Comprehensive Inorganic Chemistry III. Volume 7: Inorganic Electrochemistry [7, 3 ed.] 9780128231449

Table of contents :
Cover
Half Title
Comprehensive Inorganic Chemistry III. Volume 7: Inorganic Electrochemistry
Copyright
Contents of Volume 7
Editor Biographies
Volume Editors
Contributors to Volume 7
Preface
Vol. 1: Synthesis, Structure, and Bonding in Inorganic Molecular Systems
Vol. 2: Bioinorganic Chemistry and Homogeneous Biomimetic Inorganic Catalysis
Vol. 3: Theory and Bonding of Inorganic Non-molecular Systems
Vol. 4: Solid State Inorganic Chemistry
Vol. 5: Inorganic Materials Chemistry
Vol. 6: Heterogeneous Inorganic Catalysis
Vol. 7: Inorganic Electrochemistry
Vol. 8: Inorganic Photochemistry
Vol. 9: NMR of Inorganic Nuclei
Vol. 10: X-ray, Neutron and Electron Scattering Methods in Inorganic Chemistry
7.01. Introduction: Inorganic electrochemistry
Abstract
References
7.02. Status of Li(Na)-based anionic redox materials for better batteries
Content
Abstract
7.02.1 Introduction
7.02.2 The journey of anionic redox chemistry
7.02.2.1 The birth of insertion chemistry based on dichalcogenides
7.02.2.2 Anion–cation redox competition: Ligand hole chemistry and anion polymerization
7.02.2.3 Oxygen redox activity in LiCoO2
7.02.2.4 Li2MnO3-based compounds
7.02.2.5 Model Li-rich or Na systems
7.02.2.6 Theoretical progresses
7.02.2.7 Practical issues: Sluggish kinetics, voltage hysteresis, and voltage fade
7.02.2.8 Sulfides: There and back again
7.02.3 Fundamentals behind anionic redox
7.02.3.1 Band structure descriptions
7.02.3.2 Anionic redox activity from O 2p “NB states” or “lone-pair states"
7.02.3.3 Charge-transfer (Δ) vs. Mott–Hubbard (U) classification
7.02.3.4 Reductive coupling mechanism
7.02.3.5 Anionic activity from O (2p)-M(nd) π-type interaction
7.02.3.6 The nature of oxidized O2-: Electron holes, O-O dimers, trapper O2, molecules, and oxygen release
7.02.4 Anionic redox opening a new rich materials chemistry
7.02.4.1 Increasing the Li/M and O/M ratio in layered rock-salt compound
7.02.4.2 Playing with the alkaline ion: From Li to Na
7.02.4.3 Ligand manipulation: From oxides to sulfides/selenides
7.02.4.4 Cation disorder, superstructure, and structural dimensionality
7.02.5 Practical issues and their fundamental understandings
7.02.5.1 Chemical and electrochemical irreversibility
7.02.5.2 Voltage fade
7.02.5.3 Voltage hysteresis
7.02.6 Conclusions and outlook
Acknowledgment
References
7.03. Electrode materials for reversible sodium ions de/intercalation
Content
Abstract
7.03.1 Introduction
7.03.2 Positive electrode materials
7.03.2.1 Layered oxides of transition metals
7.03.2.2 Oxoanion-based compounds
7.03.2.2.1 NaMPO4: Maricite and triphylite
7.03.2.2.2 NASICON-structured electrode materials
7.03.2.2.3 Alluaudites
7.03.2.2.4 Pyrophosphates
7.03.2.2.5 Mixed phosphates
7.03.2.3 Mixed-anion positive electrode materials
7.03.2.3.1 Na3V2(PO4)2(O,F)3
7.03.2.3.2 Na2MPO4F (M = Mn, Fe, Co)
7.03.2.3.3 AVPO4F (A – alkali metal): Tavorite and KTiOPO4 structure types
7.03.2.4 Compounds with other oxoanions and mixed anion groups: Silicates, carbonate-phosphates, Prussian Blue analogs
7.03.3 Negative electrode materials
7.03.3.1 Carbon-based negative electrode materials
7.03.3.1.1 Approaches for describing mechanisms of sodium ion storage
7.03.3.2 Titanium-based materials
7.03.3.3 Alloys
7.03.3.3.1 Group 14
7.03.3.3.1.1 Silicon-based electrode materials
7.03.3.3.1.2 Sn- and Pb-based electrode materials
7.03.3.3.2 Group 15
7.03.4 Summary and outlook
Acknowledgment
References
7.04. Electrode materials for K-ion batteries
Content
Abstract
7.04.1 Introduction to K-ion battery
7.04.2 Positive electrode materials
7.04.2.1 Layered oxides as positive electrode materials
7.04.2.1.1 Classification of layered structures
7.04.2.1.2 Stable structure types of layered AxMO2
7.04.2.1.3 Single transition metal oxides
7.04.2.1.4 P2- and P3-type binary and ternary transition-metal systems
7.04.2.2 Prussian blue analogues
7.04.2.2.1 Prussian blue analogues as electrode materials
7.04.2.2.2 Li, Na, and K insertion into Prussian blue analogues
7.04.2.2.3 Prussian blue analogues for K-ion batteries
7.04.2.2.4 Structural evolution during Kþ insertion
7.04.2.2.5 Particle size and anion vacancy effect on the electrochemical performance
7.04.2.3 Polyanionic compounds as positive electrode materials
7.04.2.3.1 KTiOPO4-type structure materials
7.04.2.3.2 KxMP2O7 (M = Fe, Mn, and V)
7.04.2.3.3 K3V2(PO4)3 and K3V2(PO4)2F3
7.04.3 Negative electrode materials
7.04.3.1 Carbon materials
7.04.3.1.1 K intercalation into graphite
7.04.3.1.2 Electrochemical properties of graphite
7.04.3.1.3 Hard and soft carbon
7.04.3.2 K Alloys and other potassiatable compounds
7.04.3.2.1 Alkali metal alloy materials and compounds for Li-, Na-, and K-ion batteries
7.04.3.2.2 Group 14 elements and compounds
7.04.3.2.3 Group 15 elements and compounds
7.04.3.3 Transition metal oxides as negative electrode materials
7.04.3.3.1 Ti, Mo, and Nb oxides
7.04.3.3.2 Transition metal oxides based on conversion reaction
7.04.3.4 Transition metal chalcogenides
7.04.3.4.1 3d transition metal dichalcogenides (TiS2 and VS2)
7.04.3.4.2 4d and 5d transition metal dichalcogenides (MoS2, MoSe2, and WS2)
7.04.3.4.3 Metal sulfides based on conversion or conversion-alloying reactions
7.04.4 Summary and perspective
References
7.05. Charge transfer through interfaces in metal-ion intercalation systems
Content
Abstract
7.05.1 Introduction
7.05.2 Interphases and electron transfer in metal-ion batteries
7.05.3 Charge transfer kinetics in ion intercalation processes
7.05.3.1 Heterogeneous charge transfer mechanism
7.05.3.2 Heterogeneous charge transfer kinetics
7.05.3.3 Phenomenological description of ion intercalation kinetics
7.05.4 Experimental studies of charge transfer kinetics
7.05.4.1 Electrochemical signatures of rate-limiting steps
7.05.4.2 Electrochemical impedance spectroscopy of LIBs
7.05.4.3 Apparent activation energies for interfacial charge transfer
7.05.5 Modeling charge transfer in metal-ion batteries
7.05.5.1 Modeling methods
7.05.5.2 Construction of slabs with interfaces
7.05.5.2.1 Stability of interface
7.05.5.2.2 Solid-solid interfaces
7.05.5.2.3 Solid-liquid interfaces
7.05.5.3 Solid/solid interfaces
7.05.5.3.1 Structure and energetics of solid/solid interfaces on the anode
7.05.5.3.2 Structure and energetics of solid/solid interfaces on the cathode
7.05.5.4 Metal-ion diffusion inside inorganic SEI phases
7.05.5.5 Charge transfer across solid/solid interfaces
7.05.5.5.1 Li+ transfer across Li/iSEI interfaces
7.05.5.5.2 Li+ transfer across graphite/iSEI interfaces
7.05.5.5.3 Li+ transfer across electrode/iSEI and iSEI/iSEI interfaces
7.05.5.6 ET during Li+ electrodeposition
7.05.5.7 Modeling of liquid solutions
7.05.5.8 First stages of liquid electrolyte decomposition
7.05.5.9 Structure and charge transfer at the anode/liquid electrolyte interfaces
7.05.5.9.1 Li+ transfer through the electrolyte/oSEI interface
7.05.5.9.2 Li+ transfer through the oSEI/iSEI interface
7.05.5.9.3 Li+ transfer through the anode/electrolyte interface
7.05.5.10 Structure and charge transfer at cathode/liquid electrolyte interface
7.05.6 Concluding remarks
Acknowledgments
References
7.06. Transition metal hexacyanoferrates as catalysts for (bio)sensors
Content
Abstract
7.06.1 Introduction
7.06.2 Structure and electroactivity
7.06.3 Synthesis
7.06.4 Transition metal hexacyanoferrates as electrocatalysts of hydrogen peroxide reduction
7.06.5 Prussian Blue based (bio)sensors
7.06.6 Prussian Blue based nanozymes and their applications
7.06.7 Conclusion
References
7.07. Conventional and less conventional solution-based synthesis of battery materials: Cathodes, anodes and electrolytes
Content
Abstract
7.07.1 Introduction
7.07.2 Synthesis routes for inorganic materials
7.07.2.1 Different applications require different synthesis routes
7.07.2.2 Aqueous sol(ution)-gel synthesis
7.07.2.3 (Co)precipitation
7.07.2.3.1 Precipitation
7.07.2.3.2 Co-precipitation
7.07.2.4 Hydrothermal and solvothermal synthesis routes
7.07.2.5 Combustion synthesis
7.07.3 Synthesis of battery materials
7.07.3.1 LiMn2O4 (LMO) cathode material
7.07.3.2 Li2MnO3 (Li-rich LMO) cathode material
7.07.3.3 LiNi0.5Mn1.5O4 (LNMO) cathode material
7.07.3.4 LiNi-x-yMnxCoyO2 (NMC or NCM) cathode materials
7.07.3.5 Li4Ti5O12 (LTO) anode material
7.07.3.6 Sulfide solid-state electrolytes
7.07.3.6.1 Binary Li2S - P2S5-type electrolytes
7.07.3.6.2 Ternary sulfide electrolytes
7.07.4 Surface modification
7.07.4.1 Strategies for surface modification
7.07.4.2 Synthesis of surface modified lithium-ion battery cathode materials
7.07.4.2.1 The behavior of metal oxides dispersed in aqueous solution
7.07.4.2.2 Interfacial deposition mechanisms of ionic species on solid surfaces
7.07.4.2.3 Surface modification through ‘deposition-precipitation’
7.07.4.2.4 Surface modification by hydro- or solvothermal synthesis
7.07.4.2.5 Surface modification by sol(ution)-gel synthesis
7.07.4.2.6 Post-synthesis thermal treatment
7.07.5 Conclusion
References
7.08. Nanostructured materials for electrochemical capacitors
Content
Abstract
7.08.1 Introduction to electrochemical capacitors
7.08.1.1 Examples of devices
7.08.1.2 Energy storage of batteries vs. electrochemical capacitors
7.08.1.3 Electrochemical capacitor applications
7.08.2 Electrochemical capacitor mechanisms
7.08.2.1 Electric double layer capacitance
7.08.2.2 Pseudocapacitance
7.08.3 Electrochemical capacitor materials and applications
7.08.3.1 Activated carbon
7.08.3.2 Manganese oxides
7.08.3.3 Insertion-type materials
7.08.3.4 Other materials
7.08.4 Future directions and opportunities
References
7.09. Development of polyanionic sodium-ion battery insertion materials
Content
Abstract
7.09.1 Introduction
7.09.2 Phosphate class of polyanionic cathodes
7.09.2.1 NASICON-type phosphates
7.09.2.2 Olivine phosphates
7.09.2.3 Pyrophosphates
7.09.2.4 Metaphosphates
7.09.2.5 Alluaudite phosphates
7.09.3 Sulfate class of polyanionic cathodes
7.09.3.1 Bisulfates
7.09.3.2 Alluaudite sulfates
7.09.4 Other polyanionic cathodes
7.09.4.1 Borates
7.09.4.2 Silicates
7.09.4.3 Alluaudites
7.09.5 Mixed polyanionic cathodes
7.09.5.1 Fluorophosphates
7.09.5.1.1 Vanadium-based fluorophosphates
7.09.5.1.2 Other fluorophosphates
7.09.5.2 Mixed phosphates [(PO4)(P2O7)]
7.09.5.2.1 Na4M3(PO4)2P2O7 (M = Mn, Fe, Co, Ni)
7.09.5.2.2 Na7V4(P2O7)4PO4
7.09.5.3 Fluorosulfates
7.09.5.4 Hydroxyosulfates
7.09.5.5 Carbonophosphates
7.09.5.6 Nitridophosphates
7.09.5.7 Oxalate derivatives
7.09.5.8 Phosphosulfates
7.09.5.9 Phosphonitrates
7.09.6 Conclusions
Acknowledgement
References
7.10. Electrode materials viewed with transmission electron microscopy
Content
Abstract
7.10.1 Introduction
7.10.2 Transmission electron microscopy techniques in brief
7.10.2.1 TEM data visualization, manipulation and treatment
7.10.2.2 (S)TEM image simulation
7.10.3 Electron beam damage in transmission electron microscopy
7.10.4 Electron diffraction techniques for metal-ion battery electrodes
7.10.5 Imaging of the local crystal and defect structure
7.10.5.1 Point defects and order-disorder in electrode materials
7.10.5.2 Planar defects
7.10.5.3 Phase boundaries, grain boundaries and surfaces
7.10.5.4 Imaging in 3D
7.10.6 Spectroscopy with electrons
7.10.7 In situ and operando observations of electrochemical reactions
7.10.8 Conclusions and outlook
Acknowledgement
References
7.11. Chemistry of Li-air batteries
Content
Abstract
7.11.1 Introduction
7.11.2 Positive electrode
7.11.2.1 Discharge process
7.11.2.1.1 General reaction pathway
7.11.2.1.2 Superoxide anion formation and solvation. Surface and solution-mediated mechanism
7.11.2.1.3 Fundamental aspects of lithium peroxide crystallization
7.11.2.1.4 Lithium peroxide deposition in porous electrode
7.11.2.2 Charge process
7.11.2.3 Heterogeneous ORR/OER catalysts
7.11.3 Negative electrode
7.11.3.1 Many shapes of lithium
7.11.3.2 Mechanisms of morphological instability
7.11.3.3 Steps toward uniform deposition
7.11.3.3.1 SEI design
7.11.3.3.2 Electrolyte design
7.11.3.3.3 Electrode design
7.11.3.3.4 Alternative anode materials
7.11.4 Reactions with reactive oxygen species
7.11.4.1 Electrolyte decomposition
7.11.4.2 Carbon electrode degradation
7.11.5 Redox mediators
7.11.5.1 Basic principles
7.11.5.2 Redox mediators for charge
7.11.5.3 “Shuttle effect”
7.11.5.4 Redox mediators for discharge
7.11.5.5 Bifunctional and dual mediators
7.11.6 Concluding remarks and future prospects
References
7.12. Mineral inspired electrode materials for metal-ion batteries
Content
Abstract
7.12.1 Introduction
7.12.2 Phosphate minerals in pegmatites
7.12.3 Olivine-type cathode materials LiMPO4
7.12.3.1 Trihylite LiFePO4
7.12.3.2 LiMPO4 (M = Mn, Co, Ni)
7.12.4 NASICON–type materials, structurally related to kosnarite, KZr2(PO4)3
7.12.4.1 Li3V2(PO4)3
7.12.4.2 Na3V2(PO4)3
7.12.4.3 A3Ti2(PO4)3, A = Li, Na
7.12.5 Electrode materials structurally related to natisite, Na2TiSiO5
7.12.5.1 Na3V2(PO4)2F3
7.12.5.2 Na3V2(PO4)2(O,F)3
7.12.6 Tavorite based electrode materials
7.12.6.1 LiVPO4Y (Y = F, O)
7.12.6.2 LiMPO4F (M = Fe, Ti)
7.12.6.3 Fluoride sulfates LiMSO4F
7.12.7 Electrode materials structurally related to katiarsite, KTiOAsO4
7.12.7.1 KVPO4+δdF1-δ
7.12.7.2 KTiPO4+δF1-δ
7.12.8 Concluding remarks
7.12.9 Appendix
7.12.9.1 The list of mentioned minerals
Acknowledgment
References
7.13. Computational design of materials for metal-ion batteries
Content
Abstract
7.13.1 Introduction. Metal-ion batteries - state of the art and the role of computational design
7.13.2 Materials and main characteristics of batteries
7.13.2.1 Classification of ion conducting materials
7.13.2.2 Battery Characteristics
7.13.3 Computational design of ion conducting materials
7.13.3.1 Prerequisites for high ionic conductivity in solids
7.13.3.2 Simulation techniques
7.13.3.2.1 Geometrical/topological analysis
7.13.3.2.2 Bond valence sum energy modeling
7.13.3.2.3 Classical molecular dynamics and kinetic Monte Carlo simulations
7.13.3.2.4 Density functional theory calculations
7.13.3.3 Software for modeling ion conducting materials
7.13.3.3.1 Software for geometrical/topological analysis
7.13.3.3.2 Software for BVSE modeling
7.13.3.3.3 Software for classical MD and KMC simulations
7.13.3.3.4 Software for DFT modeling
7.13.3.4 Databases of ion conducting materials
7.13.4 Modeling versus experiment: A comparison
7.13.5 Conclusions
References
Relevant Websites
7.14. Lithium sulfur batteries: Electrochemistry and mechanistic research
Content
Abstract
7.14.1 Lithium–sulfur batteries
7.14.1.1 Operational principles
7.14.1.2 Problems and challenges
7.14.1.3 Common materials and parameters
7.14.1.3.1 Cathodes
7.14.1.3.2 Anodes
7.14.1.3.3 Electrolyte
7.14.1.3.4 Separator
7.14.1.4 Characterization techniques used in Li-S battery mechanistic research
7.14.1.4.1 Electrochemical techniques
7.14.1.4.2 Modeling
7.14.1.4.3 Chromatographic analysis
7.14.1.4.4 Liquid electrolyte physicochemical properties determination
7.14.1.4.5 X-ray techniques
7.14.1.4.6 Optical spectroscopy
7.14.1.4.7 Microscopy
7.14.1.4.8 Nuclear magnetic and electron paramagnetic resonance spectroscopy
7.14.1.5 Li-S battery mechanism
7.14.1.6 Summary
References
7.15. Fundamentals and applications of enzymatic bioelectrocatalysis
Content
Abstract
7.15.1 Introduction to bioelectrocatalysis
7.15.2 Principles of bioelectrocatalysts
7.15.2.1 Enzymes
7.15.2.1.1 Chemical nature of enzymes
7.15.2.1.2 Enzyme specificity
7.15.2.1.3 Principles of enzyme catalysis
7.15.2.1.4 Fundamentals of enzyme kinetics
7.15.2.1.5 Factors impacting enzyme activity
7.15.2.1.6 Enzyme classifications: Metalloenzyme and non-metalloenzymes
7.15.2.2 Enzymatic bioelectrocatalysts
7.15.2.3 Enzyme cascades
7.15.2.4 Enzyme engineering
7.15.2.4.1 Rational design of proteins
7.15.2.4.2 Directed evolution
7.15.3 Electron transfer mechanisms
7.15.3.1 Mediated electron transfer (MET)
7.15.3.2 Direct electron transfer (DET)
7.15.4 Electrodes, electrode materials, and bioelectrocatalyst-electrode connections
7.15.4.1 Electrodes and electrode materials
7.15.4.1.1 High surface area electrodes
7.15.4.1.2 Nanostructured electrodes
7.15.4.2 Enzyme bioelectrocatalyst-electrode connections and immobilization strategies
7.15.5 Applications
7.15.5.1 Biosensors
7.15.5.1.1 Principles
7.15.5.1.2 Characterization of the analytical performance of electrochemical enzymatic biosensors
7.15.5.1.3 Electron transfer in electrochemical enzymatic biosensors
7.15.5.1.3.1 First-generation biosensors
7.15.5.1.3.2 Second-generation biosensors
7.15.5.1.3.3 Third generation biosensors
7.15.5.1.4 Electrochemical enzymatic biosensing applications
7.15.5.1.4.1 Biosensors for environmental sensing
7.15.5.1.4.2 Biosensors for analysis of food and beverage quality
7.15.5.1.4.3 Biosensors for clinical sensing and medical diagnostics
7.15.5.1.4.4 Wearable electrochemical biosensors
7.15.5.1.4.5 Self-powered biosensors
7.15.5.2 Enzymatic fuel cells
7.15.5.2.1 Principles
7.15.5.2.2 Electrochemical methods for characterization of enzymatic fuel cells
7.15.5.2.2.1 Open circuit voltage, polarization and power curves, and power generation
7.15.5.2.2.2 Cyclic voltammetry
7.15.5.2.2.3 Rotating disk electrode voltammetry
7.15.5.2.2.4 Electrochemical impedance spectroscopy
7.15.5.2.3 Applications of enzymatic biofuel cells
7.15.5.3 Bioelectrosynthesis
7.15.5.3.1 Principles
7.15.5.3.2 Characterization of enzymatic electrosynthesis performance
7.15.5.3.3 Advances in enzymatic electrosynthesis
7.15.5.3.3.1 Cofactor regeneration
7.15.5.3.3.2 Enzymatic bioelectrochemical CO2 conversion
7.15.5.3.3.3 Enzymatic electrochemical N2 reduction
7.15.5.3.3.4 Enzymatic electrochemical H2 production
7.15.6 Conclusion
Acknowledgments
References
7.16. Benchmarking in electrocatalysis
Content
Abbreviations
Abstract
7.16.1 Introduction
7.16.2 Factors affecting the rate of an electrocatalytic reaction
7.16.2.1 Overpotential and its nature
7.16.2.2 Charge-transfer overpotential for a single electron step
7.16.2.3 Multi-electron processes
7.16.2.4 Influence of adsorption
7.16.2.5 Double-layer effects
7.16.2.6 Mass-transport effects
7.16.2.7 Ohmic effects
7.16.2.8 Faradaic efficiency (FE)
7.16.3 Electrocatalytic materials
7.16.4 Evaluation of the real surface area of electrocatalysts
7.16.4.1 Non-electrochemical methods
7.16.4.1.1 Microscopy and diffraction techniques
7.16.4.1.2 Gas-phase adsorption techniques
7.16.4.2 Electrochemical methods
7.16.4.2.1 Hydrogen adsorption/desorption
7.16.4.2.2 Formation/reduction of surface (hydr)oxides
7.16.4.2.3 CO stripping and metal UPD
7.16.4.2.4 Red-ox pseudocapacitance
7.16.4.2.5 Double layer capacitance
7.16.5 Experimental measurement of the electrocatalytic activity
7.16.5.1 Measurements in a 3-electrode liquid electrolyte cell
7.16.5.1.1 The rotating disk electrode (RDE)
7.16.5.1.1.1 Brief description of the RDE
7.16.5.1.1.2 Application of the RDE for studies of planar electrodes
7.16.5.1.1.3 Application of the RDE for studies of modified electrodes
7.16.5.1.1.4 Application of the RDE for studies of nanoparticles and porous materials
7.16.5.1.2 The rotating ring-disk electrode (RRDE) and other dual-electrode methods
7.16.5.2 Measurements with a membrane-electrode assembly (MEA)
7.16.5.3 Other methods
7.16.5.3.1 Half-cell measurements with gas-diffusion electrodes (GDE)
7.16.5.3.2 Microelectrode techniques
7.16.6 Activity metrics
7.16.6.1 Overpotential at a defined current density
7.16.6.2 Onset potential and half-wave potential
7.16.6.3 Kinetic current density
7.16.6.4 Exchange current density
7.16.6.5 Turnover frequency
7.16.6.6 Tafel slope
7.16.6.7 Faradaic efficiency
7.16.7 Best practice for evaluating performance of electrocatalysts
7.16.7.1 Catalytic layer (CL)
7.16.7.1.1 Ascertaining materials stability
7.16.7.1.2 Electrode conditioning and pretreatment
7.16.7.1.3 Homogeneity of the thin catalyst film
7.16.7.1.4 Conductivity of the catalytic layer (CL)
7.16.7.1.5 Co-catalytic materials
7.16.7.1.6 Catalyst loading and utilization
7.16.7.2 Choice of the supporting electrolyte
7.16.7.3 Choice of the electrochemical cell, counter and reference electrode
7.16.7.4 Measurement protocol
7.16.7.5 Data analysis and reporting
7.16.8 Summary and outlook
Acknowledgments
References

Citation preview

COMPREHENSIVE INORGANIC CHEMISTRY III

COMPREHENSIVE INORGANIC CHEMISTRY III EDITORS IN CHIEF

Jan Reedijk Leiden Institute of Chemistry, Leiden University, Leiden, the Netherlands

Kenneth R. Poeppelmeier Department of Chemistry, Northwestern University, Evanston, IL, United States

VOLUME 7

Inorganic Electrochemistry VOLUME EDITORS

Artem Abakumov Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia

Keith Stevenson Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia

Evgeny Antipov Chemistry Faculty, Lomonosov Moscow State University, Moscow, Russia

Amsterdam • Boston • Heidelberg • London • New York • Oxford Paris • San Diego • San Francisco • Singapore • Sydney • Tokyo

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge MA 02139, United States Copyright Ó 2023 Elsevier Ltd. All rights reserved No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers may always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN 978-0-12-823144-9

For information on all publications visit our website at http://store.elsevier.com

Publisher: Oliver Walter Acquisitions Editors: Clodagh Holland-Borosh and Blerina Osmanaj Content Project Manager: Pamela Sadhukhan Associate Content Project Manager: Abraham Lincoln Samuel Designer: Victoria Pearson Esser

CONTENTS OF VOLUME 7 Editor Biographies

vii

Volume Editors

ix

Contributors to Volume 7

xv

Preface

xix

7.01

Introduction: Inorganic electrochemistry Artem M Abakumov, Evgeny V Antipov, and Keith J Stevenson

1

7.02

Status of Li(Na)-based anionic redox materials for better batteries Biao Li and Jean-Marie Tarascon

6

7.03

Electrode materials for reversible sodium ions de/intercalation Aleksandr Sh Samarin, Ivan A Trussov, and Stanislav S Fedotov

46

7.04

Electrode materials for K-ion batteries Tomooki Hosaka, Kei Kubota, and Shinichi Komaba

83

7.05

Charge transfer through interfaces in metal-ion intercalation systems Dmitry A Aksyonov and Victoria A Nikitina

128

7.06

Transition metal hexacyanoferrates as catalysts for (bio)sensors Maria A Komkova and Arkady A Karyakin

172

7.07

Conventional and less conventional solution-based synthesis of battery materials: Cathodes, anodes and electrolytes D De Sloovere, B Joos, F Ulu, SK Mylavarapu, AS Kelchtermans, R Bolia, T Vranken, A Paulus, MK Van Bael, and A Hardy

186

7.08

Nanostructured materials for electrochemical capacitors Ran Ding, Matthew Chagnot, Saeed Saeed, and Veronica Augustyn

225

7.09

Development of polyanionic sodium-ion battery insertion materials Shashwat Singh, Sai Pranav Vanam, Shubham Lochab, Maximilian Fichtner, and Prabeer Barpanda

241

7.10

Electrode materials viewed with transmission electron microscopy Elena D Orlova, Anatolii V Morozov, and Artem M Abakumov

272

7.11

Chemistry of Li-air batteries Alina Inozemtseva, Alexey Rulev, Tatiana Zakharchenko, Valerii Isaev, Lada Yashina, and Daniil Itkis

324

v

vi

Contents of Volume 7

7.12

Mineral inspired electrode materials for metal-ion batteries Nellie R Khasanova, Oleg A Drozhzhin, Olga V Yakubovich, and Evgeny V Antipov

363

7.13

Computational design of materials for metal-ion batteries Artem A Kabanov, Yelizaveta A Morkhova, Iliya A Bezuglov, and Vladislav A Blatov

404

7.14

Lithium sulfur batteries: Electrochemistry and mechanistic research Robert Dominko, Sara Drvaric Talian, and Alen Vizintin

430

7.15

Fundamentals and applications of enzymatic bioelectrocatalysis Olja Simoska, Yoo Seok Lee, and Shelley D Minteer

456

7.16

Benchmarking in electrocatalysis Elena R Savinova and Alexandr G Oshchepkov

492

EDITOR BIOGRAPHIES Editors in Chief Jan Reedijk Jan Reedijk (1943) studied chemistry at Leiden University where he completed his Ph.D. (1968). After a few years in a junior lecturer position at Leiden University, he accepted a readership at Delft University of Technology in 1972. In 1979 he accepted a call for Professor of Chemistry at Leiden University. After 30 years of service, he retired from teaching in 2009 and remained as an emeritus research professor at Leiden University. In Leiden he has acted as Chair of the Department of Chemistry, and in 1993 he became the Founding Director of the Leiden Institute of Chemistry. His major research activities have been in Coordination Chemistry and Bioinorganic Chemistry, focusing on biomimetic catalysis, molecular materials, and medicinal inorganic chemistry. Jan Reedijk was elected member of the Royal Netherlands Academy of Sciences in 1996 and he was knighted by the Queen of the Netherlands to the order of the Dutch Lion (2008). He is also lifetime member of the Finnish Academy of Sciences and Letters and of Academia Europaea. He has held visiting professorships in Cambridge (UK), Strasbourg (France), Münster (Germany), Riyadh (Saudi Arabia), Louvain-la-Neuve (Belgium), Dunedin (New Zealand), and Torun (Poland). In 1990 he served as President of the Royal Netherlands Chemical Society. He has acted as the Executive Secretary of the International Conferences of Coordination Chemistry (1988–2012) and served IUPAC in the Division of Inorganic Chemistry, first as a member and later as (vice/past) president between 2005 and 2018. After his university retirement he remained active as research consultant and in IUPAC activities, as well as in several editorial jobs. For Elsevier, he acted as Editor-in-Chief of the Reference Collection in Chemistry (2013–2019), and together with Kenneth R. Poeppelmeier for Comprehensive Inorganic Chemistry II (2008–2013) and Comprehensive Inorganic Chemistry III (2019-present). From 2018 to 2020, he co-chaired the worldwide celebrations of the International Year of the Periodic Table 2019. Jan Reedijk has published over 1200 papers (1965–2022; cited over 58000 times; h ¼ 96). He has supervised 90 Ph.D. students, over 100 postdocs, and over 250 MSc research students. Kenneth R. Poeppelmeier Kenneth R. Poeppelmeier (1949) completed his undergraduate studies in chemistry at the University of Missouri (1971) and then volunteered as an instructor at Samoa CollegedUnited States Peace Corps in Western Samoa (1971–1974). He completed his Ph.D. (1978) in Inorganic Chemistry with John Corbett at Iowa State University (1978). He joined the catalysis research group headed by John Longo at Exxon Research and Engineering Company, Corporate Research–Science Laboratories (1978–1984), where he collaborated with the reforming science group and Exxon Chemicals to develop the first zeolite-based light naphtha reforming catalyst. In 1984 he joined the Chemistry Department at Northwestern University and the recently formed Center for Catalysis and Surface Science (CCSS). He is the Charles E. and Emma H. Morrison Professor of Chemistry at Northwestern University and a NAISE Fellow joint with Northwestern University and the Chemical Sciences and Engineering Division, Argonne National Laboratory. Leadership positions held include Director, CCSS, Northwestern University; Associate Division Director for Science, Chemical Sciences and Engineering Division, Argonne National Laboratory; President of the Chicago Area Catalysis Club; Associate Director, NSF Science and Technology Center for Superconductivity; and Chairman of the ACS Solid State Subdivision of the Division of Inorganic Chemistry. His major research activities have been in Solid State and Inorganic Materials Chemistry focusing on heterogeneous catalysis, solid state chemistry, and materials chemistry. His awards include National Science Council of Taiwan Lecturer (1991); Dow Professor of Chemistry (1992–1994); AAAS Fellow, the American Association for the Advancement of Science (1993); JSPS Fellow, Japan Society for the Promotion of Science (1997); Natural Science Foundation of China Lecturer (1999); National Science Foundation Creativity Extension Award (2000

vii

viii

Editor Biographies

and 2022); Institut Universitaire de France Professor (2003); Chemistry Week in China Lecturer (2004); Lecturer in Solid State Chemistry, China (2005); Visitantes Distinguidos, Universid Complutenses Madrid (2008); Visiting Professor, Chinese Academy of Sciences (2011); 20 years of Service and Dedication Award to Inorganic Chemistry (2013); Elected foreign member of Spanish National Academy: Real Academia de Ciencia, Exactas, Fisicas y Naturales (2017); Elected Honorary Member of the Royal Society of Chemistry of Spain (RSEQ) (2018); and the TianShan Award, Xinjiang Uygur Autonomous Region of China (2021). He has organized and was Chairman of the Chicago Great Lakes Regional ACS Symposium on Synthesis and Processing of Advanced Solid State Materials (1987), the New Orleans National ACS Symposium on Solid State Chemistry of Heterogeneous Oxide Catalysis, Including New Microporous Solids (1987), the Gordon Conference on Solid State Chemistry (1994) and the First European Gordon Conference on Solid State Chemistry (1995), the Spring Materials Research Society Symposium on Environmental Chemistry (1995), the Advisory Committee of Intense Pulsed Neutron Source (IPNS) Program (1996–1998), the Spring Materials Research Society Symposium on Perovskite Materials (2003), the 4th International Conference on Inorganic Materials, University of Antwerp (2004), and the Philadelphia National ACS Symposium on Homogeneous and Heterogeneous Oxidation Catalysis (2004). He has served on numerous Editorial Boards, including Chemistry of Materials, Journal of Alloys and Compounds, Solid State Sciences, Solid State Chemistry, and Science China Materials, and has been a co-Editor for Structure and Bonding, an Associate Editor for Inorganic Chemistry, and co-Editor-in-Chief with Jan Reedijk for Comprehensive Inorganic Chemistry II (published 2013) and Comprehensive Inorganic Chemistry III (to be published in 2023). In addition, he has served on various Scientific Advisory Boards including for the World Premier International Research Center Initiative and Institute for Integrated Cell-Material Sciences Kyoto University, the European Center SOPRANO on Functional Electronic Metal Oxides, the Kyoto University Mixed-Anion Project, and the Dresden Max Planck Institute for Chemistry and Physics. Kenneth Poeppelmeier has published over 500 papers (1971–2022) and cited over 28000 times (h-index ¼ 84). He has supervised over 200 undergraduates, Ph.D. students, postdocs, and visiting scholars in research.

VOLUME EDITORS Risto S. Laitinen Risto S. Laitinen is Professor Emeritus of Chemistry at the University of Oulu, Finland. He received the M.Sc and Ph.D. degrees from Helsinki University of Technology (currently Aalto University). His research interests are directed to synthetic, structural, and computational chemistry of chalcogen compounds focusing on selenium and tellurium. He has published 250 peer-reviewed articles and 15 book chapters and has edited 2 books: Selenium and Tellurium Reagents: In Chemistry and Materials Science with Raija Oilunkaniemi (Walther de Gruyter, 2019) and Selenium and Tellurium Chemistry: From Small Molecules to Biomolecules and Materials with Derek Woollins (Springer, 2011). He has also written 30 professional and popular articles in chemistry. He is the Secretary of the Division of Chemical Nomenclature and Structure Representation, International Union of Pure and Applied Chemistry, for the term 2016–2023. He served as Chair of the Board of Union of Finnish University Professors in 2007–2010. In 2017, Finnish Cultural Foundation (North Ostrobothnia regional fund) gave him an award for excellence in his activities for science and music. He has been a member of Finnish Academy of Science and Letters since 2003.

Vincent L. Pecoraro Professor Vincent L. Pecoraro is a major contributor in the fields of inorganic, bioinorganic, and supramolecular chemistries. He has risen to the upper echelons of these disciplines with over 300 publications (an h-index of 92), 4 book editorships, and 5 patents. He has served the community in many ways including as an Associate Editor of Inorganic Chemistry for 20 years and now is President of the Society of Biological Inorganic Chemistry. Internationally, he has received a Le Studium Professorship, Blaise Pascal International Chair for Research, the Alexander von Humboldt Stiftung, and an Honorary PhD from Aix-Maseille University. His many US distinctions include the 2016 ACS Award for Distinguished Service in the Advancement of Inorganic Chemistry, the 2021 ACS/SCF FrancoAmerican Lectureship Prize, and being elected a Fellow of the ACS and AAAS. He also recently cofounded a Biomedical Imaging company, VIEWaves. In 2022, he was ranked as one of the world’s top 1000 most influential chemists.

ix

x

Volume Editors

Zijian Guo Professor Zijian Guo received his Ph.D. from the University of Padova and worked as a postdoctoral research fellow at Birkbeck College, the University of London. He also worked as a research associate at the University of Edinburgh. His research focuses on the chemical biology of metals and metallodrugs and has authored or co-authored more than 400 peer-reviewed articles on basic and applied research topics. He was awarded the First Prize in Natural Sciences from Ministry of Education of China in 2015, the Luigi Sacconi Medal from the Italian Chemical Society in 2016, and the Outstanding Achievement Award from the Society of the Asian Biological Inorganic Chemistry in 2020. He founded Chemistry and Biomedicine Innovation Center (ChemBIC) in Nanjing University in 2019, and is serving as the Director of ChemBIC since then. He was elected to the Fellow of the Chinese Academy of Sciences in 2017. He served as Associated Editor of Coord. Chem. Rev and an editorial board member of several other journals.

Daniel C. Fredrickson Daniel C. Fredrickson is a Professor in the Department of Chemistry at the University of WisconsinMadison. He completed his BS in Biochemistry at the University of Washington in 2000, where he gained his first experiences with research and crystals in the laboratory of Prof. Bart Kahr. At Cornell University, he then pursued theoretical investigations of bonding in intermetallic compounds, the vast family of solid state compounds based on metallic elements, under the mentorship of Profs. Stephen Lee and Roald Hoffmann, earning his Ph.D. in 2005. Interested in the experimental and crystallographic aspects of complex intermetallics, he then carried out postdoctoral research from 2005 to 2008 with Prof. Sven Lidin at Stockholm University. Since starting at UW-Madison in 2009, his research group has created theory-driven approaches to the synthesis and discovery of new intermetallic phases and understanding the origins of their structural features. Some of his key research contributions are the development of the DFT-Chemical Pressure Method, the discovery of isolobal bonds for the generalization of the 18 electron rule to intermetallic phases, models for the emergence of incommensurate modulations in these compounds, and various design strategies for guiding complexity in solid state structures.

Patrick M. Woodward Professor Patrick M. Woodward received BS degrees in Chemistry and General Engineering from Idaho State University in 1991, an MS in Materials Science, and a Ph.D. in Chemistry from Oregon State University (1996) under the supervision of Art Sleight. After a postdoctoral appointment in the Physics Department at Brookhaven National Laboratory (1996–1998), he joined the faculty at Ohio State University in 1998, where he holds the rank of Professor in the Department of Chemistry and Biochemistry. He is a Fellow of the American Chemical Society (2020) and a recipient of an Alfred P. Sloan Fellowship (2004) and an NSF Career Award (2001). He has co-authored two textbooks: Solid State Materials Chemistry and the popular general chemistry textbook, Chemistry: The Central Science. His research interests revolve around the discovery of new materials and understanding links between the composition, structure, and properties of extended inorganic and hybrid solids.

Volume Editors

xi

P. Shiv Halasyamani Professor P. Shiv Halasyamani earned his BS in Chemistry from the University of Chicago (1992) and his Ph.D. in Chemistry under the supervision of Prof. Kenneth R. Poeppelmeier at Northwestern University (1996). He was a Postdoctoral Fellow and Junior Research Fellow at Christ Church College, Oxford University, from 1997 to 1999. He began his independent academic career in the Department of Chemistry at the University of Houston in 1999 and has been a Full Professor since 2010. He was elected as a Fellow of the American Association for the Advancement of Science (AAAS) in 2019 and is currently an Associate Editor of the ACS journals Inorganic Chemistry and ACS Organic & Inorganic Au. His research interests involve the design, synthesis, crystal growth, and characterization of new functional inorganic materials.

Ram Seshadri Ram Seshadri received his Ph.D. in Solid State Chemistry from the Indian Institute of Science (IISc), Bangalore, working under the guidance of Professor C. N. R. Rao FRS. After some years as a Postdoctoral Fellow in Europe, he returned to IISc as an Assistant Professor in 1999. He moved to the Materials Department (College of Engineering) at UC Santa Barbara in 2002. He was recently promoted to the rank of Distinguished Professor in the Materials Department and the Department of Chemistry and Biochemistry in 2020. He is also the Fred and Linda R. Wudl Professor of Materials Science and Director of the Materials Research Laboratory: A National Science Foundation Materials Research Science and Engineering Center (NSF-MRSEC). His work broadly addresses the topic of structure–composition– property relations in crystalline inorganic and hybrid materials, with a focus on magnetic materials and materials for energy conversion and storage. He is Fellow of the Royal Society of Chemistry, the American Physical Society, and the American Association for the Advancement of Science. He serves as Associate Editor of the journals, Annual Reviews of Materials Research and Chemistry of Materials.

Serena Cussen Serena Cussen née Corr studied chemistry at Trinity College Dublin, completing her doctoral work under Yurii Gun’ko. She then joined the University of California, Santa Barbara, working with Ram Seshadri as a postdoctoral researcher. She joined the University of Kent as a lecturer in 2009. She moved to the University of Glasgow in 2013 and was made Professor in 2018. She moved to the University of Sheffield as a Chair in Functional Materials and Professor in Chemical and Biological Engineering in 2018, where she now serves as Department Head. She works on next-generation battery materials and advanced characterization techniques for the structure and properties of nanomaterials. Serena is the recipient of several awards including the Journal of Materials Chemistry Lectureship of the Royal Society of Chemistry. She previously served as Associate Editor of Royal Society of Chemistry journal Nanoscale and currently serves as Associate Editor for the Institute of Physics journal Progress in Energy.

xii

Volume Editors

Rutger A. van Santen Rutger A. van Santen received his Ph.D. in 1971 in Theoretical Chemistry from the University of Leiden, The Netherlands. In the period 1972–1988, he became involved with catalysis research when employed by Shell Research in Amsterdam and Shell Development Company in Houston. In 1988, he became Full Professor of Catalysis at the Technical University Eindhoven. From 2010 till now he is Emeritus Professor and Honorary Institute Professor at Technical University Eindhoven. He is a member of Royal Dutch Academy of Sciences and Arts and Foreign Associate of the United States National Academy of Engineering (NAE). He has received several prestigious awards: the 1981 golden medal of the Royal Dutch Chemical Society; in 1992, the F.G. Chiappetta award of the North American Catalysis Society; in 1997, the Spinoza Award from the Dutch Foundation for Pure and Applied Research; and in 2001, the Alwin Mittasch Medal Dechema, Germany, among others. His main research interests are computational heterogeneous catalysis and complex chemical systems theory. He has published over 700 papers, 16 books, and 22 patents.

Emiel J. M. Hensen Emiel J. M. Hensen received his Ph.D. in Catalysis in 2000 from Eindhoven University of Technology, The Netherlands. Between 2000 and 2008, he worked at the University of Amsterdam, Shell Research in Amsterdam, and Eindhoven University of Technology on several topics in the field of heterogeneous catalysis. Since July 2009, he is Full Professor of Inorganic Materials and Catalysis at TU/e. He was a visiting professor at the Katholieke Universiteit Leuven (Belgium, 2001–2016) and at Hokkaido University (Japan, 2016). He is principal investigator and management team member of the gravitation program Multiscale Catalytic Energy Conversion, elected member of the Advanced Research Center Chemical Building Blocks Consortium, and chairman of the Netherlands Institute for Catalysis Research (NIOK). Hensen was Head of the Department of Chemical Engineering and Chemistry of Eindhoven University of Technology from 2016 to 2020. Hensen received Veni, Vidi, Vici, and Casmir grant awards from the Netherlands Organisation for Scientific Research. His main interests are in mechanism of heterogeneous catalysis combining experimental and computation studies. He has published over 600 papers, 20 book chapters, and 7 patents.

Artem M. Abakumov Artem M. Abakumov graduated from the Department of Chemistry at Moscow State University in 1993, received his Ph.D. in Chemistry from the same University in 1997, and then continued working as a Researcher and Vice-Chair of Inorganic Chemistry Department. He spent about 3 years as a postdoctoral fellow and invited professor in the Electron Microscopy for Materials Research (EMAT) laboratory at the University of Antwerp and joined EMAT as a research leader in 2008. Since 2015 he holds a Full Professor position at Skolkovo Institute of Science and Technology (Skoltech) in Moscow, leading Skoltech Center for Energy Science and Technology as a Director. His research interests span over a wide range of subjects, from inorganic chemistry, solid state chemistry, and crystallography to battery materials and transmission electron microscopy. He has extended experience in characterization of metal-ion battery electrodes and electrocatalysts with advanced TEM techniques that has led to a better understanding of charge–discharge mechanisms, redox reactions, and associated structural transformations in various classes of cathode materials on different spatial scales. He has published over 350 papers, 5 book chapters, and 12 patents.

Volume Editors

xiii

Keith J. Stevenson Keith J. Stevenson received his Ph.D. in 1997 from the University of Utah under the supervision of Prof. Henry White. Subsequently, he held a postdoctoral appointment at Northwestern University (1997– 2000) and a tenured faculty appointment (2000–2015) at the University of Texas at Austin. At present, he is leading the development of a new graduate level university (Skolkovo Institute for Science and Technology) in Moscow, Russia, where he is Provost and the former Director of the Center for Energy Science and Technology (CEST). To date he has published over 325 peer-reviewed publications, 14 patents, and 6 book chapters in this field. He is a recipient of Society of Electroanalytical Chemistry Charles N. Reilley Award (2021).

Evgeny V. Antipov Evgeny V. Antipov graduated from the Department of Chemistry at Moscow State University in 1981, received his Ph.D. in Chemistry in 1986, DSc degree in Chemistry in 1998, and then continued working at the same University as a Researcher, Head of the Laboratory of Inorganic Crystal Chemistry, Professor, Head of Laboratory of fundamental research on aluminum production, and Head of the Department of Electrochemistry. Since 2018 he also holds a professor position at Skolkovo Institute of Science and Technology (Skoltech) in Moscow. Currently his research interests are mainly focused on inorganic materials for application in batteries and fuel cells. He has published more than 400 scientific articles and 14 patents.

Vivian W.W. Yam Professor Vivian W.W. Yam is the Chair Professor of Chemistry and Philip Wong Wilson Wong Professor in Chemistry and Energy at The University of Hong Kong. She received both her B.Sc (Hons) and Ph.D. from The University of Hong Kong. She was elected to Member of Chinese Academy of Sciences, International Member (Foreign Associate) of US National Academy of Sciences, Foreign Member of Academia Europaea, Fellow of TWAS, and Founding Member of Hong Kong Academy of Sciences. She was Laureate of 2011 L’Oréal-UNESCO For Women in Science Award. Her research interests include inorganic and organometallic chemistry, supramolecular chemistry, photophysics and photochemistry, and metal-based molecular functional materials for sensing, organic optoelectronics, and energy research. Also see: https://chemistry.hku.hk/wwyam.

xiv

Volume Editors

David L. Bryce David L. Bryce (B.Sc (Hons), 1998, Queen’s University; Ph.D., 2002, Dalhousie University; postdoctoral fellow, 2003–04, NIDDK/NIH) is Full Professor and University Research Chair in Nuclear Magnetic Resonance at the University of Ottawa in Canada. He is the past Chair of the Department of Chemistry and Biomolecular Sciences, a Fellow of the Royal Society of Chemistry, and an elected Fellow of the Royal Society of Canada. His research interests include solid-state NMR of low-frequency quadrupolar nuclei, NMR studies of materials, NMR crystallography, halogen bonding, mechanochemistry, and quantum chemical interpretation of NMR interaction tensors. He is the author of approximately 200 scientific publications and co-author of 1 book. He is the Editor-in-Chief of Solid State Nuclear Magnetic Resonance and Section Editor (Magnetic Resonance and Molecular Spectroscopy) for the Canadian Journal of Chemistry. He has served as the Chair of Canada’s National Ultrahigh-Field NMR Facility for Solids and is a past co-chair of the International Society for Magnetic Resonance conference and of the Rocky Mountain Conference on Magnetic Resonance Solid-State NMR Symposium. His work has been recognized with the Canadian Society for Chemistry Keith Laidler Award and with the Gerhard Herzberg Award of the Canadian Association of Analytical Sciences and Spectroscopy.

Paul R. Raithby Paul R. Raithby obtained his B.Sc (1973) and Ph.D. (1976) from Queen Mary College, University of London, working for his Ph.D. in structural inorganic chemistry. He moved to the University of Cambridge in 1976, initially as a postdoctoral researcher and then as a faculty member. In 2000, he moved to the University of Bath to take up the Chair of Inorganic Chemistry when he remains to the present day, having been awarded an Emeritus Professorship in 2022. His research interests have spanned the chemistry of transition metal cluster compounds, platinum acetylide complexes and oligomers, and lanthanide complexes, and he uses laboratory and synchrotron-based X-ray crystallographic techniques to determine the structures of the complexes and to study their dynamics using time-resolved photocrystallographic methods.

Angus P. Wilkinson

Angus P. Wilkinson completed his bachelors (1988) and doctoral (1992) degrees in chemistry at Oxford University in the United Kingdom. He spent a postdoctoral period in the Materials Research Laboratory, University of California, Santa Barbara, prior to joining the faculty at the Georgia Institute of Technology as an assistant professor in 1993. He is currently a Professor in both the Schools of Chemistry and Biochemistry, and Materials Science and Engineering, at the Georgia Institute of Technology. His research in the general area of inorganic materials chemistry makes use of synchrotron X-ray and neutron scattering to better understand materials synthesis and properties.

CONTRIBUTORS TO VOLUME 7 Artem M Abakumov Skolkovo Institute of Science and Technology, Moscow, Russia Dmitry A Aksyonov Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia Evgeny V Antipov Skolkovo Institute of Science and Technology, Moscow, Russia; and Lomonosov Moscow State University, Moscow, Russia Veronica Augustyn Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC, United States Prabeer Barpanda Faraday Materials Laboratory (FaMaL), Materials Research Centre, Indian Institute of Science, Bangalore, India; Helmholtz Institut Ulm (HIU), Helmholtzstraße, Ulm, Germany; and Institute of Nanotechnology, Karlsruhe Institute of Technology (KIT), Karlsruhe, Germany Iliya A Bezuglov Samara Center for Theoretical Materials Science (SCTMS), Samara State Technical University, Samara, Russia Vladislav A Blatov Samara Center for Theoretical Materials Science (SCTMS), Samara State Technical University, Samara, Russia R Bolia Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium Matthew Chagnot Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC, United States

D De Sloovere Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Ran Ding Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC, United States Robert Dominko National Institute of Chemistry, Ljubljana, Slovenia Oleg A Drozhzhin Lomonosov Moscow State University, Moscow, Russia Sara Drvaric Talian National Institute of Chemistry, Ljubljana, Slovenia Stanislav S Fedotov Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia Maximilian Fichtner Helmholtz Institut Ulm (HIU), Helmholtzstraße, Ulm, Germany; and Institute of Nanotechnology, Karlsruhe Institute of Technology (KIT), Karlsruhe, Germany A Hardy Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Tomooki Hosaka Department of Applied Chemistry, Tokyo University of Science, Tokyo, Japan Alina Inozemtseva Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia

xv

xvi

Contributors to Volume 7

Valerii Isaev Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia

Shubham Lochab Faraday Materials Laboratory (FaMaL), Materials Research Centre, Indian Institute of Science, Bangalore, India

Daniil Itkis Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia

Shelley D Minteer Department of Chemistry, University of Utah, Salt Lake City, UT, United States

B Joos Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Artem A Kabanov Samara Center for Theoretical Materials Science (SCTMS), Samara State Technical University, Samara, Russia; and Samara Branch of P.N. Lebedev Physical Institute of the Russian Academy of Sciences, Samara, Russia Arkady A Karyakin Chemistry faculty of M.V. Lomonosov Moscow State University, Moscow, Russia AS Kelchtermans Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Nellie R Khasanova Lomonosov Moscow State University, Moscow, Russia Shinichi Komaba Department of Applied Chemistry, Tokyo University of Science, Tokyo, Japan Maria A Komkova Chemistry faculty of M.V. Lomonosov Moscow State University, Moscow, Russia Kei Kubota Department of Applied Chemistry, Tokyo University of Science, Tokyo, Japan Yoo Seok Lee Department of Chemistry, University of Utah, Salt Lake City, UT, United States Biao Li Chimie du Solide-Energie, UMR 8260, Collège de France, Paris, France; and Réseau sur le Stockage Electrochimique de l’Energie (RS2E), Amiens, France

Yelizaveta A Morkhova Samara Center for Theoretical Materials Science (SCTMS), Samara State Technical University, Samara, Russia; and Institute of Solid State Chemistry and Mechanochemistry SB RAS, Novosibirsk, Russia Anatolii V Morozov Skolkovo Institute of Science and Technology, Moscow, Russia SK Mylavarapu Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Victoria A Nikitina Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia Elena D Orlova Skolkovo Institute of Science and Technology, Moscow, Russia Alexandr G Oshchepkov Institut de Chimie et Procédés pour l’Energie, l’Environnement et la Santé, UMR 7515 CNRSUniversity of Strasbourg, Strasbourg Cedex, France; and Boreskov Institute of Catalysis, Novosibirsk, Russia A Paulus Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Alexey Rulev Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia Saeed Saeed Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC, United States Aleksandr Sh Samarin Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia

Contributors to Volume 7

Elena R Savinova Institut de Chimie et Procédés pour l’Energie, l’Environnement et la Santé, UMR 7515 CNRSUniversity of Strasbourg, Strasbourg Cedex, France; and Boreskov Institute of Catalysis, Novosibirsk, Russia Olja Simoska Department of Chemistry, University of Utah, Salt Lake City, UT, United States Shashwat Singh Faraday Materials Laboratory (FaMaL), Materials Research Centre, Indian Institute of Science, Bangalore, India Keith J Stevenson Skolkovo Institute of Science and Technology, Moscow, Russia Jean-Marie Tarascon Chimie du Solide-Energie, UMR 8260, Collège de France, Paris, France; Réseau sur le Stockage Electrochimique de l’Energie (RS2E), Amiens, France; and Sorbonne Université, Paris, France Ivan A Trussov Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Moscow, Russia F Ulu Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium

xvii

Sai Pranav Vanam Faraday Materials Laboratory (FaMaL), Materials Research Centre, Indian Institute of Science, Bangalore, India MK Van Bael Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Alen Vizintin National Institute of Chemistry, Ljubljana, Slovenia T Vranken Hasselt University, Institute for Materials Research (IMO-IMOMEC), DESINe, Diepenbeek, Belgium; IMEC division IMOMEC, Diepenbeek, Belgium; and EnergyVille, Genk, Belgium Olga V Yakubovich Lomonosov Moscow State University, Moscow, Russia Lada Yashina Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia Tatiana Zakharchenko Lomonosov Moscow State University, Moscow, Russia; and N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia

PREFACE Comprehensive Inorganic Chemistry III is a new multi-reference work covering the broad area of Inorganic Chemistry. The work is available both in print and in electronic format. The 10 Volumes review significant advances and examines topics of relevance to today’s inorganic chemists with a focus on topics and results after 2012. The work is focusing on new developments, including interdisciplinary and high-impact areas. Comprehensive Inorganic Chemistry III, specifically focuses on main group chemistry, biological inorganic chemistry, solid state and materials chemistry, catalysis and new developments in electrochemistry and photochemistry, as well as on NMR methods and diffractions methods to study inorganic compounds. The work continues our 2013 work Comprehensive Inorganic Chemistry II, but at the same time adds new volumes on emerging research areas and techniques used to study inorganic compounds. The new work is also highly complementary to other recent Elsevier works in Coordination Chemistry and Organometallic Chemistry thereby forming a trio of works covering the whole of modern inorganic chemistry, most recently COMC-4 and CCC-3. The rapid pace of developments in recent years in all areas of chemistry, particularly inorganic chemistry, has again created many challenges to provide a contemporary up-to-date series. As is typically the challenge for Multireference Works (MRWs), the chapters are designed to provide a valuable long-standing scientific resource for both advanced students new to an area as well as researchers who need further background or answers to a particular problem on the elements, their compounds, or applications. Chapters are written by teams of leading experts, under the guidance of the Volume Editors and the Editors-inChief. The articles are written at a level that allows undergraduate students to understand the material, while providing active researchers with a ready reference resource for information in the field. The chapters are not intended to provide basic data on the elements, which are available from many sources including the original CIC-I, over 50-years-old by now, but instead concentrate on applications of the elements and their compounds and on high-level techniques to study inorganic compounds. Vol. 1: Synthesis, Structure, and Bonding in Inorganic Molecular Systems; Risto S. Laitinen In this Volume the editor presents an historic overview of Inorganic Chemistry starting with the birth of inorganic chemistry after Berzelius, and a focus on the 20th century including an overview of “inorganic” Nobel Prizes and major discoveries, like inert gas compounds. The most important trends in the field are discussed in an historic context. The bulk of the Volume consists of 3 parts, i.e., (1) Structure, bonding, and reactivity in inorganic molecular systems; (2) Intermolecular interactions, and (3) Inorganic Chains, rings, and cages. The volume contains 23 chapters. Part 1 contains chapters dealing with compounds in which the heavy p-block atom acts as a central atom. Some chapters deal with the rich synthetic and structural chemistry of noble gas compounds, low-coordinate p-block elements, biradicals, iron-only hydrogenase mimics, and macrocyclic selenoethers. Finally, the chemistry and application of weakly coordinating anions, the synthesis, structures, and reactivity of carbenes containing non-innocent ligands, frustrated Lewis pairs in metal-free catalysis are discussed. Part 2 discusses secondary bonding interactions that play an important role in the properties of bulk materials. It includes a chapter on the general theoretical considerations of secondary bonding interactions, including halogen and chalcogen bonding. This section is concluded by the update of the host-guest chemistry of the molecules of p-block elements and by a comprehensive review of closed-shell metallophilic interactions. The third part of the Volume is dedicated to chain, ring and cage (or cluster) compounds in molecular inorganic chemistry. Separate

xix

xx

Preface

chapters describe the recent chemistry of boron clusters, as well as the chain, ring, and cage compounds of Group13 and 15, and 16 elements. Also, aromatic compounds bearing heavy Group 14 atoms, polyhalogenide anions and Zintl-clusters are presented. Vol. 2: Bioinorganic Chemistry and Homogeneous Biomimetic Inorganic Catalysis; Vincent L. Pecoraro and Zijian Guo In this Volume, the editors have brought together 26 chapters providing a broad coverage of many of the important areas involving metal compounds in biology and medicine. Readers interested in fundamental biochemistry that is assisted by metal ion catalysis, or in uncovering the latest developments in diagnostics or therapeutics using metal-based probes or agents, will find high-level contributions from top scientists. In the first part of the Volume topics dealing with metals interacting with proteins and nucleic acids are presented (e.g., siderophores, metallophores, homeostasis, biomineralization, metal-DNA and metal-RNA interactions, but also with zinc and cobalt enzymes). Topics dealing with iron-sulfur clusters and heme-containing proteins, enzymes dealing with dinitrogen fixation, dihydrogen and dioxygen production by photosynthesis will also be discussed, including bioinspired model systems. In the second part of the Volume the focus is on applications of inorganic chemistry in the field of medicine: e.g., clinical diagnosis, curing diseases and drug targeting. Platinum, gold and other metal compounds and their mechanism of action will be discussed in several chapters. Supramolecular coordination compounds, metal organic frameworks and targeted modifications of higher molecular weight will also be shown to be important for current and future therapy and diagnosis. Vol. 3: Theory and Bonding of Inorganic Non-molecular Systems; Daniel C. Fredrickson This volume consists of 15 chapters that build on symmetry-based expressions for the wavefunctions of extended structures toward models for bonding in solid state materials and their surfaces, algorithms for the prediction of crystal structures, tools for the analysis of bonding, and theories for the unique properties and phenomena that arise in these systems. The volume is divided into four parts along these lines, based on major themes in each of the chapters. These are: Part 1: Models for extended inorganic structures, Part 2: Tools for electronic structure analysis, Part 3: Predictive exploration of new structures, and Part 4: Properties and phenomena. Vol. 4: Solid State Inorganic Chemistry; P. Shiv Halasyamani and Patrick M. Woodward In a broad sense the field of inorganic chemistry can be broken down into substances that are based on molecules and those that are based on extended arrays linked by metallic, covalent, polar covalent, or ionic bonds (i.e., extended solids). The field of solid-state inorganic chemistry is largely concerned with elements and compounds that fall into the latter group. This volume contains nineteen chapters covering a wide variety of solid-state inorganic materials. These chapters largely focus on materials with properties that underpin modern technology. Smart phones, solid state lighting, batteries, computers, and many other devices that we take for granted would not be possible without these materials. Improvements in the performance of these and many other technologies are closely tied to the discovery of new materials or advances in our ability to synthesize high quality samples. The organization of most chapters is purposefully designed to emphasize how the exceptional physical properties of modern materials arise from the interplay of composition, structure, and bonding. Not surprisingly this volume has considerable overlap with both Volume 3 (Theory and Bonding of Inorganic NonMolecular Systems) and Volume 5 (Inorganic Materials Chemistry). We anticipate that readers who are interested in this volume will find much of interest in those volumes and vice versa Vol. 5: Inorganic Materials Chemistry; Ram Seshadri and Serena Cussen This volume has adopted the broad title of Inorganic Materials Chemistry, but as readers would note, the title could readily befit articles in other volumes as well. In order to distinguish contributions in this volume from

Preface

xxi

those in other volumes, the editors have chosen to use as the organizing principle, the role of synthesis in developing materials, reflected by several of the contributions carrying the terms “synthesis” or “preparation” in the title. It should also be noted that the subset of inorganic materials that are the focus of this volume are what are generally referred to as functional materials, i.e., materials that carry out a function usually through the way they respond to an external stimulus such as light, or thermal gradients, or a magnetic field.

Vol. 6: Heterogeneous Inorganic Catalysis; Rutger A. van Santen and Emiel J. M. Hensen This Volume starts with an introductory chapter providing an excellent discussion of single sites in metal catalysis. This chapter is followed by 18 chapters covering a large part of the field. These chapters have been written with a focus on the synthesis and characterization of catalytic complexity and its relationship with the molecular chemistry of the catalytic reaction. In the 1950s with the growth of molecular inorganic chemistry, coordination chemistry and organometallic chemistry started to influence the development of heterogeneous catalysis. A host of new reactions and processes originate from that time. In this Volume chapters on major topics, like promoted Fischer-Tropsch catalysts, structure sensitivity of well-defined alloy surfaces in the context of oxidation catalysis and electrocatalytic reactions, illustrate the broadness of the field. Molecular heterogeneous catalysts rapidly grew after high-surface synthetic of zeolites were introduced; so, synthesis, structure and nanopore chemistry in zeolites is presented in a number of chapters. Also, topics like nanocluster activation of zeolites and supported zeolites are discussed. Mechanistically important chapters deal with imaging of single atom catalysts. An important development is the use of reducible supports, such as CeO2 or Fe2O3 where the interaction between the metal and support is playing a crucial role.

Vol. 7: Inorganic Electrochemistry; Keith J. Stevenson, Evgeny V. Antipov and Artem M. Abakumov This volume bridges several fields across chemistry, physics and material science. Perhaps this topic is best associated with the book “Inorganic Electrochemistry: Theory, Practice and Applications” by Piero Zanello that was intended to introduce inorganic chemists to electrochemical methods for study of primarily molecular systems, including metallocenes, organometallic and coordination complexes, metal complexes of redox active ligands, metal-carbonyl clusters, and proteins. The emphasis in this Volume of CIC III is on the impact of inorganic chemistry on the field of material science, which has opened the gateway for inorganic chemists to use more applied methods to the broad areas of electrochemical energy storage and conversion, electrocatalysis, electroanalysis, and electrosynthesis. In recognition of this decisive impact, the Nobel Prize in Chemistry of 2019 was awarded to John B. Goodenough, M. Stanley Whittingham, and Akira Yoshino for the development of the lithium-ion battery.

Vol. 8: Inorganic Photochemistry; Vivian W. W. Yam In this Volume the editor has compiled 19 chapters discussing recent developments in a variety of developments in the field. The introductory chapter overviews the several topics, including photoactivation and imaging reagents. The first chapters include a discussion of using luminescent coordination and organometallic compounds for organic light-emitting diodes (OLEDs) and applications to highlight the importance of developing future highly efficient luminescent transition metal compounds. The use of metal compounds in photo-induced bond activation and catalysis is highlighted by non-sacrificial photocatalysis and redox photocatalysis, which is another fundamental area of immense research interest and development. This work facilitates applications like biological probes, drug delivery and imaging reagents. Photochemical CO2 reduction and water oxidation catalysis has been addressed in several chapters. Use of such inorganic compounds in solar fuels and photocatalysis remains crucial for a sustainable environment. Finally, the photophysics and photochemistry of lanthanoid compounds is discussed, with their potential use of doped lanthanoids in luminescence imaging reagents.

xxii

Preface

Vol. 9: NMR of Inorganic Nuclei; David L. Bryce Nuclear magnetic resonance (NMR) spectroscopy has long been established as one of the most important analytical tools at the disposal of the experimental chemist. The isotope-specific nature of the technique can provide unparalleled insights into local structure and dynamics. As seen in the various contributions to this Volume, applications of NMR spectroscopy to inorganic systems span the gas phase, liquid phase, and solid state. The nature of the systems discussed covers a very wide range, including glasses, single-molecule magnets, energy storage materials, bioinorganic systems, nanoparticles, catalysts, and more. The focus is largely on isotopes other than 1H and 13C, although there are clearly many applications of NMR of these nuclides to the study of inorganic compounds and materials. The value of solid-state NMR in studying the large percentage of nuclides which are quadrupolar (spin I > ½) is apparent in the various contributions. This is perhaps to be expected given that rapid quadrupolar relaxation can often obfuscate the observation of these resonances in solution. Vol. 10: X-ray, Neutron and Electron Scattering Methods in Inorganic Chemistry; Angus P. Wilkinson and Paul R. Raithby In this Volume the editors start with an introduction on the recent history and improvements of the instrumentation, source technology and user accessibility of synchrotron and neutron facilities worldwide, and they explain how these techniques work. The modern facilities now allow inorganic chemists to carry out a wide variety of complex experiments, almost on a day-to-day basis, that were not possible in the recent past. Past editions of Comprehensive Inorganic Chemistry have included many examples of successful synchrotron or neutron studies, but the increased importance of such experiments to inorganic chemists motivated us to produce a separate volume in CIC III dedicated to the methodology developed and the results obtained. The introduction chapter is followed by 15 chapters describing the developments in the field. Several chapters are presented covering recent examples of state-of-the-art experiments and refer to some of the pioneering work leading to the current state of the science in this exciting area. The editors have recognized the importance of complementary techniques by including chapters on electron crystallography and synchrotron radiation sources. Chapters are present on applications of the techniques in e.g., spin-crossover materials and catalytic materials, and in the use of time-resolved studies on molecular materials. A chapter on the worldwide frequently used structure visualization of crystal structures, using PLATON/PLUTON, is also included. Finally, some more specialized studies, like Panoramic (in beam) studies of materials synthesis and high-pressure synthesis are present. Direct observation of transient species and chemical reactions in a pore observed by synchrotron radiation and X-ray transient absorption spectroscopies in the study of excited state structures, and ab initio structure solution using synchrotron powder diffraction, as well as local structure determination using total scattering data, are impossible and unthinkable without these modern diffraction techniques. Jan Reedijk, Leiden, The Netherlands Kenneth R. Poeppelmeier, Illinois, United States March 2023

7.01

Introduction: Inorganic electrochemistry

Artem M. Abakumov, Evgeny V. Antipov, and Keith J. Stevenson, Skolkovo Institute of Science and Technology, Moscow, Russia © 2023 Elsevier Ltd. All rights reserved.

References

4

Abstract This introductory chapter provides a vision of crystallographer, inorganic chemist and electrochemist on inorganic electrochemistry as a self-standing research discipline along with the outline of the content of Volume 7 “Inorganic Electrochemistry”.

The volume title “Inorganic Electrochemistry” was perhaps chosen as a general place holder for such a broad and now interdisciplinary area that bridges several fields across chemistry, physics and material science. Perhaps this topic is best associated with the book “Inorganic Electrochemistry: Theory, Practice and Applications” by Piero Zannello that was intended to introduce inorganic chemists to electrochemical methods for study of primarily molecular systems including metallocenes, organometallic and coordination complexes, metal complexes of redox active ligands, metal-carbonyl clusters and proteins.1 In 2013, two of us published an Edition of Comprehensive Inorganic Chemistry (CIC II) with the major focus on modern solid state inorganic chemistry.2 This collection stressed “composition-structure-property” relationships with special attention to solid state materials, coordination and organometallic compounds, porous and nanomaterials, biomaterials, chemistry of the elements etc., with the support of computational chemistry on crystal and electronic structures and chemical bonding, and close relation to various fields of material science. Now we see the impact of inorganic chemistry on the field of material science that has opened the gateway for inorganic chemists to use more applied methods to the broad areas of electrochemical energy storage and conversion, electrocatalysis, electroanalysis, and electrosynthesis during past few decades. In recognition of this decisive impact, the Nobel Prize in Chemistry of 2019 has been awarded to John B. Goodenough, M. Stanley Whittingham and Akira Yoshino for the development of the lithium-ion battery. One of the Laureates was actually recognized for creating the commercially viable lithium-ion battery (i.e., for the design of an electrochemical system per se), whereas two othersdfor creating innovative intercalation-type cathode materials which enable these batteries to achieve the competitive energy density.3 The interest in (electro)chemical intercalation was stemmed from the possibility to tune electronic structure and physical properties by introducing the guest cations into the host structure, and chemical intercalation into transition metal sulfides was the prerequisite for S. Whittingham to employ TiS2 as a lithium battery positive electrode (cathode).4 Then this concept has been developed by J. Goodenough based on the fundamental differences in band structure of sulfides and oxides, enabling for the latter a possibility to extract electrons from the low-lying d-levels of transition metals thus increasing the cell voltage and, hence, energy density.5,6 This has been deeply rooted in such fundamental concepts of inorganic chemistry as difference in electronegativity of the elements and ionicity/covalency of chemical bonding.7,8 Exponentially growing number of studies on metal-ion batteries greatly shifted a conventional focus of electrochemistry from charge/mass transfer at the electrode/electrolyte interface (i.e. at the surface or subsurface areas of the electrode) towards the processes occurring inside the volume of the electrode. Indeed, the key performance parameters of the metal-ion batteries, such as capacity, energy and power density, largely depend on electronic capacity and energy positioning of cationic band(s) with respect to the Fermi level of the anode and top of the anionic band, availability and population of crystallographic sites containing mobile guest cations (mostly s-metals of the Group I and II of the Periodic Table of the Elements), energy barriers for the guest cation diffusion, electronic conductivity and exact (de)intercalation mechanisms. All these topics belong to the scope of inorganic and solid-state chemistry (Fig. 1). Accordingly, the characterization techniques which have been employed by inorganic and solid-state chemists for decades became extremely relevant for scrutinizing the electrode materials and their transformations. Crystal structures and phase transitions are studied with X-ray, neutron and electron diffraction, local structure changes and defects are visualized with advanced transmission electron microscopy and probed with nuclear magnetic resonance and Mössbauer spectroscopy, valence states of cations and anions are measured with X-ray absorption and electron energy loss spectroscopies, whereas surface characterization is provided with X-ray photoelectron spectroscopy. These methods deliver huge pile of data, which has to be correlated with the electrochemical responses measured with cyclic voltammetry, galvanostatic cycling with potential limitation, chronopotentio- and chronoamperometric methods. This inspired the development of in situ and operando techniques enabling simultaneous collection of two (or more) very different datasets.9–13 At the same time, establishing the “composition-structure-property” relationships is impossible without extensive use of various computational tools greatly contributing to inorganic electrochemistry of both molecular and solid state systems. Crystal space partitioning visualizes cavities and tunnels which can be occupied by the guest cations and through which the cations can migrate.14 Pauling’s rule on local charge neutrality formulated in 1929 gave rise to bond valence sum method semi-quantitatively correlating the formal valence of cation or anion with the bond lengths of the closest coordination environment.15 This correlation is utilized for computing the guest cation migration paths and estimating the associated energy barriers.16

Comprehensive Inorganic Chemistry III, Volume 7

https://doi.org/10.1016/B978-0-12-823144-9.00174-6

1

2

Introduction: Inorganic electrochemistry

Fig. 1 The “composition-structure-property” triad in metal-ion batteries. Reprinted from Abakumov, A. M.; Fedotov, S. S.; Antipov, E. V.; Tarascon, J.-M. Nat. Commun. 2020, 11, 4976.

Electronic structure, (de)intercalation potentials and migration barriers can be retrieved more precisely from density functional theory calculations, whereas structure transformations, short- and long-range atomic displacements are assessed with molecular dynamics.17 Nowadays, machine learning and artificial intelligence approaches come to the scene further empowering the computational inorganic electrochemistry.18 Functional electrochemical properties coupled with deep understanding of the electrode material’s crystal and electronic structure based on a variety of characterization methods and supported with computational tools create exceptionally rich research landscape where electrochemistry, inorganic and solid-state chemistry act synergistically delivering enormous added value. The benefit which these fields gain from mutual exchange of ideas, methods and approaches is, however, not always recognized and accepted by conventional electrochemists who stand against risky hybridization of electrochemistry with “material science”.19 It must be admitted that this view is not entirely without merit. Chemists still poorly understand the (electro) chemical reactions at the surface of electrocatalysts compared to their bulk properties, thus oversimplifying the electrocatalytic reactions confining them to mechanistic descriptors,20,21 whereas the actual catalytic activity is often related to totally different structures and compositions arising at the surface through a dynamical interaction with the electrolyte.22–25 This puts particular emphasis on mastering the electrochemical experiments followed by correct interpretation of the results based on rigorous use of rich theoretical background of electrochemistry.26–28 In order to optimize the properties of the solid state electrode materials, inorganic electrochemistry utilizes the approaches similar to those which have been widely applied in solid state inorganic chemistry and physics for high temperature superconducting cuprates and colossal magnetoresistive manganites. The important difference, however, remains. Most commonly, in cuprates and manganites the electronic band filling has been varied mostly by heterovalent substitution at the electronically inactive sites which play a role of charge reservoir, such as La2-xAxCuO4 or La1-xAxMnO3 (Addivalent alkali-earth element). This way the electron/hole concentration was changed in a controlled manner resulting in optimal charge carrier doping to maximize the critical temperature of superconducting transition,29 or metal/insulator behavior was changed along with charge and/or magnetic ordering by enabling transition metal d-orbitals of appropriate symmetry to participate in exchange interactions.30,31 This approach is, however, not very common in the intercalation-type cathode materials for metal-ion batteries. For instance, in the rock-salt-type layered LiMO2 oxides (Mdtransition metal), the 2Co3þ / Ni2þ þ Mn4þ substitution does not formally change the overall number of electrons available for charge compensation of Li deintercalation as two Co3þ/Co4þ redox pairs are replaced with the Ni2þ/Ni3þ and Ni3þ/Ni4þ pairs. At the same time, the electrochemically inactive Mn4þ cations provide low-lying d-states capable of tight binding with the oxygen p-orbitals thus increasing the structure stability and enlarging the range of accessible Li concentrations.6 In the olivine-type LiFePO4 cathode the Mn2þ and/or Co2þ for Fe2þ substitution also does not change the formal electron count, but creates additional d-bands which are lower in energy than the d-band of Fe2þ that increases the Liþ deintercalation potential and energy density.32 Thus, the doping goal is merely confined to creating extra energy levels rather than changing the electron concentration in the existing ones. In the electrocatalysts of the oxygen evolution and oxygen reduction reactions the main goal of doping confines to tuning both the electron count and cationic d-band positioning with respect to the anionic p-band, at the same time

Introduction: Inorganic electrochemistry

3

maintaining good electric conductivity and reasonable stability towards the electrolyte at the conditions of the electrochemical reaction.33 The catalysts comprising two transition metals often exceed in activity the catalysts containing these metals separately due to a formation of dynamically stable active-site/host pair (for instance, Fe/Ni) at the electrode/electrolyte interface.34 Mixing anions instead of cations also became a fruitful strategy in inorganic chemistry, resulting in a range of materials, such as superconductors, (photo)catalysts, thermoelectrics, ferroelectrics, topological insulators.35 In inorganic electrochemistry mixing anions has been widely employed for altering the covalency/ionicity of bonding between electroactive transition metals and anions in order to increase the Mn þ/M(n þ 1)þ redox potential, so as the compounds traditionally considered for negative electrode (anode, with low potential vs. A/Aþ, Adalkali metal) can be turned to a positive electrode material (cathode, with moderately high potential vs. A/Aþ).36 Breaking local symmetry around transition metal cation by introducing two different anions into its coordination environment triggers involvement of the anion p-orbitals into the redox reaction by switching the nature of the redox-active electronic states from localized non-hybridized lone pairs to the hybridized mixed d-p states near the Fermi level.37 Additional benefit of mixing anions comes from expanding structural diversity of intercalation-type cathodes by getting inspiration from the mineral world, in which Nature provides the guides for chemists (such as carbonates-phosphates sidorenkite and bonshtedtite or fluoridesulfates tavorite and triplite).38–41 Although the structural and electronic re-organizations inside the electroactive materials are vitally important for certain electrochemical applications, the processes at the electrode/electrolyte interface should not be neglected. Inorganic electrochemistry greatly benefits from mastering various approaches for surface modification, which are, however, not necessarily nesting directly in the nature of the electrochemical reactions. Conductive coatings (most often made from carbon) are necessary for enabling the polyanion cathode materials (i.e., phosphates, sulfates, silicates, borates etc.) to deliver their anticipated energy and power, as these materials are typically good electrical insulators.42 Such coatings promote transport of electrons between the cathode particles, but, from another side, might block the alkali cation transfer through the electrode/electrolyte interface.43 Protective coatings on the layered oxide Li(Ni,Mn,Co)O2 cathodes for Li-ion batteries prevent parasitic redox reactions with organic electrolyte at high charge potential and improve safety.44,45 Decorating the grain boundaries in the agglomerates of the Li(Ni,Mn,Co)O2 nanoparticles with thin layers of Li-ion conductors helps improving their mechanical stability and prolong the cathode lifetime.46,47 Such coatings, possessing reasonable Li-ion conductivity and poor electronic conductivity, might increase the electric resistance of the cathode. Strict requirement of balancing various functionalities poses a challenge for chemists to master the synthesis techniques resulting in the desired chemical composition, thickness and uniformity of the coatings. Surface engineering becomes particularly important in heterogeneous catalysis because the surface of the catalyst serves as the main playground where the (electro)chemical reaction occurs. Except of the straightforward goal of maximizing the specific surface area to increase the number of catalytic centers per unit mass, additional degrees of freedom are nested in identifying the most active crystallographic surfaces, chemical ways of controlling the crystal growth to maximize their presence and stabilizing the stepped vicinal surfaces which carry highly active undercoordinated atomic species.48,49 Motion of ions through solids is also inseparable part of inorganic electrochemistry being important for both intercalation-type electroactive materials and solid electrolytes. Importance of alkali cation diffusion for the metal-ion battery cathodes and anodes is obvious, but, in fact, transition metal cations also do not stand still while the battery goes through charge-discharge cycles. Being only partially reversible, such detrimental movements are responsible for gradual capacity and voltage fade in layered oxide cathodes, particularly those which are enriched with Li.50 Deciphering the complexity of these local structure rearrangements requires crystal chemistry considerations, such as commensurability between the size of anionic cavity and ionic radius of transition metal in a particular oxidation state, appropriateness of certain coordination number and geometry for a given electronic state of transition metal, flexibility and distortion of the oxygen surroundings, which might depend on the involvement of oxygen into the redox reactions.51 Crystal chemistry is a tool to either arrest cation migration, for instance by increasing the difference between the alkali metal and transition metal sites with Na for Li replacement,52 or to improve its reversibility, as exemplified by introducing “hexagonal” layers into the “cubic” close packing of the oxygen atoms of layered oxide cathodes.53 Suppressing migration barriers for cations or anions in solid electrolytes, depending on application ranging from all-solid-state batteries to solid oxide fuel cells, is a longstanding challenge for chemists. Various crystal chemistry measures have been applied to lower these barriers: selecting the right basic structure with direct hops between sites with low barriers (i.e. body-centered cubic), expansion of the diffusion channels, switching to more polarizable anionic framework (i.e. from oxides to sulfides), coupling cation migration to rotations of polyanionic fragments (so called “paddle-wheel” mechanism) improved alkali-ion conductivity in solid electrolytes. Solid state ionics manipulates concentration and nature of point defects, controls phase transitions via doping to stabilize high-temperature highsymmetry disordered phases with increased ionic conductivity, and pays particular attention to sintering ceramics with optimized properties of grain boundaries and interfaces.54,55 Besides a general requirement for solid electrolyte to be chemically and electrochemically stable towards the electrode materials, every particular application implies additional complications. For instance, in high-temperature solid oxide fuel cells it stems from the necessity to adjust the thermal expansion coefficients of the solid electrolyte and electrodes to maintain the cell integrity.56 In all-solid-state batteries, it is a demand to withstand lithium dendrite growth and propagation.57 The design and manufacturing of the electroactive and electrolyte materials appears to be a great challenge for chemists. Inorganic electrochemistry employs a plethora of advanced synthesis methodsd(co)precipitation at ambient and hydro(solvo)thermal conditions (incl. microwave-assisted hydrothermal synthesis), sol-gel, spray drying and pyrolysis, mechanochemical synthesis, carbothermal reduction and combustion, to name just a few.58 Chemical and electrochemical ion exchange or low-temperature strategies provide a pathway towards metastable compounds which cannot be prepared via a direct reaction. Thin layers are produced

4

Introduction: Inorganic electrochemistry

with dip and spin coating and vacuum-based techniques, such as pulsed laser deposition (PLD) and atomic layer deposition (ALD).59 A selected method of material preparations should provide right chemical composition and crystal structure, chemical homogeneity and phase purity, nanostructuring and agglomeration which assure morphology, surface area, particle size distribution and porosity desired for a particular application. On top of that, more sophisticated approaches are often required to achieve competitive electrochemical performancedsurface coatings, core-shell and gradient structures, composites with modified intergranular and interfacial layers. The synthesis methods targeting mass production must be scalable, reliable, inexpensive, resource- and energy-efficient.60 The collection of the chapters in the Volume 7 reflects, to some degree subjectively, the aforementioned diverse nature of inorganic electrochemistry. Chemistry, crystallography and electrochemistry of positive electrode (cathode) materials for post-lithium energy storage technologies, namely sodium-ion and potassium-ion batteries, are reflected in the chapters by Samarin, Trussov & Fedotov (see chapter 7.03) and Hosaka, Kubota & Komaba (see chapter 7.04), respectively. Polyanion cathode materials for metalion batteries are viewed at two different angles. Singh, Vanam, Lochab, Fichtner & Barpanda provide an overview of the trends in design principles, crystal structures, future challenges and perspectives of polyanion cathodes for sodium-ion batteries, with the particular accent on mixed polyanionic materials (see chapter 7.09). Khasanova, Drozhzhin, Yakubovich & Antipov establish the intriguing connections between minerals and metal-ion battery cathodes (see chapter 7.12). Li & Tarascon take the reader into a journey through a fascinating world of the metal-ion battery cathode materials demonstrating oxidation of the anionic sublattice, the topic which provides excellent opportunity for discussing the role of chemical bonding, crystal and electronic structures in reversibility of the anionic redox and in overcoming various difficulties in its practical realization (see chapter 7.02). Nikitina & Aksenov in their chapter displace the accent towards the surface of the electrodes in metal-ion intercalation system being focused at rigorous consideration of charge transfer across solid/solid and solid/liquid interfaces (see chapter 7.05). Comparative discussion of various computational approaches for screening and computational design of fast metal-ion conductors is given in the chapter by Kabanov, Morkhova, Bezuglov & Blatov (see chapter 7.13). Beyond metal-ion batteries, alternative electrochemical energy storage technologies and related materials are represented in the chapters by Ding, Chagnot, Saeed & Augustyn (see chapter 7.08), Inozemtseva, Rulev, Zakharchenko, Isaev, Yashina & Itkis (see chapter 7.11) and Dominko, Talian & Vizintin (see chapter 7.14). Komkova & Karyakin provide a look at the hexacyanoferrate coordination compounds as the catalysts for biosensors (see chapter 7.06). Simoska, Lee & Minteer survey bioelectrocatalysts and their applications for electrochemical enzymatic biosensors, biofuel cells, and bioelectrosynthesis (see chapter 7.15). Detailed theoretical and practical methodologies of measuring, analyzing and reporting electrocatalytic activities while avoiding common pitfalls are given in the chapter by Savinova & Oshchepkov (see chapter 7.16). Introduction into solution-based synthesis techniques of electrode materials and electrolytes is provided by De Sloovere, Joos, Ulu, Mylavarapu, Kelchtermans, Bolia, Vranken, Paulus, Van Bael & Hardy (see chapter 7.07). Finally, diffraction-based operando investigation techniques and transmission electron microscopy studies on electrode materials are outlined in the chapter by Orlova, Morozov & Abakumov (see chapter 7.10). Enjoy reading these stories as they represent forefront efforts to fully understand the complexities of such inorganic electrochemical systems!

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24.

Zanello, P. Inorganic Electrochemistry: Theory, Practice and Application, The Royal Society of Chemistry, 2003. Reedijk, J., Poeppelmeier, K., Eds.; Comprehensive Inorganic Chemistry II, Elsevier, 2013. The Royal Swedish Academy of Sciences. The Nobel Prize in Chemistry 2019. Press Release, The Royal Swedish Academy of Sciences, 2019. Whittingham, M. S. Chem. Rev. 2004, 104, 4271–4301. Manthiram, A. Nat. Commun. 2020, 11, 1550. Goodenough, J. B.; Kim, Y. Chem. Mater. 2010, 22, 587–603. Melot, B. C.; Tarascon, J.-M. Acc. Chem. Res. 2013, 46, 1226–1238. Abakumov, A. M.; Fedotov, S. S.; Antipov, E. V.; Tarascon, J.-M. Nat. Commun. 2020, 11, 4976. Lin, F.; Liu, Y.; Yu, X.; Cheng, L.; Singer, A.; Shpyrko, O. G.; Xin, H. L.; Tamura, N.; Tian, C.; Weng, T.-C.; Yang, X.-Q.; Meng, Y. S.; Nordlund, D.; Yang, W.; Doeff, M. M. Chem. Rev. 2017, 117, 13123–13186. Wu, Y.; Liu, N. Chem 2018, 4, 438–465. Yuan, Y.; Amine, K.; Lu, J.; Shahbazian-Yassar, R. Nat. Commun. 2017, 8, 15806. Balagurov, A. M.; Bobrikov, I. A.; Samoylova, N. Y.; Drozhzhin, O. A.; Antipov, E. V. Russ. Chem. Rev. 2014, 83, 1120–1134. Pecher, O.; Carretero-González, J.; Griffith, K. J.; Grey, C. P. Chem. Mater. 2017, 29, 213–242. Meutzner, F.; Nestler, T.; Zschornak, M.; Canepa, P.; Gautam, G. S.; Leoni, S.; Adams, S.; Leisegang, T.; Blatov, V. A.; Meyer, D. C. Phys. Sci. Rev. 2018, 4, 20180044. Brown, I. D. Chem. Rev. 2009, 109, 6858–6919. Fedotov, S. S.; Kabanova, N. A.; Kabanov, A. A.; Blatov, V. A.; Khasanova, N. R.; Antipov, E. V. Solid State Ion. 2018, 314, 129–140. Urban, A.; Seo, D.-H.; Ceder, G. NPJ Comput. Mater. 2016, 16002. Schmidt, J.; Marques, M. R. G.; Botti, S.; Marques, M. A. L. NPJ Comput. Mater. 2019, 5, 83. Tsirlina, G. A. J. Solid State Electrochem. 2020, 24, 2187–2188. Hong, W. T.; Risch, M.; Stoerzinger, K. A.; Grimaud, A.; Suntivich, J.; Shao-Horn, Y. Energ. Environ. Sci. 2015, 8, 1404–1427. Song, J.; Wei, C.; Huang, Z.-F.; Liu, C.; Zeng, L.; Wang, X.; Xu, Z. J. Chem. Soc. Rev. 2020, 49, 2196–2214. Baeumer, C.; Li, J.; Lu, Q.; Liang, A. Y.-L.; Jin, L.; Martins, H. P.; Duchon, T.; Glöß, M.; Gericke, S. M.; Wohlgemuth, M. A.; Giesen, M.; Penn, E. E.; Dittmann, R.; Gunkel, F.; Waser, R.; Bajdich, M.; Nemsák, S.; Tyler Mefford, J.; Chueh, W. C. Nat. Mater. 2021, 20, 674–682. Shen, T.-H.; Spillane, L.; Vavra, J.; Pham, T. H. M.; Peng, J.; Shao-Horn, Y.; Tileli, V. J. Am. Chem. Soc. 2020, 142, 15876–15883. Lopes, P. P.; Chung, D. Y.; Rui, X.; Zheng, H.; He, H.; Farinazzo Bergamo Dias Martins, P.; Strmcnik, D.; Stamenkovic, V. R.; Zapol, P.; Mitchell, J. F.; Klie, R. F.; Markovic, N. M. J. Am. Chem. Soc. 2021, 143, 2741–2750.

Introduction: Inorganic electrochemistry 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60.

5

Samira, S.; Hong, J.; Camayang, J. C. A.; Sun, K.; Hoffman, A. S.; Bare, S. R.; Nikolla, E. JACS Au 2021, 1, 2224–2241. Talaie, E.; Bonnick, P.; Sun, X.; Pang, Q.; Liang, X.; Nazar, L. F. Chem. Mater. 2017, 29, 90–105. Nikitina, V. A.; Vassiliev, S. Y.; Stevenson, K. J. Adv. Energy Mater. 2020, 1903933. Nikitina, V. A. Curr. Opin. Electrochem. 2021, 29, 100768. Rybicki, D.; Jurkutat, M.; Reichardt, S.; Kapusta, C.; Haase, J. Nat. Commun. 2016, 7, 11413. Tokura, Y.; Nagaosa, N. Science 2000, 288, 462–468. Salamon, M. B.; Jaime, M. Rev. Mod. Phys. 2001, 73, 583–628. Zhou, F.; Cococcioni, M.; Kang, K.; Ceder, G. Electrochem. Commun. 2004, 6, 1144–1148. Grimaud, A.; Hong, W. T.; Shao-Horn, Y.; Tarascon, J.-M. Nat. Mater. 2016, 15, 121–126. Chung, D. Y.; Lopes, P. P.; Martins, P. F. B. D.; He, H.; Kawaguchi, T.; Zapol, P.; You, H.; Tripkovic, D.; Strmcnik, D.; Zhu, Y.; Seifert, S.; Lee, S.; Stamenkovic, V. R.; Markovic, N. M. Nat. Energy 2020, 5, 222–230. Kageyama, H.; Hayashi, K.; Maeda, K.; Attfield, J. P.; Hiroi, Z.; Rondinelli, J. M.; Poeppelmeier, K. R. Nat. Commun. 2018, 9, 772. Fedotov, S. S.; Luchinin, N. D.; Aksyonov, D. A.; Morozov, A. V.; Ryazantsev, S. V.; Gaboardi, M.; Plaisier, J. R.; Stevenson, K. J.; Abakumov, A. M.; Antipov, E. V. Nat. Commun. 2020, 11, 1484. Leube, B. T.; Robert, C.; Foix, D.; Porcheron, B.; Dedryvère, R.; Rousse, G.; Salager, E.; Cabelguen, P.-E.; Abakumov, A. M.; Vezin, H.; Doublet, M.-L.; Tarascon, J.-M. Nat. Commun. 2021, 12, 5485. Hautier, G.; Jain, A.; Chen, H.; Moore, C.; Ong, S. P.; Ceder, G. J. Mater. Chem. 2011, 21, 17147–17153. Ati, M.; Sathiya, M.; Boulineau, S.; Reynaud, M.; Abakumov, A.; Rousse, G.; Melot, B.; Van Tendeloo, G.; Tarascon, J.-M. Understanding and Promoting the Rapid Preparation of the Triplite-Phase of LiFeSO4F for Use as a Large-Potential Fe Cathode. J. Am. Chem. Soc. 2012, 134, 18380–18387. Senthilkumar, B.; Murugesan, C.; Sharma, L.; Lochab, S.; Barpanda, P. Small Methods 2018, 1800253. Yakubovich, O.; Khasanova, N.; Antipov, E. Minerals 2020, 10, 524. Masquelier, C.; Croguennec, L. Chem. Rev. 2013, 113, 6552–6591. Iarchuk, A. R.; Nikitina, V. A.; Karpushkin, E. A.; Sergeyev, V. G.; Antipov, E. V.; Stevenson, K. J.; Abakumov, A. M. ChemElectroChem 2019, 6, 5090–5100. Wang, K.-X.; Li, X.-H.; Chen, J.-S. Adv. Mater. 2015, 27, 527–545. Lee, K. T.; Jeong, S.; Cho, J. Acc. Chem. Res. 2013, 46, 1161–1170. Yan, P.; Zheng, J.; Liu, J.; Wang, B.; Cheng, X.; Zhang, Y.; Sun, X.; Wang, C.; Zhang, J.-G. Nat. Energy 2018, 3, 600–605. Savina, A. A.; Orlova, E. D.; Morozov, A. V.; Luchkin, S. Y.; Abakumov, A. M. Nanomaterials 2020, 10, 2381. Hammer, B.; Nørskov, J. K. Adv. Catalysis 2000, 45, 71–129. Hendriksen, B. L. M.; Ackermann, M. D.; van Rijn, R.; Stoltz, D.; Popa, I.; Balmes, O.; Resta, A.; Wermeille, D.; Felici, R.; Ferrer, S.; Frenken, J. W. M. Nat. Chem. 2010, 2, 730–734. Sathiya, M.; Abakumov, A. M.; Foix, D.; Rousse, G.; Ramesha, K.; Saubanère, M.; Doublet, M. L.; Vezin, H.; Laisa, C. P.; Prakash, A. S.; Gonbeau, D.; Van Tendeloo, G.; Tarascon, J.-M. Nat. Mater. 2015, 14, 230–238. Gent, W. E.; Lim, K.; Liang, Y.; Li, Q.; Barnes, T.; Ahn, S.-J.; Stone, K. H.; McIntire, M.; Hong, J.; Song, J. H.; Li, Y.; Mehta, A.; Ermon, S.; Tyliszczak, T.; Kilcoyne, D.; Vine, D.; Park, J.-H.; Doo, S.-K.; Toney, M. F.; Yang, W.; Prendergast, D.; Chueh, W. C. Nat. Commun. 2017, 8, 2091. Wang, Q.; Mariyappan, S.; Rousse, G.; Morozov, A. V.; Porcheron, B.; Dedryvère, R.; Wu, J.; Yang, W.; Zhang, L.; Chakir, M.; Avdeev, M.; Deschamps, M.; Yu, Y.-S.; Cabana, J.; Doublet, M.-L.; Abakumov, A. M.; Tarascon, J.-M. Nat. Mater. 2021, 20, 353–361. Eum, D.; Kim, B.; Kim, S. J.; Park, H.; Wu, J.; Cho, S.-P.; Yoon, G.; Lee, M. H.; Jung, S.-K.; Yang, W.; Seong, W. M.; Ku, K.; Tamwattana, O.; Park, S. K.; Hwang, I.; Kang, K. Nat. Mater. 2020, 19, 419–427. Famprikis, T.; Canepa, P.; Dawson, J. A.; Saiful Islam, M.; Masquelier, C. Nat. Mater. 2019, 18, 1278–1291. Tuller, H. Ionic Conduction and Applications. In Springer Handbook of Electronic and Photonic Materials; Kasap, S., Capper, P., Eds., Springer International Publishing AG, 2017; pp 247–266. Istomin, S. Y.; Antipov, E. V. Russ. Chem. Rev. 2013, 82, 686–700. Liu, X.; Garcia-Mendez, R.; Lupini, A. R.; Cheng, Y.; Hood, Z. D.; Han, F.; Sharafi, A.; Idrobo, J. C.; Dudney, N. J.; Wang, C.; Ma, C.; Sakamoto, J.; Chi, M. Nat. Mater. 2021, 20, 1485–1490. Marchal, W.; De Sloovere, D.; Daenen, M.; Van Bael, M. K.; Hardy, A. Chem. A Eur. J. 2020, 26, 9070–9083. Balaish, M.; Gonzalez-Rosillo, J. C.; Kim, K. J.; Zhu, Y.; Hood, Z. D.; Rupp, J. L. M. Nat. Energy 2021, 6, 227–239. Larcher, D.; Tarascon, J.-M. Nat. Chem. 2015, 7, 19–29.

7.02

Status of Li(Na)-based anionic redox materials for better batteries

Biao Li and Jean-Marie Tarascona,b,c, a Chimie du Solide-Energie, UMR 8260, Collège de France, Paris, France; b Réseau sur le Stockage Electrochimique de l’Energie (RS2E), Amiens, France; and c Sorbonne Université, Paris, France a,b

© 2023 Elsevier Ltd. All rights reserved.

7.02.1 7.02.2 7.02.2.1 7.02.2.2 7.02.2.3 7.02.2.4 7.02.2.5 7.02.2.6 7.02.2.7 7.02.2.8 7.02.3 7.02.3.1 7.02.3.2 7.02.3.3 7.02.3.4 7.02.3.5 7.02.3.6 7.02.4 7.02.4.1 7.02.4.2 7.02.4.3 7.02.4.4 7.02.5 7.02.5.1 7.02.5.2 7.02.5.3 7.02.6 Acknowledgment References

Introduction The journey of anionic redox chemistry The birth of insertion chemistry based on dichalcogenides Anion–cation redox competition: Ligand hole chemistry and anion polymerization Oxygen redox activity in LiCoO2 Li2MnO3-based compounds Model Li-rich or Na systems Theoretical progresses Practical issues: Sluggish kinetics, voltage hysteresis, and voltage fade Sulfides: There and back again Fundamentals behind anionic redox Band structure descriptions Anionic redox activity from O 2p “NB states” or “lone-pair states” Charge-transfer (D) vs. Mott–Hubbard (U) classification Reductive coupling mechanism Anionic activity from O (2p)-M(nd) p-type interaction The nature of oxidized O2
: Electron holes, O-O dimers, trapped O2 molecules, and oxygen release Anionic redox opening a new rich materials chemistry Increasing the Li/M and O/M ratio in layered rock-salt compounds Playing with the alkaline ion: From Li to Na Ligand manipulation: From oxides to sulfides/selenides Cation disorder, superstructure, and structural dimensionality Practical issues and their fundamental understandings Chemical and electrochemical irreversibility Voltage fade Voltage hysteresis Conclusions and outlook

6 8 8 8 9 11 12 12 13 14 15 15 16 16 17 18 19 21 21 23 26 28 30 30 33 36 38 40 40

Abstract For decades, Li-ion batteries are serving as a transformative technology in propelling off-network energy use as exemplified by the prosperity of portable electronics and electric vehicles. However, they are suffering from the shortage of energy density for modern-society energy storage, as the conventional way of Li-ion batteries work solely relies on redox reaction of transition metal (cationic redox) upon Li removal/uptake. This situation has been changed since the emergence of anionic redox, i.e., anions is redox center, that offers extra capacity in Li-rich electrode materials. In this chapter, we aim to briefly introduce how anionic redox chemistry concept rises and revolutionizes the cathode materials design in Li(Na)-ion batteries. We revisit the fundamental science behind anionic redox, and indicate how these knowledge opened new dimensionalities of electrode material design. We also summarize the progresses in understanding and solving the practical roadblocks for anionic redox, and highlight the remaining issues deserving to be explored in the future.

7.02.1

Introduction

Since the Industrial Revolution in the 18th century, the wide-scale use of fossil fuel has largely reshaped people’s life and speeded up the world’s economic pace. However, simultaneously, it also brings climate change and sustainability issues that arouse many sociological and political disputes. This enforces a transformative change in the way of exploiting the energy from the nature for human beings, and more specifically, a revolution from non-renewable fossil fuel to renewable energy sources like solar, wind, bioenergy etc. The intermittent nature of renewable energy makes it essential to develop energy storage solutions to maximize the energy usage

6

Comprehensive Inorganic Chemistry III, Volume 7

https://doi.org/10.1016/B978-0-12-823144-9.00084-4

Status of Li(Na)-based anionic redox materials for better batteries

7

efficiency. On the other hand, the intelligentization of modern society results in massive demand of off-network energy use, such as portable electronics and electrical vehicles that can only be satisfied by cost-effective and safe energy storage technologies. The Li-ion technology is the one that now dominates the market due to its competitive energy density outperforming its counterparts. Having been granted with the Nobel prize in 2019, Li-ion batteries has just reached its 30th year since its first commercialization in 1991 by Sony and Asahi Kasei companies. Though being a younger technology than its predecessors, it totally revolutionized the society since its emergence and its impact never ceased increasing even nowadays. However, the current available Li-ion technology is still far from meeting the demand for a fast-developing society in the quest of higher energy density to power the electrical vehicles. The main limitation comes from cathode side for which the current available commercial materials (LiCoO2, LiFePO4, and LiMn2O4 etc.) are suffering from low specific capacities (< 180 mA h g
1) and therefore low energy densities.1,2 For long time, the community is putting enormous efforts in searching for new electrode materials with capacities surpassing those aforementioned, with, however, little progress. The difficulty is rooted in the way of how rechargeable Li-ion batteries work, as shown by the schematic in Fig. 1A, that Li ions deintercalate from cathode host and shuttle to anode host during charge

Fig. 1 (A) Schematic for the working principle of Li-ion batteries, wherein the redox reactions involving both cations and anions are highlighted with zoomed-in views. (B and C) Show the schematic crystal structures for conventional layered oxides and Li-rich layered oxides, respectively, with the electrochemical curves of their representatives, i.e., LiCoO2 and Li-rich NMC, shown in (D) and (E) correspondingly. (B–E) reproduced from Assat, G.; Tarascon, J.-M., Fundamental Understanding and Practical Challenges of Anionic Redox Activity in Li-Ion Batteries. Nat. Energy 2018, 3, (5), 373–386.

8

Status of Li(Na)-based anionic redox materials for better batteries

(contrary for discharge), with concerted electron transfer through external circuit. Associated with every Liþ extraction is the removal of one electron from the electrode (green arrows in Fig. 1A) to maintain charge neutrality with this process, that constitutes the groundwork of any electrochemical system, being termed as redox reaction. The redox center of the electrodes that either donates or uptakes the electron plays a crucial role since it determines the basic properties of electrochemical reactions, such as reaction potential, kinetics, and most importantly, the capacity. Conventionally, the redox center for cathode electrodes solely relies on cationic redox processes (i.e., transition metal (M)) (Fig. 1A, cationic redox) involving limited electron transfer and hence limited capacity, as exemplified by the conventional layered LiCoO2 (Fig. 1B and D).3 A great change has been made with the discovery of the so-called “Li-rich” compounds (Fig. 1C), in which the anions were also found to be redox active (Fig. 1A, anionic redox). Taking advantage of this anionic redox paradigm, a capacity as high as 300 mA h g
1 can be achieved in the famous Li[Li0.2Ni0.13Mn0.54Co0.13]O2 compound (Fig. 1E), called as Li-rich NMC. This finding has received worldwide resonance with the commitment of several research groups to be fully engaged in exploring anionic redox chemistry. In the recent decade, myriad compounds exhibiting anionic redox were proposed, spanning over not only Li-ion batteries, but also to Na-ion and other beyond Li-ion technologies. These findings have fertilized the field of cathode materials on both fundamental and applied levels. Therefore, it is now timely to sum up the overall anionic redox research as it is done in this chapter with emphasis on recalling main achievements and identifying the main remaining challenges. This review is structured as follows. It will start with a brief introduction on how anionic redox was proposed and how it evolved with the key advances mentioned. Afterwards, fundamental understanding of anionic redox as deduced from numerous published reports and perspectives, such as the typical band-structure descriptions, O 2p non-bonding (NB) states, charge-transfer (D) vs. Mott–Hubbard (U) classification, and so on will be presented. Next an illustration of the richness of the structural chemistry unlocked by anionic redox relying on the engineering of lone-pair states (increasing Li/M and O/M ratio), manipulating ligands (oxides to sulfides), playing with alkaline ions (from Li to Na) and characteristic structures (e.g., cation disordering and superstructure) will be given. Lastly, practical issues associated with unusual electrochemistry of anionic redox, including chemical and electrochemical irreversibility, voltage hysteresis and voltage decay, are summarized and fundamentally discussed. A commentary of the unresolved puzzles of anionic redox that remain to be further explored will close this review together with proposed perspectives on how to narrow the gap between fundamental understandings and practical utilizations. Overall, we intend this chapter to serve as a pedagogical introduction for laymen or beginners in this subject, but also to arouse interests and curiosity for those coming from other related areas so as to broaden the field of anionic redox beyond batteries.

7.02.2

The journey of anionic redox chemistry

The past decade had witnessed the booming of anionic redox as the cutting-edge research for the cathode materials of Li-ion batteries. Though the formal conceptualization of “anionic redox” was completed in recent years, the history of similar concept can be traced back to the end of 20th century. In the following, we give an overview of the several periods in which anionic redox has emerged and largely developed, as navigated by the subtitles.

7.02.2.1

The birth of insertion chemistry based on dichalcogenides

Insertion chemistry lays the foundation of Li-ion batteries and is crucial to the current success that Li-ion technology has achieved. In 1970s, Whittingham proposed the use of TiS2 as the cathode Li host based on a reversible Li intercalation/deintercalation chemistry4 (Fig. 2A), opening up the era of rechargeable Li batteries. Adopting a hexagonal close-packed sulfur slabs with Ti ions residing alternatively in between, TiS2 was an ideal Li host that can accommodate Li ions at the octahedral sites in between TiS2 slabs (Fig. 2B and C). Since then, a lot of similar layered dichacogenides or trichalcogenides, represented by MoS2, VSe2, TiS3, and so on,5–7 were found to have similar Li intercalation activity. Although the redox chemistry of these chalcogenides was not investigated at that time by the battery community, it did attract some physicists since some of these materials were showing exciting phenomena such as superconductivity, charge density waves or 2D magnetism. At the beginning, it was assumed that guest ion intercalation into layered dichalcogenides would not cause significant change to the electronic energy levels of the ions in the host. This was then disputed by J. V. McCanny, in 1979, who conducted first principle calculations on Li insertion into TiS2, showing that both titanium 4s and sulfur 3p bands are affected upon Li uptake.8 Yet the fundamental redox mechanism remains obscure due to the immaturity of experimental techniques.

7.02.2.2

Anion–cation redox competition: Ligand hole chemistry and anion polymerization

The interests of layered chalcogenides were extended to 1990s with numerous studies being implemented. Among them are the works led by J. Rouxel who proposed the concept of anion-cation redox competition in transition metal chalcogenides.9–11 The general idea, schematically shown in Fig. 3, proposed that the early transition metal based chalcogenides (e.g., ZrS2) have separated ligand sp and M nd bands, resulting in more ionic systems with valence electrons filling mostly the anion sp states.9 Moving from the left to the right hand of periodic table, the anion sp band will rise in energy and finally pour its sp electrons into empty M nd band. This leaves holes in the anion sp bands with the overall process termed as “ligand-hole chemistry”.9 Moreover, by moving too far to

Status of Li(Na)-based anionic redox materials for better batteries

9

(A)

Voltage vs. Li/Li+

2.8 2.6

Delithia

2.4

tion

2.2 2.0

Lithiati

1.8

on

1.6 1.4 0.0

0.2

TiS2

(B)

0.4

0.6

0.8

1.0

x in LixTiS2

LiTiS2

(C)

Li Ti

c

c

S b a Fig. 2

a

b

(A) Reversible Li (de)intercalation chemistry in TiS2. The crystal structure of TiS2 and LiTiS2 are shown in (B) and (C), respectively.

the right-hand of periodic table, anions will form dimers to stabilize the electron holes created due to this anion-cation redox competition.10 These pioneering understandings of the redox competition process in these layered chalcogenides, although not fully rationalized theoretically, partially shape the idea of anionic redox chemistry in batteries prevailing nowadays.

7.02.2.3

Oxygen redox activity in LiCoO2

It was several years later after the finding of dichalcogenides that the community started turning their attention towards oxides in pursuit of higher potential and hence energy density for cathode materials of Li-ion batteries. LiCoO2 was first recognized by Goodenough et al.12 in 1980 to be a good candidate of lithium cathode host as it bears similar layered structure with dichalcogenides,

Fig. 3 Anion-cation redox competition in layered transition metal chalcogenides. From the left to right panel, the rising of anionic sp band will finally results in the electron transfer from anions to empty transition metal nd band. Reproduced from Rouxel, J. Anion-Cation Redox Competition and the Formation of New Compounds in Highly Covalent Systems. Chem. A Eur. J. 1996, 2, 1053–1059.

10

Status of Li(Na)-based anionic redox materials for better batteries

except that it is in a lithiated state. It adopts a a-NaFeO2 structure with a cubic close-packed oxygen framework, as shown in Fig. 4A, that Li ions are sandwiched in between CoO2 slabs. The redox process of LiCoO2 was initially supposed to be Co3þ/Co4þ based on the high open circuit voltage of 4–5 V for LixCoO2/Li cell observed.12 However, it came out as a surprise that, after the end member CoO2 was successfully prepared electrochemically in 1996,13 Co was not found fully oxidized to 4 þ as inferred from magnetic studies,14 hence interrogating the possibility of oxygen involvement. Shortly after that, Tarascon et al. found the shortening of O-O distance in CoO2 (Fig. 4B) by synchrotron X-ray diffractions and associated it with the oxygen electron-hole chemistry,15 in line with what J. Rouxel proposed before. In that case, the redox of LiCoO2 was proposed consisting of a first step of Co3þ oxidation followed by second step of oxygen oxidation, leading to polymerization of oxygen framework in order to stabilize the CoO2 structure. The above experimental finding was also theoretically supported in earlier year by Ceder et al. with first-principles

Fig. 4 Crystal structure of (A) LiCoO2 and its delithiated end member (B) CoO2. The O-O distance within transition metal layer is marked out by dashed line. (C) Positive charge density difference map between Li(Al0.33Co0.67)O2 and (Al0.33Co0.67)O2 in a plane perpendicular to the direction of layering in the structure. Darker indicates more charge density. (D) O 1s XPS results of LixCoO2, CoO2 and CoO2 in special experimental conditions (charge neutralization) that allow to shift the components of oxygenated species formed at the interface with respect to the CoO2 component itself. It clearly demonstrates an oxidized oxygen species whose peak is located at the binding energy of 530.2 eV. (E) Normalized oxygen K-edge XAS of Li1
xCoO2. Figures adapted from: (C) Ceder, G.; Chiang, Y. M.; Sadoway, D. R.; Aydinol, M. K.; Jang, Y. I.; Huang, B., Identification of Cathode Materials for Lithium Batteries Guided by First-Principles Calculations. Nature 1998, 392 (6677), 694–696; (D) Dahéron, L.; Dedryvère, R.; Martinez, H.; Ménétrier, M.; Denage, C.; Delmas, C.; Gonbeau, D., Electron Transfer Mechanisms Upon Lithium Deintercalation From LiCoO2 to CoO2 Investigated by XPS. Chem. Mater. 2007, 20(2), 583–590; (E) Yoon, W.-S.; Kim, K.-B.; Kim, M.-G.; Lee, M.-K.; Shin, H.-J.; Lee, J.-M., Oxygen Contribution on Li-Ion Intercalation-Deintercalation in LiAlyCo1
yO2 Investigated by O K-Edge and Co L-Edge X-Ray Absorption Spectroscopy. J. Electrochem. Soc. 2002, 149(10), A1305.

Status of Li(Na)-based anionic redox materials for better batteries

11

calculations, showing that, by charge density difference analysis (Fig. 4C), electron is transferred to both metal and anions upon Li intercalation in LiCoO2 and especially Al-doped LiCoO2.16 In the following decade, more experimental efforts were devoted to unveiling the oxygen redox activity in LiCoO2 via the development of sophisticated spectroscopic techniques. As displayed in Fig. 4D and E, it become possible to successfully deconvolute O(2
n)
species out of O 1s X-ray photoelectron spectroscopy (XPS),17 or spot the broadening of pre-edge peak of O K-edge Xray absorption spectroscopy (XAS),18 hence supporting early predictions that oxygen is oxidized in LixCoO2 with electron holes being generated on O 2p band. In light of these findings, the structural instability of highly delithiated LiCoO2 could be understood by the release of O2 gas associated to large amount of holes created at O 2p states, hence explaining the encountered experimental difficulty in isolating pure CoO2. However, although concrete evidences were established for the anionic redox in LiCoO2, the underlying origin of its activity is intrinsically different from that of Li-rich compounds, as discussed in the fundamental section (Section 7.02.3).

7.02.2.4

Li2MnO3-based compounds

The anionic redox research in LiCoO2 went dim for a while as such mechanism didn’t bring any practical assets (e.g., high capacity) but only challenges associated with oxygen release and structural destabilization. However, the discovery of the so-called Li-rich layered oxides had rejuvenated the research of anionic redox as it enables abnormal reversible high capacity. This was first initiated by J. Dahn et al., who led the investigation on Ni-substituted Li2MnO3 layered compounds, Li[Li(1/3–2x/3)NixMn(2/3–x/3)]O2, which delivered anomalous 220 mA h g
1 capacity that cannot be solely covered by cationic redox.19 Its parental structure, Li2MnO3, either written as Li[Li1/3Mn2/3]O2, shares a similar layered structure with LiCoO2, but differs in a way that there is additional 1/3 Li ions replacing transition metal in Mn layer forming a honeycomb superstructure (Fig. 5A and B)dcharacteristic of socalled Li-rich layered oxides. A great deal of efforts was then heaped on substituting Li2MnO3 with Ni, Co, Fe, Cr etc. or their

Fig. 5 (A) Crystal structure of Li2MnO3 with Li layer and M layer being indicated. (B) Top view of M layer in Li2MnO3, with honeycomb ordering. (C) Reported electrochemistry of Li2MnO3 at different temperatures. (D) First-cycle electrochemistry of Li-rich NMC (Li1.2Ni0.13Mn0.54Co0.13O2), showing a staircase curve on charge but an S-shape curve on discharge. Reproduced from: (C) Robertson, A. D.; Bruce, P. G., The Origin of Electrochemical Activity in Li2MnO3. Chem. Commun. 2002, (23), 2790–2791.

12

Status of Li(Na)-based anionic redox materials for better batteries

combinations, resulting in a wide family of derivatives.20,21 Among them, one series of compounds, with compositions similar to Li1.2Ni0.13Mn0.54Co0.13O2 that termed as Li-rich NMCs, received enormous attentions due to its high-capacity over 250 mA h g
1. The success in Li2MnO3-based compounds leads to the interrogation of where the redox activity comes from. The early efforts trying to answer this question were targeted on the parental Li2MnO3 structure due to its simpler structure. The electrochemical activity of Li2MnO3 was successfully activated by the reducing particle size or cycling under higher temperature,22,23 as witnessed by the appearance of a plateau at around 4.5–4.6 V in first charge followed by an S-shape discharge curve (Fig. 5C). However, the redox mechanism of its charging process turned out to be more intricate than expected, as reflected by various proposed scenarios, such as Mn4þ/Mn5þ, Liþ/Hþ exchange, oxygen release, reversible oxygen redox, and so on.23–30 Fortunately, the situation has settled down and it is nearly commonly believed that Li extraction from Li2MnO3 is compensated by oxygen loss rather than reversible anionic redox, as evidenced by resonant inelastic X-ray scattering (RIXS) as well as operando mass spectroscopy.29,30 The signatures of 4.5 V first charging plateau of Li2MnO3 was almost entirely inherited by Ni/Co substituted derivatives, the Lirich NMC compounds, which, however, have a staircase type curve on charge (Fig. 5D) due to the presence of additional cationic redox. Yet the discharge also shows a similar S-shape curve with Li2MnO3, hence triggering a same type of debates about the redox mechanism of Li-rich NMC compounds.26,31–40 Alike before, after a long journey a consensus has been reached and it is now widely admitted that the abnormal high capacity delivered by these Li-rich NMC compounds originates from cumulated cationic and anionic redox activities. This agreement was established on some antecedent fundamental understandings of several model systems, as will be expanded next.

7.02.2.5

Model Li-rich or Na systems

Inspired by the unusual electrochemistry of Li2MnO3-based compounds, researchers started seeking for similar but simpler compounds so as to get deeper insights of the reaction mechanism. This was initially done via 4d or 5d transition metal substitution in Li2RuO3 and Li2IrO3 oxides sharing nearly identical structures with Li2MnO3. More specifically, doping Li2RuO3 with inactive spectator Sn ions led to a Li2Ru1
ySnyO3 solid solutions out of which several members exhibit impressive electrochemical behavior alike that of Li-rich NMC compound (Fig. 6A), both sharing a staircase to S-shape curve transition during first cycle.41 By unambiguously evidencing the reversible anionic redox chemistry via XPS and electron paramagnetic resonance (EPR)41 (Fig. 6B and C), the study of Li2Ru0.75Sn0.25O3 compound hence opened up a new strategy to explore model Li-rich systems aiming at fundamental understandings of anionic redox. These include the successive studies on a-Li2IrO3 for directly visualizing O-O dimers,42 b-Li2IrO3 for anionic redox activity in tridimensional structure,43 Li2Ir1
ySnyO3 on stabilization of anionic redox through M–O decoordination,44 Li3MO4 system (M ¼ Ru and Ir) in activating more anionic redox activity though the chemical tuning of the M/O ratio,45,46 and so on. Moreover, these fundamental studies enable the anionic redox topic to extend to Na-ion batteries, as represented by efforts on O3-type Na2MO3 (M ¼ Ru, Ir),47,48 Na(Li1/3M2/3)O2 (M ¼ Ir, Mn)49,50 and also P2/P3-type Nadeficient layered compounds,51–53 as well as other structures like Na2Mn3O7.54,55 Overall, most of these aforementioned electrodes were fruitful to get deeper insights towards a better understanding of the anionic redox but of little help for overcoming cost or performances bottlenecks plaguing their real-world applications.

7.02.2.6

Theoretical progresses

The lack of theoretical guidance at the early stage of anionic redox in layered oxides has somewhat limited experimental progresses. This arises from the structural and compositional complexity of typical Li-rich compound (Li-rich NMC) that is even more acute in their fully oxidized states rendering the band structure prediction by density functional theory (DFT) calculations difficult at the early stage. With the evolution of theoretical approaches relying on the DFT þ U framework,56,57 more theoretical works on interpreting redox behaviors springs up. This culminated with the unlocking of various model systems mentioned above showing structural simplicity, hence more appropriated for DFT calculations. Besides, the use of several powerful approaches based on DFT calculations largely aided the experimental findings. These including but not limited to density of states (DOS),58 crystal orbital overlap populations (COOP),59 electron localization function (ELF),60 Fukui function,61 and Bader charge analysis,42 as shown in Fig. 7. An illustrative example of theoretical study in benefiting experimental progress is the so-called “reductive coupling mechanism,”41,59 involving a ligand-to-metal charge transfer behavior that was predicted early in Li2Ru0.75Sn0.25O3 compound, although experimental evidences came later. Besides, the theoretical finding of O 2p NB states (or lone-pair states), aroused by Li-O-Li configurations due to the additional Li ion in M layer, as the source of oxygen redox activity is a crucial step forward.58 It differentiates Lirich materials from conventional layered oxides (e.g., LiCoO2) whose oxygen activities are only from M(nd)–O(2p) antibonding states. This successfully guided the subsequent findings of Mg, Zn, Al, or even cation vacancy in generating O 2p NB states for anionic   redox in sodium compounds. More impressive theoretical progresses are the predictions of the threshold h  1/3 (h denotes elec60 tron holes per oxygen atom) limit of reversible oxygen redox, together with high-valent cationic intermediate species,62,63 metaloxygen p redox process,64 and so on. Owing to these successes, theoretical calculations are becoming more indispensable nowadays than ever for anionic redox related research.

Status of Li(Na)-based anionic redox materials for better batteries

13

Fig. 6 Li2Ru0.75Sn0.25O3 model compound with anionic redox. (A) Charge and discharge curves of Li2Ru0.75Sn0.25O3 with its cycling and rate performance shown as insets. (B) O 1s XPS study showing oxygen-redox activity during cycling. (C) X-band EPR signal recorded at 4K for Li2Ru0.75Sn0.25O3 and its charged states (4 V and 4.6 V). The EPR spectra of CaO2 was included as inset for comparison. The 4.6 V spectra clearly shows a similar peroxo-like species alike CaO2, indicating the oxygen redox activity. Adapted from Sathiya, M.; Rousse, G.; Ramesha, K.; Laisa, C. P.; Vezin, H.; Sougrati, M. T.; Doublet, M. L.; Foix, D.; Gonbeau, D.; Walker, W.; Prakash, A. S.; Ben Hassine, M.; Dupont, L.; Tarascon, J. M., Reversible Anionic Redox Chemistry in High-Capacity Layered-Oxide Electrodes. Nat. Mater. 2013, 12(9), 827–835.

7.02.2.7

Practical issues: Sluggish kinetics, voltage hysteresis, and voltage fade

After the excitement of enjoying a variety of model compounds for fundamental understandings, more efforts are trickling into understanding and solving the practical issues that are hampering the real-world implementation of anionic redox. These issues include the sluggish kinetics, voltage hysteresis, voltage fade and so on, as frequently referred in recent anionic-redox papers. Among them, voltage fade and voltage hysteresis are attracting most of the attentions. Voltage fade, describing a phenomenon with gradual voltage drop and hence energy density fading along with cycling, was first noticed in the Li-rich NMC compound.65,66 Although various perspectives were proposed to account for it, such as spinel-to-layered phase transition,67 microdefects,68 redox evolution,69 and so on, a common consensus is now reached on the fact that irreversible cation migration induced by anionic redox is responsible for voltage fade.70–73 However, the reason why anionic redox ubiquitously induces voltage hysteresis, i.e., voltage gap between charge and discharge, which penalizes the energy efficiency, is still under intensive debate. Although cation migration is frequently being put forward as the reason of voltage hysteresis,71,74–77 it is being challenged by other findings, such as redox inversion,60,78 OO dimerization,79,80 and the recent proposed ligand-to-metal charge transfer.81 The intricacy of voltage hysteresis issue is rooted at the elusive thermodynamics and kinetics behind it, which should be fully mastered before giving an answer. Additionally, anionic redox also causes chemical and electrochemical irreversibilities, such as oxygen release and capacity fading, which are also receiving

14

Status of Li(Na)-based anionic redox materials for better batteries

Fig. 7 Typical approaches in DFT calculations for revealing anionic redox activity. The following figures are shown as examples. (A) DOS calculated for Li2MnO3. (B) COOP of Li2RuO3, LiRuO3 and RuO3 showing the bonding and antibonding states of Ru-O and O-O. (C) ELF of Li2MnO3. The yellow semi-spheres indicate the O 2p lone-pair states. (D) Fukui function of Na2/3Mg1/3Mn2/3O2 upon 1 electron removal (from 24 formula units supercell). (E) Charge transfer calculated from Bader charge for Ir and O during delithiation of a-Li2IrO3. Adapted from: (A) Seo, D. H.; Lee, J.; Urban, A.; Malik, R.; Kang, S.; Ceder, G., The Structural and Chemical Origin of the Oxygen Redox Activity in Layered and Cation-Disordered Li-Excess Cathode Materials. Nat. Chem. 2016, 8(7), 692–697; (B) Saubanère, M.; McCalla, E.; Tarascon, J. M.; Doublet, M. L., The Intriguing Question of Anionic Redox in High-Energy Density Cathodes for Li-ion Batteries. Energ. Environ. Sci. 2016, 9(3), 984–991; (C) Ben Yahia, M.; Vergnet, J.; Saubanere, M.; Doublet, M. L., Unified Picture of Anionic Redox in Li/Na-ion Batteries. Nat. Mater. 2019, 18(5), 496–502; (D) Vergnet, J.; Saubanère, M.; Doublet, M.-L.; Tarascon, J.-M., The Structural Stability of P2-Layered Na-Based Electrodes During Anionic Redox. Joule 2020, 4(2), 420–434; (E) McCalla, E.; Abakumov, A. M.; Saubanere, M.; Foix, D.; Berg, E. J.; Rousse, G.; Doublet, M. L.; Gonbeau, D.; Novak, P.; Van Tendeloo, G.; Dominko, R.; Tarascon, J. M., Visualization of O-O Peroxo-Like Dimers in High-Capacity Layered Oxides for Li-Ion Batteries. Science 2015, 350(6267), 1516–1521.

considerable efforts for understanding the origin and conceiving solutions. The exploration towards these practical issues is still ongoing and progressing.

7.02.2.8

Sulfides: There and back again

Given the challenges imposed by anionic redox, looking for cathodes that are free of above issues is therefore imperative. This has been addressed by various chemical strategies, such as cation doping, surface coating, heterogeneous structure, and concentration gradient, but none of them can eradicate all these issues simultaneously. This highlights the urgent need to develop new high-capacity Li-rich systems based on anionic redox. One promising way is Li-rich sulfides, simply motivated by the fact that the higher positioned and more delocalized sulfur 3p band may stimulate more anionic redox activity. This was exemplified by recent findings of Li2TiS3-based compounds, whose sulfur redox activity can be activated by doping with active M like Co2þ, Fe2þ and even Ti3þ 82–84, by making it cation-disorder or via chemical substitution at both alkaline ion (Li by Na) and the anionic levels (S by Se).85,86 The rejuvenation of sulfides chemistry enjoys the primal wisdom of stable ligand-hole chemistry in sulfides, such as TiS3 (Ti4þS2
(S22
)), Fe2þS22
and VS4 (V4þ(S22
)2), offered by the pioneering works from Rouxel and others.9,11,87 Tackling with the ligands, therefore, unlocks a new avenue for exploring anionic redox as one can envisage diverse combinations of metal and ligands (O, S, Se etc.). Moreover, Li-rich sulfides also present some interests within the field of solid state batteries that is rapidly expanding due to their more compatibility towards sulfur-based solid electrolyte. All of these are leading researchers visiting back the sulfides compounds for anionic redox nowadays.

Status of Li(Na)-based anionic redox materials for better batteries

7.02.3

15

Fundamentals behind anionic redox

Since anionic redox involves electrons of the anions, a fundamental understanding of its electronic structural origination is pivotal to guide practical utilization. The rich family of electrodes exhibiting anionic redox turns out to be a gift towards a better understanding of the science underlying anionic redox. Herein, some crucial points regarding the energy levels and band structures of anionic redox based materials are revisited.

7.02.3.1

Band structure descriptions

Since most of the electrode materials consist of MO6 octahedral units, the band structures have to be derived from the octahedral molecular orbital theory. For a conventional MO6 configuration, as shown in Fig. 8A, the metal ions have nine valence atomic orbitals, i.e., nd, (n þ 1)s and (n þ 1)p, that can interact with symmetry adapted linear combinations (SALC) of oxygen 2p atomic orbitals based on their similar energies and symmetries. Specifically, among the five M nd orbitals, the two eg orbitals will form strong s-overlap with O 2p(s) orbitals resulting in highly energy-splitted eg and eg* orbitals, whereas the three M t2g orbitals will weakly interact with O 2p (p) orbitals forming less-splitted t2g* and t2g orbitals.88 The electrons from metal nd and O 2p states can therefore be accommodated into these molecular orbitals according to their energy levels, hence generating a band structure that can be schematically shown in Fig. 8B. The antibonding (M-O)* bands are predominantly filled by metal d states due to its highlying energy, while the bonding (M-O) bands are of mainly O 2p character as oxygen 2p electrons are in low energy (Fig. 8A). Considering a typical Li deintercalation process from a conventional lithium layered rock-salt compound, the removal of Li ions are concomitantly accompanied by the removal of electrons from the highest occupied states, generally the high-lying (M-O)* states, at the vicinity of Fermi level. This can be termed as classical “cationic redox” as normally (M-O)* states are mostly dominated by M character. However, in some cases the (M-O)* states also involve a large part of O 2p character endowed by the high M-O covalence that will cause substantial O participation in redox process, as is the case of deep delithiation in LiCoO2. Nevertheless, one should note that this O redox activity doesn’t supply extra capacity, since the O contribution is established at the expense of M contribution, both are from a single (M-O)* band. This (M-O)* redox behavior is ubiquitous in a lot of rock-salt layered compounds having the M layer fully occupied by M ions, such as LiCoO2, LiNiO2, LiNi1/3Co1/3Mn1/3O2 (NMC111), Nax[Ni1/ 3Mn2/3]O2 and so on, especially when deep delithiation is applied. Though, it is noteworthy that this (M-O)* redox behavior is intrinsically different from what we formally called as “anionic redox,” the latter originating from the O 2p NB/lone-pair states, as will be elaborated next.

Fig. 8 (A) Molecular orbitals of typical MO6 octahedral (Oh symmetry) considering both s- and p-type interactions. The symmetry allowed orbital interactions between metal and ligand atomic orbitals were indicated by dashed lines, with the resulted molecular orbitals shown in the middle. Note that, to avoid any confusion, the O p non-bonding orbitals originated from SALC of O 2p atomic orbitals were not shown. The anti-bonding (M-O) states with mainly M character and the bonding (M-O) states with mainly O character were outlined with shaded region. (B) Schematic band structure of the corresponding molecular orbitals in (A). EF, fermi level.

16

Status of Li(Na)-based anionic redox materials for better batteries

7.02.3.2

Anionic redox activity from O 2p “NB states” or “lone-pair states”

To enhance the specific capacity of an electrode, one should put more Li ions per molecular weight so that more Liþ transfer per unit mass can be enabled. In this context, Li-rich materials, typically based on Li2MO3 structure that can either be written as Li[Li1/3M2/3] O2, show superiority than conventional LiMO2 compounds as the former have more accessible Li ions with less molecular weight. However, moving from LiMO2 to Li2MO3 (Li[Li1/3M2/3]O2), Li substitution in M layer unavoidably decreases M content and increases the valence of M for maintaining charge neutrality, leaving less accessible cationic redox in Li2MO3, hence calling for the activation of O redox as a compensation. However, unlike the previously referred oxygen redox from (M-O)* states, the anionic redox activity from Li2MO3 is from O 2p NB or lone-pair states that have unique structural and chemical origin.3,58,89 As shown in Fig. 9A, conventional LiMO2 has full occupation M in M layer, engendering a uniform OM3Li3 local environment and hence a typical (M-O) and (M-O)* band structure engaging all three O 2p orbitals. However, Li2MO3 enlists additional Li ions in M layer, arousing an OM2Li4 environment with a Li-O-Li configuration (Fig. 9B). In this case, the one oxygen 2p orbital pointing towards this Li-O-Li direction has negligible covalent bonding interactions with Li 2s orbitals due to their large energy difference, giving rise to pure O 2p NB states (Fig. 9B). This O 2p NB states have higher energy than the bonding (M-O) states but lower than that of antibonding (M-O)* states, and therefore can provide extra redox activity once the (M-O)* states are fully depleted. It should be noted that this “O 2p NB states” is well recognized but named differently by other researchers to be “O 2p lone-pair states”,60 or “orphaned” or “unhybridized”58 O 2p states in previous reports. Being aware of how the O 2p NB states are generated, we then can invariably envisage its presence as long as O 2p orbitals are barely involved to form molecular orbitals. In this consideration, all metals that can form peroxides, such as alkaline, alkaline earth, or divalent d10 transition metals, are able to create such O 2p NB states due to their strong reducing ability to form very ionic MeO bond with negligible orbital overlapping. This has been verified by several reports observing the anionic redox activity even in nonalkaline-ion-rich compounds, such as Na2/3[Mg0.28Mn0.72]O251 and Na2/3Mn1
yZnyO253 compounds. Moreover, cation vacancy in M layer can also activate O 2p NB states as exemplified by Na4/7[,1/7Mn6/7]O2 (Na2Mn3O7).54,55 Steered by this criteria, the number of O 2p lone pairs can be regulated via controlling the M/O ratio or structural dimensionality, such as ranging from Li2MO3 to Li3MO4 to Li4MO5., or from cation-order to cation-disorder. This consequently opens up a large avenue for searching for high-capacity electrodes by merits of anionic redox activity.

7.02.3.3

Charge-transfer (D) vs. Mott–Hubbard (U) classification

As explained above, the electrochemical activity from O 2p NB states can be formalized to be “anionic redox” from now on. One more question that may appeal to experimenters is when such O 2p NB states will participate in redox process. This calls for knowledge on the ground electronic structures of the electrode. A universal guidance can be established based on the Zaanen–Sawatzky– Allen classification,90 which enlists two parameters, d-d Coulomb interactions term (U) and charge transfer energy (D). U describes the energy needed for electron transfer between two adjacent M sites (dndn / dn þ 1dn
1) that causes the splitting of the partially filled (M-O)* band into an empty upper and filled lower Hubbard bands, termed as UHB and LHB, respectively (Fig. 10A). D stands for the energy cost for a charge transfer from occupied ligand p band to an empty M d band (dn / dn þ 1L, wherein L denotes ligand hole), as shown in Fig. 10A. The relative values of U and D therefore classify the ground electronic structures of any electrode into three types,91 the Mott-Hubbard (Fig. 10A), charge-transfer (Fig. 10C), and intermediate regimes (Fig. 10B), as detailed below.

Fig. 9 (A) Local environment and band-structure of LiMO2, with homogeneous oxygen environments that every O ion is surrounded by three Li and three M ions with only Li-O-M configurations (dashed frame outlined). (B) Local environment and band-structure of Li2MO3. Compared with LiMO2, Li2MO3 has different O environments with Li-O-Li configurations (dashed frame outlined) enabled by additional Li ions in M layer. Such Li-O-Li arouses O 2p NB states as shown by the band structure schematic.

Status of Li(Na)-based anionic redox materials for better batteries

17

Fig. 10 Three band-structure cases classified by the relative values of U and D. (A) Mott-Hubbard case with U  D, exhibiting a conventional cationic redox scenario. (B) Intermediate case where U/2 z D, in which reversible anionic redox is enabled, coupled with MO6 distortion and O-O shortening. (C) Charge-transfer case when U [ D. This case was assumed to have irreversible anionic redox resulting in O2 release. Figures reproduced from Assat, G.; Tarascon, J.-M., Fundamental Understanding and Practical Challenges of Anionic Redox Activity in Li-Ion Batteries. Nat. Energy 2018, 3, (5), 373–386.

(i) Mott-Hubbard case (U  D): The Mott-Hubbard case can be predicted when U  D, wherein the partially occupied (M-O)* states lie just below Fermi level but above that of O 2p NB band. If (M-O)* states is highly separated from O 2p NB band, a classical reversible cationic redox will take place, as shown in Fig. 10A. This case can be frequently found in many cationic redox based polyanionic compounds, such as LiFePO4 and Na3V2(PO4)2F3, that have high metal-ligand ionicity (large D) and low lying O 2p states due to the inductive effect of polyanions. (ii) Intermediate case (U/2 z D): Once the LHB band and O 2p NB band are close in energy when U/2 z D (Fig. 10B), a unique situation happens called as “reductive coupling mechanism.”41,59,91 It happens since the initial removal of electrons from LHB band (cationic redox) causes the (quasi)degeneracy of M and O states (i.e., similar energies). The following anionic redox process will proceed by lowering the structural symmetry (e.g., Jahn-teller distortion) to allow efficient orbital interactions between (M-O)* and O 2p NB states so that electrons from the very localized O 2p NB band can be removed. Accompanying with this process is the simultaneous reduction of M and partial O-O dimerization involving the ligand-to-metal charge transfer. This reductive coupling process is essential to trigger a reversible anionic redox process which can provide extra capacity, as firstly manifested in the case of Li2Ru0.75Sn0.25O3. (iii) Charge-transfer case (U [ D): The charge-transfer regime is dictated by U [ D so that the LHB band is located at a lower energy than that of O 2p NB band. In this case, O 2p NB band is the first to deplete its electrons upon Li removal, giving rise to a pure anionic redox scenario. However, since O 2p NB states are very localized, the electron removal from this states creates unstable holes that will have to recombine together forming peroxo species. This new OeO bond formation will cause the O 2p NB band splitting into s, p, p*, s* bands, with the shorter the OeO bond (or the more oxidized the (O2)n
species), the larger the amplitude of the splitting (Fig. 11). The positions of newly generated O-O p*, s* bands with respect to the UHB band further divides this charge-transfer case into three subclasses60 (Fig. 11B and D): (1) If O-O s* band lies below the UHB band of metal, the (O-O)n
species can be stabilized and the anionic redox is reversible. (2) If the UHB band lies in between O-O p* and s* bands, then (O2)2
species can form. However, in this case, a O(2p s*)-M(d) band inversion will happen and the subsequent discharge will start with M reduction at first, followed by reduction of (O2)2
species. (3) For a third case when UHB band lies below O-O p* band, (O-O)n
will be fully oxidized by transferring its p* electrons to metal, resulting in oxygen release and cation reduction, known as a reductive elimination mechanism. This whole charge-transfer framework can be well represented by the cases of d0-metal based Li-rich compounds (e.g., Li1.17Ti0.33Fe0.5O2) or metal inactive compounds (e.g., Li2MnO3), as they have no d electrons or d electrons in much lower energy than O 2p states, leaving solely O 2p states as the majority below Fermi level. It is therefore not hard to understand why, as has been proven recently by RIXS and quantitative differential electrochemical mass spectrometry (DEMS),29,30 that Li2MnO3 shows only oxygen release rather than reversible anionic redox, as it neatly falls into this framework with O-O dimers having higher antibonding states than empty Mn 3d states. With above framework, the anionic redox activity of various compounds can, therefore, be predicted and rationalized. Although the accurate prediction of their band structures through DFT method, which are functional-dependent, is precarious, this framework provides general guidance through which anionic redox activity can be wisely utilized. Yet there are also other puzzles exist for the anionic inactivity of, for instance, layered Li2TiS3 or Li2TiSe3,82,86 both belonging to charge-transfer regime with sulfur p states being the top below Fermi level. This cast doubts on the activation mechanism of anionic redox that necessitates further scrutiny.

7.02.3.4

Reductive coupling mechanism

As mentioned above, a reductive coupling mechanism describes a process in which cations will be over-oxidized first, followed by electron transfer from O NB states to cation that is associated with structural distortion (O-O dimerization etc.). This mechanism

18

Status of Li(Na)-based anionic redox materials for better batteries

Fig. 11 (A) The band-structure of a charge-transfer type compound with U [ D, with O 2p NB band lying on the top. Once electrons are removed from O 2p NB band, the peroxo-like species formation splits the band into s, p, p*, s* bands enlisting three cases depending on the magnitude of the splitting (or O-O distance), as shown in (B).

was first conceptualized at 2013 for elucidating the anionic redox activity in Li2Ru0.75Sn0.25O3 compound,41 and it consists in an electron transfer according to the schematized reaction Ru6þ-O2
/ Ru5þ-(O2)2
(Fig. 12). This process was afterwards rationalized by theoretical calculations on Li2RuO3 model structures59 and used to account for reversible anionic redox in high-covalent systems. Worth mentioning is that, similar ligand-to-metal charge transfer (LMCT) processes was also proposed in some compounds, for instance, a-Li2IrO3 (Ir> 5.5 þ-On
/ Ir5.5 þ-O(n
1)
)44 and Li2MnO3-based materials (Mn7þ-O2
/ Mn4þO(2
n)
),62,63 to account for the anionic redox-cation migration coupling behavior. Note that nomenclature causes here again a source of confusions since LMCT shares exactly a similar mechanism with reductive coupling process. Though theoretically it is well accepted by the community, the intermediated states associated to reductive coupling hasn’t been fully observed yet by experimentalists. The reason could be rooted at the instability of the highly oxidized cationic intermediates (e.g., Ru6þ, Mn7þ, Ir> 5.5 þ) that cannot be captured by any time-resolved and damage-free techniques. This obstacle has been recently overcome by the successful spotting, via Mössbauer spectroscopy, of Fe4þ intermediate and its disappearance owing to its longer lifetime.81 Nevertheless, the dynamic charge transfer process from ligand to metal still remains a hypothesis unless one can directly observe it through a unique model system, for example, system with a long-lived cationic intermediate.

7.02.3.5

Anionic activity from O (2p)-M(nd) p-type interaction

In addition to O 2p NB states activity which is well accepted by the community, some groups also proposed the redox activity from O(2p)–M(t2g) p-type interactions. The concept arises from the deduction of the molecular orbitals based on group theory if we consider the existence of p-type interactions between metal nd and O 2p states.88 Unlike layered LiMO2 compounds in which the MO6 octahedra shows a conventional Oh symmetry, the MO6 units in Li2MO3 show a C2 or C2v symmetry.92,93 This symmetry led researchers to deduce a symmetry-allowed interactions between M and O, including both s- and p-type interactions. Among them, the O 2p (p) band will overlap with M t2g orbitals to form bonding and antibonding states (t2g and t2g* in Fig. 8A, also called as b1 and b1*,92 respectively), which is the so-called p-type interaction. Similar case was also reported for Na2Mn3O7 phase with a D3d symmetrized Mn-O system.64 Nevertheless, some reports show that the symmetry adapted linear combinations of the three O 2p orbitals also generate some unmatched symmetrized orbitals that remain as non-bonding states,88,93 which is missing in other reports solely considering p-type interactions, leading to confusions that remain to be reconciled. The existence of such weak O(2p)-M(t2g) p-type interactions has also be confirmed experimentally by RIXS94 as well as highenergy X-ray Compton scattering techniques.95 It was even shown in Li-rich NMC compound that, the p interaction will become stronger upon Li extraction due to the oxidation of metal and holes created at O NB states, resulting in (s þ p) multi-orbital bonds.94 Moreover, this p-type interaction was also linked to some unusual behaviors, such as the negligible voltage hysteresis of Na2Mn3O7 phase, for which the delocalized Mn-O p states were shown to be responsible for its activity rather than O 2p

Status of Li(Na)-based anionic redox materials for better batteries

19

Fig. 12 Reductive coupling mechanism of Li2Ru0.75Sn0.25O3. (A) Schematic of two oxo-ligands coordinated to the transition metals leads to a single or double metal reduction depending on the coordination mode of the O2 moiety. (B) Representative local structures in the Ru layer after delithiation obtained from DFT calculation. The peroxo-like species with a shortened O-O distance was indicated. Reproduced from Sathiya, M.; Rousse, G.; Ramesha, K.; Laisa, C. P.; Vezin, H.; Sougrati, M. T.; Doublet, M. L.; Foix, D.; Gonbeau, D.; Walker, W.; Prakash, A. S.; Ben Hassine, M.; Dupont, L.; Tarascon, J. M., Reversible Anionic Redox Chemistry in High-Capacity Layered-Oxide Electrodes. Nat. Mater. 2013, 12(9), 827–835.

lone-pair states necessitating oxygen dimerization coupled with structural changes.64 Although intriguing, the validity of such p states redox still requires further scrutiny as divergence still exists between several reports.

7.02.3.6

The nature of oxidized O2L: Electron holes, O-O dimers, trapped O2 molecules, and oxygen release

Unlike cationic redox, anionic redox behaves much more complicated, as electron removal from O may generate electron holes that might not be stable, and that finally will coalesce to form O-O dimer, or O2 molecule trapped in the lattice or evolving out. Resolving the exact form of resultant species for anionic redox is therefore important as it determines the kinetics and reversibility of electrochemical reactions. The electron “hole” concept in oxides was already proposed earlier in some specific systems for high-temperature superconductors, notable example being the YBa2Cu3O7 wherein Cu has a “3d9L” configuration (L denotes electron hole).96 A similar “hole” concept was also extended to sulfides by J. Rouxel proposing that holes can be created at the top of sulfur sp band and stabilized without association into anionic pairs for many systems, such as CuCr2S4 (Cuþ(Cr3þ)2(S2
)3S
•, where • denotes an electron hole). These pioneering works provides hints that the electrochemical Li removal from some compounds could also be compensated by electron holes. A good example is Na2Mn3O7 compound, which was well agreed by several groups to have O electron hole redox chemistry, as confirmed by RIXS, XAS and magnetic measurements (Fig. 13A).80,97 The stabilization of electron holes in O 2p band were proposed to be rooted in either Coulombic interactions between oxidized O ions and Na vacancies or (s þ p) multiorbital MneO bond. Though convincing, it is understandable as one Na removal from Na2Mn3O7 only produce 1/7 hole per O, that is far from forming short OeO bonds. This is true as previous theory already unraveled a quasi-linear relationship between the anion-anion bond distance and formal charge per anion atom.11 In simple words, the O redox in Na2Mn3O7 could proceed by forming indiscernible O-O dimer having similar property to holes. In contrast to holes, the anion dimer formation has been well recognized by the community to account for anionic redox behavior in many compounds. The S-S dimer formation was clearly demonstrated in Li2FeS2 system in which its delithiation firstly oxidize Fe2þ to Fe3þ followed by S pairing, resulting in FeS2 with a SeS bond distance (2.12 Å) similar to that of pyrite-type FeS2 (2.18 Å).98,99 Such anionic dimer formation is less remarkable but obvious in oxides. For typical Li-rich oxides, the investigations on Li2RuO3 and Li2IrO3 based systems unambiguously disclosed the evidences of peroxide-like (O2)n
species by means of EPR, transmission electron microscopy (TEM), and Neutron diffraction in combination with consistent theoretical predictions.41–43 Though, it is noteworthy that the reported bond length of O-O dimers are in the range of 2.4–2.5 Å that is far from that of a peroxide

20

Status of Li(Na)-based anionic redox materials for better batteries

Fig. 13 (A) Electron hole formation in Na2
xMn3O7. The left panel shows the charge and discharge curves of Na2Mn3O7 in a Na half-cell. The right panel shows the O K-edge XAS of Na2
xMn3O7 from pristine to charged 4.5 and 4.7 V. A small peak arising at 527.5 eV indicates the formation of O 2p electron holes. (B) Peroxo-like species formation in both a- and b-Li2IrO3 upon delithiation. Left panel shows the enlarged ABF-STEM figure of a-LiIrO3 with O-O pairs with shortened projected distances, as labeled by the O-O dumbbells. An ABF intensity profile figure was shown below for directly visualizing the long (blue) and short (red) projected O-O distance. (C) Trapped O2 molecules in Na0.75
x[Li0.25Mn0.75]O2 by high-resolution RIXS. The region highlighted by dashed square shows the RIXS signal arising from the vibrations of the OeO bond with a fundamental vibrational frequency (1600 cm
1) that is similar to that of molecular O2 and the 1.2 Å OeO bond predicted from DFT. Figures adapted from: (A) Abate, I. I.; Pemmaraju, C. D.; Kim, S. Y.; Hsu, K. H.; Sainio, S.; Moritz, B.; Vinson, J.; Toney, M. F.; Yang, W.; Gent, W. E.; Devereaux, T. P.; Nazar, L. F.; Chueh, W. C., Coulombically-Stabilized Oxygen Hole Polarons Enable Fully Reversible Oxygen Redox. Energ. Environ. Sci. 2021, 14(9), 4858–4867; (B) McCalla, E.; Abakumov, A. M.; Saubanere, M.; Foix, D.; Berg, E. J.; Rousse, G.; Doublet, M. L.; Gonbeau, D.; Novak, P.; Van Tendeloo, G.; Dominko, R.; Tarascon, J. M., Visualization of O-O Peroxo-Like Dimers in High-Capacity Layered Oxides for Li-Ion Batteries. Science 2015, 350(6267), 1516– 1521, Pearce, P. E.; Perez, A. J.; Rousse, G.; Saubanere, M.; Batuk, D.; Foix, D.; McCalla, E.; Abakumov, A. M.; Van Tendeloo, G.; Doublet, M. L.; Tarascon, J. M., Evidence for Anionic Redox Activity in a Tridimensional-Ordered Li-Rich Positive Electrode beta-Li2IrO3. Nat. Mater. 2017, 16(5), 580–586; (C) House, R. A.; Maitra, U.; Perez-Osorio, M. A.; Lozano, J. G.; Jin, L.; Somerville, J. W.; Duda, L. C.; Nag, A.; Walters, A.; Zhou, K. J.; Roberts, M. R.; Bruce, P. G., Superstructure Control of First-Cycle Voltage Hysteresis in Oxygen-Redox Cathodes. Nature 2020, 577(7791), 502–508.

Status of Li(Na)-based anionic redox materials for better batteries

21

species (1.3–1.5 Å), hence the reason why they were normally called as peroxo-like species. This situation is apparently caused by the less oxidized O2
ions in these peroxo-like species compared with peroxides, as the bond length is nearly proportional to the formal charge per O atom.100 Further pushing the dimerization will, as instructed by previous band-structure explanation, cause oxygen molecule formation that may be trapped in the lattice or evolve out. While oxygen release has been identified very early with DEMS, the trapped oxygen molecule was validated recently by P. G. Bruce et al. in Na0.75[Li0.25Mn0.75]O2101 as well as Li-rich NMC compound102 by merits of high-resolution resonant inelastic X-ray scattering (HR-RIXS) coupled with solid-state 17O magic angle spinning (MAS) nuclear magnetic resonance (NMR) techniques. Same concept was also extended to other typical Li-rich compounds including Li2Ru0.5Sn0.5O3 and Li2Ir0.5Sn0.5O3, and also cation-disorder Li-rich oxyfluorides by the same group.103,104 This (re)formation/breaking of molecular O2 was shown to be reversible in the bulk thanks to the caging effect of cation vacancies, while it is less reversible in the near surface region that causes oxygen release accompanied with surface reconstruction. However, the possibility of oxygen loss from bulk cannot be ruled out as oxygen vacancies were reported previously, and their subsequent annihilation by migrating outward, though with slow kinetics, increases the M/O ratio called as “densification” that has been well accepted.105,106

7.02.4

Anionic redox opening a new rich materials chemistry

Both combined experimental and theoretical approaches have indicated that anionic redox activity in Li-rich layered oxides is governed by the respective positioning of the metal d-bands and the ligand non-bonding p states. Chemical insights based on simple electronegativity considerations can help us in properly manipulating band positioning, hence enabling to define the most fertile directions in achieving higher energy density materials. Some of the strategies tried are discussed below.

7.02.4.1

Increasing the Li/M and O/M ratio in layered rock-salt compounds

Layered rock-salt compounds are the most prevailing cathode hosts not only at the beginning but also nowadays. As the early members being discovered, LiCoO2 and LiNiO2 and their derivatives (NCA and NMC) have achieved a great success in various application scenarios. However, a drawback is their structural stability upon deep delithiation whose origin is nested in the large involvement of oxygen 2p character in (M-O)* that participates redox process, which creates holes at oxygen and leads to oxygen release.89 A practical way commonly to remain out of this instability regime is to limit the charge cut-off voltage below 4.4 V. Besides, Li-rich materials, especially Li2MO3 analogies, provide another way to contour this structural instability trap. Their structural stabilization arises because of the O 2p NB states, which enables reversible anionic redox unlike the (M-O)* states. As a result, such Li2MO3-based layered rock-salt structures can deliver high capacities exceeding conventional ones. Two dimensionalities can be envisaged regarding materials design when move from LiMO2 to Li2MO3. A first one consists in d electrons, as illustrated from Li2TiO3 (Ti4þ: 3d0) to Li2MnO3 (Mn4þ: 3d3) to Li2NiO3 (Ni4þ: 3d6), while another one has dwelled the replacement of 3d by 4d and to even 5d metals (Li2MnO3 (3d) / Li2RuO3 (4d) / Li2IrO3 (5d)) (Fig. 14). The use of these two degrees of freedom was a gift to both experimentalists and theorists to predict and synthesize new phases showing high-capacity electrodes. The transition metals on the left side of the periodic table, such as Ti and Zr, mostly form inactive Li2MO3 phases due to the M d0 configuration in absence of classical cationic redox, though it remains unclear why O cannot supply electrons for redox reaction. A way to get out of this trap and to trigger the anionic activity was achieved, for instance, via Fe substitution in Li2TiO3 phase,81,82 hence unveiling the crucial role of redox-active metal (e.g., Fe3þ) that brings extra d electrons. This inactiveness also extends to bulk Li2MnO3 phase, which, however, can be activated by nanosizing or high-temperature treatment.22,23 Pushing it further to the right side of transition metals, anionic redox activity becomes inherent to the synthesized phases like Li2NiO3 for instance.107 This trend might reveal the importance of charge-transfer bandgap, which decreases from Ti to Mn and to Ni as inferred from the increased M-O covalence, in favoring anionic redox. However, for active Li2MnO3 and Li2NiO3, their oxygen redox at first charge almost entirely end up with oxygen release rather than reversible anionic redox, as demonstrated in recent studies.30,107 Due to these shortcomings, Li2MO3 phases based on 3d-column M cannot work by themselves, but mostly act as parental structures in

Fig. 14 Li2MO3 electrode materials design from two dimensions based on nd orbital of M. One is by increasing the d electrons, as indicated by the horizontal arrow from d0 to d6 etc. Second is to increase the d period from 3d to 5d.

22

Status of Li(Na)-based anionic redox materials for better batteries

which cation substitutions could enable reversible anionic redox, as reflected by the substantial reversible capacities delivered by their derivatives (e.g., Li1.2Ti0.4Mn0.4O2108 and Li-rich NMC). Unlike 3d Li2MO3 phases, the 4d and 5d ones are getting more active and reversible in terms of their anionic redox capability. Two most studied compounds are Li2RuO3 and a-Li2IrO3, which are isostructural to Li2MnO3 but with redox-active transition metal (Ru4þ and Ir4þ). A lot of their derivative phases, such as Li2Ru1
ySnyO3 and Li2Ir1
ySnyO3, has been reported serving as model structures for fundamental understandings of anionic redox.41,42,44,109,110 In addition, the merits of going to 4d or 5d system also include their simpler but more robust structures that endow more convenience in determining their structural changes. This led to the success in directly observing the peroxo-like species as an evidence of anionic redox by Neutron diffraction and electron microscopy.42 Another branch of the Li2MO3 family worth mentioning is the combination of two metal or metalloid ions whose valence sum equals 8, as represented by the general formula Li4MM0 O6. This enlists several example compounds such as Li4Fe2þ 6þ Te O6,111 Li4Fe3þSb5þO6112 and Li4Ni2þTe6þO6,113 with interesting redox chemistries being reported. Bearing in mind that the capacity of the Li-rich phases is governed by the amount of Li per unit formula, Li3MO4 and Li4MO5 phases were also explored in addition to Li2MnO3. These phases are all based on rock-salt layered structure (Fig. 15A) but with different superstructure ordering, as manifested by their compositions rewritten as Li1 þ yM1
yO2, where y ¼ 1/3 for Li2MO3 and y ¼ 1/2 for Li3MO4 etc. Such increase of Li content at the expense of M content simultaneously increases the number of Li-O-Li configurations around an individual O atom, and therefore, increases the number of O 2p NB states (Fig. 15B). Following this guideline, a new Li3IrO4 phase was hence designed (Fig. 15C) showing solely anionic redox upon removing two Li ions, while the third Li extraction finally ends up with oxygen release.45 However, by limiting the charge to the removal of 2 Liþ and by extending the discharge to 1.5 V to trigger the Irþ 5 reduction, this compound was shown being able to reversibly insert 3.5 Liþ per unit

Fig. 15 Increasing Li/M and O/M ratio to trigger more O 2p NB states. (A) Typical layered rock-salt structure. The local environment of specific O ion was marked out by a frame. (B) The evolution of the O environment from (a) upon the increasing of Li/M and O/M ratio. With higher the Li/M and O/M ratio, the O 2p NB states increases, as schematically shown by the band structures below. (C) XRD and crystal structure of R-3m Li3IrO4 with random cation distribution in Ir layer. (D) Electrochemical property of Li3IrO4 tested in a Li half-cell that shows high accessible Liþ capacity (3.5 Liþ per unit formula) with considerable anionic redox activity. Adapted from (B–D) Perez, A. J.; Jacquet, Q.; Batuk, D.; Iadecola, A.; Saubanère, M.; Rousse, G.; Larcher, D.; Vezin, H.; Doublet, M.-L.; Tarascon, J.-M., Approaching the Limits of Cationic and Anionic Electrochemical Activity With the Li-Rich Layered Rocksalt Li3IrO4. Nat. Energy 2017, 2(12), 954–962.

Status of Li(Na)-based anionic redox materials for better batteries

23

formula (Fig. 15D), hence enabling to reach a capacity as high as 356 mA h g
1 for 25 cycles. Although this strategy is attractive, it should be noted that the number of compounds is getting limited moving from Li2MO3 to Li3MO4 and to Li4MO5. This is because a higher Li/M ratio requires a higher-valence M, which is getting less abundant from M4 þ to M5 þ to M6 þ, for balancing the charge. Besides, simply pushing higher the Li/M ratio is not recommended as it will impair structural stability, as more Li leads to more ionic structure wherein the MO6 would be isolated by LiO6 units (Fig. 16). Such structure with higher Li/M ratio is susceptible to anionic redox due to more O 2p NB states generated, and finally more oxygen release and structural collapse will happen. One example is Li3RuO4, whose superstructure in MO2 slab is zig-zag chains without connection together (Fig. 16), showing irreversible oxygen redox accompanied with oxygen release and Ru dissolution upon deep charge.114,115 This concern will be aggravated if we go further towards Li4MO5 and even Li5MO6 structures, as exemplified by Li4MoO5 and Li5ReO6, both showing less connected and even entirely isolated MO6 units (Fig. 16). Therefore, a tradeoff between high capacity from anionic redox and structural reversibility is required when considering this over-lithiation strategy. Additionally, the recently reported study of Co3O4loaded Li2O116 or Li2O/iridium-graphene117 nanocomposites, that turns the insertion chemistry into a conversion-type reaction functioning as Li2O (s) 4 Li2O2 (s) 4 LiO2 (s), could be an interesting approach to go beyond this limit, although it needs further evaluation. Alternatively, merging these phases having high M-O ratio with other more robust compounds forming solid solution can be a good strategy for designing new high-capacity electrodes. As first initialized by Yabuuchi et al., the combinations of Li3NbO4 and LiMO2 (M ¼ Fe or Mn) or M0 O (M0 ¼ Ni or Co) phases,118 forming fully cation-disordered compounds, are effective in triggering high reversible capacities based on contributions from both cationic and anionic redox. Among them, Li1.3Nb0.3Mn0.4O2 is the best one delivering  250 mA h g
1 reversible capacity at 50  C, and hence received enormous attentions in the following research in view of its potential practicality. Since then, more similar trials, such as Li3TaO4-, Li4MoO5-, and Li4WO5-based systems,78,79,119 were implemented to explore fundamental issues in these compounds. Apparently, a prolific materials chemistry is open towards researchers by these disparate combinations among MO, LiMO2, Li2MO3, Li3MO4, and so on, which awaits for deeper exploration.

7.02.4.2

Playing with the alkaline ion: From Li to Na

Alike Li-ion systems, the developing Na-ion technology could also be greatly benefited from electrodes with enhanced capacities, hence the recent efforts devoted to unraveling Na-ion compounds showing anionic redox activity. Worth recalling is that conventional layered O3-type NaMO2 (M ¼ Fe, Co, Mn etc. or their combinations) phases were intensively studied in parallel to Li-based ones but fell into oblivion owing to the outstanding performances of their Li counterparts.120,121 With the increased awareness of sustainability issues, the Na-ion chemistry is enjoying revival with both O3 or non-stoichiometric P2 or P3 phases.122,123 Regardless of Na stoichiometry, the anionic redox activity has been claimed in a lot of these compounds having full M occupation in M layer, namely, no O 2p NB states. Examples are O3-NaFeO2 and its derivatives, NaTi0.5Ni0.5O2, P2-Na2/3Ni1/3Mn2/3O2, P3-type Na2/ 124–126 and therefore not further discussed 3Ni0.2Mn0.8O2 and so on, with a full list of them previously enumerated in other reviews here. These compounds share the similar redox mechanism as in LiCoO2 wherein only (M-O)* states are involved for redox process, but with no extra capacity gained as early mentioned. In an attempt to get extra capacity, a more legitimate pursuit of anionic redox should rely on Na-rich compounds or analogies that have lone-pair states serving as the electron reservoir. However, unlike Li-rich oxides, placing additional Na ions in transition metal layer is a difficult task for most of the 3d systems due to the size mismatch between NaO6 and MO6. Though Na2MnO3 was mentioned in a theoretical calculation study,127 the

Fig. 16 Structural evolution for layered rock-salt compounds upon increasing Li/M and O/M ratios. For the whole structure (3D view), it remains well as rock-salt layered. For local view (M layer), it can be found that the higher the O/M (or Li/M) ratio, the less connected the MO6 units.

24

Status of Li(Na)-based anionic redox materials for better batteries

successful preparation of Na2MnO3 phase is never reported. Nevertheless, Na2TiO3 that exists under three polymorphs (a-, b-, and g-Na2TiO3),128 could be an exception but none of them are electrochemically active. By implementing the same strategy as for Lirich electrodes, namely, combining inactive b-Na2TiO3 phase with other active NaMO2 phase, single cation-disordered Na1.14Mn0.57Ti0.29O2 phase showing capacities as high as 200 mA h g
1 based on both Mn and O redox could be obtained.129 Alternatively, moving from 3d to 4d or even 5d systems is also viable for stabilizing Na-rich layered oxides. This is made possible by the larger size of the transition metal (hence larger MO6 units), which facilitates the facile accommodation of NaO6 octahedra in M layerdthe reason why the earliest anionic redox activity was reported in a 4d system, Na2Ru1
ySnyO3.130 More 4d and 5d Na-rich compounds, as represented by Na2RuO3 and Na2IrO3 (Fig. 17), mostly based on Ru and Ir, were then successively uncovered with considerable anionic redox contributions.47–49,114,131,132 Among them, one unique system, Na3RuO4, was identified to have Ru5þ/ Ru6þ redox prior to having oxygen oxidation during charge, unlike that of Li3RuO4.114 Such difference was shown to be nested in the bigger ionic size of Na than Li triggering a less coordinated Ru so that an unusual Ru6þ species can be stabilized. This highlights the difference between Na and Li ion batteries when considering the redox mechanism involving anionic redox, setting the challenges for one directly transposing the knowledge from Li to Na.

Fig. 17 Structure and electrochemistry of 4d and 5d transition metal based Na-rich compounds. (A) Crystal structure of Na2RuO3 with (B) a honeycomb ordering in Ru layer. (C) The charge and discharge curves of Na2RuO3, exhibiting both Ru and O redox separately at two plateaus. (D) Electrochemistry of Na2IrO3 with its crystal structure (isostructural with Na2RuO3) shown as insets. (E) Enlarged ABF-STEM figure of Na0.5IrO3 with distorted MO6 rings being noted, and (F) the corresponding intensity profiles highlighting short and long projected O-O distances. The short O-O distance indicates the formation of peroxo-like species and hence the evidence of anionic redox. Adapted from: (A–C) Mortemard de Boisse, B.; Liu, G.; Ma, J.; Nishimura, S.; Chung, S. C.; Kiuchi, H.; Harada, Y.; Kikkawa, J.; Kobayashi, Y.; Okubo, M.; Yamada, A., Intermediate Honeycomb Ordering to Trigger Oxygen Redox Chemistry in Layered Battery Electrode. Nat. Commun. 2016, 7, 11397; (D–F) Perez, A. J.; Batuk, D.; Saubanère, M.; Rousse, G.; Foix, D.; McCalla, E.; Berg, E. J.; Dugas, R.; KHW van den Bos; Doublet, M.-L.; Gonbeau, D.; Abakumov, A. M.; Van Tendeloo, G.; Tarascon, J.M., Strong Oxygen Participation in the Redox Governing the Structural and Electrochemical Properties of Na-Rich Layered Oxide Na2IrO3. Chem. Mater. 2016, 28(22), 8278–8288.

Status of Li(Na)-based anionic redox materials for better batteries

25

Although 4d and 5d systems are in advantage of high covalence that empowers more reversible anionic redox activity, a pressing need is to shift the research effort towards more practical compounds that are mostly based on cost-effective 3d transition metals. Implementing this is difficult due to the limited Na-rich phase in 3d system. However, this can be circumvented via substituting the transition metal in Na-deficient NaxMO2 compounds with other metal ions that do not form covalent bond with oxygen ions, such as alkaline-earth metal ions (e.g., Mg and Zn) or even vacancies, to vitalize O 2p NB states, as mentioned in an earlier section. This strategy was first initiated in an Mg-substituted P2-type compound, Na2/3[Mg0.28Mn0.72]O2, that shows anionic redox activity without causing cation migration (Fig. 18), offering a large playground for designing anionic redox based sodium ion compounds.51 Consequently, a variety of NaxM1
yAyO2 (A ¼ Li, Mg, Zn, or vacancy) compounds52–54,101,133–136 were synthesized exhibiting anionic redox, with M mostly being Mn and Ti due to their relatively higher valence (tetravalent) for charge compensation upon the introduction of low-valence A ions.

Fig. 18 Triggering anionic redox by Mg doping in sodium manganese compound. (A) P2-type structure of Na2/3[Mg0.28Mn0.72]O2 and its (B) honeycomb ordering of the transition metal layer of P2-type Na2/3[Mg0.28Mn0.72]O2 with MgMn6 units. (C) Electrochemical performance, with charge I period showing Mn redox while II showing O redox. Discharge III consists of successive O and Mn redox. (D) O redox activity characterized by RIXS. A peak emerging at 523 eV indicates the oxidation of O2
ions. Adapted from Maitra, U.; House, R. A.; Somerville, J. W.; Tapia-Ruiz, N.; Lozano, J. G.; Guerrini, N.; Hao, R.; Luo, K.; Jin, L.; Perez-Osorio, M. A.; Massel, F.; Pickup, D. M.; Ramos, S.; Lu, X.; McNally, D. E.; Chadwick, A. V.; Giustino, F.; Schmitt, T.; Duda, L. C.; Roberts, M. R.; Bruce, P. G., Oxygen Redox Chemistry Without Excess Alkali-Metal Ions in Na2/ 3[Mg0.28Mn0.72]O2. Nat. Chem. 2018, 10(3), 288–295.

26

Status of Li(Na)-based anionic redox materials for better batteries

Nonetheless, the merit of anionic redox in acquiring high capacity cannot be fully exerted in Na-deficient compounds, as the number of exchangeable Na ions, rather than electrons, is limited. This calls for searching full stoichiometry Na compounds having the composition of NaM1
yAyO2 in the same manner with Li ones. An early attempt has been reported on a theoretical prediction of O3-NaLi1/3Mn2/3O2 compound,137 analogy to Li2MnO3 (Li[Li1/3Mn2/3]O2). The first success preparation of this compound was achieved recently by carefully adjusting the synthetic conditions and stoichiometry.50 The obtained moisture-stable O3-NaLi1/
1 over 40 cycles 3Mn2/3O2 phase shows reversible anionic redox, resulting in high sustainable capacity of  190 mA h g (Fig. 19). Separately, a Ti-substituted NaLi1/3Mn2/3O2 compound, NaLi1/3Ti1/6Mn1/2O2, was equally synthesized, guided by an empirical “cationic potential” parameter for rational designing the O3 or P2 Na-layered oxides.135 These advances offer momentum for continuing seeking practical electrodes based on anionic redox for Na-ion batteries. Besides, this anionic research is also being extended to other metal-ion batteries, such as K-ion, Mg-ion, and Al-ion batteries.138–140 Altogether, anionic redox offers a new dimension to solid-state chemists for exploiting high-energy density battery systems.

7.02.4.3

Ligand manipulation: From oxides to sulfides/selenides

Another viable strategy for tuning band positioning is to act on the ligand. On the basis of electronegativity considerations (cSe < cS < cO), the highly lying p bands of sulfur or selenium are higher in energy than the ones corresponding to O (Fig. 20), and can be more readily involved in redox process than O 2p band. Besides, reversible sulfur redox is prone to be easily achieved than O redox due to the well-known stable ligand-hole chemistry in sulfides. In view of these advantages, researchers are now revisiting back to sulfides or even selenides chemistry for uncovering fruitful electrode material chemistry. The earlier reports regarding sulfide electrodes for Li batteries frequently refer to Li-free sulfides, mostly the two-dimension chalcogenides, such as TiS2, MoS2, NbS2, VSe2, and so on.2 It was shown that Li could be reversibly inserted/deinserted in these hosts, with, however, limited number of works focusing on interpreting their redox centers. For some covalent systems, such as LiMS2 phases, Li removal was shown to be charge compensated by S-3p “holes” but without forming disulfide bonds due to the sufficient cation-3d character involved.141 Similar to LiMO2, the redox activity in LiMS2 should only come from M(nd)-S(3p) antibonding states, namely, one-band (M-S)* redox, without supplying additional capacity. Alike Li-rich oxides, one may consider if it is possible to synthesize Li-rich sulfide phases to generate S 3p NB states for sulfur anionic redox. However, moving from oxides to sulfides inevitably increases the energy of p band whose electrons may finally pour into empty M d band. This sets a limit for stabilizing high-valence metal species that are especially pivotal in forming Li-rich sulfides (e.g., Li2MS3). d0 transition metals ions, such as Ti4þ, Zr4þ, and Nb5þ, are ideals for forming Li-rich sulfides as their empty d bands are positioned at very high energy. This lays the foundations of the stabilization of Li2TiS3, Li2ZrS3, and Li3NbS4 structures etc., though which are electrochemically inactive. Once shifting towards the right-hand transition metal, the drop in d-level energy leads to the lowering of cationic oxidation states through a sp-to-d electron transfer, resulting in the polymerization of sulfur and consequently explaining the unfeasibility to form Li-rich sulfides like Li2MnS3 or Li2RuS3, since S2
will be fully reduced in these cases. This explains why the right-hand transition metals generally form pyrites and marcasites or even more complex groups with polymerized sulfur frameworks (e.g., IrSe3) to stabilize the holes created at anion sp band.10 This narrows down the exploitable Li-rich sulfides to only the left side of transition metal, and more so for selenides. Even though limited by transition metals, the Li-rich sulfides chemistry can be enriched via several strategies. The first one is by introducing active transition metals in electrochemically inactive layered sulfides (Fig. 21A), such as Li2TiS3, to trigger the sulfur redox activity, as embarked by Co-substituted Li2TiS3 in obtaining Li1.2Ti0.6Co0.2S2 showing S redox activity.82 Shortly after, a similar approach has consisted in substituting Ti by Fe in Li2TiS3 for obtaining a Li1.13Ti0.57Fe0.3S2 phase having a sustainable capacity of 245 mA h g
1 owing to cumulated Fe2þ/Fe3þ and S2
/S(2
n)
redox processes83 (Fig. 21A). Other than capacity, Li1.13Ti0.57Fe0.3S2 shows advantages over the Li-rich oxides counterparts in terms of fast kinetics and negligible voltage hysteresis and voltage fade.

Fig. 19 (A) Refinement of the synchrotron XRD for NaLi1/3Mn2/3O2 compound, showing a layered rock-salt structure with honeycomb ordering in Mn layer, as inferred by the Li-Mn-Mn ordering in STEM figure (inset). (B) Electrochemical performance. (C) Anionic redox activity by hard XPS. P, pristine. 1C, 1st fully charged. Reproduced from Wang, Q.; Mariyappan, S.; Rousse, G.; Morozov, A. V.; Porcheron, B.; Dedryvere, R.; Wu, J.; Yang, W.; Zhang, L.; Chakir, M.; Avdeev, M.; Deschamps, M.; Yu, Y. S.; Cabana, J.; Doublet, M. L.; Abakumov, A. M.; Tarascon, J. M., Unlocking Anionic Redox Activity in O3-Type Sodium 3d Layered Oxides Via Li Substitution. Nat. Mater. 2021, 20, (3), 353–361.

Status of Li(Na)-based anionic redox materials for better batteries

27

Fig. 20 Schematic of the band structure evolution from oxides to sulfides and to selenides. The anion p band shows higher and higher energy due to the increasing M-L covalence (or charge transfer bandgap, as indicated by the double-headed arrow).

Based on this strategy, more and more similar Li-rich sulfide compounds (e.g., Ti3þ-doped Li2TiS384 and Fe2þ-doped Li2SnS3142) are being discovered. Alternatively, anionic redox can also be triggered by performing ligand substitution (Fig. 21C). This was shown by the recent successful synthesis of a series of Li2TiS3
ySey compounds, whose capacity peaks at 260 mA h g
1 based on Se2
/Sen
, S2
/Sm
and Ti3þ/Ti4þ redox processes, whereas the two end members, Li2TiS3 and Li2TiSe3, show negligible activity.86 A locally distorted Ti environment, which lower the symmetry of anion p bands, due to the formation of heteroleptic TiCh6 octahedra

Fig. 21 Two strategies to activate the sulfur redox activity in d0 layered chalcogenides, with Li2TiS3 shown as an example. The electrochemistry of inactive Li2TiS3 is shown in the middle (b), with its layered structure presented as inset. The sulfur anionic redox activity can be activated by cation substitution, as represented by Fe substutition generating (A) Li1.33
2y/3Ti0.67
y/3FeyS2, or by anion substitution, as represented by Se substitution forming (C) Li2TiS3
ySey. Both strategies result in considerable electrochemical activities as inferred from the electrochemistry. A schematic bandstructure evolution by these two strategies are shown below to illustrate the redox activities. Note that in Li2TiS3
ySey, a small amount of Ti3þ/Ti4þ was also triggered. Adapted from: electrochemistry in (A and B) Saha, S.; Assat, G.; Sougrati, M. T.; Foix, D.; Li, H.; Vergnet, J.; Turi, S.; Ha, Y.; Yang, W.; Cabana, J.; Rousse, G.; Abakumov, A. M.; Tarascon, J.-M., Exploring the Bottlenecks of Anionic Redox in Li-Rich Layered Sulfides. Nat. Energy 2019, 4(11), 977–987; electrochemistry in (C) Leube, B. T.; Robert, C.; Foix, D.; Porcheron, B.; Dedryvere, R.; Rousse, G.; Salager, E.; Cabelguen, P. E.; Abakumov, A. M.; Vezin, H.; Doublet, M. L.; Tarascon, J. M., Activation of Anionic Redox in d(0) Transition Metal Chalcogenides by Anion Doping. Nat. Commun. 2021, 12(1), 5485.

28

Status of Li(Na)-based anionic redox materials for better batteries

upon Se substitution, was pushed forward to account for the activation of anionic redox activity in Li2TiS3
ySey. Equally, the anionic redox chemistry shows excellent kinetics in these anionic mixing compounds along the line of what has been reported for Li2FeS2
ySey99 as well as NaCrSSe,143 as the result of shifting towards more covalent Se system. Lastly, a third strategy to trigger anionic redox activity has consisted in synthesizing, via ball milling, cation-disorder Li2TiS3 or even Li3NbS4 with considerable electrochemical activities,85,144 although the science behind the activation in these pure d0-M based chalcogenides remains unclear and deserves further investigation. However, although Li-rich sulfides outperform Li-rich oxides in terms of fast kinetics and more stable structure, their low operational potential sacrifices the energy density, which actually goes against the initial intention of anionic redox. This dilemma could be partially solved by finding a balance between oxides and sulfides, such as oxysulfides, to optimize the holistic performances, given that the difficulty of the synthesis can be overcome. This stands as a difficult challenge to solid-state chemists.

7.02.4.4

Cation disorder, superstructure, and structural dimensionality

In addition to O/M ratio, alkaline ions, and ligands, another avenue for anionic redox is to play with the nature of the structure itself, such as cation disordering, superstructure, and structural dimensionality. These variables further enhances the rich material chemistry pertaining to anionic redox. “Cation-disordered” is generally termed to describe the cubic rock-salt structures, AxM2
xO2 (A ¼ alkaline ions), wherein the cations (A and M) randomly occupy the cationic site without ordering, as shown in Fig. 22A. It is also called as disordered rocksalt structures. Initially, cation-disordered Li compounds were expected to show poor capacity owing to the Li/M mixing that fully destroys the 2D Li diffusion channels. This is true for conventional cation-disorder structures like cubic a-LiFeO2 phase, which has a poor electrochemical performance.145 However, Ceder et al. found that an over Li-stoichiometry in Li[LixM1
x]O2 (x > 0.09) could trigger a percolating 0-TM Li diffusion network (Fig. 22B), hence opening a new direction for seeking high capacity electrode materials among the rich family of cation-disordered Li-rich oxides.146 Worth noting is that in these compounds, both

Fig. 22 (A) Schematic of the typical cation-disordered rock-salt structure. A: alkaline ions. M: transition metal. O: oxygen. (B) Computed probability of finding a percolating network of 0-TM channels (inset) versus Li content (x in LixTM2
xO2) and cation mixing (TMLi layers/TMTM layers 100%). (C) XRD patterns of a series of cation-disordered Li-rich oxides based on Li3NbO4. (D) The corresponding electrochemical performances of the compounds in (C). Adapted from: (B) Lee, J.; Urban, A.; Li, X.; Su, D.; Hautier, G.; Ceder, G., Unlocking the Potential of Cation-Disordered Oxides for Rechargeable Lithium Batteries. Science 2014, 343(6170), 519–22; (C and D) Yabuuchi, N.; Takeuchi, M.; Nakayama, M.; Shiiba, H.; Ogawa, M.; Nakayama, K.; Ohta, T.; Endo, D.; Ozaki, T.; Inamasu, T.; Sato, K.; Komaba, S., High-Capacity Electrode Materials for Rechargeable Lithium Batteries: Li3NbO4-Based System With Cation-Disordered Rocksalt Structure. Proc. Natl. Acad. Sci. U. S. A. 2015, 112(25), 7650–7655.

Status of Li(Na)-based anionic redox materials for better batteries

29

excess of Li and cation-disorder will facilitate the formation of Li-O-Li configurations and therefore the O 2p NB states, arousing large anionic redox activity. A first demonstration of this activity was pioneered by Yabuuchi et al. in studying Li3NbO4-based cation-disordered compounds (Fig. 22C and D).118 This led to a burgeoning research activity mainly based on d0 transition metal ions, such as Li2TiO3-, Li4MoO5-, Li3TaO4-, Li4WO5-based oxides that however are of low electronic conductivities.78,79,108,119 Moreover, their relative lower covalence cast challenges in stabilizing the oxidized oxygen species, as inferred from DFT calculations.91 Thus, as no surprise, this family of electrodes generally exhibits lower activity especially at room temperature and in bulk-sized particles together with poorer cycling stability, compared with layered ones. A fluorination strategy was frequently pushed to alleviate the capacity degradation problem in these cation-disordered compounds, as substitution O2
with F
can lower the M valence hence triggering more cationic redox at the cost of oxygen redox.147–149 However, one overlooked aspect of anionic redox behavior in cation-disordered system is the inhomogeneous oxygen environments, which differs from the homogeneous ones in layered structure. For example, the oxygen in layered Li2MO3 is generally surrounded by two Li and four M ions (OLi2M4), whereas in full cation-disordered systems a variety of environments can be present, ranging from OM6 to OLi6 (Fig. 23), though each case has different probabilities. As the energy of O 2p NB band is only dictated by the electrostatic field exerted by its surrounding cationic charges, the more the Liþ around it, the less the electrostatic attraction, hence the higher the energy of O 2p NB electrons (Fig. 23). These O with different environments will behave differently during oxidation due to their different O 2p (NB) energy levels. However, it should be noted that some and maybe most of the cationdisordered compounds have structured diffuse scattering due to short-range order of the Li and M cations.81 This short-range order is characteristic of oxygen coordination environment that is as homogeneous as possible, such as OLi3M3 in LiMO2 or OLi2M4 in Li2MO3, in order to preserve the local charge neutrality. In these compounds, anionic redox may behave similarly with the one in layered structures.

Fig. 23 Schematic of various local O environments (OLixMy) in Li-rich cation-disordered structure. x and y indicates the number of Li and M ions surrounding the O ions. For OM6 and OLi1M5, no O 2p NB states is possible since no Li-O-Li is present. From OLi2M4 to OLi6, the decreasing number of M surrounding the O leads to weaker electrostatic attraction between O 2p charge and M, hence the higher energy of O 2p NB electrons. The end members (OLi6 and OM6) have the lowest probability to occur, but the probability increase from ends to the middle configurations (OLi4M2 and OLi3M3) according to the stoichiometry of Li/M (typical ranges from 1 to 2).

30

Status of Li(Na)-based anionic redox materials for better batteries

Cation disorder may also appear locally, such as Li/M disorder in transition metal layer, which we herein called as “intra-layer cation disorder.” Such intra-layer cation disorder will eliminate the superstructure ordering, such as LiM6 honeycomb ordering in monoclinic Li2MO3 structure, thus symmetrizing the lattice to an R-3m rhombohedral structure. In some cases, this intra-layer cation disorder was supposed to play an important role in governing the anionic redox behavior in Li-rich cathodes. This is illustrated by the recent synthesis of an intra-layer cation disordered Li2RuO3, obtained by an ion-exchange method from intra-layer Na/ Ru disordered Na2RuO3, that does not show evidence of O-O dimerization during anionic redox process, thereby suppressing the oxygen release with enhanced cycling stability.150 Notably, a similar strategy was applied for synthesizing a Li0.7[Li0.22Mn0.76]O2151 or Li0.6[Li0.2Mn0.8]O2152 phase with intra-layer Li/Mn disorder, showing high sustainable capacities approaching 300 mA h g
1 owing to successive O and Mn redox. These studies provide some hopes for finding solutions to overcome the bottlenecks of anionic redox, albeit the unclear mechanism of the merits brought by intra-layer cation disordering. Another way to diversify the Li-rich anionic redox chemistry is playing with the superstructure ordering. Since “extra Li” is pivotal for Li-rich materials to generate O 2p NB states in triggering anionic redox, where to locate these extra Li-ions is important in pursuit of more robust framework for anionic redox. The superstructure in Li-rich materials can range from the typical honeycomb ordering in most Li2MO3 compounds to zig-zag chain ordering in Li3RuO4, or Nb4O16 clusters in Li3NbO4, and even to isolated metal ions in Li5ReO6. Although this diversity is established on increasing the Li/M ratio, another lever that can modulate superstructure ordering is to combine two structures having different superlattice together, such as the combination of Li3RuO4 and Li3NbO4, hence producing helical chain ordering in Li3Ru0.7Nb0.3O4 or Li3Ru0.5Nb0.5O4 or jagged chain ordering in Li3Ru0.3Nb0.7O4 (Fig. 24A).46 The benefits of such superstructure engineering can be well reflected by the sodium layered compounds with honeycomb, ribbon, and mesh superstructure alike P2-Na0.75Li0.25Mn0.75O2, P2-Na0.6Li0.2Mn0.8O2, and Na2Mn3O7, respectively, in view of their different local structural response towards anionic redox activity (Fig. 24B).101 Specifically, due to higher Li/vacancy spreading in transition metal layer, honeycomb ordering is much more susceptible to cation migration and oxygen dimerization than ribbon and mesh orderings. The latter two cases have, therefore, less first-cycle voltage hysteresis compared to honeycomb structure. A last strategy for enriching the Li-rich chemistry is moving from two-dimension (2D) to three-dimension (3D) structure, as well represented by the transition from 2D a-Li2IrO3 to 3D b-Li2IrO3 (Fig. 25). The latter shows a high reversible anionic redox capability (2.5e
per formula unit) that is sustained upon cycling and free of cation migration and layer gliding.43 Therefore, this finding of anionic redox activity in 3D ordered network unlocks more possibilities in terms of structural dimensionality for future highcapacity cathodes design. Such 3D dimensionality effect can be equally found in b-Na1.7IrO3 compound, which can uptake nearly 1.3 Naþ based on the cumulative activity of cationic (Ir5þ/Ir4þ) and anionic ((O2)n
) networks while preserving the structure nicely.153 Overall, driven by anionic redox research, a vast number of new compounds were synthesized and unlocked for potential Li-ion cathodes. These also include other kinds of structures that are not commonly used in Li/Na electrodes, such as Li-rich anti-fluorite oxides (e.g., Li5FeO4),154 octalithium compounds Li8MO6,155 the early reported pnictogenides for anode materials,156,157 as well as some recently arisen organic frameworks.158 Apparently, thanks to anionic redox paradigm, a rich material chemistry is emerging up and more new compounds pertaining to this topic await to be explored.

7.02.5

Practical issues and their fundamental understandings

Li-rich layered oxides display outstanding energy densities associated to the emergence of anionic redox activity. However, an exceeding energy density is not sufficient by itself to warrant the practical use of these new phases in in Li-ion batteries. Indeed, several roadblocks preventing their real-world implementation need to be solved. These include chemical and electrochemical irreversibility, sluggish kinetics, voltage hysteresis, and voltage fade, which are largely counteracting the merits brought by anionic redox. A lot of studies has been taken to explore the fundamental origins of these issues while offering potential solutions to settle them, as expanded below.

7.02.5.1

Chemical and electrochemical irreversibility

For anionic redox, removing electrons from ligands requires partial ligand dimerization for stabilizing electron holes, causing a series of local structural distortions as the origins for its chemical irreversibility. This differs from cationic redox scenario simply involving bond contraction/stretching, that is less destructive to the local structure. Oxygen release is the main concern for oxygen redox, as it is not only determinative for electrochemical reversibility, but also related to safety issues by providing oxidizing agent for fire accident. The origin of oxygen release can be explicitly linked to the over-oxidation of oxygen ion during charge. However, the structural and chemical complexity of Li-rich oxides functioning through oxygen redox make it difficult to pinpoint the threshold of oxygen release during oxidation. This was recently overcome by a theoretical study coming up with the number of holes  per oxygen (h ) as an indicator for predicting the reversibility of oxygen redox.60 It was proposed that, in alkaline-rich transition  metal oxides (A-rich-TMOs), maintaining h < 1/3 is essential for reversible oxygen redox, based on a unified picture of anionic redox that helps gauge the extra anionic-redox capacities. This is quite a step forward, since based on this indicator, one can simply predict and control the reversibility of anionic redox in a given material. Nevertheless, specific considerations should be given to different Li-rich compounds that enlist different redox scenarios. They can be qualitatively classified into four categories (Fig. 26):

Status of Li(Na)-based anionic redox materials for better batteries

31

Fig. 24 (A) superstructure ordering of Li3RuyNb1
yO4 for y ¼ 1, 0.7, 0.5, 0.3, 0.1, and 0, with their Ru/Nb frameworks (MeM bonds only) shown above. (B) Superstructure ordering evolution from honeycomb, to ribbon, and to mesh, with the corresponding compounds indicated below. Reproduced from: (A) Jacquet, Q.; Perez, A.; Batuk, D.; Van Tendeloo, G.; Rousse, G.; Tarascon, J.-M., The Li3RuyNb1–yO4 (0  Y  1) System: Structural Diversity and Li Insertion and Extraction Capabilities. Chem. Mater. 2017, 29(12), 5331–5343; (B) House, R. A.; Maitra, U.; Perez-Osorio, M. A.; Lozano, J. G.; Jin, L.; Somerville, J. W.; Duda, L. C.; Nag, A.; Walters, A.; Zhou, K. J.; Roberts, M. R.; Bruce, P. G., Superstructure Control of First-Cycle Voltage Hysteresis in Oxygen-Redox Cathodes. Nature 2020, 577(7791), 502–508.

1. The first one is charge-transfer system as represented by Li2MO3 where M being inactive for redox process. In this system, oxygen solely participates redox process, with 2/3 holes per oxygen being created if Li ions are fully removed. This suggests a large irreversibility of oxygen redox that mostly ends with oxygen release, as confirmed by recent studies showing only irreversible oxygen redox in Li2MnO3.30 2. The second one is Li2M’O3 systems wherein M’ being redox active that can account one Li. This largely relieves the burden of oxygen in redox process and reduces the ho to a reversible level of 1/3, as exemplified by b-Li2IrO3 showing no oxygen release43 and the recently confirmed high reversible oxygen redox (96%) in Li2RuO3.29

32

Status of Li(Na)-based anionic redox materials for better batteries

Fig. 25 Structural dimensionality in Li-rich compounds showing anionic redox. (A) Structure of 2D a-Li2IrO3. (B) Structure of 3D b-Li2IrO3. Their electrochemical performance are shown correspondingly in (C) and (D). With a three dimensional network, the b-Li2IrO3 exhibits greater structural stability against oxygen redox and hence higher cycling stability. Figure adapted from: (B and D) Pearce, P. E.; Perez, A. J.; Rousse, G.; Saubanere, M.; Batuk, D.; Foix, D.; McCalla, E.; Abakumov, A. M.; Van Tendeloo, G.; Doublet, M. L.; Tarascon, J. M., Evidence for Anionic Redox Activity in a Tridimensional-Ordered Li-Rich Positive Electrode beta-Li2IrO3. Nat. Mater. 2017, 16(5), 580–586; (C) McCalla, E.; Abakumov, A. M.; Saubanere, M.; Foix, D.; Berg, E. J.; Rousse, G.; Doublet, M. L.; Gonbeau, D.; Novak, P.; Van Tendeloo, G.; Dominko, R.; Tarascon, J. M., Visualization of O-O Peroxo-Like Dimers in High-Capacity Layered Oxides for Li-Ion Batteries. Science 2015, 350(6267), 1516–1521.

3. The third class is the Li[LixM1
x
yM0 y]O2 system, with M being inactive while M0 being active in redox, covering most of the Lirich compounds reported, such as Li[Li0.2Ni0.13Mn0.54Co0.13]O2 or Li2Ru0.75Sn0.25O3. In this category, the active M0 redox partially replaces the O contribution, hence enabling partially reversible oxygen redox, as witnessed by numerous studies reporting reversible oxygen redox accompanied with some oxygen release at first charge.

Fig. 26

Four categories of Li-rich materials with different oxygen redox reversibility.

Status of Li(Na)-based anionic redox materials for better batteries

33

4. The last class refers to the case of Li-rich cation-disordered compounds. In this category, situation becomes more complicated due to the statistical distribution of the oxygen environment, as mentioned before in Fig. 23. For one specific O ion, the energy of its O 2p NB band is inversely proportional to the number of transition metal ions surrounding the O due to Coulombic interaction. Therefore, for the O ions with less surrounding transition metal ions, their O 2p NB bands will lie at higher energy levels (due to less electrostatic attraction) and will be more involved in redox process and are more prone to be over-oxidized and release as oxygen gas. Therefore, it is not that surprising to see oxygen release happens even at the beginning of oxygen participation in some cation-disordered compounds.159 A more realistic question regards how to mitigate or even suppress oxygen release by breaking the ho < 1/3 limit. As surface oxygen was shown to be the prime source of oxygen release, one efficient way is to anchor these surface oxygen ions by surface modification, as inspired from the trapping of oxygen molecules in lattice as a result of kinetically hindering their escaping.160 Enormous efforts were dedicated to this strategy in more practical compounds, for instance, Li-rich NMCs, by various surface modifications or heterostructure engineering, and they have been summarized in other review papers.161–163 However, homogenous coating can hardly be achieved with normal methods except for atomic layer deposition (ALD), hence calling for strategies that can thermodynamically stabilize the oxygen. This can indeed be realized by increasing the covalence between M and O, since, as suggested by DFT calculations, it preferentially promotes reversible anionic redox with higher oxygen release enthalpies.91 These two strategies (surface modification and increasing M-O covalence) is reminiscent of a recent study, which uncovers the dual roles of Ni and Co substitution in immobilizing the oxygen ions in Li2MnO3.164 This study highlights that a protective layer with less Li-rich but more Ni-rich rock-salt structure, and sole Co doping (high covalent Co-O), are both working for reducing oxygen release. A third direction could consist in acting on the ligand. Partially S substitution, namely, oxysulfides, might be able to suppress oxygen release with less voltage sacrifice, but runs into the difficulty of identifying the proper synthesis process. Alternatively, F substitution was also shown being effective in mitigating oxygen release in both layered165 and cation-disordered Li-rich oxides,147 as it lowers the valence of M and therefore releases more cationic redox capacity at the expense of oxygen redox, hence decreasing ho. Another unique feature that stands out the Li-rich layered oxides from conventional ones is the staircase to S-shape curve transition during first cycle. First spotted in Li-rich NM(C), the staircase behavior observed through the first charge was once supposed to be a two-step activation of LiMO2 and Li2MnO3 components successively, the latter being a characteristic 4.5 V flat plateau (Fig. 27A).166 It later turns out that the first slope was associated to the cationic activity of M in LiMO2 and the second plateau to reversible oxygen redox that ends with some oxygen release.39 Due to the local structural distortion induced by oxygen redox, both inter- and intra-layer irreversible cation migration are triggered during the 4.5 V O-redox plateau.77,106 Such irreversible cation migration will create different local environments that have different potentials in terms of Li (de)intercalation, hence causing an Sshape curve during discharge (Fig. 27A). What is interesting is this irreversible cation migration can be easily recovered by heating the fully discharged sample at a temperature of over 150  C, as inferred from the reemerged honeycomb superstructure peak in X-ray diffraction (XRD) after heating (Fig. 27B).106 Accompanied with this structural reordering is the reemergence of the staircase curve during charge (Fig. 27C), though which is converted to an S-shape curve again during subsequent discharge. This further consolidates the relationship between staircase to S-shape curve transition with cation migration. For practical application, such first-cycle reorganization is problematic but can be partially overcome by formation process. Though being common in some layered oxides, the staircase to S-shape transition is absent in structures like a-Li2IrO3 and Na2RuO3 due to their more robust local structures that are almost immune to cation migration. Specially noted is that a tridimensional structure, represented by b-Li2IrO3 as previously mentioned, can be more insusceptible to cation migration and hence shows better reversibility and cycling stability. This inject confidence for us to design strategies to deal with the chemical and electrochemical irreversibility.

7.02.5.2

Voltage fade

A prominent but non-negligible issue faced by Li-rich oxides is voltage fade, either called as voltage decay or voltage drop, that corresponds to a continuous decrease of the cell output voltage upon cycling (Fig. 28). It is detrimental for practical applications and even more troublesome than capacity decay since it not only causes attenuation of energy density, but also complicates the prediction of states of charge by voltage, being challenging for battery management for large-scale use. The origin of the voltage fade was early believed to be associated to the layered-to-spinel phase transition. This is because upon cycling the voltage continuously fades to values approaching the one of typical cubic-to-tetragonal phase transition reaction in spinel oxides. Such assignment was supported by TEM studies67 that unambiguously unraveled the formation of spinel phase in Li-rich NM(C) compounds (Fig. 29A) although it could not be detected by XRD due to the local-range order of spinel domain or the strain/defect effects that broadens the diffraction peaks. However, the legitimacy of this viewpoint was challenged by the absence of the typical high-voltage ( 4 V) plateaus of spinel phases. Thus, several other explanations were also brought forward. One of them proposed that, the oxygen release, which triggers the activation of lower-voltage Mn3þ/Mn4þ and Co2þ/Co3þ redox couples in addition to original ones enlisting Ni2þ/Ni4þ and O2þ/O(2
n)
, should be responsible for voltage fade (Fig. 29B).69 Another explanation, as deduced from operando three-dimensional Bragg coherent diffractive imaging, links voltage fade with the formation of dislocations that causes local strain perturbing the Li site energies (Fig. 29C).68 It was further corroborated by the ability to recover the initial voltage via heating the sample to eliminate the dislocations. Asides from the above explanations, irreversible cation migration has also been established very early by K. G. Gallagher et al. as the origin of voltage fade for Li-rich NMC materials.70 Evidence for irreversible cation migration was deduced by NMR study (Fig. 29D),71 hence enabling the direct

34

Status of Li(Na)-based anionic redox materials for better batteries

Fig. 27 (A) First-cycle reorganization of Li-rich NMC materials. Accompanying the electrochemistry, structure snapshots are shown in pristine, charged, and discharged states. For charged states, The oxygen vacancies, created due to oxygen release, are refilled implying their long-range migration from bulk to surface with subsequent annihilation. In parallel, M ions migrate to the Li vacancies in Li layer, causing the disappearance of the honeycomb ordering. In total, this was termed as “densification” as the M/O ratio is increased (see the chemical formula evolution). (B) Superstructure peak evolution before and after mild-temperature annealing. (C) Cycling curve before and after annealing. The staircase curve reemerges again after annealing, which is due to the recovery of the Mn migration, as shown by the inset. Adapted from: insets in (A–C) Yin, W.; Grimaud, A.; Rousse, G.; Abakumov, A. M.; Senyshyn, A.; Zhang, L.; Trabesinger, S.; Iadecola, A.; Foix, D.; Giaume, D.; Tarascon, J. M., Structural Evolution at the Oxidative and Reductive Limits in the First Electrochemical Cycle of Li1.2Ni0.13Mn0.54Co0.13O2. Nat. Commun. 2020, 11(1), 1252.

link between cation migration and voltage fade. Simply speaking, irreversible cation migration perturbs the Li site energies and thereby lowers the potential. Separately, this linkage between cation migration and voltage fade was also confirmed on two archetypical oxides, Li2Ru1-ySnyO3 and Li2Ru1-yTiyO3, with the former showing negligible voltage fade whereas the latter being the opposite.72 Such contrast was rationalized by scrutinizing their local structural changes upon cycling with TEM and XPS, revealing a progressive Ti migration from octahedral to tetrahedral site in Li2Ru0.75Ti0.25O3 (Fig. 29E) whereas less prominent cation migration in Li2Ru0.75Sn0.25O3. The study therefore built a robust relationship between irreversible cation migration and voltage fade and also suggested solutions to cope with this issue by using larger sized M ions, such as Sn4þ.

Fig. 28

Voltage fade in Li-rich NMC cathode.

Status of Li(Na)-based anionic redox materials for better batteries

35

Fig. 29 Several typical views on the origin of voltage fade. (A) Layered-to-spinel phase transformation was captured by TEM and was supposed to be the origin of voltage decay in Li1.2Ni0.2Mn0.6O2. (B) Redox couple evolution in Li1.2Ni0.15Co0.1Mn0.55O2 quantified by XAS. The activation of Co2þ/ Co3þ and Mn3þ/Mn4þ redox couples at low potential was believed to cause the voltage fade. (C) Dislocation defects in Li-rich layered oxide were supposed to be the cause of voltage decay. The dislocations were monitored by operando three-dimensional Bragg coherent diffractive imaging, as shown in the right panel. After heat treatment, the voltage was recovered, as indicated by the arrows in the figure at left panel. (D) Capturing the irreversible octahedral to tetrahedral TM migration in Li-rich NMC compound via 6Li NMR spectroscopy. (E) Voltage fade in Li2Ru0.75Sn0.25O3 was linked to irreversible Ti migration from octahedral site to tetrahedral site, as shown by the TEM figure in the right panel. An intensity profile along the TM column is attached below to reveal tetrahedral Ti (green). Figures adapted from: (A) Gu, M.; Belharouak, I.; Zheng, J.; Wu, H.; Xiao, J.; Genc, A.; Amine, K.; Thevuthasan, S.; Baer, D. R.; Zhang, J. G.; Browning, N. D.; Liu, J.; Wang, C., Formation of the Spinel Phase in the Layered Composite Cathode Used in Li-Ion Batteries. ACS Nano 2013, 7(1), 760–767; (B) Hu, E.; Yu, X.; Lin, R.; Bi, X.; Lu, J.; Bak, S.; Nam, K.-W.; Xin, H. L.; Jaye, C.; Fischer, D. A.; Amine, K.; Yang, X.-Q., Evolution of Redox Couples in Li- and Mn-Rich Cathode Materials and Mitigation of Voltage Fade by Reducing Oxygen Release. Nat. Energy 2018, 3(8), 690–698; (C) Singer, A.; Zhang, M.; Hy, S.; Cela, D.; Fang, C.; Wynn, T. A.; Qiu, B.; Xia, Y.; Liu, Z.; Ulvestad, A.; Hua, N.; Wingert, J.; Liu, H.; Sprung, M.; Zozulya, A. V.; Maxey, E.; Harder, R.; Meng, Y. S.; Shpyrko, O. G., Nucleation of Dislocations and their Dynamics in Layered Oxide Cathode Materials during Battery Charging. Nat. Energy 2018, 3(8), 641–647; (D) Dogan, F.; Long, B. R.; Croy, J. R.; Gallagher, K. G.; Iddir, H.; Russell, J. T.; Balasubramanian, M.; Key, B., Re-Entrant Lithium Local Environments and Defect Driven Electrochemistry of Li- and Mn-Rich Li-Ion Battery Cathodes. J. Am. Chem. Soc. 2015, 137(6), 2328–2335; (E) Sathiya, M.; Abakumov, A. M.; Foix, D.; Rousse, G.; Ramesha, K.; Saubanere, M.; Doublet, M. L.; Vezin, H.; Laisa, C. P.; Prakash, A. S.; Gonbeau, D.; VanTendeloo, G.; Tarascon, J. M., Origin of Voltage Decay in High-Capacity Layered Oxide Electrodes. Nat. Mater. 2015, 14(2), 230–238.

Altogether, this led to the consensus that irreversible cation migration causes voltage fade, as other perspectives are also highly related with cation migration. Additional studies show that irreversible cation migration is rooted in anionic redox that induces large structural distortion and hence opens a space/path for cations to migrate. This has been validated by the observation of negligible and high voltage fade in Li-rich NMC cycled with low (no anionic redox) and high (anionic redox) charge cut-off voltages, respectively.40 Moreover, one study based on Li2IrO3 and Li2Ir1-ySnyO3 also suggested that anionic redox is tightly associated with cation migration in order to trigger the decoordination of M-O bond in facilitating the formation of TM ¼ O p bond as well as short O-O species ( 1.4 Å).44 Given such linkage between cation migration and anionic redox, the latter are well acknowledged to be the ultimate origin of voltage fade. However, cation migration is not universal for all Li-rich electrodes exhibiting anionic redox. This provides implications on how to design Li/Na-rich electrodes free of voltage fade, such as the 3D b-Li2IrO3, which is intrinsically immune to cation migration and hence offers stable voltage during cycling.43 On the other hand, several strategies can be applied to counter voltage fade in terms of material design. One way is to use large sized M ions that can hardly migrate as exemplified by previously mentioned Li2Ru172 A similar strategy is to act on the layer spacing rather than cation itself to prevent the migration. This was hinted by ySnyO3. the newly reported O3-NaLi1/3Mn2/3O2 compound with no discernible voltage fade when cycled versus Na, which is due to the larger Na layer spacing that disfavors the accommodation of transition metal ions and therefore suppresses cation migration.50 Another strategy is to engineer the local structure to enable highly reversible cation migration. This is well represented by the O2-type Lix[Li0.2Ni0.2Mn0.6]O2 compound that shows remarkably mitigated voltage fade compared with its O3 counterpart.73 The improvement was shown to be nested in the altered oxygen stacking sequence that imposes strong repulsion towards the migrated cations, thereby preventing it being permanently trapped in lithium layer. Therefore, voltage fade is problematic but solutions can be found for real-world applications. Tremendous efforts have been devoted to developing strategies like surface

36

Status of Li(Na)-based anionic redox materials for better batteries

modification, bulk doping, and heterostructures to successfully mitigate voltage fade in more practical cathodes, such as Li-rich NMC oxides (see other review papers167,168).

7.02.5.3

Voltage hysteresis

Another practical issue that further counteracts the high-energy-density benefits brought by anionic redox is voltage hysteresis, that is, a prominent voltage gap or difference between charge and discharge (Fig. 30A), which penalizes the use of Li-rich electrode with low energy efficiency. This is not good for either electric vehicles or stationary energy storage, as it not only induces cost issue due to energy loss but also produces heat that causes safety concern. Therefore, it is imperative to understand its underlying origin so as to develop solutions to deal with it. Similar with voltage fade, voltage hysteresis was also first noticed in Li-rich NMC oxides, which generally exhibit a voltage gap of around 400–500 mV (Fig. 30A). However, such a voltage gap is indeed nested at interweaved cationic (Ni and/or Co) and anionic redox. By decoupling oxygen redox from cationic redox by means of voltage-window-opening experiment, it has been shown that oxygen redox actually arouses a large voltage hysteresis up to  1 V in contrast to negligible hysteresis for cationic redox (Fig. 30B and C). The large hysteretic gap for anionic redox was initially attributed, alike what has been taken for voltage fade, to cation migration that perturbs the local Li site energies and hence its (de)intercalation potential. Unlike voltage fade linked to irreversible cation migration, reversible cation migration was supposed to be responsible for voltage hysteresis (Fig. 31A and B), as it should enable a loop from which the path of charge and discharge are asymmetrical.71,74,77 Such reversible cation migration, as seen by NMR and XRD techniques, was considered to cause a structural hysteresis with a path-dependent feature,74 that is, charge to different SoC produces different discharge path and vice versa, as shown in Fig. 31C. In light of this path-dependent feature, it was concluded (A)

Voltage (V) vs. Li/Li+

5.0

Cha rge

4.5 4.0

Voltage hysteresis

Dis ch arg e

3.5 3.0 2.5 2.0 0.0

0.2

0.6

0.8

1.0

x in LixNi0.13Mn0.54Co0.13O2

(B)

(C) 1.5

2nd cycle

5.0

2.0-4.8 V 2.0-4.3 V

1.0

4.5

2.0-4.8 V 2.0-4.3 V

4.0 3.5 3.0

Ni/Co oxidation O oxidation

0.5

dQ/dV

Voltage (V) vs. Li/Li+

0.4

0.0 -0.5 Ni/Co reduction

2.5

-1.0

2.0

O reduction

-1.5 0.0

0.2

0.4

0.6

0.8

x in LixNi0.13Mn0.54Co0.13O2

1.0

2.0

2.5

3.0

3.5

4.0

4.5

5.0

+

Voltage (V) vs. Li/Li

Fig. 30 Voltage hysteresis in Li-rich NMC compound. (A) 2nd cycle of Li-rich NMC showing a prominent voltage hysteresis. (B) The cycling curve of Li1.2Ni0.13Mn0.54Co0.13O2 at different voltage windows in the 2nd cycle. (C) Corresponding dQ/dV plots of the cycling curve in (B). The Ni/Co and O redox peaks are marked out according to previous spectroscopic results. The O redox area is shaded with pink to highlight the voltage hysteresis behavior.

Status of Li(Na)-based anionic redox materials for better batteries

37

that the voltage hysteresis is nested at a thermodynamic origin, which is further corroborated by the galvanostatic intermittent titration technique (GITT) measurements that never closed at nearly zero current, hence the so-called “quasi-static” voltage hysteresis.40 Besides, cation migration may cause the change of Li site energies, as supported by some DFT calculations,77 but from synchrotron XRD performed in Li-rich NMC oxide, around 4% reversible cation migration was determined that is far from substantial to induce a large voltage hysteresis behavior. Additionally, cation migration was also recently found in NaCrSSe that shows with

Fig. 31 Several existing views on the origin of voltage hysteresis in anionic-redox compounds, including reversible cation migration (A-C), redox inversion (D), and O-O dimerization (E and F). (A) dQ/dV plot of the first cycle of Li-rich NMC compound showing a large hysteresis of O redox (red) while low voltage hysteresis of TM (or TM-O hybridized states) redox (green). (B) Evolution of the amount of TM ions in Li layer (TMLi) during first cycle. It shows around 4.7% TM reversible migration, which was supposed to cause the variation of O 2p band and hence voltage difference between charge and discharge. (C) Path-dependent feature by gradually opening the charge voltage window. (D) Redox mechanism of Li1.3Ni0.27Ta0.43O2 showing a Ni / O oxidation on charge but Ni / O reduction on discharge. (E) Schematic of charge/discharge curve with large voltage hysteresis when O-O dimerization is present, which causes different Gibbs energy change for charge and discharge, hence the voltage hysteresis (polarization). (F) Large voltage hysteresis in Li4Ni0.85WO6 compound, which was supposed to be related to the O-O dimer formation as observed from the Raman spectroscopy and confirmed by DFT. Figures adapted from: (A and B) Gent, W. E.; Lim, K.; Liang, Y.; Li, Q.; Barnes, T.; Ahn, S. J.; Stone, K. H.; McIntire, M.; Hong, J.; Song, J. H.; Li, Y.; Mehta, A.; Ermon, S.; Tyliszczak, T.; Kilcoyne, D.; Vine, D.; Park, J. H.; Doo, S. K.; Toney, M. F.; Yang, W.; Prendergast, D.; Chueh, W. C., Coupling between Oxygen Redox and Cation Migration Explains Unusual Electrochemistry in Lithium-Rich Layered Oxides. Nat. Commun. 2017, 8(1), 2091; (C) Croy, J. R.; Gallagher, K. G.; Balasubramanian, M.; Chen, Z.; Ren, Y.; Kim, D.; Kang, S.-H.; Dees, D. W.; Thackeray, M. M., Examining Hysteresis in Composite xLi2MnO3$(1–x)LiMO2 Cathode Structures. J. Phys. Chem. C 2013, 117(13), 6525–6536; (D) Jacquet, Q.; Iadecola, A.; Saubanere, M.; Li, H.; Berg, E. J.; Rousse, G.; Cabana, J.; Doublet, M. L.; Tarascon, J. M., Charge Transfer Band Gap as an Indicator of Hysteresis in Li-Disordered Rock Salt Cathodes for Li-Ion Batteries. J. Am. Chem. Soc. 2019, 141(29), 11452–11464; (E) Tsuchimoto, A.; Shi, X. M.; Kawai, K.; Mortemard de Boisse, B.; Kikkawa, J.; Asakura, D.; Okubo, M.; Yamada, A., Nonpolarizing Oxygen-Redox Capacity Without O-O Dimerization in Na2Mn3O7. Nat. Commun. 2021, 12(1), 631; (F) Taylor, Z. N.; Perez, A. J.; Coca-Clemente, J. A.; Braga, F.; Drewett, N. E.; Pitcher, M. J.; Thomas, W. J.; Dyer, M. S.; Collins, C.; Zanella, M.; Johnson, T.; Day, S.; Tang, C.; Dhanak, V. R.; Claridge, J. B.; Hardwick, L. J.; Rosseinsky, M. J., Stabilization of O-O Bonds by d0 Cations in Li4 þ xNi1-xWO6 (0 x0.25) Rock Salt Oxides as the Origin of Large Voltage Hysteresis. J. Am. Chem. Soc. 2019, 141(18), 7333–7346.

38

Status of Li(Na)-based anionic redox materials for better batteries

negligible voltage hysteresis,143 while large hysteresis was frequently observed in anion-redox compounds like Na2RuO347 and Na2/ 51 3[Mg0.28Mn0.72]O2 which are free of cation migration. It also remains to reconcile why a more reversible cation migration, found in a O2-type Lix(Li0.2Ni0.2Mn0.6)O2 compound, favors the mitigation of voltage hysteresis (asymmetry) rather than aggravating it.73 Altogether, these results negate the direct linkage between these two events, and a more sound explanation awaits to be found. Aside cation migration, anion-cation redox inversion was also put forward as the origin of voltage hysteresis. Such phenomenon, describing a cationic / anionic oxidation processes in charge while a cationic / anionic reduction in discharge, is quite unusual, as it violates the conventional wisdom of electrochemistry which enlists sequential oxidation and reduction reaction (redox couple). Nevertheless, this phenomenon has been found in several layered or cation-disordered compounds like Li-rich NMC (Fig. 30C) and Li1.3Ni0.27Ta0.43O278 (Fig. 31D), and it was explained by a band-inversion theorydoxygen dimerization that causes the O(2p) band splitting and thus leads to a higher O 2p antibonding states than unoccupied M nd states.60 Although this explanation accounts for the voltage hysteresis nicely, it still remains as a hypothesis that is hard to be proven experimentally. Besides, let’s keep in mind that band structure is a thermodynamic approach and therefore can hardly explain the kinetic effect, as shown by the GITT curves which marked a large recovery of the voltage during relaxation. Another elusive explanation recently put forward to account for such a large hysteresis in anionic redox compounds relies on reversible oxygen dimerization. It was indeed proposed that the reversible OeO bond (re)formation/breaking imposes energy penalty during charge and discharge, hence producing voltage hysteresis (Fig. 31E). This scenario was inspired from Li-driven conversion-type reactions taking place in transition metal oxides that proceed with large voltage hysteresis whose origin is supposed to be reversible MeO bond (re)formation/breaking. Evidences for this explanation in Li-rich oxides were provided by Li4 þ xNi1
xWO6 showing > 2 V voltage hysteresis with oxygen dimerization79 (Fig. 31F), as well as Na2Mn3O7 free of voltage hysteresis as no trace of oxygen dimer was captured.80 Nevertheless, it remains to establish a more robust scientific link between O-O dimerization with voltage hysteresis. At this stage, although the above explanations rely on different scientific causes, it is interesting to note that they can be unified by a framework in which the thermodynamics and kinetics of voltage hysteresis are well integrated (Fig. 32A–C) thanks to a microcalorimetry study of model Li2Ru0.75Sn0.25O3 compound.169 Here, an electrochemical and a chemical step was considered for the anionic redox process, the latter consisting of sluggish local structural reorganizations, such as cation migration, oxygen dimerization, layer gliding and so on. The chemical step with sluggish kinetics constitutes the energy penalty on which voltage hysteresis can be established, as shown by the free-energy diagram in Fig. 32C. Whichever the chemical reaction is, such framework clarifies the thermodynamic and kinetic picture of voltage hysteresis that is frequently missing in previous studies. The correctness of the above framework was recently supported by studies conducted on Li1.17Ti0.33Fe0.5O281 with the exception that the chemical step was not merely a sluggish structural reorganization, but also involves charge transferdwhich is true since electronic structure dictates local structure. It was found that, instead of following an equilibrium redox pathway involving sole O redox, Li1.17Ti0.33Fe0.5O2 prefers undergoing a non-equilibrium redox pathway enlisting Fe3þ/Fe4þ and O2
/O(2-n)- redox during charge (Fig. 32D), hence causing large voltage hysteresis. Using sensitive but harmless Mössbauer spectroscopy, this Fe4þ was demonstrated to be an intermediate state that can be eliminated by properly relaxing the electrode due to an internal ligand-tometal charge transfer process (or reductive coupling). The formation of Fe4þ was further rationalized by Marcus theory (Fig. 32E) in a way that the electron removal from O 2p lone-pair states, involving large structural distortion, is a non-adiabatic process (i.e., high kinetic barrier), whereas the electron removal from Fe3þ, with facile FeO bond contraction, is an adiabatic process (i.e., low kinetic barrier). This implies the ultimate origin of the sluggish kinetics and voltage hysteresis in Li-rich oxides is rooted in the nature of O 2p lone-pair states, which is very localized and can hardly be accessed with fast kinetics. Benefited by this finding, one can, therefore, understand why oxygen participation via (M-O)* anti-bonding states are often observed with negligible voltage hysteresis, such as the cases of LiCoO2 and Nax[Ni1/3Mn2/3]O2 phases, as oxygen redox could proceed by direct electron removal from these delocalized states that is free of reductive coupling process. The same situation happens for Na2Mn3O7 and Na0.6[Li0.2Mn0.8]O2 in which the oxygen redox takes place by involving O 2p localized holes or delocalized p states,64,101 bearing similarity with delocalized (M-O)* states. At this stage, having provided a satisfactory fundamental picture of the large voltage hysteresis in Li-rich oxides, it remains to dwell on means to mitigate or eliminate voltage hysteresis. This is not an easy task as witnessed by the limited amount of reported examples mainly focusing on moving from 3d to 4d/5d transition metal elements (Li2Ru0.75Sn0.25O3 or Li2IrO3) or replacing O by S/Se ligands (Li1.13Ti0.57Fe0.3S2 and Li2TiS3
ySey). In both cases, by increasing the metal-ligand covalence, we facilitate charge transfer and minimize the lifetime of the exited cationic intermediates. Optimism must therefore prevail after we fully uncovered the unified picture of the fundamental origin of the voltage hysteresis.

7.02.6

Conclusions and outlook

With almost a decade’s intense investigation, anionic redox has grown from its infancy to a giant that predominates today’s cuttingedge research for exploring new electrode materials for high-energy density Li-ion batteries. It’s encouraging that such a paradigm shift has spurred the discovery of a colossal amount of new phases spanning over not only Li-ion hosts but also beyond, and cultivated rich fundamental science for material chemistry as well as electrochemistry. Though making it to be practical is stumbling, the future is still full of hope for developing better electrode materials for Li(Na)-ion batteries thanks to the advances of anionic redox chemistry. To end this chapter, we would like to summarize some unanswered questions for the fundamental understandings of

Status of Li(Na)-based anionic redox materials for better batteries

39

Fig. 32 (A) Single electrochemical step of cationic redox. (B) “Square-scheme” mechanism of the anionic redox. “A” denotes the anionic species in the reduced state (for example, O2
) that on oxidation forms unstable Aþ (for example, On
with n < 2). A structural reorganization was supposed to happen spontaneously after every electrochemical step. Ar þ denotes the stabilized species after structural reorganization while Ar denotes the unstable reduced species during discharge that can further relax to A state via another structural reorganization step. (C) Energy landscape of the corresponding “square-scheme” mechanism of the anionic redox. (D) Schematic of the non-equilibrium redox path and equilibrium redox path in Li1.17Ti0.33Fe0.5O2 compound. The voltage hysteresis was linked to the charge transfer from O2
to Fe4þ that involves large structural distortion. (E) Rationalizing the charge transfer mechanism in Li1.17Ti0.33Fe0.5O2 compound. Charge transfer directly from O lone-pair states (Fe3þ-O2
/ Fe3þOn
) was supposed to be a non-adiabatic process and hard to happen, whereas the charge transfer from s-type Fe(3d)-O(2p) delocalized states is adiabatic process that can be proceeded easily. This leads to the formation of kinetically trapped Fe4þ species and then disappear slowly via a ligandto-metal charge transfer process. Figures reproduced from: (A–C) Assat, G.; Glazier, S. L.; Delacourt, C.; Tarascon, J.-M., Probing the Thermal Effects of Voltage Hysteresis in Anionic Redox-Based Lithium-Rich Cathodes Using Isothermal Calorimetry. Nat. Energy 2019, 4(8); (D–E) Li, B.; Sougrati, M. T.; Rousse, G.; Morozov, A. V.; Dedryvere, R.; Iadecola, A.; Senyshyn, A.; Zhang, L.; Abakumov, A. M.; Doublet, M. L.; Tarascon, J. M., Correlating Ligand-to-Metal Charge Transfer with Voltage Hysteresis in a Li-Rich Rock-Salt Compound Exhibiting Anionic Redox. Nat. Chem. 2021, 13(11), 1070– 1080.

anionic redox, and indicate how resolving these issues can help bridging the gap between fundamental studies and practical application. Besides, we also assess the practicality for Li-rich materials with anionic redox in Li(Na)-ion batteries against their competitors like (Ni-rich) NMC oxides, and shed new insights on future cathode design. The mysteriousness of anionic redox has been not fully wiped off even though numerous studies have been performed. Of primary importance is how anionic redox can be activated and the complication here resides in the dedicate balance between thermodynamic and kinetic aspects. One can find that, in Li2TiO3 or Li2TiS3, even the O 2p states and S 3p states predominate the vicinity of fermi level in absence of Ti 3d contribution, anionic redox cannot be activated until doping with redox-active metals or other ligands. This implies a hidden prerequisite for the activation of anionic redox in addition to the well-known band-

40

Status of Li(Na)-based anionic redox materials for better batteries

structure premise (O 2p NB states), and most probably due to some unknown kinetic parameters awaiting for further exploration. Deciphering this query will help us to smartly utilize anionic redox especially based on some assumed inactive electrodes, such as d0 or d10 transition metal based Li-rich sulfides. Besides, the activation of anionic redox was frequently linked to the so-called “reductive coupling mechanism,” or “ligand-to-metal charge transfer” named by other researchers, with, however, no full experimental evidence provided so far. A cationic intermediate species (Fe4þ) was indeed captured by Mössbauer spectroscopy in Li1.17Ti0.33Fe0.5O2 compound, but the O-to-Fe charge transfer still remains a hypothesis due to the short lifetime of Fe4þ. An answer may be provided in the future by seeking for long-lived cationic intermediates in suitable model compounds to resolve this mystery. It is crucial to fully validate the reductive coupling mechanism that is ruling anionic redox in terms of sluggish kinetics and voltage hysteresis, hence providing guidance for designing curing strategies to manage. Additional questions regard other anomalies coming along with anionic redox, such as the cationic-anionic redox inversion, need to be revisited by considering not only thermodynamics but also kinetics behind it. Application-wise, the main hope of commercializing anionic redox is being largely placed on Li-rich NMCdthe “holy grail” of cathode materials for Li-ion batteries. However, several roadblocks remain in spite of a great deal of efforts that have been made enlisting chemical-physical or surface science approaches. Nevertheless, little progress have been achieved towards accelerating their implementation in real-world applications. In parallel is the community trying to make a higher-energy density NMC oxide electrodes by putting more Ni inside, the so-called Ni-rich NMC materials, while beset by their poor thermal stabilities that cause safety concerns. The representative Ni-rich NMC811 outperforms Li-rich NMC for various aspects except for materials specific energy.3 Regardless of its lower cost and probably higher safety due to Mn rich, Li-rich NMC is still facing fierce competition against NMC811, unless its practical issues, such as voltage hysteresis and voltage fade, can be overcome. In terms of this, the future of anionic redox should rely on finding new Li-rich systems superior than Li-rich NMC, which is quite difficult, or, as we proposed here that, complementing Li-rich NMC with Ni-rich NMCddeveloping Li-rich Ni-rich NMC oxides. By this way, the merits of Li-rich NMC can be preserved while its drawbacks can be mitigated via increasing the usage of Ni at the cost of Mn, the latter being the source of voltage fade and poor kinetics. However, such strategy is rather a compromise than a solution since in this case the anion-redox contribution will be reduced while cationic redox (Ni) will increase. But at least, this combination of Li-rich and Ni-rich NMC could be an alternative choice to enable its competitiveness against Ni-rich oxides by delicately balancing the stoichiometry of Li and Ni. On the other hand, the surge of studies on Na-ion electrodes and sulfides with anionic redox is indeed being driven more by fundamental understandings rather than practical promises, although Li-rich sulfides could be important for solid-state batteries owing to their improved chemical compatibility with S-based ionic conductors (Li3PS4). Na-rich electrodes are only limited to costly 4d or 5d transition metals, whereas for 3d system it is meaningless to pursue anionic redox since they are generally under stoichiometric in Na (P2, and P3-type Na1
xMO2), and therefore less capacity that can be easily fulfilled by cationic redox. Unfortunately, when this situation is overcome as for O3-NaLi1/3Mn1/3O2, the material, besides its complex synthesis, shows poor kinetics and voltage hysteresis that make it less competitive against other sodium compounds. However, the future work can still be focused on some unique Li/Na systems represented by Na2Mn3O7 in which the cation vacancy-mediated superstructure could be a potential solution for tackling voltage hysteresis. Regarding Li(Na)-rich sulfides, although free of some practical issues like voltage hysteresis and voltage fade, they are suffering from lower working potentials and high-cost synthesis, and hence less practical than their oxide counterparts. Nonetheless, this under-explored area is still full of hope and surprise that may offer solutions in the future. In summary, the research of anionic redox is still ongoing and progresses are steadily being made. Its history has so far been rich of learning lessons starting first with the reluctance of the battery community to accept it so 10 years were lost after the first report of anionic redox back to 1999. It was overcompensated by the wonderful demonstration of the power of our community in tackling rapidly the fundamental aspects of a new paradigm in view of impactful applications. The story is not over and such a long fascinating journey further strengthens our confidence towards implementing anionic redox in real world applications. It is a difficult but not unsurmountable task that will still need few years as usually required for any research with transformational changes to impact our society.

Acknowledgment J.-M.T and B.L. acknowledge funding from the European Research Council (ERC) (FP/2014)/ERC Grant-Project 670116-ARPEMA.

References 1. 2. 3. 4. 5. 6. 7.

Tarascon, J. M.; Armand, M. Issues and Challenges Facing Rechargeable Lithium Batteries. Nature 2001, 414 (6861), 359–367. Whittingham, M. S. Lithium Batteries and Cathode Materials. Chem. Rev. 2004, 104 (10), 4271–4302. Assat, G.; Tarascon, J.-M. Fundamental Understanding and Practical Challenges of Anionic Redox Activity in Li-Ion Batteries. Nat. Energy 2018, 3 (5), 373–386. Whittingham, M. S. Electrical Energy Storage and Intercalation Chemistry. Science 1976, 192 (4244), 1126–1127. Whittingham, M. S.; Gamble, F. R. The Lithium Intercalates of the Transition Metal Dichalcogenides. Mater. Res. Bull. 1975, 10 (5), 363–371. Whittingham, M. S. The Electrochemical Characteristics of VSe2 in Lithium Cells. Mater. Res. Bull. 1978, 13 (9), 959–965. Whittingham, M. S. The Role of Ternary Phases in Cathode Reactions. J. Electrochem. Soc. 1976, 123 (3), 315–320.

Status of Li(Na)-based anionic redox materials for better batteries 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35.

36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48.

41

McCanny, J. V. A Theoretical Study of the Effects of Lithium Intercalation on the Electronic Structure of TiS2. J. Phys. C Solid State Phys. 1979, 12 (16), 3263–3276. Rouxel, J. Anion-Cation Redox Competition and the Formation of New Compounds in Highly Covalent Systems. Chem. A Eur. J. 1996, 2, 1053–1059. Jobic, S.; Brec, R.; Rouxel, J. Occurrence and Characterization of Anionic Bondings in Transition Metal Dichalcogenides. J. Alloys Compd. 1992, 178 (1), 253–283. Rouxel, J. Some Solid State Chemistry With Holes: Anion–Cation Redox Competition in Solids. Curr. Sci. 1997, 73 (1), 31–39. Mizushima, K.; Jones, P. C.; Wiseman, P. J.; Goodenough, J. B. LixCoO2 (0 100 mAh g 1 per 1 de/inserted Naþ ion), i.e. have the lowest possible molar mass; (3) exhibit a high enough operating potential within the electrochemical stability window of the electrolyte (normally higher than 2.5 V vs. Naþ/Na); (4) allow sufficient mobility of Naþ in the bulk; Additionally, such a compound is expected to possess electronic conductivity (intrinsic or grain boundary) in pristine or composite form (for instance, coated by a conductive material), and to display minimal volume variation as to not introduce severe mechanical stress. Similar to Li-based counterparts, positive electrode (cathode) materials for SIBs can be generally subgrouped as follows: (1) layered oxides with d-metal cations adopting a formal charge of þ2; (2) oxoanion-based compounds, mainly phosphates and sulfates10,11; (3) compounds, containing mixed-anion groups, for example, fluoride-phosphates, fluoride-sulfates, etc. In the following sections main classes of electrode materials are highlighted and several cases are described in detail.

7.03.2.1

Layered oxides of transition metals

The sodium layered oxides with the general formula NaxMO2 (M is a transition metal) have drawn a lot of attention due to attractive electrochemical performance, earth abundant chemistry (Fe, Mn, Ti based compounds) and easy-scalable production. A long and successful commercial history of LiCoO2 cathodes also supports a rapid development of this type of battery materials.12 The vast majority of NaxMO2 compounds crystallizes in the a-NaFeO2-type structure related to the rock salt type framework, where transition metal and Naþ ions are ordered into layers alternating along the threefold axis. Such arrangement allows tolerating a large difference between the ionic radii. The transition metal cations in the layers are in the octahedral MO6 sites edge-sharing with six other metal sites into a hexagonal pattern. Na occupies sites between transition metal sheets (Fig. 2). If a layered structure contains ions that exhibit Jahn-Teller distortions (for example Mn3þ or Ni3þ), the unit cell is capable of undergoing symmetry

Fig. 2 Schematic representation of polytypes for layered oxides. Blue bodies represent transition metal, yellow – alkali metal polyhedras respectively. The motifs of orientation of layers are designated with letters.

Electrode materials for reversible sodium ions de/intercalation

49

lowering.13 The sodium layered oxides may stabilize large variations of chemical compositions depending on the transition metal nature. The NaxMO2 compounds can be synthesized sodium-deficient down to 0.5 equivalent per formula unit. The temperature and atmosphere of the synthesis, nature of M, and the sodium content in the reaction media affect the resulting coordination of sodium being trigonal prismatic (P), octahedral (O), or tetrahedral (T) according to the notation proposed by Delmas.14 Various stacking modes of the layers bring the variety of polytypes for NaxMO2 materials (Fig. 2). The most stable among them are O3, P3 and P2, where the letter indicates a type of sodium coordination in the structure and the number designates the number of sodium layers within the unit cell (Fig. 2, here and elsewhere crystal structures were drawn using VESTA software package15). Most of the NaMO2 materials crystallize in the O3-type structure possessing hexagonal or rhombohedral symmetry, however, Ni- and Mn-based layered oxides experience a monoclinic distortion due to Mn3þ and Ni3þ being Jahn-Teller ions. The NaxMO2 compositions, where x ¼ 0.6  0.8, adopt a P2-type structure, while for x ¼ 0.5  0.6 the P3-type structures are more common. The O3, P3 and O1 type structures have the same orientation of the transition metal octahedra and therefore may convert into each other via gliding of the layers. The O3-type phases experience a reversible phase transition into P3-type when the structure loses Na during the charge process. The layers in the P2-type structure feature two different orientations of the transition metal octahedra, therefore making phase transitions into P3 and O3-type polytypes via sliding unavailable and requiring an additional 60 in-layer rotation.16 Due to the more prominent ionicity of the Na-O bond in comparison to Li-O the sodium layered oxides allow better stabilization of tetravalent transition metals in addition to trivalent. This fact is the key to the ability of Na-deficient phases to be synthesized by high-temperature annealing of V, Co and Mn precursors. Depending on the nature of the transition metal the stability regions of different polytypes related to the sodium content may vary. The synthesis temperature also affects the phase type. For example, the P3 phase decomposes above 600  C, but the O3 and P2 ones are stable in the 700–900  C range. During the charge discharge process the Na content varies inducing the phase transition O3-P3 via gliding of the layers. Such a process causes a large unit cell volume change (around 23%) leading to a rapid degradation and capacity loss during cycling.17 When sodium content in such phases drops below x ¼ 0.2–0.3 it is considered that there is the O1-type phase forming from the O3 and P3 phases and the O2-type one from the P2 phases. The explicit identification of the structure is, however, very tricky due to the formation of a large number of stacking faults in the system. The zero Na content (full deinsertion) was not achieved for the vast majority of sodium layered oxides materials, presumably due to trapping of residual Na ions between layers when they collapse.13 In general, the NaMO2 cycling curve is characterized by the presence of multiple voltage plateaus related to the ordering of Naþ ions for the specific compositions (Fig. 3). The origin of such orderings for octahedral types lies only in the strong Naþ-Naþ repulsive interactions, while in case of the prismatic environment the NaO6 and MO6 face-sharing plays a significant role on top of the Na-Na repulsion resulting in much more complicated structures. In some cases, the M-M interactions may occur leading to the formation of superstructures as for V3þ and V4þ ions.18 Such phase transitions provide a substantial change in the voltage related to the change of stabilization energy required for the ordered structure. The significant changes in the unit cell volume often cause cracking of the particles and consequent increase in the resistance of the material. Overall, phase transformations cause degradation of the sodium layered oxides and therefore should be avoided. The voltage difference between x ¼ 1 and x ¼ 0.3 for the O3-NaxCoO2 develops from  2.6 V to  4.0 V resulting in a nearly 2 V/ 1Na-extracted slope. Such a drastic voltage change is dictated by the nature of the Naþ ion: large Naþ ions tightly-packed and accommodated by octahedral voids tend to expel each other from the surroundings to diminish the coulombic interactions which are less pronounced in case of smaller Liþ in LiCoO2 that displays only 0.3 V slope upon deintercalation.19 The other important difference between the O and P phases is that in the O-type phases Naþ has to migrate via the tetrahedral site (passing via oxygen triangle) resulting in hindered diffusion. In the P-type structure the Na ion migrates via the trigonal prismatic site (passing via oxygen rectangle) that provides a lower migration barrier than that in the tetrahedral site. Such peculiarity explains much higher diffusion coefficients for the P-type phases resulting in an ability to charge/discharge at the high currents.20 NaxCoO2 was obtained in all key structure polytypes via sintering in oxygen atmosphere at different time-vs-temperature conditions depending on initial reagents type and their ratio.21 This material demonstrates a reliable reversibility during cycling both for the O3-P3 phase transition case and for the P2 phase while keeping the structural integrity in the voltage range of 2.0–4.1 V.22 The studies of the electronic properties of NaxCoO2 showed that all polytypes possess very low electronic resistance after the deintercalation.23 NaxCoO2 can be charged in the x ¼ 0.3–1.0 range accompanied by multiple phase transitions. NaxCoO2 has been studied for the effect of Mn and Fe doping. Yamada et al. did a careful investigation of the P2-Na2/3Co14þ 3þ and Mn4þ/Mn3þ redox couples increase systematically gMngO2 system and found that the average redox potential of the Co /Co with the increase of manganese content. However, the cycling stability turned out to be lower due to the accelerated surface amorphization.24 The oxidation states of Co3þ and Mn4þ in these materials were confirmed by the ESR analysis.25 Additionally, it was demonstrated by XAS studies that at low potentials Co3þ can be partially reduced to Co2þ.26 The iron-containing O3-NaCo1/ 1 at 10C).27 2Fe1/2O2 showed decent electrochemical performance at the high rates (135 mAh g NaxMnO2 was synthesized in a O3-NaMnO2 polytype with a monoclinic distortion due to Jahn-Teller effect of Mn3þ and in a P2-Na0.7MnO2 form.28 The distorted O3 polytype was shown to have different shapes of charge and discharge curves.29 The nature of such a phenomenon is in the formation of a number of ordered phases,30 which stem from the presence of both distorted Mn3þ and non-distorted Mn4þ ions causing Naþ ordering to minimize the destabilizing coulombic repulsions.31 P2-NaxMnO2 phases tend to form Mn vacancies during synthesis, but Komaba et al. managed to synthesize a stoichiometric material as for the Mn occupancy, while keeping the P2 structural polytype.32 They demonstrated that the material with a stronger Jahn-Teller distortion has a larger reversible capacity and better capacity retention, however, the reasons for that phenomenon are unclear.

50

Electrode materials for reversible sodium ions de/intercalation 5.0 5th 2nd

4.0 3.0 2.0

O3-NaFeO2

1.0 0.0 5.0 5th

4.0

2nd

3.0 2.0 1.0

O’3-NaMnO2

Voltage / V

0.0 5.0 4.0 3.0 2.0 1.0

O3-Na[Fe1/2Mn1/2]O2

5th 3rd

0.0 5.0 4.0 3.0 2.0 1.0

5th 2nd

P2-Na2/3[Fe1/2Mn1/2]O2

0.0 5.0 4.0 3.0 2.0 1.0 0.0

5th 3rd 0

P2-Na2/3MnO2 50

100

150 –1

200

Capacity / mAh g

Fig. 3 Comparison of galvanostatic charge/discharge curves of layered NaxMO2 (M is denoted on the curves respectively) oxides, adopting various polytypes (left). Morphology of particles for each sample is also compared (right). Reprinted with the permission from Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Chem. Rev. 2014, 114, 11636–11682. Copyright (2014) American Chemical Society.

The simultaneous Fe and Mn introduction results in the layered oxide with large specific capacity. The electrochemical properties of P2-Na0.67Mnx-1FexO2 can be tuned via changing the Fe/Mn ratio. A larger Mn content delivers a greater capacity, but lowers capacity retention.33,34 It was shown that 80% of the Mn content gives the best trade of specific energy/stability balance. The largest capacity was obtained for the P2-NaxFe0.5Mn0.5O2 phase – 200 mAh g 1 for the first cycle. It was shown that voltages over 4 V are detrimental to the capacity due to layers gliding when a low amount of Na remains in the structure. The further study of this phase demonstrated that replacement of a half of Fe with Ti (P2-Na0.67Mn0.8Fe0.1Ti0.1O2) results in the 130 (C/10) and 80 (1C) mAh g 1, with capacity retention over 95% for more than 50 cycles, making it a competitive Co- and Ni-free material.34,35 The interesting results were achieved for Li-doping onto the Mn site with formation of the P2-Na(Li/Mn)O2 phase in order to investigate the possible oxygen oxidation in the sample.36 The same was later done and confirmed by DFT calculations for a P3 polytype as well.37,38 210 mAh g 1 capacity was achieved due to oxygen redox corresponding to 0.72 Na extracted in the first cycle for P2-Na0.72[Li0.24Mn0.76]O2 where only Mn4þ is present. Such a structure was proven to be stable during cycling.38,39 NaNiO2 is very difficult to be synthesized Na-deficient due to the instability of Ni4þ ions. Similar to Mn-based layered oxides, NaNiO2 is an O3 polytype with a monoclinic distortion due to Ni3þ being a Jahn-Teller ion. The material exhibits several phase transitions on charge/discharge at the specific Na content of x ¼ 1/3, 2/5, 1/2, and 2/3.40,41 The last three points are related to the superstructure formation; however, their structures are not clear yet.42 The suppression of the Jahn-Teller distortion of Ni3þ was reported for the O3-Na(Ni0.6Fe0.4)O2 improving cycling stability.43 The partial substitution of Ni for Co (40%) was found to be inhibiting the intermediate phase formation: the overall deintercalation process can be described as O3 4 O3dist 4 P3dist

Electrode materials for reversible sodium ions de/intercalation

51

(“dist” here and further designates monoclinic or triclinic distortion). The detailed investigation demonstrated that Ni oxidizes prior to Co. The P3dist phase has a Co3þ/Co4þ redox pair and exhibits high electronic conductivity resulting in a very low polarization of the cell.44 The substitution of Ni for Mn eliminates the presence of Jahn-Teller ions Ni3þ and Mn3þ via their recombination into Ni2þand Mn4þ. In P2-Na2/3(Ni1/3Mn2/3)O2, the difference in the ion sizes (Ni2þ is 0.69 Å and Mn4þ is 0.53 Å) promoted the formation of the honeycomb ordering of Ni and Mn.45,46 P2-Na2/3(Ni1/3Mn2/3)O2 and O3-Na(Ni1/2Mn1/2)O2 were considered as potential practical electrodes with improved safety since the Ni4þ appears only at very high potentials. However, the P2-type phase undergoes a phase transition into an O2 polytype if charged above 4.2 V which is typically accompanied with oxygen evolution compromising its’ safety advantages.47,48 To overcome these problems the cutoff potential must be limited to 4.1 V at the cost of capacity lowering. The additional downside of the P2-type phases is a reduced capacity due to the inability to reach a fully sodiated state. Unlike the P2-phase, the O3-phase does not have such a problem and therefore may deliver a larger capacity. The O3-NaNi0.5Mn0.5O2 delivers 105 mAh g 1 at C/10 in the potential range 2.2–3.8 V.49 However, similarly to a P2 phase, the structure collapses above 4.2 V due to layers gliding.50 NaTiO2 is also shown to be electrochemically active in the Na cell. The material demonstrates evidence for Ti migration to the Na site when more than 0.3 Na are deintercalated from the structure.51 The partial replacement of Ti with Li allowed the stabilization of the P2-Na0.66(Li0.22Ti0.78)O2 material. The Ti4þ/Ti3þ redox pair is characterized by low potential allowing to cycle this material between 1.0 and 0.5 V. Half of Na can be reversibly extracted without changes in unit cell parameters therefore rendering this a zero-strain material.52 One of the most interesting examples of the Ti-based systems is the P2-Na0.66(Cr0.6Ti0.4)O2 phase. The presence of two rather different in potential redox couples (Cr3þ/Cr4þ and Ti4þ/Ti3þ) allows using this material as both anode and cathode offering a possibility for assembling symmetric cells showing high-rate performance.53 Various successful doping strategies of sodium layered oxides with Li, Mg, Cu, Zn, Fe, Ti etc. have been reported. Some of these dopants allow increasing the Na content from common x ¼ 0.67 to x ¼ 0.85 therefore increasing the capacity of P2- and O3-type materials, improving cycling performance. However, in all cases the charge above 4.2 V leads to the capacity degradation resulted from the phase transitions with layers gliding and the consequent cracks on particles.38,54–56 The electrochemically inactive elements (Ca and Mg) in this potential range were used with promising results as pillars in the Na layers to prevent the transition metal layer gliding and allow to achieve higher potentials.57,58 Various coating strategies (carbon, TiO2, Al2O3, polymers etc.) are applied to provide better cycling stability and suppress parasitic reactions of materials with electrolyte.59 The Al2O3 coating of P2-Na0.67Ni0.33Mn0.67O2 particles proved to be an efficient way to achieve flexible cathode electrolyte interphase protecting it from exfoliation during battery cycling and simultaneously improving coulombic efficiency of the material in the potential window of 2.3–4.5 V.60 The wide variety of compositions leaves the possibility to find the most appropriate electrode material to build high-performance battery cell.

7.03.2.2 7.03.2.2.1

Oxoanion-based compounds NaMPO4: Maricite and triphylite

Since the electrochemical behavior of iron-based layered oxides is quite modest and they demonstrate tendency to severely degrade during cycling, the investigation of compounds realizing other chemistries is reasonable, for example, oxoanionic compounds such as phosphates. Presented as a positive electrode material in 1997, triphyllite LiFePO4 is one of the most widely investigated compound in the battery field.61–63 It is worth mentioning that triphylite belongs to the olivine family, which naturally occurs as a mineral (Mg2SiO4). The olivine crystal structure type (Fig. 4A and B) is adopted by some Na-based phosphates, e.g. natrophilite (NaMnPO4).64 However, karenwebberite (Na(Fe2þ,Mn2þ)PO4 is unstable thermodynamically and cannot be synthesized using conventional solid-state approaches requiring high-temperature annealing.65 The diffusion of alkali metal cation along the tapes defines the electrochemical behavior of the corresponding phosphate. The crystal structure of the triphyllite is based on the hexagonal close packing (hcp) of oxygen atoms, ions of one- and double-charged metal occupy 6-coordinated positions, while P is located in tetrahedral one. If Naþ cation is introduced inside the above structure, distortion of the hcp lattice occurs (Fig. 4C). The NaFePO4 compound crystallizes in the maricite structure type (space group Pnma, also belongs to the olivine group),66 while the direct synthesis of the triphylite-structured NaFePO4 is hindered due to thermodynamic reasons. In the maricite crystal structure, the FeO6 octahedra are edge-sharing along the a axis, with the only one crystallographic position of Na atoms. Synthetic maricites (NaMPO4, M – transition metal or mixture of thereof) have been considered mostly electrochemically inactive in sodium-based cells.66,67 The origin of inactivity is that sodium atoms are isolated and accessible channels for alkali metal cation transport are absent, resulting in high energy migration barriers. Nevertheless, in 2015 Kim and co-workers stated, that maricite-structured NaFePO4 is able to demonstrate electrochemical activity.68 They claimed that during deinsertion of Naþ the following deep transformations are associated with the reverse distribution of alkali and transition metal atoms. As a result, the channels along the b axis are locked preventing Naþ ions migration. The alternative pathway which is parallel to the [011] crystal direction is characterized with small open windows (but activation energy still exceeds 1.5 eV).68 A composite consisting of nano-sized NaFePO4 (the size of primary particles is  50 nm) synthesized by a solid-state method demonstrated reversible capacity close to 140 mAh g 1 during cycling. No separated peaks can be distinguished on the galvanostatic curve and average potential is quite low - slightly higher than 2 V vs. Naþ/Na. However, the observed reversible capacity corresponds to amorphous FePO4, forming upon the first sodium deinsertion.68

52

Electrode materials for reversible sodium ions de/intercalation

Fig. 4 (A) Comparison of octahedra motifs along [010] unit cell direction of triphyllite (crystal data of LiFePO4, ICSD code number 56291, light green octahedra refer to LiO6) and (B) maricite, NaFePO4 (ICSD code number 56292); (C) typical XRD pattern of triphylite-LiFePO4, chemically delithiated FePO4 and triphylite-NaFePO4; (D) Galvanostatic curves of tr-NaFePO4 in a Na half-cell for the 1st and 200th cycles (C/10 rate). Adapted with permission from Berlanga, C.; Monterrubio, I.; Armand, M.; Rojo, T.; Galceran, M.; Casas-Cabanas, M. ACS Sustain. Chem. Eng. 2020, 8, 725– 730. Copyright (2019) American Chemical Society.

The synthesis of triphylite-type NaFePO4 (the same is relevant also for NaMnPO4) requires soft chemistry approaches, such as the ion-exchange method. During deinsertion of Liþ, LiFePO4 undergoes a transformation to a heterosite-type FePO4,69 whose framework can be filled by another species, i.e., Naþ. The substitution of Liþ for Naþ can be realized through a solid-state approach, but undesirable particle growth can occur in this case. In order to avoid this phenomenon, strong oxidizers, such as NO2BF4, can be used in a form of suspension in a solvent, for example, acetonitrile (CH3CN), to extract Liþ ions from the host with subsequent treatment with a sodium salt, such as NaI, to form tr-NaFePO4.70,71 Electrochemical ion-exchange was also performed during last decades.72 At the same time, both nitronium-containing compounds and acetonitrile are very toxic and water-sensitive substances, sodium iodide reacts with air and H2O with the subsequent oxidation-reduction processes. Such an approach is usually acceptable for research purposes only. In light of these statements, new approaches should be proposed, such as the one described by Berlanga and co-workers. It is based on sequential treating of LiFePO4 in Na2S2O8 and Na2S2O3 in water that allows obtaining trNaFePO4, demonstrating capacity of 132 mAh g 1 at C/10 rate (theoretical one is 154 mAh g 1) with good capacity retention.73 At the potentials of  3.1 and 2.95 V vs. Naþ/Na two plateaus at the charge curve can be distinguished and only one on the discharge curve near 2.80 V (Fig. 4D). Finally, dittmarite-structured compounds should be highlighted as the structurally-related templates (general formula NH4MPO4 $H2O, M ¼ Fe, Mn, Co, Ni, Cd, Mg) to triphylite. Their use as a precursor is a viable strategy to stabilize Mn-based triphylite-structured phosphates. Various features and peculiarities of the dittmarite-structured precursors synthesis (coprecipitation, solvothermal) provides vast opportunities for optimizing particles morphology.74–76

7.03.2.2.2

NASICON-structured electrode materials

An intensive study of the NASICON materials started when a significant mobility of alkali ions in its framework was discovered.77–79 The NASICON (Na Super Ionic CONductor) is a 3D framework structure of polyanionic compounds with a general formula AxM2(XO4)3, where A is an alkali metal, M is a polyvalent metal, and XO4 is an oxoanion.80,81 The main focus of the study was on the development of the solid electrolytes for Na-S batteries and only 10 years later the NASICON materials were proposed as an insertion type electrode material for battery applications.82 The typical NASICON structure can be considered as a 3D array of the so-called “lantern” units consisting of two transition metal (in case of electrode application) octahedra (MO6) interconnected via three oxoanion tetrahedra (XO4) (Fig. 5A). These units build a rigid 3D framework with a trigonal symmetry containing a system of interconnected migration tunnels for alkali ions. Depending on the amount of alkali ions in the crystal lattice and the nature of the forming ions the unit cell may distort from rhombohedral down to monoclinic or triclinic symmetry. There are two sites for alkali ions in the structure: the first is a six-coordinated site between the octahedra MO6 in the vertical column of “lanterns,” the second is an interstitial site between three trigonally arranged “lantern” units. When an alkali ion migrates through the lattice it jumps from one site to another via a hopping mechanism. The vacancies play a key role in this process. It was shown that the change of the alkali metal content significantly affects the

Electrode materials for reversible sodium ions de/intercalation

53

Fig. 5 (A) Crystal structure of the NASICON-structure (ISCD code number 248140), where VO6 octahedra (dark blue-violet), PO4 tetrahedra (light violet), and Na atoms (yellow) are depicted and the so-called “lantern” unit; (B) Charge/discharge curves for Na3þxMnxV2-x(PO4)3 (0  x  1) cycled in the 2.5–3.8 V and 2.5–4.1 V voltage windows at C/10 current rate (1st, 3rd and 4th cycles are shown). A voltage step on discharge of Na3.4Mn0.4V1.6(PO4)3 is most probably due to the side reaction with metallic Na electrode. Reprinted with permission from Zakharkin, M. V.; Drozhzhin, O.A.; Ryazantsev, S. V.; Chernyshov, D.; Kirsanova, M.A.; Mikheev, I. V.; Pazhetnov, E.M.; Antipov, E. V.; Stevenson, K.J. J. Power Sources. 2020, 470, 228–231. Copyright (2020) Elsevier.

conductivity.83 The vast majority of the NASICONs for battery electrode applications are phosphate-based, although there are reports on other anion-based compounds, for example, sulfates.84,85 The V-based NASICON – Na3V2(PO4)3 – was first reported by Porter and co-workers.83,86 It has attracted a lot of attention due to a flat plateau at a high working potential of 3.4 V vs. Naþ/Na and a theoretical capacity of 117 mAh g 1 (Fig. 5B).87 The material exhibits a reliable thermal stability up to 450  C in a charged state. Interestingly it undergoes three minor reversible symmetry changes between  30  C to 225  C weakly affecting electrochemical performance. The discharge to low potentials (< 1.5 V) brings the possibility to insert the fourth Na to the structure engaging the V2þ/V3þ redox couples and resulting in the Na4V2(PO4)3 composition, and also the fifth Na to form Na5V2(PO4)3. Such a behavior allows this material to be used on both cathode and anode sides of the battery. The anode can be reversibly cycled between 3.00 and 0.01 V with two plateaus at 1.57 and 0.28 V resulting in the capacity of 170 mAh g 1.88,89 Due to the ambivalence of electrochemical properties Na3V2(PO4)3 can be used in the symmetrical Na-ion cells. Zhang et al. reported a study of a symmetric cell with an output voltage plateau of 1.8 V, capacity of 90 mAh g 1, and stable cycling for over 280 cycles with a capacity retention of 81%.90 The significant efforts were made in modifications of this material via various doping strategies, particle structure/size control, and carbon coating. The introduction of Mg into the V site and K into the Na position has led to an improved Na diffusion in the channels due to an expansion of the unit cell delivering a more preferable rate and cycle behavior compared to the pristine phase. Structure-wise it was shown that Na can be electrochemically extracted from the interstitial site, while the six-coordinated site remains unchanged suggesting a predominant ion migration via the interstitial sites.91,92 However, it was shown that an introduction of Fe into the V site allows engaging the V4þ/V5þ redox couple and consequently accessing the six-coordinated Na site for electrochemical processes.93 As any other polyanionic compound, Na3V2(PO4)3 has low electronic conductivity and, therefore, requires extra treatment for an improvement of the electrochemical behavior. The formation of a carbon layer on the surface of the particles was found very beneficial to the electrochemical properties. The attempt of such a modification resulted in the 93 mAh g 1 capacity with an improved coulombic efficiency and cyclability.94 There are many reports on various nano-structural modifications of carbon coating. For example, carbon coating during actual synthesis was found to be very successful due to formation of very small particles (40 nm) and demonstrated 95 mAh g 1 at 5C and a capacity retention of 96.1% after 700 charge/discharge cycles.95 Others suggest various methods of reduction of organic compounds at low and high temperatures to achieve required electrochemical behavior.96 The enhanced electrochemical performance was reported for the nitrogen and boron doped carbon coating.97 It was shown that the use of graphene as a coating agent supports stable rate capability and cycling resulting in 70 mAh g 1 at 30C for over 300 cycles.98 The mixed V-Fe, Mn systems were also studied in order to overcome problems of pristine V-based compounds as well as to  while an introduction decrease the cost of pricy precursors.99 The Mn-containing analog crystallizes in the trigonal space group R3c,

54

Electrode materials for reversible sodium ions de/intercalation

of Fe causes a monoclinic distortion. The Na3FeV(PO4)3 structure demonstrates two plateaus at 3.3 and 2.5 V on discharge related to both V4þ/V3þ and Fe3þ/Fe2þ redox couples. The material showed 95% of capacity retention after 1000 cycles at 1C. On the other hand, Na3MnV(PO4)3 provides a higher potential (3.6 V) due to the Mn3þ/Mn2þ redox couple. As a downside, Mn in this material tends to move to the Na site in the crystal structure hindering electrochemical cycling.100 The increase of the cutoff voltage up to 4.3 V allows to access the V5þ/V4þ redox couple and to increase the capacity engaging extra Na to the process.101 However, it was shown that the activation of this high-voltage redox activity leads to a phase transition which is not reversible on discharge thus influencing the stability of the material.102 The introduction of Zr in place of V allows obtaining a stable high-voltage material solely depending on the Mn redox activity.103  space group with three independent Na sites. Such a configuration allows to The Na3MnZr(PO4)3 crystallizes in the trigonal R3c suppress Jahn-Teller effect of Mn3þ and, therefore, the material has a better stability than other Mn containing analogs and has only a 5.5% volume expansion resulting in a more pronounced cycling stability. So far, Na3MnZr(PO4)3 demonstrated 105 mAh g 1 at two voltage plateaus around 3.5 and 4 V. Another example of NASICON-structured materials for batteries is a titanium-based NaTi2(PO4)3. It adopts a rhombohedral  space group. NaTi2(PO4)3 has high ionic conductivity in the order of  10 6 S cm 1 providing an crystal type with the R3c þ enhanced Na migration and perspective for high-rate applications. This material is capable to reversibly insert two Na equivalents per formula unit into an interstitial alkali site resulting in the theoretical capacity of 133 mAh g 1.104 The material demonstrates a discharge plateau at 2.1 V vs. Na/Naþ corresponding to a two-phase electrochemical transition. A key advantage as an anode is safety due to the absence of the metallic Na deposition in the cell. The relatively high anode potential makes it an attractive candidate to be used in water-based systems such as NaTi2(PO4)3/Zn where it is capable of delivering a relatively high reversible capacity of 124 mAh g 1 at  0.8 V vs. Ag/AgCl reference electrode.105 Further modifications of the material with conductive coatings allowed to achieve improved electrochemical performance in terms of cycling stability and rate capabilities.106 NaTi2(PO4)3 has low electronic conductivity and therefore requires carbon coating of the particles in order to realize high performance. Covering of the particles with graphene and carbon nanotubes allowed this material to be cycled at a 50C rate for over 1000 cycles.107,108 Yang et al. reached tangible results with carbon-coated porous NaTi2(PO4)3 nanocubes cycled for 10,000 cycles at a 100C rate.109 The application of a thin carbon shell and interconnected carbon network for the NaTi2(PO4)3 modification delivered 108 mAh g 1 at 100C and a long life-time at 50C for over 6000 cycles.110 Various attempts to use graphene and graphene oxide as a coating material were found to be successful as well. The detailed studies of the Na storage mechanism in NaTi2(PO4)3 revealed that there appears an additional working plateau near 0.4 V suggesting activity of Ti3þ/Ti2þ redox couple. Potentially this material can deliver a capacity of 210 mAh g 1.111 NaTi2(PO4)3 was successfully doped with Fe resulting in Na1.5Fe0.5Ti1.5(PO4)3 and Na2FeTi(PO4)3. Na1.5Fe0.5Ti1.5(PO4)3 displays a much more complex de/intercalation mechanism and revealed an additional voltage plateau at 2.2 V related to the Fe3þ/Fe2þ redox couple.112,113 In contrast, Na2FeTi(PO4)3 demonstrates only one flat plateau on the discharge corresponding to a two-phase transition into Na2.5FeTi(PO4)3. The stability and sustainability of the NASICON family members can become a key advantage for battery application. Though there is still room for improvement, this family promises a bright future for the SIB industry.

7.03.2.2.3

Alluaudites

The alluaudite-type structured compounds only relatively recently (2010) were introduced as a class of insertion materials for SIBs.114 This structure type with the general formula A2A1M1M22(XO4)3 (A – alkali metal cation, M – metal cation or mixture thereof, and X – p-element, e.g., P, As, etc.) usually crystallizes in a monoclinic crystal system (space group C2/c).115,116 The structural framework is built on sequences of doubled octahedral dimers (M22O10) alternating with M1O6 octahedra to form twisted  (Fig. 6A). These chains are interconnected via vertices of oxoanion tetrahedra forming layers perpendicular chains parallel to ½101 to the b-axis. Additionally, these chains are bridged via oxoanion tetrahedra to form a 3D framework. Such a framework has two spacious tunnels along the c-axis to be occupied by alkali cations. The A1 cation takes a site in a wider tunnel 1, while the narrower tunnel 2 site is taken by the A2 cation. Transition metal sites tend to be always fully occupied; alkali ion sites may be occupied partially or left completely empty while keeping the structure stable. This fact allows using this framework as a positive electrode material for SIBs.117 Chemistry-wise, this family is represented mostly by Fe, Mn, Co and V phosphates and sulfates though the structural flexibility enables many other compositions. The Na ion has the most suitable ionic size to form an alluaudite structure and to be accommodated in large diffusion channels within the corresponding framework. This is the reason why most of the alluaudite minerals contain Na. So far, a large variety of alluaudite structured type cathode materials for Na-ion batteries is based on the phosphate compounds. These materials are usually prepared by solid-state, solvothermal, combustion, and sol-gel routes with annealing temperatures up to 850  C. The very first report on the use of the alluaudite NaMnFe2(PO4)3 as cathode in SIBs was published by Delmas’ group.114 The initial study showed that the morphology plays a key role in realizing the large electrochemical capacity of the material. The change of Mn:Fe ratio allowed shifting the focus of redox activity from mixed Mn/Fe to solely Fe3þ/Fe2þ resulting in 140 mAh g 1 for Na2Fe3(PO4)3.118 A further study of Fe-based alluaudites via varying the Na content showed that the change of relative occupancy of alkali sites via a solvothermal route results in the range of Na2 xFe3(PO4)3 compounds, where x ¼ 0.01  0.53.119,120 Such offstoichiometric compositions demonstrated capacities over 140 mAh g 1 at an average potential of 2.7 V and a reliable cycling behavior at high C-rates. The substitution of nearly 30% of Fe with V in Na2Fe3(PO4)3 rises the Fe3þ/Fe2þ redox potential to

Electrode materials for reversible sodium ions de/intercalation

55

Fig. 6 (A) Crystal structure of typical alluaudite representative (Na2.56Fe1.72(SO4)3, ICSD code number 243842), FeO6 octahedra are light brown, SO4 tetrahedra are deep yellow, Na atoms are yellow; (B) Galvanostatic cycling of Na2-xFe2(SO4)3 material proposed by Barpanda, Oyama, Nishimura, Chung and Yamada in 2014; (C) Capacity retention upon cycling up to 30 cycles under various rate of C/20 (2 Na in 20 h) to 20C (2 Na in 3 min). (Inset) The discharge curves of Na2-xFe2(SO4)3 as a function of rate (from C/20 to 20C). Before each discharge, the cells were charged at C/10 to 4.2 V. Reprinted from Barpanda, P.; Oyama, G.; Nishimura, S.I.; Chung, S.C.; Yamada, A. Nat. Commun. 2014, 5, 1–8. CC BY 4.0.

3.1 V. Additionally, this material demonstrated a nearly theoretical capacity even at a 5C rate accompanied with a high cycling stability over 2500 cycles.121,122 Among alluaudites, Na2Fe1.96V0.96(PO4)3 is so far considered to be the best in this group. The Mn analog of this material demonstrated both Mn3þ/Mn2þ and V4þ/V3þ to be active with an average potential 3.5 V and the discharge capacity more than 95 mAh g 1. The V-doped sulfates fit the need of high energy density cathode materials for SIBs. Another interesting example is NaCoFe2(PO4)3 showing a 2.9 V Fe3þ/Fe2þ redox potential and delivering 70 mAh g 1 capacity. This material possesses a rather low ion migration activation energy of  0.31 eV. However, this representative provides relatively low energy density mostly due to the low redox potentials.123 Although alluaudites mostly participate in intercalation processes, Na2Co2Fe(PO4)3 interaction with Naþ leads to a conversion process. Interestingly, this material exhibits electrochemical activity in both cathode (3.5 V with Fe3þ/Fe2þ and Co3þ/Co2þ redox pairs) and anode (0.6 V with conversion Fe2þ/Fe0 and Co2þ/Co0) regions.124 The general problem of phosphate-based alluaudites is lower energy density in comparison to other cathode materials. One of the ways to overcome this problem is a transition from phosphates to sulfates. The inductive effect of the SO4 group in the crystal lattice results in a higher redox potential of the active redox couple.125 For the Fe sulfate-based compounds, this rise may result in more than 1 V of the additional potential, therefore increasing the energy density by 25%. At this moment, great attention is drawn to the sulfate-based alluaudites – non-stoichiometric Na2þ2xFe2-x(SO4)3 (x  0.28) with a high average redox potential (3.8 V vs. Na/ Naþ).126 The material delivers 100 mAh g 1 capacity with three plateaus at 3.42, 3.8 and 4.04 V. Considering theoretical capacity of 120 mAh g 1 this compound may deliver more than 540 Wh kg 1 energy density making it competitive to LIB polyanion cathodes. The redox potential, electrochemical activity and discharge capacity in such sulfates can be tuned via controlling the Na and Fe

56

Electrode materials for reversible sodium ions de/intercalation

content at the synthesis stage. Though these sulfates are able to deliver rather high energy density they suffer from slow ionic and electronic conductivities resulting in the significant capacity fading during cycling. Another drawback is high sensitivity to moisture promoting easy oxidation of Fe2þ to Fe3þ. The carbon coating is supposed to mitigate both conductivity issues and sensitivity to the environment. Stoichiometric Na2Fe2(SO4)3 is supposed to have the highest energy density and to be able to undergo complete desodiation. The stabilization of the stoichiometric system is tricky since more thermodynamically stable Na6Fe(SO4)4 tends to be formed. There are reports about successful synthesis via sol-gel methods and co-precipitation.127,128 It was shown that this system may give up to 108 mAh g 1 at 3.8 V vs. Naþ/Na. High redox potential in the sulfate alluaudites originates from unique coordination of the Fe2O10 dimers tightly bridged by SO4 units that results in a short Fe-Fe interaction distance ( 3.2 Å) capable to provide a strong electrostatic repulsion for Fe3þ-Fe3þ and consequently a high redox potential. Also, it was found that complete desodiation of these materials along with retaining of the structure is impossible.125,129 In the Na2þ2xFe2-x(SO4)3 solid solutions at the limit of x ¼ 0.5, another meta-stable structure Na3Fe1.5(SO4)3 (or Na2Fe(SO4)2) appears. This structure can be synthesized directly in the “charged” state as NaFe(SO4)2 resulting in an insertion type eldfellite material which is electrochemically active in Na cells.130 The attractive feature of this material is an easiness of preparation via simple dissolution-evaporation method with subsequent sintering at moderate 300  C. Its structure consists of layers of FeO6 octahedra interconnected via SO4 tetrahedra pointing outside the layers on both sides, while Na ions are located between the layers. When additional Na is incorporated, it goes to the interstitial sites next to already present Na. The insertion goes via solid-solution mechanism. It was shown that the material can be reversibly cycled at an average voltage of 3.2 V resulting in the 80 mAh g 1 of specific capacity. The modifications via partial cationic and anionic replacements in this material were shown to be successful structurally, however, these were found to be detrimental to electrochemical performance.131,132 Although the main focus of research was placed on Fe-based sulfate alluaudites, some efforts were devoted to the Mn- and Cobased compounds. However, lack of suitable high-voltage electrolytes does not allow observing any electrochemical activity. The first principles DFT calculations provided evaluation of redox potentials in these systems as 4.4 and 5.1 V for the Mn- and Cobased ones, respectively.133,134 Finally, this group of compounds is not limited only by phosphates and sulfates. For example, Na2Mn2Fe(VO4)3 was reported, delivering 35 mAh g 1 with a certain potential for further development.135 Additionally, Gao et al. reported Na2.67Mn1.67(MoO4)3 providing 80 mAh g 1 at 3.45 V with a 2D migration pathway system.136 Researches have done a significant amount of work on these materials and much more is left opening the possibility for further development.

7.03.2.2.4

Pyrophosphates

The pyrophosphate (P2O4 7 anion) based compositions for SIBs are considered to be more stable at elevated temperatures (up to 500  C) due to a larger stability of the pyrophosphate anion.137 The first-ever report on the electrochemical activity in a pyrophosphate-type intercalation electrode in a Na-ion cell was presented by Yamaki et al., who demonstrated that (MoO2)2P2O7 can deliver 180 mAh g 1 of reversible capacity though with extremely poor rate capability.138,139 Later Yamada’s group reported a Na2FeP2O7  space group forming a 3D-framework based on the positive insertion material (Fig. 7A).140 This material crystallizes in the P 1 Fe2O11 dimers of corner-sharing FeO6 octahedra, and corner- and edge-sharing P2O7 units. Such a construction contains large ion migration tunnels along the [001] direction with room to store Na. This structure type of transition metal pyrophosphates is referred to as the rose-phase. Na2FeP2O7 is capable of extracting only a single Na ion per formula unit delivering a specific capacity of 90 mAh g 1 (Fig. 7E). The discharge voltage profile demonstrates multiple steps related to the different Na orderings.141 The calculated activation energy for Na migration was reported to be 0.5 eV proving fast ion transport.142,143 The particle surface modifications allowed achieving above 80% of capacity retention after 2000 cycles at 10C.144 However, a relatively low energy density in Na2FeP2O7 questions the possibility of its industrial application. Na2MnP2O7 has a very similar structure to the Fe-based analog, but demonstrates a much more complex charge-discharge behavior.145 0.9 Na per formula unit can be extracted in the 1.5–4.5 V voltage range without significant lattice distortion due to the Jahn-Teller effect of Mn3þ ions, which is avoided thanks to the rigidity of pyrophosphate ions. An advantage of the Mn based material is specific energy higher than that of the Fe-based one due to a higher average redox potential (3.5 vs. 2.7 V). The discharge voltage profile demonstrates a similar multistep process to the Fe-based analog. There is another modification of this material reported, b-Na2MnP2O7, with a different orientation of MnO6 edge-sharing octahedra building Mn2O10 dimers.146 This structure delivers 80 mAh g 1 capacity at the 3.3 V average voltage. Interestingly, the voltage profile does not demonstrate multistep processes, therefore, indicating a solid-solution-like mechanism of discharge. Nazar’s group reported electrochemical properties of a series of Fe-based pyrophosphates Na4-aFe2þa/2(P2O7)2 (2/3  a  7/8) in 2013.147 Such an off-stoichiometric composition allowed accessing a higher capacity of 107 mAh g 1 in comparison to pristine Na2FeP2O7. The material has a similar structure to the parent phase, however the Fe2O11 dimers are organized into higher symmetry crown-like units: Fe2P4O22 and Fe2P4O20.148 The volume difference between charged and discharged phases does not exceed 2.2% resulting in good cycling stability over long operation time.149 The similar improvement was also reported for the off-stoichiometric Mn-based structure.149 The Co-based Na2CoP2O7 structure crystallizes in the orthorhombic lattice having a poor electrochemical behavior.150 However, the synthesis strategy with the control of Na deficiency leads to the triclinic rose-phase with an average discharge potential of 4.3 V and specific capacity close to 80 mAh g 1 resulting in an improved energy density in comparison to the Fe-based analog.151

Electrode materials for reversible sodium ions de/intercalation

57

Fig. 7 (A) Crystal structure of Na2FeP2O7 (ICSD code number 237850) viewed along a-axis; (B) Crystal structure of Na4Fe3(PO4)2P2O7 (ICSD code number 236316) viewed along b-axis; (C) Crystal structure of b-NaVP2O7 viewed along c-axis, FeO6 and VO6 octahedra are light brown and dark blue-violet respectively, PO4 tetrahedra are light violet; (D) Typical CV profile of Na4M3(PO4)2(P2O7) (M ¼ Ni, Co, and Fe) and dQ/dV plot for M ¼ Mn. (E) Schematic view of typical discharge profiles at low current densities (C/10) of Na4Fe3(PO4)2P2O7 versus other polyanionic cathodes. Adapted with permission from Gezovic, A.; Vujkovic, M.J.; Milovic, M.; Grudic, V.; Dominko, R.; Mentus, S. Energy Storage Mater. 2021, 37, 243–273. Copyright (2021) Elsevier.

The other promising example of the pyrophosphate electrode materials is V-based NaVP2O7. Unlike the materials mentioned above, this compound does not exhibit a rose-structure and crystallizes in the monoclinic P21/c space group.152 It was reported to have a poor electrochemical activity due to low ionic conductivity stemming from high activation energies for Naþ migration in the structure. The modification of V3þ into V4þ results in the Na2VOP2O7 composition, which crystallizes in the layered structure formed by corner-shared P2O7 units with VO5 square pyramids and layers of Na ions (space group P4bm).153 This structure delivers 80 mAh g 1 at a high average potential of 3.8 V ascribed to the V5þ/V4þ redox couple. The second polymorph b-NaVP2O7 can be obtained via a hydrothermal route with subsequent dehydration of the intermediate NaV(HPO4)2 phase.154 It crystallizes in the KAlP2O7-type structure with the monoclinic space group P21/c. The 3D-framework of this material contains several tunnels for efficient Na migration (Fig. 7C).155 The material demonstrated 104 mAh g 1 at C/10 and average potential 3.9 V and a remarkably small volume change of  0.5% upon cycling. The cycling at 20C exhibited 90 mAh g 1. The material is active over a wide potential window of 1.0–4.4 V opening a possibility to be used in the symmetrical cells. V-based pyrophosphate Na7V3(P2O7)4 crystallizes in the monoclinic space group C2/m.154 The material has a complex quasilayered structure built upon corner-sharing octahedra VO6 with P2O7 units forming large diffusion tunnels for Na-ions. The capacity of 80 mAh g 1 was obtained at an average potential of 4.13 V related to the V4þ/V3þ redox couple, which is the highest redox potential among the V-based polyanionic compounds for sodium-ion batteries. The nature of such a phenomenon lies in the strong inductive effect of pyrophosphate units around the V ions. The rate capability of this material can be further improved via nanostructuring.156

7.03.2.2.5

Mixed phosphates

The mixed phosphates have gained a significant attention due to low barriers of Na transport within the structure, structural stability, and cheap and non-toxis components accompanied with relatively high redox potentials.157 The Na4M3(PO4)2(P2O7)

58

Electrode materials for reversible sodium ions de/intercalation

(M ¼ Fe, Mn, Co etc.) family crystallizes in the orthorhombic space group Pn21a with a slight variation of cell parameters.158,159 The structure consists of alternating double layers built upon M3P2O13 units (Fig. 7B). The MO6 octahedra are corner- and edge-sharing with the PO4 tetrahedra within the layers, which are connected via P2O7 pillar groups forming tunnels going along all three crystallographic axes. The tunnels are filled with Na and are capable of providing reversible de/insertion. Such a framework benefits the additional integrity during the electrochemical cycling due to rather small changes of the cell parameters. The Fe-based framework undergoes only a 4% volume change,160 Mn-based  7%,161 Co-based  10%,162 and Ni-based  2%.163 The transition metal occupies three octahedral sites, while Na atoms occupy four sites connected by channels. The activation barrier is low for all diffusion pathways.164,165 Sodium has a different coordination in this structure which additionally depends on the stage of deintercalation. Such variations affect the electrode potential of the material during discharge therefore resulting in multiple plateaus. The discharge potential of Na4M3(PO4)2(P2O7) varies significantly depending on the nature of the transition metal. The average redox potential of the Fe based material is around 3 V. The Mn-based compound has  3.8 V, Co based demonstrates  4.5 V, while the Ni based exhibits an average potential of  4.8 V. The evaluation of the theoretical capacity depends on the consideration of 3 or 4 Na ions to participate in the de/insertion process (Fig. 7D and E). For example, Na4Fe3(PO4)2(P2O7) should therefore result in 129 mAh g 1 (three-electron process) or 170 mAh g 1 (four-electron process). Kim et al. suggested that the only three Na de/insertion should be considered due to the fourth Na being a pillar of the framework similarly to the case of the layered Na oxides.164 The electrochemical behavior of this family strongly depends on the synthesis route. The best results for Na4Fe3(PO4)2(P2O7) are achieved when nanosized particles are tested providing up to 97 mAh g 1 at 10C166 or 79 mAh g 1 at 100C167 even after 4000–6000 cycles. However, such outstanding results are achieved via complicated preparation techniques involving formation of the desired phase directly in the mass of the current collector (carbon cloth) and, therefore, this material still needs further development. Additionally, these excellent rate capabilities are not yet accessible for the Co-, Mn- and Ni-based compounds because the capacity degrades quickly for high discharge rates. Interestingly, small substitution of Co for Al in Na4Co3(PO4)2(P2O7) is beneficial for higher rate capabilities.168 There are no reports on high rate cycling of the Ni-based compounds so far due to very high average redox potential leading to severe electrolyte degradation. The development of more stable electrolytes would help to open the full potential of this system. Unlike other examples of electrode materials for NIBs, structure modifications of Na4M3(PO4)2(P2O7) via doping with other transition metals showed relatively small effect on the electrochemical properties. The key to utilize high-rate performance in this material mostly lies in the nanostructural modifications. The formation of a conductive carbon network and use of nanoparticles of the active material less than 200 nm were found to be beneficial.161,169,170 There is a variety of carbon additives reported including nanosheets, graphene, and graphene oxide.166,167,171

7.03.2.3 7.03.2.3.1

Mixed-anion positive electrode materials Na3V2(PO4)2(O,F)3

The introduction of fluorine into sodium-vanadium phosphates produces a wide range of structural types. One of them is Na3V2(PO4)2O2-xF1þx (0  x  2) firstly reported by Park et al.,172 which in turn belongs to the structural type natisite described by Le Meins in 1999.173,174 The amount of fluorine in the structure affects the oxidation state of vanadium, such that Na3V2(PO4)2F3 contains only V3þ and Na3V2(PO4)2O2F has only V4þ ions. The V coordination in the VO4þxF2-x octahedra is also hugely affected by the vanadium electronic configuration. Moreover, the V3þ octahedra tend to be axially elongated due to a weaker F-F repulsion. Considering that the octahedra are axially directed along the c axis, this cell parameter increases with the rise of F content (Fig. 8A). All materials based on the Na3V2(PO4)2O2-xF1þx family demonstrate a close average redox potential near 3.8 V with a specific capacity around 125 mAh g 1 (Fig. 8B). The volume change during cycling was found to be rather low.172 Members of this class can be synthesized via different routes, but the key one is a rather simple solid-state reaction of NaF and proper V- and PO4-containing initial reagents. The end member Na3V2(PO4)2F3 is more attractive due to its highest operating potential ( 3.9 V) and the possibility to extract two Na ions per formula unit using the V3þ/V4þ redox couple (Fig. 8B). The detailed crystallographic studies based on the synchrotron experiments revealed that this material crystallizes in the orthorhombic space group Amam.175,176 The origin of such complexity is in the sophisticated distribution of Na ions in the ion conduction tunnels which gives rise to a series of phase transformations during the charge-discharge process. The redistribution of the Na ions upon extraction leads to four intermediate phases until the NaV2(PO4)2F3 composition is reached. While being rather complex structurally, the electrode potential profile of NaxV2(PO4)2F3 has only two relatively distinctive plateaus corresponding to de/insertion of two Na equivalents with a  0.4 V potential step.176 Tarascon’s group has attempted to access the third Na in the electrochemical reaction via formation of a disordered NaxV2(PO4)2F3 (x  0) phase with the tetragonal space group I4/mmm at charging potentials up to 4.8 V vs. Naþ/Na.177 With the support of an NMR study, it was confirmed that the third Na can be reversibly accessed during electrochemical cycling though its subsequent intercalation requires much lower electrode potentials of around 1.5 V resulting in a severe voltage hysteresis. Despite almost doubling the attainable specific capacity, the reversible uptake and release of three sodium ions on the following cycles provides only a 20% gain in specific energy, though sacrificing the cycling stability. The Na3V2(PO4)2O2F representative demonstrated two plateaus at the discharge voltage profile (4.0 and 3.6 V vs. Naþ/Na) delivering 128 mAh g 1 capacity.178 It was shown that a phase transition occurs with a significant symmetry change at 4 V. However, the volume change upon cycling does not exceed 2.7% resulting in a stable reversible electrochemical process. The step between potentials corresponds to a complete extraction of the first Na and the beginning of the second Na extraction.

Electrode materials for reversible sodium ions de/intercalation

59

Fig. 8 (A) Crystal structure of Na3V2(PO4)2F3 (space group Amam) space group, showing V2O8F3 bioctahedra, PO4 tetrahedra and the three different sites of sodium atoms (top). Top view displaying the a  b plane in the case of room temperature unit cell (solid line indicates the associated space group Amam, dotted line indicates the unit cell observed at high temperature (400 K), in the space group I4/mmm). Orange spheres represent fully occupied Na sites and white/orange spheres partially occupied ones. (B) Galvanostatic cycling of a Na3V2(PO4)2F3 in sodium half-cells at C/50 rate per exchanged ion (electrolyte 1 M NaPF6 (99%) in PC). (C) Crystal structure of Na2MnPO4F and its relationship with Na2FePO4F (ICSD code number 194076); (D) Typical view of galvanostatic cycling curves for Na2FePO4F/C in sodium half-cells. (E) Typical view of galvanostatic cycling curves for Na2MnPO4F in sodium half-cells. (A) Reprinted with permission from Bianchini, M.; Fauth, F.; Brisset, N.; Weill, F.; Suard, E.; Masquelier, C.; Croguennec, L. Chem. Mater. 2015, 27, 3009–3020. Copyright (2015) American Chemical Society. (B) Reprinted with permission from Bianchini, M.; Brisset, N.; Fauth, F.; Weill, F.; Elkaim, E.; Suard, E.; Masquelier, C.; Croguennec, L. Chem. Mater. 2014, 2, 1393–1399. doi:10.1021/cm501644g. Copyright (2014) American Chemical Society. (C, D) Reprinted with permission from Kawabe, Y.; Yabuuchi, N.; Kajiyama, M.; Fukuhara, N.; Inamasu, T.; Okuyama, R.; Nakai, I.; Komaba, S. Electrochem. commun. 2011, 13, 1225–1228. Copyright (2011) Elsevier. (E) Adapted with the permission from Gezovic, A.; Vujkovic, M.J.; Milovic, M.; Grudic, V.; Dominko, R.; Mentus, S. Energy Storage Mater. 2021, 37, 243–273. Copyright (2021) Elsevier.

7.03.2.3.2

Na2MPO4F (M ¼ Mn, Fe, Co)

The first electrochemical active member of the Na2MPO4F family was proposed by Ellis and co-workers in 2007 – this was Na2FePO4F.179–181 Iron-based compounds are attractive in terms of price and availability, the phosphate-based framework provides high thermal and cycling stability, the introduction of fluorine atoms leads to increasing of the Fe3þ/Fe2þ potential. Moreover, this chemical composition was expected to provide opportunity for realization more than one electron transfer, however, this is still a questionable issue.182 The crystal structure of Na2FePO4F (space group Pbcn) can be described as a quasi-layered system, consisting of Fe2O7F2 building blocks (which in turn are two face-sharing FeO4F2 octahedra), bonded by the PO4 units, which form infinite [FePO4F]N layers (Fig. 8).183 A carbon-covered (carbon coating from ascorbic acid pyrolyzed under Ar flow), Na2FePO4F/C

60

Electrode materials for reversible sodium ions de/intercalation

demonstrates reversible capacity up to 120 mAh g 1 with two well-distinguished plateaus at 3.06 and 2.91 V (Fig. 8D).184 For comparison, the uncoated fluoride-phosphate is almost electrochemically inactive, which is a clear evident of its poor electronic conductivity. The So-based counterpart Na2CoPO4F adopts the same crystal structure and is able to de/intercalate reversibly 1 Naþ per formula unit at an operating potential about 4.3 V vs. Naþ/Na.185 The Mn-based member, Na2MnPO4F, was initially stated to adopt a new specific crystal type.186 However, it is obvious that Na2MnPO4F is topologically related to Na2FePO4F (Fig. 8C). Considering four non-equivalent metal positions, the structure of Na2MnPO4F can be derived from Na2FePO4F through a mutual exchange of alkali and transition metal positions, for example, Na1 4 M1. For Na2MnPO4F the capacity close to 100 mAh g 1 was achieved (the theoretical is 125 mAh g 1) with an average potential about 3.5 V vs. Naþ/Na (Fig. 8E).187

7.03.2.3.3

AVPO4F (A – alkali metal): Tavorite and KTiOPO4 structure types

While NASICON and natisite structure types are the most common oxoanion electrode materials for the Na–V–P–O–F set of elements, there is a number of reports devoted to the chemical composition labelled as “NaVPO4F.” It was stated that this is a tavorite-typed cathode material (analogous to LiMPO4F (M ¼ V, Fe, Ti)188–191 for SIBs.192–197 However, the detailed analysis of the cycling curve profiles and XRD patterns indicates that this electrochemical activity can be attributed to mixtures of various NASICON-type sodium-vanadium phosphates and natisite-typed fluoride-phosphates. In 2017, Boivin et al. found specific mild hydrothermal conditions to obtain an authentic sodium-based fluoride-phosphate isostructural to LiVPO4F. For this purpose, a mixture of sodium phosphate, sodium fluoride, phosphoric acid and VCl3 were hydrothermally treated for 240  C in a PTFElined vessel for 24 h.198 The synthesis duration is especially important in this case, as attempts to decrease reaction time led to formation of Na3V2(PO4)2F3 admixture. Later, Kosova’s group demonstrated the importance of quenching and found specific conditions for a solid-state synthesis of a tavorite-structured NaVPO4F.199–201 The crystal structure of the NaVPO4F tavorite is shown in Fig. 9A. The fluoride-phosphate is built up by VO4F2 octahedra sharing common fluorine atoms, forming .–V–F–V–F–V–F–. infinite chains along the [101] direction. The octahedra are connected through PO4 tetrahedra forming .–V–O–V–O–V–O–. sequences and leading to a three-dimensional Na migration network. Sodium ions occupy hexagonal tunnels along the [ 110] direction (Fig. 9A).198 It was demonstrated that tavorite NaVPO4F is almost electrochemically inactive both in Li- and Na-based half-cells. The operating potential of the V4þ/V3þ couple cannot be determined due to very small reversible capacity, high polarization and absence of a clear voltage plateau. Boivin et al. also attempted to chemically extract Naþ using NO2BF4, but no changes in the XRD profile were observed despite the strength of the oxidizing agent. BVEL (Bond Valence Energy Landscape) calculations and VDP analysis (Voronoi-Dirichlet partition) of Naþ migration in NaVPO4F also lead to the conclusion that this compound is electrochemically inactive.200–202 Additionally, it is worth mentioning that these geometrical-topological approaches are a strong instrument complimentary to quantum chemical methods, such as

Fig. 9 (A) Ball-polyhedral representation of tavorite crystal structure. (B) Ball-polyhedral representation of KTP (AMXTO4) crystal structure with split A sites, showing continuous channels along the [100] and [010] directions. (C) Schematic view of voltammograms for Li, Na and K cations in the common reference Fcþ/Fc (ferrocenium/ferrocene) potential scale. Scan rate is 50 mV$s 1; d) Schematic view of voltammograms curves of the K0.36VOPO4 cathode in sodium half-cells; e) Galvanostatic curves for Na0.9K0.1TiOPO4 in the 1.0–2.1 V (vs. Naþ/Na). Adapted with permission from Fedotov, S.S.; Samarin, A.S.; Antipov, E.V. J. Power Sources 2020, 480, 228840. Copyright (2020) Elsevier.

Electrode materials for reversible sodium ions de/intercalation

61

DFT.203 A fluoride-sulfate counterpart, NaFeSO4F, proposed by Tripathi and co-workers can be considered as a 1D Na-ions conductor, but DFT calculations clearly reveal quite high ( 0.9 eV) activation energy barriers.204,205 Being similar to the members of the tavorite crystal family in terms of chemical composition, potassium-based compounds, such as KVPO4F, adopt KTiOPO4 (KTP) crystal type.79 During the last decades KTP-structured compounds were mainly investigated as materials for non-linear optics, namely, for the second harmonic generation.206–208 Interestingly, its mineral analog – KTiOAsO4, katiarsite – was discovered only in 2016 in Kamchatka peninsula.209,210 Members of the KTP family are described by the general formula AMXTO4, where A is a single-charged metal cation/NH4þ or mixture of those, M is a double, or higher-charged transition metal or non-metal cation or a mixture of those, X is an oxygen or fluorine atom, or OH-group, T is a p-element (P, As, S etc.).211,212 The crystal structure of KTP represents a system of helical chains of the vertex-sharing MO4X2 octahedra running along the b direction and corner-linked by the TO4 tetrahedra into a 3D-framework containing Aþ ions in cavities (Fig. 9B). If X is fluorine, a sequence of alternating XO4F2 octahedra appears – so-called cis- and trans- XO4F2. The majority of representatives belongs to the orthorhombic crystal system adopting a number of centrosymmetric and non-centrosymmetric space groups.211 The compounds of this structural type revealed a high ionic conductivity and therefore high mobility of alkali metal cations inside the host structure with values up to 10 4 S cm 1 at 298 K213–215 being similar to NASICONs (10 3  10 6 S cm 1).84,216 Recham et al. firstly investigated electrochemically a compound of this structure – it was a sulfate-based KFeSO4F,217 which demonstrates an electrochemical activity in all Li-, Na- and K- cells. It delivers a reversible capacity of  109 mAh g 1 at an average potential of 3.4 V vs. Naþ/Na in sodium half-cells in the potential range of 3.0–4.5 V. The synthesis of this fluoride-sulfate requires dry KF and absence of water during all procedures with the aim to avoid possible hydrolysis taking into account the hygroscopicity of both initial and resulting compounds. Shifting from sulfates to phosphates significantly improves stability of materials against moisture. Several years ago AVPO4X (A ¼ Li, Na, K, Rb; X ¼ F, O) were proposed as materials for alkali metal de/intercalation (Fig. 9).218,219 For example, the KVPO4F fluoride-phosphate,220 which can be synthesized by both advanced solid-state and solvothermal methods, delivers about at initial cycles 100 mAh g 1 with an average potential of 3.98 V in the 2.0–4.7 V vs. Naþ/Na potential range (Fig. 9C).221–223 Its oxygenated counterpart KVOPO4 also demonstrates reversible Naþ de/intercalation both as a cathode and anode (Fig. 9D).224 Attempts to remove full equivalent of NH4þ from NH4VOPO4 in aim to stabilize pristine KTP-typed “VOPO4” were made by Whittingham’s group. It was found that the thermal treatment leads to structure collapse and formation of (VO)2P2O7. However, “chimie douce” methods (refluxing NH4VOPO4 with 3 M NaBr in a tetraethylene glycol solution at elevated temperature) allows substituting some amount of NH4þ.225 Further investigation of soft chemistry approaches may lead into new and intriguing insights in this topic. While the potential of V4þ/V3þ couple is slightly higher that the “stability window” of most electrolytes and full removal of alkali metal cations for the V5þ/V4þ couple seems to be hindered, the future development of such compounds based on other d-metals redox active centers will be promising. Titanium-based NaxTiOPO4 (x / 1) prepared through ion-exchange from NH4TiOPO4 (synthesized under mild hydrothermal conditions at 200  C for 3 days) is an attractive phase for design of solid-electrolyte interface (SEI)-free negative electrode materials with an average sodiation potential about 1.4 V. Unfortunately, specific capacity of only about 85 mAh g 1 was achieved so far, this compound requires further investigation (Fig. 9E).226,227

7.03.2.4

Compounds with other oxoanions and mixed anion groups: Silicates, carbonate-phosphates, Prussian Blue analogs

From the crystal chemistry point of view, an orthosilicate anion group SiO44 is similar to PO43 and SO42. One of specific feature of the silicate group is its ability to catenate into (SixOy)n polyanions: di- tri-silicates etc.228–232 Nevertheless, at this moment compounds with general formula NaxMSiO4 (M – transition metal cation or their mixture) are the most investigated. Generally, their crystal structure is an ordered mixture of NaO4, MO4 and SiO4 tetrahedra (Fig. 10B and C)229,233–235 with high thermal stability due to strong covalent bonding within the Si-O units. Among all these silicates, the Fe- and Mn-based compositions are the candidates of main interest. Sodium-iron orthosilicate, Na2FeSiO4, was first reported by Kee et al. in 2016. It is worth mentioning, that the XRD pattern of Na2FeSiO4 is quite complex. At the first glance, one can suppose that Na2FeSiO4 is isostructural to Na2MnSiO4 and Na2ZnSiO4 (both structures were assigned to space group Pn, Z ¼ 2) (Fig. 10B). However, it was proven that the space group P1 (Z ¼ 1) describes this crystal structure more precisely. This composition might demonstrate attractive capacity of 276 mAh g 1 corresponding to two-electron transfer, but it is worth expecting only one-electron process delivering 138 mAh g 1. Nonetheless, it was shown that the pristine electrode material demonstrates poor electrochemical activity, which is coupled with irreversible amorphization during cycling (Fig. 10D).236 Use of carbon coating formed during thermal destruction of sucrose at 600  C under inert atmosphere improves to some extent electrochemical behavior of Na2FeSiO4 (Fig. 10D). An illustrative example of a more complex silicate is sodium-iron pyrosilicate Na2Fe2Si2O7, which was originally proposed by Yamada’s group (Fig. 10A). However, even being thoroughly mixed with a conductive carbon, the material demonstrated capacity of only about 20 mAh g 1 with a plateau centered at 2.6 V vs. Naþ/Na.232 Na2MnSiO4 (space group Pn, theoretical capacity is 278 mAh g 1 for two-electron process) was for the first time investigated as a material for SIBs by Hagiwara’s group in 2014 (Fig. 10B).237 Noteworthy, the idea of using Na2MnSiO4 in batteries in principle refers to experiments by Dunkan et al., who suggested it as a template for the synthesis of Li2MnSiO4 (space group Pn). Earlier, Li2MnSiO4 as an electrode material was also mentioned by a number of authors233,238239,240 The 3D-framework of Na2MnSiO4 is very stable and it does not collapse during substitution of Na by Li ions. At temperatures 298, 323 and 363 K Na2MnSiO4/C electrode delivers 70, 94 and 125 mAh g 1 capacity, respectively, with well-resolved plateau at  3.0 V vs. Naþ/Na and acceptable

62

Electrode materials for reversible sodium ions de/intercalation

Fig. 10 (A) Phase diagram of the Na2O-FeO-SiO2 pseudoternary system. Color map indicates expected initial charge capacity based on the one Naþ extraction referring to the Fe3þ/Fe2þ redox couple. (B) Axonometric projection of Na2MSiO4 crystal structure (M ¼ Fe, Mn), polyhedra are shown on the figure (ICSD code number 244408); (C) Crystal structure and Fe coordination in Na2Fe2Si2O7. (D) The 5 first charge/discharge curves of Na2FeSiO4 at 1/40C (1 C ¼ 138 mA$ g 1, electrolyte 1 M NaPF6 in EC:DEC) and its cyclic performance in the voltage range of 1.0–4.1 V vs. Naþ/Na. (E) Galvanostatic chargedischarge profiles of the Na3Mn(PO4)(CO3) in Na cell (cycling rate C/30, electrolyte 1 M NaPF6 in EC:DEC). (F) Crystal structure of sidorenkite, Na3Mn(PO4)(CO3), in an axonometric projection (code ICSD database 200789). (A) Reprinted with permission from Panigrahi, A.; Nishimura, S.I.; Shimada, T.; Watanabe, E.; Zhao, W.; Oyama, G.; Yamada, A. Chem. Mater. 2017, 29, 4361–4366. Copyright (2017) American Chemical Society. (B, C) Reprinted with permission from Panigrahi, A.; Nishimura, S.I.; Shimada, T.; Watanabe, E.; Zhao, W.; Oyama, G.; Yamada, A. Chem. Mater. 2017, 29, 4361–4366. Copyright (2017) American Chemical Society. (D) Reprinted with permission from Kee, Y.; Dimov, N.; Staykov, A.; Okada, S. Mater. Chem. Phys. 2016, 171, 45–49. Copyright (2016) Elsevier. (E) Reprinted with permission from Hassanzadeh, N.; Sadrnezhaad, S.K.; Chen, G. Electrochim. Acta. 2016, 208, 188–194. Copyright (2016) Elsevier.

cycling stability.237 However, possible dissolution of Mn ions (resulting in capacity fading) should not be ruled out, particles coating and electrolyte optimization could become a reasonable way for overcoming these issues.241 Silicates are earth-abundant, cheap and eco-friendly compounds, which can be regarded as a possible basis for green and truly sustainable electrode materials for batteries. Nonetheless, silicates preparation requires a tight control of the synthesis conditions,

Electrode materials for reversible sodium ions de/intercalation

63

mainly in order to avoid formation of other phases, such as, Na2SiO3 and various amorphous impurities. High hygroscopicity, poor alkali metal ion diffusion and low electronic conductivity are inherent features of silicates, which creates playground for extensive fundamental and applied research aimed at overcoming these difficulties.232 Borates can also be regarded as electrode materials, at the same time, they suffer from the same issues as silicates since borateanion (BO33) demonstrates tremendous tendency to form both crystallized and amorphous phases with highly unstable and reactive surface.242,243 Two oxoanion groups can be combined giving rise to mixed-anion framework materials. One of such examples are phosphatecarbonates belonging to the bradleyite mineral group: sidorenkite Na3Mn(PO4)(CO3)244 and bonshtedtite Na3Fe(PO4)(CO3),245 which were discovered on Khibiny massif, Kola Peninsula.239 Their crystal structures are quite similar and described by cornersharing MO6 octahedra and PO4 tetrahedra, forming layers parallel to the bc plane of the monoclinic structure (SG P21/m). The layers of carbonate groups are inlaid, sharing an edge with the octahedra (Fig. 10F) along the a axis. The investigation of these mixed phosphates was preceded by ab initio calculations, which predicted possibility of the reversible de/insertion of two Naþ at a twoelectron reaction (Mn3þ/Mn2þ and Mn4þ/Mn3þ sequentially) per formula unit with capacity of 191 mAh g 1. Later, Na3Mn(PO4)(CO3) was synthesized under mild hydrothermal conditions (120  C) and investigated in Na half-cells.240 Unfortunately, the material exhibits a considerable irreversible capacity loss, however, it is able to deliver about 100 mAh g 1 (at the 5th cycle) at an average operating potential of 3.2 V vs. Naþ/Na during subsequent cycles (Fig. 10E).246 Recently significant attention was devoted to the series of compounds known as Prussian Blue Analogs (PBAs) described by the general formula AxM(1)[M(2)(CN)6]y $ zH2O, where (M(1), M(2)) are Fe, Co, Mn, Ni, Zn etc. and A is Li, Na, K, Rb. PBAs have been known since the XVIII century but for many years were used only as pigments.247 Nowadays, hexacyanoferrates are used in a number of applications, such as biosensors for amperometric detection of various organic species,248,249 agents for removal of radioactive ions from water,250 smart glasses, and energy storage materials.251,252  space group (Fig. 11)253 The first structure of this family Fe4[Fe(CN)6]3 $xH2O was solved by Buser et al. in 1977 in the Pm3m containing iron in both þ 2 and þ 3 oxidation states (Fig. 11A). The iron atoms alternate in the cubic lattice, being interconnected by C–N groups: every Fe3þ is surrounded by six nitrogen atoms, while Fe2þ is surrounded by six carbon atoms. A specific ratio of Fe2þ and Fe3þ species (3:4) in this compound is related to the requirement of charge neutrality determining the presence of 25% of the [Fe2þ(CN)6]4 vacancies. The resulting void can be filled with H2O molecules. Two types of H2O positions can be distinguished in the crystal structure: one is located near the Fe3þ (up to six molecules, the so-called “coordinated water”), another is at the A site (the so-called “interstitial water”). A huge number of reports devoted to PBAs electrochemical properties in sodium cells was published during last decades.251 However, the most attention is focused on the iron- and manganese-based Na2-d[M1M2(CN)6] compositions proposed as highly stable positive electrodes. The deviation from 2 leads to the lattice distortion: for instance, at lower Na contents the rhombohedral symmetry is adopted (Fig. 11D). A rhombohedrally distorted PBA, Na1.92Fe[Fe(CN)6] also called Prussian White demonstrates 120 mAh g 1 discharge capacity at the 750th cycle with two well-resolved plateaus at  3.00 and  3.3 V vs. Naþ/Na corresponding to the electrochemical transitions at two different irons (Fig. 11B and C).254 The water content in PBA compounds affects the electrochemical properties - the conventional statement is that the presence of water is a negative factor. Due to the presence of H2O the charge cutoff potential is limited to the oxidative potential of water (3.94 V vs. Naþ/Na). Overcoming this value leads to formation of gases, transformations of the structure with its subsequent collapse.255 Several groups reported acceptable values (e.g. at least 70 mAh g 1) for hundreds of cycles, but the most number of cycling stable compounds are dehydrated.256 The operating potential of Mn-based PBAs is, as it is expected, usually higher (the difference is up to 0.5 V) and capacities are in the same range as for the Febased ones.252,257 Na1.24Mn[MnII(CN)6] is a rare example of MnII/MnI transition. The Mn reduction occurs in a potential range 1.2–0.6 V vs. SHE and existence of MnI was confirmed with soft X-ray spectroscopy and resonant inelastic X-ray scattering.257 PBA materials have reached commercial maturity being utilized as positive and negative electrodes in the consumer Na-ion batteries produced by Novasis, Natron,258 and recently announced by CATL.259

7.03.3

Negative electrode materials

According to chemical compositions and mechanisms of interaction with Naþ, the negative electrode materials can be grouped as follows: 1. 2. 3. 4. 5.

Metallic sodium – plating-stripping; Carbonaceous – intercalation/adsorption260; Oxides and polyanions – intercalation (solid-solution/two-phase)6; p-elements – alloying67,261; Transition metal oxides and related compounds (TMO) – intercalation/conversion.262

Intercalation-type anode materials are most widely investigated at the moment. Their achievable capacities are strictly limited by the composition and crystal structure resulting in values usually not exceeding 100–200 mAh g 1 for oxides and polyanions. For carbonaceous materials the specific capacities reach higher values of 200–350 mAh g 1. Noteworthy, alloying-type anode materials (to

64

Electrode materials for reversible sodium ions de/intercalation

Fig. 11 (A) PBAs have a face-centered cubic geometry and open-framework lattice. The green and dark-blue atoms are transition-metal ions at the R site and P site, respectively. The yellow atoms are inserting ions. Gray atoms are carbon and light-blue atoms are nitrogen. Alkali-metal ions insert into the subcubes of the lattice as the transition-metal ions change oxidation state. (B) Electrochemical performance of a R-FeHCF electrode (86 wt% active material, 7 wt% Ketjen black, 7 wt% PTFE). (C) Discharge curves of the R-FeHCF/Na half-cell at different currents. d) Schematic three-phase evolutions during cycling of rhombohedral PB. (A) Reprinted with permission from Hurlbutt, K.; Wheeler, S.; Capone, I.; Pasta, M. Joule. 2018, 2, 1950–1960. Copyright (2018) Elsevier. (B) Reprinted with permission from Wang, L.; Song, J.; Qiao, R.; Wray, L.A.; Hossain, M.A.; Chuang, Y. De; Yang, W.; Lu, Y.; Evans, D.; Lee, J.J.; Vail, S.; Zhao, X.; Nishijima, M.; Kakimoto, S.; Goodenough, J.B. J. Am. Chem. Soc. 2015, 137, 2548–2554. Copyright (2015) American Chemical Society. (C) Reprinted with permission from Wang, L.; Song, J.; Qiao, R.; Wray, L.A.; Hossain, M.A.; Chuang, Y. De; Yang, W.; Lu, Y.; Evans, D.; Lee, J.J.; Vail, S.; Zhao, X.; Nishijima, M.; Kakimoto, S.; Goodenough, J.B. J. Am. Chem. Soc. 2015, 137, 2548– 2554. Copyright (2015) American Chemical Society. (D) Reprinted from Wang, W.; Gang, Y.; Hu, Z.; Yan, Z.; Li, W.; Li, Y.; Gu, Q.F.; Wang, Z.; Chou, S.L.; Liu, H.K.; Dou, S.X. Nat. Commun. 2020, 11, 1–9. CC BY 4.0.

be discussed further) could offer even higher energy densities for the full cells.263 The other criteria for the negative electrode material of SIB can be formulated as follows: 1. Low operating potential. The ideal negative electrode material should exhibit electrochemical activity at low (below 1.0 V) potentials maximizing the cell voltage and conferring high energy density, but not too low to experience sodium plating; 2. Reversibility and cycling stability. Usually, some irreversible capacity is observed at the several initial cycles with a major contribution from the SEI formation stemming from the partial decomposition of the electrolyte with consumption of the Naþ ions. As a result, the amount of charge carriers becomes lower which typically deteriorates the electrochemical behavior of the cell. One of the ways to tackle this problem is performing electrochemical pre-sodiation of the anode material to compensate for the

Electrode materials for reversible sodium ions de/intercalation

65

Naþ loss due to SEI formation. While this method seems to be attractive for research purposes, it is not easily scalable. Another approach for overcoming this issue is to use a so-called sacrificial salt.264 During the first cycle, the anion of the sacrificial salt undergoes an entire decomposition coupled with production of gas (N2, CO or CO2) and an extra amount of sodium ions to restore their concentration in the electrolyte. Additionally, various additives such as fluorinated organics were proposed to enable the formation of a more stable and flexible SEI265; 3. Sustainability, costs, safety. Negative electrode materials should not contain possibly hazardous elements. Some carbonaceous materials seem to be prospective, however, at the same time these display low first-cycle coulombic efficiency, insufficient C-rate performance and are very sensitive to storage conditions. In light of this criteria, various iron-, and titanium-based compounds appear to be attractive.266,267

7.03.3.1

Carbon-based negative electrode materials

The vast majority of investigated composites and materials of this class are based on sp2-hybridized carbons. Although several groups reported electrochemical activity of sp3 carbons as well as in specific nanostructures or thin films for capacitors or matrices for metal-air systems, in general their applicability in conventional SIBs are doubtful. This is related to their physical properties – sp3carbons are typically wide bandgap semiconductors with rather low electronic conductivity.268 One of the most common forms of carbon in modern lithium-ion batteries is graphite.269 The crystal structure of graphite consists of hexagonal layers (also referred to as graphene layers) stacked in parallel by Van der Waals forces. Due to the low energy of this forces, layers can easily glide under the mechanical stress.270 If we take into account diffraction data (a schematic view of XRD patterns is given on Fig. 12B), one can observe the most intensive diffraction peak for graphite at about 3.35 Å d-spacing and the less intensive peaks are observed at  2 Å, giving direct information about the polymorphic modification (2H or 3R, description is given in Fig. 12B). The most intensive peak corresponds to the interplanar distance between the infinite hexagonal graphene layers. Nowadays, graphite is one of the most widespread negative electrode components for Li-ion batteries – the best commercially available graphite demonstrates capacities close to the theoretical one (372 mAh g 1). At the same time sodium ions do not perform well when used with graphite as an intercalation host. The only compound described by the formula NaC64 can be formed during intercalation resulting in a negligible capacity of 35 mAh g 1.271 The thermodynamic stability of various compounds in a sodium-graphite system is still a matter of debate. One of the conventional hypotheses is related to the following phenomena: (i) structural deformation of resulting Na-GICs (GIC – graphite intercalation compound) (ii) the difference in the binding nature of the alkali metal ions and graphite. The DFT calculations of the Li/Na/K-graphite systems are in agreement with experimental results. It was demonstrated that the deformation energy of the graphite lattice linearly increases with the size of alkali metal cations.272 Surprisingly, in the middle of 2010-s co-intercalation of Naþ-diglyme couple into graphite was discovered by Adelheim et al. The specific capacity near 100 mAh g 1 was observed for a Na ||graphite cell.273,274 Finally, in 2016 a proof of concept was presented – a full cell, consisting of Na0.7CoO2 as a cathode, graphite as anode and 1 M NaClO4 in tetraethylene glycol dimethyl ether (TEGDME) as electrolyte can be considered as a device for stationary storage applications. The ether enables the co-intercalation of solvated sodium ions by forming ternary GICs (Fig. 12C and D).275 The first report by Doeff et al. on a reaction between sodium ions and petroleum coke dates back to 1993.276 At 86  C the coke demonstrates reversible electrochemical reaction with sodium ions, forming NaC24 and delivering the capacity about 90 mAh g 1. This is a representative of so-called “soft carbon” (SC) (Fig. 13) or graphitizable carbon, which undergoes a transformation into graphite at temperatures above 2200  C.277 This graphitization process can be described as a transition from a fully turbostratic initial structure to the “hexagonal” or “cubic” close-packed stacking of graphene layers (Fig. 12A). The lamellar structure determines a presence of broad intensity maxima on the XRD pattern at angles close to those of graphite. A slight shift toward smaller angles is related to the increase of the interlayer distance and the noticeable broadening of peaks is attributed to the crystallinity loss. The interlayer reflection (101) and (012) totally disappear indicating a complete loss of the plane-to-plane coherency. Only the (100) reflection is observed with its presence pointing to the preservation of strong intralayer C-C covalent bonds remaining unstretched.278 Typically, soft carbon is derived from organic aromatic compounds, for example, pitch or tar.279 It is also well-known as a coproduct of coal purifying and plastics production.277 Soft carbon demonstrates a generally sloping potential during cycling with a lack of the low-potential plateau in comparison to hard-carbon described further. The major part of the reversible capacity is usually observed at potentials close to the metal sodium plating. This, in turn, might cause dendrites formation followed by an electrical short-circuit and battery collapse. Bommier et al. hypothesized that the irreversibility at the first cycle is associated with the intercalation of sodium ions between graphitic layers. Studies on SCs of materials demonstrated that the large irreversible capacity at first cycles is their inalienable property (Fig. 13D and E). Similar to SCs, “hard carbon” (HC) also attracted a lot of attention clearly demonstrated by the numerous research articles and reviews starting from the beginning of 2010-s.260,280,281 Hard carbon appears as a prospective object for further optimization and implementation in SIBs due to relatively high capacity (up to hundreds mAh g 1), suitable working potential (mainly below 1 V vs. Naþ/Na) and decent cycling stability. Feature of particular interest of HCs is that they can be obtained from a huge variety of organic substances: from food wastes to scarps. Oxygen-rich precursors demonstrate tendency to form HCs, for example, wood,282 cellulose,283,284 sugar,285 cotton,286 various resins,287,288 resorcinol-formaldehyde,289 etc. Almost infinite choice of precursors and

66

Electrode materials for reversible sodium ions de/intercalation

Fig. 12 (A) Hard-carbon formation scheme as a function of temperature. The circles in the polymeric-structure schemes represent moieties like functional groups and side chains. (B) Schematic view of XRD pattern (Cu X-ray tube, l ¼ 1.5406 Å) of a hard carbon, a soft carbon and a graphite from top to bottom respectively. For graphite 2H and 3R notations refer to various types of layers stacking, 2H means ABA type, 3R – to the ABCA stacking type. The ratio between peaks of 2H ((100) and (101)) and 3R ((101) and (012)) corresponds to relative amounts of these two phases (for better visibility the intensities of these peaks were multiplied by 10. Miller indexes and schematic view of carbons are shown on the figure; (C) Voltage profiles of the galvanostatic cycling tests of a Na/Graphite half-cell employing the various electrolytes (highlighted on the figure); (D) Voltage profiles at various cycle number of the Graphite | 1 M NaClO4-TEGDME | Na0.7CoO2 full-cell galvanostatically cycled at 1 C (175 mA $ g 1) within the 0.5–3.7 V voltage range at room temperature (20  C). (A) Adapted from Dou, X.; Hasa, I.; Saurel, D.; Vaalma, C.; Wu, L.; Buchholz, D.; Bresser, D.; Komaba, S.; Passerini, S. Mater. Today. 2019, 23, 87–104 with permission. Copyright (2019) Elsevier. (B–D) Reprinted with permission from Hasa, I.; Dou, X.; Buchholz, D.; Shao-Horn, Y.; Hassoun, J.; Passerini, S.; Scrosati, B. J. Power Sources 2016, 310, 26–31. Copyright (2016) Elsevier.

tuning of thermal treatment conditions provide great opportunities for further commercialization and consequent industrial production. On the other hand, composition and microstructure of the initial mixture of organic compounds might scatter a lot, creating severe challenges for large-scale reproducible production.290 Hard carbon, as it is implied by its name, is a mechanically tough form of carbon. Mechanical properties are closely related to the microstructure of HC – a high degree of cross-linking of graphene layers prevents their gliding under mechanical stress. This, in turn is directly associated with the arrangement of molecules in the initial organic precursor. In contrast to SC, hard carbon is unable to graphitize completely even at temperatures higher than 3000  C, i.e., this is a form of “non-graphitizable carbon.” HC represents a highly amorphous structure. The (002) reflection on the XRD pattern is further shifted toward larger d-spacings and is even broader, than that of SC. At the same time, the (100) reflection is quite similar to that of SC (Fig. 12B).291 This resemblance indicates that the main difference between SC and HC is in the manner of graphene layers stacking. The cross-linking of graphene layers in HC also results in a specific fine structure of voids and pores. The internal structure includes defect sites, “open” surface, and pores, which all can act as storage sites for Naþ providing a particular storage mechanism for any specifically synthesized material. Conventionally, pores in HC are categorized into mesopores – less than 50 nm, micropores – less than 2.0 nm, and ultramicropores – less

Electrode materials for reversible sodium ions de/intercalation

67

Fig. 13 HRTEM (High-resolution transmission electron microscopy) images of (A) hard carbon (synthesized from sucrose), (B) soft carbon synthesized at 900  C, (C) soft carbon synthesized at 1600  C. The graphene layers in hard carbon are curved and not aligned. These in soft carbon exhibit less curvature and are much better aligned. Upon annealing at 1600  C, the resulting soft carbon displays a quasi-graphitic structure under transmission electron microscopy, which reveals the graphitizable nature of this carbon. (D) Typical view of first cycles sodiation/desodiation potential profiles of soft carbon synthesized at 900  C and soft carbon synthesized at 1600  C at 20 mA $g 1; (E) Cyclic voltammetry curves of soft carbon synthesized at 900  C and hard carbon with a scan rate of 0.1 mV$ s 1. Reprinted with permission from Jian, Z.; Bommier, C.; Luo, L.; Li, Z.; Wang, C.; Greaney, P.A.; Ji, X. Chem. Mater. 2017, 29(5), 2314–2320. Copyright (2017) American Chemical Society.

than 1.0 nm. The difference in densities of graphite (2–2.25 g/cm3) and non-graphitizable carbons (1.4–1.7 g/cm3) additionally proves the distinct layers arrangements in these carbonaceous materials.260

7.03.3.1.1

Approaches for describing mechanisms of sodium ion storage

In most studies, the charge/discharge process in HC is divided into two stages – the first one is a sloping region above 0.1 vs. Na/Naþ and the second one is a plateau from 0.1 to  0 V vs. Naþ/Na. In general, these regions are usually assigned to the following consecutive mechanisms: “intercalation – adsorption” or “adsorption – intercalation.” The intercalation – adsorption scheme was originally proposed by Stevens and Dahn.292,293 In frame of their model, the sloping region is attributed to Naþ intercalation into pseudo-graphitic areas forming a so-called “house of cards,” and at the plateau Naþ ions adsorb in the micropores eventually forming Na nanoclusters.294–296 The reverse sequence – “adsorption – intercalation” – for the first time was described by Cao et al.297 The sloped region was attributed for adsorption of sodium ions on the surface, their interaction with defect sites and various functional groups. The plateau corresponds to further intercalation between sheets of graphene.298,299 More in-depth approaches imply dividing the curve into three or even more regions. For example, Bommier et al. suggested a change in the mechanism of “sodiation” near to the cutoff potential and supposed sodium ions adsorption on HC pore surfaces (Fig. 14B).300 The approach by Reddy et al. assumes that the sloping region corresponds to sodium adsorption on defective sites with a consequent Naþ intercalation and formation of NaC24 (Fig. 14C).280 The extended four-staged model by Weaving and Alvin et al., considers adsorption on defective sites and two steps of adsorption in micropores occurring before and after intercalation into

68

Electrode materials for reversible sodium ions de/intercalation

Fig. 14 (A) Model of non-graphitizing carbon. (B) Visual representation of the card-house model on Na-ion storage in hard carbon. The two distinct phases: intercalation inside turbostatic nanodomains and pore filling are seen. (C) Galvanostatic curve (discharge) of hard carbon by Reddy et al. The color code qualitatively represents different processes during discharge. The inset schematically depicts these processes at the atomic scale, using the corresponding color code for the respective atoms. (D) Linear sweep voltammetry (0.01 mV$ s 1) and proposed model of processes occurring during desodiation of hard carbon. Color regions of LSV plot represent different charge storage processes. (E) Voltage profile (solid line) and G-peak position (discrete points) for discharge of cycle 1, indicating regions of SEI formation (gray shading), intercalation (light pink), and pore occupancy (violet); stylized visualization (not to scale) illustrating sodiation mechanism in hard carbon: intercalation (light pink) between graphene layers occurs over the sloping region and pore occupancy (violet) occurs over the low-voltage plateau of the discharge curve. The approximate dimensions of the model in are consistent with those commonly reported for hard carbon comprising curved TNDs with crystallite size dimensions 1–2 nm (out-ofplane) and 2–4 nm (in-plane) encasing micropores with dimensions of 1000  C) cannot be neglected. If the first issue, in principle, can be overcome by the interplay between composition of electrolyte and surface design, the second one is a necessary condition for the formation of a desirable structure of carbon.390 If all these features (temperature conditions, price of pure initial reagents and electrolyte additives, and the recycling process of spent HCs) are taken into account, the question regarding the attractiveness of HCs in terms of price and sustainability principles arises. Additionally, low tap density of various forms of carbons is also an industrial challenge; (iii) Among groups 14 and 15 elements, Sn, Sb and their alloys should be mentioned. As all alloy-based electrode materials, they suffer from noticeable volume changes, but in comparison with heavy and toxic elements (As, Pb, e.g.) a further optimization of electrode does not look hopeless. Implementation of new “breathing” binders and electrolyte mixtures may probably lead to prospective outcomes; (iv) A number of polyanion compounds (NASICON-structured NaTi2(PO4)3, KTiOPO4-structured NaTiOPO4, e.g.) were investigated as negative electrode materials during last two decades. Despite the delivered capacities are smaller than those for the other classes of materials, the higher thermal and mechanical stability may also be useful properties for a specific area of applications, such as large-scale stationary storage.384

Acknowledgment This chapter is prepared with the financial support from Russian Science Foundation (RSF grant #17-73-30006).

References 1. Fotedar, V.; Mongird, K.; Koritarov, V.; Viswanathan, V.; Hadjerioua, B.; Balducci, P.; Alam, J. Energy Storage Technology and Cost Characterization Report, Pacific Northwest National Lab (PNNL), 2019. 2. Uribe, J. M.; Restrepo, N.; Guillen, M. Doc. Treb. 2021, 1. 3. Kudryavtsev, P. G. Altern. Energy Ecol. 2016, 72–88. 4. Martin, G.; Rentsch, L.; Höck, M.; Bertau, M. Energy Storage Mater. 2017, 6, 171–179. 5. Dessemond, C.; Lajoie-Leroux, F.; Soucy, G.; Laroche, N.; Magnan, J. F. Minerals 2019, 9. 6. Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Chem. Rev. 2014, 114, 11636–11682. 7. Tarascon, J.-M. Joule 2020, 4, 1616–1620. 8. Fleischmann, S.; Mitchell, J. B.; Wang, R.; Zhan, C.; Jiang, D. E.; Presser, V.; Augustyn, V. Chem. Rev. 2020, 120, 6738–6782. 9. Choi, C.; Ashby, D. S.; Butts, D. M.; DeBlock, R. H.; Wei, Q.; Lau, J.; Dunn, B. Nat. Rev. Mater. 2020, 5, 5–19. 10. Damhus, T.; Hartshorn, R. M.; Hutton, A. T. Nomenclature of Inorganic Chemistry: IUPAC Recommendations 2005, Chemistry International, 2005; pp 325–328. 11. Hartshorn, R. M.; Hellwich, K.-H.; Yerin, A.; Damhus, T.; Hutton, A. T. Pure Appl. Chem. 2015, 87, 1039–1049. 12. Goikolea, E.; Palomares, V.; Wang, S.; de Larramendi, I. R.; Guo, X.; Wang, G.; Rojo, T. Adv. Energy Mater. 2020, 10, 2002055. 13. Delmas, C.; Carlier, D.; Guignard, M. Adv. Energy Mater. 2021, 11, 2001201. 14. Delmas, C.; Braconnier, J.; Fouassier, C.; Hagenmuller, P. Solid State Ionics 1981, 3–4, 165–169. 15. Momma, K.; Izumi, F. J. Appl. Crystallogr. 2011, 44, 1272–1276. 16. Delmas, C.; Fouassier, C.; Hagenmuller, P. Phys. Bþ C 1980, 99, 81–85. 17. Gao, R.-M.; Zheng, Z.-J.; Wang, P.-F.; Wang, C.-Y.; Ye, H.; Cao, F.-F. Energy Storage Mater. 2020, 30, 9–26. 18. Guignard, M.; Didier, C.; Darriet, J.; Bordet, P.; Elkaïm, E.; Delmas, C. Nat. Mater. 2013, 12, 74–80. 19. Radin, M. D.; Hy, S.; Sina, M.; Fang, C.; Liu, H.; Vinckeviciute, J.; Zhang, M.; Whittingham, M. S.; Meng, Y. S.; Van der Ven, A. Adv. Energy Mater. 2017, 7, 1602888. 20. Ong, S. P.; Chevrier, V. L.; Hautier, G.; Jain, A.; Moore, C.; Kim, S.; Ma, X.; Ceder, G. Energy Environ. Sci. 2011, 4, 3680. 21. Fouassier, C.; Matejka, G.; Reau, J.-M.; Hagenmuller, P. J. Solid State Chem. 1973, 6, 532–537. 22. Ohzuku, T.; Ueda, A.; Nagayama, M. J. Electrochem. Soc. 1862, 1993, 140. 23. Molenda, J.; Delmas, C.; Hagenmuller, P. Solid State Ionics 1983, 9–10, 431–435. 24. Wang, X.; Tamaru, M.; Okubo, M.; Yamada, A. J. Phys. Chem. C 2013, 117, 15545–15551. 25. Carlier, D.; Cheng, J. H.; Berthelot, R.; Guignard, M.; Yoncheva, M.; Stoyanova, R.; Hwang, B. J.; Delmas, C. Dalt. Trans. 2011, 40, 9306. 26. Cheng, J.-H.; Pan, C.-J.; Lee, J.-F.; Chen, J.-M.; Guignard, M.; Delmas, C.; Carlier, D.; Hwang, B.-J. Chem. Mater. 2014, 26, 1219–1225. 27. Yoshida, H.; Yabuuchi, N.; Komaba, S. Electrochem. Commun. 2013, 34, 60–63. 28. Parant, J.-P.; Olazcuaga, R.; Devalette, M.; Fouassier, C.; Hagenmuller, P. J. Solid State Chem. 1971, 3, 1–11. 29. Ma, X.; Chen, H.; Ceder, G. J. Electrochem. Soc. 2011, 158, A1307. 30. Chen, X.; Wang, Y.; Wiaderek, K.; Sang, X.; Borkiewicz, O.; Chapman, K.; LeBeau, J.; Lynn, J.; Li, X. Adv. Funct. Mater. 2018, 28, 1805105.

Electrode materials for reversible sodium ions de/intercalation

77

31. Li, X.; Ma, X.; Su, D.; Liu, L.; Chisnell, R.; Ong, S. P.; Chen, H.; Toumar, A.; Idrobo, J.-C.; Lei, Y.; Bai, J.; Wang, F.; Lynn, J. W.; Lee, Y. S.; Ceder, G. Nat. Mater. 2014, 13, 586–592. 32. Kumakura, S.; Tahara, Y.; Kubota, K.; Chihara, K.; Komaba, S. Angew. Chemie. 2016, 128, 12952–12955. 33. Dose, W. M.; Sharma, N.; Pramudita, J. C.; Avdeev, M.; Gonzalo, E.; Rojo, T. Chem. Mater. 2018, 30, 7503–7510. 34. Arnaiz, M.; Gómez-Cámer, J. L.; Gonzalo, E.; Drewett, N. E.; Ajuria, J.; Goikolea, E.; Galceran, M.; Rojo, T. Mater. Today Proc. 2021, 39, 1118–1131. 35. Zarrabeitia, M.; Gonzalo, E.; Pasqualini, M.; Ciambezi, M.; Lakuntza, O.; Nobili, F.; Trapananti, A.; Di Cicco, A.; Aquilanti, G.; Katcho, N. A.; López del Amo, J. M.; Carrasco, J.; Muñoz-Márquez, M.Á.; Rojo, T. J. Mater. Chem. A 2019, 7, 14169–14179. 36. Yabuuchi, N.; Hara, R.; Kajiyama, M.; Kubota, K.; Ishigaki, T.; Hoshikawa, A.; Komaba, S. Adv. Energy Mater. 2014, 4, 1301453. 37. de la Llave, E.; Talaie, E.; Levi, E.; Nayak, P. K.; Dixit, M.; Rao, P. T.; Hartmann, P.; Chesneau, F.; Major, D. T.; Greenstein, M.; Aurbach, D.; Nazar, L. F. Chem. Mater. 2016, 28, 9064–9076. 38. Rong, X.; Liu, J.; Hu, E.; Liu, Y.; Wang, Y.; Wu, J.; Yu, X.; Page, K.; Hu, Y.-S.; Yang, W.; Li, H.; Yang, X.-Q.; Chen, L.; Huang, X. Joule 2018, 2, 125–140. 39. Maitra, U.; House, R. A.; Somerville, J. W.; Tapia-Ruiz, N.; Lozano, J. G.; Guerrini, N.; Hao, R.; Luo, K.; Jin, L.; Pérez-Osorio, M. A.; Massel, F.; Pickup, D. M.; Ramos, S.; Lu, X.; McNally, D. E.; Chadwick, A. V.; Giustino, F.; Schmitt, T.; Duda, L. C.; Roberts, M. R.; Bruce, P. G. Nat. Chem. 2018, 10, 288–295. 40. Braconnier, J. J.; Delmas, C.; Hagenmuller, P. Mater. Res. Bull. 1982, 17, 993–1000. 41. Han, M. H.; Gonzalo, E.; Casas-Cabanas, M.; Rojo, T. J. Power Sources 2014, 258, 266–271. 42. De Boisse, B. Universite´ de Bordeaux, 2014. 43. Wang, X.; Liu, G.; Iwao, T.; Okubo, M.; Yamada, A. J. Phys. Chem. C 2014, 118, 2970–2976. 44. Saadoune, I.; Maazaz, A.; Ménétrier, M.; Delmas, C. J. Solid State Chem. 1996, 122, 111–117. 45. Paulsen, J. M.; Dahn, J. R. Solid State Ionics 1999, 126, 3–24. 46. Lu, Z.; Dahn, J. R. J. Electrochem. Soc. 2001, 148, A237. 47. Zhang, Y.; Wu, M.; Ma, J.; Wei, G.; Ling, Y.; Zhang, R.; Huang, Y. ACS Cent. Sci. 2020, 6, 232–240. 48. Lee, D. H.; Xu, J.; Meng, Y. S. Phys. Chem. Chem. Phys. 2013, 15, 3304. 49. Komaba, S.; Yabuuchi, N.; Nakayama, T.; Ogata, A.; Ishikawa, T.; Nakai, I. Inorg. Chem. 2012, 51, 6211–6220. 50. Wang, P.-F.; You, Y.; Yin, Y.-X.; Guo, Y.-G. J. Mater. Chem. A 2016, 4, 17660–17664. 51. Maazaz, A.; Delmas, C.; Hagenmuller, P. J. Incl. Phenom. 1983, 1, 45–51. 52. Wang, Y.; Yu, X.; Xu, S.; Bai, J.; Xiao, R.; Hu, Y.-S.; Li, H.; Yang, X.-Q.; Chen, L.; Huang, X. Nat. Commun. 2013, 4, 2858. 53. Wang, Y.; Xiao, R.; Hu, Y.-S.; Avdeev, M.; Chen, L. Nat. Commun. 2015, 6, 1–9. 54. Wang, J.; Liu, H.; Yang, Q.; Hu, B.; Geng, F.; Zhao, C.; Lin, Y.; Hu, B. ACS Appl. Mater. & Interfaces 2020, 12, 34848–34857. 55. Li, Y.; Yang, Z.; Xu, S.; Mu, L.; Gu, L.; Hu, Y.-S.; Li, H.; Chen, L. Adv. Sci. 2015, 2, 1500031. 56. Jin, T.; Wang, P.; Wang, Q.; Zhu, K.; Deng, T.; Zhang, J.; Zhang, W.; Yang, X.; Jiao, L.; Wang, C. Angew. Chemie 2020, 132, 14619–14624. 57. Clément, R. J.; Billaud, J.; Robert Armstrong, A.; Singh, G.; Rojo, T.; Bruce, P. G.; Grey, C. P. Energy Environ. Sci. 2016, 9, 3240–3251. 58. Matsui, M.; Mizukoshi, F.; Imanishi, N. J. Power Sources 2015, 280, 205–209. 59. Yu, F.; Du, L.; Zhang, G.; Su, F.; Wang, W.; Sun, S. Adv. Funct. Mater. 2020, 30, 1906890. 60. Alvarado, J.; Ma, C.; Wang, S.; Nguyen, K.; Kodur, M.; Meng, Y. S. ACS Appl. Mater. & Interfaces 2017, 9, 26518–26530. 61. Yakubovich, O. V.; Simonov, M. A.; Belov, N. V. Sov. Phys. Dokl. 1977, 347. 62. Streltsov, V. A.; Belokoneva, E. L.; Tsirelson, V. G.; Hansen, N. K. Acta Crystallogr. Sect. B Struct. Sci. 1993, 49, 147–153. 63. Padhi, A. K. J. Electrochem. Soc. 1997, 144, 1188. 64. Moore, P. B. Am. Mineral. J. Earth Planet. Mater. 1972, 57, 1333–1344. 65. Vignola, P.; Hatert, F.; Fransolet, A.-M.; Medenbach, O.; Diella, V.; Andò, S. Am. Mineral. 2013, 98, 767–772. 66. Kosova, N. V.; Podugolnikov, V. R.; Devyatkina, E. T.; Slobodyuk, A. B. Mater. Res. Bull. 2014, 60, 849–857. 67. Kulova, T. L.; Skundin, A. M. Russ. Chem. Bull. 2017, 66, 1329–1335. 68. Kim, J.; Seo, D. H.; Kim, H.; Park, I.; Yoo, J. K.; Jung, S. K.; Park, Y. U.; Goddard, W. A.; Kang, K. Energy Environ. Sci. 2015, 8, 540–545. 69. Yakubovich, O. V.; Khasanova, N. R.; Antipov, E. V. Minerals 2020, 10 (6), 524. 70. Heubner, C.; Lein, T.; Schneider, M.; Michaelis, A. J. Mater. Chem. A 2020, 8, 16854–16883. 71. Le Poul, N.; Baudrin, E.; Morcrette, M.; Gwizdala, S.; Masquelier, C.; Tarascon, J. M. Solid State Ionics 2003, 159, 149–158. 72. Oh, S. M.; Myung, S. T.; Hassoun, J.; Scrosati, B.; Sun, Y. K. Electrochem. Commun. 2012, 22, 149–152. 73. Berlanga, C.; Monterrubio, I.; Armand, M.; Rojo, T.; Galceran, M.; Casas-Cabanas, M. ACS Sustain. Chem. Eng. 2020, 8, 725–730. 74. Koleva, V. G.; Boyadzhieva, T. J.; Stoyanova, R. K. Cryst. Growth Des. 2019, 19, 3744–3754. 75. Boyadzhieva, T.; Koleva, V.; Kukeva, R.; Nihtianova, D.; Harizanova, S.; Stoyanova, R. RSC Adv. 2020, 10, 29051–29060. 76. Boyadzhieva, T.; Koleva, V.; Zhecheva, E.; Nihtianova, D.; Mihaylov, L.; Stoyanova, R. RSC Adv. 2015, 5, 87694–87705. 77. Goodenough, J. B.; Hong, H.-P.; Kafalas, J. A. Mater. Res. Bull. 1976, 11, 203–220. 78. Hong, H.-P. Mater. Res. Bull. 1976, 11, 173–182. 79. Ivanov-Schitz, A. K.; Sigarev, S. E. Fiz. Tverd. tela. 1986, 28, 3528–3531. 80. Ivanov-Schitz, A. K. Solid State Ionics New Dev. BVR Chowdari. Singapore 1996, 63–81. 81. Ivanov-Schitz, A. K.; Demyanets, L. N. Crystallogr. Rep. 2003, 48. 82. Delmas, C.; Nadiri, A.; Soubeyroux, J. L. Solid State Ionics. 1988, 28–30, 419–423. 83. Porter, D. B.; Olazcuaga, R.; Delmas, C.; Cherkaoui, F.; Brochu, R.; Le Flem, G. Chem. Informationsd. 1981, 12. 84. Jian, Z.; Hu, Y. S.; Ji, X.; Chen, W. Adv. Mater. 2017, 29 (20), 1601925. 85. Chung, S.-C.; Ming, J.; Lander, L.; Lu, J.; Yamada, A. J. Mater. Chem. A 2018, 6, 3919–3925. 86. Porter, D. B.; Olazcuaga, R.; Delmas, C.; Cherkaoui, F.; Brochu, R.; Leflem, G. Rev. Chim. Miner. 1980, 17, 458–465. 87. Li, S.; Dong, Y.; Xu, L.; Xu, X.; He, L.; Mai, L. Adv. Mater. 2014, 26, 3545–3553. 88. Wang, D.; Chen, N.; Li, M.; Wang, C.; Ehrenberg, H.; Bie, X.; Wei, Y.; Chen, G.; Du, F. J. Mater. Chem. A 2015, 3, 8636–8642. 89. Jian, Z.; Sun, Y.; Ji, X. Chem. Commun. 2015, 51, 6381–6383. 90. Zhang, Y.; Zhao, H.; Du, Y. J. Mater. Chem. A 2016, 4, 7155–7159. 91. Bui, K. M.; Dinh, V. A.; Okada, S.; Ohno, T. Phys. Chem. Chem. Phys. 2015, 17, 30433–30439. 92. Song, W.; Cao, X.; Wu, Z.; Chen, J.; Huangfu, K.; Wang, X.; Huang, Y.; Ji, X. Phys. Chem. Chem. Phys. 2014, 16, 17681–17687. 93. de Boisse, B. M.; Ming, J.; Nishimura, S.; Yamada, A. J. Electrochem. Soc. 2016, 163, A1469–A1473. 94. Jian, Z.; Han, W.; Lu, X.; Yang, H.; Hu, Y.-S.; Zhou, J.; Zhou, Z.; Li, J.; Chen, W.; Chen, D.; Chen, L. Adv. Energy Mater. 2013, 3, 156–160. 95. Duan, W.; Zhu, Z.; Li, H.; Hu, Z.; Zhang, K.; Cheng, F.; Chen, J. J. Mater. Chem. A 2014, 2, 8668–8675. 96. Klee, R.; Aragón, M. J.; Alcántara, R.; Tirado, J. L.; Lavela, P. Eur. J. Inorg. Chem. 2016, 19, 3212–3218. 97. Shen, W.; Li, H.; Wang, C.; Li, Z.; Xu, Q.; Liu, H.; Wang, Y. J. Mater. Chem. A 2015, 3, 15190–15201. 98. Tao, S.; Wang, X.; Cui, P.; Wang, Y.; Haleem, Y. A.; Wei, S.; Huang, W.; Song, L.; Chu, W. RSC Adv. 2016, 6, 43591–43597. 99. Zhou, W.; Xue, L.; Lü, X.; Gao, H.; Li, Y.; Xin, S.; Fu, G.; Cui, Z.; Zhu, Y.; Goodenough, J. B. Nano Lett. 2016, 16, 7836–7841.

78 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116. 117. 118. 119. 120. 121. 122. 123. 124. 125. 126. 127. 128. 129. 130. 131. 132. 133. 134. 135. 136. 137. 138. 139. 140. 141. 142. 143. 144. 145. 146. 147. 148. 149. 150. 151. 152. 153. 154. 155. 156. 157. 158. 159. 160. 161. 162. 163. 164. 165. 166. 167. 168. 169. 170. 171. 172.

Electrode materials for reversible sodium ions de/intercalation Li, H.; Jin, T.; Chen, X.; Lai, Y.; Zhang, Z.; Bao, W.; Jiao, L. Adv. Energy Mater. 2018, 8, 1801418. Chen, F.; Kovrugin, V. M.; David, R.; Mentre, O.; Fauth, F.; Chotard, J.-N.; Masquelier, C. Small Methods. 2019, 3, 1800218. Zakharkin, M. V.; Drozhzhin, O. A.; Tereshchenko, I. V.; Chernyshov, D.; Abakumov, A. M.; Antipov, E. V.; Stevenson, K. J. ACS Appl. Energy Mater. 2018, 1, 5842–5846. Gao, H.; Seymour, I. D.; Xin, S.; Xue, L.; Henkelman, G.; Goodenough, J. B. J. Am. Chem. Soc. 2018, 140, 18192–18199. Delmas, C.; Cherkaoui, F.; Nadiri, A.; Hagenmuller, P. Mater. Res. Bull. 1987, 22, 631–639. Park, S. I.; Gocheva, I.; Okada, S.; Yamaki, J. J. Electrochem. Soc. 2011, 158, A1067. Li, X.; Zhu, X.; Liang, J.; Hou, Z.; Wang, Y.; Lin, N.; Zhu, Y.; Qian, Y. J. Electrochem. Soc. 2014, 161, A1181. Wang, L.; Wang, B.; Liu, G.; Liu, T.; Gao, T.; Wang, D. RSC Adv. 2016, 6, 70277–70283. Wu, C.; Kopold, P.; Ding, Y.-L.; van Aken, P. A.; Maier, J.; Yu, Y. ACS Nano 2015, 9, 6610–6618. Yang, G.; Song, H.; Wu, M.; Wang, C. J. Mater. Chem. A 2015, 3, 18718–18726. Jiang, Y.; Shi, J.; Wang, M.; Zeng, L.; Gu, L.; Yu, Y. ACS Appl. Mater. & Interfaces 2016, 8, 689–695. Wang, D.; Liu, Q.; Chen, C.; Li, M.; Meng, X.; Bie, X.; Wei, Y.; Huang, Y.; Du, F.; Wang, C. ACS Appl. Mater. & Interfaces 2016, 8, 2238–2246. Difi, S.; Saadoune, I.; Sougrati, M. T.; Hakkou, R.; Edstrom, K.; Lippens, P.-E. J. Phys. Chem. C 2015, 119, 25220–25234. Difi, S.; Saadoune, I.; Sougrati, M. T.; Hakkou, R.; Edstrom, K.; Lippens, P.-E. Hyperfine Interact. 2016, 237, 61. Trad, K.; Carlier, D.; Croguennec, L.; Wattiaux, A.; Ben Amara, M.; Delmas, C. Chem. Mater. 2010, 22, 5554–5562. Yakubovich, O. V.; Simonov, M. A.; Egorov-Tismenko, Y. K.; Belov, N. V. Dokl. Akad. Nauk 1977, 1123–1126. Moore, P. B. Am. Mineral. J. Earth Planet. Mater. 1971, 56, 1955–1975. Moore, P. B.; Ito, J. Mineral. Mag. 1979, 43, 227–235. Huang, W.; Li, B.; Saleem, M. F.; Wu, X.; Li, J.; Lin, J.; Xia, D.; Chu, W.; Wu, Z. Chem. Eur. J. 2015, 21, 851–860. Liu, D.; Palmore, G. T. R. ACS Sustain. Chem. & Eng. 2017, 5, 5766–5771. Walczak, K.; Kulka, A.; Molenda, J. Solid State Sci. 2019, 87, 21–26. Kim, J.; Kim, H.; Lee, S.; Myung, S.-T. J. Mater. Chem. A 2017, 5, 22334–22340. Wen, M.; Liu, X.; Zhao, Y.; Liu, S.; Liu, H.; Dong, Y.; Kuang, Q.; Fan, Q. J. Solid State Electrochem. 2018, 22, 891–898. Dwibedi, D.; Baskar, S.; Barpanda, P. ECS Trans. 2017, 80, 337. Essehli, R.; Belharouak, I.; Ben Yahia, H.; Maher, K.; Abouimrane, A.; Orayech, B.; Calder, S.; Zhou, X. L.; Zhou, Z.; Sun, Y.-K. Dalt. Trans. 2015, 44, 7881–7886. Barpanda, P.; Oyama, G.; Nishimura, S.; Chung, S.-C.; Yamada, A. Nat. Commun. 2014, 5, 1–8. Oyama, G.; Pecher, O.; Griffith, K. J.; Nishimura, S.; Pigliapochi, R.; Grey, C. P.; Yamada, A. Chem. Mater. 2016, 28, 5321–5328. Chen, M.; Cortie, D.; Hu, Z.; Jin, H.; Wang, S.; Gu, Q.; Hua, W.; Wang, E.; Lai, W.; Chen, L. Adv. Energy Mater. 2018, 8, 1800944. Jungers, T.; Mahmoud, A.; Malherbe, C.; Boschini, F.; Vertruyen, B. J. Mater. Chem. A 2019, 7, 8226–8233. Araujo, R. B.; Chakraborty, S.; Barpanda, P.; Ahuja, R. Phys. Chem. Chem. Phys. 2016, 18, 9658–9665. Singh, P.; Shiva, K.; Celio, H.; Goodenough, J. B. Energy Environ. Sci. 2015, 8, 3000–3005. Nisar, U.; Gulied, M. H.; Shakoor, R. A.; Essehli, R.; Ahmad, Z.; Alashraf, A.; Kahraman, R.; Al-Qaradawi, S.; Soliman, A. RSC Adv. 2018, 8, 32985–32991. Trussov, I. A.; Kokhmetova, S. T.; Driscoll, L. L.; Smith, R.; Berry, F. J.; Marco, J. F.; Galeyeva, A. K.; Kurbatov, A. P.; Slater, P. R. J. Solid State Chem. 2020, 289, 121395. Dwibedi, D.; Araujo, R. B.; Chakraborty, S.; Shanbogh, P. P.; Sundaram, N. G.; Ahuja, R.; Barpanda, P. J. Mater. Chem. A. 2015, 3, 18564–18571. Dwibedi, D.; Gond, R.; Dayamani, A.; Araujo, R. B.; Chakraborty, S.; Ahuja, R.; Barpanda, P. Dalt. Trans. 2017, 46, 55–63. Ben Yahia, H.; Shikano, M.; Tabuchi, M.; Belharouak, I. Inorg. Chem. 2016, 55, 4643–4649. Gao, J.; Zhao, P.; Feng, K. Chem. Mater. 2017, 29, 940–944. Barpanda, P.; Liu, G.; Ling, C. D.; Tamaru, M.; Avdeev, M.; Chung, S.-C.; Yamada, Y.; Yamada, A. Chem. Mater. 2013, 25, 3480–3487. Uebou, Y.; Okada, S.; Yamaki, J. J. Power Sources 2003, 115, 119–124. Deng, W.; Feng, X.; Xiao, Y.; Li, C. ChemElectroChem. 2018, 5, 1032–1036. Barpanda, P.; Ye, T.; Nishimura, S. I.; Chung, S. C.; Yamada, Y.; Okubo, M.; Zhou, H.; Yamada, A. Electrochem. Commun. 2012, 24, 116–119. Kim, H.; Shakoor, R. A.; Park, C.; Lim, S. Y.; Kim, J.-S.; Jo, Y. N.; Cho, W.; Miyasaka, K.; Kahraman, R.; Jung, Y.; et al. Adv. Funct. Mater. 2013, 23, 1147–1155. Clark, J. M.; Barpanda, P.; Yamada, A.; Islam, M. S. J. Mater. Chem. A 2014, 2, 11807–11812. Chen, X.; Du, K.; Lai, Y.; Shang, G.; Li, H.; Xiao, Z.; Chen, Y.; Li, J.; Zhang, Z. J. Power Sources 2017, 357, 164–172. Song, H. J.; Kim, D.-S.; Kim, J.-C.; Hong, S.-H.; Kim, D.-W. J. Mater. Chem. A 2017, 5, 5502–5510. Park, C. S.; Kim, H.; Shakoor, R. A.; Yang, E.; Lim, S. Y.; Kahraman, R.; Jung, Y.; Choi, J. W. J. Am. Chem. Soc. 2013, 135, 2787–2792. Barpanda, P.; Ye, T.; Avdeev, M.; Chung, S. C.; Yamada, A. J. Mater. Chem. A 2013, 1, 4194–4197. Ha, K.-H.; Woo, S. H.; Mok, D.; Choi, N.-S.; Park, Y.; Oh, S. M.; Kim, Y.; Kim, J.; Lee, J.; Nazar, L. F. Adv. Energy Mater. 2013, 3, 770–776. Lin, B.; Zhang, S.; Deng, C. J. Mater. Chem. A 2016, 4, 2550–2559. Li, H.; Zhang, Z.; Xu, M.; Bao, W.; Lai, Y.; Zhang, K.; Li, J. ACS Appl. Mater. & Interfaces 2018, 10, 24564–24572. Barpanda, P.; Lu, J.; Ye, T.; Kajiyama, M.; Chung, S.-C.; Yabuuchi, N.; Komaba, S.; Yamada, A. RSC Adv. 2013, 3, 3857. Kim, H.; Park, C. S.; Choi, J. W.; Jung, Y. Angew. Chemie Int. Ed. 2016, 55, 6662–6666. Kee, Y.; Dimov, N.; Staikov, A.; Barpanda, P.; Lu, Y.-C.; Minami, K.; Okada, S. RSC Adv. 2015, 5, 64991–64996. Barpanda, P.; Liu, G.; Avdeev, M.; Yamada, A. ChemElectroChem. 2014, 1, 1488–1491. Kim, J.; Park, I.; Kim, H.; Park, K.-Y.; Park, Y.-U.; Kang, K. Adv. Energy Mater. 2016, 6, 1502147. Drozhzhin, O. A.; Tertov, I. V.; Alekseeva, A. M.; Aksyonov, D. A.; Stevenson, K. J.; Abakumov, A. M.; Antipov, E. V. Chem. Mater. 2019, 31, 7463–7469. Deng, C.; Zhang, S.; Zhao, B. Energy Storage Mater. 2016, 4, 71–78. Gezovic, A.; Vujkovic, M. J.; Milovic, M.; Grudic, V.; Dominko, R.; Mentus, S. Energy Storage Mater. 2021, 37, 243–273. Sanz, F.; Parada, C.; Amador, U.; Monge, M. A.; Valero, C. R. J. Solid State Chem. 1996, 123, 129–139. Sanz, F.; Parada, C.; Rojo, J. M.; Ruíz-Valero, C. Chem. Mater. 2001, 13, 1334–1340. Kim, H.; Park, I.; Lee, S.; Kim, H.; Park, K.-Y.; Park, Y.-U.; Kim, H.; Kim, J.; Lim, H.-D.; Yoon, W.-S. Chem. Mater. 2013, 25, 3614–3622. Kim, H.; Yoon, G.; Park, I.; Park, K.-Y.; Lee, B.; Kim, J.; Park, Y.-U.; Jung, S.-K.; Lim, H.-D.; Ahn, D.; Lee, S.; Kang, K. Energy Environ. Sci. 2015, 8, 3325–3335. Zarrabeitia, M.; Jauregui, M.; Sharma, N.; Pramudita, J. C.; Casas-Cabanas, M. Chem. Mater. 2019, 31, 5152–5159. Zhang, H.; Hasa, I.; Buchholz, D.; Qin, B.; Geiger, D.; Jeong, S.; Kaiser, U.; Passerini, S. NPG Asia Mater. 2017, 9, e370. Kim, H.; Park, I.; Seo, D.-H.; Lee, S.; Kim, S.-W.; Kwon, W. J.; Park, Y.-U.; Kim, C. S.; Jeon, S.; Kang, K. J. Am. Chem. Soc. 2012, 134, 10369–10372. Wood, S. M.; Eames, C.; Kendrick, E.; Islam, M. S. J. Phys. Chem. C 2015, 119, 15935–15941. Ma, X.; Pan, Z.; Wu, X.; Shen, P. K. Chem. Eng. J. 2019, 365, 132–141. Yuan, T.; Wang, Y.; Zhang, J.; Pu, X.; Ai, X.; Chen, Z.; Yang, H.; Cao, Y. Nano Energy. 2019, 56, 160–168. Liu, X.; Tang, L.; Li, Z.; Zhang, J.; Xu, Q.; Liu, H.; Wang, Y.; Xia, Y.; Cao, Y.; Ai, X. J. Mater. Chem. A 2019, 7, 18940–18949. Pu, X.; Wang, H.; Yuan, T.; Cao, S.; Liu, S.; Xu, L.; Yang, H.; Ai, X.; Chen, Z.; Cao, Y. Energy Storage Mater. 2019, 22, 330–336. Wu, X.; Zhong, G.; Yang, Y. J. Power Sources 2016, 327, 666–674. Ryu, S.; Wang, J. E.; Kim, J.-H.; Ruffo, R.; Jung, Y. H.; Kim, D. K. J. Power Sources 2019, 444, 227274. Park, Y.-U.; Seo, D.-H.; Kim, H.; Kim, J.; Lee, S.; Kim, B.; Kang, K. Adv. Funct. Mater. 2014, 24, 4603–4614.

Electrode materials for reversible sodium ions de/intercalation 173. 174. 175. 176. 177. 178. 179. 180. 181. 182. 183. 184. 185. 186. 187. 188. 189. 190. 191. 192. 193. 194. 195. 196. 197. 198. 199. 200. 201. 202. 203. 204. 205. 206. 207. 208. 209. 210. 211. 212. 213. 214. 215. 216. 217. 218. 219. 220. 221. 222. 223. 224. 225. 226. 227. 228. 229. 230. 231. 232. 233. 234. 235. 236. 237. 238. 239.

79

le Meins, J.; Courbion, G. J. Solid State Chem. 1999, 277, 260–277. Yakubovich, O. V.; Kireev, V. V.; Mel’nikov, O. K. Crystallogr. Rep. 2000, 45, 578–584. Bianchini, M.; Brisset, N.; Fauth, F.; Weill, F.; Elkaim, E.; Suard, E.; Masquelier, C.; Croguennec, L. Chem. Mater. 2014, 26 (14), 4238–4247. Bianchini, M.; Fauth, F.; Brisset, N.; Weill, F.; Suard, E.; Masquelier, C.; Croguennec, L. Chem. Mater. 2015, 27, 3009–3020. Yan, G.; Mariyappan, S.; Rousse, G.; Jacquet, Q.; Deschamps, M.; David, R.; Mirvaux, B.; Freeland, J. W.; Tarascon, J. M. Nat. Commun. 2019, 10, 585. Guo, J.-Z.; Wang, P.-F.; Wu, X.-L.; Zhang, X.-H.; Yan, Q.; Chen, H.; Zhang, J.-P.; Guo, Y.-G. Adv. Mater. 2017, 29, 1701968. Ellis, B. L.; Makahnouk, W. R. M.; Makimura, Y.; Toghill, K.; Nazar, L. F. Nat. Mater. 2007, 6, 749–753. Kabalov, Y. K.; Simonov, M. A.; Belov, N. V. Dokl. Akad. Nauk 1974, 215, 850–853. Ellis, B. L.; Michael Makahnouk, W. R.; Rowan-Weetaluktuk, W. N.; Ryan, D. H.; Nazar, L. F. Chem. Mater. 2010, 22, 1059–1070. Fedotov, S. S.; Kabanov, A. A.; Kabanova, N. A.; Blatov, V. A.; Zhugayevych, A.; Abakumov, A. M.; Khasanova, N. R.; Antipov, E. V. J. Phys. Chem. C 2017, 121, 3194–3202. Avdeev, M.; Ling, C. D.; Tan, T. T.; Li, S.; Oyama, G.; Yamada, A.; Barpanda, P. Inorg. Chem. 2014, 53, 682–684. Kawabe, Y.; Yabuuchi, N.; Kajiyama, M.; Fukuhara, N.; Inamasu, T.; Okuyama, R.; Nakai, I.; Komaba, S. Electrochem. Commun. 2011, 13, 1225–1228. Kubota, K.; Kazuki, Y.; Yabuuchi, N.; Komaba, S. Electrochemistry 2014, 82, 909–911. Yakubovich, O. V.; Karimova, O. V.; Mel’nikov, O. K. Acta Crystallogr. Sect. C Cryst. Struct. Commun. 1997, 53, 395–397. Kim, S. W.; Seo, D. H.; Kim, H.; Park, K. Y.; Kang, K. Phys. Chem. Chem. Phys. 2012, 14, 3299–3303. Barker, J.; Gover, R. K. B.; Burns, P.; Bryan, A.; Saidi, M. Y.; Swoyer, J. L. J. Power Sources 2005, 146, 516–520. Gover, R. K. B.; Burns, P.; Bryan, A.; Saidi, M. Y.; Swoyer, J. L.; Barker, J. Solid State Ionics 2006, 177, 2635–2638. Julien, C.; Mauger, A.; Vijh, A.; Zaghib, K. Lithium Batteries: Science and Technology, Springer, 2016; pp 269, 274–278. Barker, J.; Saidi, M. Y.; Swoyer, J. L. Electrochem. Solid-State Lett. 2003, 6, 4–7. Pizarro-Sanz, J. L.; Dance, J. M.; Villeneuve, G.; Arriortua-Marcaida, M. I. Mater. Lett. 1994, 18, 327–330. Zhao, J.; He, J.; Ding, X.; Zhou, J.; Ma, Y.; Wu, S.; Huang, R. J. Power Sources 2010, 195, 6854–6859. Lu, Y.; Zhang, S.; Li, Y.; Xue, L.; Xu, G.; Zhang, X. J. Power Sources 2014, 247, 770–777. Ruan, Y. L.; Wang, K.; Song, S. D.; Han, X.; Cheng, B. W. Electrochim. Acta 2015, 160, 330–336. Jin, T.; Liu, Y.; Li, Y.; Cao, K.; Wang, X.; Jiao, L. Adv. Energy Mater. 2017, 7 (15), 1700087. Law, M.; Balaya, P. Energy Storage Mater. 2018, 10, 102–113. Boivin, E.; Chotard, J. N.; Bamine, T.; Carlier, D.; Serras, P.; Palomares, V.; Rojo, T.; Iadecola, A.; Dupont, L.; Bourgeois, L.; Fauth, F.; Masquelier, C.; Croguennec, L. J. Mater. Chem. A 2017, 5, 25044–25055. Li, L.; Xu, Y.; Sun, X.; Chang, R.; Zhang, Y.; Zhang, X.; Li, J. Adv. Energy Mater. 2018, 8, 1801064. Kosova, N. V.; Semykina, D. O. Solid State Ionics 2019, 343, 115119. Semykina, D. O.; Yakovlev, I. V.; Lapina, O. B.; Kabanov, A. A.; Kosova, N. V. Phys. Chem. Chem. Phys. 2020, 22, 15876–15884. Goubard-Bretesché, N.; Kemnitz, E.; Pinna, N. Mater. Chem. Front. 2019, 3, 2164–2174. Fedotov, S. S.; Kabanova, N. A.; Kabanov, A. A.; Blatov, V. A.; Khasanova, N. R.; Antipov, E. V. Solid State Ionics 2018, 314, 129–140. Tripathi, R.; Ramesh, T. N.; Ellis, B. L.; Nazar, L. F. Angew. Chemie – Int. Ed. 2010, 49, 8738–8742. Tripathi, R.; Gardiner, G. R.; Islam, M. S.; Nazar, L. F. Chem. Mater. 2011, 23, 2278–2284. Masse, R.; Grenier, J.-C. Bull. Mine´ralogie 1971, 94, 437–439. Bierlein, J. D.; Arweiler, C. B. Appl. Phys. Lett. 1986, 49, 917–919. Sorokina, N. I.; Voronkova, V. I. Crystallogr. Reports 2007, 52, 80–93. Plechov, P. Y.; Belakovskyi, D. I.; Pautov, L. A.; Pekov, I. V.; Kasatkin, A. B.; Agahanov, A. A.; Karpenko, V. Y.; Necrylov, N. A.; Gritsenko, Y.; Garanin, B. K. Novye Dannye o Miner. 2020, 54, 74–95. Pekov, I. V.; Yapaskurt, V. O.; Britvin, S. N.; Zubkova, N. V.; Vigasina, M. F.; Sidorov, E. G. Mineral. Mag. 2016, 80, 639–646. Fedotov, S. S.; Samarin, A. S.; Antipov, E. V. J. Power Sources 2020, 480, 228840. Ge, G. W.; Luo, W.; Zhang, X. L.; Du, X. D.; Liang, L. L. Polyhedron. 2018, 152, 55–60. Furusawa, S.; Hayasi, H.; Ishibashi, Y.; Miyamoto, A.; Sasaki, T. J. Phys. Soc. Japan 1993, 62, 183–195. Furusawa, S.; Sugiyama, H.; Itoh, F.; Miyamoto, A.; Sasaki, T. J. Phys. Soc. Japan 2000, 69, 2087–2091. Park, J. H.; Kim, C. S.; Choi, B. C.; Moon, B. K.; Seo, H. J. Appl. Phys. A Mater. Sci. Process 2004, 78, 745–748. Guin, M.; Tietz, F. J. Power Sources 2015, 273, 1056–1064. Recham, N.; Rousse, G.; Sougrati, M. T.; Chotard, J. N.; Frayret, C.; Mariyappan, S.; Melot, B. C.; Jumas, J. C.; Tarascon, J. M. Chem. Mater. 2012, 24, 4363–4370. Fedotov, S. S.; Khasanova, N. R.; Samarin, A. S.; Drozhzhin, O. A.; Batuk, D.; Karakulina, O. M.; Hadermann, J.; Abakumov, A. M.; Antipov, E. V. Chem. Mater. 2016, 28, 411–415. Fedotov, S. S.; Samarin, A. S.; Nikitina, V. A.; Aksyonov, D. A.; Sokolov, S. A.; Zhugayevych, A.; Stevenson, K. J.; Khasanova, N. R.; Abakumov, A. M.; Antipov, E. V. J. Mater. Chem. A. 2018, 6, 14420–14430. Pinna, N.; Goubard-Bretesché, N. Adv. Energy Mater. 2021, 11, 2002971. Nikitina, V. A.; Kuzovchikov, S. M.; Fedotov, S. S.; Khasanova, N. R.; Abakumov, A. M.; Antipov, E. V. Electrochim. Acta 2017, 258, 814–824. Nikitina, V. A.; Fedotov, S. S.; Vassiliev, S. Y.; Samarin, A. S.; Khasanova, N. R.; Antipov, E. V. J. Electrochem. Soc. 2017, 164, A6373–A6380. Kim, H.; Ishado, Y.; Tian, Y.; Ceder, G. Adv. Funct. Mater. 2019, 29, 1–8. Ding, J.; Lin, Y. C.; Liu, J.; Rana, J.; Zhang, H.; Zhou, H.; Chu, I. H.; Wiaderek, K. M.; Omenya, F.; Chernova, N. A.; Chapman, K. W.; Piper, L. F. J.; Ong, S. P.; Whittingham, M. S. Adv. Energy Mater. 2018, 8, 1–8. Buyuker, I.; Hidalgo, M. F. V.; Zuba, M.; Diamint, J.; Zhou, H.; Chernova, N. A.; Piper, L. J.; Whittingham, M. S. J. Electrochem. Soc. 2021, 168, 050513. Mu, L.; Ben, L.; Hu, Y. S.; Li, H.; Chen, L.; Huang, X. J. Mater. Chem. A 2016, 4, 7141–7147. Galceran, M.; Rikarte, J.; Zarrabeitia, M.; Pujol, M. C.; Aguiló, M.; Casas-Cabanas, M. ACS Appl. Energy Mater. 2019, 2, 1923–1931. Bianchini, F.; Fjellvåg, H.; Vajeeston, P. Phys. Chem. Chem. Phys. 2017, 19, 14462–14470. Nalbandyan, V. B.; Zvereva, E. A.; Shukaev, I. L.; Gordon, E.; Politaev, V. V.; Whangbo, M. H.; Petrenko, A. A.; Denisov, R. S.; Markina, M. M.; Tzschoppe, M.; Bukhteev, K. Y.; Klingeler, R.; Vasiliev, A. N. Inorg. Chem. 2017, 56, 14023–14039. Renman, V.; Valvo, M.; Tai, C. W.; Gómez, C. P.; Edström, K.; Liivat, A. Sustain. Energy Fuels 2018, 2, 941–945. Otroshchenko, L. P.; Simonov, V. I.; Belov, N. V. Dokl. Akad. Nauk 1982, 265, 76–79. Panigrahi, A.; Nishimura, S. I.; Shimada, T.; Watanabe, E.; Zhao, W.; Oyama, G.; Yamada, A. Chem. Mater. 2017, 29, 4361–4366. Duncan, H.; Kondamreddy, A.; Mercier, P. H. J.; Le Page, Y.; Abu-Lebdeh, Y.; Couillard, M.; Whitfield, P. S.; Davidson, I. J. Chem. Mater. 2011, 23, 5446–5456. Guan, W. H.; Lin, Q. Y.; Lan, Z. Y.; Pan, W. L.; Wei, X.; Sun, W. P.; Zheng, R. T.; Lu, Y. H.; Shu, J.; Pan, H. G.; Yan, M.; Jiang, Y. Z. Mater. Today Nano 2020, 12, 100098. Treacher, J. C.; Wood, S. M.; Islam, M. S.; Kendrick, E. Phys. Chem. Chem. Phys. 2016, 18, 32744–32752. Kee, Y.; Dimov, N.; Staykov, A.; Okada, S. Mater. Chem. Phys. 2016, 171, 45–49. Chen, C. Y.; Matsumoto, K.; Nohira, T.; Hagiwara, R. Electrochem. commun. 2014, 45, 63–66. Trottier, J.; Brochu, F.; Rodrigues, I.; Zaghib, K.; Mauger, A.; Julien, C. M. ECS Meet. Abstr. 2011, 580. Khomyakov, A. P. Int. Geol. Rev. 1983, 25, 368–372.

80 240. 241. 242. 243. 244. 245. 246. 247. 248. 249. 250. 251. 252. 253. 254. 255. 256. 257. 258. 259. 260. 261. 262. 263. 264. 265. 266. 267. 268. 269. 270. 271. 272. 273. 274. 275. 276. 277. 278. 279. 280. 281. 282. 283. 284. 285. 286. 287. 288. 289. 290. 291. 292. 293. 294. 295. 296. 297. 298. 299. 300. 301. 302. 303. 304. 305. 306. 307.

Electrode materials for reversible sodium ions de/intercalation Chen, H.; Hao, Q.; Zivkovic, O.; Hautier, G.; Du, L.-S.; Tang, Y.; Hu, Y.-Y.; Ma, X.; Grey, C. P.; Ceder, G. Chem. Mater. 2013, 25, 2777–2786. Law, M.; Ramar, V.; Balaya, P. J. Power Sources 2017, 359, 277–284. Barpanda, P.; Lander, L.; Nishimura, S. I.; Yamada, A. Adv. Energy Mater. 2018, 8, 1–26. Stafeeva, V. S.; Drozhzhin, O. A.; Panin, R. V.; Filimonov, D. S.; Fabrichnyi, P. B.; Yashina, L. V.; Khasanova, N. R.; Antipov, E. V. Russ. J. Electrochem. 2015, 51, 619–626. Kurova, T. A.; Shumyatskaya, N. G.; Voronkov, A. A.; Pyatenko, Y. A. Dokl. Akad. Nauk 1980, 605–607. Krivovichev, S. V.; Chernyatieva, A. P.; Britvin, S. N.; Yakovenchuk, V. N.; Krivovichev, V. G. Geol. Ore Depos. 2013, 55, 669–675. Hassanzadeh, N.; Sadrnezhaad, S. K.; Chen, G. Electrochim. Acta. 2016, 208, 188–194. Kraft, A. Bull. Hist. Chem. 2008, 33, 61–67. Karyakin, A. A. Electroanalysis 2001, 13 (10), 813–819. Karyakin, A. A.; Gitelmacher, O. V.; Karyakina, E. E. Anal. Chem. 1995, 67, 2419–2423. Vipin, A. K.; Hu, B.; Fugetsu, B. J. Hazard. Mater. 2013, 258, 93–101. Paolella, A.; Faure, C.; Timoshevskii, V.; Marras, S.; Bertoni, G.; Guerfi, A.; Vijh, A.; Armand, M.; Zaghib, K. J. Mater. Chem. A 2017, 5, 18919–18932. Hurlbutt, K.; Wheeler, S.; Capone, I.; Pasta, M. Joule 2018, 2, 1950–1960. Buser, H. J.; Ludi, A.; Schwarzenbach, D.; Petter, W. Inorg. Chem. 1977, 16, 2704–2710. Wang, L.; Song, J.; Qiao, R.; Wray, L. A.; Hossain, M. A.; De Chuang, Y.; Yang, W.; Lu, Y.; Evans, D.; Lee, J. J.; Vail, S.; Zhao, X.; Nishijima, M.; Kakimoto, S.; Goodenough, J. B. J. Am. Chem. Soc. 2015, 137, 2548–2554. Ojwang, D. O.; Svensson, M.; Njel, C.; Mogensen, R.; Menon, A. S.; Ericsson, T.; Häggström, L.; Maibach, J.; Brant, W. R. ACS Appl. Mater. & Interfaces 2021, 13, 10054– 10063. Song, J.; Wang, L.; Lu, Y.; Liu, J.; Guo, B.; Xiao, P.; Lee, J.-J.; Yang, X.-Q.; Henkelman, G.; Goodenough, J. B. J. Am. Chem. Soc. 2015, 137, 2658–2664. Firouzi, A.; Qiao, R.; Motallebi, S.; Valencia, C. W.; Israel, H. S.; Fujimoto, M.; Wray, L. A.; De Chuang, Y.; Yang, W.; Wessells, C. D. Nat. Commun. 2018, 9 (1), 1–10. Bauer, A.; Song, J.; Vail, S.; Pan, W.; Barker, J.; Lu, Y. Adv. Energy Mater. 2018, 8, 1–13. Will, L.; Huang, E. Bloomberg. https://www.bloomberg.com/press-releases/2021-07-29/catl-unveils-its-latest-breakthrough-technology-by-releasing-its-first-generation-ofsodium-ion-batteries, 2021. Dou, X.; Hasa, I.; Saurel, D.; Vaalma, C.; Wu, L.; Buchholz, D.; Bresser, D.; Komaba, S.; Passerini, S. Mater. Today 2019, 23, 87–104. Kulova, T. L.; Skundin, A. M. Russ. J. Electrochem. 2020, 56, 3–19. Fang, S.; Bresser, D.; Passerini, S. Adv. Energy Mater. 2019, 10 (1), 1902485. Zhang, H.; Hasa, I.; Passerini, S. Adv. Energy Mater. 2018, 8, 1702582. Singh, G.; Acebedo, B.; Cabanas, M. C.; Shanmukaraj, D.; Armand, M. Electrochem. Soc. 2013, 37, 61–63. Kubota, K.; Komaba, J. Electrochem. Soc. 2015, 162, A2538–A2550. Yusoff, N. F. M.; Idris, N. H.; Din, M. F. M.; Majid, S. R.; Harun, N. A.; Rahman, M. M. Sci. Rep. 2020, 10, 1–10. Cho, M. K.; Jo, J. H.; Choi, J. U.; Kim, J.; Yashiro, H.; Yuan, S.; Shi, L.; Sun, Y. K.; Myung, S. T. Nano Energy 2017, 41, 687–696. Zeng, A.; Neto, V. F.; Gracio, J. J.; Fan, Q. H. Diam. Relat. Mater. 2014, 43, 12–22. Yoshino, A. Angew. Chemie – Int. Ed. 2012, 51, 5798–5800. Seehra, M. S.; Narang, V.; Geddam, U. K.; Stefaniak, A. B. Carbon 2017, 111, 380–385. Ge, P.; Fouletier, M. Solid State Ionics 1988, 30, 1172–1175. Lenchuk, O.; Adelhelm, P.; Mollenhauer, D. Phys. Chem. Chem. Phys. 2019, 21 (35), 19378–19390. Jache, B.; Adelhelm, P. ECS Meeting Abstracts 2014, 10169–10173. Goktas, M.; Bolli, C.; Buchheim, J.; Berg, E. J.; Novák, P.; Bonilla, F.; Rojo, T.; Komaba, S.; Kubota, K.; Adelhelm, P. ACS Appl. Mater. & Interfaces 2019, 11, 32844–32855. Hasa, I.; Dou, X.; Buchholz, D.; Shao-Horn, Y.; Hassoun, J.; Passerini, S.; Scrosati, B. J. Power Sources 2016, 310, 26–31. Doeff, M. M.; Ma, Y.; Visco, S. J.; De Jonghe, L. C. Electrochem. Soc. 1993, 140, 169–170. Jian, Z.; Bommier, C.; Luo, L.; Li, Z.; Wang, C.; Greaney, P. A.; Ji, X. Chem. Mater. 2017, 29 (5), 2314–2320. Zarrabeitia, M.; Chagas, L. G.; Kuenzel, M.; Gonzalo, E.; Passerini, S.; Miguel, A. Appl. Mater. Interfaces 2019, 11 (32), 28885–28893. Stevens, D. A.; Dahn, J. R. J. Electrochem. Soc. 2001, 148, A803. Anji Reddy, M.; Helen, M.; Groß, A.; Fichtner, M.; Euchner, H. ACS Energy Lett. 2018, 3, 2851–2857. Xiao, B.; Li, X. ChemSusChem 2019, 12 (1), 133–144. Smith, A. J.; MacDonald, M. J.; Ellis, L. D.; Obrovac, M. N.; Dahn, J. R. Carbon 2012, 50, 3717–3723. Thomas, P.; Billaud, D. Electrochim. Acta 2002, 47, 3303–3307. Simone, V.; Boulineau, A.; de Geyer, A.; Rouchon, D.; Simonin, L.; Martinet, S. J. Energy Chem. 2016, 25, 761–768. Bobyleva, Z. V.; Drozhzhin, O. A.; Dosaev, K. A.; Kamiyama, A.; Ryazantsev, S. V.; Komaba, S.; Antipov, E. V. Electrochim. Acta 2020, 354, 136647. Li, Y.; Hu, Y. S.; Titirici, M. M.; Chen, L.; Huang, X. Adv. Energy Mater. 2016, 6 (18), 1600659. Zheng, T.; Liu, Y.; Fuller, E. W.; Tseng, S.; von Sacken, U.; Dahn, J. R. J. Electrochem. Soc. 1995, 142, 2581–2590. Xing, W.; Xue, J. S.; Zheng, T.; Gibaud, A.; Dahn, J. R. J. Electrochem. Soc. 1996, 143, 3482–3491. Hasegawa, G.; Kanamori, K.; Kannari, N.; Ozaki, J. I.; Nakanishi, K.; Abe, T. J. Power Sources 2016, 318, 41–48. Frida, R.; Alvin, S.; Kim, J. J. Ind. Eng. Chem. 2020, 91, 317–329. Muñoz-Márquez, M.Á.; Saurel, D.; Gómez-Cámer, J. L.; Casas-Cabanas, M.; Castillo-Martínez, E.; Rojo, T. Adv. Energy Mater. 2017, 7, 1–31. Stevens, D. A.; Dahn, J. R. J. Electrochem. Soc. 2000, 147, 1271–1273. Franklin, R. Proc. R. Soc. London 1951, 209 (1097), 196–218. Stratford, J. M.; Allan, P. K.; Pecher, O.; Chater, A.; Grey, C. P. Chem. Comm. 2016, 52 (84), 12430–12433. Morikawa, Y.; Nishimura, S.; Hashimoto, R.; Ohnuma, M. Adv. Energy Mater. 2019, 10 (3), 1903176. Morita, R.; Gotoh, K.; Fukunishi, M.; Kubota, K.; Komaba, S.; Nishimura, N.; Yumura, T.; Deguchi, K.; Ohki, S.; Shimizu, T.; et al. J. Mater. Chem. A 2016, 4, 13183–13193. Cao, Y.; Xiao, L.; Sushko, M. L.; Wang, W.; Schwenzer, B.; Xiao, J.; Nie, Z.; Saraf, L. V.; Yang, Z.; Liu, J. Nano Lett. 2012, 12 (7), 3783–3787. Qiu, S.; Xiao, L.; Sushko, M. L.; Han, K. S.; Shao, Y.; Yan, M. Adv. Energy Mater. 2017, 7 (17), 1700403. Gomez-Martin, A.; Martinez-Fernandez, J.; Ruttert, M.; Winter, M.; Placke, T.; Ramirez-Rico, J. Chem. Mater. 2019, 31, 7288–7299. Bommier, C.; Surta, T. W.; Dolgos, M.; Ji, X. Nano Lett. 2015, 15, 5888–5892. Weaving, J. S.; Lim, A.; Millichamp, J.; Neville, T. P.; Ledwoch, D.; Kendrick, E.; Mcmillan, P. F.; Shearing, P. R.; Howard, C. A.; Brett, D. J. L. Appl. Energy Mater. 2020, 3 (8), 7474–7484. Wang, J.; Polleux, J.; Lim, J.; Dunn, B. J. Phys. Chem. C 2007, 111, 14925–14931. Brezesinski, T.; Wang, J.; Polleux, J.; Dunn, B.; Tolbert, S. H. J. Am. Chem. Soc. 2009, 131 (5), 1802–1809. Yamamoto, H.; Muratsubaki, S.; Kubota, K.; Fukunishi, M.; Watanabe, H.; Kim, J.; Komaba, S. J. Mater. Chem. A 2018, 6, 16844–16848. Li, Z.; Bommier, C.; Chong, Z. S.; Jian, Z.; Surta, T. W.; Wang, X.; Xing, Z.; Neuefeind, J. C.; Stickle, W. F.; Dolgos, M.; Greaney, P. A.; Ji, X. Adv. Energy Mater. 2017, 7, 1–10. Raccichini, R.; Varzi, A.; Passerini, S.; Scrosati, B. Nat. Mater. 2015, 14, 271–279. Chandula, K.; Li, H.; Xu, L.; Yan, C. J. Energy Chem. 2020, 42, 91–107.

Electrode materials for reversible sodium ions de/intercalation 308. 309. 310. 311. 312. 313. 314. 315. 316. 317. 318. 319. 320. 321. 322. 323. 324. 325. 326. 327. 328. 329. 330. 331. 332. 333. 334. 335. 336. 337. 338. 339. 340. 341. 342. 343. 344. 345. 346. 347. 348. 349. 350. 351. 352. 353. 354. 355. 356. 357. 358. 359. 360. 361. 362. 363. 364. 365. 366. 367. 368. 369. 370. 371. 372. 373. 374. 375. 376. 377. 378.

81

Kulova, T. L.; Skundin, A. M. Elektrohimiya. 2012, 48, 362. Natarajan, S.; Subramanyan, K.; Aravindan, V. Small 2019, 15, 1–14. Deschanvres, A.; Raveau, B.; Sekkal, Z. Mater. Res. Bull. 1971, 6, 699–704. Senguttuvan, P.; Rousse, G.; Seznec, V.; Tarascon, J. M.; Palacín, M. R. Chem. Mater. 2011, 23, 4109–4111. Leonidov, I. A.; Leonidova, O. N.; Perelyaeva, L. A.; Samigullina, R. F.; Kovyazina, S. A.; Patrakeev, M. V. Phys. Solid State 2003, 45, 2183–2188. Ohzuku, T.; Ueda, A.; Yamamoto, N. J. Electrochem. Soc. 1995, 142, 1431–1435. Sun, Y.; Zhao, L.; Pan, H.; Lu, X.; Gu, L.; Hu, Y. S.; Li, H.; Armand, M.; Ikuhara, Y.; Chen, L.; Huang, X. Nat. Commun. 2013, 4, 1810–1870. Petkovich, N. D.; Rudisill, S. G.; Wilson, B. E.; Mukherjee, A.; Stein, A. Inorg. Chem. 2014, 53, 1100–1112. Opra, D. P.; Gnedenkov, S. V.; Sinebryukhov, S. L. J. Power Sources 2019, 442, 227225. Wu, L.; Buchholz, D.; Bresser, D.; Gomes Chagas, L.; Passerini, S. J. Power Sources 2014, 251, 379–385. Myung, S. T.; Takahashi, N.; Komaba, S.; Yoon, C. S.; Sun, Y. K.; Amine, K.; Yashiro, H. Adv. Funct. Mater. 2011, 21, 3231–3241. Siebert, A.; Dou, X.; Garcia-Diez, R.; Buchholz, D.; Felix, R.; Handick, E.; Greco, G.; Hasa, I.; Wilks, R. G.; Passerini, S. ACS Appl. Energy Mater. 2020, 4 (1), 164–175. Baur, W. H.; Khan, A. A. Acta Crystallogr. Sect. B Struct. Crystallogr. Cryst. Chem. 1971, 27, 2133–2139. Weyl, R. Zeitschrift fu¨r Krist. Mater. 1959, 111, 401–420. Wu, L.; Bresser, D.; Buchholz, D.; Giffi, G.; Castro, C. R.; Ochel, A.; Passerini, S. Adv. Energy Mater. 2014, 5 (2), 1401142. Lee, J.; Lee, J. K.; Chung, K. Y.; Jung, H. G.; Kim, H.; Mun, J.; Choi, W. Electrochim. Acta 2016, 200, 21–28. Wang, W.; Liu, Y.; Wu, X.; Wang, J.; Fu, L.; Zhu, Y.; Wu, Y.; Liu, X. Adv. Mater. Technol. 2018, 3, 1–20. Andersson, S.; Wadsley, A. D. Acta Crystallogr. 1961, 14, 1245–1249. Andersson, S.; Wadsley, A. D. Acta Crystallogr. 1962, 15, 194–201. Batygin, V. G. Zhurnal Neorg. Khimii 1967, 12, 1442–1452. Avdeev, M.; Kholkin, A. Acta Crystallogr. Sect. C Cryst. Struct. Commun. 2000, 56, e539–e540. Yakubovich, O. V.; Kireev, V. V. Crystallogr. Reports. 2003, 48, 24–28. Nava-Avendaño, J.; Morales-García, A.; Ponrouch, A.; Rousse, G.; Frontera, C.; Senguttuvan, P.; Tarascon, J. M.; De Dompablo, M. E. A.; Palacín, M. R. J. Mater. Chem. A 2015, 3, 22280–22286. Muñoz-Márquez, M. A.; Zarrabeitia, M.; Castillo-Martínez, E.; Eguía-Barrio, A.; Rojo, T.; Casas-Cabanas, M. ACS Appl. Mater. & Interfaces 2015, 7, 7801–7808. Zarrabeitia, M.; Castillo-Martínez, E.; López Del Amo, J. M.; Eguía-Barrio, A.; Muñoz-Márquez, M. A.; Rojo, T.; Casas-Cabanas, M. J. Power Sources 2016, 324, 378–387. Kebede, M. A.; Ezema, F. I. Electrochemical Devices for Energy Storage Applications, CRC Press, 2019. Kulova, T. L.; Fateev, V. N.; Seregina, E. A.; Grigoriev, A. S. Int. J. Electrochem. Sci. 2020, 15, 7242–7259. Chevrier, V. L.; Ceder, G. J. Electrochem. Soc. 2011, 158, 1011–1014. Jung, S. C.; Jung, D. S.; Choi, J. W.; Han, Y. J. Phys.Chem. Lett. 2014, 5 (7), 1283–1288. Witte, J.; Schnering, H. G.; Klemm, W. ZAAC – J. Inorg. Gen. Chem. 1964, 327, 260–273. Morito, H.; Yamada, T.; Ikeda, T.; Yamane, H. J. Alloys Compd. 2009, 480, 723–726. Komaba, S.; Matsuura, Y.; Ishikawa, T.; Yabuuchi, N.; Murata, W.; Kuze, S. Electrochem. Commun. 2012, 21, 65–68. Ellis, L. D.; Wilkes, B. N.; Hatchard, T. D.; Obrovac, M. N. J. Electrochem. Soc. 2014, 161, 416–421. Legrain, F.; Malyi, O. I.; Manzhos, S. Comput. Mater. Sci. 2014, 94, 214–217. Xu, Y.; Swaans, E.; Basak, S.; Zandbergen, H. W.; Borsa, D. M.; Mulder, F. M. Adv. Energy Mater. 2016, 6 (2), 1501436. Buriak, J. M.; Sayed, S. Y.; Peter Kalisvaart, W.; Luber, E. J.; Olsen, B. C. ACS Appl. Energy Mater. 2020, 3, 9950–9962. Kalisvaart, W. P.; Olsen, B. C.; Luber, E. J.; Buriak, J. M. ACS Appl. Energy Mater. 2019, 2, 2205–2213. Lu, X.; Adkins, E. R.; He, Y.; Zhong, L.; Luo, L.; Mao, S. X.; Wang, C. M.; Korgel, B. A. Chem. Mater. 2016, 28, 1236–1242. Baggetto, L.; Keum, J. K.; Browning, J. F.; Veith, G. M. Electrochem. commun. 2013, 34, 41–44. Jow, T. R.; Shacklette, L. W.; Maxfield, M.; Vernick, D. Proc. Electrochem. Soc. 1987, 134 (7), 1730–1733. Yabuuchi, N.; Matsuura, Y.; Ishikawa, T.; Kuze, S.; Son, J.-Y.; Cui, Y.-T.; Oji, H.; Komaba, S. ChemElectroChem. 2014, 1 (3), 580–589. Jain, A.; Ong, S. P.; Hautier, G.; Chen, W.; Richards, W. D.; Dacek, S.; Cholia, S.; Gunter, D.; Skinner, D.; Ceder, G.; Persson, K. APL Mater. 2013, 1, 011002. Morachevskii, A. G. Russ. J. Appl. Chem. 2018, 91, 1785–1798. Ellis, L. D.; Hatchard, T. D.; Obrovac, M. N. J. Electrochem. Soc. 2012, 159, A1801–A1805. Baggetto, L.; Ganesh, P.; Meisner, R. P.; Unocic, R. R.; Jumas, J. C.; Bridges, C. A.; Veith, G. M. J. Power Sources 2013, 234, 48–59. Stratford, J. M.; Mayo, M.; Allan, P. K.; Pecher, O.; Borkiewicz, O. J.; Wiaderek, K. M.; Chapman, K. W.; Pickard, C. J.; Morris, A. J.; Grey, C. P. J. Am. Chem. Soc. 2017, 139, 7273–7286. Nam, D. H.; Kim, T. H.; Hong, K. S.; Kwon, H. S. ACS Nano. 2014, 8, 11824–11835. Li, Z.; Ding, J.; Mitlin, D. Acc. Chem. Res. 2015, 48, 1657–1665. Wiberg, E. Elsevier, 2001. DeWitt, T. W.; Skolnik, S. J. Am. Chem. Soc. 1946, 68, 2305–2309. Hittorf, W. Ann. Phys. 1865, 202, 193–228. Bachhuber, F.; Von Appen, J.; Dronskowski, R.; Schmidt, P.; Nilges, T.; Pfitzner, A.; Weihrich, R. Angew. Chemie – Int. Ed. 2014, 53, 11629–11633. Aykol, M.; Doak, J. W.; Wolverton, C. Phys. Rev. B 2017, 95 (21), 214115. Keyes, R. W. Phys. Rev. 1953, 92, 580–584. Sun, Y.; Wang, L.; Pei, A.; He, X.; Sun, Y.; Wang, L.; Li, Y.; Li, Y.; Lee, H. R.; Pei, A.; He, X. Joule 2019, 3 (4), 1080–1093. Hasa, I.; Mariyappan, S.; Saurel, D.; Adelhelm, P.; Koposov, A. Y.; Masquelier, C.; Croguennec, L.; Casas-Cabanas, M. J. Power Sources 2021, 482, 228872. Zhou, J.; Liu, X.; Cai, W.; Zhu, Y.; Liang, J.; Zhang, K.; Lan, Y.; Jiang, Z.; Wang, G.; Qian, Y. Adv. Mater. 2017, 29, 1–7. Köpf, M.; Eckstein, N.; Pfister, D.; Grotz, C.; Krüger, I.; Greiwe, M.; Hansen, T.; Kohlmann, H.; Nilges, T. J. Cryst. Growth 2014, 405, 6–10. Sangster, J. M. J. Phase Equilibria Diffus. 2010, 31, 62–67. Dong, Y.; DiSalvo, F. J. Acta Crystallogr. Sect. E Struct. Rep. Online 2005, 61, i223–i224. Mortazavi, M.; Ye, Q.; Birbilis, N.; Medhekar, N. V. J. Power Sources 2015, 285, 29–36. Liu, H.; Hu, K.; Yan, D.; Chen, R.; Zou, Y.; Liu, H.; Wang, S. Adv. Mater. 2018, 30, 1–22. Pang, J.; Bachmatiuk, A.; Yin, Y.; Trzebicka, B.; Zhao, L.; Fu, L.; Mendes, R. G.; Gemming, T.; Liu, Z.; Rummeli, M. H. Adv. Energy Mater. 2018, 8, 1–43. Kulova, T. L.; Kudryashova, Y. O.; Gryzlov, D. Y.; Skundin, A. M.; Andreev, V. N. Patent RU 2 732 988 C1, 2020. Morachevskii, A. G. Russ. J. Appl. Chem. 2019, 92, 321–331. Liu, Y.; Xu, J.; Kang, Z.; Wang, J. Thermochim. Acta 2013, 569, 119–126. Johnson, C. E.; Fischer, A. K. J. Less-Common Metals 1970, 20, 339–344. Lim, Y. R.; Shojaei, F.; Park, K.; Jung, C. S.; Park, J.; Cho, W. I.; Kang, H. S. Nanoscale 2018, 10, 7047–7057. Ramireddy, T.; Sharma, N.; Xing, T.; Leforestier, J.; Glushenkov, A. M. ACS Appl. Mater. Interfaces 2016, 8 (44), 30152–30164. Darwiche, A.; Marino, C.; Sougrati, M. T.; Fraisse, B.; Stievano, L.; Monconduit, L. J. Am. Chem. Soc. 2012, 134 (51), 20805–20811. Li, X.; Ni, J.; Savilov, S. V.; Li, L. Chem. Eur. J. 2018, 24, 13719–13727.

82

Electrode materials for reversible sodium ions de/intercalation

379. Sottmann, J.; Herrmann, M.; Vajeeston, P.; Hu, Y.; Ruud, A.; Drathen, C.; Emerich, H.; Fjellvåsg, H.; Wragg, D. S. Chem. Mater. 2016, 28, 2750–2756. 380. Kuze, S.; Kageura, J.; Matsumoto, S.; Nakayama, T.; Makidera, M.; Saka, M.; Yamaguchi, T.; Yamamoto, T.; Nakane, K. Development of a Sodium Ion Secondary Battery, Sumitomo Chemical Co. Ltd., 2013; pp 1–13. 381. Kageura, J.; Matsumoto, S.; Kuze, S.; Yamaguchi, T. 224th ECS Meet. Abstr, 2013. 382. Chayambuka, K.; Mulder, G.; Danilov, D. L.; Notten, P. H. L. Adv. Energy Mater. 2020, 10 (38), 2001310. 383. Liu, J.; Chang, D.; Whitfield, P.; Janssen, Y.; Yu, X.; Zhou, Y.; Bai, J.; Ko, J.; Nam, K.-W.; Wu, L. Chem. Mater. 2014, 26, 3295–3305. 384. Yao, W.; Sougrati, M.-T.; Hoang, K.; Hui, J.; Lightfoot, P.; Armstrong, A. R. Chem. Mater. 2017, 29, 2167–2172.

7.04

Electrode materials for K-ion batteries

Tomooki Hosaka, Kei Kubota, and Shinichi Komaba, Department of Applied Chemistry, Tokyo University of Science, Tokyo, Japan © 2023 Elsevier Ltd. All rights reserved.

7.04.1 7.04.2 7.04.2.1 7.04.2.1.1 7.04.2.1.2 7.04.2.1.3 7.04.2.1.4 7.04.2.2 7.04.2.2.1 7.04.2.2.2 7.04.2.2.3 7.04.2.2.4 7.04.2.2.5 7.04.2.3 7.04.2.3.1 7.04.2.3.2 7.04.2.3.3 7.04.3 7.04.3.1 7.04.3.1.1 7.04.3.1.2 7.04.3.1.3 7.04.3.2 7.04.3.2.1 7.04.3.2.2 7.04.3.2.3 7.04.3.3 7.04.3.3.1 7.04.3.3.2 7.04.3.4 7.04.3.4.1 7.04.3.4.2 7.04.3.4.3 7.04.4 References

Introduction to K-ion battery Positive electrode materials Layered oxides as positive electrode materials Classification of layered structures Stable structure types of layered AxMO2 Single transition metal oxides P2- and P3-type binary and ternary transition-metal systems Prussian blue analogues Prussian blue analogues as electrode materials Li, Na, and K insertion into Prussian blue analogues Prussian blue analogues for K-ion batteries Structural evolution during Kþ insertion Particle size and anion vacancy effect on the electrochemical performance Polyanionic compounds as positive electrode materials KTiOPO4-type structure materials KxMP2O7 (M ¼ Fe, Mn, and V) K3V2(PO4)3 and K3V2(PO4)2F3 Negative electrode materials Carbon materials K intercalation into graphite Electrochemical properties of graphite Hard and soft carbon K Alloys and other potassiatable compounds Alkali metal alloy materials and compounds for Li-, Na-, and K-ion batteries Group 14 elements and compounds Group 15 elements and compounds Transition metal oxides as negative electrode materials Ti, Mo, and Nb oxides Transition metal oxides based on conversion reaction Transition metal chalcogenides 3d transition metal dichalcogenides (TiS2 and VS2) 4d and 5d transition metal dichalcogenides (MoS2, MoSe2, and WS2) Metal sulfides based on conversion or conversion-alloying reactions Summary and perspective

84 87 87 87 87 89 91 93 93 94 95 95 97 99 100 102 103 104 104 104 106 108 110 110 112 113 115 115 117 117 118 120 120 120 122

Abstract Li-ion batteries (LIBs) have the highest energy density among practical secondary batteries and are widely utilized in electronics, electric vehicles, and even stationary energy storage systems. Along with expanding their demand and application, concern about Li and Co resources is growing. Therefore, high-performance secondary batteries composed of earth-abundant elements are desired to complement LIBs. Recently, K-ion batteries (KIBs) have attracted much attention as potential alternatives to LIBs. Previous studies have developed positive and negative electrode materials for KIBs and demonstrated several unique advantages of KIBs over LIBs and Na-ion batteries (NIBs). Besides being free from any scarce/toxic elements, the low standard electrode potentials of Kþ/K electrodes lead to high operation voltages competitive to those observed in LIBs. This chapter provides an overview of the studies on electrochemical materials for KIBs and discussions on recent achievements and remaining/emerging issues. The review also includes insights into electrode reactions and solid-state ionics.

Comprehensive Inorganic Chemistry III, Volume 7

https://doi.org/10.1016/B978-0-12-823144-9.00053-4

83

84

7.04.1

Electrode materials for K-ion batteries

Introduction to K-ion battery

Li-ion batteries (LIBs), commercialized in 1991, have the highest energy density of all practical rechargeable batteries and are widely used in electronic devices, electric vehicles, and even stationary energy storage systems. With the expansion of its demand and applications, there is a growing concern about the resources of Li and Co. Therefore, rechargeable batteries made of earth-abundant elements, such as Na-ion batteries (NIBs) and Mg batteries, are extensively studied as complements to the LIBs. Recently, K-ion batteries (KIBs) have been attracting attention due to their unique advantages. Fig. 1A shows the schematic configuration of a typical KIB, with a current collector consisting of a positive electrode (e.g., Prussian blue analog (PBA)) and an anode (e.g., graphite), and an electrolyte (e.g., KPF6/ethylene carbonate (EC) :diethyl carbonate (DEC)) between the electrodes KIBs are rocking-chair batteries similar to LIBs and NIBs. During charging, Kþ ions act as ionic charge carriers and are inserted from the active material of the positive electrode into the active material of the negative electrode via the electrolyte. During discharge, the inserted Kþ ions are inserted from the active material of the positive electrode into the active material of the negative electrode via the electrolyte. During discharge, the Kþ ions inserted in the negative electrode are extracted and reinserted into the positive electrode. Thus, during the charging process, electrons pass through the external circuit from the positive electrode to the negative electrode based on the redox of the respective active materials of the positive and negative electrodes. On the other hand, electrons reversibly pass through the external circuit from the anode to the cathode during the discharge of the battery, based on the reversible oxidation and reduction of the respective anode and cathode materials. In LIBs, Li-containing and Li-free materials are used for the positive and negative electrodes, respectively. Therefore, similarly, in KIB, K-containing and K-free materials are used as the active materials for the positive and negative electrodes, respectively. These positive and negative electrodes are sandwiched between insulating separators immersed in a nonaqueous electrolyte solution or, in the case of allsolid-state KIBs, solid electrolytes without separators. It is also worth mentioning that K thermodynamically forms no AleK intermetallic compounds, indicating that Al foil can also be applied as a negative-electrode current-collector in such batteries.2 According to the NIB cost analysis results reported by Vaalma et al., replacing the current collectors from Cu to Al reduces the battery costs; moreover, the reduction rate is higher than that observed when replacing Li with Na.3 Although the cell composition and operating mechanism of KIBs are very similar to those of LIBs and NIBs, different carrier ions have a significant impact on battery properties and material design.

Fig. 1 (A) Schematic illustration of the cell configuration and operational mechanism of a typical K-ion battery and (B) elemental abundance in the Earth’s crust. Data derived from Ref 1.

Electrode materials for K-ion batteries

85

Fig. 1B shows the abundance of elements in the earth’s crust.1 Li, Co, Cu, and Ni, which are widely used in high-energy LIBs, are present in the crust at less than 20 ppm (0.002%), while Al, Fe, Ca, Na, K, and Mg are present at more than 1%. Therefore, unlike LIBs, K-based batteries have significant advantages for large-scale energy storage applications in terms of the abundance of elements in the battery material. In particular, in addition to the abundance of K, Kþ ions are expected to provide higher voltage operation and better rate performance than LIBs. Table 1 summarizes the characteristics of carrier ions for rechargeable batteries. First, based on Faraday’s law of electrolysis, the lower the mass-to-charge ratio of the carrier ions, the higher the weight capacity; the mass-to-charge ratio of Liþ ions is as low as 6.94, while the mass-to-charge ratios of multivalent Mg2þ, Al3þ, and Ca2þ ions are as low as 12.16, 8.86, and 20.04, respectively. On the other hand, Naþ and Kþ ions show high mass-to-charge ratios of 23.00 and 39.10, which are 3.3 and 5.6 times higher than those of Liþ ions. However, since the carrier ions are mainly contained in the cathode, it is necessary to compare the formula masses of the cathode materials and consider their specific capacities. For example, the theoretical gravimetric capacity of KCoO2 is 206 mAh g 1, assuming that one alkali metal ion per compositional formula is extracted/inserted from ACoO2 (A ¼ Li, Na, K) during charge/discharge. This is lower than the value of 274 mAh g 1 observed for LiCoO2. According to this estimate, K (de) inserted materials should exhibit as much as 75–80% of the gravimetric capacity of their corresponding Li counterparts, since the transition metal oxide framework is much heavier than these alkali metal carrier ions. Although the theoretical volume capacity of KCoO2 is lower than that of Li and Na, KCoO2 retains 65–70% of the theoretical volume capacity of LiCoO2 due to its lower density. Secondly, the operating voltage, which is the potential difference between the positive and negative electrodes, is another important factor determining the energy density of a battery. In order to achieve a high energy density in a full cell, the high and low operating potential is desired for the positive and negative electrodes, respectively. In the rocking chair-type batteries, the lower limit of available potential is determined by the standard electrode potential of the carrier ions. LIBs have been recognized to offer the highest operation voltage owing to the lowest standard electrode potential of Liþ/Li ( 3.04 V vs. SHE) in aqueous electrolytes (Table 1). The standard electrode potential is highly dependent on the solvation state of the metal ion and differs between aprotic electrolyte solutions and aqueous solutions. Potassium has a lower standard electrode potential than Liþ/Li in carbonate solvents such as propylene carbonate (PC) and ethylene carbonate (EC):diethyl carbonate (DEC).2,9 A previous study reported that the standard electrode potential of Kþ/K in PC, calculated from the standard molar Gibbs energy and entropy, was  0.09 V vs. Liþ/Li.4 The lower E0(Kþ/K) than E0(Liþ/Li) implies that the negative electrode for KIBs can operate at a lower potential cutoff than that for LIBs, allowing higher voltage KIBs. To take advantage of the lower-potential limitation, it is essential to select an appropriate negative electrode showing low operation potential without K metal plating. Therefore, it is necessary to develop battery components such as high potential cathode materials, low potential anode materials, electrolytes, binders, and current collectors in order to achieve practical high voltage battery performance. In addition to high voltage operation, high power is also expected due to the fast diffusion of Kþ ions. Among the ions listed in Table 1, the Kþ ion has the largest ionic radius, but its Stokes radius in aqueous and PC solutions is the smallest among these ions; the small Stokes radius of the Kþ ion can be attributed to the weak interaction between the Kþ ion and the solvent molecules due to the low surface charge density of the solvent molecules. As a result, the potassium ion has the fastest diffusion rate in the electrolyte solution and the highest molar conductivity (Table 1). Therefore, it can be seen that if the degree of dissociation and anion mobility are the same, the Kþ ion provides higher cationic conductivity and transport number, which is advantageous for high power battery operation. It is also known that the desolvation process at the electrode/electrolyte interface is one of the rate-determining processes in lithium-ion batteries composed of graphite negative electrodes.10 DFT calculations by Ohkoshi et al. revealed that the desolvation energy of Kþ ions for PC is much lower than that of Liþ, Naþ, and Mg2þ ions (Table 1).8 Thus, in terms of cationic conductivity and desolvation energy, KIB is expected to have better rate capability than LIB and NIB. On the other hand, the diffusion of Kþ ions at the electrode/electrolyte interface and in the electrode material, as well as the kinetics of structural changes during K extraction and

Table 1

Comparison of the physical properties of Liþ, Naþ, Kþ, and other candidates as charge carriers for rechargeable batteries.

Relative atomic mass Mass-to-charge ratio Theoretical gravimetric capacity of ACoO2a/mAh g 1 Theoretical volumetric capacity of ACoO2a/mAh cm 3 E0 (Aþ/Aaq.)/V vs. SHE E0 (Aþ/APC)/V vs. Liþ/LiPC4 Shannon’s ionic radii/Å5 Stokes radii in water/Å6 Stokes radii in PC/Å7 Limiting molar ionic conductivity in PC/S cm2 mol 17 Desolvation energy in PC/kJ mol 18 Melting point/ C a

Liþ

Naþ

Kþ

Mg2þ

Al3þ

Ca2þ

6.94 6.94 274 1378  3.04 0 0.76 2.38 4.8 8.3 215.8 180.5

23.00 23.00 235 1193  2.71 0.23 1.02 1.84 4.6 9.1 158.2 97.8

39.10 39.10 206 906 2.93 0.09 1.38 1.25 3.6 15.2 119.2 63.4

24.31 12.16 260 – 2.4 – 0.72 3.47 – – 569.4 650

26.58 8.86 268 – 1.7 – 0.54 4.39 – – – 660

40.08 20.04 242 – 2.9 – 1.06 3.10 – – – 842

Theoretical capacities of ACoO2 with the multivalent ions are estimated based on Mg0.5CoO2, Al1/3CoO2, and Ca0.5CoO2.

86

Electrode materials for K-ion batteries

insertion, can also be rate-limiting processes for KIBs. Therefore, to realize high-power KIBs, it is crucial to demonstrate a lowresistance surface layer and develop electrode materials with high rate capability. Owing to the unique characteristics of Kþ ions, such as elemental abundance, no formation of AleK intermetallic compounds, a lower standard electrode potential than that of Liþ/Li, and fast diffusion in the electrolyte, KIBs have attracted attention to realize high voltage and power-density performances. Since 2015, many K insertion materials have been evaluated in K cells. Fig. 2A and B respectively summarize the electrochemical performances of the positive and negative electrode materials, according to the results of these publications, as capacity versus average operating potential plots. The reported cathode materials (Fig. 2A) can be classified into several material groups: layered transition metal oxides, PBAs, polyanionic compounds, and organic materials. Since layered transition metal oxides are the most promising cathode materials for LIBs and NIBs, K-containing layered transition metal oxides are also of scientific interest for KIB applications. In contrast to Li- and Na-based materials, typical K-containing layered transition metal oxides, such as KxCoO2, exhibit small reversible capacity, low operating potential, and stepwise voltage fluctuations with multiphase transitions due to strong Kþ-Kþ repulsion and stabilizing Kþ/vacancy ordering.12–14 Some PBAs exhibit high capacity and operating potential because of their three-dimensional

Fig. 2 Materials overview for selected (A) positive and (B) negative electrode materials for K-ion batteries. The values are based on the potassiation (positive electrodes) and depotassiation (negative electrodes) results for practical cells reported in the literature. The average voltage was calculated from the respective potassiation and depotassiation curves at low current density using digitalization and graph analysis software. The energy densities of the full cells were calculated assuming graphite (279 mAh g 1 at 0.25 V vs. Kþ/K9) as the negative electrode material against the positive electrode materials and K 2Mn[Fe(CN)6] (155 mAh g 1 at 3.8 V vs. Kþ/K11) as the positive electrode material against the negative electrode materials.

Electrode materials for K-ion batteries

87

(3D) open-framework consisting of channels and interstitial sites suitable for the diffusion and insertion of large Kþ ions.11,15–17 Like PBA, polyanionic compounds with 3D open frameworks are potential cathode materials and tend to exhibit high operating potentials.18–21 Some organic electrodes also exhibit redox activity in K-cells, providing a large reversible capacity.22–24 As shown in Fig. 2B, the anode materials for KIBs can be classified into carbonaceous materials, K-alloy materials, oxides, chalcogenides, polyanionic compounds, and organic materials. Graphite materials, which are widely utilized in LIBs and technically developed to enhance their performance, are promising candidates as negative electrode materials for KIBs, owing to their low operating potential and sufficiently large capacity.2,25,26 Some non-graphitic carbon materials also exhibit larger capacity and better rate performance than graphite.25,27 Materials showing alloying or conversion reactions provide higher gravimetric and volumetric capacities in K-cells than carbon-based materials.28 Future KIBs will include such high-capacity materials. Transition metal oxides,29–31 chalcogenides,32–34 and polyanionic compounds35,36 are expected to be used as safe anodes with less risk of Kmetal precipitation because of their relatively high operating potential. Besides, some organic materials have a high energy density comparable to carbonaceous materials.37–39

7.04.2

Positive electrode materials

7.04.2.1

Layered oxides as positive electrode materials

Layered AMO2 (A ¼ alkali metal, M ¼ 3d transition metal) materials have been extensively studied as cathode materials for Li and Na ion batteries since the electrochemical Li and Na intercalation behavior of layered LiCoO2 and NaCoO2 was first reported in 1980.40–42 In the potassium system, early contributions to the synthesis and structural characterization of KxMO2 have been reported by Hoppe, Fouassier, Delmas, and Hagenmuller et al. in the 1960s–1970s.43,44 Since 2016, the electrochemical K intercalation reactions into the materials in aprotic K cells have been extensively studied. Here, we review the layered KxMO2 material, including early structural studies.

7.04.2.1.1

Classification of layered structures

7.04.2.1.2

Stable structure types of layered AxMO2

LiCoO2, LiNi0.8Co0.15Al0.05O2, and LiNi1/3Mn1/3Co1/3O2, which are isostructural to a-NaFeO2, have been commercialized as cathode materials for lithium-ion batteries. These materials have a rock-salt type layered structure with an R-3m space group (SG) known as the a-NaFeO2 type. The a-NaFeO2-type structure consists of MO2 slabs consisting of MO6 octahedra with shared ends and alkali metals between the slabs. When the MO2 slabs are stacked in different stacking arrangements in the c-axis direction, polymorphism emerges. A systematic notation for layered transition metal oxides containing alkali metals was proposed by Delmas et al.45,46 Fig. 3A shows a schematic diagram of a typical layered structure of AxMO2. In the O3-type (a-NaFeO2-type) structure, the MO2 slabs are stacked along the c-axis with cubic close-packed (ccp) oxygen in the AB CA BC array, and the alkali metal ions are accommodated in octahedral sites in the interslab spaces. Each hexagonal unit cell is composed of three MO2 slabs. Specifically, the “O” in the O3 type represents the octahedral sites where the alkali metal ions are accommodated, and the subsequent “3” is the number of MO2 slabs in the hexagonal unit cell. Large alkali metal ions, such as Kþ and Rbþ, generally prefer prismatic sites with larger spaces than octahedral sites, thus forming P-type layered oxides such as P3-type and P2-type. In the P3-type layered structure, alkali-metal ions occupy the prismatic sites in the interslab space between the MO2 slabs with the oxygen packing array AB BC CA along the c-axis. There are three MO2 slabs in the hexagonal unit cell with an R3m or R-3m SG. The structures of most of the O3-type NaMO2 electrodes reversibly transform into P3-type NaxMO2 electrodes during electrochemical Na extraction on charging.47 When the unit cell is distorted from the hexagonal symmetry, a prime symbol is added between the letter and the number; however, the number of MO2 slabs is counted by assuming a pseudo-hexagonal unit cell. For instance, O3type KCrO2 transforms into O’3-, P3-, and P’3-type KxCrO2 by electrochemical K extraction.48 During K extraction, the P3 type structure is reversibly formed by the gliding of the MO2 slabs from the O3 type structure without CreO bond breakage and thus, a P3-type phase is generally alkali-metal deficient. Indeed, P’3-type K0.8CrO2 can be synthesized as a K-deficient phase via a solid-state reaction.49 Compared to the P3-type phases, the P2-type phase is generally synthesized at higher temperatures.12,50 Thus, P2-type KxCoO2 is produced by calcination at 600  C, a higher temperature than the 400  C required for the synthesis of P3-KxCoO2.12 As the phase transition from P3 to P2 type phase has to be accompanied by CoeO bond breakage, the P3 to P2 phase transition does not occur during electrochemical extraction of alkali-metal ions at room temperature. In the P2-type structure, alkali-metal ions occupy the prismatic sites in the interslab space between the MO2 slabs, with the oxygen packing array AB BA AB along the c-axis. There are two MO2 slabs in the hexagonal unit cell with a P63/mmc SG.

The structural types of layered AxMO2 reported in literature are summarized in Table 2. It should be noted that the materials synthesized by a solid-state reaction are listed, and the distorted O’3-, P’3-, and P’2-type materials are included together with the undistorted types. The Li system only comprised O3-type materials synthesized with M ¼ V, Cr, Co, and Ni in LiMO2. O’3-LiMnO251,52 and O3LiFeO251,53 were synthesized by Liþ/Naþ exchange and hydrothermal reactions from O’3-NaMnO2 and O3eNaFeO2, respectively; however, they could not be synthesized directly from a solid-state reaction. On the other hand, the synthesis of O3-, P3-, and P2type NaxMO2 and KxMO2 have been reported. O3-type NaMO2 compounds with M ¼ Sc, Ti, V, Cr, Mn, Fe, Co, and Ni has been synthesized, whereas the only reported O3 types are KScO243 and KCrO2.54 The difference in the structural stability of AMO2 was explained

88

Electrode materials for K-ion batteries

Fig. 3 (A) Schematic illustrations of the crystal structures of O3-, P3-, and P2-type AxMO2. (B) Plots of the distance between the nearest-neighbor Kþ ions vs. ionic radius of the octahedral M3 þ ions in stoichiometric KMO2. The open and filled squares represent unstable (hypothetic) and stable (reported in Inorganic Crystal Structure Database (ICSD)) compounds with layered structures, respectively, while the filled circles are the stable compounds with nonlayered structures. (C) Site preference energies of Cr, Mn, Fe, Co, and Ni in the KMO2 composition. Figures (B, C) are reproduced from Ref. Kim, H.; Seo, D. H.; Urban, A.; Lee, J.; Kwon, D. H.; Bo, S. H.; Shi, T.; Papp, J. K.; McCloskey, B. D.; Ceder, G., Stoichiometric Layered Potassium Transition Metal Oxide for Rechargeable Potassium Batteries. Chem. Mater. 2018, 30(18), 6532–6539 with permission. Copyright, 2018 American Chemical Society.

via the correlation between the ionic radii of the alkali (RA þ) and transition (RM3 þ) metals53,55–57 and that between RA þ and the ionicity ratio of the A-O and M-O bonds.55,58 The ionic radius of Liþ in octahedral coordination (0.76 Å) is similar to those of the transition-metal ions5 (Table 1). Thus, lithium ions are often mixed with transition metal ions to form cation-disordered rocksalt-type (NaCl type) phases such as a-LiFeO2.59 On the other hand, the Naþ (1.02 Å) and Kþ (1.38 Å) ionic radii are larger than those of Liþ and the transition-metal ions and thus, no site-exchange occurs between the alkali- and transition-metal ions. O3-type structures are crystallized at an almost fully alkali-metal-inserted state, and electrostatically repulsive interaction between the alkali-metal ions is significant for the large ionic radius of the Kþ ion. In 2018, Kim et al. plotted the distance between the nearest-neighbor Kþ ions vs. ionic radii of M3 þ ions (Fig. 3B).48 Large M3 þ cations (e.g. Sc3þ, In3þ, Er3þ, Tl3þ, Y3þ, Pr3þ, and La3þ) can retain long Kþ-Kþ distances and stabilize the O3-type structures. In contrast, small M3 þ cations (e.g. Ti3þ, V3þ, Mn3þ, Fe3þ, Co3þ, and Ni3þ) cannot stabilize O3-type structures and display short Kþ-Kþ distances. Kim et al. also calculated the site preference energies of Cr, Mn, Fe, Co, and Ni in the stoichiometric KMO2 composition (Fig. 3C).48 They reported that small cations theoretically prefer to occupy tetrahedral and pyramidal sites as nonlayered structures of KMnO2,60 KFeO2,61 and KCoO262,63 Table 2

Structural types of layered AxMO2 (A ¼ alkali-metal, M ¼ transition metal) compounds reported in the literature as a single transition-metal system.

3d M

21Sc

22Ti

23V

24Cr

25Mn

26Fe

27Co

28Ni

29Cu

LixMO2 NaxMO2

– O3

– O3

O3 O3

– O3 P2

– O3

– –

O3

–

O3 P3

P3 P2

–

O3 O3 P3 P2 P3 P2

O3 O3

KxMO2

O3 O3 P3 P2 –

–

–

Electrode materials for K-ion batteries

89

are experimentally synthesized in stoichiometric compositions. Instead, P3- and P2-type layered KxMO2 compounds are crystallized in nonstoichiometric compositions (Table 2).44 Notably, except for KScO2, the electrochemical potassium (de)intercalation properties of all the layered KxMO2 compounds listed in Table 2 have been examined after 2017.

7.04.2.1.3

Single transition metal oxides

7.04.2.1.3.1 O3-type KxMO2 (M ¼ Sc and Cr) O3-type KScO2 was first reported by Hoppe in 1965.43 Since Sc3þ ions have the electron configuration [Ar]4s03d0, O3-type KScO2 should be electrochemically inactive as O3eNaeScO2 delivers a negligibly small capacity in a Na cell.57 The only other reported O3type KMO2 oxide is O3eKCrO2. Chromium is the only 3d transition metal that crystalizes into O3-type structures with Li, Na, and K as LiCrO2, NaCrO2, and KCrO2, respectively (Table 2). Galvanostatic charge/discharge curves of the three chromium oxides are compared in Fig. 4A–C. As reported by our group, O3eLiCrO2 delivered a negligibly small capacity in the Li cell (Fig. 4A), while O3eNaCrO2 delivered a reversible capacity of  110 mAh g 1 with a specific voltage plateau at  3 V vs. Na in a Na cell (Fig. 4B).64 In a K cell, O3eKCrO2 delivered a reversible capacity of  90 mAh g 1 with stepwise voltage variation in the range 4.0–1.5 V (Fig. 4C).48 Such stepwise voltage curves are commonly observed for layered NaMO2 in single 3d transition-metal systems57; however, more pronounced voltage steps were observed for O3eKCrO2 than for O3eNaCrO2. This stepwise voltage variation originates from phase transitions (e.g. O3 into P3, O3 into O’3, and O’3 into P’3) and discretely stabilizes structures as a result of charge ordering in the transition-metal slabs65 and ordering of the alkali-metal ions and vacancies in the interslab space.66–70 7.04.2.1.3.2 P’2- and P3-type KxMnO2 As mentioned above, the only reported O3-type KMO2 (M ¼ 3d transition metal) compounds are KScO2 and KCrO2. In contrast, nonstoichiometric KxMnO2 oxides crystalize into P3- and P2-type phases (Table 2). P3- and P2-type polymorphs are recognized as the respective low- and high-temperature phases in NaxMO2 (x  2/3).71–73 In addition, the K content x significantly influences the structural type of KxMO2 obtained via the solid-state reaction, even at the same synthetic temperature. The electrochemical K (de)intercalation properties of layered KxMnO2 was first reported by Vaalma et al. in 2016.74 According to the synthetic procedure of Kim et al.75 and Gaillot et al.,76 P’2-type K0.3MnO2 having an orthorhombic lattice with the SG Cmcm was synthesized, and its electrochemical K (de)intercalation properties were examined in K cells.74 P’2-type K0.3MnO2 delivered a reversible capacity of  70 mAh g 1 in the potential range 1.5–3.5 V vs. Kþ/K and  130 mAh g 1 in 1.5–4.5 V vs. Kþ/K in three-electrode K-half cells (Fig. 5A). On the other hand, higher upper cut-off potentials resulted in larger capacities, and the capacity retention severely deteriorated. The restricted potential range 1.5–3.4 V vs. Kþ/K led to better capacity retention (Fig. 5B). Similar to P’2-K0.3MnO2, the electrochemical K (de)intercalation properties of P3-type KxMnO2 were examined by Kim et al. in 2017,77 while Liu et al. compared those of P’2-K0.3MnO2 and P3-K0.45MnO2.78 The as-prepared P3-type K0.5MnO2, having a rhombohedral lattice with the SG R3m, comprised more potassium atoms than those of the as-prepared P’2-K0.3MnO2 and delivered a reversible capacity of  100 mAh g 1 in a K-metal half-cell in the voltage range 1.5–3.9 V vs. K (Fig. 5C).77 P3-K0.5MnO2 exhibited different capacity retention behaviors, depending on the upper cut-off voltages, as observed for P’2-K0.3MnO2. Here, acceptable cycle stability was obtained in the voltage range 1.5–3.9 V vs. K (Fig. 5D).78 P’2-K0.3MnO2 and P3-K0.3MnO2 displayed similar charge/discharge profiles with stepwise curves. Kim et al. performed in situ XRD measurements for P’3-K0.5MnO2, proving that structural changes from the P3- to the O3-type phase occurred via a two-phase reaction for 0.425 > x > 0.395 in KxMnO2 and from the O3- to O3-like-phase via a two-phase reaction for 0.364 > x > 0.316 in the K extraction process.77 In the reverse process of K intercalation during discharge, reverse structural changes were observed. This suggested a reversible K insertion/extraction into/from P3-KxMnO2 in the voltage range 1.5–3.9 V vs. K. When the in situ XRD measurements were performed in the voltage range 1.5–4.2 V vs. K, weakened intensities of the 00 l diffraction peaks were observed at the end of charging to 4.2 V with less recovery of the peak intensities during the subsequent discharging process.77 This suggested that stacking faults and defects were irreversibly introduced in the structure during K extraction at x > 0.27 in KxMnO2, resulting in less

Fig. 4 Typical galvanostatic charge and discharge curves of the (A) Li||O3-LiCrO2 cell at 20 mA g 1 between 3.0 and 4.5 V, (B) Na||O3-NaCrO2 cell at 25 mA g 1 between 2.0 and 3.6 V, and (C) K ||O3-KCrO2 cell at 5 mA g 1 between 1.5 and 4.0 V. (D) In situ XRD patterns of O3-KCrO2 between 5 and 25 . (C, D) Reproduced from Ref. Kim, H.; Seo, D. H.; Urban, A.; Lee, J.; Kwon, D. H.; Bo, S. H.; Shi, T.; Papp, J. K.; McCloskey, B. D.; Ceder, G., Stoichiometric Layered Potassium Transition Metal Oxide for Rechargeable Potassium Batteries. Chem. Mater. 2018, 30(18), 6532–6539 with permission. Copyright, 2018 American Chemical Society.

90

Electrode materials for K-ion batteries

Fig. 5 (A) Galvanostatic charge and discharge curves and (B) cycle performance of the K||P’2-K0.3MnO2 cell in different potential ranges at 0.1C (27.9 mA g 1). (C) Galvanostatic voltage–capacity profiles of P3-type K0.5MnO2 in the voltage range 3.9–1.5 V at 5 mA g 1. (D) Discharge capacity of P3-type K0.5MnO2 over 20 cycles with the two different voltage cutoffs at the current rate 5 mA g 1. (A, B) Reproduced from Ref. Vaalma, C.; Giffin, G. A.; Buchholz, D.; Passerini, S., Non-Aqueous K-Ion Battery Based on Layered K0.3MnO2 and Hard Carbon/Carbon Black. J. Electrochem. Soc. 2016, 163(7), A1295-A1299 with permission. Copyright, 2016 Electrochemical Society. (C, D) Reproduced from Ref. Kim, H.; Seo, D. H.; Kim, J. C.; Bo, S. H.; Liu, L.; Shi, T.; Ceder, G., Investigation of Potassium Storage in Layered P3-Type K0.5MnO2 Cathode. Adv. Mater. 2017, 29(37), 1702480 with permission. Copyright, 2017 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

stepwise voltage curves and capacity fading during cycles. Thus, the voltage range should be optimized to enhance the cycle stability of P’3-KxMnO2 in K batteries. 7.04.2.1.3.3 P’2- and P3-type KxCoO2 Similar to KxMnO2, KxCoO2 compounds are known to crystallize into P’2- and P3- type phases63,79 (Table 2). These compounds have been studied since the 1970s, but their electrochemical K (de)intercalation properties have not been examined until late 2010s. In 2017, K (de)intercalation properties were reported.12,13 Since layered cobalt oxides are vital materials for the application of Liand Na-ion batteries, the charge/discharge behavior of AxCoO2 compounds was compared in Li, Na, and K batteries (Fig. 6A–C). When the charge/discharge behaviors of the ACoO2 polymorphs in the Li, Na, and K cells were compared, much clearer tendencies were observed: layered AxCoO2 containing a lighter alkali-metal displayed (i) wider compositional ranges of alkali-metal extraction/insertion, (ii) higher average voltages of the A ||AxCoO2 cell in the voltage plots as a function of x in AxCoO2, (iii) smoother voltage curves with less stepwise variation, and (iv) less inclined voltage slopes. Notably, the variation in the interslab distances of K1-xCoO2 due to the extraction of the large Kþ ions is almost comparable to that of Na1-xCoO2 (Fig. 6E). The large potassium ions probably act as pillars and maintain the wide interslab space, even in the small amount of residual potassium. Goodenough et al. compared the voltage profiles of O3eLiCoO2 and P2-NaxCoO2 and proposed that the electrostatic repulsion between the alkalimetal ions (Aþ-Aþ) influences the CoeO bond-length.41 The larger sodium ions elongate the Co(III)eO bond in O3eNaCoO2 as compared to that in O3eLiCoO2, thereby providing differences in the working potentials of the redox reactions. Therefore, the low-voltage operation of P2- and P3-K1-xCoO2 would originate from the long CoeO distances. This was attributed to the large ionic size of the Kþ ions and the strong Kþ-Kþ repulsion in the interslab space. The Kþ-Kþ repulsion also stabilized the Kþ/vacancy ordering at the discrete K contents in KxMO2, resulting in the stepwise voltage curves observed for the K cells.12,80 The voltage steps are significant in the discharge voltage curve of P2-K1-xCoO2, as indicated by the arrows in Fig. 6D. Fig. 6E reveals that the changes in the lattice parameters are rather monotonous without any changes in the structural type and SG. These results suggest that the

Electrode materials for K-ion batteries

91

Fig. 6 Galvanostatic charge/discharge curves of (A) O3-, O4-, and O2-type LiCoO2 in Li cells, (B) P2-, P’3-, and O3-type NaxCoO2 in Na cells, and (C) P2- and P3-type KxCoO2 in K cells. The compositions at the end of KxCoO2 discharge were normalized to K0.333CoO2. (D) Operando in situ XRD patterns of P’2-KxCoO2 in a K cell during discharge and the corresponding voltage curve as a function of K content. (E) Interslab distances of A|| AxCoO2 (A ¼ Li, Na, and K) as a function of the alkali metal content estimated from the capacity. The K content of KxCoO2 at the voltage drop > 3 V was normalized to 0.5. (A–C) Reproduced from Ref. Kubota, K.; Dahbi, M.; Hosaka, T.; Kumakura, S.; Komaba, S., Towards K-Ion and Na-Ion Batteries as “Beyond Li-Ion”. Chem. Rec. 2018, 18(4), 459–479 with permission. Copyright, 2018 The Chemical Society of Japan & Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim. (D, E) Reproduced from Ref. Hironaka, Y.; Kubota, K.; Komaba, S., P2- and P3-KxCoO2 as an Electrochemical Potassium Intercalation Host. Chem. Commun. 2017, 53(26), 3693–3696 with permission. Copyright, 2017 The Royal Society of Chemistry.

voltage steps mainly arise as a result of Kþ/vacancy ordering.80,81 Such orderings produce voltage steps in the charge/discharge voltage curves, while fewer voltage-steps and smooth voltage curves are desired for the practical use of KxMO2. According to the experience and knowledge obtained for NaxMO2, fewer voltage-steps can be partially realized in binary, ternary, and quaternary transition-metal systems.57

7.04.2.1.4

P2- and P3-type binary and ternary transition-metal systems

The 3d transition-metal ions resulting from crystallization into O3-, P3-, and P2-type KxMO2 are very limited in the single transition-metal system (Table 2). However, the trend is quite different in the binary and ternary transition-metal (including pblock elements) systems, as observed in the NaxMO2 trends.57 Thus, many transition-metals can be accommodated in the binary and ternary systems. Since O3- and P2-type Na2/3Fe0.5Mn0.5O2 were reported as element-abundant and cobalt-free high-capacity positive electrode materials for NIBs, electrochemical K (de)intercalation properties were examined for P2-K0.7Fe0.5Mn0.5O2 nanoparticles82,83 and P3-K0.7Fe0.5Mn0.5O2 nanowires.84 All the materials were prepared by conventional solid-state reactions of nanometer-sized precursors under Ar gas atmosphere. Both the P2-K0.7Fe0.5Mn0.5O2 nanoparticles and P3-K0.7Fe0.5Mn0.5O2 nanowires (Fig. 7A) deliver reversible capacities > 170 mAh g 1 with good cycle stability in the voltage range 1.5–4.0 V. The abnormally large capacities in the potassium iron manganese oxides were due to the specific particle size and morphology. Indeed, smaller reversible capacities and poor rate performances were observed for the micrometer-sized P2-K0.65Fe0.5Mn0.5O283 and P3-K0.7Fe0.5Mn0.5O2.84 Good cycle stability in the K cells of the P2-K0.7Fe0.5Mn0.5O2 nanoparticle and P3-K0.7Fe0.5Mn0.5O2 nanowire electrodes were observed, and reversible structural changes during K (de)intercalation were confirmed by operando in situ XRD.82–84 Deng et al. also performed ex situ XRD measurements of the K0.65Fe0.5Mn0.5O2 submicron particle electrodes and observed an OP4-like phase with a smaller interslab distance than that of the pristine electrode after charging to 4.2 V.83 These results indicated a wider K-content range of KxMO2 during the charging/discharging than the expected range of 1/3  x  2/3. In addition, K ||P3-K0.54Co0.5Mn0.5O285 cell exhibited reversible charge/discharge behavior showing a capacity of > 100 mAh g 1 (Fig. 7B). Interestingly, the operando in situ XRD patterns of the P3-KxCo0.5Mn0.5O2 electrode revealed insignificant changes in the c-lattice parameter corresponding to the interslab distance and the wide K-content range 0.25  x  0.75 in KxCo0.5Mn0.5O2 during charging/discharging

92

Electrode materials for K-ion batteries

Fig. 7 (A) Initial charge and discharge curves of the P3-K0.7Fe0.5Mn0.5O2-nanowire electrode in a K cell measured at 20 mA g 1. (B) Cycle performance of the P3-K0.7Fe0.5Mn0.5O2-nanowire and P3-K0.7Fe0.5Mn0.5O2-particle electrodes measured at 500 mA g 1. (C) Comparison of the c-lattice parameters of P3-Kx[Co0.5Mn0.5]O2 obtained by first-principles calculations and the operando in situ XRD data. (D) X-ray absorption nearedge structure (XANES) spectra of Mn K-edge for P3-Kx[Co0.5Mn0.5]O2 measured at the OCV (2.8 V vs. Kþ/K) and at the end of charging/ discharging. (A) Reproduced from Ref. Wang, X.; Xu, X.; Niu, C.; Meng, J.; Huang, M.; Liu, X.; Liu, Z.; Mai, L., Earth Abundant Fe/Mn-Based Layered Oxide Interconnected Nanowires for Advanced K-Ion Full Batteries. Nano Lett. 2017, 17(1), 544–550 with permission. Copyright, 2016 American Chemical Society. (B–D) Reproduced from Ref. Choi, J.U.; Kim, J.; Hwang, J. Y.; Jo, J. H.; Sun, Y. K.; Myung, S. T., K-0.54[Co0.5Mn0.5]O2: New Cathode with High Power Capability for Potassium-Ion Batteries. Nano Energy 2019, 61, 284–294 with permission. Copyright, 2019 Elsevier Ltd.

(Fig. 7C). The binary transition metal ions would stabilize the layered structures in the wide K-content range, resulting in the large capacity > 100 mAh g 1. Choi et al. also examined the valence state and local environment of Mn and Co by X-ray absorption spectroscopy for the P3-KxCo0.5Mn0.5O2 electrodes at OCV ( 2.8 V) and after charging/discharging.85 The X-ray absorption near-edge structure (XANES) spectra at the Mn K-edge revealed that the Mn3 þ/4 þ redox reaction mainly occurred in the voltage range 1.5– 2.5 V. On the other hand, cobalt was found to contribute to the Co3 þ/4 þ redox reaction in the range 2.5–3.9 V. As manganese and cobalt displayed different redox potentials during charging/discharging of the P3-KxCo0.5Mn0.5O2 electrode, voltage versus capacity plots were also altered according to the transition metals in the MO2 slab. To realize a high energy performance (i.e. high capacity and high voltage) of layered KxMO2 as a positive electrode, the selection of transition metals is essential, and electrochemically inactive metals are also accommodated in the slab with changing electrochemical properties of KxMO2. KxMO2 layered materials generally display low-voltage redox in K cells, evidenced in the charge/discharge curves of KxCoO2 compared to those of LixCoO2 and NaxCoO2 (Fig. 6A–C). In 2018, Masese et al. reported high-voltage K intercalation performances in K2/3Ni2/3-xMxTe1/3O2 (M ¼ Mg, Co, Zn), which comprised Te6þ ([Kr]4d10) in the [Ni2/3-xMxTe1/3]O2 slab.86 The Te and Ni/M were ordered, and honey-comb structures were formed in the slab. This resulted in P2-K2/3[Ni2/3-xMxTe1/3]O2 [described as K2(Ni,M)2TeO6], which comprised a hexagonal lattice with the SG P63/mcm (No. 193; Fig. 8A and B)86 and was isostructural to Na2Ni2TeO6.87 The P2-K2/3Ni2/3Te1/3O2 electrode delivered a reversible capacity of 65 mAh g 1 (Fig. 8C), which is smaller than those of KxCrO2 (ca. 90 mAh g 1), KxMnO2 (80–100 mAh g 1), and KxCoO2 (70 mAh g 1). However, the averaged voltage of P2-K2/3Ni2/3Te1/3O2 (3.1 V) during potassiation is one of the highest among the KxMO2 layered oxides reported to date. Moreover, the P2-K2/3[Ni1/2Zn1/6Te1/3]O2 [described as K2Ni1.5Zn0.5TeO6] electrode delivered a higher reversible capacity of 75 mAh g 1 but displayed inferior capacity retention compared to that of the P2-K2/3[Ni2/3Te1/3]O2 electrode.86 The increase in the reversible capacities by replacing Ni2þ with Zn2þ is interesting because redox-inactive Zn2þ ions are exchanged for redox-active Ni2þ ions in P2-K2/3[Ni2/3Te1/3]O2. This observation indicates the contributions of the oxide ions to the redox reaction, which is supported DFT calculation and XANES spectra. In the research field of

Electrode materials for K-ion batteries

93

Fig. 8 Schematic illustrations of the (A) crystal structure of P2-type K2/3[Ni2/3Te1/3]O2 (K2Ni2TeO6) and (B) honeycomb in-plane structure of the [Ni2/3Te1/3]O2 slab. (C) Charge and discharge curves of the K||K2Ni2TeO6 cell filled with ionic liquid (0.5 M potassium bis(trifluoromethanesulfonyl) amide (KTFSA) in Pyr13TFSA) electrolyte. Galvanostatic measurements were conducted at the rate C/20 for 70 cycles at room temperature. The grey dashed line indicates the calculated voltages. (C) Reproduced from Ref. Masese, T.; Yoshii, K.; Yamaguchi, Y.; Okumura, T.; Huang, Z. D.; Kato, M.; Kubota, K.; Furutani, J.; Orikasa, Y.; Senoh, H.; Sakaebe, H.; Shikano, M., Rechargeable Potassium-Ion Batteries with Honeycomb-Layered Tellurates as High Voltage Cathodes and Fast Potassium-Ion Conductors. Nat. Commun. 2018, 9, 3823 with permission from Springer Nature. Copyright 2018 Nature Publishing Group.

NIBs, the redox of oxide ions has been reported for P2-Na2/3[Mn7/9Zn2/9]O288 and Na2/3[Mg0.28Mn0.72]O289 to enhance the reversible capacities of NIBs. We predict that studies on anion redox will extend to potassium layered oxides in the near future. In summary, the 3d transition metals leading to crystallization into O3-, P3-, and P2-type structures are limited for the layered KxMO2 compounds as compared to those of NaxMO2.57 However, multi-transition-metal systems of KxMO2 enable crystallization as observed in from the synthesis of P3- and P2-KxFe0.5Mn0.5O2, P3-KxCo0.5Mn0.5O2, and P3-KxNi1/6Co1/6Mn2/3O2 and as wellknown in NaxMO2.57 Since the Kþ-Kþ electrostatic repulsion is stronger than those of Liþ-Liþ and Naþ-Naþ in AxMO2, more stepwise voltage variation, corresponding to Kþ/vacancy ordering, is generally observed in the charge and discharge curves of KxMO2 in single transition-metal systems. Multiple transition-metal systems and partial replacement with redox-inactive metals are efficient to disturb Kþ//vacancy ordering in the interslab and reduce the voltage steps, thereby progressing toward their use in practical KIBs. Achieving both a high capacity and high average voltage, KxMO2 oxides are still challenging for their practical use as positive electrode materials for KIBs.

7.04.2.2 7.04.2.2.1

Prussian blue analogues Prussian blue analogues as electrode materials

In the case of light alkali metals such as Li and Na, the layered oxides AMO2 are attractive in terms of material science, topotactic/ reversible intercalation, the two-dimensional (2D) diffusion path of Aþ ions, and practical use in secondary batteries. On the other hand, in our opinion, open 3D framework structures would be preferable for the electrochemical insertion of larger alkali metals such as Kþ ions. Among the various 3D materials, open-channel hexacyanometalate compounds, also known as PBAs, have been studied as Li,90 Na,91 and K11,15,17 insertion hosts. The chemical formulas of PBAs are represented as AxM1[M2(CN)6]y$nH2O (0  x  2, y  1), where A is a mobile metal, and M1 and M2 can be various metals such as Ti, V, Cr, Fe, Co, Ni, Cu, and Zn. AxM1[M2(CN)6]y$nH2O is abbreviated as M1M2-PBA or AM1M2-PBA for simplicity hereafter. Fig. 9A illustrates a typical PBA crystal structure. PBAs have 3D open-frameworks in which M1N6 and M2C6 octahedra are linked via cyano ligands.92 The framework structure provides an open 3D channel available for mobile metal diffusion, which is suitable for the diffusion of large ions. Another feature of PBAs is the different spin states of the transition metals in the M1N6 and M2C6

94

Electrode materials for K-ion batteries

Fig. 9 Crystal structures of PBAs: (A) open (A0M1[M2(CN)6], where A is a mobile metal, and M1 and M2 can be various metals) and (B) dense (A2M1[M2(CN)6]) structures; the blue and pink polyhedra represent M1N6 and M2C6, while the yellow spheres represent mobile metals. (C) Galvanostatic discharge curves of LixMn[Fe(CN)6]y (LiMnFe-PBA), NaxMn[Fe(CN)6]y (NaMnFe-PBA), K1.75Mn[Fe(CN)6]0.93 (KMnFe-PBA), Na1.80Fe [Fe(CN)6]0.93 (NaFeFe-PBA), and K1.63Fe[Fe(CN)6]0.89 (KFeFe-PBA) in nonaqueous Li, Na, and K cells, respectively.

octahedra, which is explained by the ligand field theory. For transition metals such as Fe and Mn, the high-spin (HS) M1 is located in the M1N6 octahedron, while the low-spin (LS) M2 is located in the M2C6 octahedron owing to their respective weak N-coordinated and strong C-coordinated ligand fields.93 The typical PBAs possess a cubic structure (SG Fm-3m) in which all the octahedra are arranged linearly (Fig. 9A). On the other hand, the octahedra rotate cooperatively depending on the guest cations and molecules, resulting in monoclinic (SG P21/n) or rhombohedral (typically SG R-3; Fig. 9B) structures. The Na-rich phases have monoclinic structures when the water molecules are in the interstitial site (interstitial water). On the other hand, Na-rich PBAs have rhombohedral structures when they contain a negligible amount of interstitial water.92,94 Similar to the Na-poor phase, K-poor PBAs display a cubic structure. However, to the best of our knowledge, K-rich PBAs always have monoclinic structures, regardless of the drying conditions.11,17,38 This phenomenon is explained by the large ionic radii of the Kþ ions, which increase the Pauli repulsion in the lattice and prevent lattice shrinkage,94 and the formation of a rhombohedral phase with a smaller lattice volume than that of the monoclinic phase. Furthermore, most K-rich PBAs contain significantly less interstitial water11,17,38 than Na-rich PBAs, which typically include  2 mol interstitial water per formula.94 These facts indicate that Kþ ions and water molecules rarely occupy the interstitial sites of the PBA simultaneously.

7.04.2.2.2

Li, Na, and K insertion into Prussian blue analogues

A review of the history of PBAs as electrode materials for secondary batteries reveals that Li insertion/extraction into M[Fe(CN)6] (M ¼ V, Mn, Co, Ni, Cu) was first demonstrated in 1999 by Imanishi et al.90 Although some studies were reported after this initial work, little attention has been paid to PBA positive electrode materials as lithium insertion hosts.95,96 Research after 2010s has focused on PBAs as positive electrode materials for NIBs. The crystal structures and chemical compositions of Na-containing PBAs have been extensively investigated and thus, transition metals,91,97 [Fe(CN)6] vacancies,98 Na contents,93,99 and interstitial water contents94 are known to influence the Na insertion/extraction performances. Potassium insertion/extraction into/from KFe [Fe(CN)6] thin films in a nonaqueous K cell was first reported in 2004.15 After 2015, PBAs were reported as electrode materials for nonaqueous KIBs. Thus, we next discuss the influence of insertion guest ions on the redox voltage and reversible capacity of PBAs. Fig. 9C summarizes the charge/discharge curves of LixMn[Fe(CN)6]y (LiMnFe-PBA), NaxMn[Fe(CN)6]y (NaMnFe-PBA), K1.75Mn [Fe(CN)6]0.93 (KMnFe-PBA), Na1.80Fe[Fe(CN)6]0.93 (NaFeFe-PBA), and K1.63Fe[Fe(CN)6]0.89 (KFeFe-PBA) in the Li, Na, and K cells, respectively.14 The Na- and K-PBAs were synthesized by a simple precipitation method. On the other hand, LiMnFe-PBA was

Electrode materials for K-ion batteries

95

synthesized via electrochemical ionic exchange in the K cell because it is hardly obtained via a simple precipitation method. The LiMnFe-PBA, NaMnFe-PBA, and KMnFe-PBA delivered similar discharge capacities of 137, 140, and 137 mAh g 1, respectively. The Mn-Fe-based PBAs exhibited average discharge voltages of 3.5, 3.4 (3.7 V vs. Li), and 3.8 V (3.7 V vs. Li) in the Li, Na, and K cells, respectively. Therefore, the Kþ ion insertion potential into KMnFe-PBA was higher than that of its Li counterparts, even when the differences in the operating voltages due to the electrode potential at the counter electrode were considered. This high operation potential contrasts with that of the layered oxides described in the previous section. Therefore, the specific energy density of LiMnFe-PBA, NaMnFe-PBA, and KMnFe-PBA were 446, 472, and 521 Wh (kg of PBA) 1 in the Li-, Na-, and K-metal cells, respectively. These results indicated that PBAs are promising electrode materials, especially for high energy density KIBs. Thus, various PBAs have been developed for KIBs.

7.04.2.2.3

Prussian blue analogues for K-ion batteries

7.04.2.2.4

Structural evolution during Kþ insertion

Reversible K insertion into a thin film ( 1 mm) of KFe[Fe(CN)6] was first reported by Eftekhari in 2004.15 More than 10 years after publication, further studies on the electrode performance of powdery PBAs have been performed in K cells as summarized in Table 3. KFeFe-, KMnFe-, KCoFe-, KNiFe-, KCuFe-, and KTiFe-PBAs have been studied as positive electrode materials and only Co3[Co(CN)6]2 has been studied as a negative electrode material. K2Zn3[Fe(CN)6]2 was also reported as a positive electrode material. Notably, K2Zn3[Fe(CN)6]2 is not included in the narrow sense of PBAs because its structure is different. The charge/discharge curves of KFeFe-PBA, KMnFe-PBA, KNiFe-PBA, KCuFe-PBA, KCoFe-PBA, and KTiFe-PBA are illustrated in Fig. 10. KNiFe-PBA, KCuFe-PBA, and KCoFe-PBA displayed capacities in the range 30–70 mAh g 1 corresponding to a one-electron redox, attributed to the LS Fe2 þ/3 þ redox. On the other hand, Ni2þ, Cu2þ, and Co2þ were inactive. Conversely, KFeFe-PBA, KMnFe-PBA, and KTiFePBA presented high capacities in the range 130–140 mAh g 1, corresponding to a two-electron reaction, i.e. redox of Fe2 þ/ 3þ , Mn2 þ/3 þ and likely Ti2 þ/3 þ in the PBAs. In all the PBAs, a capacity in the range 30–60 mAh g 1, which appeared at 3.75– 4.0 V vs. Kþ/K, was attributed to the redox of LS FeC6. The redox potential of M1N6 varied greatly depending on the transition metal. Thus, the HS Mn2 þ/3 þ redox occurred at high potentials > 4 V vs. Kþ/K, whereas HS Fe2 þ/3 þ and Ti2 þ/3 þ occurred at values of 3.3 and  2.0 V, respectively. These results indicate that KMnFe-PBA is an attractive positive electrode material with potential to increase the energy density of KIBs. Moreover, PBAs generally show good long-term cycle stabilities. For example, KMnFe-PBA demonstrated capacity retentions > 90% after 350 cycles in 7 mol kg 1 potassium bis(fluorosulfonyl)amide (KFSA)/1,2-dimethoxyethane (DME) electrolyte.107 These reversible Kþ insertion/extraction properties were attributed to the 3D open framework structure, which provides suitable open channels and interstitial sites for Kþ ion insertion.

The structural evolution of PBA accompanying K extraction/insertion is next compared with that of Na based on the literature.11,94 Rhombohedral Na1.89Mn[Fe(CN)6]0.97$0.3H2O exhibited a reversible transition from the rhombohedral to the cubic structure in the Na content range x ¼ 2–1 in NaxMn[Fe(CN)6] accompanying a volume expansion > 20% and from the cubic to the tetragonal structure in the range x ¼ 1–0 in NaxMn[Fe(CN)6] with  10% volume shrinkage.94 The structural evolution of KMnFe-PBA has been investigated with operando XRD (Fig. 11A). Monoclinic KMnFe-PBA exhibited a reversible transition from the monoclinic to the cubic structure in the lower voltage plateau (K content range x ¼ 2–1 in KxMn[Fe(CN)6]) via a two-phase reaction and from the cubic to the tetragonal structure in the higher voltage plateau (K content range x ¼ 1–0 in KxMn[Fe(CN)6]; Fig. 11B).11 This structural evolution resulted in volume changes of 14.4 and  6.0% for the respective monoclinic to cubic and cubic to Table 3

Overview of PBAs and their electrochemical properties for K-ion batteries.

Material

Specific capacity/mAh g 1

Average discharge voltage/V vs. K

KxFe[Fe(CN)6] thin film15 rGO@KxFe[Fe(CN)6]@steel100 KxFe[Fe(CN)6]16 KFe[Fe(CN)6]101 K1.92Fe[Fe(CN)6]0.94$0.5H2O38 K1.89Mn[Fe(CN)6]0.9217 K1.75Mn[Fe(CN)6]0.93$0.16H2O11 K1.6Mn[Fe(CN)6]0.96$0.27H2O102 K1.64Fe[Fe(CN)6]0.89$0.15H2O11 K0.220Fe[Fe(CN)6]0.805$4.01H2O103 KxCo[Fe(CN)6]16 KxNi[Fe(CN)6]16 KxCu[Fe(CN)6]16 K0.3Ti0.75Fe0.25[Fe(CN)6]0.95$2.8H2O104 K2Zn3[Fe(CN)6]2105 Co3[Co(CN)6]2106

79 97 111 119 128 142 141 107 130 78 60 64 35 138 67 366

3.7 3.1 3.5 3.5 3.6 3.6 3.8 3.9 3.5 3.1 3.3 3.6 3.2 2.3 3.7 0.6

96

Electrode materials for K-ion batteries

Fig. 10 Charge/discharge curves of PBAs in nonaqueous K cells: (A) KFeFe-PBA, (B) KMnFe-PBA, (C) K-NiFePBA, (D) K-CuFePBA, (E) K-CoFePBA, and (F) K-TiFePBA. (A–B) Reproduced from Ref. Bie, X.; Kubota, K.; Hosaka, T.; Chihara, K.; Komaba, S., A Novel K-Ion Battery: Hexacyanoferrate(ii)/ Graphite Cell. J. Mater. Chem. A 2017, 5(9), 4325–4330 with permission. Copyright, 2017 The Royal Society of Chemistry. (CE) Reproduced from Ref. Wu, X.; Jian, Z.; Li, Z.; Ji, X., Prussian White Analogues as Promising Cathode for Non-Aqueous Potassium-Ion Batteries. Electrochem. Commun. 2017, 77, 54–57 with permission. Copyright, 2017 Elsevier Ltd. (F) Reproduced from Ref. Luo, Y.; Shen, B.; Guo, B.; Hu, L.; Xu, Q.; Zhan, R.; Zhang, Y.; Bao, S.; Xu, M., Potassium Titanium Hexacyanoferrate as a Cathode Material for Potassium-Ion Batteries. J. Phys. Chem. Solids 2018, 122, 31–35.

Electrode materials for K-ion batteries

97

Fig. 11 Structural evolution of KMnFe-PBA: (A) contour maps of the operando XRD of KMnFe-PBA during the first cycle in a K cell, illustrating the major diffraction peaks indexed using monoclinic (SG P21/n), cubic (SG Fm-3m), and tetragonal (SG I-4m2) phases and (B) crystal structures and phase transitions observed by electrochemical potassium extraction (charging) and insertion (discharging). Reproduced from Ref. Bie, X.; Kubota, K.; Hosaka, T.; Chihara, K.; Komaba, S., A Novel K-Ion Battery: Hexacyanoferrate(ii)/Graphite Cell. J. Mater. Chem. A 2017, 5(9), 4325–4330 with permission. Copyright, 2017 The Royal Society of Chemistry.

tetragonal transitions. A comparison of the volume changes of rhombohedral NaMnFe-PBA and monoclinic KMnFe-PBA revealed that KMnFe-PBA displayed 6% less volumetric change during K extraction/insertion, which is advantageous for cyclability.

7.04.2.2.5

Particle size and anion vacancy effect on the electrochemical performance

He and Nazar reported the synthesis of KFeFe-PBAs with different particle sizes and elucidated their electrochemical properties.108 They synthesized KFeFe-PBA with small (KFeFe-PBA-S) and large (KFeFe-PBA-L) whose particle sizes were  20 nm and > 1.5 mm, respectively. KFeFe-PBA-S presented a high discharge capacity of 140 mAh g 1 in the K cell, whereas KFeFe-PBA-L only presented a capacity of  10 mAh g 1. The low discharge capacity of KFeFe-PBA-L was attributed to the influence of Kþ ion diffusion kinetics.108 Fiore et al. synthesized KMnFe-PBAs in various conditions, i.e. different concentrations of precursors, citrate as a chelator, and KCl in the synthesis solution.109 They have also found that the larger particle size lowers the reversible capacity due to the

98

Electrode materials for K-ion batteries

kinetic effect, whereas larger particles exhibited better cyclability and Coulombic efficiency. These results differed from the Na insertion properties of PBA, whereby NaFeFe-PBA displayed a large particle size of several mm indicating high discharge capacity.93 In 2020, we have investigated the effect of particle size and [Fe(CN)6] vacancies on the K-insertion reaction. By chelate-assisted precipitation method, we synthesized two samples (KMnFe-PBA-S, KMnFe-PBA-L) whose particle size were  200 nm and > 2 mm (Fig. 12A and 15B). Both samples had negligible [Fe(CN)6] vacancy content. Moreover, [Fe(CN)6] vacancy-rich KMnFe-PBA (KMnFe-PBA-V) with large particle size (Fig. 12C) and considerable anion vacancies (15%) via the ionic exchange method from NaMnFe-PBA.110 Fig. 12D depicts charge-discharge curves of the KMnFe-PBAs at 0.1C and 5C. Compared with KMnFe-PBA-L, KMnFe-PBA-V showed a large reversible capacity of about 140 mAh g 1, which was comparable to that of KMnFe-PBA-S. In addition, KMnFe-PBA-V maintained much larger capacity at 5C than KMnFe-PBA-L and even than KMnFe-PBA-S. GITT measurements and DFT calculations suggest that the proper number of anion vacancies enhances Kþ ion diffusion. Therefore, material design based on particle size and anion vacancies is important for the high performance of PBAs as K-insertion hosts. Despite their promising electrochemical performances, PBAs exhibit two general disadvantages, namely low bulk density and electronic conductivity, which originate from their unique structure and chemical-bond character. Although the low bulk density of PBAs is an essential feature and is difficult to overcome, modification of the mixing conditions of the electrode components and additives would improve the density and homogeneity of the current distribution of the composite electrode. On the other hand, the low electronic conductivity of PBAs causes a low energy density of the electrode due to a large amount of conductive carbon in the electrode. Although carbon coating is a typical technique to improve the electronic conductivity of electrode materials, the limited thermal stability (< 400  C) of the PBAs hinders carbon coating for PBAs. Therefore, to improve the low electronic conductivity of PBAs, researchers are developing PBA carbon composite synthetic methods for NIBs.111,112 Although few such studies for KIBs have been published to date, previous success in the field of NIBs will induce further improvement of the electrochemical performance of PBAs in K cells.

Fig. 12 TEM images of (A) KMnFe-PBA-S, (B) KMnFe-PBA -L-, and (C) KMnFe-PBA-V. (D) Charge-discharge curves of KMnFe-PBA-S, L-, and V at a current density of 0.1C and 5C. Reproduced from Ref. Hosaka, T.; Fukabori, T.; Kojima, H.; Kubota, K.; Komaba, S., Effect of Particle Size and Anion Vacancy on Electrochemical Potassium Insertion of Potassium Manganese Hexacyanoferrates. ChemSusChem 2021 with permission. Copyright, 2020 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Electrode materials for K-ion batteries 7.04.2.3

99

Polyanionic compounds as positive electrode materials

As discussed above, PBAs display high working potentials > 3.5 V in K cells, values that are even higher than those of the PBAs in Li cells. Moreover, the 3D open framework structure of PBAs would be suitable for the insertion/extraction of large Kþ ions. Similar to the PBAs, various polyanionic compounds have 3D open-framework structures comprising MOx (M ¼ transition metals) and (XO4)n- (X ¼ P, S, As, Si, Mo, and W) polyhedra. This open structure of polyanionic compounds would be advantageous for fast ionic diffusion. For example, NASICON-type compounds have attracted attention as solid-state Naþ ion conductors since the first discovery of Na1 þ xZr2P3-xSixO12 (0  x  3) in 1976.113,114 Moreover, the inductive effect proposed by Manthiram and Goodenough increases the redox potential of transition metals.115,116 The higher redox-potential originates from the decrease in the covalency of the M-O bonds due to the highly covalent X-O bond, which depends on the electronegativity of X. Because of their fast ionic diffusion and inductive effect, polyanionic compounds have been studied as electrode materials for LIBs since the 1980s.117 To date, the most successful material is the triphylite LiFePO4, which displays an olivine-type structure. Since Padhi et al. realized the electrochemical lithium extraction/insertion from/into triphylite LiFePO4 in 1997,118 the synthetic methods, elemental substitutions, and lithium insertion/extraction mechanisms of triphylite LiFePO4 have been widely studied.117 Following the significant breakthrough of triphylite LiFePO4, researchers have focused on replacing PO43 with other anions and introducing F to realize both high reversible capacities and high redox voltages and developed promising polyanionic compounds, e.g. pyrophosphate Li2FeP2O7,119 phosphate fluoride Li2FePO4F,120 and sulfate fluoride LiFeSO4F.121 Interestingly, iron- and vanadium-based polyanionic compounds displayed promising electrochemical properties, whereas high-performance layered oxides for LIB mostly contained Co, Ni, and Mn. With respect to practical use, the use of costly elements, as well as the toxicity of V, would also be problematic, while abundant and non-toxic elements such as Fe and Mn are highly desired. The progressive studies on Li and Na battery fields have motivated researchers to study polyanionic compounds as positive electrode materials for high voltage KIBs. The polyanionic compounds for KIB application and their electrochemical performances are summarized in Table 4. Similar to the Li and Na cases, vanadium-based polyanionic compounds such as phosphates and fluorophosphates have been actively studied and have demonstrated good electrochemical performances.18,20,124,126 Furthermore, redoxactive Fe- and Mn-based polyanionic compounds have shown potential as potassium insertion hosts. As mentioned above, the most common polyanionic compound used as an electrode material is the triphylite LiFePO4. On the other hand, triphylite KFePO4 is not

Table 4

Overview of polyanionic compounds and their electrochemical properties for K-ion batteries.

Materials

Structure

KVPO4F18 Stoichiometric KVPO4F19 KVOPO418 KVP2O720 KTiPO4F122 K4Fe3(PO4)2(P2O7)123

KTiOPO4 (KTP)-type orthorhombic (Pna21) KTP-type orthorhombic (Pna21) KTP-type orthorhombic (Pna21) KAlP2O7-type monoclinic (P21/c) KTP-type orthorhombic (Pna21) Na4Co3(PO4)2(P2O7)-type orthorhombic (Pn21a) NASICON-type trigonal (R-3c) Orthorhombic (Cmcm) – Heterosite-type orthorhombic (Pmna) Monoclinic (P21/n) Triclinic (P-1) Layered tetragonal (P-421/m) KTP-type orthorhombic (Pna21)

K3V2(PO4)3124,125 K3V2(PO4)2F3126 Amorphous - FePO4127 Heterosite FePO421 KFePO421 KMnPO421 K2FeP2O721 Orthorhombic KFeSO4F21,128 Monoclinic KFeSO4F21,129 KFe(C2O4)F130 K2[(VOHPO4)2(C2O4)]131 KMoP2O720 KFeP2O720 KCrP2O720 K2MnP2O720,21 K2Fe2(SO4)3132 K2Cu2(SO4)3132

Layered monoclinic (C2/c) Orthorhombic (Cmc21) Triclinic (P-1) KAlP2O7-type monoclinic (P21/c) KAlP2O7-type monoclinic (P21/c) KAlP2O7-type monoclinic (P21/c) Monoclinic (P21/c) Langbeinite-type cubic (P213) Orthorhombic (P212121)

Theoretical capacity/ mAh g 1

Actual capacity/ mAh g 1

Average discharge potential/V vs. Kþ/K

131 131 133 102 133 117

90 100 80 61 95 120

4.13 4.3 4 4.15 3.1 2.8

106

54

3.5

115 178 178 141 141 175 (2-electron reaction) 128

104 156 120 Inactive Inactive 60 100

3.7 2.64 2.3

128 133 108 87 100 101 175 (2-electron reaction) 129 124

50 100 98 25 Inactive Inactive Inactive Inactive in Li cell Inactive in Li cell

2.7 3.6 3.4 2.6 3.85 2.5

100

Electrode materials for K-ion batteries

thermodynamically stable, while KFePO4 (SG P21/n) is a stable phase comprising FeO4 tetrahedra, FeO5 polyhedra, and PO4 tetrahedra.133 However, this compound was electrochemically inactive in K cells because of the difficulty in Kþ diffusion.21 Therefore, the selection of polyanionic compounds having a suitable structure for potassium insertion/extraction is critical to developing promising electrode materials.

7.04.2.3.1

KTiOPO4-type structure materials

As shown in Table 4, some KTiOPO4 (KTP)-type materials, such as KVPO4F and KVOPO4, show good electrode performance. Fedotov et al. demonstrated the electrochemical extraction of Kþ ions from KVPO4F to K0.15VPO4F, with the K0.15VPO4F electrode displaying excellent rate capabilities in the Li cells in 2016.134 Moreover, the electrochemical performances of KVPO4F18,135 and KVOPO418 in K cells were reported in 2017. Fig. 13A illustrates the crystal structures of these compounds. Both KVPO4F and KVOPO4 have a KTiOPO4 (KTP)-type structure (SG Pna21), comprising corner-sharing VO4F2/VO6 octahedra and PO4 tetrahedra. KVPO4F and KVOPO4 delivered reversible capacities of  92 and 84 mAh g 1, respectively, and their average depotassiation voltages were as high as 4 V (Fig. 13B and C). Both KVPO4F and KVOPO4 exhibited reversible structural evolution, with the lattice volume changes, after charging to 5.0 V, of 5.8% for KVPO4F and 3.3% for KVOPO4. These two compounds also exhibited excellent rate capabilities, maintaining 90.7 and 93.9%, respectively, at a 5C rate compared to the C/20 rate in K cells. A possible reason for the small volume changes and excellent rate capabilities is the open framework of their KTP-type structure. In 2018, Kim et al. reported the chemical composition effect on the electrochemical performance of KVPO4F.19 Stoichiometric KVPO4F and oxidized KVPO4.36F0.64 were synthesized by changing the calcination conditions, i.e. carbon black was used as a reducing agent only for the stoichiometric KVPO4F. Stoichiometric KVPO4F delivered a high reversible capacity of  105 mAh g 1 and the high average voltage 4.33 V with clear plateaus compared to those of the oxidized KVPO4.36F0.64. This latter compound presented a smooth charge/discharge curve likely due to anion disordering. Indeed, DFT calculations revealed that KxVPO4F exhibited a stable intermediate phase at the K amounts of 0.75, 0.625, and 0.5 during K extraction, which is consistent with the plateau potential in the charge/discharge curves and in situ XRD results of the stoichiometric KVPO4F. The ionic diffusion mechanism in AVPO4F was studied via the bond valence energy landscape (BVEL) method, DFT calculations, and electrochemical impedance measurements. The BVEL results indicated that Liþ, Naþ, and Kþ ions can diffuse into the AVPO4F

Fig. 13 (A) Crystal structure of KTP-type KVPO4F and KVOPO4. Charge/discharge profiles of (B) KVPO4F and (C) KVOPO4 at 2.0–5.0 V with 1 M KPF6/EC:PC (1:1 v/v) electrolyte. (B–C) Reproduced from Ref. Chihara, K.; Katogi, A.; Kubota, K.; Komaba, S., KVPO4F and KVOPO4 Toward 4 VoltClass Potassium-Ion Batteries. Chem. Commun. 2017, 53(37), 5208–5211 with permission. Copyright, 2017 The Royal Society of Chemistry.

Electrode materials for K-ion batteries

101

structure via three-, two-, and one-dimensional paths, respectively,135 and even Rbþ ion can diffuse via the one-dimensional path.136 Moreover, both the DFT calculations and impedance measurements revealed that AVPO4F displayed high diffusion coefficients for Kþ ions compared to those for Liþ and Naþ ions, indicating that the KTP structure is suitable for the diffusion of large ions such as Kþ ions.136 Orthorhombic KFeSO4F (o-KFeSO4F) having a KTP-type structure has also been studied as a potential positive electrode material.128 Lander et al. reported that KFeSO4F has two polymorphs with orthorhombic and monoclinic structures129 (Fig. 14A and B, respectively). As shown in Fig. 14C, BVEL calculation showed that the activation energy of potassium diffusion in o-KFeSO4F (0.40 eV) was much lower than that of m-KFeSO4F (2.83 eV).129 In 2019, we reported the electrochemical performances of oKFeSO4F and m-KFeSO4F in K cells. As expected from the BVEL results, o-KFeSO4F delivered higher reversible capacity and better rate capabilities than m-KFeSO4F (Fig. 14D).21 Moreover, the average voltage of o-KFeSO4F in a K half-cell was 3.6 V, which is a notably high Fe2þ/Fe3þ redox potential, owing to the strong inductive effects of the SO42 and F anions on the Fe2 þ/3 þ couple. Furthermore, the o-KFeSO4F electrode presented capacity retention of 86% after 100 cycles by using 5.5 mol kg 1 KFSA/diglyme electrolyte (Fig. 14E).21 This excellent cyclability in the highly concentrated electrolyte was attributed to negligible electrolyte decomposition as a result of the wide potential window.137

Fig. 14 Schematic crystal structure of (A) o-KFeSO4F and (B) m-KFeSO4F and (C) bond valence energy landscape (BVEL) of monoclinic and orthorhombic KFeSO4F. The energy values selected for the plots were 0.2 and 0.96 eV above the activation energies of the monoclinic and orthorhombic polymorphs, respectively. (D) Discharge rate capabilities of o-KFeSO4F and m-KFeSO4F with 5.5 mol kg 1 KFSA/diglyme electrolyte. The electrodes were charged at 0.1C and discharged at different current densities; 1C was defined as 128 mA g 1 from the theoretical capacity. Charge/discharge curves of the (E) o-KFeSO4F electrode in K half cells filled with 5.5 mol kg 1 KFSA/diglyme electrolyte at the current density of 12.8 mA g 1. (C) Reproduced from Ref. Lander, L.; Rousse, G.; Abakumov, A. M.; Sougrati, M.; van Tendeloo, G.; Tarascon, J. M., Structural, Electrochemical and Magnetic Properties of a Novel KFeSO4F Polymorph. J. Mater. Chem. A 2015, 3(39), 19754–19764 with permission. Copyright, 2015 The Royal Society of Chemistry. (D, E) Reproduced from Ref. Hosaka, T.; Shimamura, T.; Kubota, K.; Komaba, S., Polyanionic Compounds for Potassium-Ion Batteries. Chem. Rec. 2019, 19(4), 735–745 with permission. Copyright, 2019 The Chemical Society of Japan & Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

102

Electrode materials for K-ion batteries

In 2020, Fedotov et al. reported electrode performance of KTP-type KTiPO4F.122 The carbon-coated KTiPO4F showed a large reversible capacity of 94 mAh g 1 and good rate performance. Notably, KTiPO4F exhibited an electrode potential of 3.6 V in K cell, which is unusually high for a redox transition of Ti4 þ/Ti3 þ. The high redox potential would be due to a combination of inductive effect of PO43 and F and structural (charge/vacancy ordering, electrostatics).

7.04.2.3.2

KxMP2O7 (M ¼ Fe, Mn, and V)

In 2018, Park et al. conducted a systematic screening for K-containing polyanionic compounds as positive electrode materials.20 They searched for positive electrode material candidates from the Inorganic Crystal Structure Database (ICSD; > 187,000 compounds) by screening with composition restriction, channel tolerance, transition metal site restriction, channel type, transition metal valance, and maximum capacity from chemical composition. As a result, the redox reaction properties of the selected 10 materials were investigated with cyclic voltammetry, and the data were discussed by combining the visual observation results of color change after chemical oxidation. Fig. 15 displays the cyclic voltammograms and photographs of the synthesized materials before and after chemical oxidation. Interestingly, three kinds of pyrophosphate, namely KTiP2O7, KVP2O7, and KMoP2O7, reasonably displayed reversible redox activities as well as color changes after chemical oxidation. All the three pyrophosphates were structurally

Fig. 15 Cyclic voltammograms of 10 selected compounds in KPF6/EC:DEC tested at 20  C and the scan rate 0.1 mV s 1. Photos compare the color of the (left) as-prepared and (right) chemically oxidized samples. Reproduced from Ref. Park, W. B.; Han, S. C.; Park, C.; Hong, S. U.; Han, U.; Singh, S. P.; Jung, Y. H.; Ahn, D.; Sohn, K.-S.; Pyo, M., KVP2O7 as a Robust High-Energy Cathode for Potassium-Ion Batteries: Pinpointed by a Full Screening of the Inorganic Registry under Specific Search Conditions. Adv. Energy Mater. 2018, 8(13), 1703099 with permission. Copyright, 2018 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Electrode materials for K-ion batteries

103

related to KAlP2O7 with the SG P21/c. In addition, the charge-discharge data revealed that KTiP2O7, KMoP2O7, and KVP2O7 exhibited reversible potassium insertion, delivering reversible capacities of 22, 25, and 54 mAh g 1, respectively, when cycled at 50  C. On the other hand, the Fe- and Mn-based pyrophosphates KFeP2O7 and K2MnP2O7 were inactive (Fig. 15). The DFT calculations indicated that the redox potential of Fe3þ/Fe4þ in KFeP2O7 was > 5.0 V.20 This impractically high redox potential was attributed to the Fe3þ/Fe4þ redox center and inductive effect of the P2O4 7 unit. In contrast to KFe(III)P2O7, we reported that K2Fe(II)P2O7, comprising 2D structures, delivered a moderate reversible capacity of  60 mAh g 1 with a redox potential < 3.0 V.21 This low working voltage possibly originates from both the Fe2þ/Fe3þ redox center and the FeO4 tetrahedral environment, which leads to a weaker inductive effect compared to that observed in the FeO6 octahedral environment.138,139 Overall, among these pyrophosphates, KVP2O7 is the most promising because it presented the highest working voltage ( 4.0 V) and a relatively high reversible capacity.

7.04.2.3.3

K3V2(PO4)3 and K3V2(PO4)2F3

In addition to the good electrochemical performance of the vanadium pyrophosphate KVP2O7, some vanadium-based phosphate and phosphate fluoride compounds, such as K3V2(PO4)3 and K3V2(PO4)2F3, KVPO4F have been developed for KIBs. Electrochemical performance of NASICON-type K3V2(PO4)3 was first reported by Han et al., and K3V2(PO4)3/C nanocomposites showed 54 mAh-1 and an average operating voltage of ca. 3.5 V. Moreover, Zheng et al. reported that Rb-doped K2.95Rb0.05V2(PO4)3/C demonstrated superior rate performance, which was attributed to the higher Kþ ion diffusion coefficients for the Rb-doped material.140 According to the study by Lin et al.,126 orthorhombic K3V2(PO4)2F3 was synthesized via electrochemical ionic exchange from Na3V2(PO4)2F3 in a K cell. K3V2(PO4)2F3 presented an impressive electrochemical performance, delivering reversible capacities > 100 mAh g 1 with a working potential of  3.7 V vs. Kþ/K (Fig. 16A). Moreover, the K3V2(PO4)2F3 electrode demonstrated good capacity retention (95%) over 180 cycles. The structural evolution of K3V2(PO4)2F3 is schematically illustrated in Fig. 16B. The original structure is the same as that of Na3V2(PO4)2F3 with an SG Cmcm.141,142 Thus, the crystal structure of K3V2(PO4)2F3

Fig. 16 (A) Charge/discharge curves of K3V2(PO4)2F3 in the initial two cycles. The inset displays the corresponding dQ/dV curves. (B) Phase transformations of K3V2(PO4)2F3 during K ion insertion and extraction. The VO4F2 octahedral and PO4 tetrahedral structures are labeled in blue and purple, respectively, the K atoms are in cyan, and the white parts represent the vacancies. Reproduced from Ref. Lin, X.; Huang, J.; Tan, H.; Huang, J.; Zhang, B., K3V2(PO4)2F3 as a Robust Cathode for Potassium-Ion Batteries. Energy Storage Mater. 2019, 16, 97–101 with permission. Copyright, 2019 Elsevier Ltd.

104

Electrode materials for K-ion batteries

is built up of VO4F2 octahedra and PO4 tetrahedra, with an open framework. The structural changes during potassium insertion/ extraction, which were reversible, were accompanied by a minor volume change of 6.2%. They also revealed the transformation of K3V2(PO4)2F3 (SG Cmcm) into K2V2(PO4)2F3 (SG I4/mmm) upon charging to 4.2 V and the formation of the K1V2(PO4)2F3 phase (SG Cmc21) upon further charging up to 4.6 V.126 The structural change during potassium insertion/extraction was believed to be similar to that of Na3V2(PO4)2F3 in the Na cells.143 These results indicated that the open framework structure of Na3V2(PO4)2F3 is suitable even for topotactic potassium diffusion, although this has not been achieved in the direct synthesis of K3V2(PO4)2F3. Overall, previous studies indicated that the crystal structure plays a key factor for reversible potassium insertion and the rate capabilities. Because Kþ ions are too large to occupy the octahedral sites in heterosite FePO4, its crystallinity significantly decreased after potassium insertion. On the other hand, a few structures, such as K3V2(PO4)2F3 comprised the same open framework structure as Na3V2(PO4)2F3 and KTP-type structure, provide suitable channels to diffuse the Kþ ions and therefore exhibiting reversible potassium insertion/extraction properties and excellent rate capabilities. Most of the polyanionic compounds reported to date exhibit a lower specific energy density than that of the PBAs K 2Mn [Fe(CN)6] and K 2Fe[Fe(CN)6], owing to their limited specific capacities, while some polyanionic compounds displayed higher operation potentials. Therefore, the development of polyanionic compounds with large specific capacities (> 130 mAh g 1) as well as high operation voltages is a very attractive research target. Since different polyanionic compounds have many diverse structures compared to the structures of PBAs, the design of polyanionic compounds of suitable crystal structures will maximize both the volumetric energy density and potassium diffusion.

7.04.3

Negative electrode materials

7.04.3.1

Carbon materials

Like LIBs, electrochemical Kþ ion intercalation into carbon materials, such as graphite, soft carbon, and hard carbon, have been extensively studied. In this chapter, carbon materials as anode materials for KIB will be introduced.

7.04.3.1.1

K intercalation into graphite

Based on the development of carbon materials for LIBs, graphite is expected to be a promising candidate as negative electrode materials for KIBs because of graphite intercalation compounds (GICs) with K intercalant.144–147 Fig. 17 shows voltage profiles of electrochemical Li, Na, and K intercalation into graphite. In the Li cell, the graphite electrode provides a maximum reversible capacity of  370 mAh g 1, corresponding to reversible LiC6 formation from graphite.148 On the other hand, negligible capacitance is observed in the Na cell, since electrochemical intercalation of Na hardly occurs in defect-free graphite.149 In contrast to Na, in the K cell, K is electrochemically reversibly inserted into the graphite.2,25,26 Reversible capacities in the range 240–260 mAh g 1 are typically obtained in the K cell,2 values that are consistent with the theoretical capacity of 279 mAh g 1, by assuming reversible KC8 formation.150 After electrochemical intercalation, the electrodes of LiC6 and KC8 appeared gold (inset in Fig. 17). The study of the intercalation mechanism is of fundamental and practical importance to understand the intercalation mechanism and improve the electrochemical performance. In this section, the mechanism of K-intercalation into graphite is explained and compared to that of Li-intercalation. Graphite is composed of stacked graphene layers, where the carbon atoms are planarly bonded with sp2 hybrid bonds in six-ring configurations (Fig. 18A). The graphene layers are held by weak van der Waals interactions and are stacked in an ABAB sequence along the c-axis with an interlayer distance of 3.354 Å (Fig. 18A). When alkali metal ions are intercalated into the interlayer space, the stacking sequence of the graphene layers changes into A|A for the intercalated layers,151

Fig. 17 Typical charge/discharge curves of the graphite electrode with a sodium polyacrylate (PANa) binder in a Li cell with 1 M LiPF6/EC dimethyl carbonate (DMC; 1:1 v/v; black curve), Na cell with 1.0 M NaPF6/EC:DEC (1:1 v/v; blue curve), and K cell with 1 M KFSA/EC:DEC (1:1 v/v; red curve). Reproduced from Ref. Kubota, K.; Dahbi, M.; Hosaka, T.; Kumakura, S.; Komaba, S., Towards K-Ion and Na-Ion Batteries as “Beyond Li-Ion”. Chem. Rec. 2018, 18(4), 459–479 with permission. Copyright, 2018 The Chemical Society of Japan & Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Electrode materials for K-ion batteries

105

forming “stage n” GICs, where the nth stage represents a number (n) of graphene layers between the two closest intercalated layers. In electrochemical Li intercalation, the Liþ ions are inserted between the graphene layers, and Li-GICs are formed with stage transformations through stages 4, 3, 2, and 1.148 In stage 1, all graphene layers are alternately stacked with Li layers, forming LiC6 at the end of the electrochemical reduction process by potential scanning down to 0.0 V vs. Liþ/Li.148,152–154 The interlayer distance in stage 1 is  3.7 Å, and all the graphene layers are stacked in the AA sequence (Fig. 18B). Similarly, Kþ ions are inserted into graphite, and stage 1 is formed at the end of the electrochemical reduction process in a K cell.2,25,26 The composition of the stage-1 K-GIC is KC8, of which the alkali-metal density and ordering structure in each alkalimetal layer are different from those of LiC6. The interlayer distance between the graphene layers becomes  5.3 Å,150 which is much longer than the 3.7 Å observed in LiC6. In contrast to the hexagonal symmetry of LiC6, KC8 has orthorhombic symmetry with ordered stacking of the K layers along the c-axis. It displays the stacking sequence AaAbAgAdA, where a, b, g, and d refer to potassium layers with the same in-plane distribution of the Kþ ions but a different stacking manner (Fig. 18C).150 Similar to stage 2–4 Li-GICs (Fig. 18D), Rüdorff and Schulze prepared discrete stoichiometric compositions of KC24, KC36, and KC48 by K-vapor synthesis and determined their stage structures as stages 2, 3, and 4, respectively.145 Nixon and Parry determined the stacking sequences of stages 2–4 (Fig. 18E) and confirmed the formation of the equilibrium stages in the stage sequence stage 4 / stage 3 / stage 2 / stage 1 under potassium vapor.151 Fig. 19A illustrates the deviated charge/discharge potential curves of a K||graphite three-electrode cell in comparison with those of a Li ||graphite cell, which were operated at room temperature at a low current density of 9.3 and 12.4 mA g 1, respectively. The graphite electrode in the K cell displayed at least three reduction peaks corresponding to the potential-plateaus at 0.28, 0.20, and 0.14 V vs. Kþ/K during the K intercalation process. On the other hand, in the K deintercalation process, at least three oxidation peaks were reversibly observed at 0.29, 0.46, and 0.58 V vs. Kþ/K. The one main-redox-peak 0.14 V/0.29 V vs. Kþ/K is different from those observed for Li intercalation,155 i.e., two main-redox-peaks of 0.10 V/0.12 V and 0.07 V/0.08 V vs. Liþ/Li (Fig. 19A). Thus, the formation of the stage structures of K-GICs different from those of Li-GICs is expected for electrochemical K intercalation. The structural evolution of graphite during K insertion/extraction was elucidated using operando in situ XRD by our group14 and others.156–158 Fig. 19B displays the operando XRD patterns of graphite in a K/graphite cell during K insertion.14 The XRD patterns

Fig. 18 Schematic illustrations of crystal structures of (A) graphite having a hexagonal lattice with the SG P63/mmc, (B) LiC6 having a hexagonal lattice with the SG P6/mmm, (C) KC8 having an orthorhombic lattice with the SG Fddd. Schematic illustrations of stages of (D) Li-GICs and (E) K-GICs based on the Daumas-Hérold model.

106

Electrode materials for K-ion batteries

Fig. 19 (A) Deviated charge/discharge curves of the graphite electrode with a PANa binder in a Li cell with 1 M LiFSA/EC:DEC (1:1 v/v; red curve) and K cell with 1 M KFSA/EC:DEC (1:1 v/v; blue curve) at a current density of 12.4 and 9.3 mA g 1, respectively. (B) Operando in situ XRD patterns for a graphite electrode in a non-aqueous K cell with 1 M KFSA/EC:DEC (1:1 v/v) during K insertion at the 3rd cycle. Reproduced from Ref. Kubota, K.; Dahbi, M.; Hosaka, T.; Kumakura, S.; Komaba, S., Towards K-Ion and Na-Ion Batteries as “Beyond Li-Ion”. Chem. Rec. 2018, 18(4), 459–479 with permission. Copyright, 2018 The Chemical Society of Japan & Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

reveal stage transformations through at least stages 3, 2, and 1 during electrochemical K intercalation into graphite, and reversible stage transformations are also confirmed during K deintercalation.157,158

7.04.3.1.2

Electrochemical properties of graphite

Enhancement of the electrochemical K-intercalation properties of graphite is important for its practical use as negative electrode material in KIBs. Previous studies indicate that the electrochemical performance of graphite electrodes highly depends on the binder, electrolytes, and the nature of graphite. Table 5 listed the reported electrochemical properties of graphite electrodes. Graphite generally has large crystalline domains and a small surface area of < 20 m2 g 1. Thus, high Coulombic efficiencies > 90% are expected for graphite electrodes, even at initial charge/discharge cycles, as achieved for Li ||graphite cells. However, the initial Coulombic efficiencies of K||graphite cells are often low, except for the > 80% results reported by our group2,11,107 and Beltrop et al.160 The low initial Coulombic efficiencies would originate from electrolyte decomposition until the formation of a stable SEI on graphite, and would lead to highly resistive electrodes resulting in capacity degradation during cycles. With regards to a binder, our group has demonstrated excellent cycle stabilities with higher Coulombic efficiencies (79% and 89%) in K||graphite cells, using sodium polyacrylate (PANa) and sodium carboxymethylcellulose (CMC) binders, respectively, compared to the 59% of the graphite electrode with a poly(vinylidene fluoride) (PVdF) binder.2 Beltrop et al. also used sodium carboxymethylcellulose as a binder and achieved 81.8% efficiency in the initial cycle and excellent cycle stability during subsequent cycles.160 This group also reported a potassium-containing ionic liquid as the electrolyte.160 Electrolyte solutions are also known to influence the Coulombic efficiency and cycle stability of K/graphite cells.14 To date, high Coulombic efficiencies and excellent cycle stabilities have been demonstrated using KFSA or KTFSA-based electrolyte solutions such as 1 M KFSA/EC:DEC (1:1 v/v)2; N-butylN-methyl bis(trifluoromethanesulfonyl) amide (Pyr14TFSA) þ 0.3 M KTFSA þ 2 wt% ethylene sulfite (ES) and 7 mol kg 1 KFSA/ DME,107 and KFSA:ethylmethylcarbonate (EMC) (1:2.5, molar ratio).158 Electrolyte solutions of KPF6/EC:DEC (1:1 and 3:7 v/v) and KPF6/EC:PC (1:1 v/v) could be acceptable to identify the electrochemical activities of carbon materials in K-metal half cells159; however, high resistance was observed for a K ||graphite cell filled with 0.8 KPF6/EC:DMC (1:1 v/v) electrolyte solution.161 Moreover, low resistance of a K||graphite cell was observed with 1 M KPF6/DME as the electrolyte solution.163 The structures and compositions of GICs formed during the charge/discharge reactions in ester-based electrolytes are however different from those observed for K-GICs with carbonate-based electrolytes and potassium-vapor synthesis. Only Kþ ions are reversibly intercalated into graphite with the formation of K-GIC stages in highly-concentrated ester-based electrolytes such as 7 mol kg 1 KFSA/DME.107 On the other hand, Kþ ions are co-intercalated with the electrolyte solvent into graphite in conventional concentrations of ester-based electrolytes such as 1 M potassium trifluoromethanesulfonate (KOTf)/diethylene glycol dimethyl ether (diglyme)166 and 1 M KPF6/DME.163 In addition to the binder and electrolyte, the structural properties of graphite affect the properties. Fig. 20 compares the cycle stabilities of the reversible capacities in K||graphite cells composed of different interlayer distances of graphite samples, examined by our group in 2020.167 Graphite with a small interlayer distance (d002) of 3.355 Å exhibited the best cycle stability and the cycle retention declined in the interlayer distance order 3.355, 3.360, and 3.366 Å. These results suggested that well-crystallized graphite having an average small d002 is suitable for reversible K intercalation at room temperature. It should be noted that the slightly larger capacity than the theoretical value of 279 mAh g 1 would originate from the electric conductive agent carbon black, which was added (6 wt%) in the graphite electrodes and the capacity contribution was not counted in the gravimetric capacities of the graphite

Electrode materials for K-ion batteries Table 5

107

Electrochemical properties of K||graphite cells reported in the literature.

Product name of graphite

1st reversible capacity, Q (QMax.)/ mAh g 1

1st Coulombic efficiency/%

Cycle stability

TIMREX SLP50 (Imerys Graphite & Carbon)25 SNO3 (SCE carbon)2 -26 TIMREX KS4 (Imerys Graphite & Carbon)159 -156 TIMREX KS6L (Imerys Graphite & Carbon)160 SNO3 (SCE carbon)107

273

57.5

50.8% after 50 [email protected] A g 1 0.8 M KPF6/EC:DEC

243 (247) 207 246

58.5–88.9 73.9 66.5

99.9% after 40 cycles@25 mA g 1 1 M KFSA/EC:DEC – 0.5 M KPF6/EC:DEC 89.4% after 1 M KPF6/EC:PC or EC:DEC 200 cycles@20 mA g 1

249 (274) 270

47.8 82

12.4% after 50 [email protected] mA g 1 1 M KPF6/EC:DMC 100% after 50 cycles@50 mA g 1 Pyr14TFSA þ 0.3 M KTFSA þ2 wt% ES

259 (263)

81.5

Synthetic graphite (MTI)161 -158

300

65

170(255)

62

99.9% after 300 cycles@25 mA g 1 90.9% after 100 cycles@50 mA g 1 > 90% after 2000 cycles@93 mA g 1 96.9% after 50 [email protected]% after 700 cycles@280 mA g 1 98.4% after 100 cycles@25 mA g 1 94.7% after 1000 cycles@558 mA g 1

-162 239 Graphite (Iopsilion)163 135

57.4 69.6

SNO3 (SCE carbon)164 257 (261)

88.7

SNO3 (SCE carbon)165 260

88.6

Electrolyte

7 mol kg 1 KFSA/DME 0.8 M KPF6/EC:DEC KFSA: EMC (1:2.5, molar ratio) 0.8 M KPF6/EC:DEC 1 M KPF6/DME 1 mol kg 1 (KPF6)0.9(KFSA)0.1/ EC:DEC 1 M K[FSA]/[C3C1pyrr][FSA]

electrodes. Carbon black is an important electrode additive acting as an electric conductive agent between the particles of the active material powder. In particular, because the interlayer distance is enlarged 1.6-fold to form stage 1 KC8 by electrochemical K intercalation, carbon black addition enhances the cycle stability and rate performance of a graphite electrode. Optimization of the electrode composition is of high importance to achieve a good cycle life, high Coulombic efficiency, and good rate capability of a graphite electrode for practical KIBs.

Fig. 20 Capacity vs. cycle number plots of graphite electrodes at the graphite:carbon black (Super C45, Imerys Graphite & Carbon):CMC weight ratio 89:6:5. The graphite samples used were of different interslab distances, d002. Reproduced from Ref. Hosaka, T.; Kubota, K.; Hameed, A. S.; Komaba, S., Research Development on K-Ion Batteries. Chem. Rev. 2020, 120(14), 6358–6466 with permission. Copyright, 2020 American Chemical Society.

108

Electrode materials for K-ion batteries

7.04.3.1.3

Hard and soft carbon

Graphite is not electrochemically active for sodium insertion, despite the high capacity in the Li and K cells. On the other hand, the hard carbon electrode displayed similar electrochemical activity as those of the Li, Na, and K insertion hosts. Hard carbon has small domains of sp2 carbon layers with a large interlayer distance and micropores inside the particles168 and Liþ and Naþ ions are known to be accommodated into both the interlayer and micropore spaces.169 The charge/discharge performance of a commercially available hard carbon developed for practical LIBs, Carbotron P(J) (Kureha Corporation), was examined in Li, Na, and K cells by our group (Fig. 21A).14 The Li, Na, and K cells presented sloping curves during the charging (insertion) process, followed by plateau regions located at  0.05 V vs. Liþ/Li, 0.1 V vs. Naþ/Na, and 0.3 V vs. Kþ/K, respectively. Li and K cells deliver almost the same reversible capacities of  215 mAh g 1. On the other hand, the reversible capacity of the Na cell, which is approaching 256 mAh g 1, is higher than those of the Li and K cells. When hard carbon is exposed to air, the surface of hard carbon gradually reacts with CO2, O2, and moisture,170 resulting in large irreversible capacities at the initial cycle and large voltage-hysteresis between charge and discharge in the Li cells.171,172 The galvanostatic intermittent charge/discharge curves of hard carbon exposed to air for more than 3 months for the Li, Na, and K cells are compared in Fig. 21B–D. The Li cell displays large hysteresis between the charge and discharge open-circuit voltages, while the hysteresis in the Na cell is smaller and that in the K cell is negligible. These results suggest weaker interaction of the large Kþ ions with the functional groups formed by exposure to air, probably due to the lower Lewis acidity compared to those of the Liþ and Naþ ions.167 Low-voltage hysteresis and less air-sensitive behaviors are advantageous for the use of hard carbon for practical KIBs. The electrochemical K-insertion properties of hard carbon were first reported by Jian et al. in 2016173 and improved by using CMC as a binder in 2017.174 The group synthesized hard carbon spheres by hydrothermal reactions of sucrose at 195  C followed by thermal treatment at 1100  C under Ar atmosphere.174 The hard carbon electrode presented sloping curves in the range 1.25– 0.3 V vs. K metal and plateau regions at 0.2 and 0.33 V during the potassiation and depotassiation processes, respectively.174 The initial potassiation and depotassiation capacities were 344 and 260 mAh g 1, respectively, with an initial Coulombic efficiency of 76%, which was higher than the 49% observed when using a PVdF binder.173

(A)

(B)

2.0

Licell Nacell Kcell

2.0

discharge charge

1.0

1.0

0.5

0.5 0.0

0.0 0

50

100

150

200

250

0

300

50

100

150

200

250

Capacity / mAh g-1

Capacity / mAh g-1 (C)

(D)

2.0

discharge charge

2.0

Na cell

discharge charge

K cell

1.5 voltage / V

1.5 voltage / V

Li cell

1.5

1.5

voltage / V

Potential / V vs. A+/A

2.5

1.0

0.5

1.0

0.5

0.0

0.0 0

50

100

150

Capacity / mAh g

200 -1

250

0

50

100

150

200

250

-1

Capacity / mAh g

Fig. 21 (A) Galvanostatic charge/discharge curves of the hard carbon [Carbotron P(J), Kureha Corporation] electrodes in Li, Na, K cells at the 2nd cycle. Galvanostatic intermittent charge/discharge curves of the hard carbon electrodes in the (B) Li, (C) Na, and (D) K cells. The hard carbon sample was exposed to air for more than 3 months. Reproduced from Ref. Hosaka, T.; Kubota, K.; Hameed, A. S.; Komaba, S., Research Development on KIon Batteries. Chem. Rev. 2020, 120(14), 6358–6466 with permission. Copyright, 2020 American Chemical Society.

Electrode materials for K-ion batteries

109

As hard carbon has been extensively studied in NIBs, K-insertion properties have been examined for hard carbon materials derived from rice-starch,74 oak,175 waste-tire rubber,176 cellulose,177 skimmed cotton,178 potatoes,179 loofahs,180 and maple leaves.181 For example, Vaalma et al. synthesized hard carbon by the thermal decomposition of starch powder from rice (SigmaAldrich) at 503 K in air, followed by heating at 1300 K under argon flow (Fig. 22A).74 They also examined the influence of the addition of an electric conductive agent of carbon black (Super C65, Imerys Graphite & Carbon). The pure hard carbon electrode composed of 90 wt% hard carbon and 10 wt% PVDF (HC in Fig. 22B) and the composite electrode of hard carbon and carbon black (70 wt% HC, 20 wt% C65, and 10 wt% PVDF; HC þ CB in Fig. 22B), deliver the same reversible capacities of 200 mAh g 1 with an initial Coulombic efficiency of  50%. However, the composite electrode exhibits a better rate performance and cycle stability with capacities of 200 mAh g 1 at 0.1C after 50 cycles and 54 mAh g 1 at 10C (1C ¼ 279 mA g 1 in the literature).74 Carbon black seems to be necessary in the hard carbon electrodes as an electric conductive agent to enhance the rate capabilities and cycle stabilities of the K cells. Our group also studied the electrode performances of hard carbon in Na and K cells using hard carbons of different turbostratic stacked sheets and nano-sized pores.177,182 We demonstrated a reversible Na-insertion capacity of 350 mAh g 1 for the hard carbon prepared by the pre-treatment of cellulose at 275  C and post-treated at 1300 or 1500  C.183 Thus, we next discuss the influence of post-treatment temperatures (700–1500  C) for cellulose precursors pre-treated at 275  C. The charge and discharge curves of the cellulose-derived hard carbon electrodes (post-treatment at 700–1500  C) with a 5 wt% PANa binder are represented in Fig. 22C.177 The reversible K insertion capacities increased from 50 to 290 mAh g 1 with increasing post-treatment temperatures. The optimal hard carbon synthesized by pre-heat treatment and post-treatments of cellulose at 275  C and 1500  C, respectively, demonstrated a high initial Coulombic efficiency of 83% as well as a high reversible capacity, which was higher than the theoretical capacity of graphite (279 mAh g 1). Furthermore, good retention and higher Coulombic efficiencies

Fig. 22 (A) SEM image of a rice starch-derived hard carbon (HC) at 2kx. (B) Cycle performance of the HC, CB, and HC/CB (ratio 70:20) electrodes with 1.5 M KFSA/EC:DMC in K-metal three-electrode cells in the potential range 0.02–2.0 V vs. Kþ/K. Cycles 11–40 correspond to the C-rate test (five cycles each at 0.2, 0.5, 1, 2, 5, and 10C). (C) Charge/discharge curves of cellulose-derived hard carbons (pretreatment at 275  C and post-treatment at 700–1500  C) with 5 wt% PANa binder in K cells tested at a current rate of 25 mA g 1 under constant current–constant voltage (CC-CV) potassiation and (D) capacity retention of the hard carbon prepared at 275  C and 1300  C in the K cell. (A, B) Reproduced from Ref. Vaalma, C.; Giffin, G. A.; Buchholz, D.; Passerini, S., Non-Aqueous K-Ion Battery Based on Layered K0.3MnO2 and Hard Carbon/Carbon Black. J. Electrochem. Soc. 2016, 163(7), A1295-A1299 with permission. (C, D) Reproduced from Ref. Yamamoto, H.; Muratsubaki, S.; Kubota, K.; Fukunishi, M.; Watanabe, H.; Kim, J.; Komaba, S., Synthesizing Higher-Capacity Hard-Carbons from Cellulose for Na- and K-Ion Batteries. J. Mater. Chem. A 2018, 6(35), 16844–16848 with permission. Copyright, 2018 The Royal Society of Chemistry.

110

Electrode materials for K-ion batteries

> 99.5% were observed over 50 cycles (Fig. 22D). Compared to the graphite electrode, a hard carbon electrode has the potential to deliver larger reversible potassiation capacities. The history of the development of higher capacity hard carbon over the past decade184 indicates the creation of a new hard carbon material with increased reversible capacities in K cells for practical use in KIBs within the next decade. In the hard carbon materials for Na insertion, Kano et al. selected meso- and micro-porous carbon as a precursor and prepared hard carbon via high-temperature treatment at 2100–2400  C.185 The most porous hard carbon delivered a significantly large reversible capacity of 438 mAh g 1 with a high Coulombic efficiency of 93% at the initial cycle of a Na cell. In 2020, we demonstrated a higher Na-insertion capacity of 478 mAh g 1 for a hard carbon electrode prepared by heating a freeze-dried mixture of magnesium gluconate and glucose by a MgO-template technique.186 In both cases, the charge/discharge curves displayed a small capacity in the sloping region and a large capacity in the long plateau region at  0.1 V vs. Naþ/Na. The suitably sized micropores would enable high-density accommodation of Naþ ions in the structure. In 2019, we studied the influence of the synthesis temperature of macroporous phenolic resin-derived hard carbon on the structure and electrochemical properties in Na and K cells in order to clarify the relationship between the microstructure and electrochemical properties.187 As the heat-treatment temperature is raised from 1100 to 1500  C, the interlayer distance of stacked carbon sheets decreased, and the nanosized void size increased from XRD and SAXS results. The reversible capacity in both Na and K cells increases as the heat-treatment temperature rises. Hard carbon prepared at 1500  C in argon delivers the highest capacity of 336 mAh g 1 in the K cell (Fig. 23).

7.04.3.2

K Alloys and other potassiatable compounds

In this section, we summarize K alloy-related electrode materials, such as KeMe alloying materials and P. Group 14 and 15 elements, including metals (Sn, Pb, and Bi), metalloids (Si, Ge, As, and Sb), and polyatomic nonmetals (P), are known to form binary compounds with K, as is observed with sodium.47 These electrode materials have recently been studied as potential negative electrodes for KIBs because they tend to accommodate a larger amount of K and provide much higher capacities than carbonaceous materials.28

7.04.3.2.1

Alkali metal alloy materials and compounds for Li-, Na-, and K-ion batteries

Alloying materials and binary compounds have been extensively studied as electrode materials for LIBs and NIBs. First, we describe the theoretical specific capacity, volume capacity, operating potential, and volume change of K-alloys and binary compounds in comparison with the Li and Na counterparts. Fig. 24A compares the theoretical capacities, both specific and volumetric, of selected negative electrodes for Li, Na, and K systems. Al exhibits electrochemical alloying to form b-LiAl in Li cells,188,189 while Al does not present an electrochemical alloying reaction in Na and K cells.14 Thus, an Al electrode would never function as an active material for NIBs and KIBs. Al foil can instead be used as a negative electrode current collector for NIBs and KIBs, unlike LIBs.3,14,47 Si is a promising negative electrode material for LIBs because the electrochemical lithiation of Si to form Li15Si4 (Li3.75Si) displays a theoretical capacity of 3580 mAh g 1.188 On the other hand, NaSi and KSi are known to be the most alkali-metal-rich silicide phases of binary compounds, delivering a theoretical capacity of 954 mAh g 1.190 However, to the best of knowledge, the electrochemical formation of NaSi and KSi in Na and K cells has not been reported to date. Other group 14 elements (Ge, Sn, and Pb) are known to form 1:1 compounds as the most K-rich phases in binary compounds, delivering theoretical specific capacities of 369, 226, and 129 mAh g 1, respectively.190 These theoretical capacities are lower than their Li counterparts. In the case of the group 15 elements,

Fig. 23 Charge and discharge curves of hard carbon electrodes in K cells filled with 1 mol dm 3 KFSA/EC:DEC (1:1 v/v) tested at a current rate of 10 mA g 1. Reproduced from Ref. Kamiyama, A.; Kubota, K.; Nakano, T.; Fujimura, S.; Shiraishi, S.; Tsukada, H.; Komaba, S., High-Capacity Hard Carbon Synthesized from Macroporous Phenolic Resin for Sodium-Ion and Potassium-Ion Battery. ACS Appl. Energy Mater. 2019, 3(1), 135–140 with permission. Copyright, 2018 American Chemical Society.

Electrode materials for K-ion batteries

111

Fig. 24 (A) Theoretical specific and volumetric capacities of selected negative electrode materials for Li, Na, and K cells. The volumetric capacities were calculated at full potassiation. Note that the electrochemical formation of NaSi and KSi has not been reported though these compounds are found in the binary phase diagram. Electrochemical formation of K3P has not been reported yet while some KxP compounds including K4P3 and KP have been reported in K cells as K-rich compounds. (B) Volume expansion upon K alloying with different metals as a function of x in KxM. The NaeSb and LieSb systems are also plotted for comparison. Reproduced from Ref. Hosaka, T.; Kubota, K.; Hameed, A. S.; Komaba, S., Research Development on K-Ion Batteries. Chem. Rev. 2020, 120(14), 6358–6466 with permission. Copyright, 2020 American Chemical Society.

K3Sb and K3Bi are known as K-rich phases, which are identical to their Li and Na counterparts, with calculated capacities of 660 and 385 mAh g 1, respectively. P has the stable phase A3P in all the Li-, Na-, and KeP binary compounds, with a high specific capacity of 2597 mAh g 1.191 On the other hand, DFT calculations indicated unfavorable formation energy for K3P, suggesting that the full potassiation phase in K cells is K4P3 or KP, with specific capacities of 1154 and 865.5 mAh g 1, respectively. From these data, group 15 elements such as Sb, Bi, and P are expected to achieve specific capacities that are competitive with those of their Li and Na counterparts. The volumetric energy density of the electrode active materials should be discussed based on the negative electrode volume at full potassiation, i.e. after volume expansion. This is because the battery volume regulates the actual battery capacity, as previously discussed by Obrovac et al.192 and our group.191 In this study, we calculated the volumetric capacities of the negative electrodes based on the volume at the fully charged state. Fig. 24A displays the volumetric capacities based on the volume after a full charge. In general, K compounds tend to show large molar volumes due to the large ionic radius of the Kþ ions; therefore, the volumetric capacities of K-alloy materials are small compared to those of Li- and Na-alloy materials. For example, although Sb presented the same theoretical gravimetric capacity for the Li, Na, and K cells, its theoretical volumetric capacities were 1890, 1120, and 760 Ah L 1 for the Li, Na, and K cell, respectively. Fig. 24B illustrates the relative volume changes of the K-M systems (M ¼ Pb, Sb, Bi, Sn, and red P) with respect to the original volume of M as a function of the K content. The NaeSb and LieSb systems were also plotted for comparison. The volume of all the materials monotonically increased with increasing K content, and no significant difference in volume change was observed between these materials. The K1M compounds presented volume expansions in the range 200–250%, while the K3M compounds displayed

112

Electrode materials for K-ion batteries

volume expansions of  500%. Clearly, the volume change in the K-M systems was significantly more pronounced than that of the NaeSb and LieSb systems. The large volume change would destruct the electrode structure and cause side reactions such as continuous decomposition of the electrolyte, thereby leading to capacity degradation. To overcome these problems related to large volume changes, several strategies such as the use of nanostructured electrodes193,194 and modified electrolytes195 have been developed as further discussed in later sections. In the next section, the material properties are reviewed according to their group, i.e., group 14 (Si, Ge, Sn, and Pb) and group 15 (P, Sb, and Bi) elements.

7.04.3.2.2

Group 14 elements and compounds

This section focuses on the electrochemical performances of some group 14 elements such as Si and Sn in a K cell. Sultana et al. described the electrochemical properties of a Si/graphene composite electrode in a K cell.28 The Si/graphene electrode showed reversible capacity of only 150 mAh g 1 in K cell, while the same electrode displayed a reversible capacity > 2000 mAh g 1 in the Li cell. Moreover, obtained linear charge/discharge curves indicated that the capacitive capacity of graphene would mainly contribute to the reversible capacity, and ex situ XRD data indicated the absence of KeSi compounds. Therefore, to date, there is no evidence of electrochemical potassiation in K cells. The electrochemical properties of Sn as a K-ion negative electrode was first reported by Sultana et al. in 2016.196 In 2017, the electrochemical alloying of K and Sn was studied in more detail by Ramireddy et al., using a 1 mm Sn film electroplated on a copper foil.197 The charge/discharge curves of the Sn film (Fig. 25A) reveal a reversible capacity of 245 mAh g 1, suggesting the formation of the KSn phase. In situ synchrotron XRD was conducted during cyclic voltammetry (Fig. 25B). Upon potassiation, the Sn peak

Fig. 25 (A) Charge/discharge curves of 1 mm thick Sn film in a K cell at a current density of 25 mA g 1, (B) in situ XRD patterns corresponding to the region associated with the phase transformation of primary interest during electrochemical potassiation of Sn. (C) Initial charge/discharge curves and (D) cycle performance of Sn-based electrodes in 0.5 M KPF6/EC:DEC at current density of 25 mA g 1. (A–B) Reproduced from Ref. Ramireddy, T.; Kali, R.; Jangid, M. K.; Srihari, V.; Poswal, H. K.; Mukhopadhyay, A., Insights into Electrochemical Behavior, Phase Evolution and Stability of Sn upon K-Alloying/De-Alloying via In Situ Studies. J. Electrochem. Soc. 2017, 164(12), A2360-A2367 with permission. Copyright, 2017 Electrochemical Society. (C, D) Reproduced from Ref. Shimizu, M.; Yatsuzuka, R.; Koya, T.; Yamakami, T.; Arai, S., Tin Oxides as a Negative Electrode Material for Potassium-Ion Batteries. ACS Appl. Energy Mater. 2018, 1(12), 6865–6870 with permission. Copyright, 2018 American Chemical Society.

Electrode materials for K-ion batteries

113

intensity decreased continuously, whereas the reverse trend was observed for KSn. This suggested a two-phase reaction, from b-Sn to KSn, without an intermediate phase during potassiation. During subsequent depotassiation, no intermediate phase was observed; this was inconsistent with the fact that at least two plateaus were confirmed in the charge/discharge curves in Fig. 25A. This contradiction may be due to the formation of an amorphous and/or nano-crystalline intermediate phase. Actually, Wang et al. have discovered that amorphous K4Sn9 is formed as an intermediate phase before the formation of partially crystalline KSn as a fully potassiated phase using in situ TEM measurements.198 Moreover, Gabaudan et al. clarified that the formation of KSn at the end of the potassiation and that of the intermediate phase of K4Sn9 during potassiation using DFT calculation and 119Sn Mössbauer spectroscopy.199 Although these studies successfully demonstrated reversible potassiation/depotassiation of the Sn electrodes, the electrodes displayed severe capacity degradation. This was attributed to the large volume expansion of 210%, which resulted from the formation of the KSn phase. Therefore, some groups reported a modified electrode design, such as nanocomposite electrodes, to achieve a practical cycle performance by suppressing the loss of electrical contact and crack formation.200,201 Another strategy to improve the cycle performance is the utilization of SnO as the electrode material instead of Sn.202 The SnO electrode presented a reversible capacity of  230 mAh g 1, while SnO2 displayed a significantly lower electrochemical activity (Fig. 25C). Ex situ XRD data suggested that the metal Sn formed in the initial reduction contributed toward the reversible capacity. This synthetic process occurred via a conversion method and accompanied the formation of the KSn alloy. The alloying reaction of Sn to KSn is reversible, in contrast with the irreversible conversion reaction. Moreover, regeneration of SnO rarely occurred under upper cutoff voltages of 2.0 and 3.0 V during the depotassiation process. Therefore, the charge/discharge mechanism of SnO is very similar to that of Sn, except for the initial irreversible conversion reaction. However, the SnO electrode displayed better capacity retention with a reversible capacity of 183 mAh g 1 after 30 cycles than the Sn electrode (Fig. 25D). The good cycle performance of the SnO electrode occurred because the K2O matrix, irreversibly formed in the initial cycle, acts as a buffer to suppress the aggregation of Sn nanoparticles and relieve the stress caused by the large volume change observed in Sn. This is similar to the electrochemical lithiation of SnO203 and SiO.204

7.04.3.2.3

Group 15 elements and compounds

The group 15 elements Sb, Bi, and P are potential negative active materials, owing to their higher theoretical capacities over those of the group 14 elements. McCulloch et al. first published a paper on the electrode performance of Sb in K cells in 2015.205 The Sb/C nanocomposite electrode presented a reversible capacity of 650 mAh (g-Sb) 1, which corresponds to the theoretical capacity (660 mAh g 1) based on the formation of K3Sb (Fig. 26A). It should be noted that the carbon material in the composite material and the electrode are electrochemically active, and it may contribute up to 110 mAh g 1 for the capacity according to the control test of carbon electrode.205 Fig. 26B illustrates the operando XRD patterns of Sb during the first cycle in a K cell reported by Gabaudan et al.206 The operando XRD patterns revealed that Sb was alloyed with K to form K3Sb without an intermediate phase during the potassiation process. Moreover, the major phase was cubic K3Sb with a minor phase of hexagonal K3Sb. The charge/discharge mechanism of Sb is schematically summarized in Fig. 26C. Although antimony is one of the promising K alloy materials for application as a negative electrode for KIBs, its extensive volume change of  500% is a critical challenge that hinders practical cycle performance. As was observed for other K alloy materials, the design of electrode materials, such as nanoporous antimony,207 and Sb@C composite electrodes,208–217 has been reported to improve their cycle performance. Huang et al. synthesized tin-antimony alloy (SnSb)-graphene-carbon (SnSb-G-C) nanofibers with a porous multichannel structure via an electrospinning method and investigated their electrode performance in K cells.218 The SnSb-G-C electrodes exhibited a good cycling performance of 275 mAh g 1 over 100 cycles, which was better than those of its SnSb-carbon and graphene-carbon counterparts. This good cycle performance was attributed to the nanochannels and dispersive graphene inside the carbon matrix, which effectively buffer the structural changes and protect the SnSb particles from destruction. Bi is a potential negative electrode because of the stability of the K3Bi phase and satisfactory high theoretical specific and volumetric capacities, as previously mentioned. The electrochemical alloying of Bi in K cells was reported in 2018.195,219,220 Bulk Bi electrode presented a reversible capacity of  400 mAh g 1. Huang et al. reported a reversible dealloying/alloying process from K3Bi to Bi, via the two intermediate phases K3Bi2 and KBi.219 The estimated volume expansion of the Bi electrode to form K3Bi is as high as 500%. Actually, Huang et al. observed that the K||Bi cell exhibited rapid capacity decay to half of the initial capacity in six cycles, while the capacity was almost 0 mAh g 1 after 10 cycles when KPF6/EC:PC was used as the electrolyte solution.219 On the other hand, the addition of fluoroethylene carbonate (FEC) into the electrolyte suppressed the capacity decay, and the electrode displayed a capacity > 100 mAh g 1 after 10 cycles. Moreover, stable cycle performance was achieved by changing the electrolyte solvent from EC:PC to glymes such as DME and diglyme,195,219,220 a similar tendency as that observed for the NaeBi system.221 Lei et al. reported that the Bi electrode became a porous structure during cycling in the KPF6/DME unlike KPF6/EC:PC electrolyte. This porous structure should contribute to the good cycle performance.220 The formation of the porous structure is explained by the transfer of surface Bi atoms via the strong chemical adsorption of DME molecules on the Bi.220 The lighter group-15 element P is another potential negative electrode material with potential to realize high energy density KIBs. Phosphorus-based materials such as black phosphorus, red phosphorus, Sn4P3, and blue phosphorene have been investigated as negative electrodes for KIBs. Sultana et al. investigated black phosphorus as a negative electrode material for KIBs.222 Fig. 27A displays the charging/discharging of a black P@C composite electrode in a K cell. The composite electrode exhibited a high reversible capacity of 617 mAh g 1. The ex situ XRD results (Fig. 27B) suggested electrochemical reaction of the phosphorus electrode to form KP as the final product.222 Zhang et al. examined the negative electrode performances of red P@C and Sn4P3@C composites, which were synthesized by a facile ball-milling method.223 Fig. 27C and D respectively compare the charge/discharge curves and

114

Electrode materials for K-ion batteries

Fig. 26 (A) Charge/discharge curves of the Sb@C electrode in a K cell. (B) Operando XRD patterns of Sb electrodes recorded at C/10 rate. The patterns are recorded during the first charge (black) and discharge (red). The Be window used in the in situ electrochemical cell creates an intense (100) Bragg peak at 45.8 (together with three minor peaks at 35.7, 40.6, and 41.9 due to the impurities included in the beryllium sample). (C) Crystal structures of the phases observed during operando XRD for Sb. (A) Reproduced from Ref. McCulloch, W. D.; Ren, X.; Yu, M.; Huang, Z.; Wu, Y., Potassium-Ion Oxygen Battery Based on a High Capacity Antimony Anode. ACS Appl. Mater. Interfaces 2015, 7(47), 26158–26166 with permission. Copyright, 2015 American Chemical Society. (C) Reproduced from Ref. Gabaudan, V.; Berthelot, R.; Stievano, L.; Monconduit, L., Inside the Alloy Mechanism of Sb and Bi Electrodes for K-Ion Batteries. J. Phys. Chem. C 2018, 122(32), 18266–18273 with permission. Copyright, 2018 American Chemical Society.

cycle performances of Sn4P3@C with those of P@C and Sn@C. The Sn4P3@C electrode delivered a reversible capacity of 384.8 mAh g 1 and better cycle performance than that of the red P@C electrode (Fig. 27D). Ex situ XRD of K-inserted Sn4P3 indicated that the KeSn (K4Sn23, KSn) and KeP (K3–xP) phases were formed electrochemically. The good cycle performance would occur because these phases act as mutual buffers to alleviate the volume changes during cycling. Furthermore, carbonphosphorus composite materials have also been actively developed.213,224–228 Various K alloys and compounds have been reported for their potential application as negative electrodes, some of which (e.g., Sb, Bi, and P) have demonstrated high reversible capacities and moderately low operating potentials in K cells. Therefore, alloying negative electrodes are attractive as negative electrode materials to realize high energy density KIBs. However, the larger volume expansion during the potassiation process over that observed during the lithiation and sodiation processes resulted in pulverization of the active materials, detachment of the composite layer, and continuous electrolyte decomposition/SEI formation. Thus, the realization of practical cycle performance is still challenging for alloying materials. Previous intensive studies successfully highlighted that engineering of the electrode structure and optimization of the electrolyte are highly effective in demonstrating better cycle performances.

Electrode materials for K-ion batteries

115

Fig. 27 (A) Charge/discharge curves of the black phosphorus (BP)@C 7:3 electrode at a current density of 50 mA g 1 and (B) XRD pattern of the BP@C 1:1 electrode after the first charge and discharge. (C) Charge/discharge curves of the Sn4P3@C electrode at a current density of 50 mA g 1 and (D) cycle performance of the Sn4P3@C, Sn@C, and P@C electrodes. (A–B) Reproduced from Ref. Sultana, I.; Rahman, M. M.; Ramireddy, T.; Chen, Y.; Glushenkov, A. M., High Capacity Potassium-Ion Battery Anodes Based on Black Phosphorus. J. Mater. Chem. A 2017, 5(45), 23506–23512 with permission. Copyright, 2017 The Royal Society of Chemistry. Reproduced from Ref. Zhang, W.; Mao, J.; Li, S.; Chen, Z.; Guo, Z., PhosphorusBased Alloy Materials for Advanced Potassium-Ion Battery Anode. J. Am. Chem. Soc. 2017, 139(9), 3316–3319 with permission. Copyright, 2017 American Chemical Society.

7.04.3.3

Transition metal oxides as negative electrode materials

In the former sections, we have reviewed carbonaceous and alloying materials compatible as negative electrodes with low redox potentials. Generally, a low redox potential is advantageous to realize the high energy density required by batteries but raises concern on safety issues relating to metal plating. Therefore, researchers have been developing negative electrode materials with moderately low operation potentials. In this context, transition metal oxides have been actively studied as negative electrode materials for LIBs, NIBs, and KIBs. Li4Ti5O12 (LTO) is the most popular transition metal oxide as a negative electrode material for LIBs. LTO exhibits highly reversible charge/discharge curves with a voltage plateau at  1.5 V vs. Liþ/Li.229 The relatively high operating potential realizes highly safe LIBs at the expense of the energy density. This is because lithium dendrite formation can be satisfactorily avoided even at low temperatures and high-power operation, as compared to those of the graphite negative electrode. Therefore, the development of negative electrode materials operating in the moderate potential range 1–1.5 V vs. Kþ/K is expected to demonstrate safe and high-power KIBs. In this section, we provide an overview of the development of transition metal oxides.

7.04.3.3.1

Ti, Mo, and Nb oxides

Various titanium oxides such as K2Ti4O9,31,230,231 K2Ti8O17,29 and K2Ti6O1330,232 have been studied as negative electrode materials for KIBs as well as Li and Na systems. K2Ti4O9 is known to have a layered structure with an interlayer gap suitable to accommodate Kþ ions (Fig. 28A).233 K2Ti4O9 was studied in Li230 and K31 cells. In 2016, Kishore et al. first reported that a bulk K2Ti4O9 electrode

116

Electrode materials for K-ion batteries

exhibited reversible Kþ insertion/extraction in K cells, delivering an initial discharge capacity of 97 mAh g 1 at a current density of 30 mA g 1.31 One year later, Dong et al. synthesized K2Ti4O9 nanoribbons by the simultaneous oxidation and alkalization of MXene (Ti3C2).231 Fig. 28B displays the HRTEM image of the MXene-derived K2Ti4O9, which offers ultrathin thickness (< 11 nm), narrow nanoribbon widths (< 60 nm), and open macroporosity. Fig. 28C illustrates the charge/discharge curves of MXene-derived K2Ti4O9 at different current densities. The MXene-derived K2Ti4O9 electrode displayed a high initial reversible capacity of 151 mAh g 1 at 50 mA g 1 and good capacity retention compared to those of bulk K2Ti4O9. The better electrochemical performance was likely attributed to the nanostructure. In addition, the MXene-derived K2Ti4O9 presented a good rate capability maintaining  75 mAh g 1 at 300 mA g 1 (Fig. 28C), which is superior to that of the K2Ti4O9 synthesized by the conventional solid-state method.231 Although some previous studies did not take into account the capacity of the conductive agent, it is necessary to evaluate the contribution of conductive carbons to the capacity because conductive carbons such as AB are electrochemically active at potentials as low as 0–1 V. In addition to K2Ti4O9, layered K2Ti8O1729 and K2Ti6O1330,232 was also reported as a possible negative electrode material for KIBs. The reversible capacity and charge discharge curves of K2Ti8O17 and K2Ti6O13 are quite similar to K2Ti4O9, and nanoparticle is essential to obtain a sufficient capacity (> 100 mAh g 1). Some transition metal oxides such as MoO2234,235 and Nb2O5236 have also been reported as negative electrode materials based on the Kþ insertion reaction. The electrochemical performance of MoO2 was investigated by computational235 and experimental work.234 Liu et al. synthesized MoO2-rGO composites via a hydrothermal method followed by a heat treatment process, and its electrochemical performance was compared to that of bulk MoO2. The as-prepared MoO2-rGO composite electrode delivered initial charge and discharge capacities of 711 and 367.2 mAh g 1, respectively, in the wide potential range 0–3 V. The MoO2-rGO composite electrode displayed a reversible capacity of 218.9 mAh g 1 after 200 cycles, whereas lower reversible capacities of 146.1 and 119.6 mAh g 1 were reported after 200 cycles for bulk MoO2 and rGO, respectively.

Fig. 28 (A) Crystal structure of K2Ti4O9, (B) HRTEM images of MXene-derived K2Ti4O9, (C) charge/discharge curves cycled at different current densities of MXene-derived K2Ti4O9. (A) Reproduced from Ref. Kishore, B.; Munichandraiah, N., K2Ti4O9: A Promising Anode Material for Potassium Ion Batteries. J. Electrochem. Soc. 2016, 163(13), A2551-A2554 with permission. Copyright, 2016 Electrochemical Society. (B–C) Reproduced from Ref. Dong, Y.; Wu, Z.-S.; Zheng, S.; Wang, X.; Qin, J.; Wang, S.; Shi, X.; Bao, X., Ti3C2 MXene-Derived Sodium/Potassium Titanate Nanoribbons for High-Performance Sodium/Potassium Ion Batteries with Enhanced Capacities. ACS Nano 2017, 11(5), 4792–4800 with permission. Copyright, 2017 American Chemical Society.

Electrode materials for K-ion batteries

117

Orthorhombic niobium pentoxide (T-Nb2O5), which possesses a layered structure, was also reported as a negative electrode for KIBs.236 Li et al. synthesized T-Nb2O5 nanowires (diameter of 20–40 nm) via a hydrothermal approach.236 T-Nb2O5 electrode presented a reversible capacity of 152 mAh g 1 in a K cell based on the reversible electrochemical redox of Nb4 þ/5 þ. Overall, the titanium oxides, K2Ti4O9, K2Ti8O17, and K2Ti6O13, MoO2, and Nb2O5 have achieved acceptable reversible capacities of potassium insertion/extraction and good cycle performances. However, these oxides showed slopy charge/discharge curves in the wide potential range including  1 V, which requires protection with a passivation layer (SEI layer), resulting in complicated electrochemical properties compared to those of LTO, which require little SEI passivation. Although a lower working potential of the negative electrode favors a higher energy density, the development of oxide negative electrodes with a higher working potential of  1.6 V vs. Kþ/K, corresponding to 1.5 V vs. Liþ/Li is still a challenge to realize safe and high-power KIBs.

7.04.3.3.2

Transition metal oxides based on conversion reaction

In addition to insertion-type negative electrode material, transition metal oxides with conversion reactions237 have been studied as negative electrodes for KIBs. Sultana et al. first proposed mixed transition-metal oxides of Co3O4–Fe2O3, in which the electrode undergoes a typical conversion reaction mechanism in the K cells.238 The dispersion of Co3O4 and Fe2O3 nanoparticles with a Super P carbon matrix was prepared with a gentle ball milling method. The electrode comprised the cubic phase of Co3O4 (JCPDS No. 00043-1003, SG Fd-3m) and rhombohedral phase of hematite Fe2O3 (JCPDS No. 01-072-6226, SG R-3c). The TEM bright-field image and EDS mapping of the Co3O4-Fe2O3 sample revealed the existence of two different crystallite sizes of very fine (< 10 nm) and coarser (> 40 nm) nanoparticles, in which large Fe2O3 particles were decorated with small Co3O4 particles. The charge/discharge curve of the Co3O4-Fe2O3-C electrode is depicted in Fig. 29A. The electrode exhibited a reversible capacity > 400 mAh g 1 at the initial cycle in the potential range 0–3 V, while the discharge capacity faded to  250 mAh g 1 after 10 cycles and then stabilized at  220 mAh g 1 over 50 cycles. The ex situ XRD pattern of the electrode after full potassiation confirmed the diffraction peaks of Co, Fe, and K2O (JCPDS No. 00-027-0431), evidencing the conversion reaction of the Co3O4–Fe2O3/C electrode as follows: Fe2 O3 þ 6Kþ þ 6e %2Fe þ 3K2 O

(1)

Co3 O4 þ 8Kþ þ 8e– %3Co þ 4K2 O

(2)

However, both the Co3O4 and Fe2O3 phases also remained in the fully potassiated electrode, indicating that the Co3O4–Fe2O3/C electrode did not undergo full conversion reaction with potassium. After depotassiation to 3.0 V, the Co3O4 and Fe2O3 diffraction peaks were as predicted, while the discharge products of Co and K2O were still present, suggesting a partial irreversible conversion reaction. In 2019, Niu et al. studied the electrochemical performance of amorphous FeVO4 (a-FVO) in K cells and compared it with that of crystalline FeVO4.239 The initial potassiation and depotassiation capacities of a-FVO were 746.2 and 359.6 mAh g 1, respectively, corresponding to a Coulombic efficiency of 48.2% (Fig. 29B). A capacity of 220 mAh g 1 was maintained after 200 cycles. The crystalline samples calcined at high temperature (600, 700, and 800  C) displayed a lower reversible capacity (< 300 mAh g 1) than that of a-FVO (Fig. 29B). Ex situ XPS and TEM measurements revealed the formation of K2O, V2O3, VO2, and FeO, while the formation of Fe0 and V0 was not confirmed. From XPS and HRTEM analyses, the initial discharge reaction of FeVO4 would proceed as follows: 10Kþ þ 10e þ 4FeVO4 /4FeO þ V 2 O3 þ 2VO2 þ 5K2 O

(3)

Tan et al. reported an N-doped FLG and FexO composite synthesized via an in situ two-step calcination strategy.240 The synthesized sample (FexO@NFLG-240) exhibited a bubble-like structure packed with graphene nanosheets. The sample calcined at 240  C (FexO@NFLG-240) displayed a larger specific surface area, higher reversible capacity of 423 mAh g 1, and better cycle performance than those of the other samples synthesized at higher calcination temperatures (Fig. 29C and D). The FexO@NFLG-240 would work as a typical conversion reaction as follows: Fex O þ 2Kþ þ 2e %Fe þ K2 O

(4)

Overall, a-FVO and FexO@NFLG-240 exhibited good capacities and cycle performances; however, the potential hystereses and low Coulombic efficiencies of these materials should be improved for K-ion full cells. Since the volume change, potential hysteresis, and destruction of the electrode structure are general challenges for conversion-type negative electrodes, the design of the electrode structure, such as the use of a carbon composite, is indispensable to achieve sufficient cycle performance.

7.04.3.4

Transition metal chalcogenides

Similar to the oxides described in the last section, the transition metal dichalcogenides MX2 (M ¼ transition metal, X ¼ S, Se, and Te) have attracted attention as Kþ ion intercalation hosts, as Li intercalation into TiS2 is first demonstrated by Whittingham at 1976.241 Moreover, some dichalcogenides also exhibited conversion reactions. Fig. 30A illustrates the typical crystal structure of a transition metal dichalcogenide, namely a 2D layered structure with a wider interlayer space than those of graphite and typical transition metal layered oxides. The history of LIBs indicates that Liþ intercalation into TiS2, typical layered dichalcogenides, was

118

Electrode materials for K-ion batteries

Fig. 29 (A) Charge/discharge curves of the Co3O4-Fe2O3@C electrode at the current density 50 mA g 1. (B) Charge/discharge curves of a-FVO and FVO synthesized at different temperatures and a current density of 100 mA g 1. (C) Charge/discharge curves of FexO@NFLG-240 at 50 mA g 1, and (D) capacity retention of the samples, FexO@NFLG-240, Fe2O3@NFLG-300, and Fe2O3@NFLG-900, calcined at different temperatures for 100 cycles at 50 mA g 1. (A) Reproduced from Ref. Sultana, I.; Rahman, M. M.; Mateti, S.; Ahmadabadi, V. G.; Glushenkov, A. M.; Chen, Y., K-Ion and Na-Ion Storage Performances of Co3O4 –Fe2O3 Nanoparticle-Decorated Super P Carbon Black Prepared by a Ball Milling Process. Nanoscale 2017, 9(10), 3646–3654 with permission. Copyright, 2017 The Royal Society of Chemistry. (B) Reproduced from Ref. Niu, X.; Zhang, Y.; Tan, L.; Yang, Z.; Yang, J.; Liu, T.; Zeng, L.; Zhu, Y.; Guo, L., Amorphous FeVO4 as a Promising Anode Material for Potassium-ion Batteries. Energy Storage Mater. 2019, 22, 160–167 with permission. Copyright, 2019 Elsevier Ltd. (C, D) Reproduced from Ref. Tan, Q.; Li, P.; Han, K.; Liu, Z.; Li, Y.; Zhao, W.; He, D.; An, F.; Qin, M.; Qu, X., Chemically Bubbled Hollow FexO Nanospheres Anchored on 3D N-Doped Few-Layer Graphene Architecture as a PerformanceEnhanced Anode Material for Potassium-Ion Batteries. J. Mater. Chem. A 2019, 7(2), 744–754 with permission. Copyright, 2019 The Royal Society of Chemistry.

the prototype of electrode materials as a Liþ ion intercalation host.241 Moreover, Kþ ion intercalation into dichalcogenides, such as TiS2, VS2, MoS2, MoSe2, and WS2 have been extensively studied as well as Liþ and Naþ ion intercalation.242

7.04.3.4.1

3d transition metal dichalcogenides (TiS2 and VS2)

Tian et al. investigated Kþ ion insertion into TiS2 and the accompanying phase transition.33 Fig. 30B displays the charge/discharge curves of TiS2 and pre-potassiated K0.25TiS2. The initial potassiation capacity of TiS2 (210.7 mAh g 1) corresponded to 0.88 mol of K intercalated into 1 mol of TiS2 to form K0.88TiS2. The subsequent depotassiation capacity was only 145.8 mAh g 1, corresponding to 0.61 mol K extracted from TiS2, while the initial Coulombic efficiency was as low as 69.2%. A comparison of the charge and discharge profiles in Fig. 30B reveals that the irreversible capacity mainly originated from a plateau of 2.5–2.3 V but not from a lower potential. Thus, the initial irreversible capacity would be due to trapping of the Kþ ions in TiS2 rather than electrolyte decomposition. To bypass this irreversible capacity, they synthesized K0.25TiS2 by chemical pre-potassiation using potassium terephthalate. The K0.25TiS2 cells displayed an OCP of  2.4 V to Kþ/K (Fig. 30B), which was lower than that of bulk TiS2 cells ( 2.8 V). Thus, K0.25TiS2 successfully bypassed the first irreversible plateau of the initial charge curve. The K0.25TiS2 cell achieved an initial charge capacity of 145.0 mAh g 1 and a discharge capacity of 151.1 mAh g 1 with an initial Coulombic efficiency of 104.2%. These results revealed that the low initial Coulombic efficiency of the TiS2 electrodes for KIBs can be compensated by pre-potassiation. More interestingly, chemical pre-potassiation significantly improved the cycle performance of the TiS2 electrodes. To reveal the reason behind the

Electrode materials for K-ion batteries

119

Fig. 30 (A) Crystal structure of layered dichalcogenides. (B) Charge/discharge curves of KxTiS2 (x ¼ 0, 0.25) at a current density of 24 mA g 1 in the voltage range 1.0–3.0 V. (C) In situ XRD patterns collected during the first and second discharge/ charge processes of the K ||TiS2 cell under a current rate of 24 mA g 1. Left: image plot of the diffraction patterns for the (001) reflections during the first and second discharge/charge cycles; right: first two discharge/charge curves. (B–D) Reproduced from Ref. Tian, B.; Tang, W.; Leng, K.; Chen, Z.; Tan, S. J. R.; Peng, C.; Ning, G.-H.; Fu, W.; Su, C.; Zheng, G. W.; Loh, K. P., Phase Transformations in TiS2 During K Intercalation. ACS Energy Lett. 2017, 2(8), 1835–1840 with permission. Copyright, 2017 American Chemical Society.

improved electrochemical performance of K0.25TiS2, they conducted in situ XRD and ex situ TEM measurements. The in situ XRD patterns (Fig. 30C) showed that revealed that the peak of TiS2 (marked as 1 in Fig. 30C) was not recovered after depotassiation, revealing that the phase transition from TiS2 to K 0.25TiS2 was irreversible and that K remained after the first discharge. VS2 was also investigated as a negative electrode material for KIBs as well as LIBs and NIBs. Zhou et al. examined the electrochemical intercalation of Liþ, Naþ, and Kþ ions into a VS2 nanosheet synthesized by the solvothermal method.34 The reversible capacities of the VS2 nanosheets were > 1000 mAh g 1 in a Li cell,  700 mAh g 1 for a Na cell, and  00 mAh g 1 for a K cell. Notably, although the electrode comprised 20 wt% Super P, the influence of its reversible capacity was not taken into consideration. The VS2 electrode displayed the moderate average discharge potential 1.5 V vs. Kþ/K, which was 0.5 V lower than that of TiS2. Good rate capability was also demonstrated in the K cell.

120

Electrode materials for K-ion batteries

7.04.3.4.2

4d and 5d transition metal dichalcogenides (MoS2, MoSe2, and WS2)

In addition to the 3d transition metal dichalcogenides of TiS2 and VS2, MoS2, MoSe2, and WS2, which are common 4d and 5d transition metal dichalcogenides, have been studied as negative electrode materials for KIBs. Ren et al. demonstrated electrochemical potassium intercalation into MoS2 in 2017.32 They investigated electrochemical performance of composite electrode using commercially available hexagonal (2H) MoS2 powder. The MoS2 electrode presented a reversible capacity of  70 mAh g 1, with a voltage plateau at 1.2 V in a K cell, and a stable charge/discharge cycle over 200 cycles without significant degradation. Ex situ XRD measurements revealed that reversible transition from 2H-MoS2 to hexagonal K0.4MoS2. The volume expansion from MoS2 to K0.4MoS2 was  34%. When K0.4MoS2 was further potassiated to 0.25 V, formation of hexagonal KxS was confirmed, indicating the conversion reaction of MoS2 into Mo and KxS. Following this pioneering work, reversible electrode reaction, including the conversion reaction at the low potential of MoS2, was achieved by the design of carbon composite materials, such as MoS2@rGO243 MoS2@N-doped-C hollow tubes composite.244 Molybdenum diselenide (MoSe2) was also evaluated in K cells as a negative electrode material.245,246 Ge et al. reported that the þ K ion intercalation reaction into MoSe2 mainly proceeded at potentials > 0.53 V in the first charge cycle. On the other hand, the conversion reaction occurred at a lower potential, forming K5Se3 as the major final discharge product. In 2019, Zhang et al. investigated Kþ ion intercalation into WS2.247 The WS2 displayed a layer spacing of 6.18 Å, which is 0.5 and 0.02 Å wider than those of TiS2 and MoS2, respectively. WS2 could accommodate 0.62 Kþ per formula unit leading to a reversible capacity of 67 mAh g 1 and exhibited a voltage plateau at  1.0 V with an average operating potential of 0.72 V for Kþ/K. The plateau voltage in a K cell was lower than those of TiS2 and MoS2 as well as the plateau voltage in the Li ||WS2 and Na||WS2 cells. In situ XRD and ex situ TEM data revealed that the interlayer space expanded from 6.18 to 8.26 Å (34% expansion) and structural evolution was highly reversible. The interlayer expansion was as small as that of MoS2. As a result, WS2 demonstrated good cyclability up to 600 cycles with 89.2% capacity retention at a current rate of 0.3C. Although the gravimetric capacity is limited due to the heavy atomic weight of tungsten, the cyclability of WS2 is the most promising among the reported chalcogenides.

7.04.3.4.3

Metal sulfides based on conversion or conversion-alloying reactions

As discussed above, some transition metal dichalcogenides accommodate potassium through a combination of insertion and conversion reactions. Moreover, some metal sulfides, such as SnS2248–251 and Sb2S3 exhibit a combination of alloying and conversion reactions, thereby resulting in high theoretical capacities. In 2017, Lakshmi et al. first reported the electrochemical performance of SnS2@rGO composite material, which delivered initial charge and discharge capacities of 630 and 355 mAh g 1, respectively, in K cells.248 In 2019, Cao et al. investigated the potassium storage mechanism of nanoscale SnS2 crystals anchored on nitrogen-doped graphene nanosheets (SnS2@NC).251 In situ XRD and ex situ TEM measurements revealed the formation of K2S, K2S5, KSn, and K4Sn23 after potassiation to 0.01 V. After depotassiation to 3.0 V, a mixed phase of SnS2, SnS, and K2S was confirmed, indicating that the reaction was partly irreversible. Thus, SnS2 can deliver a large capacity through a conversion-alloying reaction. Future work should solve the remaining issues of insufficient cycle performance and large voltage hysteresis. Antimony sulfide (Sb2S3) was reported as a negative electrode material for LIBs, NIBs, and KIBs.252 Lu and Chen synthesized a Sb2S3-graphene architecture (Sb2S3-SNG) consisting of porous S, N-codoped graphene and Sb2S3 nanoparticles ( 20 nm) through a hydrothermal reaction as can be seen in the HR-TEM image of Fig. 31A.252 Fig. 31B displays the charge/discharge curve of the Sb2S3-SNG electrode in a K cell. The Sb2S3-SNG electrode exhibited a reversible capacity of 548 mAh g 1 with a good capacity retention of 89.4% over 100 cycles. Liu et al. conducted in situ synchrotron XRD measurement for a Sb2S3 electrode and revealed the conversion/alloying reaction of Sb2S3, which produced K2S3 and K3Sb as the final products via intermediate phases of Sb and K2S6.213 To summarize the electrochemical properties of chalcogenides, 2D layered dichalcogenides have wider interlayers than graphite and transition-metal-layered oxides and demonstrate highly reversible Kþ intercalation reaction with small volume expansion. Some layered dichalcogenides, such as MoS2 and WS2, are also expected to be safe high-power negative electrode materials because they exhibit potential plateaus > 1 V and good rate performances. These dichalcogenides undergo a conversion reaction accompanied by the formation of KxS at lower potential and exhibit high reversible capacity at the expense of safety and cycle stability. Moreover, some metal sulfides, such as SnS2 and Sb2S3, have been studied as alloying-conversion type negative electrodes, which provide high capacity. Stable cycling of dichalcogenide negative electrode materials, including the conversion reaction or alloyingconversion reaction, could be achieved by the design of a suitable material morphology and carbon composite.

7.04.4

Summary and perspective

We have discussed the overview of electrode material research for KIBs and the remaining issues. This section presents an overview and prospects of KIB as a next-generation battery comparable to LIB and NIB. Although the operating mechanisms of these batteries are very similar except for the carrier ions, the unique properties of Kþ ions as carrier ions characterize the attractiveness of KIBs as a promising battery complement to LIBs. The low standard electrode potential of Kþ/K (¼  0.1 V Liþ/Li in PC) is very advantageous for high voltage operation. Furthermore, Fe-Mn-based active materials such as K2Mn[Fe(CN)6] and KFeSO4F show a redox of 4 V vs. Kþ/K in K cells, allowing the demonstration of high-voltage K-ion cells free of minor and harmful elements such as cobalt, lead, and cadmium. This remarkable high-voltage operation opens up the possibility that KIBs can exhibit a gravimetric energy density

Electrode materials for K-ion batteries

121

Fig. 31 (A) HR-TEM image of the as-prepared Sb2S3-SNG composite and (B) charge/discharge curves of the Sb2S3-SNG composite electrode at a current density of 20 mA g 1. Reproduced from Ref. Lu, Y.; Chen, J., Robust self-supported anode by integrating Sb2S3 nanoparticles with S,Ncodoped graphene to enhance K-storage performance. Sci. China: Chem. 2017, 60(12), 1533–1539 with permission. Copyright, 2017 Science China Press and Springer-Verlag GmbH.

comparable to LIBs. Another potential advantage of KIBs is their high-power operation. This is because Kþ ions diffuse faster in solution than Liþ or Naþ ions, and therefore interact less with solvents and anions. In many studies, electrode materials with crystal structures suitable for Kþ ion diffusion have been developed to achieve high power operation. Previous research on cathode materials for KIBs has revealed the framework’s importance, including the crystal structure, redox centers, and anions. Although transition metal layered oxides are the most promising cathode materials for LIBs and NIBs, typical transition metal layered oxide materials exhibit low capacitance, low operating potential, and multiple phase transitions due to K þ/ vacancy ordering. Therefore, by selecting appropriate dopants, these structural difficulties can be mitigated to provide high capacity, high operating potential, and good cycle life. PBA is one of the promising groups of cathode materials in terms of energy density, cycle performance, and rate capability. Its 3D open framework structure provides suitable channels, and interstitial sites for the diffusion and insertion of large Kþ ions; similar to PBA, polyanionic compounds with 3D open framework structure may be candidates for other cathode materials due to their high ionic conductivity and high operating potential. Previous studies have revealed suitable crystal structures, including a KTiOPO4-type structure for Kþ ion diffusion. As an anode, graphite, which is widely used and technically established in LIBs, is a promising candidate for KIBs. However, early studies have reported that graphite’s electrochemical performance (e.g., cyclability and low-temperature performance) is inferior to that of Li batteries. On the other hand, recent studies have focused on the dramatic effects of electrolyte, binder, and physical properties on graphite electrodes’ cyclability. These studies have revealed that the optimal properties are different from those of the Li system. Therefore, the development of graphite materials suitable for Kþ insertion is also an interesting and practically important topic. Some non-graphite carbon materials show superior performance in terms of specific capacity and rate capability. Future comparative studies will further understand the Kþ insertion properties of carbonaceous materials and their mechanisms, which will enable the development of anode materials suitable for KIB. Transition metal oxides, chalcogenides, and polyanionic compounds are other notable groups of anode materials. Their relatively high K-intercalation potential is expected to be a safety advantage by avoiding K-metal plating and dendrite formation even in overcharged conditions and high-rate operation. Anode materials based on alloying/conversion reactions generally have higher specific power and volume capacity than other anode materials. However, most of the anode materials using alloying and conversion reactions do not have sufficient cycling stability of K-cells because the volume change during the potassiation reaction is larger than that during the lithiation and sodiation reactions of LIB and NIB materials. This large volume change is attributed to the large molar volume of Kþ ions. Therefore, the electrode structure’s design is essential to achieve satisfactory cyclicity, as evidenced by previous studies. In the future, KIBs are expected to improve further energy density, cyclicity, and high power operation, paving the way for practical applications. From a scientific point of view, the study of Kþ ion insertion hosts has enabled a better understanding of the effect of insertion cations on the crystal structure and electronic state of insertion host materials by comparing the insertion chemical properties of cations (e.g., Liþ, Naþ, Kþ, Mg2þ, Cu2þ, Al3þ). Thus, KIB’s research provides new insights into electrode reactions and solid-state ionics, and opens up new strategies for realizing next-generation batteries.

122

Electrode materials for K-ion batteries

References 1. Rudnick, R. L.; Gao, S. Composition of the Continental Crust. Treatise Geochem. 2003, 3, 1–64. 2. Komaba, S.; Hasegawa, T.; Dahbi, M.; Kubota, K. Potassium Intercalation into Graphite to Realize High-Voltage/High-Power Potassium-Ion Batteries and Potassium-Ion Capacitors. Electrochem. Commun. 2015, 60, 172–175. 3. Vaalma, C.; Buchholz, D.; Weil, M.; Passerini, S. A Cost and Resource Analysis of Sodium-Ion Batteries. Nat. Rev. Mater. 2018, 3, 18013. 4. Marcus, Y. Thermodynamic Functions of Transfer of Single Ions From Water to Nonaqueous and Mixed Solvents: Part 3-Standard Potentials of Selected Electrodes. Pure and Applied Chemistry 1985, 57 (8), 1129–1132. 5. Shannon, R. D. Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides. Acta Crystallogr., Sect. A: Cryst. Phys., Diffr., Theor. Gen. Crystallogr. 1976, 32 (5), 751–767. 6. Nightingale, E. R. Phenomenological Theory of Ion Solvation. Effective Radii of Hydrated Ions. J. Phys. Chem. 1959, 63 (9), 1381–1387. 7. Matsuda, Y. Behavior of Some Ions in Mixed Organic Electrolytes of High Energy Density Batteries. J. Electrochem. Soc. 1981, 128 (12), 2552–2556. 8. Okoshi, M.; Yamada, Y.; Komaba, S.; Yamada, A.; Nakai, H. Theoretical Analysis of Interactions between Potassium Ions and Organic Electrolyte Solvents: A Comparison with Lithium, Sodium, and Magnesium Ions. J. Electrochem. Soc. 2017, 164 (2), A54–A60. 9. Hosaka, T.; Muratsubaki, S.; Kubota, K.; Onuma, H.; Komaba, S. Potassium Metal as Reliable Reference Electrodes of Nonaqueous Potassium Cells. J. Phys. Chem. Lett. 2019, 10 (12), 3296–3300. 10. Abe, T.; Fukuda, H.; Iriyama, Y.; Ogumi, Z. Solvated Li-Ion Transfer at Interface Between Graphite and Electrolyte. J. Electrochem. Soc. 2004, 151 (8), A1120–A1123. 11. Bie, X.; Kubota, K.; Hosaka, T.; Chihara, K.; Komaba, S. A Novel K-Ion Battery: Hexacyanoferrate(ii)/Graphite Cell. J. Mater. Chem. A 2017, 5 (9), 4325–4330. 12. Hironaka, Y.; Kubota, K.; Komaba, S. P2- and P3-KxCoO2 as an Electrochemical Potassium Intercalation Host. Chem. Commun. 2017, 53 (26), 3693–3696. 13. Kim, H.; Kim, J. C.; Bo, S. H.; Shi, T.; Kwon, D. H.; Ceder, G. K-Ion Batteries Based on a P2-Type K0.6CoO2 Cathode. Adv. Energy Mater. 2017, 7 (17), 1700098. 14. Kubota, K.; Dahbi, M.; Hosaka, T.; Kumakura, S.; Komaba, S. Towards K-Ion and Na-Ion Batteries as “Beyond Li-Ion”. Chem. Rec. 2018, 18 (4), 459–479. 15. Eftekhari, A. Potassium Secondary Cell Based on Prussian Blue Cathode. J. Power Sources 2004, 126 (1-2), 221–228. 16. Wu, X.; Jian, Z.; Li, Z.; Ji, X. Prussian White Analogues as Promising Cathode for Non-Aqueous Potassium-Ion Batteries. Electrochem. Commun. 2017, 77, 54–57. 17. Xue, L.; Li, Y.; Gao, H.; Zhou, W.; Lü, X.; Kaveevivitchai, W.; Manthiram, A.; Goodenough, J. B. Low-Cost High-Energy Potassium Cathode. J. Am. Chem. Soc. 2017, 139 (6), 2164–2167. 18. Chihara, K.; Katogi, A.; Kubota, K.; Komaba, S. KVPO4F and KVOPO4 Toward 4 Volt-Class Potassium-Ion Batteries. Chem. Commun. 2017, 53 (37), 5208–5211. 19. Kim, H.; Seo, D.-H.; Bianchini, M.; Clément, R. J.; Kim, H.; Kim, J. C.; Tian, Y.; Shi, T.; Yoon, W.-S.; Ceder, G. A New Strategy for High-Voltage Cathodes for K-Ion Batteries: Stoichiometric KVPO4F. Adv. Energy Mater. 2018, 8 (26), 1801591. 20. Park, W. B.; Han, S. C.; Park, C.; Hong, S. U.; Han, U.; Singh, S. P.; Jung, Y. H.; Ahn, D.; Sohn, K.-S.; Pyo, M. KVP2O7 as a Robust High-Energy Cathode for Potassium-Ion Batteries: Pinpointed by a Full Screening of the Inorganic Registry under Specific Search Conditions. Adv. Energy Mater. 2018, 8 (13), 1703099. 21. Hosaka, T.; Shimamura, T.; Kubota, K.; Komaba, S. Polyanionic Compounds for Potassium-Ion Batteries. Chem. Rec. 2019, 19 (4), 735–745. 22. Jian, Z.; Liang, Y.; Rodríguez-Pérez, I. A.; Yao, Y.; Ji, X. Poly(anthraquinonyl sulfide) Cathode for Potassium-Ion Batteries. Electrochem. Commun. 2016, 71, 5–8. 23. Ma, J.; Zhou, E.; Fan, C.; Wu, B.; Li, C.; Lu, Z.-H.; Li, J. Endowing CuTCNQ with a New Role: A High-Capacity Cathode for K-Ion Batteries. Chem. Commun. 2018, 54 (44), 5578–5581. 24. Yang, S.-Y.; Chen, Y.-J.; Zhou, G.; Fu, Z.-W. Multi-Electron Fused Redox Centers in Conjugated Aromatic Organic Compound as a Cathode for Rechargeable Batteries. J. Electrochem. Soc. 2018, 165 (7), A1422–A1429. 25. Jian, Z.; Luo, W.; Ji, X. Carbon Electrodes for K-Ion Batteries. J. Am. Chem. Soc. 2015, 137 (36), 11566–11569. 26. Luo, W.; Wan, J.; Ozdemir, B.; Bao, W.; Chen, Y.; Dai, J.; Lin, H.; Xu, Y.; Gu, F.; Barone, V.; Hu, L. Potassium Ion Batteries with Graphitic Materials. Nano Lett. 2015, 15 (11), 7671–7677. 27. Bin, D.-S. S.; Lin, X.-J. J.; Sun, Y.-G. G.; Xu, Y.-S. S.; Zhang, K.; Cao, A.-M. M.; Wan, L. Engineering Hollow Carbon Architecture for High-performance K-ion Battery Anode. J. Am. Chem. Soc. 2018, 140 (23), 7127–7134. 28. Sultana, I.; Rahman, M. M.; Chen, Y.; Glushenkov, A. M. Potassium-Ion Battery Anode Materials Operating through the Alloying-Dealloying Reaction Mechanism. Adv. Funct. Mater. 2018, 28 (5), 1703857. 29. Han, J.; Xu, M.; Niu, Y.; Li, G.-N.; Wang, M.; Zhang, Y.; Jia, M.; Li, C. Exploration of K2Ti8O17 as an Anode Material For Potassium-Ion Batteries. Chem. Commun. 2016, 52 (75), 11274–11276. 30. Dong, S.; Li, Z.; Xing, Z.; Wu, X.; Ji, X.; Zhang, X. Novel Potassium-Ion Hybrid Capacitor Based on an Anode of K2Ti6O13 Microscaffolds. ACS Appl. Mater. Interfaces 2018, 10 (18), 15542–15547. 31. Kishore, B.; Munichandraiah, N. K2Ti4O9: A Promising Anode Material for Potassium Ion Batteries. J. Electrochem. Soc. 2016, 163 (13), A2551–A2554. 32. Ren, X.; Zhao, Q.; McCulloch, W. D.; Wu, Y. MoS2 as a Long-Life Host Material for Potassium Ion Intercalation. Nano Res. 2017, 10 (4), 1313–1321. 33. Tian, B.; Tang, W.; Leng, K.; Chen, Z.; Tan, S. J. R.; Peng, C.; Ning, G.-H.; Fu, W.; Su, C.; Zheng, G. W.; Loh, K. P. Phase Transformations in TiS2 During K Intercalation. ACS Energy Lett. 2017, 2 (8), 1835–1840. 34. Zhou, J.; Wang, L.; Yang, M.; Wu, J.; Chen, F.; Huang, W.; Han, N.; Ye, H.; Zhao, F.; Li, Y.; Li, Y. Hierarchical VS2 Nanosheet Assemblies: A Universal Host Material for the Reversible Storage of Alkali Metal Ions. Adv. Mater. 2017, 29 (35), 1702061. 35. Han, J.; Niu, Y.; Bao, S.-J.; Yu, Y.-N.; Lu, S.-Y.; Xu, M. Nanocubic KTi2(PO4)3 Electrodes for Potassium-Ion Batteries. Chem. Commun. 2016, 52 (78), 11661–11664. 36. Wei, Z.; Wang, D.; Li, M.; Gao, Y.; Wang, C.; Chen, G.; Du, F. Fabrication of Hierarchical Potassium Titanium Phosphate Spheroids: A Host Material for Sodium-Ion and Potassium-Ion Storage. Adv. Energy Mater. 2018, 8 (27), 1801102. 37. Lei, K.; Li, F.; Mu, C.; Wang, J.; Zhao, Q.; Chen, C.; Chen, J. High K-Storage Performance Based on the Synergy of Dipotassium Terephthalate And Ether-Based Electrolytes. Energy Environ. Sci. 2017, 10 (2), 552–557. 38. Liao, J.; Hu, Q.; Yu, Y.; Wang, H.; Tang, Z.; Wen, Z.; Chen, C. A Potassium-Rich Iron Hexacyanoferrate/Dipotassium Terephthalate@carbon Nanotube Composite Used for KIon Full-Cells with an Optimized Electrolyte. J. Mater. Chem. A 2017, 5 (36), 19017–19024. 39. Li, C.; Deng, Q.; Tan, H.; Wang, C.; Fan, C.; Pei, J.; Cao, B.; Wang, Z.; Li, J. Para-Conjugated Dicarboxylates with Extended Aromatic Skeletons as the Highly Advanced Organic Anodes for K-Ion Battery. ACS Appl. Mater. Interfaces 2017, 9 (33), 27414–27420. 40. Mizushima, K.; Jones, P.; Wiseman, P.; Goodenough, J. B. LixCoO2 (0< x X(1), M(2) > M(1)].65 Alluaudite phosphates assume a monoclinic framework (s.g. C2/c).

Fig. 5 (a) Structural illustration of general alluaudites, e.g. NaMnFe2(PO4)3. (FeO6 octahedra: brown, MnO6 octahedra: pink, PO4 tetrahedra: blue, Na atoms: yellow spheres). (b) Galvanostatic voltage-capacity profiles of gel combustion made NaMnFe2(PO4)3 alluaudite cathode. (c) Comparative voltage-capacity plots of alluaudites Na2Fe3-xMnx(PO4)3 (x ¼ 0–2). (d) Voltage-capacity profiles of vanadium-based alluaudite Na2Mn2V(PO4)3 working as a 3.5 V cathode. From Panel (b) Dwibedi, D., Jaschin, P.W., Gond, R., Barpanda, P., Revisiting the Alluaudite NaMnFe2(PO4)3 Sodium Insertion Material: Structural, Diffusional and Electrochemical Insights. Electrochim. Acta. 2018, 283, 850–857, with permission. Panel (c) Huang, W., Wu, Z., et al., Self-Assembled Alluaudite Na2Fe3-xMnx(PO4)3 Micro/Nanocompounds for Sodium-Ion Battery Electrodes: A New Insight Into Their Electronic and Geometric Structure. Chem. Eur. J. 2015, 21, 851, with permission. Panel (d) Zhang, P., Yang, K., Song, L., Gao, J., et al., An Alluaudite¼Type Sodium-Ion Battery Cathode Candidate Na2Mn2V(PO4)3: Crystal Growth, Preparation, Stucture and Electrochemical Properties, J. Alloy. Compd. 2019, 783, 409415

Development of polyanionic sodium-ion battery insertion materials

251

Crystal structure of phosphate alluaudite NaMnFe2(PO4)3 is shown in Fig. 5a. This structure consists of infinite chains formed by
Equivalent chains are a succession of M(2) octahedral pairs linked by M(1) octahedra that run as kinked chains parallel to [101]. connected by vertices with PO4 tetrahedra to delimit sheets oriented normal to the b–axis. The three-dimensional framework is built by bridging these equivalent chains with corner connected PO4 tetrahedra. In case of NaMnFe2(PO4)3, M1 site is statistically occupied by Fe1 and Mn1 whereas M2 site is occupied by Fe2 and Mn2. This alluaudite system offers two large open tunnels along c axis, which are occupied by Na species. NaMnFe2(PO4)3 alluaudite was introduced as a potential SIB cathode by Delmas group.66 Though it has open structure suitable for Naþ (de)insertion, meagre electrochemical activity was observed. Prepared by sol-gel method, it led to a reversible capacity below 40 mAh/g (at a slow rate of C/100) due to large micrometric particles. Following, NaMnFe2(PO4)3/MWCNTs composite was employed to control the particle size/morphology, which enhanced the electrochemical activity with capacity over 150 mAh/g. The presence of MWCNTs can improve the electronic conductivity and diffusion coefficient to activate both Fe3þ/ Fe2þ and Mn3þ/Mn2þ redox couples. Autocombustion route was later employed to form nanoscale ( 200–300 nm) alluaudite particles with porous morphology, which led to a reversible capacity over 60 mAh.g 1 involving single-phase Fe3þ/Fe2þ redox activity at 2.8 V (Fig. 5b).67 Exploring high capacity PO4-based alluaudite compounds, family of self-assembled Na2Fe3 xMnx(PO4)3 (x ¼ 0--2) insertion materials were prepared by low temperature (ca. 200  C) solvothermal route developing self-assembled monodispersed nanorods. The formation of nanoparticles led to efficient electrochemical performance in Na2Fe3 xMnx(PO4)3 alluaudite cathodes.68 Higher Mn content led to lower capacity due to redox inactive Mn3þ/Mn2þ species. The sole capacity came from Fe3þ/Fe2þ redox couple, with the highest capacity (ca. 140 mAh/g) obtained in case of Na2Fe3(PO4)3 (¼ Na0.67FePO4) alluaudite (Fig. 5c). While alluaudites works as intercalation based cathodes, Na2Co2Fe(PO4)3 was reported as the first alluaudite working as a conversion type cathode material. When this alluaudite was cycled between 3 and 0.03 V, it led to the insertion of  1 Naþ involving Fe3þ/Fe2þ redox activity around 0.9 V. It was followed by long plateau  0.6 V involving conversion reaction of constituent Fe2þ and Co2þ species to their metallic forms (Fe0/Co0). During conversion reaction, partial sodiation of Na2Co2Fe(PO4)3 led to the formation of alluaudite Na3Co2Fe(PO4)3, which can work as a 3.5 V cathode involving both Fe3þ/Fe2þ and Co3þ/Co2þ redox activity. Thus, alluaudite type Na2Co2Fe(PO4)3 can function as a dual positive/negative electrode material for SIBs.69 Various Fe-based alluaudites Na2  xFe3(PO4)3 (x ¼ 0.01–0.53) can be prepared by controlling the stoichiometry of Na sites.70,71 While they retain the overall alluaudite framework similar to Na2Fe3(PO4)3, they differ in relative occupancy of alkali sites and the final electrochemical performance. One notable example is alluaudite Na2.01Fe3(PO4)3 ( Na0.67FePO4) acting as a 2.7 V sodium battery cathode. Upon usage of carbon nanotubes, monodisperse alluaudite nanocomposite Na2.01Fe3(PO4)3/CNT can be prepared, which can deliver excellent electrochemical activity (ca. 140 mAh/g) with good cycling stability due to reduced Naþ diffusion length, improved electronic conductivity and electrode-electrolyte wetting. Similar phosphate alluaudites can be designed like Na1.86Fe3(PO4)3, Na1.702Fe3(PO4)3 and Na1.47Fe3(PO4)3 by varying the Na/Fe cationic ratio, all of them working as reversible SIB insertion materials. As these Fe-based alluaudites suffer from low voltage, in an effort to elevate the redox potential, vanadium-substituted Na2Fe1.96V0.96(PO4)3 alluaudites were realized using solid-state and sol-gel routes. In case of Na2Fe1.96V0.96(PO4)3, efficient Naþ (de)intercalation was observed involving a 3.1 V Fe3þ/Fe2þ redox potential along with very fast rate kinetics.72 In-situ XRD study revealed a single-phase redox reaction involving minimal volume change ( 1.6%) and simultaneous activity of Fe3þ/Fe2þ, V4þ/ V3þ and V3þ/V2þ redox couples. Among the PO4-based alluaudites, Na2Fe1.96V0.96(PO4)3 works as the best candidate with desirable combination of high voltage, fast rate and cycling stability. Following, the Mn-analogue Na2Mn2V(PO4)3 alluaudite was synthesized by sol-gel route. Enabling both Mn3þ/Mn2þ and V4þ/V3þ redox couples, it delivered a discharge capacity over 95 mAh/g involving a central redox potential at 3.5 V (Fig. 5d). Enabled by high redox potential, these vanadium-based phosphate alluaudites can serve as high energy density cathodes for SIBs. Pursuing other transition metal chemistry, NaCoFe2(PO4)3 was discovered with combustion synthesis route having nanoscale particles ( 300 nm) with porous morphology. Isostructural to the first alluaudite NaMnFe2(PO4)3, this open framework facilitates one-dimensional Naþ mobility with low migration barrier (ca.  0.31 eV).73 With an underlying solid-solution redox mechanism, it works as a 2.9 V Fe-based SIB cathode delivering capacity over 70 mAh/g with moderate rate kinetics. Overall, these PO4-based alluaudites suffer from low redox potential, slow rate kinetics, moderate capacity (limited below 100 mAh/g) and hence low energy density. Till date, no report exists showing multiple electron reaction. Thus, these alluaudites are far from real-life battery application.

7.09.3

Sulfate class of polyanionic cathodes

7.09.3.1

Bisulfates

Based on inductive effect, the more electronegative moieties in polyanionic compounds can upshift the redox potential of transition metals. Using Pauling’s electronegative scale, sulfur (S) forms the most electronegative anion, thus SO4-based compounds can in principle exhibit the highest redox potential.74 This point has been proved in case of LIBs with the advent of fluoro(hydroxy)sulfates and bisulfates reaching close to 4 V activity.75,76 Although the metal fluorosulfates offer high redox potential stemming from the electronegativity of both SO42 and F species, their presence (particularly F) make them prone to moisture attack and chemical

252

Development of polyanionic sodium-ion battery insertion materials

degradation. This stability issue can be partly circumvented by removing F species. Fluorine free polysulfate [e.g., bisulfate Li2Fe(SO4)2] can be designed as high voltage ( 3.8 V) cathodes for LIBs. Similar strategy can be adopted for SIBs with metal bisulfate [Na2M(SO4)2, M ¼ 3d metals] chemistry. Unfortunately, these sulfates are hygroscopic with the tendency to form hydrated bisulfates [Na2M(SO4)2$nH2O, n ¼ 0, 2, 4].77 It is difficult to isolate anhydrous Na2M(SO4)2 bisulfate. However, it can be prepared indirectly by dehydration of hydrated bisulfates [Na2M(SO4)2$4H2O / Na2M(SO4)2$2H2O / Na2M(SO4)2] (Fig. 6a). Na2M(SO4)2$4H2O (M ¼ Fe, Co) belongs to bloedite mineral family named after Na2Mg(SO4)2$4H2O bloedite system.78 They can be prepared by simple dissolution of Na2SO4 and MSO4$7H2O polyhydrate precursors in aqueous media followed by precipitation in alcohol media. These bloedites have a monoclinic structure (s.g. P21/c) built with isolated MO2(OH)4 octahedra where four O atoms are linked to 2H atoms to form four units of structural H2O. The remaining two O atoms are connected to SO4 tetrahedra so as to form M(SO4)2(H2O)4 blocks delimiting large voids locating the Naþ ions. The bloedite Na2Fe(SO4)2.4H2O was found to be redox active with a capacity of 50 mAh/g, albeit at a slow rate of C/50. With a sloping voltage profile, it works as a 3.3 V sodium insertion host. However, cathode cycling leads to steady amorphization and material instability in bloedite cathode systems. On the other hand, Na2Fe(SO4)2$2H2O belongs to kröhnkite mineral type compound, which can be prepared by dissolution and precipitation of Na2SO4 and MSO4$7H2O in distilled water without using alcohol. Kröhnkite Na2Fe(SO4)2$2H2O has a monoclinic (s.g. P21/c) framework built from isolated MO4(OH)2 octahedra with two O atoms linked to 2H atoms thereby forming two units of structural water.79 The other four O atoms are bridged to four SO4 tetrahedra to create tortuous one-dimensional Naþ diffusion pathways. It delivers reversible capacity approaching 70 mAh/g at a faster rate (C/20) with the average Fe3þ/Fe2þ redox activity located at 3.25 V. Kröhnkites are more stable than bloedites upon cycling, though both suffer from limited capacity due to weight penalty of structural water units. Higher capacity can be obtained by removing these structural water molecules. It can be achieved by dehydration of both bloedite and kröhnkite phases to get anhydrous Na2Fe(SO4)2. The direct synthesis of Na2Fe(SO4)2 invariably leads to the formation of vanthoffite Na6Fe(SO4)4 phase. Anhydrous Na2Fe(SO4)2 has a monoclinic structure with FeO6 octahedra sharing O atoms with five SO4 tetrahedra forming one dimensional Naþ diffusion pathways. Involving an Fe3þ/Fe2þ redox activity centered at 3.4 V, it exhibits a reversible capacity over 60 mAh/g. While these (de)hydrated bisulfates work as Na insertion host, they are marred by materials stability, low voltage/capacity, and inability to realize two-electron reaction (removal of 2 Naþ). Deviating from these Fe2þ-based bisulfates, Goodenough group reported the first Fe3þ-based bisulfate NaFe(SO4)2 cathode that belongs to eldfellite class of minerals having a monoclinic structure (s.g. C2/m).80 The framework is built from layers of edgesharing FeO6 octahedra abridged by SO4 tetrahedra with large voids enabling two-dimensional Naþ migration pathways (Fig. 6b). This eldfellite system delivered a reversible capacity of 80 mAh/g involving Fe3þ/Fe2þ redox activity centered at 3.2 V (Fig. 6c). The sloping voltage profile is related to single-phase redox mechanism with moderate rate kinetics. While complete sodiation (1 Na) is not attained, it can form a new polymorph of Na2Fe(SO4)2 that can be structurally different than anhydrous

(a) Na S

Na2Fe(SO4)2.4H2O Bloedite : 3.4 V cathode (b)

(d)

NaFe(SO4)2

4.0

Voltage (vs. Na/Na+)

Voltage (vs. Na/Na+)

(c)

NaM(SO4)2: Eldfellite

Na2Fe(SO4)2 anhydrous : 3.3 V cathode

Na2Fe(SO4)2.2H2O Krohnkite : 3.25 V cathode

3.0

2.0 0

20

40

60

Capaci y (mAh/ )

80

3.0

NaV(SO4)2

2.0

1.0

20

40

60

80

Capacity (mAh/g)

Fig. 6 (a) Progressive dehydration (shown by arrow marks) leading from bloedite to kröhnkite to anhydrous Na2Fe(SO4)2 bisulfate. (FeO6 octahedra: green, SO4 tetrahedra: yellow, Na atoms: yellow spheres, structural waters are highlighted by dashed marks.) (b) Structure of NaM(SO4)2 eldfellite along [010] direction. (MO6 octahedra: cyan, SO4 tetrahedra: yellow, Na atoms: green spheres). Galvanostatic voltage-capacity profiles of eldfellite (c) NaFe(SO4)2 and (d) NaV(SO4)2 cathodes. From panel (c) Singh, P., Shiva, K., Celio, H., Goodenough, J. B., Eldfellite, NaFe(SO4)2: An Intercalation Cathode Host for Low Cost Na-Ion Batteries. Energy Environ. Sci. 2015, 8, 3000–3005, with permission.

Development of polyanionic sodium-ion battery insertion materials

253

Na2Fe(SO4)2 derived by thermal treatment of bloedite/kröhnkite. Exploring eldfellite systems, isostructural V-analogue NaV(SO4)2 has been recently examined as a 2.3 V Na-insertion host based on V3þ/V2þ redox couple delivering capacity close to 70 mAh/g (Fig. 6d).81 While its energy density is poor, this system can be implemented as a potential high voltage cathode by activating the V4þ/V3þ redox couple with the use of high-voltage electrolytes. Along this line, isostructural NaM(SO4)2 (M ¼ Cr, Ti) eldfellite systems call for exploration as potential cathode in sodium-ion batteries.

7.09.3.2

Alluaudite sulfates

As discussed in Section 7.09.2.4, the PO4 class of alluaudites offer poor electrochemical performance even though alluaudites framework have suitable structural, ionic migration and electrochemical viability to work as efficient cathodes. Using the advantage of inductive effect as discussed in Section 7.09.3.1, it is enticing to design SO4-based alluaudite compositions for sodium batteries. While nature is rich with naturally occurring PO4-based alluaudite minerals, there is no report on naturally occurring sulfate alluaudite. To fill this gap, Yamada group attempted to design synthetic SO4-based alluaudites. Leading to Na2Fe2(SO4)3 as the first alluaudite-type sulfate.82 This Na-Fe-S-O alluaudite assumes a monoclinic framework (s.g. C2/c) (Fig. 7a). This Na2Fe2(SO4)3 structure is completely different from the earlier reported NASICON type AxFe2(SO4)3 built with lantern units [Fe2(SO4)3]. Instead, Na2Fe2(SO4)3 assumes alluaudite mineral-type with generic formula AA0 BM2(XO4)3 (B ¼ Na-1 site, A and A0 ¼ partially occupied Na-2 site and Na-3 sites respectively, M ¼ Fe site and X ¼ S). All Fe species occupy a crystallographically distinct site forming isolated edge-sharing FeO6-FeO6 [Fe2O10] octahedral dimers, bridged to SO4 tetrahedra strictly in corner-sharing fashion having tunnels along c axis. The Na occupies three distinct crystallographic sites (Na-1 to Na-3). While Na-1 site is fully filled with no migration tunnels, the Na-2 and Na-3 are partially occupied with large tunnels for efficient one-dimensional migration (along c-axis). It facilitates efficient one-dimensional Naþ migration in alluaudites with low energy barrier as confirmed by theoretical studies (Fig. 7b). This sulfate alluaudite works as an exceptionally high (ca. 3.8 V) voltage cathode registering the highest potential among all known Fe-based cathodes for sodium-ion batteries. This voltage is close to the 4 V LiCoO2 cathode for LIBs. Prepared by classical solid-state route, it offers rare combination of earth-abundant and economic Na-Fe-S-O elements, high redox voltage operation, excellent rate kinetics and long cycle life. Involving multiple step wise voltage profiles, it delivers a reversible capacity of 100 mAh/g with three plateaus at 3.42, 3.8, and 4.04 V (Fig. 7c). With its high voltage and theoretical capacity  120 mAh/g, this alluaudite can offer high energy density (> 540 Wh kg 1) making Na-ion batteries competitive with the state-of-the-art Liion batteries. Further, with its open tunnels, this cathode exhibits excellent rate kinetics even with the presence of large micrometric particles (Fig. 7d). These attributes have attracted wide attention for further development of SO4-based alluaudite type cathodes. While stoichiometric Na2Fe2(SO4)3 alluaudite is outstanding as cathode, its synthesis is tricky due to the formation of thermodynamically stable vanthoffite Na6Fe(SO4)4 impurity along with unreacted FeSO4.xH2O precursor. Thus, an excess amount of Na2SO4 is essential in the precursor mixture to obtain pure endmember. Also, in the Na2SO4-FeSO4 binary system, it is possible to obtain off-stoichiometric alluaudites with sodium rich and iron deficient compositions. Thus, a range of alluaudite Na2 þ 2xFe2 x(SO4)3 compositions (x ¼ 0 0.4) have been investigated. The stoichiometric Na2Fe2(SO4)3 composition is highly unstable unlike the off-stoichiometric Na2 þ 2xFe2  x(SO4)3 (x ¼ 0.25– 0.3) alluaudite products.83 Extending the offstoichiometry, even bisulfate Na2Fe(SO4)2 can be expressed as an off-stoichiometric alluaudite system Na2þ2xFe2-x(SO4)3 (with x ¼ 0.5). Depending on the sodium rich and iron deficient off-stoichiometric compositions, these alluaudites showed tunable redox potential and discharge capacity. Essentially, this off-stoichiometry limits the degree of desodiation, hence restricting the theoretical capacity to 100 mAh/g. Although the extent of off-stoichiometry can be manipulated, stabilizing stoichiometric Na2Fe2(SO4)3 alluaudite is difficult. Apart from synthesis, another key concern is the severe capacity fading due to the sluggish ionic and electronic conductivity. DFT study suggests complete desodiation of Na2Fe2(SO4)3 to form Fe2(SO4)3 involves high anisotropic distortion in FeO6 octahedra making the structure unstable.84 Also, these sulfate alluaudites are moisture sensitive with fast Fe2þ to Fe3þ oxidation tendency. This issue can be alleviated by carbon coating that can improve the conductivity and moisture resistance of sulfate alluaudites. Similar to other SO4-based cathodes, caution must be exercised to avoid waterbased routes and high-temperature annealing as the sulfate products are soluble in water (or moisture sensitive) and decompose upon heating with SOx gas evolution. Nonetheless, two-step aqueous methods, such as spray drying synthesis and Pechini route, have been successfully implemented in alluaudite formation. Playing with the stoichiometry, it is also possible to observe ambient temperature phase transformation from Pechini synthesized anhydrous bloedite Na2Fe(SO4)2 to alluaudite Na2 þ 2xFe2 x(SO4)3 by simple mechanical milling. The alluaudite Na2 þ 2xFe2 x(SO4)3 is sensitive to moisture attack leading to the formation of krohnkite Na2Fe(SO4)2.2H2O or more probably bloedite Na2Fe(SO4)2.4H2O. Hence, the krohnkite and bloedite phases can also be used as precursors to form anhydrous alluaudite phase by careful heat treatment. The most striking feature of alluaudite is its high redox voltage (ca. 3.8 V) matching close to that of LiCoO2 (4 V). The origin of this high redox potential can be related to its crystal structure.85,86 Deviating from mixed metal-based phosphate alluaudites, the sulfate alluaudites consist of single transition metal, where FeO6 octahedra share edge to form Fe2O10 dimers, which are corner shared with neighboring SO4 tetrahedra. This unique edge-sharing FeO6-FeO6 coordination leads to a small Fe  Fe interatomic distance, which is a potential reason behind the high redox potential. From thermodynamics point-of-view, the high redox voltage can be expressed as:

254 Development of polyanionic sodium-ion battery insertion materials

Fig. 7 (a) Structure of alluaudite Na2Fe2(SO4)3 (FeO6 octahedra: green, SO4 tetrahedra: yellow, Na atoms: yellow spheres). Bond valence site generated Naþ migration pathways showing one-dimensional channels along c axis. (b) Corresponding Naþ migration barrier predicted by DFT calculations. (c) Voltage-capacity profiles (at C/20) showing reversible Naþ (de)insertion with average redox potential at 3.8 V. (inset) corresponding dQ/dV profiles showing three sets of redox peaks. (d) Capacity as a function of cycling rate. (inset) Comparative voltage-capacity profiles at various rates. From Barpanda, P., Oyama, G., Nishimura, S., Chung, S.C., Yamada, A., A 3.8 V Earth-Abundant Sodium Battery Electrode. Nat. Commun. 2014, 5, 4358, with permission.

Development of polyanionic sodium-ion battery insertion materials

255

E ¼ DGo =nF ¼ ðxGo Na þ Go Host  Go NaxHost Þ=nF where n ¼ number of electrons, G ¼ Gibbs free energy, and F ¼ Faraday constant. In case of alluaudites, GoHost  GoNaxHost is quite large as the sodiated state is stable (low GoNaxHost), whereas the desodiated phase is highly unstable (high GoHost predicted from DFT studies), thereby upshifting the redox voltage. In addition, the edge sharing FeO6-FeO6 geometry offers strong Fe3þ-Fe3þ repulsion in the charged state that can lead to high redox potential. Following the discovery of Fe-based alluaudite, isostructural 3d transition metal-based alluaudite analogues have been discovered. For example, low temperature solid-state route (350  C for 12 h) has been employed to isolate Na2 þ 2xMn2 x(SO4)3 (x ¼ 0.22) and Na2 þ 2xCo2 x(SO4)3 (x ¼ 0.16) by using Na2SO4 and MnSO4.H2O and CoSO4$7H2O precursors, respectively.87,88 These alluaudite phases generally adopt monoclinic C2/c structure. However, the Co-based phase displays polymorphism depending on the synthesis method. When prepared via (dry) solid-state route, it crystallizes in the monoclinic C2/c framework. However, when formed via (wet) Pechini method, it adopts anhydrous bloedite type C2/c structure. Independent of the synthesis routes, the Ni-based analogue crystallizes in the anhydrous bloedite type C2/c structure. Formation of alluaudite type Na2 þ2xNi2 x(SO4)3 remains a challenge, which can act as high-voltage cathode. Despite having open framework, the Mn- and Co-based alluaudites are found to be electrochemical inactive due to the lack of stable high-voltage electrolytes, although DFT calculations predict them as high-voltage cathodes with average redox potential of Mn-, Co- and Ni-based alluaudite phases around 4.4, 5.1 and 5.2 V (vs. Na/Naþ) respectively.84 In pursuit of improved electrochemical activity, the solid solution family of Fe-Co, Co-Mn, Mn-Fe alluaudites have been prepared. Incorporation of Mn or Co in solid solution with Fe can further enhance the overall redox potential although with lower electrochemical performance due to inactivity of Mn and Co species in these alluaudites. Among them, mixed cobalt-manganese based alluaudite sulfate, Na2 þ 2d(Co0.63Mn0.37)2 d(SO4)3, has been shown to work as a high voltage (ca. > 4.0 V) cathode. The simultaneous participation of cobalt and manganese ions and hence the redox activity was said to originate from the occupancy of the 8f crystal site by Co and Mn, leading to the lower extent of cationic deficiency and hence favorable Naþ migration. Moving on, in an effort to invoke heteropolyanionic substitution strategy, PO4-SO4 solid-solution alluaudite-type cathode materials, e.g., NaxFey(PO4)3 z(SO4)z (0  z  3) have been reported.89 With the aim to produce materials with additional degree of freedom, this polyanionic substitution strategy in alluaudite provided the first material platform for a polyanionic solid solution system with distinct function as an intercalation electrode with tunable battery performance. As a whole, sulfate class of alluaudites form a milestone cathode system as earth-abundant and economically viable high-voltage insertion cathodes for sodium-ion batteries.

7.09.4

Other polyanionic cathodes

7.09.4.1

Borates

Among the polyanionic species, borates BO33 stand out as the lightest polyanionic units that in principle can deliver the highest theoretical capacity. Also, the borates can exist in widely diverse structures having various B-O coordination polyanions (e.g., BO33, BO45, B2O44, B3O63, B5O105 etc.). While LiMBO3 (M ¼ Fe, Mn, Co) borates have been reported as high capacity (ca. 200 mAh/ g) cathodes for LIBs,90 the borate insertion host for SIBs are rare. One such example is pentaborate Na3MB5O10 (M ¼ Fe, Co).91 Prepared by solid-state route, while Na3FeB5O10 assumes an orthorhombic structure (s.g. Pbca), the Na3CoB5O10 has a monoclinic structure (s.g. P21/n). These pentaborates are built from MO4 tetrahedra bridged with all four vertices to B5O105 units. The pentaborate B5O105 units are formed from BO4 tetrahedra connected to three trigonal-planar BO3. The MO4-B5O10 network forms layers stacked accommodating Naþ ions between the layers. The Fe-analogue Na3FeB5O10 works as a 2.5 V cathode, while the Na3CoB5O10 is electrochemically inactive. This low redox potential stem from weaker inductive effect of BO33 vis-à-vis PO43 and SO42 units. These pentaborates suffer from low electronic conductivity, large polarization, low ionic conductivity, and poor insertion activity. Similar to the case of LIBs, borate chemistry suffers from low energy density and rate kinetics making them unsuitable for sodium-ion batteries.

7.09.4.2

Silicates

Considering elemental abundance, silicate SiO44 based polyanionic cathodes (e.g., various polymorphs of Li2MSiO4) have been investigated for Li-ion batteries. Silicates have lower molecular weight so can give high theoretical capacity with the possibility of two-electron reaction. However, they exhibit lower redox potential than PO4/SO4 due to weaker inductive effect. Extending the silicate chemistry to SIBs, Na2MSiO4 (M ¼ Mn, Fe, Co) phases have been reported as cathode candidates for Na-ion batteries. Na2MnSiO4 has a monoclinic structure (s.g. Pn) having isolated MnO4 tetrahedra connected to SiO4 tetrahedra having channels for Naþ ions along c-axis.92 Employing a Na[FSA]–[C3C1pyrr][FSA] (FSA ¼ bis(fluorosulfonyl)amide anion and C3C1pyrr ¼ N-methyl-Npropylpyrrolidinium cation) ionic liquid electrolyte at 25–90  C, Na2MnSiO4 worked as a 3 V insertion host with a capacity of 70 mAh/g, but with poor cycling stability. Use of 5% vinylene carbonate additive in the electrolyte led to the formation of passivation film on cathode mitigating Mn dissolution to electrolyte, thereby activating over 1 electron reaction in Na2MnSiO4 with reversible capacity approaching 210 mAh/g. The Co-analogue Na2CoSiO4 exhibits polymorphism assuming either orthorhombic (s.g. Pbca) or monoclinic (s.g. Pn) structure depending on the annealing conditions. While in monoclinic structure, Co and Si

256

Development of polyanionic sodium-ion battery insertion materials

have distinct sites, they share one crystallographic site in orthorhombic structure. These Co-silicates offer reversible capacity over 100 mAh/g and 125 Ah/g for monoclinic and orthorhombic phases respectively. Fe-based silicates can be touted as economic cathodes due to their elemental abundance. Na2FeSiO4 has a monoclinic structure (s.g. P1), which becomes irreversibly amorphous upon cycling. It is possible to isolate cubic polymorph of Na2FeSiO4 (s.g. F-43 m). While the monoclinic phase is obtained by heating at 230  C for 7 days followed by annealing at 300  C for 2 days, the cubic polymorph can be obtained by heating at 120  C for 10 h in ethylene glycol followed by heating at 500  C for 10 h. Cubic Na2FeSiO4 works as a 1.9 V insertion cathode involving Fe3þ/Fe2þ redox couple delivering a capacity of 106 mAh/g with good structural stability and cycling stability.93,94 Exploring silicate chemistry further, Na2Fe2Si2O7 orthosilicate has been reported as a potential cathode. Na2Fe2Si2O7 crystallizes into a monoclinic (s.g. P21/n) unit cell having 1D chains, built with corner- and edge-sharing distorted FeO5 bipyramids and FeO4 tetrahedra, which are further connected through Si2O7 units.95 Although having open frameworks with diffusion channels, Na2Fe2Si2O7 exhibited a meagre capacity of 20 mAh/g (theoretical capacity: 164.5 mAh/g) that can be due to poor electronic/ionic conductivity and surface side reactions. While silicate chemistry indeed offers rich structural diversity, similar to LIBs, they are not suitable for practical SIB application due to material instability, poor electrochemical activity and rate kinetics.

7.09.4.3

Alluaudites

Similar to the case of PO4 alluaudites (Section 7.09.2.4) and SO4 alluaudites (Section 7.09.3.2), it is possible to extend alluaudite chemistry to other anionic systems. While the naturally occurring alluaudite minerals are based on PO4 and AsO4 systems, synthetic alluaudites with VO4/WO4/MoO4 etc. can be designed in principle delivering variable redox potential arising from inductive effect. In this spirit, the first alluaudite type vanadates (Fe-based) Na0.70(Na0.70Mn0.30)(Fe3þ/Fe2þ)2Fe2þ(VO4)3 and (Co-based) Na2(Fe/Co)2Co(VO4)3 have been reported.96,97 The V5þO4 moieties have larger unit cell volume due to the large size of V5þ resulting in large tunnels for Liþ and Naþ migration. Till date, poor electrochemical activity is noticed for vanadate alluaudites (e.g., Na2Mn2Fe(VO4)3 offers capacity of 35 mAh/g).98 On the other hand, Na2.67Mn1.67(MoO4)3 has been reported as the first MoO4-based alluaudite cathode for SIBs made by sol-gel synthesis with final (air) annealing at 650  C.99 This Na-rich alluaudite system has two completely filled (Na1, Na2) and one partially filled (Na3) sites; all Na sites being six coordinated (Fig. 8a). It offers two-dimensional migration pathways as per bond valence sum calculation (Fig. 8b).100 Thus, Na2.67Mn1.67(MoO4)3 is a rare alluaudite offering two-dimensional Naþ migration pathways and exhibits efficient Naþ (de) insertion leading to capacity over 80 mAh.g 1 with a redox plateau  3.45 V (Fig. 8c). It forms a unique alluaudite with the capacity stemming from both cationic Mn3þ/Mn2þ redox ( 3.5 V) and anionic Mo6þ/Mo5þ redox (2.4 V) activity. It calls for exploration of various other members of alluaudites Na2 þ 2xM2 x(XO4)3 (where M ¼ 3d metals, X ¼ Mo, V, W, As) with scope to (i) discover new SIB cathode materials, (ii) tune the redox potential with polyanionic moieties, and (c) realize high capacity by combining both cationic (M) and anionic (X) redox activity.

7.09.5

Mixed polyanionic cathodes

As shown in Sections 7.09.2–7.09.4, indeed, polyanionic compounds offer rich structural diversity and playground to unveil sodium-based cathode insertion materials. Starting from 1989, polyanionic systems gained prominence with materials like Fe2(SO4)3. The polyanionic materials chemistry can be further enlarged by combining various polyanionic moieties to form mixed-polyanionic cathode compounds.30 The first such case (fluorophosphates, e.g., NaVPO4F) was reported in 2003. This section elaborates the discovery and developments of various class of mixed-polyanionic sodium battery insertion materials.

Fig. 8 (a) Crystal structure of molybdate alluaudite Na2.67Mn1.67(MoO4)3 cathode (MnO6 octahedra: pink, MoO4 tetrahedra: green, Na atoms: yellow spheres). (b) Corresponding bond valence sum map showing Naþ migration pathways showing one-dimensional channels along c axis. (c) Voltagecapacity profiles (at C/10) showing reversible Naþ (de)insertion with average redox potential at 3.45 V. From Gao, J., Zhao, P., Feng, K., Na2.67Mn1.67(MoO4)3: A 3.45 V alluaudite-type Cathode Candidate for Sodium-Ion Batteries. Chem. Mater. 2017, 29, 940–944, with permission.

Development of polyanionic sodium-ion battery insertion materials 7.09.5.1 7.09.5.1.1

257

Fluorophosphates Vanadium-based fluorophosphates

Introducing electronegative fluorine (F) anions in sodium vanadium phosphates brings in new fluorophosphate class of materials. The vanadium based fluorophosphates have three-dimensional (3D) sodium-ion intercalating structure and competent electrochemical properties with high theoretical capacity and energy density (> 500 Wh kg 1). The presence of high ionicity F anions in the anion sub-lattice structure with strong covalency stabilizes the antibonding V5þ/V4þ state through the strong inductive effect, which increases the operating voltage. The strong phosphate-metal bond (P-O) in fluorophosphates enhances the structural stability apart from reducing the menace of oxygen evolution.101 The stable (PO4)3 polyanion can reduce the volume expansion in the progression of Na (de)intercalation relative to transition metal oxides.102 Sodium vanadium fluorophosphates with the chemical formula NaVPO4F, Na3V2(PO4)2F3 and Na3(VO1 xPO4)2F1 þ 2x (0  x  1) have been widely investigated as positive electrode materials for sodium-ion batteries (NIBs) as described below. 7.09.5.1.1.1 NaVPO4F Sodium vanadium fluorophosphate, NaVPO4F (NVPF), can exist in two polymorphs: a high temperature tetragonal phase and a low-temperature monoclinic phase. The tetragonal phase is isostructural to Na3Al2(PO4)3F2 with an I4/mmm symmetry. It has a three-dimensional structure built from VO4F2 octahedra abridged to PO4 tetrahedra rendering open channels locating the Naþ ions.103 Barker group first prepared the tetragonal NaVPO4F phase by conventional solid-state synthesis.104 When implemented in full cell against hard carbon anode, it delivered a discharge capacity of 82 mAh/g at an average cell voltage of 3.7 V (vs. Na/ Naþ) involving V4þ/V3þ redox activity with distinct two-step (dis)charge profiles. The monoclinic polymorph of NaVPO4F (symmetry C2/c) was later reported (Fig. 9a). However, its exact crystal structure is yet to be solved. Following, a tavorite type NaVPO4F was reported in which the oxidation state of vanadium can be slightly higher than V3þ (Fig. 9b). Though it has open channels, it offers poor electrochemical activity (only 15% Naþ can be taken out). Many synthesis techniques have been adopted to synthesize NaVPO4F. Balaya group synthesized monoclinic NaVPO4F using soft template method using V2O3 and V2O5 as precursor. NVPF synthesized by utilizing V2O5 as precursor delivered a discharge capacity of 121 mAh/g at 1C current rate with a discharge plateau at 3.33 V. It delivered excellent cyclability retaining 81% of discharge capacity at the end of 10,000 cycles running at 10C rate. Negligible volumetric changes were observed by ex situ field emission scanning electron microscopy. This system suffers from low electronic conductivity that can be improved by carbon coating to obtain superior electrochemical performance. Following, a bottom-up synthesis route based on hydrogen bonds was explored to synthesize NaVPO4F/C cathode material. The resulting nano-sized three-dimensional coral-like structures exhibited excellent rate capability and cycling stability. A capacity of 88 mAh g 1 was obtained at 50C rate, with 70% capacity retention after 2500 cycles at 5C rate (Fig. 9c).

Fig. 9 Crystal structure of (a) monoclinic NaVPO4F and (b) triclinic tavorite NaVPO4F. (c) Cycling stability of NaVPO4F upto 2500 cycles at 5C rate showing 70% of capacity retention. (d) (Dis)charge profile of monoclinic NaVPO4F at 0.5C rate when cycled between 2.0 and 4.3 V. (e) In-siu XRD pattern of NaVPO4F during (dis)charge. From panels (a, d, e) Ling, M., Zhang, H., et al., Superior Na-Storage Performance of Molten State Blending Synthesized Monoclinic NaVPO4F Nanoplates for Na-Ion Batteries. J. Mater. Chem. A., 2018, 6, 24201–24209, with permission. Panel (b) Boivin, E., Croguennec, L., et al., Vanadyl-Type Defects in Tavorite-Like NaVPO4F: From the Average Long Range Structure to Local Environments. J. Mater. Chem. A. 2017, 5, 25044–25055, with permission. Panel (c) Feng, P. Jiang, K., et al., A 3D coral-like Structured NaVPO4F/C Constructed by a Novel Synthesis Route as High-Performance Cathode Material for Sodium-Ion Battery. Chem. Engg. J., 2018, 353, 25–33, with permission.

258

Development of polyanionic sodium-ion battery insertion materials

Monoclinic NaVPO4F/C prepared by molten-state-blending technique showed high crystallinity, high thermal stability, and good electron/Naþ transport.105,106 When cycled between 2 and 4.3 V, discharge capacity of  130 mAh/g was obtained at 0.5C rate with a voltage plateau at 3.4 V (Fig. 9d). The reversible structural change during (de)intercalation was examined using in-situ XRD study (Fig. 9e). It exhibited a discharge capacity of 112.1 mAh/g at fast rate of 30C. It offered stable cycling performance for 1500 cycles at 20C rate with capacity fading of just 0.0064% per cycle. While carbon coating by various synthesis protocols can enhance the electronic conductivity and hence final capacity, it is also possible to incorporate small amount of metal-ion (Cr3þ, Al3þ) into the V sites to improve the cyclability of monoclinic NaVPO4F phase.106 7.09.5.1.1.2 Na3V2(PO4)2F3 As discussed in Section 7.09.2.1, NASICON type compounds are widely explored for sodium-ion batteries, where Na3V2(PO4)3 is a key member. Substituting one PO43 units by three fluorine atoms (for charge balance), Na3V2(PO4)2F3 can be designed as a vanadium based fluorophosphate having high theoretical capacity involving three electron transfer reaction. Due to multiple electron reaction, it has theoretical energy density (507 Wh kg 1) close to commercial LiFePO4.107 It assumes a tetragonal framework (s.g. P42/mnm) built with V2O8F3 bioctahedra linked to PO4 tetrahedra units to form a three-dimensional network with Naþ diffusion pathways along (110) and (100) directions. While the presence of F enhances the voltage, the PO4-group renders structural/ thermal stability and long term cyclability.108 The structure of Na3V2(PO4)2F3 varies with temperature, forming more symmetric tetragonal framework at higher temperature with complete randomness on Na sites. Na3V2(PO4)2F3 works as an SIB cathode involving two voltage plateaus at 3.7 and 4.2 V. During (de)sodiation, multiple phases appear at different states (Fig. 10a).109 Interestingly, only one of these phases exhibit solid-solution process in the range of x ¼ 1.8 to 1.3 during Naþ deinsertion, while the completely discharged product has a Cmc21 symmetry. Assuming 3 electron transfer, Na3V2(PO4)2F3 can deliver a theoretical discharge capacity of 256 mAh/g but only 1 electron transfer capacity can be practically realized. The capacity can be improved by doping of electrochemically inactive Ga3þ ions in the V3þ sites. The resulting Na3GaV(PO4)2F3 delivered a capacity of 141 mAh/g.110 In an effort to improve the energy density of Na3V2(PO4)2F3 by activating the third Naþ ion present in the structure, Tarascon group have demonstrated the removal of third Naþ ion from the pristine structure during charge leading to formation of

Fig. 10 (a) Galvanostatic intermittent titration curves of Na3V2(PO4)2F3 during the first cycle between Na3V2(PO4)2F3 and Na2V2(PO4)2F3 and the second one is galvanostatic cycling between Na2V2(PO4)2F3 and NaV2(PO4)2F3. The single-phase compositions are highlighted by colored circles. (b) First cycle activation of Na3V2(PO4)2F3 charged till 4.8 V and then the charge is controlled by limiting Dx(Na) ¼ 2, 2.25, 2.5, 2.75, 3.0 extracted followed by discharge till 1 V at C/10 rate. (c) Intermediate crystal structures of NaxV2(PO4)2F3 at different states of charge. From panel (a) Broux, T. Croguennec, L., et al., High Rate Performance for Carbon-Coated Na3V2(PO4)2F3 in Na-Ion Batteries. Small Methods, 2019, 3, 1800215, with permission. Panel (b and c) Yan, G., Tarascon, J. M., et al., Higher Energy and Safer Sodium Ion Batteries Via an Electrochemically Made Disordered Na3V2(PO4)2F3 Material. Nat. Commun., 2019, 10, 585, with permission.

Development of polyanionic sodium-ion battery insertion materials

259

a disordered phase of tetragonal symmetry.111 It can uptake three Naþ ions during subsequent discharge when cycled in voltage range of 1–4.8 V with the last sodium ion being inserted at 1.6 V (vs. Na/Naþ) (Fig. 10b). During this process, Naþ distribution varies widely during different states of charge with rich structural diversity (Fig. 10c). The electrochemical performance of Na3V2(PO4)2F3 is limited by low intrinsic electrical conductivity and large particle size, which can be improved by carbon-coating to increase electronic conductivity, particle downsizing to reduce diffusion length, and/or alkali/metal-ion doping to broaden the diffusion pathway in the structure. Carbon-coated Na3V2(PO4)2F3 has been shown to exhibit a discharge capacity of 130 mAh/g with cycling stability over 3000 cycles. 7.09.5.1.1.3 Na3(VO1 xPO4)2F1 þ 2x Oxygen substitution in Na3V2(PO4)2F3 can produce a new family of vanadium oxyfluorophosphates Na3(VO1 xPO4)2F1 þ 2x (0  x  1) having both V3þ and V4þ (VO2þ) states and the intermediate V3 þ/4 þ mixed valence phases. Among them, the Na3(VO)2(PO4)2F is attractive with its high energy density and stable cycling. Na3(VO)2(PO4)2F has a tetragonal framework (s.g. I4/mnm or P42/mnm).112–114 Na3V2(PO4)2F3 and Na3(VO)2(PO4)2F have similar frameworks where one of the F atoms gets replaced by an oxygen atom. Na3(VO)2(PO4)2F exhibits two step voltage plateaus at 3.6 and 4.0 V as per Eqs. (1) and (2) with a capacity of 87 mAh/g (at C/100). Plateau _ : Na3 V 2 O2 ðPO4 Þ2F4Nax V 2 O2 ðPO4 Þ2F þ xNaþ þ xe ð2  x  3Þ

(1)

Plateau __ : Na2 V 2 O2 ðPO4 Þ2F4Nax V 2 O2 ðPO4 Þ2F þ xNaþ þ xe ð1  x  2Þ

(2)

Various compositions of Na3(VO1 xPO4)2F1 þ 2x cathode materials can be prepared. With good control over particle size and carbon coating, high capacity (ca.  100 mAh/g) and cycle retention can be obtained. The Naþ-ion (de)intercalation mechanism in Na3(VO)2(PO4)2F involves both solid-solution and two-phase activity.115,116 The redox behavior can vary with the amount of F. Na3(VO1 xPO4)2F1 þ 2x has remarkable energy densities of  520 Wh kg 1 with ultrahigh stability among known SIB cathodes. Even though V has multiple oxidation states, high-voltage (de)sodiation is difficult to realize due to electrolyte instability. All these vanadium based fluorophsophates suffer from low intrinsic electronic conductivity that can be improved by carbon additives (CNTs, graphene oxide, rGO, conducting carbons etc.), surface coatings, metal substitution and nanoscaling. Na3(VO)2(PO4)2F can be implemented in full cell with hard carbon giving specific capacity of 120 mAh/g with average working voltage of 3.1 V.117

7.09.5.1.2

Other fluorophosphates

While vanadium-based compounds offer structural diversity and high voltage cathode performance, the quest for economic and non-toxic cathodes has led to the exploration of other earth-abundant transition metals-based fluorophosphates. Exploiting the inductive effect of F species, these fluorophosphates (Na2MPO4F, M ¼ Fe, Co, Mn, Ni) can be potential SIB cathodes. Economic Fe-based fluorophosphate Na2FePO4F and Na2CoPO4F were first reported by Nazar group.118,119 Isostructural to Na2FePO4OH hydroxyphosphate, they assume orthorhombic structure (s.g. Pbcn) those consist of face-sharing MO4F2 octahedra connected via bridging F-atoms giving rise to M2O7F2 bioctahedra units (Fig. 11a and b).120 These bioctahedra units are connected by PO4 tetrahedra to form infinite MPO4F slabs having Naþ ions in two distinct crystallographic sites. Using nanoscale Na2FePO4F prepared by ionothermal synthesis, Tarascon group reported its electrochemical activity with an initial discharge capacity of 120 mAh/g (1 electron theoretical capacity 124 mAh/g 1).121 Variety of synthesis protocols can be used to control the particle size, morphology, and apply carbon coating to enhance the Naþ (de)insertion activity in this system with an average Fe3þ/Fe2þ redox potential located at 3 V (Fig. 11d). A careful look shows the presence of two sets of redox plateaus during cycling of Na2FePO4F involving solid-solution mechanism.122,123 It can be due to different site energy of Naþ cations leading to different insertion voltage and an ordering behavior. While in principle Na2FePO4F can yield two-electron reaction (Fe2þ to Fe4þ), the high potential of Fe4þ/Fe3þ redox couple make it difficult to extract the second Naþ out of the structure. Therefore, the capacity of Na2FePO4F is limited. Isostructural to Na2MgPO4F and Na2FePO4F, Komaba group first reported the electrochemical activity of solid-state prepared Na2CoPO4F.124 It works as a high-voltage (ca. 4.4 V) sodium insertion material with a charge capacity of 213 mAh/g corresponding to extraction of 1.74 Naþ. However, a discharge capacity of only 71 mAh/g was observed with huge polarization. Large irreversible capacity loss was attributed to decomposition of electrolytes at higher voltages, which can be minimized by adding 2% of fluoroethylene carbonate (FEC) additive in electrolytes. Following, Yang group reported spray-drying synthesis of Na2CoPO4F with discharge capacity of 107 mAh/g at an average Co3þ/Co2þ redox potential of 4.5 V (Fig. 11e).125 However, severe capacity fading was observed in NCPF due to electrolyte decomposition at higher voltages. Going by the trend, the isostructural Na2NiPO4F can give even higher redox voltage (ca. > 5 V as per DFT calculations). However, both synthesis of pure end-member and its electrochemical performance has not been achieved till date due to formation of Ni-based impurity and non-availability of highvoltage electrolytes respectively.119 Deviating from all previous orthorhombic Na2MPO4F systems, Na2MnPO4F assumes a monoclinic structure with P21/n symmetry (Fig. 11c).121 It consists of corner-shared MnO4F2 octahedra connected via F-atoms to form Mn2O8F2 chains. These chains are interlinked via PO4-group giving rise to a three-dimensional framework. Based on poor Na-ion diffusion kinetics, this material was found to be electrochemically inactive despite of having an open pathway for diffusion. To realize electrochemical activity, carbon-coated nano-sized Na2MnPO4F was found to be active at 60  C with a discharge capacity of 98 mAh g 1, albeit

260

Development of polyanionic sodium-ion battery insertion materials

Fig. 11 Structural illustration of fluorophosphates (a) orthorhombic Na2FePO4F (FeO4F2 octahedra: blue, PO4 tetrahedra: pink, Na atoms: yellow and orange spheres), (b) orthorhombic Na2CoPO4F (CoO4F2 octahedra: brown, PO4 tetrahedra: blue, Na atoms: yellow and green spheres), (c) monoclinic Na2MnPO4F (MnO4F2 octahedra: pink, PO4 tetrahedra: blue, Na atoms: yellow spheres), Galvanostatic potential-capacity profiles of (d) Na2FePO4F (inset: corresponding dQ/dV profiles and impedance spectra at different voltages), (e) Na2CoPO4F and (f) Na2MnPO4F. From panel (e) Zou, H. Yang, Y., et al., Spray-Drying Synthesis of Pure Na2CoPO4F as Cathode Material for Sodium Ion Batteries. ECS Electrochem. Lett., 2015, 4, A53, with permission. Panel (f) Lin, X. Yang, Y., et al., Exploiting Na2MnPO4F as a High-Capacity and Well-Reversible Cathode Material for Na-Ion Batteries. RSC. Adv., 2014, 4, 40985–40993, with permission.

with poor cycling stability (Fig. 11f).126 Overall, Na2MPO4F class of fluorophosphates can deliver high voltage operation with suitable electrolytes. Among them, the iron analogue Na2FePO4F forms the most feasible SIB cathode insertion material.

7.09.5.2 7.09.5.2.1

Mixed phosphates [(PO4)(P2O7)] Na4M3(PO4)2P2O7 (M ¼ Mn, Fe, Co, Ni)

Based on the corner or edge-sharing coordination of PO4 anions, diverse structural building blocks can be conceived with diverse orientation (or bending) to form polyphosphate units like pyrophosphates (P2O7), metaphosphates (P3O9) and so on. Combination of different PO4-based building blocks, mixed phosphates can be designed with diverse structural and electrochemical properties. When isolated PO43 units are combined with P2O74 pyrophosphate units, mixed phosphates are formed having a general formula of Na4M3(PO4)2P2O7 (M ¼ Mn, Fe, Co, Ni).127 These isostructural Na4M3(PO4)2P2O7 polyanionic frameworks assume an orthorhombic framework (s.g. Pn21a). Its structure is built from a three dimensional (3D) network of [M3P2O13]N blocks parallel to the bc plane (Fig. 12a). These [M3P2O13]N blocks are connected by pyrophosphate (P2O74) groups offering long tunnels along b axis favoring Naþ diffusion. This structure has four different interlinked Na sites. Na1 and Na4 are located along the bc plane made by NaO6 octahedra, whereas Na2 and Na3 sites have seven coordinated NaO7 polyhedra and NaO6 octahedra along the a-axis. Bond valence sum and DFT calculations point out that Naþ can move sinusoidally along the b-axis (Fig. 12b), similar to the case of Liþ migration in triphylite LIB cathode LiFePO4.128 The Na4M3(PO4)2P2O7 mixed phosphate materials are attractive as cathode materials for SIBs due to their high theoretical capacity (ca. 170 mAh g 1) (extraction/insertion of 3 Naþ), open frameworks with three-dimensional Naþ diffusion pathways as revealed by Bond Valence Sum analysis (BVS), high redox potentials induced by inductive effect of PO4 and P2O7 anions, broad selectivity in forming solid solution phases of central transition metal and easy scalable synthesis. Kang group first demonstrated the electrochemical activity of solid-state prepared Na4Fe3(PO4)2P2O7.129,130 It delivered about 88% of the 1 electron theoretical capacity involving reversible removal of 3 Naþ with an average Fe3þ/Fe2þ redox potential centered at 3.2 V (vs. Na/Naþ). Implementing solution combustion route, carbon coated Na4Fe3(PO4)2P2O7 was shown to deliver capacity over 100 mAh/g at a fast rate of C/10 with a stepwise voltage profile (Fig. 12d). The P2O7 dimer undergoes distortion during the cycling involving a solid-solution (single-phase) (de)sodiation mechanism with minimal volume change (ca. 4%). Isostructural Na4Co3(PO4)2P2O7 has also been unveiled as a suitable SIB cathode exhibiting a reversible capacity of 95 mAh/g in the high potential region of 4.1–4.7 V involving a multiple step wise voltage profile (Fig. 12e).131 It retained 84% of reversible capacity over 100 cycles keeping the 4.5 V (vs. Na/Naþ) redox potential. Na4Co3(PO4)2P2O7 mixed phosphate cathode can be

Development of polyanionic sodium-ion battery insertion materials

(a)

(b)

M(1)O6

(c)

Na4

Na1

[PO4]3-

261

NaO6 Na2

Na Fe P O

M(2)O6

Na3

c

b

[P2O7]4-

Na4Fe3(PO4)2P2O7

(d) 4

NaO5

M(3)O5

b

Poten al (V vs. Na/Na+)

NaO7

Na4Co3(PO4)2P2O7

(e) 5

a

Na4Mn3(PO4)2P2O7

(f)

4

4 4

3

3 3

2

3 1 cycle

2

1 2

cycles 10-1

0

20

40

60

80

100 120

Capacity (mAh/g)

Na7V4(P2O7)4PO4

(g)

5

0

20

40

60

80

100 120

Capacity (mAh/g)

120

90

60

30

Capacity (mAh/g)

0

0

20

40

60

80

Capacity (mAh/g)

Fig. 12 (a) Structural illustration of family of mixed phosphates Na4M3(PO4)2P2O7 (MO6 octahedra: brown, MO5 pentahedra: brown, PO4 tetrahedra: pink, Na atoms: yellow), (b) Bond valence sum map showing sinusoidal Na4 and Na1 diffusion pathways along b-direction, (c) Structural illustration of family of mixed phosphates Na7V4(P2O7)4PO4 (VO6 octahedra: blue, PO4 tetrahedra: pink, Na atoms: yellow), Galvanostatic potential-capacity profiles of mixed phosphate cathodes (d) Na4Fe3(PO4)2P2O7, (e) Na4Co3(PO4)2P2O7, (f) Na4Mn3(PO4)2P2O7 and (g) Na7V4(P2O7)4PO4. From panel (e) Nose, M. Iba, H., et al., Na4Co3(PO4)2P2O7: A Novel Storage Material for Sodium-Ion Batteries. J. Power Sources, 2013, 234, 175–179, with permission. Panel (f) Kim H. Kang, K., et al., Anomalous Jahn-Teller Behavior in a Manganese-Based Mixed-Phosphate Cathode for Sodium Ion Batteries. Energy Environ. Sci. 2015, 8, 3325–3335, with permission. Panel (g) Lim, S. Y. Choi, J. W., et al., Role of Intermediate Phase for Stable Cycling of Na7V4(P2O7)4PO4 in Sodium Ion Battery. PNAS, 2014, 111, 599–604, with permission.

used against hard carbon anode to form full cell delivering excellent capacity and cyclability. Exploring mixed phosphate family further, Kang group reported the Mn analogue Na4Mn3(PO4)2P2O7 realizing  100% theoretical capacity with excellent cycling stability at a rate of C/20 (Fig. 12f). With its high redox potential (Mn3þ/Mn2þ redox potential at 3.84 V), it can lead to high energy density (416 Wh kg 1).132 The unique Jahn–Teller distortion in this material initiates Naþ diffusion channels as well as provides high-power capability and cycling stability. It also involves multiple step voltage profile mostly dominated by solid-solution redox mechanism with small volume change (about 7%) suitable for long term cycling stability. The Ni analogue Na4Ni3(PO4)2P2O7 was the last member in this family that was found to be moderately active. It was found to involve competitive two-phase reactions followed by single phase reaction by adjusting residual Naþ sites uniformly. The structural rearrangement during Naþ extraction can lead to more stable phase, and reversible cycling. Unlike other phosphates, Na4Ni3(PO4)2P2O7 registers a rare case of electrochemically active Ni-based polyanionic cathode for SIBs.133 The mixed phosphate Na4M3(PO4)2P2O7 (M ¼ Mn, Fe, Co, Ni) family of cathodes uniquely offer a system with all members being electrochemically active. With stepwise voltage profiles involving complex (de)sodiation mechanism, they deliver high voltage operation (3.2–4.8 V) that can arise due to inductive effect due to encompassing of M octahedra by the electron-drawing P2O7 groups the edge-sharing geometry of MO6 polyhedra leading to strong M3 þ–M3 þ repulsion in the desodiated structure that destabilizes the charged state thereby enhancing the redox voltage. As all members are active, it is also possible to form various mixed metal mixed phosphates (e.g., Na4Fe3 xMnx(PO4)2P2O7, Na4Co x yMnxNiy(PO4)2P2O7) to obtain SIB cathode systems with tuneable voltage and high energy-density.134

7.09.5.2.2

Na7V4(P2O7)4PO4

Crystal structure and cation transport of Na7M(P2O7)4PO4 (M ¼ Fe, Al, Cr) were first reported by Bretey group.135 Isostructural Na7V4(P2O7)4PO4 mixed phosphate with a tetragonal structure (s.g. P-421c) was later reported.136 It has a three-dimensional skeleton of [V4(P2O7)4PO4]N built from pyrophosphate P2O7 group with two adjacent VO6 octahedra and PO4 tetrahedron with four adjacent VO6 octahedra in corner sharing fashion. Three Naþ ions occupied three distinct interstitial sites with suitable diffusion pathways (Fig. 12c). Involving a biphasic V4þ/V3þ redox activity centered at 3.88 V (vs. Na/Naþ), Na7V4(P2O7)4PO4 exhibited a reversible capacity approaching 80 mAh/g (Fig. 12g). It formed Na5V3.5 þ4(P2O7)4PO4 as an intermediate phase utilizing V4þ/ V3þ redox reaction, resulting in two plateaus in (dis)charge curves. The existence of intermediate phases favors better kinetics by lowering down the reaction barriers between two end members. With conductive carbon coating, Na7V4(P2O7)4PO4 exhibited excellent cycling stability retaining 94% capacity over 800 cycles with high-rate kinetics. Vanadium-based Na7V4(P2O7)4PO4 mixed

262

Development of polyanionic sodium-ion battery insertion materials

phosphate forms a high-voltage SIB cathode with excellent electrochemical performance, which can be improved by optimization of morphology (e.g., one-dimensional nanowires, three-dimensional hybrid foams), particle downsizing and carbon coating.

7.09.5.3

Fluorosulfates

As described in Section 7.09.3, as per inductive effect principle, sulfate class of cathodes can exhibit the highest redox potential among all polyanionic families. The inductive effect can be enhanced by adding electronegative F species to SO42 to form fluorosulfate mixed polyanionic systems. This concept has been demonstrated by Tarascon group developing high voltage (ca. 3.6– 3.9 V) LiMSO4F cathodes for LIBs.75,137,138 Extending the fluorosulfate chemistry to Na-based systems, analogous NaMSO4F materials can be prepared by low temperature (T < 300  C) solid-state and non-aqueous solvothermal topotactic reaction between NaF and MSO4$H2O monohydrate precursors. With the large ionic size of Naþ, NaMSO4F fluorosulfates adopt a maxwellite-type monoclinic structure (s.g. C2/c).139,140 The maxwellite-type NaMSO4F structure is built with corner-sharing MO4F2 octahedral chains abridged by SO4 tetrahedra along [001] having a distinct site for Naþ (Fig. 13a). NaMSO4F fluorosulfate family has two anomalies: (a) observation of Jahn-Teller instability in NaCuSO4F with highly distorted CuO4F2 octahedra, and (b) deviation from maxwellite to triplite structure in case of NaMnSO4F. While the NaMSO4F phases are electrochemically inactive, the Fe-analogue NaFeSO4F shows slight degree of Naþ (de)insertion with a distinct plateau at 3.5 V (vs. Na/Naþ) (Fig. 13b). Atomistic modeling revealed three root causes behind the poor electrochemical activity: (a) existence of one-dimensional zigzag Naþ diffusion along [101] direction vs. the three-dimensional Liþ diffusion in LiMSO4F, (ii) large volume change during (de)sodiation (ca. 16%), and (iii) high activation energy barrier (ca. 0.9 eV) hindering Naþ diffusion.141 Despite several cathode optimization (e.g., particle downsizing), there is no improvement in their electrochemical activity. In search of suitable Naþ insertion host, potassium fluorosulfate analogues KMSO4F were prepared by topotactic reaction of KF with MSO4$H2O monohydrates. As the cationic size of Kþ is large, it develops different structure than Li- and Na-analogues.142,143 The KMSO4F fluorosulfates assume KTiOPO4 (KTP)-type orthorhombic structure (s.g. Pna21) with distorted MO4F2 octahedral chains abridged by SO4 tetrahedra, where the F species have an alternate cis and trans positions (Fig. 13c). The buckling of these chains along [011] and [0-11] directions develop large cavities hosting the Kþ ions. Chemical oxidation of KFeSO4F led to an orthorhombic FeSO4F polymorph (s.g. Pnna) with large open cavities. This FeSO4F favors efficient Naþ (de)intercalation delivering a capacity over 120 mAh/g involving a flat step-wise voltage profile centered at 3.5 V (Fig. 13d). While maxwellite NaFeSO4F is electrochemically inactive, (de)potassiated KFeSO4F can work as a high voltage SIB cathode.

(c)

(a)

NaFeSO4F 5 4 3 2 0.92

0.96

x in NaxFeSO4F

KFeSO4F

(d)

Potential (V vs. Na/Na+)

Potential (V vs. Na/Na+)

(b)

1.0

4

3

2 0.2

0.4

0.6

0.8

x in KxFeSO4F

Fig. 13 Structural illustration of fluorosulfates: (a) maxwellite NaFeSO4F (FeO4F2 octahedra: green, SO4 tetrahedra: yellow, Na atoms: yellow), (c) KTP-type KFeSO4F (FeO4F2 octahedra: green, SO4 tetrahedra: yellow, K atoms: violet). Galvanostatic potential-capacity profiles showing Naþ (de) insertion into (b) maxwellite NaFeSO4F and (d) KTP-type KFeSO4F cathodes. From Barpanda, P., Sulfate Chemistry for High-Voltage Insertion Materials: Synthetic, Structural and Electrochemical Insights. Isr. J. Chem., 2015, 55, 537–557, with permission.

Development of polyanionic sodium-ion battery insertion materials

263

The presence of SO4 and F units make these fluorosulfates prone to moisture attack leading to the formation of hydrated derivatives. Such hydrated fluorosulfates NaMSO4F$2H2O can be prepared by dissolution and precipitation route assuming an uklonskovite structure (monoclinic, s.g. P21/m) built from parallel chains of MO4F2 octahedra where two O atoms are bridged to H atoms to form structural H2O units.144 Although having open cavities favoring Naþ migration, these bihydrated fluososulfate NaMSO4F$2H2O materials are found to be electrochemically inactive.

7.09.5.4

Hydroxyosulfates

The moisture resistance in sulfate compounds can be improved by switching from F to OH. This substitution can lead to hydroxysulfates (NaMSO4.OH) class of mixed polyanionic insertion materials. Nonetheless, synthesis of hitherto unknown NaMSO4OH family leads to mixture of impurity phases. Rather, various other hydroxysulfate phases have been reported like Na0.84Fe2.86(SO4)2(OH)6 (jarosite), NaFe3(SO4)2(OH)6 (natrojarosite), Na2FeOH(SO4)2.H2O (metasideronatrite), Na2FeOH(SO4)2.3H2O (sideronatrite) and Na2Cu(SO4)2.xH2O (sideronkite). Pralong group reported NaFe3(SO4)2(OH)6 as the first hydroxysulfate based SIB cathode materials prepared by precipitation of solution containing Na2SO4 and Fe2(SO4)3.nH2O.145,146 This jarosite compound assumes a trigonal structure with R-3m symmetry having layered framework built from corner-sharing FeO4(OH)2 octahedral chains. These chains are in turn connected by SO4 tetrahedra to form hexagonal voids accommodating Naþ ions. This host framework has ample cavities to insert multiple alkali ions. It can reversibly uptake of 2 Naþ (per f.u.) leading to capacity of 120 mAh/g with an average Fe3þ/Fe2þ redox potential of 2.72 V. Reversible cycling was observed involving sloping voltage profiles with monophasic redox mechanism. Interestingly, Naþ insertion transforms the crystalline jarosite into an amorphous Na3Fe3(SO4)2(OH)6 phase indexed to a hexagonal cell. The jarosite building layers [Fe3(SO)4(OH)6] are very thin, which can be distorted, corrugated and separated from neighboring layers during sodiation while keeping all the polydehral features and interatomic bonds intact. In contrast, desodiation relaxes the local strains leading back to initial jarosite crystalline framework. While far from practical application due to low energy density, the jarosite hydroxysulfate system paves way for exploration of amorphous polyanionic phases as possible insertion host for SIBs.

7.09.5.5

Carbonophosphates

In search of new battery cathode materials, it is possible to employ classical solid-state chemistry knowledge along with mineralogical database and computational tools (e.g., Materials Project). Using computational search of materials database, Ceder group unveiled carbonophosphates (Na3MPO4CO3, M ¼ Mn, Fe) family of mixed polyanionic cathode materials for Na-ion batteries.147 Similar to BO3 moieties, CO3 polyanion moieties exist in triangular planar units, which can be combined with tetrahedral PO4 to stabilize diverse crystal structure. Prepared by hydrothermal route, Na3MnPO4CO3 assumed a monoclinic framework (s.g. P21/m) related to sidorenkite mineral structure. Its structure is built as a layered system consisting of isolated MnO6 octahedra connected to four PO4 tetrahedra by corner-sharing manner, while sharing one edge with planar CO3 units (Fig. 14a). With the presence of both corner-sharing and edge-sharing coordination, the MnO6 octahedra are highly distorted. The Naþ species occupy two distinct crystallographic sites with feasible migration pathways. Sidorenkite Na3MnPO4CO3 works as an SIB cathode involving two distinct plateaus corresponding to Mn3þ/Mn2þ and Mn4þ/ Mn3þ redox activity with an average potential around 3.7 V (vs. Na/Naþ). While it is redox active, it can (de)intercalate Naþ only at a slow rate (ca. C/100) delivering a stable capacity over 125 mAh/g (theoretical capacity ¼ 191 mAh/g) (Fig. 14b). Although the structure has open pathways for three-dimensional Naþ migration, similar to many Mn-based cathodes, it inherently suffers from low electronic conductivity that limits its activity. Here, carbon coating on sidorenkite can be performed using carbon nanostructures (graphene oxide, CNTs etc) improve the capacity over 176 mAh/g. This study has been extended to isostructural

Fig. 14 (a) Structural illustration of sidorenkite class of carbonophosphate Na3MnPO4CO3 along (left) ab plane and (right) ac plane (MnO6 octahedra: red, PO4 tetrahedra: blue, CO3 triangular units: black, Na atoms: green and yellow), (b) Corresponding galvanostatic potential-capacity profiles showing Naþ (de)insertion at a rate of C/100. From Chen, H. Ceder, G., et al., Sidorenkite (Na3MnPO4CO3): A New Intercalation Cathode Material for Na-Ion Batteries. Chem. Mater. 2013, 25, 2777–2786, with permission.

264

Development of polyanionic sodium-ion battery insertion materials

bonshtedtite mineral type Na3FePO4CO3 carbonophosphate with nano-platelet morphology.148 With nanoscale morphology, this Fe-based cathode can (de)intercalate Naþ with reversible capacity  100 mAh/g involving both Fe3þ/Fe2þ and Fe4þ/Fe3þ activity with an average redox potential of 2.6 V (vs. Na/Naþ). Carbonophosphates can be considered as economic cathode with high thermal/chemical stability with potential to involve 2-electron reaction leading to high capacity. Nonetheless, effort should be geared to improve their conductivity, rate kinetics and cycling stability.

7.09.5.6

Nitridophosphates

As discussed in Section 7.09.2.4, metaphosphate frameworks can work as SIB insertion systems. Combining metaphosphate (P3O9) with electronegative nitrogen (N), nitridophosphate class of mixed polyanions can be designed with potential high voltage operation owing to inductive effect and high energy density due to the high mass per charge ratio. Two such SIB cathodes, Na3Ti(PO3)3N and Na3V(PO3)3N (and its solid-solution) have been reported till date.149–151 This nitridophosphates have a CUBICON structure, i.e., cubic system (s.g. P213) built with MO6 octahedra (M ¼ Ti, V) corner-sharing with three neighboring PO3N to form (PO3)3N6 building blocks. Constituent Naþ occupy three distinct sites with open migration pathways (Fig. 15a). The Ti-analogue Na3Ti(PO3)3N delivers capacity of 67 mAh/g (theoretical capacity ¼ 74 mAh/g) involving Ti4þ/Ti3þ redox activity centered at 2.7 V (Fig. 15b). It involves single-phase redox mechanism having negligible volume change (z0.5%). This near zero strain is similar to many Ti-based insertion materials like Li4Ti5O12. Compared to other Ti-based systems involving Ti4þ/Ti3þ redox activity, Na3Ti(PO3)3N exhibits the highest Ti redox potential (ca. 2.7 V) among all known Ti-based compounds, which can arise from its unique structure and inductive effect of N. While the voltage is high, this nitridophosphate compound has limited net theoretical capacity (and hence energy density) with weight penalty of metaphosphate units. One approach to enhance the energy density is by realizing high-voltage operation involving transition metal with multiple redox activity. In this spirit, multivalent transition metal (V) based nitridophosphate Na3V(PO3)3N has been reported. Isostructural to the Ti analogue, it offers highvoltage (ca. 4 V) Naþ (de)intercalation involving monophasic V4þ/V3þ redox activity with minimal volume change (0.25%) (Fig. 15c). Though it offers high-voltage operation and possible multiple electron reaction, its capacity is limited below 50 mAh/g. With its tricky synthesis involving ammonolysis and limited capacity, the nitridophosphate cathodes offer poor scope for practical application.

7.09.5.7

Oxalate derivatives

In an effort to produce sustainable cathodes, there has been attempts to design oxalate (C2O4) based compounds. One such organic-based polyanionic system is iron-oxalate K4Na2[Fe(C2O4)2]3$2H2O. Its open framework structure has FeO6 octahedra connected to six oxalate (C2O4) units via oxygen atoms. While the first charge triggers depotassiation creating open voids, the subsequent cycles undergo reversible Naþ (de)insertion along one-dimensional channels involving redox plateau at 2.7 V (vs. Na/Naþ) with a moderate capacity of 50.2 mAh/g.152 The presence of multiple oxalate and structural water involves huge weight penalty limiting its capacity. Following this study, oxalate type mixed polyanion Na2Fe(C2O4)F2 cathode was synthesized via a hydrothermal route. This fluoro-oxalate structure assumes a monoclinic structure (s.g. C2/c) built with FeO4F2 octahedra, which are connected via their oxygen vertices to oxalate units. Each oxalate group in turn is linked to three FeO4F2 octahedra forming zig-zag chains of Fe(C2O4)F 2 units along the b-axis. (Fig. 16a). The Na atoms are located between the chains. This system works as an SIB cathode with two Fe3þ/Fe2þ redox plateaus on charge (3.3 and 3.6 V vs. Na/Naþ) and two on discharge (2.95 and 3.25 V vs. Na/Naþ) with large polarization.153 The presence of F species can be the reason behind this high voltage. Nonetheless, with the weight penalty of oxalate units, it offers a discharge capacity of 70 mAh/g corresponding to the reversible extraction of 0.56 Naþ.

Fig. 15 (a) Structural illustration of nitridophosphate Na3M(PO3)3N (M ¼ Ti, V) (MO6 octahedra: blue, PO4 tetrahedra forming trimers of PO3N: green, Na atoms: gray, pink and cyan). Galvanostatic potential-capacity profiles showing Naþ (de)insertion into (b) Na3Ti(PO3)3N and (c) Na3V(PO3)3N cathodes. From panel (a and b) Liu, J., Khalifah, P., et al., Ionic Conduction in Cubic Na3TiP3O9N, a Secondary Na-Ion Battery Cathode With Extremely Low Volume Change. Chem. Mater., 2014, 26, 3295–3305, with permission. Panel (c) Reynaud, M. Casas-Cabanas, M., et al., Sodium Vanadium Nitridophosphate Na3V(PO3)3N as a High-Voltage Positive Electrode Material for Na-Ion and Li-Ion Batteries. Electrochem. Commun. 2017, 84, 14–18, with permission.

Development of polyanionic sodium-ion battery insertion materials

265

Fig. 16 Structural illustration of oxalate class of mixed polyanionic cathodes (a) Na2Fe(C2O4)F2 (FeO4F2 octahedra: blue, C2O4 units shown by red dashed line, Na atoms: green) and (b) Na2Fe2(C2O4)3.2H2O (FeO6 octahedra: green, C2O4 units shown by red dashed line, Na atoms: blue). The corresponding potential-capacity profiles showing Naþ (de)insertion are shown in the bottom panel. From panel (a) Yao, W. et al., Armstrong, A.R., Na2Fe(C2O4)F2: A New Iron-Based Polyoxyanion Cathode for Li/Na Ion Batteries. Chem. Mater. 2017, 29, 2167–2172, with permission. Panel (b) Yao, W. et al., Armstrong, A.R., Na2Fe2(C2O4)2.2H2O: An Iron-Based Positive Electrode for Secondary Batteries. Chem. Mater. 2017, 29, 9095–9101, with permission.

Pursuing oxalate chemistry, a hydrated oxalate Na2Fe2(C2O4)3$2H2O has been reported by hydrothermal synthesis.154 This hydrated oxalate has a monoclinic structure (s.g. P21/c) having FeO6 octahedra chains that are connected by C2O4 oxalate units (Fig. 16b). It can reversible uptake Naþ involving multiple stepwise voltage profiles with an average Fe3þ/Fe2þ redox potential of 3.3 V giving rise to near theoretical capacity (ca. 117 mAh/g) and high energy density (ca. 325 Wh kg 1). It benchmarks among the best of Fe-based cathode materials for SIBs. The oxalate chemistry is relatively less explored with room to study various carboxylate-based compounds (e.g., malonate Na2Fe(H2C3O4)2$2H2O, Fe2(C2O4)3$4H2O etc), other 3d metal based compositions and possible exploitation of anionic redox activity involving oxalate (C2O4) units.

7.09.5.8

Phosphosulfates

While PO4-based compounds offer chemical/thermal stability, the SO4-based materials can induce high-voltage operation. Combining these units, phosphosulfate class of mixed polyanion cathode can be developed. Goodenough group used this approach to unveil NaFe2PO4(SO4)2 isostructural to hexagonal NASICON Na3Zr2PO4(SiO4)2 structure.31,155 As NASICON framework offers fast ion migration, various NASICON type cathodes and their derivatives have been examined for decades (Section 7.09.2.1). The NaFe2PO4(SO4)2 system can be synthesized by solid-state or solvothermal route with annealing temperature as
consisting low as 100  C. Due to the presence of Fe3þ species, it is high stable in air. This system has a hexagonal structure (s.g. R3c) of FeO6 octahedra corner shared with PO4 and SO4 tetrahedra units (Fig. 17a). Bond valence sum calculations predicts low migration barrier for Naþ diffusion. Indeed, it works as a sodium insertion host involving a single-phase redox mechanism locating the

266

Development of polyanionic sodium-ion battery insertion materials

(a)

(b)

Voltage (V)

Potential (V vs. Na/Na+ )

4.5 4.0 3.5 3.0 2.5 2.0

cycles 20-1

1.5 1.0 0.5 0.0 0

20

40

60

80

Capacity (mAh g-1)

100

120

Capacity (mAh/g) Fig. 17 (a) Structural illustration of NASICON type NaFe2PO4(SO4)2 mixed polyanionic cathodes. (FeO6 octahedra: brown, PO4 and SO4 tetrahedra: yellow, Na atoms: blue) Naþ migration pathways are shown. (b) Voltage-capacity profiles showing Naþ (de)insertion involving 3 V Fe redox activity.

average Fe3þ/Fe2þ redox potential located at 3.2 V (Fig. 17b). It can deliver discharge capacity of 100 mAh/g with moderate rate kinetics and cycling stability.

7.09.5.9

Phosphonitrates

Parallel to phsophosulfates in earlier section, combining PO4 with NO3 units, phosphonitrate class of compounds can be realized. In this direction, phosphonitrates AFePO4NO3 (A ¼ NH4, Li, K) have been reported.156 It must be noted that synthesis of NO3 based compounds comes with high risk of explosion (upon annealing over 200  C). Thus, large-scale production is not practical. AFePO4NO3 stabilizes into a triclinic (s.g. P-1) framework built with distorted FeO6 octahedra sharing an edge with NO3 groups. Overall structure is very close to layered carbonophosphates due to presence of triangular planar NO3 units similar to BO3 units. While NaFePO4NO3 is difficult to prepare, NH4FePO4NO3 analogue can be prepared and can reversibly (de)intercalate Naþ delivering a reversible capacity of 80 mAh/g (at a rate of C/50) involving an average Fe2þ/Fe3þ redox potential at 2.7 V (vs. Na/Naþ). While this cathode can offer 2 electron reaction, these phosphonitrates are not practical due to difficulty in synthesis, low voltage operation, sloping voltage profile with large hysteresis, and poor rate kinetics.

7.09.6

Conclusions

Sodium-ion batteries indeed form a practical post-Li-ion battery systems that has seen enormous research and development leading to its commercialization (circa 2021). Sodium insertion materials offer rich structural diversity due to the sharp contrast in size of Naþ and the 3d transition metal ions, which leads to the stabilization of a larger number of functional structures. These frameworks offer one dimensional to three-dimensional diffusion pathways favoring efficient Naþ migration. While oxide cathodes offer large capacity, polyanionic cathodes cater large database of insertion systems with many outperforming their oxide counterparts particularly concerning structural stability, high operating voltage, and operational safety. As noted in this chapter, a wide array of simple and mixed polyanionic compounds have been reported as potential cathodes for sodium-ion batteries. Their electrochemical performance (voltage, capacity) and energy density values are summarized in Fig. 18. While they can lead to high redox potential, many of them invariably suffer from limited theoretical capacity due to the weight penalty of polyanion units. Also, low electronic conductivity is a common problem in many polyanionic cathodes. This issue can be mitigated by strategies such as surface carbon coating, composite formation with carbon nanostructures, controlling particle size and morphology, iso/aliovalent metal ion doping and so on. These methods have been successfully employed in LIBs and can be extended to SIBs. One niche feature of polyanionic cathode systems is the possibility of multiple electron reaction. However, it is not a trivial feature as extraction of more than one Naþ (corresponding to multiple redox change in constituent transition metal species) can lead to significant change in bond length triggering polyhedral distortion which can be irreversible and can destroy the framework. Also, the extraction of second Naþ is often involved with very high diffusion barrier and occur at very high voltage beyond the safe operating electrolytes. In principle, multiple electron reaction can be realized by implementation of suitable electrolytes stable at high voltage. These electrolytes can also impart long cycling stability and safe operation. In this direction, various combination of electrolytes as well as superconcentrated electrolytes offer room for exploration. Among the wide gamut of polyanionic cathodes with various active redox centers, while Ti-based systems work at low voltage, Vand Co-based cathodes offer high-voltage operation close to electrolyte stability border  4.5 V. Among the large number of polyanionic compounds identified so far, vanadium-based phosphates and fluorophosphates such as Na3V2(PO4)3 and Na3V2(PO4)2F3

Development of polyanionic sodium-ion battery insertion materials

267

Fig. 18 Voltage-capacity plot summarizing the electrochemical activity of polyanionic cathode materials for sodium-ion batteries: (left) various phosphate containing cathodes and (right) various non-phosphate class of cathodes.

offer the best electrochemical performance and high energy density. Their industrial deployment is yet hampered by the toxicity and cost of vanadium. Also, phosphate class of materials like pyrophosphate Na2FeP2O7 and mixed phosphates Na4Fe3(PO4)2P2O7 can work as economic 3 V cathodes with fast rate kinetics, thermal safety, and cyclability. Sulfate based alluaudite Na2Fe2(SO4)3 forms yet another milestone material consisting of earth-abundant elements giving rise to the highest redox potential operation and hence high energy density. However, these sulfate cathodes warrant careful synthesis and storage/handling due to their hygroscopic nature. Since the inception of sodium insertion materials in 1980s, the research on SIBs has seen unprecedented growth and success leading to its commercialization by CATL.157 SIBs can ably cater small-scale transportation (e.g., Faradion in UK and Tiamat in France have demonstrated Na-ion cells in electric bikes) and mostly stationary energy storage. For practical cell assembly and usage, parallel effort should be geared towards development of anodes, binders, and electrolytes. In this SIB sector, polyanionic cathodes can potentially occupy some market share. This chapter is an effort to summarize the research work on polyanionic cathode materials for sodium-ion batteries, which offers rich playground for discovery of novel cathode materials for potential battery applications.

Acknowledgement S.S., S.P.V. and S.L. contributed equally to this work. The authors acknowledge the financial support from the Technology Mission Division (Department of Science and Technology, Govt. of India) under the Materials for Energy Storage (MES-2018) program (DST/TMD/MES/2k18/207). P.B. is grateful to Alexander von Humboldt Foundation (Bonn, Germany) for a 2022 Humboldt fellowship for experienced researchers. The crystal structures were illustrated using Vesta-3 software.158

References 1. Smalley, R. E. Future Global Energy Prosperity: The Terawatt Challenge. MRS Bull. 2005, 30, 412–417. 2. Gur, T. M. Review of Electrical Energy Storage Technologies, Materials and Systems: Challenges and Prospects for Large-Scale Grid Storage. Energ. Environ. Sci. 2018, 11 (10), 2696–2767. 3. Jurasz, J.; Canales, F. A.; Kies, A.; Guezgouz, M.; Beluco, A. A Review on the Complementarity of Renewable Energy Sources: Concept, Metrics, Application and Future Research Directions. Solar Energy. 2020, 195, 703–724. 4. Tarascon, J. M.; Armand, M. Issues and Challenges Facing Rechargeable Lithium Batteries. Nature 2001, 414 (6861), 359–367. 5. Yang, Z.; Zhang, J.; Kinter-Meyer, M. C. W.; Lu, X.; Choi, D.; et al. Electrochemical Energy Storage for Green Grid. Chem. Rev. 2011, 111 (5), 3577–3613. 6. Larcher, D.; Tarascon, J. M. Towards Greener and More Sustainable Batteries for Electrical Energy Storage. Nat. Chem. 2015, 7, 19–29. 7. Whittingham, M. S. Lithium Batteries and Cathode Materials. Chem. Rev. 2004, 104 (10), 4271–4302. 8. Nitta, N.; Wu, F.; Lee, J. T.; Yushin, G. Li-Ion Battery Materials: Present and Future. Mater. Today 2015, 18 (5), 252–264. 9. Etacheri, V.; Marom, R.; Elazari, E.; Salitra, G.; Aurbach, D. Challenges in the Development of Advanced Li-Ion Batteries: A Review. Energ. Environ. Sci. 2011, 4, 3243–3262. 10. Manthiram, A. A Reflection on Lithium-ion Battery Cathode Chemistry. Nat. Commun. 2020, 11, 1550. 11. Abakumov, A.; Fedotov, S. S.; Antipov, E.; Tarascon, J. M. Solid State Chemistry for Developing Better Metal-Ion Batteries. Nat. Commun. 2020, 11, 4976. 12. Choi, J. W.; Aurbach, D. Promise and Reality of Post-Lithium-Ion Batteries with High Energy Densities. Nat. Rev. Mater. 2016, 1, 16013. 13. Usiskin, R.; Lu, Y.; Popovic, J.; Law, M.; Balaya, P.; et al. Fundamentals, Status and Promise of Sodium-Based Batteries. Nat. Rev. Mater. 2021, 6, 1020–1035. 14. Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research Development on Sodium-Ion Batteries. Chem. Rev. 2014, 114 (23), 11636–11682.

268

Development of polyanionic sodium-ion battery insertion materials

15. Kundu, D.; Talaie, E.; Duffont, V.; Nazar, L. F. The Emerging Chemistry of Sodium Ion Batteries for Electrochemical Energy Storage. Angew. Chem. Int. Ed. 2015, 54 (11), 3431–3448. 16. Kubota, K.; Komaba, S. Review- Practical Issues and Future Perspectives for Na-Ion Batteries. J. Electrochem. Soc. 2015, 162 (14), A2538–A2550. 17. Hwang, J. Y.; Myung, S. T.; Sun, Y. K. Sodium-Ion Batteries: Present and Future. Chem. Soc. Rev. 2017, 46 (12), 3529–3614. 18. Fouassier, C.; Matejka, G.; Reau, J. M.; Hagenmuller, P. Sur de Nouveaux Bronzes Oxygenes de Formule NaxCoO2 (x 1). Le Systeme Cobalt-Oxygene-Sodium. J. Solid State Chem. 1973, 6 (4), 532–537. 19. Mizushima, K.; Jones, P. C.; Wiseman, P. J.; Goodenough, J. B. LixCoO2 (0 < x < 1): A New Cathode Material for Batteries of High Energy Density. Mater. Res. Bull. 1980, 15 (6), 783–789. 20. Wang, P. F.; You, Y.; Yin, Y. X.; Guo, Y. G. Layered Oxide Cathodes for Sodium-Ion Batteries: Phase Transition, Air Stability, and Performance. Adv. Energy Mater. 2018, 8 (8), 1701912. 21. Manthiram, A.; Goodenough, J. B. Lithium Insertion into Fe2(SO4)3 Frameworks. J. Power Sources 1989, 26 (3–4), 403–408. 22. Padhi, A. K.; Nanjundaswamy, K. S.; Goodenough, J. B. Phospho-Olivines as Positive-Electrode Materials for Rechargeable Lithium Batteries. J. Electrochem. Soc. 1997, 144 (4), 1188–1194. 23. Yamada, A.; Chung, S. C.; Hinokuma, K. Optimized LiFePO4 for Lithium Battery Cathodes. J. Electrochem. Soc. 2001, 148 (3), A224–A229. 24. Masquelier, C.; Croguennec, L. Polyanionic (Phosphates, Silicates, Sulfates) Frameworks as Electrode Materials for Rechargeable Li (or Na) Batteries. Chem. Rev. 2013, 113 (8), 6552–6591. 25. Manthiram, A.; Goodenough, J. B. Lithium Insertion into Fe2(MoO4)3 Frameworks: Comparison of M ¼ W with M ¼ Mo. J. Solid State Chem. 1987, 71 (2), 349–360. 26. Padhi, A. K.; Nanjundaswamy, K. S.; Masquelier, C.; Okada, S.; Goodenough, J. B. Effect of Structure on the Fe3þ/Fe2þ Redox Couple in Iron Phosphates. J. Electrochem. Soc. 1997, 144 (5), 1609–1613. 27. Padhi, A. K.; Manivannan, V.; Goodenough, J. B. Tuning the Position of the Redox Couples in Materials with NASICON Structure by Anionic Substitution. J. Electrochem. Soc. 1998, 145 (5), 1518–1520. 28. Goodenough, J. B.; Kim, Y. Challenges for Rechargeable Li Batteries. Chem. Rev. 2010, 22 (3), 587–603. 29. Barpanda, P.; Lander, L.; Nishimura, S.; Yamada, A. Polyanionic Insertion Materials for Sodium-Ion Batteries. Adv. Energy Mater. 2018, 8 (17), 1703055. 30. Senthilkumar, B.; Murugesan, C.; Sharma, L.; Lochab, S.; Barpanda, P. An Overview of Mixed Polyanionic Cathode Materials for Sodium-Ion Batteries. Small Methods. 2019, 3 (4), 1800253. 31. Avdeev, M. Crystal Chemistry of NaSICONs: Ideal Framework, Distortion, and Connection to Properties. Chem. Mater. 2021, 33 (19), 7620–7632. 32. Goodenough, J. B.; Hong, H. Y. P.; Kafalas, J. A. Fast Naþ-Ion Transport in Skeleton Stuctures. Mater. Res. Bull. 1976, 11 (2), 203–220. 33. Gopalakrishnan, J.; Rangan, K. K. Vanadium Phosphate (V2(PO4)3): A Novel NASICON-Type Vanadium Phosphate Synthesized by Oxidative Deintercalation of Sodium from Sodium Vanadium Phosphate (Na3V2(PO4)3). Chem. Mater. 1992, 4 (4), 745–747. 34. Delmas, C.; Olazcuaga, R.; Cherkaoui, F.; Brochu, R.; Leflem, G. A New Family of Phosphates with the Formula Na3M2(PO4)3 (M ¼ Ti, V, Cr, Fe). C R Seances Acad Sci Ser C 1978, 287 (5), 169–171. 35. Chotard, J.-N.; Rousse, G.; David, R.; Mentré, O.; Courty, M.; et al. Discovery of a Sodium-Ordered Form of Na3V2(PO4)3 Below Ambient Temperature. Chem. Mater. 2015, 27 (17), 5982–5987. 36. Plashnitsa, L. S.; Kobayashi, E.; Noguchi, Y.; Okada, S.; Yamaki, J. Performance of NASICON Symmetric Cell With Ionic Liquid Electrolyte. J. Electrochem. Soc. 2010, 157 (4), A536–A543. 37. Jian, Z.; Zhao, L.; Pan, H.; Hu, Y.-S.; Li, H.; et al. Carbon Coated Na3V2(PO4)3 as Novel Electrode Material for Sodium Ion Batteries. Electrochem. Commun. 2012, 14 (1), 86–89. 38. Lim, S. Y.; Kim, H.; Shakoor, R. A.; Jung, Y.; Choi, J. W. Electrochemical and Thermal Properties of NASICON Structured Na3V2(PO4)3 as a Sodium Rechargeable Battery Cathode: A Combined Experimental and Theoretical Study. J. Electrochem. Soc. 2012, 159 (9), A1393–A1397. 39. Jian, Z.; Han, W.; Lu, X.; Yang, H.; Hu, Y.-S.; et al. Superior Electrochemical Performance and Storage Mechanism of Na3V2(PO4)3 Cathode for Room-Temperature Sodium-Ion Batteries. Adv. Energy Mater. 2013, 3 (2), 156–160. 40. Saravanan, K.; Mason, C. W.; Rudola, A.; Wong, K. H.; Balaya, P. The First Report on Excellent Cycling Stability and Superior Rate Capability of Na3V2(PO4)3 for Sodium-Ion Batteries. Adv. Energy Mater. 2013, 3 (4), 444–450. 41. Jian, Z.; Yuan, C.; Han, W.; Lu, X.; Gu, L.; et al. Atomic Structure and Kinetics of NASICON NaxV2(PO4)3 Cathode for Sodium-Ion Batteries. Adv. Funct. Mater. 2014, 24 (27), 4265–4272. 42. Ren, W.; Yao, X.; Niu, C.; Zheng, Z.; Zhao, K.; et al. Cathodic Polarization Suppressed Sodium-Ion Full Cell with a 3.3 V High-Voltage. Nano Energy 2016, 28, 216–223. 43. Ren, W.; Zheng, Z.; Xu, C.; Niu, C.; Wei, Q.; et al. Self-Sacrificed Synthesis of Three-Dimensional Na3V2(PO4)3 Nanofiber Network for High-Rate Sodium-Ion Full Batteries. Nano Energy 2016, 25, 145–153. 44. Zhou, W.; Xue, L.; Lu, X.; Gao, H.; Li, Y.; et al. NaxMV(PO4)3 (M ¼ Mn, Fe, Ni) Structure and Properties of Sodium Extraction. Nano Lett. 2016, 16 (12), 7836–7841. 45. Gao, H.; Seymour, I. D.; Xin, S.; Xue, L.; Henkelman, G.; et al. Na3MnZr(PO4)3: A High-Voltage Cathode for Sodium Batteries. J. Am. Chem. Soc. 2018, 140 (51), 18192– 18199. 46. Avdeev, M.; Mohamed, Z.; Ling, C. D.; Lu, J.; Tamaru, M.; et al. Magnetic Structures of NaFePO4 Maricite and Triphylite Polymorphs for Sodium Ion Batteries. Inorg. Chem. 2013, 52 (15), 8685–8693. 47. Oh, S. M.; Myung, S. T.; Hassoun, J.; Scrosati, B.; Sun, Y. K. Reversible NaFePO4 Electrode for Sodium Secondary Batteries. Electrochem. Commun. 2012, 22, 149–152. 48. Moreau, P.; Guyomard, D.; Gaubicher, J.; Boucher, F. Structure and Stability of Sodium Intercalated Phases in Olivine FePO4. Chem. Mater. 2010, 22 (14), 4126–4128. 49. Lu, J.; Chung, S. C.; Nishimura, S.; Yamada, A. Phase Diagram of Olivine NaxFePO4 (0 < x < 1). Chem. Mater. 2013, 25 (22), 4557–4565. 50. Kim, J.; Seo, D.-H.; Kim, H.; Park, I.; Yoo, J.-K.; et al. Unexpected Discovery of Low-Cost Maricite NaFePO4 as a High-Performance Electrode for Na-Ion Batteries. Energ. Environ. Sci. 2015, 8 (2), 540–545. 51. Gutierrez, A.; Kim, S.; Fister, T. T.; Johnson, C. S. Microwave-Assisted Synthesis of NaCoPO4 Red-Phase and Initial Characterizations as High Voltage Cathode for Sodium-Ion Batteries. ACS Appl. Mater. Interfaces 2017, 9 (5), 4391–4396. 52. Barpanda, P.; Nishimura, S.; Yamada, A. High-Voltage Pyrophosphate Cathodes. Adv. Energy Mater. 2012, 2 (7), 841–859. 53. Honma, T.; Togashi, T.; Ito, N.; Komatsu, T. Fabrication of Na2FeP2O7 Glass-Ceramics for Sodium-Ion Battery. J. Cerma. Soc. Jpn. 2012, 120 (1404), 344–346. 54. Kim, H.; Shakoor, R. A.; Park, C.; Lim, S. Y.; Kim, J.-S.; et al. Na2FeP2O7 as a Promising Iron-Based Pyrophosphate Cathode for Sodium Rechargeable Batteries: A Combined Experimental and Theoretical Study. Adv. Funct. Mater. 2013, 23 (9), 1147–1155. 55. Barpanda, P.; Ye, T.; Nishimura, S.; Chung, S. C.; Yamada, Y.; et al. Sodium Iron Pyrophosphate: A Novel 3.0 V Iron-Based Cathode for Sodium-Ion Batteries. Electrochem. Commun. 2012, 24, 116–119. 56. Barpanda, P.; Liu, G.; Ling, C. D.; Tamaru, M.; Avdeev, M.; et al. Na2FeP2O7: A Safe Cathode for Rechargeable Sodium-Ion Batteries. Chem. Mater. 2013, 25 (17), 3480–3487. 57. Ha, K. H.; Woo, S. H.; Mok, D.; Choi, N.-S.; Park, Y.; et al. Na4-aM2 þ a/2(P2O7)2 (2/3  a  7/8, M ¼ Fe, Fe0.5Mn0.5, Mn): A Promising Sodium Ion Cathode for Na-Ion Batteries. Adv. Energy Mater. 2013, 3 (6), 770–776. 58. Barpanda, P.; Ye, T.; Avdeev, M.; Chung, S. C.; Yamada, A. A New Polymorph of Na2MnP2O7 as a 3.6 V Cathode Material for Sodium-Ion Batteries. J. Mater. Chem. A 2013, 1 (13), 4194–4197.

Development of polyanionic sodium-ion battery insertion materials

269

59. Kim, H.; Park, C. S.; Choi, J. W.; Jung, Y. Defect-Controlled Formation of Triclinic Na2CoP2O7 for 4 V Sodium-Ion Batteries. Angew. Chem. Int. Ed. 2016, 55 (23), 6662–6666. 60. Barpanda, P.; Lu, J.; Ye, T.; Kajiyama, M.; Chung, S. C.; et al. A Layer-Structured Na2CoP2O7 Pyrophosphate Cathode for Sodium-Ion Batteries. RSC Adv. 2013, 3 (12), 3857–3860. 61. Gond, R.; Meena, S. S.; Yusuf, S. M.; Shukla, V.; Jena, N. K.; et al. Enabling the Electrochemical Activity in Sodium iron Metaphosphate [NaFe(PO3)3] Sodium Battery Insertion Materials: Structural and Electrochemical Insights. Inorg. Chem. 2017, 56 (10), 5918–5929. 62. Gond, R.; Rao, R. P.; Praong, V.; Levedev, O. I.; Adams, S.; et al. Cubic Sodium Cobalt Metaphosphate [NaCo(PO3)3] as a Novel Cathode Material for Sodium-Ion Batteries. Inorg. Chem. 2018, 57 (11), 6324–6332. 63. Dwibedi, D.; Barpanda, P.; Yamada, A. Alluaudite Battery Cathodes. Small Methods 2020, 4 (7), 2000051. 64. Fisher, D. J. Alluaudite. Am. Mineral. 1955, 40 (11 12), 1100–1109. 65. Moore, P. B. Crystal Chemistry of the Alluaudite Structure Type: Contribution to the Paragenesis of Pegmatite Phosphate Giant Crystals. Am. Mineral. 1971, 56 (11–12), 1955–1975. 66. Trad, K.; Carlier, D.; Croguennec, L.; Wattiaux, A.; Amara, M. B.; et al. NaMnFe2(PO4)3 Alluaudite Phase: Synthesis, Structure, and Electrochemical Properties as Positive Electrode in Lithium and Sodium Batteries. Chem. Mater. 2010, 22 (19), 5554–5562. 67. Dwibedi, D.; Jaschin, P. W.; Gond, R.; Barpanda, P. Revisiting the Alluaudite NaMnFe2(PO4)3 Sodium Insertion Material: Structural, Diffusional and Electrochemical Insights. Electrochim. Acta 2018, 283, 850–857. 68. Huang, W.; Li, B.; Saleem, M. F.; Wu, X.; Li, J.; et al. Self-Assembled Alluaudite Na2Fe3-xMnx(PO4)3 Micro/Nanocompounds for Sodium-Ion Battery Electrodes: A New Insight Into Their Electronic and Geometric Structure. Chem. A Eur. J. 2015, 21 (2), 851–860. 69. Essehli, R.; Belharouak, I.; Ben Yahia, H.; Maher, K.; Abouimrane, A.; et al. Alluaudite Na2Co2Fe(PO4)3 as an Electroactive Material for Sodium Ion Batteries. Dalton Trans. 2015, 44 (17), 7881–7886. 70. Huang, W.; Zhou, J.; Li, B.; An, L.; Cui, P.; et al. A New Route Toward Improved Sodium Ion Batteries: A Multifunctional Fluffy Na0.67FePO4/CNT Nanocactus. Small 2015, 11 (18), 2170–2176. 71. Liu, D.; Palmore, G. T. R. Synthesis, Crystal Structure, and Electrochemical Properties of Alluaudite Na1.702Fe3(PO4)3 as a Sodium-Ion Battery Cathode. ACS Sustain. Chem. Eng. 2017, 5 (7), 5766–5771. 72. Kim, J.; Kim, H.; Lee, S.; Myung, S. T. Development of a New Alluaudite-Based Cathode Material with High Power and Long Cyclability for Application in Na Ion Batteries in Real-Life. J. Mater. Chem. A 2017, 5 (42), 22334–22340. 73. Dwibedi, D.; Gond, R.; Barpanda, P. Alluaudite NaCoFe2(PO4)3 as a 2.9 V Cathode for Sodium-Ion Batteries Exhibiting Bifunctional Electrocatalytic Activity. Chem. Mater. 2019, 31 (18), 7501–7509. 74. Barpanda, P. Pursuit of Sustainable Iron-Based Positive Insertion Materials for Sodium-Ion Batteries: Two Case Studies. Chem. Mater. 2016, 28 (4), 1006–1011. 75. Barpanda, P.; Ati, M.; Melot, B. C.; Rousse, G.; Chotard, J. N.; et al. A 3.90 V Iron-Based Fluorosulphate Material for Lithium-Ion Batteries Crystallizing in the Triplite Structure. Nat. Mater. 2011, 10 (10), 772–779. 76. Reynaud, M.; Ati, M.; Melot, B. C.; Sougrati, M. T.; Rousse, G.; et al. Li2Fe(SO4)2 as a 3.83 V Positive Electrode Material. Electrochem. Commun. 2012, 21, 77–80. 77. Barpanda, P. Sulphate Chemistry for High-Voltage Insertion Materials: Synthetic, Structural and Electrochemical Insights. Isr. J. Chem. 2015, 55 (5), 537–557. 78. Reynaud, M.; Rousse, G.; Abakumov, A. M.; Sougrati, M. T.; Van Tendeloo, G.; et al. Design of New Electrode Materials for Li-Ion and Na-Ion Batteries From the Bloedite Mineral Na2Mg(SO4)2.4H2O. J. Mater. Chem. A 2014, 2 (8), 2671–2680. 79. Barpanda, P.; Oyama, G.; Ling, C. D.; Yamada, A. Krohnkite-Type Na2Fe(SO4)2.2H2O as a Novel 3.25 V Insertion Compound for Na-Ion Batteries. Chem. Mater. 2014, 26 (3), 1297–1299. 80. Singh, P.; Shiva, K.; Celio, H.; Goodenough, J. B. Eldfellite, NaFe(SO4)2: An Intercalation Cathode Host for Low-Cost Na-Ion Batteries. Energ. Environ. Sci. 2015, 8 (10), 3000–3005. 81. Singh, S.; Singh, D.; Ahuja, R.; Fichtner, M.; Barpanda, P. Eldfellite NaV(SO4)2 as a Versatile Cathode Insertion Host for Li-Ion and Na-Ion Batteries. J. Mater. Chem. A 2022. https://doi.org/10.1039/D2TA03673H. In Press. 82. Barpanda, P.; Oyama, G.; Nishimura, S.; Chung, S. C.; Yamada, A. A 3.8 V Earth-Abundant Sodium Battery Electrode. Nat. Commun. 2014, 5, 4358. 83. Oyama, G.; Nishimura, S.; Suzuki, Y.; Okubo, M.; Yamada, A. Off-Stoichiometry in Alluaudite-Type Sodium Iron Sulfate Na2 þ 2xFe2-x(SO4)3 as an Advanced Sodium Battery Cathode Material. ChemElectroChem 2015, 2 (7), 1019–1023. 84. Araujo, R. B.; Chakraborty, S.; Barpanda, P.; Ahuja, R. Na2M2(SO4)3 (M ¼ Fe, Mn, co and Ni): Towards High Voltage Sodium Battery Applications. Phys. Chem. Chem. Phys. 2016, 18 (14), 9658–9665. 85. Oyama, G.; Pecher, O.; Griffith, K. J.; Nishimura, S.; Pigliapochi, R.; et al. Sodium Intercalation Mechanism of 3.8 V Class Alluaudite Sodium Iron Sulfate. Chem. Mater. 2016, 28 (15), 5321–5328. 86. Nishimura, S.; Suzuki, Y.; Lu, J.; Torri, S.; Kamiyama, T.; et al. High-Temperature Neutron and X-Ray Diffraction Study of Fast Sodium Transport in Alluaudite-Type Sodium Iron Sulfate. Chem. Mater. 2016, 28 (7), 2393–2399. 87. Dwibedi, D.; Araujo, R. B.; Chakraborty, S.; Shanbogh, P.; Sundaram, N.; et al. Na2.44Mn1.79(SO4)3: A New Member of Alluaudite Family of Insertion Compound for Sodium Ion Batteries. J. Mater. Chem. A 2015, 3 (36), 18564–18571. 88. Dwibedi, D.; Gond, R.; Dayamani, A.; Araujo, R. B.; Chakraborty, S.; et al. Na2.32Co1.84(SO4)3 as a New Member of Alluaudite Family of High-Voltage Sodium Battery Cathode. Dalton Trans. 2017, 46 (1), 55–63. 89. Lu, J.; Nishimura, S.; Yamada, A. Polyanionic Solid-Solution Cathodes for Rechargeable Batteries. Chem. Mater. 2017, 29 (8), 3597–3602. 90. Barpanda, P.; Dwibedi, D.; Ghosh, S.; Kee, Y.; Okada, S. Lithium Metal Borate (LiMBO3) Family of Insertion Materials for Li-Ion Batteries: A Sneak Peak. Ionics 2015, 21 (7), 1801–1812. 91. Strauss, F.; Rousse, G.; Sougrati, M. T.; Corte, D. A. D.; Courty, M.; et al. Synthesis, Structure, and Electrochemical Properties of Na3MB5O10 (M ¼ Fe, Co) Containing M2 þ in Tetrahedral Coordination. Inorg. Chem. 2016, 55 (24), 12775–12782. 92. Law, M.; Ramar, V.; Balaya, P. Na2MnSiO4 as an Attractive High Capacity Cathode Material for Sodium-Ion Battery. J. Power Sources 2017, 359, 277–284. 93. Kee, Y.; Dimov, N.; Staykov, A.; Okada, S. Investigation of Metastable Na2FeSiO4 as a Cathode Material for Na-Ion Secondary Battery. Mater. Chem. Phys. 2016, 171, 45–49. 94. Li, S.; Guo, J.; Ye, Z.; Zhao, X.; Wu, S.; et al. Zero-Strain Na2FeSiO4 as Novel Cathode Material for Sodium-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8 (27), 17233– 17238. 95. Panigrahi, A.; Nishimura, S.; Shimada, T.; Watanabe, E.; Zhao, W.; et al. Sodium Iron(II) Pyrosilicate Na2Fe2Si2O7: A Potential Cathode Material in the Na2O-FeO-SiO2 System. Chem. Mater. 2017, 29 (10), 4361–4366. 96. Benhsina, E.; Assani, A.; Saadi, M.; El Ammari, L. Crystal structure of (Na0.70)(Na0.70,Mn0.30)(Fe3þ,Fe2þ)2Fe2þ(VO4)3, A Sodium-, Iron- and Manganese-based Vanadate with the Alluaudite-type Structure. Acta Crystallogr. 2016, E72, 220–222. 97. Hadouchi, M.; Assani, A.; Saadi, M.; El Ammari, L. The Alluaudite-Type Crystal Structures of Na2(Fe/Co)2Co(VO4)3 and Ag2(Fe/Co)2Co(VO4)3. Acta Crystallogr. 2016, E72, 1017–1020. 98. Ben Yahia, H.; Shikano, M.; Tabuchi, M.; Belharouak, I. Synthesis, Crystal Structure, and Properties of the Alluaudite-Type Vanadates Ag2–xNaxMn2Fe(VO4)3. Inorg. Chem. 2016, 55 (9), 4643–4649. 99. Gao, J.; Zhao, P.; Feng, K. Na2.67Mn1.67(MoO4)3: A 3.45 V Alluaudite-Type Cathode Candidate for Sodium-Ion Batteries. Chem. Mater. 2017, 29 (3), 940–944.

270

Development of polyanionic sodium-ion battery insertion materials

100. Serdtsev, A.; Medvedeva, N. I. Ab Initio Insights into Na-Ion Diffusion and Intercalation Mechanism in Alluaudite NaxMn2(MoO4)3 as Cathode Material for Sodium-Ion Batteries. J. Alloys Compd. 2019, 808, 151667. 101. Gover, R. K. B.; Withers, N. D.; Allen, S.; Withers, R. L.; Evans, J. S. O. Structure and Phase Transitions of SnP2O7. J. Solid State Chem. 2002, 166 (1), 42–48. 102. Zheng, J. C.; Han, Y.; Zhang, B.; Shen, C.; Ming, L.; et al. Electrochemical Properties of VPO4/C Nanosheets and Microspheres as Anode Materials for Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2014, 6 (9), 6223–6226. 103. Le Meins, J. M.; Crosnier-Lopez, M. P.; Hemon-Ribaud, A.; Courbion, G. Phase Transitions in the Na3M2(PO4)2F3 Family (M ¼ Al3þ, V3þ, Cr3þ, Fe3þ, Ga3þ): Synthesis, Thermal, Structural, and Magnetic Studies. J. Solid State Chem. 1999, 148 (2), 260–277. 104. Barker, J.; Saidi, M. Y.; Swoyer, J. L. A Sodium-Ion Cell Based on the Fluorophosphate Compound NaVPO4F. Electrochem. Solid St. 2002, 6 (1), A1–A4. 105. Zhuo, H.; Wang, X.; Tang, A.; Liu, Z.; Gamboa, S.; et al. The Preparation of NaV1-xCrxPO4F Cathode Materials for Sodium-Ion Battery. J. Power Sources 2006, 160 (1), 698–703. 106. Liu, Z. M.; Wang, X. Y.; Wang, Y.; Tang, A.; Yang, S. Y.; et al. Preparation of NaV1-xAlxPO4F Cathode Materials for Application of Sodium-Ion Battery. Trans. Nonferrous Met. Soc. Chin. 2008, 18 (2), 346–350. 107. Shakoor, R. A.; Seo, D. H.; Kim, H.; Park, Y. U.; Kim, J.; et al. A Combined First Principles and Experimental Study on Na3V2(PO4)2F3 for Rechargeable Na Batteries. J. Mater. Chem. 2012, 22 (28), 20535–20541. 108. Jiang, T.; Wei, Y. J.; Pan, W. C.; Li, Z.; Ming, X.; et al. Preparation and Electrochemical Studies of Li3V2(PO4)3/Cu Composite Cathode Material for Lithium Ion Batteries. J. Alloys Compd. 2009, 488 (1), L26–L29. 109. Yan, G.; Mariyappan, S.; Rousse, G.; Jacquet, Q.; Deschamps, M.; et al. Higher Energy and Safer Sodium Ion Batteries Via an Electrochemically Made Disordered Na3V2(PO4)2F3 Material. Nat. Commun. 2019, 10, 585. 110. Matts, I. L.; Dacek, S.; Pietrzak, T. K.; Malik, R.; Ceder, G. Explaining Performance-Limiting Mechanisms in Fluorophosphate Na-Ion Battery Cathodes through Inactive Transition-Metal Mixing and First-Principles Mobility Calculations. Chem. Mater. 2015, 27 (17), 6008–6015. 111. Broux, T.; Fauth, F.; Hall, N.; Chatillon, Y.; Bianchini, M.; et al. High Rate Performance for Carbon-Coated Na3V2(PO4)2F3 in Na-Ion Batteries. Small Methods. 2019, 3 (4), 1800215. 112. Sauvage, F.; Quarez, E.; Tarascon, J. M.; Baudrin, E. Crystal Structure and Electrochemical Properties Vs. Naþ of the Sodium Fluorophosphate Na1.5VOPO4F0.5. Solid State Sci. 2006, 8 (10), 1215–1221. 113. Massa, W.; Yakubovich, O. V.; Dimitrova, O. V. Crystal Structure of a New Sodium Vanadyl(IV) Fluoride Phosphate Na3{V2O2F[PO4]2}. Solid State Sci. 2002, 4 (4), 495–501. 114. Tsirlin, A.; Nath, R.; Abakumov, A.; Fukurawa, Y.; Johnston, D.; et al. Phase Separation and Frustrated Square Lattice Magnetism of Na1.5VOPO4F0.5. Phys. Rev. B. 2011, 84, 014429. 115. Sharma, N.; Serras, P.; Palomares, V.; Brand, H. E. A.; Alonso, J.; et al. Sodium Distribution and Reaction Mechanisms of a Na3V2O2(PO4)2F Electrode during Use in a SodiumIon Battery. Chem. Mater. 2014, 26 (11), 3391–3402. 116. Park, Y. U.; Seo, D. H.; Kim, H.; Kim, J.; Lee, S.; et al. A Family of High-Performance Cathode Materials for Na-Ion Batteries, Na3(VO1-xPO4)2 F1 þ 2x (0  x  1): Combined First-Principles and Experimental Study. Adv. Funct. Mater. 2014, 24 (29), 4603–4614. 117. Xu, M.; Wang, L.; Zhao, X.; Song, J.; Xie, H.; et al. Na3V2O2(PO4)2F/Graphene Sandwich Structure for High-Performance Cathode of a Sodium-Ion Battery. Phys. Chem. Chem. Phys. 2013, 15 (31), 13032–13037. 118. Ellis, B. E.; Makahnouk, W. R. M.; Makimura, Y.; Toghill, K.; Nazar, L. F. A Multifunctional 3.5 V Iron-Based Phosphate Cathode for Rechargeable Batteries. Nat. Mater. 2007, 6, 749–753. 119. Ellis, B. E.; Makahnouk, W. R. M.; Rowan-Weetaluktuk, W. N.; Ryan, D. H.; Nazar, L. F. Crystal Structure and Electrochemical Properties of A2MPO4F Fluorophosphates (A ¼ Na, Li; M ¼ Fe, Mn, Co, Ni). Chem. Mater. 2010, 22 (3), 1059–1070. 120. Avdeev, M.; Ling, C. D.; Tan, T. T.; Li, S.; Oyama, G.; et al. Magnetic Structure and Properties of the Rechargeable Battery Insertion Compound Na2FePO4F. Inorg. Chem. 2014, 53 (2), 682–684. 121. Recham, N.; Chotard, J. N.; Dupont, L.; Djellab, K.; Armand, M.; et al. Ionothermal Synthesis of Sodium-Based Fluorophosphate Cathode Materials. J. Electrochem. Soc. 2009, 156 (12), A993–A999. 122. Sharma, L.; Nayak, P.; de la Llave, E.; Chen, H.; Adams, S.; et al. Electrochemical and Diffusional Investigation of Na2Fe(II)PO4F Fluorophosphates Sodium Insertion Material Obtained from Fe(III) Precursor. ACS Appl. Mater. Interfaces 2017, 9 (4), 34961–34969. 123. Smiley, D. L.; Goward, G. R. Ex Situ 23Na solid-State NMR Reveals the Local Na-Ion Distribution in Carbon-Coated Na2FePO4F during Electrochemical Cycling. Chem. Mater. 2016, 28 (21), 7645–7656. 124. Kubota, K.; Yokoh, K.; Yabuuchi, N.; Komaba, S. Na2CoPO4F as a High-Voltage Electrode Material for Na-Ion Batteries. Electrochemistry 2014, 82 (10), 909–911. 125. Zou, H.; Li, S.; Wu, X.; McDonald, M. J.; Yang, Y. Spray-Drying Synthesis of Pure Na2CoPO4F as Cathode Material for Sodium Ion Batteries. ECS Electrochem. Lett. 2015, 4 (6), A53–A55. 126. Wu, X.; Zheng, J.; Gong, Z.; Yang, Y. Sol-Gel Synthesis and Electrochemical Properties of Fluorophosphates Na2Fe1-xMnxPO4F/C (x ¼ 0, 0.1, 0.3, 0.7, 1) Composite as Cathode Materials for Lithium Ion Battery. J. Mater. Chem. 2011, 21 (46), 18630–18637. 127. Sanz, F.; Parada, C.; Rojo, J. M.; Ruiz-Valero, C. Synthesis, Structural Characterization, Magnetic Properties, and Ionic Conductivity of Na4MII3(PO4)2(P2O7) (MII ¼ Mn, Co, Ni). Chem. Mater. 2001, 13 (4), 1334–1340. 128. Wood, S. M.; Eames, C.; Kendrick, E.; Islam, M. S. Sodium Ion Diffusion and Voltage Trends in Phosphates Na4M3(PO4)2P2O7 (M ¼ Fe, Mn, Co, Ni) for Possible High-Rate Cathodes. J. Phys. Chem. C 2015, 119 (28), 15935–15941. 129. Kim, H.; Park, I.; Seo, D. H.; Lee, S.; Kim, S. W.; et al. New Iron-Based Mixed-Polyanion Cathodes for Lithium and Sodium Rechargeable Batteries: Combined First Principles Calculations and Experimental Study. J. Am. Chem. Soc. 2012, 134 (25), 10369–10372. 130. Kim, H.; Park, I.; Lee, S.; Kim, H.; Park, K. Y.; et al. Understanding the Electrochemical Mechanism of the New Iron-Based Mixed-Phosphate Na4Fe3(PO4)2(P2O7) in a Na Rechargeable Battery. Chem. Mater. 2013, 25 (18), 3614–3622. 131. Nose, M.; Nakayama, H.; Nobuhara, K.; Yamaguchi, H.; Nakanishi, S.; et al. Na4Co3(PO4)2P2O7: A Novel Storage Material for Sodium-Ion Batteries. J. Power Sources 2013, 234, 175–179. 132. Kim, H.; Yoon, G.; Park, I.; Park, K. Y.; Lee, B.; et al. Anomalous Jahn-Teller Behavior in a Manganese-Based Mixed-Phosphate Cathode for Sodium Ion Batteries. Energ. Environ. Sci. 2015, 8 (11), 3325–3335. 133. Zhang, H.; Hasa, I.; Buchholz, D.; Qin, B.; Geiger, D.; et al. Exploring the Ni Redox Activity in Polyanionic Compounds as Conceivable High Potential Cathodes for Na Rechargeable Batteries. NPG Asia Mater. 2017, 9, e370. 134. Nose, M.; Shiotani, S.; Nakayama, H.; Nobuhara, K.; Nakanishi, S.; et al. Na4Co2.4Mn0.3Ni0.3(PO4)2P2O7: High Potential and High Capacity Electrode Material for Sodium-Ion Batteries. Electrochem. Commun. 2013, 34, 266–269. 135. de la Rochere, M.; Kahn, A.; d’Yvoire, F.; Bretey, E. Crystal Structure and Cation Transport Properties of the OrthodDiphosphates Na7(MP2O7)4PO4 (M ¼ Al, Cr, Fe). Mater. Res. Bull. 1985, 20 (1), 27–34. 136. Lim, S. Y.; Kim, H.; Chung, J.; Lee, J. H.; Kim, B. G.; et al. Role of Intermediate Phase for Stable Cycling of Na7V4(P2O7)4PO4 in Sodium Ion Battery. Proc. Natl. Acad. Sci. U. S. A. 2013, 111 (2), 599–604. 137. Recham, N.; Chotard, J. N.; Dupont, L.; Delacourt, C.; Walker, W.; et al. A 3.6 V Lithium-Based Fluorosulfate Insertion Positive Electrode for Lithium-Ion Batteries. Nat. Mater. 2010, 9, 68–74.

Development of polyanionic sodium-ion battery insertion materials

271

138. Tarascon, J. M.; Recham, N.; Armand, M.; Chotard, J. N.; Barpanda, P.; et al. Hunting for Better Li-Based Electrode Materials Via Low Temperature Inorganic Synthesis. Chem. Mater. 2010, 22 (3), 724–739. 139. Barpanda, P.; Chotard, J. N.; Recham, N.; Delacourt, C.; Ati, M.; et al. Structural, Transport, and Electrochemical Investigation of Novel AMSO4F (A ¼ Na, Li; M ¼ Fe, Co, Ni, Mn)Metal Fluorosulphates Prepared Using Low Temperature Synthesis Routes. Inorg. Chem. 2010, 49 (16), 7401–7413. 140. Reynaud, M.; Barpanda, P.; Rousse, G.; Chotard, J. N.; Melot, B. C.; et al. Synthesis and Crystal Chemistry of the NaMSO4F Family (M ¼ Mg, Fe, Co, Cu, Zn). Solid State Sci. 2011, 14 (1), 15–20. 141. Tripathi, R.; Gardiner, G. R.; Islam, M. S.; Nazar, L. F. Alkali-Ion Conduction Paths in LiFeSO4F and NaFeSO4F Tavorite-Type Cathode Materials. Chem. Mater. 2011, 23 (8), 2278–2284. 142. Recham, N.; Rousse, G.; Sougrati, M. T.; Chotard, J. N.; Frayret, C.; et al. Preparation and Characterization of a Stable FeSO4F-Based Framework for Alkali Ion Insertion Electrodes. Chem. Mater. 2012, 24 (22), 4363–4370. 143. Hosaka, T.; Kubota, K.; Hameed, A. S.; Komaba, S. Research Development on K-Ion Batteries. Chem. Rev. 2020, 120 (14), 6358–6466. 144. Ati, M.; Dupont, L.; Recham, N.; Chotard, J. N.; Walker, W.; et al. Synthesis, Structural, and Transport Properties of Novel Bihydrated Fluorosulfates NaMSO4F.2H2O (M ¼ Fe, Co and Ni). Chem. Mater. 2010, 22 (13), 4062–4068. 145. Gnanavel, M.; Pralong, V.; Lebedev, O. I.; Caignaert, V.; Bazin, P.; et al. Lithium Intercalation into the Jarosite-Type Hydroxysulfate: A Topotactic Reversible Reaction from a Crystalline Phase to an Inorganic Polymer-Like Structure. Chem. Mater. 2014, 26 (15), 4521–4527. 146. Gnanavel, M.; Lebedev, O. I.; Bazin, P.; Raveau, B.; Pralong, V. Reversible Transformation from Amorphous Na3Fe3(SO4)2(OH)6 to Crystallized NaFe3(SO4)2(OH)6 Jarosite-Type Hydroxysulfate. Solid State Ion. 2015, 278, 38–42. 147. Chen, H.; Hao, Q.; Zivkovic, O.; Hautier, G.; Du, L. S.; et al. Sidorenkite (Na3MnPO4CO3): A New Intercalation Cathode Material for Na-Ion Batteries. Chem. Mater. 2013, 25 (14), 2777–2786. 148. Huang, W.; Zhou, J.; Li, B.; Ma, J.; Tao, S.; et al. Detailed Investigation of Na2.24FePO4CO3 as a Cathode Material for Na-Ion Batteries. Sci. Rep. 2014, 4, 4188. 149. Liu, J.; Chang, D. H.; Whitfield, P.; Janssen, Y.; Yu, X. Q.; et al. Ionic Conduction in Cubic Na3TiP3O9N, a Secondary Na-Ion Battery Cathode with Extremely Low Volume Change. Chem. Mater. 2014, 26 (10), 3295–3305. 150. Kim, J.; Yoon, G.; Lee, M. H.; Kim, H.; Lee, S.; et al. New 4V-Class and Zero-Strain Cathode Material for Na-Ion Batteries. Chem. Mater. 2017, 29 (18), 7826–7832. 151. Reynaud, M.; Wizner, A.; Katcho, N. A.; Loaiza, L. C.; Galceran, M.; et al. Sodium Vanadium Nitridophosphate Na3V(PO3)3N as a High-Voltage Positive Electrode Material for Na-Ion and Li-Ion Batteries. Electrochem. Commun. 2017, 84, 14–18. 152. Wang, X.; Kurono, R.; Nishimura, S.; Okubo, M.; Yamada, A. Iron-Oxalato Framework with One-Dimensional Open Channels for Electrochemical Sodium-Ion Intercalation. Chem. A Eur. J. 2015, 21 (3), 1096–1101. 153. Yao, W.; Sougrati, M. T.; Hoang, H.; Hui, J.; Lightfoot, P.; et al. Na2Fe(C2O4)F2: A New Iron-Based Polyoxyanion Cathode for Li/Na Ion Batteries. Chem. Mater. 2017, 29 (5), 2167–2172. 154. Yao, W.; Sougrati, M. T.; Hoang, H.; Hui, J.; Lightfoot, P.; et al. Reinvestigation of Na2Fe(C2O4)3.2H2O: An Iron-Based Positive Electrode for Secondary Batteries. Chem. Mater. 2017, 29 (21), 9095–9101. 155. Shiva, K.; Singh, P.; Zhou, W.; Goodenough, J. B. NaFe2PO4(SO4)2: A Potential Cathode for a Na-Ion Battery. Energ. Environ. Sci. 2016, 9 (10), 3103–3106. 156. Asl, H. Y.; Choudhury, A. Combined Theoretical and Experimental Approach to the Discovery of Electrochemically Active Mixed Polyanionic Phosphatonitrates, AFePO4NO3 (A ¼ NH4/Li, K). Chem. Mater. 2016, 28 (14), 5029–5036. 157. https://www.catl.com/en/. 158. Momma, K.; Izumi, F. VESTA 3for Three-Dimensional Visualization of Crystal, Volumetric and Morphology Data. J. Appl. Cryst. 2011, 44, 1272–1276.

7.10

Electrode materials viewed with transmission electron microscopy

Elena D. Orlova, Anatolii V. Morozov, and Artem M. Abakumov, Skolkovo Institute of Science and Technology, Moscow, Russia © 2023 Elsevier Ltd. All rights reserved.

7.10.1 7.10.2 7.10.2.1 7.10.2.2 7.10.3 7.10.4 7.10.5 7.10.5.1 7.10.5.2 7.10.5.3 7.10.5.4 7.10.6 7.10.7 7.10.8 Acknowledgement References

Introduction Transmission electron microscopy techniques in brief TEM data visualization, manipulation and treatment (S)TEM image simulation Electron beam damage in transmission electron microscopy Electron diffraction techniques for metal-ion battery electrodes Imaging of the local crystal and defect structure Point defects and order-disorder in electrode materials Planar defects Phase boundaries, grain boundaries and surfaces Imaging in 3D Spectroscopy with electrons In situ and operando observations of electrochemical reactions Conclusions and outlook

272 273 283 283 284 285 291 292 297 301 304 305 309 315 319 320

Abstract Advanced transmission electron microscope is no more a simple “picture-taking” machine, but a powerful analytical tool combining various cutting-edge imaging and spectroscopy techniques. Nowadays comprehensive investigation of electrode materials, as well as their further development, cannot be imagined without transmission electron microscopy (TEM). Indeed, only TEM can simultaneously deliver a whole range of local structural and chemical data with sub-Å spatial and sub-eV energy resolution, providing valuable contributions to the fundamental understanding of key factors affecting the electrochemical performance of various electrode materials. This Chapter outlines the traditional and cutting-edge TEM techniques along with examples of their successful application to the electrochemical systems with the primary focus on metal-ion batteries and electrocatalysis in order to demonstrate a wide range of TEM opportunities for electrode materials exploration. Intrinsic limitations of the TEM methods and possible sources of artifacts are also discussed in the specific context of the electrochemical systems.

7.10.1

Introduction

Prospects for the “green” future based on renewable energy sources, electric transportation, hydrogen economy and declining use of fossil fuels have triggered enormous interest to electrochemical energy storage and conversion technologies. The performance of such electrochemical devices as metal-ion batteries and fuel cells largely depends on the proper design of the electrode materials that in turn calls for improved understanding of the intricate relationships between their chemical composition, crystal and electronic structure, both in the bulk and at the surface, peculiarities of charge transfer at the electrode/electrolyte interface, mechanisms of the surface catalytic reactions, degradation phenomena, ionic diffusion, reversibility of cationic (de)intercalation and related chemical and structural changes, and many other aspects that constitute the ultimate complexity of the electrochemical systems. No wonder that whole diversity of available diffraction, spectroscopic and imaging techniques has been applied alone or in combination, using ex situ or in situ/operando regimes to interrogate the electrode materials at different spatial and temporal scales. The long-range ordered crystal structures are investigated with synchrotron X-ray and neutron powder diffraction or combination of both, the local atomic arrangement is analyzed with pair distribution function analysis, nuclear magnetic resonance spectroscopy and extended X-ray absorption fine structure, the electrochemically-induced changes in the oxidation states of the elements and electronic structure are probed with X-ray absorption spectroscopy and resonant inelastic X-ray scattering, the surface states are assessed with X-ray photoelectron spectroscopy. However, none of these techniques provides such remarkable locality as can be probed with advanced transmission electron microscopy (TEM). The first transmission electron microscope has been invented by Max Knoll and Ernst Ruska in 1931, but the last two decades were marked with tremendous progress in the development of TEMs, nowadays reaching unprecedented sub-Å spatial and < 0.1 eV energy resolutions thanks to several key improvements in the electron sources, optics and detectors. Perfectioning hardware is complemented with the development of new imaging and data treatment techniques allowing visualization of the elements with low atomic number (including even hydrogen),1 atomicresolution imaging at very low electron dose, column by column collection of spectroscopic data, three-dimensional (3D) viewing

272

Comprehensive Inorganic Chemistry III, Volume 7

https://doi.org/10.1016/B978-0-12-823144-9.00031-5

Electrode materials viewed with transmission electron microscopy

273

of the elemental distribution, quantification of atomic displacements, the population of the crystallographic positions and oxidation states, crystal structure solution and refinement from diffraction data collected on submicron-sized crystals. A combination of analytical techniques for assessment of local structure and chemistry makes modern transmission electron microscope more than just a picture-taking machine, but a whole material science laboratory whose capabilities can be applied to a nanometer-sized area of the electrode.2 Regarding the ability to retrieve local information, TEM is unmatched among other characterization techniques. The goal of this chapter is to provide the material scientists and electrochemists with an overview of the advanced TEM techniques with particular emphasis on their applicability to various energy-related materials, such as metal-ion battery electrodes and fuel cell electrocatalysts. This overview is preceded with a brief survey of the transmission electron microscope structure, operation modes and implementation of various TEM methods, which will then be illustrated with concrete examples from electrochemical material science. The vast majority of the experimental examples provided in this chapter are confined to positive (cathode) or negative (anode) electrodes for lithium- or sodium-ion batteries represented with the layered complex oxides of alkali and transition metals, and with the large group of so called polyanion materials (i.e. containing phosphate, sulfate, silicate or other multiatomic anionic groups in the structure). Oxides with the perovskite-type structures are used to illustrate the ability of TEM to unravel the defect and surface structures of lithium-ion conducting solid electrolytes and electrocatalysts of the oxygen evolution and oxygen reduction reactions in low-temperature alkaline electrolyzers and fuel cells. Particular consideration will be given wherever possible to artifacts created by the interaction of high-energy electron beam with the electrode materials and electrolytes, complemented with a special section on mechanisms of electron beam damage specific to the abovementioned groups of materials. A section on electron diffraction will provide an introduction into the world of quantitative electron crystallography and its use for elucidating the crystal structures of metal-ion battery cathodes at different states of charge. Atomic-resolution imaging of point and extended defects, phase boundaries, grain boundaries and surfaces will be complemented with examples of three-dimensional (3D) visualization of the material’s microstructure. As part of TEM measurements can be conducted in situ or in operando regime inside the dedicated electrochemical cells with either liquid or solid electrolytes, offering an opportunity to look at the electrochemical systems in their native state, the advantages and limitations of such approach will also be discussed in a separate section. Finally, the abilities of spectroscopic techniques, based on characteristic X-rays and electron energy losses, to probe local chemistry, electronic structure and chemical bonding will be illustrated. Due to space limitation, we are not able to comprehensively review all the aspects of the theory and practice of the TEM techniques (which could be found in the famous textbook by D.B. Williams and C.B. Carter3) and all reported applications of TEM to the electrochemical materials, systems and problems that encounter hundreds or may be thousands of publications. The main goal is to highlight emerging possibilities delivered by TEM for enhancing our understanding of the complex electrochemical systems and facilitating the development of the advanced electrode/electrolyte materials.

7.10.2

Transmission electron microscopy techniques in brief

The basic components of the transmission electron microscope (TEM) column are reminiscent to those of the conventional optical microscope. Although manipulating electron beams and light rays requires completely different hardware, the major parts of both devices (light/electron source, illumination (condenser) system, sample stage, objective lens, intermediate and projector lenses, image-recording devices, see Fig. 1) play very similar roles.4 In contrast to the light microscope, the optical system of TEM is built of magnetic lenses, which can change their strength and focal length by changing electric current passing through the coils of the lenses and creating axially-symmetric magnetic field. The ray paths of the electrons passing through the lens (except those traveling through its center) are bend by the Lorentz force toward the main optical axis being focused at the focal plane of the lens. Electron gun provides a beam of high energy electrons accelerated in the electrostatic potential applied between the cathode, serving as an electron source, and anode. Electrons are extracted from the cathode due to thermoemission from a heated tungsten of LaB6 single crystal filament, or by field-emission due to tunneling effect when high extraction potential is applied to a very sharp tip of a tungsten needle. The wavelength of the emitted electrons can be calculated from applied accelerating voltage U as: h 1:226 l ¼ pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ¼ pffiffiffiffi 2em0 U U

(1)

where e and m0 are the electron charge and mass, respectively, U is taken in volts, and l – in nanometers. As the velocity of high energy electrons constitutes measurable fraction of the speed of light, a correction taking into account the relativistic effects at U > 100 kV is applied for a more precise calculation of the electron wavelength: 1=2 h 1:226 l ¼ sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ 9:79:107 U ffi ¼ pffiffiffiffi  U 2m0 eU 1 þ 2meU0 c2

(2)

The velocities and wavelengths of the electrons at the energies typically used in transmission electron microscopes are listed in Table 1. Due to the strong Coulomb interaction of charged electrons with the electrostatic potential of the material, the mean free path of electrons in a TEM specimen is about few tens of nm that naturally limits the thickness of the TEM specimens. As the differential cross section of the electron scattering by atom is roughly proportional to its atomic number as Z2, increasing the average atomic

274

Electrode materials viewed with transmission electron microscopy

Fig. 1 Principal components of the TEM column including (from top to bottom) electron gun, condenser system, sample stage, objective lens, projector lenses, as well as a set of apertures, electron detectors and cameras. Note that the lenses here are the magnetic lenses whose strength can be adjusted by changing electric current through the lens coils. Detector of secondary X-ray radiation and electron energy loss spectrometer are also shown. Many other essential parts of the electron optics (monochromator, beam deflectors, stigmator coils, aberration correctors, etc.) are not shown for the sake of simplicity. Modified from Inkson, B.J., Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) for materials characterization. In Materials Characterization Using Nondestructive Evaluation (NDE) Methods, Hübschen, G.; Altpeter, I.; Tschuncky, R.; Herrmann H.-G. Ed.; Elsevier, 2016; pp. 17–43. Copyright 2016 Elsevier.

Table 1

Velocities and wavelengths of the electrons at the energies typically used in transmission electron microscopes.

Electron energy (keV)

Electron velocity (108 m.s 1)

Electron wavelength (A˚)

100 120 200 300 400

1.644 1.759 2.086 2.330 2.484

0.0370 0.0335 0.0251 0.0197 0.0164

Electrode materials viewed with transmission electron microscopy

275

number of constituting elements requires progressive lowering of the sample thickness making the preparation of thin, electrontransparent specimen of ultimate importance. As proper TEM specimen preparation largely determines the quality and interpretability of the subsequent imaging, diffraction and spectroscopic measurements, a number of dedicated techniques have been developed for material’s thinning. Uniformly thin slices of matter can be prepared by mechanical cutting, polishing and final thinning with a high energy beam of Arþ ions or using the focused ion beam (FIB), where the material is milled away by high energy Gaþ ions of Xeþ ions focused onto the sample by a dedicated optical system consisting of electrostatic lenses. Depending on the nature of the sample, other techniques, such as electropolishing, drop casting, vitrification, microtome sectioning, taking surface replica can be employed. For brittle inorganic materials, such as oxide ceramics, grinding the sample into powder with a mortar and pestle, followed by dispersing the powder in some organic solvent and depositing it onto holey amorphous carbon film supported by a Cu grid usually provides crystallites with electron-transparent edges formed due to cleavage of larger crystals. This way is most commonly used for the post mortem investigation of electrode materials and requires the operations to be performed in air-free atmosphere due to sensitivity of the electrodes to oxygen, moisture and CO2 (see Section 7.10.5.1 for details). The condenser system of TEM takes the electrons from the gun and transfers them to the specimen providing illumination with either a broad parallel beam or a focused beam (often called an ‘electron probe’). In the conventional TEM mode, the condenser lenses are adjusted to illuminate the specimen with a parallel beam of electrons, typically several micrometers across (Fig. 2A). In the scanning TEM (STEM) mode, the illumination system creates a focused probe by converging the beam at the specimen into a spot of sub-nm or even sub-Å size (Fig. 2B), which scans over the specimen using the STEM scan coils (Fig. 1). If in the TEM mode the image of the whole area of interest is recorded at once, in the STEM mode image recording is a sequential process of capturing the signal point by point scanned with the electron probe. The objective lens is the strongest lens in the TEM, creating the first image of the specimen. Let’s assume that the specimen is a thin crystalline plate positioned at the object plane of the objective lens and illuminated with a parallel electron beam (Fig. 2A). The sample contains a family of crystallographic planes hkl, and their orientation with respect to the incoming electron beam satisfies the Bragg’s law (i.e. 2dhklsinq ¼ nl, dhkl is the interplanar spacing for the hkl family, q is the angle between the hkl plane and the direction of the primary electron beam, n - integer). The electrons which travel along the optical axis of the TEM column being not scattered by the specimen will be focused by the objective lens into a point at the back focal plane (main focus, blue rays in Fig. 2A). The rays diffracted by the hkl planes also form a parallel electron beam traveling at the angle 2q to the main optical axis; they will be focused into another point at the back focal plane (side focus, red rays in Fig. 2A). Thus, the spots corresponding to the 000 (transmitted electron beam) and hkl (diffracted electron beam) reflections are observed meaning that the electron diffraction (ED) pattern is formed at the back focal plane of the objective lens, which is also called a diffraction plane. The first image of the

Fig. 2 Parallel-beam (A) and convergent-beam (B) illumination modes in TEM. The rays transmitted through the specimen without scattering are colored in blue, the rays scattered by the specimen are colored in red. The angular ranges for bright field (BF), annular dark field (ADF) and high angle annular dark field (HAADF) collection areas are schematically marked.

276

Electrode materials viewed with transmission electron microscopy

specimen appears as a result of interference of the transmitted electron beam and diffracted beams (which could be many) at the image plane, conjugated with the object plane of the objective lens. By changing the strength of the projector system lenses either the diffraction pattern or specimen image can be further transferred to a registration device which can be a fluorescent viewing screen with a TV camera or electron camera based on a charge coupled device (CCD), complementary metal oxide semiconductor (CMOS) electronics or direct electron detector (Fig. 1). This way one can easily switch between the direct and reciprocal space of a crystalline specimen to observe its electron diffraction pattern or lattice image. The excitation of the projector lenses defines the magnification of the image in the direct space or magnification of the diffraction pattern in the reciprocal space, the latter is also referred to as “camera length.” Due to substantially shorter wavelength of accelerated electrons in TEM compared to that of laboratory X-ray radiation or thermal neutrons (for instance, lelectron ¼ 0.0251 Å at U ¼ 200 kV and lX-ray ¼ 1.5406 Å for CuKa1 radiation typical for lab X-ray powder diffractometers), the ED pattern of a crystalline material to the first approximation can be considered as a section of its reciprocal lattice normal to the direction of the primary electron beam (Fig. 3A). This approximation originates from the radius of the Ewald sphere r ¼ 1/l which is much larger for electrons than for X-rays so that the Ewald sphere for electrons is so flat that can

Fig. 3 Typical electron diffraction pattern demonstrating a reciprocal lattice section normal to the [001] crystallographic direction (A). The hkl indexes for the reciprocal lattice vectors along a* and b* reciprocal directions are marked. Bragg equation for electron diffraction patterns as a result of similarity of two right-angled triangles in reciprocal and direct space (B).

Electrode materials viewed with transmission electron microscopy

277

roughly be considered as a plane cutting the reciprocal lattice that also simplifies the expression of the Bragg’s law for the ED patterns. From the similarity of two right-angled triangles – one in reciprocal space with the catheti 1/l and 1/dhkl and another one in the direct space with the catheti L and rhkl – it follows that dhkl ¼ Ll/rhkl, where dhkl is the interplanar distance for the crystallographic planes hkl, rhkl is the distance between the reflections 000 and hkl measured from the ED pattern and L is an effective “camera length” (measure of the ED pattern magnification) (Fig. 3B). The geometry of the reflection net in the ED patterns and measured interplanar spacings are used for the analysis of the crystal symmetry and unit cell dimensions,5 the reflection conditions indicate presence/absence of lattice centering, glide planes and screw axes, whereas reflection intensities can be used for crystal structure solution and refinement provided that the dynamical diffraction effects are reduced or rigorously taken into account (see Section 7.10.4).6–10 As a reciprocal lattice of a crystal is a three-dimensional object, several reciprocal lattice sections are needed to enable its reconstruction that requires collection of the ED patterns along different crystallographic directions. In the most conventional way, several sub-micron sized single crystals are selected in TEM and series of the ED patterns are collected by rotating them around selected directions. The ED patterns can be registered discretely while the crystal is viewed along the most prominent low-index crystallographic directions, or in quasi-continuous manner when the crystal is rotated over a large angular range and the patterns are taken with a small angular step that is a basis of electron diffraction tomography (EDT).11 The double-tilt specimen holders are most frequently used for such electron diffraction experiments. These holders serve as two-circle goniometers, where, however, the angular ranges might be severely limited by the construction of the holder and the space limitations within the objective lens of TEM, where the holder tip is placed. The typical tilt range of a double tilt holder is  40–60 along the holder’s rod and  20– 30 along the perpendicular direction. Precise manipulation of the orientation of the specimen is achieved through a computercontrolled sample stage. Dedicated tomography holders with narrower tip allow for larger tilt angles. For air-sensitive samples (such as charged Li-ion battery cathode materials), the vacuum-tight double-tilt holders offer transfer to the TEM column under argon or vacuum enabling specimen preparation in an Ar-filled glove box (see Section 7.10.5.1). As the lateral dimensions of a typical TEM specimen are much larger than its thickness in the electron beam direction, the specimen is described by a two-dimensional distribution of the electrostatic potential V(r) projected along the incident electron beam. The incoming electron wave interacts with the projected potential V(r) of the crystal resulting in the exit wavefunction f(r), r is in the plane of the exit face (Fig. 2A). The amplitude of the diffracted wave corresponding to the reciprocal vector g is given by a Fourier transform of the exit wavefunction F(g) ¼F [f(r)]. The back focal plane visualizes the square of the Fourier transform of the object in the form of the electron diffraction pattern. The reflections in the back focal plane act as sources of secondary waves, which interfere with each other creating the image of the object at the image plane. The amplitude in the image is given by inverse Fourier transform g(r) ¼ F 1 [F(g)], thus restoring the wave function at the exit face of the crystal.12 The high resolution TEM (HR-TEM) image obtained this way is the result of interference of the transmitted beam and diffracted beams taking into account their phase shifts due to interaction with the crystal potential and due to defocus and aberrations of the objective lens. Thus, in reality, the wave function at the image plane g(r) is affected by distortions from the optical system of the microscope. Generally, the observed contrast in the HR-TEM image, being affected by the imaging parameters and aberrations, has no one-to-one correspondence with the exit wave function (Fig. 4). The relationship between the experimentally observed image contrast and the exit wave function is highly nonlinear that hampers the intuitive interpretation of the HR-TEM images. An additional complication with revealing the projected potential from the HR-TEM image is caused by the fact that the exit wave function is a product of dynamical interaction between the crystal potential and incident plane electron wave, i.e. the exit function is thickness-dependent. Thus, extensive computational modeling of the contrast behavior as a function of the imaging conditions and specimen thickness is necessary to relate the observed HR-TEM contrast to the anticipated crystal structure projection. This might pose serious obstacles if the crystal structure is not known a priori and should be retrieved from the HR-TEM images leading to a series of trial-and-error attempts to construct the plausible structure model and modify it to achieve the best correspondence between the experimental and calculated HR-TEM images. Different techniques are employed in order to decipher the crystallographic information hidden in the HR-TEM images, such as crystallographic image processing,13 taking a series of images with variable defocus (through-focus series),14 or imaging of projected crystal potential with negative spherical aberration coefficient.15 Difficulties with the direct and straightforward interpretation represent a significant disadvantage of the HR-TEM imaging that has largely been overcome with the advancements in atomic-resolution scanning transmission electron microscopy (STEM) imaging techniques.16 In the STEM mode the sample is illuminated with a focused electron beam in a form of a cone with the convergence semiangle a thus forming a very fine probe of sub-Å size at the specimen (Fig. 2B). In the back focal plane a convergent beam electron diffraction (CBED) pattern is formed which consists of the bright field disk formed by the electrons transmitted through the specimen (colored in blue in Fig. 2B) and Bragg reflections disks (colored in red in Fig. 2B) with the disk diameter proportional to a. Besides the Bragg diffraction due to coherent elastic interaction of electrons with the crystal electrostatic potential, part of the incident electron beam is scattered over high angles by the Rutherford scattering (i.e. incoherent elastic electron interaction with the nucleus) and the inelastic thermal diffuse scattering (marked in green in Fig. 2B). When the electron probe scans over the specimen, the electrons at the back focal plane are collected with the disk- or ring-shaped detectors (Fig. 1), and the registered intensity at every point of the scanned raster is plotted against the position of the probe thus forming an image. The disk-shaped bright field (BF) detector picks up the forward-scattered electrons, and the obtained image is reminiscent to that obtained in HR-TEM. For the ring-shaped detector the diameters of the disk and hole define the inner and outer collection angles and the angular range, in which the scattered electrons are recorded. These angles can be varied without changing the physical size of the detector by adjusting the

278

Electrode materials viewed with transmission electron microscopy

Fig. 4 Comparative imaging of the same structure (top) with the HR-TEM, HAADF-STEM and ABF-STEM techniques. The structure consists of heavy cation columns (Pb, Bi, shown as large circles), lighter cation columns (Fe, medium-sized circles) and light anion columns (O, small circles). HR-TEM image does not show direct correspondence with the structure projection. Only cationic columns are visible as brighter (Pb, Bi) and fainter (Fe) dots in the HAADF-STEM image. The structure projection is fully resolved in the ABF-STEM image (note the contrast is inverted with respect to HAADF-STEM image and the atomic columns appear dark).

Electrode materials viewed with transmission electron microscopy

279

“camera length” L that is the magnification degree of the diffraction pattern. The annular dark-field (ADF) detector collects the electrons scattered over small angles typically ranging to 3 except those falling into the BF cone. However, the most informative are the STEM images formed with electrons incoherently scattered over high angles and picked up by the high angle annular dark-field (HAADF) detector. In the HAADF-STEM images the contribution of the Bragg scattering is minimized and the signal strongly depends on the chemical composition. The intensity in atomic resolution HAADF-STEM images roughly scales as Zn (n < 2), where Z is the average atomic number along the projected atomic column,2,17,18 hence the term “Z-contrast” imaging. This makes the interpretation of the HAADF-STEM images more straightforward because the atomic columns appear as bright dots on a dark background and their position and intensities directly correlate with the projected electrostatic potential of the crystal (Fig. 4). The lenses are not directly involved in the image formation making HAADF-STEM contrast relatively robust to the defocus and crystal thickness variations. The spatial resolution of the HAADF-STEM technique depends on the size of the focused electron probe given by the degree of the electron source demagnification and might be greatly improved by aberration correctors for the probe-forming condenser system of the TEM. A useful measure of the spatial resolution in the STEM mode is given by full-width at half-maximum (FWHM) of the electron probe dprobe,19 which is contributed by diffraction limitations ddif and the source size dsrc limitations required to keep reasonable source brightness: dprobe ¼ d2dif þ d2src

(3)

1=4

(4)

ddif ¼ 0:4l3=4 Cs dsrc ¼

rffiffiffiffiffiffiffiffiffiffiffiffiffiffi 4I Bp2 a2

(5)

where l is the electron wavelength, Cs is a spherical aberration coefficient of the probe-forming lens, B is the source brightness and I is the beam current, a is the convergence semiangle. From one side, strong almost quadratic dependence of the signal in the HAADF-STEM images on the atomic number facilitates their direct interpretation, but from another side, it appears as a significant disadvantage rendering “light” elements, such as oxygen and lithium, virtually invisible, particularly in the presence of the atoms with large Z. Nevertheless, they can be visualized using annular bright-field (ABF) STEM imaging (Fig. 4) which utilizes contributions of both coherently and incoherently scattered electrons so that the Z-dependence of the signal scales approximately as Z1/3,20,21 allowing very light atom columns to be visible next to the columns of heavy cations.1,22 Compared to the HAADF-STEM images, the ABF-STEM contrast is inverted and the atomic columns appear dark on a bright background enabling qualitative interpretation and assignment of atomic columns. A combination of the HAADF-STEM and ABF-STEM images taken with sub-Å resolution is a very powerful tool for retrieving local structural information. However, practically-wise the ABF-STEM imaging demands the specimen to be thin being more sensitive to the thickness variation compared to HAADF-STEM that might cause severe limitations if the near-surface areas are corrupt or adopt the structure different from the bulkier part of a crystal. With the task to extract quantitative information on the atomic positions and interatomic distances from high resolution TEM images such parameter as signal-to-noise ratio comes to the forefront because the precision of the measurement of the separations between the atomic columns is inversely proportional to the square root of the number of detected electrons23: pffiffiffi 2w2 sz pffiffiffiffiffiffi; if d  2w d Ne

(6)

pffiffiffi pffiffiffi 2w sz pffiffiffiffiffiffi ; if d  2w Ne

(7)

where s is a standard deviation in the measured distance d, w is a width of the atomic column in the image (representing the spatial resolution) and Ne is a total number of electrons collected to form an image of this atomic column. Unfortunately, the electron detectors used in the HAADF- and ABF-STEM techniques collect only small fraction of scattered electrons requiring a relatively high electron dose to maintain reasonable signal-to-noise ratio. Employing advanced segmented or high-speed pixelated detectors provides a way to collect a much larger fraction of the scattered electrons thus enabling low-dose STEM techniques such as differential phase contrast (DPC) and ptychography. In the latter, a pixelated detector captures an image of CBED pattern at every point of the specimen scanned by the electron beam (Fig. 5A), in contrast to the conventional annular ring-like STEM detector which registers only the integral intensity of the electrons scattered over a certain angular range. An atomic-resolution phase image of the specimen is then reconstructed from the interference between the transmitted beam and diffracted beams. Such images were proven to show the positions of both “heavy” and “light” elements being recorded even at sub-picoampere electron currents.24,25 A phase image of the exit electron wave can also be directly recorded with integrated differential phase contrast (iDPC-STEM) technique which is based on a linear relationship between the gradient of phase shift of the electron wave and position of center of mass (COM) of the intensity of the CBED pattern. The COM position at every point of the sample is mapped with a 4-quadrant STEM detector (Fig. 5B), and 2D integration of this map provides the iDPC-STEM image reproducing the projected electrostatic potential of the specimen. The differentiated DPC image (dDPC-STEM) can be interpreted as a projected charge density map of the specimen.26 The contrast dependence on atomic number in iDPC-STEM images is linear that is more appropriate for visualizing

280

Electrode materials viewed with transmission electron microscopy

Fig. 5 CBED pattern captured with a pixelated detector (shaded in gray) at every scanned point of the specimen in a ptychography imaging (A) and displacement of the center of mass of the CBED pattern (red arrow) registered with a 4-quadrant annular detector in DPC-STEM imaging (B). The disk of the transmitted undiffracted beam is shaded in blue, the disks of the diffracted beams are shaded in red.

“light” elements than nearly quadratic dependence in the HAADF-STEM imaging.27 Particular advantage of the DPC images comes from the fact that they maintain a good signal-to-noise ratio being recorded at the electron dose at least few orders of magnitude lower than required in the conventional HAADF- and ABF-STEM techniques. Treating the (S)TEM images as a source of two-dimensional (2D) projected information of the real 3D distribution of the scattering density in the object is related to the ratio between the specimen thickness (typically below few tens of nm) and depth of field (DoF) of TEM that is the thickness range of the object which is viewed in focus. For a point object, DoF is commonly defined through a Rayleigh’s criterion of resolution, and for incoherent imaging in an aberration-corrected microscope with sub-Å spatial resolution DoF can be estimated as 2l/a2, where l is the electron wavelength, a is the convergence semiangle.28 For U ¼ 200 kV electron beam (l ¼ 0.0251 Å) and a ¼ 21 mrad as typical imaging conditions, the DoF is  11 nm meaning that a significant part of the specimen thickness is simultaneously in focus. For observation of a finite-sized object DoF can be estimated as rffiffiffiffi 3 d (8) DoFz 2 a where d is the diameter of the object.29,30 For nanoparticles with the average diameter of 5 nm at a ¼ 21 mrad, the DoF spreads over  300 nm meaning that the entire specimen is in focus. The essential 2D character of the (S)TEM images makes the assessment of 3D features cumbersome and sometimes misleading. This issue can be resolved with STEM tomography,31 although this method often requires complex and multistage preparation of dedicated specimens. Incoherent imaging conditions in the HAADF-STEM mode result in monotonic dependence of the contrast on the projected scattering density required for tomography reconstruction of a 3D object from a set of 2D projections. A STEM tomography experiment involves the acquisition of a tilt series of images, covering the highest accessible range of angles, where the images are recorded at every angle with a step of 1–2 (Fig. 6). Because of geometric restrictions caused by the specimen holder size and limited space available for its rotation in the microscope, the available tilt range is limited to  70–80 that results in undersampling of a certain region of the specimen known as “missing wedge.” The collected set of images is going through alignment (i.e. bringing them to a common coordinate system eliminating the effects of small sample movement) and 3D reconstruction with one of various mathematical algorithms (see Refs. 32, 33 for an overview) followed by visualization. The quality of the 3D reconstruction is most severely limited by the number of 2D projected images and the missing wedge artifacts, thus additional experimental efforts are invested to increase the tilt range, such as using a dedicated rod-like sample holder and needle-like specimens. In addition to the crystallographic and atomic structure information retrieved from electron diffraction and imaging, wealth of chemically-sensitive information can be extracted from spectroscopic data originating from inelastic interactions of high-energy electrons with the matter.34 Among various spectroscopy techniques, energy-dispersive X-ray (EDX) spectroscopy and electron energy loss spectroscopy (EELS) are the most widespread analytical tools implemented into TEMs. Incident electrons can interact with the atoms of the solid specimen transferring a small fraction of their energy to the electrons of the inner atomic shell thus kicking them out to the unoccupied electron states above the Fermi level (Fig. 7). The energy losses of the transmitted electrons contribute to the core-loss part of the EELS spectrum. When the hole in the inner shell is filled back by an electron from higher energy levels, a photon is emitted with the energy corresponding to the energy difference between the upper and lower levels. This characteristic X-ray radiation is measured in EDX spectroscopy (Fig. 7).

Electrode materials viewed with transmission electron microscopy

281

Fig. 6 An electron tomography experiment, including the acquisition of a tilt series within a limited angular range (A, B) (missing wedge is indicated) and 3D reconstruction through back projection of the 2D images along their acquisition directions (C). Reprinted with permission from Bals, S.; Goris, B.; Altantzis, T.; Heidari, H.; Van Aert, S.; VanTendeloo, G. C. R. Physique 2014, 15, 140–150. Copyright 2014 Elsevier.

Fig. 7

Scheme of electronic transitions generated in an atom by incident high energy electron and typical appearance of the EDX and EELS spectra.

282

Electrode materials viewed with transmission electron microscopy

Both EDX and EELS spectra can be recorded in TEM and STEM modes and used to identify the elements constituting the specimen, determine their quantities and build their spatial distribution maps.35,36 Quantification of the EDX spectra is based on the relation between the intensities of the peaks in the EDX spectrum for the elements A and B, IA and IB, and the elemental concentrations CA and CB: IA sA CA  IB sB CB

(9)

where sA and sB are the inelastic electron scattering cross-sections for the elements A and B, respectively.37 The proportionality coefficient sA/sB is called the Cliff-Lorimer factor; for every EDX spectrometer and acceleration voltage these factors are tabulated for the intensities of characteristic X-ray lines of the analyzed elements with respect to the intensity of the Si Ka line (the so-called kfactors) and can be used for standardless quantitative compositional analysis. Similar expression can also be applied to the quantification of the EELS spectra, but also the concentration of the element can be calculated directly from the intensity of the corresponding EELS edge IA and the intensity of transmitted electron beam I0: IA  NA sA I0

(10)

where NA is the number of the atoms A per unit area, and the intensity I0 can be measured directly from the zero-loss peak in the EELS spectrum (note that this formula neglects multiple electron scattering being valid only for sufficiently thin specimens). Strong confinement of the electrons to the atomic columns oriented along the beam direction enables atomic resolution elemental mapping with both spectroscopic techniques. The STEM variant of EDX spectroscopy has acquired a strong boost after the invention of advanced silicon drift (SDD) X-ray detectors with large solid collection angles. Their compact design allows placing several SDD detectors in a symmetrical manner around the sample area (Fig. 8) and achieve the total detector area of about 120 mm2 and the collection angle up to 0.9–1.2 sr enhancing the speed of data collection by a factor of up to  50 compared to a conventional Si(Li) X-ray detector.38 Windowless design of the SDD detectors improves the sensitivity to “light” elements, whereas their symmetrical arrangement around the specimen makes the EDX technique much less sensitive to the specimen tilt (Fig. 8) that is important if the chemical mapping at a certain specific orientation of the specimen with respect to the electron beam is required, or the data have to be collected over a large range of tilt angles to enable compositionally-sensitive 3D tomography reconstruction. Although due to rapidly falling X-ray fluorescence yield with lowering Z, EDX cannot be used for mapping atoms lighter than carbon, EELS can be effectively used for that. A combination of the monochromated electron source and advanced high energy resolution EELS spectrometer allows recording EELS spectra with the energy resolution below 0.1 eV (i.e. comparable to that of X-ray absorption spectroscopy) thus revealing very fine details of the electron energy loss near edge structure (ELNES) which carry information on the unoccupied local density of states. Energy position and intensity distribution in the ELNES spectra depend on the oxidation states of the transition metals and their coordination number and could be used for mapping this chemically-sensitive information at different spatial scales and even with atomic resolution.39–41 A brief overview of the vast number of TEM-based characterization techniques should be complemented with highlights of some useful software, which allow stepping forward from data acquisition toward decent analysis. The conscious and careful treatment of gathered TEM data is the key point on the way to relate the observations to the peculiarities of the material’s behavior. A wide variety of freeware is available nowadays for sustainable TEM data handling. However, one should note that the provided list of TEM-

Fig. 8 A scheme illustrating four SDD X-ray detectors arranged symmetrically around the sample and the objective lens pole pieces (A). Comparison of relative EDX count rates of the 4 SDD X-ray detector system and a single Si(Li) detector (B), with the diagrams showing the effects of detector shadowing for the four SDD X-ray detectors (upper row), and diagrams of shadowing effects for the single detector system (bottom row). Reprinted with permission from ChemiSTEM Technology-Application-Brochure, Thermo Fisher Scientific, 2013.

Electrode materials viewed with transmission electron microscopy

283

related computer tools is by far not limited to the given set of programs, as the selection reflects the personal preferences of the authors and does not include the commercially available software.

7.10.2.1

TEM data visualization, manipulation and treatment

ImageJ (Fiji, https://imagej.net) is one of the most frequently used software designed for manipulations with visual data, including TEM microphotographs.42 Different manipulations with raw images, such as enhancing contrast, reducing noise, image deblurring, subtracting background and many others could be performed with various computational approaches embedded into ImageJ.43 ImageJ is capable of performing routine tasks, for example, measuring the interplanar spacing in HR-TEM and STEM images, interatomic distances in HAADF-STEM and reciprocal lattice vectors in ED patterns, plotting intensity profiles over selected areas, estimating particles size and plotting size distribution histograms, calculating fast Fourier transforms of the images. The fact that ImageJ is an open-source image processing software makes it very appealing to the international scientific community, thus, the numerous plugins and packages, provided not only by ImageJ developers, but also by scientific groups and independent researchers, have been created recently, and are being developed now in order to facilitate scientific image analysis. The enormous number of plugins bundles a wide range of computational methods, based on different mathematical models and approximations. In the list of various plugins there are packages for automatized distance measurements procedures, some of which are grouped for the most common TEM treatment tasks, for example, the one called “TEM suite” (https://imagej.net/TEM_suite); and plugins for more specific tasks, such as numerical post-treatment aimed at retrieving phase of the exit electron wave from limited number of TEM images44 or semi-automated image denoising.45 As for quantification of the atomic-resolution HAADF-STEM images, the techniques based on model-based parameter estimation theory (see Section 7.10.5.1) are realized in the StatSTEM software (https://github.com/quantitativeTEM/StatSTEM),46 which enables accurate and precise extraction of the information on the positions and integrated intensities of atomic columns with increased speed and over a large field of view due to the optimized image segmentation and calculation algorithm. First, the column positions are found with peak finder routine, which searches for local intensity maxima in the image. Then, the intensity distribution at each atomic column is modeled as a Gaussian peak, in accordance with an algorithm, described in Ref. 47 and coordinates of the atomic columns and the total intensity of electrons, scattered by atomic columns are fitted, making it possible to quantify the number of atoms in each atomic column and their respective locations. Based on these quantitative values and indexed column positions, a 3D model, which gives an idea of how a nanoparticle might look like, could be constructed. Besides, StatSTEM is designed to calculate and map atomic displacement and the components of the strain tensor, extracting them from comparison of expected and observed atoms coordinates. The rapid development of neural networks and machine learning has brought more advanced tools for image processing. One of the most prominent machine learning methods – deep learning – which is widely deployed in areas that require precise object detection, recognition, and classification, is deployed in AtomSegNet software (https://github.com/xinhuolin/AtomSegNet).48 Due to validated deep-learning models and large training datasets, AtomSegNet allows performing robust atom segmentation and localization along with proper noise reduction and deblurring of atomic-resolution STEM images, therefore, significantly cutting the time spent on image processing, and increasing the quality and precision in detection and localization of atomic columns from STEM images. For electron diffraction patterns treatment, the Crystallographic Tool Box (CrysTBox, https://old.fzu.cz/crystbox)49 is designed to deal with reciprocal space: it could be used for the SAED and CBED patterns processing, reciprocal space visualization along specified plane, automatic identification of individual diffraction spots and corresponding sample orientation (zone axis), and determination of interplanar angles and distances with up to a few picometer precision. Moreover, due to this tool, it is possible to visualize two misoriented phases and their interface, and perform dislocations, strain, and displacement mapping in the direct space high-resolution image. Noteworthy, the application allows calculating TEM holder tilts needed to achieve the desired zone axis, which could facilitate the sample preparation via FIB, when the sample should be cut and thinned close to the particular zone axis, and help less experienced users to operate the TEM in a faster manner. There are also other programs, which complement the functionality for treatment of the ED patterns, such as ProcessDiffraction (https://www.energia.mta.hu/labar/ProcDif.htm).50 Quantitative analysis of electron diffraction tomography (EDT) experiments (see Section 7.10.4), which imply collection and processing of a relatively large set of SAED patterns, can be performed with the PETS software (http://pets.fzu.cz).51 This program is dealing with searching for diffraction spots, clusterization of the reflections from all measured frames,52 “cleaning the peak list” by replacing the individual reflections with the cluster centers, indexing the cluster set, refinement of the lattice parameters and the orientation matrix, calculation of the reflections positions and integration of reflection intensities that can be used for crystal structure solution and refinement with other software, such as Jana (http://jana.fzu.cz).53

7.10.2.2

(S)TEM image simulation

In many cases, simulation of atomic-resolution (S)TEM images becomes a crucial step in a proper interpretation of the experimental observations and confirmation of the proposed structure models. Among other options, JEMS (http://www.jems-swiss.ch, freely available in Student Edition)54 and QSTEM (https://www.physics.hu-berlin.de/en/sem/software/software_qstem)55 offer a possibility to calculate HR-TEM, HAADF- and ABF-STEM images. The JEMS software offers a possibility to utilize the Bloch wave, Howie-Whelan or multislice methods to calculate HR-TEM images, while for HAADF-STEM images it uses the frozen lattice method; it is also possible to transform unit cells and efficiently define large supercells. Simulations of the TEM and STEM images in QSTEM

284

Electrode materials viewed with transmission electron microscopy

can be done using the multislice algorithm; sophisticated structure models consisting of several crystalline grains can be constructed and atomic models can be selected to match the atomic columns in the image.

7.10.3

Electron beam damage in transmission electron microscopy

Irradiation of a TEM specimen with high-energy electrons may cause unwanted and unexpected altering of the structure and chemical composition through various mechanisms of electron beam damage. A head-on collision of the electron with transferring part of its momentum to the atomic nuclei is at the origin of the knock-on damage mechanism, typically intrinsic in electrically conducting materials, such as metals and semiconductors. Knock-on damage involves displacement of the atomic entities from their native positions causing vacancy formation or exchange of the atomic species between neighboring crystallographic sites. The probability of this displacement increases with the energy of incident electrons (i.e. with the acceleration voltage), decreases with increasing atomic number Z, and also depends on the strength of the chemical bonding. The atomic displacement does not occur below the threshold energy Eth which can be estimated from the atomic weight A of the displaced species and the displacement activation energy Ed56: ! rffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi AEd 1þ Eth ðeV Þ ¼ ð511 KeVÞ 1 (11) 561 eV In intrinsically insulating or poorly conducting materials radiolysis becomes the dominant mechanism of electron beam damage. Radiolysis breaks the chemical bonds by inelastic energy transfer between the incoming electron and the atomic electron clouds resulting in displacement and ejection of atoms.56,57 These damage mechanisms are particularly crucial for the analysis of the layered Li-rich Li4/3  xNixMn2/3  xCoxO2 battery cathode materials at different states of charge, in which the electrochemicallyinduced cationic migration is the subject of interest as it might be at the origin of voltage fade,58–60 irreversible capacity and poor rate capability,61 capacity loss,62,63 or energy loss due to voltage hysteresis between charge and discharge.64,65 Unfortunately, the beam-induced atomic displacements in these Li-rich transition metal oxides tend to mimic well the electrochemically induced displacements that may cause artifacts and misinterpretations.66 The tendency to cation displacements increases for the metal-ion battery cathode materials in the charged state, in which there are vacant sites next to the transition metal cations electrochemically created due to deintercalation of the alkali cations.67 The probability of beam-induced cation displacement is not entirely tightened to getting sufficient energy to overcome the displacement threshold, but also can be associated with simultaneous changes in the valence state of the migrating cation. Ni cations in LiNiO2 are demonstrated to be more prone to migration to the Li positions than Co in LiCoO2 because of reduction of Ni3þ to Ni2þ (being closer in size to Liþ) and concomitant oxygen release under electron beam irradiation.68 In the large family of “polyanion” metal-ion battery cathodes (the materials with the general formula 3 2 4 2 3 AMxM’1  xXyY1  y (A ¼ Li, Na, K; M, M’ ¼ V, Mn, Fe, Co, Ni; X, Y ¼ PO3 4 , VO4 , SO4 , SiO4 , CO3 , BO3 )) the damage through radiolysis is particularly severe due to their highly insulating nature. It should be noted that the beam-induced transformations and damage become even more severe if the surface of the particles is of interest, as in nanosized electrode materials for electrocatalysis. The surface atoms are loosely bound and can be sputtered off the specimen by the electron beam or displaced along the surface, and the activation energies for these displacements are generally lower than those for the displacement of atoms within a crystal.56 This causes reconfiguration of the surface, leading to unusual reconstructions propagating toward sub-surface layers.69,70 Adjusting the electron energy through changing the acceleration voltage generally leads to a trade-off between knock-on damage (more severe at higher voltages) and radiolysis (prevailing at lower voltages). If at U ¼ 200 kV acceleration voltage the damage to the LiNi0.6Co0.2Mn0.2O2 Li-ion battery cathode structure was obvious, the original structure has been retained after prolonged observations at U ¼ 80 kV.71 The aberration-corrected electron optics allows for lowering the acceleration voltage without sacrificing much spatial resolution. However, this strategy might not work properly if the positions of the Liþ cations are of interest. The threshold of knock-on displacement for the Liþ cation is at U  30 kV at which most transmission microscopes do not operate yet. The atomic displacement cross section curve for lithium demonstrates anomalous behavior reaching maximum at 60–80 kV and then gradually descending,72,73 so that observations at U ¼ 120–200 kV might be less damaging for the Li positions than those at the lower voltage. For instance, Li2O2 can withstand 4–5 times higher electron dose at 200 kV acceleration voltage rather than at 80 kV.74 Another strategy which is proven to be successful is decreasing the electron dose. If the crystal structure is of interest at the submicron spatial scale, quantitative electron diffraction (see the Section 7.10.4) is the method of choice as electron diffraction tomography data can be collected at a very low electron dose rate,  250 times lower compared to that in the conventional atomic resolution STEM imaging75 that allows for the dissipation of the extra energy transferred from the electron beam. Further decrease of the electron dose can be achieved through the utilization of highly sensitive noise-free direct electron detectors.76,77 Performing atomic resolution TEM or STEM imaging at cryogenic conditions suppresses atomic displacements and allows for increasing the critical electron dose by a factor of 3–4 that, however, requires a dedicated low-drift TEM specimen holder.78 For instance, TEM imaging at liquid nitrogen temperature allowed for resolving the atomic structure of metallic lithium dendrites and their interface with a solid electrolyte that is important for understanding the dendrite growth at the anode site as one of the fault reasons in allsolid-state metal-ion batteries.79 True breakthrough has been provided by DPC-STEM imaging which utilizes a much larger fraction of electrons interacting with the specimen compared to the HAADF- and ABF-STEM techniques and maintains a high signal-to-noise ratio even at very low electron dose.24,80

Electrode materials viewed with transmission electron microscopy

7.10.4

285

Electron diffraction techniques for metal-ion battery electrodes

Laboratory/synchrotron X-ray and neutron powder diffraction in both ex situ and operando regimes remain dominant diffraction techniques to study the crystal structures of positive (cathode) and negative (anode) electrode materials for metal-ion batteries. They are capable of delivering structural information on the materials in their pristine state as well as in different stages of (de)intercalation of the alkali cations, also in a dynamical manner while the electrochemical cell is charged or discharged. Thus, the legitimate question remains on what advantages the electron diffraction can still demonstrate on top of the well-developed and widely available powder diffraction techniques? Indeed, it is well-known that the electron diffraction intensities are strongly distorted by dynamical and multiple diffraction effects stemming from the strong interaction of electrons with the matter. This significantly complicates retrieving the structure factors from the ED experiments imposing severe constraints on structure solution and refinement while X-ray and neutron diffraction can be analyzed in the assumption of single scattering events, allowing for using wellestablished algorithms of kinematical diffraction theory for the crystal structure solution and an accurate refinement of atomic coordinates, occupancy factors and atomic displacement (thermal) parameters. From the consideration above, very strong arguments must be provided to convince the material scientists to employ electron diffraction in their research. The first such argument is again related to very strong electron scattering by matter leading to a possibility to register the ED patterns from very small, typically submicron-sized crystallites, which are generally intractable with the conventional bulk diffraction techniques. By selecting small individual single crystals constituting a polycrystalline sample and taking their ED patterns along main crystallographic directions, their reciprocal lattice can be fully reconstructed as the ED patterns can be considered as the 2D reciprocal lattice sections (see the Section 7.10.2), in contrast to powder diffraction pattern representing 1D projection of the reciprocal lattice. Thus, ED patterns provide a straightforward way to determine the crystallographic shape of the crystallites that is particularly important for the electrode materials with anisotropic diffusion pathways. For instance, the olivine-structured phosphate cathodes, such as LiFePO4, possess 1D system of alkali cation diffusion with the easiest migration path aligned along the crystallographic direction b of the orthorhombic Pnma structure.81 The crystallites with the shortest dimension along the b axis are beneficial for these cathodes to ensure a high charge/discharge rate.82 The crystallite shape can be retrieved from a rotation ED

Fig. 9 TEM images and corresponding ED patterns of a rod-like crystal in olivine-structured NaFe0.2Mn0.8PO4. The rods have the longest dimension along the [010] direction and the shortest dimension normal to the (301) crystal planes ([203] direction, 10 off the [101] direction).

286

Electrode materials viewed with transmission electron microscopy

experiment as illustrated in Fig. 9, and used further as input to revise the synthesis technique and obtain the particles of the desired shape. Analysis of qualitative and quantitative phase composition of polycrystalline samples from powder X-ray diffraction data is now a routine automated procedure based on matching the experimental diffraction pattern with the standard single-phase patterns stored in databases, and applying various quantification procedures, such as based on reference intensity ratios (RIRs), internal/external standards or Rietveld refinement. ED can complement this essentially bulk information with local visualization of phase distribution among different crystallites. In LiFePO4, the industrially most relevant polyanion cathode material, extraction/insertion of Li proceeds through a two-phase mechanism, in which the lithiated LiFePO4 and delithiated FePO4 phases co-exist in almost entire range of Li concentrations. Mapping spatial distribution of LiFePO4 and FePO4 particles at different states of charge has been performed with electron diffraction, based on computer-aided matching between the experimental nanobeam ED patterns and the kinematically-calculated patterns of the LiFePO4 and FePO4 structures for different orientations with small angular step (Fig. 10A).83–86 Additionally, the template matching provides the orientational map of the crystallites in the sample (Fig. 10B). In order to facilitate this “template matching,” the ED patterns were recorded with a precessed electron beam that diminishes distortions of the measured ED intensities caused by dynamical diffraction and multiple scattering bringing them closer to those from kinematical approximation (the details on precession electron diffraction (PED) technique can be found in Refs. 87–89). Electron diffraction is a primary tool to detect the presence of planar defects (twins and stacking faults, see Section 7.10.2 for a more detailed description) and understand their nature as each type of defect demonstrates characteristic fingerprints in the ED patterns. Twin planes separate parts of the crystal which differ in the orientation of the crystal axes; they manifest themselves by reflection splitting in the ED patterns. The signature of mirror twins is a presence of a reciprocal row of unsplit reflections along the direction perpendicular to the twin plane. The reflections in other reciprocal lattice rows are split, and the splitting magnitude increases proportionally to the distance from the unsplit row.90 Mirror-twinned domains often appear in the Li1 þ xM1-xO2 (M – transition metals) Li-ion battery cathode materials based on the “cubic” close-packed structure, where the twin planes are confined to the {111} crystal planes of the underlying face-centered cubic sublattice (Fig. 11).91 Stacking faults, antiphase boundaries, shear planes contain a translational component implying a displacement of one part of the crystal with respect to another, and, if such defects are abundant and randomly spaced, diffuse intensity lines appear in the ED patterns. These lines are orthogonal to the fault planes and pass through reciprocal lattice nodes g for which the dot product g.R s integer (R is the displacement vector at the fault plane) that gives a possibility to calculate the components of the vector R using the g.R ¼ integer for the unstreaked diffraction spots.92 This is also a very typical situation in the Li-rich Li1 þ xM1  xO2 layered oxides. The Li and M cations in the (Li1  xM2 þ x)  basic structure layers are ordered forming a “honeycomb” pattern resulting in a monoclinic C2/m superstructure on top of the R3m characteristic of the “cubic” close-packing of the oxygen layers. The adjacent “honeycomb” layers can be shifted with respect to each  structure at little energy cost giving rise to abundant stacking faults, other by symmetry-related lattice vectors of the parent R3m which produce pronounced diffuse intensity lines in the ED patterns (Fig. 12). There is an indication, exemplified with the Li2MnO3 (¼ Li(Li1/3Mn2/3)O2) material as the end-member of the Li1 þ xM1  xO2 series, that these stacking faults affect electrochemical capacity and capacity retention when the material is tested as a cathode for Li-ion batteries.93 For crystal structure analysis, electron diffraction has been used for decades as a method auxiliary to powder crystallography.94–96 Because of clear overview of the reciprocal space obtained from ED patterns taken in a single crystal manner from polycrystalline

Fig. 10 Phase map (A) and corresponding orientation map (B) in the LiFePO4/FePO4 two-phase sample with 75% delithiation degree. Reprinted with permission from Brunetti, G.; Robert, D.; Bayle-Guillemaud, P.; Rouviere, J. L.; Rauch, E. F.; Martin, J. F.; Colin, J. F.; Bertin, F.; Cayron, C. Chem. Mater. 2011, 23, 4515–4524. Copyright 2011 American Chemical Society.

Electrode materials viewed with transmission electron microscopy

287

 subcell and showing reflection splitting in the h0l, h s 3n reciprocal Fig. 11 The [010] ED pattern of Li1.2Ni0.13Mn0.54Co0.13O2 indexed in the R 3m lattice rows caused by mirror twinning of the cubic close-packed O3 R 3m structure with the (001) twin plane (A). The primitive reciprocal unit cells of two twinned variants are outlined in red and green. Twinned domains with few tens of nanometers in thickness along the c-axis are clearly visible in low magnification HAADF-STEM image (B). Reprinted from Yin, W.; Grimaud, A.; Rousse, G.; Abakumov, A.M.; Senyshyn, A.; Zhang, L.; Trabesinger, S.; Iadecola, A.; Foix, D.; Giaume, D.; Tarascon, J.-M. Nat. Commun. 2020, 11, 1252.

Fig. 12 The [100] ED pattern of Li1.2Ni0.13Mn0.54Co0.13O2 indexed in the monoclinic C2/m supercell (A). The diffuse intensity lines pass along the c* reciprocal direction leaving the unstreaked reflections 0kl, k ¼ 6n resulting in a displacement vector R ¼ [u 1/6 0] for the fault planes (the u component cannot be determined from this pattern). HAADF-STEM image demonstrates random shifts of the “honeycomb” layers (marked with the white zig-zag line) over 1/6 of the repeat period of the C2/m supercell (B).

materials, they are of indispensable help for indexing powder diffraction patterns and finding correct unit cell, determine the crystal system, lattice centering and limit the selection of possible space groups from the analysis of reflection conditions (Fig. 13). Nevertheless, the crystal structure solution (in case of hitherto unknown structures) and Rietveld refinement (in case of known good initial structure model) were performed using powder diffraction intensities due to their essentially kinematical nature. This approach has extensively been applied to crystallographic investigations of many pristine polyanion metal-ion battery cathodes helping to tackle their structure complexity (AMXO4Y (A ¼ Li, K, Rb, M ¼ Fe, V, Ti, X ¼ P, S, Y ¼ OH, F,97–101 Fe2O(SO4)2,102 Na2Co(SO4)2,103 LiNaCoPO4F,104 Li1  xFe1 þ x(PO4)1  y(OH)4y,105 Li2FeSiO4106,107). However, this scheme becomes particularly challenging being applied to the metal-ion battery electrode materials extracted from electrochemical cells at different states of charge. For casting the electrode layer, the electroactive material is usually mixed with amorphous additives, such as carbon black to improve electric conductivity and polymeric binder to assure the electrode integrity and adhesion to the metallic foil of the current collector. Scattering contribution from these additives complicates the background shape of powder diffraction patterns introducing uncertainties in background fitting with conventional polynomial functions, particularly in the regions with substantial reflection overlap. If the diffraction pattern is taken from the electrode deposited on a current collector (typically Al or Cu foil), the reflections from the latter dominate the pattern, being at the same time hard to account for because of strong non-uniaxial texture in these rolled foils. The electroactive material by itself could also consist of a mixture of phases with similar structures and close unit cell parameters that is very common for the compounds operating through a two-phase (de)intercalation mechanism. Intrinsic limitations of powder diffraction also stem from extensive reflection overlap (particularly severe for the structures with low symmetry and/or large unit cell volume, typical for the polyanion compounds), limited reciprocal space coverage and low signal-to-noise ratio (intrinsic in the lab X-ray diffractometers with conventional X-ray tubes operating with Co Ka or Cu Ka radiation with relatively long wavelengths). Quantitative ED offers an alternative possibility for structure analysis of metal-ion battery electrodes through collecting and operating with single crystal diffraction data from submicron-sized crystallites that are hardly accessible for other diffraction

288

Electrode materials viewed with transmission electron microscopy

Fig. 13 Crystal structure analysis flowchart illustrating the synergistic effect of powder diffraction and electron diffraction. Transmission electron microscope allows operating crystallites constituting the polycrystalline sample as separate single crystals obtaining series of the ED patterns representing reciprocal lattice sections and providing important input for further crystal structure solution and refinement from powder diffraction data.

methods. On top of eliminating many abovementioned disadvantages of powder diffraction, ED offers better sensitivity to “light” atoms compared to X-rays that is important for resolving the structure of intercalation-type electrodes in Li-ion batteries, in which Liþ cations can switch crystallographic positions depending on the state of charge and/or occupy several crystallographic sites with fractional occupancies. Relative preference of ED over X-rays in detectability of “light” atoms can be estimated as [Zheavy/Zlight]1/2, where Zheavy and Zlight are the atomic numbers of “heaviest” and “lightest” elements in the structure, respectively.108 Thus, in LiCoO2, a typical layered oxide cathode for Li-ion batteries, the Li cations (ZLi ¼ 3, ZCo ¼ 27) are three times more visible with electrons than with X-rays. Employing neutron powder diffraction in many cases improves visibility of the Li cations because of the absence of systematic dependence of coherent scattering lengths on atomic numbers, but usually only samples of few tens of mg are available from laboratory electrochemical cells that is not enough for high quality neutron diffraction experiments. As it was already discussed in this section and in the Section 7.10.2, the ED intensities are significantly distorted by multiplebeam dynamical scattering which is thickness-dependent and more severe for the materials containing elements with high Z. Chemists who design novel intercalation materials to be used as cathodes or anodes in metal-ion batteries always keep in mind the formula for calculating the theoretical capacity C ¼ 26.8n/M, where C is the electrochemical capacity in Ah/g, n is the number of electrons transferred in the redox reaction or the number of (de)intercalated alkali cations per formula unit and M is the molecular weight of the formula unit. A quest for better electroactive materials aims at maximum possible capacity as one of the most important goals, naturally prohibiting “heavy” high-Z elements in their chemical compositions. As a positive consequence of this chemical and practical reasoning, the metal-ion battery electrode materials generally should demonstrate a reduced probability of multiple scattering events. Additionally, multiple scattering is greatly diminished by the data collection routine employing PED and/or EDT techniques.109 The obtained quasi-kinematical intensities can be used for ab initio structure solution with the direct methods or charge flipping algorithm.110,111 In KMnCO3F, consisting from the elements with relatively low Z and with a moderately small difference in Z between “heavy” and “light” atoms (theoretical density r ¼ 3.01 g .cm 3, wED  0.34, wED ¼ [Zlight/Zheavy]0.75 is the estimate of detectability of atoms,108) the crystal structure solution with charge flipping from EDT data revealed a complete structure and correct space group, which then was confirmed with joint Rietveld refinement from synchrotron X-ray and neutron powder diffraction data.112 For Li2CoPO4F (theoretical density r ¼ 3.42 g cm 3, wED  0.19), the direct methods applied to PED data revealed the positions of all atoms except those of the Li cations, which were subsequently located using difference Fourier maps.113 In the half-charged Li0.5RhO2 with the rutile-ramsdellite intergrowth structure (theoretical density r ¼ 6.18 g cm 3, wED  0.13) the direct structure solution provided only the Rh positions, whereas the O atoms were found from difference Fourier maps, which, however, did not allow locating the Li cations due to large residual peaks of scattering density.114 The Li positions were found using direct space structure optimization with Monte Carlo-based methods employing bond valence sum, antibump

Electrode materials viewed with transmission electron microscopy

289

conditions (prevention of short interatomic distances) and quality of the EDT intensity fit as the cost functions.115 The comparison of these three cases illustrates that the incompleteness of the structure solution and further efforts required to complete the structure model clearly correlate with the tendency toward the dynamical diffraction (reflected here through the material’s density indicating the total scattering power) and the ratio between atomic numbers of the “light” and “heavy” atoms. It is a fully legitimate question on how reliable is the crystallographic information refined in a kinematic approximation (i.e. assuming Ihkl  |Fhkl |2, where Fhkl is the structure factor of the reflection hkl and Ihkl – its intensity) from the PED or/and EDT intensities? A comparison of atomic coordinates for LiFePO4 after the Rietveld refinement from laboratory powder X-ray diffraction data and EDT data (Fig. 14) provided in Table 2 demonstrates that the refined structures are very close to each other. Bond valence sums for the cations and anions calculated using interatomic distances from the EDT refinement do not differ more than by 10% with the nominal oxidation states that confirm that the obtained structure is chemically sensible. The reliability factor RF for the EDT refinement of  0.18 is, however, unusually high compared to that is generally accepted in X-ray diffraction refinements, and the observed |Fobs | and calculated | Fcalc | structure amplitudes noticeably scatter around the | Fobs | ¼ |Fcalc | straight line (Fig. 14). This indicates deviations from kinematical approximation caused by the residual dynamical diffraction effects. Ex situ EDT data on metal-ion battery cathodes at different states of charge were used so far for analyzing different structural aspects of alkali cation (de)intercalation. In the olivine-structured LiFe0.5Mn0.5PO4 cathode the occupancies of the Li sites were

Fig. 14 3D reconstruction of reciprocal space from the EDT data on LiFePO4 (A). The missing wedge due to the limitation of the rotation angle of the TEM holder is clearly visible. |Fobs |  |Fcalc | plot after the refinement of the LiFePO4 crystal structure in a kinematical approximation (B).

290 Table 2

Electrode materials viewed with transmission electron microscopy Crystallographic parameters of LiFePO4 as refined from laboratory powder X-ray diffraction (XRD) data and electron diffraction tomography (EDT) data (space group Pnma, a ¼ 10.3181(1)Å, b ¼ 6.00132(6)Å, c ¼ 4.68908(5)Å). Powder XRD

EDT

Atom

x/a

y/b

z/c

x/a

y/b

z/c

BVS

Li Fe P O1 O2 O3

0 0.2823(3) 0.0966(6) 0.092(1) 0.455(2) 0.164(1)

0 1/4 1/4 1/4 1/4 0.051(2)

0 0.974(1) 0.422(1) 0.747(3) 0.203(2) 0.285(2)

0 0.2809(6) 0.094(1) 0.094(2) 0.458(1) 0.165(1)

0 1/4 1/4 1/4 1/4 0.047(2)

0 0.972(2) 0.408(3) 0.726(5) 0.187(5) 0.296(4)

0.94(1) 2.00(4) 5.6(2) 2.0(1) 2.21(8) 2.15(5)

Bond valence sums (BVS) are calculated from the EDT atomic coordinates.

refined being in full agreement with the experimentally observed electrochemical capacity and the analysis of polyhedral distortion and the (Fe, Mn) – O bond lengths variation due to Jahn-Teller distortion intrinsic in the Mn3þ cations.116 New distorted octahedral Li sites arise in KVPO4 with the KTiOPO4-type structure after de-potassiation and subsequent lithiation in the Li-ion cell that was unequivocally demonstrated with the EDT structure solution and refinement.117 EDT revealed strong Li/Fe antisite disorder in Li2FePO4F obtained by electrochemical Na for Li replacement through the analysis of residual scattering density peaks in difference Fourier maps and subsequent refinement of the occupancy factors for the Li and Fe sites.118 Finally, it has been demonstrated from the refinement of the complex tunnel (Nax⎕1-x)5[MnO2]13 structure that the precision in atomic positions obtained from EDT data applying full dynamical scattering theory competes with that from the Rietveld refinement based on synchrotron powder X-ray diffraction data.119 These examples, also complemented with numerous success stories from other fields of crystallography, solid state chemistry and material science, led to a significant revision of the role of electron diffraction in structure analysis of polycrystalline materials making it a fully legitimate partner of powder diffraction techniques (Fig. 15). Being a low-dose technique, electron diffraction provides a unique opportunity for local structure characterization of battery electrodes in situ inside the electrochemical cell filled with liquid electrolyte. Conventional organic carbonate electrolytes readily decompose under intense electron beam120 required for atomic resolution TEM imaging causing carbon contamination which grows quickly and deteriorates the image. Moreover, electron scattering in the electrolyte layer spreads the electron beam and delocalizes the image. However, it has been demonstrated that the EDT data can be collected inside the 500 nm thick electrochemical

Fig. 15 Crystal structure analysis flowchart illustrating complementarity of powder diffraction and electron diffraction tomography. Besides basic crystallographic parameters retrieved from the 3D reconstruction of the reciprocal lattice, the initial structure model can be determined and used as input for the Rietveld refinement from powder diffraction data or finalized with the single crystal refinement rigorously taking into account dynamical diffraction.

Electrode materials viewed with transmission electron microscopy

291

cell with the electron-transparent Si3N4 windows.121 In spite of a limited angular range of  30 available with such an electrochemical holder, the collected data allowed for reliable structure refinement of the pristine and fully charged LiFePO4 (Fig. 16). Finally, it is worth mentioning that electron diffraction can reveal the structure of not only crystalline, but partially disordered and amorphous materials from the total scattering approach with the analysis of pair distribution function (PDF) delivering information on local interatomic distances and coordination numbers even if the material does not possess periodicity. PDF is calculated from a continuous scattering curve (instead of point Bragg reflections) usually measured with short X-ray or neutron wavelengths to ensure scattering at high momentum transfer, but it also can be measured with electrons (ePDF) assuming a low probability of multiple scattering.122 This approach has been employed for distinguishing the structure of amorphous delithiated FePO4 by comparing the experimentally measured ePDF curve with those calculated from possible trigonal and orthorhombic FePO4 structures revealing that the local atomic arrangement in the amorphous phase is closer to the latter.123 To summarize, the ED and its advanced varieties provide decent data for proper structure solution and various defects detection. Still, the information on the exact location of crystal structure features, such as point and extended defects and their evolution upon the electrochemical reactions are missing. Therefore, the comprehensive investigation requires the usage of other TEM techniques, which offer the opportunity to investigate crystal structure locally, so, the following section is dedicated to the detailed description of these techniques applied to the analysis of the nature, location and evolution of various crystal structure defects.

7.10.5

Imaging of the local crystal and defect structure

The degree to what the crystal and defect structure of the electrode material determines its electrochemical properties heavily depends on the application. If in catalysis the surface and subsurface structure are of primary importance, the intercalation-type electrodes rely on long-range cation diffusion (such as Liþ, Naþ or Kþ diffusion in the electrodes for metal-ion batteries) reckoning on availability of the cation migration pathways with low energy barriers. As these materials are typically polycrystalline, both intragrain diffusion through the interior of the grains and intergrain diffusion through the grain boundaries need to be considered. These issues are equally important for solid electrolytes, with, however, particular significance of the atomic structure of the electrode/electrolyte interfaces often causing high impedance of the electrochemical cell. Thus, the presence, spatial distribution, concentration and atomic structure of point and/or extended defects and grain boundaries, which can block the ionic and electronic transport, have a huge impact on the rate capability of the metal-ion batteries. In that light, TEM techniques such as atomic-resolution

Fig. 16 EDT experiment in a TEM cell with liquid electrolyte and Si3N4 windows: scheme of the cell and data collection procedure, reconstruction of 3D reciprocal space and difference Fourier map demonstrating the Li positions in the LiFePO4 structure. Reprinted with permission from Karakulina, O. M.; Demortiere, A.; Dachraoui, W.; Abakumov, A. M.; Hadermann, J. Nano Lett. 2018, 18, 6286–6291. Copyright 2018 American Chemical Society.

292

Electrode materials viewed with transmission electron microscopy

HAADF-, ABF-, iDPC-STEM and electron holography imaging, providing direct visualization of appearance and segregation of defects in the area of interest, are very powerful tools for crystal structure analysis in terms of all types of defects. When applied to the metal-ion battery materials, they enable not only the detailed analysis of relevant structural features in the pristine electrodes, but can also be conducted post-mortem delivering the information on the evolution of the crystal and defect structure depending on a number of charge/discharge cycles. Tracking this evolution can deliver valuable information on the peculiarities of charge/ discharge mechanism, voltage hysteresis, capacity and voltage fade. Capacity fade might stem not only from the crystal structure degradation, but also from mechanical deterioration of the electrodes, hence the importance of understanding the local strain development and relaxation during electrochemical cycling. Shearing of the atomic layers causing stacking faults in course of electrochemically-induced phase transitions, microcracks initiated at the extended defects and twin boundaries, formation of micropores are the subject of interest allowing material scientists to correlate the chemical composition, crystal structure, and mechanical behavior of the intercalation-type electrodes. A possibility to extract quantitative information on the occupancies of the atomic sites and atomic displacements from the atomic-resolution TEM images is an extra asset for establishing these correlations.

7.10.5.1

Point defects and order-disorder in electrode materials

Li-ion battery cathodes possessing one-dimensional channels for the Liþ cation migration have been scrutinized in detail for the role of point defects which might block these channels. In LiFePO4 with the olivine-type structure the cationic exchange between the Liþ and Fe2þ positions (the FeFe þ LiLi / FeLi0 þ LiFe• defect pairs denoted with the Kröger-Vink notations, also called anti-site defects) results in the most energetically favorable type of point defect.81 These defects are the major obstacle for the Liþ migration within the LiFePO4 structure because the [010] diffusion channels are effectively blocked by Fe2þ cations appearing in these channels due to anti-site disorder.124 The Fe atoms in the Li positions can be discerned using the HAADF-STEM imaging due to sensitivity of this imaging technique to atomic number Z and relatively large difference between the atomic numbers of Li (Z ¼ 3) and Fe (Z ¼ 26).125,126 However, the minimum concentration at which the Li4 Fe exchange can be detected from the HAADF-STEM images with the naked eye is limited to > 8%–15% of Fe in the Li position (Fig. 17) that significantly exceeds the concentration of just few percent, which is sufficient to block the diffusion channels ( 4% of antisite defects rise the activation energy of Li migration by  50%).124 Plotting the HAADF intensity profiles across the atomic columns allows for a more precise quantitative estimate of the anti-site disorder. The ratio R between integral intensities measured at the Li and transition metal (M) atomic columns in the [010] HAADF-STEM images of the Pnma LiMPO4 olivine-type structure is related to the degree of anti-site disorder x as following: R¼

4½ð1  xÞZLi þ xZM 2 ½xZLi þ ð1  xÞZM 2

(12)

where ZLi and ZM are the atomic numbers of Li and transition metal, respectively.127 An additional concern arises from attributing the observation of extra HAADF intensity at the Li positions to the particular type of defects (i.e. anti-site defects), whereas the actual defect chemistry might be more diverse and sophisticated. For example, extra iron can appear at the cost of Li vacancies in order to maintain the charge balance in a form of the FeLi0 þ VLi• defect

Fig. 17 [010] HAADF-STEM images of LiFePO4 with 8% (A) and 2% (B) of the Li/Fe antisite defects. Note weak intensity at the Li positions in A and a virtual absence of this intensity in B (the crystal structure is overlaid with the Li positions marked in green). The computer-simulated HAADFSTEM images are calculated with no anti-site defects (C) and 10% of the Li/Fe antisite defects (D). Reprinted with permission from Paolella, A.; Bertoni, G.; Hovington, P.; Feng, Z.; Flacau, R.; Prato, M.; Colombo, V.; Marras, S.; Manna, L.; Turner, S.; Van Tendeloo, G.; Guerfi, A.; Demopoulos, G.P.; Zaghib, K. Nano Energy 2015, 16, 256–267. Copyright 2015 Elsevier.

Electrode materials viewed with transmission electron microscopy

293

pair.128,129 In the hydrothermally prepared Li1  xFe1 þ x(PO4)1  y(OH)4y solid solutions Fe in the Li positions appears to compensate for extra negative charge caused by the PO43 / 4OH- substitution.105,130 This defect chemistry is drastically different from anti-site disorder, but the HAADF-STEM images of the Li1  xFe1 þ x(PO4)1  y(OH)4y solid solutions look virtually identical to the images of the Li/Fe-disordered LiFePO4. Anti-site disorder in LiFePO4 is favored by the striking similarity of the octahedral oxygen coordination of the Liþ and Fe2þ cations with the average LieO and FeeO distances of 2.150 and 2.156 Å, respectively, as the ionic radii of Liþ (0.76 Å) and Fe2þ (0.78 Å) are very close to each other. A similar situation is also realized in the layered oxide cathodes for Li-ion batteries LiNixCoyMnzO2 (NMC) with high Ni content (x  0.6, x þ y þ z ¼ 1), in which the proximity of the ionic radii of Liþ (0.76 Å) and Ni2þ (0.69 Å) is also a sufficient prerequisite for cation mixing. This disorder is present in the pristine materials and its degree depends on the Ni content and the synthesis conditions (temperature and partial oxygen pressure) defining the Ni2þ/Ni3þ ratio in the obtained compounds. The Li/Ni mixing is not only the anti-site disorder but a result of more complex defect chemistry involving changing the chemical composition along the LiMO2 / MO line (M - transition metal). In contrast to LiFePO4, NMCs possess a structure built of the layers of edge-sharing (Ni, Mn, Co)O6 octahedra with the Li cations in between, so that the Liþ diffusion system in NMCs is essentially two-dimensional. Nevertheless, the Li/M disorder also negatively impacts the electrochemical capacity and rate capability of NMCs that justifies extended investigation of this disorder with the bulk diffraction methods and locally with TEM.131–134 In contrast to X-ray and neutron powder diffraction, HAADF-STEM imaging provides the direct view on defects’ arrangement and location, distinguishing the contributions of the interior of the crystallites and their surface to the whole defect concentration (Fig. 18A–C). Semi-quantitative map of the cation disorder can be extracted from the HAADF-STEM images with the help of model-based parameter estimation.46,135,136 In this approach, the intensity distribution at the position of atomic column is modeled with a Gaussian function and the expectation of the intensity of a pixel (k, l) at the position (xk, yl) in the image can then be represented as a function fkl(q) of a set q of unknown structure parameters: 1 0  2  y 2 x Mi I X x  b þ y  b X m k l i mi C B (13) Ymi exp@  fkl ðqÞ ¼ BG þ A 2 2w i m i¼1 i

Fig. 18 [010] HAADF-STEM images of LiNi0.8Mn0.1Co0.1O2 as-synthesized cathode material taken from the interior part (A) and the surface (C) of a crystallite shown in the low magnification HAADF-STEM image (B). Note the remarkable difference in intensity associated with the Li positions indicating a significant population of this position by M cations at the surface compared to significantly lower Li/M disorder in the bulk. The [010] HAADF-STEM image of the same material (D), the color-coded distribution of the M atoms among the cationic columns (E) and the corresponding map of the Li/M disordering level throughout the image, calculated with the model-based parameter estimation method.

294

Electrode materials viewed with transmission electron microscopy

where BG is the constant background intensity, wi is a width of the Gaussian intensity distribution for the i-th type of atomic column, Ymi is the intensity of the Gaussian peak associated with the atomic column mi, bmi x and bmi y are the x and y coordinates of the atomic column mi, the index i spans from 1 to I encountering the types of atomic columns (i.e. consisting of different atomic entities or/and with different density of atoms), the index mi encounters the number of columns of the type i. The parameters BG, Y, b and w are then defined using least squares minimization of the discrepancies between the experimental and model-based image intensities. Proper quantification of scattering cross-section allows quantification of a number of atoms in projected atomic columns of Li and M and calculation of the Li/M relative disorder map (Fig. 18D–F). The majority of existing calculation approaches lie down on the direct implementation of the least squares estimator, where all the unknown parameters of the parametric model are calculated at the same time in an iterative manner, comparing this model to the experimental data at each iteration using the uniformly weighted least squares criterion.137 The uniformly weighted least squares estimates b q of the unknown parameters q are then given by the values of t that minimize the uniformly weighted least squares criterion: b q ¼ argmin t

K X L X

ðwkl  fkl ðt ÞÞ2

(14)

k¼1 l¼1

where wkl stands for the observed pixel value in an image. However, this calculation technique is computationally very timeconsuming and feasible only for a limited field of view. The alternative algorithm is to perform segmentation of the image into smaller sections containing individual atomic columns without ignoring overlap between neighboring columns, so, the image will possess superposition of the Gaussian peaks, describing individual atomic columns.46 In that way, unknown parameters of a single atomic column are evaluated at the same time, which significantly reduces calculation time for each iteration. Furthermore, the observation of the cation disordering can be logically extended to the electrode structures at different states of charge. Atomic-resolution imaging of Li-ion battery cathodes can be performed in situ in an all-solid-state battery configuration (see Section 7.10.7) as the electrochemical TEM cell with liquid electrolyte severely limits the spatial resolution. However, the vast majority of investigations are performed post mortem on the electrodes extracted from the electrochemical cells and carefully washed from the electrolyte components. Handling the charged cathode materials or discharged anode materials must be conducted with extreme care to prevent unwanted redox reactions by interacting with oxygen and moisture of the ambient air. After opening the cell in an Ar-filled glove box (Fig. 19A) and washing the electrodes with the electrolyte solvents, the electrode material is scraped

Fig. 19 Handling the electrode materials for TEM specimen preparation: the operations are conducted in an Ar-filled glove box (A), the material is grinded and dispersed in inert liquid (B) and dropped onto a TEM grid, which then is transported to the TEM in a tightly closed holder (C).

Electrode materials viewed with transmission electron microscopy

295

off the current collector and grinded in a mortar in a liquid media creating dispersion (Fig. 19B). The electrochemical stability window of the liquid for such dispersion should match the voltage used for charge/discharge of the electrode (using the same solvent as for the electrolyte is highly recommended). Few drops of the suspension are placed on a carbon-coated Cu TEM grid which is transported into the microscope column under Ar using a dedicated vacuum-tight transfer holder. Typically, such holders possess a tip with an O-ring moving in and out of the holder’s body (Fig. 19C). The holder is sealed in the glove box and opened in the airlock chamber of the TEM during evacuation. HAADF- and ABF-STEM imaging of the charged and discharged layered oxide cathodes revealed that the transition metal cation sublattice does not stay intact during Liþ extraction/insertion, but undergoes dynamical order-disorder transformations with the degree of reversibility depending on the material’s nature. HAADF-STEM images of the Li2Ru0.75Ti0.25O3 electrode charged to 4.6 V vs. Liþ/Li revealed massive migration of the transition metal cations to the Li sites causing strong cation disorder (Fig. 20A and B). This migration appears to be highly reversible as the initial layered structure is almost fully restored upon subsequent discharge to 2.0 V (Fig. 20C).58 Very local changes in the cation ordering can be spotted with TEM at a very early stage of the electrochemically-induced degradation. ABF-STEM image of the Li1.2Ni0.13Mn0.54Co0.13O2 cathode material demonstrates a variety of local defects formed at the surface and subsurface region after just two charge/discharge cycles within the 2.0–4.8 V vs Liþ/Li potential range (Fig. 21A).63 At the  1 nm thick surface layer the layered ordering of Li and M cation is lost and isotropic rock-salt-like structure is formed. In the subsurface region a migration of cations into tetrahedral interstices of underlying “cubic” close packing of the oxygen atoms is observed, along with the cation exchange between Li and M crystallographic positions. Note that in contrast to the HAADF-STEM technique, ABF-STEM imaging clearly visualizes not only “heavier” M cations, but also the positions of “lighter” Li and O atoms that in turn allows tracing oxygen coordination of the cationic species even at the sites of point defects. After 100 cycles the concentration of the cationic defects increases and they form locally-ordered domains of the spinel-type structure (Fig. 21B and C). The sensitivity of the ABF-STEM imaging to “light” elements permits tracing changes in the oxygen sublattice that is very instructive for establishing true Liþ (de)intercalation mechanism in the Li-ion battery cathodes. In the charged state of Li2IrO3 layered cathode the ABF-STEM images revealed cooperative deformation of the IrO6 octahedra with a formation of the set of short (the projected distance of 1.56(1)Å) and long (the projected distance of 1.83(1)Å) O-O separations (Fig. 22A and B).138 The O-O projected distances have been measured straight from the ABF-STEM images by plotting the inverted ABF intensity profiles and quantifying the atomic positions through fitting the profiles assuming Gaussian intensity distribution at the atomic columns (Fig. 6C). The determined projected distances were found to agree well with the results of the Rietveld refinement from neutron powder diffraction data and DFT-based calculations. Shortening of the O-O separations reflects covalent bonding formed due to partial oxidation of the oxygen atoms when the Liþ cations are extracted at high potentials that delivers extra electrochemical capacity to this class of cathode materials based on so-called “anionic” redox reactions.139,140 Although practical demands require the anionic redox to be fully reversible, the concurring full oxygen oxidation with releasing O2 often cannot be completely avoided. ABF-STEM imaging picks up oxygen vacancies formed at the early stage of the irreversible oxygen oxidation and O2 release from the Li4.27Fe0.56TeO6 cathode that correlates well with the differential electrochemical mass spectroscopy (DEMS)

Fig. 20 HAADF-STEM images of pristine Li2Ru0.75Ti0.25O3 (A), charged to 4.6 V vs Li/Liþ (B) and discharged to 2.0 V (C). The charged state is highly disordered with the transition metal cations migrating to the octahedral sites in the Li layer and also to the Li positions in the “honeycomb”ordered LiM2 layers (the insert in (B) shows the HAADF intensity profiles along the LiM2 layers in the pristine ordered state (top) and in the charged disordered state (bottom)). After discharge the layered ordered structure is restored. Reprinted from Sathiya, M.; Abakumov, A.M.; Foix, D.; Rousse, G.; Ramesha, K.; Saubanère, M.; Doublet, M.L.; Vezin, H.; Laisa, C. P.; Prakash, A. S.; Gonbeau, D.; Van Tendeloo, G.; Tarascon, J.-M. Nat. Mater. 2015, 14, 230–238.

296

Electrode materials viewed with transmission electron microscopy

Fig. 21 [010] ABF-STEM image of Li1.2Ni0.13Mn0.54Co0.13O2 after two charge/discharge cycles (A). Dashed line marks the near-surface layer with pronounced Li/M disorder resembling the rock-salt (RS) structure. The region 1 shows point defect where the cations are located at the tetrahedral voids (outlined in the corresponding enlargement) instead of their native octahedral positions. The area 2 demonstrates a migration of the M cations toward the adjacent octahedral position in the Li1  xM2 þ x layers. [010] HAADF-STEM images of the pristine (B) and 100 times cycled (C) Li1.2Ni0.13Mn0.54Co0.13O2 demonstrating increasing cation disorder and local transformation from the layered to spinel-type structure. Reprinted with permission from Pimenta, V.; Sathiya, M.; Batuk, D.; Abakumov, A.M.; Giaume, D.; Cassaignon, S.; Larcher, D.; Tarascon, J.-M. Chem. Mater. 2017, 29, 9923–9936. Copyright 2017 American Chemical Society.

data showing oxygen eruption at the first charge and CO2 release on subsequent cycles due to catalytic electrolyte oxidation (Fig. 23).141 Despite all advantages of the HAADF-STEM and ABF-STEM techniques in direct atomic structure imaging, they require a high electron dose rate to maintain an appropriate signal-to-noise ratio, which increases the risks of undesirable damages and/or artifacts (see Section 7.10.3). A high signal-to-noise ratio is becoming particularly important if quantification of interatomic distances is the target, as the precision of such quantification depends on the number of detected electrons (see Eqs. 6 and 7 in Section 7.10.2). Electron ptychography has been demonstrated to be a viable low-dose STEM technique capable of visualization of lithium, oxygen, and transition metal sublattices in the Li1.2Mn0.6Ni0.2O2 cathode material at subpicoampere electron currents.142 Moreover, ptychography also allows for postprocessing correction for aberration-induced image distortions that reduces the time necessary for the alignment at the stage of image collection thus further contributing to lowering the total electron dose. Another cutting-edge STEM technique is differential phase contrast (DPC) imaging, which possesses the visualization possibilities of ABF-STEM while requiring a few orders of magnitude lower electron dose for providing a comparable signal-to-noise ratio. Fig. 24 represents experimental dDPC-STEM projected charge density maps for LiCoO2 at 5% and 60% delithiation states. The high sensitivity of dDPCSTEM technique makes the formation of Li vacancies clearly visible even at low concentrations for Li0.95CoO2. As for Li0.4CoO2, the

Electrode materials viewed with transmission electron microscopy

297

Fig. 22 [001] HAADF-STEM and ABF-STEM images of Li2IrO3 charged to 4.6 V vs Li/Liþ (A). Enlarged ABF-STEM image where O-O pairs with short projected distances are marked with dumbbells (B). Inverted ABF intensity profiles along the O-O pairs, which demonstrate long (blue) and short (red) projected distances (C). Reprinted with permission from McCalla, E.; Abakumov, A.M.; Saubanere, M.; Foix, D.; Berg, E.J.; Rousse, G.; Doublet, M.-L.; Gonbeau, D.; Novák, P.; Van Tendeloo, G.; Dominko, R.; Tarascon, J.-M. Science 2015, 350, 1516–1521. Copyright 2015 American Association for the Advancement of Science.

ordering of Li vacancies becomes apparent. In addition, delithiation is accompanied by Co cation migration from the octahedral positions to tetrahedral voids.140

7.10.5.2

Planar defects

Planar defects are intrinsic in many metal-ion battery electrode materials, particularly in those based on anionic close packings. A variety of planar defects in solids can be rationalized if the parent defect-free structure is split into modules whose structures are related to each other by a translation over a vector R (which is most typically a rational fraction of a repeat period of the parent structure) or/and by rotation over the angle f (Fig. 25). The former causes the parent structure to be split into translational domains separated by conservative (i.e. preserving the chemical composition) interfaces if R is parallel to the fault plane (antiphase boundaries, stacking faults) or non-conservative (i.e. changing the chemical composition) interfaces if R contains a component orthogonal to the fault plane (such as crystallographic shear planes). The latter gives rise to orientational domains, such as twins. Planar defects demonstrate their characteristic features in the reciprocal space, which can be easily picked up by electron diffraction, as briefly described in Section 7.10.4. Aberration-corrected STEM imaging provides further details on the atomic structure of defects. a-Li2IrO3 with a layered structure built of alternating Li3 and (LiIr2) “honeycomb”-ordered layers can gradually be transformed into its b-polymorph by increasing the number and ordering of antiphase boundaries (APBs) which are formed by a displacement of the layers over ½ of the interlayer distance (Fig. 26). These planar defects are conservative and do not change

298

Electrode materials viewed with transmission electron microscopy

Fig. 23 DEMS data for Li4.27Fe0.56TeO6 demonstrating O2 and CO2 release during charge (A). [010] ABF-STEM image after removing 1 lithium per formula unit (B). The black arrows indicate oxygen-deficient columns, while the white arrows indicate the M cations migrating to the lithium layer. Reprinted with permission from McCalla, E.; Prakash, A. S.; Berg, E.; Saubanere, V; Abakumov, A.M.; Foix, D.; Klobes, B.; Sougrati, M.-T.; Rousse, G.; Lepoivre, F.; Mariyappan, S.; Doublet, M.-L.; Gonbeau, D.; Novak, P.; Van Tendeloo, G.; Hermann, R.P.; Tarascon, J.-M., J. Electrochem. Soc. 2015, 162, A1341–A1351. Copyright 2015 The Electrochemical Society.

the chemical composition, but they significantly alter the electrochemical behavior of Li2IrO3 as a Li-ion battery cathode. If the a-form is prone to Ir cation migration to the octahedral sites in the Li layer on charge, such migration is suppressed in the b-polymorph. Additionally, pillaring the layers through APBs prevents their shearing behavior that remarkably improves the cycling stability of the b-polymorph.138,143,144 The layered AxMO2 oxide cathodes for metal-ion batteries (M - transition metal cation(s), A - alkali cations) demonstrate rich variety of structures originating from different stacking modes of the oxygen close-packed layers, hence the stacking faults are the intrinsic features in these materials. According to commonly accepted notations, these stacking sequences are designated with a letter denoting the coordination environment of the A cation (O - octahedral, P – trigonal prismatic), and a number reflecting the number of the octahedral MO2 layers per one repeat period along the stacking direction.145 For instance, the O3 notification represents the “cubic” ABCABC close-packed stacking sequence with the octahedral coordination of the alkali cations, whereas P2 stands for prismatic coordination with the ABBA stacking sequence. The stacking sequence significantly affects the electrochemical properties of the cathode materials, as it can be altered during battery operation resulting in phase transitions accompanied by the volume change, strain formation, and sluggish kinetics.146,147 Different stacking sequences can be distinguished by tracing the displacements of the transition metal cationic layers in the HAADF-STEM images. The pristine Li1.2Ni0.13Mn0.54Co0.13O2 material has the O3 stacking, which is partially preserved in the delithiated material at 4.8 V vs. Liþ/Li. The O3 stacking is recognized in the [010] HAADF-STEM image by the lateral shift of every next

Electrode materials viewed with transmission electron microscopy

299

Fig. 24 Artificially-colored dDPC-STEM maps of Li0.95CoO2 and Li0.4CoO2 demonstrating the projected charge density distributions. Enlarged part of the dDPC image for Li0.4CoO2 also demonstrates the Co atoms migrated from their octahedrally coordinated initial positions to tetrahedral interstices. The scale bar is 0.5 nm. Reprinted from Abakumov, A. M.; Fedotov, S. S.; Antipov, E. V.; Tarascon, J.-M. Nat. Commun. 2020, 11, 4976.

layer by 1/3 of the separation between the M-cation columns along the layer plane (Fig. 27A).91 However, slabs or even whole crystallites of the O1 stacking (“hexagonal” ABAB close packing) appear due to gliding the layers upon delithiation, which can be distinguished from the original O3 stacking by the absence of the lateral displacement of the layer (Fig. 27A and B). However, often the analysis of the layer shifts from the HAADF-STEM images cannot unequivocally identify the stacking sequence. For instance, both O1 and P2 stackings demonstrate no lateral shifts of the layers (Fig. 28), and can be distinguished indirectly by analyzing the distances between layers (Fig. 28). The AxO6 trigonal prisms in the P2 stacking are more extended along the stacking direction than the octahedra in the O1 stacking in order to accommodate larger A cations and maximize the distance between the oxygen atoms that causes larger interlayer separation. An obvious drawback of the HAADF-STEM imaging in this context is its insensitivity to the oxygen atoms. Since the oxygen positions are “seen” in the ABF-STEM images, the A-cation coordination by the oxygen atoms can be directly visualized and the stacking sequences as well as stacking faults can be discerned with more certainty (Fig. 29). Twinning is much less common in the layered AxMO2 metal-ion battery cathodes compared to stacking faults. Nevertheless, abundant twinning can also influence electrochemical properties of the intercalation-type electrodes. In monoclinic layered a-NaMnO2 the randomly occurring coherent twin planes twist the layers of the MnO6 octahedra transforming flat 2D Na-ion migration pathways into zig-zag shaped ones.148 Twisting of the cationic layers is easily recognized in the HR-TEM and HAADF-STEM images (Fig. 30). Increasing the abundance of the twin planes leads to various quasi-periodic modulated sequences, and, finally to the orthorhombic b-NaMnO2 in which the twin planes pass through each MnO6 octahedron. The twins increase the desodiation potential by destabilizing the desodiated structure through oxygen overbonding at the twin planes and also impede the Na

300

Electrode materials viewed with transmission electron microscopy

Fig. 25

Planar defects in crystalline solids.

Fig. 26 Series of HAADF-STEM images illustrating a transformation of the a-Li2IrO3 layered structure into its b-polymorph as a result of accumulation and ordering of antiphase boundaries (marked with arrows at the structure sketches). The IrO6 octahedra are shown, the Li cations are hidden for clarity.

Electrode materials viewed with transmission electron microscopy

301

Fig. 27 [010] HAADF-STEM images of Li1.2Ni0.13Mn0.54Co0.13O2 charged to 4.8 V vs. Li/Liþ representing the retaining O3 stacking with a thin domain of the O1 structure inserted (A), and the extended O1 domain (B). The O3 and O1 stackings can be identified by the presence and absence of the lateral displacements of the (Li1  xM2 þ x) layers followed by white straight lines. Reprinted from Yin, W.; Grimaud, A.; Rousse, G.; Abakumov, A.M.; Senyshyn, A.; Zhang, L.; Trabesinger, S.; Iadecola, A.; Foix, D.; Giaume, D.; Tarascon, J.-M., Nat. Commun. 2020, 11, 1252.

Fig. 28 [010] HAADF-STEM images of the O1 and P2 stacking sequences with different interlayer distances along with the corresponding crystal structures projections (not in the same scale as the images).

diffusion. Similar twins have also been observed with TEM in the LiCoO2 thin films resulting in  1000 times lower Li diffusion coefficient across the twin boundary compared to the diffusion along the twin boundary.149

7.10.5.3

Phase boundaries, grain boundaries and surfaces

The interfaces between the lithiated and delithiated phases arising in course of a two-phase (de)lithiation mechanism in the Li-ion battery electrode materials provide a significant contribution to the total free energy of the electrochemical system because substantial lattice mismatch between the lithiated and delithiated phases causes strong interfacial strain. Thus, the phase nucleation and interface propagation upon charge/discharge are contributing to the explanation of such phenomena as rate-depending competition between the two-phase and solid-solution mechanisms, changing the miscibility gap due to particle size and zero-current voltage hysteresis. HAADF-STEM images, taken with short dwell time to minimize the image distortions due to specimen drift and then aligned through cross-correlation to improve signal-to-noise ratio have been used to map the strain at the interface between the pristine LiFePO4 and chemically delithiated FePO4 phases.150 Fitting the HAADF intensity distribution with a set of Gaussian functions provided the coordinates of the Fe columns and resulted in a and b lattice dimensions maps across the coherent

302

Electrode materials viewed with transmission electron microscopy

Fig. 29 ABF-STEM images of the O3 (A), P3 (B) and P2 (C) stacking sequences in the AxMO2 layered oxides. The projections of the MO6 octahedra are traced in red, the yellow lines denote the coordination polyhedra of the A cations.

Fig. 30 HR-TEM image of a-NaMnO2 with inserted corresponding HAADF-STEM image (A). The white arrows mark the twin planes; the black arrows mark an arrangement of two subsequent twin planes which does not change the orientation of the layers but leads to their displacements thus resulting in out-of-phase boundary. HR-TEM image of b-NaMnO2 (B). White and black arrowheads indicate the local domains of the a- and b-phases, respectively. The model of the twin boundary (C), the MnO6 octahedra are shown in green, the Na cations are the brown spheres. Reprinted with permission from Abakumov, A. M.; Tsirlin, A. A.; Bakaimi, I.; Van Tendeloo, G.; Lappas, A. Chem. Mater. 2014, 26, 3306–3315. Copyright 2014 American Chemical Society.

(100) interface (Fig. 31). The strain component along the a-axis, which is perpendicular to the interface, relaxes sharply without violating the coherency, but the b-axis strain component (parallel to the interface) needs at least 30 nm for relaxation. Imaging the Li sites with ABF-STEM revealed that partially delithiated phase exists at the LiFePO4/FePO4 interface.151 The empty and occupied Li sites in this intermediate phase are ordered into separate layers leading to so-called Li “staging” which spreads over  15 nm forming curved boundaries with the LiFePO4 and FePO4 phases (Fig. 32). Similar ABF-STEM observations of the interface between the delithiated and lithiated phases have been performed for the Li4Ti5O12 Li-ion battery anode also demonstrating a two-phase discharge/charge mechanism.152 The atomic structure of coherent phase and grain boundaries formed at topotactic (de)lithiation is largely inherited from the structures of the adjacent domains of the lithiated and delithiated phases. More sophisticated incoherent boundaries are formed between the grains of the material with the same structure and chemical composition but with misorientation of the lattices. Such grain boundaries are common in the layered NMC-based cathode materials as they are produced as spherical of  5– 10 mm agglomerates consisting of thousands of submicron-sized grains with numerous grain boundaries. Such incoherent grain boundaries can demonstrate complex atomic reconstructions, for instance, forming a spinel-like arrangement,153 being the seeds for crack formation in the agglomerates in course of the charge/discharge cycling. The nature of the grain boundaries and electrode-electrolyte interfaces is of high importance for all-solid-state batteries (SSBs) as the primary bottleneck of SSB’s practical implementation is limited Liþ diffusion across the grain boundaries and high interfacial

Electrode materials viewed with transmission electron microscopy

303

Fig. 31 [001] HAADF-STEM image of the coherent interface between the LiFePO4 and chemically delithiated FePO4 phases (A). Normalized maps of the lattice dimensions along the a-axis (B) and b-axis (C) showing direction-dependent strain relaxation. Reprinted with permission Nakamura, A.; Furutsuki, S.; Nishimura, S.; Tohei, T.; Sato, Y.; Shibata, N.; Yamada, A.; Ikuhara, Y. Chem. Mater. 2014, 26, 6178–6184. Copyright 2014 American Chemical Society.

Fig. 32 [010] ABF-STEM images of the LiFePO4/FePO4 phase boundary with the appearance of the “staging” phase with the layered Li-vacancy ordering (A, marked by the dashed yellow lines). The enlarged images of the pristine LiFePO4, fully delithiated FePO4 and partially delithiated “staging” phases are provided in the panels B, C, and D, respectively. Reprinted with permission from Zhu, C.; Gu, L.; Suo, L.; Popovic, J.; Li, H.; Ikuhara, Y.; Maier, J. Adv. Func. Mater. 2014, 24, 312–318. Copyright 2014 WILEY-VCH Verlag GmbH & Co.

304

Electrode materials viewed with transmission electron microscopy

resistance between the cathode and solid electrolyte. HAADF-STEM imaging of incoherent grain boundaries in Li0.33La0.55TiO3 (LLTO) solid electrolyte revealed that the majority of them demonstrate significantly lower HAADF intensity that was related to strong La depletion assuming the vast dependence of the HAADF intensity on Z (Fig. 33).154 In combination with EELS, these grain boundaries were shown to consist mostly of Ti and O and have completely different crystal structure from the bulk LLTO being closer to a binary titanium oxide than to perovskite. The lack of La and Li near these grain boundaries makes them the Litransport barriers blocking the Li diffusion. As revealed with HAADF- and ABF-STEM imaging, the La-depleted Li2TiO3-structured defects can also intergrowth coherently with LLTO forming closed loops, embracing large LLTO domains and isolating them from ionic conduction.155 Revealing the complete atomic structure of such incoherent grain boundaries is a challenging problem because they typically provide very limited number of clear projections for TEM imaging, and the grains at both sides do not obey strict epitaxial relations as in the coherent interfaces. Atomic resolution electron tomography might be an appropriate tool to reveal the atomic structure of such grain boundaries in the future. The potential power of 3D atomic reconstruction has already been demonstrated on the grain boundaries between gold particles and could be further extended toward multiatomic structures.156,157 Highly disrupted structure at the surfaces and grain boundaries of nanoparticles is a subject of great interest for electrocatalysis because asymmetric and incomplete coordination environment of transition metal cations gives rise to the centers with enhanced catalytic activity. In contrast to the HR-TEM images suffering from delocalization at surfaces and interfaces (which, however, could be diminished with aberration correction), STEM techniques provide information on the very surface atomic structure revealing the most prominent crystal planes, terminating layers, and atomic terraces. Nanoparticles of cubic Mn2O3 prepared with high surface area to be used as electrocatalyst for the oxygen reduction reaction in alkaline media (ORR, O2 þ 2H2O þ 4e / 4OH) appear to be round-shaped, but close-up HAADF-STEM images demonstrate that their surface consists of short fragments of {111} crystal planes terminated by a step of one atomic layer, thus forming numerous terraces and surfaces of any arbitrary curvature (Fig. 34).158 HAADF- and ABF-STEM images of the surface-terminating grain boundaries in LaCoO3 and LaMnO3 revealed the transition metal sites that are exceptionally active in electrocatalysis of the oxygen evolution reaction (OER, 4OH / O2 þ 2H2O þ 4e) due to Jahn-Teller-induced displacement of the transition metal cations in asymmetric oxygen coordination environment.159 Finally, the aberration-corrected HR-TEM imaging of nanosized multimetallic electrocatalysts can be employed for quantification of strain at the surface and the interior of the particles through measuring the positions of the atomic columns by fitting the intensities with the Gaussian functions (Fig. 35). The strain has been demonstrated to be an effective measure to enhance the catalytic activity by adjusting the binding strength of intermediates through varying the metal-metal separations at the surface.160,161

7.10.5.4

Imaging in 3D

As outlined in the Section 7.10.2, the (S)TEM images carry 2D projected information of the real 3D object that might be a source of confusions which, however, can be at least partially resolved with electron tomography experiments enabling retrieving 3D information. Size, shape, porosity, orientation, and spatial distribution over carbon support of monometallic and bimetallic

Fig. 33 HAADF-STEM image of a boundary between two grains in Li0.33La0.55TiO3 (A) demonstrating the presence of two regions with different contrast, denoted as Type I and Type II, respectively. Green and red arrowheads designate the alternating La-poor and La-rich layers in the Li0.33La0.55TiO3 perovskite structure, respectively. Magnified HAADF-STEM images of Type I and Type II areas are shown in (B) and (C), respectively. Reprinted with permission from Ma, C.; Chen, K.; Liang, C.; Nan, C.-W.; Ishikawa, R.; Morea, K.; Chi, M. Energy Environ. Sci. 2014, 7, 1638–1642. Copyright 2014 The Royal Society of Chemistry.

Electrode materials viewed with transmission electron microscopy

305

Fig. 34 HAADF-STEM image of the Mn2O3 nanoparticles (A), the enlarged image of a single round-shaped nanoparticle (B) and atomic-resolution image of the surface showing atomic terraces confined to {111} crystal planes of the underlying cubic Ia-3 structure (C). Reprinted with permission from Ryabova, A. S.; Napolskiy, F. S.; Poux, T.; Istomin, S.Ya.; Bonnefont, A.; Antipin, D.M.; Baranchikov, A.Ye.; Levin, E.E.; Abakumov, A.M.; Kéranguéven, G.; Antipov, E.V.; Tsirlina, G.A.; Savinova, E.R. Electrochim. Acta 2016, 187, 161–172. Copyright 2016 Elsevier Ltd.

Fig. 35 Aberration-corrected HRTEM images of a Pt–Fe nanoparticle (A) and the map of the lattice contraction relative to the bulk Pt lattice (B). Reprinted with permission from Gan, L.; Yu, R.; Luo, J.; Cheng, Z.; Zhu, J. J. Phys. Chem. Lett. 2012, 3, 934–938. Copyright 2012 American Chemical Society.

nanoparticles as electrocatalysts for various types of fuel cells have been assessed with electron tomography.162–164 Tracing degradation of both catalyst support and particles can be performed in 3D by combining the identical location technique (see Section 7.10.7) and electron tomography.165,166 Fig. 36 demonstrates the HAADF-STEM images taken at different viewing angles of the (PtNbOx)/carbon composite cathode catalyst for polymer electrolyte membrane fuel cells. The identical location tomography series were measured before and after potential cycling of the TEM grid posed on a rotating disk electrode.167 The comparative analysis of the 3D reconstructions revealed the disappearance of the NbOx crystals, reducing the size of the Pt particles and changing their locations due to overall shrinkage of the carbon support. Quantitative measurement of the particle size suggested preferential dissolution of small Pt particles. For Li-ion battery cathode materials electron tomography has been successfully applied to detect the voids and cracks arising in course of prolonged electrochemical cycling. The voids are formed due to material “densification” in course of oxygen release from the Li-rich NMC cathodes as a consequence of irreversible oxygen redox reactions. On top of small voids formed at the material synthesis, large pores were observed surrounded with the material demonstrating the spinel/rock-salt structural reconstruction confirming the densification scenario.168 In delithiated LiNiO2 heating at 500  C promotes escalated oxygen evolution followed by the formation of lamellar arrangement of voids confined to the (001) planes as visualized with electron tomography (Fig. 37). Similar to the Li-rich NMCs, the voids are surrounded with the spinel- and rock-salt type domains demonstrating the common features of the “densification” process.169

7.10.6

Spectroscopy with electrons

Energy-dispersive X-ray (EDX) spectroscopy is one of the most widespread analytical techniques enabling qualitative and quantitative assessment of chemical compositions being realized in both SEMs and TEMs. The (S)TEM-EDX spectroscopy with multiple SDD X-ray detectors in combination with high-brightness FEG electron source (see the Section 7.10.2) opens up a possibility to

306

Electrode materials viewed with transmission electron microscopy

Fig. 36 Identical location electron tomography experiment for the pristine (A) and cycled (B) (Pt-NbOx)/carbon electrocatalyst showing selected images from the tilt series and the result of 3D reconstruction with color-coded particle’s volume. Reprinted with permission from Rossouw, D.; Chinchilla, L.; Kremliakova, N.; Botton, G. A. Part. Part. Syst. Charact. 2017, 34, 1700051. Copyright 2017 WILEY-VCH Verlag GmbH & Co.)

Fig. 37 3D reconstruction of delithiated LiNiO2 after heating to 500  C (A). The open and closed voids are indicated in blue and yellow, respectively. HR-TEM image of a lamellar arrangement of voids (outlined with cyan dashed lines) surrounded with the spinel and rock-salt (RS) domains (B). Reprinted from Wang, C.; Han, L.; Zhang, R.; Cheng, H.; Mu, L.; Kisslinger, K.; Zou, P.; Ren, Y.; Cao, P.; Lin, F.; Xin, H.L. Matter 2021, 4, 1–14.

gather the characteristic X-ray spectrum from a very small sample volume thanks to greatly improved signal-to-noise ratio and data collection rate. As the electrons transmitted through properly-oriented crystalline material demonstrate strong tendency to propagate along the atomic columns parallel to the electron beam due to electron channeling effect,170 the collection of the EDX spectra from a single atomic column became possible, thus enabling atomic-resolution compositional mapping with secondary X-rays. The HAADF-STEM images and atomic-resolution EDX maps in Fig. 38 demonstrate the stepped (vicinal) structure of the round nanoparticles of the LaCoO3 ORR electrocatalyst showing that the surface is composed of short fragments of perovskite {100} layers separated by one-unit-cell steps with terminating (LaO) layer and underlying (CoO2) layers (Fig. 38A), whereas the {110} terraces are terminated with the (LaCoO) layers (Fig. 38B) thus revealing different accessibility of Co as a catalytically active center.171

Electrode materials viewed with transmission electron microscopy

307

Fig. 38 HAADF-STEM images and atomic resolution EDX compositional maps (Co – green, La – red) of the surfaces of LaCoO3 particles. The surface consisting of the (LaO)-terminated {100} terraces stepped by one perovskite unit cell (A). The surface consisting of the (LaCoO)-terminated {110} terraces (B). The inserts show enlarged images of the regions used for the EDX mapping. The steps between the terraces are marked with arrows. Reprinted with permission from Mefford, J. T.; Kurilovich, A. A.; Saunders, J.; Hardin, W.G.; Abakumov, A.M.; Forslund, R.P.; Bonnefont, A.; Dai, S.; Johnston, K.P.; Stevenson, K.J. Phys. Chem. Chem. Phys. 2019, 21, 3327–3338. Copyright 2019 Royal Society of Chemistry.

Enhanced performance of the SDD X-ray detectors and their tolerance to the specimen tilt enabled a combination of the EDX compositional mapping and 3D tomography visualization as the EDX signal also satisfies the projection requirement, demonstrating monotonic intensity dependence on thickness for the high energy X-ray analytical lines. In contrast to conventional electron tomography, in the STEM-EDX tomography the spectral image is collected at every tilt angle. Each pixel of the spectral image contains an EDX spectrum covering the analytical lines of the elements of interest, thus the required elemental distribution map can be extracted by selecting the X-ray line of the desired energy and then used for 3D reconstruction. STEM-EDX tomography is particularly useful for visualization of compositional inhomogeneities in the layered oxide cathodes for Li-ion batteries, in which several transition metals are often combined and must be uniformly distributed even at the unit cell scale to deliver optimal electrochemical performance (Fig. 39).172,173 However, one should be aware of various artifacts which appear as an unavoidable cost of high counting rate of the SDD X-ray detectors and better signal-to-noise ratio on the X-ray spectra.174,175 These artifacts might spoil the identification of trace elements and result in false conclusions on the presence of certain minor admixtures (Fig. 40). Silicon escape peaks occur when high-energy X-ray photons are absorbed by a material of the detector exciting the Si-Kab X-rays. The associated energy loss results in a new peak in the X-ray spectrum with the energy Eescape ¼ EA  ESi-K (¼ 1.74 keV), where EA is the energy of the characteristic X-ray peak of the element A which is present in the sample in high concentration. The escape peaks are usually of low intensity but they become clearly visible with the large number of photons received by the SDD detectors at high count rate. A particularly misleading situation arises if the specimen is mounted on a copper support (very common for TEM studies) which produces intense X-ray fluorescence giving rise to Cu-Ka and Cu-Kb peaks even if Cu is not present in the sample. The escape peak for Cu-Ka radiation is at Eescape ¼ 8.046–1.74 keV ¼ 6.306 keV that coincides with the position of Fe-Ka peak at 6.405 keV within the typical energy resolution of the EDX spectrometers ( 130 eV). Thus, an erroneous conclusion can be drawn on contamination of the material with Fe. Another type of artifact originates from coincidence events when two photons with the energies EA and EB enter the detector within the time range shorter than the time resolution of the detector’s inspection circuit that will be interpreted by the detector as a single photon with the energy EA þ EB. If two strong peaks with the EA and EB energies are present in the spectrum, the weaker coincidence or “sum” peaks at the EA þ EA, EB þ EB and EA þ EB energies will be generated (Fig. 40) also spoiling qualitative analysis of minor admixture elements. Electron energy loss spectroscopy (EELS) can equally be employed for atomic-resolution compositional mapping albeit being confined to thin areas of the specimen.34 STEM-EELS is beneficial for visualization of coatings on nanosized particles of the electrode materials, particularly if the coating contains “light” elements (protective oxide layers, conductive carbon coatings), and can be used for assessment of the coating integrity and thickness. The Fe and C STEM-EELS compositional maps of LiFePO4 coated via carbonization of polymerized dopamine (Fig. 41) demonstrate progressively increasing thickness of the carbon layer with the dopamine polymerization time that actually became detrimental for Liþ diffusion through thick carbon coating layer worsening the charge/discharge rate capability.176 EELS is a unique technique to map Li distribution with high spatial resolution as lithium is not accessible with the EDX analysis, in contrast to carbon or oxygen. However, practical difficulties in Li mapping arise from the location of the Li-K edge in the low energy loss region (the onset energy of 55 eV), where it overlaps with a steeply rising background due to closeness of very strong zero loss peak and with the M-edges of 3d transition metals (the onset energies: 47 eV (Ti, V),

308

Electrode materials viewed with transmission electron microscopy

Fig. 39 3D visualizations of STEM-EDX reconstructions demonstrating the distribution of Mn, Ni and O in the Li1.2Ni0.2Mn0.6O2 cathode material. Reprinted with permission from Genc, A.; Kovarik, L.; Gu, M.; Cheng, H.; Plachinda, P.; Pullan, L.; Freitag, B.; Wang., C. Ultramicroscopy 2013, 131, 24–32. Copyright 2013 Elsevier.

48 eV (Cr), 51 eV (Mn), 57 eV (Fe), 62 eV (Co), 68 eV (Ni)). The latter is a crucial obstacle for studying Li-ion battery electrodes usually comprising one or more 3d metals. In this case, extracting the Li signal requires its deconvolution from the background and other overlapping edges through the spectrum modeling and fitting.177 Tomography with STEM-EELS or energy-filtered TEM is somewhat cumbersome because of strong multiple scattering effects and background variation with thickness changing while the specimen is rotated over large angles. Nevertheless, STEM-EELS tomography was successfully applied to mapping chemical compositions of the rod-like specimens and separate nanoparticles in which the thickness variations are minimized.178,179 The great added value of EELS for studying the electrochemical reactions and electroactive materials is its ability to provide local information on the oxidation states and chemical bonding, especially for the 3d transition metal cations. The position and shape of the L2,3 (2p-3d transitions) excitation edge demonstrate a strong correlation with the formal oxidation state of the 3d metals (Fig. 42), but establishing the quantitative relationships is not straightforward and lacks universal approach. In a so-called “white line” ratio method the intensity ratio of the L3 (2p3/2-3d transition) and L2 (2p1/2-3d transition) peaks is empirically correlated with the oxidation state. The energy position of the excitation edge also demonstrates systematic shift with changing the oxidation state that is called a “chemical shift”: increasing the oxidation state displaces the excitation edge toward higher energy loss. The parametrizations of the oxidation states of 3d metals through intensities and their ratios, absolute and relative energy positions of the excitation edges are summarized in Table 3. Practical application of these equations requires exact reproduction of the measurement and data treatment conditions employed in the original works, including background subtraction, taking into account multiple scattering effects, selecting proper intensity integration windows and energy position calibration. One should also note that the shape of the L2,3 excitation edge depends not only on the oxidation state of the transition metal, but also on its coordination number.39 STEM-EELS spectral imaging can be followed by the selection of the desired energy range in the EELS spectrum where the element in the oxidation state of interest dominates and reconstructing the image in the electrons contributing to this specific energy loss that generates the elemental map sensitive to this oxidation state. Valence mapping can also be facilitated by decomposition of

Electrode materials viewed with transmission electron microscopy

309

Fig. 40 EDX spectrum taken with SDD X-ray detectors (Titan Themis Z transmission electron microscope, Super-X EDX system, 200 kV) on the cathode material LiCoO2. The top panel demonstrates full intensity scale, while scaled-up view is provided at the bottom panel illustrating a set of artifacts. The Cu-K and Pb-L,M signals arise from interfering fluorescence of the copper TEM grid used as sample support and the TEM radiation shield, respectively. Intense CoKa, CoKb and CuKa peaks generate the corresponding escape peaks with the energy offset of 1.74 keV. The Co-K and Cu-K lines produce a series of sum peaks at the high energy part of the spectrum.

the EELS spectra into contributions by multiple linear least-square (MLLS) fitting with the reference spectra measured on the materials in which the oxidation states are well defined.190 Applied to the LiNi1/3Mn1/3Co1/3O2 Li-ion battery cathode material, the EELS valence mapping reveals that even at the elevated electrode potential (4.1 V vs Liþ/Li) the  2 nm surface layer contains divalent transition metals (Fig. 43). This surface reduction region originates from the first charge contributing to lowering of Coulombic efficiency and then develops into a rock-salt type reconstruction layer. In the cathode materials with two-phase Li (de)intercalation mechanism mapping of the oxidized and reduced valence states visualizes the distribution of the delithiated and lithiated phases, such as FePO4 and LiFePO4 in a half-charged Li0.5FePO4 (Fig. 44).191

7.10.7

In situ and operando observations of electrochemical reactions

Employing the full resolving power of modern transmission electron microscopes for unraveling the electrochemically-induced chemical and structural transformations in real time would shed light on many mysterious processes occurring at the surface and in the bulk of the electrode materials. These might include surface restructuring and degradation of electro- and photocatalysts, capturing electronic structure and chemical nature of the surfaces at different real environments, tracing chemistry and structure of intercalation-type electrode materials, solid electrolyte interphase (SEI) formation and evolution as a function of applied potential and number of cycles in metal-ion batteries and many other challenging subjects.192 Unfortunately, the most serious obstacles for these observations arise from the physics of the interaction of high energy electrons with matter: 1) the specimen is immersed into deep vacuum of the TEM column that prevents using liquid electrolyte with high vapor pressure in open cells without protecting electron-transparent windows; 2) even relatively thin electrolyte layer introduces substantial scattering of electrons thus deteriorating quality of images and EELS spectra193,194; 3) most electrolytes are susceptible to electron beam damage through radiolysis introducing side reactions and contaminating the specimen by decomposition products.120,195,196 Thus, in situ and operando TEM observations in the electrochemical cells with liquid electrolyte are conducted at low electron dose being limited in spatial resolution precluding atomic-scale imaging. They allow monitoring many electrochemical phenomena, such as electroplating and stripping,193,197–201 particularly lithium electrodeposition dynamics (Fig. 45),202,203 lithium insertion and deinsertion dynamics in Li-ion battery cathodes,193 solid electrolyte interface (SEI) formation.204,205 A visualization of the

310

Electrode materials viewed with transmission electron microscopy

Counts

(E)

2.5

3.0

3.5

4.0

4.5

5.0

5.5

5.5

6.6

Thickness, nm

Counts

(G)

2.5

3.0

3.5

4.0

4.5

5.0

5.5

6.0

6.5

Thickness, nm

Counts

(H)

4.5

5.0

5.5

6.0

6.5

7.0

7.5

8.0

8.5

Thickness, nm

Fig. 41 STEM-EELS compositional maps of Fe (green) and C (red) showing conductive carbon coating on the LiFePO4 nanoparticles treated with dopamine for polymerization times of 2 h (A), 24 h (B) and 40 h (C). The corresponding coating thickness distributions are provided in the panels E, F, G, respectively. The carbon distribution in LiFePO4 after annealing with glucose is demonstrated in the panel D. Reprinted with permission from Iarchuk, A. R.; Nikitina, V. A.; Karpushkin, E. A.; Sergeyev, V.G.; Antipov, E.V.; Stevenson, K.J.; Abakumov. A.M. ChemElectroChem 2019, 6, 5090– 5100. Copyright 2019 Wiley-VCH Verlag GmbH & Co.

nanoscale Kirkendall effect is provided by operando TEM in a liquid cell performed on Ag nanocubes in course of the galvanic Au for Ag replacement according to the reaction 3Ag þ AuCl4 / Au þ 3AgCl þ Cl.206 Kirkendall effect arising from unequal diffusion rates of Au and Ag generates vacancies in the metal sublattice which accumulate into voids at the interface between two materials. The voids then evolve forming intricate hollow microstructures (Fig. 46). Interpretation of the TEM observations in liquid electrolyte, however, requires particular caution because electron beam irradiation initiates cascade of chemical transformations generating the products which might affect the nature of the studied electrochemical processes. High energy electrons excite the valence electrons of the solvent molecules breaking the chemical bonds and generating a whole range of the radical, ionic and molecular species. In water they are hydrated electrons eh, hydrogen radical H•, hydroxyl radical OH•, and H2, which further react producing H2O2, H3Oþ, HO2•, and eventually recombine back to H2O or form bubbles at sufficiently high dose rate.195,207 The computational modeling demonstrates that at uniform irradiation without concentration gradients and mass transfer, steady state concentrations of the decomposition products should be achieved, with the values depending on the electron dose rate, pH and amount of dissolved oxygen (Fig. 47). If the beam-induced decomposition of water is simulated in detail, much less is known on the radiolysis products for organic solvents, such as linear and cyclic alkyl carbonates commonly used as electrolyte solvents in metal-ion batteries. However, the most unstable component of such electrolytes are the dissolved salts, such as LiPF6, which are reduced with the solvated electrons according to the reaction LiXF6 þ 2e þ 2Liþ / XF3 þ 3LiF (X ¼ pnictogen), whereas the solvents are much less susceptible to the electron beam damage.120 The particularly stable combination was found to be a solution of lithium trifluoromethanesulfonate CF3SO3Li in dimethyl sulfoxide (DMSO) known to be an effective radical scavenger.208 Limitation in the TEM spatial resolution for an object immersed in liquid stems from the inelastic interaction of the electron beam with the liquid layer. It is largely defined by the caused energy spread and chromatic aberration of the optical system: (15) dTEM ¼ AE aCs t E2 where dTEM is the TEM spatial resolution, a is a beam convergence semi-angle, CS is a chromatic aberration coefficient, t is the thickness of the liquid layer and E is the electron beam energy. The coefficient AE defines the nature of the electrolyte and depends on

Electrode materials viewed with transmission electron microscopy

311

Fig. 42 The EELS L2,3 edges together with the O–K edges for some 3d transition metals relevant for applications in metal-ion batteries and electrocatalysis sorted by their formal oxidation states. Energy resolution is 0.9 eV. Reprinted with permission from Tan, H.; Verbeeck, J.; Abakumov, A.; Van Tendeloo, G. Ultramicroscopy 2012, 116, 24–33. Copyright 2012, Elsevier.

its molecular weight, atomic number and density.209–211 For the imaging in a STEM mode, the resolution can be considered as the smallest size of the object visible above the experimental noise and can be expressed as: rffiffiffiffiffiffi t (16) dSTEM  N0 where dSTEM is the STEM spatial resolution, t is the thickness of the liquid layer and N0 is the number of incident electrons. Additionally, the spatial resolution in the TEM and STEM modes depends on the signal-to-noise ratio, scaling as d  D-1/4 (D – electron dose rate). Thus, in both TEM and STEM cases, higher electron dose and thinner liquid cell improve the spatial resolution (another low-dose technique to retrieve the atomic-scale information on the materials in liquid cells is electron diffraction tomography described in detail in Section 7.10.4). Atomic resolution TEM imaging in liquid electrolyte has been attained only for extremely thin cell with graphene windows,212 in which applying electric bias was not possible by the time of writing this Chapter and the electrochemical reactions were mimicked by the affection of the electron beam.213 This approach utilizes the reducing power of solvated electrons and, being applied to Li-ion battery materials, enables investigating anodic processes, such as expansion of Si nanoparticles upon lithiation revealing its anisotropic character,214 or morphology evolution in conversion-type anodes.215,216 However, the interpretation of these observations and, particularly, correlations with the processes in real electrochemical cells is significantly complicated by excitations of the electronic and vibrational states of the electrode material(s) by inelastic energy transfer from high-energy electrons. The reactions starting from excited states might demonstrate significantly reduced energy barriers or even proceed through barrierless paths, thus, significantly changing the reaction kinetics. Understanding these excitations requires sophisticated time-dependent computational approaches.217 While correlating the (S)TEM observations in liquid cells with the electrochemistry measured in the conventional two-electrode cells, additional caution is required because of open circuit electrode potential negative offset caused by the electron beam, which may range from dozens to hundreds of millivolts depending linearly on the electron beam current density.218 As the liquid cells suffer from significant intrinsic limitations imposed by extensive electron scattering in relatively thick (hundreds of nm) electrolyte layer, alternative nano-electrochemical cells have been designed based on the electrolytes with low (or virtually zero) vapor pressure, namely ionic liquids and solid electrolytes. These cells do not need protecting vacuum-tight windows and can be fabricated with the slicing and shaping capabilities of the focused ion beam (FIB) technique and mounted on a TEM holder allowing for micromechanical manipulation and external electric biasing. Such experimental setup enables

Table 3 Element Ti

V

Cr

Mn

Electrode materials viewed with transmission electron microscopy Empirical relationships between the parameters of the EELS L2,3 excitation edges of transition metals and their formal oxidation states (I – intensity of the corresponding edge, E – its position along the energy loss axis, V – formal oxidation state of the transition metal). Parameters

Equation

Comments

x ¼ IL2/IL3 y ¼ Ti4 þ/STi

  y ¼ ln x 0:88ð20Þ  0:218ð9Þ 0:22ð4Þ

þ3  VTi  þ4; oxides;

x ¼ Ti4 þ/STi y ¼ IL2/IL3 x ¼ VV y ¼ EL3

y ¼ 1.963e3.446x þ 2.653

x ¼ VV y ¼ EL3

y ¼ 1.15(4)x þ 510.28(16)

x ¼ VCr y ¼ E(CrL3)  E(OK) x ¼ VCr y ¼ [EL2  EL3]/[IL2/IL3]

y ¼ 1.7(1)x þ 41.8

þ3  VTi  þ4; oxides;

y ¼ 2.5x þ 510

þ3  VV  þ 5; oxides; EL3 is measured at 10% of the maximum þ3  VV  þ 5; oxides; EL3 is measured at 10% of the maximum þ2  VCr  þ 6; oxides;

y ¼ 0.78(7)x þ 2.1(2)

þ2  VCr  þ 6; oxides, silicate, pyrophosphate; 5.0 eV integration windows centered on the maxima of the L3 and L2 edges

x ¼ VMn y ¼ EL3

y ¼ 0.715x þ 641.16

x ¼ VMn y ¼ IL3/IL2

y ¼  0.734x2 þ 4.638x  5.014

x ¼ VMn y ¼ EL2  EL3

y ¼ 2.65(6)x þ 103.45(23)

x ¼ IL3/IL2 y ¼ VMn

y ¼ 0.68x2  5.02x þ 11.27 y ¼ 0.64x2  4.83x þ 11.02

x ¼ VMn y ¼ IL3/IL2

Fe

integration windows: 455.8–456.8 eV 465.25–466.25 eV

x 0:83ð11Þ

EL3at the gravity center þ3  VMn  þ4; oxides; 4.0 eV integration windows centered on the maxima of the L3 and L2 edges þ3  VMn  þ4; oxides; energy difference between the L2 and L3 peak maxima þ2  VMn  þ4; oxides; curves without (top) and with multiple scattering correction; 4.0 eV integration windows centered on the maxima of the L3 and L2 edges



 y ¼ 24ð9Þ  e

þ2  VMn  þ4; silicates;

=

312

þ 1:64ð8Þ

x ¼ VMn y ¼ EL3

y ¼ 0.83(2)x þ 637.43(7)

x ¼ Fe3 þ/SFe y ¼ IL3/IL2

1 y ¼ 0:193ð7Þx 2 0:465ð9Þx 1 þ0:366ð3Þ

x ¼ Fe3 þ/SFe y ¼ I1/SIi

y ¼  68.4(7)x þ 83.1(4)

x ¼ Fe3 þ/SFe y ¼ IL3/IL2

y ¼ 0.74x2 þ 0.69x þ 4.22

þ2  VMn  þ4; oxides; 8.0 eV integration windows centered on the maxima of the L3 and L2 edges with multiple scattering correction þ2  VMn  þ4; oxides; EL3 is measured at 10% of the maximum þ2  VFe  þ 3; oxides, silicates; integration windows: 708.5–710.5 eV 719.7–721.7 eV þ2  VFe  þ 3; oxides, silicates; integral intensity of the first peak of the L3 edge at 707.8 eV normalized to the total L3 edge intensity þ2  VFe  þ 3; oxides, silicates; 2.0 eV integration windows centered on the maxima of the L3 edge and at þ12.8 eV off the L3 edge

References 180

181 182 40 183 184

185 186

186 187

40

40 188

188

189

Electrode materials viewed with transmission electron microscopy

313

Fig. 43 HAADF-STEM image and STEM-EELS valence maps for Ni, Mn and Co in LiNi1/3Mn1/3Co1/3O2 charged to 4.1 V vs Li/Liþ demonstrating the reduced surface layer. Reprinted with permission from Liu, H.; Bugnet, M.; Tessaro, M. Z.; Harris, K.J.; Dunham, M.J.R.; Jiang, M.; Goward, G.R.; Botton, G.A. Phys. Chem. Chem. Phys. 2016, 18, 29064–29075. Copyright 2016 Royal Society of Chemistry.

Fig. 44 ADF-STEM images of the crystals in Li0.5FePO4 and the maps of the LiFePO4 (LFP) and FePO4 phases obtained through the MLLS fitting of the STEM-EELS spectral imaging data with the LFP and FP reference spectra. Reprinted with permission from Honda, Y.; Muto, S.; Tatsumi, K.; Kondo, H.; Horibuchi, K.; Kobayashi, T.; Sasaki, T. J. Power Sources 2015, 291, 85–94. Copyright 2015, Elsevier.

monitoring of the electrochemically-induced changes in the electrode structure with atomic resolution. A formation and propagation of the step-like grain boundary between LiFePO4 and its delithiated form FePO4 have been demonstrated in the nano-sized Liion battery with LiFePO4 as a positive electrode, metallic Li as a negative electrode and a thin layer of Li2O as a solid electrolyte (Fig. 48).219 The electrode of interest can also be prepared in a form of nanowire being suitable for atomic-resolution observations, such as peeling-off atomic layers at the electrochemical Si/LixSi interface upon lithiation (Fig. 49),220 as well as monitoring of volumetric variation and expansion anisotropy, fracturing and amorphization upon lithium insertion into Si anode.221–223 The clear drawback of such in situ TEM nanowire cells is a significant difference between the electrochemical conditions compared to the real battery cell with liquid electrolyte. This is, however, not true for FIB-sliced thin-film all-solid-state battery cells cut out from the real cell stacks. Using such nanobattery setups the cathode nanocrystallization in the LCO/LLZO (LCO – LiCoO2, LLZO - Li6.75La2.84Y0.16Zr1.75Ta0.25O12 solid state electrolyte) cell upon delithiation was monitored resulting in splitting of initially

314

Electrode materials viewed with transmission electron microscopy

Fig. 45 An example of the TEM liquid cell design: schematic of the imaging geometry in a STEM mode (A), SEM images of the electrodes at the center of the chip with the patterned Ti electrode marked with blue arrows (B, C, D). Li cycling while imaging every 15 s: BF STEM image before Li deposition (E), near beginning and end of deposition cycle (F, G), and near beginning and end of stripping cycle (H, I) demonstrating disconnected “dead” Li. Reprinted with permission from Leenheer, A. J.; Jungjohann, K. L.; Zavadil, K. R.; Sullivan, J. P.; Harris, C. T. ACS Nano 2015, 9, 4379– 4389. Copyright 2015 American Chemical Society.

Fig. 46 Time series of TEM images and corresponding schematics showing the morphological evolution of an Ag nanocube during galvanic Au for Ag replacement. Yellow arrows indicate the deposited shell, green arrows indicate pores. Reprinted from Chee, S. W.; Tan, S. F.; Baraissov, Z.; Bosman, M.; Mirsaidov, U. Nat. Commun. 2017, 8, 1224.

single crystalline LCO into a set of domains separated by coherent twin and antiphase boundaries.224 Similar observations on the LNMO/LLZO (LNMO – LiNi0.5Mn1.5O4) and NMC/LLTO (NMC - LiNi0.76Mn0.14Co0.10O2, LLTO - Li0.5La0.5TiO3) stacks revealed inhomogeneous delithiation, localized migration of the transition metal cations, formation of antiphase boundaries and structural transition to a disordered rock-salt-type state (Fig. 50).225,226 Deeper insight into the factors determining ionic transport in all-solid-state batteries has been achieved from in situ scrutinizing the grain boundaries and interfaces with (S)TEM imaging, electron holography and electron energy loss spectroscopy. EELS mapping of the Li concentration as a function of state of charge in the particles of the LiNi0.85Co0.15Al0.05O2 (NCA) cathode revealed strong concentration gradients at the interfaces between the crystallites of different orientation and virtually no gradient between the crystallites of the same orientation (Fig. 51) demonstrating that the Li ion transport significantly depends on the mutual alignment of the crystal structures in the adjacent particles.227 Electron holography imaging of the LiCoO2/ Li1 þ x þ yAlyTi2ySixP3xO12 cell revealed a sharp potential drop at the cathode – solid electrolyte interface due to high interface resistance, one of the main obstacles for practical realization of all-solid-state Li batteries.228 In situ STEM-EELS chemical mapping revealed that this interface contains inactive regions consisting of LiCoO2 and Co3O4. The latter leads to the higher interfacial resistance of the Li-ion transfer because of the lack of Li-ion diffusion pathways.229 Beyond battery research, additional examples of the application of in situ and operando TEM in electrocatalysis can be found in several reviews.230,231 An elegant technique of identical location electron microscopy (IL-EM) has especially to be mentioned in the

Electrode materials viewed with transmission electron microscopy

315

Fig. 47 Concentrations of eh, H•, H2, H2O2, OH• and O2 as a function of time (deaerated water, 7.5  107 Gy/s dose rate) (A); steady-state concentrations of the same products as functions of dose rate (B); steady-state concentrations as functions of pH (C) and the ratio of concentrations in oxygen-saturated water (C(O2) ¼ 0.255 mM) and in deaerated water as functions of pH (D). Reprinted with permission from Schneider, N. M.; Norton, M. M.; Mendel, B. J.; Grogan, J.M.; Ross, F.M.; Bau, H.H. J. Phys. Chem. C 2014, 118, 22373–22382. Copyright 2014 American Chemical Society.

context of catalyst research.232 Although it is neither in situ nor operando technique, it allows for tracking local electrochemicallyinduced changes in the sample by comparing their (S)TEM images before and after electrochemical testing. The location of the area of interest is noted in the numbered or marked holes in the electrically conductive, corrosion-resistant and catalytically inactive TEM grid, which then is used as a working electrode in an electrochemical cell (Fig. 52A–G). After cyclic voltammetry measurements, the grid is mounted in the TEM holder and observations are performed exactly at the same place tracing back its location through the grid’s coordinate system. This approach became a standard method of studying degradation phenomena in various electrocatalysts, such as particle detachment from support, dissolution, dealloying, agglomeration and Ostwald ripening.233,234 Because electron scattering at the electrolyte layer or cell windows does not obstacle achieving atomic spatial resolution, tracing the structure of catalytic particles down to a single atom at the surface layers has been demonstrated being also coupled with the atom counting and strain quantification using the statistical parameter estimation theory (Fig. 52H–M, see also the Section 7.10.5.1).235

7.10.8

Conclusions and outlook

Transmission electron microscopy has already provided an invaluable contribution to unraveling the complexity of the electrochemical materials. Nevertheless, the TEM field develops very dynamically and within a short time period new and, perhaps, now unforeseen possibilities may emerge. Electron dose required for atomic-resolution imaging with a high signal-to-noise ratio will be further lowered due to progress in noiseless detectors capable of single electron counting. This will be critically important for TEM observations on fragile materials (such as metallic Li dendrites) and for the experiments in liquid cells where the sensitivity of the electrolyte to the electron beam is an issue. Another mean for reducing beam damage and unwanted artifacts is lowering the electron beam energy by decreasing acceleration voltage while keeping reasonable spatial resolution thanks to aberration correctors. Advances in aberration-corrected electron optics already enabled TEMs to operate at acceleration voltages of 15–60 kV delivering atomic resolution,236–239 and more progress in the field of low-voltage TEM is yet to come, particularly, in the area of specimen thinning to transparency for low-energy electrons. Hardware improvement will also be spread over electron sources and EELS spectrometers. Already now the attainable energy resolution of the commercial analytical TEMs is shifted to < 0.05 eV, and the limit of  0.004 eV is already reported.240 This enables retrieving finer details on the electron energy loss near-edge structure and the analysis of vibrational properties of solids that in turn calls for further progress in ab initio computational modeling of the EELS spectra which might unlock quantitative interpretation of the local spectral data.

316

Electrode materials viewed with transmission electron microscopy

Fig. 48 In situ TEM electrochemical setup (A), the morphology and orientation of the FePO4 nanosized positive electrode (B). High resolution TEM images demonstrating a step-like phase boundary between FePO4 and LiFePO4 and its migration along the [010] direction during lithiation (C, D, E). Reprinted with permission from Zhu, Y.; Wang, J. W.; Liu, Y.; Liu, X.; Kushima, A.; Liu, Y.; Xu, Y.; Mao, S.X.; Li, J.; Wang, C.; Huang, J.Y. Adv. Mater. 2013, 25, 5461–5466. Copyright 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

(S)TEM imaging has already achieved sub-Å resolution and demonstrated the ability to visualize the atomic species from the whole Periodic Table. Further progress is expected in 3D imaging with atomic resolution, which has already been demonstrated for monoatomic metallic nanoparticles,156,157,241 but the true challenge lies in reconstructing the 3D structures of the crystallites consisting of many sorts of atoms, planar defects and grain boundaries that constitute the vast majority of the electrode materials. On a larger spatial scale, precise crystallographic information averaged over sub-micron area can already be extracted ex situ or in situ with electron diffraction tomography, but there is still room for improvement related to the design of dedicated

Electrode materials viewed with transmission electron microscopy

317

Fig. 49 In situ TEM images of the pristine (A) and partially lithiated (B) silicon nanowire. High-resolution TEM image after 5 s of lithiation showing that the lithiation proceeds through peeling-off the {111} Si crystal planes (C). Reprinted with permission from Liu, X. H.; Wang, J. W.; Huang, S.; Fan, F.; Huang, X.; Liu, Y.; Krylyuk, S.; Yoo, J.; Dayeh, S.A.; Davydov, A.V.; Mao, S.X.; Picraux, S.T.; Zhang, S.; Li, J.; Zhu, T.; Huang, J.Y. Nat. Nanotech. 2012, 7, 749–756. Copyright 2012 Macmillan Publishers Limited.

electrochemical cells allowing for larger tilt angles thus increasing the reciprocal space coverage, and to application of dynamical refinement for improving the quality and reliability of the obtained structural data. Electron diffraction imaging (4D STEM) measures the diffraction pattern at every scanned point thus generating huge data sets with enciphered information on electric and magnetic fields, ferroelectric polarizations, particle orientation and strain, and even maps of the fields around single atoms.242,243 This approach will be of high demand in the quantitative microstructural analysis of the layered oxide Li-ion battery cathodes in relation to the mechanical properties of their agglomerated particles influencing the cyclic stability.244–246 Processing large datasets generated with 4D STEM calls for advanced procedures for pattern recognition and analysis of diffraction pattern positions, where the approaches based on neural networks come into play.247 The analysis with the machine learning algorithms will play an increasing role in disentangling multiple spatial and spectral physical and/or chemical components of the electron microscopy and spectroscopy data.248 Summarizing, with the expected developments in the transmission electron microscopy imaging and spectroscopic techniques, supported with quantitative data treatment and modeling, coupled with the ability to interrogate the electrochemical reactions in

318

Electrode materials viewed with transmission electron microscopy

Fig. 50 In situ ABF-STEM images of pristine and charged LiNi0.5Mn1.5O4 demonstrating the transformation from ordered to disordered structure. Reprinted from Gong, Y.; Chen, Y.; Zhang, Q.; Meng, F.; Shi, J.-A.; Liu, X.; Liu, X.; Zhang, J.; Wang, H.; Wang, J.; Yu, Q.; Zhang, Z.; Xu, Q.; Xiao, R.; Hu, Y.-S.; Gu, L.; Li, H.; Huang, X.; Chen, L. Nat. Commun. 2018, 9, 3341.

Fig. 51 Spatial distribution of Li concentration in the NCA cathode at different states of charge mapped with in situ EELS on Li-K edge. The remaining crystallite with high Li concentration after 15 h of relaxation is disconnected because of cracking along the grain boundaries. Clear Li concentration gradient is observed at the grain boundaries between the particles having different orientation (1 and 2, 3 and 4 in the 3.59–3.69 V panel (top right), the orientation is determined with electron diffraction (not shown)), whereas no Li gradient is seen between the particles 2 and 3 oriented identically. Reprinted with permission from Nomura, Y.; Yamamoto, K.; Hirayama, T.; Igaki, E.; Saitoh, K. ACS Energy Lett. 2020, 5, 2098– 2105. Copyright 2020 American Chemical Society.

Electrode materials viewed with transmission electron microscopy

319

Fig. 52 Illustration of the IL-EM approach: using the TEM grid as an electrode by connecting it through a wire (A), tweezer (B) or pressing it against glassy carbon rotating disk electrode (C); location tracking through progressively increasing magnification and tracing the grid coordinate system (D–G). An example of the IL electron microscopy: ADF-STEM images of a Pt-Co nanoparticle before (H) and after (I) electrochemical activation; the removed and deposited atomic columns are marked with yellow and green, respectively. Strain tensor components in the Pt-Co nanoparticle before (J, K) and after (L, M) electrochemical activation. (A–G) Reprinted with permission from Hodnik, N.; Cherevko, S. Curr. Opin. Electrochem. 2019, 15, 73–82. Copyright 2019 Elsevier. (H–M) Reprinted from Moriau, L.; Hrnjic, A.; Pavlisic, A.; Kamsek, A.R.; Petek, U.; RuizZepeda, F.; Sala, M.; Pavko, L.; Selih, V.S.; Bele, M.; Jovanovic, P.; Gatalo, M.; Hodnik, N. iScience 2021, 24, 102102.

their native environment, anticipated formidable results will greatly contribute to our fundamental understanding of diverse electrochemical systems.

Acknowledgement This work has been supported by the Russian Science Foundation (grant 20-13-00233).

320

Electrode materials viewed with transmission electron microscopy

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62. 63. 64.

Ishikawa, R.; Okunishi, E.; Sawada, H.; Kondo, Y.; Hosokawa, F.; Abe, E. Nat. Mater. 2011, 10, 278–281. Van Tendeloo, G.; Bals, S.; Van Aert, S.; Verbeeck, J.; Van Dyck, D. Adv. Mater. 2012, 24, 5655–5675. Williams, D. B.; Carter, C. B. Transmission Electron Microscopy: A Textbook for Materials Science, Springer Science & Business Media, LLC, 1996, 2009. Inkson, B. J. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) for materials characterization. In Materials Characterization Using Nondestructive Evaluation (NDE) Methods; Hübschen, G., Altpeter, I., Tschuncky, R., Herrmann, H.-G., Eds., Elsevier, 2016; pp 17–43. Morniroli, J. P. Ultramicroscopy 1992, 45, 219–239. Kolb, U.; Mugnaioli, E.; Gorelik, T. E. Cryst. Res. Techn. 2011, 46, 542–554. Mugnaioli, E. Acta Crystallogr. B 2015, 71, 737–739. Mugnaioli, E.; Gemmi, M. Z. Krist. 2018, 233, 163–178. Palatinus, L.; Jacob, D.; Cuvillier, P.; et al. Acta Crystallogr. A 2013, 69, 171–188. Palatinus, L.; Petrícek, V.; Correâ, C. A. Acta Crystallogr. A 2015, 71, 235–244. Gemmi, M.; Mugnaioli, E.; Gorelik, T. E.; Kolb, U.; Palatinus, L.; Boullay, P.; Hovmöller, S.; Abrahams, J. P. ACS Cent. Sci. 2019, 5, 1315–1329. Van Dyck, D. Atomic resolution electron microscopy. In Handbook of Nanoscopy; Van Tendeloo, G., Van Dyck, D., Pennycook, S. J., Eds.; vol. 1; Wiley-VCH, 2012; pp 45–79. Chapter 2. Zou, X. Microsc. Res. Tech. 1999, 46, 202–219. Zandbergen, H. W.; Van Dyck, D. Microsc. Res. Tech. 2000, 49, 301–323. Urban, K. W.; Jia, C. L.; Houben, L.; Lentzen, M.; Mi, S.-B.; Tillmann, K. Phil. Trans. R. Soc. A 2009, 367, 3735–3753. Pennycook, S. J.; Nellist, P. D. Scanning Transmission Electron Microscopy, Springer New York, 2011. Hartel, P.; Rose, H.; Dinges, C. Ultramicroscopy 1996, 63, 93–114. Abe, E. Mater. Trans. 2003, 44, 2035–2041. Nellist, P. D. Scanning transmission electron microscopy. In Science of Microscopy; Hawkes, P., Spence, J. C. H., Eds., Springer, 2007. Findlay, S. D.; Shibata, N.; Sawada, H.; Okunishi, E.; Kondo, Y.; Yamamoto, T.; Ikuhara, Y. Appl. Phys. Lett. 2009, 95, 191913. Findlay, S. D.; Shibata, N.; Sawada, H.; Okunishi, E.; Kondo, Y.; Ikuhara, Y. Ultramicroscopy 2010, 110, 903–923. Batuk, D.; Hadermann, J.; Abakumov, A. M.; Vranken, T.; Hardy, A.; Van Bael, M.; Van Tendeloo, G. Inorg. Chem. 2011, 50, 4978–4986. den Dekker, A. J.; Gonnissen, J.; De Backer, A.; Sijbers, J.; Van Aert, S. Ultramicroscopy 2013, 134, 34–43. Yang, H.; Rutte, R. N.; Jones, L.; Simson, M.; Sagawa, R.; Ryll, H.; Huth, M.; Pennycook, T. J.; Green, M. L. H.; Soltau, H.; Kondo, Y.; Davis, B. G.; Nellist, P. D. Nat. Commun. 2016, 7, 12532. Yang, H.; Jones, L.; Ryll, H.; Simson, M.; Soltau, H.; Kondo, Y.; Sagawa, R.; Banba, H.; Maclaren, I.; Nellist, P. D. J. Physics: Conf. Ser. 2015, 644, 012032. Lazic, I.; Bosch, E. G. T.; Lazar, S. Ultramicroscopy 2016, 160, 265–280. Yücelen, E.; Lazic, I.; Bosch, E. G. T. Sci. Rep. 2018, 8, 2676. Intaraprasonk, V.; Xin, H. L.; Muller, D. A. Ultramicroscopy 2008, 108, 1454–1466. Lupini, A. R.; De Jonge, N. Microsc. Microanal. 2011, 17, 817–826. Xin, H. L.; Muller, D. A. J. Electr. Microsc. 2009, 58, 157–165. Bals, S.; Goris, B.; Altantzis, T.; Heidari, H.; Van Aert, S.; VanTendeloo, G. C. R. Physique 2014, 15, 140–150. Goris, B.; Roelandts, T.; Batenburg, K. J.; Heidari, H.; Bals, S. Ultramicroscopy 2013, 127, 40–47. Alpers, A.; Gardner, R. J.; Konig, S.; Pennington, R. S.; Boothroyd, C. B.; Houben, L.; Dunin-Borkowski, R. E.; Batenburg, K. J. Ultramicroscopy 2013, 128, 42–54. Muller, D. A. Nat. Mater. 2009, 8, 263–270. D’Alfonso, A. J.; Freitag, B.; Klenov, D.; Allen, L. J. Phys. Rev. B 2010, 81, 100101. Muller, D. A.; Kourkoutis, L. F.; Murfitt, M.; Song, J. H.; Hwang, H. Y.; Silcox, J.; Dellby, N.; Krivanek, O. L. Science 2008, 319, 1073–1076. Sigle, W. Annu. Rev. Mater. Res. 2005, 35, 239–314. Schlossmacher, P.; Klenov, D. O.; Freitag, B.; von Harrach, H. S. Microscopy Today 2010, 18, 14–20. Turner, S.; Egoavil, R.; Batuk, M.; Abakumov, A. M.; Hadermann, J.; Verbeeck, J.; Van Tendeloo, G. Appl. Phys. Lett. 2012, 101, 241910. Tan, H.; Verbeeck, J.; Abakumov, A.; Van Tendeloo, G. Ultramicroscopy 2012, 116, 24–33. Tan, H.; Egoavil, R.; Beche, A.; Martinez, G. T.; Van Aert, S.; Verbeeck, J.; Van Tendeloo, G.; Rotella, H.; Boullay, P.; Pautrat, A.; Prellier, W. Phys. Rev. B 2013, 88, 155123. Schneider, C. A.; Rasband, W. S.; Eliceiri, K. W. Nat. Methods 2012, 9, 671–675. Rueden, C. T.; Schindelin, J.; Hiner, M. C.; DeZonia, B. E.; Walter, A. E.; Arena, E. T.; Eliceiri, K. W. BMC Bioinformatics 2017, 18, 529. Amandinea, V.; Cédric, M.; Sergio, M.; Patricia, D. Micron 2019, 121, 90–98. Roels, J.; Vernaillen, F.; Kremer, A.; Gonçalves, A.; Aelterman, J.; Luong, H. Q.; Goossens, B.; Philips, W.; Lippens, S.; Saeys, Y. Nat. Commun. 2020, 11, 771. De Backer, A.; van den Bos, K. H. W.; Van den Broek, W.; Sijbers, J.; Van Aert, S. Ultramicroscopy 2016, 171, 104–116. Van Aert, S.; De Backer, A.; Martinez, G. T.; Goris, B.; Bals, S.; Van Tendeloo, G. Phys. Rev. B 2013, 87, 064107. Lin, R.; Zhang, R.; Wang, C.; Yang, X.-Q.; Xin, H. L. Sci. Rep. 2021, 11, 5386. Klinger, M.; Jäger, A. J. Appl. Cryst. 2015, 48, 2012–2018. Labar, J. L. Ultramicroscopy 2005, 103, 237–249. Palatinus, L.; Brázda, P.; Jelínek, M.; Hrdá, J.; Steciuk, G.; Klementová, M. Acta Crystallogr. B 2019, 75, 512–522. Schlitt, S.; Gorelik, T. E.; Stewart, A. A.; Schömer, E.; Raasch, T.; Kolb, U. Acta Crystallogr. A 2012, 68, 536–546. Petrícek, V.; Dusek, M.; Palatinus, L. Z. Kristall. 2014, 229, 345–352. Stadelmann, P. Microsc. Microanal. 2003, 9, 60–61. Koch, C. 2002. Determination of Core Structure Periodicity and Point Defect Density Along Dislocations. Ph.D. Thesis, Arizona State University, USA. Egerton, R. F. Ultramicroscopy 2014, 145, 85–93. Jiang, N. Rep. Prog. Phys. 2016, 79, 016501. Sathiya, M.; Abakumov, A. M.; Foix, D.; Rousse, G.; Ramesha, K.; Saubanère, M.; Doublet, M. L.; Vezin, H.; Laisa, C. P.; Prakash, A. S.; Gonbeau, D.; Van Tendeloo, G.; Tarascon, J.-M. Nat. Mater. 2015, 14, 230–238. Wu, Y.; Ma, C.; Yang, J.; Li, Z.; Allard, L. F.; Liang, C.; Chi, M. J. Mater. Chem. A 2015, 3, 5385–5391. Eum, D.; Kim, B.; Kim, S. J.; Park, H.; Wu, J.; Cho, S.-P.; Yoon, G.; Lee, M. H.; Jung, S.-K.; Yang, W.; Seong, W. M.; Ku, K.; Tamwattana, O.; Park, S. K.; Hwang, I.; Kang, K. Nat. Mater. 2020, 19, 419–427. Xu, B.; Fell, C. R.; Chic, M.; Shirley Meng, Y. Energy Environ. Sci. 2011, 4, 2223–2233. Boulineau, A.; Simonin, L.; Colin, J.-F.; Bourbon, C.; Patoux, S. Nano Lett. 2013, 13, 3857–3863. Pimenta, V.; Sathiya, M.; Batuk, D.; Abakumov, A. M.; Giaume, D.; Cassaignon, S.; Larcher, D.; Tarascon, J.-M. Chem. Mater. 2017, 29, 9923–9936. Gent, W. E.; Lim, K.; Liang, Y.; Li, Q.; Barnes, T.; Ahn, S.-J.; Stone, K. H.; McIntire, M.; Hong, J.; Song, J. H.; Li, Y.; Mehta, A.; Ermon, S.; Tyliszczak, T.; Kilcoyne, D.; Vine, D.; Park, J.-H.; Doo, S.-K.; Toney, M. F.; Yang, W.; Prendergast, D.; Chueh, W. C. Nat. Commun. 2017, 8, 2091.

Electrode materials viewed with transmission electron microscopy 65. 66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116. 117. 118. 119. 120. 121. 122. 123. 124. 125.

321

Sun, G.; Yu, F.-D.; Zhao, C.; Yu, R.; Farnum, S.; Shao, G.; Sun, X.; Wang, Z.-B. Adv. Funct. Mater. 2021, 31, 2002643. Lin, F.; Markus, I. M.; Doeff, M. M.; Xin, H. L. Sci. Rep. 2014, 4, 5694. Hobbs, L. W. Ultramicroscopy 1987, 23, 339–344. Shim, J.-H.; Kang, H.; Kim, Y.-M.; Lee, S. ACS Appl. Mater. Interfaces 2019, 11, 44293–44299. Kisielowski, C.; Frei, H.; Specht, P.; Sharp, I. D.; Haber, J. A.; Helveg, S. Adv. Struct. Chem. Imag. 2016, 2, 13. Zhang, X.; Zhang, X.; Yuan, B.; Liang, C.; Yu, Y. J. Phys.: Condens. Matter 2020, 32, 413004. Shim, J.-H.; Kang, H.; Lee, S.; Kim, Y.-M. J. Mater. Chem. A 2021, 9, 2429–2437. Rossell, M. D.; Erni, R.; Asta, M.; Radmilovic, V.; Dahmen, U. Phys. Rev. B 2009, 80, 024110. Wang, F.; Graetz, J.; Moreno, M. S. et al. ACS nano 2011, 5 (2), 1190–1197. Basak, S.; Jansen, J.; Kabiri, Y.; Zandbergen, H. W. Ultramicroscopy 2018, 188, 52–58. Palatinus, L.; Brazda, P.; Boullay, P.; Perez, O.; Klementova, M.; Petit, S.; Eigner, V.; Zaarour, M.; Mintova, S. Science 2017, 355, 166–169. Georgieva, D.; Jansen, J.; Sikharulidze, I.; Jiang, L.; Zandbergen, H. W.; Abrahams, J. P. J. Instrum. 2011, 6, C01033. MacLaren, I.; Macgregor, T. A.; Allen, C. S.; Kirkland, A. I. APL Mater. 2020, 8, 110901. Tyukalova, E.; Duchamp, M. J. Phys. Mater. 2020, 3, 034006. Li, Y.; Li, Y.; Pei, A.; Yan, K.; Sun, Y.; Wu, C.-L.; Joubert, L.-M.; Chin, R.; Koh, A. L.; Yu, Y.; Perrino, J.; Butz, B.; Chu, S.; Cui, Y. Science 2017, 358, 506–510. Ilett, M.; S’ari, M.; Freeman, H.; Aslam, Z.; Koniuch, N.; Afzali, M.; Cattle, J.; Hooley, R.; Roncal-Herrero, T.; Collins, S. M.; Hondow, N.; Brown, A.; Brydson, R. Phil. Trans. R. Soc. A. 2020, 378, 20190601. Islam, M. S.; Driscoll, D. J.; Fisher, C. A. J.; Slater, P. R. Chem. Mater. 2005, 17, 5085–5092. Fisher, C. A. J.; Islam, M. S. J. Mater. Chem. 2008, 18, 1209–1215. Brunetti, G.; Robert, D.; Bayle-Guillemaud, P.; Rouviere, J. L.; Rauch, E. F.; Martin, J. F.; Colin, J. F.; Bertin, F.; Cayron, C. Chem. Mater. 2011, 23, 4515–4524. Robert, D.; Douillard, T.; Boulineau, A.; Brunetti, G.; Nowakowski, P.; Venet, D.; Bayle-Guillemaud, P.; Cayron, C. ACS Nano. 2013, 7, 10887–10894. Mu, X.; Kobler, A.; Wang, D.; Chakravadhanula, V. S. K.; Schlabach, S.; Szabó, D. V.; Norby, P.; Kübel, C. Ultramicroscopy 2016, 170, 10–18. Rauch, E. F.; Véron, M. Mater. Charact. 2014, 98, 1–9. Eggeman, A. S.; Midgley, P. A. Adv. Imaging Electron. Phys. 2012, 170, 1–63. Eggeman, A. S. Acta Crystallogr. B 2019, 75, 475–484. Eggeman, A. S.; Midgley, P. A. IUCrJ 2015, 2, 126–136. Boulesteix, C.; van Landuyt, J.; Amelinckx, S. Phys. Stat. Sol. (a) 1976, 33, 595–606. Yin, W.; Grimaud, A.; Rousse, G.; Abakumov, A. M.; Senyshyn, A.; Zhang, L.; Trabesinger, S.; Iadecola, A.; Foix, D.; Giaume, D.; Tarascon, J.-M. Nat. Commun. 2020, 11, 1252. Van Tendeloo, G.; Amelinckx, S. Phase Transitions 1998, 67, 101–135. Serrano-Sevillano, J.; Reynaud, M.; Saracibar, A.; Altantzis, T.; Bals, S.; Van Tendeloo, G.; Casas-Cabanas, M. Phys. Chem. Chem. Phys. 2018, 20, 23112–23122. Batuk, D.; Batuk, M.; Abakumov, A. M.; Hadermann, J. Acta Crystallogr. B 2015, 71, 127–143. McCusker, L. B.; Baerlocher, C. Chem. Commun. 2009, (12), 1439–1451. McCusker, L. B.; Baerlocher, C. Z. Kristollgr. 2013, 228, 1–10. Fedotov, S. S.; Luchinin, N. D.; Aksyonov, D. A.; Morozov, A. V.; Ryazantsev, S. V.; Gaboardi, M.; Plaisier, J. R.; Stevenson, K. J.; Abakumov, A. M.; Antipov, E. V. Nat. Commun. 2020, 11, 1484. Fedotov, S. S.; Samarin, A. S.; Nikitina, V. A.; Aksyonov, D. A.; Sokolov, S. A.; Zhugayevych, A.; Stevenson, K. J.; Khasanova, N. R.; Abakumov, A. M.; Antipov, E. V. J. Mater. Chem. A 2018, 6, 14420–14430. Lander, L.; Rousse, G.; Abakumov, A. M.; Sougrati, M.; van Tendeloo, G.; Tarascon, J.-M. J. Mater. Chem. A 2015, 3, 19754–19764. Ati, M.; Sathiya, M.; Boulineau, S.; Reynaud, M.; Abakumov, A.; Rousse, G.; Melot, B.; Van Tendeloo, G.; Tarascon, J.-M. J. Am. Chem. Soc. 2012, 134, 18380–18387. Subban, C. V.; Ati, M.; Rousse, G.; Abakumov, A. M.; Van Tendeloo, G.; Janot, R.; Tarascon, J.-M. J. Am. Chem. Soc. 2013, 135, 3653–3661. Sun, M.; Rousse, G.; Abakumov, A. M.; Van Tendeloo, G.; Sougrati, M.-T.; Courty, M.; Doublet, M.-L.; Tarascon, J.-M. J. Am. Chem. Soc. 2014, 136, 12658–12666. Reynaud, M.; Rousse, G.; Abakumov, A. M.; Sougrati, M. T.; Van Tendeloo, G.; Chotard, J.-N.; Tarascon, J.-M. J. Mater. Chem. A 2014, 2, 2671–2680. Fedotov, S. S.; Aksyonov, D. A.; Samarin, A. S.; Karakulina, O. M.; Hadermann, J.; Stevenson, K. J.; Khasanova, N. R.; Abakumov, A. M.; Antipov, E. V. Eur. J. Inorg. Chem. 2019, 4365–4372. Sumanov, V. D.; Aksyonov, D. A.; Drozhzhin, O. A.; Presniakov, I.; Sobolev, A. V.; Glazkova, I.; Tsirlin, A. A.; Rupasov, D.; Senyshyn, A.; Kolesnik, I. V.; Stevenson, K. J.; Antipov, E.; Abakumov, A. M. Chem. Mater. 2019, 31, 5035–5046. Nishimura, S.-I.; Hayase, S.; Kanno, R.; Yashima, M.; Nakayama, N.; Yamada, A. J. Am. Chem. Soc. 2008, 130, 13212–13213. Boulineau, A.; Sirisopanaporn, S.; Dominko, R.; Armstrong, A. R.; Bruce, P. G.; Masquelier, C. Dalton Trans. 2010, 39, 6310–6316. Vainshtein, B. K.; Zvyagin, B. B.; Avilov, A. S. Electron diffraction structure analysis. In Electron Diffraction Techniques; Cowley, J. M., Ed.; vol. 1; Oxford University Press: Oxford, 1992; pp 216–312. Chapter 6. Hadermann, J.; Abakumov, A. M. Acta Crystallogr. B 2019, 75, 485–494. Burla, M. C.; Caliandro, R.; Carrozzini, B.; Cascarano, G. L.; Cuocci, C.; Giacovazzo, C.; Mallamo, M.; Mazzone, A.; Polidori, G. J. Appl. Cryst. 2015, 48, 306–309. Palatinus, L. Acta Crystallogr. A 2004, 60, 604–610. Rousse, G.; Ahouari, H.; Pomjakushin, V.; Tarascon, J.-M.; Recham, N.; Abakumov, A. M. Inorg. Chem. 2017, 56, 13132–13139. Hadermann, J.; Abakumov, A. M.; Turner, S.; Hafideddine, Z.; Khasanova, N. R.; Antipov, E. V.; Tendeloo, V. Chem. Mater. 2011, 23, 3540–3545. Mikhailova, D.; Karakulina, O. M.; Batuk, D.; Hadermann, J.; Abakumov, A. M.; Herklotz, M.; Tsirlin, A. A.; Oswald, S.; Giebeler, L.; Schmidt, M.; Eckert, J.; Knapp, M.; Ehrenberg, H. Inorg. Chem. 2016, 55, 7079–7089. Favre-Nicolin, V.; Cerný, R. J. Appl. Crystallogr. 2002, 35, 734–743. Drozhzhin, O. A.; Sumanov, V. D.; Karakulina, O. M.; Abakumov, A. M.; Hadermann, J.; Baranov, A. N.; Stevenson, K. J.; Antipov, E. V. Electrochim. Acta 2016, 191, 149–157. Fedotov, S. S.; Khasanova, N. R.; Samarin, A. S.; Drozhzhin, O. A.; Batuk, D.; Karakulina, O. M.; Hadermann, J.; Abakumov, A. M.; Antipov, E. V. Chem. Mater. 2016, 28, 411–415. Karakulina, O. M.; Khasanova, N. R.; Drozhzhin, O. A.; Tsirlin, A. A.; Hadermann, J.; Antipov, E. V.; Abakumov, A. M. Chem. Mater. 2016, 28, 7578–7581. Mugnaioli, E.; Gemmi, M.; Merlini, M.; Gregorkiewitz, M. Acta Crystallogr. B 2016, 72, 893–903. Abellan, P.; Mehdi, B. L.; Parent, L. R.; Gu, M.; Park, C.; Xu, W.; Zhang, Y.; Arslan, I.; Zhang, J.-G.; Wang, C.-M.; Evans, J. E.; Browning, N. D. Nano Lett. 2014, 14, 1293–1299. Karakulina, O. M.; Demortière, A.; Dachraoui, W.; Abakumov, A. M.; Hadermann, J. Nano Lett. 2018, 18, 6286–6291. Gorelik, T. E.; Neder, R.; Terban, M. W.; Lee, Z.; Mu, X.; Jung, C.; Jacob, T.; Kaiser, U. Acta Crystallogr. B 2019, 75, 532–549. Zhu, C.; Mu, X.; Popovic, J.; Weichert, K.; van Aken, P. A.; Yu, Y.; Maier, J. Nano Lett. 2014, 14, 5342–5349. Hu, B.; Tao, G. J. Mater. Chem. A 2015, 3, 20399–20407. Chung, S.-Y.; Choi, S.-Y.; Yamamoto, T.; Ikuhara, Y. Phys. Rev. Lett. 2008, 100, 125502.

322

Electrode materials viewed with transmission electron microscopy

126. Paolella, A.; Bertoni, G.; Hovington, P.; Feng, Z.; Flacau, R.; Prato, M.; Colombo, V.; Marras, S.; Manna, L.; Turner, S.; Van Tendeloo, G.; Guerfi, A.; Demopoulos, G. P.; Zaghib, K. Nano Energy 2015, 16, 256–267. 127. Boulineau, A.; Gutel, T. Chem. Mater. 2015, 27, 802–807. 128. Chen, J.; Vacchio, M. J.; Wang, S.; Chernova, N.; Zavalij, P. Y.; Whittingham, M. S. Solid State Ionics 2008, 178, 1676–1693. 129. Jensen, K. M.Ø.; Christensen, M.; Gunnlaugsson, H. P.; Lock, N.; Bøjesen, E. D.; Proffen, T.; Iversen, B. B. Chem. Mater. 2013, 25, 2282–2290. 130. Aksyonov, D. A.; Varlamova, I.; Trussov, I. A.; Savina, A. A.; Senyshyn, A.; Stevenson, K. J.; Abakumov, A. M.; Zhugayevych, A.; Fedotov, S. S. Inorg. Chem. 2021, 60, 5497–5506. 131. Zhang, S. S. Energy Storage Mater. 2020, 24, 247–254. 132. Yin, L.; Li, Z.; Mattei, G. S.; Zheng, J.; Zhao, W.; Omenya, F.; Fang, C.; Li, W.; Li, J.; Xie, Q.; Erickson, E. M.; Zhang, J.-G.; Whittingham, M. S.; Meng, Y. S.; Manthiram, A.; Khalifah, P. G. Chem. Mater. 2020, 32, 1002–1010. 133. Lin, Q.; Guan, W.; Meng, J.; Huang, W.; Wei, X.; Zeng, Y.; Li, J.; Zhang, Z. Nano Energy 2018, 54, 313–321. 134. Zhu, J.; Sharifi-Asl, S.; Garcia, J. C.; Iddir, H. H.; Croy, J. R.; Shahbazian-Yassar, R.; Chen, G. ACS Appl. Energy Mater. 2020, 3, 4799–4811. 135. Bals, S.; Van Aert, S.; Van Tendeloo, G.; Ávila-Brande, D. Phys. Rev. Lett. 2006, 96, 096106. 136. De Backer, A.; van den Bos, K. H. W.; Van den Broek, W.; Sijbers, J.; Van Aert, S. J. Physics: Conf. Series 2017, 902, 012013. 137. den Dekker, A. J.; Van Aert, S.; van den Bos, A.; Van Dyck, D. Ultramicroscopy 2005, 104 (2), 83–106. 138. McCalla, E.; Abakumov, A. M.; Saubanere, M.; Foix, D.; Berg, E. J.; Rousse, G.; Doublet, M.-L.; Gonbeau, D.; Novák, P.; Van Tendeloo, G.; Dominko, R.; Tarascon, J.-M. Science 2015, 350, 1516–1521. 139. Assat, G.; Tarascon, J.-M. Nat. Energy 2018, 3, 373–386. 140. Abakumov, A. M.; Fedotov, S. S.; Antipov, E. V.; Tarascon, J.-M. Nat. Commun. 2020, 11, 4976. 141. McCalla, E.; Prakash, A. S.; Berg, E.; Saubanere, V.; Abakumov, A. M.; Foix, D.; Klobes, B.; Sougrati, M.-T.; Rousse, G.; Lepoivre, F.; Mariyappan, S.; Doublet, M.-L.; Gonbeau, D.; Novak, P.; Van Tendeloo, G.; Hermann, R. P.; Tarascon, J.-M. J. Electrochem. Soc. 2015, 162, A1341–A1351. 142. Lozano, J. G.; Martinez, G. T.; Jin, L.; Nellist, P. D.; Bruce, P. G. Nano Lett. 2018, 18, 6850–6855. 143. Pearce, P. E.; Assat, G.; Iadecola, A.; Fauth, F.; Dedryvère, R.; Abakumov, A.; Rousse, G.; Tarascon, J.-M. J. Phys. Chem. C 2020, 124, 2771–2781. 144. Pearce, P. E.; Perez, A. J.; Rousse, G.; Saubanère, M.; Batuk, D.; Foix, D.; McCalla, E.; Abakumov, A. M.; Van Tendeloo, G.; Doublet, M.-L.; Tarascon, J.-M. Nat Mater. 2017, 16, 580–586. 145. Delmas, C.; Fouassier, C.; Hagenmuller, P. Physica Bþ C 1980, 99, 81–85. 146. Hwang, J.-Y.; Myung, S.-T.; Sun, Y.-K. Chem. Soc. Rev. 2017, 46, 3529–3614. 147. Berthelot, R.; Carlier, D.; Delmas, C. Nat. Mater. 2011, 10, 74–80. 148. Abakumov, A. M.; Tsirlin, A. A.; Bakaimi, I.; Van Tendeloo, G.; Lappas, A. Chem. Mater. 2014, 26, 3306–3315. 149. Moriwake, H.; Kuwabara, A.; Fisher, C. A. J.; Huang, R.; Hitosugi, T.; Ikuhara, Y. H.; Oki, H.; Ikuhara, Y. Adv. Mater. 2013, 25, 618–622. 150. Nakamura, A.; Furutsuki, S.; Nishimura, S.; Tohei, T.; Sato, Y.; Shibata, N.; Yamada, A.; Ikuhara, Y. Chem. Mater. 2014, 26, 6178–6184. 151. Zhu, C.; Gu, L.; Suo, L.; Popovic, J.; Li, H.; Ikuhara, Y.; Maier, J. Adv. Func. Mater. 2014, 24, 312–318. 152. Lu, X.; Zhao, L.; He, X.; Xiao, R.; Gu, L.; Hu, Y.-S.; Li, H.; Wang, Z.; Duan, X.; Chen, L.; Maier, J.; Ikuhara, Y. Adv. Mater. 2012, 24, 3233–3238. 153. Sharifi-Asl, S.; Yurkiv, V.; Gutierrez, A.; Cheng, M.; Balasubramanian, M.; Mashayek, F.; Croy, J.; Shahbazian-Yassar, R. Nano Lett. 2020, 20, 1208–1217. 154. Ma, C.; Chen, K.; Liang, C.; Nan, C.-W.; Ishikawa, R.; Morea, K.; Chi, M. Energy Environ. Sci. 2014, 7, 1638–1642. 155. Zhu, F.; Islam, M. S.; Zhou, L.; Gu, Z.; Liu, T.; Wang, X.; Luo, J.; Nan, C.-W.; Mo, Y.; Ma, C. Nat. Commun. 2020, 11, 1828. 156. Wang, C.; Duan, H.; Chen, C.; Wu, P.; Qi, D.; Ye, H.; Jin, H.-J.; Xin, H. L.; Du, K. Matter 2020, 3, 1999–2011. 157. Miao, J.; Ercius, P.; Billinge, S. J. L. Science 2016, 353 (aaf2157). 158. Ryabova, A. S.; Napolskiy, F. S.; Poux, T.; Istomin, S. Y.; Bonnefont, A.; Antipin, D. M.; Baranchikov, A. Y.; Levin, E. E.; Abakumov, A. M.; Kéranguéven, G.; Antipov, E. V.; Tsirlina, G. A.; Savinova, E. R. Electrochim. Acta 2016, 187, 161–172. 159. Heo, Y.; Choi, S.; Bak, J.; Kim, H.-S.; Bin Bae, H.; Chung, S.-Y. Adv. Energy Mater. 2018, 8, 1802481. 160. Gan, L.; Yu, R.; Luo, J.; Cheng, Z.; Zhu, J. J. Phys. Chem. Lett. 2012, 3, 934–938. 161. Luo, M.; Guo, S. Nat. Rev. Mater. 2017, 2, 17059. 162. Sneed, B. T.; Cullen, D. A.; Mukundan, R.; Borup, R. L.; More, K. L. J. Electrochem. Soc. 2018, 165, F3078–F3084. 163. Grothausmann, R.; Zehl, G.; Manke, I.; Fiechter, S.; Bogdanoff, P.; Dorbandt, I.; Kupsch, A.; Lange, A.; Hentschel, M. P.; Schumacher, G.; Banhart, J. J. Am. Chem. Soc. 2011, 133, 18161–18171. 164. Sneed, B. T.; Cullen, D. A.; Reeves, K. S.; Dyck, O. E.; Langlois, D. A.; Mukundan, R.; Borup, R. L.; More, K. L. ACS Appl. Mater. Interfaces 2017, 9, 29839–29848. 165. Meier, J. C.; Galeano, C.; Katsounaros, I.; Topalov, A. A.; Kostka, A.; Schüth, F.; Mayrhofer, K. J. J. ACS Catal. 2012, 2, 832–843. 166. Hengge, K.; Gänsler, T.; Pizzutilo, E.; Heinzl, C.; Beetz, M.; Mayrhofer, K. J. J.; Scheu, C. Int. J. Hydrogen Energy 2017, 42, 25359–25371. 167. Rossouw, D.; Chinchilla, L.; Kremliakova, N.; Botton, G. A. Part. Part. Syst. Charact. 2017, 34, 1700051. 168. Hu, E.; Yu, X.; Lin, R.; Bi, X.; Lu, J.; Bak, S.; Nam, K.-W.; Xin, H. L.; Jaye, C.; Fischer, D. A.; Amine, K.; Yang, X.-Q. Nat. Energy 2018, 3, 690–698. 169. Wang, C.; Han, L.; Zhang, R.; Cheng, H.; Mu, L.; Kisslinger, K.; Zou, P.; Ren, Y.; Cao, P.; Lin, F.; Xin, H. L. Matter 2021, 4, 1–14. 170. Van Aert, S.; Geuens, P.; Van Dyck, D.; Kisielowski, C.; Jinschek, J. R. Ultramicroscopy 2007, 107, 551–558. 171. Mefford, J. T.; Kurilovich, A. A.; Saunders, J.; Hardin, W. G.; Abakumov, A. M.; Forslund, R. P.; Bonnefont, A.; Dai, S.; Johnston, K. P.; Stevenson, K. J. Phys. Chem. Chem. Phys. 2019, 21, 3327–3338. 172. Genc, A.; Kovarik, L.; Gu, M.; Cheng, H.; Plachinda, P.; Pullan, L.; Freitag, B.; Wang, C. Ultramicroscopy 2013, 131, 24–32. 173. Lin, R.; Hu, E.; Liu, M.; Wang, Y.; Cheng, H.; Wu, J.; Zheng, J.-C.; Wu, Q.; Bak, S.; Tong, X.; Zhang, R.; Yang, W.; Persson, K. A.; Yu, X.; Yang, X.-Q.; Xin, H. L. Nat. Commun. 2019, 10, 1650. 174. Newbury, D. E. Scanning 2009, 31, 91–101. 175. Ritchie, N. W. M.; Newbury, D. E.; Lindstrom, A. P. Microsc. Microanal. 2011, 17, 903–910. 176. Iarchuk, A. R.; Nikitina, V. A.; Karpushkin, E. A.; Sergeyev, V. G.; Antipov, E. V.; Stevenson, K. J.; Abakumov, A. M. ChemElectroChem 2019, 6, 5090–5100. 177. Akita, T.; Taguchi, N. Surf. Interface Anal. 2016, 48, 1226–1230. 178. Goris, B.; Turner, S.; Bals, S.; Van Tendeloo, G. ACS Nano 2014, 8, 10878–10884. 179. Jarausch, K.; Thomas, P.; Leonard, D. N.; Twesten, R.; Booth, C. R. Ultramicroscopy 2009, 109, 326–337. 180. Stoyanov, E.; Langenhorst, F.; Steinle-Neumann, G. Am. Mineral. 2007, 92, 577–586. 181. Slick, L. R.; Mane, P.; Howe, J.; Zega, T. Microsc. Microanal. 2018, 24 (S1), 2086–2087. 182. Kalavathi, S.; Amirthapandian, S.; Chandra, S.; Sahu, P. C.; Sahu, H. K. J. Phys.: Condens. Matter 2014, 26, 015601. 183. Arévalo-López, Á. M.; Alario-Franco, M.Á. Inorg. Chem. 2009, 48, 11843–11846. 184. Stoyanov, E.; Langenhorst, F. Geochemistry 2014, 74, 497–505. 185. Bourdelle, F.; Lloret, E.; Durand, C.; Airaghi, L. Phys. Chem. Miner. 2021, 48, 18. 186. Zhang, S.; Livi, K. J. T.; Gaillot, A.-C.; Stone, A. T.; Veblen, D. R. Am. Miner. 2010, 95, 1741–1746. 187. Loomer, D. B.; Al, T. A.; Weaver, L.; Cogswell, S. Am. Miner. 2007, 92, 72–79. 188. van Aken, P. A.; Liebscher, B. Phys. Chem. Miner. 2002, 29, 188–200.

Electrode materials viewed with transmission electron microscopy 189. 190. 191. 192. 193. 194. 195. 196. 197. 198. 199. 200. 201. 202. 203. 204. 205. 206. 207. 208. 209. 210. 211. 212. 213. 214. 215. 216. 217. 218. 219. 220. 221. 222. 223. 224. 225. 226. 227. 228. 229. 230. 231. 232. 233. 234. 235. 236. 237. 238. 239. 240. 241. 242. 243. 244. 245. 246. 247. 248.

323

Cavé, L.; Al, T.; Loomer, D.; Cogswell, S.; Weaver, L. Micron 2006, 37, 301–309. Liu, H.; Bugnet, M.; Tessaro, M. Z.; Harris, K. J.; Dunham, M. J. R.; Jiang, M.; Goward, G. R.; Botton, G. A. Phys. Chem. Chem. Phys. 2016, 18, 29064–29075. Honda, Y.; Muto, S.; Tatsumi, K.; Kondo, H.; Horibuchi, K.; Kobayashi, T.; Sasaki, T. J. Power Sources 2015, 291, 85–94. Yuan, Y.; Amine, K.; Lu, J.; Shahbazian-Yassar, R. Nat. Commun. 2017, 8, 15806. Holtz, M. E.; Yu, Y.; Gao, J.; Abruña, H. D.; Muller, D. A. Microsc. Microanal. 2013, 19, 1027–1035. Holtz, M. E.; Yu, Y.; Gunceler, D.; Gao, J.; Sundararaman, R.; Schwarz, K. A.; Arias, T. A.; Abruna, H. D.; Muller, D. A. Nano Lett. 2014, 14, 1453–1459. Schneider, N. M.; Norton, M. M.; Mendel, B. J.; Grogan, J. M.; Ross, F. M.; Bau, H. H. J. Phys. Chem. C 2014, 118, 22373–22382. Woehl, T. J.; Jungjohann, K. L.; Evans, J. E.; Arslan, I.; Ristenpart, W. D.; Browning, N. D. Ultramicroscopy 2013, 127, 53–63. Williamson, M. J.; Tromp, R. M.; Vereecken, P. M.; Hull, R.; Ross, F. M. Nat. Mater. 2003, 2, 532–536. Radisic, A.; Vereecken, P. M.; Hannon, J. B.; Searson, P. C.; Ross, F. M. Nano Lett. 2006, 6, 238–242. White, E. R.; Singer, S. B.; Augustyn, V.; Hubbard, W. A.; Mecklenburg, M.; Dunn, B.; Regan, B. C. ACS Nano 2012, 6, 6308–6317. Chen, X.; Noh, K. W.; Wen, J. G.; Dillon, S. J. Acta Mater. 2012, 60, 192–198. Schneider, N. M.; Park, J. H.; Grogan, J. M.; Steingart, D. A.; Bau, H. H.; Ross, F. M. Nat. Commun. 2017, 8, 2174. Leenheer, A. J.; Jungjohann, K. L.; Zavadil, K. R.; Sullivan, J. P.; Harris, C. T. ACS Nano 2015, 9, 4379–4389. Mehdi, B. L.; Qian, J.; Nasybulin, E.; Park, C.; Welch, D. A.; Faller, R.; Mehta, H.; Henderson, W. A.; Xu, W.; Wang, C. M.; Evans, J. E.; Liu, J.; Zhang, J.-G.; Mueller, K. T.; Browning, N. D. Nano Lett. 2015, 15, 2168–2173. Zeng, Z.; Liang, W.-I.; Liao, H.-G.; Xin, H. L.; Chu, Y.-H.; Zheng, H. Nano Lett. 2014, 14, 1745–1750. Unocic, R. R.; Sun, X.-G.; Sacci, R. L.; Adamczyk, L. A.; Alsem, D. H.; Dai, S.; Dudney, N. J.; More, K. L. Microsc. Microanal. 2014, 20, 1029–1037. Chee, S. W.; Tan, S. F.; Baraissov, Z.; Bosman, M.; Mirsaidov, U. Nat. Commun. 2017, 8, 1224. Grogan, J. M.; Schneider, N. M.; Ross, F. M.; Bau, H. H. Nano Lett. 2014, 14, 359–364. Woehl, T. J.; Abellan, P. J. Microsc. 2017, 265, 135–147. de Jonge, N.; Ross, F. M. Nat. Nanotech. 2011, 6, 695–704. Pu, S.; Gong, C.; Robertson, A. W. R. Soc. Open Sci. 2020, 7, 191204. de Jonge, N.; Houben, L.; Dunin-Borkowski, R. E.; Ross, F. M. Nat. Rev. Mater. 2019, 4, 61–78. Koo, K.; Park, J.; Ji, S.; Toleukhanova, S.; Yuk, J. M. Adv. Mater. 2021, 33, 2005468. Park, J.; Koo, K.; Noh, N.; Chang, J. H.; Cheong, J. Y.; Dae, K. S.; Park, J. S.; Ji, S.; Kim, I.-D.; Yuk, J. M. ACS Nano 2021, 15, 288–308. Yuk, J. M.; Seo, H. K.; Choi, J. W.; Lee, J. Y. ACS Nano 2014, 8, 7478–7485. Seo, H. K.; Park, J. Y.; Chang, J. H.; Dae, K. S.; Noh, M. S.; Kim, S. S.; Kang, C. Y.; Zhao, K.; Kim, S.; Yuk, J. M. Nat. Commun. 2019, 10, 3428. Chang, J. H.; Cheong, J. Y.; Yuk, J. M.; Kim, C.; Kim, S. J.; Seo, H. K.; Kim, I. D.; Lee, J. Y. ACS Omega 2017, 2, 6329–6336. Kalinin, S. V.; Dyck, O.; Balke, N.; Neumayer, S.; Tsai, W.-Y.; Vasudevan, R.; Lingerfelt, D.; Ahmadi, M.; Ziatdinov, M.; McDowell, M. T.; Strelcov, E. ACS Nano 2019, 13, 9735–9780. Fahrenkrug, E.; Alsem, D. H.; Salmon, N.; Maldonado, S. J. Electrochem. Soc. 2017, 164, H358–H364. Zhu, Y.; Wang, J. W.; Liu, Y.; Liu, X.; Kushima, A.; Liu, Y.; Xu, Y.; Mao, S. X.; Li, J.; Wang, C.; Huang, J. Y. Adv. Mater. 2013, 25, 5461–5466. Liu, X. H.; Wang, J. W.; Huang, S.; Fan, F.; Huang, X.; Liu, Y.; Krylyuk, S.; Yoo, J.; Dayeh, S. A.; Davydov, A. V.; Mao, S. X.; Picraux, S. T.; Zhang, S.; Li, J.; Zhu, T.; Huang, J. Y. Nat. Nanotech. 2012, 7, 749–756. Liu, X. H.; Zheng, H.; Zhong, L.; Huang, S.; Karki, K.; Zhang, L. Q.; Liu, Y.; Kushima, A.; Liang, W. T.; Wang, J. W.; Cho, J.-H.; Epstein, E.; Dayeh, S. A.; Picraux, S. T.; Zhu, T.; Li, J.; Sullivan, J. P.; Cumings, J.; Wang, C.; Mao, S. X.; Ye, Z. Z.; Zhang, S.; Huang, J. Y. Nano Lett. 2011, 11, 3312–3318. Liu, X. H.; Zhong, L.; Huang, S.; Mao, S. X.; Zhu, T.; Huang, J. Y. ACS Nano 2012, 6, 1522–1531. Wang, Z.; Gu, M.; Zhou, Y.; Zu, X.; Connell, J. G.; Xiao, J.; Perea, D.; Lauhon, L. J.; Bang, J.; Zhang, S.; Wang, C.; Gao, F. Nano Lett. 2013, 13, 4511–4516. Gong, Y.; Zhang, J.; Jiang, L.; Shi, J.-A.; Zhang, Q.; Yang, Z.; Zou, D.; Wang, J.; Yu, X.; Xiao, R.; Hu, Y.-S.; Gu, L.; Li, H.; Chen, L. J. Am. Chem. Soc. 2017, 139, 4274–4277. Li, S.; Yao, Z.; Zheng, J.; Fu, M.; Cen, J.; Hwang, S.; Jin, H.; Orlov, A.; Gu, L.; Wang, S.; Chen, Z.; Su, D. Angew. Chem. Int. Ed. 2020, 59, 22092–22099. Gong, Y.; Chen, Y.; Zhang, Q.; Meng, F.; Shi, J.-A.; Liu, X.; Liu, X.; Zhang, J.; Wang, H.; Wang, J.; Yu, Q.; Zhang, Z.; Xu, Q.; Xiao, R.; Hu, Y.-S.; Gu, L.; Li, H.; Huang, X.; Chen, L. Nat. Commun. 2018, 9, 3341. Nomura, Y.; Yamamoto, K.; Hirayama, T.; Igaki, E.; Saitoh, K. ACS Energy Lett. 2020, 5, 2098–2105. Yamamoto, K.; Iriyama, Y.; Asaka, T.; Hirayama, T.; Fujita, H.; Fisher, C. A. J.; Nonaka, K.; Sugita, Y.; Ogumi, Z. Angew. Chem. Int. Ed. 2010, 49, 4414–4417. Nomura, Y.; Yamamoto, K.; Hirayama, T.; Ohkawa, M.; Igaki, E.; Hojo, N.; Saitoh, K. Nano Lett. 2018, 18, 5892–5898. Yang, Y.; Xiong, Y.; Zeng, R.; Lu, X.; Krumov, M.; Huang, X.; Xu, W.; Wang, H.; DiSalvo, F. J.; Brock, J. D.; Muller, D. A.; Abruña, H. D. ACS Catal. 2021, 11, 1136–1178. Hwang, S.; Chen, X.; Zhou, G.; Su, D. Adv. Energy Mater. 2020, 10, 1902105. Mayrhofer, K. J. J.; Ashton, S. J.; Meier, J. C.; Wiberg, G. K. H.; Hanzlik, M.; Arenz, M. J. Power Sources 2008, 185, 734–739. Hodnik, N.; Cherevko, S. Curr. Opin. Electrochem. 2019, 15, 73–82. Meier, J. C.; Katsounaros, I.; Galeano, C.; Bongard, H. J.; Topalov, A. A.; Kostka, A.; Karschin, A.; Schuth, F.; Mayrhofer, K. J. J. Energy Environ. Sci. 2012, 5, 9319–9330. Moriau, L.; Hrnjic, A.; Pavlisic, A.; Kamsek, A. R.; Petek, U.; Ruiz-Zepeda, F.; Sala, M.; Pavko, L.; Selih, V. S.; Bele, M.; Jovanovic, P.; Gatalo, M.; Hodnik, N. iScience 2021, 24 (102102). Krivanek, O. L.; Dellby, N.; Murfitt, M. F.; Chisholm, M. F.; Pennycook, T. J.; Suenaga, K.; Nicolosi, V. Ultramicroscopy 2010, 110, 935–945. Kaiser, U.; Stöger-Pollach, M. Ultramicroscopy 2014, 145, 1–2. Bell, D. C.; Mankin, M.; Day, R. W.; Erdman, N. Ultramicroscopy 2014, 145, 56–65. Sasaki, T.; Sawada, H.; Hosokawa, F.; Sato, Y.; Suenaga, K. Ultramicroscopy 2014, 145, 50–55. Krivanek, O. L.; Dellby, N.; Hachtel, J. A.; Idrobo, J.-C.; Hotz, M. T.; Plotkin-Swing, B.; Bacon, N. J.; Bleloch, A. L.; Corbin, G. J.; Hoffman, M. V.; Meyer, C. E.; Lovejoy, T. C. Ultramicroscopy 2019, 203, 60–67. Goris, B.; De Beenhouwer, J.; De Backer, A.; Zanaga, D.; Joost Batenburg, K.; Sanchez-Iglesias, A.; Liz-Marzan, L. M.; Van Aert, S.; Bals, S.; Sijbers, J.; Van Tendeloo, G. Nano Lett. 2015, 15, 6996–7001. Müller, K.; Ryll, H.; Ordavo, I.; Ihle, S.; Struder, L.; Volz, K.; Zweck, J.; Soltau, H.; Rosenauer, A. Appl. Phys. Lett. 2012, 101, 212110. Hachtel, J. A.; Idrobo, J. C.; Chi, M. Adv. Struct. Chem. Imag. 2018, 4, 10. Xu, C.; Märker, K.; Lee, J.; Mahadevegowda, A.; Reeves, P. J.; Day, S. J.; Groh, M. F.; Emge, S. P.; Ducati, C.; Mehdi, B. L.; Tang, C. C.; Grey, C. P. Nat. Mater. 2021, 20, 84–92. Mao, Y.; Wang, X.; Xia, S.; Zhang, K.; Wei, C.; Bak, S.; Shadike, Z.; Liu, X.; Yang, Y.; Xu, R.; Pianetta, P.; Ermon, S.; Stavitski, E.; Zhao, K.; Xu, Z.; Lin, F.; Yang, X.-Q.; Hu, E.; Liu, Y. Adv. Funct. Mater. 2019, 29, 1900247. Romano Brandt, L.; Marie, J.-J.; Moxham, T.; Forstermann, D. P.; Salvati, E.; Besnard, C.; Papadaki, C.; Wang, Z.; Bruce, P. G.; Korsunsky, A. M. Energy Environ. Sci. 2020, 13, 3556–3566. Xu, W.; LeBeau, J. M. Ultramicroscopy 2018, 188, 59–69. Muto, S.; Shiga, M. Microscopy 2020, 69, 110–122.

7.11

Chemistry of Li-air batteries

Alina Inozemtsevaa,b, Alexey Ruleva,b, Tatiana Zakharchenkoa,b, Valerii Isaeva,b, Lada Yashinaa,b, and Daniil Itkisa,b, a Lomonosov Moscow State University, Moscow, Russia; and b N.N. Semenov Federal Research Center for Chemical Physics, Moscow, Russia © 2023 Elsevier Ltd. All rights reserved.

7.11.1 7.11.2 7.11.2.1 7.11.2.1.1 7.11.2.1.2 7.11.2.1.3 7.11.2.1.4 7.11.2.2 7.11.2.3 7.11.3 7.11.3.1 7.11.3.2 7.11.3.3 7.11.3.3.1 7.11.3.3.2 7.11.3.3.3 7.11.3.3.4 7.11.4 7.11.4.1 7.11.4.2 7.11.5 7.11.5.1 7.11.5.2 7.11.5.3 7.11.5.4 7.11.5.5 7.11.6 References

Introduction Positive electrode Discharge process General reaction pathway Superoxide anion formation and solvation. Surface and solution-mediated mechanism Fundamental aspects of lithium peroxide crystallization Lithium peroxide deposition in porous electrode Charge process Heterogeneous ORR/OER catalysts Negative electrode Many shapes of lithium Mechanisms of morphological instability Steps toward uniform deposition SEI design Electrolyte design Electrode design Alternative anode materials Reactions with reactive oxygen species Electrolyte decomposition Carbon electrode degradation Redox mediators Basic principles Redox mediators for charge “Shuttle effect” Redox mediators for discharge Bifunctional and dual mediators Concluding remarks and future prospects

324 327 327 327 328 330 333 334 335 336 338 339 341 341 342 342 342 342 342 343 344 344 345 346 346 346 346 352

Abstract Transition to the renewable energy sources (solar, wind energies, etc.) and electrified transport is needed to improve our environment. For these purpose long-lived, safe, and affordable rechargeable batteries are of demand. Metal-air batteries, and especially lithium-air batteries with aprotic electrolytes promise very high specific energy in comparison to the LIBs and thus gained lots of attention in recent decades. The lithium-air battery consists of a lithium anode, a porous air cathode, and an electrolyte composed of a lithium salt dissolved in nonaqueous, organic solvent. In 1996, Abraham et al. reported the first rechargeable nonaqueous lithium-O2 battery, that laid the foundation for a new field and other analogous systems such as Na-O2 and K-O2 batteries. Significant efforts were undertaken a decade ago to commercialize Li-air batteries, which failed but prompted the research aimed at a deeper understanding of the unclear underpinning chemistry and electrochemistry. Extensive studies uncovered fundamental obstacles limiting the further development of nonaqueous Li-O2 batteries, including limited capacity, electrode passivation and degradation, poor cycling and safety issues related to Li metal anode. This chapter is aimed to shed the light on the diversity of mechanisms and elementary processes occurring in this system.

7.11.1

Introduction

Nowadays fossil fuels satisfy more than 80% of the global energy demand,1 that leads to a huge release of carbon dioxide and other greenhouse gases and drives the global climate change. A third of greenhouse gases is released from the petroleum used in automobile applications. Therefore, transition to the renewable energy sources (solar, wind energies, etc.) and electrified (or hybrid) transportation systems is needed to improve our environment and energy security. For these purpose efficient energy storage devices, such as long-lived, safe, and affordable rechargeable batteries, are of demand. The utilization of the batteries in electric (EV) or

324

Comprehensive Inorganic Chemistry III, Volume 7

https://doi.org/10.1016/B978-0-12-823144-9.00055-8

Chemistry of Li-air batteries

325

hybrid electric vehicles (HEV) requires them to be light and compact, and have sufficient power and energy densities to cover most driving range scenarios for a day.2 In the 1990s, lead-acid batteries were used in electric vehicles, but their applications were limited by relatively low energy density.3 Besides, nickel-cadmium (Ni-Cd), nickel-metal hydride (Ni-MH), lithium-polymer, sodium-sulfur, and sodium-metal chloride battery systems were proposed for vehicle propulsion applications.4,5 Most of today’s EVs, especially cars, utilize lithium-ion batteries (LIBs).6 However, lithium-ion battery limits the EV’s driving range to 160 km on a single charge and the battery accounts for nearly 65% of the total cost.7 It is known that the theoretical energy density of gasoline is 13,000 Wh kg
1 and the energy density of lithium-ion batteries is around 100–200 Wh kg
1. If the tank-to-wheel energy conversion efficiency of 12.6% is taken into account, the practical energy density of gasoline is 1700 Wh kg
1, which is still much higher than that of lithium-ion batteries as shown in Fig. 1. Therefore, novel energy systems with higher energy densities are urgently needed.9 As seen from Fig. 1, metal-air batteries10,11 promise a higher specific energy in comparison to the LIBs and thus gained the most attention in recent decades.12 As metal-air batteries comprise a metal anode and an oxygen electrode, they combine the features of both batteries and fuel cells. To date, metal-air batteries using alkali metals (Li, Na, and K), alkaline earth metal (Mg), and first-row transition metals (Fe, Zn) or Al as the anode have been reported and their theoretical specific energies are presented in Fig. 2. Among them, the Li-air battery (LAB or Li-O2) system possesses the highest theoretical energy density, and is one of the most studied class of metal-air batteries. The Li-O2 battery comprises a metallic lithium negative electrode (anode), air-O2 as the positive electrode (cathode) active mass, and a Liþ-containing electrolyte solution. The cathode in this system is basically an electronically conducting porous matrix that enables the electrochemical contact between electrode, dissolved oxygen gas and Liþ ions from the electrolyte solution. A solid discharge product is deposited in the porous cathode. It primarily functions with the dissolution/deposition of lithium metal at the anode and an oxygen reduction reaction (ORR)/oxygen evolution reaction (OER) at the cathode. To date four types of Li-O2 systems under investigation exist, referred to as aprotic, aqueous, hybrid, and solid-state batteries. These types differ from each other in the electrolytes used. This, in turn, determines the specific electrochemical reactions during energy storage and release. A schematic illustration of these four types of batteries is provided in Fig. 3 including liquid electrolytes

Fig. 1 The gravimetric energy density of some representative types of primary/rechargeable batteries, metal-air batteries, H2–air fuel cell and gasoline. The theoretical values are calculated on the basis of thermodynamics of active materials.8

Fig. 2 Theoretical energy densities for different types of metal-air batteries. Reprinted with permission from Li Y, Lu J. Metal–Air Batteries: Will they be the Future Electrochemical Energy Storage Device of Choice? ACS Energy Lett. 2017, 2, 1370–1377, doi:10.1021/acsenergylett.7b00119. Copyright 2017 American Chemical Society.

326

Chemistry of Li-air batteries

Fig. 3 Schematic configurations of lithium-air battery types. Reprinted from Tan P, Jiang HR, Zhu XB, An L, Jung CY, Wu MC, et al. Advances and Challenges in Lithium-Air Batteries. Appl. Energy 2017, 204, 780–806, doi:10.1016/j.apenergy.2017.07.054. Copyright 2017, with permission from Elsevier.

(Fig. 3A–C) and solid ones (Fig. 3D). The all-solid-state lithium/air battery (Fig. 3D) was first developed by Kumar and coworkers.13,14 The electrolyte used in this case was based on lithium aluminum germanium phosphate (LAGP) mixed with polyethylene oxide (PEO). This battery operated at the temperature of 75  C and exhibited a high discharge and a low charge potentials (2.8 V and 3.6 V, respectively). A lithium phosphorous oxynitride (LiPON) electrolyte was also employed by Lu and coworkers.15,16 The scheme of aqueous Li-air battery is shown in Fig. 3B. It should be mentioned that researches on aqueous metal-air batteries appeared much earlier than lithium-ion batteries.17 The first primary zinc-air battery was designed by Maiche in 1878,18 later aqueous iron-air, aluminum-air, and magnesium-air batteries were developed in 1960s.19–21 However, most of the aqueous metal-air systems function as primary cells and can only be mechanically recharged by replacing the used metal anodes and the electrolyte solution. On the contrary, Li-O2 aqueous battery is based on reversible reaction and could be recharged electrochemically. Aqueous lithium-air batteries have many advantages compared to the non-aqueous systems, namely a high discharge potential, a high round-trip efficiency, the absence of solid discharge products that cause electrode pore clogging. In this battery, the cathode operates in an aqueous electrolyte separated from the Li anode by additional inorganic and/or organic electrolyte layers. This type of battery was first investigated in the 1970s22 where the lithium anode was immersed into a concentrated LiOH aqueous solution. During the discharge, LiOH partially soluble in the electrolyte have been formed. LiOH then precipitates with crystalline water when its concentration exceeds 5.2 M, therefore aqueous Li-air cells face an intrinsic limitation of 477 Wh kg
1 in theoretical specific energy,23 although they might compete with their non-aqueous counterpart in volumetric terms.24

Chemistry of Li-air batteries

327

Another attempt to implement an aqueous electrolyte brought the hybrid lithium-air battery into existence.25,26 In the proposed design, the anode side of the battery was filled with an aprotic solvent, whereas the air electrode was in contact with an aqueous electrolyte. The two electrolyte solutions were decoupled by a solid electrolyte membrane. The hybrid lithium-air battery avoided direct contact between the lithium and the solid electrolyte membrane and indeed resulted in improved Liþ conductivity. However, the disadvantage of such battery is complicated design and different diffusion kinetics in the two electrolyte solutions. The nonaqueous Li-O2 battery system with Li2O2 discharge product has gained attention due to its high theoretical specific energy. The battery consists of a lithium anode, a porous air cathode, and an electrolyte composed of a lithium salt dissolved in nonaqueous, organic solvent (Fig. 3A). In 1996, Abraham et al.27 reported the first rechargeable nonaqueous lithium-O2 battery, that laid the foundation for a new field and other analogous systems such as Na-O2 and K-O2 batteries. Among these, the Li-O2 battery is the most attractive because it offers the highest operational voltage ( 2.96 V) with a theoretical specific energy of  3500 Wh kg
1 (based on Li2O2formation on discharge). Significant efforts were undertaken a decade ago to commercialize Li-air batteries, which failed but prompted the research aimed at a deeper understanding of the unclear underpinning chemistry and electrochemistry. Extensive recent studies uncovered certain fundamental obstacles limiting the further development of nonaqueous Li-O2 batteries, including limited practical capacity, electrode passivation and degradation, poor cycling and also safety issues related to Li metal anode. Different aspects of the Li-O2 fundamentals are already reviewed in the excellent publications.2,12,28–35 Our chapter is focused to shed the light on the diversity of mechanisms and elementary processes occurring in this system and describes electrode processes occurring under discharge and recharge.

7.11.2

Positive electrode

7.11.2.1

Discharge process

7.11.2.1.1

General reaction pathway

Oxygen reduction reaction (ORR) mechanism was extensively studied many years before Li-O2 battery concept. In aprotic media in the absence of lithium ions oxygen is reduced to superoxide anion,36–41 according to the reaction: O2 þ e
/O2
37,41

as was confirmed by the EPR spectroscopy. Superoxide anion can be reduced further to peroxide anion at higher overpotential at the mercury electrode.40 Electrochemical and chemical reactions occurring during oxygen reduction in aprotic electrolytes in the presence of lithium ions, as well as in a Liþ-containing melt, are listed in Table 1. In the presence of lithium ions in the electrolyte, the superoxide anions form ion pairs with lithium ions according to the Eq. (1.3) in Table 1. Then Liþ$O2
ion pairs can be electrochemically reduced to Li2O2 or disproportionate to form Li2O2 and O2. The latter reaction can occur both on the electrode surface and in the electrolyte bulk.59 It was established that the first electron transfer step from carbon electrode to oxygen molecule in presence of Liþ ions does not involve oxygen chemisorption on the electrode surface43 and is analogous to the reaction 1.1 (Table 1) in aprotic media containing no Liþ ions.60,61 On the contrary, the second electron transfer (1.4) requires the preliminary adsorption step (1.2).43,45 The molecular dynamics (MD) simulations support this idea, showing that Liþ$O
2 pair could be reduced on the electrode surface according to the reaction (1.4), while the reduction of free superoxide anion is unlikely.47 The presence of superoxide species in Liþ-containing

Table 1 Product

Electrochemical and chemical processes in Liþ-containing electrolyte under ORR conditions. N

Reaction

O2 þ e
/ O2
(sol) O2
(sol) þ Liþ(sol) / Liþ$O2
(ads) Liþ(sol) þ O2
(sol) / Liþ$O2
(sol) Liþ$O2
(ads) þ Liþ þ e
/ Li2O2 2 Liþ$O2
(ads) / Li2O2 þ O2 2 Liþ$O2
(sol) / Li2O2 þ O2 O2 þ 2e
þ 2Liþ / Li2O2 O2 þ H2O þ 2e
/ HO2
þ OH
HO2
þ 2OH
þ 2Liþ / Li2O2 þ 2H2O 1.9 2Li2O2 þ H2O / 4LiOH þ H2O2 2Li2O2 þ H2O / 4LiOH þ O2 1.10 2Li2O2 þ 2Liþ þ2e
/ 2Li2O 1.11 Li2O2 þ 2Liþþ2O2
/ 2Li2O þ 2O2 1.12 2Li2O2 / 2Li2O þ O2

References

Superoxide species 1.1 1.2 1.3 Li2O2 1.4 1.5 1.6 1.7 1.8

42,43 44,45 46 43,45,47 45,46 45,46,48 49 50–52

LiOH

53–55

Li2O

56,57 58 49

328

Chemistry of Li-air batteries

electrolyte solution under ORR conditions was confirmed by UV–visible42,62 and electron spin resonance (ESR) spectroscopy.63 The surface-enhanced Raman scattering (SERS) experiments also detected the superoxide species on the electrode surface.44,45 In the presence of Liþ superoxide anions undergo disproportionation (reaction 1.6 in Table 1) with the formation of Li2O2, since solid LiO2 is thermodynamically unstable at room temperature.64 The precipitate mainly consisted of Li2O2 is formed upon the addition of KO2 to Liþ-containing solutions in various aprotic solvents.65,66 Kinetics of this reaction was studied by ESR48 and UV–visible spectroscopy.45,46 Reaction 1.6 was shown to be first-order with respect for both superoxide ions and Liþ46 Therefore, one can suggest that the rate-limiting step of the disproportionation is the formation of a contact ion pair (CIP) Liþ-O
2 (1.3). However, some works report that the superoxide concentration does not change when a lithium salt is added to its solution,42 although it might be due to the low concentration of O
2 in the experiment. Li2O2 could also be formed by direct two-electron transfer (1.7) in the molten electrolyte,49 but in solution it is unlikely due to the high energy barrier compared to one-electron transfer.67 Two-electron reduction of O2 in Liþ-containing solutions is facilitated in aqueous/nonaqueous mixtures68 or in aqueous alkaline solutions.51,52 In this case, O2 is reduced to hydroperoxide species HO2
instead of O2
, and Li2O2 is formed in chemical reactions (1.8). Typically, hydrolysis of Li2O2 takes place as described by the Eq. (1.9); however, unexpectedly it had been found that the hydrolysis can be remarkably restrained in saturated aqueous LiOH solution.52 Li2O2 undergo hydrolysis (1.9) in nonaqueous electrolyte as well when it is humid, and LiOH crystallization is observed upon discharge.53,54 Several studies also report the formation of lithium oxide under ORR conditions. Two different mechanisms have been proposed: electrochemical reduction of Li2O2 (1.10)56,57 and somproportionation of Li2O2 and the superoxide anion (1.11).58 The formation of Li2O was also observed during the discharge at 150  C in eutectic melt of LiNO3/KNO3 as an electrolyte. At this temperature D G(Li2O2) > D G(Li2O), and the main discharge product is Li2O formed via Li2O2 disproportionation (1.12). Overall, the ORR mechanism is strongly influenced by the electrolyte solvent,45,69,70 the electrolyte salt,71 as well as electrolyte additives72–74 and the presence of a catalyst on the electrode surface.75–77 Also, the cathode overpotential45,70,78 and operation temperature79,80 play an important role in oxygen reduction processes.

7.11.2.1.2

Superoxide anion formation and solvation. Surface and solution-mediated mechanism

The mechanism of Li-O2 battery discharge is strongly influenced by the electrolyte composition. Laoire et al.56 adopted the concept of hard-soft acid-base (HSAB) theory to explore various electrolyte systems to determine the solvation effects that has been invoked many times since then.45,81,82 According to the HSAB theory, Liþ is a hard acid and since superoxide is a moderately soft base, it has a low affinity for the hard acid Liþ present in Liþ-conducting electrolytes. Consequently, the superoxide formed as the first reduction product of O2 either decompose or undergo a fast second reduction to form the hard base, peroxide (O22
) which readily forms Li2O2 with a hard acid Liþ according to equations 1.4, 1.5 in Table 1. Although Liþ behaves as a hard acid, its acidity can be diminished by coordination with solvent molecule in Liþ-(solvent)n. However, Baltruschat et al. pointed out widespread misconception that interaction between soft bases and acids are always stronger than interactions between soft bases and hard acids.83,84 This is because concept of hardness/softness of an ion does not replace the concept of weak and strong Lewis acids.85 For instance, the interaction between Hþ and S2
is certainly stronger than the interaction between Hþ and Cl
despite the former being a hardsoft acid-base pair and the latter being hard-hard. Once O2 is reduced to superoxide, a strong Lewis base is present in the electrolyte. It may strongly interact with Liþ, who is a strong Lewis acid, this leads to the superoxide disproportionation. In contrast to that, the interaction between the weak Lewis acid tetrabuthylammonium cation (TBAþ) and O2
is much weaker, and thus O2
is not as much destabilized as in the case of Liþ. For strongly solvation solvents such as dimethylsulfoxide (DMSO), the Lewis acidity of the cation is lower, resulting in a weaker interaction and, thus, an increase of superoxide species lifetime. Johnson et al. studied the ORR mechanism in electrolyte solvents with different Liþ solvation ability that is thought to be connected with the solvent donor number (DN).45 Gutmann DN is enthalpy of dissolution typical Lewis acid SbCl5 in the corresponding solvent.86 Kwabi et al. showed the strict correlation between Gutmann DN and calculated solvation energy for the series of conventional Li-O2 battery electrolyte solvents.69 However, further inspection reveals that for a wider set of solvents presented in Fig. 4A the correlation between experimental solvation energy87 and Gutmann DN is not sufficiently solid. It should be taken into account that both experimental and calculated Liþ solvation energies have serious uncertainties. In high-DN solvents (DMSO, 1-methylimidazolium, etc.) Liþ is strongly coordinated by the solvent molecules, and forms solvent separated ionic pairs (SSIPs) with O2
that prevents their fast disproportionation to Li2O2. Such SSIPs diffuse from the electrode surface to the electrolyte bulk and can be detected experimentally on the ring in rotating ring-disk electrode (RRDE) voltammetry. In solution bulk the chemical disproportion occurs, and Li2O2 crystals nucleate homogeneously.88 In low-DN solvents (acetonitrile) Liþ is weakly solvated and form contact ionic pairs (CIPs) with O2
which undergo fast disproportionation (1.5). Alternatively, the second electron transfer from the electrode to superoxide species can occur (1.4), both processes giving rise to Li2O2 layer on the electrode surface.45 These processes are schematically shown in Fig. 4B. Using mixtures of solvents with different Liþ solvation abilities, it is possible to tune the amount of mobile superoxide species in the electrolyte bulk. By means of RRDE voltammetry it was demonstrated that with an increase of DMSO (high-DN solvent) concentration in the mixture with tetraethylene glycol dimethyl ether (TEGDME) (low-DN solvent), the concentration of superoxide species in the solution proportionally increases and reaches its maximum in pure DMSO.89 The stability of Liþ$O2
ion pairs can be also tuned by strong solvation of O2
. The latter is realized in solvents with high acceptor number (AN). This phenomenon was observed when water (AN ¼ 55) and methanol (AN ¼ 42) were added to the electrolyte based on 1,2-dimethoxyethane (DME) (AN ¼ 10).90 Another approach to suppress the Li-O2 CIP formation is usage of

Chemistry of Li-air batteries

Li+ solvation energy, kJ/mol

(A)

329

(B)

–450 –470 –490 –510 10

15

20

25

30

35

Donor number, kcal/mol Fig. 4 (A) Experimental Liþ solvation energy87 plotted against Gutmann solvent donor number (B) Oxygen reduction mechanistic pathways in nonaqueous Li-O2 batteries. (B) Adapted from Johnson L, Li C, Liu Z, Chen Y, Freunberger SA, Ashok PC, et al. The Rrole of LiO2 Ssolubility in O2 Rreduction in Aaprotic Ssolvents and its Cconsequences for Li–O2 Bbatteries. Nat. Chem.Nat Chem 2014, 6, 1091–1099, doi:10.1038/nchem.2101.

Li-containing salts with low dissociation level (e.g., LiNO3 instead of LiTFSI)71,91 that decreases the amount of free Liþ ions in solution; but this concept is relevant only for low-DN solvents. Alternative explanation of the higher Liþ$O
2 SSIPs stability in DMSO than in acetonitrile was proposed by Kislenko et al. Using molecular dynamics (MD) simulations, they predicted that the solvation shell of Liþ in DMSO is more stable than in acetonitrile due to the lock-and-key steric factor.92 Unfortunately, similar calculations were not performed yet for other conventional solvents to establish quantitative correlation and make general conclusions. Moreover, solvent’s DN and AN affect the value of the standard potential of lithium oxidation and oxygen reduction, respecþ tively. Kwabi et al. measured the standard potentials of the O2/O
2 and Li /Li redox in acetonitrile, DME, dimethylacetamide 69 (DMA), and DMSO (TBAClO4 was used as supporting electrolyte). Standard potential O2/TBAþ-O2
increases with increasing solvent’s AN, while Liþ/Li potential decreases with decreasing DN, which is associated with an increase in the solvation energy of Liþ and O2
ions, respectively (Fig. 5). The effect of solvent on the potential of the reaction (1.4) in Table 1 (reduction of lithium superoxide to peroxide) in DMSO mixtures in various solvents was studied in Ref. [83]. It was found that the potential is influenced by the additive solvent’s AN, but the best correlation is observed with the normalized shift of Betaine-30 absorption band (EsN) in the additive solvent. This observation is consistent with the Marcus-Hush theory which shows the major contribution of the reorganization energy to the solvent activation barrier for the electron transfer from an electrode to a reagent,93 which can be evaluated as EsN.

Fig. 5 Experimental standard O2/TBAþ-O2
and Liþ/Li redox potentials vs. Me10Fcþ/Me10Fc plotted against acceptor and donor numbers of each solvent. Reproduced from Kwabi DG, Bryantsev VS, Batcho TP, Itkis DM, Thompson CV, Shao-Horn Y. Experimental and Computational Analysis of the Solvent-Dependent O2/Liþ-O2
Redox Couple: Standard Potentials, Coupling Strength, and Implications for Lithium–Oxygen Batteries. Angew. Chem. Int. Ed. 2016, 55, 3129–3134, doi:10.1002/anie.201509143 by permission of John Wiley and Sons.

330

Chemistry of Li-air batteries

7.11.2.1.3

Fundamental aspects of lithium peroxide crystallization

The morphology of solid lithium peroxide deposited at the cathode surface or within the porous cathode is diverse. Nevertheless, the observed structures can be divided to amorphous films covering the electrode surface or nanostructured crystalline agglomerates of lithium peroxide crystals. A homogeneous film less than 10 nm thick is formed both on smooth glassy carbon electrodes in an electrolysis cell94 and on porous electrodes during discharge at high current densities.45,78 Agglomerates usually appear within the porous electrodes during the discharge at relatively low specific current densities, and demonstrate a variety of morphologies as illustrated in Fig. 6. The most frequently observed shapes are separated plates or nanowalls,38,70,85–89 full or truncated discs,45,78,95–99 toroidal78,95-97,100–106 or spherical agglomerates77,90,97,98 composed of plates. Nanowalls can be frequently found in SEM images of discharged electrodes; they may cover the electrode surface and large agglomerates. Probably they are formed during evaporation of the electrolyte. Other crystalline morphologies present different stages of agglomerate shape evolution.101,106 In the most of observed agglomerates the individual plates can be easily distinguished along their side surface95. Detailed agglomerate structure is clearly seen in side view TEM image obtained in Ref. [95] and presented in Fig. 7A. This structure shows ordering of the layers. In addition, electron diffraction pattern presented in Fig. 7C for crystal in Fig. 7B directly confirms the crystalline order. The red and blue dots show the calculated electron diffraction pattern [001] of lithium peroxide, which is in good agreement with the experimental data. It can be concluded that the toroidal particle is lithium peroxide plates, which are stacked in ordered manner although they have small angular misalignment. Periodic packing of plates is confirmed also by small angle neutron107 and X-ray scattering88 . The plate thickness is few nm and  1 nm smaller than the period of mesocrystals, pinpointing solvent inclusions in the interplate space. It should be mentioned that the similar mesocrystals with toroidal and spherical shapes were observed earlier for CaCO3108 and ZnO.109 Li2O2 crystallizes in the hexagonal symmetry (space group P63/mmc).110 The crystal structure is shown in Fig. 7E.111 According to calculations in Ref. [112] the equilibrium crystal habit deduced from Wulff theorem is an anisotropic hexagonal prism as shown

(A)

(B)

(C)

100 nm 100 1 00 нм нм (D)

(E)

(F)

100 nm Fig. 6 Diverse morphology of LI2O2 formed during Li-O2 battery discharge: nanowalls. (B) Reprinted with permission from Sharon D, Hirsberg D, Afri M, Garsuch A, Frimer AA, Aurbach D. Reactivity of Amide Based Solutions in Lithium–Oxygen Cells. J. Phys. Chem. C 2014, 118, 15207–13, doi:10.1021/jp506230v. Copyright 2014 American Chemical Society; discs - (C) Reprinted with permission from Griffith LD, Sleightholme AES, Mansfield JF, Siegel DJ, Monroe CW. Correlating Li/O2 Cell Capacity and Product Morphology with Discharge Current. ACS Appl. Mater. Interfaces 2015, 7, 7670–7678, doi:10.1021/acsami.5b00574. Copyright 2015 American Chemical Society; toroids - (D) Reprinted with permission from Fan W, Cui Z, Guo X. Tracking Formation and Decomposition of Abacus-Ball-Shaped Lithium Peroxides in Li–O2 Cells. J. Phys. Chem. C 2013, 117, 2623– 2627, doi:10.1021/jp310765s. Copyright 2013 American Chemical Society (E) Reproduced from Schwenke KU, Metzger M, Restle T, Piana M, Gasteiger HA. The Influence of Water and Protons on Li2O2 Crystal Growth in Aprotic Li-O2 Cells. J. Electrochem. Soc. 2015, 162, A573–A584, doi:10.1149/2.0201504jes under the terms of the Creative Commons Attribution Non-Commercial No Derivatives 4.0 License; sphere - (F) Reproduced from Ref. Zakharchenko TK, Sergeev AV, Bashkirov AD, Neklyudova P, Cervellino A, Itkis DM, et al. Homogeneous Nucleation of Li2O2 under Li–O2 Battery Discharge. Nanoscale 2020, 12, 4591–601, doi:10.1039/c9nr08493b with permission from The Royal Society of Chemistry.

Chemistry of Li-air batteries

(A)

(B)

(C)

(D)

331

L

H

(E)

(F)

d h l

Fig. 7 (A) Side-profile TEM images of Li2O2 disc particle and (B) bright-field TEM images of toroid particles formed during Li-O2 battery discharge. (C) Simulated Li2O2 [001] zone axis (red and blue dots) superimposed over an experimental diffraction pattern for the particle pictured in panel (B). (D) Schematic structure of the Li2O2 agglomerate. (E) Crystal structure of Li2O2. Red spheres represent oxygen atoms, green spheres represent lithium atoms. (F) Equilibrium shape of Li2O2 crystallites based on calculated surface energies and the Wulff construction. The termination and relative areal fractions of basal and prismatic facets are identified. (A–C) Reprinted with permission from Mitchell RR, Gallant BM, Shao-Horn Y, Thompson CV. Mechanisms of Morphological Evolution of Li2O2 Particles during Electrochemical Growth. J. Phys. Chem. Lett. 2013, 4, 1060–1064, doi:10.1021/jz4003586. Copyright 2013 American Chemical Society; (E) Reprinted from Qiao R, Chuang Y-D, Yan S, Yang W. Soft X-Ray Irradiation Effects of Li2O2, Li2CO3 and Li2O Revealed by Absorption Spectroscopy. PLoS One 2012, 7, e49182, doi:10.1371/journal.pone.0049182 under the terms of the Creative Commons Attribution License; (F) Reprinted with permission from Radin MD, Rodriguez JF, Tian F, Siegel DJ. Lithium Peroxide Surfaces Are Metallic, while Lithium Oxide Surfaces Are Not. J. Am. Chem. Soc. 2012, 134, 1093–1103, doi:10.1021/ja208944x. Copyright 2012 American Chemical Society.

in Fig. 7F. Both (0001) and (1
100) faces are oxygen-enriched. At the same time, detailed inspection reveals that considering more faces and different oxygen pressure produces equilibrium habit much closer to isotropic.113 Therefore, the reason of the anisotropy is not clear at the moment. Moreover, the aspect ratio l/h for individual crystallites estimated from the broadening of (100) and (101) X-ray diffraction peaks is significantly higher (up to 10),114 than the equilibrium one for Li2O2 crystal, which is equal to 4.1.112 In addition, the mean lateral size is several times larger than the crystallite size estimated from coherent scattering domain dimensions.97 Therefore, each plate is not perfectly single crystalline but composed of several crystallites of the same thickness. Upon growth larger toroids or spherical particles formed during prolonged discharge do not demonstrate the perfect alignment of nanoplates anymore; this behavior is associated with secondary nucleation on the edges of the plates supposed in [95,97,105] on the basis of XRD data demonstrating the decrease of the mean lateral size of crystallites (nanoplates) at later stages of the process. Most probably, the electric field of the electrode plays an important role in mesocrystallization process. In the absence of electric field, the disproportionation reaction in the electrolyte bulk where the source of superoxide is KO2, yield Li2O2 plates which are not assembled in mesocrystals.66 Contrary, toroid mesocrystals form during discharge even in all-solid-state cells115 in the absence of Liþ$O2
SSIPs that support the hypothesis of forming mesocrystals exclusively in electric field. Unique kinetic studies were carried out for Li-O2 system by in situ TEM116–118 and electrochemical quartz crystal microbalance,119 that shed some light on the assembly process. Narrow size distribution for Li2O2 agglomerate particles is evidently supported by the growth of existing particles rather than by formation of new ones. Linear size of the agglomerates increases with time as t1/3,116 while the mass of Li2O2 grows linearly.119 It indicates diffusion-limited regime. Based on the analysis of the Li2O2 particle size dependence on the discharge current density of lithium-oxygen cells, it was suggested that the growth of agglomerates from plates can be limited by the local mass transfer of uncharged particles (LiO2), while lithium peroxide can be formed from these particles both by disproportionation of LiO2 on the surface of already formed Li2O2, and as a result of its electrochemical reduction.99 The high stability toward disproportionation of superoxide ions formed during the ORR and their sufficient diffusivity lead to a quite high superoxide ion concentration across all the internal spaces in a porous electrode. Further disproportionation results in large initial super- saturation for lithium peroxide, and its solubility is quite low. This enables Li2O2 nanoplates to nucleate homogeneously over the whole volume of the electrolyte, which was observed experimentally by operando small-angle X-ray scattering.88 These nanoplates are assembled into mesocrystals via a non-classical crystallization mechanism both at the electrode surface and in the electrolyte bulk. Subsequently, these particles can grow further, most likely via the classical mechanism by enlargement of the building plates with the additional nucleation of new plates on the edges of the initial plates. The general scheme is presented in Fig. 8.

332

Chemistry of Li-air batteries

Fig. 8 Overall scheme of Li2O2 agglomerate formation. Reproduced from Ref. Zakharchenko TK, Sergeev AV, Bashkirov AD, Neklyudova P, Cervellino A, Itkis DM, et al. Homogeneous Nucleation of Li2O2 under Li–O2 Battery Discharge. Nanoscale 2020, 12, 4591–4601, doi:10.1039/ c9nr08493b with permission from The Royal Society of Chemistry.

The shape of agglomerates and its evolution upon discharge is more or less the same in different solvents, though their appearance is more prominent in solvents where Liþ$O2
pair is more stable. Larger particles formed in electrolytes based on solvents with higher DN45,70 and AN90 or in the electrolytes with higher ionic association strength71,91 (Fig. 9A). There are three different explanations to size increase of toroid-shape particles in electrolytes with high water content (more than 500 ppm): (1) increase of Liþ$O2
ion pair stability due to better solvation of O2
90; (2) formation of Li2O2 via H2O2 soluble in the electrolyte68; (3) increase of Li2O2 solubility in electrolytes containing traces of water and, consequently, Ostwald ripening of the particles.96 However, Li2O2 solubility in the electrolytes with water content up to 1.5% is about 10
4–10
5 N that is not a high value,88,96 and narrow size distribution of Li2O2 toroidal agglomerates during discharge indicates that Ostwald ripening is unlikely. The discharge current density plays one of the key roles in the deposition of Li2O2 during Li-O2 battery discharge. At high current densities, an amorphous film is formed. At lower current densities, toroidal crystalline agglomerates are formed45,70,78 . The mechanism that explains different nature of the deposits at different current densities and, accordingly, overvoltages, was proposed by Adams et al.78 and is schematically shown in Fig. 9B. At the high current densities, lithium superoxide is predominantly reduced electrochemically immediately after its formation due to its high concentration. At low current densities and, accordingly, at lower supersaturation, crystals are formed more slowly that leads to a larger particle size. Homogeneous nucleation and further free growth of the large Li2O2 crystallites in the electrolyte bulk occurring in high-DN solvents is the important mechanism enabling the high capacity of Li-O2 batteries. In contrast, the capacity of Li– O2 cells containing low-DN solvents is restricted by electrode surface passivation.

Fig. 9

Surface and solution growth mechanisms of Li2O2 (A) in different electrolyte solutions30 and (B) as a function of current density.78

Chemistry of Li-air batteries 7.11.2.1.4

333

Lithium peroxide deposition in porous electrode

One of the key parameters that defines the perspective of Li-O2 battery practical application of is the specific discharge capacity. It is directly related to the maximal amount of Li2O2 which can be deposited within the space of porous positive electrode. The latter, in general, is mostly determined by the following three parameters: the discharge current density, solvation properties of the electrolyte solvent, and the positive electrode structure. Li2O2 is formed inside pores of different sizes (from micro to macro). It is crucial to reveal the factors limiting the Li2O2 deposition inside porous electrode, and structural parameters determining the specific discharge capacity of lithium-oxygen cells. Generally, the experimental discharge capacity of model lithium-oxygen cells with porous electrodes is a poorly reproducible value,99 depending on the microstructure of the electrode, specific location of the carbon agglomerates relative to each other, which depends essentially on the electrode fabrication process. It was shown in120 that discharge capacity depends on the pressing of the electrode. Pore structure of the initial carbon black electrode was studied using X-ray tomography by Torayev et al.121 Based on the experimental data, 3D reconstruction of the porous electrode structure was created, for which the discharge curves were calculated. It was found that the difference in the discharge curves is due to the stochastic nature of the pore interconnections inside the electrode. During Li-O2 battery discharge, the cathode surface is covered by either Li2O2 layer or/and islands and larger species. In the case of porous electrodes, Li2O2 is deposited inside the open pores. The cell discharge is limited by two main factors that are schematically shown in Fig. 10A: pores clogging that hinders oxygen transport, and the electrode surface passivation that blocks electron transport (Li2O2 is an electron insulator,123 and 4 nm thick film is enough to fully block the charge transport94). Theoretical124 and experimental107,125 studies explain a sudden and large overpotential increase, so-called “sudden death” of LiO2 cell, by limiting oxygen diffusion to the electrode pores rather than due to the surface passivation by the insulating product. To enhance the discharge voltage, it is preferable to use cathode materials with high specific surface area. For the positive electrodes made of carbon black the discharge capacity increases with an increase of the surface area.126 However, this approach has a natural limit being the fast clogging of small pores by Li2O2 deposit. For this reason, the cells with nanoporous electrode materials demonstrate low specific capacity, since most of the surface area of such materials is related to the pores that are not involved in the discharge process (Fig. 10B).127 Moreover, in carbon black pores thinner than  10 nm are not even wetted by the electrolyte according to SANS data.107 It should be noted that for electrodes made of dense and thick layer carbon of black, most of the Li2O2 is deposited nonuniformly - only within a limited thickness from the side where electrolyte is in contact with the gaseous medium. The electrode volume located closer to the lithium anode remains unfilled with the discharge product. This is particularly noticeable at high discharge current densities.120 This observation pinpoints limited oxygen transport in the electrode volume. To achieve higher specific capacity, the electrodes with complex pore architecture were proposed. The structure of the electrode should contain macrochannels more than 1 mm in diameter that are capable for oxygen transport into the electrode depth without being clogged with the product.128,129 Oxygen availability in pores can be enhanced by partial wetting of the cathode (incomplete flooding) with an electrolyte solution for the formation of three-phase boundaries,122 and by addition of hydrophobic fluorinated compounds to the cathode material to tune electrode surface wetting.130 High efficiency was demonstrated by using fluorinated additives in the electrolyte solution for increasing the oxygen solubility.131 An alternative mechanism of “sudden death” revealed by theoretical studies is electrode surface passivation when it becomes covered by electron insulating Li2O2 layer of remarkable thickness.106 The interplay between these two mechanisms is determined by many circumstances, electrolyte choice being a dominant factor. As it was discussed in the previous section, electrolyte solvation

(B) (A)

Fig. 10 (A) Schematic illustration of the pore filling during discharge. The growing Li2O2 layer leads to cathode passivation by electrical isolation (top right) and pore blocking (bottom right).122 (B) Accommodation of Li2O2 in the pores of various sizes. (B) Reprinted from Tran C, Yang X-Q, Qu D. Investigation of the Gas-Diffusion-Electrode Used as Lithium/Air Cathode in Non-aqueous Electrolyte and the Importance of Carbon Material Porosity. J. Power Sources 2010, 195, 2057–2063. doi:10.1016/j.jpowsour.2009.10.012. Copyright 2010, with permission from Elsevier.

334

Chemistry of Li-air batteries

ability governs the pathway of Li2O2 formation – on the electrode surface or in solution. First is generally typical for low-DN solvents and the second is realized for high-DN solvents. Their balance plays an important role in cathode pore filling during the discharge. Theoretical studies132,133 typically suppose than Li2O2 is deposited on the electrode as a film by direct twoelectron reduction and do not consider solution-mediated pathway focusing mostly on the differences in oxygen solubility and diffusivity for different solvents. Such approach is useful for poorly solvating media, but significantly underestimates the capacities for electrolytes with high Liþ and/or O2
solvation ability. However, it is known that the discharge capacity of model cells with flat electrode surface increases with an increase of the solvent DN,45 and it was also confirmed for porous electrodes in two solvents with different Liþ solvation ability - DMSO and acetonitrile.107 In certain theoretical studies134,135 generation of peroxide in the electrolyte bulk was taken into account, however, the number of fitted parameters in these simulations is high due to the lack of fundamental transport characteristics of discharge intermediates in Liþ-containing electrolytes. Several works120,136,137 were aimed at experimental studies of the electrolyte solvent effect on the porous electrode capacity, and also explain the revealed differences primarily by oxygen transport properties of the solvents. It should be mentioned that the discharge capacity can be limited by the other factors such as, for instance, Liþ ion diffusion. Under certain conditions cell discharge can result in increase of the Li2O2 amount deposited from the O2-side to Li-side of electrode evidencing the limited Liþ ion access to the reaction zone.138

7.11.2.2

Charge process

Charge process in Li-O2 batteries is generally much less studied than the discharge, and has several unsolved problems related to both positive and negative electrodes. Charging includes lithium peroxide oxidation accompanied by oxygen evolution (OER), and Liþ ion transport to anode where metallic Li is deposited. The mechanism of lithium peroxide oxidation in the cathode space plays a crucial role. One of the major difficulties in Li2O2 decomposition is insulating nature of the bulk lithium peroxide. At the same time, several studies pinpoint remarkable surface conductivity of Li2O2 demonstrated by various methods. Zheng et al. observed large toroidal particle growth and decomposition in solid state Li-O2 battery in SEM chamber.115 Theoretical studies predict the oxygen-rich termination of Li2O2 crystallites, along with the fact that Li-deficient surface species lead to the formation of halfmetallic structure.112 XANES studies have revealed that Li2O2 formed upon the Li-O2 battery discharge is substoichiometric and polycrystalline.139 Furthermore, Hummelshøj et al. have found that Li2O2 becomes an electronic conductor during the charge because mobile Li vacancies can be formed at the potentials above 2.85 V vs Liþ/Li.123 Although the studies mentioned above suggest that lithium peroxide surface exhibits half-metallic states and OER can occur, in practice the large charging overpotential is observed. Charging profiles are strongly dependent on Li2O2 morphology and side reactions such as carbonate formation. This assertion is supported by DFT calculations that predict the charging overpotential < 0.2 V,140 and experimental observations of O2 evolution at potentials ca. 3 V vs Liþ/Li at early stages of charge where the latter increase in potential is believed to appear due to solid carbonates deposition.141,142 Several mechanisms of Li2O2 oxidation have been proposed so far. One-electron Liþ de-insertion with Li-deficient component (Li2
xO2) formation was first suggested theoretically and predicted 0.3–0.4 V overpotential upon charging.143 Experimental evidence was later provided using operando XRD by Ganapathy et al.104 Another mechanism is mesocrystal decomposition that considers several steps of Li2O2 oxidation, namely decomposition of amorphous Li2O2, and oxidation of underlying crystalline phase through Li-deficient solid solution mentioned above. Schematic description of discussed processes is presented in Fig. 11.

(A)

(B)

Fig. 11 Cartoon showing the mechanism of (A) E-Li2O2 (formed during discharge) and (C) C-Li2O2 (bulk crystalline Li2O2) oxidation during the charge process. Reprinted with permission from Ganapathy S, Adams BD, Stenou G, Anastasaki MS, Goubitz K, Miao X-F, et al. Nature of Li2O2 Oxidation in a Li-O2 Battery Revealed by Operando X-Ray Diffraction. J. Am. Chem. Soc. 2014, 136, 16335–16344. doi:10.1021/ja508794r. Copyright 2014 American Chemical Society.

Chemistry of Li-air batteries

335

The mechanism of oxygen evolution is not yet clear; however, one can suppose it includes disproportionation of LiO2 at the electrode surface which yields molecular oxygen and Li2O2 (similar to discharge).104 Direct two-electron oxidation was suggested to explain significant difference in charge and discharge overpotentials.44,56,144 However, low overpotential at the initial stages of charge cannot be explained by this mechanism. Besides, theoretical studies have shown that delithiation and electron extraction with Li2
xO2 formation is more kinetically feasible than direct oxidation of Li2O2. Similar to discharge, Lu et al. have investigated the role of solvent in charging process.145 They have observed LiO2 formation on surface that generates soluble superoxide species in high-DN solvents. At the same time, in low-DN solvents further oxidation with O2 evolution occurs. The solvent effect is schematically explained in Fig. 12. All models consider that Li2O2 can form large crystalline species. However, mechanistic studies of thin film decomposition should be developed as well. Film-like Li2O2 usually has amorphous structure that leads to improved oxidation kinetics and smaller overpotential, due to the better Liþ and O2
mobility.78,97,146 Charge transfer through amorphous Li2O2 can occur via both tunneling of the charged vacancies and hole polaron migration.118,147,148 Another aspect of Li2O2 decomposition process is singlet oxygen evolution. Its appearance was confirmed by operando EPR and fluorescence spectroscopy.149,150 The thermodynamic potential of singlet O2 evolution from Li2O2 is about 0.5 V higher than that for triplet O2.151 Possible formation path is disproportionation of LiO2-like species: 2LiO2 /Li2 O2 þ 1 O2 1

O2 can also be released via electrochemical oxidation of LiO2 or Li2O2.143,150,152 Singlet oxygen evolution is important not only in terms of kinetic limitations; it is a quiet reactive form of oxygen that can react with Li-O2 cell components.153

7.11.2.3

Heterogeneous ORR/OER catalysts

The heterogeneous catalyst added to the cathode composition can potentially solve various problems associated with charge and discharge. The need for the true catalysis (i.e. reducing the overpotential) on discharge is currently a matter of dispute. Generally, ORR overpotential is not very high even on uncatalyzed carbon electrodes, and Tafel analysis on a flat glassy carbon provided a high exchange current density of 1 mA/cm2.140 Moreover, the conventional catalytic model is hardly applicable for Li-O2 system, as it assumes soluble reactants, products and intermediates, which can easily access and escape from the active site. Upon Li-O2 battery discharge, however, the catalytic sites are covered by insoluble products, and catalytic surface quickly becomes deactivated.154 After that, the reaction primarily occurs on the surfaces of solid discharge products (instead of the catalyst surfaces), that results in similar discharge voltages for various catalysts at the higher depths of discharge.155,156 This hypothesis was further supported by DFT studies showing that ORR kinetics on the Li2O2 surfaces are facile.123,141 Therefore, many authors suggest that ORR electrocatalysis is unprofitable in the Li-O2 electrochemistry.30,140,154 In contrast, the charge overpotential is basically much higher (> 1 V) on uncatalyzed electrodes, and the facilitation of the reaction kinetics is important here. However, McCloskey and co-workers154 compared the charging voltages of Li-O2 batteries with XC72 carbon, Au-XC72 and MnO2-XC72 based cathodes, and observed O2 evolution by DEMS starting from the voltage of  3.0 V for all three materials.154 That indicates there is hardly any kinetic barrier for the Li2O2 decomposition even if no catalyst is introduced. Nevertheless, many groups157–161 observed the catalytic effect (e.g. the decreased overpotential and improved

Fig. 12 Proposed solvent-controlled Li2O2 decomposition. Mechanism “H” denotes high-DN solvent and “L” denotes low-DN solvent. Li2O2* denotes the Li2O2 generated by LiO2(sol) disproportionation. Reprinted from Wang Y, Lai N-C, Lu Y-R, Zhou Y, Dong C-L, Lu Y-C. A SolventControlled Oxidation Mechanism of Li2O2 in Lithium-Oxygen Batteries. Joule 2018, 2, 2364–2380, https://doi.org/10.1016/j.joule.2018.07.021. Copyright 2018, with permission from Elsevier.

336

Chemistry of Li-air batteries

rechargeability compared with a carbon-based cathode) on the charge process. At the same time, many reports stress that this apparent improvement is not connected with the catalytic effect on Li2O2 product decomposition.153,154,157,159,162–166 For some catalysts, similar or even worse O2 recovery efficiency (OER/ORR) compared to a carbon cathode was demonstrated.157,159,162,164,165 The reason is that the apparent rechargeability improvement can be attributed to the different catalystrelated effects: (1) facilitation of Li2O2 decomposition; (2) promotion of byproducts (Li2CO3, LiOH) decomposition, and (3) facilitation of electrolyte decomposition. While only the first effect is desirable and increase O2 recovery efficiency, the others decrease it, although byproducts decomposition might be helpful to reduce the electrode passivation.141,154,167 It was reported for RuO275 and NiO168 catalysts. Facilitated electrolyte decomposition inevitable leads to electrolyte consumption.154,159,164,169 The choice of the electrolyte is, therefore, of high importance. Unfortunately, many studies, especially early ones, utilize the carbonate-based or ether-based electrolytes, which are now known to be unstable in Li-O2 environment, and many catalytic materials make the electrolyte decomposition even worse. For example, McCloskey et al. studied heterogeneous catalysts including Pt, MnO2, and Au and showed that both Pt and MnO2 induced considerable decomposition and an increase in CO2 evolution, and neither Au nor MnO2 demonstrated any catalytic activity.154 In ether-based electrolytes the situation is similar. Ma et al. have demonstrated that both Pd- and Ru-based catalysts increase the degree of irreversibility in Li-O2 cells in TEGDME-based electrolyte.159 Genorio et al.170 studied oxygen reduction in TBAþ-containing dimethoxyethane at well-defined single crystal electrodes. A decrease in the electrochemical reversibility of the redox couple is observed in the order Au > Pt > Ir, and chemical analysis of the electrode surface showed formation of a solid electrolyte interphase in the case of Pt and Ir. For the metal oxides, Barile and Gewirth171 showed that Pt, Pd, and Cu(II) oxide catalysts also increased the amount of CO2 evolved and decreased O2 evolution. For that reason, recent catalytic studies generally utilize DMSO-based electrolyte which is one of the most stable among the common electrolyte solvents. Table 2 represents different inorganic catalytic materials that have been investigated, and corresponding electrolyte compositions. It should be noted that much wider range of materials, including their combinations, have been tested so far; more examples can be found in the several reviews.205–207 Table 2 only illustrate the variety of the materials and the electrolytes used. Taken the electrolyte instability problem into account, the quantification of the O2 consumption and evolution during battery cycling is the right way to judge whether the improved battery performance is a result of a true electrocatalytic effect.205,208,209 Among various types of compounds, the catalyzed Li2O2 decomposition was confirmed for nanoporous gold (NPG)192 and Rubased catalysts.167 Although the reversibility of the cells was dependent on the electrolyte composition,159,164,165,167 in DMSObased electrolyte Ru167 and NPG192 demonstrated high O2 recovery ratios (> 90%). Transition metal carbides have been also considered as a promising class of cathode materials due to the high electrical conductivity and high oxidation resistance.157,200 Bruce and co-workers reported that cells using a TiC cathode and a LiClO4/DMSO electrolyte can be cycled with  99% O2 recovery for 100 cycles.200 Nazar et al.157 later showed that TiC surfaces with a sub-nanometer TiO2/TiOC passivation film can effective catalyze Li2O2 decomposition. Nazar and co-workers158 also exploited titanium oxide phases with high electrical conductivity for Li-O2 battery applications, and found that Ti4O7-based cathodes provide an ultralow onset voltage from 3.0 V and e
/O2 ratio of 2.42 on charging. For the materials which promotion effect on Li2O2 decomposition was proved, various catalytic mechanisms have been proposed so far. For transition metal oxide catalysts, Nazar and co-workers160 suggested they reduce the binding energies of LixO2 species with the surface, that promote their mass transport on the electrode surface and facilitate both further oxidation and reduction of LixO2. Shao-Horn and co-workers supposed that on metal catalysts solid-state Li2O2 oxidation is mediated by chemical conversion of Li2O2 to a lithium metal oxide, where Li2O2/catalyst interface is important.210 To prove this idea, Lu and co-workers211 studied the electrochemical behavior of all solid-state Li-O2 batteries with Vulcan carbon and Ru-catalyzed Vulcan carbon electrodes, to eliminate the electrolyte-mediated charging route that may also contribute to the charge process.212–214 As a result, authors observed more O2 evolution at a lower overpotential using Ru catalyst compared to the bare carbon,211 that verifies the solid-state catalysis for Li2O2 oxidation. The authors211 propose that reaction intermediates such as Li2
xO2 bind to the catalyst surface, that facilitates their delithiation at lower overpotential. Although in Li-O2 system many solid catalysts enable a very low onset charging voltage at 3.0 V (such as Ru,167 NPG,192 Ti4O7,158 TiCrOx,215 a common issue is that cells gradually polarize toward 4.0 V and higher upon deeper charging, usually due to the accumulation of parasitic products at the electrode/electrolyte interface.141,155,216,217 Additionally, above 3.5 V electrolyte decomposition and carbon corrosion occur to a greater extent141,155,216–218 and cause the battery irreversibility. Summing up, minimizing side reactions at the interfaces and constantly keeping the charging voltage below 3.5 V is needed to improve the reversibility of Li-O2 batteries. A desirable solid catalyst should satisfy the following criteria: (1) be stable to the reactive oxygen species and high overpotentials upon charging; (2) does not promote the electrolyte decomposition at the battery operating potentials (typically 2.5–3.5 V); (3) have a high binding energy between OER reaction intermediates (e.g., Li2
xO2) and the catalyst surface to promote Li2O2 delithiation on charge; (4) readily decompose LiOH and Li2CO3, because the accumulation of these typical byproducts rapidly cause the cell failure.

7.11.3

Negative electrode

In contrast to the existing Li-ion batteries, Li-air batteries need metallic lithium as an anode. In addition, while in Li-ion batteries the graphite anode constitutes rather low fraction of the total battery weight, in energy-dense systems such as Li-O2 the contribution of

Chemistry of Li-air batteries Table 2

337

Summary of the inorganic catalytic materials used in nonaqueous Li-O2 batteries.

Material Metal oxides Fe2O3, Fe3O4, Co3O4, NiO, CuO, CoFe2O4 MnO2 RuO2 NiO RuO2, Co3O4, MnO2 Co3O4@Ni MCo2O4 (M ¼ Mn,Ni,Fe,Zn) FePO4 Perovskites LaNi0.9M0.1O3 (M ¼ Cu, Co) La0.75Sr0.25MnO3 La0.5Sr0.5CoO3
x La1.7Ca0.3Ni0.75Cu0.25O4 Sr2CrMoO6
d Non-noble metals CuFe Nanoporous Ni Noble metals Au, Pt Pd Nanoporous Au Ag Ru Ir Doped carbon materials N-doped carbon S-doped graphene Titanium compounds Ti4O7 TiC TiN Other binary compounds SiC B4C Mo2C MoS2

Electrolyte

References

1 M LiPF6 in propylene carbonate 1 M LiPF6 in propylene carbonate 0.5 M LiTFSI in TEGDME 1 M LiPF6 in EC:DMC 1 M LiClO4 in DMSO 1 M LiClO4 in DMSO 1 M LiTFSI in TEGDME 1 M LiTFSI in TEGDME

172,173 173,174 75 168 175 176 177–181 182

1 M LiCF3SO3 in TEGDME 1 M LiTFSI in TEGDME 1 M LiTFSI in TEGDME 1 M LiTFSI in TEGDME 0.1 M LiTFSI in TEGDME

183 184 185 186 187

0.2 M lithium triflate in PC:TFP 1 M LiTFSI in TEGDME

188 189

1 M LiClO4 in PC: DME 1 M LiCF3SO3 in TEGDME 0.1 M LiClO4 in DMSO 1 M LiCF3SO3 in TEGDME 0.5 M LiClO4 in DMSO 0.1 M LiClO4 in TEGDME/DMSO (volume ratio 1:2) 1 M LiCF3SO3 in TEGDME

151,190 191 192 193 194 195 196

1 M LiTFSI in TEGDME 1 M LiPF6 in TEGDME

197 198

1 M LiNO3 in N,N-dimethylacetamide 0.5 M LiClO4 in DMSO 1 M LiTFSI in TEGDME 1 M LiTFSI in TEGDME

199 200 157 157

0.5 M LiClO4 in DMSO LiCF3SO3/TEGDME 1 M LiCF3SO3/TEGDME EMIM-BF4

200 201,202 203 204

the anode becomes larger, so energy-dense anode materials are desired. Obviously, pure lithium metal would be the ultimate material to satisfy both requirements. The very first rechargeable lithium batteries have used metallic lithium anodes, although several issues have made them impractical. The problem was solved by replacing lithium with sp2-carbonaceous materials; now a great effort is put into returning back to metallic lithium. The use of lithium metal faces some major challenges, which don’t have a robust solution up to now. Usually, they are outlined as two problems, one of which is high reactivity of lithium, which leads to inevitable formation of a passivation layer of solid electrolyte interphase (SEI). The second is morphological instability of lithium upon electrodeposition and dissolution: lithium tends to form very sparse non-compact deposits on charge and corrosion pits on discharge. Such protrusions can detach from the electrode on discharge, forming “dead lithium”219 which leads to the capacity loss since it no longer participates in electrochemical reactions, or can reach the opposite electrode, resulting in short-circuit. In addition, such deposits have a huge surface-to-mass ratio, and the high reactivity makes it even worse, leading to irreversible consumption of both lithium and electrolyte on the SEI formation. The morphological instability issue is usually referred as “dendrite growth” problem; however, sometimes the morphology is different from dendritic, and there are some other morphological issues (e.g., pitting on discharge), so we believe that this issue is more general. The problem of morphological instability of lithium was known since the very beginning of Li-batteries era in 1970s220 when lithium was just proposed as an electrode material. Generally, lithium doesn’t deposit smoothly, but forms bulky and loose mosslike structures. Dendrite growth is a well-known case in electrodeposition of many other metals like silver or zinc, so lithium deposits were similarly considered to be dendrites. However, later research revealed that at least in some cases the morphology

338

Chemistry of Li-air batteries

of electrodeposited lithium was not dendritic: for example, Yamaki et al. used term “whiskers” to describe the morphology of electrodeposited lithium basing on observations of growth dynamics in optical microscopy and SEM.221 Further observations revealed a huge variety of different morphologies including “moss-like,” particles, whiskers. It should be noted that deposits with morphology resembling true dendrites was also observed in gel electrolyte.222 The question of whether these structures are whiskers or dendrites is not just a matter of terminology. Different morphologies might be governed by fundamentally different mechanisms, which means that approaches to suppress such morphological instability would be different as well.

7.11.3.1

Many shapes of lithium

Dendrite growth is a fairly common phenomenon in electrodeposition. By definition, a dendrite is a skeleton of a single crystal with fractal tree-like structure.223 The stem and branches are formed from specific lattice directions, so the geometry is defined by the space lattice. Dendrites usually grow due to tendency of self-amplified growth of any small protrusions: upon electrodeposition cations are consumed on the surface of an electrode and concentration gradient is formed near the electrode, so any protrusion falls in a region with higher concentration of metal ions, and deposition at the tip occurs faster. The theory of dendrite growth is wellknown and developed for other metals. Such models focus on diffusion limitations on the tip of a growing dendrite, and one of the most notable models was developed by Barton and Bockris for silver dendrites.224 Later, Aogaki and Makino pointed out that dendritic growth starts after some incubation time when ion concentration at the electrode drops to zero,225 which became known as Sand’s time.226 Chazalviel pointed out that at such conditions unscreened high electric field may even further contribute to dendrite development.227 Later such theories were adapted for lithium, with some additions taking into account plasto-elastic behavior of polymer electrolytes.228,229 However, usually the observed morphology of electrodeposited lithium is completely different from dendrites. High-resolution imaging methods such as SEM or TEM reveal that such sparse deposits are usually a combination of whiskers and particles of  100 nm size,230 with bending and branching of whiskers having a random character, unlike regular fractal branching in dendrites. The critical current density for planar deposition of lithium also doesn’t seem to exist, as whiskers start to grow at current density as low as mA/cm2, which is far below the limit, proposed in dendrite growth theories.231 One of the first detailed observations of lithium electrodeposits was performed by Yamaki et al. using SEM.221 They observed long whiskers of  100 nm diameter with small particles on top. Similar results were obtained by Steiger et al., who used both in situ optical microscopy and SEM.231 Deposits that were described as “mossy” basing on optical micrographs, were actually a combination of particles and needles of sub-micrometer size. Since lithium is very soft and fusible metal, so the electron beam damage may significantly affect the deposits, cryo-electron microscopy techniques (TEM and SEM) were utilized to overcome this issue.230,232 It has revealed that lithium whiskers are single-crystalline without any regular defects (in some cases, metal whiskers can grow around screw dislocation,233 which is not the case for lithium). The crystal orientation of such whiskers is random and constant along the whole length without any sign of specific crystallographic facets on the side surface. When optic microscopy with low magnification is used, various different morphologies are also distinguished: in addition to distinct needles,228,234 tree-like, bush-like or mossy lithium was observed (Fig. 13); however, sometimes detailed observation of such mossy morphologies using SEM reveals that it also consists of bent of fractured needles.237 Another important yet controversial topic is the growth regime of such deposits. Initially, dendrite growth implies crystallization at the tip; however, since lithium deposits are not dendrites in a strict sense, other options are possible. Yamaki et al. have performed one of the first in situ observations of lithium electrodeposition using optical microscope.221 They have noticed that lithium filaments slightly swing and rotate upon electrodeposition, which have led to conclusion that they actually grow from the base by squeezing the metal from the electrode due to accumulating stress. Later, similar observations were performed by other groups.238,239 In some experiments the tip shape of a growing whisker remains constant, also allowing to conclude that insertion of new lithium atoms occurs at the base of a whisker.240 Furthermore, Steiger et al. observed elongation of not only base of a whisker, but also some middle parts between kinks and bends.241 They proposed that lithium inserts not at the base, but right underneath the tip or a kink, which leads to apparent root growth. Some essential details to this discussion were added using in situ NMR with isotopic labelling.242 This method allows to distinguish between lithium metal of different morphology. When lithium was deposited from naturally abundant electrolyte (z 92.5% 7Li) on 6Li foil, the composition of whiskers also corresponded to natural abundance, i.e., nutrient phase. Even though it doesn’t answer the question at which spot deposition occurs, it makes clear that whiskers are composed of freshly-deposited atoms and not ones pushed from the bulk. Apparent root growth was also observed by in situ TEM.243–245 Such setup used a very thin electrochemical cell with liquid electrolyte and silicon nitride windows, allowing the electron beam to transmit the cell. These observations should be treated with caution: room temperature TEM was shown to be very damaging technique: electron beam can burn lithium whiskers230 and severely damage the SEI,244 which definitely affects the morphology of the electrodeposits. However, this method provides very high magnification for direct in situ monitoring of electrodeposits in liquid electrolyte. One of such experiments have shown that in addition to root growth, lithium can deposit relatively uniform on the surface of whiskers (i.e., at the tip and sidewalls), leading to thickening of whiskers and formation of uniform particles.245 This observation backs up a concept that growth modes can switch at different conditions such as current density and charge passed. Bei et al. proposed Sand’s capacity (a charge passed before Sand’s time) as a criterion for a mechanism change236,246: before that, lithium deposits in a form of moss, which actually consists of small root-growing whiskers. After Sand’s capacity has passed, near-electrode region becomes depleted of lithium ions, and growth mode switches to typical tip-controlled dendrite growth. Morphology of such deposits differs from typical whiskers: instead of rather thick whiskers ( 1 mm diameter), particles are

Chemistry of Li-air batteries

339

Fig. 13 Various morphologies of electrodeposited lithium: (1) mossy, bush- and tree-like lithium, deposited at different temperatures235; (2) lithium needles under cryo-SEM and TEM230; (3) e,f – lithium mossy lithium, consisting of needles, g,h – dendritic lithium, transition between mossy and dendritic deposition.236 Published by The Royal Society of Chemistry.

much thinner; however, it is impossible to observe whether it has fractal structure; in addition, such dendrite growth mode was observed at rather high current density of 50 mA/cm2, which is pretty unusual in lithium-related literature and exceeds practical needs (Fig. 14).

7.11.3.2

Mechanisms of morphological instability

Following their observations, Yamaki et al. proposed their mechanism of lithium whisker growth.221 According to this mechanism, mechanical stress is accumulated underneath the SEI during lithium deposition due to non-uniformity of this process. Eventually, SEI breaks and lithium whiskers are extruded through such breaches in form of whiskers, which leads to their observed rotation and waving. In this model lithium was treated as a Newtonian liquid, which was rather unreasonable assumption, for which this idea was criticized later229 and kind of neglected for a long time. Later, another very different mechanism was proposed by Steiger et al.

Fig. 14

Status of published efforts on the cycling of lithium metal.247

340

Chemistry of Li-air batteries

basing on the observation that sometimes whiskers grow in between kinks.241 In this concept, such kinks or tips of whiskers are actually of different chemical nature and act as catalysts of lithium reduction, the same way as a liquid drop in vapor-liquid-solid whisker growth method. This theory aligns well with the observation of some subsurface structures detected with X-ray tomography, which appear just before whisker growth. However, the nature of such structures is unclear and this observation is open for interpretations. Recently, extrusion theory got its development. Even Yamaki et al. have noticed that lithium whisker growth resembles tin whiskers, another harmful case of whisker growth.248 This should not be confused with the famous demonstration of tin dendrite growth upon electrodeposition; tin whiskers grow when thin tin layer is thermally cycled249 or is plated on the copper substrate.250 Diffusion of copper, or the difference in thermal expansion coefficients of a substrate and tin causes mechanical stress in the tin film, and this leads to whisker growth. The mechanism of tin whisker growth was developed in 2000s251,252 and recently was adapted for lithium (Fig. 15).253 Indeed, lithium and tin have a lot in common besides tendency for whisker growth: both are covered with a passivation layer (oxide on tin and SEI on Li) and both are soft metals with low melting point (Li z 181  C and Sn z 232  C). According to this mechanism, the main driving force for whisker growth is mechanical stress accumulated during the electrodeposition of lithium underneath SEI. It was estimated that this stress can be relatively high and reach  100 MPa. One of possible reasons why lithium doesn’t deposit uniformly on the surface and builds internal stress is the presence of SEI. Formation of compressive stress in electroplated metal layers is a common phenomenon,254,255 but in lithium it eventually leads to whisker growth due to low lattice energy and high mobility of atoms in the metal. The mechanism of whisker growth is schematically presented in Fig. 15. Accumulated stress leads to the transport of lithium atoms to inclined shallow grain boundaries, where they are joining the lattice of an upper grain, pushing it upwards. Stress is released at such grain boundaries, allowing further flow of atoms during the deposition. This way tin whiskers grow from existing metal grains. However, the lithium whisker thickness is much less that lithium metal grain size (typical grain size is  10 mm vs. whisker thickness  1 mm). This stage eventually ends, because grain boundaries are very complex and the fraction of grains which can freely expand up is very low; majority of grains have only a fraction of such boundaries, which bend and accumulate stress, while other boundaries hold them like roots. When this stress accumulates, grain lifting can no longer proceed, and new grains nucleate at the surface. Nucleation threshold depends strongly on the surface morphology like electrode edges, grain edges or pits,256 so in non-ideal conditions when rough lithium foil or non-lithium substrate is used the first stage is extremely short and usually unnoticed. New nuclei are actually small grains without deep grain boundaries which could hold their growth, so they start to grow from the base by the same mechanism and form actual whiskers.

Fig. 15 Mechanism of lithium whisker growth. Reprinted with permission from Rulev AA, Kondratyeva YO, Yashina LV, Itkis DM. Lithium Planar Deposition Vs Whisker Growth: Crucial Role of Surface Diffusion. J. Phys. Chem. Lett. 2020, 11, 10511–10518, doi:10.1021/acs.jpclett.0c02674. Copyright 2020 American Chemical Society.

Chemistry of Li-air batteries

341

According to this mechanism, lithium whisker growth is governed not by the processes occurring in the nutrient phase, but within the electrode. It means that any electrolyte is not even needed for whiskers to grow, which was observed in multiple experiments. When lithium was deposited by physical vapor deposition, the final morphology of lithium whiskers was almost identical to the one obtained by electrodeposition.257 In various experiments on in situ electrodeposition in SEM chamber, where a cell with solid electrolyte was used, lithium whiskers were growing on the electrode-vacuum interface258–260 or simply on the surface of solid electrolyte, where lithium was reduced directly by the electron beam.261 Notably, in this case the residual pressure in the chambers was playing the critical role: morphology changed from uniform particles to whiskers and back to uniform particles upon pressure change from 10
7 Pa to 10
5 Pa.262 In other notable experiments, lithium was deposited inside the ETEM chamber on the tip of AFM cantilever. The tip was brought to a small lithium particle and the native SEI on lithium acted as solid electrolyte, what allowed to deposit a single nucleus on the cantilever tip, which later evolved into a root-growing whisker.263,264 All these observations show that whiskers grow without any mass transfer limitations or inhomogeneities, supporting the idea that reasons lie within the electrode. Key factors are the presence of SEI and low energy of lithium lattice and thus, high mobility in solid phase. For example, when the temperature is reduced to 
10  C, portion of whiskers is much lower and lithium deposits in a form of spherical particles.253 All in all, a tendency to whisker growth seems to be an inherent feature of lithium. In any real-world conditions the metal would always tend to form whiskers. Three main processes and their interplay govern the apparent morphology: nucleation, root whisker growth and surface deposition of lithium. In most cases, different combinations of three results in different apparent morphology such as needles, mossy, tree-like and bush-like or “dendrites”. Different mechanisms may act simultaneously on different scales: on the smaller level, lithium always tends to form root-growing whiskers. At larger scale, transport limitations in the electrolyte or high electric fields after Sand’s time may come into play, defining relations between nucleation rate, whisker growth and surface deposition. This way, we believe that all of the theories above are not necessarily mutually exclusive, and morphology of electrodeposited lithium is hierarchically defined by multiple factors, that makes this problem very complex and multiscale.

7.11.3.3

Steps toward uniform deposition

Dendrite growth is a well-known issue in metal electrodeposition and for many other metals is suppressed by addition of specific levelling and brightening agents, which adsorb on specific planes blocking the development of protrusions.265 However, actual lithium dendrites start to grow at high current density exceeding practical needs, and the real problem most likely is whisker growth. Whiskers, according to the current theories, have some minimal thickness, so the short-circuiting can easily be mitigated using nanoporous separators, impermeable for thick whiskers.246 “Dead” lithium and high surface area of non-uniform deposits are much bigger problems: presence of SEI is considered to be one of the main reasons of morphological instability; in turn, the formation of SEI makes the developed morphology an actual problem, so these two phenomena are deeply interconnected. Associated problems were still not addressed properly: even though a large volume of works on suppressing morphological instability is published, for some reasons cycling parameters (like depth of cycling, current density etc.) are far below practical needs (Fig. 14), so one can assume that such efforts and approaches are not completely successful. The big problem in the efforts to achieve planar deposition is that there is no commonly accepted universal mechanism of morphological instability. Another obstacle is the high reactivity of lithium and thus sensitivity of the system to changing parameters. It was described above how the slight change of residual gas pressure changes the morphology of lithium particles in the SEM chamber; in addition, the composition and thus properties of SEI highly depend on the electrolyte composition. For this reason, it is sometimes difficult to determine which parameter of the system was actually affected by the approach and what actually leads to the morphology change. It can be demonstrated by a very illustrative case: Ding et al. have shown that addition of tiny amount of Csþ cations to the electrolyte led to complete suppression of lithium whiskers and deposition of a uniform compact layer, which was explained by the “electrostatic shield” of non-electroactive Csþ cations, which repel lithium ions from protrusions.266 In the following paper the same group observed the exactly same effect just by adding small amount of water267 . Later the concept of “electrostatic shield” was proven to be wrong,240 so the role of Csþ cations was thus very questionable. This way, it is difficult not only to verify and reproduce different approaches, but also to systematize them. Another subtle point of most approaches proposed in the literature is the fact, that they usually deal with kind of isolated lithium anode without any relation to lithium-air batteries. Operation conditions in Li-air batteries usually impose several restrictions on the system: in the cells with solid electrolyte membrane isolating the anode and the cathode space, approaches may be picked more freely. However, when it comes to the cells with non-isolated electrolyte, oxidative stability of proposed additives should also be considered. In addition, one has to take care of protecting lithium electrode from reactive species in the electrolyte like dissolved oxygen. Also, usually the electrolytes which were optimized for the cathode performance have poor reduction stability on the lithium electrode, so the need to protect lithium electrode must be taken into account. Various efforts to suppress morphological instability were comprehensively reviewed in multiple papers268–271; here we will briefly review the main directions in this field and their key points.

7.11.3.3.1

SEI design

As SEI is one of the key reasons for whisker growth, many efforts concentrate on designing SEI with desired properties. Besides, the designed SEI may protect the anode from unwanted side-reactions. Two main approaches to design specific SEI are usually utilized: using electrolyte additives which react with lithium and modify SEI in the desired way; and creating an artificial SEI on lithium before the cell assembly. Advantages of the first approach include the ability to repair SEI if it cracks during cycling and its

342

Chemistry of Li-air batteries

adaptation to the volume change of the electrode. One of the most popular SEI-modifying agents are fluorine-containing salts, which are usually used in Li-metal batteries since they create “good” stable SEI, mostly consisting of LiF. For example, adding small amounts of water was shown to significantly improve the morphology due to hydrolysis of F-containing salts and LiF generation.267 Other common SEI-modifying additives are components which form polymer interphase which, from the one side, is resistant to cracking, and from the other side, blocks whisker growth due to its elasticity.272,273 Such polymer layers can also be created ex situ prior to the cell assembly; other artificial interphases include alloys, inorganic solid electrolytes such as LIPON etc., which can be coated by various methods such as PVD, spray coating or many others.274,275

7.11.3.3.2

Electrolyte design

A similar approach is the design of the media adjacent to the electrode, i.e., the electrolyte or separator. Use of solid electrolyte membrane, beside protecting lithium anode from the reactive species on the cathode, may also prevent whisker growth. It was shown that mechanical pressure applied to lithium electrode can block whisker growth; to achieve this, the electrolyte with a modulus higher than of lithium is required.276 While polymer or, obviously, liquid electrolytes are unable to achieve it, inorganic solid electrolytes can provide such stiffness.277 Apparently, lithium whiskers usually have thickness of 100 nm – 1 mm, thus special separators with smaller pores may be effectively impermeable for whiskers protecting the cell from short-circuit, and can mechanically prevent whisker growth the same way as actual solid electrolytes.246,278 Besides, some special separators which redistribute mass-transfer or have some impact on the surface were proposed.279

7.11.3.3.3

Electrode design

One of the recent directions is design of the electrode, or, more precisely, current collector. One of notable solutions was proposed basing on the role of mechanical stress in whisker formation. To release stress in electrodeposited lithium film, authors used current collector deposited on the soft polydimethylsiloxane (PDMS) substrate. Upon deposition, stress in the lithium layer accumulates and eventually releases by wrinkling of the current collector, resulting in unstressed whisker-free lithium layer.280 In order to reduce local current density, porous or spongy lithium hosts are also proposed. Since nucleation is another important factor defining the overall morphology of lithium deposits, such hosts usually have to be lithiophilic, which makes nucleation more uniform. Lithiophilic host only allows to accommodate a small amount of lithium before whisker growth becomes possible again, so large surface area is important.281–283

7.11.3.3.4

Alternative anode materials

Although alternative anode materials don’t directly relate to metal lithium anode, they may be more probable high-density candidates for anode in Li-air batteries than metal lithium with its unsolved problems. A large variety of different materials from amorphous carbon to transition metal oxides are proposed and described in several reviews.281,284 One of the most promising candidates is silicon, which makes phase Li15Si4, which has the highest specific capacity after lithium (1866 mAh/g vs 3860 for Li and 339 for graphite; here we calculate capacity per lithiated phase unlike conventional Li-ion literature to be able to compare to metal lithium). The major challenge for silicon, as well as many other alternative materials is large volume expansion upon lithiation ( 400%), leading to cracking of material. To get around this problem, various nanostructured or composite materials are proposed.281

7.11.4

Reactions with reactive oxygen species

7.11.4.1

Electrolyte decomposition

Due to a promise of extraordinary high specific energy, numerous efforts on developing cathode materials285,286 and electrolyte compositions287,288 were performed after Abraham’s report on implementation of polymer electrolytes in Li-air batteries in 1996.27 For about 15 years the electrolyte stability issues were out of the focus until it was found that carbonate species, which are formed as a result of the alkyl carbonate-based electrolyte decomposition, prevail in the discharge products.289 From that time, electrochemical measurements were coupled with advanced characterization techniques, e.g., differential electrochemical mass spectrometry (DEMS),192,290 Fourier transform infrared spectroscopy (FTIR),290,291 X-ray diffraction (XRD),292 and nuclear magnetic resonance (NMR) spectroscopy,293 which enabled identifying the byproduct formation mechanisms. The decomposition of carbonated-based electrolytes in LABs was revealed by Freunberger et al.294 It was shown that oxygen reduction proceeds via the formation of superoxide species, which nucleophilically attack the C]O groups in organic carbonates. This reaction yields Li2CO3, C3H6(OCO2Li)2, CH3CO2Li, HCO2Li, and CO2. Further theoretical studies295 showed that such radical attack can be a dominant reaction pathway for all carbonate-based electrolytes. Many aprotic noncarbonate-based electrolyte solvents were also found to be non-stable, however, they react in a different way. Superoxide anion, as a strong base, attacks the solvent molecule, which can be represented as H-R, and abstracts a- or b-hydrogen atom from it.137,296–298 As a result, HO2
and R are formed and involved into further reactions leading to the electrolyte decomposition and various side product formation. The side reaction rate thus depends on the solvent’s acidity and O2
concentration. Solvents comprising high donor and acceptor numbers (DN and AN) were found to prolong the superoxide anion lifetime in solution and increase the O2
concentration.69 If this is accompanied by the acidic hydrogen in the solvent molecule, the latter degrades. For instance, DMSO having a high DN was reported to react with O2
yielding hydroperoxyl radical that further attacks the S atom of DMSO, which, as known,299 finally leads to LiOH. Ethers like DME with low DN and absence of strong electron-

Chemistry of Li-air batteries

343

withdrawing functional groups156 could remain stable against the superoxide radical attack for a week or more. Long exposures in battery operating conditions, however, also lead to degradation of ethers, which happens mainly during recharge cycles. Both linear (e.g., DME, TEGDME) and cyclic ethers were found to oxidize at potentials above 4 V on recharge, which was proven by numerous analytical tools including DEMS, FTIR, XRD, and NMR spectroscopy.156,290,300 Reactivity of various classes of solvents is summarized in Table 3. A number of efforts were made to develop more stable electrolyte solvents for LABs. As the acidity of a-hydrogen strongly affects the solvent stability against oxidation triggered by the superoxide attack, one of the most popular trends in development of new solvents for LAB electrolytes is the synthesis of a-proton-free compounds. For instance, 2,2,4,4,5,5-hexamethyl-1,3-dioxolane320 and 2,4-dimethoxy-2,4-dimethylpentan-3-on321 were considered as prospective candidates. Using strongly solvating media can be also effective as shown for, e.g., hexamethylphosphoramide, which was found to support nearly reversible oxygen reduction/ evolution to Li2O2.322 Another aggressive agent capable of triggering the electrolyte decomposition is highly reactive singlet molecular oxygen, which can evolve due to lithium peroxide oxidation on recharge. As the oxidation of hydrogen peroxide in alkaline media is a common way to generate singlet oxygen, yet in 2011 it was supposed that Li2O2 oxidation during LAB recharge can also result in 1O2.323 Direct experimental prove was grasped later using operando EPR.150 It was found that singlet oxygen is evolved during charge at 3.55–3.75 V vs Liþ/Li range. The process was interpreted as two-electron reduction of lithium peroxide to 1O2,150,160 however, the debates on the mechanism are ongoing. All in all, singlet oxygen participates in the side reactions not only with the electrolyte solvents, but also with other cell components. Some electrolyte salt anions can also be reactive in battery operation conditions. The side products formation upon LAB discharge was studied in solutions of LiPF6, LiBF4, LiClO4, LiBr, CF3SO3Li, LiB(C2O4)2, and LiTFSI in TEGDME as electrolytes.324,325 Combining XRD and XPS analysis of the electrodes after discharge, and NMR spectroscopy of the electrolyte solution, it was shown that LiClO4 is the most stable salt. LiBr is stable during discharge but is easily oxidized upon recharge. Using LiTFSI, CF3SO3Li j LiPF6 results in a little amount of LiF on the surface. Lithium fluoride was detected not only at the positive electrode surface, but also at the surface of Li2O2 after soaking the powder in LiPF6 solution.309 Boron-containing electrolytes LiB(C2O4)2 and LiBF4 cause the formation of Li2C2O4 and LiF as main discharge products.324

7.11.4.2

Carbon electrode degradation

Unfortunately, carbon electrodes, which are widely used as cathodes for LABs, also are involved into side reactions under LAB operation conditions. Using ex situ XPS analysis of the electrodes after discharge, and DEMS with isotope labeling, McCloskey et al.141 suggested that carbon reacts chemically with Li2O2 and O2. This reaction is favorable thermodynamically, however, the kinetics was supposed to be sluggish as the solid product covers the interface and hinders further reaction. The formation of thin Li2CO3 film was confirmed by other studies, e.g.,155; however, carbon demonstrated better stability when poorly solvating media was used (e.g., in glymes293). Later it was shown that carbon electrode oxidation to Li2CO3 more probably proceed via another pathway.326 Carbon-free solid electrolyte was used to avoid detection of the electrolyte decomposition products in the operando near-ambient pressure XPS

Table 3

The reactivity of various classes of aprotic solvents in the lithium-air batteries.

Compound class

Examples

Sulfoxides

DMSO

Alkylcarbonates

PC EC TEGDME DME

Ethers Nitriles Esters Amides Sulfones Ionic liquids

CH3CN trimethylacetonitrile g-butirolactone methylformate DMA DMF N-methylpirrolidone TMS EMI TFSI, PP13TFSI, PYR14TFSI

pKa 298

51.8 30.3 49.4 34.4 42.9 35.2

Stability

References

Relatively low reactivity toward superoxide-anion; reacts with Li2O2 in presence of superoxide species. LiOH, Li2SO4, DMSO2 byproducts are formed. Non-stable toward nucleophilic attack

301–304

Slow degradation due to proton abstraction and further oxidation; Li2CO3, OSO2Li, CH3CO2Li side products after long-term cycling. Non-stable during recharge. Relatively stable, minor reaction products were detected by in situ XPS Non-stable toward nucleophilic attack

214,218,291,300,307

Degradation due to proton abstraction and further oxidation; Li2CO3, CH3CO2Li side products after longterm cycling Stable Imidazolium-based liquids are non-stable toward superoxide; Pyrrolidinium- and piperidinium-base liquids are stable

101,312–315

294,300,305,306

295,298,308–310 295,306,311

65,311,316,317 318,319

344

Chemistry of Li-air batteries

experiment. The results indicated that during the discharge superoxide species are generated, and it is followed by conversion of superoxide into Li2O2 and simultaneously into epoxy or similar groups on the surface of a thin carbon electrode. Afterwards, the concentration of carbonate species starts to increase with consumption of the oxygen functional groups on carbon electrode, and this process occurs even at open circuit if there are residues of superoxide species. Based on these findings, it was suggested that it is superoxide anion or LiO2, which promote carbon degradation due to nucleophilic attack or electron-transfer reactions followed by oxidation with molecular oxygen. More recently, Kataev et al.327 performed model chemical experiments in clean ultra-high vacuum conditions and showed that Li2O2 synthesized in situ does not react with graphene serving as a model carbon material, although the reaction possesses negative free energy change. In contrast, mixture of peroxide and superoxide species (KO2 was used as a model) quickly reacts with carbon in presence of molecular oxygen. These results are in line with previous hypothesis on carbonate formation after glassy carbon came in contact with potassium superoxide powder, which was based on appearance of carbonate decomposition peak in cyclic voltammogram.326 Thus, an increasing number of experimental data reveals that superoxide-promoted pathway gives major contribution to carbon electrode corrosion. Comprehensive understanding of the mechanism is still missing, unfortunately. Nevertheless, both spectroscopic and electrochemical43,326,328 studies give a hint that defects in sp2-carbon lattice and oxygen functionalities on the surface speed up carbon electrode degradation. Carbon oxidation can also occur during the cell recharge. Mainly, Li2CO3 is a detectable byproduct. It is a result of carbon interaction with either singlet oxygen or superoxide radicals, which both can be generated during Li2O2 decomposition. The mechanisms are still actively discussed in the literature and summarized in recent reviews, e.g.12 Going to high potentials (over 4.5 V vs Liþ/Li) on recharge enables oxidative decomposition of Li2CO3; it leads, however, to inevitable consumption of carbon electrode material.141,326

7.11.5

Redox mediators

7.11.5.1

Basic principles

The main discharge product of non-aqueous lithium-oxygen batteries is insoluble and insulating lithium peroxide Li2O2, thus its decomposition on recharge requires drastically large overpotential, which makes the charge process extremely challenging. Li2O2 thin film or nanoparticle formation may lead to the regime in which overpotential is not too high, and the charging proceeds at applicable potentials.329 However, in order to reach high specific capacity, large Li2O2 particles formation is preferable. In the latter case, particles might be located far from the electrode surface that leads to poor charging efficiency, electrolyte decomposition and other side processes. One of the possible ways to overcome this issue is substituting direct electrochemical oxidation of Li2O2 (Reaction 1) by the chemical reaction involving so called redox mediators (RM) or electron shuttles. These are soluble electroactive species, which experience reversible electrochemical oxidation at the electrode, diffuse to the Li2O2 particles and oxidize them chemically.330 Schematic mechanism is represented by the Eqs. (1)–(5) (Fig. 16). Li2O2 – 2e
/ 2Liþ þ O2

(1)

RM – e
/ RMþ

(2)

2RMþ þ Li2O2 / O2 þ 2Liþ þ 2RM

(3)

RMþ þ Li2O2 / LiO2 þ Liþ þ RM

(4)

2LiO2 / Li2O2 þ O2

(5)

Several requirements are made for compound to act as a RM for Li-O2 battery. First of all, the equilibrium redox potential of RMþ/RM pair has to be greater than that for Li2O2 decomposition reaction E(2Li þ O2 ¼ Li2O2) of 2.96 V vs Liþ/Li. On the other hand, to lower the overpotential, E(RMþ/RM) should not be too high and should definitely not exceed electrolyte decomposition potential. One must note that, although thermodynamically 3 V is suitable potential, the actual minimal value for properly working RM has to be much higher to achieve reasonable reaction rate. Another key parameter is diffusion coefficient since relatively fast diffusion is required to maintain sufficient charging rate. Finally, RM has to be chemically inert toward oxidizing species like O2 and O2
. Schematic illustration of ORR and OER processes and three electrochemical properties affected by the RMs are presented in Fig. 17. However, not only thermodynamics but also kinetics affects RM performance. Precise study on the Li2O2 oxidation kinetics by variety of RMs revealed crucial role of RM’s kinetic parameters and shed light on the mechanism and the rate-determining step.332 It was shown that RM with relatively high redox potential may exhibit slow kinetics of Li2O2 oxidation. Among the investigated compounds, only RMs with E > 3.6 V vs Liþ/Li showed high rate constant values. Besides, it was demonstrated that RM reaction with Li2O2 is inner sphere process thus making oxidation rate depending on the RM structure.

Chemistry of Li-air batteries

345

Fig. 16 (A) Schematic illustration of ORR and OER RMs in Li-O2 batteries, (B) Three electrochemical properties affected by RMs: energy efficiency, power capacity, and coulombic efficiency. Reprinted from Ko Y, Park H, Kim B, Kim JS, Kang K. Redox Mediators: A Solution for Advanced Lithium– Oxygen Batteries. Trends Chem. 2019, 1, 349–360, doi:10.1016/j.trechm.2019.03.016. Copyright 2019, with permission from Elsevier.

Fig. 17 First-cycle load curves (constant-current discharge/charge) with and without the redox mediator. 1 M LiClO4 in DMSO at a NPG under 1 atm O2 with 10 mM TTF (blue) and without TTF (red). The rate - 0.078 mA cm
2.331

7.11.5.2

Redox mediators for charge

Organic molecules having reversible redox near 3.5 V vs Liþ/Li is a large group of potential RMs including TEMPO and its derivatives,333–336 DMPZ,333,337–340 TDPA.341 Tetrathiafulvalene (TTF) was a pioneer compound to act as a Li-O2 redox mediator. The cell containing TTF was shown to have lower overpotential (by more than 0.5 V) (Fig. 17) and stability for 100 cycles.331 However, organic RMs typically tend to be oxidized by superoxide species, although this reaction is poorly investigated. Besides, not only superoxide species, but also singlet oxygen may lead to RM degradation which was studied by Kwak et al.337 From the other hand, the advantage of organic RMs is the synthetic opportunities to modify their structure. For instance, Bergner et al. explored various TEMPO-based compounds and demonstrated not only the difference in redox potentials, but also significant deviation in oxidation stability.336 Other common compounds tested as RMs are metal-organic transition metal complexes. Iron phthalocyanine (FePc),342 cobalt tetraphenylporphyrin, phthalocyanine and bis(terpyridine) derivatives (CoTPP, CoPc, Co(Terp)2)102,343 were investigated and demonstrated similar results including decrease in overpotential by ca. 0.5 V. Organometallic compounds, similarly to organic ones, could be oxidized upon Li-O2 battery cycling. Having a transition metal core, these structures are engaged in reversible redox which potential is mainly determined by the potential of corresponding M(n þ 1)þ/Mn þ couple. This approach enables tuning the potential by the choice of suitable transition metal. Also, adjustment the ligand structure is another way to vary precisely the redox potential of the RM, and in general defines its chemical properties that was shown in the work of Matsuda et al.102 Some inorganic compounds, mainly lithium halides LiI and LiBr, were reported to be appropriate for oxidizing Li2O2.333,344–346
Upon charging, halide salts are believed to form trihalide structures (I
3 , Br3 ) that further oxidize Li2O2. Such simple compounds have their obvious advantages as stability toward oxidation by molecular oxygen and its reduced forms. However, LiBr causes

346

Chemistry of Li-air batteries

corrosion problems due to the formation of Br2 and Br3
species, and LiI exhibits slow kinetics in the reaction with Li2O2.333,347 The results of the several works devoted to various RMs for charge are summarized in Table 4.

7.11.5.3

“Shuttle effect”

Relatively large redox potential required for RM makes its reaction with metallic Li thermodynamically favorable. Since Li is a strong reducing agent, not only oxidized form of RM could be reduced by Li to original RM state lowering the capacity (Eq. 6), but also several irreversible side reactions with RM may take place leading to gradual decrease of RM efficiency.372 RMþ þ Li / RM þ Liþ

(6)

This is a fundamental problem of “classically” designed Li-O2 cells where cathode and anode share the same electrolyte. However, this problem can be treated in several ways. An obvious way to prevent RM reaction with Li anode is to separate cathodic and anodic spaces of the cell. The use of solid electrolyte membrane appears to be an effective way to suppress anode reaction with RM, to increase cycling stability and to lower the overpotential.334 More intricate ways include coating some parts of the cell: Li anode coated with ceramic-polymer composite retain sufficient stability to RM reactions.335 Similar concept of using PEDOT:PSS-modified separator was proven to suppress RM diffusion to the anode due to strong interaction of oxidized RM with negatively charged polymer groups.340 Another way to prevent RM from migrating toward the anode is linking certain active groups to the cathode surface. Immobilized TEMPO groups on the cathode showed sufficient stability compared to dissolved TEMPO, however, the decrease in overpotential was only ca. 300 mV.383

7.11.5.4

Redox mediators for discharge

The RMs can be also applied to Li-O2 discharge. Since carbon cathode materials used in Li-O2 systems are unstable toward oxidation by superoxide intermediate O2
, reducing the concentration of O2
near the carbon electrode surface may provide more stable battery performance and increase the discharge capacity. One of the suggested RMs for discharge is 2,5-ditert-butyl-1,4benzoquinone (DBBQ) proven to increase discharge capacity up to 80–100 times in low-DN solvents like TEGDME and DME.72 DBBQ is believed to coordinate Liþ cation and superoxide anion, thus Li2O2 is formed far away from the electrode surface. ORR mechanism in presence of DBBQ is shown in Fig. 18. Despite the fact that anthraquinone (AQ) has redox potential lower than ORR in aprotic media, Han et al. investigated modified AQs as RMs for oxygen reduction. The discharge capacity in TEGDME was in good correlation with the RM’s redox potentials, that opens the possibility to obtain the desired redox potential values by the proper RM’s design.384 AQs immobilized on the cathode surface were also studied and showed stable performance, increased discharge capacity and no effect on Li anode.385

7.11.5.5

Bifunctional and dual mediators

Addition of LiNO3 to the electrolyte solution or even using it as electrolyte salt was reported to have dual effect on Li-O2 cell performance. Nitrate anions induce the formation of stable SEI on metallic lithium which is accompanied by NO2 evolution. Highly reactive NO2 molecules may diffuse to the cathodic space and act as RM for Li2O2 decomposition. As a result, In314,382,386 improved Li anode stability and decreased charging overpotential were observed in LiNO3 electrolyte solution.314,382,386

7.11.6

Concluding remarks and future prospects

Li-O2 batteries are certainly a promising next step in energy storage development. However, they are still far from practical implementation, despite the tremendous efforts of scientific community. These efforts were first focused on achieving theoretical capacity. In the aprotic Li-O2 system ORR kinetics is fast, and a catalyst is not needed. Capacity is directly dependent on efficient deposition of maximal amount of Li2O2. To achieve it, the conditions of homogenous Li2O2 nucleation and free growth in the electrolyte solution (rather than insulating film deposition at the electrode surface) should be supported by using solvents that efficiently solvate Liþ and/or superoxide ions. In turn, the internal space of an electrode and the reactants mass-transfer should be organized properly. Porous electrodes of optimized microstructure, for instance, with bimodal pore distribution, are promising for this purpose. Second, the electrode passivation by the side products originated from reactions of superoxide anions with electrode and electrolyte should be avoided. These reactions (especially with the electrode material) can be essentially diminished upon LiO2 stabilization in highly solvating electrolytes, due to the fast outflow of reactive species from the electrode surface. Of course, the use of proper electrode material and stable electrolyte would be an ideal solution. However, at the moment great variety of proposed modified carbon and other cathode materials still suffer from the stability problem. The next important problem is poor rechargeability. To move toward practical lithium-air battery, the charge voltage should be brought well below 3.5 V; this is because carbon cathode corrosion, singlet oxygen generation, and associated electrolyte decomposition are all sensitive to higher voltages. To reduce charging overpotential, solid catalysts have been widely explored. However,

Table 4

The performance of various RMs for Li-O2 battery charge. Name and structure

ECV, VLi Plateau 3.5 3.7Au

tetrathiafulvalene

TTF

3.4–3.6 V

3.67 V

3.38 3.69Pt

4.5 V

1 Vvs LFP

3.5 V

TTF

3.6 3.8gc

Electrolyte

CRM, mM

Air electrode

Overpotential, mV References

Reduces overpotential by more than 500 mV with a current density up to 1 mA/cm2. It shows 2.01 e
/O2 ratio (close to theoretical) showing stability unlike Fc and TMPD. Reduces overpotential. On gold TTF reacts with LiO2 but not Li2O2. TTF unlikely works as an ideal RM at the porous carbon electrode since carbonate formation appears with and without TTF. Overpotential on second charge is greater than without TTF. TTF chemically reduces Fe3þ and induces lithiation of FePO4 in the presence of Liþ. Li2O2 dissolution with TTF is shown by STEM. Observed Li2O2 decomposition at Li2O2/electrolyte interfaces following the TTFþ diffusion forefront is specific to the presence of TTF. Reduced overpotential, and stable cycling at large capacities were demonstrated in both coin cells and pouch batteries. TTFþ removes the solid Li2O2 from the electrode surface at lower potential than required for the electrochemical oxidation of DMSO to dimethyl sulfone as shown by AFM. SECM experiment demonstrated that soluble TTFþ chemically reacts with surface species blocking the electrode consistent with EQCM mass recovery, RRDE and redox titration. CO2 evolution upon charging.

1.0 M LiClO4 in DMSO

10

Gold

>500

331

0.5 M LiClO4 in DMSO

10

C (Ketjenblack) or Gold

200–400

348

1.0 M LiTFSI in TEGDME

50

Super C65

500 on 1st 200 on 2nd

349

1.0 M LiClO4 in DMSO

50

Gold

1500

350

1.0 M LiClO4 in DMSO

50

nanoporous graphene

660

351

0.1 M LiPF6 in DMSO

1

GC disk Au ring

352

1.0 M LiTFSI in diglyme, diglyme: Pyr14TFSI or Pyr14TFSI 1.0 M LiTFSI in diglyme

50

CNT

343

10

C (Ketjenblack)

500

353

347

Improves round-trip efficiency and reduces overpotential. Showed better performance than LiI and TTF. No additional degradation products were observed.

Chemistry of Li-air batteries

3.41; 3.6 V 3.37; 3.2 V 3.15gc 3.74gc 3.7 V

Summary

(Continued)

Table 4

The performance of various RMs for Li-O2 battery charge.dcont'd

TEMPO

ECV, VLi Plateau

3.7 V

3.67 V and 4.32 V

3.6 V

TEMPO

3.7 V

3.5 V MeO Methoxy-TEMPO TEMPO DMPZ

5,10-Dihydro-5,10dimethylphenazine

3.35, 4.1Au

DMPZ

MPT

3.1–3.2 V

10-Methylphenothiazine

3.67?

3.6 V

Electrolyte

CRM, mM

Air electrode

Overpotential, mV References

Modified TEMPO was studied. Theoretical model derived. Structure correlates with redox potential. 1-Me AZADO is most efficient among studied compounds with E ¼ 3.6 V. Great cycling stability and discharge capacities were achieved using high concentrations of RM with solid electrolyte. Solid electrolyte also suppresses side reaction on air electrode. Conducting polymer layer suppresses Li reaction with RM. Suppression of the self-discharge of RMþ by anode protection can provide stable cycling while maintaining redox mediation at the O2 electrode. Larger capacities available with dual mediators. Observed