Advances in Sustainable Energy: Policy, Materials and Devices [1 ed.] 3030744051, 9783030744052

Table of contents :
Preface
Contents
About the Editors
Abbreviations
Chapter 1: Promising Clean Energy Development: Practice, Challenges, and Policy Implications
1.1 Fossil Fuels: A Prolog
1.1.1 Life-Cycle Assessment: Common Approaches
1.1.2 Role of Energy Storage
1.2 The Electron Economy
1.2.1 Public Policy and Discourse
1.2.2 The Hydrogen Economy
1.2.3 The Photon Economy
1.2.4 The Circular Economy
References
Chapter 2: Applications and Fundamentals of Photocatalysis with Solar Energy
2.1 Introduction
2.2 Historical Background
2.3 Applications
2.3.1 Environmental Protection
2.3.2 Clean Energy Production: Water Splitting and Reduction of Carbon Dioxide
2.3.3 Sterilization and Disinfection
2.4 Reaction System of Photocatalysis
2.4.1 Basic Principles of Photocatalysis
2.4.2 Photophysical Properties of the Semiconductor
2.4.3 Band Structure of Semiconductor
2.4.4 Optical Absorption Spectrum of Semiconductor
2.4.5 Emission Spectrum of the Semiconductor: Recombination
2.5 Potential and Charge Distribution Across the Semiconductor-Electrolyte Interface
2.5.1 Fermi Level
2.5.2 Space Charge Layer and Band Bending
2.5.3 Flat-Band Potential
2.5.4 Quasi-Fermi Level
2.5.5 Surface States
2.6 Strategies for Improving the Performance of Photocatalysts
2.6.1 Doping
2.6.2 Heterojunction/Homojunction
2.6.3 Semiconductor-Semiconductor Heterojunction
2.6.4 Metal-Semiconductor (M-S) Heterojunction
2.6.5 Carbon-Semiconductor Heterojunction
2.6.6 Homojunction
2.7 Morphology Control
2.8 Loading Cocatalysts
2.9 Conclusions and Future Trends
References
Chapter 3: Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells: Design, Fabrication, and Performance
3.1 Introduction to Nucleic Acid Hybrid Materials for Clean Energy
3.2 Optoelectronic Properties of Nucleic Acid
3.2.1 Structure and Conformation
3.2.1.1 Molecular Basis and Double-Helix Structure
3.2.1.2 Higher-Order Structures
3.2.2 Optoelectronic Properties of Nucleic Acid
3.2.2.1 Conductivity, Magnetism, and Spintronics
3.2.2.2 Optical Properties
3.3 Design Principles and Fabrication Techniques
3.3.1 Design Principles
3.3.2 Fabrication Techniques
3.4 Enhancement of Performance
3.5 Summary and Outlook
References
Chapter 4: First-Principles Calculations for the Interfaces of Perovskite Solar Cells
4.1 Introduction
4.2 Methods
4.3 CH3NH3PbI3 Phase Transformation Under Pressure
4.3.1 Lattice Structures of CH3NH3PbI3 Under Pressure
4.3.2 Projected Band Structures for Three Phases of CH3NH3PbI3
4.3.3 Variation of the Bandgap of CH3NH3PbI3 with Pressure
4.3.4 Optical Properties of CH3NH3PbI3 Under Pressure
4.4 Perovskite Solar Cell Materials Upon Doping: Ag-Doped MAPbI3: MAPb1-xAgxI3
4.4.1 Structural Properties
4.4.2 Electronic Properties
4.4.3 Optical Properties
4.5 Mn-Doped CsPbI2Br: Structural, Electronic, and Optical Properties
4.5.1 Structural Properties
4.5.2 Electronic Properties
4.5.3 Optical Properties
4.6 Surface in Perovskite Solar Cell Materials: PCBM-Adsorbed MAPbI3 Surface
4.6.1 Structure of PCBM-Adsorbed Surface
4.6.2 Adsorption Energy of PCBM on MAPbI3 Surface
4.6.3 Electronic Properties of PCBM-Adsorbed MAPbI3 Surface Model
4.6.4 Optical Absorption Properties of PCBM-Adsorbed MAPbI3 Surface
4.7 PEA+-Adsorbed MAPbI3 Surface
4.7.1 Structure of PEA+-Adsorbed Surface
4.7.2 Adsorption Energy of PEA+ on MAPbI3 Surface
4.7.3 Electronic Properties of PEA+-Adsorbed MAPbI3 Surface Model
4.8 Interfaces in Perovskite Solar Cells Materials: MAPbI3/WZ-ZnO Interface
4.8.1 Interface Structure
4.8.2 Electronic Properties
4.8.3 Optical Properties
4.9 MAPbI3/NiO Interface
4.9.1 Interface Structural
4.9.2 Electronic Properties
4.10 Au/MAPbI3 Interface
4.10.1 Interface Structure
4.10.2 Electronic Properties
4.11 MAPbI3/SnO2 Interface
4.11.1 Interface Structural
4.11.2 Electronic Properties
4.11.3 Optical Properties
4.12 CsPbI3/SnO2 Interface
4.12.1 Interface Structural
4.12.2 Electronic Properties
4.12.3 Optical Properties
4.13 CsPbI2Br/SnO2 Interface
4.13.1 Interface Structure
4.13.2 Electronic Properties
4.14 Conclusions
References
Chapter 5: Clean Hydrogen Production Technologies
5.1 Introduction
5.2 Hydrogen Production by Methane Reforming Techniques
5.2.1 Steam Reforming of Methane (SRM)
5.2.2 Partial Oxidation of Methane (POM)
5.2.3 Auto-thermal Reforming of Methane (ATR)
5.2.4 Dry Reforming of Methane (DRM)
5.3 Hydrogen Production by Coal and Biomass Gasification Techniques
5.4 Hydrogen Production by Biological Techniques
5.5 Hydrogen Production by Photochemical Techniques
5.6 Hydrogen Production by Electrolysis of Water
5.7 Challenges and Barriers in Hydrogen Production Techniques
5.8 Conclusion
References
Chapter 6: Thermal Hydrogen Compression Based on Metal Hydride Materials
6.1 Introduction
6.1.1 Metal Hydrides for Hydrogen Compression Systems
6.2 Metal Hydride-Based Thermal Compressors
6.3 Hybrid Compression Systems
6.4 Techno-Economic Analysis of Metal Hydride Compressors
6.4.1 Metal Hydride Compressor System Techno-Economic Model
6.4.2 Techno-Economic Performance of a Mini-Channel Tube Metal Hydride Compressor
6.5 Conclusions
References
Chapter 7: Oxygen Reduction Reaction Performed by Ru-Based Catalysts
7.1 Introduction
7.2 Fundamentals of Fuel Cells
7.3 Main Components of a PEM Fuel Cell
7.4 Characteristics of Fuel Cell Discharge Evaluations
7.5 Mechanism of Methanol Oxidation
7.6 The Oxygen Reduction Reaction
7.6.1 Electrocatalysts for Oxygen Reduction
7.7 Experimental
7.7.1 Preparation of the New Ru-Based Catalysts
7.7.2 Electrochemical Characterization
7.7.3 Structural and Morphological Characterization
7.7.4 DMFC Performance Evaluation
7.8 Results and Discussion
7.8.1 Synthesis of the Ru-Based Materials and Their Physical Properties
7.8.2 Structural and Morphological Characterization of the Ru-Based Materials
7.8.3 Electrochemical Study
7.8.4 Cyclic Voltammetry
7.8.5 Linear Sweep Voltammetry (LSV): Oxygen Reduction Reaction (ORR)
7.8.6 Oxygen Reduction Reaction (ORR) in the Presence of Methanol
7.8.7 Direct Methanol Fuel Cell Evaluation
7.9 Conclusions
References
Chapter 8: Direct Catalytic Low-Temperature Conversion of CO2 and Methane to Oxygenates
8.1 Introduction
8.2 Methane and CO2: Major Natural Gas Constituents
8.3 Activation of Methane and CO2: Thermodynamic Challenges
8.4 Oxygenates
8.5 Methane and CO2 to Oxygenates
8.5.1 Role of Oxidizing Agents
8.5.2 O2 as the Oxidizing Agent
8.5.3 NO as an Oxidizing Agent
8.5.4 H2O2 as an Oxidizing Agent
8.5.5 CO2 as an Oxidizing Agent
8.6 Activating Methane with CO2: Thermodynamics
8.7 Metals as Candidate Catalysts
8.8 Oxygenates from CO2 and Methane
8.9 Catalytic Deactivation
8.10 Future Perspectives
8.11 Conclusions
References
Chapter 9: Heat Transfer Analysis in Solar Thermal Collectors
9.1 Introduction
9.2 Classification of Solar Collectors
9.3 Solar Flat Plate Collectors
9.4 Evacuated Tube Collectors (ETC)
9.5 Solar Concentrating Collectors
9.6 Solar Distillation
9.7 Solar Pond
9.8 Solar Dryer
9.9 Solar Refrigeration System
9.10 Conclusions
Nomenclature
References
Chapter 10: Heat Transfer Fluids in Concentrating Solar Power Systems: Principle and Practice
10.1 Introduction: The Case for Concentrating Solar Power (CSP)
10.2 Why Use CSP to Generate Electricity?
10.3 CSP Engineering Considerations I and Installed Capacity
10.4 Thermodynamic Considerations I
10.5 CSP Engineering Considerations II
10.6 The Effect of Thermodynamic Factors on CSP Performance Efficiency
10.7 Heliostat and Engineering Consideration for Efficient Light Collection
10.8 Molten Salts as Heat Exchangers: An Overview
10.9 Particle Receivers as Heat Exchangers
10.10 Water as a Heat Transfer Fluid
10.11 The Figure of Merit as an Index of Performance Efficiency vs. Effectiveness
10.12 Thermal Oils
10.13 Organic Fluids
10.14 Alkali Molten Salts
10.15 Liquid Metals as Heat Transfer Fluids
10.16 Gases as Heat Transfer Fluids
10.17 Effect of Tube Length, Area, and Pressure on CSP Efficiency
10.18 Effect of Heat Engine Cycle and S-CO2 Parameters on CSP Efficiency
10.19 Conclusions
References
Chapter 11: Electrocatalysis for the Water Splitting: Recent Strategies for Improving the Performance of Electrocatalyst
11.1 Electrochemical Water Splitting
11.2 Mechanism of Water Splitting
11.2.1 Oxygen Evolution Reaction
11.2.2 Hydrogen Evolution Reaction
11.3 Thermodynamics of Electrochemical Water Splitting
11.3.1 Electrode Reactions
11.4 Electrode Potential and Function of Electrocatalyst
11.5 Function of Electrocatalyst
11.6 Problems Associated with Electrocatalyst
11.7 Possible Strategies to Improve the Performance of Electrocatalyst
11.7.1 Nanostructuring
11.7.2 Self-supported Electrocatalyst
11.7.3 Layered Double Hydroxide Structure
11.7.4 Introduction of Support
11.7.5 Doping of Heteroatom
11.7.6 Multi-metal Electrocatalyst
11.8 Conclusion and Future Perspective
References
Chapter 12: Future of Electrochemical Energy Storage and Its Impact on the Transition Metals
12.1 Introduction
12.2 Energy Storage Overview
12.3 Electrochemical Energy Storage
12.4 Life-Cycle Impact Analysis of Key Transition Elements
12.4.1 Cobalt
12.4.2 Copper
12.4.3 Lead
12.4.4 Manganese
12.4.5 Nickel
12.4.6 Vanadium
12.4.7 Zinc
12.5 Solar Thermal Processing: Sustainable Future
12.5.1 Zinc
12.6 Recycling End of Life Products - Circular Economy
12.7 Conclusion
References
Chapter 13: The Application in Energy Storage and Electrocatalyst of Vanadium (Based) Oxides
13.1 Introduction
13.1.1 Monovalence Vanadium Oxides
13.1.2 V2O5
13.1.3 Lithium-Ion Batteries
13.1.4 Other Metal-Ion Batteries
13.1.5 VO2
13.1.6 Lithium-Ion Batteries
13.1.7 Other Metal-Ion Batteries
13.1.8 V2O3
13.1.9 Lithium-Ion Batteries
13.1.10 Other Metal-Ion Batteries
13.2 Wadsley Phase Vanadium Oxides
13.2.1 V3O7
13.2.2 V6O13
13.3 Vanadium-Based Oxides
13.3.1 Electrocatalysis
13.4 Summary and Outlook
References
Chapter 14: Water-Stable Metal-Organic Frameworks for Water Adsorption
14.1 Fundamental Basics About Water-Stable MOFs
14.2 Prototypes of Water-Stable MOF Series
14.3 The MIL Series
14.3.1 MIL-53
14.3.2 MIL-100
14.3.3 MIL-101
14.4 The ZIF Series
14.4.1 ZIF-8
14.4.2 ZIF-68
14.4.3 ZIF-69
14.4.4 ZIF-70
14.5 The UiO Series
14.5.1 UiO-66 to UiO-68
14.6 The CAU Series
14.6.1 CAU-3
14.6.2 CAU-10
14.7 The PCN Series
14.7.1 PCN-222/MOF-545
14.7.2 PCN-224
14.7.3 PCN-228-230
14.7.4 PCN-250
14.7.5 PCN-333
14.7.6 PCN-601
14.7.7 PCN-777
14.8 The MOF-800 Series
14.8.1 MOF-801
14.8.2 MOF-841
14.9 Water-Stable MOFs for Water Adsorption
14.10 Early Stage for Structure Characterization
14.11 Investigations of Adsorption and Mechanism
14.12 Development of MOF-Based Device in Practical Applications
14.13 Outlook
References
Chapter 15: Supercapacitors: History, Theory, Emerging Technologies, and Applications
15.1 Introduction
15.2 The History of Supercapacitors
15.3 Working Principles and Classification
15.3.1 Electric Double-Layer Capacitors (EDLCs)
15.3.2 Pseudocapacitive Supercapacitors (SCs)
15.3.3 Hybrid Supercapacitors (SCs)
15.4 Electrolyte
15.4.1 Aqueous Electrolyte
15.4.2 Organic Electrolyte
15.4.3 Ionic Liquid
15.4.4 (Quasi-)Solid-State Electrolyte
15.5 Carbon Material Electrodes
15.5.1 Activated Carbon (AC)
15.5.2 Carbon Nanotubes (CNTs)
15.5.3 Graphene
15.6 Transition Metal Compounds Electrode Materials
15.6.1 Transition Metal Oxides (TMOs)
15.7 Transition Metal Carbides and Nitrides (MXene)
15.8 Other Transition Metal Compounds (Hydroxides, Sulfides, Phosphides, and Selenides)
15.9 Emerging Electrode Materials
15.9.1 Black Phosphorus (BP)
15.9.2 Metal-Organic Frameworks (MOFs)
15.9.3 Covalent Organic Frameworks (COFs)
15.9.4 Conductive Polymers (CPs)
15.10 Conclusion
References
Chapter 16: Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage
16.1 Introduction
16.2 Interlayer Structural Engineering of 2D MXene
16.2.1 Synthesis of MXenes
16.2.2 Layered Structure with an Enlarged Interlayer Spacing
16.3 Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage Applications
16.4 Supercapacitors
16.5 Li-ion Batteries/Capacitors
16.6 Na-ion Batteries/Capacitors
16.7 Other Energy Storage Systems
16.8 Conclusions and Perspectives
References
Chapter 17: The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries
17.1 Introduction
17.1.1 Overview
17.1.2 Challenges in Lithium Metal Batteries
17.1.3 Strategies to Revive LMA
17.1.3.1 The Development of 3D Micro-/Nanostructured Li host
17.1.3.2 The Development of Solid-State Electrolyte
17.1.3.3 The Development of Solid Electrolyte Interphase
17.1.4 Ex Situ-Based SEI
17.1.4.1 Physical Deposition
17.1.4.2 Chemical Deposition
17.1.5 SEI Properties and Functionality
17.1.5.1 Thickness
17.1.5.2 Transference Number
17.1.5.3 Young´s Modulus and Flexibility
17.1.5.4 Mixed Ionic/Electronic Conductor
17.1.5.5 Antioxidative
17.1.5.6 Hybrid SEI
17.1.5.7 Lithiophilic
17.1.6 Operation Under Practical Conditions
17.1.7 Conclusion and Outlook
References
Chapter 18: 3D X-Ray Characterization of Energy Storage and Conversion Devices
18.1 Introduction to X-rays
18.2 X-ray Interactions with Matter
18.2.1 Absorption
18.2.2 Scattering
18.2.3 Elastic Scattering - X-ray Diffraction
18.2.4 Total Attenuation
18.2.5 Refraction
18.2.6 Fluorescence
18.3 X-ray Sources for Characterization
18.3.1 Lab-Based X-ray Sources
18.3.2 Synchrotron Light Sources
18.4 Applications of 3D Imaging
18.4.1 Full-Field Techniques: X-ray Absorption CT
18.4.2 X-ray Phase-Contrast CT
18.4.3 Transmission X-ray Microscopy and Nano-tomography
18.4.4 X-ray Absorption Near-Edge CT
18.4.5 Diffraction Contrast Tomography
18.5 Scanning Probe Techniques
18.5.1 Coherent Diffraction Imaging
18.5.2 Ptychography
18.6 Limitations of X-rays
18.7 Future Potential and Conclusions
References
Chapter 19: In Situ Transmission Electron Microscopy for Studying Lithium-Ion Batteries
19.1 Overview
19.2 A Brief Introduction of the Lithium-Ion Battery
19.3 Material Characterization of the Li-Ion Battery Cell
19.4 Experimental Setups for In Situ Transmission Electron Microscopy (TEM)
19.4.1 Overview of TEM Methods of Battery Materials
19.4.2 In Situ TEM of Battery Materials
19.4.2.1 Open-Cell Configuration
19.4.2.2 Liquid Chemistry Open Setup
19.4.2.3 Solid Chemistry Open Setup
19.4.2.4 Advantages and Disadvantages Between Liquid and Solid Setup
19.4.2.5 Sealed Liquid-Cell Configuration
19.5 Discussion
19.6 Application of In Situ TEM to LIBs
19.6.1 Reactions at the Electrode Interface
19.6.2 Study of Electrode Material Degradation
19.6.3 Intercalation Mechanism
19.6.4 Alloying Mechanism
19.6.5 Conversion Mechanism
19.7 The Beam Effects of In Situ Liquid TEM
19.7.1 Beam Effects in Open-Cell Configurations
19.7.2 Beam Effects in Liquid-Cell Configurations
19.8 Conclusions
References
Chapter 20: Clean Coal Conversion Processes-The Present and Future Challenges
20.1 Introduction: The Case for Clean Coal
20.2 The Relationship Between Energy Generation and Environmental Pollution
20.3 The Capture of Soluble Oxides of Carbon, Nitrogen, and Sulfur (SONS/SOCNS) Using Lewis Bases
20.4 The Capture of Soluble Oxides of Nitrogen and Sulfur (SONS) Using Lewis Bases: Engineering Considerations
20.5 The Capture of Carbon Dioxide Using Lewis Bases: Energy Considerations
20.6 The Capture of Carbon Dioxide Using Membrane: Engineering Considerations
20.7 The Capture of Carbon Dioxide Using Membrane: Energy Considerations
20.8 Outlook
20.9 Conclusion
References
Chapter 21: Coal Gasification with Exergy Recuperation and CO2 Recovery
21.1 Introduction
21.2 Coal Gasification
21.3 Gasifying Agents
21.4 Different Types of Gasifier
21.5 Exergy Recuperation
21.6 Carbon Dioxide Recovery in Coal Gasification
21.7 Removal of Sulfur, Nitrogen, and Particulate Matters
21.7.1 Sulfur Removal
21.7.2 Nitrogen Removal
21.7.3 Tar Removal
21.8 Advanced Gasification Systems
21.8.1 IGCC/IGFC System
21.8.2 Advanced IGCC/IGFC Systems with Exergy Recovery
21.8.3 Fischer-Tropsch Synthesis System
21.8.4 Urea Synthesis System
21.9 Conclusions and Outlook
References
Chapter 22: Lignin and Lignocellulosic Materials: A Glance on the Current Opportunities for Energy and Sustainability
22.1 Introduction
22.2 Chapter Taxonomy
22.3 Bioresources: From Biomasses to Advanced Materials
22.4 Lignin
22.4.1 Nature and Chemical Structure
22.4.2 Lignin Processing
22.5 From Basic Material to Multifunctional Applications
22.5.1 Production of Biofuels
22.5.2 Lignin-Based Polymers
22.5.3 Lignin as a Carbon Precursor
22.6 Micro- and Nanoscale Applications
22.7 Summary and Future Perspectives
References
Chapter 23: Municipal Solid Waste Incineration and Sustainable Development
23.1 Introduction
23.2 Incineration Process
23.3 Environmental Aspects
23.3.1 Air Pollution
23.3.1.1 Pollutants
23.3.1.2 Prevention and Control Systems
23.3.2 Solid Waste Management
23.4 Economic Aspects
23.4.1 Annual Capital Costs
23.4.2 Annual Operational Costs
23.4.3 Revenues
23.4.4 Externalities
23.5 Social Aspects
23.6 Conclusion
References
Chapter 24: Microbial Fuel Cells: Design and Evaluation of Catalysts and Device
24.1 Introduction
24.2 Synthesis and Characterization of Cathode and Anode Catalysts
24.3 Review of Cathode Materials
24.4 Review of Anode Materials
24.5 Catalyst Synthesis Methods
24.5.1 Solgel Synthesis
24.5.2 Microemulsion Synthesis
24.5.3 Hydrosolvothermal Chemistry
24.5.4 Chemical Vapor Deposition
24.6 Development and Modification of Exchange Membrane
24.6.1 Cation Exchange Membrane (CEM): Synthesis and Characterization
24.6.1.1 Poly(vinylidene fluoride) (PVDF): Synthesis and Characterization
24.6.1.2 Poly(phenylene oxide) (PPO): Synthesis and Characterization
24.6.1.3 Anion Exchange Membrane (AEM): Synthesis and Characterization
24.6.1.4 Quaternized Polysulfone (QPSU): Synthesis and Characterization
24.6.1.5 Polyvinyl Alcohol (PVA): Synthesis and Characterization
24.7 Construction and Evaluation of Microbial Fuel Cells
24.7.1 Fundamental Definitions
24.7.2 Effect of Electropolymerization
24.7.3 Effect of Catalyst Modification
24.8 Understanding and Postulations of Electron Transfer Mechanisms
24.8.1 Biofilm Anode and Electron Transfer
24.8.2 Anode Microbiology
24.8.3 Cathode ORR Mechanism
24.9 Conclusion
References
Chapter 25: Observation on Comprehensive Energy Trend
25.1 Introduction
25.2 Industrial Evolution
25.3 Industry Status
25.4 Trend Observation
25.5 Future Focus
25.6 Conclusions
References
Chapter 26: Smart Energy Trend Observation
26.1 Smart Energy Trend Observation
26.1.1 Industry Evolution
26.1.2 Industry Status
26.1.2.1 Intellectualized Power Transmission and Distribution
26.1.2.2 Intelligent Power Distribution
26.1.2.3 Intelligent Power Consumption
26.1.2.4 Progress of Demonstration Projects
26.1.3 Smart Technology
26.1.3.1 Power IoT Technology
26.1.3.2 Energy Blockchain Technology
26.1.3.3 Smart Energy Technology
26.1.3.4 Smart Power Plant Technology
26.1.4 Practical Issues
26.1.4.1 Source Network Load Storage Multiphase Coordinated Dispatch Control Project
26.1.4.2 Intelligent Power Distribution House Project
26.1.4.3 Energy Blockchain Project
26.1.4.4 Megacity Grid Energy Internet Demonstration Project
26.1.5 Trend Observation
26.2 Conclusion
References
Chapter 27: Postface: Conclusion on Renewable Energy Strategies for a Sustainable Future: Part A: Role of Energy Storage
27.1 Introduction
27.2 Renewable Energy
27.3 Electrochemical Energy Storage
27.4 Concluding Remarks
27.5 Description
27.6 Key Features
27.7 Readership
References
Index

Citation preview

Yong-jun Gao Weixin Song Jingbo Louise Liu Sajid Bashir   Editors

Advances in Sustainable Energy Policy, Materials and Devices

Advances in Sustainable Energy

Yong-jun Gao • Weixin Song • Jingbo Louise Liu • Sajid Bashir Editors

Advances in Sustainable Energy Policy, Materials and Devices

Editors Yong-jun Gao Hangzhou Branch of the Zhejiang Tsinghua Yangtze River Delta Research Institute Clean Energy and Energy Conservation and Environmental Protection Centre Hangzhou City, China Jingbo Louise Liu Texas A&M University – Kingsville Kingsville, TX, USA

Weixin Song Department of Materials University of Oxford Oxford, UK Sajid Bashir Department of Chemistry Texas A&M University – Kingsville Kingsville, TX, USA

ISBN 978-3-030-74405-2 ISBN 978-3-030-74406-9 https://doi.org/10.1007/978-3-030-74406-9

(eBook)

© The Editor(s) (if applicable) and The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 This work is subject to copyright. All rights are solely and exclusively licensed by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Switzerland AG The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Preface

At the time of starting this book with Springer Nature, I lost my doctorate advisor (‘academic father’), and at the start of writing for the actual monograph, I lost my biological father from severe acute respiratory syndrome (SARS), a ribonucleic acid encoded virus with a corona geometry. At present SARS is probably the most wellknown phrase and has affected everyone in all sectors of society. “To lose one parent, Mr. Worthing, may be regarded as a misfortune; to lose both looks like carelessness.”―Oscar Wilde, The Importance of Being Earnest. Like, Mr. Wilde, I reflect on, where we are as a society. During this reflection period, when most of us are working remotely, teleconferencing, or are limited in how many people can work in any laboratory, we have to ask is Sustainable Energy Necessary? and if yes, what is the likely energy production roadmap over the next five decades? This monograph aims to address this central thematic question and is split into several themes, relating to energy generation from coal, hydrogen, thermal heat, light, electrons, and associated storage to ensure a continuous supply of electricity as demanded by industry, residual, home and businesses. The authors cover themes related to the above mission and in this preface as editors will summarize and expand the scope (1-Bashir) contained within the pages of this monograph. The first three chapters deal with photocatalysis concerning solar energy (2- Li/Kong), hybrid nature-derived mimetics (3-Bai/Ran), and the first principle-based density level theory calculations (4-Tang). The next four chapters deal with the production of hydrogen or oxygen or their utilization, with clean hydrogen production technologies (5-Abdullah), from metal hydride (6-Corgnale) and oxygen reduction using ruthenium electrocatalysts for direct methanol fuel cell devices (7-Uribe-Godínez) and direct conversion of carbon dioxide to oxygenates (8-Spivey). We return to the theme of solar energy is respect to solar thermal fluids thermodynamics (9-Kalita), modeling (10-Fox/Bashir), and electrolysis (11-Haik).

v

vi

Preface

The theme of energy storage is dealt with in the next two chapters with a focus on transition metal-based electrochemical storage (12-Huda) and an emphasis on mixed vanadium oxides (13-Li). We then explore an application of certain metal oxides in the form of metal-organic frameworks for water adsorption (14-Wang), advanced materials for supercapacitors (15-He), or metal layered two-dimensional MXenes for electrochemical storage as a departure from traditional metal oxides, addressed earlier in the monograph (16-Luo). The role of solid-electrolyte in high-density lithium-ion batteries is discussed next (17-Qiao) and the modern characterization approach to better understand energy materials is further explored as a standalone topic for X-ray energy (18-Tan/Shearing) or advanced transmission electron microscopy(19-Robertson). We then break the next segment to discuss clean coal (20-Meyer), coal and biomass gasification (21-Guan), energy from lignocellulosic materials (22-El-Azazy/Shibl), and how these materials can be used for sustainable development. Two examples are used to illustrate this approach. The role of Municipal solid waste incineration (23-Quina) and carbon sugars as fuels for microbial fuel cells, both theory and actual practice is extensively reviewed (24-Liu). The last segment policy, energy materials and electrification in china in two perspectives, one from the application of artificial intelligence in the monitoring of smart devices and applications such as electric vehicles (25-Gao) and the application of blockchain, smart devices and photovoltaic/wind/hydrogen stored energy in a delocalized grid as a framework for renewables energy (26-Gao) and the monograph ends with a general review of themes covered by one of our editors (27-Song) with the back matter subject index. All authors justify the role of energy storage or application of sustainable energy to lessen the influence of carbon dioxide that is derived from the combustion of coal to generate electricity. Fossil fuels such as coal and natural gas are the mainstay for electricity generation both within the United States and globally as observed in figure 1, however at the cost of reliable and relatively cheap electricity from coal is the expense of emissions from soluble oxides of carbon, nitrogen, and sulfur, as well as particulate matter. This has an environmental cost associated with it and will contribute towards respiratory distress, which in turn will enhance the susceptibility of a slice of the population towards SARS. Lastly, should carbon be used to generate thermal heat or used to manufacture materials, drugs, plastics, paints and other useful resources is a question that is being asked to generate a circular economy where waste is minimized and each atom is utilized, a form of conversation of matter?. The authors within this monograph highlight a few of the approaches that are currently under development at the technical level of attainment. The review and case studies are broad because the challenges are diverse, such as the development of clean biofuels for aviation, transportation and heating and cooling, waste to energy conversion technologies, a value-added material transformation from waste and life cycle assessment. On the proceedings page, we as the editorial team will further develop some areas that we consider emerging and pivotal towards the development of strategies to engineering a future with zero carbon emissions and cheaper energy resources that are sustainable.

Preface

vii

Any project, be it a greeting card, a book chapter, a review article, or a grant proposal is an intense endeavor and almost always a team effort; this monograph is no different, which we acknowledge here. Charlotte Hollingworth, for the day-today management, pre-print layout of the various manuscripts and Solomon, Cynthia, and Sharon also at Springer (Springer Science+Business Media, LLC) for taking an idea based on our work within the American Chemistry Society (ACS) division of energy and fuels (ENFL, Liu and Bashir), and support from Texas A&M UniversityKingsville. Lastly, we pay our respects to Peter J Derrick and Mohammed Bashir who died recently. A short biography can be found here for Dr. Derrick at https://doi.org/ 10.1177/1469066717739174. Thank you, Peter, and Dad your legacy will endure. Kingsville, TX, USA

Sajid Bashir

Kingsville, TX, USA

Sai Chava

Oxford, UK Hangzhou City, China Kingsville, TX, USA

Weixin Song Yong-jun Gao Jingbo Louise Liu

Contents

1

2

3

4

Promising Clean Energy Development: Practice, Challenges, and Policy Implications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sajid Bashir, Sai Chava, Weixin Song, Yong-jun Gao, and Jingbo Louise Liu

1

Applications and Fundamentals of Photocatalysis with Solar Energy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Qiuyang Huang, Dan Kong, and Yongdan Li

27

Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells: Design, Fabrication, and Performance . . . . . . . . . . . . . . . . . . . . . . Dan Bai, Huhu Feng, Xingchen Yu, Chenxin Ran, and Wei Huang

67

First-Principles Calculations for the Interfaces of Perovskite Solar Cells . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Jun-Peng An, Ying Tian, Hong-Tao Xue, Jun-Chen Li, Jun-Qiang Ren, Xue-Feng Lu, and Fu-Ling Tang

95

5

Clean Hydrogen Production Technologies . . . . . . . . . . . . . . . . . . . . 159 Mohammad Yusuf, Mohamad Sahban Alnarabiji, and Bawadi Abdullah

6

Thermal Hydrogen Compression Based on Metal Hydride Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171 Claudio Corgnale, Robert C. BowmanJr., and Theodore Motyka

7

Oxygen Reduction Reaction Performed by Ru-Based Catalysts . . . . 193 J. Uribe-Godínez and A. Altamirano-Gutiérrez

8

Direct Catalytic Low-Temperature Conversion of CO2 and Methane to Oxygenates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227 Ashraf Abedin and James J. Spivey

ix

x

Contents

9

Heat Transfer Analysis in Solar Thermal Collectors . . . . . . . . . . . . 251 Pankaj Kalita, Dudul Das, Samar Das, Rabindra Kangsha Banik, and Urbashi Bordoloi

10

Heat Transfer Fluids in Concentrating Solar Power Systems: Principle and Practice . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 279 Elise B. Fox, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

11

Electrocatalysis for the Water Splitting: Recent Strategies for Improving the Performance of Electrocatalyst . . . . . . . . . . . . . . 315 Tanveer ul Haq and Yousef Haik

12

Future of Electrochemical Energy Storage and Its Impact on the Transition Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 341 Nazmul Huda and Shahjadi Hisan Farjana

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based) Oxides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 359 Yue Wang, Fan Li, Sajid Bashir, and Jingbo Louise Liu

14

Water-Stable Metal-Organic Frameworks for Water Adsorption . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 387 Xuan Wang and Charles Lee

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 417 Yiyang Liu, Paul R. Shearing, Guanjie He, and Dan J. L. Brett

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . 451 Jianmin Luo

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 479 Rajesh Pathak, Yue Zhou, and Qiquan Qiao

18

3D X-Ray Characterization of Energy Storage and Conversion Devices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 513 Chun Tan, Andrew S. Leach, Thomas M. M. Heenan, Rhodri Jervis, Dan J. L. Brett, and Paul R. Shearing

19

In Situ Transmission Electron Microscopy for Studying Lithium-Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 545 Chen Gong, Shengda Pu, and Alex W. Robertson

20

Clean Coal Conversion Processes–The Present and Future Challenges . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 571 E. Gerald Meyer, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

Contents

xi

21

Coal Gasification with Exergy Recuperation and CO2 Recovery . . . 593 Zhongkai Zhao, Yohanes Andre Situmorang, Atsushi Tsutsumi, Xiaogang Hao, Abuliti Abudula, and Guoqing Guan

22

Lignin and Lignocellulosic Materials: A Glance on the Current Opportunities for Energy and Sustainability . . . . . . . . . . . . . . . . . . 621 Marwa El-Azazy, Sajid Bashir, Jingbo Louise Liu, and Mohamed F. Shibl

23

Municipal Solid Waste Incineration and Sustainable Development . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 653 Beatriz Sales Bandarra and Margarida J. Quina

24

Microbial Fuel Cells: Design and Evaluation of Catalysts and Device . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 681 Sajid Bashir, Shawn P. Mulvaney, William Houf, Luis Villanueva, Zhaohui Wang, Gregory Buck, and Jingbo Louise Liu

25

Observation on Comprehensive Energy Trend . . . . . . . . . . . . . . . . 765 Rui Guan, Yunchuan Wang, Sai Raghuveer Chava, Jingbo Louise Liu, Sajid Bashir, and Yong-jun Gao

26

Smart Energy Trend Observation . . . . . . . . . . . . . . . . . . . . . . . . . . 797 Ran Wei, Yong-jun Gao, Zhihua Wu, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

27

Postface: Conclusion on Renewable Energy Strategies for a Sustainable Future: Part A: Role of Energy Storage . . . . . . . . 839 Weixin Song, Yong-jun Gao, Sajid Bashir, and Jingbo Louise Liu

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 847

About the Editors

Yong-jun Gao, a Postgraduate trained and is at present the Deputy Director of Clean Energy and Energy Conservation and Environmental Protection Center, Zhejiang Yangtze River Delta. Since 2001, he has been doing research and work in the fields of energy and environmental protection. He has won the second academic prize of China Nuclear Energy Association, the third academic prize of China Electricity Council, and the second prize of the science and technology management innovation of State Power Investment Corporation Limited. He has published many articles as the first author or co-author in publications such as China Nuclear Power, Nuclear Power Engineering, China Power Enterprise Management, East China Electric Power, and Nanostructured Materials for nextgeneration Energy Storage and Conversion. He was awarded the model worker and the May Day medal by a Chinese city.

Weixin Song is an MPLS Enterprise Fellow and Postdoctoral Research Fellow at the University of Oxford. He received his Ph.D. in Material Electrochemistry from Imperial College London in 2019 funded by the President’s Ph.D. Scholarship. He has been a member of the Centre for Doctoral Training (CDT) of Advanced Characterization of Materials and the London Centre for Nanotechnology (LCN). He completed his BSc in 2012 and MSc in 2015 from Central South University, China. He has research interests in materials electrochemistry and characterization with materials xiii

xiv

About the Editors

application in energy storage, electrocatalysis, and photo electrocatalysis. He is particularly interested in the materials and interface studies using electron microscopy and spectroscopy. He has 22 first-author publications and 18 licensed patents. He has contributed chapters on one book published by Springer. He received the highest student honor of Central South University (2014), STFC Future Early Career Award (2016), Armourers & Brasiers Rolls-Royce Scheme Grant (2018), National award for outstanding selffinanced students abroad (2018), MPLS Enterprise Fellowship (2019) and Constance Fligg Tipper Centenary Memorial prize (2020).

Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing (China) in 2001. She completed her postgraduate training at the University of Calgary in Alberta (Canada) in electrochemistry analysis of nanomaterials. She is at present a Full Professor at Texas A&M University-Kingsville (TAMUK, USA) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry, a Vebleo Fellow, and a Fellow of the International Association of Advanced Materials Society. She also holds Chartered Chemist status and separately Chartered Scientists status from the Science Council, UK. She also holds DEBI Faculty Fellow status at the US Air Force Research Laboratory and is a past JSPS IF (Japan) and FFSI in Israel awardee recipient respectively. She has authored and co-authored, books, book chapters, and peer-reviewed journal articles (> 100). During her 12.5-year services at TAMUK, in chemistry, she has taught > 8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and /or participated in the projects (> 40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Program Chair, Chair-Elect, and Chair of Energy and Fuels of the American Chemical Society.

About the Editors

xv

Sajid Bashir received his Ph.D. training in matrixassisted laser desorption/ionization time-of-flight mass spectrometry from the University of Warwick (UK) in 2001 and previously graduate training in Fourier transform ion cyclotron resonance mass spectrometry from the University of New York at Buffalo (USA). He was a postgraduate research associate at Cornell University (USA) in the field of plant proteomics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a past Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has co-authored > 80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry, and also holds Chartered Chemist and separately Chartered Scientists from the Science Council (UK). He is also the American Chemical Society, Energy and Fuels Division Technical Secretary (2018–2022). During his service at TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with Risk Management. Currently, he collaborated with local law enforcement as a consultant in forensic chemistry.

Abbreviations

δ α λ Φ θ Φfb ΔT φ (A)ZIB abs AC ACN AFC AOR ASU ATR BFB BP CAES CB CBM CCS CCS CCT CE CFB CFS CH CHP CLG CNF

Delta Elevation Angle Wavelength Polarization Potential Theta Flat Band Potential Charging/ Discharging Cycle Time Latitude Of The Collector (Aqueous) Zinc Ion Batteries Absorption Activated Carbon Acetonitrile Alkaline Fuel Cell Alcohol Oxidation Reaction Air Separation Unit Auto-thermal Reforming of Methane Bubbling Fluidized Bed Black Phosphorus Compressed Air Energy Storage Conduction Band Conduction Band Minimum Carbon Capture and Storage CO2 Capture and Storage Clean Coal Technology Coulombic Efficiency Circulating Fluidized Bed Capacitive Faradic Storage Helmholtz layer capacitance Combined Heat, and Power Chemical Looping Gasifier Carbon Nanofibers xvii

xviii

CNT CNT CO CO2 COF COS CP CQD CRM CS2 CSC CSC CSP CT CTMA CV CVD DBCFB DEPG des DFT dI DL DMC DMF DMFC DOS dx E Ec EC EDC EDLC EF EMC EPRI ESPW ESR ETC EU EUV Ev FEC FOM

Abbreviations

Carbon Nanotubes Carbon Nanotubes Carbon Monoxide Carbon Dioxide Covalent Organic Frameworks Carbonyl Sulfide Conductive Polymers Carbon Quantum Dots Critical Raw Materials Carbon Disulfide Charge layer capacitance China Scholarship Council Concentrated Solar Power Computed Tomography Cetyltrimethylammonium Chloride Cyclic Voltammetry Chemical Vapor Deposition Dual-Bed Circulating Fluidized Bed Dimethyl Ether of Polyethylene Glycol Desorption Density Functional Theory Decrement of the Light Intensity Double Layer Dimethyl Carbonate N, N-Dimethylacetamide Direct Methanol Fuel Cell Density Of States Film Thickness Occupied by the electrons CB Energy Ethylene Carbonate Ethylene Dicarbonate Electric Double-Layer Capacitors Fermi Level Enterprise Management Console Electric Power Research Institute Electrochemical Stable Potential Window Equivalent Series Resistance Evacuated tube collectors European Union Extreme Ultraviolet VB Energy Fluoroethylene Carbonate Figure Of Merit

Abbreviations

FPC FST GBL G-C-S GHG GO GPL GTE H 2O H 2S HCN HDFC HER HF HHV HOMO HOR HRSG HTFT HTMM HyPr-RING hν I IBA ICT IEA IFA IGCC IGFC IHP IoT IRENA ISE k KIB KPI LAB LCI LDH LIB LMA LMB LSV LUMO

xix

Flat Plate Collectors Foundation for Science and Technology γ-ButyroLactone Gouy-Chapman-Stern Greenhouse Gases Graphene Oxides Gel Polymer Electrolyte Gross Thermal Efficiencies Water Hydrogen Sulfide Hydrogen Cyanide Housing Development Finance Corporation Hydrogen Evolution Reaction Heating Fluid Higher Heating Value Highest Occupied Molecular Orbital Hydrogen Oxidation Reaction Heat Recovery Steam Generator High-Temperature Fischer-Tropsch Synthesis High-Temperature Mixing Method Hydrogen Production by Reaction-Integrated Novel Gasification Photon Energy Light Intensity Bottom Ash Information and Communications Technology International Energy Agency Fly Ash Integrated Coal Gasification Combined Cycle Integrated Coal Gasification Fuel Cell Inner Helmholtz Plane Internet of Things International Renewable Energy Agency Inorganic Solid Electrolytes Boltzmann Constant Potassium Ion Batteries Key Performance Indicators Lead-Acid Batteries Life Cycle Impact Layered Double Hydroxides Lithium-Ion Batteries Lithium Metal Anode Lithium Metal Batteries Linear Sweep Voltammetry Lowest Unoccupied Molecular Orbital

xx

MCFC MEA MeOH METI MEXT MH MIEC MO MOF MSW Nc NEA NFCS NMP Nv OER OHP OOH OOR ORR PAFC PAH PANI PC PC PCB PCBM PCDD PCDF PCE PCM PCP PDDA PEC PEMFC PHES PM PPy PSC PTFE PVA QFL RDE

Abbreviations

Molten Carbonate Fuel Cell Membrane-Electrode Assemblies Methanol Ministry of Economy, Trade, and Industry Ministry of Education, Culture, Sport, Science, and Technology of Japan Metal Hydride Mixed Ionic/electronic Conductors Methyl Orange Metal-Organic Frameworks Municipal Solid Waste Effective Densities of CB National Energy Administration Non-Faradaic Capacitive Storage N-Methyl-2-Pyrrolidone Effective Densities of VB Oxygen Evolution Reaction Outer Helmholtz Plane Adsorbed (per) Hydroxide species Oxidation of Organic Reactant Oxygen Reduction Reaction Phosphoric Acid Fuel Cell Polycyclic Aromatic Hydrocarbons Porous Nano flower Polyaniline Electrode Propylene Carbonate Propylene Carbonate Polychlorinated Biphenyls Phenyl-C61-Butyric Acid Methyl Ester Polychlorinated Dibenzo-p-Dioxins Polychlorinated Dibenzofurans Power Conversion Efficiency Phase Change Material Porous Coordination Polymers Poly Diallyl Dimethylammonium Chloride Photoelectrochemical Proton Exchange Membrane Fuel Cell Pumped Hydro Energy Storage Particulate Matter Polypyrrole Perovskite Solar Cells Ethylene Polytetrafluoride Polyvinyl Alcohol Quasi-Fermi Level Rotating Disk Electrode

Abbreviations

RDS ROS SARS SC SDC SEI SGSP SHE SIB SNCR SOFC SOHIO SPE SRM SSE STEM STP TBCFB TES TFE TGA TiO2 TMB TMO TOC TXM UPD VB VBM VOC

xxi

Rate-Determining Step Reactive Oxygen Species Severe Acute Respiratory Syndrome Supercapacitors Static Dielectric Constant Solid Electrolyte Interphase Salt Gradient Solar Pond Standard Hydrogen Electrode Sodium-Ion Batteries Selective Noncatalytic Reduction Solid Oxide Fuel Cell Standard Oil of Ohio Solid Polymer Electrolyte Steam Reforming of Methane Solid State Electrolyte Scanning Transmission Electron Microscopy Standard Temperature and Pressure Triple-Bed Combined Circulating Fluidized Bed Total Energy Supply TetraFluoroethylene Copolymer Thermo Gravimetric Analysis Titanium Dioxide Transition-Metal Based Transition Metal Oxides Total Organic Carbon Transmission X-ray Microscopy Under-Potential Deposition Valence Band Valence Band Minimum Volatile Organic Compounds

Chapter 1

Promising Clean Energy Development: Practice, Challenges, and Policy Implications Sajid Bashir, Sai Chava, Weixin Song, Yong-jun Gao, and Jingbo Louise Liu

Dedicated to Dr. Peter J Derrick and Mr. Mohammed Bashir, Rest in Peace.

1.1

Fossil Fuels: A Prolog

To achieve this conversation of mass principle where atoms are converted from one form to another, waste heat and energy should be recaptured and retasked, and converted to value-added materials. As is known in freshmen chemistry, one mole of carbon generates one mole of carbon dioxide, using methane as an example. Coal

S. Bashir (*) Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA e-mail: [email protected] S. Chava Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA Emergent Biosolutions, San Diego, CA, USA e-mail: [email protected] W. Song Department of Materials, University of Oxford, Oxford, UK e-mail: [email protected] Y.-j. Gao Center for Clean Energy and Energy Conservation and Environmental Protection, Zhejiang Yangtze River Delta, Hangzhou, People’s Republic of China Sunshine Times Law Firm, Hangzhou, People’s Republic of China e-mail: [email protected] J. L. Liu Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA Texas A&M Energy Institute, College Station, TX, USA e-mail: [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_1

1

2

S. Bashir et al.

being a multimeric component generates multiple moles of carbon dioxide, as well as nitric oxides and oxides of sulfur: 4 C85 H5 O7 N2 S2 þ 351 O2 ! 340 CO2 þ 8 NO2 þ 8 SO3 þ 10 H2 O Combustion of bituminous coal will generate approximately 85 moles of CO2 per mole of coal or 24 megajoules per kilogram or 1000 g of CO2eq for each kWh generated, while natural gas is approximately at 500 g CO2eq per kWh. Therefore, technologies that improved combined heat and power such as recapture heat from waste incineration will result in lower consumption of coal and natural gas. In general, the electricity demand has increased annually globally by 3% although, with the more advanced infrastructure of the United States, the increase is less, there is an increased demand in energy generation from a mix of resources, principally coal and Natural gas, followed by hydro and nuclear [1]. Since 2010, there has been a renewed emphasis on a transition to renewable energy (Fig. 1.1), where the first transition was going from wood/peat moss locally to coal enabling the early societies to become more mobile. In the modern preindustrial era, coal becomes the engine for change, which enables the Romans, Greeks, and Chinese to use coal for heat and forging of weapons. The British industrial revolution was powered by coal to provide heat for steam engines, buildings, and generate electricity and manufacture steel. The second transition is in the early 1970s as the economy transitioned to more natural Gas due to the Oil

Fig. 1.1 Summary of energy generation by nuclear, petroleum distillates, and renewable resources (left-axis) and fossil fuels (Coal, and Natural Gas, right axis; [2–6])

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

3

Crisis and to lower carbon dioxide emissions. We are in the third type of transition to a more sustainable form of generating energy [7]. A levelized cost of electricity from different energy resources by the Institute of Energy Research suggests that retrofit of current coal and gas power plants to extend their useless operating lifespans by 15 years could yield near-zero carbon dioxide emissions as the most reliable and cheapest form of energy production. Under such a scenario, soluble oxide of carbon, nitrogen, and sulfur are captured and carbon dioxide is stored in geological sites suited for long-term storage [8]. The current general-purpose analyses reported in the literature fail to connect the levelized cost of energy to the reduction of particulate matter, solid oxides of carbon, nitrogen, and sulfur, and second the cost to implement these technologies and their lifecycle energy and emission [9]. This we believe is necessary before the first step before an environmental roadmap can be fully developed and deployed [10] (Fig. 1.2). In the area of energy capture and recycling, that waste to biofuels will take an increasing role in the next 50 years as a necessary aid to transition away from the utilization of coal for the generation of electricity. There are two broad reasons: One, these resources currently exist, do not need to be mined or extracted but are generated as a result of industrial, agricultural, and process chemistries, and are essential ‘lost’ to the economy. If these resources can be integrated, waste is minimized, and the extracted energy can be applied to the grid [13]. Second, the ‘lost energy’ is currently ‘replaced’ by coal and natural gas and this will enable lesser dependence on coal. Examples of atom-conversion are the generation of biofuels

Fig. 1.2 Summary of emissions of common volatiles and particulate matter (left axis) and the price of natural gas and coal on a megawatt-hour basis (right axis; [11, 12])

4

S. Bashir et al.

from nonfood feedstocks, such as nonedible seed oil, wood chips, grape marc, chicken/cow manure, fruit peel, rotten plant matter, or recycling of municipal solid waste by a thermal process, followed by aerobic respiration, fractionation using a biorefinery process to generate bioethanol and char. These could be supported using clay-based heterogeneous catalysis and a porous metal–organic framework to capture and store biogas [14–15]. Where the carbon source is lignocellulose biomass, the typical pretreatment is acid wash and digestion; alkaline hydrogen peroxide application may also achieve similar goals with lower energy balance and lesser useless by-products. Under such a scenario, lignocellulose biomass is pretreated with hydrogen peroxide [16]. As hydrogen peroxide has a greater oxidative potential than sulfuric acid, longer incubation periods may yield phenomenologically similar results in terms of delignation followed by hydrolysis described by Eq. 1.1: 

E : H2 O2 þ 2Hþ þ 2e ! 2H2 O, E ¼ 1:78 V vs:S2 O8 2 þ 2e 

! 2SO4 2 , E : 2:01 V

ð1:1Þ

The waste can also be heated after pretreatment or subjected to co-digestion through controlled pyrolysis to generate biochar. As biochar exhibits a high carbon content, its caloric values are like coal and could be used for heat and power generation or converted to use materials such as activated carbon or porous storage materials. In the area of biofuels, lignocellulosic organics from olive trees or organic manure could be treated with solid weak acids using heterogeneous catalysts and powered by concentrated solar power to generate biofuels, or if anaerobic respiration is used to generate lactic acid as a precursor for the synthesis of poly-lactic acid polymers from waste by-products, adding value back to the supply chain [17]. Like solid waste, wastewater treatment is a high-energy process particularly when secondary sludge is formed, the degree of microbial activity is still high. Wastewater microbial-rich sludge could be subjected to anaerobic respiration and converted to inorganic salts consisting of soluble oxides of nitrogen and phosphorus, which could be used as renewable fertilizers, the organic digestates as biogas.

1.1.1

Life-Cycle Assessment: Common Approaches

The life-cycle assessment should account for environmental, economic, and technical parameters in the application of any energy production platform, summarized in Fig. 1.3. Examples of environmental parameters include generation of global warming potential, acid rain and acidification potential, eutrophication potential, and cumulative energy demand. In the area of economy, the unit cost per megawatt-hour, operational life-cycle, infrastructure costs, and power plant operational costs

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

5

Fig. 1.3 Block diagram summary of factors that underpin life-cycle assessment cradle-to-grave type of analysis for energy generation (fossil fuels) and energy storage (sustainable energy resources, [18])

including the cost of decommissioning and the technical parameters are the cost of feedstocks, generating capacity, generating reliability, technological maturity of the proposed process, and transmission and storage potential. An example of the above approach is when an analysis of electrical storage is conducted on the ‘short-term’ (< 1 min), ‘medium-term’ (~ 270 min), or ‘long-term’ (> 130,000 min). By using the general equation to determine the levelized cost of electricity as described by Eqs. 1.2 and 1.3: ðCapC þ OpC  AFÞ AF  W

ð1:2Þ

1  ð1 þ itsÞ1 , its

ð1:3Þ

where AF ¼

where capital expenditures (CapC, $), operational expenditures (OpC, $) including the balance of plant cost, W is the annualized energy output from the storage system (kWh). The energy output is dependent upon the number of cycles (n)  discharge power (kW) and discharge per cycle time (ms), AF ¼ annuity factor (for l at fixed 5%), its ¼ is the cost of in-state discounts or tariffs, or subsidies ($) and l is the lifetime of the storage system (years) in increments of 1kWh and 8-216 GWh used from storage systems. The cycles for storage depend on a short–medium or long time scale, with 40 (thermal or isothermal energy storage) or 100 (adiabatic energy storage) cycles

6

S. Bashir et al.

Fig. 1.4A Summary of energy storage by the form of storage (Mechanical, electrical, chemical, or thermal in nature, [20]) Performance Costs at 100 MW 3.5

LCOE (LOG Skwh)

3

Pump Hydro Medium-term Storage

2.5

Compressed Air Energy Medium-term Storage Pump Hydro long-term Storage

2

Battery Medium-term Storage

1.5

H2/CH4 to Power Long-term Storage H2/CH4 to Power Medium-term Storage

1

Battery Short-term Storage

0.5

Compressed Air Energy Short-term Storage Pump Hydro Short-term

0 0 –0.5

50

100

150

200

250

300

350

400

450

500

H2/CH4 to Power Short-term Storage

CO2 Emissions (g CO2e kWh)

Fig. 1.4B Summary of LCOE for different systems for fast, medium, and long-term energy retrieval on a log basis with area based on modal value. (Data from Table 3.4 of [11, 19])

per day for 1 year and cost factors from [19] (Data from Table 3.4 of [19 and Appendices of 11]) and results summarized in Fig. 1.4A by energy type, Fig. 1.4B by LCOE and Table 1.1 [51].

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

7

Table 1.1 Summary of different storage systems in terms of power, specific energy, efficiency, discharge capacity, the response time (fast is seconds to minutes), lifetimes (years or discharge cycles), stability in self-discharge tendency, storage costs, power conversion system (PCS) cost, and balance of plant (BOP) associated costs [21]

Storage technology Pumped hydro storage Conventional compressed air energy storage Lead acetate (nickel cadmium, if different) Lithium ion Hydrogen Thermal Storage technology

Power (MW) 100–5000 5–3000

0–40

Specific energy (Wh/kg) 0.5–1.5 30–60

Efficiency (%) 65–80 41–75

Discharge at capacity (h) 1–24 1–24

Response time Fast Fast

70–80 (60– 80) 65–88 35–42 14–18 Storage cost ($/kWh) 5–136 2–141

1–5 (6–8) 0.017–2 12 1–24 PCS ($/kW)

¼ cycle

403–4644 432–1674

3–30 3–30

132–915 (609– 1210) 282–4104 2–15 0.059– 0.101

211–648 (281–355)

46–140 (76–130)

161–4320 540–4809 1–50 (kJ/kg/ K)

0–130 11–43 59–101

Pumped hydro storage Conventional compressed air energy storage Lead acetate (nickel cadmium, if different)

15 k 10 k

30–50 (45– 80) 75–200 400–1000 80–200 Self-discharge/day (%) Low Low

0.25–1.5 (1.5–3.0)

0.1–5 (0.2–0.3)

Lithium-ion Hydrogen Thermal

0.6–1.2 0.1–1.0 5– 15 years

1–5 Low 0.05–1

0.015–50 0–50 50–250 Lifetime (cycles, k)

¼ cycle ¼ cycle Fast BOP cost ($/kW)

Sources: For Pumped Hydro and Compressed Air Storage: [19, 20, 22–24]. For lead acetate storage: [25–29]. For lithium-ion storage: [24, 26, 27, 28, 30, 31, 32]. For hydrogen Storage: [20, 22, 24, 26, 33–35]. For thermal Storage: [24, 36–38]. For levelized costs of electricity and lifetime cost analysis: [22, 29, 39–50, 122]

The analysis using the LCOE equation demonstrates that Li-ion batteries and advanced adiabatic and isothermal compressed air energy storage as suitable energy storage platforms with low levelized costs and high reliability, whereas under a medium time scale, pumped hydro and advanced adiabatic compressed air or energy storage or thermal storage using molten salts or mineral oil and long-term power to gas to power or pumped hydro, where hydrogen or methane is used to generate electricity. The electricity would be generated using a combined cycle power plant or proton exchange membrane fuel cells [52, 53].

8

1.1.2

S. Bashir et al.

Role of Energy Storage

Energy storage is required to enable constant power load and is driven by pumped hydro as the historically largest share form of energy storage, with lithium-ion batteries being a recent addition due to the rise of both solar and wind power. No lifetime assessment cycle reviewed encompasses all the required parameters, however, parameters that drive cost and applicability are a lifetime of storage technology, the efficiency of energy retrieval, and cycle length and depend on end-usage, with lithium-ion batteries gaining a significant share in all stationary applications due to greater research and development as a result of increased uptake of solar and wind as sustainable energy resources. For long-term storage and high-need applications, pumped hydro, compressed air, hydrogen and thermal energy storage are likely to increase due to stricter emissions standards for automobiles, fuel cells driven transport is now feasible, utilizing zeolite-based heterogeneous catalysts, for example, bismuth vanadate as an intermediate bandgap photocatalyst that could be coupled solar to photocatalysis through recombination of photoinduced electrons and holes to generate hydrogen from water or coal [54, 55].

1.2

The Electron Economy

While the last two decades have seen a dramatic rise in solar and wind energy generation resources, the future most likely will look toward a hybrid system where energy storage will play an increasingly prominent role coupled with clean coal and carbon dioxide capture. Figure 1.5 shows a conceptual workflow of how this could be accomplished. The energy generation resource such as a photovoltaic array, wind-powered turbine, solar concentrator to steam to turbine, or hydrogen feed gas will need to generate electricity, store it and then allow the system to utilize the stored energy. This will be done using inverters, and heat pumps, and a direct current (DC) bus,

Fig. 1.5 Workflow of hybrid energy generation and storage [56]

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

9

using a DC-to-DC converter and DC bus connection to alternating current (AC) bus using inverters and the electricity or lead load connected to the AC bus. In this manner, energy is drawn with surplus energy used to recharge the secondary or flow batteries or electrolyzes to generate hydrogen from water, if the load is more than the immediate power, then the energy from the stored batteries can be tapped or stored hydrogen to generate electricity with the excess heat used to assist with steam generation. Currently, approximately 2=3 of the total annualized cost is due to components from the wind-powered turbine and the rest for batteries, inverters, and hydrogen storage systems, and fuel cells. The fuel cell catalyst is the next greatest expense and is likely to decrease as more research and development is geared toward the automobile and residential energy sectors, where the likelihood of supercapacitor, battery, flex-fuel, and fuel cells are used or some combination for automobile and stationary applications [57–59]. The rapid development in alkali-metal-based batteries is driven by chemical and physical characteristics of lithium and sodium and this is related to atomic radii, electron affinity, and charge density, from which energy density and kinetics of energy release are derived: e.g. 6.9 atomic mass unit, amu per mole, 0.76 Å, 180.5  C, 3.0 V, 3860 mAg.g1 and $5 k for lithium, Li vs. 23 amu, 1.06 Å, 97.7  C, 2.7 V 1160 mAg.g1 and $0.15 k for sodium, Na. The molar mass is in atomic mass units, cation radii in Angstroms, melting point in degrees Celsius, voltage relative to standard hydrogen electrode in volts, the energy density in milliampere hour per gram, and cost of carbonate salt on a ton basis in United States dollars, thousand increments, respectively. These energy parameters favor lithium-ion batteries, while environmental and economic factors favor sodium ion-based batteries [60]. A bifurcated role is likely with lithium-ion batteries being favored in automobiles where high energy density and fast kinetics are required and sodium-ion batteries for stationary energy storage where high density is not required, such as buildings or a configuration shown in Fig. 1.5 [61]. The near-future bottleneck to battery and fuel cell development is the electrocatalysts and electrode architecture, which will ultimately define high gravi- and volumetric energy densities at the solid electrolyte interphase (SEI) layer that limits spontaneous reactions at the expense of thermodynamic instability and lower cell lifetime, which may be accomplished through alloying, described in Eqs. 1.4 and 1.5 [62]: E þ yAþ þ xē ! Ay M,

ð1:4Þ

where A represents alkali metals Li or Na and E represent metalloids Si, Ge, Sn, Sb, and P, respectively, or conversion [63]: Ma Ob þ ðb:nÞA ! aM þ bAn O,

ð1:5Þ

where M.O are metal oxides and A is alkali metals (Li or Na) such as Silicon-based anodes with Li iron phosphate batteries, e.g. Li17Si4.

10

S. Bashir et al.

1.2.1

Public Policy and Discourse

The above paragraphs imply an electron-driven economy, while this is true to a greater degree today, this is not a complete characterization of the likely environment-food-water-energy nexus in the next 50 years, because, within the United States, much policy is driven at the federal level due to the muscular resources available and at the time of writing this preface, president-elect Biden has 48 senate votes most likely confirming his policies and the best-case scenario is of a clean sweep of the State of Georgia Senate races, tying the US senate at fifty senators from the Grand Old Party (Republican GOP) and the democratic party that as a slim majority within the U.S house of representatives and either a 48–52, 49–51 minority or a tie broken by the Vice-President and confirming Democratic policies on behalf of the presidency of the United States. This political reality means major policy initiates will require bipartisan support to be enacted into law. This means the Biden Green initiative on the environment that pledges great reduction in carbon dioxide emissions could be met through not an electron economy but a carbon majority and hydrogen minority-driven economy, whereas natural gas and coal-powered generation of electricity are accomplished using plant upgrades to capture carbon dioxide and conversion of biomass, coal and natural resources to hydrogen that is used to generate energy and promote fuel cell-driven transport. Here, the U.S Department of energy could utilize its public-partnership research fund to promote companies to conduct subsidized research in fuel cells-driven transport, similar to a model employed to assist E. Musk in his development of a sports electric vehicle in the early zero–zero decades. There was bipartisan support to elicit more private sector funding for technologies that have yet to reach maturity and widespread deployment but were further in the development cycle than emerging technologies, much as like the National Institute of Health (NIH) Model in support vaccine research in developing world countries related to the human immunodeficiency viruses/acquired immunodeficiency syndrome and tropical diseases that do not reach the profit margin for many US Pharma to invest in heavily. This NIH initiative has also received support from prominent foundations such as the Gates Foundation and could be a scenario whereby, the politics is removed from the policy.

1.2.2

The Hydrogen Economy

How could such an economy work in principle? Where would the hydrogen be generated? Currently, hydrogen production, approximately 48% from natural gas, 30% from Naphtha and Oils, and 18% from coal [28] or biomass. When fossil fuels are the feedstock, hydrogen production can be accomplished by hydrocarbon reforming, pyrolysis. Reforming from hydrocarbons can be accomplished by steam reforming, partial oxidation, and autothermal reforming. When generated from biomass, this may be by bio photolysis, dark or photo fermentation, if thermochemical, by pyrolysis, gasification, combustion or liquefaction, and from water splitting as indicated in Fig. 1.5. Here, electrolysis, thermolysis, and photolysis

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

11

Table 1.2 Summary of the energy density of common fuels Fuel Hydrogen Methane Ethane Butane Methanol Ethanol Gasoline (Cetane # 8– 14) Diesel (Cetane # 51) Coal Wood (C6H10O5)n Sugar (glucose)

Phase at STP Gas Gas Gas Gas Liquid Liquid Liquid

Specific energy (MJ/ Kg) 141.9 55.5 51.9 49.5 20 29.7 45.4

Energy Density (MJ/ kg) 119.9 50.0 47.8 29.7 17.1 23.3 42.6

Liquid Solid Solid Solid

46.0 15–27 15.0 15.6

43.0 12.0–36.5 9.0 23.9

Sources and Notes’: Standard Temperature and Pressure. [64–68]

could be utilized. The advantages of hydrogen are in its high energy density when compared with carbon-based fuels, as shown in Table 1.2 and workflow in Fig. 1.6. Hydrogen production by steam reformation is nickel catalytic conversion of hydrocarbons and steam to form hydrogen and carbon oxide through a series of sequential steps, which result in the synthesis of gas (syngas), water-gas shift, and methanation or gas purification. The feed is light hydrocarbons or naphtha; however, the materials need to be carbon and hydrogen to avoid poisoning of the catalyst. If the coal or feed contains sulfur, a desulfurization step is performed [69] at 3.5 MPa and 900  C temperature using a 3:5 ratio of steam to carbon [70]. The heat is recovered and the channel promotes the water gas shift reaction, where carbon monoxide reacts with steam to produce hydrogen gas and is purified through a pressure swing adsorption [71] at the expense of carbon dioxide emission, which could be captured using supercritical conditions and the liquefied carbon dioxide stored in geological reservoirs. Alternately, the carbon dioxide could be reacted with hydrogen to generate wet methane [72]. This methane is re-fed back to the reformer to generate additional hydrogen. The different gases could be separated using Pd-based Inorganic membranes with high selectivity and at lower operational temperatures (550  C) lowering the cost of hydrogen production to $2.0 kg with Carbon dioxide capture and storage [73] described by the steam reformation, watergas shift reaction, or methanation described in Eqs. 1.6, 1.7 and 1.8: Cn Hm ðgÞ þ nH2 O ðgÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:6Þ

WGS : CO ðgÞ þ H2 O ðgÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:7Þ

Methanator : CO ðgÞ þ 3H2 ðgÞ ! CH4 ðgÞ þ H2 O ðlÞ

ð1:8Þ

Partial oxidation is a similar method but using oxygen instead of air and occurs at around 1315  C for the non-catalytic process using coal, oils, or methane [71]. The

Fig. 1.6 Workflow process to produce hydrogen using ten (〇!〇) distinct processes. Each unique process and the pathway are indicated using a numbered sequence. For example, hydrogen production by Steam reforming is shown by the process 〇, membrane separation by the process ❷, coal gasification by 〇, biomass-based processes are shown by sequences 〇,〇 and〇. Thermo-photo- or electrolysis processes by sequences 〇,〇,〇, and 〇. CRA represents carbonrich acids, CBP Carbon By-products; Pe the electrical energy, HeTF Heat Transfer Fluids, HiTHeE High-temperature Heat Exchanger [68]

12 S. Bashir et al.

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

13

feed gas or materials are oxidized directly, or sulfur is removed, and oxygen is used to partially oxidize hydrocarbons and syngas in a similar workflow to the steam reformation for the catalytic process conducted at lower operating temperatures summarized in Eqs. 1.9, 1.10, 1.11 and 1.12 [74]: Catalytic Reformer : Cn Hm ðgÞ þ ½nO2 ðgÞ ! nCO ðgÞ þ ½mH2 ðgÞ

ð1:9Þ

Non  catalytic : Cn Hm ðg=lÞ þ nH2 O ðg=lÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:10Þ

WGS : CO ðgÞ þ H2 O ðg=lÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:11Þ

Methanator : CO ðgÞ þ 3H2 ðgÞ ! CH4 ðgÞ þ H2 O ðlÞ

ð1:12Þ

Because coal has lower hydrogen to carbon ratio than methane, higher pressure and temperature are required (6 MPa and 1315  C), generating non-reactant ash and more water [75]. The above two processes are endothermal and require additional energy input, for autothermal reformation; the initial oxidation is exothermal, and the thermal energy is used to drive the endothermal steam reforming process to drive hydrogen production. The oxidizer is steam, oxygen, or air as the inlet gas to the reformer to drive the reformation and oxidation reactions to occur [76], summarized in Eq. 1.13: Cn Hm ðlÞ þ ½nH2 O ðlÞ þ ¼nO2 ðgÞ ! nCO ðgÞ þ ð½n þ ½mÞH2 ðgÞ

ð1:13Þ

Because part of the thermal energy is utilized, the overall energy demand is lesser than the before processes, lowering the cost of production. Methane conversion to hydrogen could be separated using a Pd membrane and could be accomplished at lower temperatures than the non-catalytic reaction when the air was used, approximately 900  C, that has to be tuned to the operating parameters of the inorganic membrane, as a higher temperature may lead to membrane deformation and loss of selectivity of hydrogen [77]. Hydrocarbon pyrolysis is the breakdown of hydrocarbons to hydrogen via a thermocatalytic breakdown to carbon and hydrogen at the boiling point of the liquid hydrocarbons. When heavy oil hydrocarbons are utilized, the decomposition temperature is higher around 350  C to promote gasification of the hydrocarbon followed by cracking or the methane to hydrogen gas as a form of decarbonization at under 1000  C and 1 atm pressure. Pyrolysis is a gasification process and does not involve a water gas shift or carbon dioxide absorption step. The hydrogen gas is separated using Pd-Ag alloyed membrane at lower temperatures to avoid damage to the membrane, the rate of separation is dependent upon the partial pressure of the incoming gas, noting that the gas needs to be devoid of sulfur to avoid catalytic poisoning [78]. Hydrogen gas can also be formed using biomass, where the cellulose lignin-based materials can be converted to methanol, sugars that were generated from carbon dioxide, and sunlight known as photosynthesis. The carbon biomass is then

14

S. Bashir et al.

thermochemically broken via pyrolysis to generate a gasified-rich distillate in methane and carbon monoxide that can be processed to generate hydrogen. Alternately, the biomass can be liquified and combusted to ultimately generate hydrogen at 0.5 MPa and 530  C anaerobically to generate small neutral gases, hydrocarbons, and char [79] described by Eqs. 1.14, 1.15, 1.16 and 1.17: Cn Hm ðs=lÞ þ nH2 O ðlÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:14Þ

CO ðgÞ þ H2 O ðlÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:15Þ

In biomass gasification, the biomass is converted into syngas using air, oxygen, or steam at 1400  C and 32 atm: Biomass ðs=lÞ þ Air ðgÞ ! H2 ðgÞ þ CO2 ðgÞ þ CO ðgÞ þ N2 ðgÞ þ CH4 ðgÞ þ Cn Hm ðlÞ þ H2 O ðlÞ þ char ðsÞ ð1:16Þ Biomass ðs=lÞ þ Steam ðg=lÞ ! H2 ðgÞ þ CO2 ðgÞ þ CO ðgÞ þ CH4 ðgÞ þ Cn Hm ðlÞ þ H2 O ðlÞ þ char ðsÞ þ tar ðsÞ ð1:17Þ Photosynthesis from plants or fermentation by cyanobacteria, algae, and fungi can also convert biomass to sugar-rich feedstocks that can be converted to carbonrich acids (CRA) and then to hydrogen gas using light energy or in the dark using two electron reduction by a hydrogenase or nitrogenase. Some elect microbes can directly convert water to hydrogen in a process known as photolysis, summarized in Eqs. 1.18, 1.19, 1.20, 1.21 and 1.22 [80]: 2H2 O ðlÞ þ hv ! 2H2 ðgÞ þ O2 ðgÞ or 12H2 O ðlÞ þ 6CO2 ðgÞ þ hv ! C6 H12 O6 ðsÞ þ 6O2 ðgÞ and

ð1:18Þ ð1:19Þ

C6 H12 O6 ðsÞ þ 12H2 O ðlÞ þ hv ! 12Hþ ðaqÞ

þ6CO2 ðgÞ or sugars or carbon rich acids ðCRAÞ ð1:20Þ C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! 2CH3 COOH ðCRAÞ ðlÞ þ 4H2 ðgÞ þ 2CO2 ðgÞ ðfermentationÞ

ð1:21Þ

C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! CH3 ðCH2 Þ2 COOH ðCRAÞðlÞ þ 2H2 ðgÞ þ 2CO2 ðgÞ ðfermentationÞ

ð1:22Þ

Cellulose/Starch are the common bio feed stocks and fermentation is carried out at mild acidic conditions in the dark. When sunlight is used, nitrogenase can convert CRA (e.g. acetic acid) into hydrogen and carbon dioxide, summarized in Eq. 1.23 [81]:

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

CH3 COOH ðlÞ þ 2H2 O ðlÞ þ hv ! 4H2 ðgÞ þ 2CO2 ðgÞ

15

ð1:23Þ

These reactions could be accomplished using a bio-reactor, where sugars are digested by bacteria anaerobically in the dark to generate carbon-rich acids, that could be used by photosynthetic bacteria using light and water described by Eqs. 1.24 and 1.25 [82]: C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! 2CH3 COOH ðlÞ þ 2CO2 ðgÞ þ 4H2 ðgÞ

ð1:24Þ

2CH3 COOH ðsÞ þ 4H2 O ðlÞ þ hv ! 8H2 ðgÞ þ 4CO2 ðgÞ

ð1:25Þ

Using glucose as the biomass, approximately 7 mol of hydrogen could be generated per mol of glucose [83] under mild acidic conditions. The last method to generate hydrogen is direct water splitting using sunlight [84] and redox of water. As this is an endothermic reaction, energy from sunlight is used using a fuel cell/ electrolyzer design consisting of a cathode and anode and electrolyte with an external current. Here, water is electrolytically broken to hydrogen and oxygen, shown in Eq. 1.26 [85]: 2H2 O ðlÞ ! 2H2 ðgÞ þ O2 ðgÞ

ð1:26Þ

The electrolysis can be accomplished using fuel cells such as proton exchange membrane fuel cells, solid oxide electrolysis cells, where they generate proton migrates across a membrane to the cathode to generate hydrogen, leaving water and oxygen [86, 87]. In the solid oxide electrolyzer operated under alkaline conditions, water is broken down to H+ and OH, with the hydroxide traveling through the aqueous electrolyte to the anode to form oxygen [88], with the driving energy being thermal instead of electrical in a fuel cell, with hydrogen in the steam feed and using fuel cells described by Eqs. 1.27, 1.28, 1.29 and 1.30 [89]: Alkaline Solid Oxygen Electrolyzer (Anode). OH ðaqÞ ! O2 ðgÞ þ 2H2 O ðlÞ þ 4ē

ð1:27Þ

2H2 O ðlÞ þ 2ē ! 2OH  ðaqÞ þ H2 ðgÞðcathodeÞ

ð1:28Þ

Proton exchange membrane 2H2 O ðlÞ ! O2 ðgÞ þ 4Hþ ðaqÞ þ 4ē ðanodeÞ þ

4H ðaqÞ þ 4ē ! 2H2 ðgÞðcathodeÞ

ð1:29Þ ð1:30Þ

The efficiency and relative merits of hydrogen production summarized in Eqs. (1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 1.10, 1.11, 1.12, 1.13, 1.14, 1.15, 1.16, 1.17, 1.18, 1.19, 1.20, 1.21, 1.22, 1.23, 1.24, 1.25, 1.26, 1.27, 1.28, 1.29 and 1.30) are summarized in Table 1.3A. The cost of solar thermal is due to the electrolyzer unit and for solar, it is peak and off-peak costs plus the cost of the electrolyzer unit. For

16

S. Bashir et al.

Table 1.3A Hydrogen production by energy source, rate, capacity factor, and approximate cost in the year of study Electricity platform Nuclear Solar thermal Solar photovoltaic Wind BioSystem Photolysis Fermentation (Dark) Fermentation (Light)

H2 production rate (kg/day) 1000 1000 1350 35,000 Synthesis rate (mmol H2/h)

Capacity factor (%) 95 40 30 60

0.05–0.5 9.5

Hydrogen cost ($/kg) 4.15 7.00 10.50 5.25 Bioreactor volume (m3) 340–1710 8.50

0.15

750

Sources: [68, 86, 90–96]

wind, the unit can be used in the cogeneration of hydrogen and electricity or hydrogen varying the cost. The values chosen are to the nearest-semi-decade or multiples of five and are based on a standard fuel stack of 5 kW [97]. The synthesis rate for biological processes is highly dependent on the source of carbon, whether the photolysis is by direct or indirect means and the availability of oxygen. For this reason, the modal value by semi-decade is used. As the above is an energy-intensive process, currently the energy is offset from renewable resources such as wind, solar, or concentrated solar thermal. At temperature above 2500  C, water splitting is energetically favorable and could be accomplished catalytically at lower temperatures using binuclear catalysts based on copper, summarized in Eqs. 1.31, 1.32, 1.33 and 1.34 [98, 99]): 

2H2 O ðlÞ ! 2H2 ðgÞ þ O2 ðgÞ, Temp > 2500 C

ð1:31Þ

~ C CuCl2 ðsÞ þ H2 O ðgÞ ! CuO*CuCl2 ðsÞ þ 2HCl ðgÞ, T400 ~ C CuO*CuCl2 ! 2CuCl ðlÞ þ ½O2 ðgÞ, Temp500

ð1:33Þ

~ C SnO ðsÞ þ H2 O ðgÞ ! SnO2 ðsÞ þ H2 ðgÞ, Temp550

ð1:34Þ



1.2.3

ð1:32Þ

The Photon Economy

Concentrating Solar Power is one attractive vehicle by which the thermolysis of water can be accomplished. Lastly, photolysis where sunlight is absorbed by a semiconductor catalyst where photon energy above the semiconductor bandgap at the anode will result in the generation of an electron-hole pair which migrates toward

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

17

Table 1.3B Hydrogen production by thermal source, rate, catalyst, and approximate cost in the year of study Thermal source Nuclear Nuclear Solar Solar

H2 production rate (kg/day) 7000 800,000 6000 6000

Solar

6000

Catalyst CuCl SI ZnO/Zn Fe3O4/ FeO Fe2O3/ Fe3O4

Operating temperature ( C) 550 850 1730 1630

Hydrogen cost ($/kg) 2.15 2.55 8.00 8.40

1330

8.40

Sources: [68, 86, 90–96, 104–110]

the appropriate biased electrodes and the electrolyte, with the holes migrating at the anode, where water is split into protons that migrate through the electrolyte to the cathode and oxygen that remains with water, while electron flow is through an external circuit to the cathode, where protons combine to form hydrogen gas, summarized in Eqs. 1.35, 1.36 and 1.37 [100–103]: þ

2 p þH2 O ðlÞ ! ½O2 ðgÞ þ 2Hþ ðaqÞ ðanodeÞ

ð1:35Þ

2ē þ 2H þ ðaqÞ ! H2 ðgÞðcathodeÞ

ð1:36Þ

H2 O ðlÞ ! H2 ðgÞ þ ½O2 ðgÞ

ð1:37Þ

The efficiency and relative merits of hydrogen production summarized in Eqs. (1.31, 1.32, 1.33 and 1.34) is summarized in Table 1.3B: This reaction would occur at a bandgap of 1.23 eV without the application of an external bias, with the most common semiconductor photocatalysts being SiC and TiO2 [111].

1.2.4

The Circular Economy

Steam methane reforming is the most common with biomass being used more as a carbon-neutral mechanism to generate hydrogen gas by thermochemical processes where hydrogen gas is generated from pyrolysis and gasification, whereas biochemical processes utilize fermentation using bioreactors. The energy costs for these processes can be aligned using concentrated solar power to drive coal gasification and water splitting and storage that is achieved using high pressure or liquefaction using cryogenic temperatures and cooling limiting the total capacity and are summarized in Table 1.4 for relative costs. Solid-state storage using metal–organic frameworks at lower temperatures may be achieved at lower operating pressures,

18

S. Bashir et al.

Table 1.4 Summary of relative merits of each hydrogen generation process. The efficiency is the highest practical efficiency and the cost is the modal costs where the process is highly variable such as wind, solar, or biochemical by semi-decade. Nuclear thermal is always cheaper than solar thermal due to the greater energy density of nuclear fission and greater maturity of the technology Process Steam reformation Partial oxidation Autothermal reforming Pyrolysis Coal gasification Photolysis

Efficiency (%) 85

Energy source Hydrocarbons

Feedstock Methane

Hydrogen cost ($/kg) 2.30

75 75

Hydrocarbons Hydrocarbons

Methane Methane

1.65 1.50

50 30 10

Steam Hydrocarbons Solar

Biomass Coal Water and Algae/ Cyanobacteria CRA

2.05 1.60 2.15

Water

6.50

Water Water

2.60 8.50

Dark fermentation Electrolysis

80

STP (dark)

60

Thermolysis Thermolysis

45 45

Wind\Solar\ Nuclear Nuclear Solar

2.60

Notes and Sources: CRA Carbon-rich acids; Hydrocarbon are fossil fuels including natural gas and/or coal; STP Standard Temperature and Pressure; [68, 96, 104–110] Table 1.5 Summary of current storage methods for hydrogen

Process Liquid hydrogen Compressed hydrogen Activated carbon Interstitial metal hydride Salt-like metal hydride

Storage density (mat wt %)/[volumetric density, by wt%] 100/[14]

Volume density (kg/m3)/[energy release, MJ/kg H2] 70/[120]

Temperature ( C) 253

Pressure (MPa) 0.1

100/[3]

40/[120]

25

77

10.5/[6.5]

4/[3.5]

200

10

2/[3]

150/[15]

0

1

7.5/[18]

150/[37]

330

1

Sources: [112–120]

or chemical conversion to generate metal hydrides, which can be decomposed at close to standard temperature and pressure (STP) and are summarized in Table 1.5. The energy demand is unlikely to diminish in the near and intermediate future and the environmental stress from soluble oxides of carbon, nitrogen, and sulfur are unlikely to diminish, including the long-term effects of SARS [121]. As this

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

19

pathology in part is due to respiratory distress, then energy production without particulate matter and carbon dioxide would greatly reduce the risk of proclivities to being susceptible to the virus. The development of integrated fossil fuel and renewable energy production, recycling, waste management, and lifecycle assessment will need to be developed and deployed. Hydrogen gas can be generated by several abiotic and biotic methods with steam methane reformation the most common or coal gasification using thermal pyrolysis as possible new approaches using concentrated solar thermal or nuclear thermal and separation by a membrane-based reactor with combined cycle gas turbine exhausts recycling heat and exhaust gas driven by concentrated solar power. Alternately water-splitting methods may be employed and hydrogen stored using porous materials or through compression with new safety codes for its usage and a newer transportation pipeline that can potentially improve the infrastructure, create jobs, and transition to a hydrogen economy for power and heating as well as fuel cell-based transportation, with fossil fuel as feedstocks and nuclear or thermal as energy resources to drive the synthesis of hydrogen, which we believe will garnish bipartisan support at the US Senate which will need to sign off any major environmental or energy policy initiates regardless of the occupancy of the Whitehouse. Acknowledgments This work is supported by the Petroleum Research Fund of the American Chemical Society (53827-UR10) and the Robert Welch Foundation (Departmental Grant, AC-0006). We thank the program chair and Dr. E. Gerald Meyer of the ENFL America Chemical Society for the opportunity to run or moderate symposia at the technical sessions. Lastly, The leadership at Texas A&M University-Kingsville, Department, College, and University level, as well as Springer Science+Business Media, LLC technical staff for their assistance in copy editing this and other book chapters. Author Contributions S. Bashir completed the initial draft including equations relating to thermodynamics. J.L. Liu, data for figures and Sai. Chava the data for hydrogen production and figure 6, and W. Song and Y. Gao who reviewed the final draft which was submitted by S. Chava.

References 1. US Energy Information Administration, (EIA) (2020) Electricity generation, capacity, and sales in the United States. https://www.eia.gov/energyexplained/electricity/electricity-in-theus-generation-capacity-and-sales.php 2. Capuano L (2018) International energy outlook 2018 (IEO2018). US Energy Information Administration (EIA): Washington, DC, USA, 2018, 21. https://www.hidropolitikakademi. org/uploads/wp/2018/07/International-Energy-Outlook-2018-IEO2018.pdf 3. Conti J, Holtberg P, Diefenderfer J, LaRose A, Turnure JT, Westfall L (2016) International energy outlook 2016 with projections to 2040 (No. DOE/EIA-0484 (2016). USDOE Energy Information Administration (EIA), Washington, DC (United States). Office of Energy Analysis. https://www.eia.gov/outlooks/ieo/pdf/0484(2016).pdf#:~:text¼International Energy Outlook 2016 With Projections to 2040, an analytical agency within the U.S. Department of Energy

20

S. Bashir et al.

4. Sieminski A (2013) International energy outlook in 2013. US Energy Information Administration (EIA) Report Number: DOE/EIA-0484. https://www.eia.gov/pressroom/presentations/ sieminski_07252013.pdf 5. US Energy Information Administration, (EIA) (2017) Annual energy outlook 2015: with projections to 2040. https://www.osti.gov/servlets/purl/1296780 6. Newell RG, Raimi D (2020) Global energy outlook comparison methods: 2020 update. https:// media.rff.org/documents/Global_Energy_Outlook_Comparison_Methods_2020.pdf 7. Beér, J. (2020). CO2 reduction and coal-based electricity generation, Chapter 15. In R. Malhotra (Ed.), Encyclopedia of sustainability science and technology series: Volume XVIII: Fossil energy (2nd ed., pp. 427–437). https://doi.org/10.1007/978-1-4939-9763-3 8. Midilli A, Dincer I, Ay M (2006) Green energy strategies for sustainable development. Energy Policy 34(18):3623–3633 9. Perera FP (2017) Multiple threats to child health from fossil fuel combustion: impacts of air pollution and climate change. Environ Health Perspect 125(2):141–148 10. Geissdoerfer M, Savaget P, Bocken NM, Hultink EJ (2017) The circular economy–a new sustainability paradigm? J Clean Prod 143:757–768 11. Lazard.Com (2020) Levelized cost of energy and levelized cost of storage – 2020. https:// www.lazard.com/perspective/lcoe2020 12. US Environmental Protection Agency (2020) Air trends. https://www.epa.gov/air-trends 13. Ubando AT, Felix CB, Chen WH (2020) Biorefineries in circular bio-economy: a comprehensive review. Bioresour Technol 299:122585 14. Munir M, Ahmad M, Saeed M, Waseem A, Rehan M, Nizami AS et al (2019) Sustainable production of bioenergy from novel non-edible seed oil (Prunus cerasoides) using bimetallic impregnated montmorillonite clay catalyst. Renew Sust Energ Rev 109:321–332 15. Ponnusamy VK, Nguyen DD, Dharmaraja J, Shobana S, Banu JR, Saratale RG et al (2019) A review of lignin structure, pretreatments, fermentation reactions, and biorefinery potential. Bioresour Technol 271:462–472 16. Ho MC, Ong VZ, Wu TY (2019) The potential use of alkaline hydrogen peroxide in lignocellulosic biomass pretreatment and valorization–a review. Renew Sust Energ Rev 112:75–86 17. Montoneri E (2017) Municipal waste treatment, technological scale-up, and commercial exploitation: the case of bio-waste lignin to soluble lignin-like polymers. In: Food waste reduction and valorisation. Springer, Cham, pp 79–120 18. Nazir H, Batool M, Osorio FJB, Isaza-Ruiz M, Xu X, Vignarooban K et al (2019) Recent developments in phase change materials for energy storage applications: a review. Int J Heat Mass Transf 129:491–523 19. Abdon A, Zhang X, Parra D, Patel MK, Bauer C, Worlitschek J (2017) Techno-economic and environmental assessment of stationary electricity storage technologies for different time scales. Energy 139:1173–1187 20. Zhao H, Wu Q, Hu S, Xu H, Rasmussen CN (2015) Review of energy storage system for wind power integration support. Appl Energy 137:545–553 21. Riensche E, Stimming U, Unverzagt G (1998) Optimization of a 200 kW SOFC cogeneration power plant: part I: variation of process parameters. J Power Sources 73(2):251–256 22. Schoenung SM (2001) Characteristics and technologies for long-vs. short-term energy storage. United States Department of Energy. https://prod-ng.sandia.gov/techlib-noauth/accesscontrol. cgi/2001/010765.pdf 23. Denholm P, Kulcinski GL (2004) Life cycle energy requirements and greenhouse gas emissions from large scale energy storage systems. Energy Convers Manag 45(13–14):2153–2172 24. Chen H, Cong TN, Yang W, Tan C, Li Y, Ding Y (2009) Progress in electrical energy storage system: a critical review. Prog Nat Sci 19(3):291–312 25. Rydh CJ (1999) Environmental assessment of vanadium redox and lead-acid batteries for stationary energy storage. J Power Sources 80(1–2):21–29

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

21

26. Barton JP, Infield DG (2004) Energy storage and its use with intermittent renewable energy. IEEE Trans Energy Conver 19(2):441–448 27. Baker J (2008) New technology and possible advances in energy storage. Energy Policy 36 (12):4368–4373 28. Hadjipaschalis I, Poullikkas A, Efthimiou V (2009) Overview of current and future energy storage technologies for electric power applications. Renew Sust Energ Rev 13 (6–7):1513–1522 29. Zakeri B, Syri S (2015) Electrical energy storage systems: a comparative life cycle cost analysis. Renew Sust Energ Rev 42:569–596 30. Smith SC, Sen PK, Kroposki B (2008, July) Advancement of energy storage devices and applications in the electrical power system. In: 2008 IEEE power and energy society general meeting-conversion and delivery of electrical energy in the 21st century. IEEE, pp 1–8 31. Beaudin M, Zareipour H, Schellenberglabe A, Rosehart W (2010) Energy storage for mitigating the variability of renewable electricity sources: an updated review. Energy Sustain Dev 14(4):302–314 32. Liu H, Wei Z, He W, Zhao J (2017) Thermal issues about Li-ion batteries and recent progress in battery thermal management systems: a review. Energy Convers Manag 150:304–330 33. Smith W (2000) The role of fuel cells in energy storage. J Power Sources 86(1–2):74–83 34. Benitez LE, Benitez PC, Van Kooten GC (2008) The economics of wind power with energy storage. Energy Econ 30(4):1973–1989 35. Li P (2008) Energy storage is the core of renewable technologies. IEEE Nanotechnol Mag 2 (4):13–18 36. Thaker S, Oni AO, Kumar A (2017) Techno-economic evaluation of solar-based thermal energy storage systems. Energy Convers Manag 153:423–434 37. Liu M, Tay NS, Bell S, Belusko M, Jacob R, Will G et al (2016) Review on concentrating solar power plants and new developments in high-temperature thermal energy storage technologies. Renew Sust Energ Rev 53:1411–1432 38. Hidalgo DB, Alonso JG, Pérez YM (2017) Costos De Las Tecnologías De Almacenamiento De Energía Térmica. Revista Centro Azúcar 44(4):10–10 39. Brownson DA, Kampouris DK, Banks CE (2011) An overview of graphene in energy production and storage applications. J Power Sources 196(11):4873–4885 40. Flueckiger SM, Iverson BD, Garimella SV (2014) Economic optimization of a concentrating solar power plant with molten-salt thermocline storage. J Solar Energy Eng 136(1):011015 41. Hameer S, van Niekerk JL (2015) A review of large-scale electrical energy storage. Int J Hydrog Energy 39(9):1179–1195 42. Luo X, Wang J, Dooner M, Clarke J (2015) Overview of current development in electrical energy storage technologies and the application potential in power system operation. Appl Energy 137:511–536 43. Amrouche SO, Rekioua D, Rekioua T, Bacha S (2016) Overview of energy storage in renewable energy systems. Int J Hydrog Energy 41(45):20914–20927 44. Jülch V (2016) Comparison of electricity storage options using the levelized cost of storage (LCOS) method. Appl Energy 183:1594–1606 45. Amirante R, Cassone E, Distaso E, Tamburrano P (2017) Overview of recent developments in energy storage: mechanical, electrochemical, and hydrogen technologies. Energy Convers Manag 132:372–387 46. Yulong P, Cavagnino A, Vaschetto S, Feng C, Tenconi A (2017, June) Flywheel energy storage systems for power systems application. In: 2017 6th international conference on clean electrical power (ICCEP). IEEE, pp 492–501 47. Buiskikh D, Zakeri B, Syri S, Kauranen P (2018, June) Economic feasibility of flow batteries in grid-scale applications. In: 2018 15th international conference on the European energy market (EEM). IEEE, pp 1–5 48. Schmidt O, Melchior S, Hawkes A, Staffell I (2019) Projecting the future levelized cost of electricity storage technologies. Joule 3(1):81–100

22

S. Bashir et al.

49. Mostafa MH, Aleem SHA, Ali SG, Ali ZM, Abdelaziz AY (2020) Techno-economic assessment of energy storage systems using annualized life cycle cost of storage (LCCOS) and levelized cost of energy (LCOE) metrics. J Energy Storage 29:101345 50. Rahman MM, Oni AO, Gemechu E, Kumar A (2020) Assessment of energy storage technologies: a review. Energy Convers Manag 223:113295 51. Aneke M, Wang M (2016) Energy storage technologies and real-life applications–a state of the art review. Appl Energy 179:350–377 52. Sammes N (ed) (2006) Fuel cell technology: reaching towards commercialization. Springer Science & Business Media, London 53. Davison J, Mancuso L, Ferrari N (2014) Costs of CO2 capture technologies in coal-fired power and hydrogen plants. Energy Procedia 63:7598–7607 54. Hirth L (2013) The market value of variable renewables: the effect of solar wind power variability on their relative price. Energy Econ 38:218–236 55. Schröder A, Kunz F, Meiss J, Mendelevitch R, Von Hirschhausen C (2013) Current and prospective costs of electricity generation until 2050 (No. 68). DIW data documentation. https://www.econstor.eu/bitstream/10419/80348/1/757528015.pdf 56. Zhang Y, Hua QS, Sun L, Liu Q (2020) Life cycle optimization of renewable energy systems configuration with hybrid battery/hydrogen storage: a comparative study. J Energy Storage 30:101470 57. Biggar DR, Hesamzadeh MR (2014) The economics of electricity markets. Wiley, Chichester 58. Hirth L (2015) The optimal share of variable renewables: how the variability of wind and solar power affects their welfare-optimal deployment. Energy J 36(1):149–184 59. Hirth L, Ueckerdt F, Edenhofer O (2016) Why wind is not coal: on the economics of electricity generation. Energy J 37(3):1–27 60. Vaalma C, Buchholz D, Weil M, Passerini S (2018) A cost and resource analysis of sodiumion batteries. Nat Rev Mater 3(4):1–11 61. Ong SP, Chevrier VL, Hautier G, Jain A, Moore C, Kim S et al (2011) Voltage, stability, and diffusion barrier differences between sodium-ion and lithium-ion intercalation materials. Energy Environ Sci 4(9):3680–3688 62. Slater MD, Kim D, Lee E, Johnson CS (2013) Sodium-ion batteries. Adv Funct Mater 23 (8):947–958 63. Nayak PK, Yang L, Brehm W, Adelhelm P (2018) From lithium-ion to sodium-ion batteries: advantages, challenges, and surprises. Angew Chem Int Ed 57(1):102–120 64. EngineeringToolBox.Com (2003) Fuels – densities and specific volume. https://www. engineeringtoolbox.com/fuels-densities-specific-volumes-d_166.html 65. Aifantis KE, Hackney SA, Kumar RV (eds) (2010) High energy density lithium batteries: materials, engineering, applications. Wiley, Weinham 66. Dincer I (2012) Green methods for hydrogen production. Int J Hydrog Energy 37 (2):1954–1971 67. Neutrium.net (2014) Specific energy and energy density of fuels. https://neutrium.net/ properties/specific-energy-and-energy-density-of-fuels/#:~:text¼CommonFuels 68. Nikolaidis P, Poullikkas A (2017) A comparative overview of hydrogen production processes. Renew Sust Energ Rev 67:597–611 69. Balthasar W (1984) Hydrogen production and technology: today, tomorrow, and beyond. Int J Hydrog Energy 9(8):649–668 70. Ersöz A (2008) Investigation of hydrocarbon reforming processes for micro-cogeneration systems. Int J Hydrog Energy 33(23):7084–7094 71. Steinberg M, Cheng HC (1989) Modern and prospective technologies for hydrogen production from fossil fuels. Int J Hydrog Energy 14(11):797–820 72. Hidalgo D, Martín-Marroquín JM (2020) Power-to-methane, coupling CO2 capture with fuel production: an overview. Renew Sust Energ Rev 132:110057 73. Koros WJ, Fleming GK (1993) Membrane-based gas separation. J Membr Sci 83(1):1–80

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

23

74. Chen HL, Lee HM, Chen SH, Chao Y, Chang MB (2008) Review of plasma catalysis on hydrocarbon reforming for hydrogen production—interaction, integration, and prospects. Appl Catal B Environ 85(1–2):1–9 75. Damen K, van Troost M, Faaij A, Turkenburg W (2006) A comparison of electricity and hydrogen production systems with CO2 capture and storage. Part a: review and selection of promising conversion and capture technologies. Prog Energy Combust Sci 32(2):215–246 76. Yoon HC, Erickson PA (2008) Hydrogen from coal-derived methanol via autothermal reforming processes. Int J Hydrog Energy 33(1):57–63 77. Kyriakou V, Garagounis I, Vourros A, Marnellos GE, Stoukides M (2018) A protonic ceramic membrane reactor for the production of hydrogen from coal steam gasification. J Membr Sci 553:163–170 78. Gallucci F, Fernandez E, Corengia P, van Sint Annaland M (2013) Recent advances on membranes and membrane reactors for hydrogen production. Chem Eng Sci 92:40–66 79. Chen G, Andries J, Spliethoff H (2003) Catalytic pyrolysis of biomass for hydrogen-rich fuel gas production. Energy Convers Manag 44(14):2289–2296 80. Das D, Veziroǧlu TN (2001) Hydrogen production by biological processes: a survey of the literature. Int J Hydrog Energy 26(1):13–28 81. Minowa T, Ogi T (1998) Hydrogen production from cellulose using a reduced nickel catalyst. Catal Today 45(1–4):411–416 82. Yu D, Aihara M, Antal MJ Jr (1993) Hydrogen production by steam reforming glucose in supercritical water. Energy Fuel 7(5):574–577 83. Asada Y, Tokumoto M, Aihara Y, Oku M, Ishimi K, Wakayama T et al (2006) Hydrogen production by co-cultures of Lactobacillus and a photosynthetic bacterium, Rhodobacter sphaeroides RV. Int J Hydrog Energy 31(11):1509–1513 84. Zhang Y, Ying Z, Zhou J, Liu J, Wang Z, Cen K (2014) Electrolysis of the Bunsen reaction and properties of the membrane in the sulfur–iodine thermochemical cycle. Ind Eng Chem Res 53(35):13581–13588 85. Bhandari R, Trudewind CA, Zapp P (2014) Life cycle assessment of hydrogen production via electrolysis–a review. J Clean Prod 85:151–163 86. Rozendal RA, Hamelers HV, Euverink GJ, Metz SJ, Buisman CJ (2006) Principle and perspectives of hydrogen production through bio catalyzed electrolysis. Int J Hydrog Energy 31(12):1632–1640 87. ezzahra Chakik F, Kaddami M, Mikou M (2017) Effect of operating parameters on hydrogen production by electrolysis of water. Int J Hydrog Energy 42(40):25550–25557 88. Rossmeisl J, Logadottir A, Nørskov JK (2005) Electrolysis of water on (oxidized) metal surfaces. Chem Phys 319(1–3):178–184 89. David M, Ocampo-Martínez C, Sánchez-Peña R (2019) Advances in alkaline water electrolyzers: a review. J Energy Storage 23:392–403 90. Benemann JR (1997) Feasibility analysis of photobiological hydrogen production. Int J Hydrog Energy 22(10–11):979–987 91. Call D, Logan BE (2008) Hydrogen production in a single chamber microbial electrolysis cell lacking a membrane. Environ Sci Technol 42(9):3401–3406 92. Das D, Veziroglu TN (2008) Advances in biological hydrogen production processes. Int J Hydrog Energy 33(21):6046–6057 93. Kalinci Y, Hepbasli A, Dincer I (2009) Biomass-based hydrogen production: a review and analysis. Int J Hydrog Energy 34(21):8799–8817 94. Holladay JD, Hu J, King DL, Wang Y (2009) An overview of hydrogen production technologies. Catal Today 139(4):244–260 95. Zhu J, Zäch M (2009) Nanostructured materials for photocatalytic hydrogen production. Curr Opin Colloid Interface Sci 4(14):260–269

24

S. Bashir et al.

96. Bartels JR, Pate MB, Olson NK (2010) An economic survey of hydrogen production from conventional and alternative energy sources. Int J Hydrog Energy 35(16):8371–8384 97. Park S, Shao Y, Liu J, Wang Y (2012) Oxygen electrocatalysts for water electrolyzers and reversible fuel cells: status and perspective. Energy Environ Sci 5(11):9331–9344 98. Steinfeld A (2002) Solar hydrogen production via a two-step water-splitting thermochemical cycle based on Zn/ZnO redox reactions. Int J Hydrog Energy 27(6):611–619 99. Steinfeld A (2005) Solar thermochemical production of hydrogen––a review. Sol Energy 78 (5):603–615 100. Ashokkumar M (1998) An overview of semiconductor particulate systems for photoproduction of hydrogen. Int J Hydrog Energy 23(6):427–438 101. Liao CH, Huang CW, Wu J (2012) Hydrogen production from semiconductor-based photocatalysis via water splitting. Catalysts 2(4):490–516 102. Zhang HL, Baeyens J, Degrève J, Cacères G (2013) Concentrated solar power plants: review and design methodology. Renew Sust Energ Rev 22:466–481 103. Zhang X, Peng T, Song S (2016) Recent advances in dye-sensitized semiconductor systems for photocatalytic hydrogen production. J Mater Chem A 4(7):2365–2402 104. Ewan BCR, Allen RWK (2005) A figure of merit assessment of the routes to hydrogen. Int J Hydrog Energy 30(8):809–819 105. Utgikar V, Thiesen T (2006) Life cycle assessment of high-temperature electrolysis for hydrogen production via nuclear energy. Int J Hydrog Energy 31(7):939–944 106. Balat M (2008) The potential importance of hydrogen as a future solution to environmental and transportation problems. Int J Hydrog Energy 33(15):4013–4029 107. Charvin P, Stéphane A, Florent L, Gilles F (2008) Analysis of solar chemical processes for hydrogen production from water splitting thermochemical cycles. Energy Convers Manag 49 (6):1547–1556 108. Orhan MF, Dincer I, Naterer GF (2008) Cost analysis of a thermochemical Cu–Cl pilot plant for nuclear-based hydrogen production. Int J Hydrog Energy 33(21):6006–6020 109. Roeb M, Neises M, Monnerie N, Call F, Simon H, Sattler C et al (2012) Materials-related aspects of thermochemical water and carbon dioxide splitting: a review. Materials 5 (11):2015–2054 110. Romero M, Steinfeld A (2012) Concentrating solar thermal power and thermochemical fuels. Energy Environ Sci 5(11):9234–9245 111. Mishra G, Parida KM, Singh SK (2015) Facile fabrication of S-TiO2/β-SiC nanocomposite photocatalyst for hydrogen evolution under visible light irradiation. ACS Sustain Chem Eng 3 (2):245–253 112. Yartys VA, Lototsky MV (2004) An overview of hydrogen storage methods. In: Hydrogen materials science and chemistry of carbon nanomaterials. Springer, Dordrecht, pp 75–104 113. Rashidi AM, Nouralishahi A, Khodadadi AA, Mortazavi Y, Karimi A, Kashefi K (2010) Modification of single-wall carbon nanotubes (SWNT) for hydrogen storage. Int J Hydrog Energy 35(17):9489–9495 114. Lototskyy M, Yartys VA (2015) Comparative analysis of the efficiencies of hydrogen storage systems utilizing solid-state H storage materials. J Alloys Compd 645:S365–S373 115. Lototskyy MV, Tolj I, Pickering L, Sita C, Barbir F, Yartys V (2017) The use of metal hydrides in fuel cell applications. Prog Nat Sci Mater Int 27(1):3–20 116. Lai Q, Sun Y, Wang T, Modi P, Cazorla C, Demirci UB et al (2019) How to design hydrogen storage materials? Fundamentals, synthesis, and storage tanks. Adv Sustain Syst 3(9):1900043 117. Yartys VA, Lototskyy MV, Akiba E, Albert R, Antonov VE, Ares JR et al (2019) Magnesium based materials for hydrogen-based energy storage: past, present, and future. Int J Hydrog Energy 44(15):7809–7859 118. Aziz M, Wijayanta AT, Nandiyanto ABD (2020) Ammonia as effective hydrogen storage: a review on production, storage, and utilization. Energies 13(12):3062

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

25

119. Hirscher M, Yartys VA, Baricco M, von Colbe JB, Blanchard D, Bowman RC Jr et al (2020) Materials for hydrogen-based energy storage–past, recent progress, and future outlook. J Alloys Compd 827:153548 120. Yao B, Kuznetsov VL, Xiao T, Jie X, Gonzalez-Cortes S, Dilworth JR et al (2020) Fuels, power, and chemical periodicity. Phil Trans R Soc A 378(2180):20190308 121. Moelling K, Broecker F (2020) Air microbiome and pollution: composition and potential effects on human health, including SARS coronavirus infection. J Environ Public Health 2020:1646943 122. Ueckerdt F, Hirth L, Luderer G, Edenhofer O (2013) System LCOE: What are the costs of variable renewables?. Energy 63:61–75

Sajid Bashir received his Ph.D. in Analytical Chemistry from the University of Warwick, England, in 2001. He was a postgraduate research associate at Cornell University to research in the field of plant genetics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has coauthored more than 80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry and Chartered Chemist & Chartered Scientists of Science Council. During his service in TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with risk management. Currently, he collaborated with local law enforcement as a consultant. Sai Raghuveer Chava received his M.S. in Chemistry from Texas A&M University-Kingsville in 2010. He was a Graduate Research Assistant at NNTRC, Kingsville who researched snake venoms. Currently, he is working as a Scientist on Vaccine research focusing on the analytical method development, formulation, and process-related problems using the analytical data. He has also worked on ophthalmic formulations, process characterization, method developments, and method validations. He has coauthored five journal articles and a book chapter and peerreviewed journal articles. He is an elected member of the Royal Society of Chemistry and Sigma Xi.

Dr. Weixin Song is an MPLS Enterprise Fellow and Postdoctoral Research Fellow at the University of Oxford. He received his Ph. D. in Material Electrochemistry from Imperial College London in 2019 funded by the President’s Ph.D. Scholarship. He has been a member of the Centre for Doctoral Training (CDT) of Advanced Characterization of Materials and the London Centre for Nanotechnology (LCN). He completed his B.Sc. in 2012 and M.Sc. in 2015 from Central South University, China. He has research interests in materials electrochemistry and characterization with materials application in energy storage, electrocatalysis, and photo electrocatalysis. He is particularly interested in the materials and interface studies using electron microscopy and spectroscopy. He

26

S. Bashir et al. has 22 first-author publications and 18 licensed patents. He has contributed chapters on one book published by Springer. He received the highest student honor of Central South University (2014), STFC Future Early Career Award (2016), Armourers & Brasiers Rolls-Royce Scheme Grant (2018), National award for outstanding self-financed students abroad (2018), MPLS Enterprise Fellowship (2019), and Constance Fligg Tipper Centenary Memorial prize (2020). Yong-jun Gao Postgraduate, Deputy Director of Clean Energy and Energy Conservation and Environmental Protection Center, Zhejiang Yangtze River Delta. Since 2001, he has been doing research and work in the fields of energy and environmental protection. He has won the second academic prize of China Nuclear Energy Association, the third academic prize of China Electricity Council, and the second prize of the science and technology management innovation of State Power Investment Corporation Limited. He has published many articles as the first author or co-author in publications such as China Nuclear Power, Nuclear Power Engineering, China Power Enterprise Management, East China Electric Power, and Nanostructured Materials for Next-Generation Energy Storage and Conversion. He was awarded the model worker and the May Day medal by a Chinese city. Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing in 2001. She is a Full Professor at Texas A&M UniversityKingsville (TAMUK) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry and a DEBI Faculty Fellow at the US Air Force Research Laboratory. She has authored and co-authored books, book chapters, and peer-reviewed journal articles (> 100). During her 12.5-year services in TAMUK, she taught more than 8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and/or participated in the projects (> 40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Chair-Elect of Energy and Fuels, American Chemical Society.

Chapter 2

Applications and Fundamentals of Photocatalysis with Solar Energy Qiuyang Huang, Dan Kong, and Yongdan Li

2.1

Introduction

This section will give a brief introduction to photocatalysis, mainly on the historical background, applications, and reaction systems of photocatalysis.

2.2

Historical Background

The photoinduced redox reactions on semiconductors can be traced back to the 1930s when scientists reported that TiO2 could react with organic dyes [1]. Limited to the development of social needs, such phenomena did not attract much attention at Q. Huang Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), State Key Laboratory of Chemical Engineering, Tianjin Key Laboratory of Applied Catalysis Science and Technology, School of Chemical Engineering and Technology, Tianjin University, Tianjin, China D. Kong (*) Department of Chemical and Metallurgical Engineering, School of Chemical Engineering, Aalto University, Aalto, Finland Department of Materials Science and Physics, China University of Mining and Technology, Xuzhou, China e-mail: [email protected] Y. Li Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), State Key Laboratory of Chemical Engineering, Tianjin Key Laboratory of Applied Catalysis Science and Technology, School of Chemical Engineering and Technology, Tianjin University, Tianjin, China Department of Chemical and Metallurgical Engineering, School of Chemical Engineering, Aalto University, Aalto, Finland © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_2

27

28

Q. Huang et al.

that time. With the development of human society, energy demand has been increasing dramatically. The negative effects caused by the large-scale consumption of fossil fuels also become a huge challenge such as climate change and air pollution. The chemical influence of air pollution affects the hydrologic cycle, thus contributing to water pollution and food security [2]. In the early 1970s, the outbreak of the energy crisis made people realize that it was urgent to find alternative sustainable resources to replace fossil fuels [3]. In 1972, Fujishima and Honda reported that TiO2 decomposed water to produce hydrogen under ultraviolet light irradiation [4], which immediately became a research hotspot because it provided a novel solution to directly convert solar energy to hydrogen, which is clean, renewable, and abundant on the earth. The research on the decomposition of toxic inorganic pollutants in water by photocatalysis was reported by Frank and Bard in 1977, where cyanide was decomposed in the presence of aqueous TiO2 suspensions [5]. Later on, it was discovered that TiO2 and some other semiconductors could catalyze the oxidation of organic compounds into CO2 and H2O under illumination, and similar effect on toxic inorganic compounds, such as chromium species, nitrites, and NOx [6]. In summary, photocatalysis technology is a promising alternative to protect the environment and relieve the energy crisis. In recent years, great progress has been made on semiconductor photocatalysis. Furthermore, advanced equipment and research methods have been developed and applied to the field of photocatalysis, which promotes the rapid development of both theoretical research and industrial technology.

2.3

Applications

The photocatalysis is widely applied in a variety of fields and products in the environmental and energy fields, including air and water purification, hydrogen evolution, and carbon dioxide conversion into gaseous hydrocarbons, sterilization, and disinfection.

2.3.1

Environmental Protection

Humankind has been facing up a situation of severe environmental pollution. Many kinds of industrial and domestic wastewater and gases containing toxic pollutants have been discharged to the environment. Due to the complex composition, low concentration, and poor degradation efficiency of many compounds, conventional removal methods have not been able to achieve the desired results. Photocatalytic technology provides a new way to deal with environmental pollution. It is an ideal solution because it utilizes solar energy and brings little secondary pollution. Furthermore, the photocatalysts are inexpensive and easy to produce. In 1976, Carey et al. [7] proposed that TiO2 could decompose chlorinated

2 Applications and Fundamentals of Photocatalysis with Solar Energy

29

biphenyl under ultraviolet light. Since then, more and more photocatalysts have been developed for the degradation of organic/inorganic pollutants to protect the environment. The photocatalysts such as TiO2 can oxidize hydrocarbons [8, 9], halogenated hydrocarbons [10, 11], organic acids [12–14], organic dyes [15–17], and nitrogenous organic compounds [18, 19] into CO2 and H2O. Inorganic pollutants [20–22] such as Cr6+ can also be reduced by the photocatalytic reaction. The developed countries have tried to apply photocatalytic technology to solve environmental pollution problems [23, 24]. Spain’s PSA Center built the first industrial-scale wastewater treatment reactor to deal with the resin plant wastewater containing various organic pollutants such as phenol, ethylene glycol, and styrene. The photocatalytic technology is also applied to air cleaners and purification of nuclear-contaminated soil in America. Similarly, in Japan, the photocatalyst has been deposited onto the floodlights of the tunnels to decompose NOx and CO in exhausted gases emitted by the vehicles into N2 and CO2. Besides, the photocatalysts have been used as indoor formaldehyde cleaner.

2.3.2

Clean Energy Production: Water Splitting and Reduction of Carbon Dioxide

Hydrogen is considered to be one of the alternatives to fossil energy, due to its high calorific value, pollution-free combustion, and recyclability. At present, hydrogen can be produced by fossil fuel reforming, electrolysis of water, microbial fermentation [25]. However, fossil fuel reforming can bring pollution to the environment and its efficiency is low. The cost of water electrolysis is a high and similar problem with the microbial fermentation method. Using solar energy to split water is regarded as a promising method owing to its low cost, nonpollution, and sustainable utilization. Fujishima and Honda’s work in 1972 is the milestone of photoelectrochemical (PEC) water splitting. Many factors affect the efficiency of PEC water splitting including the recombination of photogenerated carriers, sluggish surface reaction kinetics, and poor light absorption [26]. Carbon dioxide emissions have increased dramatically due to the use of fossil fuels, which is mainly responsible for the greenhouse effect [27]. Therefore, reducing CO2 emissions is one of the most challenging tasks. In recent years, many scientists are working on techniques to reuse CO2 also called carbon capture and storage (CCS) [28], such as the CO2 geological storage, wet scrubbing with aqueous amine solutions to remove CO2 after combustion. The main problem is how to convert CO2 into useful chemicals and avoid the regeneration of CO2 with low energy consumption. CO2 can act as a carbon source for many organic products. However, CO2 is very stable under normal circumstances and thus large extra energy is needed to reduce it. The photocatalytic reduction of CO2 into chemicals has to attract increasing attention around the global [29]. The low light conversion efficiency is still a challenge.

30

2.3.3

Q. Huang et al.

Sterilization and Disinfection

Bacteria in air or water are harmful to human health, so photocatalytic disinfection has also attracted people’s attention. The interaction between the photocatalysts and the bacteria produces active superoxide radicals and hydroxyl radicals, which can penetrate the cell wall of the bacteria and destroy the cell membranes to prevent the transport of the film-forming substances and block the respiratory system [30]. For example, the tiles decorated with TiO2 photocatalyst have the functions of bactericidal action, especially aimed at Escherichia coli, Pseudomonas aeruginosa, Staphylococcus aureus, Enterococcus faecium, etc. [31]. Such tiles could be widely used in sterilization in the daily environment (hospitals, pharmaceutical workshops, etc.). Many antibacterial tiles with photocatalyst TiO2 developed by several companies such as Toto (Japan) have been put into the market [32].

2.4

Reaction System of Photocatalysis

Up to now, there are two systems in the photocatalytic reaction: the semiconductor powder suspension system and the photoelectrochemical (PEC) cell [33]. For the former, the photocatalysts are directly dispersed in the solution, which is convenient and simple to operate. However, the photocatalysts are easy to aggregate and hard to be recovered. The latter is to fix the photocatalyst onto a conductive substrate to construct an electrode so that the oxidation and reduction reactions occur at the anode and the cathode, respectively, which facilitates the separation of the products. Meanwhile, the recombination of photogenerated carriers can be suppressed by the applied bias voltage.

2.4.1

Basic Principles of Photocatalysis

Photocatalysis belongs to the heterogeneous reaction that involves a solid phase of semiconductor and a liquid phase of electrolyte. The intrinsic photophysical properties of semiconductors play an important part in photocatalysis. This section will discuss the basic theory of photocatalysis in terms of the photophysical properties of semiconductors and the structure of the semiconductor-electrolyte interface.

2.4.2

Photophysical Properties of the Semiconductor

Semiconductors refer to materials having conductivity between conductors and insulators, which is closely related to the band structure. Illumination usually can

2 Applications and Fundamentals of Photocatalysis with Solar Energy

31

enhance the conductivity of the semiconductor. This section will introduce the photophysical properties of semiconductors and the general laws of semiconductor-light interactions.

2.4.3

Band Structure of Semiconductor

The photocatalytic reaction involves the migration and reaction of photoinduced carriers in the semiconductors. The energy band structure of the semiconductor is of great significance. The energy band structure of the semiconductors is composed of the valence band (VB), conduction band (CB), and the forbidden band (the gap between VB and CB) [34]. Figure 2.1 shows the energy band of a semiconductor at a certain temperature. The actual semiconductors always have complex band structures such as impurity levels (Fig. 2.2). They can bind electrons and generate localized electron states and thus introducing impurity levels in the forbidden band. According to the valence difference, the impurity level can be divided into donor level and acceptor level; the location is generally close to the bottom of CB for the former and to the top of VB for the latter. The relevant semiconductors are called n-type and p-type semiconductors, respectively. The excitation of electrons from the donor level to CB is much easier than that from VB, that is, the electrical conduction often depends on the electrons excited by the donors to CB. It is also easier for the electron to be excited from VB to the acceptor level than that to CB. The donor and acceptor levels are also possible traps for the photogenerated charge carriers, which can be divided into shallow and deep traps. The shallow trap level approaches either the bottom of CB (for n-type) or the top of VB (for p-type), whereas the deep trap level nearly at the center of the Fig. 2.1 The energy band of a semiconductor

Fig. 2.2 Band structure of n-type and p-type semiconductors. The Ed is referred to as the donor level in the n-type semiconductor and the Ea is referred to as the acceptor level in the p-type semiconductor

32

Q. Huang et al.

forbidden band. The deep trap level can become the recombination center, reducing the lifetime of the photogenerated charge carrier. The surface state energy levels also locate in the forbidden band and are also the recombination center.

2.4.4

Optical Absorption Spectrum of Semiconductor

When a semiconductor is illuminated by the light, the electron in the VB absorbs enough photon energy, passes over the forbidden band, and enters CB, becoming a free electron and simultaneously leaving a hole in the valence band; an electron-hole pair is generated. This process is called intrinsic absorption, which occurs only when the photon energy (hν) is greater than or equal to the forbidden bandgap (Eg) [35]. The corresponding characteristic wavelength (λ0) is the absorption edge. hν0 ¼

hc ¼ Eg λ0

ð2:1Þ

1240 Eg

ð2:2Þ

λ0 ¼

The electron transition related to intrinsic absorption can be divided into direct and indirect transitions, respectively. For direct transition, the lowest energy point of the potential energy surface of CB is vertically at the highest point of that of VB. When the photon energy absorbed by the electron is greater than the energy gap, a vertical transition occurs from VB to CB. For the indirect transition, the potential energy surface of CB drifts relative to that of VB. The electron not only absorbs photons but also exchanges a certain amount of vibration energy with the crystal lattice, that is, release or absorb a phonon according to the law of conservation of momentum. Since the energy of those extra transitions is small, the energy of indirect transition is still close to the energy gap. Figure 2.3 shows direct and indirect transitions. The electrons transfer from a low energy state to a high energy state due to the absorbed photons. The absorption coefficient (α) is used to describe the absorption intensity. Assuming that the light passes through a film with a thickness of dx, the decrement of the light intensity, dI, is proportional to the film thickness dx and the light intensity I: dI ¼ αIdx

ð2:3Þ

I ¼ I 0 exp ðαlÞ

ð2:4Þ

that is

2 Applications and Fundamentals of Photocatalysis with Solar Energy

33

Fig. 2.3 (a) Direct transitions and (b) indirect transitions. Point A is the top of VB and point B(B0 ) is the bottom of CB. The hν  hν’, neglecting the phonon energy, but the wave vectors of electrons are different between point B and B0 after photoexcitation.

The light intensity I decay as the light propagates through the semiconductor medium. In addition to intrinsic absorption, exciton absorption, free carrier absorption, impurity absorption, etc. also exist and their absorption wavelengths appear in the long-wavelength region of the intrinsic absorption band. However, those absorptions are weak, so the absorption spectrum of the semiconductor is still dominated by the intrinsic absorption.

2.4.5

Emission Spectrum of the Semiconductor: Recombination

Light emission is the opposite of light absorption. For light absorption, electrons in the low energy level absorb a suitable photon and jump to the high energy level; while electrons in the high energy level can go back to the low energy level with a photon emitting during the light emission. In a balanced semiconductor, light absorption and emission occur alternately. When there are excess electrons and holes, additional radiation would be generated resulting from their recombination, and thus, the radioluminescence happens. The recombination of photogenerated electron-hole pairs in semiconductors can be classified into direct and indirect recombination. The direct recombination means the electrons jump directly back to VB from CB and recombine with the photogenerated hole. The indirect recombination means the photogenerated electron is captured by the trap level first and then leaps to VB from the trap level to recombine with the holes. The photocatalytic effect stems from the photogenerated electron-hole pairs. After the electron-hole pairs are generated, they separate and migrate to the surface of the semiconductor to carry out the redox reactions. The

34

Q. Huang et al.

Fig. 2.4 Processes involved during photocatalysis. (i) Photoexcitation to generate charge carriers; (ii) charge carrier diffusion; (iii) bulk recombination; (iv) oxidation or reduction; (v) surface recombination

charge carrier recombination can occur during this process. Figure 2.4 shows the excitation and recombination process. After the charge carriers reach the surface of the semiconductors, whether they take part in the redox reactions depends on the positions of CB and VB of the semiconductor and the redox potential of the chemicals adsorbed on the surface of the semiconductor. Taking the PEC water-splitting reaction as an example, the conduction band potential of the semiconductor should be more negative than the hydrogen electrode potential, and the valence band potential to be more positive than the oxygen electrode potential. Many semiconductors (e.g., TiO2, ZnO, Fe2O3, CdS) can be used as photocatalysts for water splitting due to their suitable band structure. Figure 2.5 shows the band positions of several semiconductors in contact with aqueous electrolyte at pH ¼ 13.

2.5

Potential and Charge Distribution Across the Semiconductor–Electrolyte Interface

Photocatalysis is a solid-liquid heterogeneous reaction, and the interface properties will change after contacting the two phases. This section will describe the related properties of charges and potential at the interface.

2.5.1

Fermi Level

In solid-state physics, the Fermi level (EF) is the smallest possible increment of ground state energy caused by the addition of a particle in a system of noninteracting

2 Applications and Fundamentals of Photocatalysis with Solar Energy

35

Fig. 2.5 Band positions of several semiconductors in contact with aqueous electrolyte at pH ¼ 1

fermions, that is, the chemical potential of the fermion system in the ground state at the absolute zero temperature, or the maximum energy of a single fermion in the ground state, or the highest energy level occupied by the fermion [36]. Based on the Fermi-Dirac statistics, the probability that an electronic state with an energy state of E is occupied by the electrons is: f ðE Þ ¼

1 exp

EE F kT




þ1

ð2:5Þ

where k is the Boltzmann constant. The Fermi level plays an important role in determining the energy of the entire system, as well as the distribution of the charge carriers. According to the statistical theory, the Fermi level is the chemical potential of the system. The Fermi level of an intrinsic semiconductor can be calculated according to Eq. (2.6): 1 1 N E F ¼ ðE v þ Ec Þ  kTln c 2 2 Nv

ð2:6Þ

where Ec, Ev is referred to CB and VB energy, respectively; Nc, Nv referred to the effective densities of CB and VB, respectively. For the intrinsic semiconductor, its valence band is filled with valence electrons, while the conduction band is empty. Correspondingly, the Fermi level is in the

36

Q. Huang et al.

middle of the bandgap. For the n-type semiconductor, the more donors giving electrons to the conduction band are incorporated, the closer the Fermi level is to the bottom of CB. For p-type semiconductors, the higher the acceptor concentration is, the closer the Fermi level is to the top of VB. When the doping concentration is high enough to some extent, the Fermi level may enter CB or VB. The Fermi level of the solution can be obtained analogized to the solid Fermi level. When a semiconductor is immersed into a solution containing a redox couple, the oxidation state can accept electrons to become a reduction state, as the following equation shows. Ox þ ne Ð Red

ð2:7Þ

The Fermi level of the solution is equal to the electronic potential. The equilibrium of the solid-liquid system is reached when the Fermi levels or the electrochemical potentials of the solid and the solution are equal. The electronic potential (μe) is equal to the difference between the electrochemical potentials of Red (μ red) and Ox (μ ox), which can be described as: μox þ nμe ¼ μred

ð2:8Þ

The equilibrium redox potential of the standard hydrogen scale is given by the Nernst equation: φ ¼ φo 

RT C ln ox nF C red

ð2:9Þ

Combining eqs. (2.8) and (2.9): 1 RT C ln ox μe ¼ ðμred  μox Þ ¼ φo  n nF C red

E e ¼ eμe ¼ eφo 

kT C ln ox n Cred

ð2:10Þ

ð2:11Þ

Thus, the Fermi level of the solution can be described as follows: ) F redox ¼ F o 

kT C ln ox n C red

ð2:12Þ

where Fredox is the Fermi level of the solution. Cox and Cred are the concentration of the oxidation state and reduction state, respectively. Fo is the energy when both the oxidation state and reduction state are under equilibrium.

2 Applications and Fundamentals of Photocatalysis with Solar Energy

2.5.2

37

Space Charge Layer and Band Bending

If a semiconductor and a solution containing a redox couple have different Fermi levels, the electrons will migrate at the interface, resulting in different concentration of charge carriers at the surface and in the bulk. When the Fermi levels of the semiconductor and the solution are the same, the system reaches equilibrium. Meanwhile, the semiconductor band bending occurs at the surface, that is, the positions of the energy bands are different from those in the bulk. The region of the bent surface band is called the space charge layer. An n-type semiconductor will be used to illustrate the formation of a space charge layer, while the case of a p-type semiconductor is just the opposite. Figure 2.6 illustrates the formation of three types of space charge layers of an ntype and a p-type semiconductor. If the Fermi levels of the semiconductor and the solution are equal, the electrons would not flow at the interface, thus no space charge layer forms at the semiconductor surface. If the Fermi level of the n-type semiconductor is lower than that of the solution, the electrons will migrate from the solution to the surface of the semiconductor until the Fermi levels are equal. Extra electrons gather on the surface of the semiconductor, causing the band to bend downward relative to the bulk. An accumulation layer is thus formed. On the contrary, if the Fermi level of the n-type semiconductor is higher, then the electrons transfer to the solution. The concentration of electrons at the surface is lower than that in the bulk. The surface energy band of the semiconductor will bend upward relative to the bulk phase to form a depletion layer. Based on the depletion layer, electrons continue to transfer from the semiconductor to the solution so that the concentration of majority carriers (electron) on the surface of the semiconductor is lower than the minority carriers (hole). In this case, an inversion layer is formed. For p-type semiconductor, if the Fermi layer of the semiconductor is higher than that of the solution, an accumulation layer is formed. Otherwise, an depletion layer is formed.

Fig. 2.6 Space charge layer models for n- and p-type semiconductors in contact with solutions

38

2.5.3

Q. Huang et al.

Flat-Band Potential

The flat band potential is an important feature of the semiconductor/electrolyte system. It determines the semiconductor energy band position. When a semiconductor is immersed in the electrolyte, the band bending at the surface of the semiconductor occurs due to the difference of the Fermi levels of the semiconductor and the solution. A space charge layer is formed on the side of the semiconductor and a Helmholtz layer on the side of the electrolyte, as a result of the band bending on the surface of the semiconductor. If a voltage bias is applied to the semiconductor for polarization, the Fermi level of the semiconductor will be changed. The voltage bias to flatten the energy band at the surface is called the flat band potential (φfb). The flat band potential can be obtained by measuring the differential capacitance of the semiconductor space charge layer. The capacitance of the system includes the space charge layer capacitance (CSC) of the semiconductor and the Helmholtz layer capacitance (CH) of the solution in tandem. Generally, CSC is much larger than CH, so the system capacitance is approximately equivalent to CSC. CSC can be tuned by changing the polarization potential (φ). According to the Mott–Schottky equation, the relationship between CSC and V can be described as:   1 2 kT ¼ ΔφSC  CSC qεo εeN D e

ð2:13aÞ

where ΔφSC is the potential drop of the band bending, that is, ΔφSC ¼ φ  φfb. ε, εo are the dielectric constant of the semiconductor and the vacuum, respectively. ND is the doping concentration. It indicates that the reciprocal of CSC is linearly related to the polarization potential. The flat potential and doping concentration can be calculated from the x intercept and slope of 1/Csc ~ ΔφSC, respectively.

2.5.4

Quasi-Fermi Level

The electrons in the semiconductor have the same Fermi level under the thermal equilibrium state. However, if a disruption of such illumination is added to the semiconductor, the equilibrium is broken to generate nonequilibrium carriers. Because the timescale of carrier transfer within a band is much smaller than that of the electron-hole recombination across the bandgap, it is assumed that the carriers in the valence band or the conduction band are still in the equilibrium state, while the valence band and the conduction band themselves are not. Here, a local Fermi level, quasi-Fermi level (QFL), is introduced. The nonequilibrium between the valence band and the conduction band is manifested by different QFLs. QFL of the conduction band is also called the electron quasi-Fermi level, EFn. Correspondingly, the quasi-Fermi level of the valence band is called the hole quasi-Fermi level, EFp. The QFLs for electrons and holes are defined as below:

2 Applications and Fundamentals of Photocatalysis with Solar Energy

n NC p ¼ Ev þ kTln NV

E Fn ¼ Ec þ kTln E Fp

39

ð2:13bÞ ð2:14Þ

where n and p are the concentration of photostationary-state electron and hole, respectively. In general, QFL would deviate from the equilibrated Fermi level with increased concentrations of carriers. Besides, the QFL of the majority carrier deviates much less than that of the minority carrier. Figure 2.7 shows the QFL for an n-type semiconductor under illumination. EFn is close to EF, while EFp deviates from EF seriously and draws near the VB. The position of EFp at the surface is determined by the illumination intensity and the kinetics of interfacial charge transfer.

2.5.5

Surface States

The semiconductor surface is not regular because the ideal periodic lattice is interrupted at the surface. The atoms on the surface have unpaired electrons, that is, dangling bonds, which are capable of exchanging electrons or holes with the bulk, as well as adsorbing external species. These electronic energy levels are referred to as surface state energy (Ess) and they play an important role in the semiconductorliquid interface. Taking an n-type semiconductor as an example, it is used as a photoanode of a PEC cell [37, 38]. As illustrated in Fig. 2.8, the photoexcited holes in the valence band should diffuse to the surface to oxidize Red to Ox in the solution. However, they can be trapped by the surface states and recombine with the electrons. To avoid the surface state trapping, the position of the Fermi level in the Fig. 2.7 Quasi-Fermi level of an n-type semiconductor under illumination

40

Q. Huang et al.

Fig. 2.8 The downward shift of the Fermi level at the surface alters the electron occupation of surface states, and the corresponding change in surface charge leads to Fermi level pinning

semiconductor is adjusted by extra bias to regulate the occupation of Ess. Initially, the Fermi level of electrons (EFn) is located above the Ess so the surface states are full of electrons and no redox reaction happens. With more positive potential added to the n-type semiconductor, the EFn shifts downward until it is equal to Ess. Meanwhile, the Fermi level of holes (EFp) matches with the Fredox and the holes can initiate the oxidation reaction. Finally, the EFn is smaller than Ess when bias is positive enough, leading to empty surface states and increased photocurrent.

2.6

Strategies for Improving the Performance of Photocatalysts

There are many types of semiconductor materials, for example, elementary substances represented by Si, transition metal oxides represented by TiO2, compounds of IIB group elements and VIA group elements represented by CdS, and compounds of IIIA-VA group elements represented by GaAs. With the development of photocatalysis research, many complex compounds such as titanates (NaTaO3 [39, 40]), vanadates (BiVO4 [41, 42]), tantalites (SrTiO3 [43–45]), tungstates (Bi2WO6 [46, 47]), oxyhalides (BiOCl [48, 49]), and solid solutions (Bi2InNbO7 [50]) have been constructed. Although these semiconductors have been developed for PEC water splitting and water pollutant decomposing, they still suffer from low efficiencies, which originates from poor light absorption, severe photogenerated carrier recombination, slow surface reaction kinetics, etc. Therefore, it is of great importance to design the photocatalysts rationally. This section will focus on four strategies to modify the semiconductors in terms of doping, heterostructure construction, morphology control, and cocatalyst loading.

2.6.1

Doping

Doping is an effective way to change the physicochemical properties of the semiconductors. The addition of foreign elements to the semiconductor crystal lattice can

2 Applications and Fundamentals of Photocatalysis with Solar Energy

41

cause lattice defects, thereby affecting the energy band structure and the movement and distribution of photogenerated carriers. Typically, doping can serve the purposes as below: (i) Doping can introduce some impurity levels in the forbidden band, which makes it possible for the semiconductors with a wide bandgap to absorb longer wavelength photons in the visible light. (ii) The dopants can participate in the capture and release of photogenerated electrons and holes, control the diffusion of carriers during the photocatalytic reaction process. (iii) Doping can increase the carrier density and improve the electrical conductivity of the semiconductors. (iv) Doping affects the shape-controlled growth of semiconductors. Many methods can be applied for doping, such as hydrothermal methods [51, 52], ion thermal diffusion from the substrate [53, 54], dip coating [55], impregnation, high-energy ion implantation [56, 57]. For wide bandgap semiconductors, such as TiO2 with a bandgap of ~3.2 eV, they usually suffer from limited light absorption. Dopants have been introduced to overcome this issue. Lee et al. [52] prepared transitions metal-doped TiO2 by the hydrothermal reaction. The UV–visible spectroscopy analysis showed that Fe3+-, Cr4+-, Ni2+-doped samples can significantly enhance the light absorption in the visible region. Accordingly, the doped TiO2 achieved higher benzene photodecomposition efficiency in the artificial sunlight, especially the Cr4+-modified sample. Nonmetallic elements can also tune the TiO2 energy band. Asahi et al. [58] reported an N-doped TiO2 whose absorption edge was shifted to the visible light region. Due to the narrowed bandgap, it achieved photodegradation of methylene blue and acetaldehyde under visible light. Other nonmetallic elements such as C [59], S [60, 61] have also been studied. It should be noted that the impurity level should be tuned properly in the bandgap, otherwise it will become the recombination center. This is closely related to the concentration, distribution, and d-electron configuration of the dopants. Choi et al. [62] studied the effects of 21 kinds of metal ions doped TiO2 used for chloroform oxidation and carbon tetrachloride reduction. The Fe3+, Mo4+, Ru3+, Os3+, V4+, and Rh3+ ions with a doping concentration of 0.1–0.5% can significantly increase the TiO2 photocatalytic activity, whereas the Li+, Mg2+, A13+, and Ga3+ ions have a negative influence. The ions with positive effect can act as both electron and hole traps and the trapped charge carriers have a sufficiently long lifetime to reach the surface by detrapping and electron tunneling. Hematite (Fe2O3) can respond to visible light due to its narrow bandgap (1.9–2.2 eV), which is an advantage relative to TiO2. However, Fe2O3 has a poor conductivity due to the highly localized small polarons, leading to serious carrier bulk recombination. To enhance electrical conductivity, many elements have been incorporated such as Sn [51], Ti [54], Al [63], Cr [64], Si [65], P [66]. Ling et al. prepared Sn-doped Fe2O3 by hydrothermal method and Sn was introduced by a facile dip coating procedure. The Mott–Schottky plots showed that the donor density

42

Q. Huang et al.

was much higher after ex situ doping [67]. Cesar et al. [68] prepared Si-doped Fe2O3 and found that, in addition to the increased conductivity, Si doping reduced the particle size which could reduce the diffusion length for carriers from the bulk to the surface, further decreasing the recombination. The distribution of dopants in lattices is critical for the photocatalytic activity. Luo et al. [66] proposed a P gradient-doped Fe2O3 nanoarray by controlling the annealing time after a facile dipping. From Fig. 2.9a, b, the Mott–Schottky plots indicated that whether gradient doping or homodoping was helpful to increase the carrier density. However, the graded distribution of the P element made the band bending widen and accelerated the separation of photogenerated carrier compared to that of homogeneous distribution so as to enhance photocurrent. Besides introducing foreign elements, native defects like vacancies or interstitials also influence the carrier densities of semiconductors. Oxygen vacancy is very common in metal oxides and useful in tuning the electrical conductivity and photocatalytic activity [69–72]. For example, high-temperature calcination is generally needed for the iron oxide materials prepared by the hydrothermal method. Xiao et al. found that rapid cooling caused a partial reduction of the surface of iron oxide (Fig. 2.10a), that is, oxygen vacancies were introduced, which contributes to increased carrier densities (Fig. 2.10b), and thus, better PEC water-splitting performance [73]. Similar to doping, the fabrication of a solid solution that incorporates a certain component into a crystal usually through a solid reaction, and the product retains the structure of the original one, is also an approach to regulate the crystal properties. There are many multicomponent semiconductors reported for photocatalysis, for example, GaN:ZnO [74–76], AgSrNbTiOx [77, 78], Cu3MS4 (M ¼ V, Nb, Ta) [79], ZnxCd1-xS [80, 81], CdxIn1-xS [82], In–Ni–Ta–O–N [83], GaP–ZnS [84]. The bandgap of multicomponent semiconductors can be adjusted by changing the stoichiometry of the different components. Li et al. [80] found that the absorption band

Fig. 2.9 (a) Mott–Schottky plots of bare Fe2O3, homo-P:Fe2O3, and grad-P:Fe2O3 photoanodes. (b) I–V characteristics of bare Fe2O3, homo-P:Fe2O3, grad-P:Fe2O3, homo-P:Fe2O3/Co-Pi, and grad-P:Fe2O3/Co-Pi photoanodes measured in 1 M KOH solution under AM 1.5G illumination. (Reprinted with permission from [65], copyright 2017 Royal Society of Chemistry)

2 Applications and Fundamentals of Photocatalysis with Solar Energy

43

Fig. 2.10 (a) High-resolution XPS spectra of the Fe 2p. (b) Mott–Schottky plots. RC referred to rapid cooling, NC referred to natural cooling, SC referred to slowing cooling. (Reprinted with permission from [72], copyright 2017 Elsevier B.V.)

edge was red shifted continuously as the Cd content increased in the ZnxCd1-xS solid solution as shown in Fig. 2.11A, consistent with the color change from white to yellow. The sample with a Zn/Cd molar ratio of 1:1 exhibited the best H2 evolution performance. Another typical example is the GaN:ZnO solid solution. The bandgaps of ZnO and GaN are too wide to absorb visible light, while their solid solutions show a narrowed bandgap for the visible light absorption reflected from the UV–visible absorption spectrum in Fig. 2.11B [74]. Although separated InN or GaZnON shows no water-splitting activity, the Ni-InN/GaZnON solid solution exhibits overall water splitting under visible light and good photocatalytic stability [85].

2.6.2

Heterojunction/Homojunction

As a photocatalyst, semiconductors usually should have a suitable bandgap to ensure light absorption, efficient photogenerated carrier separation, and resistance to photocorrosion. Also, the energy levels of the conduction/valence band are proper to drive the redox reaction. However, no single semiconductor has met all of the above conditions so far. Correspondingly, the integration of two or more materials, that is, design heterojunction/homojunction, is a solution. A homojunction is from the interface of the same semiconductor materials but with different doping. On the contrary, a heterojunction is from the interface of different semiconductors (semiconductor-semiconductor heterojunction) or that of metals and semiconductors (metal-semiconductor heterojunction). Both can be built to improve the photocatalytic performance.

44

Q. Huang et al.

Fig. 2.11 (A) UVvisible diffuse reflection spectra of the Zn1xCdxS (x ¼ 0, 0.1, 0.3, 0.5, 0.7, 0.9, and 1.0) samples. Reprinted with permission from [79], copyright 2013 Royal Society of Chemistry. (B) UV–visible diffuse reflectance spectra: (a) GaN, (b) GaN:ZnO (Zn 3.4 atom %), (c) GaN: ZnO (Zn 6.4 atom %), (d) GaN:ZnO (Zn 13.3 atom %), and (e) ZnO. Reprinted with permission from [73], copyright 2005 American Chemical Society

2.6.3

Semiconductor-Semiconductor Heterojunction

Regardless of the influence of surface states, the energy band of heterojunction is shown in Fig. 2.12. When two n-type semiconductors (SC 1 and SC2) with different energy band positions and Fermi level contact with each other, the electrons flow from the side with higher Fermi level (SC1) to the other side with lower Fermi level (SC 2). A built-in electric field is constructed at the interface, whose electric field direction is from SC1 to SC2. When the Fermi levels of the two semiconductors reach the same, the electrons will stop flowing and the heterojunction is formed. The charge layers are formed on both sides of the semiconductor contact surface. If both the semiconductors are n-type, a depletion layer will be formed on the surface of SC1 who loses electrons and an accumulation layer will be formed on the surface of SC2 who accepts excess electrons. The potential difference between the two sides is the contact potential difference which is equal to the difference between the two Fermi levels. Due to the redistribution of electrons at the interface, the carrier concentration of the semiconductor surface is different from that of the bulk, resulting in band bending. An ideal semiconductor-semiconductor composite system should satisfy the following conditions: (a) the energy band structures of the coupled semiconductors can be matched to each other so that the band bending at the interface can force photogenerated electron-hole pairs to separate; (b) the ohmic contact is necessary, so the charge can move smoothly at the interface; (c) after the charge is separated at the interface, the semiconductor as an electron acceptor should have excellent electron transport capability, and the other semiconductor as electron donor should be in favor of hole transport.

2 Applications and Fundamentals of Photocatalysis with Solar Energy

45

Fig. 2.12 Heterojunctions between two semiconductors. (a) n-n heterojunction and (b) p-n heterojunction

Wide bandgap semiconductors have limited absorption of sunlight, but combining them with narrow bandgap semiconductors can solve this problem. TiO2 is a typical wide bandgap semiconductor and hardly responds to visible light. Many semiconductors (metal oxides: Bi2WO6 [86], WO3 [87, 88], Fe2O3 [89–91]; sulfides: CdS [92, 93], MoS2 [94]) have been integrated with TiO2 and successfully improved the visible light absorption. Aside from visible light activity, a reasonable heterojunction can also promote the separation of photoinduced carriers due to the built-in electric field, such as p-n junctions. Dai et al. [95] deposited BiOI on the TiO2 nanotube wall. A p-n heterojunction was formed at the interface. As shown in Fig. 2.13a, such an internal field facilitated the excited electrons on the conduction band of the p-type BiOI to transfer to that of n-type TiO2, while the holes remained in the BiOI valence band, suppressing the carrier recombination. Similarly, MoS2 nanosheet-coated TiO2 nanobelts could split water to produce hydrogen and decompose the organic dye Rhodamine B [96] due to the energy band matching of MoS2 and TiO2, which not only increases the light absorption, but also inhibits the photogenerated carrier recombination. Cong et al. [97] deposited iron oxide nanoparticles onto TiO2 nanotubes by electrodeposition. The modified TiO2 nanotubes could decompose phenol under visible light, and the decomposition efficiency was greatly improved owing to the formation of heterojunction. Graphitic carbon nitride (g-C3N4) is a metal-free semiconductor and has attracted great attention since it was applied in photocatalysis considering its suitable bandgap for visible light absorption, high stability in an aqueous solution, and nontoxicity [98–100]. It can also be used for heterojunction formation with other semiconductors like TiO2 [101–103], ZnO [104–107], CdS [108, 109], and Fe2O3 [110–112]. When g-C3N4 is loaded on a semiconductor, the light absorption and the photogenerated carrier separation efficiency can be significantly improved if band structures are matched. The absorbance intensity of the g-C3N4–ZnO composite photocatalysts

46

Q. Huang et al.

Fig. 2.13 (a) Schematic illustration for the formation of the p-n junction between p-type BiOI and n-type TiO2 under visible light irradiation. Reprinted with permission from [94], copyright 2011 American Chemical Society. (b) UV–vis diffuse reflectance spectra of g-C3N4, ZnO, and g-C3N4– ZnO composites. Reprinted with permission from [103], copyright 2012 Royal Society of Chemistry

became stronger and the absorption edge shifted to the lower energy displayed in Fig. 2.13b. Accordingly, an enhanced photodegradation of methyl orange (MO) and p-nitrophenol was obtained [104].

2.6.4

Metal-Semiconductor (M-S) Heterojunction

The heterojunction can also be formed between metals and semiconductors, generally based on the n-type semiconductors. When the metal is in close contact with the n-type semiconductor, a heterojunction, also called Schottky junction, with an energy barrier can be created. Metals and n-type semiconductors have different Fermi levels. Usually, the work function of the metal (φm) is higher than that of the semiconductor (φs). Therefore, the electrons will continuously transfer from the semiconductor to the metal until a new equilibrium if the two are closely integrated. The metal acts as an electron collector, and the excess positive charge will accumulate on the surface of the n-type semiconductor so that the bands will be bent upward and a depletion layer is generated (Fig. 2.14). M-S heterojunction has been employed in photocatalysis, especially for Ag [113– 115], Au [116–118] nanoparticles with the property of localized surface plasmon resonance. Intrinsically they have strong absorption for visible light and can promote the charge transfer with semiconductors where plasmonic nanoparticles are loaded on. Wu et al. [119] produced Ag/TiO2 nanotubes. After the Ag nanoparticles were deposited, the visible light absorption was enhanced. Other noble metals like Pd [120] and Pt [121, 122] also appear in M-S heterojunction for photocatalysis. They can act as electron collectors to accelerate charge carrier separation at the surface of semiconductors [123, 124]. Owing to the high cost of noble metals, cheaper alternatives should be developed. Some earth abundant metals, such as Ni [125], Co [126], Cu [127] can play a similar role. Chen et al. [128] demonstrated that the H2

2 Applications and Fundamentals of Photocatalysis with Solar Energy

47

Fig. 2.14 Schematic of Schottky junction formed between metals and n-type semiconductors

production rate over 1.2 wt% Ni-loaded CdS is nearly twice higher than that of 1.2 wt% Pt-loaded CdS under visible light.

2.6.5

Carbon-Semiconductor Heterojunction

In addition to the non-noble metals, carbon materials such as graphite, carbon nanotubes (CNTs), and carbon quantum dots (CQDs) are good alternatives to noble metals. Besides excellent conductivity, graphite and CNTs can supply a large surface area. Yu et al. [129] prepare mesoporous multiwalled CNTs/titanium dioxide. Only the chemically bonded composites could facilitate the separation of photogenerated electron-hole pairs and decrease their recombination rate and thus significantly enhanced photocatalytic activity for degrading acetone in the air under UV irradiation. Li et al. [130] synthesized graphene nanosheets decorated CdS clusters through a solvothermal method. These nanosized composites exhibited a high H2 production rate owing to the improved carrier transport by graphene. For CQDs, they stand out due to the size-dependent photoluminescence and the up-conversion photoluminescence, that is, CQDs can help harvest more photons. Yu et al. [131] reported that the absorption band edge of the CQDs-sensitized TiO2 composites was extended to the visible light region. The CQDs serve as electron reservoirs to trap the photoinduced electrons and enhance the separation of photoinduced electron-hole pairs. Accordingly, the photocatalytic activity of H2 production was improved. A variety of heterojunctions can be combined to modify the photocatalysts, which means the semiconductor-semiconductor heterojunctions and M-S heterojunctions can be produced in tandem. Zou’s group fabricated a ternary MoS2-graphene/ ZnIn2S4 composite, which showed a higher H2 photocatalytic evolution rate than Pt-loaded ZnIn2S4 photocatalyst [132]. A schematic was exhibited in Fig. 2.15a. It was believed that the graphene served as an electron transport bridge between the MoS2 and the ZnIn2S4, resulting in reduced charge recombination. Zhang et al.

48

Q. Huang et al.

deposited Ag nanoparticles to the interface of nanoscale ZnO/CdS, obtaining a ZnO/Ag/CdS nanocomposite. The modification with Ag nanoparticles significantly enhances the light absorption and facilitates the separation of photogenerated carriers through the localized surface plasma resonance (Fig. 2.15b) [133]. Li’s group constructed a novel nanocomposite NiO/CdS@ZnO photoanode for water splitting. The well-matched band structures between the composite components significantly promoted the transfer of charge carriers as evidenced by the decreased photoluminescence intensity [134].

Fig. 2.15 (a) Schematic illustration for photocatalytic hydrogen generation over MoS2-graphene/ ZnIn2S4 under visible light irradiation. Reprinted with permission from [131], copyright 2016 Elsevier B.V. (b) The schematic diagrams of the reaction process and carriers transport in ZnO/Ag/ CdS composite photoanode. (Reprinted with permission from [132], copyright 2014 Elsevier B.V.)

2 Applications and Fundamentals of Photocatalysis with Solar Energy

2.6.6

49

Homojunction

Doping can regulate the energy band structure of a semiconductor, which makes it possible to form a homojunction between different phases with the same chemical composition. Abdi et al. [135] constructed heterojunctions in series based on W-doped BiVO4 to suppress the photocarriers recombination. They fabricated the gradient-doped W:BiVO4 by spray pyrolysis. Two phases with different doping concentrations form a homojunction, and the photogenerated carriers can be separated and diffused to the surface under the driving force of the built-in electric fields in series. Similar junctions were applied to the ZnO nanostructure. Xiao et al. produced a ZnO:Ga/ZnO photoanode by a two-step hydrothermal reaction. The as-prepared ZnO:Ga/ZnO composite with nearly the same lattice structure in the two phases exhibits enhanced light absorption and improved electrical conductivity [136].

2.7

Morphology Control

The morphology of photocatalysts plays an important role in photocatalytic performance. A sufficient surface area is necessary to ensure light absorption and reactive sites. The photocatalysts should grow along special crystal planes due to the crystal anisotropy. The microstructure determines the transfer distance of the carriers from the excitation site to the surface, which should match the lifetime of the photogenerated hole-electron pairs to inhibit the recombination [137]. In such a context, nanostructures are desirable to achieve high performance. Accordingly, various nanostructures such as zero-dimensional (0D), one-dimensional (1D), twodimensional (2D), and three-dimensional (3D) have been proposed. 0D nanostructure – 0D nanostructures are typically nanoparticles, like nanospheres, core-shell structures, hollow spheres. Chen et al. produced submicrometer TiO2 hollow spheres (Fig. 2.16a) by a template-free solvothermal route. The hollow nanostructures exhibited good photocatalytic activity on the degradation of phenol, even superior to Degussa P25 [138]. In general, the smaller the size of the nanoparticle, the more specific surface area can be provided. However, when it is smaller enough to exhibit the quantum confinement effect, the absorption edge will move to the blue side with reduced light absorption. The band bending, a built-in field for the photon-generated carrier separation also disappears [139]. Also, if such nanoparticles are immobilized on a conductive substrate, the recombination of photogenerated carriers will be enhanced by the grain boundaries [140]. 1D nanostructure – Compared to the 0D nanostructures, the 1D nanostructures not only have a large surface area but also can provide an orthogonal direction for

50

Q. Huang et al.

Fig. 2.16 Examples of 0D and 1D nanostructures: (a) TiO2 nano hollow spheres. Reprinted with permission from [137], copyright 2011 American Chemical Society. (b) TiO2 nanotubes. Reprinted with permission from [147], copyright 2001 Materials Research Society. (c) Zr-doped Fe2O3 nanotubes. (Reprinted with permission from [148], copyright 2017 Wiley-VCH)

charge carrier separation that significantly enhances electron/hole collection at the back-contact in a PEC cell. Based on the high length-to-diameter ratio for the 1D nanostructures, light absorption and scattering are significantly enhanced. 1D nanostructures such as nanotubes (NTs) [59, 141, 142], nanorods (NRs) [143– 145], and nanowires (NWs) [67, 146] have been fabricated by hydrothermal reaction [53], electrodeposition [64], chemical vapor deposition [147], etc. Titanium dioxide nanotubes (Fig. 2.16b) are usually prepared by one-step anodization of titanium foil, which was first reported in 2001 by Gong’s group [148]. Recently, Li et al. [149] prepared Zr-doped Fe2O3 nanotubes with an outer diameter of ca. 100 nm (as displayed in Fig. 2.16c) through a solid-state reaction at a high temperature of 800 C. Such a nanostructure of Zr-doped Fe2O3 provided a shortened charge transfer distance and an enlarged depletion layer at both inner and outer surfaces when contacting with the electrolyte. 2D Nanostructure – 2D nanostructures include nanosheets, nanoflakes, a nanobelt, etc. They exhibit relatively large lateral size and an ultrathin thickness confined to the atomic scale regime that collectively leads to surface areas typically larger than those of 1D nanostructures. Moreover, it is easier to control the exposed active facets in 2D nanostructures. By precise control of the growth conditions, some nanosheets provide almost 100% exposed active sites, such as the (110) facet exposed Fe2O3 nanosheets (Fig. 2.17a, b) obtained through a silica hydrogel-mediated dissolution-recrystallization process [150]. Such ultrathin Fe2O3 nanosheets exhibited superior degradation photoreactivity of an organic contaminant in comparison with the α-Fe2O3 nanoparticles. The impressive enhancement was attributed to a special lamellar nanostructure with high surface area and improved conductivity. Another way to synthesize Fe2O3 with 2D nanostructure has been demonstrated by Yang’s group. They oxidized the iron substrate under the thermal flow of O2 and the Fe2O3 nanobelts were formed on the surface [151].

2 Applications and Fundamentals of Photocatalysis with Solar Energy

51

3D Nanostructure – It is also called a hierarchical nanostructure. It consists of a conspicuous backbone with nanobranches which are usually nanoparticles, nanowires/rods/tubes, and nanosheets. Thus, the 3D nanostructure often resembles nanotrees or nanoflowers with significantly increased surface area and a much larger number of active sites, owing to the high arrangement provided by the branches [137]. Hierarchical nanostructures can either be composed of a single material or a combination of two or more. For only one material, Liu et al. [152] reported that a flower-like CdS was fabricated through a low-temperature mixed solvothermal strategy. The flower structure constructed of several assembled nanosheets with a highly preferred (002) facet is in favor of light absorption and charge transfer. Thus, the flower-like CdS displayed better performance in photocatalytic hydrogen evolution than nanobelt CdS. Wang and coworkers fabricated ZnO nanoflowers (Fig. 2.17c, d) assembled by ZnO nanorods [153]. It was demonstrated that flowers structured ZnO could be controllable by simply adjusting the basicity in the solution. The as-prepared ZnO nanoflowers had a higher content of oxygen vacancies to serve as active centers, leading to enhanced photocatalytic activity. Cauliflower-type structured Fe2O3 with a minimum feature size of 5–10 nm (Fig. 2.17e, f) was produced by atmospheric pressure chemical vapor deposition [68, 154]. It was found that silicon doping played an important part in the formation of such a fine nanostructure. Hierarchical nanostructures composed of two or more materials are more advantageous. The backbone and branches can be made up of different compositions, thus it is possible to construct heterojunctions: WO3-Fe2O3 [155, 156], CdS-TiO2 [157, 158], ZnO-Fe2O3 [159, 160], and Fe2O3-TiO2 [161–163] assemblies have been reported. Han et al. [164] fabricated a 3D hierarchical hetero-nanostructures composed of thin α-Fe2O3 nanoflakes branched on TiO2 nanotubes, which exhibited enhanced activity in photoelectrochemical water splitting and photocatalytic degradation of an organic dye compared with pure TiO2 nanotubes due to the optimized light absorption and the enhanced separation of photogenerated charge carriers at the interface. Yang’s group constructed tree-like heterostructures comprising Si nanowires as backbones and TiO2 nanorods as branches as showed in Fig. 2.18 [165]. The nanotree heterostructure displayed much higher photoactivity than the configuration where a TiO2 thin film is partially covered onto Si nanowire due to light absorbers with large surface area and complementary electrical properties of two materials.

2.8

Loading Cocatalysts

The photocatalytic effect of semiconductors originates from the photoexcited holes with oxidizing ability and electrons with reducing ability. However, photogenerated carriers on the surface tend to recombine rather than react due to the surface states.

52

Q. Huang et al.

Fig. 2.17 Examples of 2D and 3D nanostructures: (a, b) Fe2O3 nanosheets. Reprinted with permission from [149], copyright 2017 Elsevier B.V. (c, d) ZnO nanoflowers. Reprinted with permission from [152], copyright 2017 Elsevier B.V. (e, f) Fe2O3 cauliflower-type nanostructure. Reprinted with permission from [67], copyright 2006 American Chemical Society

2 Applications and Fundamentals of Photocatalysis with Solar Energy

53

Fig. 2.18 Examples of hierarchical nanostructures: (a) Schematic of the Si-TiO2 tree-like heterarchical structure. (b, c) SEM images of Si-TiO2 nanotrees. (Reprinted with permission from [164], copyright 2017 American Chemical Society)

Also, slow surface reaction kinetics accelerates the surface recombination, which severely inhibits the efficiency of the photocatalytic reaction. A common and often useful strategy to solve those problems is to integrate a cocatalyst with a photocatalyst. Generally, suitable cocatalysts can accelerate the surface reaction kinetics and inhibit the surface electron-hole recombination, that is, the cocatalysts can capture the holes/electrons from the semiconductors and then transfer them to the liquid, providing a shortcut for the surface reaction. Some semiconductors are electrochemically unstable in the electrolyte environment or easy to be corroded by the photoinduced holes. The cocatalysts can act as physically protective layers or can extract photogenerated carriers from the semiconductors and help them participate in the reaction rather than accumulate on the semiconductor surface to prevent semiconductors from photocorrosion [166]. For PEC water splitting, cocatalysts can be derived from electrocatalysts such as metal phosphates/phosphides [167–169], metal oxides [170–172], (oxy)-hydroxides [173–175]. They can be prepared by atomic layer deposition (ALD), electrochemical deposition, spin coating, dip casting, etc. Cobalt phosphate (Co-Pi) is one of the typical electrocatalyst [176]. Many groups have successfully integrated Co-Pi onto photocatalysts [73, 177, 178], and achieved significantly improved PEC water-splitting performance. However, there are different perspectives in terms of the role of Co-Pi. Klahr et al. [179] deduced that Co-Pi could collect the photogenerated holes from the Fe2O3 photoanode. Co(III) species in Co-Pi film was oxidized to Co(IV) which is more beneficial for water oxidation, leading to reduced resistance and accelerated surface reaction kinetics (Fig. 2.19a, b). However, based on the impedance and PEC measurement, Carroll et al. [180] considered that slow charge transfer between the solid-liquid interfaces limited the photocurrent of Co-Pi/Fe2O3 photoanode while a reduced hole-electron recombination was displayed compared to bare Fe2O3. And the latter was more dominant to obtain a net enhancement of water-oxidation efficiency. In short, those different viewpoints probably are originated from the differences in the loading amount of Co-Pi cocatalysts and the microstructure of the underlying photoanodes [181]. Besides Co-Pi, Fe-Pi is also an efficient cocatalyst for Fe2O3 photoanode.

54

Q. Huang et al.

Fig. 2.19 (a) Nyquist plots measured under the illumination of a bare hematite electrode (red circles) and with 15 mC cm2 CoPi catalyst (green triangles) measured at 1.25 V versus RHE. (b) Diagram of a CoPi-coated hematite electrode under illumination and applied bias. Reprinted with permission from [178], copyright 2012 American Chemical Society. (c) Applied voltage versus densities of surface states (DOS) of the electrodes. Solid lines correspond to Lorentzian fit results. (d) Proposed kinetic scheme of the hole transfer pathways from the VB of Fe2O3 to the surface states for the Fe2O3 and FeCoW/Fe2O3 photoanodes under illumination at a bias of 1.0 V versus RHE. Shaded and blank areas represent electron filled or empty states, respectively. (Reprinted with permission from [190], copyright 2017 Elsevier B.V.)

When the surface of Fe2O3 was transformed into amorphous iron phosphate, the oxygen atoms are covalently fixed in phosphate. Accordingly, the oxygen vacancies are decreased and the surface states are effectively suppressed, leading to improved photocatalytic activity [167]. Some metal oxides also have similar passivation effect [182, 183]. Formal et al. [183] deposited an ultrathin alumina overlayer on hematite by ALD. The onset potential of the modified Fe2O3 photoanode was reduced by 100 mV due to the suppression of surface states. Many excellent catalysts (e.g., Co- and Ni-based oxides and (oxy)-hydroxides) for electrocatalytic oxygen evolution reaction (OER) can be employed as cocatalysts for photoanode. Zhan et al. [184] deposited CoOx on the WO3 surface and obtained an increased PEC water-splitting efficiency. The cobalt ion valence cycle (Co3+/Co4+) suppressed the surface carrier recombination and enhanced the reaction kinetics. A similar CoOx/WO3 structure was fabricated by Ding’s group [185]. They attributed the improved charge transfer kinetics and Faradaic efficiency to the formation of pn heterojunctions and the increased oxidation selectivity caused by CoOx.

2 Applications and Fundamentals of Photocatalysis with Solar Energy

55

Also, CoOx can protect the photoanode as well. Higashi et al. [186] claimed that the stable photocurrent of modified TaON was due to the highly dispersed CoOx nanoparticles that efficiently collected photogenerated holes to prevent TaON from self-oxidative deactivation. A similar effect was found in the NiOx layer-coated Si photoanode [187]. Du et al. [188] reported that NiFeOx-decorated Fe2O3 showed an excellent turn-on voltage of 0.61 V versus reversible hydrogen electrode (RHE). The author explained the negative shift of turn-on voltage compared to bare Fe2O3 from the thermodynamics rather than the kinetics, that is, the NiFeOx layer increased the photovoltage of the Fe2O3 photoanode. An amorphous NiOOH layer on the Fe2O3 photoanode led to ca. 150 mV cathodic shift in turn-on voltage [189]. The photogenerated holes could be trapped by Ni2+/Ni3+ pairs in Ni(OH)2/NiOOH before the surface states, followed by detrapping to oxidize water, resulting in improved photocurrent onset potential when the bias was low. When the bias was up to 1.25 V versus RHE, the oxidation process (Ni3+/Ni4+) occurred which was slow, though Ni4+ could oxidize water fast. A similar effect was found in Ni(OH)2 [190]. Li’s group reported a FeCoW oxyhydroxide gel modified α-Fe2O3 photoanode. The onset voltage of the composite photoanode was a shift to the cathode side by 170 mV. The FeCoW oxyhydroxide gel acted as a hole collector to inhibit surface recombination of Fe2O3 as well as shifted the surface state to a higher energy position to obtain a reduced turn-on voltage (Fig. 2.19c, d) [191].

2.9

Conclusions and Future Trends

Photocatalysis is a green technology because it employs inexhaustible solar energy to realize energy conversion without any toxic and harmful by-products, beneficial to both environmental protection and energy development. However, up to now, no semiconductor can perfectly achieve high solar conversion efficiency. Over the past few decades, many different strategies have been successfully proposed, and significant progress has been made to modify photocatalysts. To develop high-performance photocatalysts, attention can be paid to investigating the synergistic effects of the multimethod combination on the activity for photocatalysis in future researches. Further insight should be taken into detailed thermodynamics and kinetics of electron-hole transfer and reactions at the interface, which is crucial to design more efficient solar-to-fuel conversion systems. Although there are still many challenges ahead, it is believed that in the future, photocatalytic technology will penetrate every aspect of our daily life and provide more convenience for human beings with the unremitting efforts of researchers. Acknowledgments This work is partially supported by Prof. Lingbo Luise Liu. The authors also gratefully acknowledge the helpful comments and suggestions of the reviewers, which have improved the presentation.

56

Q. Huang et al.

Authors’ Contribution Qiuyang Huang wrote the paper. Dan Kong polished the paper and discussed with Yongdan Li the content of this work and finalized the paper. All authors have approved the final version of the manuscript.

References 1. Hashimoto K, Irie H, Fujishima A (2005) TiO2Photocatalysis: a historical overview and future prospects. Jpn J Appl Phys 44(12):8269–8285 2. Lelieveld J, Klingmüller K, Pozzer A, Burnett RT, Haines A, Ramanathan V (2019) Effects of fossil fuel and total anthropogenic emission removal on public health and climate. Proc Natl Acad Sci 116(15):7192–7197 3. Gratzel M (2001) Photoelectrochemical cells. Nature 414(6861):338–344 4. Fujishima A, Honda K (1972) Electrochemical photolysis of water at a semiconductor electrode. Nature 238(5358):37–38 5. Frank SN, Bard AJ (1977) Heterogeneous photocatalytic oxidation of cyanide ion in aqueous solutions at titanium dioxide powder. J Am Chem Soc 99(1):303–304 6. Hoffmann MR, Martin ST, Choi W (1995) Environmental applications of semiconductor photocatalysis. Chem Rev 95(1):69–96 7. Carey JH, Lawerence J, Tosine HM (1976) Photocatalytic treatment of polychloribated biphenyls on TiO2. Bull Environ Contam Toxicol 16:696–701 8. Matthews R (1987) Carbon dioxide formation from organic solutes in aqueous suspensions of ultraviolet-irradiated TiO2. Effect of solute concentration. Aust J Chem 40(4):667–675 9. Vidal A, Herrero J, Romero M, Sanchez B, Sanchez M (1994) Heterogeneous photocatalysis: degradation of ethylbenzene in TiO2 aqueous suspensions. J Photochem Photobiol 79 (3):213–219 10. Lichtin NN, Dong J, Vijayakumar K (1992) Photopromoted TiO2-catalyzed oxidative decomposition of organic pollutants in water and the vapor phase. Water Qual Res J Can 27 (1):203–210 11. Sabin F, Türk T, Vogler A (1992) Photo-oxidation of an organic compound in the presence of titanium dioxide: determination of the efficiency. J Photochem Photobiol A 63(1):99–106 12. Kormann C, Bahnemann D, Hoffmann MR (1991) Photolysis of chloroform and other organic molecules in aqueous titanium dioxide suspensions. Environ Sci Technol 25(3):494–500 13. Carraway ER, Hoffman AJ, Hoffmann MR (1994) Photocatalytic oxidation of organic acids on quantum-sized semiconductor colloids. Environ Sci Technol 28(5):786–793 14. Ollis DF, Hsiao C-Y, Budiman L, Lee C-L (1984) Heterogeneous photoassisted catalysis: conversions of perchloroethylene, dichloroethane, chloroacetic acids, and chlorobenzenes. J Catal 88(1):89–96 15. Reeves P, Ohlhausen R, Sloan D, Pamplin K, Scoggins T, Clark C, Hutchinson B, Green D (1992) Photocatalytic destruction of organic dyes in aqueous TiO2 suspensions using concentrated simulated and natural solar energy. Sol Energy 48(6):413–420 16. Matthews RW (1989) Photocatalytic oxidation and adsorption of methylene blue on thin films of near-ultraviolet-illuminated TiO2. J Chem Soc Faraday Trans 85(6):1291–1302 17. Matthews RW (1991) Photooxidative degradation of colored organics in water using supported catalysts. TiO2 on the sand. Water Res 25(10):1169–1176 18. Matthews R, Abdullah M, Low G-C (1990) Photocatalytic oxidation for total organic carbon analysis. Anal Chim Acta 233:171–179 19. Grätzel CK, Jirousek M, Grätzel M (1990) Decomposition of organophosphorus compounds on photoactivated TiO2 surfaces. J Mol Catal 60(3):375–387 20. Milis A, Peral J, Domenech X, Navio J (1994) Heterogeneous photocatalytic oxidation of nitrite over iron-doped TiO2 samples. J Mol Catal 87(1):67–74

2 Applications and Fundamentals of Photocatalysis with Solar Energy

57

21. Sabate J, Anderson M, Aguado M, Giménez J, Cervera-March S, Hill CG Jr (1992) Comparison of TiO2 powder suspensions and TiO2 ceramic membranes supported on the glass as photocatalytic systems in the reduction of chromium (VI). J Mol Catal 71(1):57–68 22. Aguado M, Giménez J, Cervera-March S (1991) Continuous photocatalytic treatment of Cr (VI) effluents with semiconductor powders. Chem Eng Commun 104(1–3):71–85 23. Fujishima A, Hashimoto K, Watanabe T (1999) TiO2 Photocatalysis: fundamentals and applications. Bkc. Inc., Tokyo, p 176 24. Jones AP (1999) Indoor air quality, and health. Atmos Environ 33(28):4535–4564 25. Turner JA (2004) Sustainable hydrogen production. Science 305(5686):972–974 26. Walter MG, Warren EL, McKone JR, Boettcher SW, Mi Q, Santori EA, Lewis NS (2010) Solar water splitting cells. Chem Rev 110(11):6446–6473 27. Rubin E, De Coninck H (2004) IPCC special report on carbon dioxide capture and storage. Cambridge University Press, Cambridge; TNO: Cost Curves for CO2 Storage, Part 2005, 2, 14 28. Gibbins J, Chalmers H (2008) Carbon capture and storage. Energy Policy 36(12):4317–4322 29. Habisreutinger SN, Schmidt-Mende L, Stolarczyk JK (2013) Photocatalytic reduction of CO2 on TiO2 and other semiconductors. Angew Chem Int Ed 52(29):7372–7408 30. Vatansever F, de Melo WCMA, Avci P, Vecchio D, Sadasivam M, Gupta A, Chandran R, Karimi M, Parizotto NA, Yin R, Tegos GP, Hamblin MR (2013) Antimicrobial strategies centered around reactive oxygen species - bactericidal antibiotics, photodynamic therapy, and beyond. FEMS Microbiol Rev 37(6):955–989 31. Kuhn KP, Chaberny IF, Massholder K, Stickler M, Benz VW, Sonntag HG, Erdinger L (2003) Disinfection of surfaces by photocatalytic oxidation with titanium dioxide and UVA light. Chemosphere 53(1):71–77 32. Zhu Y, Yao W, Zong R (2014) Photocatalysis: application on environmental purification and Green energy. Chemical Industry Press 33. Chen X, Zhang Z, Chi L, Nair AK, Shangguan W, Jiang Z (2016) Recent advances in visiblelight-driven Photoelectrochemical water splitting: catalyst nanostructures and reaction systems. Nano Micro Lett 8(1):1–12 34. Smith, R. A., Semiconductors. 1978 35. Grundmann M (2010) Physics of semiconductors, vol 11. Springer 36. Liu E, Zhu B, Luo J (2014) The physics of semiconductors, 7th edn. Electronic Industry Press 37. Steier L, Herraiz-Cardona I, Gimenez S, Fabregat-Santiago F, Bisquert J, Tilley SD, Graetzel M (2014) Understanding the role of underlayers and overlayers in thin film hematite photoanodes. Adv Funct Mater 24(48):7681–7688 38. Klahr B, Gimenez S, Fabregat-Santiago F, Hamann T, Bisquert J (2012) Water oxidation at hematite Photoelectrodes: the role of surface states. J Am Chem Soc 134(9):4294–4302 39. Kato H, Kudo A (2001) Water splitting into H2 and O2 on alkali tantalate photocatalysts ATaO3 (A¼ Li, Na, and K). J Phys Chem B 105(19):4285–4292 40. Kato H, Asakura K, Kudo A (2003) Highly efficient water splitting into H2 and O2 over lanthanum-doped NaTaO3 photocatalysts with high crystallinity and surface nanostructure. J Am Chem Soc 125(10):3082–3089 41. Ng YH, Iwase A, Kudo A, Amal R (2010) Reducing graphene oxide on a visible-light BiVO4 photocatalyst for an enhanced photoelectrochemical water splitting. J Phys Chem Lett 1 (17):2607–2612 42. Malathi A, Madhavan J, Ashokkumar M, Arunachalam P (2018) A review on BiVO4 photocatalyst: activity enhancement methods for solar photocatalytic applications. Appl Catal A Gen 555:47–74 43. Konta R, Ishii T, Kato H, Kudo A (2004) Photocatalytic activities of noble metal ion-doped SrTiO3 under visible light irradiation. J Phys Chem B 108(26):8992–8995 44. Kato H, Kudo A (2002) Visible-light-response, and photocatalytic activities of TiO2 and SrTiO3 photocatalysts codoped with antimony and chromium. J Phys Chem B 106 (19):5029–5034

58

Q. Huang et al.

45. Yu K, Zhang C, Chang Y, Feng Y, Yang Z, Yang T, Lou L-L, Liu S (2017) Novel threedimensionally ordered macroporous SrTiO3 photocatalysts with remarkably enhanced hydrogen production performance. Appl Catal B Environ 200:514–520 46. Tang J, Zou Z, Ye J (2004) Photocatalytic decomposition of organic contaminants by Bi 2 WO 6 under visible light irradiation. Catal Lett 92(1–2):53–56 47. Zhang L, Man Y, Zhu Y (2011) Effects of Mo replacement on the structure and visible-lightinduced photocatalytic performances of Bi2WO6 photocatalyst. ACS Catal 1(8):841–848 48. Zhang K-L, Liu C-M, Huang F-Q, Zheng C, Wang W-D (2006) Study of the electronic structure and photocatalytic activity of the BiOCl photocatalyst. Appl Catal B Environ 68 (3–4):125–129 49. Yu L, Zhang X, Li G, Cao Y, Shao Y, Li D (2016) Highly efficient Bi2O2CO3/BiOCl photocatalyst based on heterojunction with enhanced dye-sensitization under visible light. Appl Catal B Environ 187:301–309 50. Zou Z, Ye J, Arakawa H (2001) Substitution effects of In3+ by Al3+ and Ga3+ on the photocatalytic and structural properties of the Bi2InNbO7 photocatalyst. Chem Mater 13 (5):1765–1769 51. Annamalai A, Shinde PS, Jeon TH, Lee HH, Kim HG, Choi W, Jang JS (2016) Fabrication of superior alpha-Fe2O3 nanorod photoanodes through ex-situ Sn-doping for solar water splitting. Sol Energy Mat Sol C 144:247–255 52. Lee K, Lee NH, Shin SH, Lee HG, Kim SJ (2006) Hydrothermal synthesis and photocatalytic characterizations of transition metals doped nano TiO2 sols. Mater Sci Eng B Adv 129 (1–3):109–115 53. Ling Y, Wang G, Wheeler DA, Zhang JZ, Li Y (2011) Sn-doped hematite nanostructures for Photoelectrochemical water splitting. Nano Lett 11(5):2119–2125 54. Annamalai A, Shinde PS, Subramanian A, Kim JY, Kim JH, Choi SH, Lee JS, Jang JS (2015) Bifunctional TiO2 underlayer for alpha-Fe2O3 nanorod based photoelectrochemical cells: enhanced interface and Ti4+ doping. J Mater Chem A 3(9):5007–5013 55. Annamalai A, Lee HH, Choi SH, Lee SY, Gracia-Espino E, Subramanian A, Park J, Kong K-J, Jang JS (2016) Sn/be sequentially co-doped hematite Photoanodes for enhanced Photoelectrochemical water oxidation: effect of Be2+ as co-dopant. Sci Rep 6:23183 56. Zhou X, Haeublein V, Liu N, Nhat Truong N, Zolnhofer EM, Tsuchiya H, Killian MS, Meyer K, Frey L, Schmuki P (2016) TiO2 nanotubes: nitrogen-ion implantation at Low dose provides Noble-metal-free photocatalytic H-2-evolution activity. Angew Chem Int Ed 55(11):3763–3767 57. Ghicov A, Macak JM, Tsuchiya H, Kunze J, Haeublein V, Frey L, Schmuki P (2006) Ion implantation and annealing for an efficient N-doping of TiO2 nanotubes. Nano Lett 6 (5):1080–1082 58. Asahi R, Morikawa T, Ohwaki T, Aoki K, Taga Y (2001) Visible-light photocatalysis in nitrogen-doped titanium oxides. Science 293(5528):269–271 59. Park JH, Kim S, Bard AJ (2006) Novel carbon-doped TiO2 nanotube arrays with high aspect ratios for efficient solar water splitting. Nano Lett 6(1):24–28 60. Ohno T, Akiyoshi M, Umebayashi T, Asai K, Mitsui T, Matsumura M (2004) Preparation of S-doped TiO2 photocatalysts and their photocatalytic activities under visible light. Appl Catal A Gen 265(1):115–121 61. Lv Y, Ding Y, Zhou J, Xiao W, Feng Y (2009) Preparation, characterization, and photocatalytic activity of N, S-Codoped TiO2Nanoparticles. J Am Chem Soc 92(4):938–941 62. Choi W, Termin A, Hoffmann MR (1994) The role of metal ion dopants in quantum-sized TiO2: correlation between photoreactivity and charge carrier recombination dynamics. J Phys Chem C 98(51):13669–13679 63. Kleiman-Shwarsctein A, Huda MN, Walsh A, Yan Y, Stucky GD, Hu Y-S, Al-Jassim MM, McFarland EW (2010) Electrodeposited aluminum-doped alpha-Fe2O3 Photoelectrodes: experiment and theory. Chem Mater 22(2):510–517

2 Applications and Fundamentals of Photocatalysis with Solar Energy

59

64. Kleiman-Shwarsctein A, Hu Y-S, Forman AJ, Stucky GD, McFarland EW (2008) Electrodeposition of alpha-Fe2O3 doped with Mo or Cr as photoanodes for photocatalytic water splitting. Adv Funct Mater 112(40):15900–15907 65. Cesar I, Sivula K, Kay A, Zboril R, Graetzel M (2009) Influence of feature size, film thickness, and silicon doping on the performance of nanostructured hematite Photoanodes for solar water splitting. J Phys Chem C 113(2):772–782 66. Luo ZB, Li CC, Liu SS, Wang T, Gong JL (2017) Gradient doping of phosphorus in Fe2O3 nanoarray photoanodes for enhanced charge separation. Chem Sci 8(1):91–100 67. Li M, Yang Y, Ling Y, Qiu W, Wang F, Liu T, Song Y, Liu X, Fang P, Tong Y, Li Y (2017) Morphology and doping engineering of Sn-doped hematite nanowire Photoanodes. Nano Lett 17(4):2490–2495 68. Cesar I, Kay A, Martinez JAG, Gräetzel M (2006) Translucent thin-film Fe2O3 photoanodes for efficient water splitting by sunlight: nanostructure-directing effect of Si-doping. J Am Chem Soc 128(14):4582–4583 69. Wang Y, Cai J, Wu M, Chen J, Zhao W, Tian Y, Ding T, Zhang J, Jiang Z, Li X (2018) Rational construction of oxygen vacancies onto tungsten trioxide to improve visible light photocatalytic water oxidation reaction. Appl Catal B Environ 239:398–407 70. Pu A, Deng J, Li M, Gao J, Zhang H, Hao Y, Zhong J, Sun X (2014) Coupling Ti-doping and oxygen vacancies in hematite nanostructures for solar water oxidation with high efficiency. J Mater Chem A 2(8):2491–2497 71. Huang H, Hou X, Xiao J, Zhao L, Huang Q, Chen H, Li Y (2018) Effect of annealing atmosphere on the performance of TiO 2 nanorod arrays in photoelectrochemical water splitting. Catal Today 330:189–194 72. Yang T-Y, Kang H-Y, Sim U, Lee Y-J, Lee J-H, Koo B, Nam KT, Joo Y-C (2013) A new hematite photoanode doping strategy for solar water splitting: oxygen vacancy generation. Phys Chem Chem Phys 15(6):2117–2124 73. Xiao J, Huang H, Huang Q, Zhao L, Li X, Hou X, Chen H, Li Y (2017) Suppressing the electron-hole recombination rate in hematite photoanode with a rapid cooling treatment. J Catal 350:48–55 74. Maeda K, Takata T, Hara M, Saito N, Inoue Y, Kobayashi H, Domen K (2005) GaN: ZnO solid solution as a photocatalyst for visible-light-driven overall water splitting. J Am Chem Soc 127(23):8286–8287 75. Ohno T, Bai L, Hisatomi T, Maeda K, Domen K (2012) Photocatalytic water splitting using modified GaN: ZnO solid solution under visible light: long-time operation and regeneration of activity. J Am Chem Soc 134(19):8254–8259 76. Yan S, Wang J, Gao H, Wang N, Yu H, Li Z, Zhou Y, Zou Z (2013) Zinc gallogermanate solid solution: a novel photocatalyst for efficiently converting CO2 into solar fuels. Adv Funct Mater 23(14):1839–1845 77. Wang D, Kako T, Ye J (2008) Efficient photocatalytic decomposition of acetaldehyde over a solid-solution perovskite (Ag0. 75Sr0. 25)(Nb0. 75Ti0. 25) O3 under visible-light irradiation. J Am Chem Soc 130(9):2724–2725 78. Wang D, Kako T, Ye J (2009) New series of solid-solution semiconductors (AgNbO3) 1 x (SrTiO3) x with modulated band structure and enhanced visible-light photocatalytic activity. J Phys Chem C 113(9):3785–3792 79. Ikeda S, Aono N, Iwase A, Kobayashi H, Kudo A (2019) Cu3MS4 (M¼ V, Nb, Ta) and its solid solutions with sulvanite structure for photocatalytic and Photoelectrochemical H2 evolution under visible-light irradiation. ChemSusChem 12(9):1977–1983 80. Li Q, Meng H, Zhou P, Zheng Y, Wang J, Yu J, Gong J (2013) Zn1–x Cd x S solid solutions with controlled bandgap and enhanced visible-light photocatalytic H2-production activity. ACS Catal 3(5):882–889 81. Zhang L, Jiang D, Irfan RM, Tang S, Chen X, Du P (2019) Highly efficient and selective photocatalytic dehydrogenation of benzyl alcohol for simultaneous hydrogen and benzaldehyde production over Ni-decorated Zn0. 5Cd0. 5S solid solution. J Energy Chem 30:71–77

60

Q. Huang et al.

82. Dan M, Prakash A, Cai Q, Xiang J, Ye Y, Li Y, Yu S, Lin Y, Zhou Y (2019) Energy-bandcontrolling strategy to construct novel CdxIn1-xS solid solution for durable visible light photocatalytic hydrogen sulfide splitting. Sol RRL 3(1):1800237 83. Li Y, Jiang S, Xiao J, Li Y (2014) Photocatalytic overall water splitting under visible light over an In–Ni–Ta–O–N solid solution without an additional cocatalyst. Int J Hydrog Energy 39 (2):731–735 84. Shenoy J, Hart JN, Grau-Crespo R, Allan NL, Cazorla C (2019) Mixing thermodynamics and photocatalytic properties of GaP–ZnS solid solutions. Adv Theory Simul 2(3):1800146 85. Hou X, Jiang S, Li Y, Xiao J, Li Y (2015) Ni-doped InN/GaZnON composite catalyst for overall water splitting under visible light irradiation. Int J Hydrog Energy 40 (45):15448–15453 86. Xu QC, Wellia DV, Ng YH, Amal R, Tan TTY (2011) Synthesis of porous and visible-light absorbing Bi2WO6/TiO2 heterojunction films with improved photoelectrochemical and photocatalytic performances. J Phys Chem C 115(15):7419–7428 87. Papp J, Soled S, Dwight K, Wold A (1994) Surface acidity and photocatalytic activity of TiO2, WO3/TiO2, and MoO3/TiO2 photocatalysts. Chem Mater 6(4):496–500 88. Kwon YT, Song KY, Lee WI, Choi GJ, Do YR (2000) Photocatalytic behavior of WO3-loaded TiO2 in an oxidation reaction. J Catal 191(1):192–199 89. Palanisamy B, Babu C, Sundaravel B, Anandan S, Murugesan V (2013) Sol-gel synthesis of mesoporous mixed Fe2O3/TiO2 photocatalyst: application for degradation of 4-chlorophenol. J Hazard Mater 252:233–242 90. Zhang X, Lei L (2008) Preparation of photocatalytic Fe2O3–TiO2 coatings in one step by metal-organic chemical vapor deposition. Appl Surf Sci 254(8):2406–2412 91. Liu J, Yang S, Wu W, Tian Q, Cui S, Dai Z, Ren F, Xiao X, Jiang C (2015) 3D flowerlike α-Fe2O3@ TiO2 core-shell nanostructures: general synthesis and enhanced photocatalytic performance. ACS Sustain Chem Eng 3(11):2975–2984 92. Li G-S, Zhang D-Q, Yu JC (2009) A new visible-light photocatalyst: CdS quantum dots embedded mesoporous TiO2. Environ Sci Technol 43(18):7079–7085 93. Yu Z, Meng J, Xiao J, Li Y, Li Y (2014) Cobalt sulfide quantum dots modified TiO2 nanoparticles for efficient photocatalytic hydrogen evolution. Int J Hydrog Energy 39 (28):15387–15393 94. He H, Lin J, Fu W, Wang X, Wang H, Zeng Q, Gu Q, Li Y, Yan C, Tay BK (2016) MoS2/TiO2 edge-on Heterostructure for efficient photocatalytic hydrogen evolution. Adv Energy Mater 6 (14):1600464 95. Dai G, Yu J, Liu G (2011) Synthesis and enhanced visible-light Photoelectrocatalytic activity of pn junction BiOI/TiO2 nanotube arrays. J Phys Chem C 115(15):7339–7346 96. Zhou W, Yin Z, Du Y, Huang X, Zeng Z, Fan Z, Liu H, Wang J, Zhang H (2013) Synthesis of few-layer MoS2 Nanosheet-coated TiO2 Nanobelt Heterostructures for enhanced photocatalytic activities. Small 9(1):140–147 97. Cong Y, Li Z, Zhang Y, Wang Q, Xu Q (2012) Synthesis of alpha-Fe2O3/TiO2 nanotube arrays for photo electro-Fenton degradation of phenol. Chem Eng J 191:356–363 98. Wen J, Xie J, Chen X, Li X (2017) A review on g-C3N4-based photocatalysts. Appl Surf Sci 391:72–123 99. Cao S, Yu J (2014) G-C3N4-based photocatalysts for hydrogen generation. J Phys Chem Lett 5 (12):2101–2107 100. Ong W-J, Tan L-L, Ng YH, Yong S-T, Chai S-P (2016) Graphitic carbon nitride (g-C3N4)based photocatalysts for artificial photosynthesis and environmental remediation: are we a step closer to achieving sustainability? Chem Rev 116(12):7159–7329 101. Jo W-K, Natarajan TS (2015) Influence of TiO2 morphology on the photocatalytic efficiency of direct Z-scheme g-C3N4/TiO2 photocatalysts for isoniazid degradation. Chem Eng J 281:549–565

2 Applications and Fundamentals of Photocatalysis with Solar Energy

61

102. Yu J, Wang S, Low J, Xiao W (2013) Enhanced photocatalytic performance of direct Z-scheme g-C3N4-TiO2 photocatalysts for the decomposition of formaldehyde in the air. Phys Chem Chem Phys 15(39):16883–16890 103. Wang Y, Yang W, Chen X, Wang J, Zhu Y (2018) Photocatalytic activity enhancement of core-shell structure g-C3N4@TiO2 via controlled ultrathin g-C3N4 layer. Appl Catal B Environ 220:337–347 104. Sun J-X, Yuan Y-P, Qiu L-G, Jiang X, Xie A-J, Shen Y-H, Zhu J-F (2012) Fabrication of composite photocatalyst g-C3N4-ZnO and enhancement of photocatalytic activity under visible light. Dalton Trans 41(22):6756–6763 105. Liu W, Wang M, Xu C, Chen S (2012) Facile synthesis of g-C3N4/ZnO composite with enhanced visible light photooxidation and photoreduction properties. Chem Eng J 209:386–393 106. Wang Y, Shi R, Lin J, Zhu Y (2011) Enhancement of photocurrent, and photocatalytic activity of ZnO hybridized with graphite-like C 3 N 4. Energy Environ Sci 4(8):2922–2929 107. Xiao J, Zhang X, Li Y (2015) A ternary g-C3N4/Pt/ZnO photoanode for efficient photoelectrochemical water splitting. Int J Hydrog Energy 40(30):9080–9087 108. Zhang J, Wang Y, Jin J, Zhang J, Lin Z, Huang F, Yu J (2013) Efficient visible-light photocatalytic hydrogen evolution and enhanced photostability of core/shell CdS/g-C3N4 nanowires. ACS Appl Mater Interfaces 5(20):10317–10324 109. Ge L, Zuo F, Liu J, Ma Q, Wang C, Sun D, Bartels L, Feng P (2012) Synthesis and efficient visible-light photocatalytic hydrogen evolution of polymeric g-C3N4 coupled with CdS quantum dots. J Phys Chem C 116(25):13708–13714 110. She X, Wu J, Xu H, Zhong J, Wang Y, Song Y, Nie K, Liu Y, Yang Y, Rodrigues MTF (2017) High-efficiency photocatalytic water splitting using 2D α-Fe2O3/g-C3N4 Z-scheme catalysts. Adv Energy Mater 7(17):1700025 111. Jiang Z, Wan W, Li H, Yuan S, Zhao H, Wong PK (2018) A hierarchical Z-scheme α-Fe2O3/gC3N4 hybrid for enhanced photocatalytic CO2 reduction. Adv Mater 30(10):1706108 112. Liu Y, Yu Y-X, Zhang W-D (2014) Photoelectrochemical study on charge transfer properties of nanostructured Fe2O3 modified by g-C3N4. Int J Hydrog Energy 39(17):9105–9113 113. Hirakawa T, Kamat PV (2005) Charge separation and catalytic activity of Ag@ TiO2 coreshell composite clusters under UV-irradiation. J Am Chem Soc 127(11):3928–3934 114. Jang JS, Yoon KY, Xiao X, Fan F-RF, Bard AJ (2009) Development of a potential Fe2O3based photocatalyst thin film for water oxidation by scanning electrochemical microscopy: effects of Ag Fe2O3 nanocomposite and Sn doping. Chem Mater 21(20):4803–4810 115. Chen Y, Huang W, He D, Yue S, Huang H (2014) Construction of Heterostructured g-C3N4/ Ag/TiO2 microspheres with enhanced Photocatalysis performance under visible-light irradiation. ACS Appl Mater Interfaces 6(16):14405–14414 116. Li H, Bian Z, Zhu J, Huo Y, Li H, Lu Y (2007) Mesoporous Au/TiO2 nanocomposites with enhanced photocatalytic activity. J Am Chem Soc 129(15):4538–4539 117. Cao S-W, Yin Z, Barber J, Boey FY, Loo SCJ, Xue C (2011) Preparation of Au-BiVO4 heterogeneous nanostructures as highly efficient visible-light photocatalysts. ACS Appl Mater Interfaces 4(1):418–423 118. Li P, Wei Z, Wu T, Peng Q, Li Y (2011) Au-ZnO hybrid nanopyramids and their photocatalytic properties. J Am Chem Soc 133(15):5660–5663 119. Wu F, Hu X, Fan J, Liu E, Sun T, Kang L, Hou W, Zhu C, Liu H (2013) Photocatalytic activity of Ag/TiO2 nanotube arrays enhanced by surface plasmon resonance and application in hydrogen evolution by water splitting. Plasmonics 8(2):501–508 120. Sreethawong T, Yoshikawa S (2005) Comparative investigation on photocatalytic hydrogen evolution over Cu-, Pd-, and Au-loaded mesoporous TiO2 photocatalysts. Catal Commun 6 (10):661–668 121. Melvin AA, Illath K, Das T, Raja T, Bhattacharyya S, Gopinath CS (2015) M-Au/TiO2 (M ¼ Ag, Pd, and Pt) nano photocatalyst for overall solar water splitting: role of interfaces. Nanoscale 7(32):13477–13488

62

Q. Huang et al.

122. Hu Y-S, Kleiman-Shwarsctein A, Forman AJ, Hazen D, Park J-N, McFarland EW (2008) Pt-doped α-Fe2O3 thin films active for photoelectrochemical water splitting. Chem Mater 20 (12):3803–3805 123. Lian Z, Wang W, Li G, Tian F, Schanze KS, Li H (2017) Pt-enhanced mesoporous Ti(3+)/ TiO2 with rapid bulk to surface electron transfer for photocatalytic hydrogen evolution. ACS Appl Mater Interfaces 9(20):16959–16966 124. Tong T, Zhu B, Jiang C, Cheng B, Yu J (2018) Mechanistic insight into the enhanced photocatalytic activity of single-atom Pt, Pd or Au-embedded g-C 3 N 4. Appl Surf Sci 433:1175–1183 125. Loget G, Fabre B, Fryars S, Mériadec C, Ababou-Girard S (2017) Dispersed Ni nanoparticles stabilize silicon photoanodes for efficient and inexpensive sunlight-assisted water oxidation. ACS Energy Lett 2(3):569–573 126. Ueda K, Minegishi T, Clune J, Nakabayashi M, Hisatomi T, Nishiyama H, Katayama M, Shibata N, Kubota J, Yamada T (2015) Photoelectrochemical oxidation of water using BaTaO2N photoanodes prepared by particle transfer method. J Am Chem Soc 137 (6):2227–2230 127. Bui D-N, Mu J, Wang L, Kang S-Z, Li X (2013) Preparation of Cu-loaded SrTiO3 nanoparticles and their photocatalytic activity for hydrogen evolution from methanol aqueous solution. Appl Surf Sci 274:328–333 128. Chen S, Chen X, Jiang Q, Yuan J, Lin C, Shangguan W (2014) Promotion effect of nickel loaded on CdS for photocatalytic H2 production in lactic acid solution. Appl Surf Sci 316:590–594 129. Yu J, Ma T, Liu S (2011) Enhanced photocatalytic activity of mesoporous TiO2 aggregates by embedding carbon nanotubes as an electron-transfer channel. PCCP 13(8):3491–3501 130. Li Q, Guo B, Yu J, Ran J, Zhang B, Yan H, Gong JR (2011) Highly efficient visible-lightdriven photocatalytic hydrogen production of CdS-cluster-decorated graphene nanosheets. J Am Chem Soc 133(28):10878–10884 131. Yu H, Zhao Y, Zhou C, Shang L, Peng Y, Cao Y, Wu L-Z, Tung C-H, Zhang T (2014) Carbon quantum dots/TiO2 composites for efficient photocatalytic hydrogen evolution. J Mater Chem A 2(10):3344–3351 132. Yuan Y-J, Tu J-R, Ye Z-J, Chen D-Q, Hu B, Huang Y-W, Chen T-T, Cao D-P, Yu Z-T, Zou Z-G (2016) MoS 2 -graphene/ZnIn 2 S 4 hierarchical microarchitectures with an electron transport bridge between light-harvesting semiconductor and cocatalyst: a highly efficient photocatalyst for solar hydrogen generation. Appl Catal B Environ 188:13–22 133. Zhang X, Li Y, Zhao J, Wang S, Li Y, Dai H, Sun X (2014) Advanced three-component ZnO/Ag/CdS nanocomposite photoanode for photocatalytic water splitting. J Power Sources 269:466–472 134. Li Y, Zhang X, Jiang S, Dai H, Sun X, Li Y (2015) Improved photoelectrochemical property of a nanocomposite NiO/CdS@ZnO photoanode for water splitting. Sol Energy Mater Sol Cells 132:40–46 135. Abdi FF, Han L, Smets AHM, Zeman M, Dam B, van de Krol R (2013) Efficient solar water splitting by enhanced charge separation in a bismuth vanadate-silicon tandem photoelectrode. Nat Commun 4:2195 136. Xiao J, Hou X, Zhao L, Li Y (2016) A conductive ZnO: Ga/ZnO core-shell nanorod photoanode for photoelectrochemical water splitting. Int J Hydrog Energy 41 (33):14596–14604 137. Kment S, Riboni F, Pausova S, Wang L, Wang LY, Han H, Hubicka Z, Krysa J, Schmuki P, Zboril R (2017) Photoanodes based on TiO2 and alpha-Fe2O3 for solar water splitting – the superior role of 1D nano architectures and combined heterostructures. Chem Soc Rev 46 (12):3716–3769 138. Shang S, Jiao X, Chen D (2012) Template-free fabrication of TiO2 hollow spheres and their photocatalytic properties. ACS Appl Mater Interfaces 4(2):860–865

2 Applications and Fundamentals of Photocatalysis with Solar Energy

63

139. Wang W, Li G, Xia D, An T, Zhao H, Wong PK (2017) Photocatalytic nanomaterials for solardriven bacterial inactivation: recent progress and challenges. Environ Sci Nano 4(4):782–799 140. Sivula K, Le Formal F, Graetzel M (2011) Solar water splitting: progress using hematite (alpha-Fe2O3) Photoelectrodes. ChemSusChem 4(4):432–449 141. Kisumi T, Tsujiko A, Murakoshi K, Nakato Y (2003) Crystal-face and illumination intensity dependences of the quantum efficiency of photoelectrochemical etching, in relation to those of water photooxidation, at n-TiO2 (rutile) semiconductor electrodes. J Electroanal Chem 545:99–107 142. Su J, Zhu L, Chen G (2016) Ultrasmall graphitic carbon nitride quantum dots decorated selforganized TiO2 nanotube arrays with highly efficient photoelectrochemical activity. Appl Catal B Environ 186:127–135 143. Guo M, Yang C, Zhang M, Zhang Y, Ma T, Wang X, Wang X (2008) Effects of preparing conditions on the electrodeposition of well-aligned ZnO nanorod arrays. Electrochim Acta 53 (14):4633–4641 144. Liu B, Aydil ES (2009) Growth of oriented single-crystalline rutile TiO2 Nanorods on transparent conducting substrates for dye-sensitized solar cells. J Am Chem Soc 131 (11):3985–3990 145. Shen S, Guo P, Wheeler DA, Jiang J, Lindley SA, Kronawitter CX, Zhang JZ, Guo L, Mao SS (2013) Physical and photoelectrochemical properties of Zr-doped hematite nanorod arrays. Nanoscale 5(20):9867–9874 146. Ling Y, Wang G, Reddy J, Wang C, Zhang JZ, Li Y (2012) The influence of oxygen content on the thermal activation of hematite nanowires. Angew Chem Int Ed 51(17):4074–4079 147. Wu JJ, Liu SC (2002) Low-temperature growth of well-aligned ZnO nanorods by chemical vapor deposition. Adv Mater 14(3):215–218 148. Gong D, Grimes CA, Varghese OK, Hu W, Singh R, Chen Z, Dickey EC (2001) Titanium oxide nanotube arrays prepared by anodic oxidation. J Mater Res 16(12):3331–3334 149. Li C, Li A, Luo Z, Zhang J, Chang X, Huang Z, Wang T, Gong J (2017) Surviving hightemperature calcination: ZrO2-induced hematite nanotubes for Photoelectrochemical water oxidation. Angew Chem Int Ed 56(15):4150–4155 150. Shao P, Ren Z, Tian J, Gao S, Luo X, Shi W, Yan B, Li J, Cui F (2017) Silica hydrogelmediated dissolution-recrystallization strategy for the synthesis of ultrathin alpha-Fe2O3 nanosheets with highly exposed (110) facets: a superior photocatalyst for degradation of bisphenol. Chem Eng J 323:64–73 151. Wen XG, Wang SH, Ding Y, Wang ZL, Yang SH (2005) Controlled growth of large-area, uniform, vertically aligned arrays of alpha-Fe2O3 nanobelts and nanowires. J Phys Chem B 109(1):215–220 152. Liu Y, Ma Y, Liu W, Shang Y, Zhu A, Tan P, Xiong X, Pan J (2018) Facet, and morphology dependent photocatalytic hydrogen evolution with CdS nanoflowers using a novel mixed solvothermal strategy. J Colloid Interface Sci 513:222–230 153. Wang Y, Li X, Wang N, Quan X, Chen Y (2008) Controllable synthesis of ZnO nanoflowers and their morphology-dependent photocatalytic activities. Sep Purif Technol 62(3):727–732 154. Kay A, Cesar I, Graetzel M (2006) New benchmark for water photooxidation by nanostructured alpha-Fe2O3 films. J Am Chem Soc 128(49):15714–15721 155. Sivula K, Formal FL, Grätzel M (2009) WO3Fe2O3 Photoanodes for water splitting: a host scaffold, guest absorber approach. Chem Mater 21(13):2862–2867 156. Lei R, Ni H, Chen R, Zhang B, Zhan W, Li Y (2017) Hydrothermal synthesis of WO3/Fe2O3 nanosheet arrays on iron foil for photocatalytic degradation of methylene blue. J Mater Sci Mater Electron 28(14):10481–10487 157. Meng A, Zhu B, Zhong B, Zhang L, Cheng B (2017) Direct Z-scheme TiO2/CdS hierarchical photocatalyst for enhanced photocatalytic H-2-production activity. Appl Surf Sci 422:518–527 158. Baker DR, Kamat PV (2009) Photosensitization of TiO2 nanostructures with CdS quantum dots: particulate versus tubular support architectures. Adv Funct Mater 19(5):805–811

64

Q. Huang et al.

159. Xie J, Zhang L, Li M, Hao Y, Lian Y, Li Z, Wei Y (2015) Alpha-Fe2O3 modified ZnO flowerlike microstructures with enhanced photocatalytic performance for pentachlorophenol degradation. Ceram Int 41(8):9420–9425 160. Liu Y, Sun L, Wu J, Fang T, Cai R, Wei A (2015) Preparation and photocatalytic activity of ZnO/Fe2O3 nanotube composites. Mater Sci Eng B Adv 194:9–13 161. Li Y, Wei X, Zhu B, Wang H, Tang Y, Sum TC, Chen X (2016) Hierarchically branched Fe2O3@TiO2 nanorod arrays for photoelectrochemical water splitting: facile synthesis and enhanced photoelectrochemical performance. Nanoscale 8(21):11284–11290 162. Liu J, Yang S, Wu W, Tian Q, Cui S, Dai Z, Ren F, Xiao X, Jiang C (2015) 3D flowerlike α-Fe2O3@TiO2 Core-Shell nanostructures: general synthesis and enhanced photocatalytic performance. ACS Sustain Chem Eng 3(11):2975–2984 163. Chai X, Zhang H, Pan Q, Bian J, Chen Z, Cheng C (2019) 3D ordered urchin-like TiO2@Fe2O3 arrays photoanode for efficient photoelectrochemical water splitting. Appl Surf Sci 470:668–676 164. Han H, Riboni F, Karlicky F, Kment S, Goswami A, Sudhagar P, Yoo J, Wang L, Tomanec O, Petr M, Haderka O, Terashima C, Fujishima A, Schmuki P, Zboril R (2017) Alpha-Fe2O3/ TiO2 3D hierarchical nanostructures for enhanced photoelectrochemical water splitting. Nanoscale 9(1):134–142 165. Liu C, Tang J, Chen HM, Liu B, Yang P (2013) A fully integrated nanosystem of semiconductor nanowires for direct solar water splitting. Nano Lett 13(6):2989–2992 166. Ran J, Zhang J, Yu J, Jaroniec M, Qiao SZ (2014) Earth-abundant cocatalysts for semiconductor-based photocatalytic water splitting. Chem Soc Rev 43(22):7787–7812 167. Hu Z, Shen Z, Yu JC (2016) Covalent fixation of surface oxygen atoms on hematite photoanode for enhanced water oxidation. Chem Mater 28(2):564–572 168. Eftekharinia B, Moshaii A, Dabirian A, Vaughan NS (2017) Optimization of charge transport in a Co–Pi modified hematite thin film produced by scalable electron beam evaporation for photoelectrochemical water oxidation. J Mater Chem 5(7):3412–3424 169. Zhong DK, Choi S, Gamelin DR (2011) Near-complete suppression of surface recombination in solar photoelectrolysis by “Co-Pi” catalyst-modified W: BiVO4. J Am Chem Soc 133 (45):18370–18377 170. Tilley SD, Cornuz M, Sivula K, Graetzel M (2010) Light-induced water splitting with hematite: improved nanostructure and iridium oxide catalysis. Angew Chem Int Ed 49 (36):6405–6408 171. Hajibabaei H, Schon AR, Hamann TW (2017) Interface control of Photoelectrochemical water oxidation performance with Ni1–x Fe x O y modified hematite Photoanodes. Chem Mater 29 (16):6674–6683 172. Mao Y, Cheng Y, Wang J, Yang H, Li M, Chen J, Chao M, Tong Y, Liang E (2016) Amorphous NiO electrocatalyst overcoated ZnO nanorod photoanodes for enhanced photoelectrochemical performance. New J Chem 40(1):107–112 173. Wang G, Ling Y, Lu X, Zhai T, Qian F, Tong Y, Li Y (2013) A mechanistic study into the catalytic effect of Ni(OH)2 on hematite for photoelectrochemical water oxidation. Nanoscale 5 (10):4129–4133 174. Kim JY, Youn DH, Kang K, Lee JS (2016) Highly conformal deposition of an ultrathin FeOOH layer on a hematite nanostructure for efficient solar water splitting. Angew Chem Int Ed 55(36):10854–10858 175. Tang F, Cheng W, Su H, Zhao X, Liu Q (2018) Smoothing surface trapping states in 3D corallike CoOOH-wrapped-BiVO4 for efficient Photoelectrochemical water oxidation. ACS Appl Mater Interfaces 10(7):6228–6234 176. Kanan MW, Nocera DG (2008) In situ formation of an oxygen-evolving catalyst in neutral water containing phosphate and Co2+. Science 321(5892):1072–1075 177. Abdi FF, van de Krol R (2012) Nature and light dependence of bulk recombination in Co-Picatalyzed BiVO4 photoanodes. J Phys Chem C 116(17):9398–9404

2 Applications and Fundamentals of Photocatalysis with Solar Energy

65

178. Li Y, Chen H (2016) Facile fire treatment of nanostructured hematite with enhanced photoelectrochemical water splitting performance. J Mater Chem A 4(39):14974–14977 179. Klahr B, Gimenez S, Fabregat-Santiago F, Bisquert J, Hamann TW (2012) Photoelectrochemical and impedance spectroscopic investigation of water oxidation with “Co-Pi”-coated hematite electrodes. J Am Chem Soc 134(40):16693–16700 180. Carroll GM, Gamelin DR (2016) Kinetic analysis of photoelectrochemical water oxidation by mesostructured Co-Pi/alpha-Fe2O3 photoanodes. J Mater Chem A 4(8):2986–2994 181. Carroll GM, Zhong DK, Gamelin DR (2015) Mechanistic insights into solar water oxidation by cobalt-phosphate-modified α-Fe2O3 photoanodes. Energy Environ Sci 8(2):577–584 182. Hisatomi T, Le Formal F, Cornuz M, Brillet J, Tetreault N, Sivula K, Graetzel M (2011) Cathodic shift in onset potential of solar oxygen evolution on hematite by 13-group oxide overlayers. Energy Environ Sci 4(7):2512–2515 183. Le Formal F, Tetreault N, Cornuz M, Moehl T, Graetzel M, Sivula K (2011) Passivating surface states on water splitting hematite photoanodes with alumina overlayers. Chem Sci 2 (4):737–743 184. Zhan F, Liu W, Li W, Li J, Yang Y, Li Y, Chen Q (2016) Efficient solar water oxidation by WO 3 plate arrays film decorated with CoO x electrocatalyst. Int J Hydrog Energy 41 (28):11925–11932 185. Huang J, Zhang Y, Ding Y (2017) Rationally designed/constructed CoOx/WO3 anode for efficient photoelectrochemical water oxidation. ACS Catal 7(3):1841–1845 186. Higashi M, Domen K, Abe R (2012) Highly stable water splitting on oxynitride TaON photoanode system under visible light irradiation. J Am Chem Soc 134(16):6968–6971 187. Sun K, McDowell MT, Nielander AC, Hu S, Shaner MR, Yang F, Brunschwig BS, Lewis NS (2015) Stable solar-driven water oxidation to O2(g) by Ni-oxide-coated silicon Photoanodes. J Phys Chem Lett 6(4):592–598 188. Du C, Yang X, Mayer MT, Hoyt H, Xie J, McMahon G, Bischoping G, Wang D (2013) Hematite-based water splitting with low turn-on voltages. Angew Chem Int Ed 52 (48):12692–12695 189. Malara F, Minguzzi A, Marelli M, Morandi S, Pesaro R, Dal Santo V, Naldoni A (2015) Alpha-Fe2O3/NiOOH: An effective Heterostructure for Photoelectrochemical water oxidation. ACS Catal 5(9):5292–5300 190. Wang G, Ling Y, Lu X, Zhai T, Qian F, Tong Y, Li Y (2013) A mechanistic study into the catalytic effect of Ni(OH)2 on hematite for photoelectrochemical water oxidation. Nanoscale 5 (10):4129–4133 191. Xiao J, Huang H, Huang Q, Li X, Hou X, Zhao L, Ma R, Chen H, Li Y (2017) Remarkable improvement of the turn-on characteristics of a Fe2O3 photoanode for photoelectrochemical water splitting with coating a FeCoW oxy-hydroxide gel. Appl Catal B Environ 212:89–96

Ms. Qiuyang Huang received her bachelor’s degree from Chongqing University in 2016. Currently, she is a Ph.D. candidate under the supervision of Prof. Yongdan Li at the School of Chemical Engineering and Technology at Tianjin University, China. Her Ph.D. research is focused on the photoelectrochemical water splitting into hydrogen.

66

Q. Huang et al.

Professor Dan Kong obtained her Ph.D. degree in Chemical Engineering in 2019 from the University College London, with Professor Junwang Tang. She was a Postdoctoral Researcher at Aalto University in Finland and researched the field of photocatalysis and photoelectrocatalysis. Currently, she is an Associate Professor at the China University of Mining and Technology in China. Dan’s research focuses on the application of low-cost and tunable polymeric semiconductors for solar energy conversion such as overall water splitting and CO2 conversion. Dan’s interests also include efficient microwave-assisted synthetic techniques, photoelectrocatalysis, and mechanism investigation. Professor Yongdan Li received his Ph.D. degree in 1989 from the Industrial Catalysis Program of Tianjin University, China, with Professor Liu Chang. He spent one year at the University of Twente as a visiting researcher and one and half a year in DCPRENSIC in INPL in Nancy as a post-doc. After that, he got an associate professorship at Tianjin University and after another year, he was promoted to a full professor there. He served as the Chair of the Industrial Catalysis Program and the Chairman of the Department of Catalysis Science and Technology in Tianjin until 2017. In June 2017, he was appointed as the Tenured Full Chair Professor of Industrial Chemistry at the School of Chemical Engineering, Department of Chemical & Metallurgical Engineering of Aalto University, Finland.

Chapter 3

Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells: Design, Fabrication, and Performance Dan Bai, Huhu Feng, Xingchen Yu, Chenxin Ran, and Wei Huang

3.1

Introduction to Nucleic Acid Hybrid Materials for Clean Energy

Coping with the increasingly urgent and complex issues of energy supply for industrial manufacture, transportation, communication, finance, security and biosecurity, information storage and processing, and so forth, how to provide energy sustainably has become a serious problem. For those of us who live in the present and being weaned on petroleum fuels with limited resources, it is necessary to develop materials and technologies for renewable and environmentally friendly energy. Solar power is one of the most popular choices among the sources of sustainable energy. If we can efficiently harvest and utilize the sunlight that falls on the Earth, humanity will have almost infinite energy. However, there are still problems to tackle in the design of materials and devices in solar cells. First generation of silicon-based dye-sensitized solar cells receives sunlight by

D. Bai (*) · C. Ran (*) · W. Huang Frontiers Science Center for Flexible Electronics (FSCFE), Institute of Flexible Electronics (IFE), MIIT Key Laboratory of Flexible Electronics (KLoFE), Northwestern Polytechnical University, Xi’an, Shaanxi, China Xi’an Institute of Biomedical Materials and Engineering (IBME), Xi’an Key Laboratory of Biomedical Materials and Engineering (KLBME), Northwestern Polytechnical University (NPU), Xi’an, Shaanxi, China Research and Development Institute of Northwestern Polytechnical University in Shenzhen, Xi’an, Shaanxi, China e-mail: [email protected]; [email protected] H. Feng · X. Yu Frontiers Science Center for Flexible Electronics (FSCFE), Institute of Flexible Electronics (IFE), MIIT Key Laboratory of Flexible Electronics (KLoFE), Northwestern Polytechnical University, Xi’an, Shaanxi, China © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_3

67

68

D. Bai et al.

chromophore pigments to generate electricity. Silicon solar cells are good batteries with high efficiency of power generation and environmental stability. One of the practical problems with the silicon device is its net weight. An average 4 kW power generation solar cell system builds on the rooftop of a house, for instance, a weight of about 250 kg–400 kg [1, 2]. Besides labor costs, the architectural scaffold for solar cell installation must be designed strong enough to withstand the weight of the solar cell. In the age information computing has posted Moore’s law, the current technologies based on silicon and 2D carbon materials tend to mature and saturate, and the performances of current photoelectronic devices are gradually facing their physical limits. Considering the manufacturing cost alone, even the cost of making silicon solar cells themself is becoming cheaper; without lightweight materials the installation costs would still be above the social welfare coverage. Various approaches have been put into practice for making solar cells more compatible and easier to blend into daily life. With the advancement of biotechnology and bioengineering, biomacromolecules have come to attract attention as next-generation materials for photoelectronic devices. Biomolecule-based molecular devices can be miniaturized and integrated with less energy consumption technologies, therefore attracting much attention as a next-generation intelligent material for electronic energy devices, as a choice of carbon and silicon [3, 4]. Biomaterials are abundant in nature and their production is environmentally friendly. Synthetic biomolecular systems with rational design create intelligent systems with advanced functionalities. In recent years, nucleic acids from a natural source and chemical synthesis are taken up as a biomolecular material that contributes to the development of the sustainable energy field. Nucleic acids have been long utilized as potent material in the field of bionanotechnology, for their synthetic feasibility, versatile functionality, high performance, and environmental friendliness. The unique photoelectronic properties of nucleic acids have attracted much research attention in the application areas of electronic energy transfer, information storage, and logic circuit systems beyond the scope of their biological functionalities [5–7].

3.2

Optoelectronic Properties of Nucleic Acid

3.2.1

Structure and Conformation

3.2.1.1

Molecular Basis and Double-Helix Structure

Nucleic acids are one of nature’s most abundant and renewable biomaterials. Natural DNA/RNA and artificial XNA are substances responsible for coding and transportation of information. In recent years, quantification has come to be carried out in various fields with advances in technology and increasing needs in the industry. Photovoltaic cells using nucleic acid hybrid materials and the manufacture techniques were developed [8, 9]. The development of biomaterial-based devices and

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

69

sensors using nucleic acids can achieve high performance without compromising environmental friendliness. In this regard, nucleic acid hybrid materials are considered as a competent scaffold for hosting functional optoelectronic molecules and materials in the designated platform [10, 11]. Nucleic acid materials are easy to obtain from biomass with low cost and easily decomposed without pollution. Typical nucleic acids form in B/D/Z helical structures following the condition of surrounding microenvironments (Fig. 3.1). The exhibited peculiar double-helix structure consisted of four nucleobases in pairs: adenine (A) with thymine (T) and guanine (G) with cytosine (C). The backbones joined internally by the ester bonds were made of sugar and phosphate groups. The diameter of each helix is about 2 nm and the distance between the two base pairs is 3.4 nm [11, 12]. The base pairs are linked together by strong hydrogen bonds. The macromolecule of nucleic acids exhibits a net negative charge because of the outside phosphate groups, which may compensate by sodium ions and nonlocalized counterions that move freely along the macromolecular chain surface.

Fig. 3.1 The B/D/Z type of nucleic acid helical structures with their parameters. Reprint permission acquired from Chem. Soc. Rev. (2016), the royal society of chemistry publishing [11]

70

3.2.1.2

D. Bai et al.

Higher-Order Structures

The functionalization of biopolymers involves the control of their higher-order structures. Furthermore, the transitions of the higher-order structure extend beyond single strand, and double-helix formation plays important role in advanced biomolecular functions such as signaling and material transport. Nucleic acids are considered a cationic biopolymer because of their phosphate backbones. When nucleic acids form a higher-order structure, the electrostatic interactions are strongly affected [13, 14]. On the other hand, Nucleic acids can also be used to assemble a stable complex with the copolymer electrolyte. Nucleic acid strands can favorably define the orientation of photovoltaic donor-acceptor pairs in spatial organization. The preferred inhibition of exciton recombination could be promoted with nucleic acid incorporated with ionic surfactants such as cetyltrimethylammonium chloride (CTMA) and sensitizer molecules (Figs. 3.1 and 3.2), in the various sub-nanostructure such as films, fibers, and meshes [15–18]. On the other hand, nucleic acids could be functionalized by the intercalation interaction between guest molecules with the minor and/or major grooves in weak van der Waals forces. The formation of covalent bonding between guest molecules with nucleic acids is also possible via a randomized doping process. Of course, chemical synthesis and modification of native nucleic acids could also create a copolymer-like complex with guest molecules aligned in a designated position and orientation [19–21]. Modulating the folding and assembling process is the key to achieve designated functionalities when using biopolymer materials such as nucleic acids. Besides the innate biological activities, the opposite nucleic acids with artificial materials can manipulate the higher-order structures of biopolymers such as nucleic acids. Generally, the electrolyte complexes are formed from nucleic acids and polycation compactly in aqueous solvents. Copolymer with moieties of rich hydrophilic side chain has been found to form a completely soluble complex with nucleic acids (Fig. 3.3). In such cationic codecity, electrostatic repulsion between nucleic acids and their surrounding sphere was alleviated and stabilized the double-helix and triple-stranded nucleic acids. The various approaches in organic synthetic chemistry, photochemistry, and molecular biology have been used to stabilize the higher-order structure of nucleic acid hybrids, such as the synthesis of artificial xeno nucleic acids (XNA) or the use of nucleic acid-binding ligands. In the formed electrolyte complex of nucleic acids, structural transition catalysts activate the release and revalidation of nucleic acid base pairs. Higher-order structure assembly and functional operation of nucleic acids can be adapted in ionic peptide materials as well. The macromolecules of nucleic acid strands could act as electronic donoracceptor pairs, which are electrostatically adsorbed around calcium-titanium crystals rather than forming 2D/3D heterolytic structures. Surface-immobilized nucleic acid molecules could eliminate the ion vacancy on the crystal boundary and regulate the electronic structure of the main 3D crystal, thus modulating the carrier composite channel [22, 23]. For instance, molecular docking calculations of the interactions

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

71

Fig. 3.2 The nucleic acid helical structure incorporated with conjugated cationic surfactant by ion-exchange reaction forms copolymers

between nucleic acid and chromophore compounds such as metal complexes have been proven to adopt compact conformation to bind the major groove of the helix structure [24, 25]. The chelating ligands in metal complexes were positioned at the bottom of the major groove, surrounded by the nucleotides, forming stable hydrophobic bindings. Detailed analysis showed that hydrogen bond lengths between the nucleotide with the metal complexes of the helical nucleic acid were 2.5–2.8 Å [26] (Fig. 3.4). This case proved that a wide range of polyaromatic donor-acceptor molecules including metal complexes could interact with nucleic acids. The mechanism proposed was that the helical structure of nucleic acid could affect the exciton

Fig. 3.3 Nucleic acid hybrid materials in the arrangement of spatial organization incorporated with photovoltaic donor-accepter molecules

Fig. 3.4 Detailed and total view of iridium (III) metal complex interacting with nucleic acid double-helix strands. Reprint permission was acquired from Elsevier [26]

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

73

coupling between donor-acceptor pairs, reducing the process of charge recombination, because the helical arrangement of nucleic acids was nearby of the exciton diffusion length [26–29]. To alternate, the structure of polyion backbone and hydrophilic side chains of nucleic acids and ionic copolymers such as Poly(L-lysine)-g-dextran (PLL-g-dex) have been used to form a soluble inter-polyelectrolyte complex with nucleic acid strands to enhance their photoelectronic performance [30].

3.2.2

Optoelectronic Properties of Nucleic Acid

3.2.2.1

Conductivity, Magnetism, and Spintronics

Nucleic acid was known to be a potent material for photoelectronic devices. The high-speed charge transfers caused by charge hopping between nucleic acid base pairs are suitable for making molecular wires and semiconductors. The stacking structure and the mutual action between the base pairs in nucleic acids form the overlapping p-orbital regions. As required in photoelectronic devices, the insulator/ conductor layers were often made from cyclic compounds with double or triple bonds of p-orbital overlapping with π-conjugated areas. Therefore, electrons can be moved via the π cloud and create a tunable forbidden gap in insulators for optimized semiconductor properties [31, 32]. Modulations on the molecular array orientations in nucleic acids are feasible because of their native double-helix structure and the spontaneous higher-order spatial conformations (Fig. 3.5). As many desired functionalities in optoelectronic devices rely on carrier transportation, carrier mobility (carrier movement speed divided by the electric field strength) is a critical factor. Using nucleic acid materials with self-assembled orientation due to their chemical structural characteristics, it is easy to modulate the electronic carrier mobility of current flow and response speed for practical implementation [33, 34]. The dynamics of nucleic acids in vivo could be described by either soliton or localized modulated waves termed breather wave (isolated wave). These terahertz waves are generated due to nucleic acid folding that resonates with biomacromolecules and organelle membrane activities. In terahertz wave 1 THz (=1012 Hz) corresponds to the wavelength of 300μm. The frequency band of terahertz waves has a characteristic absorption frequency in the molecular interactions such as hydrogen bonds and van der Waals force [35–37]. Biomacromolecule systems such as nucleic acids with higher-order structures dissolved in an aqueous solution are expected to be able to obtain such information about their molecular networks. The quantities of the plasticizer incorporated in nucleic acids were also considered as an important factor in the ionic conductivity values in the strands. In an aqueous solution, the polyelectrolyte properties of nucleic acids are more pronounced because of the presence of the diffusion of ions and their counterions [38–41].

74

D. Bai et al.

Fig. 3.5 Synthetic nucleic acid as hole transport materials incorporated in electronic circuits. (Reprint permission acquired from Journal of the American Chemistry Society) [34]

Nucleic acid was firstly considered as materials without magnetism until the report on λ bacteriophage genome DNAs has shown paramagnetism at low temperature [39]. In the lyophilized dry sample of nucleic acids, magnetism property was not found. Magnetism in nucleic acids was considered to originate via the base pairs interacting with their surrounding water molecules, where the free π-electron with antiparallel or parallel spins becomes the source of spin magnetism. With both magnetic and semiconductor properties, nucleic acids are expected to be potent materials for making semiconductor devices [41–43]. Utilizing both the electronic charge and spin of nucleic acids, the technologies of electronics, semiconductor, and magnetics could be integrated into the emerging new field of spintronics and create devices with new features and improved performance. Current spintronic devices use mainly inorganic materials based on metals and silicon. By using nucleic acid materials, which are based on light elements of hydrogen and carbon, it is possible to develop a new device with little loss of deflection of spin [18, 44, 45].

3.2.2.2

Optical Properties

Production and optimization of next-generation display materials rely on a wide range of research fields, mainly from the basic optical properties of liquid crystal applications. To put a dielectric liquid crystal into practical use, it is necessary to eliminate defects in molecular orientation structures to stabilize their electric field response [46–48]. Using materials with high-speed responsiveness enables beautiful video display and significantly lowers energy consumption [49]. Apart from the ferroelectric liquid crystals and liquid crystal polymer network systems, research to explore the relationship between the liquid crystal phase and materials from living organisms was also popular in recent years [50]. The study of the mechanism between the fluidity of liquids and the anisotropic formation of crystals also plays an important role in understanding the formation of life forms. New materials with advanced techniques are required, to optimize the properties desired for practical

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

75

display applications (driving voltage, response time, contrast ratio, etc.) [46]. Nucleic acids were recently been carried under investigation as potent display material for their uniform molecular orientation and unique optical properties. There are various substances such as nucleic acids, peptides, proteins, virus particles, and other amphipathic moieties such as surfactants able to take the liquid crystal state. Liquid crystal was usually termed as a substance form having an orientation order found in a group of organic molecular components. Although inorganic component-based liquid crystals were hardly known, layered inorganic crystals such as graphite could also take information in the liquid crystal state. Therefore organic-inorganic hybrid material-based liquid crystals have also been investigated in recent years [51, 52]. As natural material from living organisms, the nucleic acid has high self-assembly, which makes it easier to modulate in nanoscale structures. Bottom-up nanotechnology applied to nucleic acid enables many functionalized materials with low energy costs. By the manner of molecular array alignment, conventional liquid crystal was classified into three phases, as their opacities increase in the order of nematic phase < smectite phase < cholesteric phase. In the nematic phase, the long axes of the molecules are parallel arranged in one direction within a column, and their ends were staggered at random intervals. In the smectite phase, rodlike molecules are arranged in planar layers, without regularity in the sequence in the layer; the molecules were also distributed in one direction (Fig. 3.6) [53]. In the cholesteric phase, the arrangement was complex, in which the molecules were aligned in layers; concerning the layer above and below, each layer rotates to a certain degree to give a spiral structure. The structure of the cholesteric phase is very much close to the natural structure of the nucleic acid helix, which makes the unique properties. Recently, blue phase liquid crystals with optically isotropic birefringent properties attracted much research attention. The structures of blue phase liquid crystals were characterized in form of double-twisted cylinders, which are arranged periodically in space. Thermotropic liquid crystal is represented by the exclusion volume effect due to anisotropic molecular shape; the riotropic liquid crystal appears due to microphase separation due to differences in intermolecular interactions. The conditions of the liquid crystal phase were classified into two types: thermotropic (temperature transition) and lyotropic (concentration transition). Thermotropic liquid crystal is intended to change the phase only by heat and pressure, such as thermoplastics. The lyotropic liquid crystal consists of a multicomponent phase change depending on the temperature and component composition, which is common in biological tissues such as cell membranes [54]. There are many exemplary liquid crystal states in living organisms. For instance, the lamellar structure is formed from components such as sphingolipids between cells of the stratum corneum of the living body. In addition to the function of preventing moisture evaporation from the body, the cholesteric phase of sphingolipids crucially modulates lateral microdomain organization for functionalities such as color expression, making the body optically active. On the other hand, the nerve pulses in the brain were modulated by the insulation effect of smectic phase phospholipids of the parent medium molecule is made, various functions have

76

D. Bai et al.

Fig. 3.6 Arrangement of a nucleic acid molecular array for display devices. Reprinted permission was acquired from Elsevier [53]

been realized. When an aqueous solution of nucleic acids was made into a certain concentration, it could effectuate optical property peculiar to the cholesterol phase [55, 56]. In the research of liquid crystal conditions and thermal stability, there are still unsolved questions. Through the utilization of nucleic acid as functional hybrid materials, methods and mechanisms could be established for liquid crystal material creation, phase structure, and property evaluations. Also, the self-assembly structure of the liquid crystal phases is very similar to the native structure of nucleic acids; the structure-activity relationship study is also interesting [57]. From the viewpoint of using natural biological self-assembly structure with the mechanisms of liquid crystal states, materials found in living organisms and functionalized synthetic materials can be used for each other, reciprocally. Nucleic acids are an ideal material to form a liquid crystal state because of their characteristic helix structure. When the base-pair molecules in nucleic acids deviate from the natural spherical symmetry, nucleic acid strands take a rodlike-shaped structure that satisfies the liquid crystal conditions. The feature helix structure was

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

77

formed by mainly the hydrophobic interactions in vivo, during which the spheres become amphiphilic [58]. Current liquid crystalline materials with long rodlike molecular arrays assemble randomly and often suffer from the disadvantageous phase separation beyond a certain concentration. On the contrary, the helical structures of nucleic acids are stable, not becoming anisotropic even in dilute solution. In the case of dsDNA, the double-helix structure is stabilized by hydrogen bonds of complementary base pairs facing each other. Using nucleic acids, it is possible to adapt their native molecular orientation to create advanced functionalities and go beyond the functions of their sugar-phosphate backbone and base-pair molecules [59].

3.3 3.3.1

Design Principles and Fabrication Techniques Design Principles

Conventional photoelectronic energy conversion devices are fabricated by depositing organic photoactive materials on inorganic substrates, and these devices require many energy costs in their synthesis, functionalization, and production. Notably, biomacromolecules, such as nucleic acids with built-in structure and functionalities in themselves, could achieve comparable or even superior optoelectronic performance with the advantages of environmentally benign and low-cost production. Therefore, the design and fabrication of the nucleic acid-based photovoltaic device are highly expected in the field. To investigate the key principles in designing nucleic acid material-based photoelectronic devices, the metrics of sensitivity, stability, reproducibility, and lifetime should be taken into consideration before putting them into practical use. Through defect heterogeneity elimination and adjustment of the energy bandwidth, the optimized nucleic acid hybrid materials can achieve higher radiation efficiency and much akin to n-type properties (Fig. 3.7) [60]. In this case, the trade-off between short-circuit current density (JSC) and fill factor (FF) does not necessarily be compromised, as long as the prepared composition of hybrid material film achieves the appropriate circuit voltage (VOC) and minimum deficit voltage difference with steady power outputs [61]. During the working of a photovoltaic cell, photonic energy is absorbed by the designated donor molecule to promote the excitation of the ground-state electrons from the highest occupied molecular orbital (HOMO) to the lowest unoccupied molecular orbital (LUMO) across the bandgap of the donor molecule. Subsequently, photoexcited electron-hole pairs are generated. In the process of energy redistribution, the excited electrons can either recombine with the hole or diffuse to the donoracceptor interface, where the separation of Coulomb-bounded electron-hole pairs into free electrons and holes is possible when the LUMO of the acceptor is below the LUMO of the donor [62, 63]. The free electrons then may transfer across the energy barrier between donor-acceptor regions toward the cathode region, while the free

78

D. Bai et al.

Fig. 3.7 Mechanistic illustration of nucleic acid hybrid material used in perovskite. Reprint permission acquired from American Chemical Society [60]

holes transport toward the anode region. During this process, the minimized energy level mismatch is required for the free charges transferring efficiently across the Schottky barriers at the interfaces. Most importantly, controlling the phase separation degree of donor-acceptor inter-network within the diffusion distance of exciton has been recognized as the most essential strategy to promote efficient charge transfer. Therefore, it is crucial to precisely modulate the donor-acceptor molecules using organized platforms, such as nucleic acids [64, 65].

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

79

In a typical nucleic acid-based photovoltaic cell, the nucleic acid layer often lies between the anode layer and the cathode layer either in direct contact or indirect contact (Fig. 3.7). Electron/hole transport or hole/electron blocking layers may also include in the device structure as buffer layers to ensure efficient charge collection at the interfaces. The main function of the nucleic acid layer is to orchestrate a plurality of donor-acceptor molecular pairs within the confined space in a certain orientation [66]. The high load of oriented donor-acceptor pairs would make high electron mobility in a preferential direction, allowing the photovoltaic cells to convert photon energy into electric energy with high efficiency. Also, the innate features of nucleic acids, such as self-assembly and self-replication in well-defined nanoscale geometry and macroscale morphologies, make them suitable for efficient optoelectronic interactions with donor-acceptor pairs through various mechanisms. The unique abilities of nucleic acid hybrid materials, therefore, have been considered as preferred platforms for their interactions with photovoltaic dye molecules, chromophores, and donor-acceptor pairs, which promote the directed orientation and prevent the aggregation of these molecules simultaneously [67, 68].

3.3.2

Fabrication Techniques

Perovskite solar cells are based on thin-film photovoltaic technology, which requires the fabrication of high-quality perovskite thin film, where the photovoltaic performance of perovskite solar cells is largely correlated to the quality of perovskite thin film. For example, the heterogeneity polycrystalline of the perovskite film may severely affect the transportation of photogenic carriers [69, 70]. Moreover, the lifetime of photo-generated excitons, the mobility of charge carriers, and the carrier diffusion lengths also determine the power conversion efficiency of perovskite solar cells. Notably, film fabrication techniques play a key role in obtaining high-quality perovskite thin film. The main techniques used for thin-film formation include spin coating, droplet coating, inkjet coating, roll-to-roll printing, and vacuum coating (Fig. 3.8). All of these techniques can be used to prepare the semiconductor film for thin film-based optoelectronic devices, and they possess their features and merits [71]. The spin coating technique is suitable for the preparation of polymer film and solution gel film because of its simple operability, simple equipment, and easy preparation in the ambient atmospheric environment [72]. Also, it only requires a relatively simple material with a certain viscosity and solubility. Moreover, the film can be simply prepared by putting the substrate on the spin coater, dropping the viscous solution in the center of the substrate, and starting the spin coater. The thickness of the prepared film can be easily adjusted by optimizing the speed, time, and acceleration of the spin coater, concentration of the solution, volume of the spincoated solution, solvent volatilization speed, and adhesion of the solution to the substrate [73].

80

D. Bai et al.

Fig. 3.8 Manufacturing technologies for solar cell devices using hybrid materials

The drop coating technique is processed by dropping a certain amount of liquid on the surface of the substrate, where the liquid diffuses to the entire surface. Then, the liquid on the substrate naturally volatilizes in the solvent atmosphere or atmospheric atmosphere, finally forming a thin film, as shown in Fig. 3.8 The method employs the principle of molecular self-assembly, in which internal molecules rearrange themselves during volatilization to form a thin film. The key point of controlling the characteristics of the film is the selection of solvent [74, 75]. Solvents with different boiling points have different volatilization rates, which will lead to the change of the morphology and the crystallinity of the prepared thin film. The advantage of this preparation method is obvious that it is easy to prepare thin films in an atmospheric environment. However, it also has disadvantages, such as poor homogeneity and high surface roughness of the deposited thin film [76]. In addition to the solution-based techniques, the physical or chemical deposition of thin film in a vacuum environment is called vacuum coating [77]. This technique is suitable to deposit semiconductor materials that show poor solubility in organic solvents. As a result, a vacuum evaporation method is an alternative option [78]. In the vacuum evaporation process, the solid semiconductor is firstly heated above its sublimation temperature to promote its transition into gas in a vacuum environment. The semiconductor molecules will uniformly despite on the substrate placed above the evaporation source. By controlling the heating temperature and deposition velocity, the order degree, thickness, and carrier transport capacity of the deposited

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

81

film can be well-controlled. The advantage of this technique is that it can prepare a uniform thin film with less impurity and high purity, but the disadvantage is that the operational procedure is complicated [79–81]. In recent years, printing and lithography, such as inkjet printing and roll-to-roll (R2R) printing, have become increasingly significant to fabricate photovoltaic and electronic devices. Inkjet coating technique is processed by ejecting a fixed amount of liquid material on the substrate from an ink storage chamber through a nozzle in the form of droplets [82]. The ejected region of the substrate forms a pattern under the combined action of gravity and air resistance, and the solidification of droplets can be realized using solvent evaporation. In this process, the viscosity and surface tension of the solution are the main factors affecting the quality of the thin film printed by inkjet printing. The advantages of inkjet printing are simple operation, high feasibility, low manufacturing cost, high operational flexibility, and mature technology, which can effectively control the deposition of materials and the formation of thin films [83]. However, its disadvantage is that the development of inkjet-based manufacturing technology is time-consuming. Compared to inkjet printing, R2R printing can also be used to efficiently produce low-cost thin film. Patterned ink on the substrate gradually forms a thin film with pressure as the roll begins to turn. This method is suitable for mass production and can reduce labor costs [84, 85]. For solution-based techniques, NAs could play a crucial role because of their tunable solubility in various solvents. As water-soluble and hydrophilic material, NAs could dissolve in water-immiscible solvents, such as acetonitrile and alcohol. Upon modification with polyethylene glycol or hydrogels, NAs have also been reported to dissolve in most organic solvents. Therefore, the introduction of NAs in solution-processed thin-film techniques could extend the solvent system from organic solvent to water, lowering the cost with environmentally friendly green chemistry. Moreover, for the packaging of advanced semiconductors, nucleic acid materials are also compatible with inkjet printing electronic technologies, which allows the device to process at ambient temperature and atmospheric pressure. As a platform in the film-forming reaction, the introduction of nucleic acid has been proven to influence the deposition process by tuning the growth rate of the thin film. As a result, the size and uniformity of the film can be fundamentally controlled. Furthermore, the presence of nucleic acid could also reduce the complex permittivity and dielectric loss of the deposited film, hence improving the electrical performance and power conversion efficiency of the corresponding electronic devices. In the manufacturing process, it is also possible to develop nucleic acid hybrid materials for optoelectronic devices without the need for a microfabrication process or extreme ultraviolet (EUV) lithography machines. For example, laser oscillation requires a resonator molecular system with a precise periodic structure to reflect circularly polarized light of a particular wavelength, and the native structure of nucleic acid themselves could serve as a distributed feedback-type laser oscillator. Also, nucleic acid-based liquid crystals could be used for the wavelength of the light that is confined spontaneously via the helical structures [86, 87]. Although lowering the energy threshold for oscillation of the liquid crystal laser is still a challenge, nucleic

82

D. Bai et al.

acids are considered a promising material for semiconductor laser devices. In practical use, nucleic acid-based cholesteric liquid crystal devices could generate electric current driven by the redistribution of the liquid crystal generated by photoelectronic excitation [88]. Also, nucleic acid materials are compatible with flexible electronic devices. Overall, the merits of NAs and their hybrid materials discussed above enable their great potential in the application of next-generation optoelectronic devices.

3.4

Enhancement of Performance

In general, organic semiconductors have difficulty obtaining materials with high purity due to the difficulty of the synthetic procedure. The performance of photoelectronic devices is easily affected by impurities [89–91]. Oxygen in the atmosphere may also deteriorate the electrical properties. The chemical and biological synthesis of nucleic acids is relatively easy and low cost, with high purity. Furthermore, through nucleic acid sequence programming, it is possible to develop nucleic acids in a three-dimensional higher-order structure accompanying new functionalities. In accord with the conductivity, magnetism, spintronics, and optical properties described in the previous section, using nucleic acid hybrid materials, the current limitation problems of photoelectronic devices and the efficiency of energy conversion could be tackled [92]. The integration of nucleic acid hybrid materials with advances in photoelectronic devices is challenging. Recent developments of perovskite solar cells (PSCs) expanded the fabrication and application possibilities of devices from the siliconbased solar cells. Perovskite solar cells can be made by applying the materials and technologies of flexible electronics. The thickness of the perovskite itself is about 500 nm (nanometer) and about 1/100th of a single hair. Since the weight of a solar cell mainly depends on the device substrate where the materials are applied to, perovskite solar cells have the potential to be produced on top of platforms that are lightweight with flexible texture, such as wrapping film or paper. Future solar cells could be made in the shape of a vinyl sheet, put on the roof like a curtain. Also, it is possible to use transparent conductive electrodes in the making of translucent solar cells [93]. The thickness of the perovskite can be tuned so that the window glass of a house or building can be made into a solar cell. The emerging print full electronic technologies allow simple and fast manufacture of solar cells anytime, anywhere. Therefore, installation costs of solar cells could be considerably reduced, easy to install to a home in populated areas. The term perovskite was originated from the mineral stone vastly distributed in the Ural Mountains of Russia. Generally, their chemical formula is represented by the composition formula of ABX3, where cation is placed in A and B and anions are placed in X (Fig. 3.9). For instance, the natural combination of perovskite is A: Ca; B: Ti; X:O. Most of the current organic-metal halide-based perovskite materials were modified with other components, which are actively used for applied research

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

83

Fig. 3.9 (Upper). Crystal structure of ABX3-type hybrid halide perovskite. The organic and inorganic monovalent cation is A (black), divalent metal cation B (red), and halogen anion X (blue); (lower) flexible perovskite solar cells. Reprint permission acquired from the American Chemistry Society [60]

on solar cells, consisting mainly of crystals from the combination of A: monovalent organic cations of CH3NH31+(MA1+), NH2CH ¼ NH21+(FA1+), Cs+, etc.; B: divalent metal cation of Pb2+, Sn2+, etc.; and X: Cl/Br/I and other halogen anions. By far perovskite is the most popular substrate for solar cell devices due to its tuneable photoelectric properties, including suitable bandgaps, light absorption coefficient, high photoluminescence quantum yield, narrowband emission, relatively small absorption length, long diffusion length, and low-cost solution processability [94, 95]. Organic metal halide perovskite solar cells were born in 2012 with the development of semiconductor materials. The conversion efficiency of the initial light energy as a photosensitized solar cell using this material as a photosensitized object was only 3.2%, which was not comparable to 26% of silicon-based systems. After a series of improvements, the PCE of perovskite solar cells (PSCs) has now exceeded 25%, and it is thought that it has the potential to exceed the records kept by silicon solar cells. Compared with silicon substrates, perovskite solar cells have received extensive attention and in-depth developments from researchers around the world in recent years [96]. When perovskite solar cells first appeared, light to energy conversion efficiency (power conversion efficiency, PCE) was low and durability was also a problem. For instance, in earlier research, the PCE was modulated by adjusting the perovskite crystal structure, blending ratio of methylammonium ion (MA+), inorganic lead ion (Pb2+), and halogen anion iodine ion (I). In the past decade, the power conversion efficiency of Pb-based perovskite solar cells has increased from 3.80% to 25.2%, and the photoelectric detection efficiency of Pb-based photodiode array detectors has increased to 1015 jones. Also, the ultrahigh

84

D. Bai et al.

photoluminescent quantum yield and high Seebeck coefficient value of approximately 1000μV/K were obtained from the Pb-based monocrystalline and Pb-based perovskite quantum dots, respectively. Despite such attractive properties in the photoelectric and thermal electrical application, the toxicity of Pb2+ in their preparation, utilization, and disposal causes many further environmental problems, which is one of the bottlenecks in its practical application. Because the Ca-Ti film can be prepared by the liquid phase method at low temperature (100  C), it is compatible with the processing of flexible electronics, especially suitable for preparation into flexible battery devices [97]. In particular, flexible devices can be produced in large areas at low cost through the roll-to-roll process. Therefore, the development of high-efficiency flexible PSCs has gradually become one of the main topics in the field of sustainable energy research. The characteristic foldable, lightweight, and robust flexible devices are very suitable for integrated design with architecture construction and logistics. The substrate of PSCs generally consists of two types, namely, flexible metal foil and flexible polymer substrate. Polyethylene terephthalate (PETE), polyethylene naphthalate (PEN), and other plastic substrates have the properties of low cost, light quality, good light transmission, and suitable mechanical performance. Yet their low melting temperature limits the carrier transport layer preparation process. The polymer substrate of flexible perovskite battery research hence has been focused on the techniques to prepare high-efficiency carrier transport layer film under low-temperature conditions. In contrast, the foil-type material has excellent ductility, conductivity, and high-temperature resistance, so the vast majority of the hightemperature preparation carrier transport layer process can be directly applied to the metal foil substrate. However, the opacity of the foil substrate requires the device’s back electrode to be transparent, which puts forward new requirements for the preparation process of the transparent electrode. Most flexible perovskite batteries now were metal foil substrate-based, which requires the development of highly efficient transparent electrodes. At the same time, the commonly used transparent conductive material indium tin oxide (In2O3-SnO2, ITO) film is very easy to crack under bending conditions, resulting in a sharp increase in resistance, hence causing severe degradation of the entire device performance Therefore, to improve the flexibility of the device, the development of the more mechanical performance of the transparent conductive electrode is an important direction for breakthrough material and technology development [98–100]. One of the unsolved issues is the material used in perovskite solar cells. Currently in piezoelectric ceramic transducer (PbZrO3, PbTiO3), lead is often used. Since Pb is toxic, it becomes a big problem when the PSCs decomposed some lead melts into the environment. However, the amount of lead per unit area used in PSCs is about the same as the amount contained in 1-cm-thick soil in nature. To be specific, the amount of lead contained in a 1 m2 area of PSCs is approximately 0.4 g. Conventional perovskite solar cells have made it easy to achieve more than 20% energy conversion efficiency by containing cesium and rubidium, but as rare earth elements, they were expensive. On the other hand, by finding a method of utilizing naturally abundant raw materials, PSCs could be made at low cost with higher performance. Also, there

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

85

is an advantage that the manufacture of bioinspired raw material is easy. Besides the flexible substrate and electrodes, the enhancement of efficiency relies heavily on the materials of the carrier transport layer. Enhancement of power conversion efficiency (PCE) has been a worldwide race for the performance of solar cells, emerging the highest record of >40% until the year 2020. Various technologies have stimulated the rising PCE in organic cells and dye-sensitized cells record of >14% up to date. Our goal is to produce perovskite solar cells with an energy conversion efficiency of more than 25%, which would contribute to the developing countries suffering from poverty by bringing solar cells to everyday life [1, 2]. The economic pressure would be alleviated once the electric power charge is reduced for industry and home. Apart from the problem of environmental stability, the current challenge in the research of solar cells is the significantly reduced hysteresis behaviors. The reason for the hysteresis occurrence in solar cells is still not fully understood. Hysteresis is commonly referred to as a historical phenomenon or historical effect, in which past events affect the status quo and represent a state that has not been undone. In the case of solar cells, during the power generation process, the current-voltage characteristics under the light irradiation should usually take the same value when the voltage is raised and lowered [101–103]. For perovskite solar cells, the hysteresis has also been reported to originate from the photocurrent density-voltage (J-V) responses. The perovskite material causing a large capacitive effect was also believed that hysteresis depends on. Ion migrations across the device and the effects of electronic defects during measurements have been ascribed to the hysteresis behavior of currentvoltage (I-V) curves in terms of mechanisms. Many cases of hysteresis also reported the different current values detected by the direction of the voltage sweep in solar cells. For electrical hysteresis, when measuring the I-V characteristics by voltage scanning on its terminals, there is a difference between the forward scanning I-V curve and the reverse scanning. The two scanned I-V curves do not coincide with each other even for a relatively low scanning rate of 100 mV/s. Hysteresis behavior was found to be correlated with the scanning rate of solar cell voltage, as the voltage scanning rate increases the hysteresis which becomes pronounced. The discovery of hysteresis motivated research on hybrid materials [103–107]. The advantage of this phenomenon is that the output value of power is stable considering the practical use of solar cells. On the other hand, from the viewpoint of basic research, it has been found that the hysteresis phenomenon could be improved by material design. Modulation of structure-activity relationship provides a great hint to think about the mechanism by which hysteresis occurs, to produce hysteresis-free PSCs with synergetic performance optimization before the commercialization [107–111]. Additionally, using nucleic acid hybrid materials has been proven to increase the photostability of the dye sensitizers and also improve the efficiency of the photovoltaic cell. Unlike conventional polymers, nucleic acids complexed with donoracceptor molecules can impose a defined spatial organization and fixed orientation [111–113]. Nucleic acid material formed with ionic surfactant or lipids with the ionic head group could form regular arrangements such as lamellar structures or parallel aligned layers that alternate through the self-assembly process. The solubility of

86

D. Bai et al.

nucleic acid materials in aqueous and various solvents improved their manufacturing processability with environmental friendliness.

3.5

Summary and Outlook

In resource-limited areas where fuel transportation or grid connection is limited by high cost, it is crucial to develop solar-generated electricity. The absence of environmental pollution during use and minimal maintenance are also key performance indicators for the technology and materials of solar power conversion devices in remote areas such as satellites, ocean vessels, mountains, and deserts. As described in this chapter, nucleic acid hybrid materials can assist the improvement of efficient photovoltaic cell performance attribute to their optoelectronic properties and their flexibility to interact with donor-acceptor molecular pairs. The advantage of manufacturing process is being solution-processable, thus reducing the cost while enhancing the durability, using nucleic acid as hybrid biomaterial with green technology. Acknowledgments This work is financially supported by the Science Technology and Innovation Commission of Shenzhen Municipality (深创-20205354JCYJ20190806153018791), Natural Science Foundation of Zhejiang Province (LGF19H200005), Japan-China Medical Association (国卫-2018920), the Frontiers Science Center for Flexible Electronics (FSCFE) and Shaanxi Institute of Biomedical Materials and Engineering (SIBME) is further acknowledged in providing space, equipment, and services. Author Contribution The research work referred to in part in the Results section was undertaken by DB, CXR (who undertook the experimental design, data analysis, and discussion; unpublished data excluded). DB, HHF, and RCX wrote the first manuscript of this book chapter and completed the editing of subsequent drafts.

References 1. Photovoltaics Report (2020) Fraunhofer Institute for Solar Energy Systems 2. IRENA (2019) Future of solar photovoltaic: deployment, investment, technology, grid integration, and socio-economic aspects (A global energy transformation). International Renewable Energy Agency, Abu Dhabi 3. Singh TB, Sariciftci NS, Grote JG (2009) Bio-organic optoelectronic devices using DNA. In: Organic electronics. Springer, pp 73–112 4. Gomez EF, Steckl AJ (2018) Engineering DNA and nucleobases for present and future device applications. In: Green materials for electronics. Wiley-VCH, pp 191–233 5. Sotzing GA, Ner YJ (2010) Nucleic acid-based photovoltaic cell. In: Google Patents 6. Wang Y, Yan X, Dong R (2014) Organic memristive devices are based on silver nanoparticles and DNA. Organ Electron 15(12):3476–3481 7. Lin Y-T, Lin T-Y, Hung Y-C (2014) Bistable memory device based on DNA biopolymer nanocomposite. Paper presented at the Organic Photonics VI

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

87

8. Pawlicka A, Firmino A, Vieira D, Sentanin F, Grote J, Kajzar F (2009) Gelatin-and DNA-based ionic conducting membranes for electrochromic devices. Paper presented at the Optical materials in defense systems technology VI; 9. Pawlicka A, Firmino A, Sentanin F, Sabadini R, Jimenez D, Jayme C, Rau I (2015) Bio-inspired materials for electrochemical devices. Paper presented at the optics and photonics for counterterrorism, crime fighting, and defence XI; and optical materials and biomaterials in security and defence systems technology XII 10. Sariciftci NS (2019) From organic electronics to bio-organic electronics. Nonlinear Opt 50 11. Wood Bayden R (2016) The importance of hydration and DNA conformation in interpreting infrared spectra of cells and tissues.[J]. Chem Soc Rev 45:1980–1998 12. Pawlik G, Mitus AC, Mysliwiec J, Miniewicz A, Grote JG (2010) Photochromic dye semiintercalation into DNA-based polymeric matrix: Computer modeling and experiment. Chem Phys Lett 484(4–6):321–323 13. Rau I, Grote JG, Kajzar F, Pawlicka A (2012) DNA–novel nanomaterial for applications in photonics and electronics. C R Phys 13(8):853–864 14. Reddeppa M, Mitta SB, Park B-G, Kim S-G, Park SH, Kim M-D (2019) DNA-CTMA functionalized GaN surfaces for NO2 gas sensor at room temperature under UV illumination. Org Electron 65:334–340 15. Reddy MSP, Park H, Lee J-H (2018) Residue-and-polymer-free graphene transfer: DNA-CTMA/graphene/GaN bio-hybrid photodiode for light-sensitive applications. Opt Mater 76:302–307 16. Rujoiu TB, Petris A, Vlad V, Rau I, Manea A-M, Kajzar F (2015) Lasing in DNA–CTMA doped with Rhodamine 610 in butanol. Phys Chem Cheml Phys 17(19):13104–13111 17. Son W-H, Reddy MSP, Choi S-Y (2014) Hydrogenated amorphous silicon thin-film solar cell with a buffer layer of DNA-CTMA biopolymer. Modern Phys Lett B 28(13):1450107 18. Nakamura K, Sagara A, Kobayashi N (2011) Photophysical properties of lanthanide (III) chelates-doped DNA-CTMA complex. Paper presented at the Nanobiosystems: processing, characterization, and applications IV 19. Zhou J, Wang ZY, Yang X, Wong C-Y, Pun EY (2010) Fabrication of low-loss, single-modechannel waveguide with DNA-CTMA biopolymer by multistep processing technology. Opt Lett 35(10):1512–1514 20. Yaney PP, Ouchen F, Grote JG (2013) Preliminary electrical and Raman characterization of DNA-CTMA films fabricated by drawn down bar. Paper presented at the Nanobiosystems: processing, characterization, and applications VI 21. Zhao T, Lin C, Zhu X, Zheng H, Li L, Huang B, Liang L (2017) Luminescence properties of Eu(TTA)3(H2O)2 doped with DNA-CTMA and its application in fluorescent inkjet ink. Springer, Singapore, pp 933–939 22. Chopade P, Dugasani SR, Jeon S, Jeong J-H, Park SH (2019) White light emission produced by CTMA-DNA nanolayers embedded with a mixture of organic light-emitting molecules. RSC Adv 9(54):31628–31635. https://doi.org/10.1039/C9RA05834F 23. Dugasani SR, Gnapareddy B, Kesama MR, Ha TH, Park SH (2018) DNA and DNA–CTMA composite thin films embedded with carboxyl group-modified multi-walled carbon nanotubes. J Indus Eng Chem 68:79–86 24. Ibisate M, Galisteo-López JF, Esteso V, López C (2013) FRET-mediated amplified spontaneous emission in DNA–CTMA complexes. Adv Opt Mater 1(9):651–656 25. Lazar CA, Kajzar F, Mihaly M, Rogozea AE, Petcu AR, Olteanu NL, Rau I (2016) Novel materials based on DNA-CTMA and lanthanide (Ce3+, Pr3+). Biopolymers 105(9):613–617 26. Bai D, Tian Y, Chen K, Zhang X, Wang F, Cheng Y, Dong X (2020) Novel 6F-iridium(III) complexes as potent theranostic agents: hypoxia probing, radiosensitization, and antiviral functionalities. Dyes Pigments 182:108635 27. Wang Z, Li P, Guo Q, Yang Z (2015) Luminescence and energy transfer of tunable emission phosphor Ca2PO4Cl: Ce3+, Mn2+. Spectrochim Acta Part A Mol Biomol Spectrosc:137, 871–876

88

D. Bai et al.

28. Liang L, Mitsumura Y, Nakamura K, Uemura S, Kamata T, Kobayashi N (2016) Temperature dependence of transfer characteristics of OTFT memory based on DNA-CTMA gate dielectric. Organ Electron 28:294–298 29. Long W, Zou W, Li X, Jiang W, Li X, Chen J (2014) Tracing photon transmission in dye-doped DNA-CTMA optical nanofibers. Opt Exp 22(6):6249–6256 30. Wahl GM, Stern M, Stark GS (1979) Efficient transfer of large DNA fragments from agarose gels to diazo benzyloxy methyl-paper and rapid hybridization by using dextran sulfate. Proc Natl Acad Sci U S A 76(8):3683–3687 31. Kelley SO, Barton JK (1999) Electron transfer between bases in double helical DNA. Science 283(5400):375–381 32. Kwon Y-W, Choi DH, Jin J-I (2012) Optical, electro-optic and optoelectronic properties of natural and chemically modified DNAs. Poly J 44(12):1191–1208 33. Nakatani K, Dohno C, Saito I (2000) Modulation of DNA-mediated hole-transport efficiency by changing superexchange electronic interaction. J Am Chem Soc 122(24):5893–5894 34. Okamoto A, Kamei T, Saito I (2006) DNA hole transport on an electrode: application to effective photoelectrochemical SNP typing. J Am Chem Soc 128(2):658–662 35. Mohiuddin M, Kumar B, Haque S (2017) Biopolymer composites in photovoltaics and photodetectors. In: Biopolymer composites in electronics. Elsevier, pp 459–486 36. Bonea A, Rău I, Ţane A, Brodeală A, Svasta P (2011) Organic capacitive structures using biopolymers. Paper presented at the 2011 IEEE 17th international symposium for design and Technology in Electronic Packaging (SIITME) 37. Firmino A, Grote J, Kajzar F, Rau I, Pawlicki A (2011) Application of DNA in electrochromic cells with switchable transmission. Nonlin Opt Quant Opt 42:181–201 38. Gomez EF (2015) Investigation of DNA Nucleobases for bio-organic light emitting diodes. University of Cincinnati 39. Apalkov V, Berashevich J, Chakraborty T (2010) Unique magnetic signatures of mismatched base pairs in DNA. J Chem Phys 132(8):85102–85102 40. Nakamae S, Cazayous M, Sacuto A, Monod P, Bouchiat H (2005) Intrinsic low-temperature paramagnetism in B-DNA. Phys Rev Lett 94(24):248102 41. Zhou B (2013) Optical properties of DNA biomaterial and application to UV-photoconductors: State University of New York at Buffalo 42. D'Souza SM, Ramees M, Sarojini B (2018) Studies on the properties of Nano metal oxides doped DNA-CTMA Matrix. Mater Today Proc 5(1):2582–2587 43. Manea A-M, Rau I, Kajzar F, Meghea A (2013) Fluorescence, spectroscopic, and NLO properties of green tea extract in deoxyribonucleic acid. Opt Mater 36(1):140–145 44. Marin CA, Kajzar F, Manea-Saghin A-M (2018) Linear and nonlinear properties of novel materials based on DNA containing [2.2] paracyclophane moiety. Opt Mater 75:547–553 45. Ouchen F, Venkat N, Johnson LE, Benight SJ, Yaney PP, Lesko A et al (2013) Effect of UV-crosslinking of DNA-CTMA biopolymer on its electrical and optical properties. Paper presented at the Nanobiosystems: processing, characterization, and applications VI 46. Reddy MSP, Puneetha P, Lee Y-W, Jeong S-H, Park C (2017) Effect of illumination and frequency-dependent series resistance and interface state densities on the electrical properties of DNA-CTMA/p-GaN bio-hybrid Schottky photodiode. Polym Test 59:107–112 47. Dancus I, Vlad V, Petris A, Rau I, Kajzar F, Meghea A, Tane A (2013) Nonlinear optical properties of Rh610 sensitized DNA-CTMA characterized by Z-scan. Paper presented at the ROMOPTO 2012: tenth conference on optics: micro-to Nanophotonics III 48. Lazar CA, Kajzar F, Rau I, Puntus L, Manea A-M (2016) Synthesis, linear and nonlinear optical properties of DNA-CTMA/europium (III) complex. Synth Met 221:120–126 49. Derkowska B, Krupka O, Smokal V, Sahraoui B (2011) Optical properties of oxazolone derivatives with and without DNA–CTMA. Opt Mater 33(9):1429–1433 50. Lazar CA, Rau I, Manea A-M (2017) Fluorescence, optical absorption and third-order nonlinear optical properties of terbium (III) complex embedded into DNA-CTMA. J Lumin 182:59–64

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

89

51. Naskar S, Saurabh S, Jang YH, Lansac Y, Maiti PK (2020) Liquid crystal ordering of nucleic acids. Soft Matter 16(3):634–641 52. Shen J, He F, Chen L, Ding L, Liu H, Wang Y, Xiong X (2017) Liquid crystal-based detection of DNA hybridization using surface-immobilized single-stranded DNA. Microchim Acta 184 (9):3137–3144 53. Liao S, Ding H, Wu Y, Wu Z, Shen G, Yu R (2016) Label-free liquid crystal biosensor for L-histidine: A DNAzyme-based platform for small molecule assay. Biosens Bioelectron 79:650–655 54. Wang L, Urbas AM, Li Q (2018) Nature-inspired emerging chiral liquid crystal nanostructures: from molecular self-assembly to DNA mesophase and nanocolloids. Adv Mater 32:1801335 55. Kim H, An Z, Jang CH (2018) Label-free optical detection of thrombin using a liquid crystalbased aptasensor. Microchem J 141:71–79 56. Khan M, Khan AR, Shin JH, Park SY (2016) A liquid-crystal-based DNA biosensor for pathogen detection. Sci Rep 6:22676 57. Verma I, Sidiq S, Pal SK (2017) Poly (l-lysine)-coated liquid crystal droplets for sensitive detection of DNA and their applications in the controlled release of drug molecules. ACS Omega 2(11):7936–7945. Kim HJ, Jang CH (2018) Imaging DNA single-strand breaks generated by reactive oxygen species using a liquid crystal-based sensor. Anal Biochem, 556:1–6 58. Nguyen DK, Jang CH (2020) Label-free liquid crystal-based biosensor for the detection of dopamine using DNA aptamer as a recognition probe. Anal Biochem:113807 59. Dierking I, Al-Zangana S (2017) Lyotropic liquid crystal phases from anisotropic nanomaterials. Nano 7(10):305 60. Hou M, Zhang H, Wang Z, Xia Y, Chen Y, Huang W (2018) Enhancing efficiency and stability of perovskite solar cells via a self-assembled dopamine interfacial layer. ACS Appl Mater Interfaces 10(36):30607–30613 61. Hou Y, Wang K, Yang D, Jiang Y, Yennawar N, Wang K, Priya S (2019) Enhanced performance and stability in DNA-perovskite heterostructure-based solar cells. ACS Energy Lett 4(11):2646–2655 62. Jayme CC, Kanicki J, Kajzar F, Nogueira AF, Pawlicka A (2016) Influence of DNA and DNA-PEDOT: PSS on dye-sensitized solar cell performance. Mol Cryst Liq Cryst 627 (1):38–48 63. Lee KW, Kim KM, Lee J, Amin R, Kim B, Park SK, Kim HJ (2011) A two-dimensional DNA lattice implanted polymer solar cell. Nanotechnology 22(37):375202 64. Elfwing A, Cai W, Ouyang L, Liu X, Xia Y, Tang Z (2018) DNA based hybrid material for interface engineering in polymer solar cells. J Appl Mater Interfaces 10(11):9579–9586 65. Bakr ZH, Wali Q, Fakharuddin A, Schmidt-Mende L, Brown TM, Jose R (2017) Advances in hole transport materials engineering for stable and efficient perovskite solar cells. Nano Energy 34:271–305 66. Yusoff AR, Kim J, Jang J, Nazeeruddin MK (2016) New Horizons for Perovskite solar cells employing DNA-CTMA as the hole-transporting material. ChemSusChem 9(13):1516–1516 67. Firmino A, Grote J, Kajzar F, M’Peko J-C, Pawlicka A (2011) DNA-based ionic conducting membranes. J Appl Phys 110(3):033704 68. Yumusak C, Singh TB, Sariciftci N, Grote J (2009) Bio-organic field-effect transistors based on crosslinked deoxyribonucleic acid (DNA) gate dielectric. Appl Phys Lett 95(26):341 69. Zeyada H, El-Nahass M, El-Shabaan M (2016) Photovoltaic properties of the 4H-pyrano [3, 2-c] quinoline derivatives and their applications in organic-inorganic photodiode fabrication. Synth Met 220:102–113 70. Paulson B, Shin I, Jeong H, Kong B, Khazaeinezhad R, Dugasani SR et al (2018) Optical dispersion control in surfactant-free DNA thin films by vitamin B 2 doping. Sci Rep 8(1):1–10 71. Meena MK, Sinhamahapatra A, Kumar A (2019) Superhydrophobic polymer composite coating on glass via spin coating technique. Colloid Polym Sci 297:1499–1505

90

D. Bai et al.

72. Melville OA, Lessard BH, Bender TP (2015) Phthalocyanine-based organic thin-film transistors: A review of recent advances. ACS Appl Mater Interfaces 7:13105–13118 73. Echelmeier A, Kim D, Cruz Villarreal J, Coe J, Quintana S, Brehm G et al (2019) 3D printed droplet generation devices for serial femtosecond crystallography enabled by surface coating. J Appl Crystallogr 52(5) 74. Hong SC, Lee G, Ha K, Yoon J, Ahn N, Cho W et al (2017) Precise morphology control and continuous fabrication of perovskite solar cells using a droplet-controllable electrospray coating system. ACS Appl Mater Interfaces 9(9):7879–7884 75. Mansour AM, Yahia IS, El Radaf IM (2018) Structural, electrical, and photovoltaic properties of PbSb2S5/n-Si heterojunction synthesized by vacuum coating technique. Mater Res Exp 5 (7):076406 76. Huang CY, Ko CC, Chen LH, Huang CT, Tung KL, Liao YC (2018) A simple coating method to prepare superhydrophobic layers on ceramic alumina for vacuum membrane distillation. Sep Purif Technol 198:79–86 77. Vasilev DD, Malevannaya EI, Moiseev KM (2017) Vacuum coating system for deposition of superconducting W x Si (1–x) ultrathin films used in single-photon detectors. J Phys Conf Ser 872(1):012027. IOP Publishing 78. Jiang D, Park SY (2016) Light-driven 3D droplet manipulation on flexible optoelectronic wetting devices fabricated by a simple spin-coating method. Lab Chip 16(10):1831–1839 79. Kelso MV, Mahenderkar NK, Chen Q, Tubbesing JZ, Switzer JA (2019) Spin coating epitaxial films. Science 364(6436):166–169 80. Zheng J, Zhang M, Lau CFJ, Deng X, Kim J, Ma Q, Ho-Baillie AW (2017) Spin-coating free fabrication for highly efficient perovskite solar cells. Sol Energy Mater Sol Cells 168:165–171 81. Liu X, Tan X, Liu Z, Ye H, Sun B, Shi T et al (2019) Boosting the efficiency of carbon-based planar CsPbBr3 perovskite solar cells by a modified multistep spin-coating technique and interface engineering. Nano Energy 56:184–195 82. Hart AS, Andersen TR, Griffith MJ, Fahy A, Vaughan B, Belcher WJ, Dastoor PC (2019) Roll-to-roll solvent annealing of printed P3HT: ICXA devices. RSC Adv 9:42294–42305 83. Homenick CM, James R, Lopinski GP, Dunford J, Sun J, Park H, Malenfant PR (2016) Fully printed and encapsulated SWCNT-based thin-film transistors via a combination of R2R gravure and inkjet printing. ACS Appl Mater Interfaces 8(41):27900–27910 84. Dou B, Whitaker JB, Bruening K, Moore DT, Wheeler LM, Ryter J et al (2018) Roll-to-roll printing of perovskite solar cells. ACS Energy Lett 3(10):2558–2565 85. Gong C, Tong S, Huang K, Li H, Huang H, Zhang J, Yang J (2020) Flexible planar Heterojunction Perovskite solar cells fabricated via sequential roll-to-roll microgravure printing and slot-die coating deposition. Solar RRL 4(2):1900204 86. Nechyporchuk O, Yu J, Nierstrasz VA, Bordes R (2017) Cellulose nanofibril-based coatings of woven cotton fabrics for improved inkjet printing with potential in e-textile manufacturing. ACS Sustain Chem Eng 5(6):4793–4801 87. Li R, Liu J, Shi A, Luo X, Lin J, Zheng R et al (2019) A facile method to modify polypropylene membrane by polydopamine coating via inkjet printing technique for superior performance. J Colloid Interface Sci 552:719–727 88. Li R, Fan H, Shen L, Rao L, Tang J, Hu S, Lin H (2020) Inkjet printing assisted fabrication of polyphenol-based coating membranes for oil/water separation. Chemosphere 250:126236 89. Spaeth HD (2012) DNA-enhanced efficiency and luminance of organic light-emitting diodes. University of Cincinnati 90. Liu K, Zheng L, Ma C, Göstl R, Herrmann A (2017) DNA–surfactant complexes: selfassembly properties and applications. Chem Soc Rev 46(16):5147–5172 91. You H, Spaeth H, Linhard V, Steckl A (2009) Role of surfactants in the interaction of dye molecules in natural DNA polymers. Langmuir 25(19):11698–11702 92. Yang C, Tang A, Li X, Zhang F, Teng F (2014) Electrochemistry of Cu (I) doped CdS nanoparticles hosted by DNA–CTMA in aqueous electrolyte. Mater Chem Phys 147 (3):1074–1078

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

91

93. Anton AM, Pădurețu C-C, Marin CA, Simion A-M, Simion C, Kajzar F (2019) Photoresponsive natural materials. Mol Cryst Liq Cryst 695(1):37–44 94. Bhatt R, Sreedhar B, Padmaja P (2017) Chitosan supramolecular cross-linked with trimesic acid–facile synthesis, characterization, and evaluation of adsorption potential for chromium (VI). Int J Biol Macromol 104:1254–1266 95. Aga RS, Telek BA, Lombardi JP, Heckman EM, Bartsch CM (2013) Dielectric tunability of DNA biopolymer films with varying amounts of hexadecyltrimethylammonium chloride. J Electron Mater 42(3):463–469 96. Ileana R, Francois K (2014) Biopolymers for application in photonics. 4 97. Irimia-Vladu M, Sariciftci NS, Bauer S (2011) Exotic materials for bio-organic electronics. J Mater Chem 21(5):1350–1361 98. Minami H, Nakamura K, Kobayashi N (2018) Enantioselective luminescence enhancement of chiral Ru (phen) 32+ complexes by interaction with deoxyribonucleic acid. J Nanophoton 12 (3):033005 99. Kim JY, Nagamani S, Liu L, Elghazaly AH, Solin N, Inganäs O (2020) A DNA and self-doped conjugated polyelectrolyte assembled for organic optoelectronics and bioelectronics. Biomacromolecules 21(3):1214–1221 100. Mindroiu M, Manea A-M, Rau I, Grote JG, Oliveira HC, Pawlicka A, Kajzar F (2013) DNA-and DNA-CTMA: novel bio-nano materials for application in photonics and electronics. Paper presented at the ROMOPTO 2012: tenth conference on optics: micro-to nanophotonics III 101. Norwood RA, Thomas J, Peyghambarian N, Wang J, Li L, Ouchen F, Grote J (2010) Hybrid DNA materials for energy storage. Paper presented at the Nanobiosystems: processing, characterization, and applications III 102. Organick L, Ang SD, Chen Y-J, Lopez R, Yekhanin S, Makarychev K, Strauss K (2018) Random access in large-scale DNA data storage. Nat Biotechnol 36(3):242–248. https://doi. org/10.1038/nbt.4079 103. Kumar J, Ouchen F, Smarra DA, Subramanyam G, Grote JG (2015) DNA based electrolyte/ separator for lithium battery application. Paper presented at the Nanobiosystems: processing, characterization, and applications VIII 104. Lazar CA, Kajzar F, Mihaly M, Pirvu C, Petcu AR, Olteanu NL, Rau I (2016) DNA based materials doped with praseodymium (III) hydroxide nanoparticles. Opt Mater 56:3–7 105. Liang L, Fu Y, Li L, Zheng H, Wei X, Wei Y, Kobayashi N (2018) RNA–CTMA dielectrics in organic field-effect transistor memory. Appl Sci 8(6):887 106. Liang L, Fu Y, Wang D, Wei Y, Kobayashi N, Minari T (2018) DNA as functional material in organic-based electronics. Appl Sci 8(1):90 107. Mamangun DMD, Santana JL, Ouchen F, Grote JG, Sotzing GA (2014) Orthogonal alignment of DNA using hexafluoroisopropanol as a solvent for film castings. RSC Adv 4 (75):39798–39801 108. Massin J, Parola S, Andraud C, Kajzar F, Rau I (2013) Enhanced fluorescence of isophorone derivatives in DNA based materials. Opt Mater 35(10):1810–1816 109. Nakamura K, Minami H, Sagara A, Itamoto N, Kobayashi N (2018) Enhanced red emissions of europium(iii) chelates in DNA-CTMA complexes. J Mater Chem C 6(16):4516–4522. https://doi.org/10.1039/c8tc00255j 110. Steckl A, Spaeth H, You H, Gomez E, Grote JJO, News P (2011) DNA as an optical material. 22(7):34–39 111. Sun Q, Subramanyam G, Dai L, Check M, Campbell A, Naik R et al (2009) Highly efficient quantum-dot light-emitting diodes with DNA-CTMA as a combined hole-transporting and electron-blocking layer. ACS Nano 3(3):737–743 112. Nizioł J, Śniechowski M (2014) DNA-CTMA complex for applications in organic electronics: thermal annealing phenomena. Paper presented at the 2014 16th international conference on transparent optical networks (ICTON)

92

D. Bai et al.

113. Sznitko L, Cyprych K, Szukalski A, Miniewicz A, Rau I, Kajzar F, Mysliwiec J (2013) Lasing and random lasing based on organic molecules. Paper presented at the optics and photonics for counterterrorism, crime fighting, and defence IX; and optical materials and biomaterials in security and defence systems technology X

Dan Bai acquired her Ph.D. in Molecular Photonics from Newcastle University, UK, in the year 2013, and was appointed as a lecturer and then associated professor in the School of Medicine, Xi’an Jiaotong University, in the year 2014–2019, followed by a current post in Institute of Flexible Electronics, Northwestern Polytechnical University. As PI and co-PI, Dan’s recent publications include the papers in Dyes & Pigments, Molecular TherapyNucleic Acids, JMCC, Biotech.&Bioeng., and ACS Appl.Mat.Int. Appl.OrgMet.Chem., alongside two book chapters; she received several grants with a sum of over 2 million including the ones from the Natural Science Foundation of China and the Japan-China Medical Association. Her group now focuses on the design and creation of molecular systems for photonic and optoelectronic applications. Huhu Feng is a first-year graduate student majoring in materials and chemistry at the Institute of Flexible Electronics, Northwestern Polytechnical University. He is interested in the interdisciplinary research of biohybrid materials for optoelectronic devices. Currently, he has four papers in press.

Xingchen Yu is currently a junior undergraduate student at the Queen Mary Engineering School, Northwestern Polytechnical University. He was awarded the Outstanding Student (2019), Outstanding Youth League Members (2020), and other distinguished achievements.

3 Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells. . .

93

Chenxin Ran received his B.S. in Applied Chemistry from the Xidian University in 2009 and obtained his Ph.D. in Electronic Science and Technology from the Xi’an Jiaotong University in 2016 under Professor Minqiang Wang’s supervision. He worked as a visiting scholar from 2014 to 2015 with Professor Liming Dai at Case Western Reserve University. He worked as a postdoc fellow in Xi’an Jiaotong University with Professor Zhaoxin Wu from 2016 to 2020. Now he works as an associate professor at Northwestern Polytechnical University. His research interests include photophysics, structural control, and property optimization of Pb/Pb-free-based perovskite materials and their optoelectronic applications.

Academician Huang Wei received his BSc in Chemistry from Peking University in 1983 and his MSc and Ph.D. in Physical Chemistry from the same university. He did his postdoctoral research in the Department of Chemistry with the National University of Singapore (NUS) where he participated in the foundation of the Institute of Materials Research and Engineering (IMRE) since 1995. In 2001, he became Chair Professor at Fudan University, where he founded and chaired the Institute of Advanced Materials (IAM). In June 2006, he was appointed as the Deputy President of the Nanjing University of Posts and Telecommunications, where he initiated the Institute of Advanced Materials (IAM) and the Key Laboratory for Organic Electronics and Information Displays (KLOEID). In November 2011, he was elected as Academician of the Chinese Academy of Sciences (CAS). In July 2012, he was appointed as the President of the Nanjing Tech University (NanjingTech, 2011 University) and founded the Institute of Advanced Materials (IAM) and the Key Laboratory of Flexible Electronics (KLOFE) with Nanjing Tech. In 2013, he was elected as Director-General of Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM). In 2015, he founded the International Research Center of Flexible Electronics (CoFE) and the Joint International Laboratory of Flexible Electronics (LoFE). In October 2016, he led the construction of the National Innovation Talent Cultivation Base of Flexible Electronics (National 111 project). In April 2017, he was appointed as the Executive Vice President of Northwestern Polytechnical University. In October 2017, he founded the Institute of Flexible Electronics (IFE) at Northwestern Polytechnical University.

Chapter 4

First-Principles Calculations for the Interfaces of Perovskite Solar Cells Jun-Peng An, Ying Tian, Hong-Tao Xue, Jun-Chen Li, Jun-Qiang Ren, Xue-Feng Lu, and Fu-Ling Tang

4.1

Introduction

Simulation is becoming an important tool for the research of solar cells, which is a useful supplement to experimental research. First-principles calculation based on density functional theory (DFT) is performed on solar cells to understand the relationship between material structure and device performance. DFT [1–3] has become an effective method in fields of physics, quantum chemistry, and computational materials science, which has reasonably predicted the structure and properties of many substances. After several years of development, the all-solid-state perovskite solar cells have achieved a power conversion efficiency of 25.6% [4], which proves that their usage prospect is huge. Perovskite solar cells consist of multilayer structures [5]: substrate material | electron transport layer | light absorption layer | hole transport layer | metal electrode. The key material to a complete device is the perovskite layer, also known as a light-absorbing layer. It has the general chemical formula ABX3, where A is CH3NH3+, HN ¼ CHNH3+, or HC (NH2)2+; B is generally occupied by Pb or Sn; and X is for I, Cl, Br, or I1-xClx [6, 7]. MAPbI3 is a typical organic-inorganic perovskite. Since elemental substitution can be carried out at MA, Pb, or I, other organic-inorganic perovskites are considered derivatives of MAPbI3. The temperature and pressure can change the distance between atoms and change the bonding mode of the materials, thereby causing the phase change of the CH3NH3PbI3 structure. Many x-ray diffraction (XRD) [8–11] results indicate that the pressure will cause the CH3NH3PbI3 phase change from the tetragonal structure

J.-P. An · Y. Tian · H.-T. Xue · J.-C. Li · J.-Q. Ren · X.-F. Lu · F.-L. Tang (*) State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Department of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou, People’s Republic of China e-mail: [email protected]; tfl@lut.edu.cn © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_4

95

96

J.-P. An et al.

to the orthorhombic structure (space group Imm2). Jiang et al. [12] used DFT to study the interaction between the high-pressure cubic structure of CH3NH3PbI3 and the bandgap. Moreover, the charge density and band structure of orthorhombic CH3NH3PbI3 at different pressures were studied by Kong et al. [10]. Our work used DFT to study the structure, electronic, and optical properties of CH3NH3PbI3 at the highest pressure value of 0.4 GPa. Nowadays, the main problem of this MAPbX3 perovskite material is that Pb is highly toxic and it is not friendly to the environment. Hence, from the long-term project to protect the environment, the promotion and application of such lightabsorbing materials will inevitably be restricted. This requires us to develop less or free-lead perovskite materials. In recent research on element doping of organicinorganic perovskites, In, Ag, Al, and Na are used to replace Pb [13–15]. We used first principles to study the changes of energy bands and optical properties caused by Ag+ doping for Pb2+. For the doping research for perovskite batteries, Ag has excellent performance among similar substitution elements (Sr, Sn, In, and so on) and is an environmentally friendly material. It has a similar ion radius (129 pm) compared with Pb2+ (133 pm) [16], so it will not greatly change the crystal structure. Recently, pure inorganic cesium lead halide perovskite CsPbX3 (X ¼ Cl, Br, and I) has been prepared due to its good thermal stability and high charge carrier mobility. Among them, cubic (α)-CsPbI3 with the bandgap 1.73 eV is the most suitable one for the photovoltaic industry. Unfortunately, in a humid environment even at room temperature, the α-CsPbI3 phase is unstable, making it transform into a nonperovskite orthorhombic phase (δ) with lower photovoltaic efficiency” ([17, 18] with permission). The work by Bai et al. [19] confirmed that Mn2+ ions will be inserted into the gaps of the CsPbI2Br lattice during film growth, thereby reducing the recombination loss and increasing the hole extraction efficiency. After adopting the interstitial Mn-doped CsPbI2Br film, the best device efficiency record of this solar cell was broken in 2018, reaching 13.47%. According to the experimental results of CsPbI2Br thin film, it was found that Mn doping seems to have substitution and interstitial doping structures. Understanding the mechanism of Mn doping will provide meaningful help for energy band engineering and the preparation of stable CsPbI2Br perovskites. In addition to the material properties of each layer of the solar cell, the surface and interface structure and their electronic properties are also important to photovoltaic properties. A tremendous number of researchers have pursued higher power conversion efficiency (PCE) through experiments of surface (interface) regulation [20– 23]. In perovskites, defects in the bulk phase, grain boundaries, surface, and interface will much affect surface charge recombination mechanism and the carrier transport, and this hinders the further improvement of the efficiency of polymer solar cells (PSCs). Dangling bonds on the surface result in new electronic states (energy levels) in the bandgap [24]. The interface or surface states are one of the reasons why perovskite batteries have lower conversion efficiency and poor stability under this

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

97

crystal plane orientation. Therefore, it is also the main method to improve device stability and battery efficiency by adopting an effective passivation method. Therefore, within this chapter, the lattice structures and electrical properties of phenethyl ammonium (PEA+) adsorbed-MAPbI3 (110) surface, WZ-ZnO (100)/ MAPbI3 (112) interface, WZ-ZnO/MAPbI3 interface, NiO (110)/MAPbI3 (100) interface, MAPbI3/Au interface, SnO2 (110)/MAPbI3 (100) interface, CsPbI3/SnO2 interface, and CsPbI2Br/SnO2 interface are investigated. For these materials and surfaces or interfaces in them, their lattice structures the surface energy, or interface energies are provided. We calculated the density of states (DOS), charge density difference, Bader charge, and optical properties (including absorption coefficient, refractive indices, and extinction coefficient). Our purposes are to obtain quantitative atomic and electronic information in the battery materials and then to understand the relationship between composition, structure, and performance.

4.2

Methods

In this chapter, most of the calculation work was performed in the Vienna ab initio simulation package (VASP) [25, 26] “based on first-principles density functional theory (DFT) with the Perdew-Burke-Ernzerhof (PBE) [27] version of the generalized gradient approximation (GGA) and the more complex Heyd-ScuseriaErnzerhof (HSE06) [28] hybrid functional. The projector augmented wave (PAW) [29] method was applied to describe the pseudopotential. The electronic configurations are [Xe]6s1, [Xe]5d106s26p2, [Kr]4d105s25p5, [Ar]3d104s24p5,” ([30] with permission) [Kr]4d105s25p2, [Ar]3d54s2, [Kr]4d105s1, O2s22p4, N2s22p3, C2s22p2, H1s1 for cesium, plumbum, iodine, bromine, stannum, manganese, silver, oxygen, nitrogen, carbon, and hydrogen, respectively. The lattice structure optimization adopts the conjugate gradient (CG) [31] method. The interface binding energy, band structure, and density of states (DOS) in the systems are obtained based on the tetrahedral method of Bloch correction [32].

4.3 4.3.1

CH3NH3PbI3 Phase Transformation Under Pressure Lattice Structures of CH3NH3PbI3 Under Pressure

At room temperature and atmospheric pressure, CH3NH3PbI3 is tetragonal, but it may be the orthorhombic structure or the cubic structure under high pressure. Figure 4.1 shows the different models of tetragonal [33–35], cubic [12], and orthorhombic structures [10] of CH3NH3PbI3.

98

J.-P. An et al.

Fig. 4.1 Lattice structures of CH3NH3PbI3: (a) tetragonal I4/mcm phase, (b) cubic Im3 phase, and (c) orthorhombic Imm2 phase at atmospheric and high pressures. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.2 The projected energy band structures: the tetragonal phase (a), cubic phase (b), and orthorhombic phase (c). The energy orbitals of I-5p, Pb-6s, and Pb-6p are represented by green, blue, and red circles, respectively. (Ref. [36] with permission. Copyright Elsevier Press)

4.3.2

Projected Band Structures for Three Phases of CH3NH3PbI3

“The projected band diagram aims to find out the contribution of the atomic orbit to the energy band near the Fermi level according to the projection of the energy band on each atom and each orbit and its ratio. CH3NH3PbI3’s three-phase projection band structure is shown in Fig. 4.2, and the Fermi levels in Fig. 4.2 are set to 0 eV. The energy bands formed by different orbits are represented by different colors, and each point represents the proportion of orbit projection” ([36] with permission). “Figure 4.2 shows that the conduction band minimum (CBM) of these three structures of CH3NH3PbI3 is contributed by Pb-6p orbitals and valence band maximum (VBM) by I-5p and Pb-6s orbitals. For the two high-pressure phases, the contribution of the Pb-6s orbitals in the cubic phase is greater than the orthorhombic phase. The CBM and VBM of the three structures shown in Fig. 4.2 are at point G, which shows that CH3NH3PbI3 is a direct bandgap semiconductor material regardless of pressure” ([36] with permission).

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

99

Under normal temperature and pressure, the calculated bandgap of the tetragonal structure is “1.76 eV, which is larger than the experimental value of 1.62 eV [8]. The calculation results show that the bandgap values of the three structures at the phase transition point are higher than the corresponding experimental values [10], and the error range of the two is 8%. In the 3 to 0 eV part, the bonding orbitals of Pb-6s and I-5p” ([36] with permission) determine the large fluctuation of the energy band, while in the range of 1.5–5 eV, the antibonding orbitals of Pb-6p and I-5p contribute to the conduction bands.

4.3.3

Variation of the Bandgap of CH3NH3PbI3 with Pressure

We not only investigated “the energy band structure of CH3NH3PbI3 at the phase transition point but also analyzed the effect of different pressure values near the phase transition point on the bandgap. It can be seen from Fig. 4.3 that before the phase transition” ([36] with permission), the bandgap of the tetragonal structure decreases with increasing pressure, while at the phase transition point, the bandgap suddenly increases. After the phase transition occurs, the band gaps of the two possible high-pressure structural phases are larger than those of the normal-pressure tetragonal phase, and after the pressure is gradually increased, the band gaps of the two possible high-voltage structural phases gradually decrease.

Fig. 4.3 Bandgap change of CH3NH3PbI3with pressure. (Ref. [36] with permission. Copyright Elsevier Press)

100

J.-P. An et al.

Fig. 4.4 CBM and VBM of three CH3NH3PbI3 phases under pressure. (Ref. [36] with permission. Copyright Elsevier Press)

To further study the change of the bandgap of three CH3NH3PbI3 phase structures under high pressure, we give the displacements of CBM and VBM of each phase structure under pressure, as shown in Fig. 4.4. “Before the phase transition, the increasing pressure causes the CBM of the CH3NH3PbI3 tetragonal phase to move downward and its VBM as a whole tends to decrease. Therefore, the pressure before the phase change introduces the bandgap of the tetragonal phase to gradually decrease. At the phase transition point of 0.4 GPa, the CBMs of the two highpressure phases move upward, while the corresponding VBMs move downward, increasing the bandgap of the two structures at the phase transition point” ([36] with permission). After the phase change point, the CBM of the two high-pressure phases tends to decrease as a whole and their VBM tends to increase as the pressure increases, indicating that the bandgap gradually decreases under the effect of pressure. Some studies [8] found that the increase and decrease of the bandgap with the change of pressure are mainly due to the octahedron structural deformation of [PbI6]4. The CH3NH3+ organic cation located in the gap of the octahedron does not significantly change the bandgap. However, the deformation mechanism of [PbI6]4 octahedrons and the principle that the distortion structure changes the bandgap are still unclear. In some literatures [37–41], the bond lengths and their average value are used to describe the change of the degree of polyhedron distortion or the symmetry change of the crystal structure. Based on this, we calculate the “average Pb-I bond length in [PbI6]4 octahedrons and the octahedron deformation degree σ of [PbI6]4 to determine the structure deformation and analyze the reason why pressure affects the bandgap” ([36] with permission). Although the “[PbI6]4 octahedrons in CH3NH3PbI3 shrinks at the same pressure, there is no significant difference in its [PbI6]4 octahedrons between atmospheric pressure and the phase transition point of each phase. Figure 4.5 shows [PbI6]4

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

101

Fig. 4.5 The octahedrons [PbI6]4 in CH3NH3PbI3 at the phase transition point: (a) tetragonal phase, (b) cubic phase, and (c) orthorhombic phase. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.6 The change of average length of Pb-I bonds (a) of [PbI6]4 octahedrons and the standard deviation σ (b) with pressure. (Ref. [36] with permission. Copyright Elsevier Press)

octahedrons with three-phase structures at the phase transition point” ([36] with permission). We use the average length of the Pb-I bonds in [PbI6]4 octahedrons in Fig. 4.6a to study and analyze the relationship between the optimized bond length and pressure in the structures of CH3NH3PbI3. We found that the gradual increase in pressure reduces the average length of all Pb-I bonds in the three [PbI6]4 octahedrons, which shows that the pressure causes the distance between Pb and I atoms to

102

J.-P. An et al.

shorten and causes [PbI6]4 octahedrons to shrink slightly. This may be the reason for the bandgap reduction of these three structures with pressure. Besides, to quantify “the dispersion of a set of data, the standard deviation σ of Pb-I bonds in [PbI6]4 octahedrons can be expressed, which can be obtained according to the following Eq. (4.1) to characterize the degree of different structures’ deformation of the octahedron. rffiffiffiffi N 1X σ¼ ðx  μÞ N i¼1 i

ð4:1Þ

In Eq. (4.1), xi represents the Pb-I bond length, N is the number of Pb-I bonds, and the average length of Pb-I bonds is expressed in μ. The lower the σ value, the closer the Pb-I bond length in [PbI6]4 octahedron to its average bond length, indicating the smaller degree of distortion of the octahedron. The higher the σ value, the longer the Pb-I bond lengths in a wider range of values, meaning greater distortion of the octahedron” ([36] with permission). Figure 4.6b shows σ, the change of [PbI6]4 octahedron with the pressure in three CH3NH3PbI3 structures. According to Fig. 4.6b, in the possible cubic and orthorhombic phases, the distortion degrees σ of [PbI6]4 octahedron is greater than that of atmosphericpressure tetragonal phase [PbI6]4 octahedron distortion. From Fig. 4.3, we find that the band gaps of the atmospheric-pressure tetragonal phase are smaller than those of the two high-pressure phases. Therefore, the pressure promotes the phase change of CH3NH3PbI3, and the enhancement of its bandgap is due to the increase of the distortion degree σ of [PbI6]4 octahedron. That is to say, for the three structural phases of CH3NH3PbI3, the higher the degree of distortion σ, the larger the bandgap.

4.3.4

Optical Properties of CH3NH3PbI3 Under Pressure

Figure 4.7 shows the optical properties of the three CH3NH3PbI3 at the phase transition point. The real ε1(ω) and imaginary ε2(ω) of the complex dielectric function are shown. The photon energy range is 0–23 eV, and function ε1(ω) can characterize the electronic polarizability. In Fig. 4.7a, when the energy is zero, the static dielectric constants of the cubic phase, tetragonal phase, and orthorhombic phase are 4.32, 4.45, and 4.5, respectively. Subsequently, these values gradually increased, reaching the maximum at 8.28 at 2.4 eV for the tetragonal phase, the maximum at 8.34 at 2.4 eV for the cubic phase, and the maximum at 7.75 at 2.42 eV for the orthorhombic phase. Their value gradually decreases below zero. It can be seen in Fig. 4.7b that these three structures have two main complex dielectric function peaks. The tetragonal phase shows peaks at 3.05 eV and 6.58 eV, respectively, with ε2(ω) “values of 8.49 and 3.02. For the cubic phase, the two peaks are 7.86 and 2.63, while the orthorhombic phase has peaked at 3.03 eV and 6.22 eV. The intensity of the peaks of the cubic phase is the lowest among the three phases, and the threshold energy of these three phases is close to 1.8 eV” ([36] with permission).

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

103

Fig. 4.7 The complex dielectric function real: ε1(ω) (a) and imaginary ε2(ω) parts (b) for three structures of CH3NH3PbI3 at the phase transition point. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.8 The extinction coefficient k(ω) (a) and refractive index n(ω) (b) at the phase transition point (0.4 GPa). (Ref. [36] with permission. Copyright Elsevier Press)

Figure 4.8a shows three structures’ extinction coefficient k(ω) at the phase transition point. Specifically, the “variation trends of k(ω) with light energy are similar to the imaginary parts ε2(ω). Figure 4.8b shows the refractive indexes n(ω). When photons penetrate the material, they slow down due to their interaction with electrons, so the refractive index is greater than 1. The static refractive index n(0) of the cubic phase, tetragonal phase, and orthorhombic phase are 2.08, 2.11, and 2.12, respectively. The refractive index of the tetragonal phase reaches a maximum of

104

J.-P. An et al.

Fig. 4.9 The optical reflectivity R(ω) (a) and absorption coefficient α(ω) (b) at the phase transition point (0.4 GPa). (Ref. [36] with permission. Copyright Elsevier Press)

2.95 at 2.53 eV. For the other two possible high-pressure phases, the maximum refractive indexes are 2.86 and 2.96, respectively” ([36] with permission). Figure 4.9a shows the optical reflectivity R(ω) at the phase transition point. When the energy is zero, the tetragonal phase has the reflective value “12.7%, and the maximum reflectivity value is about 59.4%” ([36] with permission) occurring at an energy of 7.86 eV. For the orthorhombic phase and the cubic phase, the reflectivity at energy zero is 12.9% and 12.3%, respectively, while the maximum values are 57.4% and 57%, respectively. The maximum reflectivity appears in the “light energy region of 7–8 eV” ([36] with permission). The absorption coefficients α(ω) of the three crystal structures of CH3NH3PbI3 “at the phase transition point are shown in Fig. 4.9b. The material absorbs incident light after the incident light energy exceeds 1.8 eV. As the incident light increases to 2.0 eV, the absorption coefficient of this type of material increases significantly. For these three structures, when the photon energy is greater than 23 eV, the absorption coefficient tends to be almost zero and no longer absorbs light. For the tetragonal phase and cubic phase, the maximum absorption peak appears at 6.80 eV, while for the orthorhombic phase, the maximum absorption peak appears at 6.68 eV” ([36] with permission).

4.4 4.4.1

Perovskite Solar Cell Materials Upon Doping: Ag-Doped MAPbI3: MAPb1-xAgxI3 Structural Properties

For studying the optical and electronic properties of the MAPb1-xAgxI3 system, we optimized the cubic MAPbI3 cell. The optimized lattice parameters are 6.33 Å, which is in good agreement with the experiment value of 6.295 Å [42–45]. A larger

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

105

2  4  5 calculation model is built to obtain a low concentration doped structure with lattice parameters a ¼ 12.660 Å, b ¼ 25.320 Å, and c ¼ 31.650 Å.

4.4.2

Electronic Properties

“Figure 4.10 shows the band structure and total density of states (DOS) of pure MAPbI3. From Fig. 4.10a, we see that the MAPbI3 is a direct bandgap semiconductor material. Although density functional theory (DFT, such as GGA + PBE) often underestimates the bandgap of MAPbI3 perovskites [46–48], our calculated bandgap 1.64 eV is consistent with the simulated results (1.68 eV) [49] and experimental reports (1.60 eV) [50, 51]. Figure 4.10b shows the VBM, mainly contributed by I-5s and Pb-6s orbitals, while the CBM is mainly contributed by I-5p and Pb-6p orbitals” ([52] with permission). The highest experimental Ag doping concentration is 5% [16]. Figure 4.11 shows the band structure and DOS of MAAgI3. Figure 4.11a shows few-electron states nearby the Fermi level, and MAAgI3 has a certain metallicity. We analyzed the total DOS and some DOS of MAAgI3. Figure 4.11b shows the electronic states, mainly attributed by Ag-4d and I-5p orbits contribution nearby the Fermi level. The orbital contribution of I-5p could be introduced by lattice distortion because the ionic radius of Ag+ (129 pm) is less than that of Pb2+ (133 pm). To evaluate MAPb1-xAgxI3’s stability, we calculated the formation of the energy of the MAPb1-xAgxI3 system by Eq. (4.2). E¼

  1 ½Etot ðMAPbI3 Þ þ xEðAgÞ  Etot ðMAPb1x Agx I3 Þ  xE ðPbÞ x

ð4:2Þ

Fig. 4.10 The band structure (a) and density of states (b) of single-cell CH3NH3PbI3. (Ref. [52] with permission. Copyright American Institute of Physics)

106

J.-P. An et al.

Fig. 4.11 The band structure (a) and density of states (b) of single-cell CH3NH3AgI3. (Ref. [52] with permission. Copyright American Institute of Physics)

Equation (4.2) [53] represents the total energy difference of the system before and after the doping of Pb by Ag. “Dividing by x shows the average formation energy of each Ag atom. Etot(MAPbI3) is the total energy of the undoped lattice, Etot(MAPb1xAgxI3) is the total energy of the supercell cell after doping, and x is the number of Ag substitution atoms. E(Ag) and E(Pb) are the corresponding energies of a single Ag and Pb atom, respectively. The formation of energy changes with the doping concentration in Fig. 4.12. With increasing x, the formation energy increases. The formation energy for Ag substitution is negative, indicating that the process will absorb energy as it progresses. Based on the above relationship, the optical band gap of MAPb1-xAgxI3 varies with doping concentration x in Fig. 4.13. The optical bandgap of MAPbI3 is 1.56 eV, and it agrees well with the experimental values (1.51 eV) [35, 54]. The optical bandgap of MAPb0.975Ag0.025I3 (1.57 eV) increases slightly with doping a small amount of Ag. When the Ag adding concentration increased from 0% to 30%, the MAPb1-xAgxI3 optical band gap increases from 1.56 eV to about 3.11 eV. The literature [16, 55] gave that if x < 0.025, the MAPb1-xAgxI3 power conversion efficiency is better. Then, we will show the MAPb1-xAgxI3 photoelectric properties with the low concentration doping (x  0.1)” (Copyright [52] with permission). We calculate the MAPb1-xAgxI3 system’s total DOS and the partial DOS to investigate the electronic structure changes in Fig. 4.14. Figure 4.14a shows that the total DOS of MAPb1-xAgxI3 system varies with x ¼ 0, 0.025, 0.05, 0.075, and 0.100. When x reaches 2.5%, the MAPb0.975Ag0.025I3’s total DOS is almost unchanged. “Comparing the total DOS of MAPb0.950Ag0.050I3 and those of MAPb0.975Ag0.025I3 in Fig. 4.14a, we find that the valence band moves to a higher energy level, and the energy enters the valence band with x reaches 2.5%, which shows MAPb0.95Ag0.05I3 is a p-type conductor. The CBM is almost unchanged. The bandgap is relatively decreased by 0.013 eV. From the total DOS of MAPb0.975Ag0.025I3 to the MAPb0.9Ag0.1I3 in Fig. 4.14a, the total DOS of MAPb1-xAgxI3 is almost unchanged with x increasing” ([52] with permission).

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

107

Fig. 4.12 The total energy of the supercell CH3NH3Pb1-xAgxI3 (x ¼ 0.025, 0.05, 0.075, 0.10, 0.15, 0.20, 0.25, 0.30). (Ref. [52] with permission. Copyright American Institute of Physics)

Fig. 4.13 The optical band gap of the CH3NH3Pb1xAgxI3 with doping concentration. (Ref. [52] with permission. Copyright American Institute of Physics)

Figure 4.14a shows MAPbI3’s total DOS. Combining with Fig. 4.14b, we can find that I-5p and Pb-6s orbits determine pure MAPbI3’s VBM on about 0.21 eV. The MAPbI3’s CBM is controlled by Pb-6p orbits. To study the MAPbI3’s electronic properties variation with Ag doping concentration, the MAPbI3’s supercell was doped with a different number of Ag atom atoms, and their corresponding “doping concentration x increases from 2.5% to 10%. Partial DOS is analyzed for different doping concentrations x. According to Fig. 4.14b, c, the electron orbital contribution of the Pb and I atom is almost unchanged at 2.5% concentration. The slight difference comes from the contribution of the Ag-4d orbit. Figure 4.14a shows that the MAPb0.950Ag0.050I3’s valence band shifts toward high-energy levels at x ¼ 5% concentration. In Fig. 4.14d, there is a slight local phenomenon on the Pb-6p, I-5p, and Ag-4d orbital from 3.2 to 0 eV” ([52] with permission). By analyzing Fig. 4.14d–f, it can be seen that as the doping concentration x increases, the valence band energy from Ag atoms moves to a lower energy

108

J.-P. An et al.

Fig. 4.14 The total density of states of (a) CH3NH3Pb1-xAgxI3; (b) the partial density of states of CH3NH3PbI3, (c) CH3NH3Pb0.975Ag0.025I3, (d) CH3NH3 Pb0.950Ag0.050I3, (e) CH3NH3 Pb0.925Ag0.075I3, and (f) CH3NH3Pb0.900Ag0.100I3. (Ref. [52] with permission. Copyright American Institute of Physics)

level. The conduction band introduced by Ag atoms hardly changes, and the bandgap slightly increases as the doping concentration x increases from 5% to 10%.

4.4.3

Optical Properties

The “optical properties of MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) system are investigated through the absorption coefficient α, dielectric function ε, reflectivity R, extinction coefficient k, and refractive indices n. In Fig. 4.15a, it can be seen that there are four main absorption peaks at 1.56, 2.19, 2.81, and 6.26 eV, respectively. The first absorption peak is contributed by the electronic transition from the

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

109

Fig. 4.15 Imaginary part (a) and real part of dielectric function (b) of CH3NH3Pb1-xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

I-5p or Pb-6s orbital (the valence band) to the Pb-6p orbital (the conduction band) “([52] with permission). The second absorption peak is contributed by the electronic transition from the Pb-6p orbital to the Pb-6p orbital. The third absorption peak is contributed by the I-5p or Pb-6p orbit. The above analysis also explains the absorption peaks in absorption coefficients and refractive index. “The MAPb1-xAgxI3 (x ¼ 0.025, 0.050, 0.075, 0.100) system has three main absorption peaks. Three main absorption peaks have the position of nearby 2.19, 2.81, and 6.26 eV, and with increasing doping concentration (x  0.1), their absorption intensity decreases. This shows that the doping of Ag elements has a certain effect on the absorption strength” ([52] with permission). Figure 4.15b shows the real part of MAPb1-xAgxI3’s dielectric function (x ¼ 0, 0.025, 0.050, 0.075, 0.100). For undoped MAPbI3, it has “good dielectric behavior when the energy range is less than 8.76 eV but shows a certain metal character in the energy range of larger than 8.76 eV. MAPbI3’s high-frequency dielectric constant is 7.19, which agrees very well with the experimental values of 7.1 [56], 6.6 [57], 5.5 [58], 5.6 [59], and 6.5 [60, 61]. The high-frequency dielectric constant is 6.76 at x ¼ 2.5%, lower than that of pure MAPbI3. MAPb1-xAgxI3’s highfrequency dielectric constant (x ¼ 0.050, 0.075, 0.100) is higher than that of pure MAPbI3: 7.6, 10.61, and 7.24, respectively. We note that MAPb0.925Ag0.075I3 has the highest high-frequency dielectric constant. We think that the introduction of Ag elements has a certain impact on it” ([62] with permission). Figure 4.16 shows MAPb1-xAgxI3’s “absorption coefficients (x ¼ 0, 0.025, 0.05, 0.075, 0.100) system. For a solar cell material, the absorption coefficient is a critical parameter of the light absorption layer. The larger the absorption coefficient, the thinner the light absorption layer. From Fig. 4.16, as the Ag doping concentration x increases, the absorption coefficient decreases slightly in the visible light region. Besides, if where the light absorption layer itself is thin, the electron-hole pairs excited by light can be efficiently extracted from the outside, which would achieve a high conversion efficiency. It can be noted that the absorption coefficient rises earlier nearby the basic absorption edge, and a few micron thicknesses of the light absorption layer” ([52] with permission) can fully absorb sunlight.

110

J.-P. An et al.

Fig. 4.16 Absorption coefficient of CH3NH3Pb1xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

As shown in Fig. 4.16, in the five models, MAPb1-xAgxI3’s absorption peaks and absorption intensities vary greatly “in the energy range of less than 5.57 eV. The absorption peak is relatively broad and strong in the range of 2-7 eV. The intensity of the absorption peaks begins to decrease when the energy is greater than 7 eV. This shows that the material has a certain range of light absorption. Beyond this range, light absorption will not occur. MAPb1-xAgxI3 (x ¼ 0.025, 0.050, 0.075, 0.100) model at 1.56 eV has no absorption peak. The absorption peak of MAPb0.9Ag0.1I3 (x ¼ 0.1) model is lower than that of the MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.050, 0.075) model in the range of 2-2.83 eV, but the intensity of the absorption peak is higher than that of MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.050, 0.075) model at 3.09 eV. This means that it can generate electron-hole pairs under low-energy irradiation” ([52] with permission). Figure 4.17a shows MAPb1-xAgxI3’s reflectance spectrum (x ¼ 0, 0.025, 0.05, 0.075, 0.10) model. The “absorption peaks decrease as doping concentration increases. The absorption peak is mainly introduced by the imaginary part of the dielectric function in the range of 1-10 eV” ([62] with permission). Figure 4.17b, c shows MAPb1-xAgxI3’s refractive index and extinction coefficient. The relevant data are all from the dielectric function. Now, MAPb1-xAgxI3’s data of the reflectivity, refractive index, and extinction coefficient in the experiment are relatively rare. We hope that our calculation data would be a certain reference for the experiment.

4.5 4.5.1

Mn-Doped CsPbI2Br: Structural, Electronic, and Optical Properties Structural Properties

At room temperature, “CsPbI3 has a cubic crystal structure with a space group of Pm3m (No. 221), as shown in the unit cell structure in Fig. 4.18a. We found that the lattice constants obtained from the experiment are a ¼ b ¼ c ¼ 6.289 Å, α ¼ β ¼ γ ¼

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

111

Fig. 4.17 Reflectivity R (a), refractive indices n (b), and extinction coefficient k (c) of CH3NH3Pb1xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

Fig. 4.18 Lattice structure: CsPbI3 (a) and CsPbI2Br (b). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

90 [64]. In our work, the lattice constants of CsPbI3 from first-principles calculations are 6.408 Å, similar to other theoretical calculation results (6.394 Å) [65]. Based on the optimized CsPbI3 lattice model, CsPbI2Br’s lattice model is constructed by substitution doping (see Fig. 4.18b). After optimization, the lattice constant of CsPbI2Br is a ¼ b ¼ c ¼ 6.297 Å. Compared with CsPbI3, the lattice of CsPbI2Br shrinks because the radius of the Br atom is smaller than that of the I atom” ([63] with permission). The “2  2  2 CsPbI2Br supercell is composed of 40 atoms. When one of the Pb atoms is replaced by an Mn atom, the doping concentration is 12.5%, and the lattice parameters reduce from 12.595 Å to 12.392 Å” ([63] with permission). Compared with Pb atoms’ radius, the reason for the lattice shrinkage of the lattice is the smaller radius of Mn atoms [66]. When Mn atoms occupy one of the octahedral gaps, the doping concentration is 2.5%, and then the lattice parameters rise to 12.619 Å. Therefore, only the octahedral gap of the CsPbI2Br supercell is taken into consideration for interstitial Mn doping. Figure 4.19 shows the most stable structure of Mn substitutional doping and Mn interstitial doping in CsPbI2Br. “To evaluate the thermodynamic stability of Mn-doping in CsPbI2Br, we obtain its binding energy by Eq. (4.3)” [67]:

112

J.-P. An et al.

Fig. 4.19 Mn substitutional doped CsPbI2Br (a) and Mn interstitial doped CsPbI2Br (b). (Ref. [63] with permission. Copyright Royal Society of Chemistry)



Pb I Br Mn E tot  nCs E Cs atom  nPb E atom  nI E atom  nBr E atom  nMn E atom Eb ¼ ðnCs þ nPb þ nI þ nBr þ nMn Þ

 ð4:3Þ

where Eb and Etot are the binding energy and “total energy of Mn-doped CsPbI2Br, EA atom is the energy of free atoms i (i ¼ Cs, Pb, I, Br, Mn), and ni is the number of atoms i in the supercell. The calculated binding energies of Mn substitutional and interstitial doped CsPbI2Br are 2.5761 eV/atom and 2.5789 eV/atom, respectively. The negative binding energies and the small energy difference of 2.8 meV indicate that they are two stable doping structures. We can understand the fact that the Mn substitutional and interstitial doped CsPbI2Br had been prepared by experiments [68–70]. Mn interstitial doped CsPbI2Br seems slightly more stable due to its lower binding energy” ([63] with permission).

4.5.2

Electronic Properties

Figure 4.20a shows that the CBM and “the VBM of CsPbI2Br are at the G point. This indicates that CsPbI2Br is a direct bandgap semiconductor material, and it has a bandgap of 1.42 eV. The VBM of CsPbI2Br is mainly contributed by the I-5p and Br-4p orbitals, while the CBM is mainly contributed by the Pb-6p orbital. The band structure of Mn substitutional doped CsPbI2Br is denser than that of CsPbI2Br (Fig. 4.20b). Also, several impurity bands are appearing near the Fermi level, caused by Mn-3d orbitals. As the Mn-3d orbital is half-filled, these impurity bands are also called intermediate bands (IBs). The IBs can act as effective stepping stones to help transport valence electrons into the conduction bands, that is, the existence of IBs helps additional electronic transitions from VB to IB and from IB to CB, apart from the direct electronic transition from VB to CB. IB reduces the photon absorption energy threshold and increases the photocurrent, while the output voltage of the solar cell is not limited by this threshold. CsPbI2Br has a relatively wide bandgap which is

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

113

Fig. 4.20 Band structures and TDOS: (a) undoped, (b) Mn substitutional doped CsPbI2Br, and (c) Mn interstitial doped CsPbI2Br. (Ref. [63] with permission. Copyright Royal Society of Chemistry)

not suitable for direct use as a light-absorbing material for solar cells, and Mn-doped CsPbI2Br can become a promising one for high-efficiency solar cells [71]. Also, there are many novel intermediate-band photovoltaic materials, for example, Ti or Cr-substituted CuGaS2 [72, 73] and Mn-doped AlGaN [74], which have been successfully fabricated experimentally by doping the transition elements. CBM and VBM at the G-point indicate that Mn substitutional doped CsPbI2Br is still a direct bandgap semiconductor. The Mn interstitial doped CsPbI2Br’s Fermi level passes through CBM, and the impurity level introduced by Mn doping is located above the bandgap and near the bottom of the conduction band (Fig. 4.20c). This shows its characteristics of an n-type semiconductor. Several flat bands near the Fermi level are expected to produce a sharp peak in the DOS, mainly from the Mn-3d orbital” ([63] with permission). The bonding characteristics of CsPbI2Br and Mn doping CsPbI2Br (100) planes were shown in Fig. 4.21. In Fig. 4.21a, the charge density around Pb atoms has a strong locality and the bonding characteristics are ionic. Stronger than the covalent bond, the bonding feature between Pb atoms and I or Br atoms is a mixed one. In Fig. 4.21b, the bond length of Mn-I or Br becomes shorter, which is beneficial to the charge transfer between them. The Mn substitutional doped CsPbI2Br has similar binding characteristics as with CsPbI2Br. In Fig. 4.21c, Mn atoms enter the octahedral gap of CsPbI2Br, I, and Br atoms in the PbI4Br2 octahedron are attracted by the Mn atom to cause charge change. The enhanced binding capacity indicates that the doping enhances electron transport and conductivity. PbI4Br2’s octahedron is found deformed.

4.5.3

Optical Properties

The response of ion and electron motion to dipole displacement in the external electric field can be expressed by the dielectric constant [75]. Comparing the DOS and the imaginary part of the dielectric function, we can reveal the interband transition process of excited electrons. The imaginary part of the dielectric function of CsPbI2Br has

114

J.-P. An et al.

Fig. 4.21 The charge densities on the (100) plane of the undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

Fig. 4.22 Real and imaginary parts of the dielectric function: undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

three obvious peaks (Fig. 4.22a): the peak at 2.84 eV mainly comes from “the VBM (I-5p and Br-4p) to the CBM (Pb-6p) (see Fig. 4.20a). Figure 4.22b, c shows, in the high-energy range (>2.884 eV), the imaginary part of the dielectric function of Mn substitutional and interstitial doped CsPbI2Br is the same as that of the CsPbI2Br. In the low-energy range (0–2 eV), the imaginary part of the dielectric function of CsPbI2Br is different from that of Mn-doped CsPbI2Br. The imaginary part of the dielectric function of Mn-doped CsPbI2Br produces a new peak at 0.32 eV with a peak value of 0.25, while Mn-doped CsPbI2Br has a new peak at 0.39 eV with a peak value of 2.25, mainly derived from the electron transition between the VBM and the

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

115

Fig. 4.23 Absorption coefficient spectrum: undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

impurity band (see Fig. 4.20c). The static dielectric constant (SDC) is the refractive index measured at a frequency higher than the lattice vibration frequency [76] and can be obtained from the real part of the dielectric function. Figure 4.22 shows that the SDC of the undoped, Mn substitutional, and interstitial doped CsPbI2Br are 4.97, 5.99, and 10.27, respectively” ([63] with permission). As shown in Fig. 4.23a, as the “photon energy greater than 1.36 eV, the absorption coefficient of the undoped CsPbI2Br increases to the maximum at 6.65 eV. The absorption spectra of Mn-doped CsPbI2Br in Fig. 4.23b, c show that the light absorption capacity is the same in the high-energy range (>1.28 eV). And in the low-energy range (0–1.28 eV), the light absorption capacity of Mn-doped CsPbI2Br is stronger than the other two. We found that an absorption peak appears at 0.61 eV, which means that the photon energy of 0.61 eV can be absorbed, thereby promoting the transition of electrons from the impurity band near the Fermi level to the CBM” ([63] with permission). The impurity band of Mn-doped III-V compounds had similar reports [77–79].

4.6 4.6.1

Surface in Perovskite Solar Cell Materials: PCBM-Adsorbed MAPbI3 Surface Structure of PCBM-Adsorbed Surface

Figure 4.24a shows the surface adsorption model: the upper half is the adsorbed phenyl-C61-butyric acid methyl ester (PCBM) molecule, and the lower half is the

116

J.-P. An et al.

MAPbI3 (100) plane. The structure contains 332 atoms, including 122 H, 18 N, 90 C, 66 I, 24 Pb, and 2 O atoms. Since the adjacent unit cells will have periodic effects, we add 30 Å vacuum layers along the c-axis to the adsorption system. During optimization for the structure, the positions of the four layers of atoms away from the surface of MAPbI3 are fixed, while the remaining atoms and PCBM molecules are relaxed. The single-point calculation determines the lowest energy point of the adsorption structure at different adsorption distances. In Fig. 4.25, when PCBM adsorbs to the MA+ site (see Fig. 4.24b) and the adsorption distance is 3.708 Å, the structure has the lowest energy. With structure optimization, the atom position and bond length change as shown by the red arrow in Fig. 4.24a. Comparing it with the clean MAPbI3 surface, the Pb and I atoms in the first layer move down 0.017-0.03 Å and 0.3550.437 Å along the Z-axis, respectively. The C-N atomic bond length in the second layer reduces from 1.510 Å to 1.489 Å. The I and Pb atoms in the third layer move slightly. Also, the C–C bond length in the PCBM molecular cage increased slightly, and the C-C-C angle connecting the two branches to C60 increases by 3 . This is

a

Vacuum

b

Realxed

surface layer1 layer2 layer3 layer4

Fixed

layer5 layer6 layer7 +

MA

C

I

H

Pb

O

adsorption sites: 1 above-I 2 above-Pb

z

3 bridge x

4 above-MA+

Fig. 4.24 Adsorption structure (a) and sites (b) of PCBM on MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

–1693.85

–1693.86 Total energy/eV

Fig. 4.25 Total energy at different adsorption distances. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

117

–1693.87

–1693.88

–1693.89 (3.708,-1693.895)

–1693.90 3.50 3.55 3.60 3.65 3.70 3.75 3.80 3.85 3.90 Molecule - surface distace/Å

because PCBM molecules adsorbed on the surface cause periodic damage, also cause interaction between atoms.

4.6.2

Adsorption Energy of PCBM on MAPbI3 Surface

Adsorption energy determines the stability of the adsorption system [80]. The adsorption energy of PCBM molecule on the surface of MAPbI3 can be calculated by Eq. (4.4):   E ¼ E atom=MAPbI3ð100Þ  E MAPbI3ð100Þ  NE atom =N

ð4:4Þ

In the above expression, Eatom/MAPbI3 (100) and EMAPbI3 (100) represent the total energy of the system before and after adsorption, respectively. Eatom represents the total energy of PCBM, and N (where N ¼ 1) represents the number of molecules adsorbed atoms. The calculated adsorption energy is 0.87 eV/PCBM molecule. It is a large adsorption binding capacity and the PCBM molecule has good stability on the surface” (Copyright [18] with permission).

4.6.3

Electronic Properties of PCBM-Adsorbed MAPbI3 Surface Model

Figure 4.26 shows the total DOS and (local density of states) LDOS of the adsorption surface. The bandgap is about 0.9 eV when the surface adsorbs PCBM, and the system shows semiconductor characteristics. Moreover, no new electronic state is generated near the Fermi level 0.3 to 0.2 eV, which is conducive to generate photogenerated carriers. The new surface state has a deleterious effect on the solar

118

J.-P. An et al.

A- Total B- Layer1 C- Layer2 D- Layer3 E- Layer4 F- Layer5

200

100

DOS

150

50 0 A B C D E F –5

–4

–3

–2

–1

0

1

2

3

4

5

Energy/eV

Fig. 4.26 The total DOS (a) and the LDOS (b–f) of the PCBM-adsorbed MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

cell’s filling factor and photovoltaic conversion efficiency. Comparing the LDOS of the first layer MAPbI3 (100) in Fig. 4.26b and the LDOS of the fifth layer in Fig. 4.26f (property is similar to the body), no new electrons appear near the Fermi level state. In summary, PCBM is a material suitable for electronic transmission materials. By calculating the differential charge, the charge transfer of the system is visually displayed. The calculation by Eq. (4.5) is as follows [81]: Δρ ¼ ρMAPbI3=PCBM  ρMAPbI3  ρPCBM

ð4:5Þ

where ρMAPbI3/PCBM is the total charge density of the structure, while ρMAPbI3 and ρPCBM are the charge density of MAPbI3 or PCBM separately in the same adsorption system. In other words, when calculating the charge density of MAPbI3, the PCBM in the model will be deleted as a vacuum, and vice versa. The differential charge of this system adsorbed by PCBM is shown in Fig. 4.27, and green and yellow indicate the decreasing (negative) or increasing (positive) charge density, respectively. The right part of Fig. 4.27 is the charge transfer distribution, and the left part is partially enlarged. The atoms closer to the surface have a “higher electron density, and the farther away from the surface, the smaller the electron transfer. The redistributed electrons promote the bonding. Specifically, for PCBM molecules, the number of electrons of C atoms and H atoms near I atoms decreases, while the number of electrons of C atoms and H atoms near Pb atoms increases. In the first layer of MAPbI3, the number of electrons on both I and Pb atoms increases. This difference in electron transfer also indicates the good bonding ability between the PCBM and MAPbI3 atoms” (Copyright [18] with permission).

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

119

Fig. 4.27 The differential charge of the adsorbed surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

4.6.4

Optical Absorption Properties of PCBM-Adsorbed MAPbI3 Surface

The absorption coefficient spectra of the bulk MAPbI3, MAPbI3 (100), and “MAPbI3 surface adsorbed by PCBM are given in Fig. 4.28. There are optical band gaps on all model surfaces. When the photon energy value is at 3.75 eV, the absorption coefficient of the bulk MAPbI3 reaches its maximum peak, and its absorption edge is 1.85 eV. In the photon energy range of 1.5-2.5 eV, the absorption coefficients of the MAPbI3 (100) and the adsorbed surface have a consistent change trend. But the MAPbI3 surface adsorbed by PCBM shows a lower absorption coefficient when the photon energy is less than 3.0 eV. It shows a higher light absorption intensity in the energy range of 3.15–5.0 eV.” (Copyright [18] with permission). Our work provides additional information for experimental research and guides the design of experimental research, and may help design the efficiency of perovskite-type solar cells to seek effective control methods.

4.7 4.7.1

PEA+-Adsorbed MAPbI3 Surface Structure of PEA+-Adsorbed Surface

PEA+-adsorbed MAPbI3 (110) surface contains 84 atoms in total (Fig. 4.29). Passivated H atoms were added to the left of the model to eliminate the effects of dangling bonds of surface atoms. We added a 10 Å vacuum layer to avoid the interaction for the periodic unit cell. The single-point energy calculation method obtains the optimal adsorption distance of the passivating agent (the distance from the N atom in PEA+ to the I atom in the center of the MAPbI3 surface), and then we determine the most stable adsorption structure. The relationship between the total

120 1.2x104

Naked MAPbI3 surface PCBM-adsorbed MAPbI3 surface

1.0x104

MAPbI3 bulk

Absorption/cm–1

Fig. 4.28 The absorption coefficient spectrum of the MAPbI3 bulk, clean MAPbI3 surface, and PCBM-adsorbed MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

J.-P. An et al.

4

8.0x10

6.0x104 4.0x104 2.0x104 0.0 0

1

2

3 Energy/eV

4

5

Fig. 4.29 The structure of PEA+-adsorbed MAPbI3 (110) and the changes in the positions of atoms after optimization. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

energy and the adsorption distance is shown in Fig. 4.30. When PEA+ is 3.393 Å from the center I atom, the total energy of the entire system is 389.377 eV, which is the lowest. In the optimization of the model shown in Fig. 4.30, the first and second layers of atoms near the surface of MAPbI3 (110) and PEA+ are relaxed, and the remaining layers of atoms away from the surface of MAPbI3 (110) are fixed to simulate the bulk. “The red arrow in Fig. 4.30 shows the direction of movement of the atomic position after the relaxation. In the first layer, the I atoms close to the adsorption surface approach the center, while the I atoms far away from the surface move in the opposite direction. The values of movement are 0.204 Å for the former and 0.022 Å for the latter. Also, the dPb-I bond lengths are 3.160 Å and 3.179 Å, which are shortened by 0.025 Å and 0.007 Å compared with before optimization. Compared with the clean surface, I atoms move in the opposite direction, relating to the adsorption of the passivating agent on the surface. In the second layer, the I and

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

121

Fig. 4.30 The total energy change of the model under different adsorption distance. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

Pb atoms moved by 0.057 Å and 0.008 Å along the Z-axis in the negative and positive directions, respectively. The average bond length of dPb-I formed is 3.1755 Å. The changes of bond angle and bond length in the MA+ cation in this layer are greater than the corresponding changes in the first layer. After the relaxation, the bond length (angle) of PEA+ also increased or decreased. The most significant is that the C–N bond length near the surface increases by 0.018 Å, and the N-C-C bond angle increases by 2.078 ” (Copyright [18] with permission). For the MAPbI3 part, the Pb atoms, cations, and I atoms at the central position are adsorbed by the passivating agent PEA+, and a larger movement will occur. Due to the periodic destruction of the lattice structure, the atoms closer to the surface move more significantly during relaxation.

4.7.2

Adsorption Energy of PEA+ on MAPbI3 Surface

The adsorption energy of PEA+ on the surface of MAPbI3 (110) characterizes the stability of the adsorbed substance on the surface (Eq. 4.4) [80], which is 3.898 eV. The adsorption binding shows that the adsorbent has good stability on this surface.

122

4.7.3

J.-P. An et al.

Electronic Properties of PEA+-Adsorbed MAPbI3 Surface Model

According to the density of states, the passivation of the electronic states on the surface of MAPbI3 (110) by PEA+ can be analyzed. Figure 4.31a shows the total density of states (DOS) of the adsorption system. PEA+ has greatly reduced electronic states present on the surface of MAPbI3, and its DOS value has decreased from 11.5 to about 6.5. Also, PEA+ has the function of adjusting the bandgap while maintaining the semiconductor characteristics of MAPbI3. Compared with the band gap value of the clean MAPbI3 (110) surface, the surface value of adsorbed PEA+ will be significantly reduced. The adsorption of PEA+ does not introduce a new surface state, and the passivation effect is good as shown in Fig. 4.31b. Comparing the local density of clean surfaces, it can be seen from Fig. 4.31c–f that the addition of adsorbed molecules makes the first to fourth layers have no peak value nearby the Fermi. In particular, the passivation effect of PEA+ on the first and third layers is the most obvious. As mentioned earlier, because the bottom passivation H introduces some electronic states, the fifth layer of atoms still has small peaks near the Fermi level 0.5 to 0.3 eV. Figure 4.31i–n is the partial wave state densities of C, N, I, and Pb atoms near the surface and PEA+. Combining with Fig. 4.31i, l and n, at 1.5–2.5 eV from the Fermi level, the C-2p electron orbit in PEA+, and the N-2s electron orbit in MAPbI3, Pb-6p electron orbit exists hybridization. Figure 4.31i, m shows that there are hybridizations between the C-2p electron orbit in PEA+ and the I-5p electron orbit in MAPbI3. The passivation effect of the adsorption of PEA+ on the surface of MAPbI3 (110) is good on its surface state, and the electron orbital hybridization of atoms near the surface also promotes the interaction between the atoms. Figure 4.32 shows the differential charge of PEA+ adsorbed on the surface of MAPbI3 (110). The decrease (negative) or increase (positive) of the charge density is

Fig. 4.31 Adsorption structure of PEA+-MAPbI3: total state density (a), local state density of PEA+ (b), local state density maps of the first to fifth layers (e–g); localization of the passivated H layer domain state density (h); C, N, atomic partial wave state density (i–j) of PEA+ and C, N, I, Pb atom partial wave state density (k–n) at the surface. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

123

Fig. 4.32 Differential charge of PEA+-MAPbI3 adsorption structure. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

represented by cyan and yellow, respectively. Figure 4.32a shows the charge transfer as a whole, and Fig. 4.32b is a partially enlarged view. As can be seen from Fig. 4.32, charge transfer occurs in the entire system, but charge transfer between atoms near the surface is more obvious. The farther from the surface, the less the electron transfer. The electron density at the surface is the highest, and the charge is transferred to both sides of the surface. Electron redistribution, which promotes the bonding between the atoms on the surface, indicates that the interaction between the atoms on the surface is strong. According to Fig. 4.32b, in the first layer of MAPbI3, the number of electrons on the I atom increases, and the number of electrons on the Pb atom decreases. For the PEA+ part, the number of electrons of C atoms near the surface decreases, while the number of electrons on N atoms increases.

4.8 4.8.1

Interfaces in Perovskite Solar Cells Materials: MAPbI3/ WZ-ZnO Interface Interface Structure

The interface model of WZ-ZnO (100)/MAPbI3 (112) is shown in “Fig. 4.33. The interface is composed of two parts, on the left are six monoatomic layers in the WZ-ZnO (100) plane, and on the right are six diatomic layers of the MAPbI3 (112) plane” ([84] with permission). To avoid the influence of surface dangling bonds on the interface performance, 12 and 4 passivation H were added on both sides of WZ-ZnO (100) and MAPbI3 (100) surfaces, respectively. Moreover, a ¼ 10.69 Å, b ¼ 9.35 Å, and c ¼ 51.14 Å are the lattice parameters of the WZ-ZnO (100)/ MAPbI3 (112) interface. “There are 36 Zn atoms, 36 O atoms, 6 N atoms, 6 C atoms,

124

J.-P. An et al.

Fig. 4.33 WZ-ZnO (100)/MAPbI3 (112) interface model. (Ref. [83] with permission. Copyright Springer Nature Group)

Fig. 4.34 The total energy of the interface varies with the interface distance. (Ref. [83] with permission. Copyright Springer Nature Group)

6 Pb atoms, 18 I atoms, 36 H atoms, and 16 passivated H atoms in the interface model. To eliminate the adverse effects caused by the periodic boundary conditions of atoms at the surface, a vacuum layer of 30 Å was added to the model” ([84] with permission). Also, we determined a reasonable interface distance by a single-point calculation. As shown in Fig. 4.34, when the interface distance is 2.582 Å, the model has the lowest total energy of 636.05 eV. The interface binding energy can measure the stability of the WZ-ZnO (100)/ MAPbI3 (112) interface. We use Eq. (4.4) to calculate the interface binding energy. The interface binding energy is 0.164 J/m2. This result is smaller than the values of other systems in the literature, such as WZ-ZnO/CdS (0.61 J/m2) [85], Cu2ZnSnS4/ WZ-ZnO (0.21 J/m2) [86], and WZ-CuInS2/WZ-CdS (0.68 J/m2) [87]. This shows that WZ-ZnO/MAPbI3 is a weakly bonded interface.

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

4.8.2

125

Electronic Properties

The electrical properties of the WZ-ZnO (100)/MAPbI3 (112) interface were investigated, specifically, the TDOS, PDOS, the LDOS, and differential charge density of the interface. The TDOS of the WZ-ZnO (100)/MAPbI3 (112) interface is shown in Fig. 4.35a; it is easy to see that there is a new electronic state near the Fermi level. This is because the formation of the heterojunction causes the periodical structure of the crystal to be destroyed so that an additional electron energy level is generated between the conduction band and the valence band, which is also called the interface state. Figure 4.35b, c is the LDOS of the fifth and first layers of the WZ-ZnO (100) plane. Because the fifth layer is far away from the interface, it is less affected by periodicity and dangling bonds. And its properties are closest to the bulk phase, so it is used to analogize the bulk phase WZ-ZnO. From the LDOS in the first layer of the WZ-ZnO (100) plane, a new energy density state is found at 0.5 eV to 0.3 eV near the Fermi level. Figure 4.35d, e is the LDOS of the fifth and first layers of the MAPbI3 (112) plane. Similarly, we use the fifth layer away from the interface to analogize the bulk phase. After a comparative analysis, “we found that new electronic states also appeared at the first layer of MAPbI3 (112), from 1.0 eV to 0 eV near the Fermi level. The above analysis shows that the first layers of WZ-ZnO (100)

Fig. 4.35 The total density of states of WZ-ZnO (100)/MAPbI3 (112) interface (a); the local density of states of WZ-ZnO layer 5 (b) and layer 1 (c); the local density of states of MAPbI3 layer 5 (d) and layer 1 (e); the partial density of states of Zn (f), O (g), CH3NH3 (h), I (i), Pb (j) on the interface layer 1. (Ref. [83] with permission. Copyright Springer Nature Group)

126

J.-P. An et al.

and MAPbI3 (112) have interface states at 1.0 eV to 0 eV and 0.5 eV to 0.3 eV, respectively” ([88] with permission). The PDOS of the interface system is helpful for further understanding of the WZ-ZnO (100)/MAPbI3 (112) interface. Figure 4.35f–j is the PDOS of Zn, O, I, Pb, CH3NH3 atoms at this interface. Figure 4.35f, g shows those near the Fermi level 0.5 eV to 0.3 eV, in which electronic states appear on the WZ-ZnO (100) surface side mainly from O-2p orbital contributions. Also, the orbital hybridization phenomenon exists in the Zn-3d and O-2p electron orbits at 6 eV to 4 eV near the Fermi level. Figure 4.35h, j shows that near the Fermi level, new electronic states appear on the MAPbI3 (112) surface side mainly from the contributions of I-5p and Pb-6s electron orbitals. According to Fig. 4.35f, j, at the Fermi level, 3 eV to 0 eV, the O-2p electron orbit, I-5p electron orbit, Pb-6s electron orbit, and Pb-6p electron orbit on the MAPbI3 | WZ-ZnO interface also had orbital hybridization. The PDOS of all atoms contributing to the new electronic state at the interface is shown in Fig. 4.36, which is the partial wave state density of Zn1-Zn6, Ox (x ¼ 1–6), and I1-I3 atoms. From Fig. 4.35c, f, it can be seen that from 1.0 eV to 3.0 eV from the Fermi level, the new electronic state on the surface of WZ-ZnO (100) is contributed by Zn atoms on the first layer. Combined with the analysis of Fig. 4.36a–f, it is mainly derived from the 4s orbital of Zn2 atom, Zn4 atom, and Zn6 atom, and this is because these three atoms are affected by the H atom in

Fig. 4.36 The partial density of states of Zn1 (a), Zn2 (b), Zn3 (c), Zn4 (d), Zn5 (e), Zn6 (f), Ox (x ¼ 1–6) (g) atoms on the WZ-ZnO layer 1, I1 (h), I2 (i), I3 (j) atoms on the MAPbI3 layer 1. (Ref. [83] with permission. Copyright Springer Nature Group)

4 First-Principles Calculations for the Interfaces of Perovskite Solar Cells

127

Fig. 4.37 The density charge difference (a) and partially enlarged view (b) of the WZ-ZnO (100)/ MAPbI3 (112) interface. (Ref. [83] with permission. Copyright Springer Nature Group)

CH3NH3 ion [89]. From Fig. 4.36g, it is found that the contribution of six O atoms at different positions at the interface to the interface state is the same. Figure 4.36h–j clearly shows that the 5p orbital of the I3 atom contributes more to the interface state, while the 5p orbital of the I1 and I2 atoms contributes to the interface state on a comparable level. To study the electronic properties of the WZ-ZnO (100)/MAPbI3 (112) interface more comprehensively, we calculated the differential charge density of the interface system according to Eq. (4.5). Figure 4.37a qualitatively describes the redistribution of charges in the WZ-ZnO (100)/MAPbI3 (112) interface, and Fig. 4.37b is a partially enlarged view. We use yellow to indicate electron increase and bluegreen to indicate electron loss. From Fig. 4.37, it is found that the charge transfers at the interface will promote the bonding near the interface. Moreover, the closer the atoms to the interface, the greater the degree of charge transfer.

4.8.3

Optical Properties

As a material of the light absorption layer, MAPbI3 is a necessary one to understand the optical properties of its interface. If the energy of the photon is greater than the bandgap of the material, the valence band electrons are excited to the conduction band by absorbing the photon, and light absorption may occur. From Fig. 4.38, “the optical band gap of bulk MAPbI3 is about 1.81 eV, which is close to the reference values 1.62 eV [90] and 1.65 eV [91]. The absorption edge of bulk MAPbI3 is calculated to be approximately 1.6 eV, which is higher than the experimental value of 1.5 eV and the calculated value based on SOC-GW ( n > 9) vanadium oxides, exhibit valences of 4+, +5, and +3, +4 respectively. For the

Carbon composite; hierarchical structure Carbon-based composite; hierarchical structure

Graphene-based composite and special morphology

Carbon composite and special morphology Graphene coated and special morphology Carbon composite and special morphology Carbon composite and super hierarchical

Dandelion-like V2O3/C V2O3@NSCNF

V2O3/NG

Porous shuttle-like V2O3/C G-V2O3

Hierarchical porous metallic V2O3@C

V2O3@PNCNFs

V2O3/rGO

V2O3/C

Strategy Carbon composite and special morphology Carbon composite and special morphology Graphene-based composite

Sample V2O3@C HS

Table 13.4 Some applications of V2O3 in this field

Cathode/ AZIBs

Li-S batteries Anode/KIBs

Anode of LIBs and SIBs Anode/SIBs

Anode/LIBs

Anode/LIBs

Anode/LIBs

Anode/LIBs

Application Anode/LIBs

Electrospinning and heat treatment Hydrothermal calcination

Plasma-enhanced CVD

Calcination of MOFs

Hydrothermal and annealing Hydrothermal; electrospinning and calcination Hydrothermal and calcination

Hydrothermal and annealing Hydrothermal

Method Hydrothermal

181mAh/g@2A/g, 68% retention after 1000 cycles 1000mAh/g@2C/g, ~50% retention after 1000 cycles ~230mAh/g@50 mA/g, 95.8% retention after 500 cycles ~200mAh/g@4A/g, 90% retention after 4000 cycles

435mAh/[email protected]/g, after 250 cycles (LIB); 154mAh/[email protected]/g after 500 cycles (SIB)

Performance 583mAh/g@2A/g no obvious decrease after 800 cycles 732mAh/g@100 mA/g, 95% retention after 100 cycles 675mAh/[email protected]/g, no obvious decrease after 300 cycles 315mAh/g@2A/g, 94% retention after 1000 cycles 800mAh/[email protected]/g; no obvious decrease after 100 cycles

[104]

[103]

[102]

[101]

[100]

[99]

[98]

[97]

[83]

References [96]

13 The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . . 369

370

Y. Wang et al.

Magnéli vanadium oxides, there are few reports on the application of energy storage and electrocatalysis. So in this section, we will introduce Wadsley vanadium oxides in the aspect. Wadsley phase vanadium oxide is a series of mixed-valence vanadium oxides. There are mainly three types, namely V3O7, V4O9, and V6O13. All of them are single- and double-layered compounds. V3O7 and V6O13 have some reports in the field of energy storage such as supercapacitors [105, 106] and metal-ion batteries [107], while V4O9 is rarely studied [108]. Its main application strategy in electrochemistry is to prepare Wadsley phase vanadium oxide with certain morphology or pore structure by hydrothermal method (partial post-heat treatment) [105] or to combine it with carbon carrier for application [109].

13.2.1 V3O7 Li et al. [110] fabricated a well crystalline V3O7H2O by a high-temperature mixing method (HTMM). To compare with V3O7H2O, two types of VO2 were synthesized in the same way. All the vanadium oxides were applied in cathode material of lithium-ion batteries and the test results showed that V3O7H2O has the best performance: 250 mAh/g at a current density of 100 mA/g and 80% retention after 100 cycles test. The authors also discussed the reason for the good performance of V3O7H2O in the view of crystallography. According to the results of XRD refinement, the crystal model was obtained and suggested that well crystalline was benefited in the electrical conductivity of the material. With the presence of the crystal water, the distance of the crystal plane became larger, the lithium-ion could migrate more freely between the interlayer (V3O7H2O), compared with the diffusion in tunnels (two types of VO2), and it is conducive to the lithium-ion deintercalation performance. After that, Liu et al. [111] further carried out with surface carbon coating on the material and obtained the V3O7H2O@C by the same method and found that this material could reverently insert 3.3 moles of lithium ions per mole and had stable cycling performance under different rate conditions. Shen et al. [109] composited V3O7H2O with rGO by hydrothermal method and obtained rod-shaped nanomaterials (as shown in Fig. 13.4). This material was applied in the cathode material of aqueous zinc-ion batteries and the Zn anode also was decorated by rGO. The initial capacity was about 250 mAh/g at a current density of 300 mA/g, the retention was 79% after 1000 cycles. The authors suggested that the electrode combined with rGO could not only inhibit the dendrite growth of the anode but also had better ion deintercalation performance.

13.2.2 V6O13 Like V3O7 and V4O9, V6O13 is a mixed-valence vanadium oxide of +4 and + 5. It belongs to a monoclinic crystal system and also has a single- to double-layered

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

371

Fig. 13.4 (a) XRD data of V3O7H2O/rGO. (b) SEM image of V3O7H2O/rGO nanobelts. (c) Typical diffraction pattern and (d) representative TEM image of a single nanoribbon, revealing a width of 100–200 nm. (e) Lattice-resolved HRTEM image [from the region highlighted in (d)]. (f) SEM image and corresponding elemental mapping of (g) carbon, (h) oxygen, and (i) vanadium [109]

alternating layered structure, so it also can have the performance of deintercalation metal ions [6]. The lithium-ion deintercalation reaction is as follows: V6 O13 þ x Liþ þ x e ¼ Lix V6 O13 : Since V6O13 has the highest oxygen potential in the series of vanadium oxide (VO2.16), more lithium ions can be embedded. In the earlier study, Murphy et al. [112] obtained the maximum embedded lithium-ion content of 8 mol per mole of V6O13, with a theoretical specific capacity of about 420 mAh/g, and energy density of 890 Wh/kg was determined [113]. They are higher than the current commercial lithium-ion batteries cathode materials [33], so V6O13 is also one of the potential cathode materials of lithium-ion batteries [114]. Although V6O13 has great application potential, there are still some obstacles for the material which restrict its application. Firstly, V6O13 has a low MIT temperature [115] and it is metallic at room temperature [116]. During the discharge process, the material would be embedded with lithium ions and forms LixV6O13, which greatly reduces the electrical conductivity of the material and thus limits its performance. Secondly, in the charging and discharging process, due to the removal of lithium

372

Y. Wang et al.

ions, the material is expected to undergo an irreversible phase change but also lattice distortion. These changes would contribute to a volume change of the cell up to 13.6% [117]. This volume effect would also degrade the performance of the material. Therefore, in the present research, the morphology control, doping and surface modification of the material, has become the main strategy for the electrochemical application of the material. Tian et al. [118] prepared a flower-like NH4V4O10 precursor by hydrothermal method, and then prepared ultra-thin pre-lithified V6O13 nanosheet through low-temperature exfoliation and chemical lithium intercalation. This material was applied to the cathode material of LIBs. It could deliver the first lap capacity of 300 mA/h (100 mA/g), and reach 180mAh/g at the current density of 1000 mA/g. The retention was 98% after 150 cycles. Through the electrochemical impedance and in situ conductance test, it was found that the pre-lithiated V6O13 (12 Ω) can enhance the electrical conductivity compared with the non-pre-lithiated material (27 Ω), thus reducing the charge transfer resistance and enhancing the material performance. Ding et al. [119] synthesized a 3D nano-textile-structured V6O13 by the simple liquid redox self-assemble method. With the vanadyl(IV) sulfate [VOSO4 (H2O) where 0  x  6] as oxidant used in varying amounts, different diameter of V6O13 nano-textiles was obtained and used as cathode materials of lithium-ion batteries. The first cycle capacity was 360 mAh/g at the current density of 20 mA/g, and it delivered a capacity of 215 and 134 mAh/g at the current density of 200 mA/g and 500 mA/g, respectively. The retention was about 95% after 100 laps, and the energy density could reach about 780 Wh/kg. The authors suggested the reason for this performance was that the nano-textiles have a conductive structure, which could enhance electron and ion transport. He et al. [120] prepared a Cu-doped V6O13 by a simple hydrothermal method and obtained a belt-like CuxV5-xO13. The test result showed that the Cu0.1V5.9O13 has the best performance: the first cycle capacity was about 380 mAh/g at the rate of 0.1C and 82% retention after 50 cycles. Lai et al. [121] synthesized a flower-like V6O13 which was assembled by ultrathin nanosheet and used as the cathode material of aqueous zinc-ion batteries. It showed an outstanding performance: 395 mAh/g at the current density of 100 mA/g and 300 mAh/g at 5 A/g, after 1000 times charge and discharge; the retention is 87%. The authors thought that crystal water could enhance the intercalating chemical process. Other applications of Wadsley phase vanadium oxides in this field are shown in Table 13.5.

13.3

Vanadium-Based Oxides

Vanadium oxides also are with the adsorption of some species, such as OH[17], etc. Thus they have the potential application of electrocatalysis; however, the catalytic process is complex which includes the adsorption, reduction/oxidation, desorption steps, and so on. So it is difficult for the single vanadium oxides to be

Cathode/LIBs Cathode of AZIBs Cathode of AZIBs

Graphene composite and special morphology Special morphology

Special morphology

Special morphology

Special morphology

Carbon coated and special morphology Mn-doped

– –

V3O7H2O

V6O13NS

β-V6O13

CNPs-V6O13

Mn0.02V5.98O13

V6O13 V6O13

Cathode/LIBs

Cathode/LIBs

Cathode/LIBs

Microwave solvothermal

The cathode of ZIBs and AZIBs Cathode of AZIBs

Hydrothermal Hydrothermal and annealing

Hydrothermal

Hydrothermal

Hydrothermal

In situ electrochemical oxidation Hydrothermal

Hydrothermal

Method High-temperature mixing hydrothermal method Hydrothermal

Cathode of AZIBs

Cathode/LIBs

V3O7H2O/ rGO V3O7H2O

H2V3O8

Application Cathode/LIBs

Strategy Carbon composite and special morphology Special morphology

Sample V3O7H2O@C

Table 13.5 Some applications of Wadsley phase vanadium oxides in this field Performance 262mAh/[email protected]/g, 94% retention after 100 cycles 250mAh/[email protected] A/g, 90% retention after 50 cycles 250mAh/[email protected]/g 79% retention after 1000 cycles 275mAh/g@3A/g, 80% retention after 200 cycles (AZIB) 171mAh/g@5A/g, 85% retention after 1000 cycles 330mAh/g@42 mA/g, 70% retention after 50 cycles 335mAh/g@50 mA/g, 73% retention after 75 cycles 288mAh/g@1C, 82% retention after 100 cycles 350mAh/[email protected], 95% retention after 50 cycles 206mAh/g@10A/g after 3000 cycles 245mAh/g@4A/g, no obvious decrease after 2000 cycles

[129] [130]

[128]

[127]

[126]

[125]

[124]

[123]

[109]

[122]

References [111]

13 The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . . 373

374

Y. Wang et al.

used as an electrocatalyst. They are usually combined with other transition metal (hydro) oxides or noble metals to form the vanadium-based oxides which could be applied in energy electrocatalysis [131–133]. This is what we will introduce in this section.

13.3.1 Electrocatalysis As transition metal oxides, vanadium oxide and vanadium-based oxides also have some reports on energy electrocatalysis application; in this section, we will introduce the application in water splitting [131], oxygen reduction reaction (ORR) [132], and small molecule fuel’s oxidation catalysis [133] of vanadium-based oxide briefly. Shi et al. [131] compared the reports of VS2 and NiV-layered double hydroxide (NiVLDH) and suggested that the oxygen-group compounds of vanadium had certain effects on hydrogen evolution. Then the VOOH catalyst was synthesized by the hydrothermal method (as shown in Fig. 13.5). The catalyst was used as an electrocatalyst for water splitting under the alkaline system. The excellent performance showed that the HER overpotential was 164 mV; OER was 270 mV, and the overall reaction voltage was 1.62 V. The sustainable test was kept for 25 h, which implied excellent durability. Mu et al. [132] prepared spinel CoV2O4 by cation exchange method and used an oxygen reduction reaction electrocatalyst. The performance not only showed that under alkaline conditions, the half-wave potential could reach 0.82 V vs. RHE, which is better than cobalt oxide CoO and Co3O4 alone, almost comparable to

Fig. 13.5 VOOH hollow nanospheres (sample III) for overall water splitting in 1 M KOH. (a) Schematic of the overall water splitting setup. (b) Polarization curve of the VOOH || VOOH catalyst along with Ni foam || Ni foam and Pt/C || RuO2 for comparison. Inset shows the stability test at a constant current density of 50 mAcm2. Catalyst loading: about 0.8 mgcm2 [131]

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

375

commercial carbonized platinum, but also had better durability. Subsequently, it was applied to the cathode catalyst of zinc-air batteries. It was found that the power density could reach 380 mW/cm2, and the voltage was stable after continuous testing for 150 h under the current density of 20 mA/ cm2. The authors believed that this spinel structure could enhance the electrical conductivity of the V-V chain, thus significantly improve the activity of cobalt oxide. Amiddedin et al. [134] synthesized Pt/VOx-MWCNTs by adding the amorphous VOx to Pt-MWCNTs by microwave-assist method and applied in the electrocatalyst of methanol oxidation reaction (MOR) in the acidic system. Compared with the Pt-MWCNTs, this catalyst improved the activity of MOR obviously. The authors suggested that the existence of VOx could help to generate more -OH, and thus improved oxidation reaction of absorbed CO and enhanced the performance of catalyst eventually. Other applications of vanadium-based oxides in electrocatalysis are shown in Table 13.6.

13.4

Summary and Outlook

This chapter mainly introduced the application of vanadium (based) oxides in energy storage and electrocatalysis, mainly focusing on metal-ion batteries and water splitting, oxygen reduction reaction, and catalytic oxidation of small molecular fuels. At present, there are many reports on the preparation, performance, and energy storage or catalytic mechanism of most vanadium oxides. However, there are also some deficiencies, which are summarized as follows. The application of vanadium oxide could be catalogued in two areas. Structural, where the vanadium oxide usually possesses layered structure which intercalates metal ions. Under these conditions, the modified oxide could be used in applications related to ion insertion such as metal-ion batteries or supercapacitors. The other areas are related to the materials adsorption characteristics. Here, the adsorption of vanadium oxide to species, such as -OH, are evaluated in terms of material electrochemistry and include: 1. Prepare vanadium (based) oxides with special morphology or hierarchical structure to increase the activity area. 2. Improve the electrical conductivity of prepared materials: generally, vanadium oxide is a semiconductor or poor electrical conductivity at room temperature, which restrict its performance. Therefore, researchers often composite it with carbon- or graphene-based materials and dope with other elements to improve the electrical conductivity. 3. Composite with other metal or metal oxides and synthesize the vanadium-based oxide electrocatalyst like NiVLDH, etc. using the synergistic effect between elements or compounds. Moreover, in the research progress of vanadium oxide in this field, there are still the following shortcomings that have not been addressed:

V2O5/Ni (OH)2 Pt/V2O5-C

CoV1.5Fe0.5O4

Pt loaded

Ni(OH)2 composite

Special morphology Ni, NiS composite and special morphology Ni composite, hierarchical structure NiS composite, hierarchical structure Fe doped

V10O24nH2O VOx / Ni3S2@NF Ni/V2O3

VOx/NiS/Ni

Strategy Ni composite

Sample Ni/V2O3

Water splitting MOR (acidic)

OER, ORR

OER

Microwave assist, and reduction

Thermal decomposition Hydrothermal

Hydrothermal, calcination Hydrothermal

HER

HER OER

Method Hydrothermal, calcination Hydro-thermal Hydrothermal

Application HER

Table 13.6 Some applications of vanadium-based oxides electrocatalysis

η ¼ 330 mV (50 mA/cm2); Tafel slope ¼ 121 mV/dec @1MKOH η ~ 300 mV, Tafel slope ¼ 38 mV/dec (OER), E1/ 2 ~ 0.7 V vs. RHE η ~ 40 mV, Tafel slope ¼ 44 mV/dec (HER@1 M KOH); 1.59 V for water splitting@6 M KOH 17.4 mA/cm2 at forward sweep peak(Pt@C:12.25)

η ¼ 118 mV; Tafel slope ¼ 101 mV/dec @1MKOH η ¼ 358 mV (100 mA/cm2), Tafel slope ¼ 82 mV/dec @1MKOH η ¼ 22 mV (20 mA/cm2); @1MKOH

Performance η ¼ 44 mV; Tafel slope ¼ 38 mV/dec @1MKOH

[133]

[139]

[138]

[137]

[17]

[136] [63]

References [135]

376 Y. Wang et al.

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

377

1. The phase transition problem of vanadium oxide which is used in electrode materials of metal-ion batteries has not been solved, which greatly restricts the large-scale application of vanadium oxides. 2. The influence of the change of valence of vanadium in vanadium oxide on catalysis when they are used for electrocatalysts. 3. The development of other mixed-valence vanadium oxides, such as V4O9 in Wadsley phase vanadium oxides, and Magnéli phase vanadium oxides are still rarely reported. In particular, a series of studies on Magnéli phase vanadium oxides are still limited to thermodynamic synthesis. The electrochemical applications for the preparation of special morphology have not yet been reported. Acknowledgments F. L. and Y. W. sincerely thank the National Nature Science Foundation of China (NSFC 52074016) for the financial support. J. L. Liu and S. B. would also like to acknowledge the support from the National Science Foundation (NSF-MRI, CBET 0821370) and R. Welch Foundation (AC-0006) from the Texas A&M University-Kingsville. Author Contribution F. L. and Y. W. collectively conceived this topic. J. L. Liu completed the literature survey and edited the manuscript. S. B. checked the data and the hkl assignment(s). Y. W. generated the data table and wrote the first draft of the paper. All authors reviewed the final edit and approved the final version of the manuscript.

References 1. Wang Y, Takahashi K, Lee K, Cao G (2006) Nanostructured vanadium oxide electrodes for enhanced lithium-ion intercalation. Adv Funct Mater 16(9):1133–1144 2. Chao D, Zhu C, Xia X, Liu J, Zhang X, Wang J, Liang P, Lin J, Zhang H, Shen ZX, Fan HJ (2015) Graphene quantum dots coated VO2 arrays for highly durable electrodes for Li and Na ion batteries. Nano Lett 15(1):565–573 3. Hu P, Yan M, Zhu T, Wang X, Wei X, Li J, Zhou L, Li Z, Chen L, Mai L (2017) Zn/V2O5 aqueous hybrid-ion battery with high voltage platform and long cycle life. ACS Appl Mater Interfaces 9(49):42717–42722 4. Cao AM, Hu JS, Liang HP, Wan LJ (2005) Self-assembled vanadium pentoxide (V2O5) hollow microspheres from nanorods and their application in lithium-ion batteries. Angew Chem Int Ed 44(28):4391–4395 5. Mai L, Xu L, Han C, Xu X, Luo Y, Zhao S, Zhao Y (2010) Electrospun ultralong hierarchical vanadium oxide nanowires with high performance for lithium-ion batteries. Nano Lett 10 (11):4750–4755 6. Chernova NA, Roppolo M, Dillon AC, Whittingham MS (2009) Layered vanadium and molybdenum oxides: batteries and electrochromics. J Mater Chem 19(17):2526 7. Pan A, Zhang J-G, Nie Z, Cao G, Arey BW, Li G, Liang S-q, Liu J (2010) Facile synthesized nanorod-structured vanadium pentoxide for high-rate lithium batteries. J Mater Chem 20 (41):9193–9199 8. Chao D, Xia X, Liu J, Fan Z, Ng CF, Lin J, Zhang H, Shen ZX, Fan HJ (2014) A V2O5/ conductive-polymer core/shell nanobelt array on three-dimensional graphite foam: a high-rate, ultrastable, and freestanding cathode for lithium-ion batteries. Adv Mater 26(33):5794–5800 9. Zeng L, Zheng C, Xi J, Fei H, Wei M (2013) Composites of V2O3-ordered mesoporous carbon as anode materials for lithium-ion batteries. Carbon 62:382–388

378

Y. Wang et al.

10. Jiang L, Qu Y, Ren Z, Yu P, Zhao D, Zhou W, Wang L, Fu H (2015) In situ carbon-coated yolk-shell V2O3 microspheres for Lithium-ion batteries. ACS Appl Mater Interfaces 7 (3):1595–1601 11. Wang Y, Zhang HJ, Admar AS, Luo J, Wong CC, Borgna A, Lin J (2012) Improved cyclability of lithium-ion battery anode using encapsulated V2O3 nanostructures in wellgraphitized carbon fiber. RSC Adv 2(13):5748–5753 12. Li X, Fu J, Pan Z, Su J, Xu J, Gao B, Peng X, Wang L, Zhang X, Chu PK (2016) Peapod-like V2O3 nanorods encapsulated into carbon as binder-free and flexible electrodes in lithium-ion batteries. J Power Sources 331:58–66 13. Tsang C, Manthiram A (1997) Synthesis of nanocrystalline VO2 and its electrochemical behavior in lithium batteries. J Electrochem Soc 144(2):520–524 14. Liu H, Wang Y, Wang K, Hosono E, Zhou H (2009) Design and synthesis of a novel nanothorn VO2(B) hollow microsphere and their application in lithium-ion batteries. J Mater Chem 19(18):2835–2840 15. Liu P, Zhu K, Gao Y, Luo H, Lu L (2017) Recent progress in the applications of vanadiumbased oxides on energy storage: from low-dimensional nanomaterials synthesis to 3D micro/ nano-structures and free-standing electrodes fabrication. Adv Energy Mater 7(23):1700547 16. Xing X, Liu R, Cao K, Kaiser U, Streb C (2019) Transition-metal oxides/carbides@carbon nanotube composites as multifunctional electrocatalysts for challenging oxidations and reductions. Chem Eur J 25(47):11098–11104 17. Chen Y, Rao Y, Wang R, Yu Y, Li Q, Bao S, Xu M, Yue Q, Zhang Y, Kang Y (2020) Interfacial engineering of Ni/V2O3 for hydrogen evolution reaction. Nano Res 13 (9):2407–2412 18. Tepavcevic S, Xiong H, Stamenkovic VR, Zuo X, Balasubramanian M, Prakapenka VB, Johnson CS, Rajh T (2012) Nanostructured bilayered vanadium oxide electrodes for rechargeable sodium-ion batteries. ACS Nano 6(1):530–538 19. Zhu Y-H, Zhang Q, Yang X, Zhao E-Y, Sun T, Zhang X-B, Wang S, Yu X-Q, Yan J-M, Jiang Q (2019) Reconstructed orthorhombic V2O5 polyhedra for fast ion diffusion in K-ion batteries. Chem 5(1):168–179 20. Deng L, Niu X, Ma G, Yang Z, Zeng L, Zhu Y, Guo L (2018) Layered potassium vanadate K0.5V2O5 as a cathode material for nonaqueous potassium ion batteries. Adv Funct Mater 28 (49) 21. Yan M, He P, Chen Y, Wang S, Wei Q, Zhao K, Xu X, An Q, Shuang Y, Shao Y, Mueller KT, Mai L, Liu J, Yang J (2018) Water-lubricated Intercalation in V2O5 center dot nH(2)O for highcapacity and high-rate aqueous rechargeable zinc batteries. Adv Mater 30(1) 22. Yang Y, Tang Y, Fang G, Shan L, Guo J, Zhang W, Wang C, Wang L, Zhou J, Liang S (2018) Li+ intercalated V2O5 center dot nH(2)O with enlarged layer spacing and fast ion diffusion as an aqueous zinc-ion battery cathode. Energy Environ Sci 11(11):3157–3162 23. Zhou J, Shan L, Wu Z, Guo X, Fang G, Liang S (2018) Investigation of V2O5 as a low-cost rechargeable aqueous zinc ion battery cathode. Chem Commun 54(35):4457–4460 24. Wang Z, Yu K, Gong S, Du E, Zhu Z (2020) Vanadium based carbide-oxide heterogeneous V2O5@V2C nanotube arrays for high-rate and long-life lithium-sulfur batteries. Nanoscale 12 (36):18950–18964 25. Li W, Hicks-Garner J, Wang J, Liu J, Gross AF, Sherman E, Graetz J, Vajo JJ, Liu P (2014) V2O5 polysulfide anion barrier for long-lived Li-S batteries. Chem Mater 26(11):3403–3410 26. Guo Y, Zhang Y, Zhang Y, Xiang M, Wu H, Liu H, Dou S (2018) Interwoven V2O5 nanowire/ graphene nanoscroll hybrid assembled as efficient polysulfide-trapping-conversion interlayer for long-life lithium-sulfur batteries. J Mater Chem A 6(40):19358–19370 27. Chen Z, Augustyn V, Wen J, Zhang Y, Shen M, Dunn B, Lu Y (2011) High-performance supercapacitors based on intertwined CNT/V2O5 nanowire nanocomposites. Adv Mater 23 (6):791 28. Qu Q, Zhu Y, Gao X, Wu Y (2012) Core-shell structure of polypyrrole grown on V2O5 nanoribbon as high-performance anode material for supercapacitors. Adv Energy Mater 2 (8):950–955

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

379

29. Wang L, Shu T, Guo S, Lu Y, Li M, Nzabahimana J, Hu X (2020) Fabricating strongly coupled V2O5@PEDOT nanobelts/graphene hybrid films with high areal capacitance and facile transferability for transparent solid-state supercapacitors. Energy Storage Materials 27:150–158 30. Pan AQ, Wu HB, Zhang L, Lou XW (2013) Uniform V2O5 nanosheet-assembled hollow micro flowers with excellent lithium storage properties. Energy Environ Sci 6(5):1476–1479 31. Liu J, Xia H, Xue D, Lu L (2009) Double-shelled nanocapsules of V2O5-based composites as high-performance anode and cathode materials for Li-Ion batteries. J Am Chem Soc 131 (34):12086 32. Yue Y, Liang H (2017) Micro- and Nano-structured vanadium pentoxide (V2O5) for electrodes of lithium-ion batteries. Adv Energy Mater 7(17):1602545 33. Liu X, Zeng J, Yang H, Zhou K, Pan D (2018) V2O5-based nanomaterials: synthesis and their applications. RSC Adv 8(8):4014–4031 34. Zhu Y, Yang M, Huang Q, Wang D, Yu R, Wang J, Zheng Z, Wang D (2020) V2O5 textile cathodes with high capacity and stability for flexible lithium-ion batteries. Adv Mater 32(7) 35. Ji Y, Fang D, Wang C, Zhou Z, Luo Z, Huang J, Yi J (2018) Cobalt-doped V2O5 nanowire arrays on Ti foil for enhanced lithium-ion storage. J Alloys Compd 742:567–576 36. Yu H, Zeng J, Hao W, Zhou P, Wen X (2018) Mo-doped V2O5 hierarchical nanorod/ nanoparticle core/shell porous microspheres with improved performance for the cathode of the lithium-ion battery. J Nanoparticle Res 20(5) 37. Wang L, Wang Y, Zhu X, Zhao Y (2019) Conventional synthesis of V2O5/graphite composites with enhanced lithium-ion storage. J Alloys Compd 792:418–423 38. Wang Z, Huang H, Zeng L, Wang Y, Lv L, Dai C, Mao W, Chen X, Fang D (2019) In-operando deformation studies on the mechano-electrochemical mechanism in free-standing MWCNTs/V2O5 lithium-ion battery electrode. Electrochim Acta 305:101–115 39. Cao F, Pan GX, Zhang YJ (2018) Synthesis of V2O5/C core/shell arrays on graphene foam for electrochemical energy storage. Mater Res Bull 99:479–484 40. Chen M, Liang X, Yin J, Chen Q, Xia X (2018) Graphene foam supported V2O5/N-C core/ shell arrays as advanced cathode for lithium-ion storage. J Alloys Compd 735:2022–2029 41. Zhang X, Liu X, Yang C, Li N, Ji T, Yan K, Zhu B, Yin J, Zhao J, Li Y (2019) A V2O5nanosheets-coated hard carbon fiber fabric as high-performance anode for sodium-ion battery. Surface Coatings Technol 358:661–666 42. Vishnuprakash P, Nithya C, Premalatha M (2019) Exploration of V2O5 nanorod@rGO heterostructure as potential cathode material for potassium-ion batteries. Electrochim Acta 309:234–241 43. Wang X, Ma L, Sun J (2019) Vanadium pentoxide Nanosheets in-situ spaced with acetylene black as cathodes for high-performance zinc-lon batteries. ACS Appl Mater Interfaces 11 (44):41297–41303 44. Cao L, Kou L, Li J, Huang J, Yang J, Wang Y (2018) Nitrogen-doped carbon-coated V2O5 nanocomposite as cathode materials for lithium-ion battery. J Mater Sci 53(14):10270–10279 45. Gao X-T, Liu Y-T, Zhu X-D, Yan D-J, Wang C, Feng Y-J, Sun K-N (2018) V2O5 nanoparticles confined in three-dimensionally organized, porous nitrogen-doped graphene frameworks: flexible and free-standing cathodes for high-performance lithium storage. Carbon 140:218–226 46. Dou Y, Liang X, Gao G, Wu G (2018) Template-free synthesis of porous V2O5 yolk-shell microspheres as cathode materials for lithium-ion batteries. J Alloys Compd 735:109–116 47. Yue Y, Juarez-Robles D, Mukherjee PP, Liang H (2018) Superhierarchical nickel-vanadia nanocomposites for Lithium storage. Acs Applied Energy Materials 1(5):2056–2066 48. Tian B, Tang W, Su C, Li Y (2018) Reticular V2O5 center dot 0.6H(2)O Xerogel as cathode for rechargeable potassium ion batteries. ACS Appl Mater Interfaces 10(1):642–650 49. Ke Y, Wang S, Liu G, Li M, White TJ, Long Y (2018) Vanadium dioxide: the multistimuli responsive material and its applications. Small 14(39):1802025

380

Y. Wang et al.

50. Driscoll T, Kim H-T, Chae B-G, Kim B-J, Lee Y-W, Jokerst NM, Palit S, Smith DR, Di Ventra M, Basov DN (2009) Memory metamaterials. Science 325(5947):1518–1521 51. Seo M, Kyoung J, Park H, Koo S, Kim H-S, Bernien H, Kim BJ, Choe JH, Ahn YH, Kim H-T, Park N, Park QH, Ahn K, Kim D-s (2010) Active terahertz nanoantennas based on VO2 phase transition. Nano Lett 10(6):2064–2068 52. Liu M, Hwang HY, Tao H, Strikwerda AC, Fan K, Keiser GR, Sternbach AJ, West KG, Kittiwatanakul S, Lu J, Wolf SA, Omenetto FG, Zhang X, Nelson KA, Averitt RD (2012) Terahertz-field-induced insulator-to-metal transition in vanadium dioxide metamaterial. Nature 487(7407):345–348 53. Zhu Z, Evans PG, Haglund RF Jr, Valentine JG (2017) Dynamically reconfigurable meta device employing nanostructured phase-change materials. Nano Lett 17(8):4881–4885 54. Yang S, Gong Y, Liu Z, Zhan L, Hashim DP, Ma L, Vajtai R, Ajayan PM (2013) Bottom-up approach toward single-crystalline VO2-graphene ribbons as cathodes for ultrafast lithium storage. Nano Lett 13(4):1596–1601 55. Wang W, Jiang B, Xiong W, Sun H, Lin Z, Hu L, Tu J, Hou J, Zhu H, Jiao S (2013) A new cathode material for super-valent battery based on aluminum ion intercalation and deintercalation. Sci Rep 3 56. Wang Q, Xu J, Zhang W, Mao M, Wei Z, Wang L, Cui C, Zhu Y, Ma J (2018) Research progress on vanadium-based cathode materials for sodium-ion batteries. J Mater Chem A 6 (19):8815–8838 57. Ding J, Du Z, Gu L, Li B, Wang L, Wang S, Gong Y, Yang S (2018) Ultrafast Zn2+ intercalation and deintercalation in vanadium dioxide. Adv Mater 30(26) 58. Song Y, Zhao W, Kong L, Zhang L, Zhu X, Shao Y, Ding F, Zhang Q, Sun J, Liu Z (2018) Synchronous immobilization and conversion of polysulfides on a VO2-VN binary host targeting high sulfur load Li-S batteries. Energy Environ Sci 11(9):2620–2630 59. Wang S, Liao J, Yang X, Liang J, Sun Q, Liang J, Zhao F, Koo A, Kong F, Yao Y, Gao X, Wu M, Yang S-Z, Li R, Sun X (2019) Designing a highly efficient polysulfide conversion catalyst with paramontroseite for high-performance and long-life lithium-sulfur batteries. Nano Energy 57:230–240 60. Wang H, Yi H, Chen X, Wang X (2014) One-step strategy to three-dimensional graphene/VO2 nanobelt composite hydrogels for high-performance supercapacitors. J Mater Chem A 2 (4):1165–1173 61. Zheng J, Zhang Y, Wang Q, Jiang H, Liu Y, Lv T, Meng C (2018) Hydrothermal encapsulation of VO2(a) nanorods in amorphous carbon by carbonization of glucose for energy storage devices. Dalton Trans 47(2):452–464 62. Huang D, Li S, Zhang X, Luo Y, Xiao J, Chen H (2018) A novel method to significantly boost the electrocatalytic activity of carbon cloth for oxygen evolution reaction. Carbon 129:468–475 63. Niu Y, Li W, Wu X, Feng B, Yu Y, Hu W, Li CM (2019) Amorphous nickel sulfide nanosheets with embedded vanadium oxide nanocrystals on nickel foam for efficient electrochemical water oxidation. J Mater Chem A 7(17):10534–10542 64. Xia C, Lin Z, Zhou Y, Zhao C, Liang H, Rozier P, Wang Z, Alshareef HN (2018) Large Intercalation Pseudocapacitance in 2D VO2 (B): Breaking through the Kinetic Barrier. Adv Mater 30(40) 65. Niu C (2018) Al-doped VO2(B) nanobelts as a cathode material with enhanced electrochemical properties for lithium-ion batteries. Funct Mater Lett 11(04):1850068 66. Liu X, Li B, Yao L, Wang P, Kong X, Zhang Y (2019) Flower-like VO(2)(B)@C structure: high rate capacity and stability as lithium-ion batteries. J Nanosci Nanotechnol 19 (7):4052–4057 67. Ma Y, Li W, Ji S, Zhou H, Li R, Li N, Yao H, Cao X, Jin P (2017) Fabrication of bristlegrasslike VO2 (B)-ZnO hetero-architectures as anode materials for lithium-ion batteries. Mater Res Express 4(8)

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

381

68. Hu X, Zhao Z, Wang L, Li J, Wang C, Zhao Y, Jin H (2019) VO2 (a)/graphene nanostructure: stand up to Na ion intercalation/deintercalation for enhanced electrochemical performance as a Na-ion battery cathode. Electrochim Acta 293:97–104 69. Zhao L, Wei Q, Huang Y, Luo R, Xie M, Li L, Mai L, Wu F, Chen R (2019) Pseudocapacitive graphene-wrapped porous VO2 microspheres for ultrastable and ultrahigh-rate sodium-ion storage. ChemElectroChem 6(5):1400–1406 70. Ren G, Zhang R, Fan Z (2018) VO2 nanoparticles on edge oriented graphene foam for high rate lithium-ion batteries and supercapacitors. Appl Surf Sci 441:466–473 71. Liu P, Xu Y, Zhu K, Bian K, Wang J, Sun X, Gao Y, Luo H, Lu L, Liu J (2017) Ultrathin VO2 nanosheets self-assembled into 3D micro/nano-structured hierarchical porous sponge-like micro-bundles for long-life and high-rate Li-ion batteries. J Mater Chem A 5(18):8307–8316 72. Jayalakshmi T, Nagaraju K, Nagaraju G (2018) Enhanced lithium storage of mesoporous vanadium dioxide(B) nanorods by reduced graphene oxide support. J Energy Chem 27 (1):183–189 73. Nizar SAS, Ramar V, Venkatesan T, Balaya P, Valiyaveettil S (2018) Enhanced electrochemical performance of W incorporated VO2 nanocomposite cathode material for lithium battery application. Electrochim Acta 282:480–489 74. Chen M, Liang X, Wang F, Xie D, Pan G, Xia X (2019) Self-supported VO2 arrays decorated with N-doped carbon as an advanced cathode for lithium-ion storage. J Mater Chem A 7 (12):6644–6650 75. Xiao B, Zhang B, Zheng J-c, Tang L-b, An C-s, He Z-j, Tong H, Yu W-j (2018) Nano-micro structure VO2/CNTs composite as a potential anode material for lithium-ion batteries. Ceram Int 44(11):13113–13121 76. Li W, Huang J, Cao L, Li X, Chen S, Feng L (2020) Polycrystalline VO2(M) with welldispersed crystalline zones for enhanced electroactivity of lithium-ion batteries. J Alloys Compd 812 77. Guo W, Wang Y, Li Q, Wang D, Zhang F, Yang Y, Yu Y (2018) SnO2@C@VO2 composite hollow nanospheres as an anode material for lithium-ion batteries. ACS Appl Mater Interfaces 10(17):14993–15000 78. Su H, Chang K, Ma Y, Yang D, Wang C, Zhang K, Li D, Zhang J (2019) Hierarchical flowerlike structures composed of cross-shaped vanadium dioxide nanobelts as superior performance anode for lithium and sodium ions batteries. Appl Surf Sci 480:882–887 79. Schwingenschlogl U, Eyert V (2004) The vanadium Magneli phasesV(n)O(2n-1). Ann Phys 13 (9):475–510 80. Sun Y, Jiang S, Bi W, Wu C, Xie Y (2011) Highly ordered lamellar V2O3-based hybrid nanorods towards superior aqueous lithium-ion battery performance. J Power Sources 196 (20):8644–8650 81. Han C, Liu F, Liu J, Li Q, Meng J, Shao B, He Q, Wang X, Liu Z, Mai L (2018) Facile template-free synthesis of uniform carbon-confined V2O3 hollow spheres for stable and fast lithium storage. J Mater Chem A 6(15):6220–6224 82. Luo H, Wang B, Wang F, Yang J, Wu F, Ning Y, Zhou Y, Wang D, Liu H, Dou S (2020) Anodic oxidation strategy toward structure-optimized V2O3 cathode via electrolyte regulation for Zn-ion storage. ACS Nano 14(6):7328–7337 83. Meng L, Guo R, Li F, Ma Y, Peng J, Li T, Luo Y, Li Y, Sun X (2020) Facile synthesis of flocklike V2O3/C with improved electrochemical performance as an anode material for Li-ion batteries. Energy Technol 8(3) 84. Xia X, Chao D, Zhang Y, Zhan J, Zhong Y, Wang X, Wang Y, Shen ZX, Tu J, Fan HJ (2016) Generic synthesis of carbon nanotube branches on metal oxide arrays exhibiting stable highrate and long-cycle sodium-ion storage. Small 12(22):3048–3058 85. Zhang L, Liu Y, Zhao Z, Jiang P, Zhang T, Li M, Pan S, Tang T, Wu T, Liu P, Hou Y, Lu H (2020) Enhanced polysulfide regulation via porous catalytic V2O3/V8C7 heterostructures derived from metal-organic frameworks toward high-performance Li-S batteries. ACS Nano 14(7):8495–8507

382

Y. Wang et al.

86. Zhang R, Han J, Zheng B, Shi X, Asiri AM, Sun X (2019) Metal-organic framework-derived shuttle-like V2O3/C for electrocatalytic N-2 reduction under ambient conditions. Inorg Chem Front 6(2):391–395 87. Xu H, Liu L, Gao J, Du P, Fang G, Qiu H-J (2019) Hierarchical nanoporous V2O3 Nanosheets anchored with alloy nanoparticles for efficient electrocatalysis. ACS Appl Mater Interfaces 11 (42):38746–38753 88. Zhang XY, Yuan H, Mao F, Wen CF, Zheng LR, Liu PF, Yang HG (2019) Boosting alkaline hydrogen evolution electrocatalysis over metallic nickel sites through synergistic coupling with vanadium Sesquioxide. ChemSusChem 12(23):5063–5069 89. Pan X, Ren G, Hoque MNF, Bayne S, Zhu K, Fan Z (2014) Fast Supercapacitors Based on Graphene-Bridged V2O3/VOx Core-Shell Nanostructure Electrodes with a Power Density of 1 MW kg(1). Adv Mater Interfaces 1(9) 90. Zheng J, Zhang Y, Jing X, Liu X, Hu T, Lv T, Zhang S, Meng C (2017) Synthesis of amorphous carbon coated on V2O3 core-shell composites for enhancing the electrochemical properties of V2O3 as a supercapacitor electrode. Colloids Surf A Physicochem Eng Asp 518:188–196 91. Zhang D, Li G, Li B, Fan J, Liu X, Chen D, Li L (2019) A facile strategy to fabricate V2O3/ porous N-doped carbon nanosheet framework as high-performance anode for lithium-ion batteries. J Alloys Compd 789:288–294 92. Li Y, Zhang S, Wang S, Leng J, Jiang C, Ren X, Zhang Z, Yang Y, Tang Z (2019) A multishelled V2O3/C composite with an overall coupled carbon scaffold enabling ultrafast and stable lithium/sodium storage. J Mater Chem A 7(33):19234–19240 93. Gao S, Zhang D, Zhu K, Tang JA, Gao Z, Wei Y, Chen G, Gao Y (2017) Flexible V2O3/ carbon nano-felts as a free-standing electrode for high-performance lithium-ion batteries. J Alloys Compd 702:13–19 94. Yang L, Zhang Z, Xia L, Zhao Y, Li F, Zhang X, Wei J, Zhou Z (2019) Integrated insights into Na+ storage mechanism and electrochemical kinetics of ultrafine V2O3/S and N co-doped rGO composites as anodes for sodium-ion batteries. J Mater Chem A 7(39):22429–22435 95. Tan S, Jiang Y, Wei Q, Huang Q, Dai Y, Xiong F, Li Q, An Q, Xu X, Zhu Z, Bai X, Mai L (2018) Multidimensional synergistic nanoarchitecture exhibiting highly stable and ultrafast sodium-ion storage. Adv Mater 30(18):1707122 96. Wang J, Liu Z, Yang W, Han L, Wei M (2018) A one-step synthesis of porous V2O3@C hollow spheres as a high-performance anode for lithium-ion batteries. Chem Commun 54 (53):7346–7349 97. Liu X, Zhang D, Li G, Xue C, Ding J, Li B, Chen D, Li L (2018) In situ synthesis of V2O3 nanorods anchored on reduced graphene oxide as a high-performance lithium-ion battery anode. Chemistryselect 3(43):12108–12112 98. Xun L, Gao S, Xu Y, Cheng X, Zhang X, Zhao H, Huo L (2018) Synthesis of dandelion-like V2O3/C composite with bicontinuous 3D hierarchical structures as an anode for highperformance lithium-ion batteries. Ceram Int 44(12):14128–14135 99. Gou W, Kong X, Wang Y, Ai Y, Liang S, Pan A, Cao G (2019) Yolk-shell structured V2O3 microspheres wrapped in N, S co-doped carbon as pea-pod nanofibers for high-capacity lithium-ion batteries. Chem Eng J 374:545–553 100. Zhang J, Li Q, Liao Z, Wang L, Xu J, Ren X, Gao B, Chu PK, Huo K (2018) In situ synthesis of V2O3-intercalated N-doped graphene nanobelts from VOx-amine hybrid as highperformance anode material for alkali-ion batteries. ChemElectroChem 5(10):1387–1393 101. Cai Y, Fang G, Zhou J, Liu S, Luo Z, Pan A, Cao G, Liang S (2017) Metal-organic frameworkderived porous shuttle-like vanadium oxides for sodium-ion battery application. Nano Res 11 (1):449–463 102. Song Y, Zhao W, Wei N, Zhang L, Ding F, Liu Z, Sun J (2018) In-situ PECVD-enabled graphene-V2O3 hybrid host for lithium-sulfur batteries. Nano Energy 53:432–439

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

383

103. Jin T, Li H, Li Y, Jiao L, Chen J (2018) Intercalation pseudocapacitance inflexible and selfstanding V2O3 porous nanofibers for high-rate and ultra-stable K ion storage. Nano Energy 50:462–467 104. Ding Y, Peng Y, Chen S, Zhang X, Li Z, Zhu L, Mo L-E, Hu L (2019) Hierarchical porous metallic V2O3@C for advanced aqueous zinc-ion batteries. ACS Appl Mater Interfaces 11 (47):44109–44117 105. Zhao D, Zhu Q, Chen D, Li X, Yu Y, Huang X (2018) Nest-like V3O7 self-assembled by porous nanowires as an anode supercapacitor material and its performance optimization through bonding with N-doped carbon. J Mater Chem A 6(34):16475–16484 106. Li H-Y, Wei C, Wang L, Zuo Q-S, Li X, Xie B (2015) Hierarchical vanadium oxide microspheres forming from hyperbranched nanoribbons as remarkably high-performance electrode materials for supercapacitors. J Mater Chem A 3(45):22892–22901 107. Xu N, Ma X, Wang M, Qian T, Liang J, Yang W, Wang Y, Hu J, Yan C (2016) Stationary full Li-ion batteries with interlayer-expanded V6O13 cathodes and lithiated graphite anodes. Electrochim Acta 203:171–177 108. Pang H, Dong Y, Ting SL, Lu J, Li CM, Kim DH, Chen P (2013) 2D single- or double-layered vanadium oxide nanosheet assembled 3D micro flowers: controlled synthesis, growth mechanism, and applications. Nanoscale 5(17):7790–7794 109. Shen C, Li X, Li N, Xie K, Wang JG, Liu X, Wei B (2018) Graphene-boosted, highperformance aqueous Zn-ion battery. ACS Appl Mater Interfaces 10(30):25446–25453 110. Li L, Liu P, Zhu K, Wang J, Liu J, Qiu J (2015) A general and simple method to synthesize well-crystallized nanostructured vanadium oxides for high-performance Li-ion batteries. J Mater Chem A 3(18):9385–9389 111. Liu P, Bian K, Zhu K, Xu Y, Gao Y, Luo H, Lu L, Wang J, Liu J, Tai G (2017) Ultrathin nanoribbons of in situ carbon-coated V3O7H2O for high-energy and long-life Li-ion batteries: synthesis, electrochemical performance, and charge–discharge behavior. ACS Appl Mater Interfaces 9(20):17002–17012 112. Murphy DW, Christian PA, Disalvo FJ, Carides JN, Waszczak JV (1981) Lithium incorporation by V6O13 and related vanadium (+4, +5) oxide cathode materials. J Electrochem Soc 128 (10):2053–2060 113. West K, Zachauchristiansen B, Jacobsen T, Atlung S (1985) V6O13 as cathode material for lithium cells. J Power Sources 14(1–3):235–245 114. West K, Zachauchristiansen B, Ostergard MJL, Jacobsen T (1987) Vanadium-oxides as electrode materials for rechargeable lithium cells. J Power Sources 20(1–2):165–172 115. Choi S, Son J, Oh J, Lee J-H, Jang JH, Lee S (2019) Sharp contrast in the electrical and optical properties of vanadium Wadsley (VmO2m+1, m > 1) epitaxial films selectively stabilized on (111)-oriented Y-stabilized ZrO2. Phys Rev Mater 3(6) 116. West K, Zachauchristiansen B, Jacobsen T (1983) Electrochemical properties of non-stoichiometric V6O13. Electrochim Acta 28(12):1829–1833 117. Meng W, Pigliapochi R, Bayley PM, Pecher O, Gaultois MW, Seymour ID, Liang H-P, Xu W, Wiaderek KM, Chapman KW, Grey CP (2017) Unraveling the complex delithiation and lithiation mechanisms of the high capacity cathode material V6O13. Chem Mater 29 (13):5513–5524 118. Tian X, Xu X, He L, Wei Q, Yan M, Xu L, Zhao Y, Yang C, Mai L (2014) Ultrathin pre-lithiated V6O13 nanosheet cathodes with enhanced electrical transport and cyclability. J Power Sources 255:235–241 119. Ding Y-L, Wen Y, Wu C, van Aken PA, Maier J, Yu Y (2015) 3D V6O13 Nanotextiles assembled from interconnected nanogrooves as cathode materials for high-energy lithium-ion batteries. Nano Lett 15(2):1388–1394 120. He J, Long F, Peng D, Wu X, Mo S, Zou Z (2017) Ribbon-like Cu doped V6O13 as cathode material for high-performance lithium-ion batteries. J Wuhan Univ Technol Mater Sci Ed 32 (6):1397–1401

384

Y. Wang et al.

121. Lai J, Zhu H, Zhu X, Koritala H, Wang Y (2019) Interlayer-expanded V6O13nH2O architecture constructed for an advanced rechargeable aqueous zinc-ion battery. ACS Appl Energy Mater 2(3):1988–1996 122. Sarkar S, Bhowmik A, Pan J, Bharadwaj MD, Mitra S (2016) Preparation, structure study, and electrochemistry of layered H2V3O8 materials: high capacity lithium-ion battery cathode. J Power Sources 329:179–189 123. Kundu D, Hosseini Vajargah S, Wan L, Adams B, Prendergast D, Nazar LF (2018) Aqueous vs. nonaqueous Zn-ion batteries: consequences of the desolvation penalty at the interface. Energy Environ Sci 11(4):881–892 124. Cao Z, Chu H, Zhang H, Ge Y, Clemente R, Dong P, Wang L, Shen J, Ye M, Ajayan PM (2019) An in situ electrochemical oxidation strategy for the formation of nanogrid-shaped V3O7H2O with enhanced zinc storage properties. J Mater Chem A 7(44):25262–25267 125. Zou Z, Cheng H, He J, Long F, Wu Y, Yan Z, Chen H (2014) V6O13 Nanosheets synthesized from ethanol-aqueous solutions as high energy cathode material for lithium-ion batteries. Electrochim Acta 135:175–180 126. Zhang Y, Meng C (2015) Facile one-pot hydrothermal synthesis of belt-like β-V6O13 with rectangular cross-sections for Li-ion battery application. Mater Lett 160:404–407 127. Li S, Zou Z, Wu X, Zhang Y (2019) Solvothermal preparation of carbon-coated V6O13 nanocomposite as cathode material for lithium-ion battery. J Electroanal Chem 846:113173 128. He J, Long F, Zou Z, Wang W, Fu Z (2014) Hydrothermal synthesis and electrochemical performance of Mn-doped V6O13 as cathode material for lithium-ion battery. Ionics 21 (4):995–1001 129. Shan L, Zhou J, Zhang W, Xia C, Guo S, Ma X, Fang G, Wu X, Liang S (2019) Highly reversible phase transition endows V6O13 with enhanced performance as aqueous zinc-ion battery Cathode. Energy Technol 7(6):1900022 130. Shin J, Choi DS, Lee HJ, Jung Y, Choi JW (2019) Hydrated intercalation for highperformance aqueous Zinc ion batteries. Adv Energy Mater 9(14):1900083 131. Shi H, Liang H, Ming F, Wang Z (2017) Efficient overall water-splitting electrocatalysis using lepidocrocite VOOH hollow nanospheres. Angew Chem Int Ed 56(2):573–577 132. Mu C, Mao J, Guo J, Guo Q, Li Z, Qin W, Hu Z, Davey K, Ling T, Qiao SZ (2020) Rational design of spinel cobalt vanadate oxide Co2VO4 for superior electrocatalysis. Adv Mater 32 (10):1907168 133. Maiyalagan T, Khan FN (2009) Electrochemical oxidation of methanol on Pt/V2O5–C composite catalysts. Catal Commun 10(5):433–436 134. Nouralishahi A, Khodadadi AA, Rashidi AM, Mortazavi Y (2013) Vanadium oxide decorated carbon nanotubes as promising support of Pt nanoparticles for methanol electro-oxidation reaction. J Colloid Interface Sci 393:291–299 135. Zhou P, Lv X, Gao Y, Cui Z, Liu Y, Wang Z, Wang P, Zheng Z, Dai Y, Huang B (2019) Enhanced electrocatalytic HER performance of non-noble metal nickel by the introduction of divanadium trioxide. Electrochim Acta 320:134535 136. Dey KK, Jha S, Kumar A, Gupta G, Srivastava AK, Ingole PP (2019) Layered vanadium oxide nanofibers as impressive electrocatalyst for hydrogen evolution reaction in acidic medium. Electrochim Acta 312:89–99 137. Chai Y-M, Zhang X-Y, Lin J-H, Qin J-F, Liu Z-Z, Xie J-Y, Guo B-Y, Yang Z, Dong B (2019) Three-dimensional VOx/NiS/NF nanosheets as an efficient electrocatalyst for oxygen evolution reaction. Int J Hydrog Energy 44(21):10156–10162 138. Chakrapani K, Bendt G, Hajiyani H, Lunkenbein T, Greiner MT, Masliuk L, Salamon S, Landers J, Schlögl R, Wende H, Pentcheva R, Schulz S, Behrens M (2018) The role of composition of uniform and highly dispersed cobalt vanadium iron spinel nanocrystals for oxygen electrocatalysis. ACS Catal 8(2):1259–1267 139. Meena A, Ha M, Chandrasekaran SS, Sultan S, Thangavel P, Harzandi AM, Singh B, Tiwari JN, Kim KS (2019) Pt-like hydrogen evolution on a V2O5/Ni(OH)2 electrocatalyst. J Mater Chem A 7(26):15794–15800

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

385

Yue Wang received his Master’s degree in physical chemistry at the Beijing University of Technology in 2020. He is currently a Ph.D. candidate under Prof. Fan Li’s supervise. His research interests focus on the thermodynamic, kinetic, synthesis, and application in energy electrocatalysis (MOR, ORR and related electrochemistry) of mixed-valence transition metal oxides, such as vanadium oxides, titanium oxides, tungsten oxides, and bandgap materials. He aims to study the thermodynamic, kinetic mechanism of synthesized these variable valence oxides and their applications in energy electrochemical.

Fan Li received his B.S. in Physical Chemistry in Metallurgy from the University of Science and Technology Beijing (USTB) in 1996. He received his M.E. in Chemical Technology from the Institute of Process Engineering, Chinese Academy of Sciences (IPE, CAS). He received his Ph.D. in Materials Science and Engineering from Royal Institute of Technology (KTH), Sweden 2007, and Physical Chemistry in Metallurgy from the University of Science and Technology Beijing (USTB) in 2008. He is a full professor at the Department of Chemistry and Chemical Engineering at Beijing University of Technology (BJUT). Dr. Li is focusing on the electrochemical properties of the transition metal oxides with mixed-valence states loaded Nobel metals catalyst, and applied them in polymer electrolyte membrane fuel cell. The preparation and characterization of nano-structured electro-catalyst with low content noble metals is his main study field. Sajid Bashir received his Ph.D. in Analytical Chemistry from the University of Warwick, England, in 2001. He was a postgraduate research associate at Cornell University to research in the field of plant genetics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has co-authored >80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry, and Chartered Chemist & Chartered Scientists of Science Council. During his service in TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with Risk Management. Currently, he collaborated with local law enforcement as a consultant.

386

Y. Wang et al.

Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing in 2001. She is a Full Professor at Texas A&M UniversityKingsville (TAMUK) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry and a DEBI faculty fellow at the US Air Force Research Laboratory. She has authored and co-authored, books, book chapters, and peer-reviewed journal articles (>100). During her 12.5-year services in TAMUK, she taught >8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and/or participated in the projects (>40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Chair-Elect of Energy and Fuels, American Chemical Society.

Chapter 14

Water-Stable Metal-Organic Frameworks for Water Adsorption Xuan Wang and Charles Lee

14.1

Fundamental Basics About Water-Stable MOFs

The study of metal-organic frameworks (MOFs) has boomed in the past two decades, and MOFs are such attractive advanced materials due to their crystallinity, porosity, tunability, and ease of functionality. In the earliest stages of development, many MOFs show significant instability in the presence of a small amount of water content, which is common in real-life applications. While the relative labile coordination between metal ions and organic linkers limits the stability of MOF, a strong electrostatic interaction resulted from high positive charge density (Z/r), like Ti4+, Zr4+, Al3+, Cr3+, Fe3+ ions, and carboxylate linkers, prevents the framework from the attack of any guest species. Among all the water-stable MOF series, the MIL series, the ZIF series, the UiO series, the CAU series, the PCN series, and the MOF-800 series stand out due to their robustness in the water environment. The underlying basics that the formation of a stable framework relies on strong coordination bonds aligns with the theory of Pearson’s Hard Soft Acid-Base (HSAB) (Fig. 14.1) [1, 2]. Synthetic cases of stable MOFs have been reported through direct synthesis and post-synthetic metathesis, in which the latter process involves the transformation from labile coordination to strong coordination by partial or complete replacement of the metal ions [3–5]. Another factor needing significant consideration for good stability is the rigidity of the framework. In general, an increasing length of organic linkers will decrease the rigidity of the framework and thus compromise the stability of the framework. Moreover, the decoration of hydrophobic functional groups on the surfaces or interfaces provides such immunity for water attacks [3, 6, 7]. More detailed discussions on the synthesis of stable MOFs could be found in several recent reviews [2, 3].

X. Wang (*) · C. Lee Texas A&M Higher Education at McAllen, Texas A&M University, McAllen, TX, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_14

387

388

X. Wang and C. Lee

Fig. 14.1 The hard-soft acid-base theory directs the construction of stable MOFs and some representative MOF examples

Stable MOFs have been widely used in all applications, namely gas sorption and separations, catalysis, energy generation, biomedical applications, and environmental science [8–18]. Water-stable MOFs represent a subclass of stable MOFs, in which the framework display resistance to water. Those MOFs advance other functional materials by their potentials in many practical applications especially in the presence of water [19, 20]. This chapter will introduce the structural fundamentals of the selected prototype water-stable MOF series and then present the development of the water adsorption using water-stable MOFs as adsorbents.

14.2

Prototypes of Water-Stable MOF Series

There has been an ever-growing effort to synthesize and utilize water-stable MOFs in their related applications, and many interesting examples of those MOFs have continuously been reported, especially through isoreticular chemistry or generation of composites displaying increased stability. In this section, several series of stable MOFs—the Materials of Institut Lavoisier (MIL) series, the zeolitic imidazolate framework (ZIF) series, the University of Oslo (UiO) series, the Christian-AlbrechtsUniversität zu Kiel, (CAU) series, the porous coordinate network (PCN) series, and the MOF-800 series—are selected, and some representative examples of each series that are of interest are presented in detail.

14.3

The MIL Series

The MIL series, developed by the Materials of Institut Lavoisier, is one novel category of water-stable MOFs. Credited to be one of the more diverse MOF series in existence, the discussion relating to the MIL series will be held to water-stable MIL MOFs, particularly Al-MIL-53, Fe-MIL-100, and MIL-101-NH2.

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

389

14.3.1 MIL-53 Al-MIL-53 ([Al(OH)(COO)2]n) is a hydrothermally-synthesized water-stable MOF, formed using the BDC linkers and Al3+ metal units [21]. Constructed from the fusion of trans chains of octahedra units of AlO4(OH)2 and BDC linkers [22], this MOF maintains a nanoporous framework with a pore size of 8.5 Å (Fig. 14.2). Additionally, Al-MIL-53 demonstrates remarkable thermal stability, as its Tdecomp is approximately 500
C. Unlike other MOFs, Al-MIL-53 is novel in its ability to “breathe” after being exposed to hydration and dehydration experiments. As the diffusion of water enters and exits the MOF, the interactions between the corner-sharing octahedra units and the water molecules via hydrogen bonding will fluctuate in such a manner that it will reduce the angle size of the BDC-Al-BDC (defined in the study as angle a) by 44%, and follow by the breaking of the hydrogen bonds, which will revert to the original angle size. More detailed structural changes will be discussed in the next section. Al-MIL-53 has a Langmuir surface area, and a BET surface area is 1590 and 1181 m2 g1, respectively. Because of its novel design and stability, Al-MIL-53 has been used in a broad spectrum of water-based applications.

14.3.2 MIL-100 Another novel member of the MIL series is Fe-MIL-100 [23]. First synthesized in 2007, the MOF is the result of a chemical marriage between Fe(III) and benzene1,3,5-tricarboxylate (BTC) as the inorganic units [24]. Fe-MIL-100 crystallizes into a hybridized super tetrahedra cubic system. As an isostructural complement to Cr-MIL-100, it has the structural capabilities to stabilize into two types of cages, Fig. 14.2 Structure demonstrations of MIL-53 from the infinite chains of corner-sharing octahedral metal units (blue polyhedra) connected through BDC (Reprinted with permission from Millange et al. [22]. Copyright @ 2010 American Chemical Society)

390

X. Wang and C. Lee

Fig. 14.3 Structure of Fe-MIL-100 consisted of trimers of the Fe SBUs (blue polyhedral) and H3BTC with two types of cages and pentagonal and hexagonal windows (Reprinted with permission from Wang et al. [24]. Copyright @ 2015 American Chemical Society)

each with varying apertures (ca. 25 and 29 Å) and varying window pore sizes (ca. 5.5 and 8.6 Å, respectively) (Fig. 14.3). The N2 adsorption experiments indicate this water-stable MOF has a Langmuir surface area greater than 2800 m2•g1. The thermal stability of Fe-MIL-100 is comparable, as its Tdecomp is approximately 270
C.

14.3.3 MIL-101 Additionally, Al-MIL-101-NH2, another member of the MIL series, has also attracted many interests in a variety of applications. This MOF is synthesized in solvothermal conditions by the addition of Al3+ ions and H2BDC-NH2 [23]. As a reticular structure to Cr-MIL-191 and Fe-MIL-101, the formation of Al-MIL-101NH2 requires aminoterephthalate ligands, while the previous experience using terephthalic acid as a linker yielded Al-MIL-53. It is important to note that the formation of Al-MIL-101-NH2 requires aminoterephthalate ligands, as previous experiments using terephthalic acid as a linker yielded Al-MIL-53. Furthermore, the utilization of these rigid linkers in conjunction with Al3+ octahedral clusters creates super-tetrahedral building units. Thus, this allows for the formation of two cage types, each with quasi-spherical mesoporous qualities [24]. One cage type, defined as a medium cavity, exhibits a 1.2 nm pentagonal window type and has 12 pentagonal faces, and the other cage type, named large cavity, exhibits a 1.6 nm hexagonal window type with 16 pentagonal face types. The MOF has thermal stability like its Cr-MIL-101reticular counterpart; its Tdecomp is approximately 650 K. The BET surface area for this MOF is approximately 2100 m2•g1.

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

14.4

391

The ZIF Series

The ZIF series demonstrates an alternate method of utilizing the HSAB theory to develop water-stable MOFs. As previously mentioned, stable MOFs are usually a combination of either hard acids and hard bases or soft acids and soft bases. The ZIF series are generally constructed from a soft acidic metal ion, notably Zn2+, and soft basic azolate ligands, such as imidazolates and tetrazoles. ZIF-8, ZIF-68, ZIF-69, and ZIF-70 are chosen for our discussion of water-stable MOFs.

14.4.1 ZIF-8 By undergoing solvothermal reactions, ZIF-8 utilizes imidazolate linkers and Zn2+ ions to form a 3D structure maintaining a sod topology [25]. Furthermore, the MOF exhibits permanent porosity with a BET surface area of 1630 m2•g1 and a Langmuir surface area of 1810 m2•g1. ZIF-8 is also credited with having a comparable pore size of ~11.6 Å, along with a small aperture, with the calculated size of 3.4 Å [26] (Fig. 14.4). The water-stable MOF is also characterized to have excellent thermal stability with Tdecomp above 500
C. Additionally, when exposed to various environments, especially aqueous environments, ZIF-8 demonstrated remarkable water-stability by maintaining its structure in a heated water environment.

14.4.2 ZIF-68 Another subclass of the ZIF series is ZIF-68-70 [27]. ZIF-68 also utilizes solvothermal conditions as a means of fusing bIM ligands with Zn2+ ions. The resulting chemical marriage forms a gme topological 3D structure with its largest Fig. 14.4 The single crystal structure of ZIF-8 with large cages (in blue) (Reprinted with permission from Gallaba et al. [26]. Copyright @ 2016 American Chemical Society)

392

X. Wang and C. Lee

Fig. 14.5 The network of ZIF-68 (Reprinted with permission from Van der et al. [28]. Copyright @ 2015 American Chemical Society)

cage exhibiting approximately 24 vertices [28] (Fig. 14.5). Additionally, ZIF-68 has a 10.2 Å pore size and 7.5 Å apertures. Concerning stability, the MOF has remarkable thermal and chemical stability. Specifically, the TGA analysis shows ZIF-68 has a Tdecomp of 390
C, and exposure to various environments, notably heated benzene, methanol, and water revealed that the MOF maintained its structure. The porosity of the MOF demonstrated a Langmuir surface area of 1220 m2•g1.

14.4.3 ZIF-69 In addition to ZIF-68, an isoreticular structure in the ZIF series is ZIF-69 [27]. This MOF shows similar characteristics to ZIF-68. Specifically, both share a similar topology (gme) and the number of vertices in the largest cage. They also share a Tdecomp of 390
C. However, in comparison to the pore size, ZIF-69 has a smaller pore size, measuring at 7.2 Å. Additionally, the MOF has a smaller Langmuir surface area of 1070 m2•g1. As previously addressed with ZIF-68, this MOF shares similar thermal and chemical stabilization qualities. Particularly, when exposed to the same environments and stimuli, the structural integrity of ZIF-69 was maintained.

14.4.4 ZIF-70 ZIF-70 is an isoreticular structure of ZIF-68 and ZIF-69 [27]. As addressed prior, many of the same qualities exhibited by the other two ZIF structures, including topology and the number of vertices present in the largest cage size, are exhibited in this structure. Specifically, some of the differences between them include pore size and Langmuir surface area, which are 15.9 Å and 1970 m2•g1, respectively.

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

14.5

393

The UiO Series

Arguably one of the most recognizable series in water-stable MOF development is the UiO series, pioneered by the University of Oslo. The UiO series demonstrates the effectiveness of the HSAB theory; specifically, the UiO series enhances the coordination bond strength between the carboxylate-based ligands and the hard acid nature of the Zr4+ ions.

14.5.1 UiO-66 to UiO-68 This combination results in the formation of two incredibly stable frameworks, particularly in water environments. Since all three MOFs are reported with a similar structure as well as other properties, we will discuss them together. The UiO series, namely UiO-66, UiO-67, and UiO-68, can be assembled through Zr ions and ditopic acids with different lengths to form two types of cages (Fig. 14.6), the smaller deemed as a super-tetrahedron and the larger deemed as a super-octahedron [29]. UiO-66 is the classic representation of the entire series. UiO-66 (Zr6(p.3-O)4(μ3OH)4(COO)12) is a cubic close-packed MOF built from Zr4+ ions and 1,4-benzenedicarboxylate (BDC) units through solvothermal reactions. Notably, UiO-66 is characterized by its incredible stability, partly due to its 12-coordinated system. In the 3D structure, UiO-66 has a window opening of 6 Å, as well as a Langmuir surface area of 1187 m2•g1. Furthermore, as the most restrictive MOF in the UiO series, it was experimentally derived to have a temperature of decomposition (Tdecomp) of 540
C. UiO-67 utilizes 4,40 -biphenyldicarboxylate (BPDC) linkers and Zr4+ ions, whose use increases the Langmuir surface area to 3000 m2•g1. Because of the significantly larger surface area, UiO-67 boasts an 8 Å cage size. UiO-68 is constructed out of Zr4+ ions and terphenyldicarboxylate (TPDC) units. It

Fig. 14.6 Two types of cages in the UiO series: the super-tetrahedron (a) and super-octahedron (b) (Reprinted with permission from Cavka et al. [29]. Copyright @ 2008 American Chemical Society)

394

X. Wang and C. Lee

has a Langmuir surface area of 4170 m2•g1, consequently allowing for an expanded cage size of 10 Å.

14.6

The CAU Series

Another class of representative MOFs is named after Christian Albrechts Universitat (CAU), the German university that housed the MOFs discoverers, the Stock group [30]. Among all the reported porous MOFs, our interests focus on CAU-3 and CAU-10.

14.6.1 CAU-3 The high-throughput synthesis of Al ions with nonfunctionalized H2BDC produces another framework, CAU-3, in a highly diluted basic environment [31]. Under the optimized synthetic condition, the researchers obtained another two isoreticular MOFs with H2BDC-NH2 and 2,6-naphthalene dicarboxylic acid (H2NDC). Overall, CAU-3 is an fcu net that is built from 12-connected [Al12(OCH3)24(COO)12] SBUs and BDC linkers. This network includes strongly distorted tetrahedral and octahedral cavities with a pore size of 10 and 11 Å, respectively. The BET surface area of the frameworks has been measured at 1250 m2•g1 (CAU-3-BDC-NH2), 1550 m2•g1 (CAU-3-BDC), and 2320 m2•g1 (CAU-3-NDC).

14.6.2 CAU-10 The Stock group later successfully obtained a series of isoreticular Al-based MOFs, classified as CAU-10-X (X ¼ H, CH3, OCH3, NO2, NH2, OH at the 5-position of the benzene ring of 1,3-dibenzo dicarboxylates) through high-throughput synthesis [32]. The CAU-10-X series are microporous frameworks with square-shaped, sinusoidal channels with a diameter of 7 Å.

14.7

The PCN Series

Porous Coordination Network (PCN) represents MOFs synthesized by the Zhou Group. One of their primary research focuses is to synthesize stable MOFs with large pores. Aligning the purposes of discussion for this chapter, PCN-222 (may also be referred to as MOF-545), PCN-224, PCN-228-230, PCN-250, PCN-333, PCN-601, and PCN-777 are reviewed below.

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

395

14.7.1 PCN-222/MOF-545 PCN-222 ([Zr6(μ3-OH)8(OH)8(TCPP4)2]) [33], also reported as MOF-545 [34], is a solvothermally synthesized MOF utilizing metal-chelated meso-tetrakis (4-carboxypheny) porphyrin (M-TCPP) units and Zr4+ ions. Structurally, the MOF resembles 3D needle-like crystals with a BET surface area of 2220 m2•g1 and hexagonal channels with a diameter of 3.7 nm [35] (Fig. 14.7). Additionally, the MOF crystallizes in the P6/mmm space group. PCN-222 is incredibly stable in a variety of environments, especially in acidic environments, as demonstrated by an isostructural compound, PCN-222(Fe).

14.7.2 PCN-224 Following suit, PCN-224, [Zr6(μ3-OH)8(H2O/OH)12]2[TCPP-M]3n, is another water-stable MOF harnessing the abilities of square planar M-TCPP units and the six-connected Zr6 clusters through the linker-elimination strategy [1] (Fig. 14.8). Like PCN-222, it formulates a 3D crystal in a space group of m3m, thus creating a structural design that facilitates the formation of channels with a diameter of 19 Å. In addition to its 3D design, the MOF has a BET surface area of 2600 m2 g1, arguably the highest surface area reported by porphyrin-based MOFs, and a pore volume of 1.59 cm3•g1.

14.7.3 PCN-228-230 Using MOF topology as the guiding criterion, the Zhou group further elongated the pore aperture through linker extensions of the porphyrinic units and created a series of porphyritic MOFs, classified as PCN-228-230 [36]. The BET surface areas range from 4510 m2•g1 (PCN-228), 4619 m2•g1 (PCN-229), and 4455 m2•g1

Fig. 14.7 Crystal structure of PCN-222. The TCPP (yellow) is connected to four 8-connected Zr6 clusters (green cuboid) with a twisted angle to generate a 3D network with 1D large channels (Reprinted with permission from Zhang et al. [35]. Copyright @ 2016 American Chemical Society)

396

X. Wang and C. Lee

Fig. 14.8 Microscope images of (a) PCN-224(Ni) and (b) PCN-224(no metal). Crystal structure, structural components, and underlying network topology of PCN-224(Ni): (c) the 6-connected D3d symmetric Zr6 in PCN-224. (d) Tetratopic TCPP ligands (violet square) with twisted dihedral angles generate a framework with 3-D nanochannels (e, f). (Reprinted with permission from Feng et al. [1]. Copyright @ 2013 American Chemical Society)

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

397

(PCN-230) along with an increasing pore size ranging from 2.5 to 3.8 nm. Notably, PCN-230 displays excellent stability under the treatment of aqueous acidic and basic solutions with pH levels ranging from 0 to 12.

14.7.4 PCN-250 To further the development of chemically- and thermally-stable MOFs, Feng and coworkers developed a kinetically tuned dimensional augmentation synthetic route with the use of a preformed [Fe2X (μ3-O) (CH3COO)6] (M could be Fe2+/3+, Co2+, Ni2+, Mn2+, Zn2+) building block [37]. In this approach, the use of premade secondary building units (SBUs) simplified the crystal growth process by easing the structure reorganization between metal clusters and organic linkers. As a result, the researchers reported 34 different MOFs from 30 different multitopic ligands (ditopic, tritopic, and tetratopic carboxylic acids) and mixed ligands in the presence of a modulator acetic acid. Among all the MOFs, PCN-250, constructed from the 6-connected [Fe2X (μ3-O)] cluster and the azobenzene tetracarboxylic acid, stands out for the ultrastability. Interestingly, it was also reported that variation in the arrangement of linkers produced two framework isomers [38] (Fig. 14.9). Despite the introduction of a softer Lewis acid M(II) in the cluster, PCN-250 remains robust after undergoing treatment of glacier acid and aqueous solution with a pH ranging from 1 to 11 lasting over 24 h. Moreover, the chemical stability test approves that PCN-250(Fe2Co) remains intact after being exposed to the water treatment more than 6 months.

Fig. 14.9 Structure representation of two framework isomers for PCN-250 (Reprinted with permission from Yuan et al. [38]. Copyright @2017 Elsevier Inc.)

398

X. Wang and C. Lee

14.7.5 PCN-333 After carefully considering the pore shape and symmetry of MOFs with large pore size, the Zhou group applied the ligand extension to construct PCN-332 and PCN-333 [39], which are isoreticular to MIL-100 [21]. While both MOFs display mesoporous cages, PCN-333 boasts the largest pore size of 55 Å and the highest void volume of 3.84 cm3 g1 (Fig. 14.10). Slight modifications to the framework result in exceptional water stability, as both Al-PCN-333 and Fe-PCN-333 maintain both the crystallinity and porosity upon the contact of either pure water or aqueous solution with pH from 3 to 9.

14.7.6 PCN-601 Contrary to the acid resistance of high-valent metal-based MOFs, many carboxylatebased MOFs collapse in diluted basic solutions. The Zhou group adopted the pyraolzate motif in the organic linker to increase the basic resistance of the framework [40]. PCN-601 was constructed from the Ni ions and 5,10,15,20-tetra (1H-pyrazole-4-yl)-porphyrin (H4TPP) ligands (Fig. 14.11). Nitrogen isotherms indicate that PCN-601 has a BET surface area of 1309 m2•g1 with the simulated pore size of 1.1 nm. The stability tests reveal that PCN-601 could survive not only in water but also in a saturated NaOH solution at 100
C.

Fig. 14.10 Structure drawing of PCN-332 and PCN-333 (Reprinted with permission from Feng et al. [39]. Copyright @ 2015 Macmillan Publishers Limited)

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

399

Fig. 14.11 PCN-601 is an ftw-a network (a, d), which was 12-connected [Ni8(OH)4(H2O)2 pyrozolate12] nodes shown as the blue polyhedron (b, e) and 4-connected planar TPP4 ligand as the red plane (c f) (Reprinted with permission from Wang et al. [40]. Copyright @2016 American Chemical Society)

14.7.7 PCN-777 Under careful structural rationalization, the Zhou group reported another Zr-based MOF, PCN-777 [41]. This framework results from 6-connected Zr6 clusters and 4,40 ,400 -s-triazine-2,4,6-triyl- tribenzoate with the aids of a competing reagent of trifluoroacetic acid (Fig. 14.12). Moreover, PCN-777 poses a mesoporous cage of 38 Å with a BET surface area of 2008 m2•g1 and a void volume of 2.28 cm3 g1 [42]. Also, PCN-777 is immune to the treatment of pure water and aqueous solutions of pH ranging from 3 to 11.

14.8

The MOF-800 Series

Dr. Omar Yaghi from the University of California-Berkeley is one of the earliest pioneers who initiated the study of MOFs [43] as well as establish the isoreticular chemistry [44–46]. His group has reported many milestone frameworks in MOF research. Their water-stable MOF-800 series is another novel class of Zr (IV)-based MOFs [47]. Synthesized in 2014, each MOF in the series are similar in nature with some synthetic and architectural differences. However, every MOF in the series naturally exhibits water absorption properties as well as efficient usability. For the discussion related to this family, emphasis will be made on MOF-801 and MOF-841 (Fig. 14.13).

400

X. Wang and C. Lee

Fig. 14.12 The coordination of 6-connected Zr6 units and organic linker TATB gives rise to the PCN-777 with large mesoporous cages (Reprinted with permission from Balestri et al. [42]. Copyright @2017 American Chemical Society)

14.8.1 MOF-801 Using solvothermal conditions, MOF-801 ([Zr6O4(OH)4(fumarate)6]∙6(H2O)) is synthetically derived using fumaric acid [47]. The MOF is characterized as the smallest member of the MOF-800 series, but it is novel in its construction. MOF-801 is reported with the synthetic capability to crystallize into a large crystal type (MOF-801-SC) or a microcrystalline powder (MOF-801-P). As the consequence of increasing amounts of defects in the powder format, MOF-801-P demonstrates a BET and Langmuir surface area of 990 and 1070 m2 g1, respectively. Conversely, MOF-801-SC has a measured BET and Langmuir surface area of 690 and 770 m2•g1, respectively. MOF-801 can exhibit two tetrahedral cavity diameters of 5.6 and 4.8 Å along with an octahedral diameter of 7.4 Å. Regardless of its size, MOF-801 uses Zr6O4(OH)4(CO2) n SBUs, which allows it to maintain a permanently porous structure.

14.8.2 MOF-841 The last MOF of discussion for this series, MOF-841 is the major product of the synthesis involving the production of the by-product MOF-812 [47]. Like the previously discussed MOF, MOF-841 (Zr6O4(OH)4(MTB)2(HCOO)4(H2O)2)

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

401

Fig. 14.13 Zr6O4(OH)4(-CO2)n SBUs with different coordination numbers listed under each of the SBUs as n-c (left column) are connected with organic linkers (middle) to form MOFs (right column) of various topologies (three-letter codes in the right column) (Reprinted with permission from Furukawa et al. [47]. Copyright @ 2014 American Chemical Society)

402

X. Wang and C. Lee

undergoes its reaction in solvothermal conditions, utilizing H4MTB organic linkers in developing the crystal structure. There is only one cage type that has a measured diameter of 11.6 Å in this framework. MOF-841 has a measured BET surface area of 1390 m2∙g1and a Langmuir surface area of 1540 m2∙g1. Besides the aforementioned MOF series, there are many other water-stable MOFs (JUC-62 [48], Bio MOF [49], ZJU-101 [50], SLUG-21 [51], SLUG-35 [52], FIR-53 [53], NU-1000 [54], and other series based on the same framework) that are interesting but not discussed here.

14.9

Water-Stable MOFs for Water Adsorption

In the second section, we described that a high surface area is one of the most advantageous properties for MOFs. This becomes particularly important when MOFs possess a high water capacity. Comparing to many commercial water adsorbents, like silica and zeolites, some MOFs display high water uptake in the region of low pressures. Besides, studies have approved that the guest molecules are attracted through physical adsorption [55–57], which ease the regeneration process and ultimately promise MOF reusability. Furthermore, finely tuning the hydrophilicity of the framework through linker design has a significant difference in many other guest molecules [57–59]. Since the evaluation of the use of MOFs in water-related applications is vast and complex, there have been many comprehensive reviews on other water-related applications, such as water separation and purification [60– 62]. Since we have a broad interest in the water crisis, we herein aim at the relationship between pure water and MOFs, where water either plays an essential role in the structure of the framework or is attracted to the framework through hostguest interactions. We will describe these studies based on developmental history, from the early stages in the structural characterization, to comprehensive investigations on adsorption and mechanisms, finally the most recent fabrication of MOF-based water harvester in real-world applications.

14.10

Early Stage for Structure Characterization

One of the very first examples of using water-stable MOFs could date back to 2002 [63]. The researchers noticed that the flexible MIL-53 was adopted to the guest molecular by a reversible swelling through atomic displacement, known as the “breathing effect” [21, 63] (Fig. 14.14). The heating at 275
C of the as-synthesized Al-MIL-53 yielded the anhydrous form Al-MIL-53ht at high temperature, in which the dimensions 7.3  7.7 Å2 of the original tunnels changed to channels with dimensions of 8.5  8.5 Å2 after removing the protonated and disordered BDC molecules. Notably, this transformation between the two forms is reversible. Upon the adsorption of one water molecule per Al at room temperature,

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

403

Fig. 14.14 Views of the 3D structure of Al-MIL-53 showing the channel system: (a) as-synthesized Al-MIL-53as, in which the channels are occupied by free disordered H2BDC molecules; (b) calcined form Al-MIL-53ht with empty channels; (c) room temperature form (Reprinted with permission from Serre et al. [63]. Copyright @2002 American Chemical Society)

Al-MIL-53lt was observed to change channel dimension, creating a channel with a dimension of 2.6  13.6 A2. The IR, NMR, and modeling data revealed that the water hydrogen bonding between hydration and dehydration was responsible for the structural switch between Al-MIL-53lt and Al-MIL-53ht. Four years later, Millange and coworkers demonstrated that the Cr-MIL-53 featured a similar breathing effect to the Al-MIL-53 [21] during the hydration and dehydration process [64]. However, unlike its counterparts, Fe-MIL-53 remained in a narrow pore form at high water loading through a metastable anhydrous intermediate Fe-MIL-53int. The Devautour-Vinot group further investigated the energetic features of the breathing effect of Cr-MIL-53, Fe-MIL-53, and V-MIL-47 [65]. The TGA data revealed that all the MIL materials could be dehydrated near 370 K, which was much lower than that of commercial adsorbent zeolites. Both TGA and DSC results suggested a two-step process comprised the water adsorption.

14.11

Investigations of Adsorption and Mechanism

With the development of characterizing technology and more collaborations between scientific fields, more insights about the adsorption have been discovered [66]. Kaskel and researchers compared the water adsorption isotherms of several benchmark MOFs—HKUST-1, MIL-100, MIL-101, ZIF-8, and DUT-4. Both HKUST-1 and DUT-4 collapsed upon the long-term water contact, while ZIF-8 and MIL-101 were retained; this indicates the importance of water stability when using MOFs as water sorbents. The hydrophobic characteristics of ZIF-8 were substantiated by the increasing water adsorption at increasing pressures. Particularly, the MOF initially showed insignificant water adsorption up to the relative pressure of 0.6, then a slight increase at p/p0 of 0.8, and a steep increase for the water condensation in the high-pressure region. MIL-100 and MIL-101 both displayed hysteresis loops in their water sorption isotherms. Additionally, quantitative analysis revealed that the calculated heat of adsorption for MIL-100 and MIL-101 was 48.83

404

X. Wang and C. Lee

Fig. 14.15 Bridging effect: a sketch of the formation of hydrogen bridging in the pores of the adsorbent and the resulting voids. (a) Adsorption of single water molecules to the hydrophilic centers. (b) Formation of water clusters due to hydrogen bonding. (c) Filled pore with free voids between the water agglomerates. M and L denote the metal cluster and the organic linker, respectively. Black dots depict the water molecules (Reprinted with permission from Küsgens et al. [66] Copyright @ 2009 Elsevier Inc.)

and 45.13 kJ/mol, respectively. These values were close to the energy required to generate hydrogen bridges between single water molecules. The researchers explained the different sorption behavior for the five MOFs in terms of bridging effect; the adsorbed water molecules on the first layer would generate more hydrogen bridges with additional water molecules, which ultimately formed water clusters in the pore (Fig. 14.15). In the same year, Henninger, Janiak, and coworkers reported a rare water-stable Ni-based MOF, SIE-1, which showed a remarkable loading spread of 280 g/kg at 150
C for the duration of over 10 cycles [67]. Besides previous structural studies of flexible MIL-53, the Bein research group investigated the water sorption behavior for another flexible Fe-MIL-88B with the aid of in situ X-ray diffraction (XRD) [68]. The data revealed that the relative vapor pressure increased during the adsorption; they found the lattice parameter c slightly decreased, while the other parameter increased (Fig. 14.16). These changes in crystal parameters contributed to a more than 40% increase in the cell volume, which ultimately resulted in a two-step adsorption process with a large hysteresis in the isotherms (Fig. 14.17). Since the MIL-53 network has been so attractive to many scientists, Stock and coworkers grafted five functional groups (–Cl, –Br, –CH3, –NO2, (OH)2) to the terephthalate linkers on Al-MIL-53 [69]. Among the five MOFs, the hydroxy MOF analog displayed a wide hysteresis loop. This phenomenon was translated into a two-step adsorption process. Adsorption was initiated by a fast uptake at low pressure, with increasing vapor pressure leading to pore opening within the framework and, consequently, taking up additional water molecules into the pores. To study an additional aspect of the MIL series, four isoreticular MIL-101 MOFs were evaluated by the Kitagawa group [70]. They studied the influence of functional groups (–H, –NO2, –NH2, –SO3H) on their water sorption behaviors. The adsorbed water amounts ranged from 0.8 to 1.2 g g1 for the MOFs. Strong interactions

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

405

Fig. 14.16 Complete diffraction pattern under p/p0 (left) and schematic of change of lattice parameters a and c at the relative pressure of water of 0 and 0.85 (right) (Reprinted with permission from Scherb et al. [68]. Copyright @ 2010 The Royal Society of Chemistry) 300 Adsorbed Volume / cm3g–1

Fig. 14.17 Water sorption isotherms of activated MIL-88B with hysteresis (Reprinted with permission from Scherb et al. [68]. Copyright @ 2010 The Royal Society of Chemistry)

adsorption desorption

250 200 150 100 50 0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 p/p0 (H2O)

between framework and water molecules were observed in the cases of the sulfonic and amino group with evidence of the increase of isosteric heats of adsorption of water. Tuning the pendant groups on the linkers based on hydrophobicity was approved to be an effective way to change the water sorption uptake. A similar evaluation of the influence of functionalized MOFs on water adsorption was reported by Walton and coworkers [71]. They studied the water adsorption for a series of isoreticular UiO-66 MOFs, with amino, nitro, methoxy, and naphthyl substitutes. UiO-66-1,4-Naphyl displayed the lowest water uptake, as evidenced by a very small hysteresis when compared to those of the other four materials. This could be explained by the inhibition of the water filling caused by the hydrophobic nature of the naphthyl ring. Interestingly, the UiO-66-OMe displayed the highest

406

X. Wang and C. Lee

water uptake, which was certainly not expected due to the methoxy groups being a high electron-donating functional group. The authors explained that the methoxy group acted as a directing agent for the pore filling of water, which resulted in more efficient water packing in the pore. In another Walton study, the researchers evaluated the water adsorption for seven classic MOFs—HKUST-1, Mg-MOF-74, UiO-66, UiO-66-NH2, DMOF-1, DMOF1-NH2, and UMCM-1—and the subsequent structural analysis using PXRD and nitrogen isotherms [72]. The Zn-based MOFs (DMOF-1, DMOF-1-NH2, and UMCM-1) lost crystallinity after the water vapor absorption, while others retained good crystallinity according to the PXRD patterns. Moreover, only the UiO-66 and the amino analog displayed negligible loss of surface area. In the same year, Chang and coworkers demonstrated a series of elaborate comparisons of the water sorption properties of mesoporous Cr-MIL-101 and Fe-MIL-100 to a selection of commercial water adsorbents (Zellites NaX, silicoaluminophosphate SAPO-34, and silica gel) in the context of water uptake, dehydration conditions, and adsorption sites [73]. The water sorption isotherms at 30
C at a relatively high partial pressure indicated the water uptakes of the MIL MOFs were 2.5–5.0 times higher than those of the commercial adsorbents. Additionally, most of the pore volume could be dehydrated at 30
C under vacuum, which was significantly lower than the high-temperature requirement of dehydration for the commercial adsorbents. The desorption profiles indicated that the desorption rates of these two MOFs were distinctively higher than those of commercial adsorbents. Moreover, the enthalpy profiles measured by microcalorimeter revealed two adsorption regions, in which the region below 0.3 g g1 corresponded to the formation of adsorbed water monolayer from the unsaturated metal sites and pore surface, which was similar to those of for NaX and SAPO-34, and the region above 0.3 g g1 was ascribed to the capillary condensation in the mesopores. Despite cases supporting carboxylate-based MOFs, Dinca and coworkers reported the water stable pyrazole-based MOFs, Zn(NDI-X), where the substituent X could be –H, –NHEt, or –SEt. This series of MOFs were mesoporous MOFs with ~16 A-wide channels and constructed from the 4-coordinated zinc chain with the functionalized dipyrazolate linkers [74]. Due to the largely hydrophobic surface of the framework, all three MOFs displayed Type V isotherms, which are described as little water adsorption at a low vapor pressure ( p/p0 ¼ 0–0.4) but a sharp increase in the region of high pressure. To increase hydrophilicity on the surface of the MOF, the researchers post-synthetically modified the surface of Zn(NDI-SEt) with more polar and hydrophilic sulfoxide(-SOEt) and sulfone(-SO2Et) groups in the presence of excess amounts of dimethyldioxirane. As a result, the sharp increase in water uptake began at p/p0 at 0.2–0.3 and 0.3–0.4 for Zn(NDI-SOEt) and Zn(NDI-SO2Et), respectively. This work demonstrated the use of a modular synthetic approach in the heating and cooling process in real devices. In another study about the CAU-10-X series, pronounced differences were observed in the water adsorption isotherms at 25
C as shown in Fig. 14.18 [32]. The amine and hydroxyl functional groups provided additional adsorption sites by forming strong hydrogen bonding with the adsorbed water molecules. The

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

407

Fig. 14.18 Adsorption isotherms of CAU-10-X at 25
C (Reprinted with permission from Reinsch et al. [32]. Copyright @ 2013 American Chemical Society)

dehydration could be completed by vacuum at room temperature without any structural change. In Stock’s study, CAU-10-H was later investigated using in situ powder XRD (PXRD) during the hydration/dehydration process [75]. The PXRD pattern showed a reversible structural change upon the water sorption process at ~20% RH. After 100 cycles of water adsorption, the water capacity of the CAU-10-H remained as 0.34 g g1 without any visible loss. The framework retained its crystallinity and its integrity after 700 adsorption/desorption cycles, even when exposed to water. With the expectation of improved stability between water sorption cycles, Henninger, Janiak, and coworkers deposited the microporous aluminum fumarate (μp-AF) onto aluminum metal sheets via the thermal gradient approach [76]. The bulk of μp-AF displayed a maximum water uptake of ~0.45 g g1 with a small hysteresis between the relative pressures of 0.2 and 0.35 in the water adsorption isotherms. The water capacity of the μp-AF coating was 0.37 g g1, even after 4500 cycles of water adsorption and desorption. After the modulator-regulated mature synthesis of Zr-based MOFs, the Yaghi group extensively studied the water adsorption of several water-stable MOFs [47]. They compared 20 MOFs, 10 of which were Zr-based MOFs, along with an additional three benchmark porous materials, hypothesizing that the hydrophilicity of organic linkers would dictate the efficacy of MOF water adsorption at low pressure. However, the pore size played a primary role in the overall water uptake. Among all the porous materials, MOF-801 and MOF-841 were the best performers in terms of water adsorption, ease of regeneration, and recyclability. The adsorption study indicated that the powder form of MOF-801 showed a higher affinity for water in comparison to the single crystal form, which was ascribed to the existence of more crystal defects in the powder MOF. The heat of adsorption of water for MOF-801-P

408

X. Wang and C. Lee

and MOF-841 was calculated to be approximately 60 and 50 kJ mol1, respectively. In this study, temperature tended to have little effect on the maximum water uptake of those materials. The single-crystal XRD and neutron diffraction studies also demonstrated that the primary water adsorption sites in low-pressure environments in MOF-801 were the tetrahedral cavities, formed by hydrogen bonding between the μ3-OH in the Zr SBUs and adsorbed water molecules. As the pressure increased, more guest water molecules began to populate into the octahedral cavities by forming new hydrogen bonds with the adsorbed water molecules in the primary binding sites. Farrusseng and coworkers also rationalized the mechanism of water adsorption for 15 MOFs in terms of the Henry constant, the pressure when pore filling occurs, and the maximum water uptake [77]. The mesoporous MIL-101 showed a total capacity of ~0.9 g g1 with hysteresis in the water adsorption isotherm due to the irreversible capillary condensation. The Cr-MIL-101-NH2 displayed a wider hysteresis than the Cr-MIL-101 and Cr-MIL-101-NO2 and a larger amount of water adsorbed due to the more hydrophilic nature of amine groups. Such hysteresis phenomena were absent in the water adsorption isotherms of the microporous MOFs. The breathable behavior of the Al/Ga-MIL-53 series enclosed a hysteresis loop, which resulted from the greater flexibility. Using the same topology as the frameworks, the researchers learned that the linker functionalization of amino groups increased the Henry constants as it decreased the pore filling pressure. Also, they found that the water adsorption of a mesoporous MOF was initially predominated by the physical adsorption of water that occurred at the metal cluster sites at lower pressure but changes to capillary condensation at higher pressure, as evidenced by hysteresis loops. The ideal adsorbents for heat pumps and chillers should have high water uptakes at low pressures with robust stabilities [78]. Bearing this in mind, Chang, Serre, and coworkers designed a water-stable Al-based MOF, MIL-160, with a 2,5-furan dicarboxylic acid, in which two oxygen atoms were utilized to increase the hydrophilicity of the framework. The AlO6 octahedral chain coordinated to the carboxylates from the linkers, forming square-shaped sinusoidal 5 Å-wide channels. The furan-rich MIL-160 endorsed the framework with a relatively higher adsorption enthalpy of 54 kJ/mol. Both the computational simulation and proton MAS NMR results indicated the primary adsorption sites at the hydroxyl groups on the chains. Furthermore, a second NMR peak rising from the weak interactions between the furan oxygen and the adsorbed water molecules was also observed. MIL-160 possesses a high water capacity of 235 g kg1, in which most of the water molecules were adsorbed at p/p0 < 0.18 without any hysteresis and could be removed below 373 K. Thanks to these features, they have helped MIL-160 to be a promising alternative as a water adsorbent in solar cooling applications. No weight loss was noticed after 10 cycles of water adsorption and desorption using MIL-160.

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

14.12

409

Development of MOF-Based Device in Practical Applications

Due to the high-water capacity, excellent reusability, and ease synthesis of MOF-801, Yaghi and Wang group further developed a water harvester system (Fig. 14.19), which consisted of copper foam, a layer of MOF-801 powder, and a condenser interfaced with a thermoelectric cooler [79]. Vapor from the environmental humidity was adsorbed into the MOF layer through physical diffusion, followed by a low-grade heat from the sunshine on the absorber driven by the desorption process, and concluding with those vapors being further condensed into the water collector. Their MOF-based device demonstrated the capability to capture 2.8 L of water per kg of MOF daily at 20% humidity level from the atmosphere in ambient conditions with sunlight at a flux of less than 1 sun (1 kW/m2). Overall, this design not only maximized the yield by applying the sunlight-driven vapor-desorption but also required no additional energy input. After discovering the potential of MOFs in the application of water production, the same research group implemented the MOF-based device in the practical water capture from desert air [80]. An industry scale of MOF (~1.2 kg MOF-801) was used to build the device, but no liquid water had been collected at the Arizona desert testing spot until the device was slightly modified to use the exterior insulation at a tilted position. This device produced 100 g of water per kg of MOF per one full cycle with no energy input. After carefully evaluating the parameters that affected the harvesting efficiency, they developed a next-generation device with MOF-303, which was constructed from an infinite aluminum-oxo chain and 1H-pyrazole-3,5dicarboxylate. MOF-303 was selected for its unique structure that not only facilitated

Fig. 14.19 The MOF-based water-harvesting prototype

410

X. Wang and C. Lee

a water capacity of 0.48 g.g1 but also displayed excellent water stability and good temperature response. As expected, the new-generation device gave a remarkable increase of 114% water production. After demonstrating an effective proof of concept for the water harvesting device, the Yaghi group and Wang group further optimized the engineering aspect of the MOF-based water harvester by coating the MOF layer black with an OTTI-aerogel cover as the optically transparent thermal insulator [81]. With this advancement, they made full advantage of the day-time heat to adsorb the water vapor and the night-time radiative cooling to increase the effective relative humidity (RH) during the swing from 10% (daytime at desert areas) to 20–40% (night time). The thermal efficiency for solar input to water conversion was reported to be ~14%, and the highfidelity computational simulations predicted that this device could deliver ~0.25 L water/kg MOF-801 for a single day cycle. Later, the same groups compared the kinetics on water sorption of MOF-303 to the other four commercial sorbents [82]. The kinetic study of MOF-303 displayed not only showed the fastest water uptakes at 30% and 40% but also the fastest dehydration within minutes upon mild heating. Based on the findings, the Yaghi and Wang collaboration improved their adsorbent-assisted water harvesting cycle from once a day to multiple cycles with the aid of solar power. This enhancement endorsed the MOF-303-based device with the capability to generate 1.3 L per kg of MOF every day in an arid environment of 27
C and 32% RH. In the three-day collection tested in the Mojave Desert, this device produced 0.7 L per kg of MOF per day. To avoid the irreversible capillary condensation of water in a large pore, the Dinca group took into consideration the critical diameter of water capillary action (20.76 Å) in the selection of MOFs for water capture [83]. MOFs with pore size below this critical diameter were predicted to be more likely to display continuous pore filling of water [71]. A series of mesoporous MOFs, M2Cl2(BTDD) (M ¼ Mn, Co, Ni), with an inherited 22 Å wide windows, were studied for the water isotherms. The Co-based MOF turned out to be a superior water absorber with the total water uptake of 0.968 g g1 at 94% RH with no hysteresis of adsorption as well as a calculated heat of adsorption of water at ~55 kJ/mol in zero coverage. Similar to the MOF-801 filled water harvester, this Co-based MOF was projected to capture 0.87 L water per kg of MOF, which agreed with the simulation results (0.82 g g1 MOF). Furthermore, the simulation also predicted that the water capture for this device only dropped by 5.1% over 6 cycles. They also found that this material had a cooling capacity of 400 kWh∙m3 per cycle, which corresponded to a 20
C temperature lift. Maurin, Chang, Serre, and collaborators presented a stable MOF called MIP-200 (MIP ¼ Materials of Institute of Porous Materials from Paris), which consisted of Zr6 oxo clusters and tetracarboxylic linkers [84]. MIP-200 was synthesized under a topology-guided strategy and a Kagome-type framework with both hexagonal (~13 Å in diameter) and triangular channels (6.8 Å in diameter). Both the water adsorption isotherms and TGA profile revealed MIP-200 adsorbed water in a two-step process, with the majority of adsorbed water being removed below 65
C. Negligible differences in weight loss were noticed in the 50 cycle-water

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

411

sorption tests. Also, the GCMC simulation showed that the hydroxyl group on the adjacent Zr6 SBUs coordinated to the adsorbed water molecules in both channels, resulting in the MOF exhibiting high hydrophilicity. Initially, the simulations predicted a simulated heat of adsorption at ~70 kJ/mol, which was evidenced by a steep increase in low relative pressure ( p/po < 0.1). However, as a network of hydrogen bonds formed during the pore filling stage, before the saturation stage at p/po ¼ 0.15–2.0, the adsorption of enthalpy was ascribed to be 55 kJ/mol. No decrease in the working capacity for gravimetric water adsorption was noted in the cyclability tests of MIP-200. Conclusively, MIP-200 demonstrated excellent chemical and mechanical stability, which shows its tremendous promise in its application for adsorption-driven refrigeration. Another MOF, CAU-23, was shown to exhibit properties that would make it suitable for application in adsorption-driven chillers, substantiated by the collaboration of the Henning, Maurin, Zou, and Stock groups [85]. CAU-23 is constructed from the aluminum-oxo chain with the TDC linkers, resulting in square channels with a side length of 7.6 Å. CAU-23 possesses a water adsorption capacity of 0.375 g g1 at p/po ¼ 0.3 and demonstrates a cooling temperature of 10
C. Like other Al-based MOFs, the adsorption sites are located at the μ-OH groups of metal SBUs, in which the formation of hydrogen bonds initiates the adsorption process. Following this, the water molecules occupied the channels until saturation at p/ po ¼ 0.3 by forming networks of hydrogen bonds between the adsorbed water molecules. After 5000 water sorption cycles, CAU-23 retained favorable crystallinity and water capacity. Additionally, CAU-23 displayed a coefficient of performance for cooling of 0.8, even at a very low-desorption temperature of ~50
C.

14.13

Outlook

The chapter has provided a brief overview of the fundamental basics of water-stable MOFs in terms of their synthesis and structural features as well as their applications in water storage. After the extensive development of MOFs in the last two decades, both the stability and the liability of MOFs have been utilized in various fields, with new synthetic strategies for synthesizing water-stable MOFs being discovered in recent years. Therefore, the study of these materials in all aspects is quite limited. Further investigations are needed promptly to fully utilize their potentials in energyrelated applications, especially in the area of water. Acknowledgments The authors thank Texas A&M Higher Education Center and College of Science of Texas A&M University for their additional support. Author contributions X. W. planned and prepared the manuscript and C. L. assisted in writing part of section 2.

412

X. Wang and C. Lee

References 1. Feng D, Chung W-C, Wei Z, Gu Z-Y, Jiang H-L, Chen Y-P et al (2013) Construction of ultrastable porphyrin Zr metal-organic frameworks through linker elimination. J Am Chem Soc 135:17105–17110 2. Yuan S, Feng L, Wang K, Pang J, Bosch M, Lollar C et al (2018) Stable metal-organic frameworks: design, synthesis, and applications. Adv Mater 30:1704303 3. Ding M, Cai X, Jiang H-L (2019) Improving MOF stability: approaches and applications. Chem Sci 10:10209–10230 4. Kalaj M, Cohen SM (2020) Postsynthetic modification: an enabling technology for the advancement of metal-organic frameworks. ACS Cent Sci 6:1046–1057 5. Lollar CT, Qin J-S, Pang J, Yuan S, Becker B, Zhou H-C (2018) Interior decoration of stable metal-organic frameworks. Langmuir 34:13795–13807 6. Makal TA, Wang X, Zhou H-C (2013) Tuning the moisture and thermal stability of metalorganic frameworks through incorporation of pendant hydrophobic groups. Cryst Growth Des 13:4760–4768 7. Yang C, Kaipa U, Mather QZ, Wang X, Nesterov V, Venero AF et al (2011) Fluorous metalorganic frameworks with superior adsorption and hydrophobic properties toward oil spill cleanup and hydrocarbon storage. J Am Chem Soc 133:18094–18097 8. Bigdeli F, Lollar CT, Morsali A, Zhou H-C (2020) Switching in metal-organic frameworks. Angew Chem Int Ed 59:4652–4669 9. Burtch NC, Heinen J, Bennett TD, Dubbeldam D, Allendorf MD (2018) Mechanical properties in metal-organic frameworks: emerging opportunities and challenges for device functionality and technological applications. Adv Mater 30:1704124 10. Canivet J, Fateeva A, Guo Y, Coasne B, Farrusseng D (2014) Water adsorption in MOFs: fundamentals and applications. Chem Soc Rev 43:5594–5617 11. Dissegna S, Epp K, Heinz WR, Kieslich G, Fischer RA (2018) Metal-organic frameworks: defective metal-organic frameworks (Adv. Mater. 37/2018). Adv Mater 30:1870280 12. Feng M, Zhang P, Zhou H-C, Sharma VK (2018) Water-stable metal-organic frameworks for aqueous removal of heavy metals and radionuclides: a review. Chemosphere 209:783–800 13. Jiao L, Wang Y, Jiang H-L, Xu Q (2018) Metal-organic frameworks as platforms for catalytic applications. Adv Mater 30:1703663 14. Kalmutzki MJ, Diercks CS, Yaghi OM (2018) Metal-organic frameworks for water harvesting from air. Adv Mater 30:1704304 15. Lu K, Aung T, Guo N, Weichselbaum R, Lin W (2018) Nanoscale metal-organic frameworks for therapeutic, imaging, and sensing applications. Adv Mater 30:1707634 16. Simon-Yarza T, Mielcarek A, Couvreur P, Serre C (2018) Drug delivery: nanoparticles of metal-organic frameworks: on the road to in vivo efficacy in biomedicine (Adv. Mater. 37/2018). Adv Mater 30:1870281 17. Woellner M, Hausdorf S, Klein N, Mueller P, Smith MW, Kaskel S (2018) Adsorption and detection of hazardous trace gases by metal-organic frameworks. Adv Mater 30:1704679 18. Zhao X, Wang Y, Li D-S, Bu X, Feng P (2018) Metal-organic frameworks for separation. Adv Mater 30:1705189 19. Qin J-S, Yuan S, Lollar C, Pang J, Alsalme A, Zhou H-C (2018) Stable metal-organic frameworks as a host platform for catalysis and biomimetics. Chem Commun 54:4231–4249 20. Wang C, Liu X, Keser Demir N, Chen JP, Li K (2016) Applications of water-stable metalorganic frameworks. Chem Soc Rev 45:5107–5134 21. Loiseau T, Serre C, Huguenard C, Fink G, Taulelle F, Henry M et al (2004) A rationale for the large breathing of the porous aluminum terephthalate (MIL-53) upon hydration. Chem Eur J 10:1373–1382 22. Millange F, Guillou N, Medina ME, Férey G, Carlin-Sinclair A, Golden KM et al (2010) Selective sorption of organic molecules by the flexible porous hybrid metalorganic framework MIL-53(Fe) controlled by various hostguest interactions. Chem Mater 22:4237–4245

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

413

23. Horcajada P, Surblé S, Serre C, Hong D-Y, Seo Y-K, Chang J-S et al (2007) Synthesis and catalytic properties of MIL-100(Fe), an iron(iii) carboxylate with large pores. Chem Commun 27:2820–2822 24. Wang D, Wang M, Li Z (2015) Fe-based metal-organic frameworks for highly selective photocatalytic benzene hydroxylation to phenol. ACS Catal 5:6852–6857 25. Park KS, Ni Z, Côté AP, Choi JY, Huang R, Uribe-Romo FJ et al (2006) Exceptional chemical and thermal stability of zeolitic imidazolate frameworks. Proc Natl Acad Sci 103:10186 26. Gallaba DH, Albesa AG, Migone AD (2016) Evidence of gate-opening on xenon adsorption on ZIF-8: an adsorption and computer simulation study. J Phys Chem C 120:16649–16657 27. Banerjee R, Phan A, Wang B, Knobler C, Furukawa H, Keeffe M et al (2008) High-throughput synthesis of Zeolitic imidazolate frameworks and application to CO2 capture. Science 319:939 28. Van der Perre S, Bozbiyik B, Lannoeye J, De Vos DE, Baron GV, Denayer JFM (2015) Experimental study of adsorptive interactions of polar and nonpolar adsorbates in the zeolitic imidazolate framework ZIF-68 via pulse gas chromatography. J Phys Chem C 119:1832–1839 29. Cavka JH, Jakobsen S, Olsbye U, Guillou N, Lamberti C, Bordiga S et al (2008) A new zirconium inorganic building brick forming metal-organic frameworks with exceptional stability. J Am Chem Soc 130:13850–13851 30. Ahnfeldt T, Guillou N, Gunzelmann D, Margiolaki I, Loiseau T, Férey G et al (2009) [Al4(OH) 2(OCH3)4(H2N-BDC)3]x H2O: a 12-connected porous metal-organic framework with an unprecedented aluminum-containing brick. Angew Chem Int Ed 48:5163–5166 31. Reinsch H, Feyand M, Ahnfeldt T, Stock N (2012) CAU-3: a new family of porous MOFs with a novel Al-based brick: [Al2(OCH3)4(O2C-X-CO2)] (X ¼ aryl). Dalton Trans 41:4164–4171 32. Reinsch H, van der Veen MA, Gil B, Marszalek B, Verbiest T, de Vos D et al (2013) Structures, sorption characteristics, and nonlinear optical properties of a new series of highly stable aluminum MOFs. Chem Mater 25:17–26 33. Feng D, Gu Z-Y, Li J-R, Jiang H-L, Wei Z, Zhou H-C (2012) Zirconium-metalloporphyrin PCN-222: mesoporous metal-organic frameworks with ultrahigh stability as biomimetic catalysts. Angew Chem Int Ed 51:10307–10310 34. Morris W, Volosskiy B, Demir S, Gándara F, McGrier PL, Furukawa H et al (2012) Synthesis, structure, and metalation of two new highly porous zirconium metal-organic frameworks. Inorg Chem 51:6443–6445 35. Zhang G-Y, Zhuang Y-H, Shan D, Su G-F, Cosnier S, Zhang X-J (2016) Zirconium-based porphyrinic metal-organic framework (PCN-222): enhanced photoelectrochemical response and its application for label-free phosphoprotein detection. Anal Chem 88:11207–11212 36. Liu T-F, Feng D, Chen Y-P, Zou L, Bosch M, Yuan S et al (2015) Topology-guided design and syntheses of highly stable mesoporous porphyrinic zirconium metal-organic frameworks with high surface area. J Am Chem Soc 137:413–419 37. Feng D, Wang K, Wei Z, Chen Y-P, Simon CM, Arvapally RK et al (2014) Kinetically tuned dimensional augmentation as a versatile synthetic route towards robust metal-organic frameworks. Nat Commun 5:5723 38. Yuan S, Sun X, Pang J, Lollar C, Qin J-S, Perry Z et al (2017) PCN-250 under pressure: sequential phase transformation and the implications for MOF densification. Joule 1:806–815 39. Feng D, Liu T-F, Su J, Bosch M, Wei Z, Wan W et al (2015) Stable metal-organic frameworks containing single-molecule traps for enzyme encapsulation. Nat Commun 6:5979 40. Wang K, Lv X-L, Feng D, Li J, Chen S, Sun J et al (2016) Pyrazolate-based porphyrinic metalorganic framework with extraordinary base-resistance. J Am Chem Soc 138:914–919 41. Feng D, Wang K, Su J, Liu T-F, Park J, Wei Z et al (2015) A highly stable zeotype mesoporous zirconium metal-organic framework with ultralarge pores. Angew Chem Int Ed 54:149–154 42. Balestri D, Roux Y, Mattarozzi M, Mucchino C, Heux L, Brazzolotto D et al (2017) Heterogenization of a [NiFe] hydrogenase mimic through simple and efficient encapsulation into a mesoporous MOF. Inorg Chem 56:14801–14808 43. Rowsell JLC, Yaghi OM (2004) Metal-organic frameworks: a new class of porous materials. Microporous Mesoporous Mater 73:3–14

414

X. Wang and C. Lee

44. Doonan CJ, Morris W, Furukawa H, Yaghi OM (2009) Isoreticular metalation of metalorganic frameworks. J Am Chem Soc 131:9492–9493 45. Eddaoudi M, Kim J, Rosi N, Vodak D, Wachter J, Keeffe M et al (2002) Systematic design of pore size and functionality in isoreticular MOFs and their application in methane storage. Science 295:469 46. Furukawa H, Go YB, Ko N, Park YK, Uribe-Romo FJ, Kim J et al (2011) Isoreticular expansion of metal-organic frameworks with triangular and square building units and the lowest calculated density for porous crystals. Inorg Chem 50:9147–9152 47. Furukawa H, Gándara F, Zhang Y-B, Jiang J, Queen WL, Hudson MR et al (2014) Water adsorption in porous metal-organic frameworks and related materials. J Am Chem Soc 136:4369–4381 48. Xue M, Zhu G, Li Y, Zhao X, Jin Z, Kang E et al (2008) Structure, hydrogen storage, and luminescence properties of three 3D metalorganic frameworks with NbO and PtS topologies. Cryst Growth Des 8:2478–2483 49. Mon M, Lloret F, Ferrando-Soria J, Martí-Gastaldo C, Armentano D, Pardo E (2016) Selective and efficient removal of mercury from aqueous media with the highly flexible arms of a BioMOF. Angew Chem Int Ed 55:11167–11172 50. Zhang Q, Yu J, Cai J, Zhang L, Cui Y, Yang Y et al (2015) A porous Zr-cluster-based cationic metal-organic framework for highly efficient Cr2O72 removal from water. Chem Commun 51:14732–14734 51. Fei H, Paw UL, Rogow DL, Bresler MR, Abdollahian YA, Oliver SRJ (2010) Synthesis, characterization, and catalytic application of a cationic metalorganic framework: Ag2 (4,40 -bipy)2(O3SCH2CH2SO3). Chem Mater 22:2027–2032 52. Fei H, Han CS, Robins JC, Oliver SRJ (2013) A cationic metal-organic solid solution based on Co(II) and Zn(II) for chromate trapping. Chem Mater 25:647–652 53. Fu H-R, Xu Z-X, Zhang J (2015) Water-stable metal-organic frameworks for fast and high dichromate trapping via single-crystal-to-single-crystal ion exchange. Chem Mater 27:205–210 54. Ahn S, Thornburg NE, Li Z, Wang TC, Gallington LC, Chapman KW et al (2016) Stable metalorganic framework-supported niobium catalysts. Inorg Chem 55:11954–11961 55. Choi I-H, Yoon SB, Jang S-Y, Huh S, Kim S-J, Kim Y (2019) Gas sorption properties of a new three-dimensional in-ABDC MOF with a diamond net. Front Mater 6. https://doi.org/10.3389/ fmats.2019.00218 56. Li H, Wang K, Sun Y, Lollar CT, Li J, Zhou H-C (2018) Recent advances in gas storage and separation using metal-organic frameworks. Mater Today 21:108–121 57. Mason JA, Veenstra M, Long JR (2014) Evaluating metal-organic frameworks for natural gas storage. Chem Sci 5:32–51 58. Hu Z, Wang Y, Shah BB, Zhao D (2019) CO2 capture in metal-organic framework adsorbents: an engineering perspective. Adv Sustain Syst 3:1800080 59. Lu W, Wei Z, Gu Z-Y, Liu T-F, Park J, Park J et al (2014) Tuning the structure and function of metal-organic frameworks via linker design. Chem Soc Rev 43:5561–5593 60. Li J, Wang H, Yuan X, Zhang J, Chew JW (2020) Metal-organic framework membranes for wastewater treatment and water regeneration. Coord Chem Rev 404:213116 61. Wang C, Kim J, Malgras V, Na J, Lin J, You J et al (2019) Water purification: metal-organic frameworks and their derived materials: emerging catalysts for a sulfate radicals-based advanced oxidation process in water purification (small 16/2019). Small 15:1970085 62. Wang H, Zhao S, Liu Y, Yao R, Wang X, Cao Y et al (2019) Membrane adsorbers with ultrahigh metal-organic framework loading for high flux separations. Nat Commun 10:4204 63. Serre C, Millange F, Thouvenot C, Noguès M, Marsolier G, Louër D et al (2002) Very large breathing effect in the first nanoporous chromium(III)-based solids: MIL-53 or CrIII(OH) {O2CC6H4CO2}{HO2CC6H4CO2H}xH2Oy. J Am Chem Soc 124:13519–13526 64. Millange F, Guillou N, Walton RI, Grenèche J-M, Margiolaki I, Férey G (2008) Effect of the nature of the metal on the breathing steps in MOFs with dynamic frameworks. Chem Commun 130(39):4732–4734

14

Water-Stable Metal-Organic Frameworks for Water Adsorption

415

65. Devautour-Vinot S, Maurin G, Henn F, Serre C, Devic T, Férey G (2009) Estimation of the breathing energy of flexible MOFs by combining TGA and DSC techniques. Chem Commun 21:2733–2735 66. Küsgens P, Rose M, Senkovska I, Fröde H, Henschel A, Siegle S et al (2009) Characterization of metal-organic frameworks by water adsorption. Microporous Mesoporous Mater 120:325–330 67. Henninger SK, Habib HA, Janiak C (2009) MOFs as adsorbents for low-temperature heating and cooling applications. J Am Chem Soc 131:2776–2777 68. Scherb C, Koehn R, Bein T (2010) Sorption behavior of an oriented surface-grown MOF-film studied by in situ X-ray diffraction. J Mater Chem 20:3046–3051 69. Biswas S, Ahnfeldt T, Stock N (2011) New functionalized flexible Al-MIL-53-X (X ¼ –Cl, – Br, –CH3, –NO2, –(OH)2) solids: syntheses, characterization, sorption, and breathing behavior. Inorg Chem 50:9518–9526 70. Akiyama G, Matsuda R, Sato H, Hori A, Takata M, Kitagawa S (2012) Effect of functional groups in MIL-101 on water sorption behavior. Microporous Mesoporous Mater 157:89–93 71. Cmarik GE, Kim M, Cohen SM, Walton KS (2012) Tuning the adsorption properties of UiO-66 via ligand functionalization. Langmuir 28:15606–15613 72. Schoenecker PM, Carson CG, Jasuja H, Flemming CJJ, Walton KS (2012) Effect of water adsorption on retention of structure and surface area of metal-organic frameworks. Ind Eng Chem Res 51:6513–6519 73. Seo Y-K, Yoon JW, Lee JS, Hwang YK, Jun C-H, Chang J-S et al (2012) Energy-efficient dehumidification over hierarchically porous metal-organic frameworks as advanced water adsorbents. Adv Mater 24:806–810 74. Wade CR, Corrales-Sanchez T, Narayan TC, Dincă M (2013) Postsynthetic tuning of hydrophilicity in pyrazolate MOFs to modulate water adsorption properties. Energy Environ Sci 6:2172–2177 75. Fröhlich D, Henninger SK, Janiak C (2014) Multicycle water vapor stability of microporous breathing MOF aluminum isophthalate CAU-10-H. Dalton Trans 43:15300–15304 76. Jeremias F, Fröhlich D, Janiak C, Henninger SK (2014) Advancement of sorption-based heat transformation by a metal coating of highly-stable, hydrophilic aluminum fumarate MOF. RSC Adv 4:24073–24082 77. Canivet J, Bonnefoy J, Daniel C, Legrand A, Coasne B, Farrusseng D (2014) Structure-property relationships of water adsorption in metal-organic frameworks. New J Chem 38:3102–3111 78. Cadiau A, Lee JS, Damasceno Borges D, Fabry P, Devic T, Wharmby MT et al (2015) Design of hydrophilic metal-organic framework water adsorbents for heat reallocation. Adv Mater 27:4775–4780 79. Kim H, Yang S, Rao SR, Narayanan S, Kapustin EA, Furukawa H et al (2017) Water harvesting from air with metal-organic frameworks powered by natural sunlight. Science 356:430 80. Fathieh F, Kalmutzki MJ, Kapustin EA, Waller PJ, Yang J, Yaghi OM (2018) Practical water production from desert air. Sci Adv 4:eaat3198 81. Kim H, Rao SR, Kapustin EA, Zhao L, Yang S, Yaghi OM et al (2018) Adsorption-based atmospheric water harvesting device for arid climates. Nat Commun 9:1191 82. Hanikel N, Prévot MS, Fathieh F, Kapustin EA, Lyu H, Wang H et al (2019) Rapid cycling and exceptional yield in a metal-organic framework water harvester. ACS Cent Sci 5:1699–1706 83. Rieth AJ, Yang S, Wang EN, Dincă M (2017) Record atmospheric fresh water capture and heat transfer with a material operating at the water uptake reversibility limit. ACS Cent Sci 3:668–672 84. Wang S, Lee JS, Wahiduzzaman M, Park J, Muschi M, Martineau-Corcos C et al (2018) A robust large-pore zirconium carboxylate metal-organic framework for energy-efficient watersorption-driven refrigeration. Nat Energy 3:985–993 85. Lenzen D, Zhao J, Ernst S-J, Wahiduzzaman M, Ken Inge A, Fröhlich D et al (2019) A metal– organic framework for efficient water-based ultra-low-temperature-driven cooling. Nat Commun 10:3025

416

X. Wang and C. Lee Xuan Wang is an instructional assistant professor at Texas A&M Higher Education at McAllen, where she holds appointments in the College of Science at Texas A&M University. Her research interests include the synthesis and development of metalorganic frameworks and porous coordination polymers for chemical sensing and the application in the environmental science field. Her academic training includes undergraduate study at Oklahoma Christian University (B.S., 2011) and graduate study under the direction of Professor Hong-Cai Zhou at Texas A&M University (Ph.D., 2016). She began an academic career as a visiting assistant professor at Colorado State University-Pueblo after graduation.

Mr. Charles Lee is an undergraduate student at Texas A&M Higher Education Center at McAllen. He is pursuing a degree in Biomedical Science. An emerging physician-scientist, Charles’s developing research interests combine biochemical and chemical principles in drug discovery and translational medicine applications. Outside of school, Charles is an accomplished equestrian competing with his beloved American Quarter Horses.

Chapter 15

Supercapacitors: History, Theory, Emerging Technologies, and Applications Yiyang Liu, Paul R. Shearing, Guanjie He, and Dan J. L. Brett

15.1

Introduction

With the development of a global economy, rapid population increase, and the implications of global warming, traditional energy sources will not be able to meet the demand and increasing deployment of renewable energy and transition of electrochemical power systems for vehicle propulsion calls for alternative methods of energy storage [108]. It is particularly important to seek environmentally friendly and sustainable energy resources, such as solar, wind, and geothermal energy; however, an urgent issue is to achieve efficient storage of these renewables [21, 55]. Currently, the development of novel electrochemical energy storage devices, including batteries, supercapacitors (SCs), and fuel cells, is being highly valued by researchers and enterprises. During the past three decades, the applications of rechargeable batteries have surged in many fields, from mobile electronic devices to grid-scale energy storage systems [58]. For example, lithium-ion batteries (LIBs) can achieve a high energy density of ~180 Wh
kg1 (Fig. 15.1), and this continues to

Y. Liu Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK P. R. Shearing · D. J. L. Brett Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK The Faraday Institution, Didcot, UK e-mail: [email protected] G. He (*) Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK School of Chemistry, University of Lincoln, Lincoln, UK e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_15

417

418

Y. Liu et al.

Fig. 15.1 Ragone plot comparing electrochemical energy conversion and storage devices

increase through rapid advancement of the field. However, for many applications, their power density is unsatisfactory, which leads to a slow charging rate. Meanwhile, the shortcomings of LIBs, including high cost, insufficient lithium reserve, and potential safety issues, may retard LIB deployment and their long-term prospects. In addition to ultra-high power density (10 ~ 100 kW kg1) compared to other energy conversion and storage devices, SCs have merits including operation over a wide range of temperatures (40 ~ 80  C), high efficiency, and fast charge/discharge rates (in seconds) [3, 4, 34]. Meanwhile, compared with some commercial technologies, such as fuel cells, SCs possess a much lower cost and longer cycle lives (> 100,000 cycles). As energy storage devices, the properties of SCs sit between traditional capacitors and rechargeable batteries. As shown in Fig. 15.2, supercapacitors can be used as both quick-start power supplies for electrical vehicles and balanced power supplies for lifting devices; they can also be used as traction energy for hybrid electric vehicles, internal combustion engines, and trackless vehicles, as well as power supplies for other equipment. It is worth mentioning that SCs are of great significance in plug-in hybrid vehicles. Generally, it is difficult to simultaneously deliver high energy density, power density, and long cycle life in a single power system, since the rechargeable batteries must make a trade-off among these properties. The adoption of auxiliary energy systems using SCs can address this challenge: the main energy system (rechargeable batteries) with high energy density provides a longer driving range, while the auxiliary energy system (SCs) with the high power density and cycle life provides the short-term auxiliary power in scenarios such as acceleration and slope climbing. To provide the necessary and comprehensive understanding for new researchers in the fields related to SCs, this chapter provides a concise introduction and discussion covering: (a) the development history, and fundamental theories and principles

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

419

Fig. 15.2 Application scenarios of supercapacitors

(Sect. 15.1); (b) electrolytes (Sect. 15.2); (c) electrode candidates, including carbon materials (Sect. 15.3); transition metal compounds (Sect. 15.4) and novel electrode materials (Sect. 15.5); and (d) conclusion and perspectives (Sect. 15.6).

15.2

The History of Supercapacitors

The “Leyden jar” is the earliest report of a capacitor. Invented in 1746 by Prof. Pieter Van Musschenbroek at the University of Leiden, it comprises a glass jar filled with water into which a brass rod is placed [84]. An early “Leyden jar” can be charged to a high voltage of 20,000 ~ 60,000 V and has a typical capacitance of 1 nF per pint of size. However, SCs represent an energy storage device positioned between traditional capacitors and rechargeable batteries, which possess high power densities and relatively high energy. In 1853, German physicist Hermann Von Helmholtz put forward a theoretical model of an interfacial double layer (DL): under the action of a specific potential, two layers of charge with the same amount but the opposite charges will be generated at the interface between electrode materials and electrolyte, thus forming an electrical double layer [35]. Based on insight from this theory and research on porous carbon electrode materials for fuel cells and rechargeable batteries, H. Becker from General Electric developed a “low voltage electrolytic capacitor with porous carbon electrodes” [11]. Although the underlying mechanisms for energy storage in porous carbon electrodes were not clear, the apparatus demonstrated a high energy

420

Y. Liu et al.

density similar to rechargeable batteries at that time. Inspired by Becker’s research, Donald L. Boos from Standard Oil of Ohio (SOHIO) further developed this technology and made the first attempt to market SCs [14]. Then, SOHIO licensed this technology to NEC Corporation (Nippon Denki Kabushiki-gaisha), who first successfully commercialized this electric double-layer capacitor (EDLC) as backup power for computer memory [7, 8]. However, the market of first-generation EDLCs expanded slowly, due to their unsatisfactory energy density, low operating current, and high internal resistance. To address these challenges, Trasatti and Buzzanaca found that ruthenium oxides (RuO2) demonstrated outstanding capacitive performance, which set off a wave of research on pseudocapacitors based on metal oxides as active materials [83]. Between 1975 ~ 1980, Brian E. Conway carried out extensive fundamental studies and commercial efforts on electrochemical capacitors using RuO2 electrodes, and significantly enriched the knowledge of the electrochemical capacitor field. He redefined the term “supercapacitor” and proposed the capacitance comes from two parts: (a) the electrical charges stored in the Helmholtz electric double layer and (b) the results of Faradic reactions with “pseudocapacitance” charge transfer, the mechanisms of which include redox reactions, intercalation, and electrosorption [19]. Since then, pseudocapacitive SCs using RuO2 as electrodes have been used in various applications ranging from automobiles to the military. Since the 1990s, a series of inexpensive transition metal (e.g. Mn, Ni, Co, and V) oxides, conductive polymers, and other electrode materials have also been extensively studied.

15.3

Working Principles and Classification

As shown in Fig. 15.3, the structure of SCs consists of four components: cathode, anode, electrolyte, and separator. Defined by the different types of electrode materials used, supercapacitors can be divided into three categories: electric double-layer

Fig. 15.3 Typical structure of supercapacitors

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

421

Fig. 15.4 Types of supercapacitors. (a) EDLCs; (b) pseudocapacitive SCs (pseudocapacitors); (c) hybrid SCs

capacitors (EDLCs, Fig. 15.4a), pseudocapacitive SCs (pseudocapacitors, Fig. 15.4b), and hybrid SCs (Fig. 15.4c).

15.3.1 Electric Double-Layer Capacitors (EDLCs) As shown in Fig. 15.4a, charge storage in EDLCs is achieved by accumulating charges at the interface between electrodes and the electrolyte. Under the action of an external electric field, equal amounts of charges carried by anions and cations move to the cathode and anode, respectively, thus forming a potential difference. When a charge accumulation occurs on the surface of an electrode, ions with opposite charges in the nearby electrolyte migrate to the electrode and accumulate on its surface, thus forming an electric double layer. When the external electric field is withdrawn, the electric double layer still exists and stabilizes the voltage due to the attraction between opposite electric charges. After the supercapacitor is connected to an external electrical appliance and a closed circuit is formed, the charged ions absorbed on the electrodes will migrate directionally and form a current in the external circuit until the electrolyte returns to electrically neutral. The charge and discharge of an EDLC are highly, or even completely, reversible. This is due to the fast ion electric absorption and desorption process and almost no chemical reactions involved [97]. Since 1853, a series of theories and models were developed, among which the most representative ones are the Helmholtz model, Gouy-Chapman (G-C) model, and the Gouy-Chapman-Stern (G-C-S) model. The central concept of the Helmholtz model is that opposite charges are equally distributed on both sides of the interface; furthermore, this structure can be equivalent to a flat capacitor (Fig. 15.5a).

422

Y. Liu et al.

Fig. 15.5 Models of the EDL. (a) Helmholtz model; (b) G-C model; (c) G-C-S model; (d) modified G-C-S model

Therefore, the relationship between the unilateral charge density (σ) and the potential between two layers (V ) can be determined by Eq. (15.1) [35]: σ¼

εr ε0 V d

ð15:1Þ

where εr and ε0 are the electrolyte and vacuum dielectric constant, respectively, and d represents the effective thickness of the electrical double layer. Subsequently, the capacitance (Cd) of the capacitor can be expressed as Cd ¼

∂σ εr ε0 ¼ d ∂V

ð15:2Þ

So far, the Helmholtz model has successfully abstracted a general electrochemical scene into two basic formulas. However, there is an apparent defect in this model: it can be inferred from Eq. (15.2) that Cd is a constant value, while in the experimental observation, Cd will be affected by factors including electrolyte concentration and relative potential [65]. Two apparent trends can be observed in a typical capacitance curve in the double-layer region (Fig. 15.6): (a) Cd presents a V-shaped symmetric distribution relative to the potential V and (b) the value of Cd enlarges with the increase of electrolyte concentration. To optimize the Helmholtz model, the G-C model (Fig. 15.5b) was proposed by introducing a diffuse layer model of the DL [100]. Near the electrode, the charges are strictly distributed on the surface of the electrode; while in the electrolyte, the charges will diffuse into the bulk solution far from electrode–electrolyte interfaces, due to the interaction between different ions. Therefore, when the potential difference between two sides of the interfaces is large, more ions will be compressed near the electrode; while the ions can also reach charge equilibrium with the electrodes in a small space when the electrolyte concentration is high. The improvement of the G-C model changes the effective thickness (d ) in the Helmholtz model from a constant to a variable, thus interpreting the trend of Cd more appropriately.

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

423

Fig. 15.6 Typical capacitance vs. a potential difference (Helmholtz model) in the double-layer region [10]

Nevertheless, the G-C model still deviates from the actual situation. A typical G-C model for Cd prediction is confronted with three major issues: (a) at both ends of the curve, Cd approaches infinity, which is contrary to the actual tests; (b) the predicted value of Cd is much higher than the measured value; (c) the G-S model fails for highly charged DLs. In the G-C model, the deviation of Cd prediction occurs because the charge is abstracted into a mass point, which is a standard method in physics. Under the large potential difference, these “point charges” will be infinitely close to the electrode surface; therefore, the distance between positive and negative charges will approach zero, resulting in Cd approaching infinity. Later, Otto Stern proposed the Gouy-Chapman-Stern (G-C-S) model by combining the Helmholtz model and G-C model and considering the size of ions (Fig. 15.5c) [78]. In the G-C-S model, there are two ion distribution regions at the electrode–electrolyte interface: some ions adhere to the electrode as suggested by Helmholtz, forming a Helmholtz layer, while others form the Gouy-Chapman diffusion layer. In the diffusion layer, the capacitance (Cdiff) is produced by the action of thermal motion of electrolyte ions; in the Stern layer, ions are absorbed on the electrode surface and result in a constant capacitance (CH), which is independent of the potential difference. Therefore, the capacitance of the whole double layer (Cdl) can be expressed by Eq. (15.3): 1 1 1 ¼ þ C dl CH Cdiff

ð15:3Þ

According to Eq. (15.3), the value of Cdl is always smaller than that for CH and Cdiff. When the potential difference is low, the value of Cdiff is very small, and Cdl is mainly affected by Cdiff and has a V-shaped curve versus potential difference; when under high potential difference, the value of Cdiff is large, its contribution to the Cdl is negligible. In this case, the value of Cdl approaches CH. However, there remain issues with the G-C-S model, such as a) the ions in the electrolyte will be surrounded by solvent ions to form solvated ions; b) if the adsorption force at the interface

424

Y. Liu et al.

surpasses the electrostatic force, then even the similar charges can be stable at the interface. In 1947, David C. Grahame modified the G-C-S model and suggested that some ions or uncharged substances could penetrate the Stern layer [29]. In the modified GC-S model, the term “specifically absorbed ions” refer to the ions losing their solvation shell and directly contacting with the electrode when approaching the electrode. As shown in Fig. 15.5d, the modified model suggested the existence of three regions, including the inner Helmholtz plane (IHP), outer Helmholtz plane (OHP), and diffuse layer. The IHP refers to the closest distance of specifically adsorbed ions, and the OHP refers to the distance of non-specifically absorbed ions. Although the theories and models are increasingly accurate and sophisticated, debate and further exploration continues. Some recent studies revealed that the capacitance of EDLCs is also related to the active specific surface and porosity of electrode materials, wettability between electrode and electrolyte and the acidity and alkalinity of the electrolyte solution, etc [2, 42, 87]. Currently, most of the electrode materials suitable for EDLCs are carbon materials with excellent electrical conductivity, such as graphene, porous carbon, and carbon nanotubes [37, 67, 80].

15.3.2 Pseudocapacitive Supercapacitors (SCs) The energy storage of pseudocapacitive SCs is realized through the highly reversible chemisorption or redox reactions involving electron transfer, which will trigger underpotential deposition (UPD) on the surface or the (quasi-) two-dimensional space in the bulk phase of the electrochemically active materials. Like rechargeable batteries, the Faradaic process occurs on the surface of electrodes for SCs. The electrode reactions in the rechargeable battery involve the transfer of localized valence electrons, which follows the Nernst equation. In contrast, charge storage in pseudocapacitive SCs is achieved by the transfer of delocalized valence electrons, and the corresponding electrode potential is proportional to the electric charge passing through the electrode. A typical charging and discharging process for a transition metal oxide in the aqueous electrolyte are shown as follows: In acidic electrolyte: MOx + H+ + e , MOx  1(OH) In alkaline electrolyte: MOx + OH  e , MOx(OH) Besides, adding redox reactive ions into the electrolyte can also increase pseudocapacitance [50]. For instance, Ren et al. applied a redox-active electrolyte (1 M H2SO4 electrolyte adding 0.8 M Fe2+/Fe3+ additives) in a supercapacitor with a porous nanoflower polyaniline (PANI) electrode and reported a high capacitance (1062 F g1 at 2 A g1), energy density, and cycling stability (93% at 5 A g1 over 10,000 cycles) [70]. In the as-assembled SC, two redox systems are originating from: (a) the conjugated double bonds in the polymer networks (Eqs. (15.4) and (15.5)) and (b) a Fe2+/Fe3+ active redox couple (Eq. (15.6)).

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

425

   Pxþ m Ax þ xe , Pm þ xA

ð15:4Þ

þ  þ Pyþ m C y þ ye , Pm þ yC

ð15:5Þ

Fe3þ þ e , Fe2þ

ð15:6Þ

where Pm refers to PANI, m is the polymerization degree; A and C+ refer to the anions and cations, respectively. Generally, pseudocapacitive SCs possess superior specific capacitance and energy density compared to EDLCs. However, the bulk phase of electrode materials or electrolyte composition will change, with the occurrence of redox reactions, especially the pseudocapacitive effect on the surface of the electrode. Therefore, material fatigue damage is inevitable during long-term charge and discharge cycles, resulting in poor cycle life. Therefore, the lifespan of pseudocapacitors is shorter than that of EDLCs, while still longer than the batteries. Early studies on pseudocapacitive electrodes focused on transition metal oxides (TMOs); however, the low electrical conductivity and unsatisfactory Faradaic charge storage capability of TMOs limited their application. One common approach to gain a high-performance pseudocapacitive electrode is to grow a layer of active material with the capability of pseudocapacitive or battery-type charge storage on the surface of EDLC materials. This method has been reported in many studies. At present, the materials widely studied as the electrode candidates for pseudocapacitive SCs include hetero-atom doped (N, P, O, S, B, etc.) carbon materials, transition metal oxides/hydroxides (MnO2, NiOOH, etc.), two-dimensional transition metal carbides and nitrides (MXene), and conductive polymer materials (polyaniline, polypyrrole, etc.), among others.

15.3.3 Hybrid Supercapacitors (SCs) According to Chen et al., there are three types of charge storage mechanisms: non-Faradaic capacitive storage (NFCS) for EDLC storage, capacitive Faradic storage (CFS) for pseudocapacitive storage, and non-capacitive Faradaic storage (NCFS) for battery-type storage [97]. The effective combination of an EDLC electrode and a pseudocapacitive or battery electrode is a promising methodology to improve the energy density, power density, and cycle life of SCs, as well as significantly broadening the operation potential window [17]. In early studies, hybrid SCs were given various names, such as asymmetric SCs and hybrid ion capacitors, which may lead to confusion. Therefore, based on Chen’s definition, this chapter classifies different types of hybrid SCs based on the charge storage mechanisms of the electrodes (Fig. 15.1 and Table 15.1) [97].

426

Y. Liu et al.

Table 15.1 Classification of supercapacitors based on the charge storage mechanisms EDLCs

Pseudocapacitve

Hybrid SCs

Rechargeable

SCs

Batteries

Electrode 1

NFCS

CFS

NFCS

NFCS

CFS

NCFS

Electrode 2

NFCS

CFS

CFS

NCFS

NCFS

NCFS

15.4

Electrolyte

The electrolyte is one of the critical components of SCs and provides ionic conductivity to facilitate charge compensation on each electrode. One of the substantial challenges for SCs is the unsatisfactory energy density: for a typical EDLC, the energy density is less than 5 W h kg1; for pseudocapacitive SCs or hybrid SCs, their energy density is commonly less than 15 W h kg1 [105]. This energy density cannot fully satisfy current commercial and industrial demands. The relationships among energy density of SCs (E), capacitance (C), and the operation voltage (V ) can be expressed by Eq. (15.7): 1 E ¼ CV 2 2

ð15:7Þ

According to Eq. 15.7, increasing either capacitance or both capacitances and operation voltage is an effective method to enhance the energy density of SCs. In addition to the energy density, the power density (P) is also closely related to the operating voltage, as shown in Eq. (15.8): P¼

1 V2 4W TS Rcell

ð15:8Þ

where Rcell relates to the equivalent series resistance (ESR) of SCs (Ω). It is worth noting that increasing the operating voltage will be more effective than increasing the capacitance to enhance both the energy and power density because the energy/ power density is proportional to the square of the operation voltage. Within the stable electrochemical voltage range of the electrode materials, the operation voltage is mostly dependent on the electrochemical stable potential window (ESPW) of the electrolyte. Figure 15.7 demonstrates the typical electrochemical stable potential window of commonly used electrolytes, including aqueous electrolyte, organic electrolyte, ionic liquid, and (quasi-)solid-state electrolyte. The application of electrolytes with higher ionic conductivity is essential to reduce the ESR of SCs (Rcell). Generally, the ionic conductivity (σ, S cm1) can be expressed by Eq. (15.9):

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

427

Fig. 15.7 Electrochemical stable potential window of different types of electrolytes

σ¼

  nðZeÞ2 D E exp  a kT kT

ð15:9Þ

where n refers to the concentration of the carrier, Ze is the charge of the carrier (for protons, Z ¼ 1), D relates to the self-diffusion coefficients, and Ea refers to the activation energy of ion diffusion. Adopting electrolytes with high self-diffusion coefficients and low activation energy, as well as increasing the number of charge carriers, will enhance the ionic conductivity. In addition to a wide potential window and high ionic conductivity, an ideal electrolyte possesses the following characteristics: (a) high (electro)chemical stability and wide operation range; (b) well-matched with other SC components (electrodes, current collectors, separators, etc.); (c) low flammability, volatility, and toxicity; (d) low cost; and (e) environmentally friendly nature. It is very challenging to satisfy these requirements simultaneously because different types of electrolytes possess their intrinsic strength and weakness. The following content in this section provides an overview of the different types of electrolytes to construct a basic understanding of the development of the electrolytes for SCs.

15.4.1 Aqueous Electrolyte Generally, aqueous electrolytes possess significant merits such as low cost and high ionic conductivity (e.g. ~ 0.8 S cm1 for 1 M H2SO4 at 298.15 K) [28]. Meanwhile, the small hydrated ion (2.80 ~ 4.28 Å) of the aqueous electrolyte can quickly immerse in the micropores of the electrode, so that the large size distribution and specific surface area of the electrode materials can be effectively used [105]. The aqueous electrolyte can be divided into three categories: acidic, alkaline, and neutral electrolytes. Currently, the most commonly used acidic, alkaline, and neutral

428

Y. Liu et al.

aqueous electrolytes are H2SO4 solution, KOH solution, and alkali metal salt solution (e.g. Na2SO4, KCl), respectively [1, 46, 86]. Compared with neutral aqueous electrolytes, strong acid and alkaline aqueous electrolytes provide higher ionic conductivity and lower internal resistance; however, their high corrosiveness may lead to the structural instability of SCs, which increase the risk of electrolyte leakage and cause severe environmental pollution. However, aqueous electrolytes are a relatively unpopular choice for current commercial SCs, due to a critical issue: the narrow ESPW limited its application under high operation voltage. This is because the ESPW for pure water is only 1.23 V vs. standard hydrogen electrode (SHE); under high operating voltage, the hydrogen evolution reaction (HER) and oxygen evolution reaction (OER) can occur, leading to water splitting. Although the ESPW of aqueous electrolytes can be elevated via adding salt/acid/alkali or applying “water-in-salt” aqueous electrolytes, the effects were quite limited (< 2.2 V vs. SHE) [25].

15.4.2 Organic Electrolyte Organic electrolytes currently dominate the commercial SC market due to several merits, including a) high ESPW (2 ~ 4 V vs. SHE), b) wide operating condition range (temperature and humidity), c) low corrosiveness to SC components, and d) high electrochemical stability. Additionally, the use of organic electrolytes allows low-cost materials (e.g. Al) to be used in current collectors and packages. However, organic electrolytes also face some substantial challenges. Firstly, compared to the aqueous electrolytes, organic electrolytes generally possess larger ionic radii, lower ionic conductivity, and higher internal resistance, which demands electrode materials with larger pore size and leads to the low surface area utilization of the electrode materials. Meanwhile, higher voltage and current density may lead to a lower ion concentration in the organic electrolyte, thus resulting in low efficiency and potential safety issues. Therefore, although organic electrolytes possess a high ESPW (2 ~ 4 V vs. SHE), their operation voltage is generally limited to 2.5 ~ 2.8 V. Besides, SCs using organic electrolytes require a complicated fabrication process in a tightly controlled environment due to the volatility, flammability, and toxicity of the organic electrolyte. Take lithium-ion-based SCs for example, the commonly used organic electrolyte salts include lithium chloride/perchlorate, lithium bis(trifluoromethanesulfonyl)imide (LiTFSI), lithium hexafluorophosphate (LiPF6), and tetraethylammonium tetrafluoroborate (TEABF4). The widely studied solvents are ethylene carbonate (EC), propylene carbonate (PC), dimethyl carbonate (DMC), ethylene dicarbonate (EDC), acetonitrile (ACN), N, N-dimethylacetamide (DMF), γ-butyrolactone (GBL), etc [69, 72, 98]. A single solvent system may have several issues. For instance, Japan has banned the use of commercial electrolytes based on ACN solvent in SCs, due to their low flash point and high toxicity [76]. To noticeably change both the viscosity and ionic conductivity of organic electrolytes, research has focused on

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

429

solvents based on a combination of different organic solvents or additives, such as PC-TMC, EC-DMC, and EC-EDC [82, 102]. The further development of electrolyte salts and solvents mainly focuses on improving ionic conductivity and reducing viscosity, to enable the electrolytes to maintain excellent electrochemical performance under extreme conditions (e.g. high voltage and low temperature) [56].

15.4.3 Ionic Liquid Ionic liquid electrolytes are salts that are liquid within a specific temperature range (< 373.15 K). Ionic liquids are considered as suitable alternative electrolytes owing to the following advantages: (a) high tunability due to the variety in anions and cations, (b) high thermal and (electro)chemical stability, and (c) negligible volatility and non-flammability. Many SCs based on ionic liquid electrolytes can achieve operation voltages above 3 V; under this voltage, some commercial electrolytes (e.g. ACN and PC) may face severe electrochemical decomposition issues [52, 105]. Researchers can effectively regulate and optimize ionic liquid electrolytes to satisfy the demands of applications in SCs, in terms of operation condition (temperature, voltage) and internal resistance [26]. Unfortunately, ionic liquid electrolytes are currently unavailable for commercial use, resulting from their unsatisfactory characteristics such as low ionic conductivity, undesirable viscosity, and high cost. Even for ionic liquids with relatively high ionic conductivity, such as 1-butyl-3-methylimidazolium tetrafluoroborate ([EEIM] [BF4], 0.014 S cm1 at 298.15 K), their ionic conductivity is still much lower than that for organic electrolytes (e.g. ~ 0.059 S cm1 for TEABF4/ACN at 298.15 K) and aqueous electrolytes (e.g. ~ 0.8 S cm1 for 1 M H2SO4 at 298.15 K) [28, 41, 72]. Besides, compared to aqueous and organic electrolytes, the viscosity (40 ~ 220 cp) of ionic electrolytes is much higher, which leads to poor electrochemical performance under high current densities [18]. Recent studies suggested that SCs using ionic liquid electrolyte can be improved based on two key approaches: (a) modification of anions or cations to possess similar structure to dispersants and prevent cations from aggregation, thus further improving the ionic conductivity of the electrolytes [47]; and (b) further optimization of the pore structure of electrode materials and wettability between electrodes and electrolytes [24].

15.4.4 (Quasi-)Solid-State Electrolyte With the development of wearable and printable electronic devices, flexible electrochemical energy storage devices have received increasing attention [21]. Besides, to serve as ionically conductive media, (quasi-)solid-state electrolytes also work as

430

Y. Liu et al.

separators, thus simplifying the fabrication process of SCs without electrolyte leakage issues. At present, major types of (quasi-)solid-state electrolytes are based on polymer electrolytes, which can be divided into three categories: inorganic solid electrolytes (ISEs), solid polymer electrolytes (SPEs), and gel polymer electrolytes (GPLs). ISEs are a particular type of inorganic materials in the crystalline or glass state, which conduct ions by diffusion through the lattice [9]; SPEs are composed of polymers and solvent-free salts, which conducts ions through the polymer chain [61]; in contrast, GPEs consist of polymers, aqueous electrolytes, or conductive salts dissolved in solvents [57]. Generally, only GPEs can be applied in flexible devices, because ISEs and SPEs are not bendable and possess almost no flexibility. Among these three electrolytes, GPEs are the most commonly used electrolytes for (quasi-)solid-state SCs because of their relatively high ionic conductivity (103 ~ 102 S cm1) compared to ISEs (104 ~ 102 S cm1) and SPEs (105 ~ 104 S cm1) [12, 22, 57]. However, GPEs are confronted with certain limitations: (a) their weak mechanical strength may lead to the internal short circuit, and (b) the presence of water narrows the operation voltage window. Besides, a common challenge for (quasi-)solid-state electrolytes is their limited contact surface area with the electrode materials, especially for porous nanomaterials, which will lower the utilization of the electrode active materials. This issue may increase the equivalent series resistance (ESR) of SCs and lead to a reduction in rate performance and poor specific capacitances. At present, the development of (quasi-)solid-state electrolytes is not mature, and no such electrolyte can satisfy all of the following key requirements: (a) high ionic conductivity, (b) high thermal and (electro)chemical stability, and (c) sufficient mechanical strength.

15.5

Carbon Material Electrodes

Various types of carbon materials are currently used for SC electrodes, due to their low cost, high surface area, excellent electrical conductivity, etc. Since it was first used in SCs in 1957, researchers have made numerous efforts to enhance the electrochemical performance of carbon materials, which mainly focus on surface modification, pore structure and surface area optimization, hetero-atom doping, etc. [32, 53]. This section will provide a brief introduction covering the primary carbon materials, including activated carbon (AC), carbon nanotubes (CNTs), and graphene.

15.5.1 Activated Carbon (AC) AC is the earliest widely used active electrode material for SCs, owning to the large specific surface area (> 1000 m2 g1), high electrical conductivity, and low cost

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

431

[100]. AC can be obtained through the carbonization of carbon-rich organic precursors under the protection of inert atmosphere (N2, Ar), and physical or chemical activation. These precursors originate from a wide range of sources, including (a) natural renewable resources (e.g. cellulose, coconut shells, discarded rubber, and even human hair) [69], (b) fossil fuels and their derivatives (e.g. coal, asphalt, and coke) [5], and (c) synthetic precursors (e.g. polymers, metal-organic frameworks, and covalent organic frameworks) [88]. The mainstream activation methods include physical activation, chemical activation, and microwave activation. Physical activation usually takes place at 600 ~ 1200  C in CO2 or a water vapor atmosphere [44]. Usually, it can be divided into two steps: a) precursor pyrolyzed at 400 ~ 900  C in an inert gas atmosphere, and b) oxidation at 350 ~ 1000  C to increase the porosity and specific surface area [73]. The chemical activation only contains one step: the precursors are mixed with chemical substances such as strong base (e.g. KOH, NaOH), strong acid (e.g. H2SO4, H3PO4), carbonate (e.g. K2CO3, Na2CO3), chloride (e.g. ZnCl2, FeCl2), and activated at 400 ~ 900  C [48]. The microwave activation causes internal dipole rotation and ion transfer by heating the precursor with microwave radiation, thus generating the activated carbon. All three activation methodologies have their advantages and disadvantages. For instance, the significant merits of chemical activation are lower operating temperature, less time consumption, and higher specific surface area (more suitable for EDLCs applications). Although some chemical agents (e.g. KOH) can provide hierarchically porous structure and high specific surface area (> 3000 m2 g1), their potential hazards and toxicity limited further applications [85]. Additionally, heat is transferred from the surface of the precursors to the interior during physical and chemical activation, thus inevitably creating a temperature gradient. This temperature gradient will result in the amorphous and heterogeneous microstructure of the prepared AC. In contrast, microwave activation transfers heat from the interior to the surface, thus allowing for volumetric heating and delivering energy evenly to the bulk materials. In practical applications, it is necessary to combine multiple activation methods to integrate their advantages and avoid shortcomings. Besides activation methods, the chemical and physical properties of the AC are also related to many other factors, such as the composition of the precursor, activation temperature, and activation time. For instance, a longer activation time may lead to higher porosity and broader pore size distribution [33]. The heteroatom doping and pore structure optimization are the most common strategies for the performance enhancement of porous carbon-based SCs. However, the current understanding of how these two strategies influence the ion dynamics, charge storage mechanisms, and electrochemical performance is still minimal. Many experimental and computational studies have demonstrated that hetero-atom doping in carbon materials can induce higher charge delocalization and donor density near Fermi levels, expand interlayer spacing, and enhance the wettability of active materials [68, 93]. However, the effect of different elements on the properties of carbon materials may vary due to the differences in atomic size and electronegativity. In general, the mainstream views on the mechanisms and functionalities of

432

Y. Liu et al.

Fig. 15.8 Carbon materials with different hetero-atom doping. (a) pristine, (b) O-doping, (c) N-doping, (d) P-doping, (e) B-doping

widely studied doping elements (O, N, P, and B) are (a) O-doping (Fig. 15.8b) can effectively change the surface wettability and provide more active sites; (b) N-doping (Fig. 15.8c) can distort the structure of carbon materials, thus providing additional defects and active sites; and (c) similar to nitrogen, P-doping (Fig. 15.8d) can distort the carbon structure to a greater extent due to the larger radius. As a result, it will achieve an electron-donating ability, better electron delocalization, more defects, and higher wettability; d) B-doping (Fig. 15.8e) will result in a shift of Fermi level to the conduction band, thus modulating the electronic structure of carbon materials and providing a higher electrical conductivity [13, 43]. Also, compared to single heteroatom doping, some recent studies have shown that double or multiple hetero-atom doping further optimizes the nano-/microstructure of the electrodes, thus significantly enhancing the electrochemical performance [49].

15.5.2 Carbon Nanotubes (CNTs) Both single-walled CNTs (SWCNTs) and multi-walled CNTs (MWCNTs) are promising candidates as electrode materials, due to their unique structure, excellent electrical conductivity, and good stability (chemical, electrochemical, and thermal). Although CNTs have a large theoretical specific surface area (1315 m2 g1) and high electrical conductivity (350 ~ 55,000 S cm1), their specific capacitance is only 20 ~ 80 F g1, because there are relatively few micropores that facilitate charge transfer [51, 66]. For MWCNTs, although their volume of micropores can be increased by activation, the specific capacitance is still inferior to that of AC.

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

433

There are many reports of CNT-based electrode candidates with remarkable specific capacitance; however, a critical obstacle to commercialization is their low packing density (100 ~ 250 kg m3). Densification of CNT-based electrodes to an industrial level without severely compromising ion transport is a necessary consideration to realize commercial SCs. Wei et al. proposed a strategy to shorten the transport distance of electrolyte ions via hierarchical structures, which is constructed by using 90% granulated double-walled CNTs (DWCNTs) or triple-walled CNTs (TWCNTs) and 10% MWCNTs (10μm long) as the linker (Fig. 15.9) [104]. After CO2 etching, the as-fabricated membrane with hierarchical structure demonstrates remarkable packing density (420 kg m3) and high specific surface area (871 m2 g1). Owning to the excellent mechanical properties and high electrical conductivity, CNTs are often used as supporting substrate materials to grow pseudocapacitive or EDLC electrode materials. Huang and co-workers designed and synthesized the amorphous MnO2@MWCNTs fiber, in which MnO2 was incorporated into MWCNTs fiber uniformly [75]. The as-fabricated fiber-based supercapacitor exhibits a high power density (1.5 mWh cm3) and a remarkable specific capacitance of 8 F g1 at 1 A cm3 (~12.9 times higher than that of MWCNT fiber-based devices, 0.62 F g1). The excellent electrochemical properties are as follows: (a) compared with crystalline MnO2, the amorphous MnO2 possesses a highly disordered structure, which is more conducive to rapid ion-diffusion and fast electrode kinetics; (b) the uniformly distributed MnO2 nanoparticles inhibited the stacking of MWCNTs, further facilitating the ion transportation; and (c) the wellaligned MWCNT network offers a higher electron transport performance.

MWCNTs

Granulated DWCNTs/TWCNTs DWCNTs CO2

850°C TWCNTs

Acid washing

Dispersion & filtration

Granulated DWCNTs/TWCNTs

MWCNT network

Fig. 15.9 Schematic diagram of the fabrication of CNT membrane consisted of dominant granulated DWCNTs/TWCNTs and long MWCNTs as the backbone. Reproduced with permission from Elsevier [104]

434

Y. Liu et al.

15.5.3 Graphene Since Geim and Novoselov separated graphene from graphite by the mechanical stripping method in 2004, graphene has attracted extensive attention from researchers in various fields due to its excellent electrochemical properties, strong thermal and (electro)chemical stability, high electron mobility (15,000 cm2 V1 s1), large specific surface area (2630 m2 g1), “intriguing” thermal conductivity (~5000 W m1 K1), and “muscular” mechanical strength (~ 1 TPa) [64, 71]. Graphene is a two-dimensional planar structure material composed of sp2 hybridized monolayer carbon atoms arranged in a honeycomb structure and can be reciprocally converted with CNTs, fullerenes, and other carbon materials. There are many methods to obtain graphene, including chemical vapor deposition (CVD), carbon dioxide reduction, supersonic spray, ion implantation, and laser treatment. [80]. Chemically exfoliating graphite to graphene oxides (GOs), followed by a controlled reduction of GOs to graphene via reductants (e.g. hydrazine hydrate), is generally considered to be the most efficient and economical method [36]. Theoretically, the specific capacitance of single-layer graphene is ~550 F g1, when the G entire surface area is fully utilized [91]. However, the presence of π–π bonds can cause severe aggregation, thus reducing specific surface area and suppressing the actual capacitance. Besides, there are many reports of the incorporation of graphene-based and pseudocapacitive materials as a promising strategy to elevate the electrochemical performance of electrodes [16]. However, increasing active materials loading may inhibit electron transfer and ion diffusion, resulting in a severe decline in the capacitance of SCs [38]. To address both of these challenges, Yao and colleagues synthesized a graphene aerogel with ultra-high MnO2 loading (182.2 mg cm2) via 3D printing (Fig. 15.10) and achieved a record-high areal capacitance of 44.13 F cm2 [96]. It is worth noting that the capacitance of this unique 3D structure is not limited by ion diffusion at high active materials loading, which is impossible for conventional bulk electrodes.

15.6

Transition Metal Compounds Electrode Materials

According to the definition, the theoretical specific capacitance (CT) of transition metal compounds (TMCs) can be determined by (Eq. 15.10): CT ¼

Q It ¼ mV mV

ð15:10Þ

where I refers to the applied current during the charge/discharge (A), m is the total weight of the electrode,t is the discharge time, and V is the operating voltage window. Also, CT can be expressed based on the molecular weight of TMCs (M, g mol1) as Eq. (15.11):

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

435

Fig. 15.10 Schematic diagram of the fabrication of 3D-printed graphene aerogel/MnO2 electrode. Reproduced with permission from Elsevier [96]

CT ¼

nF MV

ð15:11Þ

where n refers to the number of electrons transferred during the redox reaction, and F is Faraday’s constant. Therefore, the TMCs with lower molecular weight and more electron transfer usually have a higher specific capacitance. This section will provide a brief introduction to several extensively studied transition metal compound electrode candidates: transition metal oxides (TMOs), MXene, and other transition metal compounds (hydroxides, sulfides, phosphides, and selenides).

15.6.1 Transition Metal Oxides (TMOs) RuO2, as a precious metal oxide pseudocapacitive electrode with high specific capacitance and power density, is the first material studied and applied on SCs for applications in national defense and aerospace [83]. Although RuO2 is one of the electrode materials that possess the highest electrochemical performance at present, its negative characteristics such as high cost (~ 7140 $ per ton for ruthenium), toxicity, and need to use with strong acid electrolyte (H2SO4) has limited its civil use [7, 8, 101]. Therefore, some inexpensive and environmentally friendly TMO pseudocapacitive electrode materials (e.g. MnO2, Fe2O3) have received considerable attention in recent years [54]. TMO electrodes are confronted with two major challenges: (a) the unsatisfactory electrical conductivity results inactive materials being only partially utilized; (b) during the charging and discharging cycles, the recurring redox reactions on the surface of the active materials lead to poor cycling stability [95].

436

Y. Liu et al.

MnOOH

MnOOH

Beta-MnO2 MnOOH Birnessite

MnOOH/Beta-MnO2/Birnessite

Beta-MnO2 Birnessite

Beta-MnO2/Birnessite

MnOOH Beta-MnO2 Birnessite

Fig. 15.11 Schematic diagram of the synthesis of β-MnO2/parallel birnessite core/shell nanorod. Reproduced with permission from the American Chemical Society [107]

MnO2 has been widely studied for SCs, due to factors that include high theoretical capacitance (1370 F g1), low cost, and environmentally friendly properties [59]. Zhu et al. synthesized a β-MnO2/birnessite core-shell hybrid TMO, which using β-MnO2 as the core and highly ordered birnessite sheets (Fig. 15.11) [107]. The utilization of the MnO6 unit in the electrode was improved because the parallel highly ordered shell structure provides an effective transport channel for the ions in the electrolyte. In the 1 M Na2SO4 electrolyte, the as-assembled asymmetric SC exhibits an impressive specific capacitance of 657 F g1 at 0.1 A g1 (based on parallel birnessite) and a high energy density of 17.6 Wh kg1. Also, it is an effective methodology to composite TMOs with carbon materials, because (a) the charge transfer resistance during the (dis)charging processes can be effectively reduced and (b) the addition of carbon materials can enhance the electrical double-layer capacitance. Besides, some other strategies are promising to enhance electrochemical performance, such as nanostructure regulation, defects engineering, and heteroatom doping [45, 108].

15.7

Transition Metal Carbides and Nitrides (MXene)

First reported in 2011, Gogotsi et al. prepared two-dimensional layered Ti3C2 by hydrofluoric acid and ultrasonic method; a series of novel two-dimensional transition metal carbides and nitrides (MXene) have been developed and applied as SC electrode materials [62]. MXenes are defined as Mn + 1AXn, where M represents transition metals, A refers to elements in IIIA or IVA group, X represents carbon or nitrogen, and n ¼ 1, 2, or 3 [63]. In recent years, the MXene has drawn increasing attention in energy storage applications, owing to its excellent electrochemical performance originated from characteristics including high electrical conductivity, unique structure, and the feasibility of electron transfer due to the variable oxidation state of transition metal M [40].

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

437

Further increasing the interlayer spacing of MXene will promote ion (de)insertion kinetics and offer more effective surface exposure to electrolytes, thus enhancing the specific capacitance. The synthetic methods can affect the structure of MXene, for instance, etching with suitable concentration (< 6 M) hydrofluoric acid (HF) and the morphologies of MXene transferred into multilayer lamellas [39]. Compared with high concentration HF, the diluted HF allows the presence of more high-mobility water molecules between the voids of MXene, thus resulting in a larger permanent interlayer spacing and higher capacitance. Then, to obtain a graphene-like monolayer MXene, several methods are developed for the exfoliation of multilayered HF etched-MXene, such as scotch tape delamination, high power ultrasonication, and chemical intercalation delamination [40].

15.8

Other Transition Metal Compounds (Hydroxides, Sulfides, Phosphides, and Selenides)

The low electrical conductivity of TMOs limits their rate performance and energy density under high current densities. By replacing the anions (O) in TMOs with elements such as S, P, and Se, the electrical conductivity of TMCs can be significantly improved. A series of TMCs beyond TMOs have been developed and applied to SCs, such as CuS, MoS2, and TiS2. In recent years, transition metal sulfides with the structure of MCo2S4 (M ¼ Ni, Zn, and Cu) have attracted particular attention, due to their faster electron transport, higher electrochemical activity, and specific capacitance [30]. Chen and colleagues prepared graphene-coated NiCo2O4 core-shell materials (NiCo2O4@G) via a hydrothermal method, which demonstrates a high specific capacitance (1432 F g1 at 1 A g1) and a stable cycling performance (83.4% capacitance retention after 5000 cycles) [99]. The authors suggested that the functionalities of the graphene shell can be summarized in two aspects: (a) mitigate the potential volume change during the operation and (b) protect NiCo2S4 from dissolution.

15.9

Emerging Electrode Materials

In addition to traditional electrode materials introduced above, a variety of novel materials with excellent electrochemical properties have emerged in recent years. This section provides a brief introduction of several representative materials with great potential for use as electrodes for both EDLCs and pseudocapacitive SCs, including black phosphorus, metal-organic frameworks (MOFs), covalent organic frameworks (COFs), and conductive polymers (CPs).

438

Y. Liu et al.

15.9.1 Black Phosphorus (BP) As early as 1914, black phosphorus (BP) was synthesized by heating white phosphorus under high pressure (12,000 atmospheres); however, it was not until 2014 that it was rediscovered as a member of two-dimensional materials [92]. Similar to graphite, bulk BP has a black appearance, layered structure, high mechanical strength (~ 94 GPa), good density (2.69 g cm3), and electrical conductivity (~100 Sm1) (Fig. 15.12a) [79]. Also, the lamellar BP (phosphorene) has a graphene-like lattice of interlinked six-membered rings where each atom is bonded to three other atoms (Fig. 15.12b), thus forming an orthorhombic pleated honeycomb structure [15]. Phosphorene is prepared in a similar way to graphenes, such as scotch-tape delamination, liquid exfoliation in organic solvents (e.g. acetone, N-methyl pyrrolidone), and high-energy ball milling. [90]. At present, there are few studies based on BP. Through a modified electrochemical approach under ambient atmosphere, Yu et al. synthesized a 3D BP sponge based all-solid-state SC (Fig. 15.12c), which comprise ultrathin (< 4 nm), large (> 10μm) and high-quality (unoxidized) BP nanosheets as the basic unit. During the cyclic voltammetry test, the capacitance is 80 F g1 and 28 F g1 at the scanning rate of 10 mV s1 and 100 mV s1, respectively [89]. With in-depth study, researchers

Fig. 15.12 (a), (b) Structure of black phosphorus (BP); (c) schematic diagram of the synthesis of 3D BP sponge all-solid-state SCs. Reproduced with permission from the Royal Society of Chemistry [89]

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

439

attempt to exfoliate arsenic from the same main group with phosphorus as the electrodes for SCs, which demonstrates an impressive specific capacitance of 1578 F g1 at 14 A g1 [60].

15.9.2 Metal-Organic Frameworks (MOFs) Metal-organic frameworks (MOFs), also known as porous coordination polymers (PCPs), are organic–inorganic hybrid materials with intramolecular pores by selfassembly of organic bridging ligands and metal ions or clusters through coordination bonds. Compared with conventional porous materials, MOF materials have a series of characteristics that are conducive to the application of SCs: (a) higher porosity (e.g. 94%) and specific surface area (e.g. 7140 m2 g1 for NU-110E), due to the longer organic bridging ligands; (b) controllable chemical and physical properties and porous structure, owning to the diversity of skeletons and the high modifiability of organic ligands and metal centers; (c) the highly crystallized compounds facilitate the determination of the precise spatial structure by single crystal or polycrystalline diffraction [23, 27]. However, MOFs and their derivatives also have some common weaknesses, such as unsatisfactory (electro)chemical stability and relatively low electrical conductivity (107 ~ 102 S cm1) [94]. Currently, the following types of materials applied in SCs can be derived from MOFs: bimetallic oxides, metal oxide composites, carbon composites, and highly conductive materials. Some MOFs with good electrical conductivity can be directly used as electrode materials. By using Ni3(2,3,6,7,10.11-hexaiminotriphenylene)2 (Ni3(HITP)2) as the sole electrode for EDLCs (Fig. 15.13), Dennis et al. reported the first example of SCs entirely from MOFs without conductive additives or binders [74]. Some characteristics of Ni3(HITP)2 are conducive to its applications in energy storage, such as high electrical conductivity (50 S cm1), high specific surface area (630 m2 g1), and regular open channels (1.5 nm one-dimensional cylindrical channels). However, there is a dilemma that higher porosity tends to result in relatively lower electrical conductivity. To enhance electrical conductivity without sacrificing porosity, a practical methodology is to composite MOFs with carbon materials [103]. Besides, using MOFs as precursors, the bimetallic oxides can be efficiently designed and prepared through carbonization. Compared with single transition metal oxides, transition bimetallic oxides are equipped with superior electrochemical performances, such as (a) the co-existence of two different cations in a single crystal structure produces more electrons than a single crystal, thus enhancing the electrical conductivity; (b) more redox-active sites can be provided by the synergistic effects of transition bimetallic oxides [6].

440

Y. Liu et al.

Fig. 15.13 (a) Molecular structure of Ni3(HITP)2; (b) relative size of pore and electrolyte in the ideal Ni3(HITP)2 [74]

15.9.3 Covalent Organic Frameworks (COFs) As a novel electrode candidate for SCs, the earliest covalent organic framework (COF) was designed and synthesized by Yaghi and co-workers in 2005 [20]. Like MOFs, the COFs are equipped with high specific surface area, controllable porous structure, and excellent molecular designability. COFs can easily introduce the heteroatoms into the frameworks and change the physical and chemical properties of the materials, such as boron condensation (Fig. 15.14a) and triazine-based trimerization (Fig. 15.14b). Although COFs can be used as electrodes for both EDLCs and pseudocapacitive SCs, their poor electrical conductivity retards their deployment in SCs. Banerjee and co-workers successfully synthesized a redox-active TpOMe-DAQ thin sheet via an imine condensation, based on 2,4,6-trimethoxy-1,3,5-benzene-tri carbaldehyde (TpOMe) and 2,6-diaminoanthraquinone (DAQ) [31]. The pristine TpOMe-DAQ electrode achieved an excellent areal capacitance of 1600 mF cm2 at 10 mA cm2 and cyclic stability (> 100,000 cycles) in the 3 M H2SO4 aqueous electrolyte, originated from the existence of interlayer hydrogen bonds and the redox activity due to the quinone/hydroquinone transformation (DAQ amine).

15.9.4 Conductive Polymers (CPs) Conductive polymers (CPs) or intrinsically conductive polymers (ICPs) are electrically conductive organic polymers. Generally, the single and double bonds appear alternately in the backbone of CPs, thus forming conductive conjugated π-π bonds. When conductive polymers are used as electrode materials for SCs, they generally possess the following advantages: (a) high electrical conductivity

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

441

Fig. 15.14 COFs synthesis. (a) Boron condensation; (b) triazine-based trimerization

(102 ~ 105 S cm1), (b) wide voltage windows, (c) controllable electrochemical performance, (d) high theoretical specific capacitance, (e) excellent stability and reversibility, and (f) environmentally friendly nature [81, 106]. As shown in Fig. 15.14, the extensively studied CPs include poly(thiophene)s (PTH), polyaniline (PANI), polypyrrole (PPy), poly (p-phenylene vinylene) (PPV), and poly(3,4-ethylene dioxythiophene) (PEDOT) [77]. Among these CPs, PANI, PTH, PPy, and their derivatives are the most widely applied pseudocapacitive electrode materials for SCs. CPs store charges mainly through the redox reaction between electrons in the material and ions in the electrolyte. Generally, CPs possess an excellent microporous structure, good wettability, and a much higher electrical conductivity compared to TMCs. However, it is still essential to further enhance their electrical conductivity, since the electrochemical reactions occur throughout the CPs-based electrode during charging and discharging. Meanwhile, CPs may swell and shrink during the operation, which leads to unsatisfactory electrochemical cycle lives. To solve these issues, there are many promising strategies: (a) incorporating carbon nanomaterials to reduce the probability of deformation and further improve the electrical conductivity; (b) optimizing the structure and morphology of electrode materials to increase the specific surface area and shorten the electron transport pathways; and (c) assembling the hybrid SCs using EDLC electrode as the anode and CP as the cathode, to extend the device cycle life.

15.10

Conclusion

This chapter provides a succinct introduction and discussion of the development history, fundamental mechanisms and theories, and extensively studied electrolytes and electrodes of SCs. Although numerous encouraging achievements have been

442

Y. Liu et al.

attained in this field, it is still at an early stage to realize the employment of highperformance SCs. The development and perspectives of SCs can be briefly summarized as follows: (a) The main factors affecting the performance of SCs by electrolytes are operation range (temperature and voltage), ion mobility, ionic conductivity, thermal and (electro)chemical stability, and compatibility with electrodes. Although electrolytes for SCs have achieved much encouraging progress in recent years, the inherent challenges for each type of electrolyte have not yet been solved. Currently, there is no such electrolyte that can meet all requirements. For example, aqueous electrolytes have high ionic conductivity, while suffering from their narrow ESPV. More efforts are required in energy storage mechanism studies, advanced characterization techniques, and optimization of the electrode–electrolyte interface. (b) Further optimization of conventional materials or development of novel materials is the key to improve the (electro)chemical performance of SCs. An ideal electrode candidate is expected to be highly conductive, low cost, favorable structure, high specific surface area, high (electro)chemical and thermal stability, and environmentally friendly properties; materials with all these conducive characteristics are very challenging to develop. Currently, many effective strategies have been developed, such as structural regulation, hetero-atom doping, and composite materials engineering. (c) Compared with EDLCs and pseudocapacitive SCs, hybrid SCs combine the characteristics of the rechargeable batteries and capacitors, which is promising to achieve a high energy density and power density simultaneously. Current studies on hybrid SCs are limited to the laboratory stage, which is mainly focusing on the selection and optimization of the electrolyte and electrode materials. This chapter suggested that the design and optimization of new cell configuration, as well as the further tests and verification based on industrial scenarios and application requirements, are the research direction for hybrid SCs development. Acknowledgments The authors would like to thank the Engineering and Physical Sciences Research Council (EPSRC, EP/V027433/1, EP/533581/1), the Royal Society (RGS\R1\211080; IEC\NSFC\201261) and Faraday Institution (EP/S003053/1) Degradation project (FIRG001) for financial support. Author Contribution Mr. Yiyang Liu visualized and wrote the original draft. Prof. Paul R. Shearing wrote, reviewed, and edited the manuscript. Dr. Guanjie He conceived, supervised, wrote, reviewed, and edited the manuscript. Prof. Dan J. Brett helped in funding acquisition, supervision, writing, review, and editing.

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

443

References 1. Abbas Q, Pajak D, Frackowiak, E. & Béguin, F. (2014) Effect of binder on the performance of carbon/carbon symmetric capacitors in salt aqueous electrolyte. Electrochim Acta 140:132–138. https://doi.org/10.1016/j.electacta.2014.04.096 2. Abouelamaiem DI, He G, Neville TP, Patel D et al (2018a) Correlating electrochemical impedance with hierarchical structure for porous carbon-based supercapacitors using a truncated transmission line model. Electrochim Acta 284:597–608. https://doi.org/10.1016/j. electacta.2018.07.190 3. Abouelamaiem DI, He G, Parkin I, Neville TP et al (2018b) Synergistic relationship between the three-dimensional nanostructure and electrochemical performance in biocarbon supercapacitor electrode materials. Sust Energy Fuels 2(4):772–785. https://doi.org/10.1039/ c7se00519a 4. Abouelamaiem DI, Rasha L, He G, Neville TP et al (2018c) Integration of supercapacitors into printed circuit boards. J Energy Storage 19:28–34. https://doi.org/10.1016/j.est.2018.06.016 5. Ahmadpour A, Do DD (1996) The preparation of active carbons from coal by chemical and physical activation. Carbon 34(4):471–479. https://doi.org/10.1016/0008-6223(95)00204-9 6. An C, Zhang Y, Guo H, Wang Y (2019) Metal oxide-based supercapacitors: progress and prospectives. Nanoscale Advances 1(12):4644–4658. https://doi.org/10.1039/c9na00543a 7. Anon (n.d.-a) Ruthenium. Precious Metals Management. [Online]. https://pmm.umicore.com/ en/prices/ruthenium/. Accessed 29 Oct 2020. 8. Anon (n.d.-b) The charge of the ultra-capacitors. IEEE Spectrum. [Online]. https://spectrum. ieee.org/transportation/advanced-cars/the-charge-of-the-ultra-capacitors. Accessed 23 Oct 2020b. 9. Bachman JC, Muy S, Grimaud A, Chang HH et al (2016) Inorganic solid-state electrolytes for Lithium batteries: mechanisms and properties governing ion conduction. Chem Rev 116(1):140–162. https://doi.org/10.1021/acs.chemrev.5b00563 10. Bard AJ, Faulkner LR (2003) Electrochemical methods: fundamentals and applications, 2nd edn. Wiley, New York 11. Becker HJ (1954) United States Patent Office low volTAGE electrolytic capacitor 12. Bi H, Sui G, Yang X (2014) Studies on polymer nanofibre membranes with optimized coreshell structure as outstanding performance skeleton materials in gel polymer electrolytes. J Power Sources 267:309–315. https://doi.org/10.1016/j.jpowsour.2014.05.030 13. Bi Z, Huo L, Kong Q, Li F et al (2019) Structural evolution of phosphorus species on graphene with a stabilized electrochemical interface. ACS Appl Mater Interfaces 11(12):11421–11430. https://doi.org/10.1021/acsami.8b21903 14. Boos DL (1970) Electrolytic capacitor having carbon paste electrodes. 15. Brown A, Rundqvist S (1965) Refinement of the crystal structure of black phosphorus. Acta Crystallogr 19(4):684–685. https://doi.org/10.1107/s0365110x65004140 16. Cao X, Zheng B, Shi W, Yang J et al (2015) Reduced graphene oxide-wrapped MoO3 composites prepared by using metal-organic frameworks as a precursor for all-solid-state flexible supercapacitors. Adv Mater 27(32):4695–4701. https://doi.org/10.1002/adma. 201501310 17. Chen GZ (2017) Supercapacitor and supercapattery as emerging electrochemical energy stores. Int Mater Rev 62(4):173–202. https://doi.org/10.1080/09506608.2016.1240914 18. Chen Y, Zhang X, Zhang D, Yu P et al (2011) High-performance supercapacitors based on reduced graphene oxide in aqueous and ionic liquid electrolytes. Carbon 49(2):573–580. https://doi.org/10.1016/j.carbon.2010.09.060 19. Conway BE (1991) Transition from “supercapacitor” to “battery” behavior in electrochemical energy storage. J Electrochem Soc 138(6):1539–1548. https://doi.org/10.1149/1.2085829 20. Côté AP, Benin AI, Ockwig NW, O’Keeffe M et al (2005) Chemistry: porous, crystalline, covalent organic frameworks. Science 310(5751):1166–1170. https://doi.org/10.1126/science. 1120411

444

Y. Liu et al.

21. Dong H, Li J, Zhao S, Zhao F et al (2020) An anti-aging polymer electrolyte for flexible rechargeable zinc-ion batteries. J Mater Chem A. https://doi.org/10.1039/d0ta07086f 22. Famprikis T, Canepa P, Dawson JA, Islam MS et al (2019) Fundamentals of inorganic solidstate electrolytes for batteries. Nat Mater 18(12):1278–1291. https://doi.org/10.1038/s41563019-0431-3 23. Farha OK, Eryazici I, Jeong NC, Hauser BG et al (2012) Metal-organic framework materials with ultrahigh surface areas: is the sky the limit? J Am Chem Soc 134(36):15016–15021. https://doi.org/10.1021/ja3055639 24. Fedorov MV, Kornyshev AA (2014) Ionic liquids at electrified interfaces. Chem Rev 114(5):2978–3036. https://doi.org/10.1021/cr400374x 25. Fic K, Lota G, Meller M, Frackowiak E (2012) Novel insight into the neutral medium as an electrolyte for high-voltage supercapacitors. Energy Environ Sci 5(2):5842–5850. https://doi. org/10.1039/c1ee02262h 26. Freemantle M (1998) Designer solvents. Chem Eng News 76(13):32–37. https://doi.org/10. 1021/cen-v076n013.p032 27. Furukawa H, Go YB, Ko N, Park YK et al (2011) Isoreticular expansion of metal-organic frameworks with triangular and square building units and the lowest calculated density for porous crystals. Inorg Chem 50(18):9147–9152. https://doi.org/10.1021/ic201376t 28. Galiński M, Lewandowski A, Stepniak I (2006) Ionic liquids as electrolytes. Electrochim Acta 51(26):5567–5580. https://doi.org/10.1016/j.electacta.2006.03.016 29. Grahame DC (1947) The electrical double layer and the theory of electrocapillarity. Chem Rev 41(3):441–501. https://doi.org/10.1021/cr60130a002 30. Guan BY, Yu L, Wang X, Song S et al (2017) Formation of onion-like NiCo2S4 particles via sequential ion-exchange for hybrid supercapacitors. Adv Mater 29(6). https://doi.org/10.1002/ adma.201605051 31. Halder A, Ghosh M, Khayum AM, Bera S et al (2018) Interlayer hydrogen-bonded covalent organic frameworks as high-performance supercapacitors. J Am Chem Soc 140(35):10941–10945. https://doi.org/10.1021/jacs.8b06460 32. Hao H, Wang J, Lv Q, Jiao Y et al (2020) Interfacial engineering of reduced graphene oxide for high-performance supercapacitor materials. J Electroanal Chem 878:114679. https://doi. org/10.1016/j.jelechem.2020.114679 33. He X, Geng Y, Qiu J, Zheng M et al (2010) Effect of activation time on the properties of activated carbons prepared by microwave-assisted activation for electric double-layer capacitors. Carbon 48(5):1662–1669. https://doi.org/10.1016/j.carbon.2010.01.016 34. He G, Ling M, Han X, Abou El Amaiem DI et al (2017) Self-standing electrodes with coreshell structures for high-performance supercapacitors. Energy Storage Mater 9:119–125. https://doi.org/10.1016/j.ensm.2017.07.005 35. Helmholtz H (1853) Ueber einige Gesetze der Vertheilung elektrischer Ströme in körperlichen Leitern mit Anwendung auf die thierisch-elektrischen Versuche. Ann Phys 165(6):211–233. https://doi.org/10.1002/andp.18531650603 36. Hirata M, Gotou T, Horiuchi S, Fujiwara M et al (2004) Thin-film particles of graphite oxide 1: high-yield synthesis and flexibility of the particles. Carbon 42(14):2929–2937. https://doi.org/ 10.1016/j.carbon.2004.07.003 37. Hong S, Lee J, Do K, Lee M et al (2017) Stretchable electrode based on laterally combed carbon nanotubes for wearable energy harvesting and storage devices. Adv Funct Mater 27(48):1704353. https://doi.org/10.1002/adfm.201704353 38. Hu L, Chen W, Xie X, Liu N et al (2011) Symmetrical MnO2-carbon nanotube-textile nanostructures for wearable pseudocapacitors with high mass loading. ACS Nano 5(11):8904–8913. https://doi.org/10.1021/nn203085j 39. Hu M, Hu T, Li Z, Yang Y et al (2018) Surface functional groups and interlayer water determine the electrochemical capacitance of Ti3C2 T x MXene. ACS Nano 12(4):3578–3586. https://doi.org/10.1021/acsnano.8b00676

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

445

40. Hu M, Zhang H, Hu T, Fan B et al (2020) Emerging 2D MXenes for supercapacitors: status, challenges, and prospects. Chem Soc Rev 49(18):6666–6693. https://doi.org/10.1039/ d0cs00175a 41. Huddleston JG, Visser AE, Reichert WM, Willauer HD et al (2001) Characterization and comparison of hydrophilic and hydrophobic room temperature ionic liquids incorporating the imidazolium cation. Green Chem 3(4):156–164. https://doi.org/10.1039/b103275p 42. Ibrahim Abouelamaiem D, Mostazo-López MJ, He G, Patel D et al (2018) New insights into the electrochemical behavior of porous carbon electrodes for supercapacitors. J Energy Storage 19:337–347. https://doi.org/10.1016/j.est.2018.08.014 43. Inagaki M, Toyoda M, Soneda Y, Morishita T (2018) Nitrogen-doped carbon materials. Carbon 132:104–140. https://doi.org/10.1016/j.carbon.2018.02.024 44. Jagtoyen M, Derbyshire F (1998) Activated carbons from yellow poplar and white oak by H3PO4 activation. Carbon 36(7–8):1085–1097. https://doi.org/10.1016/S0008-6223(98) 00082-7 45. Ji J, Zhang LL, Ji H, Li Y et al (2013) Nanoporous Ni(OH)2 thin film on 3d ultrathin-graphite foam for asymmetric supercapacitor. ACS Nano 7(7):6237–6243. https://doi.org/10.1021/ nn4021955 46. Jin Z, Yan X, Yu Y, Zhao G (2014) Sustainable activated carbon fibers from liquefied wood with controllable porosity for high-performance supercapacitors. J Mater Chem A 2(30):11706–11715. https://doi.org/10.1039/c4ta01413h 47. Kunze M, Jeong S, Paillard E, Schönhoff M et al (2011) New insights to self-aggregation in ionic liquid electrolytes for high-energy electrochemical devices. Adv Energy Mater 1(2):274–281. https://doi.org/10.1002/aenm.201000052 48. Laušević Z, Apel PY, Krstić JB, Blonskaya IV (2013) Porous carbon thin films for electrochemical capacitors. Carbon 64:456–463. https://doi.org/10.1016/j.carbon.2013.07.098 49. Lee YG, An GH (2020) Synergistic effects of phosphorus and boron co-incorporated activated carbon for ultrafast zinc-ion hybrid supercapacitors. ACS Appl Mater Interfaces 12(37):41342–41349. https://doi.org/10.1021/acsami.0c10512 50. Lee J, Tolosa A, Krüner B, Jäckel N et al (2017) Asymmetric tin-vanadium redox electrolyte for hybrid energy storage with nanoporous carbon electrodes. Sustainable Energy and Fuels. 1(2):299–307. https://doi.org/10.1039/c6se00062b 51. Lekawa-Raus A, Patmore J, Kurzepa L, Bulmer J et al (2014) Electrical properties of carbon nanotube-based fibers and their future use in electrical wiring. Adv Funct Mater 24(24):3661–3682. https://doi.org/10.1002/adfm.201303716 52. Lewandowski A, Olejniczak A, Galinski M, Stepniak I (2010) Performance of carbon-carbon supercapacitors based on organic, aqueous, and ionic liquid electrolytes. J Power Sources 195(17):5814–5819. https://doi.org/10.1016/j.jpowsour.2010.03.082 53. Li B, Dai F, Xiao Q, Yang L et al (2016) Nitrogen-doped activated carbon for a high energy hybrid supercapacitor. Energy Environ Sci 9(1):102–106. https://doi.org/10.1039/c5ee03149d 54. Li Y, Xu J, Feng T, Yao Q et al (2017) Fe2O3 nanoneedles on ultrafine nickel nanotube arrays as efficient anode for high-performance asymmetric supercapacitors. Adv Funct Mater 27(14). https://doi.org/10.1002/adfm.201606728 55. Li J, McColl K, Lu X, Sathasivam S et al (2020) Multi-scale investigations of δ-Ni 0.25 V 2 O 5
nH 2 O cathode materials in aqueous zinc-ion batteries. Adv Energy Mater 10(15):2000058. https://doi.org/10.1002/aenm.202000058 56. Liang Y, Liang F, Zhong H, Li Z et al (2013) An advanced carbonaceous porous network for high-performance organic electrolyte supercapacitors. J Mater Chem A 1(24):7000–7005. https://doi.org/10.1039/c3ta11051f 57. Liang S, Yan W, Wu X, Zhang Y et al (2018) Gel polymer electrolytes for lithium-ion batteries: fabrication, characterization, and performance. Solid State Ionics 318:2–18. https:// doi.org/10.1016/j.ssi.2017.12.023

446

Y. Liu et al.

58. Lu X, Bertei A, Finegan DP, Tan C et al (2020) 3D microstructure design of lithium-ion battery electrodes assisted by X-ray nano-computed tomography and modeling. Nat Commun 11(1):1–13. https://doi.org/10.1038/s41467-020-15811-x 59. Lv Z, Luo Y, Tang Y, Wei J et al (2018) Editable supercapacitors with customizable stretchability based on mechanically strengthened ultralong MnO2 nanowire composite. Adv Mater 30(2). https://doi.org/10.1002/adma.201704531 60. Martínez-Periñán E, Down MP, Gibaja C, Lorenzo E et al (2018) Antimonene: a novel 2D nanomaterial for supercapacitor applications. Adv Energy Mater 8(11):1702606. https://doi. org/10.1002/aenm.201702606 61. Mindemark J, Lacey MJ, Bowden T, Brandell D (2018) Beyond PEO—alternative host materials for Li+conducting solid polymer electrolytes. Prog Poly Sci 81:114–143. https:// doi.org/10.1016/j.progpolymsci.2017.12.004 62. Naguib M, Kurtoglu M, Presser V, Lu J et al (2011) Two-dimensional nanocrystals produced by exfoliation of Ti 3AlC 2. Adv Mater 23(37):4248–4253. https://doi.org/10.1002/adma. 201102306 63. Naguib M, Halim J, Lu J, Cook KM et al (2013) New two-dimensional niobium and vanadium carbides as promising materials for li-ion batteries. J Am Chem Soc 135(43):15966–15969. https://doi.org/10.1021/ja405735d 64. Novoselov KS, Geim AK, Morozov SV, Jiang D et al (2004) Electric field in atomically thin carbon films. Science 306(5696):666–669. https://doi.org/10.1126/science.1102896 65. Ojha K, Arulmozhi N, Aranzales D, Koper MTM (2020) Double layer at the Pt(111)–aqueous electrolyte Interface: potential of zero charge and anomalous Gouy–chapman screening. Angew Chem Int Ed 59(2):711–715. https://doi.org/10.1002/anie.201911929 66. Peigney A, Laurent C, Flahaut E, Bacsa RR et al (2001) Specific surface area of carbon nanotubes and bundles of carbon nanotubes. Carbon 39(4):507–514. https://doi.org/10.1016/ S0008-6223(00)00155-X 67. Peng L, Fang Z, Zhu Y, Yan C et al (2018) Holey 2D nanomaterials for electrochemical energy storage. Adv Energy Mater 8(9). https://doi.org/10.1002/aenm.201702179 68. Peng H, Yao B, Wei X, Liu T et al (2019) Pore and heteroatom engineered carbon foams for supercapacitors. Adv Energy Mater 9(19):1803665. https://doi.org/10.1002/aenm.201803665 69. Qian W, Sun F, Xu Y, Qiu L et al (2014) Human hair-derived carbon flakes for electrochemical supercapacitors. Energy Environ Sci 7(1):379–386. https://doi.org/10.1039/c3ee43111h 70. Ren L, Zhang G, Yan Z, Kang L et al (2017) High capacitive property for supercapacitor using Fe3+/Fe2+ redox couple additive electrolyte. Electrochim Acta 231:705–712. https://doi.org/ 10.1016/j.electacta.2017.02.056 71. Saito Y, Luo X, Zhao C, Pan W et al (2015) Filling the gaps between graphene oxide: a general strategy toward Nanolayered oxides. Adv Funct Mater 25(35):5683–5690. https://doi.org/10. 1002/adfm.201501358 72. Sevilla M, Fuertes AB (2014) Direct synthesis of highly porous interconnected carbon nanosheets and their application as high-performance supercapacitors. ACS Nano 8(5):5069–5078. https://doi.org/10.1021/nn501124h 73. Sevilla M, Mokaya R (2014) Energy storage applications of activated carbons: supercapacitors and hydrogen storage. In: Energy and environmental science. 21 March 2014 Royal Society of Chemistry. pp. 1250–1280. DOI:https://doi.org/10.1039/c3ee43525c. 74. Sheberla D, Bachman JC, Elias JS, Sun CJ et al (2017) Conductive MOF electrodes for stable supercapacitors with high areal capacitance. Nat Mater 16(2):220–224. https://doi.org/10. 1038/nmat4766 75. Shi P, Li L, Hua L, Qian Q et al (2017) Design of amorphous manganese oxide@multiwalled carbon nanotube Fiber for robust solid-state supercapacitor. ACS Nano 11(1):444–452. https:// doi.org/10.1021/acsnano.6b06357 76. Simon P, Gogotsi Y (2013) Capacitive energy storage in nanostructured carbon-electrolyte systems. Acc Chem Res 46(5):1094–1103. https://doi.org/10.1021/ar200306b

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

447

77. Snook GA, Kao P, Best AS (2011) Conducting-polymer-based supercapacitor devices and electrodes. J Power Sources 196(1):1–12. https://doi.org/10.1016/j.jpowsour.2010.06.084 78. Stern O (1924) ZUR THEORIE DER ELEKTROLYTISCHEN DOPPELSCHICHT. Z Elektrochem Angew Phys Chem 30(21–22):508–516. https://doi.org/10.1002/BBPC. 192400182 79. Sun J, Zheng G, Lee HW, Liu N et al (2014) Formation of stable phosphorus-carbon bond for enhanced performance in black phosphorus nanoparticle-graphite composite battery anodes. Nano Lett 14(8):4573–4580. https://doi.org/10.1021/nl501617j 80. Tan YB, Lee JM (2013) Graphene for supercapacitor applications. J Mater Chem A 1(47):14814–14843. https://doi.org/10.1039/c3ta12193c 81. Tang H, Wang J, Yin H, Zhao H et al (2015) Growth of polypyrrole ultrathin films on mos2 monolayers as high-performance supercapacitor electrodes. Adv Mater 27(6):1117–1123. https://doi.org/10.1002/adma.201404622 82. Tian S, Qi L, Yoshio M, Wang H (2014) Tetramethylammonium difluoro(oxalato)borate dissolved in ethylene/propylene carbonates as electrolytes for electrochemical capacitors. J Power Sources 256:404–409. https://doi.org/10.1016/j.jpowsour.2014.01.101 83. Trasatti S, Buzzanca G (1971) Ruthenium dioxide: a new interesting electrode material. Solidstate structure and electrochemical behavior. J Electroanal Chem 29(2):A1–A5. https://doi. org/10.1016/S0022-0728(71)80111-0 84. van Bunge W, Krop H, Leeuwenburgh B, van Ruler H, et al. (2003) The dictionary of seventeenth and eighteenth-century Dutch philosophers 85. Wang J, Kaskel S (2012) KOH activation of carbon-based materials for energy storage. J Mater Chem 22(45):23710–23725. https://doi.org/10.1039/c2jm34066f 86. Wang Q, Yan J, Wang Y, Wei T et al (2014) Three-dimensional flower-like and hierarchical porous carbon materials as high-rate performance electrodes for supercapacitors. Carbon 67:119–127. https://doi.org/10.1016/j.carbon.2013.09.070 87. Wang F, Wu X, Yuan X, Liu Z et al (2017) Latest advances in supercapacitors: from new electrode materials to novel device designs. Chem Soc Rev 46(22):6816–6854. https://doi.org/ 10.1039/c7cs00205j 88. Wang J, Wang Y, Hu H, Yang Q et al (2020) From metal-organic frameworks to porous carbon materials: recent progress and prospects from energy and environmental perspectives. Nanoscale 12(7):4238–4268. https://doi.org/10.1039/c9nr09697c 89. Wen M, Liu D, Kang Y, Wang J et al (2019) Synthesis of high-quality black phosphorus sponges for all-solid-state supercapacitors. Mater Horizons 6(1):176–181. https://doi.org/10. 1039/c8mh00708j 90. Wu S, Hui KS, Hui KN (2018) 2D black phosphorus: from preparation to applications for electrochemical energy storage. Adv Sci 5(5). https://doi.org/10.1002/advs.201700491 91. Xia J, Chen F, Li J, Tao N (2009) Measurement of the quantum capacitance of graphene. Nat Nanotechnol 4(8):505–509. https://doi.org/10.1038/nnano.2009.177 92. Xia F, Wang H, Jia Y (2014) Rediscovering black phosphorus as an anisotropic layered material for optoelectronics and electronics. Nat Commun 5(1):1–6. https://doi.org/10.1038/ ncomms5458 93. Xie L, Su F, Xie L, Guo X et al (2020a) Effect of pore structure and doping species on charge storage mechanisms in porous carbon-based supercapacitors. Mat Chem Front 4(9):2610–2634. https://doi.org/10.1039/d0qm00180e 94. Xie LS, Skorupskii G, Dincǎ M (2020b) Electrically conductive metal-organic frameworks. Chem Rev 120(16):8536–8580. https://doi.org/10.1021/acs.chemrev.9b00766 95. Yang Y, Li L, Ruan G, Fei H et al (2014) Hydrothermally formed three-dimensional Nanoporous Ni(OH)2 thin-film supercapacitors. ACS Nano 8(9):9622–9628. https://doi.org/ 10.1021/nn5040197 96. Yao B, Chandrasekaran S, Zhang J, Xiao W et al (2019) Efficient 3D printed Pseudocapacitive electrodes with ultrahigh MnO 2 loading. Joule 3(2):459–470. https://doi.org/10.1016/j.joule. 2018.09.020

448

Y. Liu et al.

97. Yu L, Chen GZ (2016) Redox electrode materials for supercapatteries. J Power Sources 326:604–612. https://doi.org/10.1016/j.jpowsour.2016.04.095 98. Yu X, Ruan D, Wu C, Wang J et al (2014) Spiro-(1,10 )-bipyrrolidinium tetrafluoroborate salt as high voltage electrolyte for electric double-layer capacitors. J Power Sources 265:309–316. https://doi.org/10.1016/j.jpowsour.2014.04.144 99. Yu F, Chang Z, Yuan X, Wang F et al (2018) Ultrathin NiCo2S4@graphene with a core-shell structure as a high-performance positive electrode for hybrid supercapacitors. J Mater Chem A 6(14):5856–5861. https://doi.org/10.1039/c8ta00835c 100. Zhang L, Zhao XS (2009) Carbon-based materials as supercapacitor electrodes. Chem Soc Rev 38(9):2520–2531. https://doi.org/10.1039/b813846j 101. Zhang J, Ma J, Zhang LL, Guo P et al (2010) Template synthesis of tubular ruthenium oxides for supercapacitor applications. J Phys Chem C 114(32):13608–13613. https://doi.org/10. 1021/jp105146c 102. Zhang F, Zhang T, Yang X, Zhang L et al (2013) A high-performance supercapacitor-battery hybrid energy storage device based on graphene-enhanced electrode materials with ultrahigh energy density. Energy Environ Sci 6(6):1623–1632. https://doi.org/10.1039/c3ee40509e 103. Zhang C, Xiao J, Lv X, Qian L et al (2016) Hierarchically porous Co3O4/C nanowire arrays derived from a metal-organic framework for high-performance supercapacitors and the oxygen evolution reaction. J Mater Chem A 4(42):16516–16523. https://doi.org/10.1039/c6ta06314d 104. Zheng C, Qian W, Cui C, Zhang Q et al (2012) Hierarchical carbon nanotube membrane with high packing density and tunable porous structure for high voltage supercapacitors. Carbon 50(14):5167–5175. https://doi.org/10.1016/j.carbon.2012.06.058 105. Zhong C, Deng Y, Hu W, Qiao J et al (2015) A review of electrolyte materials and compositions for electrochemical supercapacitors. Chem Soc Rev 44(21):7484–7539. https://doi.org/10.1039/c5cs00303b 106. Zhou C, Zhang Y, Li Y, Liu J (2013) Construction of high-capacitance 3D CoO@Polypyrrole nanowire array electrode for aqueous asymmetric supercapacitor. Nano Lett 13(5):2078–2085. https://doi.org/10.1021/nl400378j 107. Zhu S, Li L, Liu J, Wang H et al (2018) Structural directed growth of ultrathin parallel birnessite on β-MnO2 for high-performance asymmetric supercapacitors. ACS Nano 12(2):1033–1042. https://doi.org/10.1021/acsnano.7b03431 108. Zou R, Xu K, Wang T, He G et al (2013) Chain-like NiCo2O4 nanowires with different exposed reactive planes for high-performance supercapacitors. J Mater Chem A 1(30):8560–8566. https://doi.org/10.1039/c3ta11361b

Mr. Yiyang Liu received the B.Eng degree in chemical engineering from the University of Nottingham and the MS degree from University College London in 2019 and 2020, respectively. He is currently a Ph.D. candidate in the Electrochemical Innovation Lab, Department of Chemical Engineering, University College London (UCL), London, UK, under the supervision of Prof. Dan Brett and Dr. Guanjie He. His current research focuses on the material design and synthesis for the application in hybrid battery systems.

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications

449

Paul Shearing is a professor in chemical engineering at University College London where he holds The Royal Academy of Engineering Chair in Emerging Battery Technologies. His research interests cover a broad range of electrochemical engineering themes with a particular interest in the relationship between performance and microstructure for energy materials.He co-directs UCL’s Electrochemical Innovation Lab and leads the UK STFC Global Challenge Network in Batteries and Electrochemical Devices. He was a founding investigator of the UK’s Faraday Institution, where he chairs the Training and Diversity Panel and is PI of the LiSTAR program investigating Li-sulfur battery technologies. Dr. Guanjie He is a senior lecturer in chemistry at the University of Lincoln, and an Honorary Lecturer in the Department of Chemistry and Department of Chemical Engineering at the University College London (UCL). His research focused on materials for electrochemical energy storage and conversion applications, especially electrode materials in aqueous electrolyte systems. He has published over 90 papers in peer-reviewed journals (citation > 2500, h-index of 27). He is serving as a guest editor for BioMed Research International, Frontiers in Materials, Frontiers in Chemistry, and Frontiers in Bioengineering and Biotechnology, a topic editor for Coatings, and an editorial board member for Batteries. He received his Ph.D. degree from the Chemistry Department, UCL, under the supervision of Prof. Ivan Parkin. Before this, he received his BSc in College of Materials Science and Engineering, Donghua University. During 2018–2019, he worked in Electrochemical Innovation Lab in the Department of Chemical Engineering, UCL, as a research fellow with Prof. Dan Brett and Prof. Paul Shearing. Prof. Dan Brett is a professor in chemical engineering at University College London (UCL) and an expert in electrochemical materials science and technology development. He received a Ph.D. from Imperial College London and a Bachelor of Science from King’s College London. His research is in the area of electrochemical engineering and technology. This includes electrochemical energy conversion and storage (fuel cells, batteries, supercapacitors, electrolyzers); and electrochemical sensors; electroanalysis, hybrid vehicles, and micro-generation technologies. He specializes in developing novel diagnostic techniques for the study of high- and low-temperature fuel cells and their materials. He is also active in modeling, instrumentation development, engineering design, device fabrication, materials development, and techno-economic analysis of electrochemical energy conversion technologies. He has published >400 peer-reviewed journal papers (citation >10000, h-index of 52) and is in the top 40 published academics in the world over the last 10 years in electrochemical power sources (combined field of fuel cells, batteries, and supercapacitors – Scopus). He was awarded the 2009 De Nora Foundation Prize and the 2011 Baker Medal (Institute of Civil Engineers) for his work on fuel cells.

Chapter 16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage Jianmin Luo

16.1

Introduction

With the ever-increasing depletion of fossil-fuel resources and energy consumption requirements, sustainable energy technologies are urgently needed. Therefore, considerable attention has been focused on the development of novel materials for highperformance electrochemical energy storage applications. Two-dimensional (2D) nanomaterials including graphene [1], hexagonal boron nitrides [2], transition metal dichalcogenides (TMDs) [3], black phosphorus (BP) [4], and metal-organic frameworks (MOFs) [5] have been widely researched and regarded as the promising candidates for electrochemical energy storage over the past decade owing to their attractive properties. 2D transition metal carbides and nitrides (MXenes), fast-growing nanomaterials in the 2D materials family, were first discovered by exfoliating from threedimensional (3D) MAX phases (M is the early transition metal, A is III or IV A-group element, X is a carbon (C) or nitrogen (N)) with hydrofluoric acid (HF) [6]. During the etching process, A layers in the MAX phase are etched and hydrophilic groups (OH, O, and/or F groups) are bonded to the outer M layers. In their general formula of Mn + 1XnTx, Tx stands for the surface hydrophilic groups [7]. Up to now, more than 30 types of MXenes, including Ti3C2, Ti2C, Ti4N3, and Mo2C, with a variety of compositions and structures have been successfully synthesized, and dozens more have been explored by computational methods [8– 11]. MXenes have a unique combination of physical and chemical properties, including high electrical conductivity, mechanical properties, enabling stable colloidal solutions in water, and efficient absorption of electromagnetic waves, which have led to a large number of applications ranging from energy storage [12–21],

J. Luo (*) Thayer School of Engineering, Dartmouth College, Hanover, NH, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_16

451

452

J. Luo

electromagnetic interference shielding [22], environment and catalysis [23–26], medicine [27], and electronics [28–30]. Besides, the unique characteristics of MXenes, including large redox-active surface area, high electrical conductivity, rich surface chemistry, and tunable layered structure, have endowed them as promising candidates for electrochemical energy storage applications [31]. However, similar to other 2D materials, the electrochemical energy storage performances of MXenes are still restricted by their trend to be restacked and agglomerated, which limits the ionic and electronic transport in electrode materials [32]. For the sake of the full utilization of their electrochemical performance, strategies including introducing interlayer spacers between MXene layers and creating porous structures of MXene were proposed [33–35]. Since MXenes have energy storage space between layers, if the interlayer space can be rationally exploited and utilized, the overall electrochemical performance of MXene for energy storage will be enhanced [36]. Therefore, interlayer structural engineering of MXene with enlarged interlayer spacing is highly necessary. Also, there have some published works that summarized the synthesis techniques, properties, and applications of MXenes, less attention has been paid to the interlayer structural engineering and regulation mechanisms of MXene-based architectures [11, 37–41]. Consequently, it is highly desirable to have a comprehensive understanding of interlayer structural engineering of MXene-based nanomaterials for electrochemical energy storage applications. In this chapter, we focus on the interlayer structural engineering of MXenes, and provide an in-depth discussion regarding the recent layered structure design with expanded and engineered interlayer structures and their electrochemical energy storage applications. The effect of different intercalation agent types and interlayer structural engineering methods related to electrochemical energy storage are discussed (Fig. 16.1). In the first section, we briefly summarize the synthesis techniques of MXenes and go over the different types of layered structure design of MXenes by selecting different intercalation agents. Besides, synthesis-dependent structural properties of MXenes with enlarged and engineered interlayer space are also investigated. In the second section, a series of interlayer structure designs along with the electrochemical energy storage performance is systematically overviewed. In the last section, the challenges and prospects of the interlayer structural engineering of MXene-based nanostructure are addressed. Accordingly, the investigation of interlayer structural engineering of MXene would be greatly conducive to the further optimization of MXene for electrochemical energy storage applications.

16.2

Interlayer Structural Engineering of 2D MXene

Before the layered structure design of MXene, different synthesis techniques are needed to prepare MXene with high quality. The intrinsic properties of MXenes are closely related to their synthesis techniques [39]. Therefore, synthesis conditions can

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

453

Fig. 16.1 The overview representative interlayer structural design of MXenes, synthetic method, and their electrochemical energy storage applications

directly influence the layered structure design of MXenes and their properties and energy storage performances. At the beginning of this section, we briefly summarize the synthesis techniques of MXenes and illuminate the relationship between synthesis techniques and the structural properties of MXenes. Electrodes made from pristine MXene with small interlayer spacing exhibit unsatisfied electrochemical performance [31]. Significantly, MXene with expanded and engineered interlayer structures for excellent storage capability is confirmed [42–44]. Therefore, in the following, we review the different types of layered structure design of MXenes by selecting different intercalation agents, including heteroatoms, alkali metal ions, other metal ions, polymers, larger cations, and 1D/2D/3D nanomaterials. Importantly, the synthesis-dependent structural properties of MXenes with expanded and engineered interlayer structures are also studied.

16.2.1 Synthesis of MXenes In general, two primary strategies, wet etching synthesis and non-etching synthesis, have been explored to synthesize MXenes [11]. Wet etching synthesis refers to a method of extraction of MnXn
1 (n ¼ 2, 3, 4) from their layered precursors MAX by removing the A atoms’ layers [8]. MAX phase is a large family of ternary carbides and nitrides. Until now, more than 70 different kinds of MAX phases have been reported, and the possible constituent atoms are shown in Fig. 16.2a [37].

454

J. Luo

Fig. 16.2 (a) Elements in periodic tables used to synthesize MAX phases. (b) Three typical structures of MAX phases (M4AX3, M3AX2, M2AX) and the corresponding MXenes. (c) The molecular structure model of single layer M3X2Tx. (d) Publication trends in MXene and MXene for electrochemical energy storage (Source: Web of Science. Search index: [topic ¼ MXene or MXene energy storage])

Three typical structure MAX phases (M4AX3, M3AX2, and M2AX) and their corresponding products MXenes are shown in Fig. 16.2b. After removing the A atom layers in MAX by an aqueous fluoride-containing acidic solution, different

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

455

terminal groups would be bonded with interfacial M atoms (Fig. 16.2c). For example, the HF etching method was first used to delaminate Ti3AlC2 MAX in 2011, and many hydrophilic groups (–OH, –F) were bonded with interfacial Ti atoms [6]. It should be noted that different etching conditions, including different etchants concentration and etching temperature, can strongly influence the properties and quality of as-synthesized MXene [39]. So far, various types of etchants (HF, HCl + LiF, NaOH, NH4HF2, and TMAOH) have been used and some novel routes have been developed [6, 45–47]. Besides, a non-etching method such as chemical vapor deposition (CVD) should also be possible for MXene synthesis [48]. In the previous research results, ultrathin 2D α-Mo2C orthorhombic crystals with up to 100-μm lateral size can be produced by using methane as the precursor and copper foil above a molybdenum foil as the substrate. MXenes with a large lateral size and few defects would be synthesized by this method, which facilitates the study of their intrinsic properties [48]. In this part, we focus on the two etching synthesis techniques of MXenes by fluoride-containing acidic solutions, which are the most widely used methods for MXenes synthesis. HF etching method is one of the most common synthesis methods for synthesizing MXenes due to its flexibility to selectively etch away the A layers from most of the precursor MAX phases [6]. Taking Ti3AlC2 as an example, the reaction process can be described by the following equations [6]. Ti3 AlC2 þ 3HF ¼ Ti3 C2 þ AlF3 þ 3=2H2

ð16:1Þ

Ti3 C2 þ 2H2 O ¼ Ti3 C2 ðOHÞ2 þ H2

ð16:2Þ

Ti3 C2 þ 2HF ¼ Ti3 C2 F2 þ H2

ð16:3Þ

The etching conditions for MAX phases vary from one to another, which depends on the compositions, structure, atomic bonding, and particle size of the material [39]. Based on the previous research results, larger n in Mn + 1CnTx requires stronger etching and longer etching time [39]. After HF etching, multilayered MXene is often obtained. Therefore, to obtain separated MXene sheets, typical intercalation and delamination processes are required [49]. In the intercalation process, organic molecules (e.g., dimethyl sulfoxide (DMSO) and tetrabutylammonium hydroxide (TBAOH)) are usually used to expand the interlayer spacing of multilayered MXene [47, 49–51]. Also, sonication treatment is often required in the followed delamination process. Besides, the disadvantage of the HF etching method is the highly corrosive nature of HF, which would lead to the formation of rich structural defects in the obtained MXene [39]. The hazardous environmental and safety concern also exists during the utilization process of HF. To avoid the use of highly toxic and dangerous HF, researchers have made a lot of effort to explore safer ways to synthesize MXenes. In addition to HF, modified fluoride-based acid based on a mixture solution of a strong acid and a fluoride salt can also be used to selectively remove A atoms layers from MAX precursor [45]. It can be found that strong acids can react with fluoride salts to form in situ HF, which

456

J. Luo

not only can selectively etch A atoms, but lead to the intercalation of cations (e.g., Li+, Na+, and K+) and water between MXene layers. Therefore, the interlayer spacing of MXene increases, and the interaction between MXene layers weakens after using this in situ HF method. It should be noted that both the concentration of fluoride salts and strong acid can influence the quality and size of obtained MXene flakes [39]. When using the minimally intensive layer delamination (MILD) method (12 M LiF/9 M HCl), the simple washing process alone is enough to delaminate multilayered Ti3C2 MXene, synthesizing single- and few-layer MXene with few defect and larger lateral size [52]. Besides, difluorides such as KHF2 or NH4HF2 can also be used to selectively etch the Ti3AlC2 MAX phase [53]. In general, compared to HF etching, modified fluoride-based acid etching exhibits different properties: (1) Synthesized MXenes have larger lateral sizes and fewer defects; (2) Larger interlayer spacing of the obtained MXenes facilitates the delamination; (3) Higher content of –O terminal groups and less content of –F terminal groups.

16.2.2 Layered Structure with an Enlarged Interlayer Spacing It is well known that the performances of energy storage applications are highly dependent on the electrode materials. As a result, the design and synthesis of highquality materials become an essential step before constructing high-performance electrodes. Until now, MXenes have been widely studied for energy storage applications with the enhanced electrochemical performance [11]. The layered structure of MXene with enlarged interlayer spacing is demonstrated with increased energy storage space for charge-carrying ions intercalation and storage as well as excellent ion transport property due to shorter diffusion pathways [31]. These favorable characteristics endow MXenes with much improved electrochemical performance, which facilitates their practical applications for energy storage applications. In this following section, we will go over the recent progress in layered structure design of MXenes with enlarged interlayer spacing by selecting different intercalation agents, including heteroatoms, alkali metal ions, other metal ions, polymers, larger cations, and 1D/2D/3D nanomaterials. After HF or in situ HF etching, MXenes have a highly negatively charged interface and form stable colloidal solutions in water as confirmed by zeta potential measurements [43, 51]. Therefore, cations can be spontaneously intercalated between MXene layers by electrostatic interaction, increasing the interlayer spacing of MXene [43, 50, 54–56]. It is noteworthy that the new chemical bonding between intercalant and MXene can be formed during the intercalation process of intercalants, which can induce the modification of MXene interfacial properties [44]. Gogotsi et al. studied the intercalation behavior of cations (e.g., Li+, Na+, K+, NH4+, Mg2+, and Al3+) from aqueous salt solutions between Ti3C2 MXene layers [54]. The results show that the intercalation of the above cations is spontaneous, and

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

457

the intercalation behavior depends on pH and the nature of the cations. When Ti3C2 MXene is treated by potassium hydroxide (KOH) and ammonium hydroxide (NH4OH) solutions, the interlayer spacing of Ti3C2 can be increased from 1 to 1.27 nm. Besides, the alkaline solution treatment can not only increase the interlayer spacing of MXene but also change the terminal groups in the interfaces between MXene layers. Peng and co-workers used an NaOH solution to treat Ti3C2 [23]. After alkalization intercalation, Ti3C2 MXene exhibits unique heavy metal ion (e.g., Pb2+) adsorption behavior. The NaOH treatment process is conducive to the transformation of –F to –OH/–ONa groups. Besides, the surface [Ti-O]
H+ groups exhibit strong metal–ligand interaction with Pb2+, and the [Ti-O]
H+ and the [Ti-O]
Na+ sites are propitious to the enhancement of Pb2+ uptake. Except for the abovementioned alkali metal ions, other metal ions (e.g., Sn4+ and 2+ Co ) can also be intercalated into the interlayer of MXene [44, 57]. Luo et al. first synthesized Sn4+ pillared Ti3C2 MXene by the use of alkali metal ions (e.g., Li+, Na+, and K+) to pre-intercalate between Ti3C2 layers, followed by the Sn4+ pillaring process based on the ion-exchange interaction between Sn4+ and pre-intercalated alkali metal ions [44]. After Sn4+ intercalation, the interlayer spacing of Ti3C2 increases from 0.98 to 1.28 nm. Note that the strong Ti-O-Sn bonding can be confirmed after the intercalation of Sn4+ due to the formation of an inner sphere complex on the interface of Ti3C2. By the same method, Song and co-workers fabricated Co2+ intercalated V2C MXene (V2C@Co) [57]. After Co2+ intercalation, the interlayer spacing of V2C increases from 0.74 to 0.95 nm by the formation of V–O–Co bonds. Although metal ions can be successfully intercalated into the interlayer of MXene by electrostatic interaction or ion exchange interaction, the added values of interlayer spacing are limited owing to the small size of metal ions [43]. Luo and co-workers used large volume cationic surfactants as the intercalation agent to intercalate between MXene layers to increase the interlayer spacing of MXene (Fig. 16.3a) [43]. Significantly, the interlayer spacing of MXene can be fine-tuned by creating pillared structures based on the spontaneous intercalation of different size of cationic surfactant (dodecyltrimethylammonium bromide (DTAB), tetradecyltrimethylammonium bromide (TTAB), cetyltrimethylammonium bromide (CTAB), stearyl trimethylammonium bromide (STAB), and dioctadecyldimethylammonium chloride (DDAC)) under different temperatures (30–70  C) (Fig. 16.3b, c) [43]. When using cetyltrimethylammonium bromide (CTAB) as the intercalation agent under the temperature of 40  C, the interlayer spacing of Ti3C2 increases from 0.98 to 2.23 nm. Since CTAB has a long-chain hydrophobic tail (Fig. 16.3d), the possible position of CTA+ in the interlayer space of Ti3C2 is shown in Fig. 16.3e [43]. When using the larger size of stearyl STAB as the intercalation agent at 50  C, the interlayer spacing of Ti3C2 can increase to 2.71 nm, a nearly 177% value-added compared with that of pristine Ti3C2 (0.98 nm). More importantly, ion storage sites for other cations can be provided after the intercalation of cationic surfactant between MXene layers. According to the ion-exchange mechanism, different pillared structure MXene can be fabricated by a facile cationic surfactant pre-pillaring followed by the cations pillaring method. Based on this method, Luo et al. successfully synthesized Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2) by CTAB pre-pillaring and subsequent Sn4+ pillaring process (Fig. 16.3a) [43]. Scanning

458

J. Luo

Fig. 16.3 (a) Schematic illustration of the preparation of Sn4+ pillared Ti3C2 (CTAB-Sn(IV) @Ti3C2) by CTAB pre-pillaring process followed by a method of Sn4+ pillaring. (b) XRD patterns of cationic surfactants pre-pillaring Ti3C2 at 40  C. (c) Interlayer spacing of cationic surfactants pre-pillared Ti3C2 at different treatment temperatures (30–70  C). (d) Schematic drawing about the structure of CTAB. (e) Possible position of CTA+ in the interlayer space of Ti3C2 MXene. (f) STEM image of CTAB-Sn(IV)@Ti3C2 and corresponding elemental mapping of Ti, C, and Sn (insets). Reproduced with permission-Copyright 2016, American Chemical Society [43]. (g) Schematic illustration of pyrrole polymerization using MXene. (h) Cross-sectional TEM image of aligned polypyrrole chains (bright layers) between MXene sheets (darker layers). Reproduced with permission-Copyright 2016, John Wiley and Sons [60]. (i) Schematic illustration of the preparation of S atoms intercalated Ti3C2 (CT-S@Ti3C2). (j) The change of interlayer spacings about S atoms intercalated Ti3C2 under different annealing temperatures. STEM image of CT-S@Ti3C2-450 and corresponding elemental mappings of Ti, C, and S. Reproduced with permission-Copyright 2019, John Wiley and Sons [36]. (k) Schematic illustration of the layer-by-layer self-assembly MXenebased multilayer films onto planar substrates. Reproduced with permission-Copyright 2019, Springer Nature [62]

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

459

transmission electron microscopy (STEM) image of CTAB-Sn(IV)@Ti3C2 combined with elemental mapping of Ti, C, and Sn confirms the successful intercalation of Sn4+ between Ti3C2 MXene layers (Fig. 16.3f). In addition to the metal ions and cationic surfactants, other large ions such as amine ions can also be used as the intercalation agents. Gogotsi et al. and co-workers proposed an isopropylamine (i-PrA) intercalation method to delaminate Nb2C MXene [50]. When placing Nb2C MXene in an i-PrA solution, the R-NH3+ cations in the solution can spontaneously intercalate between MXene layers due to the electrostatic interaction, and the corresponding interlayer spacing of Nb2C increases from 1.04 to 2.27 nm. The successful intercalation of R-NH3+ cations can effectively weaken the interlayer interactions, which helps to the delamination of Nb2C. Additionally, some polymers have been confirmed to be the promising intercalation agents for MXenes [58–60]. Gogotsi’s group developed a facile strategy for the in situ polymerization of pyrrole confined between Ti3C2 MXene layers (Fig. 16.3g) [60]. The hydrogen bonding between the anion and the surface of the electrode contributes to the self-align of polypyrrole (PPy) between the MXene layers; Ti3C2 MXene has a strongly pronounced acidic character after HF etching, which can interact with pyrrole and result in a protonated pyrrole molecule. The protonated pyrrole molecule continuously reacts with unprotonated pyrrole, forming a long chain of PPy (chain propagation). Therefore, no oxidant is needed in the polymerization process of the pyrrole in this work. Cross-sectional transmission electron microscopy (TEM) images of the PPy/Ti3C2 film in Fig. 16.3h show that PPy chains are not only intercalated but also well aligned between the MXene layers, producing a periodic pattern. Besides, the formed PPy becomes conductive after doping with the fluorine from the termination of the MXene layers. After PPy intercalation, the interlayer spacing of Ti3C2 increases from 1.4 to 2.4 nm. This increased value is comparable to that of large size cations. Except for oxidant-free polymerization, electrochemical polymerization can also be used to synthesize PPy intercalated MXene. Zhi’s group proposed an electrochemical polymerization method to intercalate PPy between Ti3C2 layers [58]. During the 60 s electrochemical polymerization process, PPy first intercalates in the interlayers of multilayered Ti3C2 and then wraps the Ti3C2 particles. After PPy intercalation, the interlayer distance of Ti3C2 increases from 0.94 to 0.96 nm. Besides, Ling et al. used a simple mixture method to mix polydiallyldimethylammoniumchloride (PDDA) and polyvinyl alcohol (PVA) with Ti3C2 MXene to produce Ti3C2/polymer composites [59]. The successful intercalation and confinement of the above polymers in the interlayers of Ti3C2 can increase the interlayer spacing of Ti3C2. The as-fabricated composites are flexible and high electrical conductivities. Moreover, facile precursor intercalation and the annealing process are effective ways to fabricate heteroatom intercalated MXene, which were confirmed to be one of the MXene-based nanostructures with enlarged interlayer spacing [36, 61]. Tao’s group successfully synthesized sulfur (S) atoms intercalated Ti3C2 MXene by a facile CTAB pretreatment, thermal diffusion with elemental S, and finally annealing and carbon disulfide (CS2) washing process (Fig. 16.3i) [36]. After CTAB pretreatment, the interlayer spacing of Ti3C2 increases from 0.98 to 2.23 nm,

460

J. Luo

which helps to the intercalation of elemental S between Ti3C2 layers. The intercalated elemental S can react with the interfacial titanium (Ti) atoms to form Ti-S bonds with expanded interlayer spacing during the annealing process. CS2 is used to wash the residual elemental S. After annealing at 500  C, the interlayer spacing of S atoms intercalated Ti3C2 (CT-S@Ti3C2) with the help of CTAB pretreatment, 1.37 nm is larger than that the cases without CTAB pretreatment (0.98 nm) and CTAB pretreatment (1.15 nm), which indicates the successful intercalation of S atoms between Ti3C2 layers (Fig. 16.3j). The STEM image of CT-S@Ti3C2 combined with elemental mapping of S confirms the successful intercalation of S atoms between Ti3C2 MXene layers (Fig. 16.3j). Wang and co-workers fabricated nitrogen (N)-doped Ti3C2 MXene (N-Ti3C2) by annealing the Ti3C2 under an ammonia (NH3) atmosphere [61]. After annealing treatment by NH3, N is introduced to substitute C atoms in the MXene, resulting in the increased interlayer spacing of Ti3C2 from 0.96 to 1.23 nm at 200  C. In addition to the above methods, the layer-by-layer assembly can also be used for constructing MXene-based nanostructures with enlarged interlayer spacing. Compared with liquid phase pre-pillaring and pillaring method, precursor intercalation and annealing/polymerization method, single layer MXenes are often required in layer-by-layer assembly method, which may need an extra exfoliation process. However, the layer-by-layer assembly can be used to fabricate free-standing electrodes, and it can assemble MXene with different dimensional nanomaterials, which has excellent flexibility. Hamedi et al. used a layer-by-layer self-assembly method to fabricate pillared Ti3C2 MXene/tris(2-aminoethyl) amine (TAEA) multilayers (Fig. 16.3k) [62]. The obtained pillared structure is highly ordered, and this structure reinforces the interconnection between MXene and TAEA. Notably, the TAEApillared MXene multilayers exhibit a high electronic conductivity of 7.3*104 Sm
1. The high electronic conductivity is ascribed to the face-to-face quasi-intimate interface contact between MXene and TAEA. Moreover, Gogotsi et al. prepared a flexible and free-standing sandwich-like Ti3C2 MXene/carbon nanotube (CNT) paper by layer-by-layer assembly method [33]. After CNT intercalation, the interlayer spacing of Ti3C2 increases from 1.21 to 1.42 nm. Furthermore, onionlike carbon, reduced graphene oxide (rGO), activated carbon, and hard carbon can also be fabricated by assembling with MXene [63–65]. The resulting MXene-based papers have good electrical conductivities, high surface areas accessible to ions, and mechanical robustness. Given the abovementioned samples, we summed up the state-of-the-art research works about designing MXene-based nanostructures with enlarged interlayer spacing via the introduction of intercalation agents. In general, the intercalation agents can be classified into the following categories: (1) Heteroatoms; (2) Alkali metal ions; (3) Non-alkali metal ions; (4) Polymers; (5) Large cations; (6) 1D/2D/3D nanomaterials. The main fabrication methods of MXene-based nanostructures with enlarged interlayer spacing, including (1) liquid phase pre-pillaring and pillaring; (2) precursor intercalation and annealing/polymerization; (3) layer-by-layer assembly. After the intercalation of intercalation agents, MXene-based nanostructure with enlarged interlayer spacing shows obvious characteristics for energy storage

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

461

applications: (1) large interlayer space for charge-carrying ions intercalation and storage; (2) excellent ion transport due to short diffusion pathways; (3) limitation of volume expansion of active materials by the confinement effect of a 2D structure during cycling. Besides, by carefully selecting the types of surfactant molecules, pillared MXene structures with controllable interlayer spacings from 1 to 2.7 nm can be prepared to match the sizes of the ions in the electrolyte, maximizing the electrochemical performance [43]. Furthermore, this fabrication method of pillared MXene is conducive to the controllable modification of the surface chemistry about MXene, which is one of the main challenges in designing MXene-based nanostructure [31].

16.3

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage Applications

MXene has a wide chemical and structural variety, which makes them competitive with other 2D materials for energy storage applications. After HF or in situ HF etching, MXenes with a unique layered structure and interfacial characteristics can be obtained. Taking the case of Ti3AlC2, after HF etching, the corresponding interlayer spacing of Ti3C2 is 0.98 nm, and the –OH and/or –F groups present at the surface of the Ti3C2 layers. MXene has interlayer space for cations intercalation and storage. If the interlayer space of MXenes can be further increased and utilized, the electrochemical performance of MXene-based electrodes will be further enhanced [43]. Theoretical studies confirm that lithium (Li) atoms can directly be adsorbed on the O terminated MXene monolayers (Fig. 16.4a), and the theoretical Li capacity of Ti3C2O2 is 268 mAh g
1 based on the configuration of Ti3C2O2Li2 [66]. The valence electron localization function (ELF) of Ti3C2O2 clearly shows the electron transfer between C and Ti and between Li and O (Fig. 16.4a). There are few electrons localized above the adsorbed Li layer. Importantly, the extra Li layers adsorption is also studied (Fig. 16.4b) [66]. With the extra Li layers, the interaction between the different Li layers is visible in the ELF, which indicates that the extra Li layer can bind to the already lithiated MXene. The structure of MXene with additional Li layers is stable (Fig. 16.4b) [66]. In the MXenes, the first extra Li layer forms about 2.8 Å above the lithiated surface. Besides, the second and third extra Li layers require 2.3 Å for each layer. Therefore, a larger space is needed to accommodate extra Li layers, which indicates MXene with enlarged interlayer spacing can store extra Li layers. Also, theoretical Na storage capacities of Ti3C2, Ti3C2F2, and Ti3C2O2, and change of Ti3C2 MXene volume after Na adsorption are explored (Fig. 16.4c–e) [42]. The results show that enlarged interlayer spacing of bare and O-functionalized Ti3C2 MXene enables stable multilayer adsorption and therefore significantly enhances their theoretical capacities. Furthermore, two layers

462

J. Luo

Fig. 16.4 (a) Side view of lithiated monolayers of Ti3C2O2 and valence ELF of (110) sections of lithiated Ti3C2O2Li2 monolayer. (b) Side view of an extra metallic Li layer adsorbed on top of lithiated Ti3C2O2Li2 monolayers and valence ELF of (110) sections of lithiated Ti3C2O2Li2 monolayer with an extra Li layer. Reproduced with permission-Copyright 2014, American Chemical Society [66]. (c) Geometric structures for bilayer adsorption on both sides of bare and O-functionalized Ti3C2 MXenes. (d) Capacity as a function of x in Ti3C2T2Nax with T ¼ bare, F, and O. (e) Percentage changes of volume as a function of x in Ti3C2T2Nax with T ¼ bare, F, and O. Reproduced with permission-Copyright 2016, American Chemical Society [42]. (f) HAADF image of Ti3C2X electrode upon Na intercalation with the cutoff potential of 0 V. Reproduced with permission-Copyright 2015, American Chemical Society [67]

of sodium (Na) intercalation between Ti3C2 MXene layers have been confirmed in the HAADF image of the Ti3C2X electrode (Fig. 16.4f) [67]. In general, the design of MXene-based nanostructure with enlarged interlayer spacing provides a new route to improve their energy storage performance. In this section, a series of interlayer structure design along with the electrochemical energy storage performance are systematically overviewed, including Li-ion batteries (LIBs), Na-ion batteries (SIBs), supercapacitors, Li/Na-ion capacitors (LICs/SICs), Na metal batteries (SMBs), potassium-ion batteries (PIBs), and magnesium ion batteries (MIBs).

16.4

Supercapacitors

Supercapacitors are regarded as one type of very promising energy storage technology since they possess higher power densities than batteries while their energy densities are superior to conventional capacitors [68]. Supercapacitors could play a role to bridge the energy and power gap between conventional capacitors and batteries [68]. MXenes’ unique structure renders them attractive for supercapacitors

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

463

applications [69], including (1) A conductive inner transition metal carbide layer enables fast electron supply to electrochemically active sites; (2) A transition metal oxide-like surface is redox-active; (3) A 2D morphology and pre-intercalated water enable fast ion transport. Therefore, MXenes are promising candidates as electrodes for energy storage applications. Gogotsi et al. demonstrate both the spontaneous and electrochemical intercalation of cations, including Na+, K+, NH4+, Mg2+, and Al3+, from aqueous salt solutions between 2D Ti3C2 MXene layers (Fig. 16.5a) [54]. After cations intercalation, the interlayer spacing of Ti3C2 MXene increases. In basic solutions, such as KOH and NaOH, the binder-free Ti3C2 paper exhibits highly flexible and yields volumetric capacitance of up to 350 F cm
3, showing notable intercalation capacitances (Fig. 16.5b, c). Besides, the fact that a variety of ions, including Na+ and Al3+, can be intercalated between the MXene layers may also enable MXene use in batteries as well as in metal-ion capacitors (battery–supercapacitor hybrids) [54]. However, the sizes of metal ions are not large enough. If larger size molecules are used as intercalation agents, the interlayer spacing of MXene can be further increased, and the corresponding electrochemical performance can be further improved [43]. Yang and Gogotsi et al. used the mechanical shearing of a discotic lamellar liquid crystal phase of Ti3C2 to assemble vertically aligned 2D Ti3C2 (Fig. 16.5d) [70]. Notably, a non-ionic surfactant, hexamethylene glycol monododecyl ether (C12E6), was introduced to enhance molecular interactions between the vertically aligned MXene nanosheets, therefore increasing the interlayer spacing to 5.8 nm (Fig. 16.5e). Compared with the traditional electrode that has limited ion transport in thick films, the vertical alignment of MXene flakes enables directional ion transport, which can lead to thickness-independent electrochemical performances in thick films. The rate performance of the vertically aligned MXene films decreases only slightly when the thickness of film increases from 40 to 200μm, especially for the scan rates below 2000 mV s
1 (Fig. 16.5f) [70]. Therefore, the vertical alignment of functional nanomaterials with enlarged interlayer spacing through the manipulation of their liquid crystal mesophase confirmed here provides a new and powerful technique to construct advanced architectures with exceptional performance. Besides, polymers can also be used as intercalation agents between MXene layers. Ling et al. mixed a charged poly-diallyl dimethylammonium chloride (PDDA) or polyvinyl alcohol (PVA) and Ti3C2 MXene to produce Ti3C2/polymer composites (Fig. 16.5g) [59]. The obtained polymers intercalated MXene composite films that have excellent flexibility, good tensile and compressive strengths, and electrical conductivity. The successful intercalation of polymers between Ti3C2 layers, which increases the interlayer spacing of MXene and can be confirmed by the high-resolution transmission electron microscopy (HRTEM) (Fig. 16.5h). Used as electrodes for supercapacitors, typical cyclic voltammograms (CVs) of Ti3C2, Ti3C2/ PDDA, and Ti3C2/PVA-KOH films are shown in Fig. 16.5i. Compared with Ti3C2 film, the enhancement in the capacitance of Ti3C2/PVA-KOH film can be ascribed to the intercalation of PVA between MXene flakes with enlarged interlayer space, improving access to deep trap sites [59]. In general, the composite films exhibit impressive volumetric capacitance values as high as 528 F cm
3 at 2 mV s
1 and

464

J. Luo

Fig. 16.5 (a) Schematic illustration of the intercalation of cations between Ti3C2Tx MXene layers. (b) SEM image of the Ti3C2Tx MXene paper electrode. (c) CV curve of Ti3C2Tx paper in the KOH electrolyte. (d) Ion transport in vertically aligned Ti3C2Tx MXene films. (e) SEM image of vertical nanosheets on the horizontally aligned MXene current collector. (f) Rate performance of MXLLC films at scan rates ranging from 10 to 100,000 mV s
1. Reproduced with permission-Copyright 2019, Springer Nature [70]. (g) Schematic illustration of MXene/PVA film with outstanding properties. (h) TEM images of Ti3C2Tx/PVA. (i) CV curves of electrodes were obtained at a scan rate of 2 mV s
1. (j) Schematic illustration of the structure of polypyrrole-Ti3C2Tx. (k) Crosssectional TEM image of polypyrrole-Ti3C2Tx. (i) Rate performance of tested compositions and comparisons of their capacitances with previously reported Ti3C2Tx electrode. Reproduced with permission-Copyright 2016, John Wiley and Sons [60]

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

465

306 F cm
3 at 100 mV s
1. Moreover, Boota et al. synthesized polypyrrole (PPy) intercalated Ti3C2 composite by simultaneous intercalation, alignment, and metalfree polymerization of pyrrole between Ti3C2 layers (Fig. 16.5j, k) [60]. When tested as supercapacitor electrodes, a PPy/Ti3C2 composite yields a high volumetric capacitance of ~1000 F cm
3 and capacitance retention of 92% after 25,000 cycles (Fig. 16.5l). The high electrochemical performance and excellent cycling performance of PPy/Ti3C2 composite are ascribed to the highly conductive nanostructure of the PPy/Ti3C2 film in which the aligned PPy chains between the metallic Ti3C2 layers provide low electronic resistance and create aligned channels for facile ionic transport. 2D conductive Ti3C2 layers may protect the intercalated PPy chains and minimize their degradation. Besides, other intercalation agents, such as N heteroatom, activated carbon, CNT, and rGO, can also be used to design MXene-based nanostructure with enlarged interlayer spacing [33, 61, 63, 64]. The above MXene-based electrodes all exhibit high capacitances and perform well at high rates due to additional and fast diffusion paths for electrolyte ions as well as large interlayer space for cations intercalation.

16.5

Li-ion Batteries/Capacitors

There are continuous efforts in the pursuit of high-performance Li-ion batteries/ capacitors with high energy/power density to meet the increasing demand for modern electronics and transportation [43]. In the family of MXenes, Ti3C2 is the most studied material as anode for LIBs [12, 44, 71]. Naguib et al. first used Ti3C2 MXene as the anodes for LIBs [71]. At a current density of 1 C, a stable capacity of 110 mAh g
1 can be obtained after 80 cycles. When the current density increases to 10 C, a capacity of 70 mAh g
1 can still be obtained for the Ti3C2 electrode. Although the capacity of Ti3C2 MXene is relatively low at the initial stage of research, there is a huge room for improvement. Luo et al. first synthesized the Sn4+ pillared MXene (PVP-Sn(IV)@Ti3C2) by the use of alkali metal ions to pre-pillar Ti3C2 MXene, followed by the Sn4+ pillaring process (Fig. 16.6a) [44]. By the ion-exchange interaction between Sn4+ and pre-intercalated alkali metal ions under the help of polyvinylpyrrolidone (PVP), Sn4+ can be successfully intercalated between Ti3C2 layers. The intercalation of Sn4+ in the interlayer of Ti3C2 MXene increases the interlayer spacing of Ti3C2 from 0.98 to 1.28 nm by forming strong Ti-O-Sn bonding. The SEM image shows that Sn4+, in the form of Sn (IV) nanocomplex, is uniformly distributed on the Ti3C2 matrix (Fig. 16.6b). Used as the anodes for LIBs, the PVP-Sn(IV)@Ti3C2 nanocomposites exhibit a superior reversible volumetric capacity of 1375 mAh cm
3 at 100 mA g
1 after 100 cycles (Fig. 16.6c). Even at a high current density of 3 A g
1, a stable specific capacity of 504.5 mAh cm
3 can be retained. The excellent electrochemical performance can be ascribed to the “pillar effect.” During the lithiation process, Li+ can alloy with the intercalated Sn4+ to form Li-Sn alloy, causing the volume expansion, which can effectively prop the MXene layers open and endow the MXene with increased

466

J. Luo

Fig. 16.6 (a) Schematic illustration of the structure of Sn4+ pillared Ti3C2 MXene (PVP-Sn(IV) @Ti3C2). (b) SEM image of PVP-Sn(IV)@Ti3C2. (c) Rate performance of PVP-Sn(IV)@Ti3C2. Reproduced with permission-Copyright 2016, American Chemical Society [44]. (d) Schematic illustration of the intercalation of Co2+ between V2C MXene layers. (e) SEM image of Co2+ intercalated V2C (V2C@Co). (f) Rate performance of Co2+ intercalated V2C. Reproduced with permission-Copyright 2018, John Wiley and Sons [57]. (g) Schematic illustration of the structure of Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2). (h) STEM image of CTAB-Sn(IV)@Ti3C2. Inset is the elemental mapping of Sn. (i) Cycling performance of CTAB-Sn(IV)@Ti3C2. Reproduced with permission-Copyright 2016, American Chemical Society [43]. (j) Schematic illustration of the sandwich-like MXene/CNT papers. (k) Cross-sectional SEM image of Nb2CTx/ CNT composite paper. (l) Cycling stability of the Nb2CTx/CNT paper electrode at different cycling rates. Reproduced with permission-Copyright 2015, John Wiley and Sons [50]

interlayer space for Li+ intercalation, namely the “pillar effect” [44]. Besides, Co2+ can also be intercalated between the MXene layers [57]. Similarly, Song et al. fabricated Co2+ intercalated V2C MXene (V2C@Co) (Fig. 16.6d) [57]. After Co2+

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

467

intercalation, strong V–O–Co bonds can be formed between Ti3C2 MXene layers. The SEM image of V2C@Co displays no obvious layer stripes, mainly attributes to the intercalation of Co2+ (Fig. 16.6e). Owing to the unique layered structure and pseudocapacitive contribution of Co species, the obtained V2C@Co exhibits improved Li-ion capacity and delivers a high capacity of 1117.3 mAh g
1 at 0.1 A g
1 and a significantly ultralong cycling stability over 15,000 cycles (Fig. 16.6f) [57]. Although alkali metal ion can pre-pillar and increase the interlayer spacing of MXene, the increments of interlayer spacing are limited due to their small size. Luo et al. selected different sizes of cationic surfactants to intercalate between MXene layers and controlled the interlayer spacing of MXene from 1 to 2.71 nm [43]. The intercalated cationic surfactants in the interlayer of MXene can provide the ion-exchange sites for other metal ions. Tao et al. successfully fabricated Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2) by using liquid-phase CTAB pre-pillaring and subsequent Sn4+ pillaring process (Fig. 16.6g) [43]. STEM image of CTAB-Sn(IV)@Ti3C2 confirms the successful intercalation of Sn4+ between Ti3C2 layers (Fig. 16.6h). As anode for lithium-ion storage, the obtained CTABSn(IV)@Ti3C2 electrode delivers a high reversible capacity of 765 mAh g
1 at 0.1 A g
1 after 100 cycles, which exhibits a higher capacity of CTAB intercalated Ti3C2 (248 mAh g
1) (Fig. 16.6i). The excellent Li-ion storage performance of CTAB-Sn(IV)@Ti3C2 is ascribed to the “pillar effect.” Besides, the CTAB-Sn(IV) @Ti3C2 anode exhibits good Li+ storage kinetics. Coupling the CTAB-Sn(IV) @Ti3C2 anode with commercial AC cathode, the assembled LICs exhibits a high energy density of 239.50 Wh kg
1 [43]. Besides, Gogotsi’s group used a layer-by-layer method to assemble a CNT/Nb2C MXene composite “paper” electrode (Fig. 16.6j) [50]. THE cross-sectional SEM image of Nb2C/CNT composites exhibits an open structure (Fig. 16.6k). After CNT intercalation between Nb2C layers, the obtained Nb2C/CNT papers exhibit good electrical conductivity, high surface areas accessible to ions, and mechanical robustness, suggesting their promising application for Li-ion storage. The obtained CNT/MXene composite “paper” electrode exhibits a high capacitance of 325 F cm
3 as an electrode for LICs (Fig. 16.6l). Also, other metal ions and nanomaterials can be used as the intercalation agents to build MXene-based nanostructures with enlarged interlayer spacing and used as electrodes for LIBs and LICs [72, 73]. Compared with aqueous electrolytes, the size of the charge carrying ions in the organic electrolyte is larger [31]. Therefore, constructing MXene-based nanostructures with enlarged interlayer spacing is crucially important to maximize their electrochemical performance.

16.6

Na-ion Batteries/Capacitors

Na-ion storage is a promising alternative to Li-ion storage due to the similar working mechanism and significant advantages as for elemental abundance and low cost of Na resources [36, 74–76]. However, compared with the smaller radius of Li ions

468

J. Luo

Fig. 16.7 (a) The structure of MXene/CNT composite. (b) SEM image of porous Ti3C2Tx/CNTSA. (c) Volumetric and gravimetric rate performances. Reproduced with permission-Copyright 2019, John Wiley and Sons [36]. (d) The structure of S atoms pillared Ti3C2 MXene (CT-S@Ti3C2). (e) STEM image of CT-S@Ti3C2-450. Inset is the elemental mapping of S. (f) Rate performance of CT-S@Ti3C2-450. Reproduced with permission-Copyright 2015, Springer Nature [17]. (g) Schematic illustration of the reaction mechanism of Ti2CTx by electrochemical activation. (h) TEM image of activated Ti2CTx after the first CV. (i) Rate capability for Ti2CTx with other reported electrodes. Reproduced with permission-Copyright 2016, Elsevier [79]

(0.76 Å), Na ions have a much larger radius (1.02 Å), which makes it difficult to find appropriate host materials for Na ions storage due to the large volume expansion and sluggish reaction kinetics. Therefore, seeking novel electrode materials for Na-ion storage is still desirable. Fortunately, MXene is reported to be a promising anode material for SIBs and SICs [36]. Heteroatoms intercalated MXene can be fabricated with enhanced Na-ion storage performance. Luo et al. successfully intercalated S atoms into the interlayer of Ti3C2 MXene (CT-S@Ti3C2) by a facile CTAB pretreatment, thermal diffusion with elemental S, and subsequent annealing process (Fig. 16.7a) [36]. After annealing under the Ar atmosphere, S atoms pillared Ti3C2 with interlayer-expanded structure via Ti-S bonding can be developed. The interlayer spacing of CT-S@Ti3C2 is about 1.37 nm, which is larger than that of CTAB pretreated Ti3C2 after annealing under

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

469

the same temperature (1.15 nm). The successful intercalation of S atoms between Ti3C2 MXene layers can be confirmed by the STEM image and the corresponding elemental mapping results of CT-S@Ti3C2 (Fig. 16.7b). Used as the anode for Na-ion storage, the CT-S@Ti3C2 delivers a high capacity of 550 mAh g
1 at 0.1 A g
1v (120 mAh g
1 at 15 A g
1) and excellent cycling stability over 5000 cycles at 10 A g
1 (Fig. 16.7c). The excellent electrochemical performance is ascribed to the strong Ti-S bonds, which not only can react with the intercalated Na that contribute to the “pillar effect” but also can provide stable double-layer Na ions adsorption that is confirmed by DFT calculation [36]. Furthermore, the obtained CT-S@Ti3C2 exhibits fast Na ion storage kinetics, which has been confirmed by the kinetic analysis. A high energy density of 263.2 Wh kg
1 can be obtained for the assembled SICs by coupling CT-S@Ti3C2 anode with commercial AC cathode. In addition to the introduction of heteroatoms, alkali metal ions can also be intercalated between the MXene layers and alter their electrochemical behaviors [17, 77]. Yamada et al. realized the usage of MXene as the electrode in a nonaqueous Na+ electrolyte. The results show that after the initial Na+ intercalation and activation processes, the Ti2C MXene electrode exhibits pseudocapacitor behavior (Fig. 16.7d) [17]. After the initial Na+ intercalation, the interlayer spacing of Ti2C increases from 0.8 to 1.0 nm, which can be confirmed by the TEM image (Fig. 16.7e). Compared with the rate capability of Ti2C with other electrode materials, such as hard carbon, expanded graphite, and P2-Na0.66[Li0.22Ti0.78]O2, Ti2C delivers higher capacity at higher current density. Therefore, Ti2C is a high-performance electrode material with a high capacity, stability, and high power for Na-ion storage. Luo et al. studied the Na+ storage behaviors of different alkali metal ions (e.g., Li+, Na+, K+) pillared Ti3C2 MXene [78]. Meanwhile, they annealed the Na+ pillared Ti3C2 MXene at different temperatures (e.g., 450 and 700  C) to regulate the terminal groups in Ti3C2 (Fig. 16.7f). The results show that Na+ intercalation during the pillaring process can increase the number of active sites for Na+ storage, and it also can decrease the Na+ diffusion barrier. In general, the results confirm that the existence of –OH groups can provide active sites for Na+ storage, while decrease the Na+ storage kinetics [78]. Compared with the most used liquid phase intercalation method to build MXenebased nanostructure with enlarged interlayer spacing, layer-by-layer assembly is more flexible because many nanomaterials can be used as intercalation agents by this method. Xie et al. prepared Ti3C2 MXene/CNT films by self-assembly of negatively charged Ti3C2 MXene and positively charged CNTs as intercalation spacers (Fig. 16.7g) [79]. In the SEM image of Ti3C2/CNT, an open sheet arrangement and CNTs can be seen sandwiched between Ti3C2 sheets, indicating the successful placement of CNTs as intercalation spacers in the Ti3C2 paper (Fig. 16.7h). The 2D/1D hybridization between MXene nanosheets and CNTs reduces the stacking of nanosheets produces hierarchical films with a porous structure, thereby improving the accessibility of MXene nanosheets to the electrolyte. Used as free-standing electrodes for Na-ion storage, the porous MXene/CNT papers exhibit a high volumetric capacity of 421 mAh cm
3 at 20 mA g
1 (Fig. 16.7i) [79].

470

16.7

J. Luo

Other Energy Storage Systems

Except for supercapacitors, LIBs, LICs, SIBs, and SICs, MXene-based nanostructures with enlarged interlayer spacing can also be used as the electrodes for other energy storage applications, including SMBs, PIBs, and MIBs [65, 80, 81]. Luo et al. used the liquid phase CTAB pre-pillaring and Sn2+ pillaring method to fabricate Sn2+ pillared Ti3C2 MXene (CT-Sn(II)@Ti3C2) and serve as a stable matrix for Na metal anode [80]. After Sn2+ intercalation, the interlayer spacing of Ti3C2 MXene increases from 0.98 to 1.9 nm. The intercalated Sn2+ in this work not only induces Na to nucleate and grow between Ti3C2 MXene layers but also endows Ti3C2 with larger interlayer space to accommodate the deposited Na under the effect of “pillar effect,” thus resulting in the uniform Na deposition [80]. Xu et al. used Ti3C2 MXene as a multifunctional binder to fabricate a free-standing and flexible hard carbon film electrode by a simple vacuum-assisted filtration strategy [65]. In the MXene-bonded hard carbon film, hard carbon film particles, as the active materials, are embedded in the 3D conductive network of MXene sheets. MXene nanosheets with enlarged interlayer spacing can effectively stabilize the electrode structure and accommodate the volume expansion of hard carbon particles during the K+ intercalation. Used as the anode for PIBs, the MXene-bonded hard carbon anode delivers a high reversible capacity of 210 mAh g
1 at the current density of 50 mA g
1 after 100 cycles with a capacity retention of 84%. Besides, Yan et al. used CTAB to intercalate between Ti3C2 MXene layers and increased the interlayer spacing of Ti3C2 [81]. The obtained CTAB intercalated Ti3C2 MXene exhibits a high Mg2+ storage capability. DFT simulations verify that the intercalated CTA+ cations reduce the diffusion barrier of Mg2+ on the MXene surface. Consequently, the MXene electrode exhibits a desirable volumetric specific capacity of 300 mAh cm
3 at 50 mA g
1.

16.8

Conclusions and Perspectives

MXenes have been widely applied in the field of electrochemical energy storage due to their high electrical conductivity, large redox-active surface area, rich surface chemistry, and tunable structures. Demonstrably, in recent years, there has been a rapid increase in the number of publications on MXenes and MXene for electrochemical energy storage applications (Fig. 16.2d). In this chapter, we first briefly summarize the synthesis techniques of MXenes, mainly engross in the HF/in situ HF etching synthesis. Different synthesis techniques used to synthesize MXene directly influence the structure and properties of

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

471

the obtained MXenes. In the research process of layered structure MXene for electrochemical energy storage, a better comprehension of the layered structure and interfacial chemistry of MXene is requisite. Up to now, the interfacial structure design of MXene-nanostructures with enlarged interlayer spacing can be classified according to different intercalation agents, including (1) heteroatoms; (2) alkali metal ions; (3) non-alkali metal ions; (4) polymers; (5) large cations; (6) 1D/2D/ 3D nanomaterials. The main fabrication methods of MXene-based nanostructures with enlarged interlayer spacing, including (1) liquid phase pre-pillaring and pillaring; (2) precursor intercalation and annealing/polymerization; (3) layer-bylayer assembly. After the intercalation of intercalation agents, MXene-based nanostructure with enlarged interlayer spacing shows obvious characteristics for energy storage applications: (1) large interlayer space for charge-carrying ions intercalation and storage; (2) excellent ion transport due to short diffusion pathways; (3) limitation of volume expansion of active materials by the confinement effect of a 2D structure during cycling. Besides, by carefully selecting the types of surfactant molecules, pillared MXene structures with controllable interlayer spacings from 1 to 2.7 nm can be prepared to match the sizes of the ions in the electrolyte, maximizing the electrochemical performance. Furthermore, this fabrication method of pillared MXene is conducive to the controllable modification of the surface chemistry about MXene, which is one of the main challenges in designing MXene-based nanostructure. Interlayer structural engineering of 2D MXene, synthesis methods, and their electrochemical energy storage applications are summarized in Table 16.1. Although the research of designing MXene-based nanostructure for electrochemical energy storage applications has been widely studied, some remaining challenges still need to be overcome for the future development of MXene. For future research, the following aspects are particularly important. (1) Achieving controllable and uniform terminations on the interface of MXene is necessary. (2) A thorough understanding and systematical study on the properties of intercalation agents (e.g., ionic, molecular species), including transport, stability, and bonding environment, and their effect on the electrochemical and physicochemical properties of layered structure MXenes is necessary. (3) A basic understanding of the property of active centers (interfacial transition metal layers and terminations) for ion transport and intercalation, and the energy storage mechanism by the cooperation between theoretical calculation and experiment lays solid foundations for the development of MXenes in electrochemical energy storage.

Polymer (PVA, PDDA, PPy)

Metal ions (Sn4+, Co2+, Sn2+)

Alkali metal ions

Intercalation agent types Heteroatoms (N, S)

PPy/l-Ti3C2

PPy-Ti3C2Tx

Ti3C2Tx/ PDDA Ti3C2Tx/PVA

CT-Sn(II) @Ti3C2

Multilayer Ti3C2 PVP-Sn(IV) @Ti3C2 V2C@Co

Materials N-Ti3C2Tx CT-S@Ti3C2

Liquid phase mixture (RT) Liquid phase mixture (RT) Precursor intercalation and in situ polymerization Electrochemical polymerization

Methods NH3 annealing (200  C) Precursor intercalation and annealing (450  C) Liquid phase intercalation (RT) Intercalation by ion exchange (RT) Intercalation by ion exchange (RT) Intercalation by ion exchange (RT)

0.96

Supercapacitors

Supercapacitors

Supercapacitors

2.25 2.40

Supercapacitors

1.84

Na metal matrix

Li-ion batteries

0.95 1.90

Li-ion batteries

Applications Supercapacitors Na-ion capacitors Supercapacitors

1.28

1.25

Interlayer spacing 1.23 1.37

59

530 F cm
3 at 2 mV s
1

35 mF cm
2

~1000 F cm
3 with capacitance retention of 92% after 25,000 cycles

59

58

60

80

57

44

54

References 61 36

1375 mAh cm
3 (635 mAh g
1) at 0.1 A g
1 1117.3 mAh g
1 at 0.1 A g
1, 199.9 mAh g
1 at 20 A g
1 High coulombic efficiency over 500 cycles (up to 10 mA cm
2 and 5 mAh cm
2) ~296 F cm
3 at 2 mV s
1

Performance 192 F g
1 in 1 M H2SO4 550 mAh g
1 at 0.1 A g
1 and ~120 mAh g
1 at 15 A g
1 ~350 F cm
3 after 10 K cycles

Table 16.1 Interlayer structural engineering of 2D MXene, synthesis methods, and their electrochemical energy storage applications

472 J. Luo

1D/2D nanomaterials

Large cations (CTA+, STA+, DDA+, amine ions)

MoS2-in-Ti3C2

Nb2CTx + iPrA + CNTs Ti3C2Tx/ MWCNT MXene-rGO

DDAC@Ti3C2

STAB-Sn(IV) @Ti3C2

CTAB-Sn(IV) @Ti3C2

Electrostatic selfassembly Annealing (500  C)

Liquid phase pre-pillaring and pillaring (40  C) Liquid phase pre-pillaring and pillaring (50  C) Liquid phase intercalation (70  C) Liquid phase mixture (RT) Lay-by-layer assembly –

2.10

1.22

1.67

1.42

Li/Na-ion batteries

Supercapacitors

Supercapacitors

Li-ion batteries

Li-ion capacitors

2.71

2.27

Li-ion capacitors

2.23

43

~450 mAh g
1 at 1 A g
1

63

1040 F cm
3 at 2 mV s
1

72

33

435 F cm
3 at 2 mV s
1

340 mAh g
1 at 20 A g
1 for Li+ storage; 310 mAh g
1 at 1 A g
1 for Na+ storage

50

~420 mAh g
1 at 0.5 C

–

43

765 mAh g
1 at 0.1 A g
1

16 Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . . 473

474

J. Luo

Acknowledgments I wish to thank the support of the Thayer School of Engineering, Dartmouth College, 14 Engineering Drive, Hanover, New Hampshire. Author Contribution The article was conceived and written by JL.

References 1. Novoselov KS, Geim AK, Morozov SV, Jiang D, Zhang Y, Dubonos SV, Grigorieva IV, Firsov AA (2004) Electric field effect in atomically thin carbon films. Science 306:666–669 2. Song L, Ci LJ, Lu H et al (2010) Large scale growth and characterization of atomic hexagonal boron nitride layers. Nano Lett 10:3209–3215 3. Tan CL, Cao XH, Wu XJ et al (2017) Recent advances in ultrathin two-dimensional nanomaterials. Chem Rev 117:6225–6331 4. Sun J, Lee HW, Pasta M, Yuan HT, Zheng GY, Sun YM, Li YZ, Cui Y (2015) A phosphorenegraphene hybrid material as a high-capacity anode for sodium-ion batteries. Nat Nanotechnol 10:980–985 5. Li H, Eddaoudi M, O’Keeffe M, Yaghi OM (1999) Design and synthesis of an exceptionally stable and highly porous metal-organic framework. Nature 402:276–279 6. Naguib M, Kurtoglu M, Presser V, Lu J, Niu JJ, Heon M, Hultman L, Gogotsi Y, Barsoum MW (2011) Two-dimensional nanocrystals produced by exfoliation of Ti3AlC2. Adv Mater 23:4248–4253 7. Naguib M, Mashtalir O, Carle J, Presser V, Lu J, Hultman L, Gogotsi Y, Barsoum MW (2012) Two-dimensional transition metal carbides. ACS Nano 6:1322–1331 8. Naguib M, Mochalin VN, Barsoum MW, Gogotsi Y (2014) 25th anniversary article: MXenes: a new family of two-dimensional materials. Adv Mater 26:992–1005 9. Urbankowski P, Anasori B, Makaryan T et al (2016) Synthesis of two-dimensional titanium nitride Ti4N3 (MXene). Nanoscale 8:11385–11391 10. Halim J, Kota S, Lukatskaya MR et al (2016) Synthesis and characterization of 2D molybdenum carbide (MXene). Adv Funct Mater 26:3118–3127 11. Anasori B, Lukatskaya MR, Gogotsi Y (2017) 2D metal carbides, and nitrides (MXenes) for energy storage. Nat Rev Mater 2:16098 12. Sun DD, Wang MS, Li ZY, Fan GX, Fan LZ, Zhou AG (2014) Two-dimensional Ti3C2 as anode material for Li-ion batteries. Electrochem Commun 47:80–83 13. Tang Q, Zhou Z, Shen PW (2012) Are MXenes promising anode materials for Li-ion batteries? Computational studies on electronic properties and Li storage capability of Ti3C2 and Ti3C2X2 (X ¼ F, OH) monolayer. J Am Chem Soc 134:16909–16916 14. Dong YF, Wu ZS, Zheng SH, Wang XH, Qin JQ, Wang S, Shi XY, Bao XH (2017) Ti3C2 MXene-derived sodium/potassium titanate nanoribbons for high-performance sodium/potassium ion batteries with enhanced capacities. ACS Nano 11:4792–4800 15. Yang Q, Huang ZD, Li XL et al (2019) A wholly degradable, rechargeable Zn-Ti3C2 MXene capacitor with superior anti-self-discharge function. ACS Nano 13:8275–8283 16. Zhang CF, Kremer MP, Seral-Ascaso A, Park SH, McEvoy N, Anasori B, Gogotsi Y, Nicolosi V (2018) Stamping of flexible, coplanar micro-supercapacitors using MXene inks. Adv Funct Mater 28:1705506 17. Wang X, Kajiyama S, Iinuma H, Hosono E, Oro S, Moriguchi I, Okubo M, Yamada A (2015) Pseudocapacitance of MXene nanosheets for high-power sodium-ion hybrid capacitors. Nat Commun 6:6544 18. Liang X, Garsuch A, Nazar LF (2015) Sulfur cathodes based on conductive MXene nanosheets for high-performance lithium-sulfur batteries. Angew Chem Int Ed Engl 54:3907–3911

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

475

19. Wang J, Tang J, Ding B et al (2017) Hierarchical porous carbons with layer-by-layer motif architectures from confined soft-template self-assembly in layered materials. Nat Commun 8:15717 20. Li J, Yuan XT, Lin C, Yang YQ, Xu L, Du X, Xie JL, Lin JH, Sun JL (2017) Achieving high pseudocapacitance of 2D titanium carbide (MXene) by cation intercalation and surface modification. Adv Energy Mater 7:1602725 21. Bak SM, Qiao RM, Yang WL, Lee S, Yu XQ, Anasori B, Lee H, Gogotsi Y, Yang XQ (2017) Na-ion intercalation, and charge storage mechanism in 2D vanadium carbide. Adv Energy Mater 7:1700959 22. Shahzad F, Alhabeb M, Hatter CB, Anasori B, Hong SM, Koo CM, Gogotsi Y (2016) Electromagnetic interference shielding with 2D transition metal carbides (MXenes). Science 353:1137–1140 23. Peng QM, Guo JX, Zhang QR, Xiang JY, Liu BZ, Zhou AG, Liu RP, Tian YJ (2014) Unique lead adsorption behavior of activated hydroxyl group in two-dimensional titanium carbide. J Am Chem Soc 136:4113–4116 24. Ying YL, Liu Y, Wang XY, Mao YY, Cao W, Hu P, Peng XS (2015) Two-dimensional titanium carbide for efficiently reductive removal of highly toxic chromium(VI) from water. ACS Appl Mater Interfaces 7:1795–1803 25. Seh ZW, Fredrickson KD, Anasori B, Kibsgaard J, Strickler AL, Lukatskaya MR, Gogotsi Y, Jaramillo TF, Vojvodic A (2016) Two-dimensional molybdenum carbide (MXene) as an efficient electrocatalyst for hydrogen evolution. ACS Energy Lett 1:589–594 26. Ran JR, Gao GP, Li FT, Ma TY, Du AJ, Qiao SZ (2017) Ti3C2 MXene co-catalyst on metal sulfide photo-absorbers for enhanced visible-light photocatalytic hydrogen production. Nat Commun 8:13907 27. Lin H, Wang XG, Yu LD, Chen Y, Shi JL (2017) Two-dimensional ultrathin MXene ceramic nanosheets for photothermal conversion. Nano Lett 17:384–391 28. Xiao X, Urbankowski P, Hantanasirisakul K et al (2019) Scalable synthesis of ultrathin Mn3N2 exhibiting room-temperature antiferromagnetism. Adv Funct Mater 29:1809001 29. Zha XH, Zhou J, Zhou YH, Huang Q, He J, Francisco JS, Luo K, Du SY (2016) Promising electron mobility and high thermal conductivity in Sc2CT2 (T ¼ F, OH) MXenes. Nanoscale 8:6110–6117 30. Yasaei P, Hemmat Z, Foss CJ et al (2018) Enhanced thermal boundary conductance in few-layer Ti3C2 MXene with encapsulation. Adv Mater 30:1801629 31. Simon P (2017) Two-dimensional MXene with controlled interlayer spacing for electrochemical energy storage. ACS Nano 11:2393–2396 32. Zhao MQ, Xie XQ, Ren CE, Makaryan T, Anasori B, Wang GX, Gogotsi Y (2017) Hollow MXene spheres, and 3D macroporous MXene frameworks for Na-ion storage. Adv Mater 29:1702410 33. Zhao MQ, Ren CE, Ling Z, Lukatskaya MR, Zhang CF, Van Aken KL, Barsoum MW, Gogotsi Y (2015) Flexible MXene/carbon nanotube composite paper with high volumetric capacitance. Adv Mater 27:339–345 34. Zhao MQ, Torelli M, Ren CE, Ghidiu M, Ling Z, Anasori B, Barsoum MW, Gogotsi Y (2016) 2D titanium carbide, and transition metal oxides hybrid electrodes for Li-ion storage. Nano Energy 30:603–613 35. Ren CE, Zhao MQ, Makaryan T, Halim J, Boota M, Kota S, Anasori B, Barsoum MW, Gogotsi Y (2016) Porous two-dimensional transition metal carbide (MXene) flakes for highperformance Li-ion storage. ChemElectroChem 3:689–693 36. Luo JM, Zheng JH, Nai JW et al (2019) Atomic sulfur covalently engineered interlayers of Ti3C2 MXene for ultra-fast sodium-ion storage by enhanced pseudocapacitance. Adv Funct Mater 29:1808107 37. Yu H, Wang YH, Jing Y, Ma JM, Du CF, Yan QY (2019) Surface modified MXene-based nanocomposites for electrochemical energy conversion and storage. Small 15:1901503

476

J. Luo

38. Xiong DB, Li XF, Bai ZM, Lu SG (2018) Recent advances in layered Ti3C2Tx MXene for electrochemical energy storage. Small 14:1703419 39. Alhabeb M, Maleski K, Anasori B, Lelyukh P, Clark L, Sin S, Gogotsi Y (2017) Guidelines for synthesis and processing of two-dimensional titanium carbide (Ti3C2TX MXene). Chem Mater 29:7633–7644 40. Tang X, Guo X, Wu WJ, Wang GX (2018) 2D metal carbides, and nitrides (MXenes) as highperformance electrode materials for lithium-based batteries. Adv Energy Mater 8:1801897 41. Fang R, Lu C, Chen A et al (2019) 2D MXene-based energy storage materials: interfacial structure design and functionalization. ChemSusChem 12:1–12 42. Yu YX (2016) Prediction of mobility, enhanced storage capacity, and volume change during sodiation on interlayer-expanded functionalized Ti3C2 MXene anode materials for sodium-ion batteries. J Phys Chem C 120:5288–5296 43. Luo JM, Zhang WK, Yuan HD et al (2017) Pillared structure design of MXene with ultralarge interlayer spacing for high-performance lithium-ion capacitors. ACS Nano 11:2459–2469 44. Luo J, Tao X, Zhang J, Xia Y, Huang H, Zhang L, Gan Y, Liang C, Zhang W (2016) Sn4+ ion decorated highly conductive Ti3C2 MXene: promising lithium-ion anodes with enhanced volumetric capacity and cyclic performance. ACS Nano 10:2491–2499 45. Ghidiu M, Lukatskaya MR, Zhao MQ, Gogotsi Y, Barsoum MW (2014) Conductive two-dimensional titanium carbide ‘clay’ with high volumetric capacitance. Nature 516:78–81 46. Li TF, Yao LL, Liu QL et al (2018) Fluorine-free synthesis of high-purity Ti3C2Tx (T¼OH, O) via alkali treatment. Angew Chem Int Ed Engl 57:6115–6119 47. Xuan JN, Wang ZQ, Chen YY et al (2016) Organic-base-driven intercalation and delamination for the production of functionalized titanium carbide nanosheets with superior photothermal therapeutic performance. Angew Chem Int Ed Engl 55:14569–14574 48. Xu C, Wang LB, Liu ZB, Chen L, Guo JK, Kang N, Ma XL, Cheng HM, Ren WC (2015) Large-area high-quality 2D ultrathin Mo2C superconducting crystals. Nat Mater 14:1135–1141 49. Mashtalir O, Naguib M, Mochalin VN, Dall’Agnese Y, Heon M, Barsoum MW, Gogotsi Y (2013) Intercalation and delamination of layered carbides and carbonitrides. Nat Commun 4:1716 50. Mashtalir O, Lukatskaya MR, Zhao MQ, Barsoum MW, Gogotsi Y (2015) Amine-assisted delamination of Nb2C MXene for Li-ion energy storage devices. Adv Mater 27:3501–3506 51. Naguib M, Unocic RR, Armstrong BL, Nanda J (2015) Large-scale delamination of multilayers transition metal carbides and carbonitrides “MXenes”. Dalton Trans 44:9353–9358 52. Lipatov A, Alhabeb M, Lukatskaya MR, Boson A, Gogotsi Y, Sinitskii A (2016) Effect of synthesis on quality, electronic properties and environmental stability of individual monolayer Ti3C2 MXene flakes. Adv Electron Mater 2:1600255 53. Karlsson LH, Birch J, Halim J, Barsoum MW, Persson POA (2015) Atomically resolved structural and chemical investigation of single MXene sheets. Nano Lett 15:4955–4960 54. Lukatskaya MR, Mashtalir O, Ren CE et al (2013) Cation intercalation and high volumetric capacitance of two-dimensional titanium carbide. Science 341:1502–1505 55. Ghidiu M, Kota S, Halim J, Sherwood AW, Nedfors N, Rosen J, Mochalin VN, Barsoum MW (2017) Alkylammonium cation intercalation into Ti3C2 (MXene): effects on properties and ion-exchange capacity estimation. Chem Mater 29:1099–1106 56. Overbury SH, Kolesnikov AI, Brown GM, Zhang ZY, Nair GS, Sacci RL, Lotfi R, van Duin ACT, Naguib M (2018) Complexity of intercalation in MXenes: destabilization of urea by two-dimensional titanium carbide. J Am Chem Soc 140:10305–10314 57. Wang CD, Xie H, Chen SM et al (2018) Atomic cobalt covalently engineered interlayers for superior lithium-ion storage. Adv Mater 30:1802525 58. Zhu MS, Huang Y, Deng QH et al (2016) Highly flexible, freestanding supercapacitor electrode with enhanced performance obtained by hybridizing polypyrrole chains with MXene. Adv Energy Mater 6:1600969

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

477

59. Ling Z, Ren CE, Zhao MQ, Yang J, Giammarco JM, Qiu JS, Barsoum MW, Gogotsi Y (2014) Flexible and conductive MXene films and nanocomposites with high capacitance. Proc Natl Acad Sci U S A 111:16676–16681 60. Boota M, Anasori B, Voigt C, Zhao MQ, Barsoum MW, Gogotsi Y (2016) Pseudocapacitive electrodes produced by oxidant-free polymerization of pyrrole between the layers of 2D titanium carbide (MXene). Adv Mater 28:1517–1522 61. Wen YY, Rufford TE, Chen XZ, Li N, Lyu MQ, Dai LM, Wang LZ (2017) Nitrogen-doped Ti3C2Tx MXene electrodes for high-performance supercapacitors. Nano Energy 38:368–376 62. Tian WQ, VahidMohammadi A, Wang Z, Ouyang LQ, Beidaghi M, Hamedi MM (2019) Layer-by-layer self-assembly of pillared two-dimensional multilayers. Nat Commun 10:2558 63. Yan J, Ren CE, Maleski K, Hatter CB, Anasori B, Urbankowski P, Sarycheva A, Gogotsi Y (2017) Flexible MXene/graphene films for ultrafast supercapacitors with outstanding volumetric capacitance. Adv Funct Mater 27:1701264 64. Yu LY, Hu LF, Anasori B, Liu YT, Zhu QZ, Zhang P, Gogotsi Y, Xu B (2018) MXene-bonded activated carbon as a flexible electrode for high-performance supercapacitors. ACS Energy Lett 3:1597–1603 65. Sun N, Zhu Q, Anasori B, Zhang P, Liu H, Gogotsi Y, Xu B (2019) MXene-bonded flexible hard carbon film as anode for stable Na/K-ion storage. Adv Funct Mater 29:1906282 66. Xie Y, Naguib M, Mochalin VN et al (2014) Role of surface structure on Li-ion energy storage capacity of two-dimensional transition-metal carbides. J Am Chem Soc 136:6385–6394 67. Wang XF, Shen X, Gao YR, Wang ZX, Yu RC, Chen LQ (2015) Atomic-scale recognition of surface structure and intercalation mechanism of Ti3C2X. J Am Chem Soc 137:2715–2721 68. Simon P, Gogotsi Y, Dunn B (2014) Where do batteries end and supercapacitors begin? Science 343:1210–1211 69. Lukatskaya MR, Kota S, Lin ZF et al (2017) Ultra-high-rate pseudocapacitive energy storage in two-dimensional transition metal carbides. Nat Energy 2:17105 70. Xia Y, Mathis TS, Zhao MQ, Anasori B, Dang A, Zhou ZH, Cho H, Gogotsi Y, Yang S (2018) Thickness-independent capacitance of vertically aligned liquid-crystalline MXenes. Nature 557:409–412 71. Naguib M, Come J, Dyatkin B, Presser V, Taberna P-L, Simon P, Barsoum MW, Gogotsi Y (2012) MXene: a promising transition metal carbide anode for lithium-ion batteries. Electrochem Commun 16:61–64 72. Ma K, Jiang H, Hu YJ, Li CZ (2018) 2D nanospace confined synthesis of pseudocapacitancedominated MoS2-in-Ti3C2 superstructure for ultrafast and stable Li/Na-ion batteries. Adv Funct Mater 28:1804306 73. Zhang CF, Park SN, Seral-Ascaso A, Barwich S, McEvoy N, Boland CS, Coleman JN, Gogotsi Y, Nicolosi V (2019) High capacity silicon anodes enabled by MXene viscous aqueous ink. Nat Commun 10:849 74. Meng RJ, Huang JM, Feng YT et al (2018) Black phosphorus quantum dot/Ti3C2 MXene nanosheet composites for efficient electrochemical lithium/sodium-ion storage. Adv Energy Mater 8:1801514 75. Li WH, Yang ZZ, Li MS, Jiang Y, Wei X, Zhong XW, Gu L, Yu Y (2016) Amorphous red phosphorus embedded in highly ordered mesoporous carbon with superior lithium and sodium storage capacity. Nano Lett 16:1546–1553 76. Chen Z, Augustyn V, Jia XL, Xiao QF, Dunn B, Lu YF (2012) High-performance sodium-ion pseudocapacitors based on hierarchically porous nanowire composites. ACS Nano 6:4319–4327 77. Okubo M, Sugahara A, Kajiyama S, Yamada A (2018) MXene as a charge storage host. Acc Chem Res 51:591–599

478

J. Luo

78. Luo JM, Fang C, Jin CB et al (2018) Tunable pseudocapacitance storage of MXene by cation pillaring for high-performance sodium-ion capacitors. J Mater Chem A 6:7794–7806 79. Xie XQ, Zhao MQ, Anasori B et al (2016) Porous heterostructured MXene/carbon nanotube composite paper with high volumetric capacity for sodium-based energy storage devices. Nano Energy 26:513–523 80. Luo JM, Wang CL, Wang H, Hu XF, Matios E, Lu X, Zhang WK, Tao XY, Li WY (2019) Pillared MXene with ultralarge interlayer spacing as a stable matrix for high-performance sodium metal anodes. Adv Funct Mater 29:1805946 81. Xu M, Lei SL, Qi J, Dou QY, Liu LY, Lu YL, Huang Q, Shi SQ, Yan XB (2018) Opening magnesium storage capability of two-dimensional MXene by intercalation of cationic surfactant. ACS Nano 12:3733–3740

Jianmin Luo received his PhD in 2018 from the Zhejiang University of Technology in China. He is currently serving as a postdoctoral associate in the Thayer School of Engineering at Dartmouth College. His research is mainly focused on the design of 2D nanomaterials for energy storage and conversion. Dr. Luo has published more than 30 papers in journals (such as Nano Letters, ACS Nano, Advanced Functional Materials, Chemical Society Review) with a total citation of over 2000 and an H-index of 20. He also peer reviews Journals including ACS Nano, Advanced Functional Materials and Small, etc. Besides, Dr. Luo holds two issued Chinese patents as well.

Chapter 17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries Rajesh Pathak, Yue Zhou, and Qiquan Qiao

17.1

Introduction

17.1.1 Overview Energy is a key to drive today’s world, from consumer applications such as portable smartphones and laptops to industrial applications such as transportation, large-scale energy storage micro-grids, and medical equipment [1–3]. To fulfill the everincreasing energy demand, fossil fuels are being widely consumed. However, the use of fossil fuels is responsible for global climate change, fluctuation in the price, and concern about the availability of limited fossil fuels shortly. To address these concerns, widespread renewable energy sources such as solar, wind, and hydropower need to be deployed. For the continuous supply of energy, the energy storage device is aiding the integration of renewable energy into the electricity system [4– 8]. As a result, in the upcoming decades, fossil fuels can become history, and electric energy can drive the whole world.

R. Pathak Department of Electrical Engineering and Computer Science, South Dakota State University, Brookings, SD, USA Applied Materials Division, Argonne National Laboratory, Lemont, IL, USA e-mail: [email protected] Y. Zhou Department of Electrical Engineering and Computer Science, South Dakota State University, Brookings, SD, USA e-mail: [email protected] Q. Qiao (*) Mechanical and Aerospace Engineering, Syracuse University, Syracuse, NY, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_17

479

480

R. Pathak et al.

The lithium (Li)-ion batteries (LIBs) are considered a reliable energy storage device due to their high energy density, power density, and long cycle life. LIBs are widely used in a variety of applications such as cellphones, laptops, electric vehicles, and hybrid electric vehicles. The LIBs are gradually reaching the upper limit of specific energy density (~ 300 Wh kg
1). For the practical applications of high energy battery applications such as long drive range electric vehicles with a single charge, a high specific energy density of more than 500 Wh kg
1 is needed. This requires aggressive electrochemical chemistries between the conversion reaction or high-voltage/high-capacity intercalation cathode and Li as an anode. Li-metal is an ideal anode material for rechargeable batteries owing to its high theoretical specific capacity of 3860 mAh g
1, low density of 0.53 g cm-3, and extremely low redox potential of
3.04 V vs. standard hydrogen electrode [9–12]. Thus, it is imperative to replace the conventional graphite anode with a lithium metal anode (LMA). Figure 17.1a shows that the use of lithium metal as an anode can provide high specific capacity and high operational voltage, leading to the high energy density of the batteries.

17.1.2 Challenges in Lithium Metal Batteries The direct use of Li metal as a node has shown a great challenge owing to the highly reactive nature, the infinite volume change of Li, and the formation of undesired lithium dendrite growth and unstable fragile/unstable SEI. Firstly, the hyperactive lithium spontaneously reacts with electrolyte, active material from the cathode, and other side reactions. The irreversible reaction consumes electrolyte, lithium, and active materials leading to electrolyte dry-out and quick capacity fading [8, 13– 18]. Secondly, the infinite volume expansion of Li leads to the pulverization of the anode, severe lithium corrosion, and formation of dead lithium. The loss of active materials, undesired side products, and the formation of inactive dead lithium increase the impedance of the cell. Thirdly, the nonuniform electric field due to the rough lithium surface induces the hot spots for Li nucleation which continues to grow Li dendrite with increased cycling. The needle or whisker-shaped dendrites can pierce the separator causing the short circuit or can catch fire. Also, the high surface area of Li dendrite growth and continuous formation/deformation of SEI consumes much electrolyte and lithium causing the electrolyte dry-out and low CE. Numerous efforts have been done for dense, reversible, and dendrite-free Li deposition. Figure 17.1b shows the volume expansion issues associated with the lithium metal anode during plating/stripping cycles. The volume expansion issues could lead to further corrosion, cracks on the electrode responsible for further Li dendrite growth. Figure 17.1c demonstrates the formation of Li dendrites and dead Li during plating/ stripping cycles using Li metal as the anode. Thus, it is essential to protect the LMA to buffer the volume expansion and suppress Li dendrite growth in LMBs.

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

481

Fig. 17.1 Importance and challenges in using lithium metal as the anode. (a) Capacity and voltage comparison of typical electrodes. The inset shows the specific energy density comparison between LiBs and LMBs. (Reproduced from Ref. [19] with permission from John Wiley and Sons). (b) The pulverization of LMA due to infinite volume change in a typical Li stripping/plating cycle. (Reproduced from Ref. [20] with permission from National Academy of Sciences). (c) The formation of dendritic, mossy, and dead Li during the plating/stripping cycle. (Reproduced from Ref. [13] with permission from Elsevier)

17.1.3 Strategies to Revive LMA Substantial efforts have been done by many researchers as demonstrated in Fig. 17.2. to address the inherent issues related to the LMA to achieve their potential applications. For guiding the uniform Li deposition and improved cycling performance, the common approaches are (1) the development of lithiophilic and conductive threedimensional (3D) micro-/nanostructured framework to guide the uniform Li deposition and accommodate the volume expansion, (2) the engineering of an artificial protective layer on top of Li by ex situ/in situ which can greatly inhibit the side

482

R. Pathak et al.

Fig. 17.2 Strategies to address the inherent problem of LMA

reactions, (3) the replacement of inflammable liquid electrolyte with higher Young’s modulus solid-state electrolyte, and (4) other strategies such as modification of separator and defining the battery testing conditions.

17.1.3.1

The Development of 3D Micro-/Nanostructured Li host

The engineering of 3D porous metal-based or carbon-based Li hosts has proven to be a successful technique to accommodate the Li for dendrite-free Li deposition. Figure 17.3a, b shows the schematic representation of the Li-ion flux distribution and the Li plating on the planar and freestanding copper nanowire (CuNW). The uniform distribution of the electric field due to the nanowires leads to the uniform Li-ion flux and homogeneous Li deposition [21]. Besides, the high surface area of the interconnected CuNW framework significantly lowers the effective current density. Moreover, the CuNW accommodates the Li inside the porous nanostructure to impede the Li dendrite growth. In contrast, the planar Li has an uneven hot spot which creates a nonuniform electric field and results in the nonuniform Li-ion flux distribution and inhomogeneous Li deposition. The commonly used metallic or carbon-based porous frameworks show a poor affinity for Li metal. Thus, the poor Li wettability/lithiophobic porous framework suggests high nucleation

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

483

Fig. 17.3 3D nano-/microstructured Li host. (a, b) Schematic representation of the Li-ion flux distribution and Li-plating behavior on planar and CuNW framework. (Reproduced from Ref. [21] with permission from American Chemical Society). (c) Schematic illustration of Li plating on VA-CuO NSs-Cu. (Reproduced from Ref. [22] with permission from John Wiley and Sons). (d) Schematic illustration of Li plating on CNF-Cu. (Reproduced from Ref. [11] with permission from John Wiley and Sons). (e) Schematic illustration of Li plating on ultrafine AgNP decorated CNF. (Reproduced from Ref. [26] with permission from John Wiley and Sons). (f) Schematic illustration of fabricating Li/c-MOF composite. (g, h) Voltage-capacity profile of Li deposition on C-MOF at 10 μA cm
2 with and without lithiophilic Zn clusters. (Reproduced from Ref. [27] with permission from John Wiley and Sons). (i) Schematic representation of fabricating Li/porous carbon composite. (Reproduced from Ref. [20] with permission from National Academy of Science)

484

R. Pathak et al.

overpotential, voltage hysteresis, and unfavorable Li deposition. As a result, the Li plating/stripping cycles show to lower the energy barrier, provide the nucleation seed for Li growth, and guide the uniform Li deposition in a porous framework; the use of lithiophilic coating or decoration with heterogeneous seed growth is essential. This significantly reduces the nucleation overpotential and voltage hysteresis. To further enhance the Li deposition, the thin layer of vertically aligned (VA) lithiophilic CuO nanosheets (NSs) was grown on the Cu [22]. The lithiophilic CuO-NSs grown on Cu as shown in the schematic Figure 17.3c assure the fast Li-ion diffusion, and the Cu current collector guarantees the fast electron transport. The high Li-ion affinity of lithiophilic CuO promotes steady Li nucleation and deposition. During Li plating/stripping, the reversible conversion reaction facilitates fast Li storage. The reversible conversion (Cu2O, Cu3N, Ni3N) and alloy (ZnO, MgO, Si3N4, Sn, Au, Si, Ag) forming lithiophilic materials have successfully reported enhanced dendrite-free Li deposition, improving the battery cycling performance. Despite the attractive features of the metal-based current collector, the high cost, high mass density (Cu, Ni, etc.), and the complex synthesis process of 3D structure hinder their further practical applications. In contrast, the carbon is naturally abundant, has a lower cost and lower density, is easier for developing a porous framework, and thus is considered as a promising candidate for Li host. The emerging porous carbons such as hollow carbon spheres, carbon nanofibers (CNFs), graphitized CNFs, carbon nanotubes (CNTs), and carbonized MOFs have been demonstrated feasible as Li host owing to their good conductivity, large surface area, flexibility, and sturdiness [23, 24]. The high surface area or voids provide sufficient space for Li deposition, and the flexibility can buffer the volume expansion of Li. Chen et al. designed the vertically aligned CNFs on Cu foil for uniform and dendrite-free Li plating as shown in Figure 17.3d. The conductivity difference between the electron-conducting current collector and less conducting carbonaceous interlayer facilitates the deposition of Li underneath the interlayer [11, 12]. The interlayer should have higher electrolyte wettability, lithiophilicity, and Li-ion diffusivity. The semiconducting interlayer does not favor the Li deposition on top of it as a separate film during initial Li plating. Instead, Li-ion diffuses and starts depositing from the bottom. Recently, mixed ionic conductors/interlayers which allow fast ionic and electronic transport have also attracted great attention in regulating the dendrite-free Li deposition. To avoid the cost and lower the density of the batteries, free-standing porous carbonaceous materials with enhanced conductivity are supposed to be efficient Li hosts. Moreover, the lithiophobic carbon cannot uniformly guide the initial Li nucleation and leads to higher nucleation and voltage overpotential. Thus, the lithiophilic decoration on such porous carbon materials is one of the promising strategies [25]. Yang et al. reported the silver nanoparticles (AgNPs) anchored CNF by treating the silver acetate-CNF by electric joule heating [26]. As lithiophilic heterogeneous silver nanoparticles (AgNPs) have considerable solubility in Li, the Li deposition on Ag-anchored CNF shows zero nucleation overpotential and significantly lowers the overpotential. Consequently, the uniform Li deposition was

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

485

achieved with ultrafine AgNP decorated CNF as shown in Figure 17.3e. However, the in situ electrochemical platings of Li in a porous framework may not have uniform spatial Li deposition control and are complex to disassemble and clean before using the Li/porous structured composite while assembling the full cell battery. The direct infusion of molten Li in such a porous framework for making composites is considered a considered suitable approach. Zhu et al. reported the infusion of molten Li into the porous, and lithiophilic Zn cluster confined carbonized metal organic framework c-MOF [27]. Figure 17.3f shows the illustration of fabricating Li/c-MOF composite. The pre-planted uniformly dispersed Zn as nucleation seeds guides the uniform Li deposition due to the thermodynamic matching with Li. Zn has appreciable solubility with Li and also forms Li-Zn alloy. As a result, the nucleation barrier in c-MOF with the Zn cluster shows an eliminated nucleation barrier compared to the c-MOF without the Zn cluster as shown in Figure 17.3g, h. Liang et al. also reported the infusion of Li into the 3D conducting carbon scaffold coated with lithiophilic Si as shown in Figure 17.2i. The Li/C composite anode demonstrated smooth and uniform confinement of Li inside the porous matrix. Besides, the Li/C composite anode showed a flat and stable voltage plateau because the Li is plated or stripped from the Li itself, which minimizes the nucleation barrier initiating from the lattice mismatch. The lithiophilic decoration on a porous nano-/ microstructured framework not only ensures uniform lithium nucleation but also reduces the effective current density and the degree of interface fluctuations and confines the Li deposition/dendrites within the porous scaffold. Further development of the SEI layer on such a porous framework can be the next leap to push the CE closer to the real battery operation. However, the quality of SEI plays a critical role in the effective and stable operation of LMBs, as several requirements for an ideal SEI need to be satisfied [28]. To further avoid the risk of high temperature during molten infusion of Li into the porous free-standing electrode, the mechanical process of the press has been applied to make the Li/C composite anode. The development of an artificial SEI on such porous Li host/current collector greatly enhances the cycling stability.

17.1.3.2

The Development of Solid-State Electrolyte

The use of flammable liquid electrolyte in LMBs not only compromises the energy density but also stimulates continuous side reactions with LMA, resulting in low Coulombic efficiency (CE) [29]. Thus, it is highly recommended to utilize the solidstate electrolyte (SSE) for addressing the aforementioned issues of LMA and liquid electrolyte. The design of an ideal SSE should consider ionic conductivity, electro/ chemical and mechanical stability, being ecofriendly, and low cost. The SSEs with high Young’s modulus is found to repress the Li dendritic growth and obtain better Li utilization. Moreover, SSE can be integrated with low cost and high capacity cathodes such as sulfur and metal sulfides. The SSE can intrinsically mitigate the shuttle effect and address the metal dissolution of cathode materials. To date,

486

R. Pathak et al. 0

75Li2S 25P2S5 glass

Li7P3S11

Li9.54Si1.74P1.44S11.7 CL0.3

log [ s (S cm–1)]

–2

1M LiPF6 EC/DMC 1M LiTFSI EC/DMC

Li6PS5Cl PEO-LiTFSI-1%LGPS-10%SN E-Li3PS4

Li10GP2S12

80Li2S 20P2S12 glass

–4 PEO-LiTFSI

–6

LiPON

70Li2S 30P2S5 glass

Li7La3Zr2O12

LiTi2Al(PO4)3

–8 2.0

2.5

3.0

3.5

4.0 4.5 1000/T (K)

5.0

5.5

6.0

Fig. 17.4 The Arrhenius plot of various SSEs compared to the organic liquid electrolyte (Reproduced from Ref. [30] with permission from John Wiley and Sons)

different types of SSEs such as oxides, sulfides, Li-rich antiperovskites, thin films, and polymers have been developed and investigated. Polymer SSE is a promising candidate because of their facile fabrication process, ultrathin nature, flexibility, high level of safety, and electrochemical stability towards Li metal anode/cathode but suffers from oxidation and thermal stability. The sulfide-based SSE has high conductivity (~ 10
2 S cm
1) which is comparable to the organic liquid electrolyte as shown in Fig. 17.4 but suffers from chemical instability. Similarly, the oxide-based SSE shows higher stability toward anode or cathode but suffers from high interfacial incompatibilities and requires high processing costs and complex device integration. In SSEs, ions migrate utilizing ionic motion across the crystal lattice. Thus, they are expected to higher transference number (~1) compared to the organic liquid electrolyte (~0.4–0.5). To address the issues with SSE, scientists are continuously investigating the facile, low-cost, and scalable method.

17.1.3.3

The Development of Solid Electrolyte Interphase

The surface layer between the LMA and the organic liquid electrolyte plays an important role in determining the morphology of Li deposition. The morphology of Li deposition influences battery safety and performance. The SEI inhibits the direct physical contact between the LMA and the solvent of electrolyte. Beyond that, SEI can regulate the Li-ion distribution from the bulk electrolyte to the LMA. The SEI can be formed in two ways: (1) in situ and (2) ex situ. The in situ formation of SEI is during the battery operation, and the ex situ SEI can be generated before the battery operation.

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

487

The common solvent used in the nonaqueous liquid electrolyte of lithium batteries includes ester and ether. The ester or carbonate electrolyte has comparatively higher oxidation stability but suffers from limited compatibility with LMA. In contrast, the ether or glyme-based electrolyte has comparatively lower reactivity with LMA, providing higher CE and lower voltage hysteresis, but suffers limited oxidation stability. The common Li salts for the nonaqueous liquid electrolyte includes lithium hexafluorophosphate (LiPF6), bis(trifluoromethanesulfonyl)imide (LiTFSI), lithium bis(fluorosulfonyl)imide (LiFSI), etc. The deposition of insulating products on the LMA during the initial cycles due to the side reaction between the liquid electrolyte and hyperactive LMA which is thermodynamically unstable (Gibbs energy
8.670 kJ/mol) in organic solvents/electrolytes or from the non-faradic adsorption and the faradic electrochemical reaction is called SEI as shown in Figure 17.5a. The reduction of Li salt and solvent in the electrolyte (usually

Fig. 17.5 The schematic illustration of the mechanism in SEI. (a) The mechanism of SEI formation due to the direct contact of bulk Li electrode and liquid electrolyte. (Reproduced from Ref. [42] with permission from John Wiley and Sons). (b) The schematic demonstration of ideal SEI. The schematic demonstration on SEI formation in (c) the Peled model, (d) Mosaic model, and (e) Coulombic interaction model, respectively. (Reproduced from Ref. [43] with permission from John Wiley and Sons)

488

R. Pathak et al.

below 1.0 V vs. Li+/Li) and the hyperactive lithium decompose, leading to the formation of SEI. This SEI layer (30–50 nm) passivates the lithium anode, preventing the further corrosion or consumption of lithium with the electrolyte. Ideally, the SEI should possess (1) high chemical stability, (2) high mechanical stability, (3) sufficient Li-ion diffusion and electron flow prevention, (4) flexibility to buffer Li volume expansion, and (5) charge/discharge reversibility as shown in Figure 17.5b. There are different models to explain the SEI formation mechanism such as the Peled model, Mosaic model, and Coulombic interaction mechanism model. Peled model (Figure 17.5c) tells that the SEI is formed from the surface reaction by the reduction of electrolyte components [31]. Mosaic model (Figure 17.5d) tells that the insoluble and insulating multiphase discharge products generated during the side reaction between the electrolyte and LMA get deposited on negatively charge LMA and the deposited layer is called mosaic layer [32]. The mosaic layers with grains and grain boundaries offer Li-ion transport across it. Coulombic interaction model (Figure 17.5e) tells that the positively charged and partially positively charged Li-ions bring into the line as head and foot, respectively [33]. Thevenin and Muller brought up several updated models which include the polymer electrolyte layer model and compact stratified layer model [34, 35]. The naturally formed SEIs are very fragile, have low Li-ion conductivity and insufficient chemical and mechanical stability, and cannot completely passivate the LMA. Such SEI cannot buffer the volume expansion issues and has weak anchoring strength with the underlying bulk lithium electrode. The continuous formation and deformation of such unstable SEI consumes both liquid electrolyte and LMA and increases the impedance of the battery, leading to premature battery death. Besides, the SEI cannot suppress lithium dendrite growth which can penetrate the separator challenging the safety of the battery. To acquire SEI with nearly ideal properties, as mentioned above, the use of various additives, salts, solvents, and fillers and optimization of the electrolyte concentration have been reported [36–40]. For the reduction, these components are generally chosen in such a way that the lower unoccupied molecular orbital (LUMO) is lower than the electrolyte and its components. Although dendrite-free Li deposition and substantial improvements in the battery cycling have been achieved, the formation of in situ SEI consumes a superfluous amount of electrolyte and bulk Li electrode [41]. Moreover, there is no calculated control on the dimension, distribution, SEI components, and thickness of in situ formed SEI. The ex situ-based method is another way to develop the SEI with more powerful control on the dimension, distribution, SEI components, and thickness. This chapter provides an overview of the efficient ex situ-derived SEI for the protection of LMA. In this chapter, the results of the studies on ex situ SEI that have been carried out to date are pulled out and summarized. Specifically, this chapter discusses the engineering process and efficient working mechanism of the ex situ artificial SEI, which can overcome the challenges associated with in situ-derived SEI. The development of ex situ SEI with sufficient Li-ion conductivity, large flexibility, high adhesion, or anchorage affinity with Li metal surface and high mechanical and chemical stability can be a promising technique. These approaches can potentially lead to much higher

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

489

efficient SEI, based on the simultaneous resolution of the above challenges hindering the battery performance. This study will allow estimating the hands-on metrics and ideal materials, in an attempt to revive the LMBs for dendrite-free Li deposition, improved cycling performance, and high energy density batteries. Finally, the perspectives and critical challenges needed to address the issues in ex situ-formed SEI for further improvement of stable and robust SEI are concluded.

17.1.4 Ex Situ-Based SEI The adding of an artificial layer on top of LMA to protect it from liquid electrolyte before the battery assembling or battery operation is considered as the process of ex situ-based SEI formation. As this is an external process, the electrolyte and plated lithium will not be consumed as in the in situ SEI formation process. Besides, there is an effective control on the thickness, dimension, and distribution, which play a significant role in the battery performance. The strong adhesion on the bulk Li electrode prevents the peel off or corrosion of SEI during Li-plating/Li-stripping cycles. The nonuniform electric field distribution in the bulk Li electrode, due to the rough, crack surface, leads to the nonuniform Li-ion flux. The nonuniform Li-ion flux deposits the Li nonuniformly, which continues to accumulate with longer charge/discharge cycles, acting as a hot spot for Li nucleation. In general, there are two methods to generate ex situ-based SEI, which include (1) physical deposition method and (2) chemical deposition method. The solution, gas, or any other separate solid layer can be utilized to develop the ex situ-based SEI. The smooth surface of SEI leads to the uniform Li-ion flux, improving the Li deposition behavior for longer Li-plating/Li-stripping cycles. Despite the attractive performance of the battery and controlled Li dendrite growth, many of the reports are concentrated on the single aspects of SEI. For example, the reports have shown that improved performance can be attributed to only high Young’s modulus or high Li-ion conductivity. It is imperative to think and consider all the ideal properties of SEI that can significantly contribute to the improved cycling performance of the battery. Besides, the material manufacturing process, cost, safety, and effectiveness are to be deemed.

17.1.4.1

Physical Deposition

The construction of SEI on top of LMA physically adding a layer is considered as a physical deposition. The physical interlayer can be from a few nanometer thicknesses to several micrometers. Depending upon the thickness, the adhesiveness of the interlayer is very important. For example, for the ultrathin film, high adhesivity is required. Otherwise, there are chances of SEI being peeled off and degraded. The separate thick (μm-range) interlayer has also been used to create a physical barrier between the LMA and liquid electrolyte. Such kind of thick interlayer does not need to bear high adhesivity. For example, glass fiber cloth [44], 3D conductive stainless

490

R. Pathak et al.

steel fibrous metal felt [45], the monolayer of nanostructured/interconnected hollow carbon nanospheres [46], and carbon paper interlayer [47] have been employed as an interlayer to protect the LMA. However, the use of an additional thick extra layer on top of LMA compromises the energy density of the battery. Various methods have been employed to develop the ex situ-based SEI by a physical deposition method such as spin coating, radio-frequency (R-F) sputtering, direct-current (DC) sputtering thermal evaporation, chemical vapor deposition (CVD), pulsed laser deposition (PLD), atomic layer deposition (ALD), drop-casting, or doctor blading techniques. While bridging the academic research to industry, the cost is an important factor. The ultrahigh vacuum needing instruments are expensive and inefficient for large-scale production. The facile and low cost, a solution-based method such as dip coating, drop a cast, and doctor blading can be more practical for large scale productions. In our previous work, the ex situ SEI was developed by depositing the ultrathin bilayer of graphite and SiO2 on the lithium metal [12]. Figure 17.6a shows the Torr combination system for R-F sputtering. Figure 17.6b, c shows the topography SEM images of Li before and after the physical deposition of graphite/SiO2 bilayer on the surface of Li. The ultrathin bilayer of graphite/SiO2 demonstrated outstanding Li-plating/Li-stripping cycles and full cell performance. The graphite provides flexibility and electrically connects the plated Li with the bulk Li underneath the SEI, which reduces the chances of dead Li formation. The high Li-ion conductivity, high mechanical strength, and reversible lithiation/delithiation properties of SiO2 facilitate fast Li-ion transport, suppress the Li dendrite growth, and store Li or slow down the Li dendrite growth, respectively. The R-F deposition of an artificial SEI layer produces a smooth and high-quality film, enabling uniform Li-ion flux on the surface of the Li anode. The rate of deposition and the carrier gas or reactive gas can be easily tuned as per the requirement. Lee et al. used the doctor blading techniques to develop the composite protective layer (CPL) on the LMA as shown in Figure 17.6d [48]. The coated CPL with a thickness of 25 μm using the slurry of Al2O3, PVDF-HFP, DMF, and carbonate electrolyte showed improved long-term battery performance compared to the bare Li. The CPL provides a physical barrier between the Li and liquid electrolyte, enables uniform Li-ion flux and sufficient Li-ion transport, and mechanically suppresses the Li dendrite growth. Figure 17.6e, f shows the surface topography and cross-sectional SEM images of CPL protected Li metal. Table 17.1 summarizes the materials and methods used for the SEI development and the corresponding electrochemical cell test performance achieved with that SEI.

17.1.4.2

Chemical Deposition

The chemical reaction between the hyperactive LMA and the material for SEI formation leads to the strong anchorage between the formed SEI and the LMA. The high anchorage affinity or adhesivity of SEI components leads to the formation

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

491

Fig. 17.6 Schematic of the physical deposition of SEI. (a) Schematic of the R-F sputtering to deposit the target materials on the Li metal surface in the Ar-environment. (b–c) Surface topography of Li before and after the physical deposition of bilayer graphite-SiO2 on the Li metal surface. (Reproduced from Ref. [12] with permission from John Wiley and Sons). (d) Schematic illustration of the CPL preparation by doctor blading method. (e, f) The surface topography and cross-sectional SEM image of the CPL deposited Li metal anode. (Reproduced from Ref. [48] with permission Elsevier)

of dense SEI. The thickness of such chemically formed SEI also varies from a few nanometers to several micrometers. The chemical reaction can be allowed to happen between the Li metal and precursors such as a solution-based, gas-based, or solid-based. The solution-based

492

R. Pathak et al.

Table 17.1 Summary of the artificial SEI engineered by a surface coating method SEI components The ultrathin bilayer of graphite and SiO2 Li3PO4 Carbon Cu3N Al2O3 MoS2 Al2O3 Al2O3 Double-layer nanodiamond Al2O3 + PVDF-HFP Cu3N + SBR ETPTA+DADMA-TFSI PEDOT-co-PEG Perfluorinated ionomer layer Langmuir-Blodgett Hollow carbon nanosphere Double-layer nanodiamond Trifluoroethanol (TFEA)-modified eggshell membrane (TESM) Metal chloride perovskite film MASnCl3 and MAPbCl3

Method R-F sputtering (40 nm)

Ref. [12]

Magnetron sputtering (30 nm) Magnetron sputtering (80 nm) R-F magnetron sputtering (115 nm) R-F magnetron sputtering (20 nm) R-F magnetron sputtering (20 nm) ALD (14 nm) ALD (~2–3 nm) MPCVD (300 nm) Doctor blade (25 μm) Doctor blade/drop-casting (1 μm) Doctor blade/UV irradiation (1 μm) Immersing LMA into the polymer and nitromethane solution (10 μm) Solvent-casting (20 nm – 9 μm) Roll press (0–3 μm) Pressing (wall thickness – 20 nm) MPCVD (300 nm) Interfacial layer (90 μm)

[49] [50] [51] [52] [53] [54] [55] [56] [48, 57] [58] [59] [60]

Spin coat-solid state transfer-pressing (1 μm and 2 μm)

[61] [62] [46] [56] [63] [64]

reaction includes doctor blading, spin coating, dripping, drop-casting, dip coating, etc. The various parameters such as exposure time, temperature, the concentration of the solution, and other environmental conditions determine the properties (thickness, SEI components, morphology, etc.) of SEI. In our previous work related to the solution-based chemical deposition, we treated the Li metal with the electrolyte containing SnF2 [11]. The bare Li shows the dendritic Li growth, but 3 wt% SnF2pretreated Li shows the dendrite-free Li deposition as illustrated in Figure 17.7a, b. The smooth surface with an artificial SEI thickness of 25 μm (Figure 17.7c) showed outstanding battery performance. Also, in our previous work related to the gas-based chemical deposition, we treated the Li metal with the N2-assisted plasma [13]. Figure 17.7d shows the experimental setup for N2 plasma, and Figure 17.7e shows the cross-sectional SEM image of Li anode with 2 min N2 plasma. The optimized thickness of 8 μm achieved with 2-min N2 plasma showed excellent electrochemical battery performance (Table 17.2).

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

493

Fig. 17.7 Chemical surface treatment of LMA. (a, b) Schematic of the dendritic Li growth and smooth Li deposition in bare Li and SnF2-pretreated Li, respectively. (c) The surface topography SEM image of Li and inset shows its cross-sectional images. (Reproduced from Ref. [11] with permission from Nature Springer). (d) Experimental setup for the developing Li3N on Li by the plasma process. (e) Corresponding cross-sectional images after 2 min of N2 plasma. (Reproduced from Ref. [13] with permission from Elsevier)

17.1.5 SEI Properties and Functionality The conventional approaches were focused only on one of the properties of SEI which cannot fully stabilize the SEI. Consequently, long-term battery performance and fast charge/discharge rates couldn’t be possible. Thus, the advanced strategies of developing SEI need to focus on all the basic requirements to achieve an ideal SEI. Figure 17.3b shows the properties of an ideal SEI which includes high Li-ion conductivity, high mechanical and chemical stability, flexibility, and reversibility. Here the representative properties of the SEI are discussed.

17.1.5.1

Thickness

The various parameters such as deposition time, the concentration of the precursor, and the fabrication method play an important role in determining the morphology and thickness of SEI. The thickness of SEI can be optimized by taking different parameters such as Li-ion diffusivity and mechanical and chemical stability into considerations. For thickness optimization, various structural and corresponding chemical characterizations can be carried out. Structural characterizations such as XRD, Raman, SEM, BET, and XPS and the electrochemical cell performance such as Li plating/stripping, Li deposition morphology, voltage profile analysis, and EIS measurement help to optimize the thickness of SEI. The higher thickness of SEI can suppress the growth of Li dendrites; however, it might sluggish the Li-ion transport. In contrast, the lower thickness of SEI may not fully protect the LMA. The excess

494

R. Pathak et al.

Table 17.2 Summary of the artificial SEI engineered by surface chemistry method SEI components/thickness of SEI Li3PO4 (50 nm) Cu + LiF (2 μm) Organic/inorganic film (75 nm) Li-rich metal alloy/Li-halides (< 10 μm) Ge, GeOx, Li2CO3, LiOH, LiCl, and Li2O (1.5 μm) LiF nanoparticles (
) MIEC LLTO (8.2 μm) Al-Li alloy/LiCl (10 μm) Cuboid LiF (77.4 μm) LiF, Sn, Sn-Li alloy (25 μm) [LiNBH]n chains (140–160 nm) Li2S (25 nm) Li2S/Li2Se (2 μm) LiF layer (40 nm) LiF coating (380 nm) LiF thin film (40 nm) Li3N (8 μm) Li-O-Si linkage (
)

Materials/method to pretreat the LMA A solution of polyphosphoric acid (PPA) and dimethylsulfoxide (DMSO) Dripping (CuF2 + DME) solution on LMA Soaking the Li metal in the pure solvent of fluoroethylene carbonate (FEC) MClx solution in tetrahydrofuran (THF) for 20 sec

Ref. [65]

Immersing LMA organic GeCl4 – THF steam (1.5 μm)

[69]

Immersing Cu into aqueous solution of LiPF6 Toluene in between LLTO and LMA Soaking LMA in AlCl3-ionic liquid Immersing LMA into NH4HF2 + DMSO Drop-cast SnF2 containing carbonate electrolyte Dehydrogenation reaction between LMA and ammonia borane-NH3BH3 (AB) + THF (Sulfur/240  C) gas – LMA (SeS2/150  C) gas – LMA Nitrogen trifluoride/175  C gas – LMA Fluoropolymer/350 Cgas – LMA/175  C Freon R134a/150  C gas – LMA Nitrogen gas-enhanced plasma – LMA Immersing LMA into tetraethoxysilane for 5 min

[70] [71] [72] [73] [11] [74]

[66] [67] [68]

[75] [76] [77] [78] [79] [13] [80]

use of chemicals or materials to construct SEI compromises the energy density of LMBs. Thus, the materials with low mass density with a high Li-ion diffusion coefficient are to be chosen. The SEI thickness ranging from ultrathin few nanometers to thick micrometer size has shown improvement in the electrochemical battery performance.

17.1.5.2

Transference Number

The high Li-ion transference number (TLi+) increases the energy density and charge rates in LMBs. Recently, scientists are motivated to develop the SEI through the surface coating of polymer which uniquely is cation conducting. The low Li-ion conductivity and low TLi+ of SEI limit the practical application of LMBs at high current density. The enhancement of Li-ion conductivity and Li-ion TLi+ is challenging work. The use of cation conducting membrane as SEI can significantly reduce the loss caused by ion polarization. Tu et al. reported SEI deposited on the LMA surface by the solvent-casting method of perfluorinated ionomer [61]. The ionomer SEI with a thickness of 20 nm demonstrated ionic conductivity of

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

495

Fig. 17.8 (a, b) Comparison of Li-ion conductivity and transference number of different electrolyte systems from the previous reports. (Reproduced from Ref. [61] with permission from John Wiley and Sons). (c) Schematic of armored MIEC SEI developed by the displacement reaction between CuF2 and Li, showing the high modulus, ionic conductivity, and surface energy. (Reproduced from Ref. [66] with permission from John Wiley and Sons). (d) Comparison between bare Li foil and rAGA-Li anodes toward water stability. (Reproduced from Ref. [102] with permission from John Wiley and Sons). (e, f) Antioxidative characteristic of pSEI-Li and bare Li using contact angle measurement, surface tarnish test, and Raman spectroscopy. (Reproduced from Ref. [59] with permission from John Wily and Sons)

~1*10
3 S cm
1 at room temperature and high TLi+ ~ 0.9. Figure 17.8a, b shows the summary and comparison of ionic conductivity and lithium transference number of various electrolytes reported previously. The gray color symbolizes the ionic conductivity measured above room temperature. Solid and open symbols indicate the solid-state and liquid electrolyte. For composite or hybrid electrolytes, half-filled symbols are used. Cho et al. demonstrated ultraviolet (UV)-polymerized ethoxylated trimethylolpropane triacrylate (ETPTA) and diallyldimethylammoniumbis (trifluoromethanesulfonyl)-imide (DADMA-TFSI) as an organic/inorganic mimic SEI [59]. Such SEI demonstrated a high ionic conductivity of 1.2*10
3 S cm
1 and TLi+ of 0.69, which can be attributed to the fast Li-ion transport provided by the inorganic layer. The lone pair electrons (such as C-N bond) and amine polar function

496

R. Pathak et al.

(such as amine group) in MOF successfully inhibit the anion migration which results in the improved TLi+ [81, 82]. Besides, the anion blocking or negatively charged group such as –SO3
, -COO
, and PO32
containing cation-selective ionomer membrane SEI lessens the chances of direct contact of solvent or anions of electrolyte with LMA. As a result, the TLi+ increases which prolonged the growth of dendritic Li. Moreover, the polar functional bonds on the surface of MOF effectively screen the ions and adsorb only Li-ion, hindering the migration of anion. This results in the improvement of TLi+ [81].

17.1.5.3

Young’s Modulus and Flexibility

The mechanical stability and flexibility are very important properties of SEI for stable long-term cycling performance in LMBs. The hostless LMA experiences repeated volumetric change which can cause serious mechanical cracks and repeated breakdown/repair on the SEI during Li plating/stripping. Such cracks behave as hot spots for Li deposition, which eventually cannot effectively guide the uniform and dendrite-free Li deposition. Therefore, the high mechanical strength and flexibility of SEI can suppress the Li dendrite growth and buffers the volume expansion issues for maintaining the interface stabilities and SEI structural integrities. Figure 17.6c shows the schematic illustration of an armored SEI (Cu and LiF) formed by the displacement reaction between CuF2 and Li. The Cu atoms in the SEI improve the Li-ion conductivity by offering more diffusion space. Besides, the high surface energy of such SEI realizes uniform Li-ion distribution, high ionic conductivity, and sufficient Young’s modulus to suppress the growing Li dendrites. Previous studies have suggested that a moderate elastic modulus above 3.0 GPa is sufficient enough for increasing the SEI failure time and suppressing the Li dendrite growth [65, 83, 84]. If Young’s modulus is high enough to improve or maintain the structural uniformity in SEI, the stability of SEI will be enhanced. Extensive research has highlighted the benefit of developing SEI with high mechanical strength, some of which are summarized in Table 17.3. The infinite volume expansion issues of Li can create cracks and repeated formation/deformation of SEI. However, the flexible SEI can accommodate the interfacial fluctuations and further prevent the SEI from corrosion and peeling off. Lie et al. designed 20 nm Li polyacrylic acid (LiPAA) as a self-adapting interface fluctuation SEI by drop-casting the solution containing LiPAA in DMSO [89]. The main organic functional group COOLi on the LiPAA anode not only provides flexibility to the surface but also protects the LMA from the air environment. The use of flexible SEI components such as styrene-butadiene rubber (SBR) [58], thinfilm carbon, or graphite coating on Li is also effective in addressing the infinite relative electrode dimension change during cycling [12, 90]. The mechanically robust metal-organic framework (MOF) SEI cemented by polymer to modify its flexibility is of great interest to develop stable SEI [81]. The use of 3D porous Li hosts such as layered reduced graphene oxide (rGO) with nanoscale interlayer gaps [91], hollow carbon nanospheres [46], and nanostructured unstacked graphene drum

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

497

Table 17.3 Ex situ SEI components with Young’s modulus value SEI components PPA-Li Double-ND Single crystal ND/ND-Li composite LiF + PVDF-HFP LiF + Sn + Li-Sn Graphite + SiO2 Li3N Cu + LiF A single atomic layer of h-BN Al2O3 Al2O3 + PVDF-HFP Nafion + LiCl Li2CO3 + LiF LiDFBOP + FEC derived SEI

Modulus (GPa) ~10.0–11.0 >200 ~170.0/30.0 6.7 55.6 10.7 48.0 12.9 ~1 Tpa 6.4–14.3 3.4 6.1 7.0 64.9

Average thickness ~ 50 nm ~ 300 nm 216 μm 12 μm 25 μm 40 nm 8 μm 2 μm 1 nm 20 nm 25 μm 2.5 μm 50 nm –

Ref. [65] [56] [85] [66] [11] [12] [13] [66] [86] [52] [48] [87] [67] [88]

[92] is another strategy to address the volume change problem of LMA. Besides, the use of ceramics materials is very promising interfacial layers to buffer the volumetric expansion of the anode [93–96]. The SEI components which offer the synergy of high mechanical strength and sufficient flexibility are proven to be effective in the stabilization of SEI [58, 97]. The dual-layered organic/inorganic film was also constructed by soaking the Li metal in the pure solvent of fluoroethylene carbonate (FEC) for different time durations [67]. The electrostatic attraction between the positively charged Li atom and negatively charged F atom can break the C-F bond of FEC to form LiF, and, subsequently, the organic components such as CH2CHOCO2Li (ROCO2Li) and CH2CHOLi (ROLi) also formed. The organic layer provides the flexibility, and the inorganic layer provides high mechanical strength to suppress the Li dendrite growth. The use of graphene oxide (GO) enables high mechanical strength and flexibility, which can tolerate the interfacial fluctuations [98].

17.1.5.4

Mixed Ionic/Electronic Conductor

The SEI with high Li-ion conductivity and poor electronic conductivity allows sufficient Li-ion diffusion but prevents the flow of electrons. The low electronic conductivity enables the Li deposition underneath the SEI (bottom to top) and suppresses the growth of Li dendrites. The lower electronic conductivity of the SEI compared to the current collector prevents the deposition of Li on top of SEI. Further, weak binding and the poor Li wettability of such SEI provide the space between the SEI and the current collector and promotes the Li deposition underneath the SEI protective layer, respectively. Nevertheless, the protective layer with high electronic conductivity which can facilitate fast diffusion of Li-ion on the surface

498

R. Pathak et al.

and within interlayer gaps has also been considered effective in stabilizing the SEI [91, 92]. The conductive protective framework is supposed to lower the large electrical impedance and strong voltage hysteresis at high current rates. Besides, the mixed Li-ion conductors can also facilitate the high Li-ion migration across the SEI. Li2S has been established as an efficient protective layer owing to its high conductivity (~10
5 S cm
1). Liu et al. designed mixed Li-ion conductors Li2S/ Li2Se where Li2Se shows higher Li-ionic conductivity than Li2S [76]. Recently, the mixed ionic/electronic conductors (MIEC) which enable rapid transport of both Li-ions and electrons have attracted great attention [71, 99]. The MIEC can store the Li at the grain boundary regions of ionic and electronic conductors, ionic conductor regulates the Li-ion, and electronic conductor transfers the electron rapidly. The electronic conductor may also store the Li by forming a Li-rich alloy. For example, in the hybrid interphase of C60 and Mg, Li and Mg form a Li-Mg alloy [100].

17.1.5.5

Antioxidative

The water penetration inside the battery during a rainy day or while washing a car can cause potential safety threats. When Li is in contact with moisture, it gets oxidized and evolves unwanted gas, deteriorating the battery performance and challenging the safety issues. The water repellent hydrophobic SEI layer can shield the LMA from such moisture exposure. Dong et al. engineered the LMA passivation SEI layers using a vertically aligned graphene framework with a hydrophobic roof of graphene tiles [101]. Figure 17.6d shows the Li metal without any protection (left side) and with reduced accordion-like graphene oxide array (rAGA) protection (right side). The reaction between Li and H2O can produce H2 gas and cause corrosion in the Li metal. An air- or waterproof graphene tile roof house like hierarchical architecture, combined with a vertically aligned graphene microstructure channel, substantially protects the LMA and stores excessive Li in its porous structure. Liao et al. demonstrated the treatment of LMA by using GeCl4 which can generate Ge, GeOx, Li2CO3, LiOH, LiCl, and Li2O on Li surface, allowing the stable symmetric and full cell Li-O2 cycling performance at a relative humidity of 45% [69]. There are numerous reports on the use of hydrophobic or water-resistive SEI to protect the LMA. The other report to protect the LMA from the moisture is to use organic components that are hydrophobic or water-resistive [59]. Figure 17.8e–g shows the antioxidative behavior of the printable solid electrolyte interphase (pSEI) deposited on Li under different humid conditions. The bare Li quickly oxidizes when in contact with the water inset of Fig. 17.8f. In contrast, the pSEI-Li has a higher contact angle of 65 indicating higher water-repelling capability. The surface of bare Li quickly tarnished into black upon exposure to a humid environment. In contrast, pSEI-Li maintained higher stability against the humid environment of ~50% relative humidity. The inset shows the photographic image of pSEI-Li and bare Li after exposure to humid air for 10 and 40 s. Figure 17.8g shows the Raman spectra of bare Li and pSEI-Li before and after 120 s exposure to humid air. The distinct and intense peak

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

499

of LiOH and Li2CO3 at 657 and 1060 cm
1 were observed for bare Li. In contrast, the peaks were very weak from the antioxidant pSEI-coated Li.

17.1.5.6

Hybrid SEI

The hybrid SEI which can simultaneously store Li by the reversible plating/stripping and redox reaction mechanism is considered effective in guiding the uniform Li deposition. The components with high Li-ion conductivity allow sufficient Li-ion diffusion, and the chemically active components react with the hyperactive dendritic Li and slow down the rate of dendrite growth during plating/stripping cycles. The SEI components of the hybrid SEI can reversibly store Li by lithiation/delithiation reaction mechanism [12, 53]. In our previous studies, the bilayer of 20 nm graphite and 20 nm SiO2 was deposited on Li foil by the R-F sputtering method. The film deposited at room temperature via magnetron sputtering does not show any diffraction peak in the XRD spectrum, indicating the amorphous phase [12, 49]. The amorphous thin film of SiO2 reacts with plated Li forming Li2Si2O5, Li4SiO4, and Li15Si4. Figure 17.9a, b shows the core level Si 2p and O 1 s peak of as-deposited SiO2, after 50th plating and after 50th stripping. The binding energy (BE) of corelevel Si 2p and O 1 spectra at 103.3 eV and 532.71 eV indicates the deposition of SiO2. The shifting of BE for both Si 2p and O 1 s toward a lower value indicates the lithiation of SiO2. Moreover, the shifting of BE for both Si 2p and O 1 s toward the original (as deposited) value indicates the delithiation/recovery of SiO2. Thus, the phenomenon of redox chemical reaction between the plated Li and thin-film SiO2 suggests the storage of Li which slows down the growth of Li dendrite. The deposition of the protective layer increases the impedance for Li-ion transport. The lithiation in such surface coatings during Li-plating/Li-stripping cycles improves the electronic conductivity, resulting in higher conductivity and lower impedance. Cha et al. sputtered 10-nm-thick semiconducting 2H-MoS2, which transforms to 1 T-MoS2 when the number of intercalated Li per MoS2 exceeds 0.4. The MoS2 film uniformity and high ionic conductivity were achieved with the galvanostatic electroplating. The SEI developed by the surface chemistry route with a chemically active SEI component could also store Li by reversible alloy formation [11, 68, 103]. The reduction of metal chlorides, fluorides, or some similar salts by the Li metal can produce the Li-rich alloy and Li-halides. The formation of Li-rich ion conductive alloy as an SEI component enables fast Li-ion migration as they have a higher Li-ion diffusion coefficient than Li metal. Besides, the resistive Li-halides prevent the reduction of Li-ion on the surface, allow Li deposition underneath the protective layer, and inhibit the electron flow, implying the suppression of Li dendrite growth. In our work, an electrolyte containing SnF2 was drop-cast on the surface of Li to form LiF, Sn, and Li-Sn alloy. The chemically active Sn reversibly stores Li by the alloy formation [11]. Figure 17.9c, d shows the cyclic voltammetry (CV) measurement of the bare Li and artificially fluorinate hybrid SEI with an optimized 25- μmthick (AFH-25) symmetrical cell. The straight line in bare Li indicates the Li-plating/

500

R. Pathak et al.

Fig. 17.9 (a, b) The XPS analysis of SiO2-hybrid SEI at different conditions. (Reproduced from Ref. [12] with permission from). (c, d) The first cycle CV measurement of symmetrical cells with bare Li and AFH-25. (e) Plating/stripping electrochemical performance of bare Li and AFH-25 symmetrical cell at 0.5 mA cm
2. (Reproduced from Ref. [11] with permission from)

Li-stripping mechanism. In contrast, AFH-25 symmetrical cell shows the reversible redox peaks at ~
0.12/0.12 v, indicating the lithiation/delithiation of active Sn. The AFH-25 symmetrical cell has the longer stable plating/stripping cycles with reduced nucleation overpotential and voltage hysteresis as shown in Fig. 17.9e. The inset of Fig. 17.9e shows the plating/stripping voltage profile for the first five cycles at a current density of 0.5 mA cm
2 to achieve a capacity of 1 mAh cm
2. The lower

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

501

nucleation overpotential and lower voltage hysteresis were observed for AFH-25 Li compared to the bare Li symmetrical cell. Besides the Li-plating and reaction mechanism to store excess Li, the SEI can also be hybrid in terms of SEI components. For example, the use of organic and inorganic SEI components has successfully stabilized the in situ- and ex situ-based SEI [58, 66, 67]. The flexibility of organic layer or components such as PVDF-HFP, PEO, ROCO2Li, and ROLi (where R represents an alkyl group and styrenebutadiene rubber of SEI buffers the volume expansion of Li during Li plating / stripping. To add up the mechanical strength to maintain the structural integrity within the SEI, the inorganic components such as (LiF, Li2CO3, Al2O3, SiO2).

17.1.5.7

Lithiophilic

The strategy of using lithiophilic decorations or coatings which undergoes conversion and/or alloy reaction has greatly improved the Li deposition behavior even at higher plating/stripping rates. Various lithiophilic coatings (such as metal oxides, metal nitrides, and metal sulfides) or seed decoration such as heterogeneous metallic nanoparticles within the SEI enhances the Li-ion adsorption. The lithiophilic coating/seed decoration guides the initial Li deposition that results in the lower nucleation overpotential and lower voltage overpotential [25]. However, the conversion reaction with lithiophilic metal oxides may lower the CE in the initial cycles due to the formation of lithiophilic Li2O. The presence of electrochemically electronic metal conductors in the SEI facilitates uniform Li-ion transport and a fast electron route. For example, the lithiophilic P2O7 and P3O9 serve as favorable sites for the adsorption of Li atoms due to the presence of a free electron pair on oxygen [62]. The other lithiophilic polar functional groups such as (C¼N, C-N, C¼O, O-H) on the 3D oxidized polyacrylonitrile (PAN) nanofiber interlayer serve as a Li nucleation site which prevents the Li deposition at the hot spots [104]. The higher affinity of the polar functional group confirms better electrode/electrolyte contact. The formation of Li-rich alloy and lithiophilic Li2O during alloying reaction and conversion reaction, respectively, can further enhance the interface stability. Besides, the development lithiophilic-lithiophobic gradient on the SEI also enhances the stability and guides the uniform Li deposition [22, 105, 106]. The lithiophobic upper layer of SEI and lithiophilic lower layer of SEI in dual-layered SEI prevent the Li deposition on the top and allow fast Li-ion transport underneath the SEI. Lithiophilic materials help to strongly capture the Li-ion, reducing the chances of dead Li formation. Lithiophilic materials suppress the growth of Li dendrites.

17.1.6 Operation Under Practical Conditions Many of the researchers are focused on the battery material level, regardless of energy density. To achieve a high energy density needed for practical applications of

502

R. Pathak et al.

Fig. 17.10 (a, b) Schematic of Li protection by PAA-capped matrix and the battery performance comparison between bare Li and PAA@Li under lean electrolyte of 15 μL mAh-1 and cathode mass loading of ~12 mg cm
2. (Reproduced with permission from [Ref. 107]) (c, d) Schematic of Li protection by PVA-modified SEI and the battery performance using bare Li and PVA-protected Li under lean electrolyte of 7.5 μL mAh
1 and cathode mass loading of ~10 mg cm-2. (Reproduced with permission from [Ref. 109]). (e, f) Schematic of Li protection by nanoporous Cu and the battery performance of (Reproduced with permission from [Ref. 108])

LMBs, the crucial cell design principles of considering high mass loading of the cathode, lean electrolyte, and lean lithium play an important role. For example, the long cycling stability with the use of excessive or flooded electrolyte has already been reported. Nevertheless, the crucial practical condition of using lean electrolyte in the cell design principle, which has been overlooked, is a key factor that impacts the energy density of the battery. Feng et al. demonstrated a capped 3D Cu foam by coating the polyacrylic acid (PAA) layer on lithiophilic CuO-Cu foam (PAA@Li matrix) as shown in Fig. 17.10a to offer low interfacial impedance, homogeneous Li-ion flux for smooth and uniform Li deposition, and low nucleation barrier [107]. The chemically and mechanically stable PAA coating, just like a cap, protects the Li and stabilizes the interface for longer cycling stability under high current density (350 h at 5 mA cm
2 for 5 mAh cm
2). The full cell (NMC811/PAA@Li matrix) tested under the electrolyte-to-capacity ratio of 15 μL mAh
1 with high

AU1

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

503

cathode mass loading of 12 mg cm
2 exhibited capacity retention of 75.3% after 200 cycles at 2C as shown in Fig. 17.10b. However, the full cell using bare Li anode showed early capacity decay after 20 cycles. Zhao et al. paired the polyvinyl alcohol (PVA) polymer-protected Li as shown in Fig. 17.10c and NCM as a cathode which demonstrated a stable cycling performance with ~74.5% capacity retention after 120 cycles and an average CE of 99.3% under a lean electrolyte condition (7.5 μL mAh
1) at 0.3 C as shown in Fig. 17.10d [108]. Designing the nanoporous Cu layer on the LMA surface as shown in Fig. 17.10e demonstrated the cycling performance of more than 100 cycles with 7.5 μL mAh
1 in Cu-Li/NCM523 full cell as shown in Fig. 17.10f [109]. Gao et al. used reactive polymer composite (RPC)-poly(vinyl sulfonyl fluoride-ran-2-vinyl-1,3-dioxolane) (P(SF-DOL)), Li fluoride nanoparticles, and graphene oxide sheets to design a stable SEI which suppresses the electrolyte consumption [98]. The RPC-derived SEI-containing polymers, nanoscale LiF particles, and GO nanosheets passivate the Li surface. This leads to the control of the side reaction between the liquid electrolyte and LMA, which has been confirmed from the lower concentration of Li (19.6%) and F (11.2%) in RPC-derived SEI compared to conventional electrolyte-derived SEI (37.3% of Li and 33.6% of F). As a result, the RPC-derived SEI effectively retains 77% of the electrolyte after 180 cycles, but without ex situ SEI it only retains 41% after 50 cycles. Besides, RPC-derived SEI demonstrates high mechanical strength and flexibility due to the presence of GO nanosheets. The polymer-inorganic SEI enabled a high CE of 99.1% at a deposition capacity of 4 mAh cm
2 and stable cycling performance over 200 cycles in 4 V Li/LiNi0.5Co0.2Mn0.3O2 cell under lean electrolyte (7 μl mAh
1), limited Li, and high areal capacity of 3.4 mAh cm
2 conditions. Cha et al. also studied the LMB performance at the lean electrolyte condition (3 g Ah
1, ~ 12 μL) by protecting the LMA through R-F sputtered MoS2 deposition [53]. Yin et al. demonstrated the use of metal chloride perovskite thin films such as MASnCl3 and MAPBCl3 on top of Li metal to shield the LMA from liquid electrolyte [64]. The high symmetry of perovskite provides vertical channels for Li-ion transport, and both the intercalated Li-ions which form Li-M alloy and insulating LiCl layer are beneficial for guiding Li deposition. The perovskite framework only allows Li-ion transport but not the solvent molecules. Such metal chloride perovskite-based interfacial-protected Li, at a lean electrolyte condition of 20 μL mAh
1, enabled a LiCoO2-Li cell with an area capacity of 2.8 mAh cm
2 for more than 100 cycles with a capacity retention of 85% at 0.5C.

17.1.7 Conclusion and Outlook LMA has been considered as a promising anode material in next-generation lithium metal batteries. LMA can deliver a high energy density when paired with conversion and intercalation cathodes such as sulfur, oxygen, metal fluoride, lithium cobalt oxide, lithium nickel manganese cobalt oxide, etc. Despite attractive features, LMA suffers from inherent issues of infinite volume expansion and its hyperactive nature,

504

R. Pathak et al.

leading the challenges in battery safety and performance. The solution to the LMA/liquid electrolyte interface incompatibilities and instabilities can bring the LMA into practical existence. The design of stable and robust SEI is the key to achieve a practical application of LMA. The development of SEI by physical or chemical deposition method has led to the dendrite-free Li metal deposition, resulting in improved electrochemical battery performance. However, the fundamental understanding of the SEI formation, the structural and electrochemical properties, and the regulation or stabilization to revive the LMA are still inadequate. Several advanced experimental and theoretical techniques have been utilized to understand the fundamental properties of SEI formation and their regulation. The continuous research and development of nanoscience and nanotechnology have led to improving the interfacial instabilities of electrode/electrolyte interphase. The theoretical calculations [110–113] and advanced characterization [114–117] have successfully disclosed the mystery of correlation between electrode/ liquid electrolyte interface chemistry and the electrochemical battery performance. Despite noteworthy advancement, there are several unresolved arguments about the formation, mechanism, and impact of SEI and challenges that remain to be discovered in the field of SEI engineering. Based on the above advanced and impactful reports, characterizations, simulations, and our thinking, we believe that this book chapter will catch additional attention in the upcoming research community focusing on the development of high-quality SEI to revive the LMA. The ultrastable and robust SEIs pave the pathway toward the practical applications of rechargeable LMBs including Li-S batteries, Li-O2 batteries, and other high-capacity/high-voltage conversion cathodes. Acknowledgments We acknowledge the financial support from the SDBOR competitive grant, NSF MRI (1428992), EDA University Center Program (ED18DEN3030025), NSF IUCRC (#1841502). Author Contributions This paper was entirely written by RP. The final draft was reviewed QQ. Conflicts of Interest There are no conflicts to declare.

References 1. Van Noorden R (2014) The rechargeable revolution: a better battery. Nature News 507 (7490):26 2. Albertus P, Babinec S, Litzelman S, Newman A (2018) Status, and challenges in enabling the lithium metal electrode for high-energy and low-cost rechargeable batteries. Nat Energy 3 (1):16–21 3. Xu K (2019) A Long journey of Lithium: from the big bang to our smartphones. Energy Environ Mater 2(4):229–233 4. Zhao Z, Pathak R, Wang X, Yang Z, Li H, Qiao Q (2020) Sulfiphilic FeP/rGO as a highly efficient sulfur host for propelling redox kinetics toward the stable lithium-sulfur battery. Electrochim Acta 137117

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

505

5. Adhamash E, Pathak R, Chen K, Rahman MT, El-Magrous A, Gu Z, Lu S, Qiao Q, Zhou Y (2020) High-energy plasma activation of renewable carbon for the enhanced capacitive performance of supercapacitor electrode. Electrochim Acta 362:137148 6. Adhamash E, Pathak R, Qiao Q, Zhou Y, McTaggart R (2020) Gamma-radiated biochar carbon for improved supercapacitor performance. RSC Adv 10(50):29910–29917 7. Zhao Z, Yi Z, Li H, Pathak R, Yang Z, Wang X, Qiao Q (2020) Synergetic effect of spatially separated dual co-catalyst for accelerating multiple conversion reaction in advanced lithiumsulfur batteries. Nano Energy 105621 8. Pathak R (2020) Lithium-ion battery and beyond: oxygen vacancy creation in tungsten trioxide and surface modification of lithium metal 9. Armand M, Tarascon J-M Building better batteries. Nature 2008, 451(7179):652–657 10. Lin D, Liu Y, Cui Y (2017) Reviving the lithium metal anode for high-energy batteries. Nat Nanotechnol 12(3):194 11. Pathak R, Chen K, Gurung A, Reza KM, Bahrami B, Pokharel J, Baniya A, He W, Wu F, Zhou Y (2020) Fluorinated hybrid solid-electrolyte-interphase for dendrite-free lithium deposition. Nat Commun 11(1):1–10 12. Pathak R, Chen K, Gurung A, Reza KM, Bahrami B, Wu F, Chaudhary A, Ghimire N, Zhou B, Zhang WH (2019) Ultrathin bilayer of graphite/SiO2 as a solid interface for reviving Li metal anode. Adv Energy Mater 9(36):1901486 13. Chen K, Pathak R, Gurung A, Adhamash EA, Bahrami B, He Q, Qiao H, Smirnova AL, Wu JJ, Qiao Q (2019) Flower-shaped lithium nitride as a protective layer via facile plasma activation for stable lithium metal anodes. Energy Storage Mater 18:389–396 14. Chen K, Pathak R, Gurung A, Reza KM, Ghimire N, Pokharel J, Baniya A, He W, Wu JJ, Qiao QQ (2020) A copper-clad lithiophilic current collector for dendrite-free lithium metal anodes. J Mater Chem A 8(4):1911–1919 15. Pathak R, Gurung A, Elbohy H, Chen K, Reza KM, Bahrami B, Mabrouk S, Ghimire R, Hummel M, Gu Z (2018) Self-recovery in Li-metal hybrid lithium-ion batteries via WO 3 reduction. Nanoscale 10(34):15956–15966 16. Pathak R, Chen K, Bahrami B, Rahman T, Lu H, Zhou Y, Qiao Q (2019) Capacity Revival of Tungsten trioxide Anode Material in Lithium-Ion Battery. In: 2019 IEEE International Conference on Electro Information Technology (EIT). IEEE, pp 1–5 17. Chen K, Pathak R, Bahrami B, Rahman MT, Lu H, Zhou Y, Qiao Q (2019) SnO 2 nanoparticles embedded biochar as anode material in lithium-ion batteries, 2019 IEEE international conference on electro information technology (EIT). IEEE, pp 1–4 18. Pathak R, Chen K, Zhou Y, Qiao Q (2019) Advances in conversion-type Li-metal fluoride battery: a mini-review 19. Guo Y, Li H, Zhai T (2017) Reviving lithium-metal anodes for next-generation high-energy batteries. Adv Mater 29(29):1700007 20. Liang Z, Lin D, Zhao J, Lu Z, Liu Y, Liu C, Lu Y, Wang H, Yan K, Tao X (2016) Composite lithium metal anode by melt infusion of lithium into a 3D conducting scaffold with lithiophilic coating. Proc Natl Acad Sci 113(11):2862–2867 21. Lu L-L, Ge J, Yang J-N, Chen S-M, Yao H-B, Zhou F, Yu S-H (2016) Free-standing copper nanowire network current collector for improving lithium anode performance. Nano Lett 16 (7):4431–4437 22. Zhang H, Liao X, Guan Y, Xiang Y, Li M, Zhang W, Zhu X, Ming H, Lu L, Qiu J (2018) Lithiophilic-lithiophobic gradient interfacial layer for a highly stable lithium metal anode. Nat Commun 9(1):1–11 23. Ye H, Xin S, Yin YX, Guo YG (2017) Advanced porous carbon materials for high-efficient lithium metal anodes. Adv Energy Mater 7(23):1700530 24. Zhang A, Fang X, Shen C, Liu Y, Zhou C (2016) A carbon nanofiber network for stable lithium metal anodes with high Coulombic efficiency and long cycle life. Nano Res 9 (11):3428–3436

506

R. Pathak et al.

25. Pathak R, Zhou Y, Qiao Q (2020) Recent advances in Lithiophilic porous framework toward dendrite-free Lithium metal anode. Appl Sci 10(12):4185 26. Yang C, Yao Y, He S, Xie H, Hitz E, Hu L (2017) Ultrafine silver nanoparticles for seeded lithium deposition toward stable lithium metal anode. Adv Mater 29(38):1702714 27. Zhu M, Li B, Li S, Du Z, Gong Y, Yang S (2018) Dendrite-free metallic Lithium in Lithiophilic carbonized metal-organic frameworks. Adv Energy Mater 8(18):1703505 28. Zhai P, Wei Y, Xiao J, Liu W, Zuo J, Gu X, Yang W, Cui S, Li B, Yang S (2020) In situ generation of artificial solid-electrolyte interphases on 3D conducting scaffolds for highperformance Lithium-metal anodes. Adv Energy Mater 10(8):1903339 29. Gurung A, Pokharel J, Baniya A, Pathak R, Chen K, Lamsal BS, Ghimire N, Zhang W-H, Zhou Y, Qiao Q (2019) A review on strategies addressing interface incompatibilities in inorganic all-solid-state lithium batteries. Sustain Energy Fuels 3(12):3279–3309 30. Zhang Q, Cao D, Ma Y, Natan A, Aurora P, Zhu H (2019) Sulfide-based solid-state electrolytes: synthesis, stability, and potential for all-solid-state batteries. Adv Mater 31(44):1901131 31. Peled E (1979) The electrochemical behavior of alkali and alkaline earth metals in nonaqueous battery systems—the solid electrolyte interphase model. J Electrochem Soc 126(12):2047 32. Kanamura K, Shiraishi S, Takehara ZI (1996) Electrochemical deposition of very smooth lithium using nonaqueous electrolytes containing HF. J Electrochem Soc 143(7):2187 33. Ein-Eli Y (1999) A new perspective on the formation and structure of the solid electrolyte Interface at the graphite anode of Li-ion cells. Electrochem Solid-State Lett 2(5):212 34. Thevenin J, Muller R (1987) Impedance of lithium electrodes in a propylene carbonate electrolyte. J Electrochem Soc 134(2):273 35. Thevenin J (1985) Passivating films on lithium electrodes. An approach using electrode impedance spectroscopy. J Power Sources 14(1–3):45–52 36. Li G, Gao Y, He X, Huang Q, Chen S, Kim SH, Wang D (2017) Organosulfide-plasticized solid-electrolyte interphase layer enables stable lithium metal anodes for long-cycle lithiumsulfur batteries. Nat Commun 8(1):1–10 37. Zheng J, Engelhard MH, Mei D, Jiao S, Polzin BJ, Zhang J-G, Xu W (2017) Electrolyte additive enabled fast charging and stable cycling lithium metal batteries. Nat Energy 2(3):1–8 38. Qian J, Henderson WA, Xu W, Bhattacharya P, Engelhard M, Borodin O, Zhang J-G (2015) High rate and stable cycling of lithium metal anode. Nat Commun 6(1):1–9 39. Weber R, Genovese M, Louli A, Hames S, Martin C, Hill IG, Dahn J (2019) Long cycle life and dendrite-free lithium morphology in anode-free lithium pouch cells enabled by a dual-salt liquid electrolyte. Nat Energy 4(8):683–689 40. Zhang XQ, Cheng XB, Chen X, Yan C, Zhang Q (2017) Fluoroethylene carbonate additives to render uniform Li deposits in lithium metal batteries. Adv Funct Mater 27(10):1605989 41. Suo L, Xue W, Gobet M, Greenbaum SG, Wang C, Chen Y, Yang W, Li Y, Li J (2018) Fluorine-donating electrolytes enable highly reversible 5-V-class Li metal batteries. Proc Natl Acad Sci 115(6):1156–1161 42. Yan C, Xu R, Xiao Y, Ding JF, Xu L, Li BQ, Huang JQ (2020) Toward critical electrode/ electrolyte interfaces in rechargeable batteries. Adv Funct Mater 1909887 43. Cheng XB, Zhang R, Zhao CZ, Wei F, Zhang JG, Zhang Q (2016) A review of solid electrolyte interphases on lithium metal anode. Adv Sci 3(3):1500213 44. Cheng XB, Hou TZ, Zhang R, Peng HJ, Zhao CZ, Huang JQ, Zhang Q (2016) Dendrite-free lithium deposition induced by uniformly distributed lithium ions for efficient lithium metal batteries. Adv Mater 28(15):2888–2895 45. Lee H, Song J, Kim Y-J, Park J-K, Kim H-T (2016) Structural modulation of lithium metalelectrolyte interface with three-dimensional metallic interlayer for high-performance lithium metal batteries. Sci Rep 6(1):1–10 46. Zheng G, Lee SW, Liang Z, Lee H-W, Yan K, Yao H, Wang H, Li W, Chu S, Cui Y (2014) Interconnected hollow carbon nanospheres for stable lithium metal anodes. Nat Nanotechnol 9 (8):618–623

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

507

47. Zhao Y, Sun Q, Li X, Wang C, Sun Y, Adair KR, Li R, Sun X (2018) Carbon paper interlayers: a universal and effective approach for highly stable Li metal anodes. Nano Energy 43:368–375 48. Lee H, Lee DJ, Kim Y-J, Park J-K, Kim H-T (2015) A simple composite protective layer coating that enhances the cycling stability of lithium metal batteries. J Power Sources 284:103–108 49. Wang L, Wang Q, Jia W, Chen S, Gao P, Li J (2017) Li metal coated with amorphous Li3PO4 via magnetron sputtering for stable and long-cycle life lithium metal batteries. J Power Sources 342:175–182 50. Zhang Y, Liu X, Bai W, Tang H, Shi S, Wang X, Gu C, Tu J (2014) Magnetron sputtering amorphous carbon coatings on metallic lithium: towards promising anodes for lithium secondary batteries. J Power Sources 266:43–50 51. Li Q, Pan H, Li W, Wang Y, Wang J, Zheng J, Yu X, Li H, Chen L (2018) Homogeneous interface conductivity for lithium dendrite-free anode. ACS Energy Lett 3(9):2259–2266 52. Wang L, Zhang L, Wang Q, Li W, Wu B, Jia W, Wang Y, Li J, Li H (2018) Long lifespan lithium metal anodes enabled by Al2O3 sputter coating. Energy Storage Mater 10:16–23 53. Cha E, Patel MD, Park J, Hwang J, Prasad V, Cho K, Choi W (2018) 2D MoS 2 as an efficient protective layer for lithium metal anodes in high-performance Li–S batteries. Nat Nanotechnol 13(4):337–344 54. Kozen AC, Lin C-F, Pearse AJ, Schroeder MA, Han X, Hu L, Lee S-B, Rubloff GW, Noked M (2015) Next-generation lithium metal anode engineering via atomic layer deposition. ACS Nano 9(6):5884–5892 55. Kazyak E, Wood KN, Dasgupta NP (2015) Improved cycle life and stability of lithium metal anodes through ultrathin atomic layer deposition surface treatments. Chem Mater 27 (18):6457–6462 56. Liu Y, Tzeng Y-K, Lin D, Pei A, Lu H, Melosh NA, Shen Z-X, Chu S, Cui Y (2018) An ultrastrong double-layer nanodiamond interface for stable lithium metal anodes. Joule 2 (8):1595–1609 57. Lee DJ, Lee H, Kim YJ, Park JK, Kim HT (2016) Sustainable redox mediation for lithiumoxygen batteries by a composite protective layer on the lithium-metal anode. Adv Mater 28 (5):857–863 58. Liu Y, Lin D, Yuen PY, Liu K, Xie J, Dauskardt RH, Cui Y (2017) An artificial solid electrolyte interphase with high Li-ion conductivity, mechanical strength, and flexibility for stable lithium metal anodes. Adv Mater 29(10):1605531 59. Cho SK, Kim HI, An JW, Jung K, Bae H, Kim JH, Yim T, Lee SY (2020) Printable solid electrolyte interphase mimic for Antioxidative Lithium metal electrodes. Adv Funct Mater 2000792 60. Ma G, Wen Z, Wang Q, Shen C, Jin J, Wu X (2014) Enhanced cycle performance of a Li–S battery based on a protected lithium anode. J Mater Chem A 2(45):19355–19359 61. Tu Z, Choudhury S, Zachman MJ, Wei S, Zhang K, Kourkoutis LF, Archer LA (2017) Designing artificial solid-electrolyte interphases for single-ion and high-efficiency transport in batteries. Joule 1(2):394–406 62. Kim MS, Ryu J-H, Lim YR, Nah IW, Lee K-R, Archer LA, Cho WI (2018) Langmuir– Blodgett artificial solid-electrolyte interphases for practical lithium metal batteries. Nat Energy 3(10):889–898 63. Ju Z, Nai J, Wang Y, Liu T, Zheng J, Yuan H, Sheng O, Jin C, Zhang W, Jin Z (2020) Biomacromolecules enabled dendrite-free lithium metal battery and its origin revealed by cryo-electron microscopy. Nat Commun 11(1):1–10 64. Yin Y-C, Wang Q, Yang J-T, Li F, Zhang G, Jiang C-H, Mo H-S, Yao J-S, Wang K-H, Zhou F (2020) Metal chloride perovskite thin-film based interfacial layer for shielding lithium metal from liquid electrolyte. Nat Commun 11(1):1–9 65. Li NW, Yin YX, Yang CP, Guo YG (2016) An artificial solid-electrolyte interphase layer for stable lithium metal anodes. Adv Mater 28(9):1853–1858

508

R. Pathak et al.

66. Xu R, Zhang XQ, Cheng XB, Peng HJ, Zhao CZ, Yan C, Huang JQ (2018) Artificial soft–rigid protective layer for dendrite-free lithium metal anode. Adv Funct Mater 28(8):1705838 67. Yan C, Cheng XB, Tian Y, Chen X, Zhang XQ, Li WJ, Huang JQ, Zhang Q (2018) Duallayered film protected lithium metal anode to enable dendrite-free lithium deposition. Adv Mater 30(25):1707629 68. Liang X, Pang Q, Kochetkov IR, Sempere MS, Huang H, Sun X, Nazar LF (2017) A facile surface chemistry route to a stabilized lithium metal anode. Nat Energy 2(9):17119 69. Liao K, Wu S, Mu X, Lu Q, Han M, He P, Shao Z, Zhou H (2018) Developing a “waterdefendable” and “dendrite-free” lithium-metal anode using a simple and promising GeCl4 pretreatment method. Adv Mater 30(36):1705711 70. Zhang XQ, Chen X, Xu R, Cheng XB, Peng HJ, Zhang R, Huang JQ, Zhang Q (2017) Columnar lithium metal anodes. Angew Chem Int Ed 56(45):14207–14211 71. Yan J, Yu J, Ding B (2018) Mixed ionic and electronic conductor for Li-metal anode protection. Adv Mater 30(7):1705105 72. Lu Z, Li W, Long Y, Liang J, Liang Q, Wu S, Tao Y, Weng Z, Lv W, Yang QH (2020) Constructing a high-strength solid electrolyte layer by in vivo alloying with Aluminum for an ultrahigh-rate Lithium metal anode. Adv Funct Mater 30(7):1907343 73. Yuan Y, Wu F, Bai Y, Li Y, Chen G, Wang Z, Wu C (2019) Regulating Li deposition by constructing LiF-rich host for dendrite-free lithium metal anode. Energy Storage Mater 16:411–418 74. Wang Z, Wang Y, Zhang Z, Chen X, Lie W, He YB, Zhou Z, Xia G, Guo Z (2020) Building artificial solid-electrolyte interphase with uniform intermolecular ionic bonds toward dendritefree Lithium metal anodes. Adv Funct Mater 2002414 75. Chen H, Pei A, Lin D, Xie J, Yang A, Xu J, Lin K, Wang J, Wang H, Shi F (2019) Uniform high ionic conducting lithium sulfide protection layer for stable lithium metal anode. Adv Energy Mater 9(22):1900858 76. Liu F, Wang L, Zhang Z, Shi P, Feng Y, Yao Y, Ye S, Wang H, Wu X, Yu Y (2020) A mixed Lithium-ion conductive Li2S/Li2Se protection layer for stable Lithium metal anode. Adv Funct Mater 30(23):2001607 77. He M, Guo R, Hobold GM, Gao H, Gallant BM (2020) The intrinsic behavior of lithium fluoride in solid electrolyte interphases on lithium. Proc Natl Acad Sci 117(1):73–79 78. Zhao J, Liao L, Shi F, Lei T, Chen G, Pei A, Sun J, Yan K, Zhou G, Xie J (2017) Surface fluorination of reactive battery anode materials for enhanced stability. J Am Chem Soc 139 (33):11550–11558 79. Lin D, Liu Y, Chen W, Zhou G, Liu K, Dunn B, Cui Y (2017) Conformal lithium fluoride protection layer on three-dimensional lithium by nonhazardous gaseous reagent freon. Nano Lett 17(6):3731–3737 80. Umeda GA, Menke E, Richard M, Stamm KL, Wudl F, Dunn B (2011) Protection of lithium metal surfaces using tetraethoxysilane. J Mater Chem 21(5):1593–1599 81. Fan L, Guo Z, Zhang Y, Wu X, Zhao C, Sun X, Yang G, Feng Y, Zhang N (2020) Stable artificial solid electrolyte interphase films for lithium metal anode via metal-organic frameworks cemented by polyvinyl alcohol. J Mater Chem A 8(1):251–258 82. Liu W, Mi Y, Weng Z, Zhong Y, Wu Z, Wang H (2017) Functional metal-organic framework boosting lithium metal anode performance via chemical interactions. Chem Sci 8 (6):4285–4291 83. Monroe C, Newman J (2003) Dendrite growth in lithium/polymer systems: a propagation model for liquid electrolytes under galvanostatic conditions. J Electrochem Soc 150(10): A1377 84. Shen X, Zhang R, Chen X, Cheng XB, Li X, Zhang Q (2020) The failure of solid electrolyte interphase on Li metal anode: structural uniformity or mechanical strength? Adv Energy Mater 10(10):1903645

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

509

85. Zhang W, Fan L, Tong Z, Miao J, Shen Z, Li S, Chen F, Qiu Y, Lu Y (2019) Stable Li-metal deposition via a 3D Nanodiamond matrix with ultrahigh Young’s Modulus. Small Methods 3 (11):1900325 86. Xie J, Liao L, Gong Y, Li Y, Shi F, Pei A, Sun J, Zhang R, Kong B, Subbaraman R (2017) Stitching h-BN by atomic layer deposition of LiF as a stable interface for lithium metal anode. Sci Adv 3(11):eaao3170 87. Li S, Fan L, Lu Y (2019) Rational design of robust-flexible protective layer for safe lithium metal battery. Energy Storage Mater 18:205–212 88. Shi P, Liu F, Feng Y, Zhou J, Rui X, Yu Y (2020) The synergetic effect of Lithium Bisoxalatodifluorophosphate and Fluoroethylene carbonate on dendrite suppression for fast charging Lithium metal batteries. Small 2001989 89. Li NW, Shi Y, Yin YX, Zeng XX, Li JY, Li CJ, Wan LJ, Wen R, Guo YG (2018) A flexible solid electrolyte interphase layer for long-life lithium metal anodes. Angew Chem 130 (6):1521–1525 90. Zhang SS, Fan X, Wang C (2018) Efficient and stable cycling of lithium metal enabled by a conductive carbon primer layer. Sustain Energy Fuels 2(1):163–168 91. Lin D, Liu Y, Liang Z, Lee H-W, Sun J, Wang H, Yan K, Xie J, Cui Y (2016) Layered reduced graphene oxide with nanoscale interlayer gaps as a stable host for lithium metal anodes. Nat Nanotechnol 11(7):626–632 92. Zhang R, Cheng XB, Zhao CZ, Peng HJ, Shi JL, Huang JQ, Wang J, Wei F, Zhang Q (2016) Conductive nanostructured scaffolds render low local current density to inhibit lithium dendrite growth. Adv Mater 28(11):2155–2162 93. Ostadhossein A, Kim S-Y, Cubuk ED, Qi Y, Van Duin AC (2016) Atomic insight into the lithium storage and diffusion mechanism of SiO2/Al2O3 electrodes of lithium-ion batteries: ReaxFF reactive force field modeling. J Phys Chem A 120(13):2114–2127 94. Wang Y-Q, Gu L, Guo Y-G, Li H, He X-Q, Tsukimoto S, Ikuhara Y, Wan L-J (2012) RutileTiO2 nanocoating for a high-rate Li4Ti5O12 anode of a lithium-ion battery. J Am Chem Soc 134(18):7874–7879 95. Sim S, Oh P, Park S, Cho J (2013) Critical thickness of SiO2 coating layer on Core@ Shell bulk@ nanowire Si anode materials for Li-ion batteries. Adv Mater 25(32):4498–4503 96. Zhang Y, Li Y, Wang Z, Zhao K (2014) Lithiation of SiO2 in Li-ion batteries: in situ transmission electron microscopy experiments and theoretical studies. Nano Lett 14 (12):7161–7170 97. Yang Q, Li W, Dong C, Ma Y, Yin Y, Wu Q, Xu Z, Ma W, Fan C, Sun K (2020) PIM-1 as artificial solid electrolyte interphase for stable lithium metal anode in high-performance batteries. Journal of Energy Chemistry 42:83–90 98. Gao Y, Yan Z, Gray JL, He X, Wang D, Chen T, Huang Q, Li YC, Wang H, Kim SH (2019) Polymer–inorganic solid–electrolyte interphase for stable lithium metal batteries under lean electrolyte conditions. Nat Mater 18(4):384–389 99. Guo W, Liu S, Guan X, Zhang X, Liu X, Luo J (2019) Mixed ion and Electron-conducting scaffolds for high-rate Lithium metal anodes. Adv Energy Mater 9(20):1900193 100. Xu Q, Lin J, Ye C, Jin X, Ye D, Lu Y, Zhou G, Qiu Y, Li W (2020) Air-stable and dendritefree Lithium metal anodes enabled by a hybrid interphase of C60 and mg. Adv Energy Mater 10(6):1903292 101. Dong L, Nie L, Liu W (2020) Water-stable Lithium metal anodes with ultrahigh-rate capability enabled by a hydrophobic Graphene architecture. Adv Mater 1908494 102. Dong L, Nie L, Liu W (2020) Water-stable Lithium metal anodes with ultrahigh-rate capability enabled by a hydrophobic Graphene architecture. Adv Mater 32(14):1908494 103. Tu Z, Choudhury S, Zachman MJ, Wei S, Zhang K, Kourkoutis LF, Archer LA (2018) Fast ion transport at solid-solid interfaces in hybrid battery anodes. Nat Energy 3(4):310–316 104. Liang Z, Zheng G, Liu C, Liu N, Li W, Yan K, Yao H, Hsu P-C, Chu S, Cui Y (2015) Polymer nanofiber-guided uniform lithium deposition for battery electrodes. Nano Lett 15 (5):2910–2916

510

R. Pathak et al.

105. Zheng H, Zhang Q, Chen Q, Xu W, Xie Q, Cai Y, Ma Y, Qiao Z, Luo Q, Lin J (2020) 3D lithiophilic–lithiophobic–lithiophilic dual-gradient porous skeleton for highly stable lithium metal anode. J Mater Chem A 8(1):313–322 106. Cheng Y, Ke X, Chen Y, Huang X, Shi Z, Guo Z (2019) Lithiophobic-lithiophilic composite architecture through co-deposition technology toward high-performance lithium metal batteries. Nano Energy 63:103854 107. Feng Y, Zhang C, Jiao X, Zhou Z, Song J (2020) Highly stable lithium metal anode with nearzero volume change enabled by capped 3D lithophilic framework. Energy Storage Mater 25:172–179 108. Zhao Y, Wang D, Gao Y, Chen T, Huang Q, Wang D (2019) Stable Li metal anode by a polyvinyl alcohol protection layer via modifying solid-electrolyte interphase layer. Nano Energy 64:103893 109. Guo F, Wu C, Chen S, Ai X, Zhong F, Yang H, Qian J (2020) Flaky and dense Lithium deposition enabled by a Nanoporous copper surface layer on Lithium metal anode. ACS Materials Letters 2(4):358–366 110. Zhou Y, Su M, Yu X, Zhang Y, Wang J-G, Ren X, Cao R, Xu W, Baer DR, Du Y (2020) Realtime mass spectrometric characterization of the solid–electrolyte interphase of a lithium-ion battery. Nat Nanotechnol 15(3):224–230 111. Wu H, Chan G, Choi JW, Ryu I, Yao Y, McDowell MT, Lee SW, Jackson A, Yang Y, Hu L (2012) Stable cycling of double-walled silicon nanotube battery anodes through solidelectrolyte interphase control. Nat Nanotechnol 7(5):310–315 112. Wang L, Menakath A, Han F, Wang Y, Zavalij PY, Gaskell KJ, Borodin O, Iuga D, Brown SP, Wang C (2019) Identifying the components of the solid–electrolyte interphase in Li-ion batteries. Nat Chem 11(9):789–796 113. Gu Y, Wang W-W, Li Y-J, Wu Q-H, Tang S, Yan J-W, Zheng M-S, Wu D-Y, Fan C-H, Hu W-Q (2018) Designable ultra-smooth ultra-thin solid-electrolyte interphases of three alkali metal anodes. Nat Commun 9(1):1–9 114. Zachman MJ, Tu Z, Choudhury S, Archer LA, Kourkoutis LF (2018) Cryo-STEM mapping of solid-liquid interfaces and dendrites in lithium-metal batteries. Nature 560(7718):345–349 115. Xu Y, Wu H, He Y, Chen Q, Zhang J-G, Xu W, Wang C (2019) Atomic to Nanoscale origin of Vinylene carbonate enhanced cycling stability of Lithium metal anode revealed by Cryotransmission Electron microscopy. Nano Lett 20(1):418–425 116. Ma Q, Zeng XX, Yue J, Yin YX, Zuo TT, Liang JY, Deng Q, Wu XW, Guo YG (2019) Viscoelastic and nonflammable Interface design–enabled dendrite-free and safe solid Lithium metal batteries. Adv Energy Mater 9(13):1803854 117. Zhang L, Yang T, Du C, Liu Q, Tang Y, Zhao J, Wang B, Chen T, Sun Y, Jia P (2020) Lithium whisker growth and stress generation in an in situ atomic force microscope–environmental transmission electron microscope set-up. Nat Nanotechnol 15(2):94–98

Rajesh Pathak received his MSc. in Physics from the GoldenGate International College, Tribhuvan University, Nepal (2009–2011), and received his PhD from South Dakota State University (2016–2020). Currently, he is a postdoctora; appointee in the Applied Materials Division at Argonne National Laboratory, Lemont, USA. He has coauthored more than 45 peerreviewed journal articles including the leading journals such as Nature Communications, 11(1), 1–10, and Advanced Energy Materials, 9(36), 1901486. Currently, his research focuses on the development of a solid electrolyte-electrode interface (SEI) between lithium metal anode and the electrolyte (liquid/solid state) and electrode property analysis to develop next-generation high energy density lithium metal batteries.

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries

511

Yue Zhou received his Ph.D. in electrical engineering at Pennsylvania State University in 2015. He was a postdoctorate associate at the Massachusetts Institute of Technology who researched the field of energy storage. Currently, he is an assistant professor at South Dakota State University directing the research on lithium metal batteries, sodium-ion batteries, and supercapacitors.

Qiquan Qiao received his Ph.D. from the Virginia Commonwealth University in 2006. He was a postdoctorate associate at the University of Florida. Currently, he is a full professor at South Dakota State University who directed the research in the field of photovoltaics, lithium metal/ion batteries, sensors, micro- or nanomanufacturing/fabrication, food-energy-water (FEW) sustainability, and precision agriculture technologies. He has directed and participated in many projects supported by NSF, NASA, USAID, EDA, 3M, Agilent, Raven Industries, etc. He has published more than 200 peer-reviewed papers in leading journals including Science, Nature Communications, Energy and Environmental Science, Journal of the American Chemical Society, Advanced Materials, Advanced Energy Materials, Advanced Functional Materials, Nanoscale, Joule, ACS Energy Letters, Nano Energy, etc.

Chapter 18

3D X-Ray Characterization of Energy Storage and Conversion Devices Chun Tan, Andrew S. Leach, Thomas M. M. Heenan, Rhodri Jervis, Dan J. L. Brett, and Paul R. Shearing

18.1

Introduction to X-rays

X-rays form part of the electromagnetic spectrum with wavelengths in the range of ca. 0.1 Å to 100 Å, equivalent to energies between ca. 0.1 keV and 100 keV. At the lower energy end of the spectrum are ‘soft’ X-rays with energies up to around 5 keV, whilst ‘hard’ X-rays typically have energies beyond this. Due to their higher energy compared to visible light, they possess a valuable property – the ability to penetrate matter and hence reveal the inner structure of materials that are not optically transparent. The principal interaction mechanisms of X-rays with matter will be discussed in detail in the following section, and these can be harnessed to probe many material properties. Advances in synchrotron and laboratory-based radiation sources have led to numerous X-ray characterization methods being developed, including X-ray absorption spectroscopy (XAS) [1], X-ray diffraction (XRD) [2], and transmission X-ray microscopy (TXM) [3]. These complementary techniques, spanning multiple time and length scales, provide a wealth of information about the electronic states (XAS), crystalline ordering (XRD), and microstructures (TXM) of materials and are being widely adopted to investigate the mechanisms behind the operation and degradation of electrochemical devices, including lithium-based batteries, fuel cells, and redox flow batteries. Many of these X-ray-based spectroscopy and imaging methods can be extended through a process known as tomography to capture a spatially resolved map of the measured property in 3D, although in practice the beam, sample, and detector configurations of some techniques are easier to implement for tomography than

C. Tan (*) · A. S. Leach · T. M. M. Heenan · R. Jervis · D. J. L. Brett · P. R. Shearing Electrochemical Innovation Lab, Department of Chemical Engineering, University College London, London, UK The Faraday Institution, Quad One, Harwell Science and Innovation Campus, Didcot, UK e-mail: [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_18

513

514

C. Tan et al.

others. Regardless, there is a strong motivation for 3D studies of the physical and chemical states of the electrodes within electrochemical devices due to the inherently heterogeneous structures and processes occurring within. Examples of applications include micro-and nano-computed tomography (CT) studies, which have already been widely conducted on a broad range of electrochemical devices to visualize factors such as mechanisms of electrode degradation. Adapted from work on other porous media, parameters such as tortuosity, porosity, and pore and particle size distributions [4–7] may also be extracted from real electrodes, which are valuable metrics, given that they are typically indicative of electrochemical performance. This chapter will begin with an introduction to X-ray interactions with matter (for a more detailed treatment of the topic, the reader is directed to references [8–11], followed by a discussion of the current state of lab-based and synchrotron light sources, including details of their operation and potential future developments. With this in mind, the reader will be guided through most X-ray-based 3D characterization methods that have been applied to energy storage and conversion devices, including more novel techniques that are still under development or are in the early stages of application to this field. The limitations of using X-rays to study energy storage and conversion materials and devices will also be discussed, and how alternative particle sources (e.g. neutrons, electrons, and muons) may provide additional insight, or be more applicable for certain materials. Finally, considering the current knowledge gaps within the literature, the potential direction of future research avenues will be discussed.

18.2

X-ray Interactions with Matter

Like any other form of electromagnetic radiation, X-ray photons can be decomposed into an oscillating set of perpendicular electrical and magnetic fields and exhibit waveparticle duality (i.e. having characteristics of both particles and waves). Both energy and momentum are contained in an electromagnetic wave (or correspondingly, a photon traveling at the speed of light, c [m s
1]) and its energy, E [J s], is inversely proportional to its wavelength, λ [m], with the proportionality constant being the Planck constant h [11], and is given by the following equation: E¼

hc λ

ð18:1Þ

When X-rays propagate through a medium, various interactions with the atoms of the medium result in the absorption, scattering, refraction, and reflection of the incident radiation, as with any other form of electromagnetic radiation. The overall attenuation through absorption or scattering of X-ray photons incident on an object is governed by several principal photon interaction mechanisms, as summarized in Fig. 18.1. An excellent account of these photon interaction mechanisms and their relative significance throughout the electromagnetic spectrum is

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

515

Fig. 18.1 X-ray interaction pathways resulting from an incident X-ray beam on an object, adapted with permission from Banhart [11]

provided at length by Evans [10] or more recently by Banhart [11]. To summarize, the principal photon interaction mechanisms are • • • •

The photoelectric effect. Coherent scattering (Rayleigh/Thomson). Incoherent scattering (Compton). Electron–positron pair production (in the fields of the nucleus and atomic electrons). • Photonuclear absorption (Fig. 18.2).

18.2.1 Absorption The photoelectric absorption of X-rays, as shown in Fig. 18.3, is the dominant effect contributing to the attenuation of incident radiation within the X-ray energy range used in most imaging applications and occurs when an incident X-ray photon interacts with a bound electron in an atom. The probability of an electron occupying a space is generally higher for lower shell numbers (e.g. shell 1 or K), compared to higher shells that are further from the nucleus (e.g. shell 7 or Q) as they have smaller orbital volumes; consequently, the photoelectric effect is more likely to occur for the most tightly bound electrons (i.e. the inner K- and L-shells). This, however, is providing that the incident photon energy is equal to, or exceeds, the binding energy of that electron [10]. If it does, the incident photon is completely absorbed, and the electron is released as a photoelectron with an energy equal to the incident photon less the electron’s binding energy.

516

C. Tan et al.

Fig. 18.2 Relative importance of principal photon interaction mechanisms that dominate at different photon energies, where at σ ¼ τ, the relative contributions between coherent and incoherent scattering and photoelectric absorption are equal, and at σ ¼ κ, the relative contributions between coherent and incoherent scattering and pair production are equal. Adapted with permission from Evans [10]

Fig. 18.3 Various scattering effects of an atom on an incident X-ray photon [11]

18.2.2 Scattering Different types of X-ray scattering will be introduced here, which can be broadly categorized into elastic and inelastic scattering. During an elastic scattering event, there is no net transfer of energy, but the direction of propagation may be modified; during an inelastic scattering event, there is a transfer of energy, meaning the scattered X-ray is of lower energy than the incident X-ray. X-ray diffraction and fluorescence will be discussed here as examples of elastic and inelastic scattering, respectively. Other types of scattering are currently not used for imaging, such as Compton scattering, resonant inelastic X-ray scattering (RIXS), and X-ray Raman scattering.

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

517

18.2.3 Elastic Scattering – X-ray Diffraction X-ray diffraction has been one of the most widely used techniques for measuring the structure of condensed matter. Its development has been the cornerstone of other solid-state measurements and a lot of our understanding of chemical bonding. Diffraction occurs when X-rays are elastically (meaning no energy is transferred) scattered off of an atom. When this atom is part of an ordered structure as found in many minerals, salts, inorganic compounds, or alloys, the diffracted X-rays interact with each other. When the spacing of the atoms, energy of the X-rays, and angle of incidence of the X-ray abide by Bragg’s law, given by Eq. 18.2, there is constructive interference of the diffracted X-rays and a peak will show on the diffraction pattern. At other atom spacings or incident angles, the interference is destructive, and this makes up the background noise to the diffraction pattern. nλ ¼ 2d sinθ

ð18:2Þ

1:2398 λ

ð18:3Þ

E¼

where E is energy (eV) and λ is photon wavelength (μm). A schematic of X-ray interaction with an atomic lattice leading to constructive interference of scattered radiation (diffraction), according to Bragg’s law is illustrated in Fig. 18.4 and defined in Eq. 18.2, where λ is wavelength of the X-ray, d is the atomic spacing, and θ is the incident angle of the X-ray and Eq. 18.2 for photon energy (in electronvolts). The patterns collected during these measurements are ‘fingerprints’ for unique crystal structures, linked to the size and shape of the unit cell. Every possible arrangement of atoms within a unit cell is given a unique identifier based on its symmetry, known as a space group (there are 230 unique space group types), and results in a distinct diffraction pattern. The shape of the unit cell determines the arrangement of the peaks in the diffraction pattern, whilst the size of the unit cell determines the relative separation of those peaks. This means that materials with the same space group give the same general diffraction pattern, although they may differ in peak position if the size of the unit cells (or inter-atom spacing) is different. There are databases of known structures to compare unknown samples, and also mixtures of phases and/or compounds can be distinguished by comparing the intensity of peaks. With high-resolution patterns, subtleties such as

Fig. 18.4 Schematic of Bragg’s law

518

C. Tan et al.

strain on the unit cell can be determined with a fitting of the shape of the diffraction peaks, known as refinement.

18.2.4 Total Attenuation Each of the principal photon interaction mechanisms can be quantified in terms of cross-sections [cm2] that define the probability of an interaction between a photon and an atom within an area traversed by the incident beam, namely, the atomic photoelectric cross-section, τ, for the photoelectric effect; the coherent scattering cross-section, σ coherent, the incoherent scattering cross-section, σ incoherent; the pair production cross-section, κ, for electron positron pair production; and the photonuclear absorption cross-section, σ photonuclear, for the photonuclear absorption effect. These interactions all result in the attenuation of the incident radiation, and therefore to determine the total attenuation of material, each of the contributions can be combined by summation to yield a total cross-section, σ tot [cm2], as defined by Hubbell [12]: σ tot ¼ σ coherent þ σ incoherent þ τ þ κ þ σ photonuclear

ð18:4Þ

At the range of photon energies used in X-ray imaging (between 1 and 100 keV), the photoelectric effect dominates, and both electron–positron pair production and photonuclear absorption are negligible. Therefore, the κ and σ photonuclear terms can be neglected as they only become significant at energies greater than ca. 5 MeV [10, 12], and the total cross-section can be simplified to σ tot ¼ σ coherent þ σ incoherent þ τ

ð18:5Þ

From the density of the material ρ, total cross-section per atom, σ tot, and the atomic mass of the element of interest, Ar, the linear attenuation coefficient, μ (cm
1) or mass attenuation coefficient, μ/ρ (cm2 g
1), can therefore be calculated, as defined by Hubbell [13]: μ N ¼ σ tot A ρ uAr

ð18:6Þ

The total cross-section per atom, σ tot, accounts for each contribution from the principal photon interactions with the material, uAr is the molecular weight of the material, and NA is Avogadro’s constant [13]. This relation is useful because theoretical estimates of the total cross-section are closely aligned to experimentally measured values [12], enabling material identification given a known density. Assuming a monochromatic and collimated X-ray beam of known energy, E, transmission, T, is related to the linear attenuation coefficient, μ, and thickness of a material, t, by the exponential relationship known as the Beer–Lambert law, as

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

519

displayed in Eq. 18.7. In turn, transmission, T, can be determined by measuring the intensities of both the incident beam, I0, and the transmitted beam, I, as it is simply the ratio between the two, as given by Eq. 18.8. T ¼ exp ð
μt Þ T¼

I I0

ð18:7Þ ð18:8Þ

Transmission is a useful concept in practical terms as the intensities of both incident, and transmitted X-ray beams can be measured using scintillators and cameras in X-ray imaging systems, where the incident beam can be measured without the sample in the field-of-view (known variously as the reference, flatfield, light-field, or open-beam image, etc.).

18.2.5 Refraction The wave-like properties of X-ray photons also result in the refraction of incident radiation when passing through a medium, as visible light would when passing through an optically transparent material. The ratio of the phase velocity of electromagnetic radiation within a material to the speed of light in a vacuum is known as the complex refractive index of the material, n. This contains both real (1
δ) and imaginary (β) components, which, respectively, describe the phase shift and attenuation of the incident radiation [14]: n ¼ 1
δ
iβ

ð18:9Þ

Within the range of X-ray energies where the photoelectric effect dominates, but away from the absorption edges of the elements within the medium of interest, both terms can be shown to reduce to [14] r c λ2 2π λ β ðλÞ ¼ μ ðλÞ 4π δ ð λ Þ ¼ ρe

ð18:10Þ ð18:11Þ

Where ρe is the electron density, rc is the classical electron radius, λ is the wavelength of the incident radiation, and μ is the wavelength (or energy)-dependent attenuation coefficient. It is important to emphasize that these equations are only valid at X-ray energies away from the absorption edge because photon interactions

520

C. Tan et al.

around the absorption edge result in a jump in the linear attenuation coefficient that is not accounted for in the term ‘μ’.

18.2.6 Fluorescence Incident beams (e.g. X-ray or electron) can eject atomic electrons and produce fluorescence X-rays. Fluorescence occurs when the hole left by the ejected electron is filled by an electron from a higher orbital, e.g. a K-shell hole populated by an L-shell electron. The energy at the orbital closer to the nucleus is lower; consequently, the surplus energy must be ejected as radiation in the form of an X-ray photon that is characteristic of the material. These types of photons are described as characteristic because their energy is equal to the difference between the two shell energies (e.g. between L and K), which in turn are determined by the atomic number of the element. Fluorescence detectors typically output spectra of photon intensities for a range of emitted X-ray energies. If fluorescence is occurring, several sharp intensity peaks are observed around energies that are characteristic of the material’s well-defined electron orbital energies. These peaks can be distinguished according to the electron transition that has occurred, i.e. which outer shell the electron traveled from to fill the lower shell hole. For instance, the L!K shell transition is generally denoted as Kα, and the M!K is noted as Kβ; and if the hole resides in the L shell, a transition from the M shell would be denoted Lα, and so on. Once the peaks have been allocated to corresponding transitions, the intensity of the characteristic emission can be correlated to the amount of each element, thus the material composition. However, although the range of Kα values for the periodic table as a whole is broad, between 0.1–100 keV, for many useful elements values can be relatively low, e.g. Si Kα is ca. 1.7 keV, and therefore highly susceptible to attenuation, making the detection of fluorescence signals from low Z elements increasingly difficult for increasing sample thickness [15].

18.3

X-ray Sources for Characterization

Two main phenomena are exploited for the generation of X-ray photons used in scientific experiments: the first involves bombarding an anode with electrons accelerated from a cathode in X-ray tubes; and the second is the release of radiation from the acceleration of electrons using magnets at particle accelerators [11]. The former is more prevalent in laboratories, where the anode target (often a metal such as tungsten or copper) produces X-ray photons via two main interaction pathways [16]: Bremsstrahlung radiation that has originated from the deceleration of high-speed electrons by the electric fields of the target nuclei, producing a continuous spectrum of photon energies; or characteristic radiation that produces peaks of

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

521

high photon intensities at discrete energies specific to the target element, via fluorescence [15]. The accelerating voltage of an X-ray tube can be considered equivalent to the kinetic energy of the electrons and dictates the highest possible energy of an X-ray photon; i.e., a 120 kV tube voltage results in the emission of X-ray photons of no more than 120 keV in energy [16]. The second origin of X-ray photons often comes from a particular type of particle accelerator: a synchrotron. These consist of many straight sections connected into a near-circular, polyhedral storage ring. Charged sub-atomic particles are injected into this ring and then circulated under a high vacuum. The electrons are decelerated in a controlled manner using magnetic fields to produce X-rays of a particular energy. Several large-scale synchrotron facilities exist today with the primary purpose of serving as light sources for experimentalists in fields ranging from biology to materials science and engineering. ‘Brilliance’ is often used to quantify beam quality, and unlike the continuous spectrum produced in lab sources, synchrotrons produce high brilliance beams; i.e., they produce a high number of photons per unit time, per unit area, with low divergence, and a high degree of monochromaticity. Moreover, beams can be further differentiated according to geometry; depending on the nature of the light source and source optics, the resulting X-ray illumination may have a pencil-, parallel-, fan-, or cone-beam geometry, as illustrated in Fig. 18.5.

18.3.1 Lab-Based X-ray Sources A lab-based X-ray source typically produces a divergent beam with a cone geometry, as presented in Fig. 18.5. This results in the formation of a penumbra and geometric magnification, M, of the object depending on the distance between the object, source, and detector [17]. The geometric magnification of the object is given by M ¼ (R1 + R2)/R1, where R1 is the source-to-sample distance, and R2 is the sampleto-detector distance, also shown in Fig. 18.5. By reducing the spot size of the beam with a microfocus source, the geometric penumbra can be limited to some extent. A comparison between different spot sizes can be seen in Fig. 18.6. The polychromatic nature of the X-ray beam generated with lab-based sources leads to undesirable effects such as beam-hardening because materials tend to attenuate X-rays at lower photon energies more than X-rays at higher photon energies [19], resulting in the transmitted beam having a different spectral composition from the incident beam [11]. To minimize this effect, a combination of source filters and software correction can be used. Despite these inherent shortcomings, the main advantages of lab-based X-ray sources are the significantly lower cost and greater accessibility when compared to large-scale facilities.

522

C. Tan et al.

Fig. 18.5 Pencil-, parallel-, and cone-beam configurations [17], and the penumbra effect arising from the cone beam

18.3.2 Synchrotron Light Sources As previously mentioned, the advantages of synchrotron radiation sources are generally due to their favorable beam brilliance; the beams produced are collimated (parallel), high-flux, and tunable to specific photon energies using complex magnets and monochromator crystals [20]. The parallel beam geometry (Fig. 18.5) of a synchrotron source simplifies 3D reconstruction and removes the issue of penumbra effects. Furthermore, the high flux significantly reduces acquisition times, to the extent that experiments may instead be time-limited by stage motor movements and detector readout frequency [17]. For comparison, a typical acquisition that may require hours with lab-based micro-computed tomography (micro-CT) may be completed within minutes for synchrotron micro-CT. Thus, the high temporal resolution achievable at synchrotron sources opens vast possibilities for operando tomography on energy storage materials, where microstructural changes in the material may occur within relatively short timescales. Finally, the highly tunable X-ray energies available with synchrotron radiation enable numerous possibilities

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

523

Fig. 18.6 Reduction in penumbra effects with a microfocus X-ray source (left) (right) compared to X-ray source with larger spot size (right), reproduced with permission from Withers [18]

for performing measurements around the absorption edges of different elements, enabling the study of electrochemical changes in energy materials [21]. The production of X-ray photons at synchrotron light sources was described in simple terms in the previous section, although the differences in configuration between synchrotron light sources and between beamlines at the same light source vary immensely. For end-users, the experimental hutch is where measurements are performed, although many components exist between this hutch and the storage ring to produce and refine the spatial and temporal coherence of the X-ray beam that is required to be compatible with the characterization method of interest. Close to the main storage ring is the ‘front end’ where X-rays are produced using insertion devices (ID), such as electron wigglers and undulators, or sometimes bending magnets, and may then be focused with a focusing mirror. For spectroscopy and scanning microscopy applications, a ‘pencil’ beam is commonly used, with a well-defined spot size. At X-ray imaging and microscopy end stations, a full-field beam is typically used, the size of which can be controlled by adjustable slits along the beam path. The energy of the beam can be selected using monochromators, although in some cases demanding high throughput, a ‘pink’ or ‘white’ beam is used due to the significantly higher flux attainable when using a wider spectral range. Whilst the spatial divergence of the beam is generally small, especially for beamlines located further away from the storage ring (i.e. a parallel geometry may be approximated), further focusing is possible with X-ray optics such as Fresnel zone plates or Kirkpatrick–Baez (KB) mirrors. The sample is typically mounted on a stage, which may be capable of up to six degrees of spatial freedom to align the sample with the beam and detector: X, Y, and Z translation steps, angular rotation, pitch, and yaw. Downstream from the sample,

524

C. Tan et al.

detectors are used to record changes to the beam resulting from its propagation through the sample, by measuring either the transmitted beam or its scattered components. The use of evacuated or helium-filled flight tubes within the beam path may be necessary for measurements conducted with ‘soft’ X-rays due to the significance of air attenuation. Many types of detectors exist; the most common type contains a scintillator made of X-ray fluorescent phosphors, such as CdWO4, that converts the incident X-ray beam into visible light. The image formed on the scintillator may then be magnified through coupled optics or captured directly by a charge-coupled device (CCD) within a camera. Whilst scintillator-based detectors measure incident X-ray radiation indirectly by conversion to visible light and offer the highest spatial resolution, they often have a limited dynamic range and introduce noise into an image from various sources (including electronic and readout noise). Photon counting detectors, such as photomultiplier tubes and semiconductor-based hybrid detectors, as offered by Medipix and PILATUS, have been developed to directly measure ionizing radiation, counting single photons that hit the detector pixel. With these detectors, dynamic range is practically unlimited and the images collected are generally noise-free, although it is still possible to saturate detector pixels and potentially damage them if the incident photon flux is too high. Furthermore, there are limits on the spatial resolution achievable due to the pixel size and pitch on these detector arrays. Most of the control, acquisition, and reconstruction software at synchrotron light sources are customized open-source code, running through either a command line or a graphical user interface. This is necessary due to the highly novel nature of each experiment and the numerous possible permutations in the configuration of the beamline. Despite the obvious advantages of using synchrotron light sources for 3D X-ray studies, several practical limitations exist that confine their broader adoption; these include high capital and running costs and limited accessibility, although extensive user programs exist that ensure affordable or free access for academic research. Additionally, the high flux of the synchrotron X-ray sources may induce dosedependent radiation damage to electrodes within in-situ cells due to material dissolution or degradation of polymers commonly used within energy storage materials [22].

18.4

Applications of 3D Imaging

The various beam geometries achievable within a laboratory and at synchrotron light sources have been discussed in the previous section. Cone- and parallel-beam geometries are most frequently used in full-field imaging techniques, where the sample is captured in a ‘single shot’ within a certain field-of-view. The pencilbeam geometry is mainly used for scanning probe imaging, where the sample is rastered across the beam whilst measurements are made, with each point corresponding to a different spatial position. Finally, novel ‘lens-less’ techniques,

18

3D X-Ray Characterization of Energy Storage and Conversion Devices

525

such as coherent diffraction imaging and ptychography, and their 3D reconstruction, will be discussed.

18.4.1 Full-Field Techniques: X-ray Absorption CT As described previously, X-ray absorption imaging (and CT) detects changes in the imaginary part of the complex refractive index of a material, and within the range of energies typically used in X-ray imaging, photoelectric absorption accounts for a large fraction of the total attenuation. Measurements obtained by absorption imaging provide information about the mass density of material through the relationship between transmission and the mass attenuation coefficient, which permits the visualization of the physical structure of a material. With a priori knowledge of the sample composition, some inferences can be made about the presence or absence of phases throughout a sample, provided there is sufficient contrast. X-ray absorption CT is a relatively mature field stemming from medical and industrial diagnostics. Its application to energy materials, whilst more recent, has proved compelling and particularly suited to the heterogeneous nature of electrode materials – enabling a more informed insight not only into the diagnosis of failure mechanisms in commercial devices but also in developing novel battery chemistries such as Li-S and Li-air through visualization of the phenomena occurring within the materials. The main advantage of tomographic techniques lies in the ability to spatially distinguish between features that may be super-positioned in a throughplane 2D image or a subsurface layer in a 2D surface image. Tomographic techniques encompass a wide range of length scales, enabling the visualization of energy storage devices and materials from the system and pack level, to individual particles, and even to individual grains within a single particle. One of the first few applications of X-ray CT on battery systems was the use of an industrial CT scanner on NiMH spacecraft batteries [23]. The use of X-ray tomography for the 3D reconstruction of battery electrodes was pioneered by Shearing et al. in 2010, who performed high-resolution X-ray micro-tomography (micro-CT) on the graphite negative electrode of a Li-ion battery using a Gatan X-ray ultramicroscope [4]. The voxel dimensions obtained were 480 nm using 41.4 magnification, producing a total reconstructed sample volume of 43  348  478μm3, from which microstructural information such as porosity, tortuosity, surface area, connectivity of pores, and pore and particle size distribution were extracted [4]. Haibel et al. performed the first in-situ micro-CT of a battery on alkaline zinc cells at a synchrotron facility and, in addition to quantifying the particle size and spatial distribution of zinc particles, attempted to correlate changes in the volume fraction of the cathode material during discharge to the mean density of the material and hence the chemical processes that occur which result in changes in mean density [24]. Since then, various research groups have applied micro- and nano-CT to diagnose cell failure mechanisms (in terms of microstructural degradation such as cracks and dislocations formed) [25]; determined the morphological evolution

526

C. Tan et al.

(in terms of microstructural parameters and the contact area between the different phases, etc.) of a wide range of Li-ion battery cathodes and anodes [26–30] with cycling; studied the current collectors used in batteries [31]; and characterized the presence and structures of lithium metal at the anode (such as the formation of lithium dendrites and other electrodeposited microstructures using synchrotron radiation sources) [32, 33]. Not limited to battery studies, absorption CT has been used to study the evolving microstructure of solid oxide fuel cell materials and multiple length scales, from the overall geometry of the cell to the macropore structure, and to the microstructure of the composite electrodes themselves (where the contrast between the electrolyte, catalyst, and pore phases is required to characterize the triple-phase boundary density and therefore performance of the material [34]). Although the feature size of polymer electrolyte membrane fuel cell catalysts is too small to resolve with X-ray imaging (