Nanomaterials for electrochemical energy storage devices 9781119510000, 1119510007, 9781119510048, 111951004X

Table of contents :
Cover
Title Page
Copyright Page
Contents
Preface
Part 1: General Introduction to Battery and Supercapacitor, Fundamental Physics Characterization Techniques
1 Electrochemistry of Rechargeable Batteries Beyond Lithium-Based Systems
1.1 Lithium-Based Batteries
1.1.1 Lithium Primary Batteries
1.1.2 Lithium Metal-Based Secondary Batteries
1.1.3 Polymer Electrolyte-Based Lithium Batteries
1.1.4 Lithium-Ion Batteries
1.1.5 Advances in Li-Ion Batteries
1.1.6 Beyond Lithium-Based Systems
1.2 Cathodes for Na-Ion Batteries
1.2.1 Transition Metal Oxides
1.2.1.1 Single Metal Oxides 1.2.1.2 Multi-Metal Oxides1.2.2 Polyanionic Compounds
1.2.3 Fluorides
1.2.4 Metal Hexacyanometalates
1.2.5 Organic Compounds
1.3 Anodes for Na-Ion Batteries
1.3.1 Carbon-Based Electrodes
1.3.2 Alloy Electrodes
1.3.3 Phosphorous, Phosphides, and Nitrides
1.3.4 Sulfides and Selenides
1.3.5 Phosphates
1.3.6 Organic Materials
1.3.7 Oxides
1.3.8 Sodium-Sulfur Batteries
1.3.9 Na-Air Batteries
1.4 Potassium Batteries
1.4.1 Potassium-Ion Batteries
1.4.1.1 Electrolytes
1.4.1.2 Cathode Materials
1.4.1.3 Anode Materials
1.4.2 Potassium-Sulfur Batteries
1.4.3 Potassium-Air Batteries 1.5 Mg-Based Rechargeable Batteries1.6 Conclusions
References
2 Li-Ion Battery Materials: Understanding From Computational View-Point
2.1 Cathode
2.1.1 Cluster Expansion
2.1.1.1 LiTi2O4
2.1.1.2 LiTiS2
2.1.1.3 LiMn2O4
2.1.1.4 LixCoO2
2.1.1.5 Li(Ni0.5Mn0.5)O2
2.1.2 Phase Stability with Gas-Phase Evolution
2.1.3 Solid State Diffusion
2.1.3.1 LiTi2O4
2.1.3.2 LiTi2S4
2.1.3.3 LiFePO4
2.1.3.4 LiCoO2
2.1.3.5 Lithium Mobility in Layered Transition Metal Oxides
2.1.4 Prediction of New Materials and Combinatorial Chemistry
2.1.4.1 Phosphates
2.1.4.2 Metal Mixing in Olivines
2.2 Anode 2.2.1 Phase Transitions in Graphite2.2.2 Fracture in Graphite
2.2.3 Diffusion in Graphene
2.2.4 Lithiation of Silicon Anodes
2.3 Electrolyte
2.3.1 Solid Electrolyte Interphase
2.3.2 Cathode Side Effects of Electrolyte
2.3.3 Solid State Electrolytes
2.3.3.1 LGPS Family
2.3.3.2 Diffusion in Solid Electrolytes --
Case of LGPS
2.4 Conclusions
Acknowledgment
References
Part 2: Battery: Anode, Cathode and Non-Li-Ion Batteries
3 Nanostructured Anode Materials for Batteries (Lithium Ion, Ni-MH, Lead-Acid, and Thermal Batteries
3.1 Introduction
3.2 Li-Ion Batteries 3.2.1 Electrochemistry of Lithium Ion Batteries3.2.2 Compatibility of Electrode Materials with the Electrolyte
3.2.3 Anode Materials for LIBs
3.2.3.1 Lithium Metal
3.2.3.2 Intercalation/De-Intercalation Materials
3.2.3.3 Alloying/De-Alloying Materials
3.2.3.4 Conversion Type Anode Materials
3.3 Nickel Metal Hydride Batteries
3.3.1 Mechanism of Ni-MH Battery Operation
3.3.2 Anode Materials
3.3.2.1 Rare Earth-Based AB5 Alloys
3.3.2.2 Ti and Zr-Based AB2 Type Alloys
3.3.2.3 Mg Based Alloys
3.3.2.4 Rare Earth-Mg-Ni-Based Superlattice Alloys

Citation preview

Nanomaterials for Electrochemical Energy Storage Devices

Scrivener Publishing 100 Cummings Center, Suite 541J Beverly, MA 01915-6106 Publishers at Scrivener Martin Scrivener ([email protected]) Phillip Carmical ([email protected])

Nanomaterials for Electrochemical Energy Storage Devices

Edited by

Poulomi Roy CSIR-Central Mechanical Engineering Research Institute, Durgapur, India and

Suneel Kumar Srivastava School of Energy Science and Engineering, Indian Institute of Technology, Kharagpur, India

This edition first published 2019 by John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, USA and Scrivener Publishing LLC, 100 Cummings Center, Suite 541J, Beverly, MA 01915, USA © 2020 Scrivener Publishing LLC For more information about Scrivener publications please visit www.scrivenerpublishing.com. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording, or otherwise, except as permitted by law. Advice on how to obtain permission to reuse material from this title is available at http://www.wiley.com/go/permissions. Wiley Global Headquarters 111 River Street, Hoboken, NJ 07030, USA For details of our global editorial offices, customer services, and more information about Wiley products visit us at www.wiley.com. Limit of Liability/Disclaimer of Warranty While the publisher and authors have used their best efforts in preparing this work, they make no representations or warranties with respect to the accuracy or completeness of the contents of this work and specifically disclaim all warranties, including without limitation any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives, written sales materials, or promotional statements for this work. The fact that an organization, website, or product is referred to in this work as a citation and/or potential source of further information does not mean that the publisher and authors endorse the information or services the organization, website, or product may provide or recommendations it may make. This work is sold with the understanding that the publisher is not engaged in rendering professional services. The advice and strategies contained herein may not be suitable for your situation. You should consult with a specialist where appropriate. Neither the publisher nor authors shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. Further, readers should be aware that websites listed in this work may have changed or disappeared between when this work was written and when it is read. Library of Congress Cataloging-in-Publication Data ISBN 978-1-119-51003-1 Cover image: Pixabay.com Cover design Russell Richardson Set in size of 11pt and Minion Pro by Manila Typesetting Company, Makati, Philippines Printed in the USA 10 9 8 7 6 5 4 3 2 1

Dedicated to Our Parents

Contents Preface

xvii

Part 1: General Introduction to Battery and Supercapacitor, Fundamental Physics Characterization Techniques 1 1 Electrochemistry of Rechargeable Batteries Beyond Lithium-Based Systems Brij Kishore, Shyama Prasad Mohanty and Munichandraiah Nookala 1.1 Lithium-Based Batteries 1.1.1 Lithium Primary Batteries 1.1.2 Lithium Metal-Based Secondary Batteries 1.1.3 Polymer Electrolyte-Based Lithium Batteries 1.1.4 Lithium-Ion Batteries 1.1.5 Advances in Li-Ion Batteries 1.1.6 Beyond Lithium-Based Systems 1.2 Cathodes for Na-Ion Batteries 1.2.1 Transition Metal Oxides 1.2.1.1 Single Metal Oxides 1.2.1.2 Multi-Metal Oxides 1.2.2 Polyanionic Compounds 1.2.3 Fluorides 1.2.4 Metal Hexacyanometalates 1.2.5 Organic Compounds 1.3 Anodes for Na-Ion Batteries 1.3.1 Carbon-Based Electrodes 1.3.2 Alloy Electrodes 1.3.3 Phosphorous, Phosphides, and Nitrides 1.3.4 Sulfides and Selenides

3

4 4 5 5 6 8 9 9 9 12 16 17 21 21 22 23 23 25 26 27

vii

viii

Contents 1.3.5 Phosphates 1.3.6 Organic Materials 1.3.7 Oxides 1.3.8 Sodium–Sulfur Batteries 1.3.9 Na-Air Batteries 1.4 Potassium Batteries 1.4.1 Potassium-Ion Batteries 1.4.1.1 Electrolytes 1.4.1.2 Cathode Materials 1.4.1.3 Anode Materials 1.4.2 Potassium–Sulfur Batteries 1.4.3 Potassium–Air Batteries 1.5 Mg-Based Rechargeable Batteries 1.6 Conclusions References

2 Li-Ion Battery Materials: Understanding From Computational View-Point Jishnu Bhattacharya 2.1 Cathode 2.1.1 Cluster Expansion 2.1.1.1 LiTi2O4 2.1.1.2 LiTiS2 2.1.1.3 LiMn2O4 2.1.1.4 LixCoO2 2.1.1.5 Li(Ni0.5Mn0.5)O2 2.1.2 Phase Stability with Gas-Phase Evolution 2.1.3 Solid State Diffusion 2.1.3.1 LiTi2O4 2.1.3.2 LiTi2S4 2.1.3.3 LiFePO4 2.1.3.4 LiCoO2 2.1.3.5 Lithium Mobility in Layered Transition Metal Oxides 2.1.4 Prediction of New Materials and Combinatorial Chemistry 2.1.4.1 Phosphates 2.1.4.2 Metal Mixing in Olivines 2.2 Anode 2.2.1 Phase Transitions in Graphite 2.2.2 Fracture in Graphite

29 29 30 33 35 38 39 40 40 41 43 43 44 49 50 67 67 68 70 73 74 77 80 80 84 86 87 93 94 98 102 102 107 113 113 115

Contents ix 2.2.3 Diffusion in Graphene 2.2.4 Lithiation of Silicon Anodes 2.3 Electrolyte 2.3.1 Solid Electrolyte Interphase 2.3.2 Cathode Side Effects of Electrolyte 2.3.3 Solid State Electrolytes 2.3.3.1 LGPS Family 2.3.3.2 Diffusion in Solid Electrolytes – Case of LGPS 2.4 Conclusions Acknowledgment References

Part 2: Battery: Anode, Cathode and Non-Li-Ion Batteries 3 Nanostructured Anode Materials for Batteries (Lithium Ion, Ni-MH, Lead-Acid, and Thermal Batteries Surendra K. Martha and Liju Elias 3.1 Introduction 3.2 Li-Ion Batteries 3.2.1 Electrochemistry of Lithium Ion Batteries 3.2.2 Compatibility of Electrode Materials with the Electrolyte 3.2.3 Anode Materials for LIBs 3.2.3.1 Lithium Metal 3.2.3.2 Intercalation/De-Intercalation Materials 3.2.3.3 Alloying/De-Alloying Materials 3.2.3.4 Conversion Type Anode Materials 3.3 Nickel Metal Hydride Batteries 3.3.1 Mechanism of Ni-MH Battery Operation 3.3.2 Anode Materials 3.3.2.1 Rare Earth-Based AB5 Alloys 3.3.2.2 Ti and Zr-Based AB2 Type Alloys 3.3.2.3 Mg Based Alloys 3.3.2.4 Rare Earth–Mg–Ni-Based Superlattice Alloys 3.3.2.5 Ti–V-Based Multicomponent Multiphase Alloys

118 122 125 126 130 131 131 135 140 141 141

145 147 148 149 149 151 153 153 156 168 176 180 181 183 184 185 185 186 187

x

Contents 3.4 Lead-Acid Batteries 3.4.1 Operating Principle 3.4.2 Negative Electrodes of Lead-Acid Batteries 3.4.2.1 Preparation of Negative Electrode 3.4.2.2 Sulfation 3.5 Thermal Batteries 3.5.1 Anode Materials for Thermal Batteries 3.5.1.1 Ca-Based Anodes 3.5.1.2 Mg and Al-Based Anodes 3.5.1.3 Li Anode 3.5.1.4 Li–Al Anodes 3.5.1.5 Li–Si Anode References

4 Nanostructured Cathode Materials for Li-/Na-Ion Aqueous and Non-Aqueous Batteries Farheen N. Sayed, Ganguli Babu and P. M. Ajayan 4.1 Introduction 4.1.1 Li+ vs. Na+ ion Batteries 4.1.2 Aqueous vs. Non-Aqueous Electrolyte 4.2 Background of Cathode Materials 4.3 Important Types of Cathode (Class) with Different Electrolytes 4.3.1 Li-ion based Nano Cathodes with Aqueous Electrolyte 4.3.2 Li-ion based Nano Cathodes with Non-Aqueous Electrolyte 4.3.3 Na+ ion based Nano Cathodes with Aqueous Electrolyte 4.3.4 Na+ ion based Nano Cathodes with Non-Aqueous Electrolyte 4.4 Methods to Prepare Nanostructured Cathodes 4.4.1 Solid-State Protocols 4.4.2 Sol–Gel Synthesis 4.4.3 Combustion Method 4.4.4 Hydrothermal Route 4.5 Future Aspects References

187 189 190 190 193 201 203 203 204 204 204 205 207 231 232 234 235 238 240 240 244 248 249 254 256 257 259 260 262 263

Contents xi 5 Polymer-Assisted Chemical Solution Method to Metal Oxide Nanoparticles for Lithium-Ion Batteries Di Huang and Hongmei Luo 5.1 Introduction 5.2 Carbon-Based Composites 5.3 Polymer-Assisted Chemical Solution Method 5.4 Oxygen Deficiency 5.5 Summary and Future Perspectives References

271 272 273 276 284 284 286

6 Li–Air: Current Scenario and Its Future 291 Saravanan Karuppiah, Remith Pongilat and Kalaiselvi Nallathamby 6.1 Introduction: Why Lithium–Air Batteries? 291 6.2 General Characteristics 296 6.2.1 Types of Lithium–Air Batteries 297 6.3 Chemistry and Mechanism 299 6.3.1 Oxygen Reduction Reaction (ORR), Oxygen Evolution Reaction (OER), and the Catalysts 301 6.4 Critical Challenges 309 6.4.1 Electrolytes 310 6.4.2 Decomposition of Electrolyte During Discharge 310 6.4.3 Passivation and Blockage of Oxygen Diffusion 314 6.4.4 Large Polarization 314 6.4.5 Lithium Dendrite Formation 315 6.4.6 Electrocatalysis 316 6.4.7 Rate Capability 317 6.4.8 Energy and Power Density 317 6.4.9 Volume Changes 318 6.5 Non-Aqueous Li/Air Systems 318 6.5.1 Electrochemistry of Oxygen Reduction and Oxidation in Non-Aqueous System 318 6.5.2 Technical Challenges in NLAS 322 6.5.2.1 Designing of Air Cathode/Oxygen Transport 322 6.5.2.2 Effective Loading of Catalysts 323 6.5.2.3 Slow Kinetics of Oxygen Reactions/Deposition of Solid Insulating Products 323 6.5.2.4 Decomposition of Non-Aqueous Electrolytes/Effect of Possible Side Reactions 323 6.5.2.5 Lithium Dendrite Formation and Side 324 Reactions of Li with H2O and Air

xii

Contents 6.5.3 Electrocatalysts for NLAS 6.5.3.1 Carbon Based Materials 6.5.3.2 Metal and/or Metal Oxides 6.5.3.3 Composite Materials 6.5.3.4 Other Cathode Materials 6.5.4 Electrolytes Deployed in Non-Aqueous Li–Air Cells 6.5.4.1 Alkyl Carbonates 6.5.4.2 Esters 6.5.4.3 Ethers 6.5.4.4 Nitriles 6.5.4.5 Amides 6.5.4.6 DMSO 6.5.4.7 Sulfones 6.5.4.8 Ionic Liquids 6.5.5 Morphology of the Deposited Products 6.6 Aqueous Lithium–Air System 6.6.1 Approaches for the Formation of Water Stable Lithium Metal 6.6.1.1 Solid Electrolyte 6.6.1.2 Stability of Solid Electrolyte— Why Do We Need Buffer Layer? 6.6.1.3 Buffer Layer 6.6.2 Catholytes 6.6.2.1 Acidic Catholyte 6.6.2.2 Alkaline Catholyte 6.6.3 Catalysts for Acidic and Alkaline System 6.6.4 Managing the Precipitation of LiOH.H2O 6.6.5 Hybrid Lithium–Air Battery 6.7 Applications 6.8 Future of Lithium–Air Systems References

7 Sodium-Ion Battery Anode Stabilization Prasit Kumar Dutta, Arnab Ghosh and Sagar Mitra 7.1 Introduction 7.2 History of NIB 7.3 Operational Principle 7.4 Types of Storage Mechanisms 7.5 Issues and Challenges in a NIB 7.6 Brief Updates on Cathode and Anode Materials Research

324 324 332 336 338 339 339 340 340 340 341 341 341 342 343 345 346 346 350 351 355 355 358 359 359 363 364 365 367 377 377 378 381 382 384 386

Contents xiii 7.6.1 Cathode Materials 7.6.1.1 Classification of Layered Structures 7.6.1.2 O3-Type Layered NaFeO2 7.6.1.3 O3-, P3-, and P2-Type NaxCoO2 7.6.1.4 Sodium Vanadium Phosphate, Na3V2(PO4)3 7.6.1.5 Emerging Cathodes 7.6.2 Anode Materials 7.6.2.1 Carbon-Based Systems 7.6.2.2 Ti-Based Oxide Anodes 7.6.2.3 Alloy Anodes 7.6.3 Room-Temperature Sodium–Sulfur (RT Na–S) Battery 7.6.4 Electrolyte Modification 7.7 Problems in a NIB on Anode Stabilization 7.7.1 Problems with Conductive Additive 7.7.2 Cyclic Voltammetry Study with Conductive Additive 7.7.3 Ex Situ SEM Studies 7.7.4 Solving the Conductive Carbon and Electrolyte Interface 7.8 Few Solutions for Future 7.8.1 In Situ Raman Mapping 7.8.2 In Situ FTIR 7.8.3 In Situ Synchrotron XRD Coupled with DFT Analysis 7.8.4 SIMS-TOF 7.8.5 In Situ TEM Coupled with DFT Analysis 7.8.6 STEM-HAADF and EELS 7.8.7 Time-Lapse Tomography of Volume Expansion 7.9 Perception References 8 Polymer-Based Separators for Lithium-Ion Batteries J. C. Barbosa, C. M. Costa and S. Lanceros-Méndez 8.1 Introduction 8.2 Polymer Types and Characteristics 8.3 Separator Types 8.3.1 Solvent Casting 8.3.2 Electrospun Separator Membranes 8.3.3 Surface Modification 8.3.4 Coating Process 8.3.5 Natural and Biopolymers

387 388 389 391 392 392 393 394 395 396 400 404 405 407 409 410 411 412 413 415 416 417 417 419 420 421 422 429 429 431 433 433 437 441 443 450

xiv

Contents 8.4 Summary and Outlook Acknowledgments List of Symbols and Abbreviations References

Part 3: Supercapacitor: Pseudocapacitor, EDLC 9 Nanostructured Carbon-Based Electrodes for Supercapacitor Applications Sanjit Saha and Tapas Kuila 9.1 Introduction 9.2 Scope of the Chapter 9.3 Charge Storage Mechanism of Carbonaceous Materials 9.4 Nanostructured Carbonaceous Materials 9.4.1 Activated Carbon 9.4.1.1 Activated Carbon as Supercapacitor Electrode 9.4.1.2 Doping of Activated Carbon as Supercapacitor Electrode 9.4.2 Graphene 9.4.2.1 Graphene as Supercapacitor Electrode 9.4.3 Carbon Nano Tube (CNT) 9.4.3.1 CNT Supercapacitor 9.4.3.2 Functionalized CNT Supercapacitor 9.5 Nanostructured Carbon-Based Supercapacitor Device 9.5.1 Carbon-Based Redox Electrode in ASC Device 9.5.2 Carbon-Based Negative (EDLC) Electrode in ASC Device 9.5.3 Different Carbon-Based ASC Device 9.5.4 Carbon-Based Printed Supercapacitor Device 9.6 Conclusions References 10 Nanostructured Metal Oxide, Hydroxide, and Chalcogenide for Supercapacitor Applications Poulomi Roy, Shipra Raj and Suneel Kumar Srivastava 10.1 Introduction 10.2 Materials Architecture and Electrode Designing 10.3 Materials 10.3.1 Metal Hydroxides 10.3.1.1 Mononuclear Metal Hydroxides 10.3.2 Layered Double Hydroxides (LDHs) 10.3.3 Layered Triple Hydroxides (LTHs)

451 452 452 454

467 469 470 471 471 473 475 476 482 483 484 497 498 500 503 504 504 505 506 508 508 521 522 524 526 526 526 532 534

Contents xv 10.4 Metal Oxide 10.4.1 Binary Metal Oxides 10.4.2 Ternary Metal Oxide 10.4.3 Quaternary Metal Oxide 10.5 Metal Chalcogenides 10.5.1 Binary Metal Chalcogenides 10.5.2 Ternary Metal Chalcogenides 10.5.3 Quarternary Metal Chalcogenides 10.6 Summary and Future Perspective References 11 Polymer-Based Flexible Electrodes for Supercapacitor Applications Syam Kandula, Nam Hoon Kim and Joong Hee Lee 11.1 Introduction 11.2 Pure Conducting Polymers (PCs) 11.2.1 Polyaniline (PANI) 11.2.2 Polypyrrole (PPy) 11.2.3 Poly(3,4-ethylenedioxythiophene) (PEDOT) 11.3 Conducting Polymer Composites (CPCs) 11.3.1 PANI-Based Binary Composites 11.3.1.1 PANI- and Carbon-Based Binary Composites 11.3.1.2 PANI and Metal Oxide/Metal Sulfide Based Binary Composites 11.3.1.3 PANI-Based Ternary Composites 11.3.2 PPy-Based Binary Composites 11.3.2.1 PPy- and Carbon-Based Binary Composites 11.3.2.2 PPy and Metal Oxide/Metal SulfideBased Binary Composites 11.3.3 PEDOT-Based Binary Composites 11.4 Conclusions and Perspective References

535 535 542 543 544 545 552 553 555 558 573 574 575 576 577 578 579 580 580 590 593 595 598 609 614 616 619

Part 4: Outlook and Conclusion

625

Outlook and Conclusion

627

Index

629

Preface The rising energy demand and scarcity of fossil fuels are major concerns these days and have drawn much attention to utilizing renewable energy sources and better storage of energy. To address this energy demand, large part of research has been focused in developing high performance energy storage devices, and extensive studies have been carried out on electrodes, electrolytes, as well as separators used in different types of batteries and supercapacitors. The optimal design of electrodes is considered as one of the main routes to achieve high performance of the device. In this regard, nanostructured materials have drawn considerable attention due to their special properties of high surface areas providing large number of active sites for electrochemical action and low charge diffusion path length compared to their bulk counterparts. Various materials, from transition metal oxides, chalcogenides to carbonaceous materials to polymers with their unique properties, tailored morphology, and size can be used as effective electrode materials for different storage mechanism. The present book aims to cover all these aspects and discusses recent achievements and various device assembles. The book comprises three sections in which the first section discusses about the fundamental physics, characterization techniques related to different types of batteries and supercapacitors and the computational view-point to understand the lithium ion battery mechanism. The next section elaborately reviews different types of batteries and their performances reported till date. The third section is focused on supercapacitor based on different nanostructured materials and their performances as reported by various groups globally. With all these components and knowing the importance of the field, it is anticipated that the content will further supplement better understanding and offers current trends of different device assembly to current ongoing works. Dr. Poulomi Roy and Prof. Suneel Kumar Srivastava August 2019 xvii

Part 1 GENERAL INTRODUCTION TO BATTERY AND SUPERCAPACITOR, FUNDAMENTAL PHYSICS CHARACTERIZATION TECHNIQUES

Poulomi Roy and Suneel Kumar Srivastava (eds.) Nanomaterials for Electrochemical Energy Storage Devices, (1–66) © 2020 Scrivener Publishing LLC

1 Electrochemistry of Rechargeable Batteries Beyond Lithium-Based Systems Brij Kishore, Shyama Prasad Mohanty and Munichandraiah Nookala* Department of Inorganic and Physical Chemistry, Indian Institute of Science, Bangalore, India

Abstract The state-of-art, Li-ion battery is the most preferred system among electrochemical energy conversion devices in recent years. The other battery systems based on Li such as Li-S and Li-O2 are either at commercialization or prototype stage. The specific energy of systems based on Li is the greatest among all known battery systems. However, the global raw material resources of Li are limited (0.007% in earth crust), and they are unevenly distributed on the earth. As a consequence, it is likely that there can be lithium crisis in the future affecting the production of Li-based rechargeable batteries for large scale applications such as electrical vehicles. Rechargeable batteries based on Na, Mg and K are expected to be viable substitutes for Li based batteries. Research activities are in progress on several electrode materials, which provide high specific capacity, cycling stability, long cycle-life, etc. The cathode materials for Na-ion batteries include layered sodium transition metal oxides, sodium transition metal polyanions such as phosphates, pyrophosphates, and fluorophosphates. Among the anodes, the most studied materials are hard carbons, low potential transition metal oxides and phosphates, and alloys of Sn, Sb, Ge, etc. The research activities for K-ion batteries are still in infancy, and K-S and K-O2 have attracted attention in recent years. The cathode materials of interest are Prussian green and Prussian blue. Several other materials analogous to Li and Na based materials may soon pick up as research interest. For the anode materials, carbon and potassium titanates are good contenders. The chapter reviews major advances in Li-based systems, detailed studies on electrode materials for emerging Na- and K-based systems, and also Mg-based rechargeable batteries. Keywords: Cathode, anode, electrolyte, Li-based batteries, Na-based batteries, K-based batteries, Mg-based batteries *Corresponding author: [email protected] Poulomi Roy and Suneel Kumar Srivastava (eds.) Nanomaterials for Electrochemical Energy Storage Devices, (3–66) © 2020 Scrivener Publishing LLC

3

4

Nanomaterials for Electrochemical Energy Storage Devices

Ever since the demonstration of Voltaic piles by Alessando Volta in 1800, batteries have been under use for a variety of applications. Their importance has gained rapid strides with the advancement of electronics and also due to dwindling resources for fossil fuels. Although rechargeable batteries such as lead-acid batteries have been widely used over more than 150 years, the recently invented Li-ion batteries have started dominating the battery market. In the present chapter, the research progress made in recent years on future battery systems beyond Li-based batteries is reviewed.

1.1 Lithium-Based Batteries The position of lithium in periodic table and in electrochemical series projects it as an excellent electrode material for battery application. Lithiumbased cells can deliver high voltage (~4 V), high volumetric as well as gravimetric energy density and hence occupy a leading position in usage for portable electronic devices. The evolution of Li metal-based cells started with primary batteries and extended to secondary batteries unsuccessfully. At present, lithium–sulfur and lithium–air batteries employing Li metal are envisioned as future high energy density batteries.

1.1.1 Lithium Primary Batteries A lithium primary battery utilizes Li metal as the anode [1]. The electrochemical reactions are irreversible. Energy density of 200 Wh kg−1 or 400 Wh l−1 can be obtained from Li-based primary batteries. Advantages of such batteries include a wide operating temperature (70 to −40°C) and good shelf life. Depending on the cathode and electrolyte used, the Li primary batteries are classified into different categories. One type is soluble cathode based battery where the cathode is either liquid (SOCl2) or gas (SO2). An organic solvent along with a Li salt serves as the electrolyte. In SOCl2 based battery, SOCl2 performs the role of solvent and a porous carbon at cathode provides the framework for electrochemical reaction. The following reactions occur, which lead to generation of electrical energy.

2Li + 2SO2 4Li + 2SOCl2

Li2S2O4 4LiCl + S+SO2

(1.1) (1.2)

Another type is solid cathode based battery where materials such as CuO, MnO2, FeS2, (CF)n, etc., have been utilized. Such active materials are combined

Electrochemistry of Rechargeable Batteries

5

with conducting carbon and binder to form the cathode and a Li salt based organic electrolyte is used. The cell reaction with CuO cathode is as follows:

2Li + CuO

Li2O + Cu

(1.3)

Replacing the organic electrolyte with any of the solid lithium ion conducting material forms the solid electrolyte based battery. Materials such as -Al2O3, LiI, AgI, LiPON (Li0.39N0.02O0.47P0.12), polymer electrolyte, etc., have been used.

1.1.2 Lithium Metal-Based Secondary Batteries The pioneering work by Whittingham resulted in fabrication of lithium metal-based rechargeable batteries with a layered compound as the cathode [2]. TiS2, which has a layered structure was utilized as the cathode material. Lithium intercalates into the structure forming lithium titanium disulfide. The cell reaction is represented as follows: Discharge

x Li + TiS2

Charge

LixTiS2

(1.4)

Although the secondary batteries with Li metal were rechargeable initially, they exhibited poor cycle life and safety issues. Lithium metal reacted with the organic electrolyte and dendrites were formed during cycling, which caused capacity fading, decrease in cycling efficiency, increase in cell volume, internal shorting, thermal run away, cell explosion, etc. Consequently, Li metal-based rechargeable cells in non-aqueous electrolytes were not successful in applications.

1.1.3 Polymer Electrolyte-Based Lithium Batteries In order to overcome issues such as volatilization of organic electrolyte, and reaction with Li metal and dendrites growth of Li on cycling, solid polymer electrolytes (SPE) were developed for Li metal-based rechargeable batteries. A SPE is composed of a Li salt and a solvating solid polymer. Typically, a SPE-based cell is composed of Li metal anode, intercalation cathode, and SPE in between. Solid polymer electrolyte also serves as the separator. Composite cathodes are formed using the active material, conducting carbon and SPE. Polyethylene oxide (PEO) with a Li salt such as LiClO4 forms a SPE. Using PEO–LiClO4 electrolyte, Osaka et al. have demonstrated that

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Nanomaterials for Electrochemical Energy Storage Devices

SPE inhibited the growth of dendrites [3]. Electrolyte with proper mechanical strength can impose stack pressure on the lithium anode and inhibit dendrite growth [4]. Solid polymer electrolyte-based cells have been cycled for over 500 cycles at different temperatures [5]. Cells have low selfdischarge and excellent efficiency. Despite of such advantages they possess poor ionic conductivity at ambient temperature. In order to enhance the conductivity, approaches like polymer-in-salt electrolyte or addition of metal oxides such as Al2O3, SiO2, TiO2, etc., to SPEs have been adapted [6]. Gel polymer electrolytes (GPE) are also utilized where an organic solvent is present in the polymer electrolyte as a plasticizer [6]. Such electrolytes improve the ionic conductivity and interfacial contact between electrolyte and electrode. In order to overcome the issue of development of concentration gradient of salt during cycling single Li+ ion conducting polymer electrolytes were also studied [6].

1.1.4 Lithium-Ion Batteries A major breakthrough in the area of Li-based batteries was achieved by replacing Li metal by carbon as the anode resulting in ‘Li-ion battery’. Li-ion battery having petroleum coke and LiCoO2 as anode and cathode, respectively, was commercialized in 1991. Since then, Li-ion batteries have dominated the market for portable electric devices. Typically, a Li-ion cell (Figure 1.1) consists of intercalation compounds as anode and cathode along with organic solvents based electrolytes. To serve as an electrode material, the intercalation compound should have appropriate lithium chemical potential, high electronic

e–

e– –

Discharge

e–

+

Load

Charge

e–

Charge Anode

Cathode

Li+

Discharge

Li+

LixC6

Electrolyte

Li1-xCoO2

Figure 1.1 Schematic of Li-ion cell. Reprinted with permission [8].

Electrochemistry of Rechargeable Batteries

7

conductivity and stability (thermal and chemical). Operating principle of a Li-ion cell with graphite anode and LiCoO2 cathode is as follows: Discharge

At cathode, x Li+ + Li1–xCoO2 + xe–

Charge

LiCoO2

(1.5)

Discharge

At anode, LiC6

Charge

Li1–xC6 + xLi+ + xe–

(1.6)

Discharge

Overall cell reaction, LiC6 + Li1–xCoO2

Charge

Li1–xC6 + LiCoO2

(1.7) Most of the cathode materials are compounds of transition metals.On removal of Li+ ion, the transition metal ions get oxidized to higher oxidation state while retaining the crystal structure. LiCoO2, a layered material was demonstrated by Goodenough et al. in 1980 [7]. It has -NaFeO2 rock salt structure with Li+ and Co3+ ions occupying alternate layers of the structure. Up to 0.5 mol of Li can be extracted from LiCoO2 without causing structural instability. The material is expensive due to the presence of Co and it suffers from structural instability with an increase in charging voltage. Dissolution of Co into electrolyte is also another issue. LiNiO2 is isostructural with LiCoO2, less expensive and has higher energy density but with poor stability because Ni gets displaced into the Li layer on cycling and affects lithiation/de-lithiation process [9]. LiMnO2 is another layered material with monoclinic crystal structure [9]. De-lithiation results in the formation of spinel LiMn2O4 which results in capacity decay. LiMn2O4 has been explored as low cost material but lithium insertion/de-insertion at 3 V leads to poor cycle life due to Jahn– Teller effect of Mn3+ [9]. LiNi1/3Mn1/3Co1/3O2 or NCM333 is isostructural to LiCoO2, the role of Ni, Co, and Mn are to deliver high capacity, cycle stability, and thermal stability, respectively. Stable capacity of 150 mAh g−1 has been obtained in NCM333 based cell operating between 2.5 and 4.2 V [10]. Olivine LiFePO4 was first reported by Goodenough et al., which possessed high capacity, thermal stability and low cost [11]. Unfortunately, the material has low electrical conductivity (10−9 S cm−1) and poor lithium diffusivity. Coating of a conducting layer on the particle and reduction in particle size have been attempted to overcome these issues. Lithium metal as the anode is the ultimate choice due to its high capacity. However, the issues like dendrites formation, reaction with electrolyte and poor stability are present in Li metal based cells. Graphite is a widely accepted anode material for Li-ion batteries. Alloys of lithium with Al, Si,

8

Nanomaterials for Electrochemical Energy Storage Devices

Sn, etc., have been studied as anode materials. Volume expansion during lithiation leads to the formation of cracks and loss of active material in such alloys based cells [9]. Amorphous silicon has shown promising results as an anode material [9]. Among the electrolytes, LiPF6 in a mixture of carbonate based solvents is the most common for Li-ion batteries. The solvent plays an important role in formation of a stable solid electrolyte interface (SEI) and the effect of additives to electrolyte has been studied to improve the properties of SEI. Several review articles on Li-on batteries are available which can provide deeper understanding of the system along with developments [9, 10].

1.1.5 Advances in Li-Ion Batteries Materials capable of storing more than one lithium per transition metal atom are categorized as ‘lithium rich compounds’. Research activity in this field has gained momentum as such materials can deliver high discharge capacity and they are considered as the next generation electrode materials. Li2MnO3 is one such material with layered structure similar to LiCoO2 [9]. Due to the inability of Mn4+ to oxidize to Mn5+, the compound is electrochemically inactive and provides lower capacity than expected. Partial substitution of lithium by proton has resulted in a high initial discharge capacity but with poor cycle life. Solid solutions of Li2MnO3 with LiMO2 (M=Mn, Ni, Co or Fe) are of great interest as high discharge capacity and cycling stability are achieved at a high voltage [9, 10]. Enhancement in energy density of Li-based cells is possible by replacement of the insertion-type cathode with conversion-type cathode. Sulfur and oxygen are the two electrode materials in this regard, which can be combined with Li resulting in high energy Li-S and Li-O2 systems, respectively. Sulfur has a high theoretical capacity of 1672 mAh g−1 in Li-S cells with Li2S as the end product of the discharge process [12]. In addition, sulfur is cheap and abundant. A typical Li-S cell consists of a sulfur cathode, Li foil as the anode and an organic electrolyte containing Li salt. Sulfur based cathode consists of sulfur or sulfur composite, conducting carbon and a binder on aluminum substrate. The average cell voltage is 2.15 V and theoretical gravimetric and volumetric energy densities are 2500 Wh kg−1 and 2800 Wh l−1, respectively [12]. Lithium-air system has a very high theoretical energy density of 11680 Wh kg−1, which projects it as a potential competitor to gasoline [13]. Lithium–air cell consists of Li metal as the anode, electrolyte, and a porous cathode, which would allow the reduction of O2. The following reactions occur during the discharge–charge processes.

Electrochemistry of Rechargeable Batteries

2Li + O2 Li2O2

Li2O2 (Discharge) 2Li + O2 (Charge)

9

(1.8) (1.9)

1.1.6 Beyond Lithium-Based Systems Since the time of commercialization, Li-ion batteries have dominated for portable devices in consumer market. The need for electric or hybrid electric vehicles in near future has prompted to search for storage systems with energy density greater than that of Li-ion system. Hence, Li–S and Li–air systems are expected to be potential candidates, if stability issues in these systems are solved. Cost and abundance of Li are concerns, which would inhibit bulk production and utilization of Li-based batteries at a reasonable cost. Only 20 ppm Li is available in the earth crust and it is unevenly distributed across the globe with South America being rich [14]. In contrast, each of Na, K, and Mg are more than 2% in the earth crust and also widely distributed, hence, they are cheap (Table 1.1). Standard electrode potentials of Na and K are close to Li [14]. Hence, Na and K are suitable in alkali metal ion batteries. Mg has advantage of high melting point, higher volumetric capacity than Li and also its ionic radius is comparable to Li [14]. Even though the standard electrode potential of Mg is less than Li, the advantages project Mg as a cheap and potential alternative to Li in batteries. Na, K, and Mg based batteries are discussed in detail in the following sections.

1.2 Cathodes for Na-Ion Batteries A wide range of compounds are studied as cathode materials, namely, oxides, compounds with polyanions such as phosphates, pyrophosphates, fluorophosphates, carbonophosphates, fluorides, hexacyanoferrates, and more recently organic polymers for Na-ion batteries (Figure 1.2). Each type of compounds has its own set of advantages and disadvantages. The following sections provide some details of these compounds.

1.2.1 Transition Metal Oxides Layered LiMO2 (M = V, Cr, Mn, Fe, Co, Ni) compounds have been extensively studied for Li-ion batteries. Following a similar trend, research activities on NaMO2 compounds have gained attention, in particular, on viable Na-intercalation cathodes. However, Li+ ion being smaller in size (0.76 Å),

Lithium (Li) 3 6.941 181 1347 0.53 157 90 340 5.392

1 0.2 0.007 3860

Properties

Atomic number

Mass number

Melting point (°C)

Boiling point (°C)

Density (g/cc)

Atomic radius (pm)

Crystal radius (pm)

Hydrated radius (pm)

Ionization potential (eV)

Oxidation state

Abundance in sea water (ppm)

Abundance in earth crust (%)

Energy density (mAh g−1)

Table 1.1 Properties of Li, Na, K, and Mg.

1165

2.6

10800

1

5.139

276

116

191

0.97

883

98

22.990

11

Sodium (Na)

687

2.4

400

1

4.341

232

152

235

0.86

774

64

39.098

19

Potassium (K)

2205

2.4

1300

2

2nd: 15.035

1st: 7.646

418

86

160

1.74

1091

650

24.305

12

Magnesium (Mg)

10 Nanomaterials for Electrochemical Energy Storage Devices

Electrochemistry of Rechargeable Batteries 5

Na3NiZr(PO4)3 α–NaFeO2 NaLi0.2Ni0.25Mn0.75O2 Na3V2(PO4)3 P2-Na2/3Ni1/3Mn2/3O2

Voltage (V vs. Na+/Na)

4

aniline-nitroaniline copolymer

Na4Fe(CN)6/C

3

2

NaMnO2 NaFePO4 NaMn0.44O2

P2-NaxCoO2 Na1-xNi0.5Mn0.5O2 NaTi2(PO4)3

Carbon-based materials

0

100

FeS

Sn/C Na2Ti3O7 Na2C8H4O4

Li4Ti5O12

1

0

11

200

SnSb/C Sb/C

300 400 500 600 Capacity (mAh/g)

P/C Sb2O4

700

800 1800

2000

Figure 1.2 The relationship between capacity and voltage for electrode materials for Na-ion batteries. Reprinted with permission [15]. MeO2 Layer Na ions

A B

A B

C A B C

c

b

B C C A

(1/3, 2/3, 0)

A B

a

A B

O3 type

P3 type

oxygen

Octahedral site

O2 type

A B

A B

A C

B A

(1/3, 2/3, 0)

A B

A B

A C

B A

Prismatic site

P2 type

Figure 1.3 Classification of Na-M-O layered materials with sheets of edge-sharing MO6 octahedra and phase transition processes induced by sodium extraction. Reprinted with permission [17].

12

Nanomaterials for Electrochemical Energy Storage Devices

usually occupies the octahedral site in the crystal lattice whereas the same is not entirely applicable for a larger Na+ ion (1.02 Å) [16]. Depending on the occupancy of alkali metal ion, the crystal structure can be On, (n = 1, 2, 3) where Na is octahedrally coordinated by oxygen, and n refers to the repeated period of transition metal stacking within a unit cell (Figure 1.3). Na metal can also occupy trigonal prismatic sites and these structures are noted as Pn. Phase transition can occur between On and Pn via gliding of MO2 sheets, which has been observed to occur at room temperature [17]. The P2 phase is generally accepted as a better cathode material because of high diffusion rate of Na+ ion and prohibited slab-gliding. Here, Na occupies the trigonal prismatic sites of ABBA oxygen stacking sequence. One of the most significant efforts to improve the electrochemical performance of layered oxides is to stabilize the inter slab space by partial metal substitution.

1.2.1.1 Single Metal Oxides The research for Na layered oxides of Mn, Ni, Cr, and Co dates back to 1980s [18]. NaxCoO2 bronze was studied as a cathode material for Na-ion battery [19]. Depending on the value of x in the range 0.4 ≤ x ≤ 1.0, NaxCoO2 is known to exist in four different phases, 0.55 ≤ x ≤ 0.60 (P’3); 0.64 ≤ x ≤ 0.74 (P2); x = 0.77 (O’3) and x = 1 (O3). P2 bronze provided better cycle life and energy efficiency, and it was first studied by Shacklette et al. [20]. It exhibited intercalation/de-intercalation of 0.6 Na and offered a theoretical energy density of 270 Wh kg−1 when charged to 3.5 V and 440 Wh kg−1 when charged to 4.0 V, based on a discharge composition of Na0.93CoO1.965. NaCrO2 and NaNiO2 electrodes showed only a small amount of reversible insertion/de-insertion ( 0.2 Na) resulting in low specific capacity [18]. Recently, however, when NaCrO2 was revisited by Komaba et al., a higher reversible capacity of 120 mAh g−1 ( 0.5 Na) was obtained with the average voltage around 3.0 V vs. Na+/Na [21]. A high discharge capacity of 90–120 mAh g−1 was maintained over 50 cycles. Sodium insertion in three different vanadium oxides (a channel structure β-NaxV2O5, and two layered structures (α-V2O5 and Na1+xV3O8)) was studied by West et al. [22]. Both Na1+xV3O8 and β-NaxV2O5 degraded gradually upon cycling, whereas α-V2O5 changed to a new phase after first discharge and exhibited excellent capacity retention thereafter upon cycling. O3-NaVO2 electrodes studied by Didier et al. delivered a specific capacity of 120 mAh g−1 corresponding to 0.5 Na reversible at 1.4–2.5 V [23]. Single crystalline bilayered V2O5 synthesized by solvothermal method was recently reported with a nanobelt morphology [24]. At a specific current of 80 mA g−1, the first discharge capacity of 206 mAh g−1 with a voltage plateau

Electrochemistry of Rechargeable Batteries

13

of 2.4 V was obtained. Based on the high discharge capacity, V2O5 was anticipated to uptake approximately 2 Na+ to form Na2V2O5. The electrode delivered a discharge capacity of 170 mAh g−1 after 100 cycles. Breaking the general notion that highly crystalline materials are preferred for cathode, Uchaker et al. recently demonstrated that amorphous V2O5 synthesized by a combination of sol–gel process and electrochemical deposition exhibited better electrochemical properties than its crystalline counterpart [25]. When tested at C/10 rate, amorphous V2O5 showed a discharge capacity of 241 mAh g−1, whereas its crystalline form showed 120 mAh g−1. Nanowire morphology of NaV3O8 was prepared by hydrothermal method which contained water of crystallization [26]. The sample was annealed at 400°C for 3 h in air to remove the intercalated water molecules. The as prepared sample with water of crystallization exhibited a higher first reversible capacity (173 mAh g−1) than the annealed sample (146 mAh g−1) at a specific current of 10 mA g−1. Although the annealed sample showed low discharge capacity, the cycle life was much better (91% retention after 50 cycles) than the as prepared sample (52% retention). The improved cycle life was attributed to contraction of cell volume by removal of water of crystallinity during annealing. Yabuuchi et al. revisited α-NaFeO2 and showed that during galvanostatic charge–discharge cycling, when the upper cut of 3.4 V vs. Na+/Na was maintained, the cell exhibited good capacity retention with reversible specific capacity of 80 mAh g−1 [27]. On increasing the upper voltage limit, the charge capacity increased but the discharge capacity decreased significantly. The XRD of the cycled electrodes revealed significant irreversible structural changes caused due to the migration of Fe into the inter slab spacing at high voltage. Manganese based cathodes are popular due to high theoretical capacity and low cost. These oxides are known to crystallize in different polymorphs depending on the way MnO6 octahedra are inter-linked [28]. In α-MnO2, these octahedra are linked at the corners to form (2 X 2) and (1 X 1) channels, which facilitate diffusion along the c-axis. Su et al. synthesized α- and β-MnO2 nanorods by hydrothermal route [29]. The β-MnO2 delivered a higher discharge capacity (298 mAh g−1) than α-MnO2 (278 mAh g−1) at 20 mA g−1. Kishore et al. synthesized amorphous MnO2 by redox reaction between Mn7+ and Mn2+ at room temperature [30]. They annealed the sample at various temperatures to get highly crystalline α-MnO2. The as prepared sample had a specific surface area of 184 m2 g−1 and mesopores of 3.5 nm. The amorphous sample delivered first discharge capacity of 194 mAh g−1 at 50 mA g−1, whereas the highly crystalline sample (heated at 800°C) showed only 22 mAh g−1. The cycle life was poor for the as prepared sample

14

Nanomaterials for Electrochemical Energy Storage Devices

whereas the crystalline sample showed good cyclability. Recently, Yuan et al. studied the dynamics of sodiation/de-sodiation by in situ TEM and in situ synchroton X-ray absorption spectroscopy in α-MnO2 [31]. During the first sodiation, α-MnO2 survives the Na+ insertion in its tunnel structure when Mn4+ reduces to Mn3.5+. It is followed by fast tunnel degradation with the formation of intermediate phase Na0.5MnO2 and final conversion to Mn2O3 and Na2O. During the first de-sodiation and subsequent cycles, partial conversion reaction occurs between Na0.5MnO2 and Mn2O3. During the subsequent cycles of sodiation, a strong destructive effect due to bigger Na+ size and strong interaction with the host structure lead to fast degradation of tunnels and complete phase conversion to Mn2O3. Pre-sodiated NaxMnO2 bronze was first studied by Doeff et al. for secondary Na and Li polymer battery cathodes [32]. A reversible intercalation up to 0.6 Na+ or Li+ per Mn at moderate current densities of 0.1–0.05 mA cm−2 delivered specific capacity of 160–180 mAh g−1 in a SPE at 85°C. Ma et al. studied low temperature α-NaMnO2, with O3 layered monoclinic structural distortion, which intercalated 0.8 Na providing a discharge capacity of 185 mAh g−1 at C/10 rate with 71% retention after 20 cycles [33]. By ex situ XRD studies of pristine α-NaMnO2 and partially charged NaMnO2, they confirmed that a two phase transformation occurred due to Na vacancy ordering or modification of O3 stacking by oxygen layer. Similarly, β-NaMnO2 was examined by Billaud et al., which had orthorhombic structure containing MnO2 sheets [34]. A high reversible capacity of 190 mAh g−1 was obtained at C/20 rate with good rate capability. When a cell was cycled at 2C rate, it offered the first discharge capacity of 142 mAh g−1 with 70 % retention after 100 cycles. Single crystalline Na0.7MnO2 nanoplates synthesized by hydrothermal route was investigated by Su et al., and it delivered a reversible capacity of 163 mAh g−1 with satisfactory cycle life and high rate capability [35]. The improved performance was attributed to (100) facet, which facilitated fast Na+ intercalation/de-intercalation during discharge/charge process. Pre-sodiated manganese oxide Na0.44MnO2 (also known as Na4Mn9O18) was investigated because of its large-sized tunnel for Na+ insertion [36]. It crystallized into orthorhombic lattice with two different environments for Mn ions. All Mn4+, half of Mn3+ were in octahedral sites and the other half of Mn3+ in square pyramidal environment. The interlinking of them resulted in two types of tunnels where Na+ ions were present. Na1 and Na2 were present in large S-shaped tunnels, and Na3 was found in smaller tunnel as shown in Figure 1.4. Sauvage et al. prepared pure Na0.44MnO2 sample by solid-state method, which exhibited a reversible capacity of 80 mAh g−1 at C/10 rate in

Electrochemistry of Rechargeable Batteries

15

Na3 Na2 Na1

a

c b

Figure 1.4 Structure of Na0.44MnO2 perpendicular to the ab plane. Reprinted with permission [36].

the voltage range 2.0–3.8 V vs. Na+/Na [36]. They also concluded that only Na1 and Na2 present in large S-shaped tunnels were reversibly accessible. Cao et al. synthesized single crystalline Na4Mn9O18 nanowires by a polymer-pyrolysis method [37]. The nanowire electrode material after calcination at 750°C delivered a discharge capacity of 128 mAh g−1 at C/10 rate with capacity retention of 77% after 1000 cycles at C/2 rate. The material was capable of inserting and extracting four Na+ ions during cycling inducing a chemical transformation from Na6Mn9O18 to Na2Mn9O18. The slightly higher value than the theoretical capacity 121 mAh g−1 suggested that Na+ ions occupying the smaller tunnels could also be extracted and contributed to reversible capacity. Recently, Dai et al. synthesized rod like Na0.44MnO2 via polyvinylpyrrolidone-combustion method, and the sample synthesized at 900°C showed specific discharge capacity of 123 mAh g−1 at C/5 rate [38]. The capacity retention after 100 cycles was 88% at 1C rate and remained at 83% when the cell was cycled for another 700 cycles at 10C rate. The reversible capacity at 20C rate was 99 mAh g−1. This high capacity at exceptionally high C rate was attributed to good solid-state diffusion of Na+ ion, and the diffusion coefficient was 3 x 10−12 cm2 s−1 as calculated from potential intermittent titration technique.

16

Nanomaterials for Electrochemical Energy Storage Devices

The influence of morphology on the performance of layered NaxMnO2+Z was studied by Bucher et al. [39]. Hollow sphere and flake like NaxMnO2+Z were synthesized. At a specific current of 50 mA g−1, the microspheres showed a reversible capacity of 94 mAh g−1 after 100 cycles, whereas flakes delivered 73 mAh g−1. The good cycle life was attributed to the better accommodation of volume changes during insertion/de-insertion by microsphere morphology, better contact with added carbon and good electrolyte wettability leading to higher electrode/electrolyte contact.

1.2.1.2 Multi-Metal Oxides Multi-metal oxides are interesting to get high capacity, fairly flat voltage profile and higher operating voltage by utilizing various properties of individual transition metals. In some cases, various metals such as Li, Mg, Ca, Ti, Te, etc., are doped, which occupy the transition metal layer and provide structural stability during insertion/extraction. Yabuuchi et al. reported a mixed transition metal oxide comprising of Fe and Mn, i.e., P2-Na2/3Mn1/2Fe1/2O2 and O3-Na2/3Mn1/2Fe1/2O2 synthesized by using solid-state method [40]. The P2 type delivered a reversible capacity of 190 mAh g−1 at C/20 rate with electrochemically active Fe3+/Fe4+ redox couple, whereas O3 type delivered 100–110 mAh g−1. Komaba et al. reported O3-NaMn0.5Ni0.5O2, which delivered a discharge capacity of 185 mAh g−1 between 2.0 and 4.5 V, but reversibility was poor due to significant expansion of inter slab space [41]. On changing the voltage window to 2.2–3.8 V, a reversible capacity of 105 mAh g−1 at 1C rate and 125 mAh g−1 at C/50 rate with 75% retention after 50 cycles were observed. Ex situ XRD studies revealed the existence of several phases O3, O’3, P3, P’3, and P’’3 during cycling and X-ray absorption spectroscopy confirmed that Mn4+ was inactive and Ni2+ contributed to capacity. Cu substituted P2-Na0.67CuxMn1-xO2 (x = 0, 0.14, 0.25, 0.33) was prepared by Kang et al. via sol–gel method [42]. The increase in Cu content resulted in a decrease in capacity and an increase in potential due to a shift in the reaction from Mn3+/Mn4+ to Cu2+/Cu3+ couple. Cycle life and rate performance were also enhanced by Cu substitution as it stabilized the crystal structure. Capacity of more than 90 mAh g−1 was obtained at 12C rate with 70% retention after 500 cycles. Substitution of Na, with Ca (NaxCayCoO2) was studied by Han et al., where bivalent Ca occupied some of the prismatic sites of Na [43]. This resulted in lattice contraction because of small size of Ca2+ ions leading to enhanced structural stability and suppression of abrupt phase transitions during cycling. Na0.60Ca0.07CoO2 showed better cycle life performance as the capacity retention was 96% after 60 cycles at C/10 rate.

Electrochemistry of Rechargeable Batteries

17

Al doping was carried out along with transition metal mixing of Mn and Ni to form P2-Na0.60[Ni0.22Al0.11Mn0.66]O2 [44]. The material when subjected to electrochemical investigation at C/10 and 5C rates delivered reversible capacity of 250 and 150 mAh g−1, respectively, in the voltage range 1.5– 4.6 V. The values reported here are higher than that of Hasa et al., where a similar composition was reported except for Al, and the metal was Fe [45]. At a specific current of 15 mAh g−1, the P2-Na0.60[Ni0.22Fe0.11Mn0.66]O2 delivered 234 mAh g−1 reversible capacity. Wang et al. prepared Na[Fe1/3Ni1/3Ti1/3]O2 by solid-state method, which delivered a discharge capacity of 117 mAh g−1 with good rate capability and cycle life at a specific current of 10 mA g−1 [46]. In situ XRD studies suggested that the phase changed from hexagonal O3 to hexagonal P3 during charging and ex situ X-ray absorption spectroscopy concluded that both Fe2+/3+ and Ni2+/4+ were active redox centers. Na analogs of commercial Li[Ni1/3Mn1/3Co1/3]O2 were evaluated by Sathiya et al. [47]. The material could cycle only 0.5 Na delivering a capacity of 120 mAh g−1 at C/10 rate in the voltage range 2.00–3.75 V. Several phase transformations and solid solutions regions were identified by in situ XRD studies. The structural evolution during de-intercalation followed the sequence O3 O1 P3 P1 with an increase in lattice parameter ‘c’. Similarly, P2-Na2/3[Mn1/3Fe1/3Co1/3]O2 synthesized by solid-state method delivered a reversible capacity of 173 mAh g−1 with an average voltage of 3.0 V in the potential range 1.5–4.5 V [48]. Yuan et al. studied P2-Na0.67Mn0.65Fe0.35-xNixO2 prepared by sol–gel method [49]. For the sample with x = 0, the material Na0.67Mn0.65Fe0.35O2 exhibited a discharge capacity of 204 mAh g−1 at C/20 rate. When a small amount of Ni was added (x = 0.15), Na0.67Mn0.65Fe0.20Ni0.15O2 delivered a discharge capacity of 208 mAh g−1 with improved cycling stability, which was attributed to Ni substitution which alleviated the Jahn–Teller distortion of Mn3+ ions. Oh et al. studied the effect of Li substitution by comparing Na  (Ni0.25Fe0.25Mn0.50)2 and Na[Li0.05(Ni0.25Fe0.25Mn0.50)0.95]O2 synthesized by co-precipitation method [50]. The Li-substituted sample exhibited a reversible capacity of 180 mAh g−1 at C/10 rate. Li helped to stabilize the structure by arresting the migration of Fe3+ from the transition metal layer to Na-layer during cycling. A full cell study with hard carbon anode and Li-substituted cathode delivered a capacity of 177 mAh g−1 (based on cathode mass) at C/10 rate with 76% capacity retention after 200 cycles at C/2 rate.

1.2.2 Polyanionic Compounds Polyanionic compounds have been extensively investigated as they provide several advantages [51, 52]. The crystal structures have open channels

18

Nanomaterials for Electrochemical Energy Storage Devices

for alkali metal ion to diffuse in and out. The operational voltage can be easily influenced by tailoring the specific redox couple by local polyanion environment. These compounds have high thermal stability due to strong covalent bonding of oxygen atoms in polyanionic polyhedral. NaFePO4 is an interesting cathode with the theoretical capacity of 154 mAh g−1 and an operating voltage of 2.9 V vs. Na+/Na based on the single electron reaction of Fe3+/Fe2+ couple [53, 54]. The crystal structure consists of a hexagonal close pack of oxygen array where Na and Fe atoms occupy half of octahedral sites and P atoms one-eighth of tetrahedral sites. The difference in ionic size and charge of Na and Fe leads to either olivine phase or maricite phase (thermodynamically stable). Olivine phase is electrochemically active, but the practical capacity is much lower due to poor electron conductivity and 1D diffusion channel. Moreover, olivine NaFePO4 shows a large non-reversible charge/discharge process due to cell mismatch between NaFePO4 and FePO4. Maricite NaFePO4 is regarded as electrochemically inactive due to the lack of cationic transport channel. Recently Kim et al. showed that nanoparticles could be active due to enhancement of Na mobility by structural transformation [54]. They prepared it by solid-state method followed by ball milling with carbon to get 50 nm particles coated with carbon to enhance the electronic conductivity. A discharge capacity of 142 mAh g−1 was reported at C/20 rate and a capacity retention of 95% was obtained after 200 cycles. Even hollow amorphous NaFePO4 nanospheres reported by Li et al. exhibited a specific discharge capacity of 152 mAh g−1 at C/10 rate, which retained a capacity of 144 mAh g−1 after 300 cycles [55]. NASICON, a Na super ionic conductor, was originally studied as a solid electrolyte that allows fast Na-ion conduction through the empty spaces in its crystal structure for high temperature Na-S batteries [56]. The presence of strong framework with large Na diffusion channel has prompted researchers to exploit it as Na-ion battery cathode. In Na3V2(PO4)3, PO4 tetrahedra share corners with VO6 octahedra to form 3D skeleton of (V2P3O12) unit with two different oxygen environments for Na atom, Na(1), which is sixfold coordinated and Na(2), which is eightfold coordinated [57]. Only two-third of Na+ ions can be extracted resulting in a theoretical capacity of 117 mAh g−1. Electrochemical performance of Na3V2(PO4)3 can be improved by carbon coating or embedding it in a carbon matrix. Saravanan et al. prepared carbon-coated Na3V2(PO4)3 with 6 wt% carbon and reported a reversible capacity of 114 mAh g−1 at 1C rate, as well as long term cycle life with 50% capacity retention after 30,000 cycles at 40C rate [58]. Cationic substitution was also carried out to enhance the electrochemical performance. Partial substitution of V with Fe, Mg, or Al

Electrochemistry of Rechargeable Batteries

19

resulted in improvement in performance due to the effective promotion in the oxidation of V to the pentavalent state or due to enhancement of ionic and electronic conductivity [59–61]. Pyrophosphates Na2MP2O7 (M = Fe, Mn, Co) can crystallize into different polymorphic forms such as triclinic, tetragonal, and orthorhombic structures depending on the type of transition metal and synthesis conditions [62]. In all these polymorphs, there exists tunnels (1D for triclinic and 3D for tetragonal/orthorhombic) containing Na+ ions with good Na+ mobility and structural stability owing to the presence of P2O74 anions, which make them promising materials for Na-ion battery cathode [62]. Triclinic Na2FeP2O7 and Na2MnP2O7 have been studied by Barpanda et al., and reversible capacities of 80–90 mAh g−1 have been demonstrated [63, 64]. The operating voltage was higher for Na2MnP2O7 (3.6 V) than for Na2FeP2O7 (3.0 V). Orthorhombic Na2CoP2O7 delivered a reversible capacity of 80 mAh g−1 and an average operating voltage of 3.0 V vs. Na+/ Na [65]. Tetragonal Na2(VO)P2O7 synthesized using solid-state route delivered a discharge capacity of 80 mAh g−1 with an average operating voltage of 3.8 V vs. Na+/Na based on V5+/V4+ redox chemistry [66]. Na4M3(PO4)2P2O7 is an interesting polyanion-based compound with mixed phosphate and pyrophosphate groups as it has four different Na+ sites located in a 3D ion channel. As a cathode material, Na4Co3(PO4)2P2O7 delivered a de-sodiation capacity of 95 mAh g−1 with negligibly small capacity fade after 100 charge–discharge cycles with redox potential in the range of 4.1–4.7 V vs. Na+/Na [67]. When the material was doped with a small amount of Mn and Ni (Na4Co2.4Mn0.3Ni0.3(PO4)2P2O7), it showed a discharge capacity of 103 mAh g−1 at a specific current of 850 mA g−1 (5C rate) [68]. This was because all three transition metals participated in the redox process, and hence a high operating voltage of 4.5 V vs. Na+/Na was observed. Carbonophosphates are class of compounds with general formula A3M(CO3)(PO4) (A = Li, Na and M = Co, Mn, Ni, Fe) and sidorenkite structure. The MO6 octahedra and PO4 tetrahedra share each corner, and CO3 triangular planer groups share an oxygen edge with the MO6 octahedra. This forms a double layer, which accommodates Na atoms at two different interstitial sites. Na(1) sites coordinate with seven oxygen atoms, and Na(2) sites coordinate with six oxygen atoms [69, 70]. Chen et al. investigated Na3MnCO3PO4 and reported a discharge capacity of 125 mAh g−1 at C/100 rate corresponding to insertion of 1.3 Na per formula against the theoretical capacity 191 mAh g−1 [69]. Recently, Wang et al. reported a discharge capacity of 176 mAh g−1, reaching 92% of theoretical capacity [70]. The increased capacity was attributed to an increase in use of carbon

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black (60 vol%), which provided a continuous network for utilization of all Na3MnCO3PO4 particles by overcoming electronic resistance. Fluorophosphates are classified broadly into two types depending on their composition, Na2MPO4F (M = Fe, Co, Mn) and Na3(VOx)2(PO4)2F3-x (O ≤ x ≤ 1). Sodium–iron/sodium–cobalt fluorophosphates possess a 2D layered structure whereas sodium–manganese fluorophosphate has a 3D tunnel structure [71, 72]. This leads to different electrochemical properties where iron and cobalt phases exchange one electron at 3.0 and 4.0 V vs. Na+/ Na, whereas manganese compound has shown poor electrochemical activity. Pure phase of Na2CoPO4F/C nanocomposite was synthesized by Zou et al. by spray drying and high temperature sintering which delivered a discharge capacity of 107 mAh g−1 with voltage plateau at 4.3 V vs. Na+/Na [71]. Komaba et al. prepared carbon coated Na2FePO4F and Na2Fe0.5Mn0.5PO4F by solid-state method where ascorbic acid was used as the carbon source [73]. The parent sample (without Mn) with 1.3 wt% carbon delivered a discharge capacity of 110 mAh g−1 at C/20 rate with two well defined plateaus at 3.06 and 2.91 V due to the formation of intermediate Na1.5FePO4F phase. In contrast, Na2Fe0.5Mn0.5PO4F was electrochemically less active, and ball milling with 6 wt% ascorbic acid as carbon source reduced the particle size and provided good carbon coating. This optimized sample delivered a discharge capacity of 110 mAh g−1 at C/20 rate. The interesting feature was the presence of three distinct plateaus, two plateaus due to iron fluorophosphate and another one assigned to Mn2+/Mn3+ redox pair. The average operating voltage was 3.53 V vs. Na+/Na. Lin et al. synthesized Na2MnPO4F/C nanocomposite via spray drying followed by a high temperature sintering [74]. The material when tested at specific current of 6.2 mA g−1 at 30°C, the first discharge capacity of 106 mAh g−1 was obtained, but the capacity decreased to 63 mAh g−1 after 50 cycles. However, at 55°C, a discharge capacity of 178 mAh g−1 was obtained with 75% capacity retention after 20 cycles. NaVPO4F synthesized by sintering NaF and VPO4 at 750°C was reported by Barker et al. as a cathode material for a 3.7 V Na-ion cell with a hard carbon anode [75]. The material crystallized into a tetragonal structure with I4/mmm space group, which was later suggested as a monoclinic structure with C2/C space group by Zhuo et al. [76]. Recently, graphene modified NaVPO4F was revisited by Ruan et al., which showed a reversible capacity of 121 mAh g−1 with 98% capacity retention after 50 cycles at C/20 rate [77]. Na3V2(PO4)2F3 is structurally made up of [V2O8F3] bi-octahedral and [PO4] tetrahedral units [72]. The oxygen sharing between these two units leads to the formation of channels along a and b directions where Na is located. Chihara et al. studied Na3V2(PO4)2F3 and obtained a specific

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discharge capacity of 95 mAh g−1 at 4C rate with two distinct plateaus at 3.65 and 4.10 V vs. Na+/Na [78]. It was also observed that only two-thirds of the Na+ ions could intercalate/de-intercalate owing to the presence of thermodynamically stable intermediate phase of NaV2(PO4)2F3. Recently Jin et al. prepared carbon wrapped multilayer Na3(VO)2(PO4)2F nanocubes embedded in graphene using hydrothermal method followed by annealing in Ar atmosphere using sucrose as the carbon source [79]. A discharge capacity of 137 mAh g−1 was obtained at C/20 rate, and plateaus were observed at 3.5 and 3.9 V vs. Na+/Na. When fluorine in Na3V2(PO4)F3 was partially substituted by oxygen, the lattice parameters and cell volume of Na3(VO)x(PO4)2F3-x decreased due to lattice shrinkage as V4+ are replaced by larger V3+ ions [80]. Hence, the charge/discharge plateaus decreased due to weakened inductive effect from fluorine.

1.2.3 Fluorides Dimov et al. investigated NaMF3 (M = Mn, Fe, Ni) series as fluorine based cathodes for Na-ion batteries and found only NaFeF3 as electrochemically active [81]. Fe-F bonds in FeF3 are strongly polarized and electrons are therefore localized in FeF3, and hence the charged species have very low electrical conductivity. FeF6 octahedra share all corners, forming a framework structure with open channels for Na intercalation/de-intercalation. Yamada et al. synthesized NaFeF3 of different particle sizes (10–600 nm) using liquid-phase synthesis [82]. They observed a dependency of discharge capacity on particle size at a current rate of C/10 or higher, while the discharge capacities obtained at C/100 rate between 1.5 and 4.5 V vs. Na+/Na were almost the same in 170–180 mAh g–1 range regardless of the particle size of NaFeF3.

1.2.4 Metal Hexacyanometalates Metal hexacyanometalates AxMM’(CN)6 (A = Na, K; M and M’ = Mn, Fe, Co, Ni) have cubic structure with metal ions (M or M’) situated at the corners of cube and cyanide groups bridging along the cube edges [83, 84]. The metal ions are octahedrally coordinated by the nitrogen or carbon end of the cyanide group. The structure allows reversible insertion/de-insertion of alkali metal ions due to the open cubic framework and favorable interstitial sites. These can be synthesized easily using a solution method from low-cost cations sources. Goodenough et al. synthesized Prussian blue (PB) and its analogs with different transition metals Fe, Mn, Ni, Cu, and Zn, at room temperature

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by solution method [83]. They demonstrated a stable reversible capacity of 100 mAh g−1 for 30 cycles for KFe2(CN)6. They further extended the work where K+ ion was exchanged with Na+ ion in PB and also replaced one of the Fe atoms with Mn leading to the formation of two sodium manganese hexacyanoferrates with different structures, rhombohedral Na1.72MnFe(CN)6 and cubic Na1.4MnFe(CN)6 [84]. Both the materials showed a reversible specific capacity greater than 120 mAh g−1 at C/20 rate. The rhombohedral structure showed fast capacity fade for initial few cycles and then stabilized, whereas cubic structure has negligible capacity fade over 30 charge– discharge cycles.

1.2.5 Organic Compounds Organic compounds are under study as cathode materials owing to their low cost, designability, and recyclability. The currently investigated organic compounds for Na-ion battery consist of carbonyl derivatives, polymers, etc. These cathode materials are categorized as cation-insertion type where reversible Na+ ion de-intercalation/intercalation is accompanied by electrochemical evolution of functional groups or anion-insertion type, which involves the incorporation/release of electrolyte anions in flexible host matrixes irrespective of the cation’s nature [85–90]. Commercial disodium rhodizonate (Na2C6O6) was studied as cationinsertion cathode material for Na-ion battery in 1 M NaClO4 in propylene carbonate by Chihara et al. [85]. A discharge capacity of 270 mAh g−1, corresponding to more than two Na+ ion insertions was obtained when cycled between 1.5 and 2.9 V vs. Na+/Na at a specific current of 18 mA g−1. Capacity fade owing to the dissolution of de-sodiated Na2C6O6 was noticed, and it was more pronounced when the upper potential limit was more than 3.2 V. Commercial 3,4,9,10-perylene-tetracarboxylic acid dianhydride (PTCDA) (C24H8O6) showed a reversible capacity of 140 mAh g−1 at a specific current of 10 mA g−1 with capacity retention of 77% after 195 cycles [86]. Similarly, when PTCDA-based polyimide containing dianhydride and alkyl chains were studied, a discharge capacity of 148 mAh g−1 at C/10 rate with 91% capacity retention after 400 cycles was obtained [87]. The cause for a good cycle stability was attributed to the presence of alkyl chain, which suppressed dissolution of the organic material in the electrolyte. Among the anion-insertion compounds, an aniline–nitroaniline copolymer (P(AN-NA)) and polytriphenylamine (PTPAm) were reported, and both these compounds exhibited intercalation/de-intercalation of [PF6]− anion [88, 89]. P(AN-NA) exhibited a reversible capacity of 180 mAh g−1,

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with capacity retention of 96% (50 cycles) in NaPF6/EC-DEC-DMC electrolyte [88]. In the case of (PTPAm), a high operating potential of 3.6 V with a rechargeable capacity of 98 mAh g−1 was obtained [89]. Sakaushi et al. prepared bipolar porous organic electrode (BPOE), which was a polymeric framework consisting of benzene rings and triazine rings in two dimensional structure, using ionothermal synthesis where a mixture of p-dicyanobenzene and ZnCl2 was heated at 900°C for 40 h in a quartz ampule [90]. The material could function by both the anion insertion mechanism at potential window of 4.1–2.8 V and as cation insertion mechanism between 2.8 and 1.3 V. A discharge capacity of 55 mAh g−1 was obtained at a specific current of 10 mA g−1 during the anion insertion mechanism in high potential window range whereas, a discharge capacity of 185 mAh g−1 was obtained for cation insertion mechanism. This suggests that the former mechanism involves storage of [ClO4]− on the surface of BPOE like pseudo-faradaic reaction, and the latter could undergo insertion/de-insertion of Na+ ion into the BPOE like a faradaic reaction.

1.3 Anodes for Na-Ion Batteries A wide range of materials is studied as possible anodes for Na-ion batteries, which includes, carbonaceous materials, sodium alloys, metal oxides, metal sulfides, phosphorous, and phosphides, etc. (Figure 1.2). Each type of material has its own set of advantages and disadvantages for applications in Na-ion energy storage systems. The following sections detail some of them.

1.3.1 Carbon-Based Electrodes Graphite is a well-studied anode for Li-ion batteries, which forms graphite intercalation compound (GIC), LiC6 and delivers a discharge capacity of 372 mAh g−1 [91]. However, the earlier studies by Doeff et al. proved that Na+ ion due to its large ionic radius forms NaC70 with a specific capacity of 31 mAh g−1 [91]. A solution to this low capacity could be envisaged by expanding the interlayer distance in graphite so that it could accommodate the larger size of Na+. Wen et al. prepared expanded graphite from graphite oxide (GO) by thermal reduction at 600°C (Figure 1.5) [92]. The expanded graphite when tested for de-intercalation/intercalation of Na+ ion delivered a discharge capacity of 284 mAh g−1 at a specific current of 20 mA g−1 and stable cycling up to 2000 cycles. Another approach used was to intercalate Na+ ion along with the electrolyte [93]. This concept was demonstrated by Adelhelm et al.

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(a)

Graphite

(b)

=C

=O

Graphite oxide

=H (c)

Expanded graphite

Figure 1.5 (a) Na+ cannot be intercalated into graphite because of the small interlayer spacing. (b) Electrochemical intercalation of Na+ into GO is enabled by the enlarged interlayer distance because of oxidation. However, the intercalation is limited by steric hindering from large amounts of oxygen-containing groups. (c) A significant amount of Na+ can be intercalated into EG owing to suitable interlayer distance and reduced oxygencontaining groups in the interlayers. Printed with permission [92].

in ether electrolyte. A ternary GIC with formula of Na(diglyme)2C20 was formed, and a discharge capacity of 100 mAh g−1 was achieved at a specific current of 37 mA g−1 with a stable cycle life for 1000 cycles. The coulombic efficiency of the first cycle was 70%, which was greater than many hard carbons used for Li+ ion intercalation. Also, high sodiation potential of 0.7 V vs. Na+/Na was considered safe as it avoided low potential Na plating. The authors found by XRD calculations that the volume expansion during intercalation was only 15%. In addition to the research on graphite, several amorphous and nanostructured carbons are under investigation. Dahn et al. reported hard carbon derived from glucose which delivered a discharge capacity of 300 mAh g−1 and proposed house of cards mechanism similar to that of Li+ ion storage in hard carbon [94]. They formulated that in the voltage region of 1.0–0.2 V vs. Na+/Na, storage occurs between the parallel/nearly parallel plane of graphitic structure whereas at potentials below 0.2 V, Na+ storage occurs in the nanovoids of hard carbon which exist between randomly stacked graphitic nano domains. A commercial hard carbon was investigated by Komaba et al. as a potential anode in a full cell Na-ion battery where NaNi0.5Mn0.5O2 was used as the cathode. The coulombic efficiency was 78%, and discharge capacity of 250 mAh g−1 was reported at a current rate of 25 mA g−1 [95]. Cao et al. prepared hollow carbon nanowires from hollow polyaniline nanowire precursor by pyrolysis. Upon testing it for Na-ion battery anode, it delivered a discharge capacity of 251 mAh g−1, and the capacity retention of 82% was maintained even after 400 cycles at a specific current of 50 mA g−1 [96]. At high current rate of 500 mA g−1, a discharge capacity of 149 mAh g−1 was obtained. They also carried out theoretical calculations and predicted that

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Na+ ion insertion was facilitated into graphitic layers when the interlayer distance was greater than 0.37 nm. The energy barrier for Na+ ion insertion dropped markedly to 0.053 eV from 0.12 eV in graphite with interlayer distance 0.34 nm. This has promoted researchers to prepare amorphous carbon with larger interlayer distance. More recently, Wang et al. prepared hard carbon with d-spacing in the range of 0.39–0.43 nm from kelp, by carbonizing at different temperatures in the range 900–1600°C under Ar for 2 h [97]. The interlayer spacing decreased and turbostatic nanocrystallite size increased with an increase in carbonization temperature. The sample prepared at 1300°C showed the best result of a de-sodiation capacity of 334 mAh g−1 at a specific current of 25 mA g−1. They also found that it was hard for Na+ ions to intercalate in kelp derived hard carbon to a large amount where the oxygen containing defects were large. Furthermore, there was a critical content of oxygen containing groups below which the Na+ intercalation became difficult. On similar lines, doping with nitrogen, fluorine or sulfur was carried out and the doped carbons were studied as anode material in Na-ion battery [98–100]. Wang et al. prepared N-doped carbon and demonstrated a discharge capacity of 259 mAh g−1 at 0.2 A g−1 in contrast to pure sample which showed only 170 mAh g−1 [98]. Similarly, fluorine doped carbon particles (F-CP) derived from lotus petioles showed a capacity of 230 mAh g−1 at a specific current of 50 mA g−1, which outperformed the capacity obtained (149 mAh g−1) from banana peel carbon (BPC) prepared by similar method except for the fluorine doping step. The capacity retention was better for F-CP (99%) than BPC (72%) after 200 cycles [99]. Qie et al. prepared sulfur doped carbon with enlarged interlayer distance and concluded that a large interlayer distance with low specific surface area are desirable to achieve high performance from carbon based anode materials for Na-ion battery [100]. They achieved first cycle capacity of 384 mAh g−1, which stabilized at 322 mAh g−1 after 10th cycle and maintained 303 mAh g−1 after 700 cycles at a specific current of 0.5 A g−1.

1.3.2 Alloy Electrodes Alloy anodes have the advantage of high capacity due to the combination of an element with more than one sodium in the process of alloy formation. However, the electrodes suffer from volume changes upon cycling. Si based alloys which are widely investigated for Li-ion cells are not suitable for Na-ion cell as redox potential of Si is close to that of Na/Na+. Alloys based on Sn, Sb, and Ge have been investigated for Na-ion cell [101]. Darwiche et al. provided insights into the phase transformation during

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sodiation/de-sodiation process with micrometric Sb based anode [102]. During sodiation, amorphous NaxSb phase was formed as intermediate which was converted to a mixture of cubic and hexagonal Na3Sb phases. Apart from providing mechanistic view, they also demonstrated that the addition of fluoroethylene carbonate as additive to the electrode improved stability due to the formation of SEI layer. Using 20 nm sized Sb crystals, excellent rate capability was achieved in Na-ion cell by He et al. [103]. Wang et al. performed in situ TEM studies on nanosized Na-ion battery with Sn nanoparticles as anode and concluded that at first sodiation of Sn occurred through a two phase process involving 56% volume change followed by a single phase process resulting in a total change in volume of 420% [104]. Single phase process was able to accommodate such huge change in volume without causing fracture of the particle. Sb-C nanofibers prepared through electrospinning technique were successful in enhancing the conductivity along with managing mechanical stress [105]. Upon subjecting the cell to cycling at C/3 rate, 90% of capacity was retained after 400 cycles. Nam et al. prepared porous Sb/Cu2Sb electrode by electrodeposition and achieved a capacity of 485 mAh g−1 after 120 cycles with 97% coulombic efficiency at C/10 rate [106]. They proposed the formation of crystalline Na3Sb through NaxSb intermediate along with copper particles during sodiation. Reverse process occurred during de-sodiation and Cu2Sb was formed. The matrix of copper particles reduced the adverse effect of volume change due to cycling.

1.3.3 Phosphorous, Phosphides, and Nitrides Phosphorous has a redox potential of 0.4 V vs. Na/Na+. In addition, with Na3P as the final product, theoretical capacity becomes 2596 mAh g−1 [107]. Hence, phosphorous is a potential anode material for Na-ion battery. However, phosphorous based electrodes suffer from poor electrical conductivity and large volume change during cycling. In order to mitigate these issues, phosphorous is combined with conducting and porous materials, which can increase the conductivity of electrode and accommodate the volume change. Kim et al. prepared amorphous red phosphorous-carbon composite by ball milling and achieved a reversible capacity of 1890 mAh g−1 and a stable performance over 30 cycles [107]. At a specific current of 2.86 A g−1, capacity of 1540 mAh g−1 was achieved. Graphene-phosphorous hybrid prepared by Song et al. resulted in a stable anode in Na-ion battery for 60 charge–discharge cycles [108]. A strategy of forming NaxPy intermediates instead of Na3P was successful to minimize volume expansion and a stable performance over 200 cycles [109]. Using vapor condensation

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technique, Zhu et al. prepared red phosphorous-single walled carbon nanotubes composite and were able to demonstrate 80% capacity retention after 2000 cycles in a Na-ion battery [110]. In order to achieve intimate contact between the components of anode, Song et al. chemically linked phosphorous to carbon nanotubes and this hybrid was bonded to a crosslinked binder [111]. Cells fabricated with such electrodes provided stable capacity over 100 cycles. Apart from phosphorous, metal phosphides were evaluated as anodes for Na-ion battery. Fullenwarth et al. studied NiP3 as anode in Li-ion and Na-ion batteries [112]. A reversible capacity of 1000 mAh g−1 was achieved and 11% loss in capacity was observed after 15 cycles. Walter et al. examined nanocrystals of transition metal phosphide as negative electrode material in Na-ion batteries [113]. Highest capacity was obtained from CuP2 but capacity fade was rapid in all materials. Agglomeration of metal particles and formation of unstable SEI layer were major causes for capacity fade. Kim et al. succeeded in overcoming these issues in Sn4P3 anode as Sn particles were embedded in amorphous phosphorous during cycling [114]. Addition of fluoroethylene carbonate resulted in formation of stable SEI layer and enhanced the capacity and stability over 100 cycles. A few reports have appeared on nitrides as promising anode materials for Na-ion battery [113, 114]. In metal nitrides, conversion reaction occurs at the surface of particles leading to the formation of metal nanoparticles and sodium nitride. Li et al. utilized Ni3N nanoparticles as anode material and were able to achieve capacity of 126 mAh g−1 after 30 cycles [115]. Amorphous carbon nitride and zinc oxide composite delivered a stable capacity over 2000 cycles [116].

1.3.4 Sulfides and Selenides Layered MoS2 was first reported for Na-ion battery by Park et al., with an initial capacity of 190 mAh g−1 and retention of 85 mAh g−1 after 100 cycles [117]. Wang et al. probed into the phase transition and stability of MoS2 [118]. Upon intercalation of sodium into MoS2, 2H-MoS2 formed 1T phase of NaxMoS2 (x < 1.5), which was partially converted in to 1T-MoS2 upon extraction of sodium. For intercalation of >1.5 Na+ ion per MoS2, NaxS, and Mo were formed. MoS2 structure could not be retained and NaxS took part in further electrochemical reactions. A composite of free standing few layers of MoS2 and reduced graphene oxide was studied. For, 60% MoS2 in the composite, capacity of 338 mAh g−1 was achieved [119]. Apart from MoS2, several other layered sulfides were studied as anodes for Na-ion battery. Xie et al. synthesized SnS2 nanoplatelet–graphene nanocomposite and

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achieved a stable capacity over 60 cycles by preventing pulverization of electrode due to volume change upon cycling and also an improvement in electron transport [120]. Porous Co3S4 nanosheets embedded in graphene sheets retained 71% of initial capacity after 50 cycles at 0.5 A g−1 [121]. The cause for stable capacity was fast sodium ion insertion/extraction kinetics and prevention of aggregation of nanoparticles as well as pulverization. Recently, V5S8–graphite hybrid nanosheets based electrode demonstrated a stable capacity over 500 cycles at 1 A g−1 [122]. Walter et al. evaluated nanocrystals of several metal sulfides including ternary metal sulfides (Cu2ZnSnS4) and demonstrated that FeS2 was superior to the others with a capacity >400 mAh g−1 retained over 600 cycles at a specific current of 1 A g−1 [123]. Among metal dichalcogenides, selenides have also been evaluated as anode materials for Na-ion battery. Yolk-shell structured MoSe2 (Figure 1.6) had higher capacity than MoO3 and stable over 50 cycles [124]. In order to achieve enhancement in conductivity and stability of the electrodes, carbon coatings were adapted. Carbon coated Sb2Se3 delivered stable capacity greater than 600 mAh g−1 over 100 cycles at a specific current of 0.2 A g−1 [125]. Xie et al. synthesized C-MoSe2/reduced graphene oxide composite by hydrothermal process and achieved stable performance in Na-ion battery over 350 cycles [126]. After the first cycle, the electrode material

500 nm Figure 1.6 FESEM image of a fractured MoSe2 microsphere showing the yolk and shell. Reprinted with permission [124].

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turned into a mixture of Mo and Se or polyselenide on reduced graphene oxide.

1.3.5 Phosphates NASICON type NaTi2(PO4)3 provides a theoretical capacity of 133 mAh g−1 by uptake of 2 Na+ ions per formula unit forming Na3Ti2(PO4)3 due to the Ti3+/Ti4+ redox couple [14]. An operating voltage of 2.1 V was achieved due to the presence of phosphate and Ti3+/Ti4+ redox couple which was greater than most of the anode materials. Park et al. studied NaTi2(PO4)3 as anode in Na-ion batteries and compared the performance in aqueous and organic electrolytes [127]. Superior capacity and stability was achieved in the aqueous electrolyte based cell due to higher conductivity of the electrolyte than organic electrolyte. Senguttuvan et al. utilized Ti3+/Ti4 and Ti2+/ Ti3+ redox couples as positive and negative electrodes, respectively, in a full titanium based symmetric cell with Na3Ti2(PO4)3 achieving an average voltage of 1.7 V in an organic electrolyte [128]. Jiang et al. embedded NaTi2(PO4)3 in carbon network which enhanced electrical conductivity of the electrode and structural stability during cycling [129]. The cell was cycled for 6000 cycles at 50C rate with stable performance. Carbon coated Na3V2(PO4)3 nanoparticles embedded in graphene sheet prevented agglomeration of particles upon cycling and also prevented pulverization of electrode due to volume changes [130]. As a result, cells retained 80% capacity after 1500 cycles at 40C rate. Anode comprising of tin phosphate based glass particles provided a capacity of 320 mAh g−1 with efficiency of 100% after 10 cycles [131]. During sodiation, the nanoparticles of Sn alloy with Na were embedded in phosphate matrix which provided stability to the electrode. NASICON-type material with Na replaced by Mg has also been evaluated as anode for Na-ion batteries. Mg0.5Ti2(PO4)3 provided an initial capacity of 97 mAh g−1 at C/20 rate which decreased to 50 mAh g−1 after 18 cycles [132]. Recently, Zhao et al. coated carbon on Mg0.5Ti2(PO4)3 particles to improve the electrical conductivity [133]. Capacity of 268 mAh g−1 was achieved at a specific current of 20 mA g−1 and the performance was stable over 50 cycles.

1.3.6 Organic Materials Organic electrode materials are comparatively less explored than inorganic materials. They offer the advantage of structural flexibility but solubility in the electrolyte and poor kinetics are challenges. Disodium terephthalate [134] (Na2C8H4O4) studied as anode for Na-ion cell resulted in a reversible

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capacity of 250 mAh g−1. Capacity retention was improved by coating the electrode with a thin layer of Al2O3. Park et al. examined disodium terephthalate and its derivatives for Na-ion cell [135]. Except with the nitro derivative, stable performance was achieved with other materials. Wang et al. designed sodium 4,4’-stilbene-dicarboxylate with an extended -conjugation, which resulted in fast charge/discharge cycling and a capacity of 72 mAh g−1 at a specific current of 10 A g−1 [136]. Croconic acid disodium salt was also studied as anode for Na-ion battery [137]. A reduction in particle size increased the performance and also stability. In order to overcome the problem of pulverization, graphene oxide was wrapped over the particles. The electrodes prepared with such wrapped particles delivered better stability than the electrodes prepared with both pristine and smaller particles. All organic Na-ion battery with energy density of 65 Wh kg−1 demonstrated the viability of organic electrode materials for Na-ion battery [138].

1.3.7 Oxides A majority of the materials explored for anodes of Na-ion batteries is oxides of transition metals. Xiong et al. prepared amorphous TiO2 nanotubes by electrochemical anodization and investigated their application as anode in a Na-ion battery [139]. The performance of electrodes comprising of nanotubes with inner diameter >80 nm was superior to tubes with inner diameter 300oC) is well established in energy storage. It consists of both Na and S in molten states, which are separated by ceramic electrolyte ( -alumina). High energy density of 760 Wh kg−1 is achieved with Na2S3 as the discharge product [157]. However, the system faces issues of safety and reliability as it involves high temperature and molten materials. Hence, room temperature Na–S battery is under investigation in order to make the system safe and reliable. Considering Na2S as the final discharge product, theoretical energy density of ambient Na–S battery is 1230 Wh kg−1 [158]. However, dissolution of polysulfides and poor conductivity of sulfur are problems similar to Li-S cells. Several studies are reported on preparation of composites of sulfur and utilization of them to attain high capacity with stable cycle life. In 2007, polyacrylonitrile (PAN)-sulfur composite in carbonate based liquid electrolyte provided a discharge capacity of 1455 mAh g−1 and stability over 18 cycles [159]. Sulfurized PAN nanofibers were used as electrodes free from binder, conducting agent and current collector. Discharge capacity of 342 mAh g−1 was obtained initially and retained 266 mAh g−1 after 200 cycles [158]. Binding sulfur to carbonized PAN nanofibers resulted in capacity retention up to 500 cycles at 1C rate [160]. Recently, sulfur covalently bonded to carbonaceous structure has been used to overcome the formation of high order polysulfide formation. Even though the content of sulfur was only 18% in the composite, highly stable performance over 900 cycles with 0.053% capacity fading per cycle was achieved [161]. Doping of hetero atoms into carbon nanostructures could confine the polysulfides. Metal-organic framework derived nitrogen doped porous carbon was able to trap polysulfides successfully and provide stability over 250 cycles at C/5 rate [162]. In another approach N and S doped porous carbons were able to retain polysulfides, and 80% of capacity was retained over 10,000 cycles at 4.6 A g−1 [163]. The concept of adding a functional layer between the cathode and separator (known as ‘interlayer’) or modification of separator was introduced in Na-S cells similar to Li-S cells. Excellent cycling stability has been achieved by inserting carbon nanofoam as interlayer and discharging the cell up to 1.8 V [164]. Interlayer prevented shuttling of polysulfides between the electrodes. Sodiated-Nafion coated on porous polypropylene separator when used in Na-S cell enhanced the capacity retention due to improved Na+ ion conductivity and retarded polysulfide migration [165]. In another

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study, activated carbon nanofiber coated on sodiated Nafion was used as the separator, which improved the Na+ ion conductivity but suppressed the polysulfide migration due to structural and electronic effects. With Na2S coated on activated carbon nanofiber cathode, sodium metal anode and liquid electrolyte, such a separator facilitated an initial discharge capacity of 680 mAh g−1 and stability over 100 cycles [166]. Sulfur (S8) was studied extensively as the active cathode material but few studies on other molecules were also reported. Complete conversion to Na2S was achieved by using small sulfur molecules (S2-4), which yielded a high discharge capacity of 1610 mAh g−1 [167]. Yu et al. fabricated Na-S cell with liquid sodium polysulfide catholyte and multiwalled carbon fabric electrode [168]. They demonstrated that irreversible transitions between high order polysulfides and low order polysulfides or Na2S resulted in capacity fading. Stable capacity was maintained by cycling cell between sulfur and high order polysulfide (Figure 1.8). Na2S instead of sulfur was used as the active material and a Na2S-carbon nanotube fabric electrode was fabricated [169]. The cell delivered an initial discharge capacity of 660 mAh g−1 at C/10 rate and retained 560 mAh g−1 after 50 cycles. For the first time, Na-S full cell using hard carbon anode was reported by Kohl et al. [170]. Addition of fluoroethylene carbonate to electrolyte during sodiation enhanced the formation of SEI. With Na2S/P2S5 as additive in an

120

Upper: 2.8/1.8 V

Capacity retention, %

100 80 Entire: 2.8/1.2 V 60

Lower: 2.2/1.2 V

40 20 0

0

20

40 60 Cycle number

80

100

Figure 1.8 Capacity retention as a function of cycle number for the Na/dissolved sodium polysulfide cells operated within different voltage windows. Reprinted with permission [168].

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35

ether-based electrolyte, a discharge capacity of 980 mAh g−1 was obtained and 200 mAh g−1 was retained after 1000 cycles. Research on ambient Na-S battery is still at early stage. Several promising results have been reported and few have been discussed here. The state-of-the-art of Na-S battery along with scope for further improvements has been reported by Manthiram and Yu [171]. For portable energy storage applications, weight of sodium as compared to lithium is an issue. However, there is immense potential in the system to compete with other batteries for stationary applications.

1.3.9 Na-Air Batteries Research on Na-O2 batteries is fuelled as an alternative to Li-O2 battery. The first rechargeable Na-O2 battery was demonstrated by Peled et al. [172] in 2011, where the battery was operated at a temperature higher that the melting point of Na. The reason for using a temperature of 100°C was cited to prevent the dendrite formation (a major problem for using alkali metal as anode, in both Li and Na systems) and the parasitic reactions caused due to the presence of moisture. Later in 2012, Sun et al. [173] demonstrated the room temperature Na-O2 cell. Similar to Li-O2 cells, the room temperature Na-O2 cells are plagued with several problems [174,175]. The formation of dendrite is one such major problem which not only leads to the failure of the cell by internal short circuit but also causes the drying of the cell because electrolyte is consumed in the formation of SEI on new dendrites. The process also leads to the corrosion of Na metal anode. Another major problem is a high reactivity of Na metal toward dissolved O2 or moisture. The superoxide anion (O2 ) formed during the discharge process is a good oxidizing agent and is known to attack the electrolyte and binder. The decomposed electrolyte/binder acts as a source for several species which in turn starts a chain of several parasitic reactions, some leading to higher charging potentials. Furthermore, the instability of the polymeric binder also leads to the disintegration of the air electrode leading to the failure of the cell. A proper air electrode or gas diffusion layer (GDL) is also important. It should possess proper porous structure with appropriate pore volume and pore size distribution. Although the above challenges are important and need to be addressed, an immediate attention is needed to study the reaction mechanism of Na-O2 cell. The current literature is filled with ambiguity in the discharge product [176]. Several research groups have claimed NaO2 as the final discharge product while some have reported Na2O2. There are reports on formation of Na2CO3, Na2O2.2H2O, and NaOH as well. It was

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Nanomaterials for Electrochemical Energy Storage Devices

found that the formation of NaO2 was via one electron transfer process at very low charging potential when compared to Li-O2 cells, whereas Na2O2 had a charging potential very similar to Li-O2 cell where Li2O2 was the discharged product. The reaction mechanism for a Na-O2 cell can be described by the following reactions. At the negative electrode:

Na

Na+ + e–

(1.12)

At the positive electrode:

O2 e

O2

(1.13)

O2 2 e

O22

(1.14)

Na + O2

NaO2 E =2.27 V

(1.15)

2 Na + O2

Na2O2 E = 2.33 V

(1.16)

Overall:

From the standard cell potentials, it is clear that Na2O2 is the thermodynamically stable product and Gibbs free energy for formation is −449.7 kJ mol−1. The Gibbs free energy of formation for NaO2 is −437.5 kJ mol−1, and the difference of 12.2 kJ mol−1 translates into a difference of 0.07 V between the cell potentials. This also shows that both the products are energetically similar. Theoretical calculations have indeed shown that crystallite size, oxygen partial pressure, temperatures, etc. are determining factors for the formation of the stable product. Ceder et al. [177] used first principle calculations to find out the thermodynamic stability of a variety of oxides of sodium as a function of oxygen partial pressure, temperature, and particle size of the discharge product. They predicted that Na2O2 is the stable product, which was supported by more negative Gibbs free energy of formation. They also concluded that NaO2 is thermodynamically stable only at partial pressures greater than 8.5 atm. They presented the phase diagram as a function of particle size and oxygen partial pressure for Na2O2 and NaO2. Based on the surface energy value which was lower for NaO2, it was

Electrochemistry of Rechargeable Batteries

37

concluded that formation of small crystallites of NaO2 was more favorable. The critical nucleation energy barrier was smaller for NaO2 when compared with Na2O2. The phase diagram data were supplemented by the findings of Lee et al. [178] where at standard experimental conditions of 298 K and PO2 = 1 atm, the phase stability of discharge products of Li-O2 and Na-O2 cells were calculated. Li2O2 and NaO2 were found to be stable discharge products of Li-O2 and Na-O2 cells, respectively. The material of air electrode needs to be very stable in an oxidizing environment of O2 and O22 ions and also it should have large surface area with appropriate pore volume. Large surface area provides high utilization of the catalyst loaded and the porous structure ensures the diffusion of O2 into the cell which enhances the formation of discharge product and decomposition of the discharged product during charging. Several carbon-based materials have been studied for this purpose. Kim et al. [179] used commercially available Ketjen Black carbon as air electrode and obtained a specific discharge capacity of 2800 mAh g−1 in a carbonate based electrolyte with Na2CO3 as the discharge product, whereas 6000 mAh g−1 was obtained in ether based electrolyte with Na2O2 as the discharge product. Diamond like carbon obtained by radio frequency sputtering exhibited a specific discharge capacity of 3600 mAh g−1 at 1/60 C rate and Na2O2 and Na2CO3 were found as the discharged products [173]. Graphene nanosheets of surface area 83 m2 g−1 demonstrated a discharge capacity of 9268 and 1110 mAh g−1 when cycled galvanostatically at specific currents of 200 and 1000 mA g−1, respectively [180]. A systematic study of correlation between the discharge capacity and surface area was carried out by Yadegari et al. [181] where a series of controlled porosity and surface area was obtained by applying heat treatment to commercial carbon black. The results exhibited a linear correlation of discharge capacity with specific surface area of the air electrode in the mesoporous range. The use of bifunctional catalyst is believed to promote the oxygen evolution reaction (OER) and oxygen reduction reaction (ORR) activity of the air electrode in Li-O2 cells. Similarly, transition metals, their oxides and composites with reduced graphene oxide have been studied as bifunctional catalyst. α-MnO2 nanowire has been used by Rosenberg and Hintennach [182]. They showed a very large first discharge capacity of 2056 mAh g−1 but suffered a 59% decrease in capacity after 2 cycles. Ni foam @NiCo2O4 nanosheet electrode was demonstrated by Liu et al. [183], which showed first discharge capacity of 1185 mAh g−1 and maintained a capacity of 401 mAh g−1 after 10 cycles. Surender et al. [184] employed Ag-RGO as the bifunctional catalyst to enhance the ORR and OER processes and

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Nanomaterials for Electrochemical Energy Storage Devices

demonstrated a specific discharge capacity of 0.54 mAh cm−2 at a specific current of 0.1 mA cm−2. A cycle life for about 30 cycles with coulombic efficiency more than 90% was demonstrated. The catalyst lowered the difference between the charge–discharge plateaus. The stability of electrolyte and binder is also important for the functioning of Na-O2 cells [185]. It has been reported that ether based electrolytes are more stable when compared with the usual carbonate based electrolytes in Li-ion batteries [186]. The carbonate based electrolytes are susceptible to nucleophilic attack by O2 anion. The presence of functional groups on the backbone of polymeric binders can also be attacked by strong nucleophiles such as O 2 or O22 produced during the ORR. This would decompose the binder and disintegrate the air electrode. The complexity of the discharged product would increase due to the change in electrochemical pathways due to presence of various disintegrated products. It has been noticed that trace amount of water has the ability to convert the discharge product from NaO2 to Na2O2.2H2O. Similarly, Na2CO3 has also been observed as the discharge product in some studies. Researchers have narrowed that polyethylene, polypropylene and fully fluorinated polymers such as polytetrafluoroethylene and Nafion are stable binders to be used to Na-O2 cells. The formation of discharge product is dependent on various parameters and various parasitic reactions also arise from the instability of the electrolyte and binder. It is also important to choose a proper catalyst for ORR and OER along with high surface area carbon with proper porosity as the air electrode. The charging overpotential can be reduced by controlling the discharge product as NaO2 and this would also improve the cycle life. All the above approaches are expected to lead toward a realistic Na-O2 prototypes.

1.4 Potassium Batteries Potassium-based systems include aqueous and non-aqueous K-ion batteries, K-O2 batteries and K-S batteries (room temperature and moderately high temperature) [187]. These systems are gaining importance primarily due to the shift towards systems using earth abundant resources which would reduce the cost and available for long periods. Hence, K-based battery systems are expected to be attractive in future over the well-known Li-based systems which have limited reserves and also involved with socio-politico factors, primarily for large scale stationary applications.

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39

Though the obvious alternate for lithium would be the next element below it in the periodic table, i.e., Na, but K can in principle compete with both of them and have several advantages as highlighted below (i)

Cell voltage: It is known that down the group, the reduction potential becomes more negative with exception being Li+/Li due to the high de-solvation energy in aqueous electrolyte [188]. But in non-aqueous carbonate based electrolytes, this behavior is not so profound. It has been proven by both theory [189] and experiment [190] that K+/K reduction potential is more negative than Li+/Li. Komaba et al. showed that the difference is about 0.15  V, which means K-based non-aqueous batteries would deliver a high cell voltage than Li-based by about 0.15 V [190]. (ii) Abundance and low cost: Potassium is the 5th abundant element in the earth crust occupying 2.4% (Na being 2.6% and Li 0.007%) [191]. Its abundance in sea water is 400 ppm, which is 2000 times higher than Li. Moreover, it is widely distributed over the globe unlike lithium, which is mainly present in Argentine, Bolivia, and Chile. More than 70% of the world reserves of Li are present in these countries. Hence, K-based batteries would be an attractively cheaper option for stationary applications. (iii) Electrochemical behavior: Potassium ion being larger in size than Li+ ion, the volume changes occurring during intercalation/deintercalation would be more in electrode materials which are analogous to Li-ion batteries. But as the reactivity of K is higher, several new classes of compounds with favorable properties can be studied. Moreover, it has been confirmed that K+ ion has high diffusion coefficient, high ionic conductivity and straight forward electrochemical behavior.

1.4.1 Potassium-Ion Batteries In view of the attractive electrochemistry of K, several electrode materials for both cathode and anode are under investigation for aqueous and non-aqueous batteries (Figure 1.9) [192]. The effect of binder and electrolyte are also under study. Some details are discussed below.

40

Nanomaterials for Electrochemical Energy Storage Devices 5 Prussian Blue analogues Layered oxides Polyanion compounds Organic materials Other intercalation compounds Alloying materials Conversion materials Carbon

KVPO4F KFeFe(CN)6 FeSO4F KMnFe(CN)6 KNiFe(CN)6 4 PTCDA

Voltage (V vs. K)

K3V2(PO4)3

FePO4

K0.6CoO2

3

K0.3MnO2 K0.7Fe0.5Mn0.5O2

K2C6O6 KTi2(PO4)3

2

PAQS MoS2 1 K2Ti8O17 K2Ti4O9 0

0

F-Graphene

Graphite Sn/C

Co3O4-Fe2O3/C

K2TP

Graphite (cointercalation)

100

200

Sn4P3/C 300

400

500

Capacity (mAh g–1)

Figure 1.9 The relationship between capacity and voltage of electrode materials for K-ion batteries. Reprinted with permission [192].

1.4.1.1 Electrolytes The choice of electrolyte for K-ion batteries has been guided by the electrolytes used for non-aqueous Li-ion batteries. Accordingly, the solvent used is a mixture of alkyl carbonates such as ethylene carbonate, propylene carbonate, diethyl carbonate and dimethyl carbonate. The salts are potassium hexafluorophosphate [191], potassium tetrafluoroborate [193], and potassium bis(fluorosulfonly)imide [190] in the concentration range 0.5–1.0 M. Potassium nitrate is the most commonly used salt for aqueous electrolyte [194].

1.4.1.2 Cathode Materials The first cathode material for K-ion batteries is reported by Eftekhari [193], where PB was synthesized electrochemically and cells were tested in non-aqueous electrolyte. At C/10 rate, a discharge capacity of 78 mAh g−1 was obtained, which was 90% of the theoretical capacity and 88% capacity retention was observed after 500 cycles. In a recent study, K0.3MnO2

Electrochemistry of Rechargeable Batteries

41

was synthesized by thermal decomposition of KMnO4 [195]. The material was tested as cathode in various voltage windows. The variations in specific capacity and cycle life study over 50 cycles were reported. A discharge capacity of 136 mAh g−1 with 58% retention after 50 charge–discharge cycles in the voltage window of 4.0–1.5 V vs. K/K+, 100 mAh g−1 with 73% retention for 4.0–2.0 V and 70 mAh g−1 with 91% retention in the window of 3.4–1.5 V were observed. A full cell study using a composite of hard carbon synthesized by thermal decomposition of starch powder with carbon black as the anode and K0.3MnO2 as the cathode in the voltage range of 3.40–0.50 V was reported. The first cycle discharge capacity was 90 mAh g−1 with cycling efficiency of 99%. The capacity retention was 50% after 100 cycles. In another study, nanostructured iron((III)oxyhydroxide/(VI) oxide) [196] was tested as the cathode material for Li-, Na-, and K- ion cells and discharge capacities of 320, 218, and 150 mAh g−1, respectively, were obtained. Equal importance was given to cathode materials for aqueous K-ion batteries and metal hexacyanometallates [194, 197, 198] were investigated for this purpose. Nickel hexacyanoferrate [194] nanoparticles co-precipitated from nickel nitrate and potassium ferricyanide delivered a discharged capacity of 58 mAh g−1 at 0.83C rate with a stable capacity retention over 1000 cycles. Similarly, Prussian green (PG) and PB were also tested and a superior performance of PG was explained as due to smaller particles (size 50–75 nm) and lower diffusion length leading to enhanced capacity utilization than bigger particles of PB (2–10 μm) [197]. PG exhibited a discharge capacity of 121 mAh g−1, whereas it was 54 mAh g−1 for PB when tested in 1 M KNO3 electrolyte using Ag/AgCl, KCl (saturated) as the reference electrode. A complete cell using KCrIIICrII(CN)6 as anode and KCrIIIFeII(CN)6 as the cathode with a cell voltage of 1.5 V was also reported [198].

1.4.1.3 Anode Materials Graphite is a commercially successful anode material for Li-ion battery and hard carbons with d-spacing greater than 0.37 nm are extensively studied for Na-ion battery. It is known that K+ ions can easily intercalate in graphite as the formation enthalpies of KCx is sufficiently negative to form thermodynamically stable GIC [199]. Carbon nanofibers [200], commercial and nanocrystalline graphite [190, 201–203], reduced graphene oxide [204], and hard carbons [205] have been investigated as potential K-ion anode. Several in situ studies such as transmission electron microscopy [200] and Raman spectroscopy [203] and ex situ X-ray diffraction [190, 201, 204] techniques have been

42

Nanomaterials for Electrochemical Energy Storage Devices

employed to study the different stages of GIC formed during intercalation/ deintercalation of K+ ion. In situ TEM studies of nanofibers showed that discharge capacity of 200 mAh g−1 at a specific current of 50 mA g−1 was contributed by the outer layer of nanofibers, which was highly disordered than the crystalline inner layer [200]. Graphite is reported to provide a discharge capacity of 273 and 244 mAh g−1 in two separate studies conducted by Jian et al. [201] and Komaba et al. [190]. Jian et al. [201] confirmed the formation of various GIC such as KC36, KC24, and KC8 on intercalation of K+ ion by ex situ XRD studies. Komaba et al. [190] demonstrated the effect of different binders such as sodium polyacrylate (PANa), sodium carboxymethylcellulose (CMCNa), or poly(vinylydenefluoride) (PVDF) on the first cycle coulombic efficiency (irreversible capacity loss), where maximum efficiency of 89 % was observed for CMCNa. The ex situ XRD provided evidence for the formation of KC8 GIC [190, 201]. By analysis of in situ Raman spectroscopy of few layer graphene obtained by chemical vapor deposition, Share et al. reported the formation of ordered intercalation compounds and also transition from GIC VI (KC72)–GIC II (KC24)–GIC I (KC8) [203]. RGO film was reported to deliver a specific discharge capacity of 200 mAh g−1 at 5 mA g−1 [204]. Hard carbon microspheres obtained by hydrothermal treatment of sucrose exhibited a specific discharge capacity of 262 mAh g−1 at C/10 rate with 83% capacity retention after 100 charge/discharge cycles [205]. It also showed a plateau voltage of 0.2 V vs. K/K+, which is an advantage in preventing plating of K in deep discharge cycles. KTi2(PO4)3, a polyanionic compound, was synthesized using hydrothermal route followed by subsequent annealing [206]. The nanocubes obtained were tested as anode material and a specific discharge capacity of 74.5 mAh g−1 was obtained when cycled in the range of 1.2–2.8 V vs. K/K+. It exhibited a poor cycling stability, which was improved by carbon coating using cane sugar. In another report, K2Ti8O17 was synthesized using a similar procedure and specific discharge capacity of 182 mAh g−1 at 20 mA g−1 was reported in the voltage range 0.01–3.00 V, which decreased to 111 mAh g−1 after 50 cycles [207]. Kishore et al. [191] have synthesized rod like morphology of K2Ti4O9 using solid state synthesis, where K2CO3 and TiO2 were taken in stoichiometric ratio and heated in a furnace at 900°C for 4 h under air ambient. A discharge capacity of 97 mAh g−1 was reported at a specific current of 30 mA g−1. Though 10 wt% of conducting carbon acetylene black was used during electrode preparation, it was confirmed that most of the capacity was obtained from the Ti+3/Ti+4 redox couple. The intercalation of K+ into the electrode material was also

Electrochemistry of Rechargeable Batteries

43

quantified by EDAX. Other anodes such as Ti3C2MXene with theoretical capacity of 192 mAh g−1 and Sn-C composite, which stores K+ ion by alloying were also studied [208, 209]. SN-C composite exhibited a discharge capacity of 150 mAh g−1 but cycle life was poor.

1.4.2 Potassium–Sulfur Batteries Zhao et al. [210] have reported room temperature K-S battery where they have used a composite of CMK-3 (highly ordered mesoporous carbon material) and sulfur composite as a cathode material. The CMK-3 improved the electronic conductivity of sulfur and also reduced the volume expansion (major problem associated with Li-S or Na-S batteries). The addition of polyaniline (PANI) was shown to improve the cycle life by reducing the solvation of polysulfide anions. CMK-3/sulfur composites were prepared by melt-diffusion process and PANI was coated by in situ polymerization in an acidic medium using ammonium persulfate as the oxidizing agent. K2S3 was confirmed to be the final discharge product and the composite with 40 wt% S gave a specific discharge capacity of 512 mAh g−1 at a specific current of 50 mA g−1. The specific discharge capacity reduced to 202 mAh g−1 after 50 cycles. When the composite with 40 wt% of S and CMK-3 was given a coating of PANI, the cycling stability improved and a discharge capacity of 329 mAh g−1 was obtained after 50 cycles.

1.4.3 Potassium–Air Batteries Potassium–air batteries have their advantages over Li-air or Na-air because in principle, K-air cell can be fabricated where the cathode, i.e., oxygen is used directly from air as K does not react with N2, which is the major component. Another advantage offered is in terms of overpotential. It has been demonstrated that K-air cell has low overpotential in comparison with Li-air or Na-air cells as the voltage gap between oxygen reduction and oxidation in K electrolyte is smaller owing to the formation of kinetically and thermodynamically stable superoxide as the discharge product [211]. In Li-air cell, the final discharge product is Li2O2, which is formed via thermodynamically unstable LiO2 as smaller Li+ is stabilized by smaller O22− according to hard soft acid base theory as reported by Ren and Wu [211]. A new concept of K+-oxygen [212] battery has been introduced recently owing to the safety risks associated with using an alkali metal as the anode. In this regard, intercalating anode such as graphite or carbon (ordered or disordered) is a good candidate material due to its excellent

44

Nanomaterials for Electrochemical Energy Storage Devices

performance. Other elements such as antimony, which store K via alloying mechanism, were also investigated.

1.5 Mg-Based Rechargeable Batteries The advantages expected from Mg metal as battery anode were visualized long ago. High theoretical potential (−2.37 V) and also high specific capacity (2.2 Ah g−1) are attractive features. Although the atomic mass of Mg (24.3) is greater than that of Li (6.9), the two electrons transfer electroxidation of Mg provides attractive theoretical capacity. Additionally, the raw material resources of Mg are abundant (104 time of Li) and therefore it is inexpensive. Furthermore, Mg can be handled in ambient atmosphere without any need for inert atmosphere environment. However, corrosion of Mg in ambient atmosphere results in the formation of an oxide film on the surface. The intrinsic corrosion-resistance of Mg is due to its surface oxide film, the removal of which accelerates its corrosion rate in aqueous solutions. The first ever known use of Mg for battery application was in 1865 by Bultinck, who observed an increase in cell voltage by replacing Zn in an voltaic pile by Mg [213]. Subsequently, Mg was used in Daniell cell, sea water activated cell, dry cell, etc. The Mg dry cell consists of Mg as the anode and MnO2 as the cathode in a suitable electrolyte similar to Zn/MnO2 dry cell. Although Mg/MnO2 cell is superior to Zn/MnO2 dry cell in terms of capacity, energy density, power density, cell voltage, shelf-life, etc., it is plagued with voltage delay action and perforation of partially discharged cells. Because of these problems, Mg/MnO2 dry cell has not become a commercial reality for consumer applications. Studies related to the development of secondary Mg batteries are limited. Investigations on rechargeable Mg batteries are interesting on many accounts in comparison with Li batteries, namely: (i) the ionic radii of Li+ and Mg2+ are 68 and 65 pm, respectively, which are comparable in magnitude, hence easy replacement of Li+ ions by Mg2+ ions in insertion compounds is possible; (ii) Mg metal is more stable than Li metal, i.e., Mg can be handled safely in oxygen and humid atmospheres unlike Li, which requires an Ar or He atmosphere of high purity; consequently, safety problems arising from a rechargeable Mg battery are minimal; (iii) global raw material resources are plentiful, and hence, Mg is cheaper than Li, the resources of which are limited and localized. Owing to these merits, investigations on the electrochemistry of Mg-based rechargeable battery systems assume significant importance.

Electrochemistry of Rechargeable Batteries

45

Only a few investigations on rechargeable Mg batteries have been reported in the literature. A brief review of those studies is given below. The reports include the electrochemistry of individual electrode materials, the electrolyte and cell studies. The kinetics of Mg electrode reaction, namely Mg 2 2e- Mg in thionly chloride solutions has been studied by Peled and Stranze [214]. Pulse currents were passed through the Mg electrode and potential/time transients recorded. Differential capacitance and reaction resistance were measured. It was concluded that the Mg electrode in thionyl chloride electrolyte was always covered by a passivating layer. This layer could not be completely removed and therefore no film-free surface could be obtained. However, it was stated that the thickness of the film could be reduced by anodic pulses. Electron transfer from Mg to Mg2+ ions in the electrolyte had a low probability as the thickness of the passivating layer was greater than 20 Å. The reaction resistance of the deposition-dissolution process of Mg metal was mainly determined by the ionic conductivity of the passivating layer. The above studies were extended further to investigate the surface electrochemical properties of Mg by Meitav and Peled [215]. Details of the formation of the passivating layer on Mg at open-circuit potential and during anodic and cathodic polarization were investigated. It was shown that the passivating layer on Mg in thionyl chloride electrolyte formed spontaneously. The reaction resistance was measured from galvanostatic micropolarization in steady state and ac impedance experiments. The differential capacitance was also measured. By assuming an appropriate value for the dielectric constant of the surface-passivating layer, its thickness was calculated. It was shown that the thickness increased from 10 Å to 100 Å in about 100 h of immersion of the Mg electrode in the electrolyte. Anticipating the use of SPEs for rechargeable Mg batteries, PEOMg(ClO4)2 solid films were studied by Patrick and co-workers [216]. Differential scanning calorimetry, ac conductivity, and galvanic cell studies were reported. The specific conductivity followed a non-Arrhenius relationship with temperature. Galvanic cells consisting of Mg anode, Mg2+ ion conducting polymer electrolyte and different cathode materials were fabricated and discharge tests carried out. In order to explore the possible cathode materials that could reversibly intercalate Mg2+ ions, Novak and co-workers studied V2O5 [217] and NaV3O8-type bronzes [218]. In acetonitrile-based electrolytes, specific charge up to 380 mAh g−1 was obtained, but the value decreased rapidly during cycling. With a similar aim, a cation-deficient mixed metal oxide Mn2.15Co0.37O4 were studied for Mg2+ ions insertion by Sanchez and

46

Nanomaterials for Electrochemical Energy Storage Devices

Pereira-Ramos [219]. The electrochemical behavior was studied by cyclic voltammetry and charge/discharge cycling. Specific capacity values in the range of 60 to 30 mAh g−1 were reported. A critical review of various aspects pertaining to secondary Mg batteries was published by Novak and co-workers [220]. Reversibility of electrochemical deposition of Mg from Grignard reagents was investigated for the application of Mg as the negative electrode in rechargeable batteries [221]. Ethyl magnesium bromide in tetrahydrofuran was used as an electrolyte to follow the cathodic deposition of Mg on Ni, Cu, Au, and Ag substrates. Au and Ag substrates were shown to be better than Ni and Cu. Polyether-based polymer electrolytes for Mg battery application were investigated by Acosta and Morales [222]. The polymer electrolytes were characterized by X-ray photoelectron spectroscopy, Fourier transform infrared spectroscopy, differential scanning calorimetry, and ac impedance spectroscopy. High conductivity values were obtained at temperatures around 100°C. A cross-linked polymethacrylate with repeated ethylene oxide units as a solid polymer medium for a Mg salt was investigated by Yoshimoto and co-workers [223]. The ionic conductivity was shown to depend on the kind of dissolved Mg salt. From the direct current (dc) polarization studies, it was shown that Mg2+ ions were mobile in the polymeric system. Research leading to the construction of an ambient temperature, rechargeable Mg battery was discussed by Gregory and co-workers [224]. Experiments were performed based on several organic electrolytes and positive electrodes capable of reversible intercalation of Mg2+ ions. It was reported that Grignard solutions obtained from magnesium halide-organoboron complexes containing aluminum halides were suitable for electrodeposition and anodic dissolution of Mg. A reaction mechanism involving organomagnesium monovalent ion as an intermediate was proposed. Many transition metal oxides, sulfides, and borides were tested for reversible intercalation of Mg2+ ions and with sufficient capacity. The technical feasibility of a secondary battery was highlighted. In conclusion, the authors stated that although there were many similarities in the non-aqueous electrochemistry of Mg and Li, there were also a number of significant differences which made the construction of a secondary Mg battery somewhat difficult. Aurbach and co-workers [225, 226] studied various aspects of the electrolyte and positive electrode materials leading to the development of a coin type secondary Mg cell. A magnesium organo haloaluminate salt in tetrahydrofuran as an electrolyte was employed. The existence of reversibility

Electrochemistry of Rechargeable Batteries

47

of Mg reaction was shown by cyclic voltammetry and electrochemical quartz crystal microbalance. A sulfide of molybdenum was identified as an appropriate host material for reversible intercalation of Mg2+ ions. Coin type cells comprising Mg anode and the cathode were assembled, and several charge/discharge cycles were demonstrated in the voltage range of 1.0 to 1.3 V. There are a few reviews on rechargeable Mg batteries which appeared in the literature in the recent past [227–229]. Saha et al. reviewed extensively on various cell components [227]. When a rechargeable Mg cell consisted of Mg as the negative electrode, the electrolyte needed to facilitate reversible electrochemical deposition of Mg during charging process and oxidation of Mg leading to dissolution of Mg2+ ions into the electrolyte during discharge process. Electrolytes based on non-aqueous solvents such as propylene carbonate, acetonitrile, etc., consisting of Mg salts such as Mg(ClO4)2 were found to be ineffective for electrodeposition of Mg by reduction of Mg2+ ions. Due to this problem coupled with the presence of surface film on Mg metal, which was formed inherently when fresh Mg surface was exposed to even a non-aqueous electrolyte, common electrolytes suitable for this application were absent. Anodic polarization of Mg in some mixed organic electrolyte solutions using solvents such as formamide, acetonitrile, propylene carbonate, tetrahydrofuran, dimethoxyethane, etc., indicated dissolution of Mg at low overpotentials [230]. However, there was no experimental evidence for the electrodeposition of Mg. Magnesium trifluoromethanesulfonate dissolved in dimethylacetamide was found to be a better electrolyte, but then its properties were still not suitable for battery application. It was known that Mg can be efficiently deposited using Grignard reagents in ethereal solvents [231, 232]. Gummow and He showed that dense films of Mg could be deposited using various Grignard’s reagents in ethereal solvents by pulse electrodeposition [233]. The electrodeposited Mg films from Grignard’s reagents were free from dendrites. It was demonstrated that ethereal solutions containing magnesium organochloro aluminate compounds based on the reaction products of some Lewis bases with a variety of Lewis acids dissolved in tetrahydrofuran were electrochemically stable and Mg was electrodeposited reversibly. Cycling efficiency of Mg was more than 95% in some electrolytes. Polymer electrolytes consisting of Mg salts were also studied for rechargeable Mg battery application. Several amidomagnesium chloride salts dissolved in polyethylene oxide and tetrahydrofuran were studied and it was shown that reversible magnesium electrodeposition occurred at potentials below −1.90 V Ag/AgCl, Cl− reference [234]. Polyethylene oxide-magnesium triflate based SPE were investigated

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using propylene carbonate and ethylene carbonate as plasticizers, but the ambient temperature ionic conductivity was low (1 x 10−5 S cm−1) [235]. But the ionic conductivity increased when GPE employing PAN, magnesium triflate, propylene carbonate and ethylene carbonate were prepared. The ionic conductivity at 20°C was 1.8 x 10−3 S cm−1, and it followed Arrhenius behavior for an optimized composition of the GPE. Polyvinylidene fluoride based electrolytes were better than the PAN-based electrolyte films in conductivity and mechanical strength. Mg/GPE/MnO2 cells were assembled and cycled. A maximum discharge capacity of 90 mAh g−1 for about 30 charge–discharge cycles was obtained. However, the discharge capacity decreased due to problems associated with the passivation of Mg surface on repeated charge-discharge cycling of the cells. Magnesium ion conducting GPE with magnesium triflate in an ionic liquid were also reported [236]. The films were free standing with a high ionic conductivity (1 x 10−3 S cm−1), a wide potential window (4 V) and a high thermal stability. The research on positive electrode materials of rechargeable Mg batteries follows similar to the materials used for Li-ion battery, which operate on the principle of intercalation/deintercalation phenomena. The material has to allow Mg2+ ions to undergo insertion and de-insertion without damaging the electrode. Several oxides, phosphates, and sulfides including Co3O4, V2O5, Mg0.5Ti2(PO4)3, and TiS2 were studied [228]. Electrochemical insertion and de-insertion of Mg2+ ions were shown to occur, although the electrochemical activity was poor in some materials. Chevrel phases, MxMo6T8 (M: metal; T: S, Se, Te) are known to be capable of intercalating multivalent cations. Mo6T8 is known to accommodate two Mg2+ ions by exchange of four electrons within Mo6 cluster compounds. The fast intercalation of Mg2+ ions with its inherent high electronic conductivity makes the chevrel phases as preferable positive electrode materials for reversible Mg-ion batteries. However, the voltage of Mg2+ ion insertion/extraction is about 1.1 V vs. Mg/Mg2+, and therefore the cell voltage is low. It was shown that MgxMo6T8 electrodes were composed of single phases for x = 0, 1, and 2 and a mixture of two phases for x = 0.5 and 1.5. Chevrel phase MgxMo6S8 was considered as the best Mg intercalation cathode material in terms of energy density. However, its selenide counterpart has high Mg2+ ion mobility in the entire intercalation range. Chevrel phase Cu2Mo6S8 was synthesized by a high temperature route, Cu was partially leached out by an acid treatment and subjected to electrochemical studies of Mg2+ ion intercalation/deintercalation processes [237]. The cyclic voltammograms corresponded to three pairs of cathodic and anodic peaks related to Mg2+ ion insertion/extraction and simultaneous Cu+ ion extraction/reinsertion phenomena. Galvanostatic charge/discharge cycling of CuxMo6S8 yielded

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49

a stable discharge capacity of about 100 mAh g−1 at C/6 rate. Several other materials such as TiS2, V2O5, V6O13, γ-MnO2, β-MnO2, λ-MnO2, WO3, MoO3, etc., were studied [228] for reversible insertion/de-insertion of Mg2+ ion. However, the cycling behavior was poor. Other cathode materials investigated for Mg-ion battery included MgCoSiO4, WSe2, MoS2, etc. Regarding the anode materials, Mg metal gained interest. However, the metal is plagued with problems related to the formation of protective layer on its surface and poor cycling performance. Similar to the investigations on different negative electrode materials of Li-ion batteries, studies were carried out on carbons, silicon, nanosized tin, bismuth, antimony, bismuth–antimony alloy, etc. Bismuth–antimony alloys provided a high discharge capacity (~300 mAh g−1) in the initial stages of cycling, but the capacity decreased to about 215 mAh g−1 after 100 cycles [238]. The theoretical capacity of Bi is 385 mAh g−1 due to the formation of Mg3Bi2 phase. However, the volume expansion is close to 100% when Mg3Bi2 phase is formed from elemental bismuth. The research and development of rechargeable Mg battery are important because this battery system is a potential alternate to Li-ion batteries. Although, there are considerable limitations of Mg batteries, the resources of Mg are plentiful unlike the scarce resources of Li. The prototype or pre-commercial rechargeable batteries known in the literature are based on Mg metal as the negative electrode and chevrel phase as the positive electrode. However, the cell voltage is only 1.0 V resulting in a low energy density (about 60 Wh kg−1). Extensive research and development efforts are required for realisation of attractive Mg-ion rechargeable batteries.

1.6 Conclusions In the two centuries old history of batteries, successful development of Li-ion batteries in the past quarter century is a commendable achievement. While the application growth is anticipated in an exponential path, the resources of Li are dwindling as also the other metals such as Co useful for manufacture of Li-ion batteries. The foreseen Li-crisis, perhaps similar to oil-crisis, needs a solution from naturally abundant alternate elements such as Na, K and Mg. In the present chapter, several compounds studied for rechargeable Na-ion, K-ion. and Mg-ion batteries are reviewed. The intercalation/de-intercalation mechanisms, and the ionic size dependent issues associated with greater radius of Na+, K+, and Mg2+ ions need clear understanding in identifying appropriate electrode materials for Na-, K-, and Mg-ion rechargeable cells. Although compounds studied till now are

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evaluated for their properties, full cell studies are scarcely reported. Such studies are likely to lead to fabrication of prototype batteries and eventually result in development of commercial level batteries. Another area of interest is to explore aqueous based Na-, K-, and Mg-ion batteries. In these cases, the performance of batteries would be inferior to non-aqueous based batteries. Nevertheless, they would be useful for stationary applications and perhaps substitutes for low energy density rechargeable batteries involving toxic elements presently under use.

References 1. Reddy, T., Linden’s Handbook of Batteries, 4th Edition, McGraw-Hill Education, New York, 2010. 2. Whittingham, M.S., Electrical Energy Storage and Intercalation Chemistry. Science, 192, 1126, 1976. 3. Osaka, T., Homma, T., Momma, T., Yarimizu, H., In situ observation of lithium deposition processes in solid polymer and gel electrolytes. J. Electroanal. Chem., 421, 153, 1997. 4. Kerr, J.B., Polymeric Electrolytes: An Overview, in: Lithium Batteries: Science and Technology, G.A. Nazri and G. Pistoia (Eds.), pp. 575–622, Springer US, Boston, MA, 2003. 5. Gauthier, M., Belanger, A., Kapfer, B., Vassort, G., Armand, M., Solid Polymer Electrolyte Lithium Batteries, in: Polymer Electrolyte Reviews-2, J.R. Maccallum and C.A. Vincent (Eds.), pp. 285–332, Elsevier Applied Science, London and New York, 1989. 6. Long, L., Wang, S., Xiao, M., Meng, Y., Polymer electrolytes for lithium polymer batteries. J. Mater. Chem. A, 4, 10038, 2016. 7. Mizushima, K., Jones, P.C., Wiseman, P.J., Goodenough, J.B., LixCoO2 (02600°C. Soft carbon or graphitizable carbon is prepared from the organic precursors by pyrolysis at a temperature of >700°C. Hard carbons are also prepared from the organic precursor at a temperature range of 700–1500°C, where systems are highly cross linked, which restricts the transformation process to go through the fluid state [4, 30]. a) Graphite The possibility of using graphite as an anode in LIBs by describing the reversible intercalation of lithium in graphite at a potential of 0.2 V (vs. Li/Li+) was reported by Rachid Yazami during 1980s. The theoretical capacity of graphite is 372 mAh g−1 by considering the intercalation mechanism of one lithium per six carbons to form LiC6 as shown

(a)

(b)

(c)

Figure 3.6 The structure of different types of carbon: (a) graphite, (b) graphitizable carbon, and (c) non-graphitizable carbon.

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Nanomaterials for Electrochemical Energy Storage Devices in Eq.  3.4 [31]. This report encouraged SONY Company, Japan to launch a commercial LIB with LiCoO2 as cathode and graphite as anode. Since then, lot of modifications have been come across the carbon-based anodes like graphite, hard carbons, graphene, carbon nanotube and fibers, etc.

Charge

C 6 +Li + +e-

Discharge

LiC 6

(3.4)

Graphite has a layered structure with covalently bonded carbons arranged in a honeycomb lattice constitute an individual layer, called graphene. The individual layers in graphite are separated by a distance of 0.335 nm, which helps in the facile intercalation of lithium. The stable cycling performance with high Columbic efficiency and less volume change (12%) of graphite during lithium insertion helps it to attain unprecedented attention as anode material in the commercial LIB market [32]. The voltage profiles during lithiation and de-lithiation of graphite shows three distinct voltage plateaus [31], indicating the presence of three ranges of compositions (LiC36, LiC18 and LiC6) as shown in Figure 3.7. Graphite electrode can accommodate one lithium for every six Carbon atoms. When graphite electrodes are lithiated, solid-electrolyte 0.4

Volts vs Li/Li+

0.3

0.2

Stage 1’

Stage 4

Stage 4

Stage 2 Stage 2

Stage 1

0.1

0 0

0.2

0.4 0.6 x in LixC6

0.8

1.0

Figure 3.7 Voltage profile of graphite during lithiation and de-lithiation (at C/50 rate) showing three distinct voltage plateaus (Reproduced with permission from Ref. [31]).

Nanostructured Anode Materials for Batteries interphase (SEI) is formed at the electrode surface during first cycle at about 0.25 V due to solvent and electrolyte salt that is electrochemically reduced to oligomers and inorganic crystals like Li2CO3, LiF, Li2O, etc., on the graphite surface as shown in Figure 3.8. A typical SEI, which allow ions to pass but electronically insulating, should have a thickness of 30–50 nm. A good SEI acts as a barrier between the electrode and electrolyte interface and helps to protect further degradation of graphite electrode. Subsequently, SEI plays a major role for kinetics and ionic transport of lithium ions. Many other attempts in modifying the graphite electrode by introducing defects, functional groups, high volume voids (micropores and macropores), etc., are reported to improve its electrochemical performance. The mild oxidation of artificial graphite is reported as an effective strategy to introduce nanochannels or micropores for the facile lithium intercalation and thereby the enhanced electrochemical performance [19]. Further, this strategy can reduce the issue of electrolyte decomposition to a certain extent due to the formation of dense oxide layers. A report shows the mild oxidation of graphite with an aqueous solution of (NH4)2S2O8 delivers a reversible capacity of 355 mAh g−1 [34]. Other than the

Typically 30–50 nm

Li2O LiF Li CO Polyolephines 2 3 2O

L 2O Li LiF

Sem e ica icarr r-

Graphite Particle

Li2CO3 Li2O

LiF rica s Semnate bo

LiF

Li+

Li2CO O3 bo bonate bon nate ates ates

Li2O Li Li2O

L 2CO Li O3

Polyolephines

SEI Interface: Lithium Intercalation into Graphite

Electrolyte

Figure 3.8 Schematic of formation of SEI on graphite electrode during lithiation (Reproduced with permission from Ref. [33]).

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Nanomaterials for Electrochemical Energy Storage Devices strong oxidizing agent such as (NH4)2S2O8, the use of HNO3, Ce(SO4)2, and H2O2 are also reported [34]. The development of metal or metal oxide composites of graphite is another strategy to enhance the electrochemical performance of graphite. The synergistic performance of the graphite and the metal/metal oxide component of the composite toward lithiation/de-lithiation can improve the overall performance of material. The composite of graphite with silver is reported to have enhanced capacity and C-rate performance due to increase in conductivity and formation of additional lithium storage sites in the composites compared with the graphite anode [35]. The composites of graphite with many other metals (Ni, Sn, Al, Zn, etc.) [35–38] and metal oxides (NiO, SnO2, Fe2O3, etc.) [36, 39] are reported to have better electrochemical performance than the conventional graphite anode. Further, the surface modification of graphite through the development of polymer or different type of carbon coating is another strategy to improve the overall performance of the anode. Accordingly, graphite modified with conductive polymer coatings like polypyrrole [40], polythiophene [41], polyaniline [42], etc., are reported as effective anode materials for LIBs. b) Non-graphitic carbon Due to their typical structure, non-graphitic carbon can offer additional sites for the lithium ion storage. The insertion of lithium ion in soft/hard carbon depends on the precursor materials as well as the heat treatment process. The average interlayer distance of hard and soft carbon deviated from that of graphite as follows, graphite—0.335 nm, soft carbon—0.375 nm, and hard carbon—0.38–0.523 nm [4, 43]. In general, hard carbon can deliver capacity of 500 mAh g−1 or more in the potential range of 0.01 to 1.5 V vs. Li/Li+. In 1991, Kureha Corporation (Japan) commercially used hard carbon as the negative electrode materials in the first built lithium-ion battery. In analogy to graphite, an irreversible capacity of about ~40% is observed in hard carbons [44], which can be attributed to the SEI formation during first cycle at about 0.75 V and side reactions between electrolytes with functional groups present on the carbon surfaces. However, the reversible capacities are found to be stable in the subsequent cycles. A typical charge/discharge

Nanostructured Anode Materials for Batteries curves of hard carbon derived from coal tar pitch is shown in Figure 3.9 [44]. Diverse nanostructured carbon materials have been developed using several precursors and synthetic strategies. In recent times, biomass-derived non-graphitic carbonaceous materials have drawn huge attention due to its low cost, environmental friendliness and accessibility [45]. These materials can be generated from a wide variety of sources such as cellulose, sweet potato, peanut shells, lignin, cornstalk, bamboo chopsticks, sea shells, corn starch, pinecone hull, banana skins, rice husk, cherry stones, and ox horns, etc. [45–47]. Biomass derived carbons and composites have delivered very good capacity, cycling stability with improved C-rate performance. These are the reasons that considerable attention has been put forward to develop cost-effective, high performance non-graphitic anode materials for LIB application. Heteroatom such as B, N, S, Si and P [47–52] used as dopant to substitute carbon atoms in the carbon lattice to change the electronic distribution in the system, which can enhance the electrochemical performance. E.g., boron (also Sulfur) acts as electron acceptor in the carbon host and accordingly, the electronic and ionic conductivity of the overall matrix as well as the intercalation potential is shifted, structure also

2.0

2nd D

1st D

Voltage (V)

1.5

1.0

0.5 2nd C 1st C

0.0 0

100 200 300 400 500 600 700 800 900 Specific capacity (mAh g–1)

Figure 3.9 The typical charge/discharge profile of non-graphitized (hard) carbons prepared at 700°C (a) first cycle, (b) second cycle (Reproduced with permission from Ref. [44]).

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Nanomaterials for Electrochemical Energy Storage Devices becomes more disorder [50]. Nitrogen can also be incorporated into carbon, can act as electron donor and shows the opposite effect to that of B-doping [50, 52]. Charge and discharge curves of pyrolytic carbon in comparison N and S doped carbons are given in Figure 3.10. c) Carbon nanotube CNTs have gained huge interest because of its unique structural, electrical, mechanical and electronic properties since its discovery by Iijima in 1991 [53]. Li+ can intercalate either on the outer surface or the inner surface of CNTs. Due to the flexible structure and morphology, CNTs can offer a stable capacity by eliminating the pulverization issue in the electrode. Literature reports suggest that CNTs can deliver capacity less than graphite as well as higher capacities [54–56]. The reversible electrochemical insertion of Li+ into single walled CNTs (SWCNT) and multi-walled CNTs (MWCNT) host demonstrate the capacity of 80–600 mAh g−1 [55–59]. Leo et al. reported SWCNTs produced by laser vaporization can deliver capacities > 1050 mAh g−1 [57]. Achieving high coulombic efficiency with CNTs remains challenging because of the presence of large structure defects and high voltage hysteresis during charge and discharge. Besides, production cost of pure CNTs restricts its use in practical systems. d) Graphene Graphene, which is single layer of graphite, honeycomb-like structure of sp2 carbons bonded into two-dimensional sheets of carbon. Graphene has gained huge attention in LIBs due to high electronic conductivity as well as minimal volume (b) 3.5 3.0 2.5 2.0 1.5 C7.3N 1.0 C12.5N 0.5 Pyrolytic carbon 0 0 400 600 800 200 Capacity (mAh/g) Voltage (V)

Voltage (V)

(a) 3.5 3.0 2.5 2.0 1.5 1.0 0.5 0 0

C28.0S C31.9S Pyrolytic carbon

200

400

600

800

Capacity (mAh/g)

Figure 3.10 Charge and discharge curves of pyrolytic carbon in comparison with (a) CxN (x = 12.5 and 7.3) and (b) CxS (x = 31.9 and 28.0) electrodes (Reproduced with permission from Ref. [51]).

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changes during intercalation/de-intercalation of lithium in comparison with the other alloy-based anode materials which can deliver high capacity ranging between 780 mAh g−1 to ~1000 mAh g−1 [60–64]. These values are associated with the interaction between lithium and carbon ending up Li2C6 stoichiometry and LiC2 stoichiometry. Graphene can offer high reversible capacity, good C-rate performance as well as good cycling stability due to its one-atom-thick 2D structure, large surface area, electrical and thermal conductivity, and superior mechanical properties. Disordered graphene sheets prepared by reduction of graphite oxide (GO) by various methods, such as low temperature reduction, chemical reduction by hydrazine was electrochemically tested and demonstrated the gravimetric capacity in the range between 790 and 1050 mAh g−1 [63]. This high capacity can be explained due to presence of active sites and defects, which provides the space for electrochemical insertion/ de-insertion of lithium ion. However, these materials suffer from high irreversible capacity due to SEI formation in the 1st lithiation process. Few-layer graphene, with surface area more than 490 m2 g−1 showed the reversible capacities of 1200 mAh g−1 during initial charge–discharge process, and with the capacity retention of 848 mAh g−1 at the end of 40th cycles (Figure 3.11) [64].

4.0 3.5 discharge-1st charge-1st discharge-2nd charge-2nd discharge-3rd charge-3rd discharge-4th charge-4th discharge-5th charge-5th

3.0

Voltage/V

2.5 2.0 1.5 1.0 0.5 0.0 0

500

1000

1500

2000

Capacity/mA h g–1

Figure 3.11 First five discharge/charge profiles of graphene sheets at the current density of 100 mA g−1 (Reproduced with permission from Ref. [64]).

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3.2.3.2.2 Titanium Dioxide The structural stability, good physico-chemical properties and reasonable cost of Titanium dioxide (TiO2) making it attractive to be used as a promising anode material for LIBs. TiO2 has theoretical capacity of 335 mAh g−1 along with other properties such as low volume change of 850°C [64]. Initially rutile phase was considered to be electrochemically inactive toward lithium at room temperature. Hence, rutile TiO2 was used as anode material at 120°C by Macklin et al. [74] for the first time. Later nanosized rutile phase was reported to be active even at room temperature and attained a capacity of 200 mAh g−1 was obtained over 20 cycles by reducing the particle size to 15 nm [75]. Eventually, capacity was reduced to 50 mAhg−1 when size was increased to 300 nm [75]. So, the improvement in the capacity is results of the nanosized material as well as high surface area. Figure 3.13 shows the charge–discharge curve of anatase and rutile nanoparticles having different sizes. The excellent cyclability observed for the nanosized materials can be attributed to the higher reversibility of the

Nanostructured Anode Materials for Batteries (a)

(b)

165

(c) Ti Ti Ti O O

O Ti

Ti

Ti O

O

Ti

Ti

O

O O Ti

Ti O Ti

O

O Ti

Ti

Ti Ti

Ti

O

Ti

O O Ti Ti

O O

Ti O

Ti O

O

Ti O

Ti O

O

O

O

O O

O

Ti O

O O Ti

O

O

Ti

O O Ti Ti

Figure 3.12 The crystal structures of TiO2 nanoparticles; (a) rutile, (b) anatase, and (c) brookite (Reproduced with permission from Ref. [72]).

3.5

3.5

Voltage (Li+/Li)

3.0 2.5 2.0 A6 A15 A30

1.5 1.0

Voltage (V vs. Li+/Li)

0.1 A/g 3.0 2.5 2.0

R15 R30 R300

1.5 1.0

0

50

100

150

Capacity (mAh/g) (a)

200

250

0

50

100 150 200 250 300 350 400

Specific capacity (mAh/g) (b)

Figure 3.13 (a) Voltage profile of anatase nanoparticles with diameters 6 (A6), 15 (A15), and 30 (A30) nm cycled at 0.1 A g−1, and (b) the initial capacity profiles of rutile nanoparticles with diameters 15 (R15), 30 (R30), and 300 (R300) nm, cycled at 0.05 A g−1 (Reproduced with permission from Ref. [77]).

Li ions inserted in both the bulk and surface of the fine rutile particles and also to the better accommodation of the volume changes upon Li+ insertion/extraction [75, 76]. The conductive phases such as carbon, polymers, and metals additives have been incorporated into the TiO2 material to further improve its capacity retention and rate capability. Generally, TiO2 is reported to be more electrochemically active in its anatase phase than rutile phase. According to the electrochemical intercalation/de-intercalation of lithium as given in Eq. 3.5, micro-sized anatase TiO2 can intercalate up to 0.5 Li with negligible volume change (~4%) [78]. Many strategies like synthesizing materials with different

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morphologies and by reducing particle size into nanoscale, incorporating carbonaceous materials, etc., are reported to be effective to improve the overall performance of anatase TiO2 [77]. In particular, it has been observed that when particle size is moving toward nano level, Li-ion diffusion is enhanced with short path length which in order facilitates the lithium ion intercalation/deintercalation. In terms of lithium intercalation efficiency, TiO2(B) is the best among all other polymorphs of TiO2 due to its large unit cell (monoclinic structure) with more open crystal structure (long a-axis of 1.216 nm) [79]. Marchand et al. in 1980 first reported the synthesis of TiO2 (B), which involves high-temperature solid-state reaction between TiO2 and K2CO3 to form K2Ti4O9, later hydrolysis of these materials leads to TiO2 (B) [80]. The unique open crystal structure of TiO2 (B) allows to accommodate a theoretical maximum of 1.25 Li per formula unit, which is even larger as compared to anatase (1 Li per formula unit). The open channels along the b-axis facilitate the Li-ion diffusion and thereby improves the kinetics of Li-ion insertion and C-rate performance. Hence, the modifications in structure and morphology of TiO2 (B) to obtain facile Li-ion diffusion pathways can offer better results in terms of the electrochemical performance. Further to improve power density as well cycle life, hybrid of TiO2 nanoparticles with nanostructure carbon matrix such as CNT, graphene has been explored [70, 71, 81]. Although TiO2 shows good cycle life (over 1000 cycles) and C-rate performance, its practical applications as anode for LIBs are limited due to low capacity and high operating plateau voltage.

3.2.3.2.3 Lithium Titanium Oxide (LTO) Like TiO2, the lithium titanate, Li4Ti5O12 (LTO) is also a promising anode material due to its flat plateau voltage of 1.55 V vs. Li/ Li+, and theoretical capacity ~175 mAh g−1. Unlike the conventional graphite anode, LTO does not form SEI during lithiation and serves as a zero strain insertion material with good cycling stability over 1000 cycles [82–85]. It has a cubic spinel structure and can be expressed as Li4/3Ti5/3O4 or in the spinel form (Li[Li1/3Ti5/3]O4). The electrochemical lithium intercalation reaction in LTO is given in Eq. 3.6.

Li 4 Ti5O12 +3Li + +3e-

Li 7 Ti5O12

(3.6)

Although the irreversible capacity of LTO is negligible due to the absence of any side reactions with the electrolyte, the low electrical conductivity of the material limits its electrochemical performance as anode for LIBs. The

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various research reports on modifying the LTO shows that the morphology and particle size plays a crucial role in the electrochemical performance. Accordingly, nano-sized LTO powders and three-dimensionally ordered macroporous (3DOM) LTO are reported to have capacities of 125 mAh g−1 and 145 mAh g−1, respectively [85]. The increased capacity of 3DOM compared to the nano-sized LTO is attributed to the porous morphology that helps in facile diffusion of lithium ions [85–87]. The possibility of using rutile TiO2 coating on LTO as a carbon-free nanocoating to improve the lithium intercalation/de-intercalation kinetics is reported as an efficient strategy by Wang et al. [81]. The bead-like structure of Li4Ti5O12 having 0.3 mm size produced from Li2CO3 and anatase TiO2 precursors by using solid-state synthesis shows excellent electrochemical performance with a capacity retention of 80% at 10 C rate [82]. Similarly, Wang et al. [81] also reported the enhanced performance of the Li4Ti5O12 synthesized through solid state synthesis with a capacity retention of 98.25% at 10 C rate. Many synthesis routes like sol–gel, microwave, molten-salt, hydro-thermal methods are reported to enhance the performance of LTO anode by tailoring its morphology and particle size [82–87]. Surface modification of the LTO by developing surface coatings such as metal oxides and carbons (carbon nanotubes, graphene, etc.) are reported to be a promising strategy to overcome the inherent limitations of LTO as anode material for LIBs [88, 89]. The carbon coated Li4Ti5O12 nanorods synthesized using hydrothermal method shows 92.7 mAh g−1 more capacity than the pristine material with a good cycling stability [88]. The carbon coating achieved from petroleum pitch precursor on to Li4Ti5O12 is reported to have 95% capacity retention even after 1000 cycles [89]. Further, graphene-coated Li4Ti5O12 is reported to have better initial discharge capacity (168 mAh g−1) and cycle life than the pristine one. The strategy of modifying the Li4Ti5O12 by using metal coating (e.g.: Ag) [90] and metal oxide coatings (e.g., Fe2O3, CuO, ZrO2, etc.) [91–93] are observed to be effective in enhancing the conductivity and thereby the overall performance. The composite obtained by the activation of Li4Ti5O12 using 3% Ag coating shows a capacity of 131 mAh g−1 at 30 C rate [94]. Although coatings can improve the surface conductivity, they are not effective in enhancing the conductivity of the crystal lattice. Hence, the idea of substituting a small percentage of Ti3+, Li+ ions with other metal ions were attained much attention to enhance the internal electronic conductivity of LTO. In this regard, magnesium (Mg) [95], Strontium (Sr) [94] and copper (Cu) [94] doping is reported to be effective in improving the electrical conductivity, decreasing the charge transfer resistance and

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improving the diffusion coefficient of Li+, respectively. Apart from these metal ions, many other cation and anion doping are also reported to be effective in improving the intrinsic conductivity and thereby the overall performance of the LTO as anode material for LIBs. Apart from the individual cation and anion doping, co-doping of the anion and cation into the LTO structure such as Al3+ and F− [96], and La3+ and F− [97] are also reported to have favorable results in terms of overall enhancement in the electrochemical performance. LTO is generally used along with olivine cathodes such as LiFePO4 or LiMnPO4 to deliver high C rate capacities and fast charging of LIBs.

3.2.3.3 Alloying/De-Alloying Materials Alloy anodes have been emerged as the most promising anode materials for LIBs owing to their high specific capacities (Figure 3.14) and safe operating potentials. The elements and their compounds (alloys or intermetallics) that can alloy and de-alloy with lithium electrochemically at room temperature are coming under this type of anode materials. Although, many elements such as Si, Sn, Sb, Al, Mg, Bi, In, Zn, Pb, Ag, Pt, Au, Cd, As, Ga, and Ge are electrochemically active toward lithium, the elements such as Si, Sn, Sb, Al, and Mg are only studied extensively due to their natural abundance, cost effectiveness and environmentally benign nature [98–102].

4500 4000

M° + xLi + xe–

LixM

R. Huggins (20 years)

Capacities (mAh/g)

3500 3000

Staggering capacity gains over C?

2500 Poor cycle life owing to large volume swings (up to 300%)

2000 1500 1000 500 0 In

C

Bi

Zn Te

Pb Sb Ga Sn

Al

As Ge

Si

Figure 3.14 A comparison of the specific capacities of different types of alloying and de-alloying type anode materials.

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The reaction mechanism of alloy anode (element of intermetallics) can be of three different types: i) insertion reaction leading to the formation of solid solution, ii) addition reaction and iii) displacement reaction [102, 103]. The reaction taking place during alloying and dealloying of Li (lithiation and de-lithiation) in the first two types is shown in Eq. 3.7.

Li+xM

LiM x

(3.7)

In insertion and addition type, the Li ions are added into the anode material phase (M), without displacing any elements from the material. However, the insertion reaction is a topotactic reaction leading to the formation of a solid solution after lithiation, without any change in phase or structure of the anode material. The insertion type material maintains same phase structure before (M) and after lithiation (LiMx). Whereas, materials following addition type reaction undergoes phase and structural changes during lithiation. The phase structure of the lithiated anode material (LiMx) is different from the parent material (M). The reaction of Li with Mg and amorphous Si are good examples of insertion type alloying/dealloying mechanism leading to the formation of solid-solutions. At the same time, due to the limited solubility of Li in Sn, Sb, crystalline Si, etc., follows addition type mechanism during lithiation/de-lithiation [102–105]. In displacement reaction, Li reacts with one component of the compound MNy (alloy or intermetallic) and the other component (N) is displaced from the parent material during lithiation. The displaced element (N) can be active or inactive toward Li. If N is inactive (e.g., Cu6Sn5, CrSb2, etc.), the mechanism of displacement reaction during alloying is represented in Eq. 3.8.

Li+xMN y

LiM x +xyN

(3.8)

As represented in Eq. 3.8, some of the displacement reactions are irreversible and the displaced component does not participate in the further alloying/dealloying processes but acts as a buffering matrix. So the alloying/dealloying after displacement of the inactive element becomes addition reaction as represented in Eq. 3.7. If the displacing element (N) is also active toward Li (e.g., SnSb, Mg2Si, etc.), the displaced element also undergoes addition type reaction with Li at potential lower than that of M [106, 107]. However, the active/active alloy compounds show displacement reactions are also not reversible after many cycles and the less active material becomes inactive buffering matrix in later cycles.

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Although, the specific capacity of these alloy based anodes like Si (4200 mAh g−1) [108], Ge (1600 mAh g−1) [109], Sn (994 mAh g−1) [110], etc., are more than that of graphite (372 mAh g−1) and LTO (175 mAh g−1), the major drawbacks of poor cycling stability and large irreversible capacities at the initial cycles limit their wide application. The poor cycle life of the alloy anodes is due to their large volume changes (up to 300%) during lithiation/de-lithiation, which can result in pulverization of the active alloy materials and lose of electrical contact. The cause of large irreversible capacity for alloy anodes are due to the loss of active material through large volume changes and pulverization, formation of SEI, permanent trapping of Li ions in the metal/alloy matrix, reaction of Li with surface oxide layer, aggregation of alloy particles, etc. [100–107]. Several strategies have been reported to surmount the major challenges of capacity fade during cycling and large irreversible capacity of alloy anodes by reducing the detrimental effects of large volume change and side reaction with the electrolyte. Examples of some of the alloying/dealloying type anode materials are discussed here.

3.2.3.3.1 Silicon Silicon is one of the most promising anode materials for LIBs due to its unique advantages like high theoretical capacity (4200 mAh g−1), natural abundance (second most abundant element), non-toxicity, low cost, etc. [111, 112]. Although silicon offers a high theoretical capacity at its fully lithiated state (Li4.4Si), its wide practical applications are limited due to the large volume change (~400%) [113] during charge–discharge cycling leading to pulverization of the electrode [114]. The volume change and pulverization can result in the gradual loss of electrical contact between the active material (silicon) and current collector, and thereby leading to rapid capacity fade and failure of the cell. The poor electronic conductivity is another limitation of silicon anode that impedes its rate capability. During initial cycles, a thick SEI layer is formed on the surface of Si anode due to the decomposition of salts and electrolyte reduction. The formation of thick SEI layer in Si anode increases the cell resistance and thereby reduces the overall performance of the cell. Owing to the attractive properties of Si as anode material for LIBs, lot of research efforts have been devoted to address the inherent issues of silicon and to enhance its electrochemical performance. The development of nanostructured Si anode was the first among them to address the issue of pulverization by providing enough free space for volume expansion during lithiation/de-lithiation. Moreover, the nanostructured electrode can ensure

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better specific capacity and rate capability by providing good electronic contact and facile lithium ion diffusion. The nanostructured Si having different morphologies such as nanowires [115, 116], nanotubes [117] and interconnected hollow nanospheres [118], and different nanoscale designs in the form of thin films [119, 120], vesical [121], and nanoparticles [122] are reported to be very effective in addressing the issue of volume change and thereby to circumvent the pulverization of the electrode. The nanoengineering of developing an electrode having silicon nanowires grown directly on current collector is reported to deliver a charge capacity (4,277 mAh g−1) almost equal to the theoretical charge capacity (4,200 mAh g−1) [115]. The unique electrode architecture of silicon nanowires provides 1D electronic pathways for efficient charge transfer and also eliminates the prevailing issue of pulverization. The nanoengineered electrode architecture of Si nanowires is attracted as the most promising anode material for LIBs with high discharge capacity (3,124 mAh g−1) and cycling stability [115]. The incorporation of carbon materials into silicon and thereby to develop silicon–carbon (Si–C) composite is attracted as an effective strategy to address the inherent issue of low conductivity and volume change of silicon-based electrodes. Composites of polymers like polystyrene, polyvinyl alcohol, polyvinyl chloride, resorcinol-formaldehyde resin, etc., and different forms of carbons like graphite, carbon nanotubes, graphene, etc. [123–129] were used as carbon matrix in the composite to mitigate the volume change of silicon. Other than Si–C composite anodes, composites of silicon with metals like silver (Ag) [130] and copper (Cu) [131], etc., to get Si–M (M = Ag, Cu) composite is also reported as effective strategy to improve the conductivity of the electrode. However, the capacity retention and cycle life of Si–M composite electrodes are not impressive because of the poor contact of silicon particles with metal up on cycling due to huge volume change. Although dense films of silicon were used earlier as anodes for LIBs, it was replaced later due to the inferior performance due to volume expansion and pulverization. As pulverization is the major failure mode of silicon-based anodes, the mass of the active material and the adherence of the active material to the substrate are crucial in the performance of the anode. Accordingly, an amorphous silicon thin film electrode having a thickness of 200 nm fabricated through magnetron sputtering is reported to have a capacity of 3000 mAh g−1, achieved as a result of the loading density optimization and intact electrical contact [132]. The development of electrode having core-shell of silicon and carbon is reported as the most effective strategy to improve the overall

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electrochemical performance rather than developing a conventional Si–C composite [121]. The enhanced performance of this electrode design is attributed to the presence of a complete but thin carbon layer over the silicon that can improve conductivity and provide mechanical stability against volume change. Among the various electrode architecture of silicon reported, binderless electrode designs have attained much attention due to its enhanced electrochemical performance. The removal of inactive binder component can reduce electrode resistance and hence can improve overall electrode performance. Beside, removal of binder which contribute to extra weight can increase specific energy of the cell. A binder-free and additive free silicon anodes developed by pressure-embedding of Si nanoparticles onto the surface of copper foil current collector shows an initial discharge capacity of 1250 mAh g−1 (C/10 rate) (Figure 3.15), and retains more than 650 mAh g−1 (C/2) capacity even after 500 cycles [133]. A comparison of the electrochemical performance of conventional and pressure-embedded Si nanoparticle electrodes are given in Figure 3.15. Though the conventional Si nanoparticle electrode shows high initial capacity (~3000 mAh g−1) than the pressure-embedded one, the capacity retention and cycling stability are more for the pressure-embedded electrode that achieved by overcoming the issue of pulverization. The pressure-embedding of Si nanoparticles resulted in an electrode with evenly distributed nanoparticles having enough free spaces in between to accommodate the volume change, and thereby effectively alleviated the issue of pulverization.

1.5

E/V

1.0 Lithiation-conventional Delithiation-conventional Lithiation-Si-NPs on copper Delithiation-Si-NPs on copper

0.5

0.0 0

500 1000 1500 2000 2500 3000 Capacity/mAh g–1

Figure 3.15 Galvanostatic charge–discharge voltage profile for the composite (Si 50 wt.%, Binder 10 wt.%, Carbon 40 wt.%) and pressure embodied Si–NPs on copper foil current collector electrode during first cycle at C/10 rate (Reproduced with permission from Ref. [133]).

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Voltage (V) vs. Li/Li+

2.5

173

1st lithiation 1st delithiation

2.0

2nd lithiation 2nd delithiation

1.5

1

1.0

2

0.5

2

1

0.0 0

500

1000

1500

2000

2500

Capacity (mAh g–1)

Figure 3.16 Voltage profiles during lithiation and de-lithiation process for the freestanding electrode annealed at 1000°C (Reproduced with permission from Ref. [133]).

The 2D structured silicon thin films and nanosheets are reported as efficient anodes for LIBs due to their advantages of small volume change and faster lithium diffusion. Other than 1D and 2D structured Si electrodes, 3D electrode architectures also attained much attention as promising anodes for LIBs. The 3D Si–Al alloy electrode developed on copper foam exhibits enhanced performance as anode material due to the microporous structure of copper foam that facilitates the electrode/electrolyte interaction and provides free space to accommodate the volume change. In another report, 3D Si–C free-standing electrode developed on carbon fiber current collector shows a reversible capacity of ~1000 mAh g−1 at 5C rate over 100 cycles (Figure 3.16) [133]. The electrode architecture provides a new strategy of electrode development by replacing the conventional organic binder and copper current collector by a carbonized p-pitch and carbon fiber, respectively. If the issues of pulverization of Silicon based anodes can be mitigated, these anodes could find a room for the next generation anodes for high energy density LIBs.

3.2.3.3.2 Germanium-Based Materials Ge is an extensively studied alloy type anode material having high specific capacity (1600 mAh g−1). Ge has higher intrinsic electronic conductivity (104 times) than Si, narrow band gap (0.67 eV) and faster Li+ diffusion rate [134, 135]. However, like other alloy anodes, Ge based materials are also vulnerable to high volume expansion (~300%) and thereby prone to pulverization during lithiation/de-lithiation [134]. Similar to Si, the

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limitations of Ge bulk materials were mitigated to certain extend through the development of nanostructures such as nanoparticles, nanowires, and nanotubes. Further, the issue of large volume change was addressed by developing highly porous 3D structures and by using carbon as a buffer material [136–138]. The hybrid composite materials of Ge were found to have noticeable variation in performance than the conventional materials.

3.2.3.3.3 Sn-Based Materials The excellent electrochemical alloying behavior of Sn with Li can be evident from the equilibrium phase diagram of Li–Sn with a number of intermetallic compositions like Li22Sn5, Li7Sn2, Li3Sn, Li5Sn, LiSn, Li2Sn5, etc. [139]. Hence, the specific capacity of the electrode can vary up to 994 mAh g−1 (for Li22Sn5) depending on type of intermetallic formed during lithiation. The large volume change during lithiation/de-lithiation reactions is the major challenge in the commercialization of Sn based anodes. The volume change during lithiation arises mainly due to the sharp density difference between Sn (7.365 g cm−3) and Li (0.534 g cm−3), which leads to the formation of intermetallics having higher Li content [140]. The first strategy followed to reduce the large volume change during cycling was to reduce the particle size of the anode material. The nanosized Sn particles showed better performance in comparison to bulk material but the cycling performance was not sufficient. Moreover, the irreversible capacity in the initial cycles were large due to increase in the amount of Li consumed for SEI formation with decrease in particle size. Including a second phase like material such as carbon in the matrix, volume change is reduced during cycling. The effect of many such second phase materials like disordered carbon [141], graphite [142], semi-amorphous carbon [143], CNT [144], TiO2 [145], semi-amorphous Cu [146], etc., have been reported. Sn based intermetallics and their composites are also reported as good anode materials for LIBs. The widely studied intermetallics to achieve better performance than the parent Sn metal includes Cu6Sn5 [147], SnSb [148], CoSn3 [149], CoSn [150], Co3Sn2 [151], CoSn2 [152], Ni3Sn4 [150, 151], Sn2Mn [152], LaSn3 [153], Ag3Sn [154], SnAg4 [154], Mo3Sn [155], TixSny [156], Mg2Sn [157], SnMn3C [152, 157], Sn2Fe [158, 159], FeSn [159], Fe2Sn3 [158], Fe3Sn5 [159], V2Sn3 [152], etc. The intermetallics containing inactive component can act as a buffering agent toward volume change as schematically illustrated in Figure 3.17. In the case of SnSb, which contains two active elements toward Li was found to form different lithiated phases such as Li3Sb and Li22Sn5. When Li3Sb is formed, the Sn phase acts as a buffer to volume expansion [161].

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Problem of Sn anode Drastic volume change during charge and discharge 4.4Li+

4.4Li+ Li4.4Sn

Sn Charge

Sn Discharge

Poor cycle durability

breakup

Alloying with inactive metal with lithium 4.4Li+ Ni-Sn alloy

4.4Li+ Li4.4Sn

Charge

Ni-Sn alloy Discharge Stable structure

Ni rich phase Long cycle durability

Figure 3.17 Schematic of the volume change in Sn and Sn alloy containing inactive component Ni (Sn–Ni alloy) (Reproduced with permission from Ref. [160]).

Based on the effective performance of intermetallics, in 2005 Sony corporation commercialized Sn–Co based alloy (ratio of Sn:Co is ~1.1:1 mol, with possible titanium of ~5%) anodes for LIBs. This novel system was able to reduce the large volume expansion during lithiation/de-lithiation and thereby enhancing the performance. The ternary alloy systems like Sn–Sb–Ni, Sn–Sb–Cu, etc., were also reported as anode materials for LIBs. The composites of the binary alloys with carbon like Sn–M–C (M = Fe, Cu, Co, Ni) were also reported to address the issues related to Sn base anodes.

3.2.3.3.4 Tin Oxide The SnO2 attained significant attention as anode material for LIBs due to high theoretical capacity (783 mAh g−1). SnO2 is an example for alloy anode materials showing displacement type reaction mechanism [162, 163]. The electrochemical alloying reaction involves a partially irreversible step as in the general Eq. 3.8, followed by the alloying/de-alloying reaction of Sn and Li as presented in Eq. 3.7 [162]. Although SnO2 shows a high theoretical capacity of 1491 mAh g−1, it shows high irreversibly capacity (~50%), vulnerable to large volume changes (>200%) and thereby capacity fade during cycling. To reduce the problems of volume change, porous nanostructures,

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nanocomposites, and hollow nanostructures of SnO2 have been reported. The presence of pores in these structures can balance the volume change during alloying and dealloying to certain extent. Different C based composites of SnO2 like carbon coated SnO2, SnO2/ carbon nanofibers, SnO2/carbon nanoparticles, SnO2/CNT, SnO2/graphene, etc., are also reported as efficient anode materials for LIBs [164–166]. The C based composites are found to show good performance in reversible capacity, Coulombic efficiency and cycling life. Moreover, different morphologies of SnO2 structures like nanowires, nanotubes, nanorods, nanoboxes, nanosheets, etc., shows better cycling stability due to its reduced volume change and mechanical stress.

3.2.3.4 Conversion Type Anode Materials The conversion type materials store lithium through reversible replacement reactions between Li+ and transition metal cations and offers high theoretical specific capacities. A wide range of transition metal oxide, sulfides, selenides, fluorides, nitrides, and phosphides have been reported as conversion type anode materials for LIB applications (Table 3.1) [167–171]. The mechanism of conversion reaction involves a redox reaction between Li+ and transition metal compounds (MaXb, where, M = Mn, Fe, Table 3.1 Conversion type anode materials for lithium ion batteries. Theoretical capacity Active anode material (mAh g−1) a. Metal oxides 500–1200 (Fe2O3, Fe3O4, CoO, Co3O4, MnxOy, Cu2O/CuO, NiO, Cr2O3, RuO2, MoO2/MoO3, etc.) [167–200] b. Metal phoshides/ 500–1800 sulfides/nitrides (MXy; M = Fe, Mn, Ni, Cu, Co, etc., and X = P, S, N) [29, 103, 167, 172, 201–210]

Advantages Common issues High capacity, High Low Coulombic energy, Low cost, efficiency, Environmentally Unstable SEI compatibility formation, Large potential hysteresis, Poor cycle life High specific Poor capacity capacity, Low retention, operation Short cycle potential, and life, High cost Low polarization of production than counter oxides

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Co, Ni, Cu, etc., and X = O, S, Se, F, N, P, etc.) leading to the formation and decomposition of a lithium binary compound (LinX) with high theoretical specific capacity as given in Eq. 3.9 [168]. A schematic of the mechanism of conversion reaction is shown in Figure 3.18.

M a X b +(b.n)Li + +(b.n)e-

aM+bLin X

(3.9)

The transition metal compounds having reaction potential (which is mainly determined by the ionization of M–X bond) in the range between 0.5 to 1.0 V vs. Li/Li+ can act as good anode material for LIBs. Although the formation of LinX is a thermodynamically feasible reaction, the reverse conversion reaction of electrochemically inactive LinX by bulk material M is difficult to happen. The reversibility of the reaction through the decomposition of the LinX matrix surrounded by the SEI layer, highly depends on the formation of electroactive M nanoparticles [168, 172]. The voltage hysteresis during charge/discharge process is an another issue in conversion type anode materials. It mainly originates from the extensive structural rearrangement during lithiation/de-lithiation leading to internal heat evolution, resulting low roundtrip energy efficiency. The voltage hysteresis also shows specificity toward the nature of the anionic species present in the conversion type anode material, and it varies in the order fluorides> oxides>sulfides>nitrides>phosphides [172–174]. Moreover, the poor rate capability and fast capacity fade due to low intrinsic conductivity and pulverization during repeated cycling are also the limitation of conversion type of anode materials. Apart from all these issues, the highly reactive M nanoparticles formed during lithiation can bring out some partially reversible side reactions through the decomposition of the electrolyte [171, 172]. A considerable effort has been devoted in the past two decades to circumvent the issues related to the conversion type anode materials. The

M nanocrystals (2~8 nm) LinX

MaXb nanoparticles 1st lithiation

LinX SEI layer

Delithiation

Current collector

Current collector

MaXb

M + LinX

Transition metal compounds (MaXb)

MaXb nanocrystals

Metal (M)

Lithiation

Current collector

MaXb + irreversible LinX LinX

Irreversible LinX

SEI layer

Figure 3.18 Schematic of the mechanism of conversion reaction in transition metal compounds (Reproduced with permission from Ref. [172]).

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strategy of bringing down the electrode material to nanoscale range is reported to be effective in enhancing the electrochemical activity toward the decomposition of LinX and intrinsic conductivity of electrodes. Some of the conversion type anode materials for LIB are given below.

3.2.3.4.1 Iron Oxide Iron oxides such as hematite ( -Fe2O3) and magnetite (Fe3O4) are reported to be effective conversion type anode material for LIBs with good theoretical capacities of 1007 and 926 mAh g−1, respectively [174]. The natural abundance, nontoxicity and low cost of iron oxides made it more attractive for LIB applications. However, the issues like poor cycling performance due to low electrical conductivity, low Li+ diffusion, high volume change, iron aggregation (poor reversibility of conversion phases), etc., during the charge/discharge process are the main limitations of these materials [174, 175]. Several approaches like developing iron oxide nanomaterials with different size, shape, porosity, by developing carbon based nanocomposites, 3D electrode architectures, etc., have been adopted to surmount these identified drawbacks of iron oxide [174–178]. For example, the -Fe2O3 hollow spheres synthesized by a facile quasiemulsion template method [179] shows a capacity of 700 mAh g−1 with a cycling stability more than 100 cycles. Similarly, different morphologies of Fe2O3 such as nanorods (300 to 500 nm), nanotubes (200 to 300 nm) and their carbon coated derivatives were also reported as anode materials with good capacity and cycling stability over 50 cycles [180]. The carbon coated crystalline Fe3O4 nanowires (20 to 50 nm) showed a reversible capacity of 830 mAh g−1 over 50 cycles. The spindle-like -Fe2O3 from an iron based metal organic structures showed good reversible capacity of 911 mAh g−1 for more than 50 cycles at 0.2C rate and a capacity of 424 mAh g−1 at 10C rate [181]. The nanocomposites of Fe3O4 cores and porous carbon-silicate layers synthesized through aerosol assisted process followed by vapor coating was found to show quite stable capacity with coulombic efficiency close to 100% over 50 cycles. The porous Fe3O4 nanotubes prepared using microporous organic nanotubes as template were also found to show excellent electrochemical performance with high capacities of 918 and 882 mAh g−1 at current densities of 500 and 1000 mA g−1, respectively [182]. Charge– discharge voltage profiles of Fe3O4 nanoparticles at a current density of 50 mA g−1 is shown in Figure 3.19 [183]. All these reports shows that the nanocomposite electrodes of low cost iron oxide with highly conductive carbon additive can be promising alternative to graphite anodes for LIB applications.

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3.0 Charge

Voltage (V)

2.5 2.0

1st 2nd 3rd 10th

1.5 1.0 0.5 Discharge

0.0

200 400 600 800 1000 1200 1400 Specific capacity (mAh/g)

0

Figure 3.19 Charge–discharge voltage profiles of Fe3O4 nanoparticles at a current density of 50 mA g-1 (Reproduced with permission from Ref. [183]).

3.2.3.4.2 Cobalt Oxide The cobalt oxides such as Co3O4 and CoO are reported as anode material for LIBs with theoretical capacities of 890 and 715 mAh g−1, respectively [184, 185]. The different forms of these oxides like nanosheets, nanocubes, nanowires, nanotubes, etc., have been developed as effective anode materials [186–190]. The pure phase CoO octahedral nanocages with edge length 100 to 200 nm, [191] shows reversible capacity of 474 mAh g−1 at 5  C rate. Different structures of CoO deliver different electrochemical performances (Figure 3.20). The excellent performance of CoO nanocages are attributed to the presence of huge voids which can

3.0

Voltage/V

2.5

CoO-1 CoO-2 CoO-3

2.0 1.5 1.0 0.5 0.0 0

300 600 900 1200 Specific Capacity/mAhg–1

1500

Figure 3.20 (a) Galvanostatic charge/discharge curves of different CoO samples such as CoO octahedrons (CoO-1), CoO octahedral nanocages (CoO-2), and CoO hollow microspheres (CoO-3) at 0.2C rate (Reproduced with permission from Ref. [191]).

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accommodate the large volume changes during cycling. The meso-crystalline Co3O4 nanoplates [188] shows reduced initial irreversible capacity along with stable cycling performance, discharge capacity of 1015 mAh g−1 over 30 cycles at 0.2C rate. The amazing performance of the Co3O4 nanoplates is ascribed to the 2D structure of the nanoplates in combination with its highly porous structure to overcome the issue of volume change during cycling [192]. In view of mitigating problems of volume change and agglomeration of cobalt oxide during lithiation/de-lithiation various composites of cobalt oxides have been developed. Carbon based nanocomposites such as CoO/ graphene and Co3O4 nanoparticles anchored on graphene were reported with high capacities of 1015 (over 520 cycles) and 935 mAh g−1 (over 30 cycles), respectively [192–195]. Apart from the oxides of iron and cobalt, various other metal oxides showing conversion type reaction mechanism with Li such as NiO, MnOx. CuOx, MoOx, CrOx, etc., have also been reported as good anode materials for LIB with reversible capacities >500 mAh g−1 [196–200].

3.3 Nickel Metal Hydride Batteries The technology of Nickel Metal Hydride (Ni-MH) batteries was developed during 1970s and 1980s, similar to sealed Ni–Cd batteries. The founder of Ovonics battery company, Stanford Ovshinsky first patented the technology in 1986 [211]. As in Ni–Cd, Ni-MH also operates by using nickel hydroxide (Ni(OH)2), a Ni based positive electrode (cathode). But, hydrogen is used as the active element, hydrogen absorbing negative electrode (anode) instead of Cd as in Ni–Cd, and alkaline potassium hydroxide as the electrolyte. The utilization of the hydrogen absorbing alloy (usually lanthanum and rare earth based alloys) as anode provides much higher energy density for Ni-MH than Ni–Cd batteries. The Ni-MH rechargeable batteries find wide use in mobile applications, stationary energy storage, transportation applications, etc. [212–214]. Ni-MH batteries have same nominal voltage (1.2 V) of Ni–Cd batteries and replaces the highly toxic Ni–Cd in many applications. They can offer high energy density close to 80 Wh kg−1, wide operation/storage temperature range and a very long service life (>5 years). But the high self-discharge rate of ~30% per month is the main disadvantage of Ni-MH batteries. Moreover, although to lesser extent, they are also susceptible to memory effect. Though Ni-MH technology is more expensive, it is more environmentally benign than Ni–Cd and Lead acid battery technologies [214].

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3.3.1 Mechanism of Ni-MH Battery Operation A typical Ni-MH battery consists of a hydrogen storage alloy, a functional intermetallic which can store hydrogen reversibly, as anode and Ni(OH)2 as cathode together with an aqueous alkaline solution of KOH (6 mol l−1) as electrolyte. The performance of the battery in terms of capacity, cycle life and reversability strongly depend on the intrinsic properties of the hydride forming alloy [213, 215]. The basic idea of Ni-MH batteries originated from the research on hydrogen storage to utilize it as an alternative energy source. Some metals/alloys/intermetallics can form hydrides reversibly and thereby store hydrogen in volumes many times higher than their own volume. Therefore, by careful selection of the alloying elements and proportions, the electrochemical hydrogenation and dehydrogenation during charging and discharging can be effectively used for battery applications [212–216]. The electrode reactions occurring at the respective electrodes of a Ni-MH battery are given in Eqns. 3.10 to 3.12. At the positive electrode:

Ni(OH)2 +OH-

Charge NiOOH+H2O+eDischarge

(3.10)

At the negative electrode:

M+H 2O+e-

Charge MH+OHDischarge

(3.11)

Charge NiOOH+MH Discharge

(3.12)

The overall reaction:

Ni(OH)2 +M

As per the overall reaction, the hydrogen moves from the positive electrode to negative electrode and in reverse reaction occurs during charge and discharge, respectively without any involvement of the electrolyte. A detailed pictorial representation of the electrochemical reactions in a Ni-MH battery is shown in Figure 3.21. Self-discharge of the battery during long time storage is a serious problem in Ni-MH batteries. The major reasons for self-discharge are decomposition of the charged metal hydroxide due to direct reduction by oxygen

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e–

e– O

M

H2O

OH Ni NiOOH

H

HO

H

Ni Ni(OH)2

OH–1

MH Negative electrode

OH

Electrolyte

Positive electrode

Figure 3.21 Schematic of the electrochemical reactions in a Ni-MH battery (Reproduced with permission from Ref. [213]).

evolution at the positive electrode (Eq. 3.13) and hydrogen oxidation at the negative electrode (Eqns. 3.14 and 3.15). These reactions can cause leakage leading to the intake of oxygen, short-circuits due to the growth of crystal bridges between the electrodes, etc. [217]. The rate of self-discharge is affected by the temperature of storage and storage time. The self-discharge characteristics of a Ni-MH battery at different temperatures is shown in Figure 3.22.

4OH-

Ni O +2H O+4e2 2

O2 +2H 2O+4eMH+OH-

(3.13)

MH 4OH-

(3.14)

M+H2O+e-

(3.15)

Although a large number of alloys can store hydrogen reversibly, every hydrogen storage alloy cannot be used as negative electrodes for Ni-MH batteries. It should satisfy certain important criteria like high reversible hydrogen storage capacity, good electrochemical catalytic properties toward hydrogen absorption and discharge, easy activation and good corrosion resistance in alkaline solution, suitable hydrogen equilibrium pressure, good charge/discharge kinetics, long cycle life, less self-discharge

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120 0°C

Capacity Retention [%]

100 80

21°C

60 45°C 40 20 0

0

5

10 15 20 25 Storage Time [days]

30

Figure 3.22 Self-discharge characteristics of a Ni-MH battery at different temperatures (Reproduced with permission from Ref. [217]).

and low cost, etc. [215–217]. The energy storage capacity of the battery is determined by the electrochemical storage capacity of the electrode, i.e. the amount of hydrogen that can be absorbed by the alloy/intermetallics negative electrode material. Moreover, after charging, the discharge pressure of the MH should be in the range of 0.1 to 1 atm at room temperature for ensuring the complete discharge of the stored hydrogen. The stability of the formed MH also plays an important role in the effective performance of the battery. The stability of MH in AB (TiFe) [218], AB2 (ZrMn2) [219, 220], and AB5 (LaNi5 or CaNi5) [221] type systems which are mainly used for Ni-MH battery applications are determined by the strength of the M–M bond between the constituent atoms. Moreover, the strength of the M–H bond in the formed metal hydride should not be too high and too low, and should be in the range of 25 to 50 kJ mol−1 to facilitate the smooth charge/discharge process [222, 223]. A typical voltage and temperature profiles of Ni-MH cell during charge at various rates is shown in Figure 3.23.

3.3.2 Anode Materials In view of the considerable importance of the hydrogen storage alloys on the performance of the Ni-MH batteries, numerous work has been reported on hydrogen storage anode materials. Among them, the rare earth based AB5 type alloys, Ti and Zr based AB2 type alloys, Mg based alloys, rare Earth–Mg–Ni based super lattice alloys, and Ti–V based alloys are of the main interest [224].

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Cell voltage / v

1.50

1.40

50

1.30

40

1.20

30

1.10 0

20

40

60

80

100

120

140

160

Temperature / °C

0.3C 0.1C 1.0C

20

Charge input / % of rated capacity

Figure 3.23 A typical voltage and temperature profiles of Ni-MH cell during charge at various rates (Reproduced with permission from Ref. [222]).

3.3.2.1 Rare Earth-Based AB5 Alloys Rare earth based AB5 alloys are the most extensively studied class of anode materials for Ni-MH batteries. In the general representation AB5, A stands for one or more rare earth elements and B can be the transition metals like Ni, Co, Mn, Al, etc. LaNi5 is a typical example for this class of materials having CaCu5 type hexagonal structure as shown in Figure 3.24. In the hexagonal

(c)

(b)

(a) La

Ni

Figure 3.24 (a) Crystal structure of LaNi5 alloy, (b) tetrahedral sites, and (c) tetrahedral sites for hydrogen storage (Reproduced with permission from Ref. [213]).

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structure of AB5, hydrogen occupies preferentially at the tetrahedral sites containing A2B2, AB3, and B4 [213, 214]. In early stages of the development of LaNi5 as electrode material, though the single LaNi5 unit can absorb six hydrogen atoms to form a hydride to give good capacity almost equivalent to its theoretical capacity, its cycling stability was very less [225]. This problem of poor cycling stability was then mitigated by partial replacement of Ni atoms by Co in LaNi5 [226, 227]. Substitutions of A and B sites in AB5 type alloys (like A = La, Ce, Pr or Nd, and B = Ni, Co, Mn, Al, Sn, or Fe) [215–217] were found to deliver capacities in the range of 300 to 320 mAh g−1.

3.3.2.2 Ti and Zr-Based AB2 Type Alloys The ability to give high energy density in Ni-MH battery applications are made based on Ti and Zr based AB2 type alloys as the second generation alloy anodes [217]. The general representation AB2 corresponds to the Laves phase structures of an intermetallic compound such as hexagonal C14 (e.g., MgZn2), cubic C15 (e.g., MgCu2), and hexagonal C36 (e.g., MgNi2). Among them, the C14 and C15 Laves phases are good hydrogen absorbers with tetrahedral interstitial sites to accommodate hydrogen (Figure 3.25) [228, 229]. But the C36 phase is not effective as a hydrogen absorber. Among numerous AB2 type hydrogen absorbing alloys, the Ti and Zr based alloys are the most effective ones for Ni-MH battery applications. Later, apart from the binary alloys, multi-element AB2 type pseudo-binary alloys containing elements such as Ti, Zr, V, Ni, Cr, Co, Mn, Al, and Fe, which can deliver higher capacities (370 to 450 mAh g−1) than AB5 type alloys became more attractive [228–231]. Hariprakash et al. synthesized a series of Zr based compounds such as Zr1.05V0.2Cr0.047Mn0.578Co0.035Ni1.196 (sample 1), Zr0.9Ti0.11V0.2Cr0.05Mn0.58Co0.042Ni1.2 (sample 2), Zr0.68Ti0.3V0.2 Cr0.02Mn0.6Co0.06Ni1.24 (sample 3) and studied the effect of Ti substitution followed by the effect of Cu powder addition. The galvanostatic discharge capacity data obtained at a discharge current density of 80 mA g−1 for various alloy electrodes is given in Figure 3.25. However, slow activation and low rate C rate capabilities of AB2 type alloy in comparison with the rare earth AB5 type alloys limited its wide applications.

3.3.2.3 Mg Based Alloys The high reversible hydrogen storage capacity, low cost and low weight of Mg based alloys make it as a promising anode material for Ni-MH rechargeable batteries. Mg and Mg2Ni alloy show high theoretical capacities of 2200 mAh g−1 and 1080 mAh g−1 for the formation of MgH2 and

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Electrode potential vs. MMO / mV

–1000

–900

–800

–700

–600 0

1

2

3

4

5

Time / h

Figure 3.25 The galvanostatic discharge capacity data obtained at a discharge current density of 80 mA g−1 for various alloy electrodes, ( ) sample 1, ( ) sample 2, ( ) sample 3, and ( ) sample 2+Cu at 25°C (Reproduced with permission from Ref. [231]).

Mg2NiH4, respectively [232–234]. However, the poor charge/discharge kinetics at room temperature and instability of the Mg based alloys hinders its practical applications. This shortcoming in Mg based alloys have been effectively addressed by amorphization and nanocrystallization approaches [213]. Thus, nano- or amorphous microstructures can provide high electrochemical capacity by facilitating hydrogen diffusion and charge transfer reactions. The performance was further enhanced by the suitable selection of alloying elements such as Sc, Ti, V, and Cr to get maximum reversible hydrogen storage capacity than the commercial anodes  [235, 236]. The Mg2Ni alloy system developed using ball milling was studied extensively as anode material Ni-MH batteries. The effect of partial substitution of other metals in Mg2Ni alloy such as Ti, Al, Co, Ce, Zr, Cr, Sn, Mn, Ca, or Y for Mg [237–239], and Co, Mn, Cu, Fe, W, Al, Pd, or C for Ni [240–242] were also reported. The strategy of partial substitution was found to be effective in enhancing the corrosion resistance and overall performance of the Mg based alloys as anodes for Ni-MH batteries. But still, the existing Mg based alloys are not satisfactory for practical applications.

3.3.2.4 Rare Earth–Mg–Ni-Based Superlattice Alloys The rare earth–Mg–Ni (R–Mg–Ni) based superlattice alloys have been emerged as the next generation anode materials for high energy density

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and high power Ni-MH batteries. In the long periodic one dimensional superstructures containing phases such as (La, Mg)Ni3 and (La, Mg)2Ni7, AB5 (CaCu5 type structures) and AB2 (Laves structures) are stacked in the ratio n:1 along the c-axis in rhombohedral or hexagonal manner results in high hydrogen storage capacity [213–215, 243]. Many approaches like composition optimization, heat treatment, ball milling, surface treatment, etc., have been reported to improve the performance of R-Mg-Ni based alloys [243–247]. Then the effect of partial replacement of Ni with other elements such as Co and Al were also found to be effective in improving the overall performance of the anode materials. These class of materials is presently used as anodes in 2500/900 series Ni–MH batteries.

3.3.2.5 Ti–V-Based Multicomponent Multiphase Alloys The high discharge capacities (>360 mAh g−1) of Ti–V based multicomponent multiphase alloys making it as the most promising negative electrode material for Ni-MH batteries. The limitations of single phase were mitigated by this approach of developing multicomponent multiphase alloys [248–250]. The improvement in the charge/discharge characteristics of V based alloys in the presence of a secondary phase such as C14 Laves phase or TiNi phase [251] was the beginning of the development of this type of materials. Further, in certain superstoichiometric Ti–V based alloy series containing C14 Laves phase and V based solid solution were found to show attractive electrochemical characteristics through the synergistic activity of these two phases [249, 250]. However, the poor cycling stability and charge/discharge kinetics of Ti–V based alloys impede its practical application. Apart from these hydrogen storage alloys, recently, cobalt based alloys especially CoB alloys have been attained significant attention due to relatively high discharge capacity (>600 mAh g−1) and good cycling stability (>100 cycles) [216]. Other than CoB alloys, the combinations such as Co–P, Co–Si, Co–BN, and Co–Si3N4 are also considered as promising negative electrodes for Ni-MH batteries, owing to their high (>450 mAh g−1) discharge capacities [216, 238].

3.4 Lead-Acid Batteries Lead-acid battery is the first rechargeable battery invented by French Physicist Raymond Gaston Plante during 1859. The battery system consists

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of spirally wound lead plates immersed in dilute sulfuric acid solution. Later in 1881, Faure and coworkers introduced new technology with pasted plates prepared by mixing of lead oxides with sulfuric acid on to the lead alloy grids [252]. This was found to give much faster better efficiency. Although lead-acid batteries have been experienced several design adjustments from time to time to allow the battery to cope with the new performance challenges as they have emerged, the fundamental electrochemistry has remained as the same even after 150 years. The basic components of a lead-acid battery and its construction are shown in Figure 3.26. The inferior properties of lead-acid batteries such as low specific energy and power characteristics compared with other rechargeable batteries are surmounted by the attractive characteristics in practical applications such as low cost, maximum recyclability (>98%) and operational safety. The operational safety is guaranteed by the use of nonflammable electrode materials and aqueous electrolyte in lead-acid batteries. Though the battery cases can catch fire, the risk of this is low compared with other batteries such as lithium-ion batteries. Further, the economical and fully established recycling process of lead-acid battery in compliance with the environmental regulations, making it more attractive for future applications compared with the other battery chemistries that need substantial improvement in the recycling process [252–257].

Electrolyte-tight sealing ring Grid plate

Positive plate pack

Positive plate Negative plate

Positive cell connection

Microporous separator

Valve adapter and valve Negative pole Negative cell conncetion Negative plate pack

Figure 3.26 The components and construction of a lead-acid battery (Reproduced with permission from Ref. [252]).

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Lead-acid batteries are available in different sizes, various designs (cylindrical, prismatic, etc.) and capacities (few Ah to 1000s of Ah) in the market. Lead-acid battery attains the large market value globally in applications like hybrid electric vehicles (HEV), industry, start–stop applications in automotive, telecommunications, etc. Presently, 70% of market value is collectively contributed by the electric bikes and automotive batteries. The worldwide market value of lead-acid batteries is 51200 million US$ in 2017 and its projected value by 2025 is 80400 million US$ [4, 256, 258, 259].

3.4.1 Operating Principle The lead-acid cell comprises of highly porous lead dioxide as positive and finely divided lead as negative active material. The electrolyte is a dilute aqueous sulfuric acid solution, takes part in the discharge process. This involvement of the electrolyte solution in the discharge and recharge reactions making it unique among other rechargeable electrochemical systems. On discharge, HSO4 ions migrate to the negative electrode and produce + H ions and lead sulfate (Eq. 3.16). At the positive electrode lead dioxide reacts with the electrolyte to form lead sulfate crystals and water (Eq. 3.17). Both electrodes are discharged to lead sulfate which is a poor conductor and the electrolyte is progressively diluted during discharge. On charge, the reverse reactions take place [252, 255]. In overall, a lead-acid cell delivers ~2.05 V as presented in Eqns. 3.16 and 3.17. At the positive electrode:

PbO2 H2SO 4 2H+ 2e

Discharge PbSO4 2H 2O (1.69 V) Charge (3.16)

At the negative electrode:

Pb2 H 2SO4

Discharge PbSO4 2H+ 2e Charge

(0.36 V)

(3.17)

Accordingly, the net discharge–charge processes in a lead-acid cell are represented by Eq. 3.18,

Pb PbO2 2H2SO4

Discharge 2PbSO4 2H 2O (2.05V ) Charge

(3.18)

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The specific gravity of the electrolyte rises during charge as the lead sulfate concentration is reduced and reduces during discharge from 1.28 (fully charged) to 1.1 (fully discharged). Along with the major electrochemical reactions discussed above, as the equilibrium potential of PbO2/PbSO4 couple is more anodic to O2/H2O couple by 0.47 V, and the equilibrium potential of Pb/PbSO4 is more cathodic to H2/H+ couple by 0.36 V there is a possibility of secondary reaction such as water decomposition beyond 2 V. Moreover, at high potential of the positive electrode (1.7 V), as the PbO2 and Pb cannot coexists due to thermodynamic reasons there is a possibility of arising grid corrosion also as a secondary reaction [252, 255–257]. O2 would evolve at the positive plate according to reaction shown in Eq. 3.19,

PbO2 + H2SO4

PbSO4 + H2O + ½ O2

(3.19)

While H2 would evolve at the negative plate according to the reaction shown in Eq. 3.20, even at open-circuit.

Pb + H2SO4

PbSO4 + H2

(3.20)

Both reactions (Eqns. 3.16 and 3.17) lead to self-discharge of the battery [252].

3.4.2 Negative Electrodes of Lead-Acid Batteries The negative electrode is one of the key components which determines the performance of lead-acid batteries. The negative electrode is composed of a metal Pb layer coated on the lead or lead alloy grid. The negative electrodes are prepared by pasting the negative active material (NAM) on lead–alloy grids (current collectors) followed by curing and formation. The performance of lead-acid batteries is highly dependent on the composition, structure, morphology, and design of the negative electrodes. The process involved in the preparation of negative electrodes is discussed below.

3.4.2.1 Preparation of Negative Electrode The negative plate paste is prepared by mixing leady oxide (85 wt. %), lignin (0.2 wt. %), barium sulfate (0.15 wt. %), Dynel fibers (0.05 wt. %), carbon black (0.1 wt. %), and 1.4 sp. gr. H2SO4 (7 wt. %) with de-ionized water (7.5 wt. %). The paste density of negative plate is maintained as 4.4 g cm−3 [260, 261]. The function of each component in the NAM is in Table 3.2.

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Table 3.2 Composition of the negative active material paste. Sl. No.

Ingredients

Composition of NAM (wt. %)

Function

1

Leady oxide

85

Active material

2

Dynel fiber

0.05

Expander, increases active material utilization

3

Carbon black

0.1

Conductive diluent

4

Lignin

0.2

Expander

5

BaSO4

0.15

Nucleating agent

7

H2SO4 (1.4 sp. gr.)

7

Electrolyte

8

Deionized water

7.5

Solvent

The ratio of each component in the NAM depends on the type of leadacid battery intended to develop. During fabrication, the NAM is pasted onto the negative grid (lead–alloy current collectors) followed by curing under controlled temperature and humidity to convert the wet paste to a dry, crack-free unformed material with good adhesion to the grid. Curing is the most time consuming technological procedure which takes about 24–72 h. During curing Pb oxidation; recrystallization of 3PbO·PbSO4·H2O (3BS), 4PbO·PbSO4 (4BS), and PbO; improvement of the paste/grid contact, and drying of the paste occurs. With an increase of curing temperature, the rates of the curing processes can be accelerated and curing time can be shortened. At temperatures >65°C, the 3BS paste can transform into 4BS paste comprising large crystals which are difficult to oxidize to PbO2 during formation [252, 260–262]. Hence, the temperature is normally maintained 1500 m2/g) activated carbons. Hence, during charge–discharge cycling of a PbC battery, the positive electrode (PbO2) undergoes charge–transfer or faradaic reactions as in conventional lead-acid batteries, while the negative electrode undergoes faradaic type charge storage reactions as in ultracapacitor are presented in Eq. 3.22. Negative electrode reaction in the PbC battery:

nC 6x- (H+ )x (a)

200 nm (c)

200 nm

nC(x-2).(H+ )x-2 6

2H+ 2e- (Discharge)

(3.22)

(b)

200 nm (d)

200 nm

Figure 3.31 Morphologies of conventional cell before cycling (a), after cycling (c), and TiO2–RGO (3:1) additive cell before cycling (b), after cycling (d) (Reproduced with permission from Ref. [307]).

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The PbC batteries can offer faster recharge rates, greater charge acceptance and longer cycle lives in deep discharge applications compared to the conventional lead-acid batteries with minimum maintenance. The costs of these batteries are the same as that of regular lead-acid batteries and are fully recyclable like lead-acid. The PbC battery has less energy density (~5 Wh/kg), but has high power density [308, 310]. Besides, in recent times, UltraBattery, which is a hybrid and long-life of lead-acid battery in a single device has been developed. It signifies a hybrid device encompasses the advantages of ultracapacitor (fast charging rates and cycle life) and lead-acid battery (energy storage efficiency) in a single unit by the use of a common electrolyte. This hybrid device sustains high energy density of lead-acid batteries with high power density of ultracapacitors in a single unit cell. UltraBattery is designed with a single PbO2 positive electrode and a twin negative electrode (one part is completely carbon and the other part is lead) in a common lead-acid battery. Although the twin electrode comprising of carbon and lead parts together acts as the negative electrode, carbon part acts as the electrode for capacitor and lead part serves as the electrode for lead-acid cell in principle. The UltraBattery does not work or give much cycle life, by a simple connection of the capacitor electrodes with the lead negative-plates, because of the inefficiency of the capacitor electrode (carbon) to share current with the lead negative plate during early stages of discharge. Hence, it may result in increased hydrogen evolution at the capacitor counterpart of the negative electrode leading to the electrolyte dry-out. To overcome these issues and to make the ultra-battery functioning smoothly, much attention is required in the modification of capacitor electrode from the fundamental side [310]. Presently the commercial UltraBatteries are developed mainly by Furukawa Battery Company of Japan and the East Penn Manufacturing Company of the United States, in partnership with CSIRO Australia [308].

3.5 Thermal Batteries A thermal battery is a non-rechargeable, primary battery that can deliver the highest level of specific power compared to other batteries available in the market. It is completely inert before being activated, by melting the solid electrolyte using thermal energy, either by mechanical or electrical ignition. The high power output of thermal battery does not come from the conversion of thermal energy to electrical energy, but from a reactive electrochemical couple. After thermal activation, the liquified electrolyte takes

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part in a non-reversible chemical reaction to produce electrical potential between the terminals. The main advantages of this primary battery are long shelf life (>20 years), less activation time (0.2 to a few seconds) and very high power density [6]. The robust design of these batteries makes it impervious to high stress, sudden shock, and sharp pressure drop. This allows their use in harsh climatic and mechanical environments. Hence, the thermal batteries find extensive use in defense applications, especially in missiles, rockets and torpedoes. The technology of thermal batteries was conceptualized and developed by Dr. Georg Otto Erb along with other German scientists during the second world war to use in rockets [6, 311]. Later the technology was brought back to the United States after the world war and developed further. A typical thermal battery consists of a stack of cells (Figure 3.32), each one having an anode, cathode, electrolyte, and a heating pellet. The electrolyte (a mixture of lithium chloride and potassium chloride) is solid, inert, and nonconductive at room temperature, which acts as a separator between the anode and cathode until activation. The inertness of the solid electrolyte gives long shelf life of more than 20 years for thermal batteries. The anode of thermal batteries was made up of calcium or magnesium initially, and presently it is replaced by lithium [6, 311–319]. The cathode is made up of chromates or sulfides. In a thermal battery, the required number of cells are connected either in series or parallel or combination of both according to the requirement of output voltage. The activation is carried out by the ignition of the heat pellet to increase the temperature and thereby to melt the solid electrolyte. As the temperature increases inside the cell, solid electrolyte melts and ion exchange takes place. Thus power in the cell is liberated. The activation time for thermal batteries ranges from 0.2 to a few seconds, and

HEAT SOURCE ANODE

ACTIVE CELL

INERT CELL

N x Cells

ELECTROLYTE CATHODE

Cell

HEAT SOURCE

(a)

(b)

Figure 3.32 Schematic of: (a) typical thermal cell components, and (b) a thermal battery containing stack of cells.

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the operating temperature is in the range of 500 to 700°C. The cell stack is thermally insulated from the case to maintain high working temperature inside the cell and also to avoid the damage of nearby components due to high temperature. The insulation plays a crucial role in the stable performance of the battery after activation by maintaining the working temperature without falling below the operating range (500 to 700°C) [6, 319]. The discharge of the battery ceases either by the exhaustion of the active material or by the solidification of the electrolyte upon decreasing the temperature below operating range. The output voltage is greatly influenced by the operating temperature, environment, activated life and the applied load or the amount of current drawn.

3.5.1 Anode Materials for Thermal Batteries In the initial stage, Ca or Mg anodes with WO3 or V2O5 cathodes and LiCl– KCl eutectic mixture as electrolyte were used in thermal batteries. However, the technology was developed without any electrochemical understanding in early stages and stopped after the report on basic electrochemistry of V2O5 in molten salts by Laitenen et al. [313]. The high solubility of V2O5 in molten electrolyte can have an adverse effect in thermal batteries by chloride oxidation in the melt. This technology was then replaced by Ca/ CaCrO4 system with LiCl–KCl eutectic as the electrolyte, without using any separator [314, 319]. The CaCrO4 cathode dissolved in the electrolyte was in direct physical contact with the Ca anode. After activation, a liquid CaLi2 alloy formed through the chemical displacement of Ca and Li + present in the molten salt acts as the real anode. The CaLi2 discharges through several stages to form Ca and it can further regenerate the CaLi2 alloy by reacting with the bulk Li+. But this technology has several disadvantages like inter-cell shorting due to liquid anode, precipitation of Ca2+ double salt in the presence of K+ (KCaCl3) to increase the cell impedance, and also the highly exothermic discharge reactions can lead to self-destruction of the battery [315]. Later this technology was replaced by the use of FeS2 (pyrite) as the cathode material [6]. The battery design was greatly simplified by the use of FeS2 cathode, as there are no chemical reactions between the electroactive components. The different anode materials used in combination with the FeS2 cathode for thermal battery applications are discussed below.

3.5.1.1 Ca-Based Anodes Initially Ca and Ca based alloys such as Ca–Al, Ca–Si, and Ca–Al–Si (mixtures of CaSi and CaAl2) were used as anodes in combination with FeS2

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cathode. However, the major disadvantage is due to the formation of liquid anode (CaLi2) in the presence of Li+ through displacement reaction [6, 316]. Moreover, as the cell emf is much less when Ca based anodes are used compared to the Li alloy anodes, Ca was replaced by Li in thermal battery applications.

3.5.1.2 Mg- and Al-Based Anodes Mg and Mg alloys such as Mg–Si, Mg–B, Mg–Al, and Mg–Cu were used as anodes in combination with FeS2 cathode for development of thermal battery [317–320]. Although the Mg/LiCl–KCl/FeS2 technology was much better than the older Ca/CaCrO4 technology, but the problems of sluggish kinetics, lower emf and capacity compared with Li alloy anode systems lead to limited development effort for Mg based anodes [6, 320]. Aluminum was also reported as anode for thermal battery applications in combination with FeS2 cathode in tetrachloroaluminate based electrolytes [321–324]. But issues due to poor rate capability, temperature limitation and low cell emf hindered its further development.

3.5.1.3 Li Anode The use of pure Li as anode in thermal batteries can offer the highest emf that can achieve ever. However, the low melting point of Li (180°C) leads to early melting of Li during thermal battery operation and hence it has to be controlled [325–327]. The use of such molten Li as anode in thermal batteries can lead to displacement reactions in the presence of K+ during discharge. This displacement reaction can result in the escape of K in the melt and to react with surrounding oxides in the insulation and thereby lowers the overall coulombic efficiency [328, 329]. Further, under severe environments, there is possibility of shorting of cells. Hence, researchers turned to the search for suitable Li alloys as anodes for thermal battery rather than using pure Li.

3.5.1.4 Li–Al Anodes The initial extensive work with FeS2 system was carried out using Li–Al alloy and thereby it became the primary anode material for all thermal batteries. Several reports and patents were issued on the preparation of Li–Al anodes for thermal battery applications [330–333]. The phases present in

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the Li–Al system were analyzed and their activity as anode material was studied. Based on these works by Moshtev et al. [334], Wen et al. [335], and Argonne National Laboratory (ANL), USA found that only β-LiAl is suitable to use as anode for thermal batteries [336]. The ternary compositions of Li–Al alloy such as Li–Al–Si and Li–Al–Fe were also studied as anodes for thermal batteries to improve the specific energy [337]. Though the specific energies of the ternary alloys were higher than Li–Al alloy, they were less than that of Li–Si anodes.

3.5.1.5 Li–Si Anode The Li–Si anode was developed to overcome the limitations of Li–Al, owing to its advantages such as better rate capabilities and multiple phases. Many research groups reported extensive study on the identification of phases in Li–Si alloy [338–344]. The studies revealed the formation of a number of compounds within the Li–Si system, and the emf vs. Li decreases with higher Li contents. On the basis of the information obtained from the discharge sequence of Li–Si anodes [341] Li13Si4 to Li7Si3 transition (44 w/o Li) was preferred as anodes for thermal battery. Though the Li–Si system shows less capacity than the Li–Al anode, the Li–Si alloy can show multiple transitions and thereby deliver high power at a higher rate. The ability to show reproducible performance and lower voltages vs. Li for Li13Si4 made it more attractive for thermal battery applications [341, 342]. Moreover, the lower specific energy and energy density for first transition of Li13Si4 to Li7Si3 than in Li–Al alloy allows it to provide ~150 mV in a cell. All these advantages of Li–Si alloy with 44 w/o Li enable it as the current preferred choice of anode material for thermal batteries. A typical cross-sectional view of a Li–Si/FeS2 thermal battery is shown in Figure 3.33. A comparison of the performance of the Li–Si/FeS2 couple in different eutectics such as LiBr–KBr–CsBr and LiCl–LiBr–LiI–KI–CsI in single cells at 250°C and 8 mA cm−2 are given in Figure 3.34. Further, to improve the overall electrochemical performance of the Li– Si binary alloy, many ternary derivatives such as Li–Si–Fe [343], Li–Si– Mg [102, 344], Li–Si–Al, and Li–Si–B were reported. Many other binary and ternary alloys of Li such as Li–Sb [345], Li–Bi [346], Li–Sn  [347], Li–B [348], Li–Ge [349], and Li–Mg–B [349, 350] have been used as anodes for thermal batteries. A comparison of the capacities and plateau voltages for various Li based alloy anodes used for thermal battery applications are given in Table 3.3.

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Nanomaterials for Electrochemical Energy Storage Devices Electric Match

Alumina-Filled Encapsulation

Fiber-Frax Wrap

Zr/BaCrO4 (Fuse Strip)

Mica Anode Collector Anode [Li(Si)] Electrolyte – Binder (Separator)

Glass Cloth Wrap

Cathode (FeS2) Cathode Collector Heat Source Pellet (Fe/KCIO4) MIN – K INSULATION

MC3907 THERMAL BATTERY Li(Si)/LiCl–KCI/FeS2 VOLUME 292(629.7)cm3 DIA 69.85(84.3)mm

WEIGHT 1.25 Kg LENGTH 76.2(106.9)mm

Figure 3.33 Schematic of the cross-sectional view of a Li–Si/LiCl–KCl/FeS2 thermal battery (Reproduced with permission from Ref. [6]).

2.2

14 (i)

12 (ii)

1.8

10

1.6

8

1.4

6

1.2

4

Total polarization / Ω

Cell voltage / V

2

2

1 (i)

(ii)

0.8 0

0.1

0.2

0.3 0.4 0.5 0.6 0.7 Capacity / Eq. Li mol–1 FeS2

0.8

0 0.9

Figure 3.34 Comparison of the performance of the Li–Si/FeS2 couple in (i) LiBr– KBr–CsBr and (ii) LiCl–LiBr–LiI–KI–CsI eutectics in single cells at 250°C and 8 mA cm−2, along with the variation in total polarization (Reproduced with permission from Ref. [319]).

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Table 3.3 Comparison of the capacities and plateau voltages for various Li based alloy anodes used for thermal battery applications [312]. Anode

Capacity (A-s g−1)

Capacity (A-s cm−3)

Discharge emf/mV vs. Li0 (T/°C)

Li–Al (20 w/o Li)

2259

3931

297 (415)

Li13Si4 (44 w/o Li) (1st transition)

1747

2411

157 (415)

Mg2–Mg–Li13Si4

3155

>4354

60 (415)

Li–B (70 w/o Li)

4800

3900

20 (500)

Li–B (80 w/o Li)

7700

5400

20 (500)

Li8Si2B (45.3 w/o Li) (1st transition)

1811

2880

157 (400)

Li9Ge4 (17.7 w/o Li) (1st transition)

1368

3420 (est.)

420 (400)

Li16Ge5 (23.4 w/o Li) (1st transition)

967

2418 (est.)

236 (400)

References 1. Chu, S. and Majumdar, A., Opportunities and challenges for a sustainable energy future. Nature, 488, 294, 2012. 2. Hall, P.J. and Bain, E.J., Energy-storage technologies and electricity generation. Energy Policy, 36, 4352, 2008. 3. Chen, H., Cong, T.N., Yang, W., Tan, C., Li, Y., Ding, Y., Progress in electrical energy storage system: A critical review. Prog. Nat. Sci., 19, 291, 2009. 4. Broussely, M., and Pistoia, G., Industrial applications of batteries: from cars to aerospace and energy storage, Elsevier, Amsterdam, UK, 2007. 5. Ratnakumar, B.V., and Smart, M.C., Aerospace applications. II Planetary exploration missions (orbiters, landers, rovers and probers). In: Industrial Applications of Batteries, pp. 327–393, Elsevier, Amsterdam, UK, 2007. 6. Guidotti, R.A. and Masset, P., Thermally activated (“thermal”) battery technology: Part I: An overview. J. Power Sources, 161, 1443, 2006. 7. Aurbach, D., Characterization of batteries by electrochemical and nonelectrochemical techniques. In: Industrial Applications of Batteries, pp. 119– 201, Elsevier, Amsterdam, UK, 2007.

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4 Nanostructured Cathode Materials for Li-/Na-Ion Aqueous and Non-Aqueous Batteries Farheen N. Sayed1*, Ganguli Babu2 and P. M. Ajayan2† 1

Solid State and Structural Chemistry Unit, Indian Institute of Science, Bengaluru, India 2 Department of Material Science and NanoEngineering, Rice University, Houston, Texas, USA

Abstract Fast charging is the main hurdle towards wide range adoption of alkali ion rechargeable batteries in electric vehicles, which is fundamentally rooted to electrode materials used in these energy storage systems. In this chapter, we have compressively reviewed nanostructured cathode materials for Li /Na-ion batteries with both aqueous and non-aqueous electrolyte. The difference between the aqueous and non-aqueous battery systems with respect to their working mechanisms, and the advancements in nano-cathode materials are discussed. The various synthesis protocols adapted to prepare cathode materials with desired nanostructures and improvements in electrochemical performance are conferred. The future prospective including considerations and limitations to enable fast charging of different class of cathode materials in nano dimensions are discussed in detail with representative studies from available literature. Keywords: Li-ion batteries, Na-ion batteries, aqueous electrolytes, non-aqueous electrolytes, cathodes, nanomaterials, specific capacity

*Corresponding author: [email protected] † Corresponding author: [email protected] Poulomi Roy and Suneel Kumar Srivastava (eds.) Nanomaterials for Electrochemical Energy Storage Devices, (231–269) © 2020 Scrivener Publishing LLC

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4.1 Introduction With the all positive response and features of rechargeable Li+ ion batteries, which are the important components of all the current electronic devices, vehicles, satellites, the first cry comes for its power capability. The energy density of batteries has grown in centuries and now reached to the 250 Wh kg−1, with strong promise of further acceleration. That means it can store more amount of energy in more efficient way. However, while talking about power performance which requires the quick charge and discharge to get the energy in a fraction of time, the supercapacitors are still the preferred choice. Hence, most of the initial commercialization of electric vehicles and other technologies involves a hybrid system of battery and supercapacitor. With increasing demand of fast charging batteries for the current fast generation, the pressure on the materials research community is building up to further accelerate the development of newer materials for energy storage requirements. Even though there are many competing technologies, like Li–S, Mg, Zn, and Li–air batteries, there is still no better alternative to Li+ ion batteries. Very close to Li ion battery chemistry is Na ion chemistry with strong potential of commercialization. The power capability can be evaluated in terms of its rate of performance. Usually, if it is mentioned on battery pack that the battery provides the energy of 140 mAh g−1, this suggests that the battery provides 140 mA g−1 of charge with 1C rate i.e. in 1 h of charge and discharge. The common trend is as the rate of charge–discharge is increased to 2 C, or 5 C, the stored charge is less and keeps on decreasing with further increase in rate. This makes batteries incompatible with high power demanding devices, which require same amount of energy to be stored and provide in reduced time. The fault lies in the internal materials chemistry of batteries. The alkali ion-rich electrode, i.e., cathode is separated from alkali ion accepting host anode by an electrolyte-soaked separator. However, in anode, the occupation of Li+/Na+ ion in the nonspontaneous step of the charged state is random, where it mostly occupies the interlayer spacing. In case of cathode materials, the Li+/Na+ ion has to travel from its original position and eventually return to same position. This shuttling and solid state diffusion of Li+/Na+ ion even in cathode material itself from bulk to surface is highly dependent on structural arrangement. Practically when the cell is charged at a higher rate, i.e. with higher current, the time interval is not sufficient enough to complete the diffusion of Li+/Na+ ion into bulk and hence the completion of insertion/de-insertion process is not attended. This leads to the loss of charge storage capacity and hence the energy density of the battery. A battery is said to be of good

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quality if it can deliver same amount of capacity as delivered at a lower rate at higher C-rates. However, if battery fails to perform in high power demand condition, it becomes replaceable with other technologies like supercapacitor. As mentioned above the major high rate capability limiting step is associated with slow solid-state diffusion of Li+/Na+ ion in the electrode materials. Hence, there has been tremendous interest in focusing on the research and development aspects of materials, specially the materials for cathode to reduce the diffusion limitations. Not only the alkali ion diffusivity is highly structure dependent, but the electron generated due to electrochemical activity of active electrode material are also slowed down while reaching to outer circuit. Hence overall efficiency of the system decreases. These two phenomena, sluggish Li+/Na+ ion diffusion and slow electron transport then lead to capacity fade in the long term, requiring the thickness of electrodes to be lesser. The widely accepted and adopted approach to achieve high rate capability is by reducing the diffusion path length for charged species via reducing the average particle size of the active cathode materials. In recent years, material science has witnessed the steep popularism for nanomaterials in which by confinement in nano-dimensions, combined properties of bulk and surface showing better host for alkali ion and electron transport are observed. Independently or in combination with particle nanosizing, surface modification and porous structure fabrication has also proven to be constructive strategy to achieve high rate capability. Surface modification can be done by range of compatible coating materials. The most common and effective being carbon coating, which increases the electronic conductivity of the material and also promote the compatibility between electrode materials and additive carbons, thus regularizing the interfacial transport of electrons and ions. Whereas, porous nature helps in faster equilibration, absorption and simultaneous release of electrolyte ions further facilitating the transport of ions in the heterogenous battery system [1]. For very small particles, the chemical potentials for active metal ions can be modified which in turn results in a changed electrode potential. The term nanomaterial could be broadly used for the nanoparticles of the size in which all the three dimensions are in nano meter range and could be of any morphologies, nanoparticle with at least two dimensions in nano range and similarly one dimensional nanoparticle with at least one dimension in nano range. The nano materials including nano spheres, nano cubes, nanorods, nano plates, and flakes have been investigated thoroughly and still being pursued [2–8].

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However, the advantages of nanosizing come with some inevitable disadvantages, the detrimental one being undesirable electrode–electrolyte interaction due to an uncompensated charge arising from high surface area. This interaction results in the formation of variable solid electrolyte interface (SEI), which has been identified as the major reason of self-discharge, poor cycling and poor safety of the alkali ion batteries.

4.1.1 Li+ vs. Na+ ion Batteries Till now in market Li+ ion battery has been an undisputed leader for all kind of electronic appliances, even though it fails to fulfill some of requirement criteria. While the economist and environmentalist are raising their concern over the use of Li+ ions which has very limited and restricted resources in selected countries, increasing competition with the other very important application as in ion exchange and nuclear, the search of an alternative leader has already started. Hence, the responsibility to run the modern energy storage devices could be shared with other potential metals. The extensive research in SIB also started at the same time as that of LIB in the 1980s however the success of later has overshadowed the potentialities and capabilities of SIB. The idea of SIBs resurfaced during 2010 and is now advancing at a rapid pace, as indicated by increased number of publications and startups. The major problem for SIB is to find a host material for Na ion with near operating voltage and capacity of Li-ion host materials. And then comes the ionic radii, where the larger Na+ radius (0.98 Å) to that of Li+ (0.69 Å) leads to the slow kinetic of Na+ insertion/de-insertion and transport in and across the host material, leading to constantly degrading specific capacity and rate capability. In case of successful Na+ insertion/de-insertion, the larger volume expansion caused also brings a change in lattice structure of the host materials, creating barrier to retrieve capacity and achieve electrochemical stability as good as in LIB. Fundamentally, they also suffer from a lower specific energy than LIBs, due to inherent lower potential and heavier weight of sodium. In the above mentioned scenario, the biggest challenge is to build affordable Na-ion host materials with a comparable high specific energy. However, SIBs are of strong interest, in view of the fact that the economically accessible global lithium resources might not be sufficient for all its applications, only fraction of which is consumer electronics, while the sodium resources are much abundant. Another advantage of using sodium is that it does not react with aluminum which is a common positive current collector for LIB, whereas lithium has tendency to irreversibly alloy with aluminum. In economic terms, manufacturers can use

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this “inertness” of aluminum to reduce the production cost by completely replacing it with copper current collector. On the anode side, copper current collector has tendency to dissolve at lower potential further limiting the anodic potentials. LIBs need to be transported at a 20–40% state of charge to prevent dissolution of the copper. However SIB can be stored under discharged state, which reduces the safety risks in transportation and storage. In electrolyte most of the common solvent complexes have similar structures for both alkali ions, with the exception of few where coordination changes from mono-dentate for Li+ ion to bidentate for Na+ ion owing to its larger radius [9]. On the other hand total binding energies reduces by ca. 20% for to smaller charge/radius ratio of Na+ ion [10]. In the charge transfer at the electrolyte–electrode interface, SIB electrolytes precedence LIB electrolytes. For these electrochemical cells, the transfer of alkali ions at the electrolyte–electrode interface can indeed be the rate determining step. In view of this Okoshi et al. [9] evaluated the de-solvation energies for various cations, including Li+ and Na+ ions, in the presence of organic solvents using DFT calculations. Smaller desolvation energy (up to 40–70 kJ mol−1) for Na+ ion as compared to Li+ ion was shown and could be successfully linked to the weaker Lewis acidity of Na+ ion.

4.1.2 Aqueous vs. Non-Aqueous Electrolyte Even though, considered passive and with the responsibility of containing and allowing the diffusion of ions of interests, Li+ and Na+ ion, the electrolyte solvent has profound effects on the electrochemical as well as thermal stability of the electrolytes. In norm the low viscosity organic solvents are mostly preferred which usually have rather high vapor pressures, creating flammable vapors at nonambient temperatures or unusual charge–discharge conditions. Most notable example is the standard organic carbonate-based electrolyte in LIB which suffers from the problems of operation at non-ambient temperatures. The other solvent candidates, polymers and ionic liquids, have no vapor pressure but both are too viscous to use for low temperature or even room temperature operations. The strongest emerging competitor for organic electrolyte is the aqueous electrolyte for both LIB and SIB. The low investment, environmental friendliness and guaranteed safety makes aqueous electrolyte promising alternative with high ionic conductivity, high rate performance and low over potentials as compared to abovementioned electrolytes.

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To understand the kinetics of the Li+ ion intercalation electrolytes, the ionic conductivity of the LiNO3 aqueous salt solution has been compared with organic electrolytes (1 M LiPF6/EC:DMC(1:1)) Figure 4.1 [11]. Where it can be concluded that given in same electrochemical conditions the impedance observed for aqueous electrolyte was less than an organic electrolyte system. For large-scale applications and also as hybrid power supplies from intermittent renewable sources, these advantages of aqueous electrolyte batteries govern the ease of fabrication as well as lower total cost. However, acidic or alkaline nature of electrolyte can cause serious deleterious corrosion. As it is known that the acidity/basicity of the electrolyte strongly affects the potential shift of system as well the stability of current collectors. Thus, neutral aqueous electrolyte with a pH value of electrolyte solution near weak alkaline and acid would be the electrolyte of choice. To be compatible with the electrode components the electrochemical stability window of electrolytes should also match with that of electrodes. To be precise, the theoretical limit of each electrolyte’s electrochemical stability window (ESW) is determined by the energy separation between the lowest unoccupied molecular orbital (LUMO) and the highest occupied molecular orbital (HOMO) of the electrolyte components. The stringent requirements for the electrolytes are that the LUMO should be less positive then anode potential and HOMO should be more negative than cathode

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Figure 4.1 Electrochemical impedance spectroscopy spectra for a thin LiFePO4 film 450 nm thick in LiNO3 salt solution and organic electrolytes [11].

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potential. If these criteria are overlooked deleterious reactions of electrolyte frequently take place precipitating insoluble products which adhere to the negative electrode surface and form a protective solid passivation layer termed as the solid electrolyte interphase, SEI. Interphase is also formed at the positive electrode, also called as cathodic SEI. These abovementioned criteria, evaluation of transport properties, the stable electrochemical window, operable temperature range, and the safety features were precisely identified by Aurbach et al. [12] as the four major figure of merits that need to be considered for any battery application. Even though there are many articles highlighting the increased operational safety and low manufacturing cost with aqueous electrolytes for both LIB and SIB [13, 14], the main drawback of aqueous electrolytes lies in its low decomposition potential. Theoretically predicted 1.23 V cannot be completely achieved experimentally. This lower electrochemical stability window restricts the higher cell voltage limit and hence lowers down the battery’s energy density. Using electrode materials possessing high hydrogen and oxygen overvoltage, battery voltages up to 2 V are attainable, as achieved in the case of lead–acid batteries. However, all recent aqueous electrolyte battery systems with voltages above 1.23  V are fundamentally unstable and will, therefore, have a certain rate of self-discharge. In principle, the electrode materials that work outside the stability range of water will tend to cause it to decompose sooner or later. Positive electrode materials which contain lithium and have potentials over the limit of water stability will react with water and absorb lithium with the generation of protons. This will decrease the pH of the water. However this is not feasible if their potentials are within the stability limits of water. The key concern here is to avoid hydrogen and oxygen evolution by using active materials with higher anodic hydrogen evolution overpotential and cathodic oxygen evolution overpotential. Among present pool of cathode materials, LiMn2O4 (LMO) qualifies very well for being within the stability range of aqueous electrolyte. For LMO the electrochemical potential is less positive than the oxygen evolution potential (~1.6 V vs. NHE) Figure 4.2 [15]. The very first introduction of aqueous rechargeable lithium battery was reported in 1995 with the use of vanadium oxide (VO2, anode) and lithium manganese oxide (LiMn2O4, cathode) in aqueous electrolyte containing lithium nitrate (LiNO3) salt [16]. For the lithium system, the increased operational stability window to >2 V has been achieved by using highly concentrated aqueous solutions of lithium salts with perfluorinated anions such as bis(trifluoromethansulfonyl)imide (TFSI) [17]. Also, for SIB,

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with ultrahigh solubility of sodium bis(fluorosulfonyl)imide (NaFSI) salt in water (up to 37m) electrochemical stability window of 2.6 V has been achieved [18]. The sodiation of the intercalation cathode material Na0.44MnO2 (NMO), is investigated in aqueous as well as non-aqueous electrolyte systems. Where the NMO samples showed better rate capability in aqueous based 0.5 M sodium sulfate electrolyte than the non-aqueous 1 M sodium perchlorate system. The apparent diffusion coefficients of Na+ ion in NMO are determined to vary drastically and be in the range of 1.08 × 10−13 –9.15 × 10−12 and 5.75 × 10−16–2.14 × 10−14 cm2 s−1 for aqueous and non-aqueous systems, respectively. The huge differences in the rate capability, as seen in Figure 4.3, are mainly attributed to magnitude of difference in the observed diffusion coefficient and charge transfer resistance from solution as well the resistance from the formed SEI layer [19].

4.2 Background of Cathode Materials The electrochemical energy storage devices with simultaneous transport of ions and electrons during usage faces many forms of resistance, as in first and foremost coming from the in built heterogenous electrochemical

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structure where electrodes are ionic as well electronic conductor but electrolyte remains to be ionic conductor. In electrodes, resistance in a material increases with dimensions and size hence affect the ion and electron kenetics. Therefore, in a battery charge transport can be improved by reducing size, i.e., nanostructuring of electrochemically active materials. As ionic insertion (discharge) and de-insertion (charge) steps, for example in cathodes in half cell configuration, can be made efficient, it will lead to an increase in both power as well as energy density. Being the most important component of batteries, the cathode materials demand more attention, owing to their irreplaceable role in the electrochemical performance for improvising specific energy, cycling life, and specific power. There has been tremendous development with this idea of achieving all the good in one system. In next few sections of this chapter, some of the representative work will be highlighted to emphasize and demonstrate the need of nanostructuring. The reduced dimensions significantly increase the rate of lithium insertion/de-insertion, as interparticle distances for lithium-ion transport shortens. The characteristic time constant for diffusion is given by general equation t = L2/D, where L is the diffusion length and D is the diffusion constant. It is obvious, the time ‘t’ for insertion decreases with the square of the particle size, as it changes from micrometer to nanometer size range. Hence there is tremendous scientific interest for functional nanostructured materials with internal or external dimensions of the order of nanometers. However, irrespective of its importance studying the effect of diffusion length on the electrochemical properties is complex as the nanostructuring of an electrode material typically results in high charged surface area and a high density of inherent defects. The emergence of these two phenomenon leads to seeding of many side reactions with unexpected change

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in energy of system and side products. Hence attention has to be given to look for optimum limit of downsizing the particles so that the benefit of reduced diffusion length can be achieved without compromising on the stability and longevity of electrode materials. In another approach with the same above-mentioned quest, Sang Bok Lee, Gary Rubloff and colleagues showed that nanofabrication and miniaturization could be used to overcome the limitations of energy storage in the form of a nanopore battery [20]. Where anode, cathode and current collector were all fabricated via ceramic membrane and called it as an ‘all-in-one’ nanopore battery. By careful control of growth conditions, the characteristic dimensions of the materials can be fabricated and regulated individually [20].

4.3 Important Types of Cathode (Class) with Different Electrolytes There are different types of cathode classes discovered and applied until now. In this section we will describe these different cathodes in nano dimensionality and their interaction with aqueous and non-aqueous electrolytes.

4.3.1 Li-ion based Nano Cathodes with Aqueous Electrolyte a. Layered oxides: The very primitive cathode material LiCoO2 belongs to the class of layered oxide materials. The inherent problem of this material is the fact that the 100% removal of the Li-ion is not possible from the structure, as it leads to the separate spinel phase formation which on discharging cannot take up the lithium into structure. Hence limiting the theoretical expected capacity to 140 mAh g−1 instead of 274 mAh g−1. Also, the high rate capability has been an issue with this material. However, it has been demonstrated that nano-LiCoO2 electrode material could deliver up to 143 mAh g−1 at 1000 mA g−1 in 0.5 M Li2SO4 aqueous electrolyte. The nano materials could further provide 133 mAh g−1 specific capacity at 10,000 mA g−1 current density, suggesting better ionic conductivity in aqueous electrolyte system [21]. During the cycle of insertion and de-insertion of ions the common feature observed is the hysteresis in charge and discharge profiles, suggesting irreversible capacity loss and polarization in system which is also reflected in charge transfer resistance values. The most notable feature was a very small difference (1.92 Ω) in the charge transfer resistance before and after a chargedischarge analysis indicating that the Li+ ions had been easily diffused during the lithiation

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de-lithiation process with 5 M LiNO3 aqueous electrolyte [22]. In the layered class of family other than LiCoO2, there are scarce of report for the LiMnO2, LiNiO2 and LiNi1/3Co1/3Mn1/3O2 (NMC, NCM) cathode materials with aqueous electrolytes. NMC electrode which is strong competitor to the LiCoO2 in commercialized Li-ion batteries is still lagging behind it in aqueous electrolyte performance. The stability of NCM in aqueous electrolytes is highly pH dependent, as investigated thoroughly in study by Aziz et al. At all pH values an oxidation peak could be observed near 0.55 V vs. SCE in first charge, corresponding to the Li+ extraction. In subsequent cycles there was a shift in the oxidation peak to higher potentials, notably at neutral (7) and alkaline (9) pH, suggesting side reactions. The stability of NCM was improved at alkaline pH of 11. However, the O2 evolution potential at pH 13 began to overlap with the redox potential of NCM. This suggests full utilization of Li+ ion extraction and insertion into the NCM electrode in aqueous electrolyte is still possible. The relative instability of NCM at low-pH is linked to H+ insertion into the NCM electrode. However, nanoporous NCM electrode with 0.5 M Li2SO4 aqueous electrolyte could deliver specific capacity of 155 mAh g–1 at 1.5 C (1 C = 160 mA g–1) and showed retention of a reversible specific capacity of 108 mAh g–1 at higher rate of 45 C. The cyclic stability and power capability were enhanced significantly for these nanoporous NCM compared to that of conventional NCM electrodes, as shown in Figure 4.4. Also, retention of more than 90% of the initial capacity after 50 cycles [23]. After 50 cycles, the discharge capacity of 65 mAh g−1 could be obtained in the charge time of 15 s, which is higher to those of the previously reported NCM nanoparticles in the aqueous electrolyte. The improvement in the rate capability and the cycling life of nanoporous NCM was attributed to its unique structure. Where the nanosized particles shortened the diffusion length of Li+ ions and this not only provided the conducting pathway for the Li+ ions into the cathode as observed by EIS but also acted as a buffer for volume change during cycling. b. Spinel: LiMn2O4 is the highly explored cathode materials with aqueous electrolyte. Since its earliest report in aqueous battery with VO2 as anode and 5 M LiNO3 aqueous electrolyte [16] there has been a continuous uprising. Also, in view of recent results where it has been shown that nanostructured materials as positive electrode materials for aqueous LIBs present relatively better electrochemical performance, nanostructuring of LiMn2O4 has gained momentum. For example, LiMn2O4 nanotube showed a superfast second-level charge and discharge capability and excellent cycling behavior because of the nano nature and preferred orientation [24].

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For the charge time of 6 s, the specific capacity of 53.9% could be retained. Further, the excellent cycling behavior is observed owing to the porous tube structure which helped to buffer the strain and stress arising due to Jahn–Teller distortion. In many of the inherent issues of LiMn2O4, like Jahn–Teller distortion, two-phase formation, crystallinity loss, and development of micro-strain had been suggested to be remarkably controlled by reducing the particle size [25]. The common issue of biphasic nature at a single point of lithiation/ de-lithiation leads to development of different domain structure. Though, in the case of the nano structure entire domain can spontaneously switch between cubic and tetragonal structures on lithiation/de-lithiation. The resultant 13% anisotropic change in lattice parameters gets adjusted by slippage at the exterior domain grain boundaries, improving performance [26] (Figure 4.5). However, the electrolyte is not mentioned in present study. Another model based study carried out on bulk LMO has confirmed that higher discharge rates induce the intercalation-induced stresses in LiMn2O4 electrodes [27]. Furthermore, stress and heat generation drastically increases with larger particle size and higher potential sweep rates [28]. c. Polyanion: The representative material from polyanion class is LiFePO4. The electrochemical oxidation of LiFePO4 in aqueous electrolyte is somehow similar to that of non-aqueous media forming FePO4. However, the reduction of FePO4 to LiFePO4 is not fully reversible since it forms a mixture of LiFePO4 and Fe3O4. All these products formed were confirmed with X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS) and

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secondary ion mass spectrometry (SIMS) [29]. Later it was shown that capacity fading was observed after repeated battery cycling with aqueous 0.5 M Li2SO4 electrolyte. The high OH− and dissolved O2 content revealed to accelerate the fading of LiFePO4. Mössbauer spectroscopy data suggested the presence of Fe3+ ion containing species in the active materials due to the irreversible secondary reactions [30]. It has been demonstrated that the capacity fading of LiFePO4 is not only due to chemical instability but also lies in its electrochemical instability. In a systematic study of LiFePO4, authors also have shown that the discharge capacity of LiFePO4 drops approximately linearly with increasing average particle size, irrespective of the presence/absence of a native carbon coating. The electrode resistance, observed from EIS, as a function of particle size followed almost the square law: Rm/dn (n = 1.994), as seen in Figure 4.6. In LiFePO4 the inherent problem reported is lesser conductivity, however, the ionic conductivity is further less than the electronic conductivity. Hence, through this study, the importance of creating more ions at the surface while reducing the particle size, termed as ionic coating is found to be of greater significance than the carbon coating resulting into better electronic conductivity [31]. However, there are many reports where optimum performance has been observed with aqueous electrolyte by surface modification with electronic conducting materials [32]. Another polyanion, Li3V2(PO4)3/C with saturated LiNO3 aqueous electrolyte, has shown potential for aqueous rechargeable lithium-ion batteries due to its high operating potential and achieved reversible capacity. The modified Li3V2(PO4)3/C could also deliver an initial capacity of 100 mAh g−1 at 5 C rates, despite an electrolyte pH condition ranging from acidic to neutral to alkaline [33].

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4.3.2 Li-ion based Nano Cathodes with Non-Aqueous Electrolyte Compared to the aqueous Li-ion battery more reports are available for the non-aqueous electrolyte cathode and also their different size effects have been deeply explored. a. Layered oxides: The layered LiCoO2 LCO, being first to be commercialized with LiPF6 in EC:DMC electrolyte and worldwide use, is also the one which has been tried to be modified in terms of its dimensionality to great extent. Variable particle sizes have been obtained by sol–gel method. The heat treatment at different temperature could also result in LiCoO2 in range of particle size (here, 30 to 50 nm). The reversible discharge capacity of 154 mAh g−1 was reported with 1 M LiPF6/EC:DEC for optimum sized LiCoO2 nanoparticles [34]. In another study, a template method was used to synthesize nano size LiCoO2 by a modified sol–gel method where P123 was used as both surfactant and chelating agent. As in above study, here also authors controlled the particle size by controlling the temperature and obtained particles in the range of 50–120 nm size. The sample with 120 nm size provided higher discharge capacity of 150 mAh g−1 with optimum cyclability compared to the samples of particles less than 100 nm size [35]. The template method also helped in modifying the surface chemistry resulting into better performance for even 120 nm particles size. However, for LiCoO2 synthesized by co-precipitation method resulting into 20–100 nm sized particles gave

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discharge capacity of 100 and 130 mAh g−1 in 1 M LiPF6/EC/DMC (1:1) [36]. The nanostructured fibers had also been used to achieve a fast-solidstate diffusion due to the shorter diffusion path for Li+ cations. These nanostructured electrodes offered an initial discharge capacity of 182 mAh g−1 compared to ~140 mAh g−1 of conventional powder and film electrodes, which strongly provide a methodology to support the construction of the microscale lithium-ion batteries [37]. Till very recently the formation of SEI layer was only envisaged to be formed over anode, however, due to nano sizing of cathode material the formation of SEI becomes inevitable driven by the high surface energy than the bulk materials. Among the publication related to SEI on cathode, Liu et al. reported the distinguished SEI layer thickness of 2–5 nm on the LiCoO2 nanoparticle which grew to be more uniform on charging/ discharging steps. On the other hand, it can be formed on the surface of nanosized particles even after just storing in dimethyl carbonate (DMC) solvent [38]. Nanoparticles of another layered material, NCM, were synthesized by Patoux and Doeff. Where precursors were calcined between temperature 600–1000 oC giving particles of average range 9–70 nm [39]. In the comparative study with solid-state synthesized bulk samples the irreversible capacity loss could be minimized for synthesized nanoparticles while cycling with 1 M LiPF6/EC:DMC (1:2). The structural analog of LiCoO2, Li2MnO3, only differs in extra lithium ions at transition metal sites. However, this compound suffers from very high irreversible loss in first discharge. It has been now accepted that partial oxygen loss and a structural reconstruction process occur on the first cycle. Where oxygen loss occurs at the interface between oxide particles and the electrolyte from crystal lattice, and therefore the particle size of active materials is the important factor. After partial oxygen loss on the first charge, tetravalent manganese ions in the ccp lattice could partly reduce to a trivalent state upon discharge. While changing the particle size and observing its effect it has been reported that a nanosized Li2MnO3 sample could deliver a large reversible capacity of 260 mAh g−1, while a micrometer-sized sample was found to be nearly electrochemically inactive, Figure 4.7 [40]. The reducing particle size alters the MeO2 inter-slab distance as well as M–M distance. Electrochemical properties of substituted Li-excess layered electrode material, Li1.2Co0.13Ni0.13Mn0.54O2, of different primary particle sizes was investigated for its phase transition occurring on continuous electrochemical cycles. Although the nanosized (  0.5) cathodes. Several materials have been identiffied as promising electrode materials for EVs such as layered mixed transition metal oxides, layered lithium rich materials. The most technologically advanced material option is layered nickel NCM with numbers indicating the decimals of the elements (e.g. NCM523: LiNi0.5Co0.2Mn0.3O2, NCM622: LiNi0.6Co0.2Mn0.2O2 and NCM811: LiNi0.8Co0.1Mn0.1O2). Traits associated with the end members are higher capacity (nickel), better rate capability (cobalt) and improved safety (manganese). Manganese with its constant oxidation state of Mn4+ during cycling can act as a structural stabilizer. While the majority capacity contribution is from the redox couples of Ni2+/Ni3+ and Ni3+/Ni4+. The low cation mixing of stoichiometric compounds could be obtained after introducing the cobalt element into the synthesis process. The capacity of 150 mAh g−1 was obtained by the battery with symmetrical NCM333 as cathode materials during the potential window of 2.5–4.3 V. Recent study demonstrated that the cost of cycle stability can bring out the high capacity

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Figure 5.8 Electrochemical analysis of LiMn2−xNixO4 (x = 0, 0.1, 0.2). (a) Voltage profile for these three samples at 1 C in the voltage range of 3.5–4.5 V; (b) Voltage profile of samples with x = 0.1 and 0.2 at 1 C in the voltage range of 3.5–4.8 V; (c) Cyclic voltammogram of these three samples; (d) Rate performance of these three samples; (e) Voltage profile of the sample with x = 0.1 at increasing C-rates of 1, 2, 10, 15, and 20 C; and (f) Cycle performance of the sample with x = 0.1 at 10 C in the voltage range of 3.5– 4.5 V. Adapted and reproduced with the permission Ref. [56], Copyright 2013, Elsevier.

of 200 mAh g−1 if the battery tests at a higher cut off voltage of 4.6 V. Since the cost for nickel, cobalt and manganese precursors do not vary significantly, an increase in nickel content in NCM is an effective approach to achieve high capacities at the same cost. Actually NCM523 has already been employed as the current generation of some EVs. Unfortunately, introducing higher nickel contents into NCM still needs to overcome

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several challenges. Highly reactive Ni4+ species, which will be predominant at the end of charge, could lead some undesired side reactions with the electrolyte solution, resulting in a consumption of active material, gas evolution and capacity fading. The decay of the high temperature stability of the material will become more serious with increasing nickel contents in NCM, leading to serious safety concerns. It is reported that the onset temperature could decrease from 250oC to ca.150oC for NCM433 to NCM811, respectively. That means that it will become more dangerous (e.g., internal short circuit) if LIB for consumer applications need to store more and more energy. Another report on NCM333 showed that the thermal stability fell from 210 to 180oC once the material was charged to 4.2 V or 4.4 V respectively. It might be attractive to reach higher capacities by increasing the operating voltage of low nickel content materials but will decrease thermal stability. The delithiated state of highly reactive Ni4+ species in NCM materials can easily convert to more stable phases under elevated temperature accompanied by the reduction of Ni4+ to the more stable Ni2+, resulting in oxygen release to ensure charge balance. The phase changes happen at higher temperatures and less rapidly or low and moderate nickel amounts (NCM433, NCM523) while the changes take place at lower temperatures with an increasing in nickel content and are accompanied by much more oxygen release. The formation of a cathode solid electrolyte interface (cSEI), which takes place at the first cycles, is also associated with the presence of reactive Ni4+ at the surface of the charged material. The unstable formed layer might persist throughout the cycling while the consumption of electrolyte solution accompanying with the formation of cSEI will result in lower coulombic effficiencies and capacity fade. Furthermore, the insulating species at the surface of NCM will hinder the lithium diffusion to NCM and subsequently increase electrical impedance. The reduction of surface Ni4+ and the oxidation of the electrolyte solution will be accompanied by gas evolution, which means that more Ni4+ presented at higher cut off voltages result in more side reactions with the electrolyte. Except for gas releasing from reaction between NCM and electrolyte solution, high nickel contents in NCM will lead to a Li/Ni cation mixing and the formation of the inactive phase on the surface will lead to the capacity fade. Even worse, cracks in the secondary particles along the grain boundaries could be observed for NCM811 during cycling, leading to a continuous increase in surface area and hence more active sites for parasitic reactions, leading to lower onset temperature and safety concern. Coating techniques which can separate the active NCM material from the electrolyte solution might be a good option to suppress the gas release and increase the thermal stability. The

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most common approaches to increase the thermal stability of active materials and suppress the gas release are wet chemical coatings (SiO2, TiO2 and Li2ZrO3 coating) which are deposited onto the active material, which can separate the reactive charged material surface from the electrolyte solution, leading to an increase in thermal stability and a suppression in gas release for NCM materials. It is a challenge to prepare uniform coating layer on active NCM materials for most wet chemical coating processes. Atomic layer deposition (ALD) technique, a better but also more expensive method, could be the relatively new, effect and alternative way to get uniform coating. ALD can controls the stepwise growth of coating layers with the desired/optimized thickness. Therefore, PACS is an alternative method to provide carbon-coated NCM nanoparticles for LIB cathodes. The carbon can prevent the agglomeration of NCM during high temperature annealing process, resulting in the smaller particle size. Carbon coating also can solve the problems that NCM cathodes are facing the performance loss as the result from the surface lithium residue, side reactions with the electrolyte and structure rearrangement during long-term cycling.

5.4 Oxygen Deficiency Most metal oxides suffer poor electronic conductivity and thereby result in the lower electrochemical performance. In addition to incorporation with conductive carbon-based materials, doping other elements and generating oxygen vacancy to improve their conductivities have also been explored. Oxygen vacancy generation in TiO2 nanotubes and MoO3−x have been demonstrated as an effective way to enhance the capacity [62, 63], since oxygen vacancy can increase the electronic conductivity and ion diffusion, leading to high capacity and rate performance and cycle durability. Our group also reported that Nb2O5−x@carbon matrix anode exhibited high rate and high durability as compared with insulator stoichiometric Nb2O5 [64].

5.5 Summary and Future Perspectives We have summarized our group recent development of anodes and cathodes with improving conductivity for high performance LIBs. Various strategies have been discussed. Carbon-based materials, e.g., graphene, CNTs and porous carbon, have been the first choice to be utilized to

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composite with the silicon, oxides or sulfides active materials in order to increase their conductivities and subsequently improve the battery performance. Besides, oxygen vacancy introduction on the active materials could also be utilized to improve the conductivities for enhancing the battery performance. To make full utilization of carbon-based materials, we have developed a so–gel strategy to uniformly coat oxides on the surface of graphene or CNTs. The strong affinity between metal ions and functional groups on graphene/reduced graphene oxide or CNTs can result in the uniform coating. However, the synthesis involves strong acids and high oxidant reagents. The electrode materials with surface carbon coating can enhance the conductivity and ion diffusion rate. Carbon-based material is used as a conductive media to connect active material nanoparticles and thus improve the battery performance by improving charge transfer and ion diffusion and maintaining structural integrity [65, 66]. Dual conductive network (carbon and CNTs or carbon and graphene composite) serve asan conductive matrix in which active materials are immersed into the matrix. This conductive network with superior conductivity can enhance the ion diffusion, charge transfer and maintain the material structure during high rate charge/discharge process. Thus, it is expected that the active materials sandwiched by dual conductive layers would further decrease the contact resistance of the electrode for better LIB performance [67, 68]. Carbon coated SnO2/graphene nanosheet is an example of using dual conductive networks to enhance the electronic conductivity [67]. SnO2 nanoparticles were wrapped by graphene nanosheets and carbon layer, leading to large reversible capacity and high cycle durability. This carbon coated SnO2/ graphene nanosheet could maintain the structure and avoid the detachment of active nanoparticles during cycling. 2D core–shell graphene@ silicon@carbon has been proposed as well as the high-performance electrode materials [68]. Our group has developed facile and one-step PACS method for making both anode and cathode materials for LIBs. A thin layer of carbon coating on the surfaces from in situ decomposition of polymers helps to connect nanoparticles to form a unique nanoparticle network. Co3O4, NiO, TiO2, Bi2O3, and V2O5 grown on nickel foam have been explored for LIB anodes without binder and extra carbon and showed higher performance. V2O5, LiMn2O4, LiMn2−xNixO4, and NCM have been studied for LIB cathodes. PACS method will be applied to grow oxide epitaxial films as LIB anode and cathode for further study the electrode/electrolyte interface and in situ FTIR and in situ Raman measurement.

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6 Li–Air: Current Scenario and Its Future Saravanan Karuppiah, Remith Pongilat and Kalaiselvi Nallathamby* CSIR-Central Electrochemical Research Institute, Karaikudi, Tamilnadu, India

Abstract Lithium-air batteries are the most promising rechargeable energy storage devices for the future society, which aims low-carbon economy and sustainability. The significant characteristic of a lithium-air battery is the exploitation of high energy density lithium metal electrode with an air cathode, resulting in huge theoretical energy density comparable to that of gasoline engines. The fundamental reaction in the battery is the electrochemical oxidation of lithium at the anode and reduction of oxygen at the cathode during discharge. However, there are numerous scientific and technical challenges, that blocks the system to solve the energy issues of the present e-society. In this chapter, we explain the classification, fundamental mechanism and recent advances in the research on lithium-air batteries, with a focus on challenges and approaches for making lithium-air batteries for large scale applications. First section of the chapter deals with the general characteristics and significance of lithium-air batteries and challenges associated with them. In the following section, recent advances made in the aqueous and non-aqueous lithium-air batteries are explained in a materials chemistry perspective, including details on electrolytes, cathode electrocatalysts and lithium metal anodes. Finally, a perspective for the future research directions are provided. Keywords: Li-air batteries, metal-air batteries, lithium anode, lithium dendrites, high energy density, electrocatalysts, electrolytes, air-cathodes

6.1 Introduction: Why Lithium–Air Batteries? Smart grid is featured as the best alternative energy storage and conversion system for the future due to the fast depleting fossil fuel reserves and *Corresponding author: [email protected] Poulomi Roy and Suneel Kumar Srivastava (eds.) Nanomaterials for Electrochemical Energy Storage Devices, (291–375) © 2020 Scrivener Publishing LLC

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alarming global warming issues [1–3]. Hydrocarbon fuel upon burning discharges immense amount of CO2 to the atmosphere and the accumulated CO2 increases the global temperature to a large extent [4]. The recently signed COP21 and COP22 agreements have driven the research community across the globe to look in to the development of indispensable strategies to address the issues due to global warming [5, 6]. According to the United Nations Climate Change Conference, an agreement to set a goal of limiting global warming to less than 2 degree Celsius (°C) has been signed by many countries and the agreement calls for zero net anthropogenic greenhouse gas emissions to be reached during the second half of the 21st century. To accomplish this target, carbon emission from gasoline engines must be reduced significantly and the only way to implement this regulation relies on the strict compliance of electric vehicles driven transportation in the place of gasoline vehicles. Such a widespread practice of e-mobility, addressing the requirements of two, three and four wheelers along with the e-market driven requirements pose unequivocal development of electrochemical energy storage systems. Because, electrochemical energy storage devices with environment friendly reaction mechanism, good round trip efficiency, better lifetime and low maintenance become the ultimate choice to offer end to end solution in this regard [7–9]. The use of portable electronic devices developing at a rapid pace, press the need for energy storage and conversion devices with high energy and power density. In this regard, lithium-ion batteries are being considered for small e-gadgets application since their energy density values are better than the Ni-MH batteries [10]. The maximum achievable energy density of lithium-ion batteries is limited by the intercalation electrodes, with an energy density value of 200 Wh kg−1, while the theoretical energy density of gasoline is 13 kWh kg−1 [11]. Many types of secondary batteries are commercialized so far and the Ragone plot describes (Figure 6.1) the comparison of specific energy and specific power of different battery chemistries. The most common batteries like Lead-acid (Pb-H), Nickel Metal hydride (NiMH) and different Li-ion based chemistries demonstrate an energy density of 2.0 V cut off is that the discharge product Li2O2 can be formed without full cleavage of O–O bond, provided a suitable catalyst is added in the cathode. Also, process of formation of Li2O2 is kinetically most favored than the formation of Li2O. Similarly, formation of Li2CO3 as the discharge product has also been reported by many researchers, indicating the involvement of passivating side reactions with the electrolyte solvents in the presence of highly reactive reduced species like O2−, O22− (Li2O2), LiO2, and LiO2 −. These parasitic side reactions and undesired products directly affect the catalytic activity and subsequently the ORR and OER mechanisms. More detailed studies on mechanism is warranted with insights substantiating the actual reaction mechanism governing the performance characteristics of Li–O2 battery. It is unclear at this stage to ascertain which reaction pathway is the major reaction or what are the possible reactions that coexist in a Li–O2 battery during discharge and charge processes. The process of formation of Li2O2 can be divided into two major classifications, as displayed in Figure 6.7 [34]. The reaction which takes place on the surface of the cathode that can be called as surface mechanism, Li+(sol)

2L

iO

Surface Mech. Li+(sol) + e– LiO2* Li2O2 (s) e–

Solution Mech.

2*

2*

Li+(sol) + O2–(sol)

LiO

Li+(sol) + O2(g)

DO2–

Li

2O 2 (s

Li tor 2 O2 oid

)+

LiO2* Li2O2 film Carbon cathode

Figure 6.7 Two general mechanisms of Li–O2 battery discharge. Reprinted from Reference [33]. Copyright 2017 American Chemical Society.

O

2 (g

)

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occurs in a poorly Lewis acid/base solvent where the solubility of LiO2 is not observed and the reaction in which LiO2 have limited solubility, forming Li2O2 through a disproportionation reaction at a nucleated Li2O2 crystal. Micrometer sized toroids are reported to be formed as a result of this reaction and will get deposited over the cathode surface. During discharge, if Li2O2 is formed through the surface mechanism, it will deposit the formed Li2O2 on the cathode surface. As a result, catalyst surface will be blocked for further discharge process and the solution reactants become inaccessible for the catalyst (Figure 6.8aI) [34]. On the other hand, if the solution mechanism is the favored route (Figure 6.8aII), the reaction will take place smoothly and the discharge will persist for longer duration. Figure 6.8c shows the fast oxygen reduction reaction activity

(a) Discharge I. Li2O2 surface growth blocks catalyst sites

Li+

II. ORR fast on carbon (panel c)

O2

LiO2,sol

Li+ O2

Low η

Li2O2 Catalyst/carbon

e–

e–

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Li2O2

Li2O2

O2

Li2O2

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log j (μA/cm2)

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0.15V charge

0

0.15V

2.0

3.0 2.5 U (V vs. Li/Li+)

0.1V 3.5

Figure 6.8 Factors affecting the effective heterogeneous electrocatalysis beyond carbon in Li–O2 batteries. Reprinted from Reference [33], Copyright 2017 American Chemical Society.

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even on carbon surfaces in the case of non-aqueous solvents in the presence of Li+ ions. The loss in polarization occurs during discharge due to the overall impedance of the cell, sluggish diffusion of oxygen (especially at high rates >1 mA/cm2) through the pores and the transport limitations of lithium ions through the SEI layer and deposited Li2O2 (trio interface) rather than due to the reduced ORR kinetics. It is also mentioned in the literature that the electrocatalyst with strong binding affinity toward Li2O and O2 will selectively activate the four electron oxygen reduction reaction pathway resulting in parasitic side reactions which will eventually degrade the electrolyte. The degradation of electrolyte results in poor OER efficiency compared to the pristine carbon cathode, resulting in poor cell rechargeability. In the case of charge reaction, the observed higher overpotential is mainly attributed to the insoluble solid Li2O2, as it is difficult to diffuse back to the catalytic active sites. In addition, due to the formation of carbonate intermediate products at the oxygen evolving Li2O2/electrolyte interface, which eventually slabs the O2 evolving surfaces, overpotential becomes unavoidable. It is further observed from the Figure 6.8c that the oxidation overpotential of peroxide is quite low and hence is not responsible for the increased overpotential during charge reaction. Y. S. Horn et al. have studied the oxygen reduction reaction (ORR) catalytic activity of polycrystalline palladium, platinum, ruthenium, gold, and glassy carbon surfaces in 0.1 M LiClO4 in 1,2-dimethoxyethane via rotating disk electrode measurements [35]. They used DME as the electrolyte solvent as it can withstand the voltage window of lithium air system— unlike carbonate based solvents. The obtained relation between catalytic activities of the cells with the discharge voltage is furnished in Figure 6.9. The Li+-ORR catalytic activity follows the order Pd > Pt > Ru ≈ Au > GC. For carbon supported metal nanoparticle, viz., Pd/C, Pt/C, Ru/C, Au/C, and VC, the ORR activity of 100mA g−1 carbon can be reached at a potential of 2.95, 2.86, 2.84, 2.76, and 2.74 V Li respectively, which is consistent with the discharge potential of lithium–air cells. The parasitic reactions of the oxygen reduction reaction intermediates with the carbonate based solvents can influence the discharge/charge voltages and the same is considered as a major hurdle in the development of active catalyst for Li+–ORR for rechargeable lithium–air batteries. Oxygen evolution reaction taking place in a rechargeable lithium–O2 cell during charge is yet another important and deciding factor as far as the performance of the cell is concerned. Similar to the discharge reaction of formation of Li2O2, the decomposition of Li2O2 to oxygen and lithium can also proceed either through a two-electron process (Li2O2 2Li+ + 2e− + O2) or  a one-electron process that involves the formation of LiO2 (Li2O2

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@ 100 mA g–1 carbon

E (V vs. Li)

2.90 Pt/C

2.85 2.80

Ru/C

2.75

Au/C

2.70

VC 50

100 –1 Q (mAh gcarbon )

150

Figure 6.9 Initial discharge profile of Li–O2 cells (assembled individually with Pd/C, Pt/C, Ru/C, Au/C, and VC) at 100 mA g−1carbon. Reprinted from Reference [35], Copyright 2011 American Chemical Society.

Li+ + e− + LiO2). Large polarization and poor cycle life in a Li–O2 cell results due to the large charging potential via sluggish oxygen evolution reaction with the non-conducting Li2O2. The development of suitable bifunctional catalyst is the best strategy to improve the performance of a lithium–air battery by reducing the charge overpotential and conversion efficiency of Li2O2 in to dissociation products. Carbon can act as ORR catalyst, but, it is not effective in oxygen evolution reaction. Carbon will react with Li2O2 discharge product and produce Li2CO3 interfacial layer thereby leading to an increased charging voltage above 4.0 V. So, in order to reduce the OER overpotential and the corresponding charging voltage, a suitable bifunctional catalyst needs to be identified, which is one of the major challenges. Carbon supported PtAu nanoparticles were shown to enhance the kinetics of OER and ORR in rechargeable Li–O2 cells (Figure 6.10) as reported by Yi Chun Lu et al. The use of PtAu nanoparticles reduces the charge voltage and enhances the round trip efficiency of the lithium–air cell [36]. Many bifunctional catalysts have been reported so far including MnO2, perovskite LaNiO3, pyrochlore Pb2Ru2O7–d, Pt/C, MnCo2O4, etc. [37–41]. Even though the precious metal based bifunctional catalysts exhibit better performance than any other catalyst materials, limited availability and high cost of these materials limit their widespread practical application in Li–O2 system. The formation of Li2O2, the insoluble discharge product on the cathode surface hinders the transport of this material to the catalytic active sites during charge and poisons OER activity. It has been demonstrated by

Li–Air: Current Scenario and Its Future 5.0

(a)

E vs. Li [VLi]

5.0

(b)

4.5

4.5

Ar-filled OER @ Pt

4.0

309

Carbon

4.0

O2-filled

3.5

3.5 3.0

3.0

PtAu/C

2.5

2.5 ORR @ Au

2.0

2.0 0

400

800

Q [mAh/gcarbon]

1200 –400

0

Q [mAh/gcarbon]

Figure 6.10 (a) Li–O2 cell discharge/charge profiles of carbon (black, 85 mA/g carbon) and PtAu/C (red, 100 mA/g carbon) in the third cycle at 0.04 mA/cm2 electrode. (b) Background measurement during charging at 100 mA/g carbon of Ar- and O2-filled cells for PtAu/C. Reprinted from Reference [36], Copyright 2010 American Chemical Society.

many research groups that the OER catalytic reaction is the main reaction which determines the charge overvoltage. But, McCloskey et al. reported that in a lithium–air cell using DME as the electrolyte solvent, the presence of electrocatalyst does not change the charge voltage. The reasons for this conclusion are, the intermediates and product (LiO2/Li2O2) are not soluble/ mobile, or their surface diffusion is not sufficiently fast over macroscopic distances to let electrocatalysis to proceed and the catalytically active sites if existing would be blocked by the insoluble nonconductive Li2O2 during discharge. In other words, the oxygen evolution reaction, oxygen reduction reaction and their effect on charge/discharge overpotentials involve complicated processes and insist the need for much more detailed studies to develop advanced bifunctional electrocatalysts for rechargeable Li−air batteries.

6.4 Critical Challenges Development of commercially viable lithium–air battery for applicability in energy grid or vehicle propulsion system requires scientific and technological breakthroughs. The underlying mechanism of the charge/ discharge: the rate limiting steps: suitable catalyst. etc should be studied in detail. Cell design and a porous cathode, which can accommodate the reaction products without any volume changes and with faster reaction kinetics (by transporting reactants to active sites and products away from

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those sites) should be developed. There are many other critical drawbacks related to electrolytes that limit the applicability of lithium–air battery for electric vehicle applications.

6.4.1 Electrolytes Volatile nature of the conventional organic based electrolytes plays a vital role in designing a rechargeable Li–Air battery since it is bestowed with an open cell structure. An electrolyte solvent with higher boiling point must be used in order to alleviate volatility issues. Even though propylene carbonate possesses many advantages like low volatility, wide electrochemical window and wide liquid temperature range, it is not stable in a system with highly reactive intermediate species like O2−, O22 , LiO2 and LiO2 [42]. Also, the performance of current lithium–air batteries depends on the potential window, which should be wide enough to withstand both high oxidation potentials and be stable to reaction with metallic lithium [43]. Similarly, low oxygen solubility and diffusivity results in poor rate performance [44, 45].

6.4.2 Decomposition of Electrolyte During Discharge The efficiency and the mechanism of a lithium–O2 battery is determined by the stability of the electrolyte solution used in the specific cell. Because, various nucleophilic oxygen species like O2−, O22 , and O2 formed during the oxygen reduction reaction will attack the electrolyte solvent and results in promoting undesirable parasitic side reactions [45]. In this regard, identification and development of suitable electrolyte for the reversible operation of an electrochemical cell is the most pressing challenge. McCloskey et al. have compared the solvent and salt combinations using DEMS (Differential Electrochemical Mass Spectrometry) and pressure decay or rise measurements, which are summarized in Table 6.2 [46]. From the table, it is understood that except for DME and DMSO, no other electrolyte is possessing an (e−/O2)chg that is close to an ideal value of 2, even though they do possess an (e−/O2)dis which is close to the ideal value of 2 and OER/ ORR is less than 0.9 for all the electrolytes after a single galvanostatic cycle. EQCM and DEMS studies suggest that DME is the best solvent with less side reactions, but, is not stable for a practical commercial battery. BD McCloskey et al. studied the mechanism of degradation of electrolytes in various carbonate and ether based electrolyte solvents under the influence of constant current cycling using LiTFSI as the electrolyte salt [47]. Figure 6.11 shows the discharge of pure DME based cells wherein

0.88

1NM3

NMP

THF

DME

CH3CNc

TGE

XC72

XC72

XC72

XC72

P50

XC72

BF4 0.75

0.72

0.58

0.48

0.51

0.33

2.04

2.05

2.01

2.01

1.96

2.14

2.05

2.30

2.33

2.71

2.33

2.59

2.80

3.35

4.44

4.05

7.04

6.41

2.65

2.59

2.71

2.59

(e−/O2)chg

0.03

0.04

0.06

0.03

0.03

0.11

0.03

0.01

1.26

0.04

0.05

0.05

0.07

CO2/ORR

0.08

0.01

0.01

0.09

0.02

0.04

0.02

0.28

0.01

0.08

0.08

0.08

0.03

H2/ORR

b

Reprinted with permission from Reference 46. Copyright 2012 American Chemical Society. DME=dimethoxyethane, THF=tetrahydrofuran, TGE=triglyme, CH3CN= acetonitrile, DMSO=dimethyl sulfoxide, NMP=N-methyl pyrrolidone, 1NM3= tri(ethylene glycol)-substituted trimethylsilane, MPP-TFSI=N-methyl-N-propylpiperidinium bis(trifluoromethylsulfonyl) imide. c Experiment performed using lithium iron phosphate (LiFePO4) as the anode. Otherwise, Li metal was used as the anode for all experiments. d LiTFSI=lithium bis(trifluoromethane sulfonyl) imide, LiBF4=lithium tetrafluoroborate, LiTrif=lithium triflate, LiBOB=lithium bis(oxalato) borate, LiClO4 = lithium perchlorate.

a

0.78

DMSO

P50

TFSI

0.36

BOB

MPP-TFSI

0.78

BF4

P50

2.00

0.77

ClO4 2.06

2.00

0.74

Trif

2.01

0.78

(e−/O2)dis

TFSI

DME

XC72

OER/ORR

Solventb

Cathode

Li Saltd

Table 6.2 Summary of DEMS results for various salt and solvent combinationsa.

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4

III I

3

2

II

IV

2.05 e–/O2 on discharge

(a) DME

Cell voltage (V)

5

4

3

2.70 e–/O2

2

(b) 1EC:1DMC

2.35 e–/O2

5

4

3

2.60 e–/O2

2

(c) 1PC:2DME 0.0

2.05 e–/O2

0.5

1.0

Charge (mAh)

Figure 6.11 Discharge behavior of pure DME based cells and comparison of performance in combination with EC and PC as electrolyte. Reprinted from Reference [47], Copyright 2011 American Chemical Society.

a flat discharge plateau at a voltage of 2.65 V is observed. The number of electrons per oxygen consumed in every single discharge is calculated to be 2.05 0.05 with a cathode weight gain of 1.0 mg (0.15 mg/mAh) (Table 6.3). By a careful examination of the electrons, consumed and the electrode weight gain measurements, it is reported that the theoretical value of formation of Li2O2 is less than the experimentally observed values, inferring the involvement of parasitic side reactions. The cathode weight gain for carbonate based electrolytes is very high compared to the ether based electrolytes which in turn could be attributed to the solvent decomposition to form solid carbonate deposits like lithium alkyl carbonates on the cathode. XRD analysis (Figure 6.12a) shows the presence of graphitic carbon peaks and Li2O2 peaks in DME based electrolyte cells, but no crystalline peaks other than carbon related peaks were seen in carbonate based

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Table 6.3 Cathode weight gain during dischargea. Discharge product

Cathode weight gain (mg/mAh)

DME discharge

1.0

PC/DME discharge

2.3

EC/DMC discharge

2.0

Li2O2

0.86

Li2CO3

1.4

Li2O

0.56

a

The standard deviation of all discharge weight gains is ~0.15 mg/mAh (based on replicate trials). b Indicates theoretical cathode weight gain if this is the only product formed during discharge, assuming that the 2e−/product molecule is formed. Reproduced with permission from reference 47. Copyright 2011 American Chemical Society.

Li2O2

20

Li2CO3 Li2CO3 Li2O2 Li2O Li2O2

(a) DME (b) 1EC:1DMC (c) 1PC:2DME

40

60 2θ (a)

80

Arbitrary counts

Log(counts), arb. units

Li2O2

DME 1EC:1DMC 1PC:2DME Neat P50 carbon paper

0

500

1000 2750 3000 Wavenumber (cm–1) (b)

Figure 6.12 (a) XRD analysis showing Li2O2 peaks in DME based electrolyte cells. (b) Raman results evidencing the formation of Li2CO3. Reprinted from Reference [47], Copyright 2011 American Chemical Society.

electrolyte cells [47]. Raman spectra (Figure 6.12b) gave a clear evidence for the formation of carbonates on the cathode surface after discharge. Even though the above studies support DME based electrolytes for Li–O2 batteries, they are unstable during recharge in the presence of Li2O2. Hence, it is clear that stability in the oxidative and reductive medium in the presence of Li2O2/LiO2 is an important requirement for a suitable electrolyte solvent/medium for better performing rechargeable lithium–air batteries.

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6.4.3 Passivation and Blockage of Oxygen Diffusion Normally, the discharge products formed in a lithium–air battery passivate the electrode surface by depositing Li2O2 on the pores thereby blocking the oxygen uptake as shown in Figure 6.13. Since the rechargeability and the energy density depend upon the fresh oxygen intake, the passivation and the blockage results in poor battery performance irrespective of the rate at which the charge/discharge cycling is performed. The reaction products like Li2O2, Li2CO3 (formed especially in carbonate based electrolytes) are electronic insulators in the bulk form. Actually, ideal electrochemical reactions take place on the surface of the porous electrode and so the passivation or blockage of pores can limit the accessible electrochemical active surface area to result in reduced discharge capacity behavior. Figure 6.13 shows the mechanism of pore blocking on the surface of porous carbon during discharge. To ensure the complete/reversible formation of Li2O2, selection of suitable carbon supports with larger pore diameter is extremely important [48].

6.4.4 Large Polarization The most critical challenge in the case of a lithium–air battery is the large charge overpotential even at very low current densities, resulting in a very low “round trip efficiency” of the cell [49]. Normal carbon supports used in conventional air batteries are good catalyst for oxygen reduction reaction, but, they are not sufficient for the oxygen evolution reactions. Many catalysts have been investigated to catalyze these reactions, such as nitrogen doped carbon and its allotropes, metal oxides, metal nitrides and precious

e– Li+

Li+ Vp

e–

O2

O2 Li+, O2

Vp: pore volume Li2O2

C-matrix

Figure 6.13 Mechanism indicating the pore blockage during cycling. Reprinted from Reference [48], Copyright 2012 Electrochemical Society.

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metals (Pt, Pd, Ru, etc.) [50–55]. Metal nanoparticles or clusters are better choices as bifunctional catalyst for lithium–air batteries as they effectively reduce the charge overpotential, but the main hurdle is their high cost. Poor electronic conductivity of Li2O2 and the resulting poor electron transport kinetics to the Li2O2/electrolyte interface and the formation of Li2CO3 may also cause a higher charge potential [56].

6.4.5 Lithium Dendrite Formation Presence of lithium metal in the lithium–air battery is associated with the unavoidable dendrite formation during repeated charge–discharge process, which is quite similar to any other lithium metal based batteries, resulting in battery failure [57]. Dendrite formation is the most common failure mechanism in a lithium based battery, which should be tackled cautiously in order to increase the stability, safety and lifetime of the battery. There are different theories and mechanisms to explain the formation of lithium dendrites. One such proposed mechanism is the SEI formation theory, which is shown in Figure 6.14. According to the theory, lithium dendrite growth is higher in places where Li+ ion conductivity is higher. The speed of the growth of lithium dendrites will be higher at inhomogeneous electrode surface, which leads to the breakage of SEI layers. These broken SEI

Li deposition Solution

Solution

Solution Li+

Li+

Li+

Li+

Li+ S.L.

Li metal

e–

Li+

Li+

Li+

S.L. Li metal

Solution Li+

S.L.

e– e– Volume changes, the surface films crack

Li deposited underneath the surface films

S.L. Li metal

Li metal

e–

Dendrite formation

Li dissolution Solution Li+

Li+

Solution Li+

Li+

S.L.

Li metal

Solution

Li+

S.L.

Li+

Solution

Li+

S.L.

S.L. Li metal

The surface films accomodate the volume changes

Low current densities

Li metal S.L.=surface layer

Li metal

The surface films are broken down and are repaired by surface reactions of Li with solution species

High current densities

Figure 6.14 Litihum dendrite formation mechanism in Li–air batteries. Reprinted from Reference [57], Copyright 2000 American Chemical Society.

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layers are prone to form lithium dendrites at higher rates. Another mechanism deals with the morphology of lithium metal surface, defect sites and grain boundaries. Third mechanism explains that the inhomogeneous current density resulted ion-concentration variation is the main reason for the formation of lithium dendrites.

6.4.6 Electrocatalysis Electrocatalyst in an air electrode plays a vital role in reducing the charge voltage in a lithium air battery. There are different materials like noble metals, transition metal oxides, carbon and its derivatives and composites involving any of these materials which are used extensively as cathode catalyst in lithium–air batteries. Until the discovery of dissociation of electrolyte by parasitic side reactions to form Li2CO3, it was believed that catalysts

U (V vs. Li/Li+)

5.0

XC72 Au MnO2/XC72

4.0

Pt/XC72 3.0

Charge

Discharge

2.0 (a) 0.0

0.5 Q (mAh)

U (V vs. Li/Li+)

5.0 Charge

1.0

XC72 Au/XC72 MnO2/XC72

4.0

Discharge

3.0 (b) 0.0

0.2 Q (mAh)

0.4

Figure 6.15 (a) Galvanostatic discharge−charge of 1 mAh discharge when using 1 M LiTFSI in organic carbonates as the electrolyte and with various nanoparticle catalyst particles on XC72 C cathode. The apparent electrocatalysis is related to the solvent decomposition. (b) Similar experiments carried out using 1 M LiTFSI in dimethyl sulfoxide. Reprinted from Reference [33], Copyright 2014 American Chemical Society.

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are responsible for the change in charge voltage plateau. McCloskey et al. reported that electrocatalysis is observed when carbonates are used as electrolyte solvents for Li−O2 whereas no other research teams observed such an electrocatalyst when relatively stable anhydrous ethers (dimethyl ether) are used as the electrolyte solvent (Figure 6.15). In a carbonate based electrolyte solvent, main charge reaction involves the parasitic reactions with the solvent to form Li2CO3. On the other hand, the OER mechanism is poisoned by the deposits of insoluble Li2O2 after the discharge cycle that blocks the active catalyst sites. Ideally, the electrocatalyst should reduce the charge overpotential for a better performing cell.

6.4.7 Rate Capability The main reactions in lithium–air cells are oxygen reduction and evolution reactions, which are the major rate limiting steps. In a non-aqueous lithium air system, high current densities of 1 mA/cm2 is only reported with carbonate based electrolytes [58]. Rate capability in carbonate based solvents is poisoned by passivation layers by electrolyte decomposition and the formation of deposits on the lithium metal surface [59]. Oxygen solubility and transport are other important parameters affecting the rate capability in a lithium–air cell [60].

6.4.8 Energy and Power Density For the development of lithium air battery for practical applications, energy and power density are important parameters. Normally, the discharge currents are kept in the range of 0.05 to 0.1 mA/g and Li–air electrode current density per gram of (carbon + catalyst) is kept at fairly high values (up to 250 mAg−1). The discharge voltage is in the range of 2.5–2.8 V. The energy density in the lithium air battery mainly depends on the weight of carbon cathode; higher loading results in lower energy capacity. When the cathode weight is less, the discharge current should also be kept at a lower value [48]. As Figure 6.16 shows, a cell producing Li2O or Li2O2 exhibits the same specific energy in the charged state, but in the discharged state, the Li2O2 cell has a lower energy per mass due to only one O2 molecule being consumed per 2 Li atoms, rather than 1/2 O2 in the case of Li2O. The specific energy calculations for the LiOH H2O and LiOH systems in the charged state depend on the assumption about where from the H2O in the discharge product originates.

Nanomaterials for Electrochemical Energy Storage Devices

14 (a) 12 10

Charged Charged w/ H2O weight Discharged

8 6 4 2 0 Li2O

Li2O2 LiOH·H2O LiOH Li/LiMO2

Active-only energy density (kWh/L)

Active-only specific energy (kWh/kg)

318

14 (b) 12 10 8 6 4 2 0 Li2O

Li2O2 LiOH·H2O LiOH Li/LiMO2

Figure 6.16 (a) Specific energy and (b) energy density values based on active materials. Reprinted from Reference [48], Copyright 2012 Electrochemical Society.

6.4.9 Volume Changes In a Li–air system, both the cathode and the lithium metal anode are prone to volume changes during cycling. In a lithium air cell, the mass of the cell increases while discharge, whereas there will be a simultaneous reduction in volume of the cell. There will also plating/stripping during cycling of the cell as the system uses pure lithium metal as the anode. So care is to be taken address those issues [61].

6.5 Non-Aqueous Li/Air Systems Recently, non-aqueous lithium air system (NLAS) is getting tremendous attention, after the first demonstration of lithium air system with a non-aqueous electrolyte by Abraham et al., in 1996 [62, 63], due to its high gravimetric (3458 Wh kg−1) and volumetric (3445 Wh L−1) energy density compared to that of conventional lithium-ion batteries [64].

6.5.1 Electrochemistry of Oxygen Reduction and Oxidation in Non-Aqueous System The general mechanism of NLAS is schematically illustrated in Figure 6.17. As shown in figure, NLAS consists of a porous cathode (air cathode), a metallic Li anode and an electrolyte between the two electrodes. During discharge, the generated Li+ moves across the electrolyte toward the porous cathode, where it combines with O2 (atmosphere) and electrons from the external circuit. Lithium atom and O2 are regenerated due to the reverse reaction

Li–Air: Current Scenario and Its Future

e–

e–

Li anode

Needs for oxygen electrode • High active area and optimized pore structure • High catalytic activity towards oxygen reactions • Stability against electrolyte decomposition • Selective O2 permeability and efficient diffusion • High loading and easy fabrication

• Safety issues • Dendritic Li formation • Low cycling efficiency • Reaction of Li anode with water • Reaction of Li anode with electrolyte

Cathode

Li+

Li Anode

Potential

Li+

Porous Cathode

Specific Capacity Capacity Retention

Li2O2

Discharge

Large overpotential

319

Recharge

Li+ O2

Li+

Electrolyte Non-aqueous electrolyte

Low cyclability

Oxygen electrolyte

• Stability agaist super radical ion • Electrolyte decomposition • Evaporation • Low O2 solubility and diffusivity

Cycle No.

• High charge potentials • Low round-trip efficiency • Slow kinetics of oxygen reaction • Formation of insoluble products

Figure 6.17 Schematic diagram and major technical challenges associated with the rechargeable non-aqueous Li–air battery (Courtesy: K.-N. Jung et al., J. Mater. Chem. A, 2016, 4, 14050).

when NLAS is charged and the generated O2 is released from the battery or retained in the reservoir [65]. Hence, the reversibility could be understood. The reaction mechanism is given below [66–72] Anode

Li

Li

e Ea0

3.05 V vs. SHE

Cathode

O2 2 Li

2e

Li2O2

Ec0

0.09 V vs. SHE

Overall battery reaction

2 Li O2

Li2O2

0 Eocv

2.96 V vs. SHE

2 Li O2

Li2O2

0 Eocv

2.91V vs. SHE

Based on the literature reports, the major discharge product formed at the air cathode is believed to be Li2O2 rather than the formation of Li2O. In other words, it is proposed from the computational study that the surface of the Li2O2 is half-metallic unlike the non-metallic and nonmagnetic surface of LiO2, which supports the Li2O2 formation in NLAS

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Nanomaterials for Electrochemical Energy Storage Devices

during discharge [73]. Moreover, the reversible formation of Li and O2 is not highly reversible from the LiO2 during charge process compared to that of Li2O2. It is widely accepted that the performance of Li–O2 system is greatly affected by the generated products during discharge and charge process. Hence, an in-depth understanding about the electrochemical reactions, i.e., ORR and OER is felt to be critical to design a new electrocatalyst for NLAS system. Basically, oxygen electrochemistry in NLAS is a very complicated process and it depends on both the electrode materials and electrolytes. Based on the literature reports, two major discharge mechanisms have been proposed [74–84]. Discharge reaction mechanism I

O2 e

O2 O2 O22

2O2 O22

2 Li

Li2O2

Discharge reaction mechanism II

O2 e O2 2 LiO2

Li

O2 LiO2 Li2O2 O2

Interestingly, supporting electrolytes with larger or smaller cation have greater influence in the discharge mechanism, whereas the anions have little influence on ORR and OER. For example, supporting electrolytes with large cations, such as tetrabutylammonium or tetraethylammonium follows reversible oxygen reduction and oxidation with one electron transfer reaction near the equivalent potential [85–87]. Here, the oxygen molecule is first reduced to form O2− and with the subsequent solvation based on TBA+ or TEA+, and O2− is oxidized at a relatively smaller overpotential. In another case, if the electrode potential is lower, the O2− is further reduced to form O22 , which can be oxidized in the higher overpotential and the same is the reason for the 2 V gap between the oxidation and reduction potential

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[70]. It is important to mention here that the noble metals (Pt, Au) as well as the carbon based electrocatalysts show identical electrocatalytic performance in non-aqueous electrolytes with larger cations [78, 88]. The possible reason behind the above mentioned mechanism is given below: When the electrolyte contains larger cation, the oxygen reduction will not involve the O–O bond cleavage mechanism, which in general is considered as a sluggish process, thus requiring noble metal catalysts. In addition, O2− is the weakly adsorbed radical which is highly soluble in electrolyte, and so the electrode under such conditions may just work as an electron transfer media only.

On the other hand, the oxygen reduction in smaller cation (Li+, Na+) containing non-aqueous electrolyte first forms O2− and it combines with the Li+, generating LiO2 on the surface of the air cathode. Further, LiO2 disproportionates in to more stable Li2O2 because LiO2 is unstable. This mechanism has been further evidenced by the in situ Surface Enhanced Raman Spectroscopy (SERS) and XRD analysis [74, 78, 60]. Similar to discharge, there are two major mechanisms proposed for the charge process also. Charge reaction mechanism I

LiO2 + Li+ + e−

Li2O2 LiO2

O2 + Li+ + e−

Charge reaction mechanism II

Li2 O2

O2 + Li+ + e−

In OER process, Li2O2 is directly oxidized to O2 and Li+ without involving LiO2 in the case of non-aqueous electrolytes containing larger cations. Apart from the above proposed mechanisms for ORR and OER, there are few other mechanisms reported in the literature. For example, the proposed ORR mechanism at the cathode side by Hummelshoj et al. is given below [69]: *denotes a surface site on Li2O2 where the growth proceeds

Li

e

O2

Li

e

LiO2

LiO2 Li 2O 2

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Oh et al. proposed a charging mechanism of the following type [89]:

O22

Li2O2

2 Li

O22

O2

O22

O2 e

e

As on date, no commonly accepted mechanism is available in the literature. Therefore, significant work is needed to understand the oxygen electrocatalysis mechanisms in nonaqueous systems, wherein various ex situ and in situ characterizations such as X-ray Diffraction (XRD), Raman, Differential Electrochemical Mass Spectrometry (DEMS), Scanning Electron Microscope (SEM), Transmission Electron Microscope (TEM), X-ray Photoelectron Spectroscopy (XPS), Fourier Transform Infrared Spectroscopy (FTIR), Nuclear Magnetic Resonance (NMR) and Atomic Force Microscopy (AFM) could be deployed.

6.5.2 Technical Challenges in NLAS 6.5.2.1 Designing of Air Cathode/Oxygen Transport Engineering of new nanostructured air cathodes plays a critical role in deciding the electrochemical performance of lithium–air system. The air cathode must be porous with sufficient porosity (to ensure the transport of all reactants viz., O2, Li+, and electrons), minimal tortuosity (to facilitate O2 transport to the active sites for the electrode reactions on the interior surface of the air cathode with minimum energy loss) and provide appropriate space (high pore volume) to accommodate the solid products (lithium oxides) during discharge. Further, it should possess adequate electrical conductivity. In addition, air cathode should ensure high electrolyte wettability to satisfy the requirements of fast ion transfer during the charge–discharge process. More importantly, cathode materials including electrocatalyst should have the ability to accelerate the kinetics of both ORR and OER. In short, the ideal cathode material with optimum structure, morphology and crystal forms provides more space for the storage of discharge products, facilitates the diffusion of oxygen and the formation of discharge product, ensures electrode wettability and enhances the catalytic performance due to the introduction of defects and vacancies [64].

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6.5.2.2 Effective Loading of Catalysts Effective loading and distribution of electrocatalysts plays the key role in minimizing the large polarization between the charge and discharge process. In other words, the enhancement in electrode kinetics could be realized by way of introducing various novel electrocatalysts (noble metal) with proper weight ratio and thereby we can reduce the energy loss attributed to the high over potential. In short, proper weight loading and percentage of the electrocatalysts can critically impact the electrochemical performance of lithium–air system.

6.5.2.3 Slow Kinetics of Oxygen Reactions/Deposition of Solid Insulating Products In general, NLAS exhibits lower round trip efficiency and poor rate capability due to the sluggish reaction kinetics associated with ORR and OER. During discharge, insoluble lithium oxides have been generated and deposited on the active sites of the cathode. Once the product is deposited on the active site of the electrode, ionic as well as electronic transport through the surface of the product is required to sustain the reversible electrode reaction. But, the poor electronic conductivity of the in situ generated discharged products will eventually terminate the electrode reactions by affecting the ion/electron transport (worsening the electronic conductivity) due to the progressive growth of the product layer. In addition, the availability of the active material should be ensured throughout the charge discharge process. In other words, the generated solid discharge product on the cathode side should not block the diffusion of O2 to the active surface of the electrode for the sustainable electrode reactions. Apart from this, the morphology of the deposited products can also play a critical role in electrochemical performance of the lithium–air system, which cannot be ignored.

6.5.2.4 Decomposition of Non-Aqueous Electrolytes/Effect of Possible Side Reactions The intermediates (O2 , O22 ,andLiO 2 / LiO 2 ) formed during discharge in lithium–air system are very reactive and they can easily decompose the electrolyte. As a result, formation of side products becomes unavoidable (Li2CO3, LiOH and lithium alkyl carbonates) rather than the formation of the desired product (Li2O2), which in turn affects the reversibility of the system in a major way.

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6.5.2.5 Lithium Dendrite Formation and Side Reactions of Li with H2O and Air It is well known that the formation of dendrite during the continuous dissolution and deposition of lithium still remains as an unaddressed challenge. In addition, the parasitic reaction of lithium with the impurities present in the electrolyte as well as the O2 (diffused through the separator) can greatly affect the reversibility and efficiency of the lithium–air system. Hence, development of high throughput air-breathing membranes (or other mechanisms) that separate O2 from ambient air is mandatory in order to avoid H2O, CO2 and other environmental contaminants from limiting the lifetime of Li–air batteries.

6.5.3 Electrocatalysts for NLAS In general, the electrocatalysts are believed to play a critical role in achieving the high power density, good round trip efficiency, long durability and improved cycling stability in lithium–air systems. Theoretically, the high specific energy of the lithium–air system is determined by the amount of product getting formed during discharge, which could be directly correlated with the electrocatalytic performance of the electrocatalysts. Practically, the obtained highest discharge capacity so far in the literature is far lower than the expected capacity value and the rate capability is also inferior compared to that of other energy conversion systems. Toward this direction, efforts are needed to develop novel electrocatalysts to realize the desirable electrochemical performances. Basically, electrocatalysts are generally classified in to the following four categories [64].

6.5.3.1 Carbon Based Materials In recent years, carbon materials have attracted tremendous attention as electrocatalyst for NLAS due to their appealing features like high surface area, desired porosity, high electronic conductivity, high chemical stability and low cost. It is important to mention here that the porosity is extremely important for application in lithium–air systems. Carbon is the best material that can provide the desired porosity and electronic conductivity. There are so many reports available about the influence of carbon materials toward the electrochemical performance of lithium–air system [90–92]. Generally, carbon materials that have been used for the NLAS can be classified in to the following three types

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1) Commercial carbon materials, 2) Functional carbon materials and 3) N-doped carbon based materials.

6.5.3.1.1 Commercial Carbon Based Materials Most of the commercial carbon based materials have been explored as electrocatalyst by various groups and the results are tabulated (Table 6.4). Interestingly, the discharge capacity of the same type of carbon is found to vary significantly even at the same current density condition. This indicates that the electrochemical performance can be influenced by some other factors in addition to the performance of the electrocatalyst. It is widely accepted that the pore volume, especially the mesopores, instead of high surface area, is the crucial factor in determining the cycling stability and rate capability in NLAS. Hence, optimization of mesopores is getting importance since too small or too big pores may lead to a less efficient usage of the available mesopore volume toward performance characteristics. Williford and Zhang [93] proposed a new concept in the design of air electrode, based on the interconnected dual pore system (one catalyzed and one noncatalyzed, or one macroporous and one mesoprous/microporous systems). Based on their modeling results, it is believed that interconnected dual pore system is playing a critical role in obtaining high energy and power density. In other words, the first pore system (mesopores) offers enough space to accommodate the discharge products and the second pore system will remain as unused/empty (macropores) for the usage of oxygen transport throughout the process. The appreciable part of the work is that the first pore system will not clog with the second pore system (the stored products in the first pore system do not clog with the second one ultimately) and this secures oxygen transport into the inner regions of the electrode and improves the utilization efficiency of the pores (Figure 6.18). Table 6.4 clearly evidences the relationship between the obtained specific discharge capacity and the specific surface area of the commercial carbon based materials. For example, carbon materials of Vulcan XC72, KB EC600JD and BP 2000 with surface area values of 240, 834 and 1509 m2 g−1 could display discharge capacities of about 183, 439, and 517 mAh gcarbon−1 respectively. Similarly, Cheng et al.’s study [93] reveal that the discharge capacities of Norit (4400 mAh gcarbon−1), acetylene (3900 mAh gcarbon−1), and super P (3400 mAh gcarbon−1) are in accordance with the relativity of their surface area, i.e., 800, 75, and 62 m2 g−1 for Norit, acetylene and super P carbon black, respectively. Even though commercial carbon based materials are being explored as electrocatalyst, there are lot of issues remain unaddressed including cycling

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(a)

(c)

Li2O

Carbon black

(b) Li2O

Voltage / V (vs Li+/Li)

2.8 2.6 2.4 2.2 2.0 pore window

0.5 mA cm–2

0.2 mA cm–2

0.1 mA cm–2

1.8 0

500

MCF-C

1000

1500

2000

2500

Specific capacity/ mAh g–1

Figure 6.18 Schematic behavior of super P carbon black (a) and MCF-C (b) after discharge (- MFC-C and … Super P) (Courtesy: X.-H. Yang et al., Electrochem. Commun., 2009, 11, 1127).

stability, low round trip efficiency, poor rate capability, poor catalytic activity and high over potential. Hence, most of the recent studies state that commercial carbons can be used preferably as conducting additive or catalyst support rather than air cathode for NLAS.

6.5.3.1.2 Functional Carbon Materials Graphene [112, 113], mesoporous carbon [72], carbon nanotubes (CNTs) [94, 119], carbon nanofibers (CNFs) [120], and carbon microfibers [121] are falling under the category of functional carbon materials and they are being considered as efficient electrocatalysts in lithium–air system due to their unique structures and higher number of defects/vacancies. In this section, few select examples on functional carbon materials are discussed. Graphene has been explored as an electrocatalyst and it delivers an extremely high specific capacity compared to that of commercial carbon based materials. The excellent electrochemical performance in terms of cycle life and round trip efficiency of graphene based materials could be correlated with its unique 3D 3-phase electrochemical area and the diffusion channels for electrolyte access and oxygen diffusion apart from the unique structural advantages. Li et al. [113], have tested graphene nanosheets (GNSs) as electrocatalyst in NLAS and realized a specific capacity of 8705.9 mAh gcarbon−1 at a current density of 75 mA gcarbon−1, which is highly superior than commercial carbon

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Table 6.4 Reported capacity behavior of commercial carbon materials in nonaqueous Li–air batteries. Carbon materials

Capacity (mAh g−1)/current density (mA cm−2)

Ref.

Super P

2120/0.05

[95]

2825/0.05

[96]

989/0.05, 528/0.2, 248/0.5

[97]

44000/0.02, 1400/0.05

[98]

1800/0.1

[99]

1736/0.1

[72]

1000/0.1

[100]

4254.7/0.1, 6587/0.15

[101]

2300/0.1

[102]

~1000/0.2

[103]

400/0.2

[50]

3400/70 mA g−1

[37]

1500/100 mA g−1

[104]

2700/0.025

[105]

850/0.05

[106]

800/0.05

[107]

1000/0.05

[108]

5813/0.1 (1.9 mg carbon loading), 3378/0.1 (4 mg carbon loading), 404/0.1 (12.2 mg carbon loading)

[109]

2600/0.1

[100]

3374.4/0.1

[101]

400/0.1

[110]

800/0.2

[50]

3000/0.2

[111]

KB EC600JD

(Continued)

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Table 6.4 Reported capacity behavior of commercial carbon materials in nonaqueous Li–air batteries. (Continued) Carbon materials

Capacity (mAh g−1)/current density (mA cm−2)

Ref.

Vulcan XC-72

1200/0.04

[92]

762/0.1

[72]

1705.7/0.1

[47]

1645/0.1

[101]

1053.8/75 mA g−1

[113]

1000/50

[74]

850/70

[114]

2000 50/0.05

[106]

1909.1/75 mA g−1

[113]

KB EC300JD

2200/0.1

[100]

Graphite

560/0.1

[1]

250/0.1

[72]

210/0.05

[115]

280/0.05

[107]

170/0.05

[108]

180/0.05

[116]

Norit carbon black

4400/70 mA g−1

[117]

Calgon activated black

80/0.05

[106]

Ensaco 250G

550/0.1

[100]

Chevron activated black

1410/0.1

[1]

Activated carbon SY TC-03

2310.9/0.1

[101]

Activated carbon M-30

2120/0.05

[118]

Denka

750/0.1

[100]

Denka

25/0.05

[106]

Super S

Black Pearls

Darco G-60

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materials (Figure 6.19). The observed increase in capacity is attributed to the presence of active sites at the edge of graphene, which is responsible for the improved electrocatalytic activity toward ORR. Interestingly, Xiao et  al., [122] reported the highest discharge capacity (15000 mAh gcarbon−1 at a current density of 0.1 mA cm−2) for graphene based electrode, using a novel hierarchical arrangement of structural and functionalized graphene sheets. Herein, oxygen spreads rapidly in the microporous channels of the hierarchically porous graphene due to the unique bimodal porous structure. In addition, isolated nanosized Li2O2 particles formed during discharge due to the defects and functional groups present on graphene, prevent the air blocking in the cathode. Properly connected nanoscale is believed to offer high density reactive sites for Li/O2 reactions and the cumulative effect results in the superior performance. Ordered mesoporous carbon (OMC, basically synthesized with a template of silica) could be explored as an electrocatalyst in NLAS due to its high specific surface area, excellent electrical conductivity and fast mass transport. For example, Sun et al. [123] demonstrated the electrocatalytic performance of the OMC in NLAS and compared the results with the commercial carbon catalyst, i.e., super P. As expected, OMC exhibits high specific capacity and low over potential compared to that of super P cathode. There are many other reports available in the literature regarding the cathode performance of OMC, synthesized using various templates. Recently, single walled and multiwalled carbon nanotubes have been investigated as air cathode for NLAS due to their appealing features such as high chemical and thermal stability, high elasticity, high tensile strength

4.8 4.4

c

b

(a)

a

Voltage / V

4.0

a b c

3.6

(b)

(c)

GNSs BP-2000 Vulcan XC-72

3.2

100 nm

100 nm

(e)

(d)

2.8 2.4 2.0

c 0

b 2000 4000 6000 8000 Specific capacity / mAh g–1

a 10000 100 nm

20 nm

Figure 6.19 Charge–discharge performance of Li–air battery with GNSs, BP-2000, and Vulcan XC-72 cathodes at a current density of 75 mA g−1 (a); SEM and TEM images of GNS electrodes before (b and c) and after (d and e) discharge (Courtesy: Y. Li et al., Chem. Commun., 2011, 47, 9438).

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and high conductivity, resulting from their unique structure. For example, pure CNT electrode could deliver a discharge capacity of about 800 mAh gelectrode−1 at a current density of 0.4 mA cm−2. In addition, CNT sponge has also been investigated as an electrocatalyst and the obtained specific capacity and discharge voltage are 6424 mAh g−1 and 2.45 V respectively at a current density of 0.05 mA cm−2. It is important to mention here that the highest discharge capacity of 34600 mAh gcarbon−1 at a current density of 500 mA gcarbon−1 reported by Chen et al. [124] using multi-walled carbon nanotube paper (synthesized by the floating catalyst method) as air cathode, leaves back an understanding that the presence of large amount of void space in the MWCNTPs are advantageous toward the storage of Li2O2 and the interpenetrating MWCNT networks facilitate facile electron transport. Carbon nanofibers [120] and carbon microfibers [121] have also been investigated as air cathode in NLAS in addition to graphene, mesoporous carbon and CNTs. A combination of CNFs and CNTs delivers an initial discharge capacity of 1500 mAh gcarbon−1 at a current density of 0.2 mA cm−2, whereas, the activated carbon microfiber (ACMF) electrode exhibits a discharge capacity of 4116 mAh gcarbon−1 and a charge voltage of 4.3 V at a current density of 0.025 mA cm−2. Of late, literature has adequate reports on the cathode performance of certain newer carbonaceous architecture such as carbon spheres, carbon nanoballs (CNB) and honeycomb-like carbon (HCC). When carbon spheres (synthesized from modified hydrothermal process) are used as air cathode, a discharge capacity of 4200 mAh gelectrode−1 in the voltage range between 2.35 to 4.35 V vs. Li/Li+ at a current density of 200 mA gelectrode−1 is realized. Nevertheless, only 13 % of the discharge capacity is obtained during charging. Kang et al. [125] reported that CNB (synthesized by solution plasma process) based air cathode can deliver a discharge capacity of 3600 mAh gcarbon−1 at a current density of about 67 mA gcarbon−1. The improved electrochemical performance in this case is associated with the high pore volume and meso–macro structure of the CNB. Apart from this, other carbon materials such as meso–macroporous carbon (synthesized by using colloidal silica as a template), diamond like carbon thin film (prepared a by radio frequency sputtering), hierarchically porous HCC (prepared using the silica templates in different diameters) and micron-sized HCC have also been employed as cathode for lithium–air system [97].

6.5.3.1.3 N-doped Carbon Materials It is generally accepted that the doping of heteroatoms in to a carbon matrix can alter the chemical and electronic nature of the material and

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responsible for the excellent electrochemical performance of the heteroatom doped carbon based electrodes. Based on the theoretical study of Yan et al. [126], doping of nitrogen in to the graphene sheet not only increases the adsorption of oxygen atoms but also decreases the energy barrier of O2 dissociation from 2.39 to 1.20 eV, leading to better catalytic activity toward ORR. Li et al. [127] has demonstrated the use of N-doped graphene nano sheet (GNS) as air cathode for the first time and realized a specific capacity of 11660, 6640 and 3960 mAh gcarbon−1 at 75, 150 and 300  mA gcarbon−1 respectively. In addition, Higgins et al. [128] reported that 20 wt.% N-doped graphene can deliver a specific capacity of 11746 mAh gcarbon−1 at a current density of 70 mA gcarbon−1, which is higher than commercial carbon electrodes (Figure 6.20). There are reports identified in the literature regarding the use of N-doped graphene, N-doped CNTs, carbon fiber and mesoporous carbon as air cathode in NLAS. For example, vertically aligned N-doped coral-like carbon fiber (VA-NCCF) arrays (synthesized by CVD method) when exploited as electrocatalyst in lithium–air system, delivers a stable limited discharge

(a)

(b)

100 nm

Pyridinic-N Graphitic-N

N-GNSs

C

GNSs 404

402

(d)

400 328 326 Binding energy / eV

O

324

N

GNSs N-GNSs

(a) 75 mA g–1 (b) 150 mA g–1 (c) 300 mA g–1

GNSs N-GNSs

3.2

Voltage / V

Intensity / a.u.

3.6

(c)

Pyrrolic-N

Intensity / a.u

100 nm

2.8 2.4 2.0

c

b

a

1.6 800

600

400

200

Bonding Energy / eV

0

0

2000

4000

6000

8000 10000 12000

Specific capacity / mAh g–1

Figure 6.20 TEM images of GNSs (a) and N-doped GNSs (b); XPS spectra of GNSs and N-doped GNSs, the inset is N 1s spectra of two samples (c); voltage profile of GNSs and N-doped GNSs electrodes at various current densities (d) (Courtesy: Y. Li et al., Electrochem. Commun., 2012, 18, 12).

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Nanomaterials for Electrochemical Energy Storage Devices

capacity of 1000 mAh gelectrode−1 at a current density of 500 mA gelectrode−1 with the very low polarization i.e., 0.3 V up to 150 cycles. Such a lower over potential could be attributed to the filling up of entire interbranch space by the special and continuous coating layer of Li2O2, which can enhance the contact between the electrode and Li2O2 compared to isolated individual particles. Further, discharge capacity of 4500 mAh gcarbon−1 with 100  mV high discharge plateau has been achieved with nitrogen-enriched mesoporous carbon (prepared by the hard template method) compared to that of BP2000 cathode (2600 mAh gcarbon−1) at a current density of 30 mA gcarbon−1 [129].

6.5.3.2 Metal and/or Metal Oxides Many researchers focus on the development of novel electrocatalysts including noble metals, metal oxides, carbides, nitrides, sulfides, and their composites due to the instability of the carbon based materials. In particular, metal and metal oxides have been deliberately studied as electrocatalyst in reducing the over potential of lithium–air system, which could be classified as (1) precious metals and/or their oxides; (2) transition metals and/or their oxides; and (3) perovskite and perovskite related oxides.

6.5.3.2.1 Precious Metals and/or Their Oxides Reports are available in the literature toward the electrocatalytic performance of the precious metal and their metal oxides in lithium–air batteries include Pt, Au, Pd, Ru, and Ir. It is generally accepted that the precious metals and their corresponding oxides are considered as the best electrocatalyst due to their excellent catalytic activity toward ORR and OER [91, 130]. The systematic investigation on the electrocatalytic performance of Pt and Au toward ORR and OER reveals that nanoparticles of Au and Pt are effective in ORR and OER respectively [91]. Interestingly, Pt–Au nanoparticles hybrid electrocatalyst supported by Vulcan XC-72 effectively decreases the over potential during charge (900 mV) and discharge (150– 360 mV), which is higher than that of pure carbon catalyst. Further, Pt–Au alloy as bifunctional electrocatalyst reduces the overpotential significantly and improves the round trip efficiency from 57% to 73% (Figure 6.10). Material loading and the structural characteristics including the size of the catalyst, composition, degree of alloying and phase segregation are found to have significant effects on the catalytic performance. For example, increasing the loading of Pt/C and changing the particle size of Au

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333

nanoparticles could improve the discharge capacity, increased discharge voltage and decreased charge voltage. In addition, fully and partially alloyed PtAu/C exhibit excellent electrochemical performance compared to that of catalyst with phase segregation. Further, Pt4Co/C, Pt2Co/C, PtCo/C and PtCo2/C (synthesized through chemical reduction method) also could be exploited as air cathode for NLAS. Various carbon supported metal and metal alloys (Pt, Pd, Ir, Ru, Pt–Pd, Pd–Ir and Pt–Ru) have also been prepared (with an impregnation reduction method) and tested as air cathode in NLAS by Ko et al. Among the chosen electrocatalysts, Ru catalyst exhibits enhanced electrochemical performance in terms of highest capacity and the lowest charge–discharge overpotential [130]. Apart from the precious metals and their alloys, their corresponding oxides have also been studied as electrocatalyst in NLAS. RuO2 and Ir2O3 exhibit an initial discharge capacity value of 317 and 345 mAh gelectrode−1 respectively, at a current density of 0.025 mA cm−2 [131].

6.5.3.2.2 Transition Metals and/or Their Oxides In addition to the precious metals and their oxides, transition metals and their corresponding oxides could also be deployed as electrocatalyst for lithium–air system to realize improved electrocatalytic performance toward ORR and OER. Among the manganese based oxides, electrolytic manganese dioxide, both α-MnO2 and β-MnO2 in bulk and nanowire forms, γ-MnO2, λ-MnO2, Mn2O3 and Mn3O4 have been reported as electrocatalyst in the literature [132]. α-MnO2 delivers the highest specific capacity of 3000 mAh gcarbon−1 at a current density of 70 mA gcarbon−1 with a discharge voltage around 2.6 V and a charge voltage at 4.0 V vs. Li/Li+. Other metal oxides including Co3O4, NiO, Fe2O3, CuO, V2O5, MoO3 and Y2O3 as air cathode catalysts for Li–air batteries have been studied and the results are tabulated (Table 6.5 and Figure 6.21). All the selected catalysts exhibit improved electrochemical performance after five cycles. There are reports available for the electrocatalytic performance of CuO, V2O5 and Co3O4 based oxides in NLAS. In addition, NiCo2O4 has been reported as cathode catalyst in lithium–air system. For example, NiCo2O4 nanoflakes exhibit a discharge capacity of 1560 mAh gcarbon−1, much higher ORR onset potential (2.9 V) and much lower overpotential when compared to those of the pure carbon cathode. Interestingly, β-FeOOH when investigated as cathode catalyst delivers a capacity of 7183 mAh gcarbon−1 along with an improved round-trip efficiency of 74.8% compared to those of pure carbon materials (62.5%) [133–139].

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Nanomaterials for Electrochemical Energy Storage Devices

Table 6.5 Comparison of electrochemical behavior of MxOy cathode catalysts. Capacity mAh g−1 Catalyst

1st cycle

5th cycle

Capacity retention per cycle

MnO2

262

653

132

Co3O4

199

304

133

NiO

298

362

134

Fe2O3

264

285

135

CuO

292

658

136

V2O5

216

829

137

MoO3

152

152

138

Y2O3

238

213

139

(b)

(c)

5

500 nm

(d)

2 μm

500 nm

(e)

Co3O4-only cathode

charge

Potential / V vs. Li/Li+

(a)

4

(j)

3 discharge

2

nanosheet

1

(f) 0

500

1000

nanoflower

1500

2000

nanoneedle

2500

3000

Specific Capacity / mAh g–1 Co3O4 600

Specific Capacity / mAh g–1 Co3O4

nanoneedle

500

(g)

2 μm

500 nm

500 nm

(h)

(i)

400 300

nanosheet

200 100

500 nm

500 nm

2 μm

0

nanoflower

(k) 10

20

30

40

Cycle Number

Figure 6.21 SEM micrographs of nanosheet (a–c), nanoneedle (d–f), and nanoflower (g–i) Co3O4; charge–discharge profiles of the carbon-free Co3O4 cathodes at a current density of 20 mA gcatalyst−1 (j); plots of capacity vs. cycle number for the Co3O4-only cathodes (k) (Courtesy: A. Riaz et al., Chem. Commun., 2013, 49, 5984).

50

Li–Air: Current Scenario and Its Future

335

6.5.3.2.3 Perovskite and Perovskite-Related Oxides

1st 2nd 3rd 4th 5th

3.0

2.0 0

4000

6000

8000 10000

Specific capacity / mA h g–1 5.0

(c)

Voltage / V vs. Li/Li+

without PNT-LSM catalyst

1 μm

100 nm

(d)

1st 2nd 3rd 4th 5th

4.0

3.0

2.0 0

(i)

100 nm

Discharge Charge Coulombic efficiency

4000 0

(h)

1

2

3 4 Cycle number

18000

40

5

Discharge Charge Coulombic efficiency

without PNT-LSM 12000

0

120 80

8000 40

4000 0

1

2

3

4

5

0

Cycle number 1050

6 Terminal voltage / V vs. Li/Li+

0.5 μm

8000 10000 2000 4000 6000 Specific capacity / mA h g–1

80 8000

SP electrode PNT-LSM/SP electrode

5

1000

4 3

950

2 1 0

without PNT-LSM catalyst 43 cycles 0

20

40

with PNT-LSM catalyst: 124 cycles 60

80 Cycle number

100

120

140

Specific capacity / mA h g–1

(g)

2000

120

with PNT-LSM 12000

Coulombic efficiency / %

with PNT-LSM catalyst 4.0

Specific capacity / mA h g–1

(f)

5.0

Specific capacity / mA h g–1

(e)

(b)

Voltage / V vs. Li/Li+

(a)

Coulombic efficiency / %

Recently, perovskite based materials are attracting much attention as electrocatalysts in NLAS due to their improved kinetics toward ORR and OER. For example, nano-sized perovskite oxides of g-La0.8Sr0.2MnO3 and s-La0.8Sr0.2MnO3 (synthesized by sol–gel and solid-state reaction method) have been investigated as cathode catalyst in lithium–air system. The observed specific discharge capacity with g-La0.8Sr0.2MnO3 cathode is 1900 mAh gcarbon−1 at a current density of 0.1 mA cm−2, which is higher than that the cathode with s-La0.8Sr0.2MnO3 (1900 mAh gcarbon−1). The surface morphology of the catalysts is believed to play a key role in achieving the enhanced electrochemical performance [140]. Further, perovskite-based porous La0.75Sr0.25MnO3 (PNT–LSM) nanotubes, La0.5Sr0.5CoO2.91 nanowires, nickel-doped lanthanum cobaltite perovskite oxides of LaNixCo1−xO3−δ (x = 0, 0.25, 0.5, 0.75 and 1), LaMn0.6Fe0.4O3 nanoparticles, Ce-incorporated LaFe0.5Mn0.5O3, Ba0.9Co0.5Fe0.4Nb0.1O3 and porous perovskite CaMnO3 have also been reported in the literature as efficient cathode catalyst for NLAS [141] (Figure 6.22). In addition to the above materials, metallic mesoporous pyrochlore, lead ruthenate and the nanocrystalline expanded pyrochlores are also used as electrocatalyst for NLAS. Metallic mesoporous pyrochlore delivers a high reversible capacity of 10400 mAh gcarbon−1 at a current density of 70 mA gcarbon−1, which is the highest value achieved among the chosen

900

Figure 6.22 FESEM (a, b) and TEM (c, d) images of PNT–LSM catalyst; cyclic performance, charge/discharge specific capacity and coulombic efficiency of Li–air battery with (e, f) and without (g, h) PNT–LSM catalyst at a current density of 0.025 mA cm−2; voltage of the terminal discharge vs. the cycle number for Li–air battery with and without PNT–LSM catalyst at a current density of 0.15 mA cm2 (i) (Courtesy: J.-J. Xu et al., Angew. Chem., Int. Ed., 2013, 52, 3887).

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Nanomaterials for Electrochemical Energy Storage Devices

cathode catalysts. The obtained high capacity value could be correlated with the engineered properties of the mesoporous pyrochlore including a high fraction of exposed oxygen vacancies, porosity that enables good diffusion to the active sites, and a nanoscale conductive network with metallic conductivity [89]. Even though a lot of reports are available for perovskite based material, the mechanism based on the substitution in cation sublattices is still unclear. Therefore, systematic study is warranted to understand the mechanism in detail and to improve the catalytic performance toward ORR as well as OER and to promote the development of Li–air batteries. In addition to the above mentioned oxide based cathode catalysts, some other oxides have also been reported in the literature. For example, ceria based catalysts (Ce1−xZrxO2 (x = 0–0.5)), cathode electrode with 50 wt.% Ce0.8Zr0.2O2 and 50 wt.% MnO2 and lithium based binary and ternary metal oxides including Li5FeO4 and Li2MnO3LiFeO2 have also been used as cathode catalysts in Li–air batteries.

6.5.3.3 Composite Materials In general, the specific capacity of the NLAS could be determined with the amount of generated discharge product during discharge and the complete conversion of the generated products into its original form during charge. Therefore, cathode catalyst should possess enough surface area to accommodate the generated solid discharge products as well as exhibit excellent bi-functional catalytic activity (toward ORR and OER). But, it is very difficult to obtain all the requirements of good cathode catalyst with the single material. Toward this direction, hybrid electrodes are expected to fulfill the above said requirements for the good cathode catalyst. Usually, metal and metal oxides combined with the carbon based materials are the good choice of candidates, where the carbon based material offers the required space to accommodate the solid products and facilitates the dispersion of catalyst particles on the surface to enhance their catalytic performance. Furthermore, carbon has some intrinsic catalytic activity toward the cathode reactions because of the vacancies and defects [142]. When carbon based composite cathode catalysts are concerned, graphene and carbon nanotubes (SWCNT and MWCNT) are the best choices available. Highly dispersed platinum nanoparticles on the graphene nanosheets exhibit a specific capacity of 4820 mAh gelectrode−1 at a current density of 70 mA gelectrode−1. Similarly, Ru–rGO and RuO2.0.64H2O–rGO (synthesized by depositing metallic ruthenium and hydrated ruthenium oxide respectively on reduced graphene oxide (rGO)) could significantly reduce the

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337

average charge potential of 4.3 V for rGO to about 3.9 and 3.7 V at 500 mA gelectrode−1 under the capacity controlled regimes of 5000 mAh gelectrode−1 [107] (Figure 6.23). In addition, the covalently coupled MnCo2O4–graphene hybrid shows a high discharge voltage at 2.95 V and a very low charge voltage at 3.75 V at a current density of 100 mA gcatalyst−1, which is the best value reported so far using the similar electrolyte condition [143]. α-MnO2 nanorods decorated on graphene sheet exhibit a discharge capacity of 11520 mAh gcarbon−1 at a current density of 200 mA gcarbon−1. Further, immobilized Co3O4 nanofibers onto nonoxidized graphene nanoflakes, Fe2O3 nanocluster-decorated graphene, La0.5Ce0.5Fe0.5Mn0.5O3 and GNSs and Zirconium doped ceria and graphene mixture (ZDC/GNSs) have also shown stable electrochemical performance as catalyst for oxygen

25 Number of particles / #

(b)

(a)

5 nm

1 μm

Ru average particle size : 2.36 nm

15 10 5 0

50 nm

(c)

20

1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 Particle size / nm

(d)

60 Number of particles / #

(e) 5 nm

1 μm

Voltage / V

Voltage / V

10

4.0 3.5 3.0 2.5 2.0 0

1000

2000 3000 4000 Capacity / mAh g–1

(h)

4.0 3.5 3.0 5th 20th

2.0 5000

Ru-rGO hybrid 500 mA/g-10 h

4.5

2.5

catalyst free rGO Ru-rGO hybrid RuO2.0.64H2O-rGO hybrid

1.5

1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 Particle size / nm

5.5

5.0

500 mA/g-10 h

4.5

1.5

20

0

1000

10th 25th

RuO2.0.64H2O-rGO hybrid 500 mA/g-10 h

4.5 4.0 3.5 3.0 2.5

15th 30th

2000 3000 4000 Capacity / mAh g–1

(i)

5.0

Voltage / V

(g)

5.0

RuO2.0.64H2O average particle size : 2.18 nm

30

5.5 5th cycle

(f)

40

0

50 nm

5.5

50

5th 20th

2.0 5000

1.5

0

1000

10th 25th

15th 30th

2000 3000 4000 Capacity / mAh g–1

5000

Figure 6.23 Microstructural analysis of Ru–rGO and RuO2.0.64H2O–rGO hybrids: SEM image and SEM-EDX (inset) of porous Ru–rGO hybrid (a); TEM images of Ru–rGO hybrid (inset: HRTEM image) (b); particle size distribution of Ru–rGO hybrid (c); SEM image and SEM-EDX (inset) of porous RuO2.0.64H2O–rGO hybrid (d); TEM images of RuO2.0.64H2O–rGO hybrid (inset: HRTEM) (e); and particle size distribution of RuO2.0.64H2O–rGO hybrid (f); charge–discharge cycles of Li–air cells using rGO, Ru– rGO hybrid, and RuO2.0.64H2O–rGO hybrid under a limited specific capacity of 5000 mAh g−1 at a current density of 500 mA g−1: voltage profiles of fifth cycle (g) and following cycles of Ru–rGO hybrid (h) and RuO2.0.64H2O–rGO hybrid (i) (Courtesy: H.-G. Jung et al., ACS Nano, 2013, 7, 3532).

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reduction in NLAS. Reduced graphene oxide–polypyrrole composite (rGO–PPy) delivers a discharge capacity of 3358 mAh g−1 at a current density of 0.3 mA cm−2 with low over potential of 1.06 V, which is lower than the pure reduced graphene oxide (rGO), i.e., 1.41 V. As we mentioned earlier, other functional carbon material such as SWCNT and MWCNT can also be used as composite materials. Based on the theoretical studies, platinum doped CNT exhibits high electrocatalytic activity toward ORR compared to those of N-doped CNTs and Pt-adsorbed CNTs. Further, increase in the concentration of Pt doping can improve the catalytic activity toward ORR. In addition, Pd coated CNTs sponge, Ru/MWCNTs, RuO2/MWCNTs, RuO2@CNT, NiO/MWCNTs, MnO2/MWCNTs, Co3O4/CNTs, α-MnO2–CNTs–CNFs and nanofibrous MnNi–CNFs have also been reported as efficient composite cathode catalyst and the performance is higher than that of the respective carbon based materials. Interestingly, other composite cathode catalysts including α-MnO2– ACM, CoO–CMK-3, carbon-sphere/Co3O4, Co3O4–RuO2–carbon sphere nanocomposite, porous carbon supported core–shelled Fe3O4–Fe nanocomposite and 3D NiCo2O4 nanowire array/carbon cloth have demonstrated as cathode catalyst for NLAS. Materials hybridized by metals and metal oxides and the compounds with transition metals coordinated to heterocyclic N (such as Fe/N/C, Co/N/C etc.) are also falling under the category of composite cathode catalyst group and exhibiting better catalytic activity toward both ORR and OER. Similarly, composites consisting of nitrogen-doped carbon materials and molybdenum nitride or Mn oxides, MnO2 nanotubes/nitrogen-doped exfoliated graphene and molybdenum nitride/N-doped carbon nanospheres have also been demonstrated as cathode catalysts in non-aqueous Li–air batteries.

6.5.3.4 Other Cathode Materials In addition to the above discussed cathode catalysts including carbon materials, metals/metal oxides and composite materials, some other materials have also been reported as cathode materials for non-aqueous Li–air batteries. For instance, sulfur-doped graphene when used as cathode catalyst, exhibits a specific capacity of 4100 mAh gcatalyst−1, which is higher than that of pristine graphene electrode at a current density of 75 mA gcatalyst−1. Fluorinated CNT also exhibits excellent electrochemical performance over the pristine CNT. Further, conjugated heterocyclic

Li–Air: Current Scenario and Its Future (a)

(b)

(c) 4.5 GPPy

X20000

1 μm

0.1 mA cm–2

AB

GPPy

4.0

Voltage/ V

TPPy

339

TPPy 3.5 3.0 2.5

X2000

2.0

100 nm

10 μm

0

300 600 900 1200 1500 1800 2100 2400

Specific capacity/ mAhg–1

(e) 4.5

0.5 mA cm–2

AB

4.0

Voltage/ V

Voltage/ V

4.5

0.1 mA cm–2 in Air

4.0

(f)

3.5 3.0 2.5

GPPy TPPy

3.5 3.0 2.5

2.0

AB

0

20

40

GPPy

60

80

TPPy

2.0

100 120 140

Specific capacity/ mAhg–1

0

300

600

900 1200 1500 1800

Specific capacity/ mAhg–1

Specific capacity/ mAhg–1

(d)

2500 0.1 mA cm–2 TPPy

2000 1500

0.5 mA cm–2 TPPy

1000 0.1 mA cm–2 AB

500 0

0.5 mA cm–2 AB

1

2

3

4

5

6

Cycle number

Figure 6.24 SEM images of TPPy (a) and GPPy (b); charge/discharge curves of AB, GPPy, and TPPy supported Li–air battery at 0.1 mA cm−2 in oxygen (c), at 0.1 mA cm−2 in argon (d), and at 0.5 mA cm−2 in oxygen (e); cycling performance of AB, GPPy, and TPPy supported Li–air battery at current densities of 0.1 and 0.5 mA cm−2, respectively, (f) (Courtesy: Y. Cui et al., Energy Environ. Sci., 2012, 5, 7893).

conductive polymers (Polypyrrole (PPy), poly(3,4-ethylenedioxythiophene) (PEDOT) and polyaniline (PANi)) have also been investigated as cathode in NLAS due to its unique metallic/semiconductor characteristics [144] (Figure 6.24).

6.5.4 Electrolytes Deployed in Non-Aqueous Li–Air Cells 6.5.4.1 Alkyl Carbonates The NLAS is initially started with the alkyl carbonates as electrolyte due to its significant characteristics such as low vapor pressure and lower viscosity. Unfortunately, the carbonate based solvents are not stable in the presence of superoxide radicals (formed during discharge process), which leads to the electrolyte decomposition by forming high molecular weight side products (called as destroyer of lithium–air battery performance). The formed side products including lithium alkyl carbonates, C3H6(OCO2Li)2, Li2CO3, HCO2Li, and CH3CO2Li are identified through the combination of XRD, FTIR and 1H NMR spectroscopy [145].

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6.5.4.2 Esters Esters have been demonstrated as electrolyte for NLAS due to their high dielectric constant and low viscosity. It is evident from the literature that the common esters such as g-butyrolactone g-butyrolactone and methyl formate are too reactive toward lithium metal due to their high vapor pressure. Further, aliphatic esters and lactones are not stable in the presence of superoxide radicals (proven through computational approach) similar to those of carbonate based electrolytes. In general, common organic esters are unstable in the presence of superoxide and could not be used in the Li–air battery [145].

6.5.4.3 Ethers Ethers and glymes are the preferred choice of stable electrolytes for non-aqueous Li–O2 system. In particular, tetraethylene glycol dimethyl ether (TEGME) is receiving tremendous attention as a popular Li–air electrolyte due to its low vapor pressure. In addition, unlike carbonate based electrolytes, these electrolytes are more stable in the presence of superoxide radicals (less susceptible against nucleophilic attack of superoxide radicals) along with the provision of higher oxidation potential of 4.5 V vs. Li/Li+. It is worth considering fluorination at the β-carbon position to improve the autoxidation stability of ethers in order to suppress C–H bond cleavage by O2, as is indicated by theoretical calculations [145].

6.5.4.4 Nitriles Mechanism of ORR has been thoroughly investigated using acetonitrile as electrolyte (containing 0.1 M n-Bu4NClO4) several years ago, which indicates the possibility of using acetonitrile as electrolyte for NLAS system. Toward this direction, the ORR and OER mechanism of the following type is involved, as understood from the in situ SERS study. The reduction of O2 in the presence of Li+ ions first forms O2 to produce the unstable LiO2, which further disproportionates to Li2O2 and O2. During charging, Li2O2 is directly decomposes in a single step to Li+ + e− + O2 and does not form LiO2. In addition, higher molecular weight aliphatic nitriles, di-nitriles, and nitriles with functionalized ether groups have also been investigated by both computational and experimental techniques. Further, trimethylacetonitrile (TMA) has been demonstrated as a relatively stable electrolyte out of the aliphatic nitrile-based solvents. Even though reports are available

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regarding nitriles as electrolytes, the mechanism and stability (other than acetonitrile) of the same are yet to be understood in the presence of superoxide ion [145].

6.5.4.5 Amides Dimethyl formamide, N,N-dimethylacetamide and N-methyl-2-pyrrolidone have been tested as electrolyte for NLAS. The main drawback of the amides in using as electrolyte for NLAS is the formation of unstable SEI on the surface of the lithium metal anode (stable SEI film is critical in obtaining stable and reversible electrochemical performance in NLAS) due to the high reactivity of N,N-dialkyl amides with lithium based negative electrodes, leading to the rapid solvent decomposition. Further, reaction between the electrolyte and the lithium metal electrode results in the formation of soluble decomposition products that are oxidized at the cathode surface upon charging. Apart from this, formation of side products (Li2CO3, HCO2Li and CH3CO2Li) are also observed upon cycling in the cathode side. The above mentioned issues are critically addressed by changing/adding different lithium salts and fluorinating the amides, resulting in improved electrolyte/lithium interface [145].

6.5.4.6 DMSO Similar to the amides, DMSO also could react with lithium metal anode resulting with the poor lithium cycling efficiency along with the decomposition of DMSO itself. Apart from this, formation of side products like DMSO2, Li2SO3 and Li2SO4 have been identified through the electrochemical quartz crystal microbalance (EQCM) measurements when carbon cathodes are used. Interestingly, reversible electrochemical performance could be obtained with DMSO electrolyte by using Au electrode [145].

6.5.4.7 Sulfones Based on the Quantum-chemical calculation results regarding the stability of sulfones toward the attack of superoxides, sulfones are getting attention as electrolyte in NLAS. But, the low melting point of the sulfones reduces its fame in considering as promising electrolyte, leading to a search in low-melting sulfones as electrolytes for NLAS [145].

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6.5.4.8 Ionic Liquids In recent years, ionic liquids are getting considerable attention as electrolyte in NLAS due to appealing features such as negligible vapor pressure, low flammability, high ionic conductivity and superior hydrophobicity, wide electrochemical potential window and less prone to O2− attack (alkyl groups bounded to the N atom are considered as poor leaving group). There are so many reports available in the literature on the electrochemical performance of ionic liquid based electrolyte in NLAS [145]. Oxygen reduction reaction mechanism in ionic liquid is investigated using two different ionic liquids, 1-ethyl-3 methyl imidazolium bis-(trifluoromethylsulfonyl) imide (EMITFSI) and 1-methyl-1-butylpyrrolidinium bis(trifluoromethanesulfonyl)imide (PYR14TFSI). Cyclic voltammetry results reveal that the formation of Li2O2 occurs through the chemical decomposition reaction of LiO2 and not from the electrochemical reduction of the same. This mechanism can be explained by the hard soft acid base (HSAB) theory of Pearson. Li+ is a hard Lewis acid in ionic liquids while O2− and O22 are soft and hard Lewis base respectively. Thus the chemical decomposition of the unstable LiO2 to form Li2O2, i.e., hard (Li+)–hard (O22 ), is favorable [145]. It is important to mention here that the pyrrolidinium and piperidinium IL-family based electrolytes are more stable against peroxide radical attack compared to that of imidazolium based ionic liquids [145, 146]. ORR

O2 + Li+ + e− 2LiO2

LiO2

Li2O2 + O2

PYR

O2 e

PYR

O2

PYR

O2

OER

LiO2 Li2O2

PYR

O2 e

O2 + Li+ + e− O2 + 2Li+ + 2e−

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Even though, ionic liquids are showing promise as electrolyte, there are so many drawbacks such as poor Li+ solubility and relatively low conductivity, low Li+ transfer numbers (0.12108) which are to be addressed before their commercial exploitation.

6.5.5 Morphology of the Deposited Products As explained earlier, the generated solid discharge product, i.e., Li2O2 can limit the electrochemical performance of the NLAS due to its insulating property (calculated band gap is 4.9 eV). In other words, the cathode reaction kinetics intimately connected with the electronic or ionic transport property of the generated bulk Li2O2 (particle size or film thickness). Therefore, the morphology and the properties of the generated Li2O2 will play a key role in obtaining excellent electrochemical performance in terms of charge–discharge polarization, discharge capacity, round-trip efficiency, rate capability and cyclability. Hence, understanding the insulative property and morphology of the Li2O2 could help in achieving the reversible electrochemical performance [64]. Based on the theoretical/mathematical and experimental measurements, it is believed that the electrical conductivity of Li2O2 could be improved by addressing the insulative properties of the same. For instance, half-metallic Li2O2 surface can mitigate the electrical passivation through the growth of Li2O2, thus responsible for the enhanced reversibility of Li2O2 during recycling of the battery. In addition, forming the polarons (both hole and electron) on the surface of the Li2O2 is expected to provide some electrical conductivity to Li2O2 due to the much higher mobility. Further, it is found that the formation of hole polarons on the surface of the Li2O2 can reduce the migration barriers. It is understood from the theoretical predictions that bulk Li2O2 has a large band gap and some unpaired spin states found on the surface of the Li2O2 could play a critical role in determining the electrical conductivity of Li2O2. It is important to mention here that the formation of superoxide-like structure of lithium peroxide clusters (due to the fact that high spin states are more stable in high stoichiometric lithium peroxide) could improve the conductivity mechanisms on the surface of Li2O2. The surface orientation of the product also plays a key role in OER mechanism. For example, it has been proven through first principle calculations that high energy surfaces are more catalytically active compared to those of abundant surfaces such as (1120) and (0001). Coming to the morphology of the generated discharge product and its influence on the electrochemical performance in NLAS, there are many

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morphologies reported in the literature including toroidal shaped [89], spherical particles [147], elongated particles [148], close-packed nanosheets [149], rough thin films (thick layer) [97], and porous ball-like morphology [150]. It is widely accepted that the mechanism of the formation of Li2O2 is as follows; the nucleation of Li2O2 starts during the initial period of discharge and grows during the discharge process. The morphology of the Li2O2 depends on the applied current density of the system. Lithium peroxide particles with several different morphologies could be observed on the surface of the discharged cathode, as shown in Figures 6.25 and 6.26. Even though many reports have been demonstrated the formation mechanism on different morphology of the Li2O2 and its influence on the electrochemical performance, a consistent understanding, particularly of their effect on OER kinetics and pathways during the recharge process, has not yet been reached. Hence, the fundamental understanding about the above mentioned issue should be required to identify an ideal reversible NLAS.

(a)

2 μm

(b) 200 nm

200 μm O2 side 40/500

(c)

(d)

500 nm

1 μm (e)

1 μm (f)

1 μm 1 μm

Figure 6.25 SEM images of various morphologies of lithium peroxide in discharged cathode. Toroidal-shaped (a), spherical particles (b), elongated particles (c), close-packed nanosheets (d), rough thin films (e), and porous ball-like (f) (Courtesy: Z. Ma et al., Energy Environ. Sci., 2015, 8, 2144).

Li–Air: Current Scenario and Its Future (a)

Pristine

(b)

Pristine

(c)

345

150 mA h g–1

200 nm 10 μm

500 nm

(d)

(e)

250 mA h g–1

500 nm

(f)

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(g)

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(h)

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(j)

(i)

500 mA h g–1

500 nm

500 nm

800 mA h g–1

500 nm

600 mA h g–1

(k) 1000 mA h g–1

(l) 1300 mA h g–1

500 nm

500 nm

Figure 6.26 Morphologic growth process of discharge products at different discharge capacities of pristine (a, b), 150 (c), 200 (d), 250 (e), 300 (f), 400 (g), 500 (h), 600 (i), 800 (j), 1000 (k), and 1300 mAh gcarbon−1 (l) (Courtesy: D. Zhai et al., J. Am. Chem. Soc., 2013, 135, 15364).

6.6 Aqueous Lithium–Air System The major advantage of aqueous lithium–air system is that the product formed during discharge reaction is soluble in water and will not passivate on the air electrode surface like non-aqueous lithium–air system [151, 152] (Figure 6.27). The reaction mechanism is as follows: Positive electrode

O2+2H2O + 4e− Negative electrode

Li

Li+ + e−

4OH−

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e–

e– Discharge O2

Li Lithium Metal

Aprotic Electrolyte

O2 Porous Air Electrode

Figure 6.27 Schematic diagram of aqueous lithium–air battery (Courtesy: Z. Ma et al., Energy Environ. Sci., 2015, 8, 2144).

Overall reaction

2 Li

1 O2 H 2O 2

2 LiOH

6.6.1 Approaches for the Formation of Water Stable Lithium Metal In general, the continuous and reversible stripping/deposition of lithium for prolonged cycles in aqueous electrolyte is the key factor in realizing the practical applications of aqueous lithium–air system. Toward this direction, solid electrolyte protection on the lithium metal has been proposed as an ideal approach to avoid the direct contact between the lithium metal and aqueous electrolyte. Table 6.6 summarizes the possible solid electrolytes for aqueous lithium air system [153–165].

6.6.1.1 Solid Electrolyte The solid electrolyte is the key in enabling the aqueous lithium air system. The important requirement of the solid electrolyte is a thin and mechanically robust membrane with high lithium-ion conductivity at the operating temperature. Furthermore, they should be chemically stable in contact with aqueous electrolytes of various pH values and lithium metal. Finally,

Type

glass ceramics

crystalline

crystalline

crystalline

crystalline

crystalline

Single crystal

Name

NASICON

NASICON

NASICON

Garnet

Perovskite

LISICON

Si wafer

yes yes

no no no

bulk 1.4 × 10−4 1 × 10−3

bulk 1.5 × 10−3 1 × 10−6 6 × 10−7

Li1.15Y0.15Zr1.85(PO4)3

Si

Li14ZnGe4O16

Li3xLa(2/3)−xZr(1/3)−2xTiO3

Li7−xLa3Zr2−xTaxO12

no

no

Stability with the lithium metal

bulk 3 × 10−3

1.3 × 10

−3

Ionic conductivity (S cm−1, RT)

Li1.3Al0.3Ti1.7 (PO4)3

Li2O–Al2O3–TiO2– P2O5

Typical composition

Table 6.6 Summary of possible solid electrolytes for hybrid and aqueous Li–air batteries.

stable in air, acids

not stable in air

stable in air

stable in air, LiCl saturated water

stable in air

stable in air

stable in air, mild acids and bases

Chemical stability

[163, 164]

[161,162]

[159, 160]

[157, 158]

[156]

[155]

[153, 154]

Ref.

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considering the industrial developments, solid-state lithium-ion conductors must be environmentally benign, non-toxic, low-cost materials and easy to fabricate [161, 165]. In recent years, NASICON type class ceramics (Li1+x+yTi2−xAlxP3−ySyO12 (LTAP) produced by Ohara Inc., Japan) have been considered as the only acceptable/promising water stable lithium ion conducting solid electrolyte due its high lithium ion conductivity, high chemical and thermal stability and excellent mechanical strength. Visco et al. [167] have proposed the first the use of composite lithium metal with a three layer protection as anode in aqueous lithium air system in 2004 with LTAP. In other words, the above mentioned composite lithium electrode consists of a lithium metal, a polymer electrolyte and a water stable lithium conducting glass ceramics, Li1+x+yTi2−xAlxP3−ySyO12 (LTAP). In addition, Li2O–Al2O3–SiO2– P2O5–TiO2–GeO2 and Li2O–Al2O3–SiO2–P2O5–TiO2 are the commercially available solid electrolytes with a conductivity of 1 × 10−4 S cm−1 and 2.5 to 4 × 10−4 S cm−1 at 25 °C, respectively [165–167]. It is worth mentioning here that the former has better mechanical strength than the latter. It is important to mention here that the lithium ion conductivity in the solid electrolyte is playing a vital role in improving the electrode performance of the aqueous lithium–air battery. Toward this direction, NASICON provides a three dimensional and open network of sites due to its corner-sharing PO4 tetrahedra and BO6 octahedra, thus responsible for the conduction pathways for various A cations. It is proposed that the substitution of trivalent cations (Al, Cr, Ga, Fe, Sc, In, Lu, Y, or La) in the octahedral sites can significantly enhance the conductivity of lithium ions. For instance, the increasing trend of the conductivity with increase in concentration of M3+ ions has been reported by Aono et al. In other words, they reported higher conductivity for all the systems when the concentration of x is 0.3 for Li1+xMxTi2−x(PO4)3 (M = Al, Sc, Y, and La) system. They concluded that the increase in the lithium ion conductivity could be correlated with the increase in the concentration of lithium ion and density of the pellet. Further, the nominal composition of Li1.3Al0.3Ti21.7(PO4)3 exhibits the highest lithium-ion conductivity (bulk conductivity 3 × 10−3 S cm−1 at 298 K) among all the above chosen systems. The partial replacement of Ti4+ with Al3+or Ga3 in LiTi2(PO4)3 system exhibits a maximum conductivity of 1.3 × 10−3 S cm−1 and 9 × 10−4 S cm−1 respectively. In addition, fast lithium ion conduction has been achieved by the heat treatment of glasses with the formula Li2O–M2O3–TiO2–P2O5 (M = Al and Ga) in class ceramics. Further, substitution of Ca2+ or Y3+ for Zr4+ in LiZr2(PO4)3 (an another system of NASICON family) can transform the structure to rhombohedral

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NASICON form, which exhibits a bulk lithium ion conductivity of above 1 × 10−4 S cm−1, which is interesting [165–167]. Li-rich garnet-type metal oxides Li5La3M2O12 (M = Nb, Ta) receive considerable attention in the race of solid electrolytes especially after the successful demonstration by Weppener’s group. In other words, facile lithium ion conduction has been demonstrated in Li7La3Zr2O12 with a high lithium ion conductivity of > 10−4 S cm−1 at 25°C. The bulk and grain boundary resistances are on the same order of magnitude, which is beneficial for ceramic solid electrolytes with a polycrystalline structure. Li6.75La3Zr1.75Nb0.25O12 and Li6.4La3Zr1.4Ta0.6O12 have also been demonstrated as solid electrolytes with the lithium ion conductivity of 8×10−4 and 1×10−3 S cm−1 respectively and the latter is considered as the highest ever lithium ion conductivity reported so far in the literature [165]. Similarly, lithium ion conductivity (1 × 10−3 S cm−1) value closer to the polymer or liquid electrolyte for solid electrolyte has been reported with the perovskite type (ABO3) based lithium lanthanum titanate (Li3xLa(2/3)−x(1/3)−2xTiO3 (0 < x < 0.16)), especially when x = 0.1. The increase in lithium ion conductivity could be attributed to the large concentration of A site vacancies, thus responsible for the movement of lithium ions between the A sites. High grain boundary resistance is the critical issue toward its usage of polycrystalline solid-electrolyte membrane regardless of its excellent bulk lithium ionic conductivity [165]. Even though the lithium super ionic conductor (LISICON, Li14ZnGe4O16. Li14ZnGe4O16 belongs to the solid solution Li2+2xZn1−xGeO4, which contains interstitial lithium ions with the range of –0.36 ≤ x ≤ + 0.87) has been considered as the choice of solid electrolyte due to its very high lithium ionic conductivity of 0.125 S cm−1 at 300°C, the low lithium ion conductivity at room temperature (1 × 10−6 S cm−1) hinders its application as a polycrystalline solid-electrolyte membrane. Further, single-crystal Si wafer has also been used as solid electrolyte due to its smaller diffusion coefficient with respect to O2 compared to that of lithium, thus prevents the gradual corrosion of lithium metal with O2. In short, the direct contact between lithium metal anode and the aqueous electrolyte has been easily avoided by the coating of water stable lithium conducting glass ceramic protective layer, which covers and isolates the lithium metal from the same. Even though LTAP is protecting the lithium metal from the aqueous electrolyte, it is quite unstable with the direct contact of lithium metal. Hence, the use buffer layer in between the lithium metal and the LTAP plate is essential in order to achieve the excellent electrochemical performance with the lithium metal as anode.

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6.6.1.2 Stability of Solid Electrolyte—Why Do We Need Buffer Layer? As mentioned earlier, NASICON type LTAP is not stable with the lithium metal due to the facile reduction of Ti4+ and cracks are also identified in the solid electrolyte. In addition, LTAP is not stable in strong acidic or alkaline medium, which have been used as catholyte in aqueous lithium–air system. For example, Manthiram et al. [165] have reported the color change in the LTAP solid electrolyte due to the contact of lithium dendrites. ABO3 type lithium lanthanum titanate (LLTO) has also been considered as an alternative choice of the solid electrolyte. The reduction of Ti4+ in LLTO with lithium metal is taking place like LTAP solid electrolyte. Further, LLTO is not stable in the low or high pH values of the electrolyte and relatively stable only in the water environment (pH = 7). LISICON type solid electrolytes are also highly reactive with lithium metal. In general, the exchange of Li+/ H+ is not favored when the lithium ions form part of the tetrahedral framework. At the same time, they easily undergo facile Li+/H+ in weak acetic acid solution when the Li+ ions are in both the framework (tetrahedral) and non-framework sites. Due to the exchange of lithium ions in the solid electrolyte with the proton from the water makes it difficult to use these kinds of solid electrolytes in aqueous Li–air system [165]. Interestingly, the need of buffer layer is eliminated with the Garnet type lithium ion conductive solid electrolytes, which is favorable for reducing the internal resistance of the system. Unfortunately, these type of lithium ion conductors are not stable with water and readily undergo Li+/H+ exchange mechanism, which hinders the use of the same as lithium ion solid electrolyte [165]. Even though single crystal Si wafer demonstrates its stability in acidic/ alkali/neutral media, lot of side reactions also occur when it is in contact with the lithium metal. Further, the crystalline nature of the Si wafer is significantly affected by the high accumulation of lithium in the silicon lattice, which is responsible for the expansion of lattice. In addition, the lithium ion conductivity value of Si wafer should be enhanced to realize the improved electrochemical performance [165]. Therefore, the stability of the solid electrolytes with lithium metal as well as aqueous electrolyte is the key issue in achieving long term and reversible cycling performance. Most of the literature reports are concentrating only on the lithium ion conductivity in the solid electrolyte rather than the stability of the same. In general, the basic requirement of the solid electrolyte is that they should be stable in aqueous solutions of various pH values. This is difficult because most oxides tend to dissolve in strongly acidic or

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alkaline solutions. In addition, the solid electrolyte is not ideally stable in contact with lithium metal. Hence, the introduction of interlayer (buffer layer) between the lithium metal and solid electrolyte is getting importance in order to increase the life span of the aqueous lithium air system. But at the same time, the introduction of buffer layer leads to a rise in the internal resistance, which could significantly affect the cell performance. In addition, the interlayer (buffer layer) should also be stable with the Li metal and should show high lithium-ion conductivity in order to realize the improved electrochemical performance.

6.6.1.3 Buffer Layer In general, the interfacial resistance between the polymer buffer layer and lithium–metal anode is large at room temperature, leading to huge overall cell resistance. However, the interfacial resistance could be greatly reduced by increasing the operating temperature to 60 °C. Thus, the aqueous Li–air batteries with a polymer–anode electrolyte are usually operated at elevated temperatures to decrease the cell resistance and make the cell operational. Some lithium-ion conductors including lithium nitride Li3N, lithium phosphorous oxynitride and poly ethylene oxide (PEO) with LiN(SO2CF3)2 (LiTFSI) have been proposed as the protective interlayer material (between lithium metal and solid electrolyte). This type of lithium electrode has been shown to be stable in the aqueous media (Figure 6.28). It is important to mention here that the interface resistance between lithium metal and polymer electrolytes in a water stable lithium electrode is the dominant part of the cell resistance and an important factor in initiating lithium dendrite formation [168–171]. Lithium

Oxygen

BaTiO3 Aqueous Electrolyte

Lithium Metal

Passivation Film

PEO 18LiTFSI Electrolyte

LTAP Solid Electrolyte

Figure 6.28 Schematic diagram of the proposed WSLE with PEO18LiTFSI–BaTiO3 buffer layer and LTAP (Courtesy: N. Imanishi et al., J. Power Sources, 2008, 185, 1392).

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Lithium nitride (Li3N) and lithium phosphorous oxynitride (LiPON) have been used as buffer layer due to their high lithium ion conductivity and stability against lithium metal. The stability and the interfacial resistance have been examined in the temperature range from 20–80°C. The impedance of the Li–Al/LiPON/LTAP/LiPON/Li–Al cell has been measured in the temperature range 25–80°C. The total cell resistance is about 8600 cm2 at room temperature and 360 cm2 at 80°C. The analysis of the impedance profiles suggests that the Li–Al/LiPON interface resistance is dominant at lower temperatures. The LTAP plate immersed in water for one month shows only a slight degradation in the conductivity. It is proposed that the LiPON protects the reaction between LTAP and Li–Al alloy [168–171]. Lithium conducting polymer electrolytes are considered as promising candidate for interlayers. Imanishi et al. [168] have proposed PEO based protecting buffer layer between the LTAP and lithium metal, i.e., PEO18LiTFSI. The composite lithium electrode exhibits stability in water, and a high interfacial resistance between the PEO18LiTFSI and the lithium metal, which is the dominant resistance in the total cell. The common EIS spectra of Li/Buffer layer/polymer electrolyte/LTAP/ Lithium salt containing electrolye/Platinum is shown in Figure 6.29. From the figure, it is apparent that the appearance of two semi circles could be correlated with the interfacial resistance, passivation resistance and charge transfer resistance. In other words, small semicircle corresponds to the

200

Initial 7 days 15 days 30 days

Zin / Ω cm2

160 120

Rb

Rf1

Rf2

Rc

CPE1

CPE2

CPE3

Zw

80 40 0 0

40

80

120 Zre / Ω cm2

160

200

Figure 6.29 Time dependent impedance behavior of Li/PEO18LiTFSI-10 wt.% BaTiO3/ LTAP/1M aqueous LiCl/Pt, air at 60°C (Courtesy: T. Zhang et al., Chem. Lett. 2011, 40, 668).

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polymer electrolyte and LTAP grain boundary resistances and the large semicircle represents the resistance of a passivation film, the interfacial resistance between the polymer electrolyte and LTAP, and the chargetransfer resistance. Hence, the interfacial resistance is playing a crucial role in realizing the improved electrochemical performance of the lithium metal as anode. Toward this direction, many attempts have been made to decrease the interfacial resistance including the addition of oxide fillers and ionic liquid with buffer layer. Herein, nanoparticle fillers are reported to improve the interfacial compatibility due to 1) a reduced reactivity of lithium–metal with the electrolyte, 2) increased stiffness and compressibility which inhibit dendrite growth and 3) even current distribution upon charge and discharge. For example, Croce et al. reported that the addition of nanosized ceramic fillers, in a PEO based electrolyte, improve the interfacial properties between the lithium electrode and the PEO electrolyte. Further, the interfacial resistance dramatically reduces from 240 cm2 to 125 cm2 by the addition of 10 wt.% nanosized BaTiO3 into the PEO18LiTFSI (Table 6.7). In addition, the total cell resistance of the Li/PEO18LiTFSI–10 wt.% BaTiO3/LTAP/aqueous 1 M LiCl/Pt air is found to be 175 cm2 at 60°C. The multilayer lithium–metal electrode demonstrates stability in an aqueous solution and exhibits a stable charge–discharge performance at a current density of up to 1.0 mA cm−2 [168]. It is important to mention here that the lithium dendrite formation has been observed in earlier days with conventional or organic solvents during repeated charge/discharge. Hence, using an organic electrolyte as a buffer layer is not a great choice and due to the same polymer electrolytes are introduced as buffer layer. Toward this direction, Brissot et al. [172, 173] have thoroughly investigated the mechanism of lithium dendrite formation in polymer electrolyte i.e., Li/polymer electrolyte/Li using in situ observation technique and simultaneous cell potential evaluation. Table shows the results of dendrite formation at the interface of lithium Table 6.7 Li/polymer interfacial resistance (at 60°C) in a Li/polymer/ Li symmetric cell relative to the average size and BET surface area of the corresponding nanofiller. PEO18LiTFSI

BaTiO3

Al2O3

SiO2

Size/nm

100

40

35

50

BET/m2 g−1

5.50

6.68

27.25

67.66

125

38

80

54

Ri/Ω cm−2

240

PP13TFSI

49

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and various electrolytes. From the table, it is clearly understood that the lithium dendrite formation has been greatly suppressed by the addition of ionic liquids. In general, dendrites will grow and penetrate this buffer layer, especially after many cycles. Studies have shown that the dendrite formation could be effectively suppressed by the incorporation of additives in the polymer electrolyte. Toward this direction, Imanishi [168] has reported that the addition of ionic liquids with the polymer electrolyte greatly suppresses the lithium dendrite formation. They found that the conductivity of PEO18LiTFSI-xPP13TFSI increases with x and the lowest interface resistance is achieved with 1.2 ≤ x ≤ 1.44. The lithium dendrite growth is reported to be suppressed by the addition of PP13FSI into PEO18LiTFSI as confirmed by the in situ optical visualization cell. However, high overpotential required for lithium deposition and the stripping reaction are observed at high current densities. The WSLE should be operated at the highest current density possible with low overpotentials to prevent the formation of lithium dendrites. Further, it is proposed that the addition of low-molecular weight plasticizers to the PEO-based electrolyte is expected to improve the interface properties between lithium metal and the polymer electrolyte and enhance the lithium ion transport number of the polymer electrolyte. Toward this direction, low molecular weight oligomer ether has been introduced as a plasticizer to enhance the transport properties of polymer electrolytes. Addition of poly(ethylene glycol) dimethyl ether (PEGDME) into propylene carbonate (PC) with LiClO4 is also reported to reduce the dendrite formation. Moreover, Gibbs activation energy for the charge transfer reaction on the lithium metal surface in PC-PEGDME-LiClO4 decreases with a decrease in the molecular weight of PEGDME in the range of 90–400. In a dedicated study, the interface resistance between lithium and polymer electrolyte and the transport properties of the PEO18LiTFSI-xTEGDME (M = 222.28 g mol−1) composite polymer electrolyte (CPE) are examined as a function of the amount of TEGDME (x). In addition, the electrochemical performance of a Li/PEO18LiTFSI-2TEGDME/LTAP/saturated LiCl aqueous solution/Pt, air cell is evaluated at 60 °C. Further, increase in the lithium transport number and decrease in the interfacial resistance have been achieved by the addition of PEGDME to polymer electrolyte PEO18LiTFSI, wherein the mean molecular weight of PEGDME is 500. Furthermore, the enhancement in oxidative stability of glyme molecules is achieved by the complex formation with LiTFSI in a molar ratio of 1:1 (Table 6.8). Interestingly, Jung et al. demonstrated a Li/tetraethylene glycol

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Table 6.8 Dendrite formation onset time and the capacity of lithium electrode with a 10 μm thick copper current collector.

Electrolyte

Temperature / °C

Current density / mA cm−2

Onset time /h

Capacity mAh g−1

PEO18LiTFSI

60

0.5

15

688

0.1

125

1025

0.5

35

1300

1.0

17

1270

PEO18LiTFSI–1.44 PP13TFSI

60

EC-DMC-EMCLiPF6

15

1.0

0.2

22

PAN-PC-EC-LiPF6 RT

1.0

1.0

100

dimethyl ether (TEGDME)-LiCF3SO3/carbon, O2 cell capable of operating over many cycles with capacity and rate values as high as 5000 mAh g−1 and 0.5 A g−1, respectively [168–171].

6.6.2 Catholytes 6.6.2.1 Acidic Catholyte In general, strong acids are not good choice of candidates for the catholyte selection since the most successful lithium ion conducting solid electrolyte LTAP is not stable in this medium. For instance, the sharp increase in the internal resistance of the LTAP has been observed after immersion of LTAP in 0.1 M HCl solution for about three weeks, which is caused by the surface corrosion of LTAP, as evidenced by SEM images (Figure 6.30). However, use of catholyte with strong acids could be enabled by the addition of imidazole, which can strongly absorb the protons in water. In other words, the pH of the system is maintained closer to the neutral value due to the trapping of the protons in the catholyte at the initial stage of the battery by imidazole–acid composite. Interestingly, a gradual release in the protons from the imidazole–acid composite occurs when all the protons are consumed upon discharge to maintain the pH of the system. In addition, imidazole is showing a better dissociation constant of 7.0, which is noteworthy. It is evident from the literature that the concentration of the

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×10k

LiCl(aq.)

1.8 μm

Basic 14.0

13.0 12.0 11.0 10.0 LiNO3(aq.)

LiOH

9.0 8.0 7.0 6.0 5.0

H2O

4.0 3.0 2.0 1.0 Acidic

HCI Acidic: Surface Decomposition Basic: High Resistance Phase Li 3PO4

Figure 6.30 SEM images of LTAP immersed in various aqueous solutions at 50°C for 3 weeks (Courtesy: T. Zhang et al., Chem. Lett. 2011, 40, 668).

imidazole should be higher than that of acid (0.01%) in order to realize better electrochemical performance. In particular, as evidenced by SEM result, no change in the structure or morphology has been observed for LTAP immersed in 6 M HCl + 6.06 M imidazole for about 2.5 months compared to that of pristine LTAP, which indicates the stability of the system [165, 174]. Further, it is interesting to mention here that the mechanical recharging of catholyte is also possible in this system. Briefly, the discharged catholyte could be easily distilled to obtain LiCl solid and imidazole + H2O liquid. The electrolysis of LiCl solid will produce lithium metal and Cl2 gas, which can be easily converted to HCl gas by reacting with H2. HCl gas can be bubbled into the recycled imidazole + H2O solution to act as the fresh catholyte (Figure 6.31). Even though imidazole is showing promising results, it is not stable at the high potential. In other words, imidazole is getting oxidized during charge before the water splitting potential. Hence, molecule with the higher oxidation potential (stability window beyond that of water) could be considered as a key role in realizing the highly reversible electrochemical system with strong acid catholyte [165]. Considering the week acid section, acetic acid and phosphoric acid could be considered as probable candidates for catholyte. It is worth mentioning here that the dissociation constant (pKa) of these acids play a crucial role in the selection of the same as catholyte for aqueous–air system. In

Li–Air: Current Scenario and Its Future H N

(b) HCl

N

H N +

pH value of catholyte

(a)

(Cl)–

357

8 pH = 5 4

1.01 3.03

6.06

1 M HCl 3 M HCl 6 M HCl

0 0

4 8 12 Concentration of Imidazole

N H

+H2 Cl2

(d)

HCl

imidazole + H2O

(c)

Electrolysis

Li

Air

LiCl + imidazole + H2O Distillation LiCl solid imidazole + H2O

Figure 6.31 (a) Illustration of the composite formed between imidazole and hydrochloric acid. (b) Calculation of the pH value of hydrochloric acid solution with various amounts of imidazole additive. (c) SEM image of LTAP after immersion in 6 M HCl + 6.06 M imidazole for 2.5 months. (d) The schematic representation of recycled imidazole buffer catholyte for hybrid Li–air batteries (Courtesy: L. Li et al., Electrochem. Commun., 2014, 47, 67).

short, the higher the dissociation constant, higher will be the pH value of the catholyte for a certain concentration. The theoretical capacity and dissociation constant of the acetic and phosphoric acid is 447, 1273 mAh g−1 (considering only one proton) and 4.7, 2.16, respectively. Even though acetic acid is bestowed with a higher pKa value, pH value of the same is not sufficient to qualify as a successful solid electrolyte. Therefore, addition of conjugate base like LiOAc is getting importance in increasing the pH value as well as decreasing the dissociation of acetic acid. Further, the highly volatile nature of acetic acid requires closed cell configuration. As far as acetic acid is concerned, the practical capacity of 250 mAh g−1 is obtained for 15 cycles, which is comparatively lower than the theoretical capacity of 447 mAh g−1 and the same could be correlated with the low utilization rate and highly volatile nature of acetic acid. The addition of conjugate base is required for the phosphoric acid also since the solid electrolyte LTAP is

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getting corroded severely with the phosphoric acid along with the increase in the grain boundary resistance [165]. Toward this direction, better results have been observed in terms of better compatibility and grain boundary resistance with the addition of LiH2PO4 as conjugate base. The improved performance could be attributed to the suppression of dissociation of phosphoric acid LiH2PO4, thus responsible for the increase in the pH value of the catholyte. However, the low solubility of phosphoric acid by LiH2PO4 in water mitigates the usage of all the three protons available in phosphoric acid and to realize the theoretical capacity of 819 mAh g−1. In other words, due to the low solubility they will deposit and clog the air electrode upon discharge. As a result, the polarization of the air electrode will increase, leading to a gradual loss of the battery efficiency.

6.6.2.2 Alkaline Catholyte For alkaline system, LiOH is considered as the most successful supporting electrolyte due to its attractive features like increasing the conductivity of the catholyte through high lithium ion concentration, enhancing the ORR activity of the noble metal catalyst and providing alkaline environment. Despite these advantages, adding high concentration of LiOH to increase the conductivity of the system leads to high alkalinity, which is not desirable. Therefore, addition of another lithium supporting salt along with LiOH is necessary to demonstrate improved electrochemical performance. For example, addition of LiClO4 could reduce the internal resistance of the cell. Apart from LiClO4, LiCl and LiNO3 have also been added as lithium supporting salt in the catholyte solution. In short, added salts should reduce the dissolution of LiOH during discharge and decrease the pH of the system, which is highly favorable toward the stability of solid electrolyte. For instance, LiCl–LiOH–H2O catholyte with saturated LiCl has been reported to reduce the LiOH dissolution due to the high concentration of lithium ions and maintains the pH of the system as 7–9, which is beneficial for LTAP [165]. Despite these advantages, several drawbacks are observed with the addition of high concentration of LiCl. In other words, the high concentration of LiCl affects the solubility of discharge product, i.e., LiOH in the catholyte, thus responsible for the clogging of air electrode upon discharge. In addition, Cl2 evolution may also take place during the high voltage charge process due to the high concentration of LiCl. Finally, the obtained practical energy density is comparably low with respect to the whole electrolyte due the high molarity (11 M) of LiCl in the saturated aqueous solution, which is considered as a dead weight in the catholyte.

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6.6.3 Catalysts for Acidic and Alkaline System In general, noble metal catalysts are the most promising candidates as elecrocatalysts in acidic medium since non-noble metal catalysts are not stable in the acidic catholyte. For example, Pt/C has been reported as the best catalyst in the acidic medium due its better ORR and OER activity. Manthiram et al. [165] reported the use of commercial Pt/C as electrocatalyst with phosphate buffer solution. Even though Pt/C exhibits excellent ORR activity in acidic medium, many critical issues associated with the same are not addressed completely. For example, while high charging voltages, dissolution and migration of Pt nanoparticles, agglomeration of Pt nanoparticles and severe corrosion of its carbon support in the acidic electrolyte are to be sorted out and solved for global acceptance. In this direction, IrO2 has been added as an additional support along with Pt/C. In general, IrO2 demonstrates excellent electro catalytic activity toward ORR in acidic medium, and hence is believed to be useful for the suppression of carbon corrosion of Pt/C catalyst. In other words, carbon corrosion is greatly suppressed by IrO2, as evident from the decrease in polarization, especially upon charging. In short, there are two reasons behind the suppression of carbon corrosion, responsible for the improved and steady state electrochemical performance. The addition of IrO2 reduces the charge overpotential of the system, which in turn reduces the carbon corrosion by keeping the potential of the system below 1.6 V NHE, the potential at which carbon corrosion occurs. Secondly, oxygen bulbs are created over the catalyst layer that will not allow the direct contact between catalyst and water molecule, without which carbon corrosion is easily avoided. In alkaline medium, non-noble metal catalysts such as carbon based materials, transition metal oxides, pervoskite oxides have been used as electrocatalysts since they are stable in alkaline medium [165].

6.6.4 Managing the Precipitation of LiOH.H2O As mentioned earlier, the discharge product (LiOH) of aqueous lithium– air system is highly soluble in water and does not create any kind of film on the positive electrode surface directly like non-aqueous lithium–air system. However, the solubility limit of the discharge product is approximately 5.25 M (at room temperature) when the product of discharge gets precipitated as LiOH.H2O beyond this solubility limit. The theoretical specific energy of the system is 430 Wh kg−1 when cycled within the solubility limit of discharge product [151]. Hence, the effective precipitation of LiOH.H2O is getting tremendous attention to qualify aqueous lithium–air

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system in the race of high specific energy storage system like non-aqueous lithium–air, lithium–sulfur etc. It is important to mention here that the formation and precipitation of LiOH.H2O do not involve any kind of electron transfer, thus responsible for the rise in the impedance and the un even distribution of discharged products, resulting from the restriction of accessible porosity. In general, the uniform distribution of discharge products on the pores is expected to provide homogeneous current distribution through the separator, thus responsible for the highly reversible electrochemical performance with long life span. In other words, the precipitation of products in certain places is due to the gravitational fields, concentration and thermal gradients, current heterogeneity, or electrolyte flow that leads to the non-uniform distribution of current, thereby resulting in the swift rise in the impedance behavior along with the reduction in the accessibility of large regions of the reservoir or electrode. In addition, uneven lithium plating/stripping would also happen due to the non-uniform current distribution during charge and discharge. Further, the non-uniform current distribution leads to the formation of dense layer of discharged product (dense layer of monohydrate) over the protective layer of positive electrode which in turn results in the increased impedance value and the early end of discharge. Hence, this leaves ample scope for the development of novel cell designs as well as the battery management system algorithms. Toward this direction, some of the companies have reported their unique cell designs with the improved reversible capacities, which are discussed in the upcoming section [151]. Poly plus company has established the unique cell design for the primary aqueous lithium–air cell (Figure 6.32), wherein the discharge products are stored in the porous reservoir of the cathode [151, 174, 175]. Porous zirconia felt reservoir has been used that expands with the solid discharge product as it precipitates and fills with a catholyte that contains hygroscopic salts, capable of absorbing water from the atmosphere. Even though the reduction in the mass of the cell (charged state) and prevention of catholyte evaporation are achieved using the above method, the salts are not appropriate for the rechargeable cells. Further, Toyota has demonstrated the reversible precipitation and dissolution of LiOH with a capacity of 200 mAh in 5 M saturated solution in symmetric cell with the use of platinum and Ohara separator [176]. Raman spectra confirms the formation of LiOH and O2 bubbles during charge process (dissolution of LiOH). In addition, they demonstrated another cell set up consisting of LaSrCoO3, carbon, PTFE on carbon paper and the anode

Li–Air: Current Scenario and Its Future

Lithium foil Interlayer Porous reservoir Case with holes for air access

361

Solid electrolyte Air electrode Anode current collector Proprietary seal

Figure 6.32 Cross section of aqueous lithium–air cell of PolyPlus (Courtesy: S. J. Visco et al., 210th Meeting of the Electrochemical Society, Cancun (2006)).

consists of Li metal in a LiTFSA/PC electrolyte with an Ohara separator between the electrodes. The consumption of LiOH during charge process (from a second cell constructed in the discharged state, with solid LiOH in the cathode) has been confirmed. Interestingly, 100 reversible charge/discharge cycles (100 cycles at 0.1 mAh cm−2, 40 cycles at 2 mAh cm−2) have been demonstrated by the EDF Company with their unique cell design (Figure 6.33), wherein issues against the storage of LiOH have been addressed to a significant extent [151, 177]. Reason behind the improved electrochemical performance could be correlated with the introduction of polymer layer on the air cathode side of the LTAP separator, thus responsible for the prevention of nucleation and preferential deposition of a resistive LiOH layer at the LTAP/reservoir interface. Further, the overpotential is found to increase exponentially with time, if the polymer layer is not incorporated. In other words, the hydroxyl ions are produced through the reaction between reduced oxygen species with water, which in turn reach the reservoir through IPN. Herein, IPN mitigates the influx of CO2, which reacts with LiOH to form Li2CO3 and leading to the degradation of the cell. The oxygen evolution electrode is embedded in the reservoir in order to enable recharge of the cell. By using a transparent EDF cell of this type, the formation of LiOH. H2O has been observed. A small amount of the LiOH.H2O (with respect to the total pore volume of reservoir) has been initially charged to 73 mAh cm−2 (for an available cell specific energy of 500 Wh/kg), 40 cycles with 2 mAh cm−2 (13 Wh kg−1) and 100 cycles with 0.1 mAh cm−2. The vertical

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Nanomaterials for Electrochemical Energy Storage Devices

Air

LiOH(aq)

Composite air electrode Oxygen evolution electrode

Air

Li

LiOH(s) Composite Lisicon separator

Figure 6.33 Schematic representation of EDF’s rechargeable aqueous Li/air cell (Courtesy: P. Stevens et al., ECS Transactions, 28, 1 (2010)).

orientation of the current EDF cell is considered as a disadvantage since the LiOH.H2O settles at the bottom of the cell and gases exit on the top of the cell, thus resulting in the non-uniform current distribution. In other words, more current flows on the top region, where the cell does not contain LiOH.H2O. This would lead to a non-uniform thickness of Li metal as a function of height in the cell. An alternative orientation of the cell, i.e., horizontal configuration is felt to be the ideal choice for LiOH.H2O to get the precipitate either on the AEM or the LTAP, resulting in more uniform current distribution. Despite this advantage, there is a difficulty for the oxygen produced at the OER electrode to exit the cell. In addition, incorporation of the OER electrode in the cell sandwich may also be more difficult [177]. Even though Toyato and EDF have demonstrated small reversible specific capacities with a small amount of solid discharge product, high specific energy with good reversibility of LiOH.H2O, remains as a challenge even today. The basic understanding about the precipitation of LiOH.H2O and its influence on the uniform distribution of current in the cell needs to be studied extensively in realizing the high specific energy aqueous lithium

Li–Air: Current Scenario and Its Future

363

air system. In addition, the stability of the porous reservoirs in terms of mechanical and electrochemical stability should be investigated for many cycles. Further, influence of current density, reservoir geometry (porosity and thickness), and cell orientation with respect to gravitational field on the product distribution requires a more comprehensive analysis. It is important to mention here that the morphology, size and porous nature of the discharge product can also influence the electrochemical performance of aqueous lithium air system. In general, during the charge process, precipitated LiOH.H2O should dissolve and LiOH to diffuse to the air cathode. The rate of particle dissolution depends upon the size and morphology of the discharge product. In other words, dissolution of bigger size LiOH particles requires longer time compared to those of smaller ones. Some novel cell designs are also reported for aqueous lithium air system, wherein external reservoir has been used to store the discharge product. In short, storing the products outside the path of ionic current flow would ensure that the pores in the positive electrode and the separator remain free for ionic transport. Further, the external storage of the product could be enabled by the flow-through cell connected with a reservoir. A uniform current density has been maintained by minimizing the LiOH concentration gradients orthogonal to the applied electric field by using sufficient high flow rates, relative to the applied current density. Recently, several promising strategies have been developed by various companies to improve the energy density and life span of the aqueous lithium–air system. In other words, adding desiccants to the discharge product and polymer protection of the reservoir to separate it from the anode protection layer and the air electrode could be correlated with the increase in energy density and durability of the aqueous lithium air system. Further, the cyclability of the system could be increased by the coating of the protective layer with lithium metal anode (water stable lithium metal anode) has also been proposed.

6.6.5 Hybrid Lithium–Air Battery The only difference between hybrid and aqueous Li–air batteries is the usage of different anode electrolyte between the solid electrolyte and lithium metal anode, as shown in Figure 6.34. In the “hybrid” Li–air battery, the anode electrolyte is a polymeric separator soaked in a liquid organic electrolyte, while in the aqueous Li–air battery, the anode electrolyte is a Li+-ion conducting polymer or ceramic. In addition, the cell voltage of the

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Nanomaterials for Electrochemical Energy Storage Devices

e–

e– Discharge O2 O2

Li Lithium Metal

Aprotic Electrolyte

Aqueous Electrolyte

Porous Air Electrode

Li+ conducting membrane

Figure 6.34 Schematic diagram of hybrid lithium–air battery (Courtesy: Z. Ma et al., Energy Environ. Sci., 2015, 8, 2144).

hybrid Li–O2–H2O cell is even higher (3.4 V), leading to a higher energy density of 3804 Wh kg−1. As a result, lot of scope remains on the development of hybrid lithium ion battery, as it is expected to offer better solution to alleviate the issues of aqueous battery in a satisfying manner [165].

6.7 Applications The main objective behind current research on Li/air batteries lies on the development of an electrochemical energy-storage technology that has a comparable power density to Li-ion batteries and a specific energy close to that of gasoline engines. The Tesla Model S, which sports a 400 km driving range with a 85 kW h−1 battery pack costs twice the price of a standard economy car. With a current cost higher than 400 $ kW h−1, electric cars have so far only entered in the high-end market, where users are willing to pay a premium. Resizing the battery pack on the other hand results in a limited driving range (typically 100–150 km) otherwise called “range anxiety” feeling. For these reasons, the need to develop energy-storage technologies enabling at least a 500 km driving range, while retaining the same battery pack volume at an affordable price, is of primary interest which in turn is the main challenge to allow widespread percolation of electric vehicle with the affordable market price.

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Sales of electric vehicles (EVs) nearly doubled in 2013, but never exceeded 100 miles on one charge. To boost the range closer to 300 miles or more, researchers are reporting new progress on a “breathing” battery that has the potential to one day replace the lithium-ion technology of today’s EVs. Imanishi et al. have demonstrated lithium–air system with the practical energy density of more than 300 Wh kg−1, which is far higher than the commercial lithium-ion battery (150 Wh kg−1). By 2022, according to the research firm Bloomberg New Energy Finance, the unsubsidized total cost of ownership of battery electric vehicles (EVs) will fall below that of an internal combustion engine vehicle. From the automotive perspective, Li–air cells could power electric cars for more than 400 miles on a single charge using a battery pack that’s a fifth of the weight of today’s EV batteries. The Li–air battery has a theoretical specific energy (energy per unit mass) of 3.5 kWh kg−1 (kilowatt hours per kilogram). By comparison, Li-ion batteries have only 105 Wh kg−1 (watt hours per kilogram) at the pack level, limiting fully electric cars to about 150 km of driving range. The energy density of gasoline is roughly 13 kWh kg−1, of which 1.7  kWh  kg−1 of energy is provided to the wheels after losses. Accounting for the weight of a full Li–air battery pack (casing, materials, etc.) the energy density will be considerably lower but estimates range up to 1.7 kWh kg−1 to the wheels. For future electrochemical energy storage, it is desirable that the battery of choice should have an energy density in excess of 500 Wh kg−1 and a cycle life over 1000 cycles at a cost lower than $100/kWh. With an energy density of 500 Wh kg−1, a battery pack including necessary ancillaries weighing no more than 300 kg would provide sufficient energy to drive EVs for at least 500 km. With such a vision only, IBM in 2009 first launched the “Battery 500” project, ambitiously aiming to develop Li−air batteries that could ensure a 500 mile (800 km) driving range. Tesla’s current growth leaves hope to realize this target in the near future.

6.8 Future of Lithium–Air Systems Together with the innovations on the air cathode, lithium–metal anode, and solid electrolyte mentioned above, as well as a flow through mode of the catholyte, the ideal aqueous Li–air batteries could realize the ultra-high energy density of the Li–O2 couple for electric vehicles and grid energy storage of electricity produced from solar or wind (Table 6.9).

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Table 6.9 Requirements for durable, high-energy automotive Li/air batteries (N = nonaqueous only, A = aqueous only). Requirement of high energy automotive lithium– air batteries Li anode

Robust and flexible containment, including flexible seals

Li protection layer(s) transport properties

Sufficient conductivity over T range of interest Negligible electronic conductivity

Manufacturability

Sufficiently thin (96% after 500 cycles. Such a high specific capacitance was suggested due to the combination of EDLC as well as Faradaic capacitance. To achieve such high performance, the combination of Ni(OH)2 with carbonaceous materials, and other metal oxides as a hybrid electrode is an efficient approach. Jang and his group [29] fabricated a hybrid electrode system composed of porous Cu nanostructure, carbon layer and Ni(OH)2 on to Ni film as a 3D Ni(OH)2/C/ Cu electrode with open structure (Figure 10.4a, b). The design of the electrode with 3D-porous or open structure provides good electrical conductivity and facilitate electron transport between the Ni(OH)2 as active materials and the current collector of the Ni-plate. The electrode offered high specific capacitance of 1860 F g−1 at 1 A g−1 current density and good cycling performance with 86.5% capacitance retention after 10,000 cycles. The asymmetric device based on the electrode exhibited the promising energy density and power density of 147.9 Wh kg−1 and 37.0 kW kg−1, respectively. They have further fabricated porous CuNi current collector in a similar way, i.e., electrodeposition method followed by deposition of Ni(OH)2 nanostructure. The electrode also exhibited outstanding performance with the specific capacitance of 3,637 F g−1 at a current density of 1 A g−1, with great cycle stability of over 80% after 10,000 cycles [26]. Chen et al. [30] combined Ni(OH)2–MnO2–RGO as hybrid spheres-based asymmetric supercapacitor, which showed satisfying specific capacitance of 1985 F g−1 and energy density of 54.0 Wh·kg−1 with good electrochemical cycling stability. Alshareef and his group [31] fabricated micropseudocapacitor based on Ni(OH)2 via top–down photolithographic process and bottom–up chemical synthesis. The fabrication of microchip was carried out with the help of photolithography on glass or PEN followed by metal deposition by sputtering and Ni(OH)2 deposition via chemical bath deposition to fabricate the electrode of total area 0.15 cm2 with 30 fingers of a typical width of 100 μm, and spacing between the fingers of 50 μm (Figure 10.4c, d). The electrode showed very high rate of redox activity up to 500 V s−1 with cell capacitance of 16 mF cm−2 and a volumetric stack

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Nanomaterials for Electrochemical Energy Storage Devices 1.8

1 A/g 2 A/g 5 A/g 10 A/g 20 A/g

(b) 1.5

(a) Voltage (V)

1.2 (d) Cu

0.6 0.3

C Ni(O

50 um

0.9

H) 2

3um

0.0

0

300

600

900 1200 1500 1800

Time (s) 10-1

(c)

(c)

Ni(OH)2μ-SCs

(d)

This work

NH(OH)2Ni/Pi/Ti

Glass

00um

1um

PEN

Ni( OH

Glass

)2

(e)

Energy density (Wh/cm3)

-2

10

10-3

4V/500 μAh Li thin-film battery

Metal oxide based μ-SCs

10-4

Carbon based μ-SCs

10-5 3V/300 μF Electrolytic capacitor

10-6 10-3

10-3

10-1

103

101

103

103

104

Power density (W/cm3)

Figure 10.4 (a) SEM image of as-synthesized 3D-Ni(OH)2/C/Cu electrode (inset: corresponding HRTEM); (b) Galvanostatic charge/discharge curves of the 3D-Ni(OH)2/C/ Cu//3D-Mn3O4/C/Cu asymmetric supercapacitor as a function of different current densities (Copyright permission taken from Ref. [29]). (c) SEM and schematic representation of Ni(OH)2-based finger electrode and its micropseudocapacitor device, (d) Ragone plot showing the comparison of energy and power density of Li thin-film batteries, electrolytic capacitors, carbon and metal oxide based μ-SCs with respect to Ni(OH)2 μ-pseudocapacitor (inset: red LED powered using μ-pseudocapacitor device) (Copyright permission taken from Ref. [31]).

capacitance of 325 F cm−3. More interestingly, the micropseudocapacitor exhibited maximum energy density of 21 mWh cm−3, which is superior to the Li-based thin film batteries. Cobal hydroxide is another promising metal hydroxide, which can be a used as alternative material owing to the environmentally benign nature and low cost. The unique structural property, tunable surface, controllability of the shape or the size makes it an excellent material electrochemical performance in alkaline solutions (Figure 10.5). At the molecular level, the similarity between pristine and charged phases of Co(OH)2 seems to be ultimately responsible for the observed enhanced performance as reported by Deng et al. [24] The theoretical capacitance value of the

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material is also found to be higher than 3000 F g−1, indicating as good electrode material for supercapacitor application. Deng and his coworkers fabricated different nanoarchitectured Co(OH)2 based electrode with nanotubular and nanoporous structure on Ni foam by electrodeposition of Cu/Ni film and successive etching of Cu [32]. The specific capacitance obtained from these electrodes was as high as 2500 F g-1 for nanotubes and 2900 F g-1 for nanoporous film at a current density of 1 A g−1 in three electrode system. In a symmetric cell arrangement of nanoporous film, remarkably high specific capacitance of 1100 F g−1 at 1 A g−1 current densities was achieved with extraordinarily high energy density of 220  W h kg−1. The device also exhibited extraordinary stability with only 13% loss after 10,000 cycles. The combination of single-layered β-Co(OH)2 with N-doped graphene as positive electrode found to be beneficial while asymmetric device is fabricated as reported by Gao et al. [33] The device exhibited high specific capacitance of 241.9 F g−1 at 1 A g−1 and good stability over 10,000 cycles with capacity retention of 93.2%. The energy density achieved by the device was significant with the value of 98.9 W h kg−1 along with the high power density of 17,981 W kg−1. Wang et al. [34] fabricated the asymmetric device combining flower-like Co(OH)2 as positive electrode and urchin-like vanadium nitride as negative electrode. Such hybrid capacitor exhibited the specific capacitance of 98.5 F g−1 at a current density of 0.5 A g−1 and showed capacity retention of 86% after 4000 cycles. It also delivered a high energy density of 22 W h kg−1 at a large power density of 15.9 kW kg−1. Iron oxide hydroxide (FeOOH) was reportedly studied as supercapacitor or battery-type anode materials due to their high theoretical capacitance, wide operating potential window, low cost and natural abundance. However, dissimilarities in the lattice structure of iron oxides/hydroxides limit the performance and importantly the cycle life [24]. The use of iron oxide as electrode material involves the following processes involving the activation into FeOOH and subsequent reversible charging discharging processes [35]:

Fe2O3 + H2O FeOOH + H2O + e–

2FeOOH (activation)

(10.2)

Fe(OH)2 + OH– (Charging–discharging)

(10.3)

A number of reports are available on the manipulation of morphology and symmetric or asymmetric assemble of material to fabricate device

Nanomaterials for Electrochemical Energy Storage Devices 1000

ar ge Ch

C0(OH)2

Di

sc

H+

ha

rg e

CoOOH c a

b

100

Batteries

10

H

NTA//CONTA

ECs Capacitors

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1e +5

Specific power (W kg–1) (d)

Ag–1

Co(OH)2 Separator

o LD Ni-C

Co(OH)2// Co(OH)2

(c) Current Collector

COSNP//COSNP

(b)

Co (HO ) ref. 2 grap 21.4 hen e 5 C ref. 15,1 o(HO) 7-20 2 ,22,4 3,44 Ni (H O ref. 24 )2 ,34

OH

H2O

(a)

Specific energy (Wh kg–1)

530

18 12 6 0

VN CO(OH)2

-6

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-12

Current Collector

-18 -1.2

V: 1.6 V -0.9 -0.6

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0.3

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Potential (V vs. SCE)

Figure 10.5 (a) Reaction model of Co(OH)2 and CoOOH phase transformation during charging–discharging process. (Reproduced permission from Ref. [24]) (b) Ragone plot comparing the performance of Co(OH)2 nanotube arrays (CONTA), sucker-like nanoporous film (COSNP) with other reported devices. (Reproduced permission from Ref. [32]). (c, d) Co(OH)2//Vanadium nitride (VN) hybrid supercapacitor: Schematic representation (c) and CV diagram of Co(OH)2 and VN in a three-electrode system. (Reproduced permission from Ref. [34]).

(Figure 10.6a, b). The CV diagram of FeOOH-based electrode indicates the materials with the characteristic of negative electrode in the device [35]. The combination of FeOOH as negative electrode and NiMoO4 as positive electrode in the asymmetric assemble delivered remarkable gravimetric and volumetric energy densities of 33.14 Wh kg−1 and 17.24 Wh l−1, respectively with robust cyclic stability over 10,000 cycles. Chen et al. [36] employed electroplating route for the synthesis of γ-FeOOH nanosheets on carbon cloth for supercapacitor application. They fabricated an asymmetric device (MnO2/CC//FeOOH/CC) at a maximum operating cell voltage of 1.85 V, which delivered maximum power density of 16000 W kg−1 at an energy density of 37.4 W h kg−1. Yu and his group [37] developed fluorine doped β-FeOOH nanorods grown on carbon cloth with high conductivity via a simple scalable

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low-temperature hydrothermal synthesis and post-annealing treatment. Such FeOOH-based supercapacitor exhibited supercapacitive property and displayed energy density of 1.85 mW h cm−3 and maximum power density of 11.11 W cm−3 with high-rate capability and long cycle life up to 5000 cycles at different current densities. Copper hydroxides have also attracted attention as another potential metal hydroxide for supercapacitor application due to their chemical stability, low cost, and environmentally friendly nature. Gurav and his co-workers [38] employed room temperature soft chemical synthesis route to grow nano-grained Cu(OH)2 thin films on glass and stainless steel substrates and used as electrode in supercapacitor application. The electrode exhibited moderately high specific capacitance of 120 F g−1 at 10 mV s−1 scan rate. In another report, 3D mesoporous Cu(OH)2 nanorods were grown on copper foam via a facile and cost-effective surface oxidation method and used as a binder-free electrode for supercapacitor application [39]. The asymmetric device combining Cu(OH)2 and activated carbon delivered high volumetric capacitance of 5.091 F cm−3 at 10 mV s−1 and good cyclic stability with capacity retention of 73.4% after 2000 cycles. The device also delivered maximum energy density and power density of 4.152 mW h cm−3 and 383.222 mW cm−3. Lei et al. [40] synthesized Cu(OH)2 nanobelt arrays by Cu-plating of dacron cloth and subsequent ammonia treatment. The electrode exhibited specific capacitance of 217 mF cm−2 at 0.5 mA·cm−2 current density in a three-electrode system and showed 90% retention in capacitance at a current density of 2 mA·cm−2 after 3000 charge/discharge cycles. The fabricated asymmetric device, Cu(OH)2/Cu/Dacron//CNF/ Dacron, delivered high areal capacitance of 195.8 mF cm−2 at 1 mA cm−2 current density (Figure 10.6c, d). It also exhibited energy density of 3.6 x 10−2 mWh·cm−2 at a power density of 0.6 mW cm−2 [40]. Fewer works are also reported based on Mn(OH)2 as supercapacitor. Anandan et al. [41] synthesized octahedral Mn(OH)2 nanoparticles in the size range of 140–200 nm using sonochemical irradiation approach. It showed specific capacitance of 127 F g−1 at current density of 0.5 mA cm−2 in the potential range from −0.1 to 0.8 V in 1 M Na2SO4 solution in a three-electrode system. As per their report, the electrochemical performance of the material originated from the reversible transition of Mn(OH)2 to MnO2 involving the occurrence of different oxidation states from Mn2+ to Mn3+ and then to Mn4+. In another study, the symmetric all-solid-state supercapacitors was fabricated based on Mn(OH)2 as electrodes and LiCl/ PVA gel as solid electrolyte [42]. Such symmetric device exhibited high volumetric capacitance of 39.3 mF cm−3 at the current density 0.3 mA cm−3 with robust cycling stability.

Nanomaterials for Electrochemical Energy Storage Devices

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Fe2O3

Fe(OH)2

Fe(OH)3

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device assembling

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PVA-KOH 100

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0 −5 −10

1st Cycle 5th Cycle 10th Cycle 20th Cycle

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FeOOH

Cu(OH)3/Cu/Dacron

(c)

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120 140 80 60 40 20 0

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Figure 10.6 (a) Models for different Fe-compounds showing the dissimilarities in crystal structure (Gold and red balls stand for Fe and O atoms, respectively) (Reproduced with permission from Ref. [24]). (b) CV diagram of Fe3O4 and its conversion to FeOOH after 1st cycle as negative electrode (Reproduced permission from Ref. [35]). (c) Areal capacitance obtained at different current densities for Cu(OH)2/Cu/Dacron//CNF/Dacron device (inset: schematic representation of the device); (d) Cycling and (e) bending stability of the device obtained from CV measurements at a scan rate of 50 mV s−1 (Reproduced permission from Ref. [40]).

10.3.2

Layered Double Hydroxides (LDHs)

Layered double hydroxides are a class of ionic lamellar compounds composed of positively charged brucite-like layers along with an interlayer region containing charge compensating anions and solvation molecules. In their structure, metal cations occupy the centers of edge sharing octahedra, and vertexes contain hydroxide ions connected to form infinite 2D sheets [43]. In general, layered double hydroxides (LDHs) are the class of two-dimensional anionic clay-like or hydrotalcite-like materials which are composed of positively charged brucite-like host layers and as well as exchangeable charge-balancing interlayer anions which are expressed as [M12 x M 3x (OH)2 ]x (A n )x/n .mH 2O [43–46]. A fraction of divalent metal ions (such as Mg2+, Fe2+, Co2+, Cu2+, Ni2+, or Zn2+) coordinated octahedrally by hydroxyl groups in the brucite like layers are uniformly replaced by trivalent metals (such as Al3+, Cr3+, Ga3+, Mn3+, or Fe3+) with the molar ratio of M3+/(M3+ + M2+), which is normally between 0.2 and 0.4 [43]. In LDHs structures, cations M2+ and M3+ occupy the octahedral holes in a

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brucite-like layer and An are anions located in the hydrated sites between the layers [45]. The layered structure and the flexibility of incorporating mixed valence transition metal ions into the LDH structure in different compositions open the enormous potential to design high-performance catalysts in both molecular and nanometer scales. Moreover, such LDH materials can be converted into the corresponding mixed metal oxides (MMOs) with a uniform M2+ and M3+ distribution and a high surface area after calcination treatment [47]. Having such unique layered structure, open structure and availability of multivalent metal cations, LDHs are considered to be very promising material for supercapacitor application. Over the past few years several efforts have been focused to tailor the physical and chemical properties of LDHs by changing the nature of metal cations, the molar ratio of M2+/M3+, the type of interlayer anions and so on to fulfill specific requirements for practical applications of LDHs in energy storage applications. A wide variety of layered transition-metal (Co, Ni, Zn) hydroxides, including layered double hydroxides (LDHs), and layered rare-earth hydroxides have recently received intense attention for the exploration of hydroxide nanosheets [47]. Yuan et al. [48] fabricated Ni foam supported Zn–Co hydroxide nanoflakes by facile solvothermal approach and used as electrode for supercapacitor application. The electrode exhibited specific capacitance of 901 F g−1 at 5 A g−1 and displayed capacity retention of about 92% after 1000 cycles. Lai et al. [49] synthesized hierarchical structures of a novel composite of nitrogen-doped carbonized bacterial cellulose@nickel-cobalt layered double hydroxide (CBC-N@LDH) via in situ oxidative polymerization, high-temperature carbonization followed by solution co-deposition process. The sample displayed high specific capacitance of 1949.5 F g−1 at current density of 1 A g−1 and retained 54.7% capacitance after 10 A g−1. It also delivered cycling stability of 74.4% retention after 5000 cycles. The asymmetric supercapacitor by assembling CBC-N@LDH as positive and CBC-N nanofibers as negative electrode materials in the potential window of 0–1.6 V delivered energy density of 36.3 W h kg−1 at power density of 800.2 W kg−1. Yang et al. [50] designed three-dimensional concentration gradient Ni–Co hydroxide nanostructures by directly growing on Ni foam using facile stepwise electrochemical deposition method as binder-free electrode for supercapacitor (Figure 10.7). The material showed specific capacitance of 1760 F g−1 at 1 A g−1 and showed rate capability above 62.5% capacitance from 1 to 100 A g−1. The asymmetric device was fabricated by assembling Ni–Co hydroxide as the cathode and activated carbon anode and it delivered a high specific capacitance of 332

Nanomaterials for Electrochemical Energy Storage Devices

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(a)

Ni hydroxide Co hydroxide

Ni foam

Ni hydroxide

x increasing

More stable

Ni1–xCox hydroxides Higher capacitance 6 M KOH

Ketjen black

Concentration gradient nanostructures

Asymmetric supercapacitors 1 A·g–1 2 A·g–1 3 A·g–1 5 A·g–1 10 A·g–1 20 A·g–1 50 A·g–1 100 A·g–1

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Figure 10.7 (a) Schematic illustration indicating the formation mechanism of 3DCGNC directly grown on Ni foam using stepwise electrochemical deposition method and the assembly of asymmetric supercapacitors using as-obtained 3DCGNC, (b) Cross-sectional SEM image of 3DCGNC deposited on flat Ni plate (inset indicates the Ni (blue) and Co (green) EDS composition distribution of the electrode), (c) galvanostatic charge– discharge profiles performed at different current densities of the asymmetric device assembled using 3DCGNC cathode and a Ketjenblack anode and (d) Cycling performance of the asymmetric device at 5 A g−1 current density (inset in (f) galvanostatic charge– discharge curves of the asymmetric device) (reproduced with permission from Ref. [50]).

F g−1 at 1 A g−1 and capacity retention of 97.8% over 20,000 cycles. It also delivered energy density and power density of 26 W h kg−1 and 133 kW kg−1.

10.3.3

Layered Triple Hydroxides (LTHs)

In recent times, very limited works based on triple layered metal hydroxides were reported. It is believed that the presence of three different metal cations may further enhance the pseudocapacitive performance. Jabeen et al. [51] synthesized Ni/Co/Al layered triple hydroxide (LTH) on Ni foam@brominated graphene hybrid by facile one-step in situ crystallization hydrothermal method and as supercapacitor electrode. They prepared

Nanostructured Metal Oxide, Hydroxide and Chalcogenide

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three different sample compositions varying the molar ratios of metal precursors. The optimized sample exhibited specific capacity of 1998 C g−1 at 6 A g−1and capacitance retention of 91% after 2000 cycles at 20 A g−1, and also obtained 100% coulombic efficiency. Chandrasekaran et al. [52] synthesized hollow Mn–Cu–Al layered triple hydroxides nanocomposite using microwave assisted technique. The material showed the specific capacitance of 1632.94 F g−1 at 4 A g−1 current density. The fabricated symmetric device based on the material delivered high capacitance of 223.89 F g−1 at 1 A g−1 and retained 92.93% of initial capacitance after 10,000 cycles in the potential window of 1.8V. It further displayed the energy density (101.75 W h kg−1) and power density (0.9 kW kg−1). The supercapacitor performances of some metal hydroxide based electrodes has been summarized in Table 10.1.

10.4 Metal Oxide Transition metal oxides are unique class of solid functional materials in which oxygen are bound to the transition metals. The crystal structures of metal oxide constitute close-pack array (ccp) structures in which oxygen anions are associated with metal cations occupying interstitial sites. Transition metal oxides usually consist of ionic bonding, show multiple oxidation states, and exhibit catalytic and semiconducting properties. With such properties and availability of metal cations, metal oxides actively participate in Faradaic redox reactions and thereby extensively used as electrode materials for supercapacitor application.

10.4.1

Binary Metal Oxides

Among various metal oxides hydrous RuO2 is the best known pseudocapacitive material exhibiting high theoretical capacitance as it possesses higher conductivity, rich redox activity and greater reversibility [12]. However, due to expensive nature the use of RuO2 are limited in energy storage application. Presently, extensive scientific research is focused on materials alternative to RuO2, which are inexpensive and abundant oxides of transition metals like, Mn, Fe, Co, Ni, Cu, etc. Manganese oxide is among the interesting class of materials, which exist in various forms such as MnO, Mn3O4, Mn2O3, MnO2 as stable oxides. According to literature studies, the crystallinity, crystal structure, morphology, and specific surface area are key parameters which decide the performance of Mn-based oxides. It is well known that high crystallinity of

Specific capacitance @ current density

230.72 F g−1 @11.25 A g−1

241.9 F g−1 @ 1 A g−1

127 F g−1 @ 5 mA cm−2

119 F g−1 @ 1 A g−1

78 F g−1 @ 2 A g−1

98.5 F g−1 @ 0.5 A g−1

5.091 F cm−3 @ 10 mV s−1

332 F g−1 @ 1 A g−1

125.2 F g−1 @ 0.88 A g−1

84.26 F g−1 @ 1 A g−1

Sample

FeOOH//NiMoO4

Co(OH)2//N-doped graphene

Ni(OH)2//AC

Ni(OH)2/UGF//a-MEGO

o-CNT-Ni(OH)2//rGO

Co(OH)2//Vanadium Nitride

Cu(OH)2//AC

Ni-Co hydroxide//AC

NixCo1−x LDHs//AC

Ni-Mn LDH/rGO//AC

10000

5000

20000

2000

4000

-

10000

1000

10000

10000

Cycles

100%

92.7%

97.8%

73.4%

86%

-

63.2%

82%

93.2%

80.8

Retention

33.8 W h kg−1

23.7 W h kg−1

70.4 W h kg−1

0.255 mW h cm−3

22 W h kg−1

35.24 W h kg−1

13.4 W h kg−1

42.3 W h kg−1

98.9 W h kg−1

104.3 W h kg−1

Energy density

[57]

[56]

[50]

[39]

[34]

[55]

[54]

[53]

[33]

[35]

Ref.

(Continued)

0.85 kW kg−1

284.2 W kg−1

133.4 kW kg−1

383.222 mW cm−3

15.9 kW kg−1

1807 W kg−1

65 W kg−1

110 W kg−1

17981 W kg−1

1.27 kW kg−1

Power density

Table 10.1 Comparative study of electrochemical performance of the metal hydroxide based supercapacitors.

536 Nanomaterials for Electrochemical Energy Storage Devices

Specific capacitance @ current density

82.5 F g−1 @ 0.5 A g−1

84.9 F g−1 @ 1 A g−1

223.89 F g−1 @ 1 A g−1

226.7 F g−1 @1 A g−1

106 F g−1 @ 6 A g−1

162 Fg−1 @ 0.1 A g−1

87.9 F g−1 @ 5mA cm−2

344.9 F g−1 @ 0.2 A g−1

Sample

Ni–Mn LDH/rGO//rGO

Co–Ni–Fe–LDH/CNFs-0.5//AC

Mn–Cu–Al–TH// Mn–Cu–Al–TH

NiCo2Al–LDH//AC

AC/CNTs//Co0.5Ni0.5(OH)2/ graphene/CNTs

Ni,Co−OH/rGO//HPC

NiCo2O4@ Co0.33Ni0.67(OH)2// CMK-3

ZnCo1.5(OH)4.5 Cl0.5·0.45H2O @ Ni foam//PVAKOH//rGO 5000

3000

17000

-

5000

10000

2000

-

Cycles

81%

82%

80%

-

92%

92.93%

82.7%

-

Retention

114.8 W h kg−1

31.2 W h kg−1

56.1 W h kg−1

41 W h kg−1

91.0 W h kg−1

101.75 W h kg−1

30.2 W h kg−1

7.8 W h kg−1

Energy density

643.8 W kg−1

396 W kg−1

76 W kg−1

4.2 kW kg−1

758.2 W kg−1

0.9 kW kg−1

800.1 W kg−1

5600 W kg−1

Power density

Table 10.1 Comparative study of electrochemical performance of the metal hydroxide based supercapacitors. (Continued)

[64]

[63]

[62]

[61]

[60]

[52]

[59]

[58]

Ref.

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material favors enhanced conductivity, but reduces the specific surface area. Therefore, choosing an optimum temperature of heat treatment is essential to achieve high electrochemical performance of Mn-based oxides [65]. Chang et al. prepared amorphous MnOx nanofibers using mixture valences of Mn3+ and Mn4+ and also investigated the effects of heat treatment on the performance of electrode material. They noted that sample annealed at 200°C obtained an increased average oxidation state of Mn and displayed the best cyclic stability. After annealing at 300°C, the capacitive performance eventually decreased due to morphological changes. However, when MnOx was annealed beyond 500°C, the pseudocapacitive behavior disappeared due to the formation of crystalline Mn2O3 [65, 66]. Manganese dioxide (MnO2) shows high theoretical capacity of 1370 F g−1, can be used for supercapacitor application. Huang et al. fabricated porous MnO2 nanotubes of 6 nm thickness using template-assisted one-step hydrothermal approach as represented in Figure 10.8. It exhibited specific capacitance of 365 F g−1 at 0.25 A g−1 in a three-electrode system. The asymmetric device was fabricated using MnO2 nanotubes as positive electrode and activated graphene as negative electrode and it delivered specific capacitance of 50 F g−1 at a current density of 0.25 A g−1, and cyclic stability with capacity retention of 76.3% after 10000 cycles. The device also delivered an energy density of 22.5 Wh kg−1 with maximum power density of 146.2 kW kg−1 [67]. Fe2O3 being one of the most abundant material is significantly advantageous due to low cost, environmental friendliness and from safety point of view. However, their intrinsic poor conductivity and aggregated morphology restricts their extensive use as electrode material. Tang and his co-workers developed Fe2O3 nanowires (NW) with 10 nm of dia on carbon fiber paper using modified hydrothermal and post calcination process. The Fe2O3 NWs exhibited specific capacitance of 908 F g−1 at 2 A g−1 in a three-electrode system and showed 90% retention in capacitance up to 10 A g−1. They further fabricated asymmetric device using Fe2O3NW as anode and NiO as cathode. The device exhibited considerably high specific capacitance of 240 F g−1 at 4 A g−1 current density and showed high energy density of 105 W h kg−1 at the power density of 1400 W kg−1, which retained up to 72.6 W h kg−1 at 12700 W kg−1 [68]. Fan et al. prepared hybrid Fe3O4@carbon nanosheets using hydrothermal process. The hybrid material as electrode showed specific capacitance of 586 F g−1 and 340 F g−1 at 0.5 A g−1 and 10 A g−1, respectively. The solid-state asymmetric supercapacitor device assembled using carbons nanosheets in situ embedded Fe3O4 composite and porous carbon delivered energy density of 18.3 W h kg−1 at power density of 351 W kg−1 [69].

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Figure 10.8 (a) Schematic illustration of synthetic procedure of porous MnO2 nanotubes, SEM images of the MnO2 nanotubes (b) detailed image, (c) side-view, (d) enlarged view, (e) low-magnification TEM image of porous MnO2 nanotube, (f) detailed images of the terminal nanosheets of the MnO2 nanotube, (g) HRTEM image of the MnO2 (inset is the corresponding SAED pattern) and (h) galvanostatic charge–discharge profiles at different current densities, (i) cycling performance of MnO2 nanotubes//AG asymmetric device at 2 A g−1 current density (inset shows the charge–discharge curves of the last 10 cycles) and (j) digital image of a red-light-emitting diode (LED) lighted by the asymmetric device MnO2 nanotubes//AG (reproduced with permission from Ref. [67]).

Nickel oxide is highly abundant, low in toxicity, and exhibit desirable electrochemical properties for energy storage devices, including supercapacitor [70]. NiO, having the theoretical capacity 2573 F g−1, mainly favored by high electrical conductivity. Kim et al. [71] synthesized NiO nanostructures of three distinct morphologies (nanoflower, nanoslice, nanoparticle) using sol gel technique and exploited their morphology

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Nanomaterials for Electrochemical Energy Storage Devices

dependant supercapacitor performance. The flower-like NiO with high surface area and open morphology showed the best supercapacitor behavior by delivering 480 F g−1 specific capacitance at 0.5 A g−1 current density. In another report, Yu et al. [72] reported the synthesis of a porous 3D Ni@NiO core-shell electrode by the activation of Ni foam under mild conditions. They first activated the Ni foam by acid treatment and exhibited areal capacitance of only 2 F cm−2 at current density of 8 mA cm−2 and offered 10-fold improvement in capacitance upon depositing NiO nanostructure on it. The material exhibited exceptionally 170% enhanced capacitance over 100000 cycles. Cobalt oxide is considered to be another promising material for supercapacitor application and thereby extensive studies have been carried out on this material. It is noteworthy to mention that the environmentally benign and low cost nature along with high theoretical specific capacitance of 3560 F g−1 make the material much desirable for supercapacitor [73]. Sun et al. [73] developed a facile and simple strategy for the synthesis of nitrogen-doped carbonaceous aerogel/cobalt oxide as composite materials through in situ encapsulating method assisted by a multilevel freeze-drying process and post carbonization technique. Such composite with highest Co3O4 content exhibited the largest specific capacitance value of 616 F g−1 at 1 A g−1 among different compositions of the composite. The asymmetric device designed by the composite material as positive and carbon aerogel as negative electrode material exhibited specific capacitance of 107 F g−1 at current density of 0.5 A g−1. The device showed the capacity retention of 92.5% after 5000 cycles and also delivered energy density of 23.75 W h kg−1 at a power density of 4.5 kW kg−1. Zhou and his co-workers [74] fabricated a supercapacitor electrode by growing well-aligned CoO nanowire array on to 3D Ni foam coated with polypyrrole. The material exhibited the synergetic contribution from CoO nanowires and conductive PPy and displayed specific capacitance of 2223 F g−1 at 1 mA cm−2 and also obtained good cycling stability (Figure 10.9). Copper oxide may also be considered as potential material for supercapacitor application, but the use remain very limited as it suffers from major drawback of low electrical conductivity and unstable cycling performance which leads to relatively poor electrochemical efficiency. Recently, some researchers have fabricated electrodes by incorporation of copper hydroxide and copper oxide into carbonaceous materials such as graphene. Recently, Ghasemi et al. [75] successfully fabricated Cu2O–Cu(OH)2– graphene nanohybrid upon stainless steel and used it as electrode material for supercapacitor application, and obtained 425 F g−1 specific capacitance

Nanostructured Metal Oxide, Hydroxide and Chalcogenide (b)

Ni foam

Current (mA/cm2)

CoO@PPy NW

AC

Cellulose paper

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Li-ion

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(e)

102 103 101 Power Density (W/kg)

104

Energy Density (Wh/cm3)

10-3

0 min CoO@PPy//AC LSG supercapacitor Li thin-film battery AC//AC

10-4

-3

-2

160

15 min

Ref.51

30 min 10

(h)

(f2)

-1

10 10 10 Power Density (W/cm3)

0

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Retention (%)

10-2 100

120 80

100%

91.5%

40

25 mA/cm2

0 0

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10000 Cycles

15000

20000

Figure 10.9 (a) Schematic representation of the asymmetric supercapacitor configuration. (b) cyclic voltammetry (CV) and (c) discharge profiles, (d) Ragone plot of the asymmetric device (EDLC and Li-ion battery are also included), (e) Volumetric energy and powder densities of device compared with other data. (f1) Image showing two supercapacitors connected in series can lighten up three LED indicators. (f2) Images of the red LED at different stages powered by the 10 s charged supercapacitors. (g) A rotating motor derived by two supercapacitors in series. (h) cycling stability of the device after 20000 cycles (reproduced with permission from Ref. [74]).

at 5 A g−1 current density. The symmetric device delivered 87% of the initial capacitance after 2000 cycles. Wang and his co-workers [76] developed electrode composed of Cu2O nanoparticles by employing oxidationassisted dealloying technique and subsequent electrochemical reactions. The electrode material displayed 210.9 F g−1 and 163.3 F g−1 at 0.5 A g−1 and 10 A g−1 current densities in three-electrode system and offered capacity retention of 94.5% after 5000 cycles at 3 A g−1 current density. The assembled asymmetric supercapacitor device showed long-term cycling stability with more than 93.3% retention after 5000 cycles. It also exhibited high energy density of 20.04 W h kg−1 and at a power density of 7.1 kW kg−1. Asen et al. [77] prepared hybrid electrode composed of rGO/ PPy/Cu2O–Cu(OH)2 on Ni foam via electrochemical deposition method.

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Nanomaterials for Electrochemical Energy Storage Devices

Such combination of materials as electrode exhibited specific capacitance of 997 F g−1 at 10 A g−1 current density and displayed capacity retention of 90% of the initial capacitance after 2000 cycles. The assembled symmetric device showed specific capacitance of 225 F g−1 at 10 A g−1 current density. It showed energy density of 20 W h kg−1 at power density of 8000 W kg−1. Based on the studies, it is understood that though copper oxide as primary material may not be beneficial but combination with other metal oxide/ hydroxides or carbonaceous material exhibit promising performance as supercapacitor.

10.4.2

Ternary Metal Oxide

The ternary transition metal oxides are cubic crystal systems possessing spinel structures. The spinel structured oxides (AB2O4), where A and B are assigned to transition metals showing +2 and +3 oxidation states [78]. According to the crystal orientation, spinel structures are mainly categorized as normal and inverse spinel structures. In normal spinel structure, all the A2+ ions occupy the tetrahedral sites while all the B3+ ions occupy the octahedral sites (example: ZnFe2O4, CdFe2O4, etc.). In case of inverse spinel structure, ½ of the B3+ ions occupy the tetrahedral sites and the remaining ½ B3+ and all A2+ ions occupy the octahedral sites (example: NiFe2O4, etc.) [79]. The spinel structured AMn2O4 (A represents Co, Ni, and Zn) series are the potential candidates which are largely studied as electrode materials for energy storage devices [78]. Up to date, ternary metal oxides based on AB2O4 with a hierarchically complex structure, like NiCo2O4 [80], CoFe2O4 [81], CuCo2O4 [82], ZnCo2O4 [83], etc., are extensively used as electrode materials for supercapacitor application. Apart from low cost and environmental benignity, the materials exhibit relatively high pseudocapacitive performance due to availability of multivalent metal cations. Some of the studies of ternary metal oxides are discussed in this section. Li et al. [84] synthesized MnCo2O4 nanosheets wrapped on a hollow activated carbon shell (C@ MnCo2O4) through facile hydrothermal method and post calcination treatment. The C@MnCo2O4 electrode material exerted higher specific capacitance than that of pure porous carbon shell and pure MnCo2O4. It also achieved the specific capacitance of 728.4 F g−1 at 1 A g−1 and displayed 95.9% of retention in capacitance after 1000 cycles. Chen et al. [85] decorated metal oxide nanoparticles (Ni, Co) on an electrospun carbon nanofiber, which initiated the surface directed growth of NiCo2O4 in the form of nanorod and nanosheet morphologies. It was believed that the metal

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nanoparticles served as a transition layer to strengthen the interface and promote charge transfer between carbon and NiCo2O4 to form NiCo2O4/ Ni/CNFs. The material was found to exhibit good specific capacitance of 781 F g−1 at 0.6 A g−1 current density. Fu et al. [86] constructed hybrid ZnCo2O4/NixCo2x(OH)6nanostructure by electrodepositing NixCo2x(OH)6 on to previously hydrothermally grown ZnCo2O4. The sample obtained 419.1 μA h cm−2 at 5 mA cm−2, nearly three times than pristine ZnCo2O4 electrode (144.5 μA h cm−2) and over four times than NixCo2x(OH)6electrode (93.5 μA h cm−2) under identical conditions. The hybrid device, ZnCo2O4/ NixCo2x(OH)6//AC displayed the capacity retention of 88.2% after 2000 cycles at 30 mA cm−2. It delivered a maximum energy density of 26.2 W h kg−1 at 511.8 W kg−1 power density (Figure 10.10).

10.4.3

Quaternary Metal Oxide

Recently, quarternary metal oxides have emerged potential as electrode material for energy storage application. Such materials are expected to offer extended synergistic properties during redox reactions. Interestingly, the existence of multi-component in the metal oxide structure favors pseudocapacitive property. However, few works related to quarternary metal oxides are reported as supercapacitor electrodes. Recently, Zhang et al. [87] prepared porous quaternary zinc–nickel–aluminum–cobalt oxide (ZNACO) nanosheets directly on Ni foam using a facile and scalable chemical bath deposition and post calcination process. The electrode displayed specific capacity of 839.2 C g−1 at 1 A g−1 and rate capability with 82% retention in capacitance from 1 A g−1 to 20 A g−1, superior to NiCo2O4, ZnCo2O4, and Co3O4 electrodes. The fabricated device, ZNACO//AC obtained the specific capacitance of 232 F g−1 at 0.5 A g−1 and also showed 90% capacity retention after 10000 cycles. It exhibited high energy density of 72.4 W h kg−1 at a power density of 533 W kg−1. Chandrasekaran et al. [88] synthesized a hollow shelled triple Mn–Cu– Al oxide by microwave-assisted successive reduction method and investigated their supercapacitor application (Figure 10.11). They noted that MnO2, CuO, and Al2O3 formed a shell layer over the core SiO2 to form uniform-sized hollow geometrical nanostructure and obtained the specific capacitance of 319.81 F g−1 at 0.5 A g−1 current density. The fabricated symmetric supercapacitor showed the specific capacitance of 131.7 F g−1 at a current density of 0.5 A g−1 and delivered 83.48% capacity retention after 3000 cycles. It showed the maximum specific energy density and power density of 62.26 W h kg−1 and 5.5 kW kg−1.

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Nanomaterials for Electrochemical Energy Storage Devices (a) Hydrothermal

Ni foam

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Figure 10.10 Schematic representation of the construction process of hierarchical ZnCo2O4@NixCo2x(OH)6x core/shell NWAs on Ni foam; (b–e) SEM images of ZnCo2O4 NWAs (b and c) and ZnCo2O4@NixCo2x(OH)6xNWAs (d and e) on Ni foam; (f and g) TEM images and EDX spectrum (h) of the ZnCo2O4@NixCo2x(OH)6x; (i) schematic illustration of the assembled hybrid supercapacitor; (j) schematic illustration of the advantages of the hierarchical ZnCo2O4@NixCo2x(OH)6xNWAs; (k) Ragone plot of the hybrid device and (l) cycling stability of the device for 2000 cycles at 30 mA cm−2 (reproduced with permission from Ref. [86]). The supercapacitor performances of some metal oxide based electrodes has been summarized in Table 10.2.

10.5 Metal Chalcogenides Transition metal chalcogenides comprising transition metals (Fe, Co, Ni, Cu, Zn) and chalcogens (S, Se, Te) are another class of materials of interest [79]. It is well known that chalcogens (S, Se, Te) exhibit lesser electronegativity than oxygen, which leads to greater electrical conductivity and improved electrochemical activity of metal chalcogenides in comparison to their oxide counterparts. Recently, electrochemically active transition

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(a)

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Figure 10.11 (a) Schematic illustration of the synthetic process of ultrathin ZNACO nanosheets grown directly on Ni foam and SEM images of (b) ZNACOprecursor nanosheets on Ni foam,inset is an image with lower magnification, (c–d) ZNACO nanosheets on Ni foam, (e) schematic illustration of the ZNACO//AC hybrid supercapacitor configuration, electrochemical performance of the ZNACO//AC hybrid supercapacitor, (f) cell capacitance vs. current density, (g) cycling stability performance at a current density of 10 Ag−1 (inset is its corresponding charge/discharge curves of the last 5 cycles and photograph of agreed round LED indicator powered by two hybrid supercapacitors (reproduced with permission from Ref. [87]).

metal sulfides (TMS), especially bimetallic nickel cobalt sulfides, have emerged highly efficient pseudocapacitive electrode materials due to good electrical conductivity of about 2 orders higher than that of the oxide counterparts [106, 107]. In addition, metal chalcogenide based electrode material displays greater flexibility and stability than metal oxides [108]. Owing to their low cost and unique electrochemical properties, such materials have drawn considerable attention as supercapacitor electrodes.

10.5.1

Binary Metal Chalcogenides

Manganese sulfide, mainly exist in the form of MnS, is the interesting class of transition metal sulfide exhibiting semi-conducting nature. In the past, MnS were used as optical materials, magnetic devices and catalysis having

195.7 F g−1 @ 0.5 A g−1

248 mF cm−2 (50 F g−1) @ 1 mA cm−2

145 F g−1 @ 1 A g−1

-

108 F g−1 @ 0.5 A g−1

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99.4 F g−1 @ 0.5 A g−1

58.5 F g−1 @ 1 A g−1

576.6 F g−1 @ 1 A g−1

CSS/Graphite/PEDOT/ MnO2//AC

NiO//rGO

NiO//rGO

CC/SnO2/MnO2//NiO/Ni foam

CNT@NiO//DGNs

CoO@ppy//AC

NiCo2O4−rGO//AC

Fe3O4@carbon nanosheets

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[95]

[69]

[94]

[74]

[93]

[92]

[91]

[90]

[89]

[67]

Ref

(Continued)

Power density (W kg−1)

Table 10.2 Comparative study of the electrochemical performance metal oxide based supercapacitor electrodes.

546 Nanomaterials for Electrochemical Energy Storage Devices

113 F g−1 @ 0.5 A g−1

232 F g−1@ 0.5 A g−1

136 F g−1 @ 0.25 A g−1

119.4 F g−1 @ 20 mA cm−2

137 F g−1 @ 1 A g−1

236 F g−1 @ 2mA cm−2

129.74 F g−1 @ 1 A g−1

113.9 F g−1 @ 1 A g−1

109.9 F g−1 @ 0.1 A g−1

186.5 F g−1@1 A g−1

Co0.45Ni0.55O–rGO//rGO

ZNACO//AC

Ni–Zn–Co oxide/hydroxide// PC

CuCo2O4//AC

CuCo2O4//AC

NiFe2O4//NiFe2O4

ZICO//NG

ZNCO//AC

MNCO//C

NixCo3−xO4//AC

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2.03 F cm (148 F g ) @ 5 mA cm−2

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95.2%

94%

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98%

86%

92.5%

84%

89.8%

96%

Retention

37.4

123.6

35.6

40.5

47

42.81

37.3

41.65

72.4

35.3

90 472

Energy density (W h kg−1)

163

264

187.6

750

333

750

1500

85

533

330

472 1445

Power density (W kg−1)

[105]

[104]

[103]

[102]

[101]

[100]

[99]

[98]

[87]

[97]

[96]

Ref

Table 10.2 Comparative study of the electrochemical performance metal oxide based supercapacitor electrodes. (Continued)

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Nanomaterials for Electrochemical Energy Storage Devices

diversity in its properties. Moreover, due to reversible redox reaction with the involvement of different oxidation states of Mn, recently its use in the field of energy storage application was also explored. Zhang et al. [109] synthesized γ-MnS/reduced graphene oxide by facile one-pot hydrothermal synthesis technique and investigated its supercapacitor application. The material showed good electrochemical characteristics with specific capacitance of 547.6 F g−1 at 1 A g−1 and displayed rate capability with 65% retention in capacitance at 20 A g−1 current density. The fabricated asymmetric supercapacitor (γ-MnS/rGO//rGO) device delivered energy density of 23.1 W h kg−1 at power density of 798.8 W kg−1. In a similar approach, Quan et al. [110] prepared α-MnS nanoparticles anchored on nitrogen-doped reduced graphene oxide (α-MnS/N–rGO) nanosheets using one-pot solvothermal approach. The electrode displayed specific capacitance of 933.6 F g−1 at 1 A g−1 current density. The asymmetric device, α-MnS/N-rGO//N-rGO delivered the specific capacitance of 77.9 F g−1 at 1 A g−1 and capacity retention of 81.7% after 2000 cycles. It further provided the energy density of 27.7 W h kg−1 at the power density of 800 W kg−1. Iron sulfide is well-known for its low cost, highly abundance nature and mainly existence in seven major phases FeS, Fe3S4, Fe1−xS, FeS, Fe1+xS, FeS2, and FeS2. Knowing their moderately high theoretical capacitance values, they have emerged as a potential candidate as supercapacitor electrodes. Recently, Karade et al. [111] synthesized FeS thin film composed of nanoflakes by SILAR method. It exhibited specific capacitance of 297 F g−1 at a current density of 0.8 mA cm−2 and in a symmetric supercapacitor assemble using two FeS thin film electrodes sandwiched with PVA–LiClO4 gel electrolyte, the specific and volumetric capacitances of 4.62 F g−1 and 65.17 mF cm−3 at a current 0.75 mA was achieved. The device also showed 91% retention in capacitance at scan rate of 100 mV s−1 after 1000 cycling test and displayed specific energy density of 2.56 W h kg−1 (36.11 μW h cm−3) and specific power density as 726 W kg−1 (10.24 mW cm−3). Nickel sulfides are inexpensive and highly abundant, mainly exist in the form of NiS, NiS2, Ni9S8, Ni3S4. They are highly redox active material due to the availability of multiple oxidation states, displaying higher electrical conductivity than nickel oxide counterparts due to low band gap. Such unique properties of nickel sulfide, makes them an ideal candidate for supercapacitor application. Flower-like Ni3S2 embedded aerogel (rGO– Ni3S2) using hydrothermal synthesis followed by high-temperature pyrolysis was reported by Lin et al. [112] The material obtained the specific capacitance of 1315 F g−1 at 1 A g−1 current density in a three electrode system and in an asymmetric assemble of rGO–Ni3S2//AC, the device exhibited specific capacitance of 104.6 F g−1 at 0.5 A g−1 with good cyclic

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stability of 85.6% retention in capacitance over 5000 cycles at 4 A g−1. It also delivered energy density of 37.19 W h kg−1 at 399.9 W kg−1 power density. In an interesting report, Zang et al. [113] synthesized Ni nanowires via chemical method and fabricated a series of hybrid combinations of Ni3S2– Ni, Ni3S2–NiS–Ni, and Ni3S2–NiS with by controlled sulfurization process. The method allowed to grow nickel sulfides directly on the conductive Ni NWs, which expected to deliver good electrochemical performances. Among these combinations, Ni3S2–NiS was found to have superior redox reactivity due to not only the availability of multivalent Ni cations but also the very high conductivity of both Ni3S2 and NiS. Though the material showed good specific capacitance of 1077.3 F g−1 at 5 A g−1, but displayed relatively poor cycling stability of 76.3% of its initial values after 10,000 cycles. On the contrary, the Ni-preserving sulfide nanowires, such as, Ni3S2–Ni and Ni3S2–NiS–Ni showed the superior cycling stability of 100% and 87.2% capacitance retention after 10,000 cycles. Han and his co-workers [114] employed facile wet-chemistry approach to design three-dimensional forest-like porous nickel sulfide nanotrees on nickel foam (NiS NTs/ Ni foam) as shown in Figure 10.12. They optimized the growth parameters and surface morphology of NTs to achieve reliable adherence and used the sample as cathode material in hybrid supercapacitors combining with activated carbon as anode. The fabricated hybrid supercapacitor, NiS NTs// AC delivered areal capacitance of 1378.1 mF cm−2 at 5 mA cm−2 current density and showed capacity retention of 87.6% after 5000 cycles at 17 mA cm−2. The device exhibited a maximum areal energy density of 0.472 mW h cm−2 and power density of 21.5 mW cm−2 in a wide potential window of 1.6 V. The device found to have good potential for practical applications as it successfully operated electrical fans and LEDs. Cobalt sulfides, in its different forms, such as, CoS, Co9S8, CoS2, Co3S4, etc., are the most electrochemically active candidate for energy storage applications. Pu et al. [115] synthesized Co9S8 nanotubes arrays directly on Ni-foam by facile two-step hydrothermal synthesis using Co(CO3)0.35Cl0.20(OH)1.10 as template precursor followed by controlled sulfurisation involving Kirkendall effect to form Co9S8 on Ni foam. The material exhibited specific capacitance of 1483 F g−1 at a current density of 24 A g−1 in a three-electrode system and in a symmetric assemble high specific capacitance of 841 F g−1 at current density of 4 A g−1 was achieved. The device delivered cyclic stability with capacity retention of 91.4% after 2000 cycles and showed energy density of 49.9 W h kg−1 at a greater power density of 900 W kg−1. The combination of cobalt sulfide with carbonaceous material, such as, MWCNT as CoS2/MWCNT is also found to be beneficial for supercapacitor application [116]. Such electrode material (20 wt% MWCNT) showed the good specific capacitance of

Nanomaterials for Electrochemical Energy Storage Devices

550 (a)

NiS coated Ni foam Oriented attachment

Sulfur source

Nickel salt

NiS thin film Ni foam

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Figure 10.12 Schematic representation of (a) the overall synthesis process using a singlestep sulfurization approach, (b) the bare nickel foam indicating the three-dimensional metal framework, (c) the nucleation and precipitation mechanism between nickel and sulfur sources, (d) the formation of forest-like NiS NTs on Ni foam, (e) cycling stability test of the hybrid device (NiS NTs//PAC), (f) photographic image of the corresponding cell setup, (g–i) real-life application of pouch-type cells operating a toy motor fan, LED and personal computer cooler fan, respectively (reproduced with permission from Ref. [114]).

1486 F g−1 at 1 A g−1 and displayed cycling stability with ~80% retention of specific capacitance after 10000 cycles. Zinc sulfides are another interesting candidate showing unique semiconducting properties have gained considerable attention as

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supercapacitor electrode material. Recently, Javed et al. synthesized ZnS nanospheres directly on carbon textile (CT) using hydrothermal method. The material (ZnS/CT) exhibited specific capacitance of 747 F g−1 at 5 mV s−1 scan rate. The fabricated symmetric device showed 540 F g−1 specific capacitance at 5 mV s−1 scan rate and showed cycling stability with 94.6% retention of initial capacitance after 5000 cycles at current density of 0.8 mA cm−2. A high energy density of 51 W h kg−1 at a power density of 205 W kg−1 was also achieved from the device. The combination of ZnS with other material like TiO2 might be beneficial to deliver good performance as can be seen for the hybrid electrode material of flower-like ZnS on TiO2 nanotubes [120]. Among other TMCs, MoS2 is one of most frequently studied materials for supercapacitor application having their unique layered structure [121]. Though the actual mechanism is still unclear, but it is believed that both EDLC as well as pseudocapacitance contribute to the total performances [122]. Bissette et al. [122] prepared solvent stabilized dispersion of MoS2, MoSe2, WS2 and TiS2 by liquid exfoliation and TMC-based electrodes were prepared with the help of binders to fabricate coin cells. Upon capacitive performance comparison, TiS2 exhibits highest performance, though the partial oxidation of TiS2 into TiO2 remains as a negative factor. The chemical exfoliation of layered TMCs, like MoS2 as an example, followed by the restacking of layers to enhance the electrode performance is a smart approach [123]. Chhowalla and his group [124] showed chemical exfoliated MoS2 nanosheets with metallic 1T phase efficiently store energy by electrochemical intercalation of ions, like H+, Li+, Na+, or K+, etc. Such mechanism led to achieve extraordinary efficiency and capacitance value up to ~700 F cm−3 in a variety of aqueous electrolytes. Interestingly, the material found suitable to achieve high operational voltage of 3.5 V using organic electrolytes. There are also some other novel materials reported on metal selenides or tellurides such as CoTe@CFP [125], Se-doped NiTe [126], NiTe/NiSe [127], Cu2Se [128], Co0.85Se [129–132], MoTe2 [133], etc., which were used as supercapacitor electrodes. Ye et al. [125] synthesized cobalt telluride (CoTe) nanosheets on carbon fiber paper (CFP) using hydrothermal method. The material delivered specific capacitance of 622.8 F g−1 at 1 A g−1. The fabricated asymmetric device, CoTe//AC obtained specific capacitance of 192.1 F g−1 at 1 A g−1 and displayed energy density of 67.0 W h kg−1 at 793.5 W kg−1 power density. Peng et al. [129] synthesized petal-like Co0.85Se nanosheets using solvothermal method. The material was assembled with N-doped porous carbon networks to fabricate asymmetric device, Co0.85Se//N-PCNs, which shows good stability with the retention value of 93.8% over 5000 cycles and displayed energy density of 21.1 W h kg−1 at 400 W kg−1 power density.

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Nanomaterials for Electrochemical Energy Storage Devices

10.5.2

Ternary Metal Chalcogenides

Ternary metal chalcogenides have drawn considerable interest in recent days for scientific studies. Among various metal chalcogenides, thiospinel (AB2S4) structured materials are largely studied as electrode materials for energy storage devices. Metal chalcogenides possessing multi-metal cations are known to be much superior over mono-metallic chalcogenides as they favor synergistic properties and thereby facilitates higher electrochemical activities. As compared with mono-metal sulfides, ternary metal sulfides such as M–Co–S (M = Ni, Cu, Zn, and Mn) delivers much higher electrochemical activity due to the availability of multiple oxidation states and smaller band-gaps favoring richer oxidation reaction. Over the past few years several reports based on ternary metal chalcogenides, such as, CuCo2S4 [134–138], MxCo3−xS4 (M = Ni, Mn, Zn) [139], Zn0.76Co0.24S [140], FeCo2S4 [141], Co0.5Ni0.5S4 [142], NixCo1−xS2 [143], CoNi2S4 [144–147], MnCo2S4 [148–151], etc., were reported. In addition, some advanced functional hybrid materials such as Zn0.76Co0.24S/NGN/ CNTs [122], CuFeS2/C [152], NiCo2S4/GR [153], Ni–Co sulfide/graphene [107], NiCo2S4 on nitrogen-doped carbon foams [154], FeNi2S4@C [155], CuFeS2/C [152], etc. Gao et al. [156] synthesized a series of NixCo1−xS1.097 (x = 0 to x = 0.48) hierarchical structures via an ion exchange method using CoS1.097 microspheres as precursors, prepared by solvothermal reaction. The composition found to have good impact in the electrochemical performances of material. Further, by modifying with rGO, the electrode Ni0.48Co0.52S1.097–rGO displayed high specific capacitance value of 1152 F g−1 at current density of 0.5 A g−1 and also showed coulombic efficiency of 98% and high rate capability with 94% retention in capacitance. Xu et al. [157] prepared nickel cobalt sulfide nanosheets directly on Ni nanowires (NWs) film substrate using facile and one-step electrodeposition technique. The material obtained volumetric capacity of 392.8 C cm−3 at 0.5 A cm−3. The fabricated asymmetric device, Ni@Ni–Co–S//rGO–CNTs showed the volumetric capacitance of 59.1 F cm−3 at 0.34 A cm−3 and displayed 90.5% retention after 10000 cycles. It showed energy density of 18.4 mW h cm−3 and power density of 254.5 mW cm−3 (Figure 10.13). Wang et al. [158] synthesized multi-layered NixZn1−xS on Ni foam using one-step hydrothermal reaction, in which they noted that Zn played a critical role in the construction of multi-layered nanostructure. The sample displayed specific capacitance of 1815 F g−1 at 1 A g−1 current density. The asymmetric device assembled with active carbon–graphene delivered specific capacitance of 127.7 F g−1 and showed no loss in specific capacitance even after 5000 charge/discharge cycles. The

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device also exhibited moderate energy density of 38.9 Wh kg−1 at 327 W kg−1 power density. Yu et al. [159] constructed double-layered MnCo2S4@Ni–Co–S nanocomposite on Ni foam by hydrothermal method and subsequent electrochemical deposition. They noted that the deposited Ni–Co–S nanosheet on MnCo2S4 increases the electro-active surface area, electrical conductivity and reinforces the mechanical stability of the electrode. The material showed high areal capacitance of 10.14 F cm−2 at 1 mA cm−2 and the asymmetric assemble, MCS@NCS//AC, delivered an areal capacitance of 1.92 F cm−2 at 1 mA cm−2 with 5000 long cycles showing 72% retention. The device also displayed maximum energy density and power density of 7.26 W h m−2 (30.5 W h kg−1) and 162.7 W m−2 (683.7 W kg−1). Hou et al. [160] prepared hollow hetero NiCo2S4/Co9S8 solvothermally and consecutive hydrothermal anion exchange reactions. The asymmetric assemble of the material as NiCo2S4/Co9S8//AC showed high specific capacitance of 107 F g−1 at 0.2 A g−1 current density along with good cyclic stability over 5000 cycles. The device also delivered energy density of 33.5 W h kg−1 at a power density of 150 W kg−1 (Figure 10.14). Like sulfides, mixed transition metal selenides, such as, (Ni–Co)0.85Se [161], Ni–Co selenide [162], [email protected] [163], Ni0.9Co1.92Se4 [164], etc., have also been reported as potential material compositions for supercapacitor applications. Gou et al. [163] fabricated uniform bimetallic ternary [email protected] nanowires in large-scale through successive cation exchange technique. Initially, NiSe nanowires were grown on nickel foam by a facile solvothermal route, which was converted into a series of ternary materials involving different proportions of Ni and Co by Co-exchange method. The asymmetric supercapacitor device based on the material, [email protected]//AC, exhibited a good specific capacitance of 86 F g−1 at a current density of 1 A g−1 and displayed cycling stability with 100% capacity retention after 2000 cycles. It also delivered energy density of 17 W h kg−1 at a high power density of 1526.8 W kg−1.

10.5.3

Quarternary Metal Chalcogenides

Few reports on quarternary metal chalcogenides, such as, Ni–Co–Fe– Sulfide [165], Cu2NiSnS4@rGO [166], Cu2ZnSnS4 [167], Ni–Co–Zn–S [168], etc., are also available in the literature for supercapacitor application. For example, Sarkar et al. [166] prepared Cu2NiSnS4 nanoparticle in situ grown on two dimensional rGOvia hydrothermal synthesis. The asymmetric device based on the material displayed good areal and volumetric

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554

(a)

(b)

(c)

500 nm

500 nm

1 μm

5μm

(d)

100 nm

Ni

(e)

load

(f) + Ni@Ni-Co-S

rGO-CNTs -

OH-

10 nm

Co

K+

S

50 nm

separator -2

Current Density/A cm 20 50 30 40

60

70

1.0

60

0.8

50 0.6

40 30

0.4

20

0.2

10 0 0.0

0.5

1.0 1.5 2.0 2.5 -3 3.0 Current Density/A cm

3.5

(h)

70

0.0 4.0

160 140 120 90.5% retention

100 1.4

80

1.2

Potential (V)

10

Capacitance Retention/%

80

0

Areal Capacitance/F cm-2

Volumetric Capacitance/F cm-3

(g)

60 40 20 0

0.0 0.2

0.4

0

2000

0

200

400 600 Time (s)

4000 6000 Cycle Number

800

8000

10000

Figure 10.13 (a) SEM images of Ni nanowires (inset. is the magnified SEM image), (b) SEM images of Ni@Ni–Co–S–3 NWs (inset. is the magnified SEM image), (c) TEM image of Ni@Ni–Co–S–3 NWs, (d) magnified TEM image (inset is the HRTEM image), (e) element mapping scan image, (f) schematic of the assembled hybrid spercapacitor (Ni@Ni–Co–S–3 positive electrode and rGO–CNTs negative electrode), (g) volumetric capacitance and area capacitance of the device as a function of current densities and (h) cycle performance of the device at 2 A cm−3 (reproduced with permission from Ref. [157]).

capacitances of 655.1 mF cm−2 and 16.38 F cm−3 at 5 mA cm−2, which retained its capacitance up to 89.2% after 2000 cycles offering energy density of 5.68 mW h cm−3 and 246.9 mW cm−3 power density. Vignesh et al. [168] synthesized spherical ball in ball archirectured Ni–Co–Zn–S using one-pot hydrothermal method. The material with such hollow architecture found to be promising showing specific capacitance of 825 F g−1 at 1 A g−1 current density. The comparative study of the supercapacitor performance by various metal chalcogenides has been compiled in Table 10.3.

Nanostructured Metal Oxide, Hydroxide and Chalcogenide

O

Co2+

EG

(I)

NiCo2S4

Co9S8

10

NCCO

(II)

20

30

40

50

JCPDS no. 73-1704

60

70

(c)

Ni

10 mV s-1 20 mV s-1

0.0

80

0.3

0.6

0.9

1.2

1.5

E (V) 1.5

(d)

(j)

S 1.2

NCCS

c

Co a b

a

0.0

(g)

(h)

400

800 Time (s)

1200

1600

120

80

100 3 A g-1

60

(k)

-1

0 Ni

S

0

40

5Ag

Co

100 nm

80 60

40 20

c

0

100

SC (F g-1)

500 nm

500 nm

2.0 A g-1 3.0 A g-1 5.0 A g-1

0.3

c

(e)

(f)

0.2 A g-1 0.5 A g-1 1.0 A g-1

0.6

S b

0.9

CE (%)

Co

200 nm

2 mV s-1 5 mV s-1

(i)

-15 -30 -45 -60

JCPDS no. 86-2273

2 Theta(deg.)

60 45 30 15 0

I (mA)

C

E (V)

O Ni2+

(b)

2-

O

Intensity(a.u.)

(a)

555

1000 2000 3000 4000 5000 Cycle number

20 0

Figure 10.14 Schematic diagram of the synthetic procedure; (b) typical WAXRD patterns of the hollow hetero-NiCo2S4/Co9S8 sub micro-spindles (inset. the digital image of powder sample); crystallographic illustration for (c) spinel NiCo2S4 and (d) cubic Co9S8; (e) FESEM image; (f–g) TEM images (inset. SAED pattern (g)); (h) elemental mapping images of Co, Ni, S; (i) CV curves at various sweep rates; (j) charge/discharge plots at a wide current density; (k) cycling performance along with CE plot (3 A g−1) at high current rates of 3 and 5 A g−1 as a function of cycle number (reproduced with permission from Ref. [160]).

10.6 Summary and Future Perspective This chapter deals with the supercapacitor applications of metal hydroxides/ oxides/chalcogenides based electrodes. The recent advancement has been much focused on designing the electrode in nanoscale level providing high surface area and easy accessibility of electrolyte to the electroactive sites. The extensive literature survey on such materials reveals the pseudocapacitive nature exhibiting greater energy density than conventional EDLCs. In order to achieve further improvement in the energy density without compromising the high power density, strategic cell voltage improvement has been considered as an effective way. Thereby, fabrication of asymmetric devices with battery-like pseudocapacitive materials with the operational

119.1 F g−1 @ 5 mVs−1

122 F g−1 @ 1 A g−1

169.4 F g−1 @ 1 A g−1

146 F g−1 @ 1 A g−1

138.4 F g−1 @ 1 A g−1

NiCo2S4//G/CS

Ni–Co–S/G//PCNS

Ni–Co–S/NF//NG

NiCo2S4/MWCNTs-5//rGO

Ni@CNTs@Ni–Co–S//CC@CNTs

128 F g @ 5 mV s

160 C g−1 @ 5mA cm−2

CuxS/CF//AC

NiCo2S4/NCF//OMC/NCF

56.6 F g−1 @ 1 A g−1

Ni3S2/CNFs//CNFs–ASC 

−1

37.1 F g−1 @ 250 mA g−1

NiS HNPs//AC-ASC

−1

Specific capacitance @ current density

Material

10000

6000

5000

70000

10000

10000

5000

2500

4000

Cycles

70.4%

82.1%

85.7%

92%

85%

78.6%

88%

97%

100%

Retention

45.5

46.5

51.8

58.1

43.2

42.3

35

25.8

11.6

Energy density

512

800

865

786

22.1

10208

2667

425

187.5

[176]

[175]

[174]

[173]

[107]

[172]

[171]

[170]

[169]

Ref.

(Continued)

Power density

Table 10.3 Comparative study of the electrochemical performance metal chalcogenide based supercapacitor electrodes.

556 Nanomaterials for Electrochemical Energy Storage Devices

Specific capacitance @ current density

123.6 F g−1 @ 0.5 A g−1

940 F g−1 @ 1.5 A g−1

150.3 F g−1 @ 0.5 A g−1

146.7 F g−1 @ 2 A g−1

124 F g−1 @ 24 mA cm−2

150 C g−1 @ 0.5 A g−1

Material

ZnS–NiS1.97//AC

Ni @ rGO–Co3S4//Ni @ rGO–Ni3S2

Zn0.76Co0.24S/NCN/CNTs//NCN/ CNTs

CuCo2S4@NiMn–LDH//AC

CuCo2S4//AC

NiCo2S4@Ni3V2O8//AC 5000

6000

10000

2000

3000

6000

Cycles

94%

94.1%

87.6%

100%

96.2%

89%

Retention

42.7

44.1

45.8

50.2

55.16

36

Energy density

200

800

1499

38.5

13000

362

Power density

[180]

[135]

[179]

[155]

[178]

[177]

Ref.

Table 10.3 Comparative study of the electrochemical performance metal chalcogenide based supercapacitor electrodes. (Continued)

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cell voltage up to 2.0 V has been attained. In recent years, several research works are also carried out in the fabrication of advanced functional materials based on transition metal hydroxide/oxide/chalcogenides and their practical application as all solid-state supercapacitor with flexibility and long cyclic stability.

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171. Xu, P., Miao, C., Cheng, K., Ye, K., Yin, J., Cao, D., Pan, Z., Wang, G., Zhang, X., High electrochemical energy storage performance of controllable synthesis of nanorod Cu1.92S accompanying nanoribbon CuS directly grown on copper foam. Electrochim. Acta, 214, 276–285, 2016. 172. Shen, L., Yu, L., Wu, H.B., Yu, X.-Y., Zhang, X., Lou, X.W.D., Formation of nickel cobalt sulfide ball-in-ball hollow spheres with enhanced electrochemical pseudocapacitive properties. Nat. Commun., 6, 6694, 2015. 173. Zha, D., Fu, Y., Zhang, L., Zhu, J., Wang, X., Design and fabrication of highly open nickel cobalt sulfide nanosheets on Ni foam for asymmetric supercapacitors with high energy density and long cycle-life. J. Power Sour., 378, 31–39, 2018. 174. Wen, P., Fan, M., Yang, D., Wang, Y., Cheng, H., Wang, J., An asymmetric supercapacitor with ultrahigh energy density based on nickle cobalt sulfide nanocluster anchoring multi-wall carbon nanotubes hybrid. J. Power Sour., 320, 28–36, 2016. 175. Peng, T., Yi, H., Sun, P., Jing, Y., Wang, R., Wang, H., Wang, X., In situ growth of binder-free CNTs@Ni–Co–S nanosheets core/shell hybrids on Ni mesh for high energy density asymmetric supercapacitors. J. Mater. Chem. A, 4, 8888–8897, 2016. 176. Laifa, S., Jie, W., Guiyin, X., Hongsen, L., Hui, D., Xiaogang, Z., NiCo2S4 Nanosheets Grown on Nitrogen-Doped Carbon Foams as an Advanced Electrode for Supercapacitors. Adv. Energy Mater., 5, 1400977, 2015. 177. Wei, C., Ru, Q., Kang, X., Hou, H., Cheng, C., Zhang, D., Self-template synthesis of double shelled ZnS-NiS1.97 hollow spheres for electrochemical energy storage. Appl. Surf. Sci., 435, 993–1001, 2018. 178. Ghosh, D. and Das, C.K., Hydrothermal growth of hierarchical Ni3S2 and Co3S4 on a reduced graphene oxide hydrogel@ Ni foam: A high-energydensity aqueous asymmetric supercapacitor. ACS Appl. Mater. Interfaces, 7, 1122–1131, 2015. 179. Lin, J., Jia, H., Liang, H., Chen, S., Cai, Y., Qi, J., Qu, C., Cao, J., Fei, W., Feng, J., Hierarchical CuCo2S4@NiMn-layered double hydroxide core-shell hybrid arrays as electrodes for supercapacitors. Chem. Eng. J., 336, 562–569, 2018. 180. Niu, L., Wang, Y., Ruan, F., Shen, C., Shan, S., Xu, M., Sun, Z., Li, C., Liu, X., Gong, Y., In situ growth of NiCo 2 S 4@ Ni 3 V 2 O 8 on Ni foam as a binder-free electrode for asymmetric supercapacitors. J. Mater. Chem. A, 4, 5669–5677, 2016.

11 Polymer-Based Flexible Electrodes for Supercapacitor Applications Syam Kandula1, Nam Hoon Kim1 and Joong Hee Lee1,2* 1

Advanced Materials Institute for BIN Convergence Technology (BK21 Plus Global Program), Department of BIN Convergence Technology, Chonbuk National University, Jeonju, Jeonbuk, Republic of Korea 2 Carbon Composite Research Centre, Department of Polymer-Nano Science and Technology, Chonbuk National University, Jeonju, Jeonbuk, Republic of Korea

Abstract In recent years, with the rapid proliferation of portable and wearable electronics, the fabrication of flexible energy storage devices (FESDs) has become one of the prime interests of current research. Among the flexible energy storage electrode materials, conducting polymers (CPs) are more significant materials, because of their unique advantages, which include ease of synthesis, economy, good conductivity, and flexibility. Combination of the CPs with either carbon-based materials or metal oxide and metal sulfide materials can greatly enhance their physicochemical properties, as well as their electrochemical properties. The current chapter summarizes current research based on innovative configurations of flexible supercapacitors (FSCs) for energy storage applications, including freestanding, interdigitated, asymmetric, and fiber-based SCs. We also demonstrate the synthesis and electrochemical applications of individual CPs—in particular, polyaniline (PANI), polypyrrole (PPy), and poly(3,4-ethylenedioxythiophene) (PEDOT), as well as their binary and ternary composites. The effects of one-dimensional materials, as well as redox electrolytes, on electrochemical properties are discussed. The realtime practical applications of these composite materials are also demonstrated. In addition, a summary of various recent flexible electrodes is also presented. Keywords: Flexible energy storage devices, supercapacitors, conducting polymers, polyaniline, polypyrrole, PEDOT, energy density, power density

*Corresponding author: [email protected] Poulomi Roy and Suneel Kumar Srivastava (eds.) Nanomaterials for Electrochemical Energy Storage Devices, (573–624) © 2020 Scrivener Publishing LLC

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11.1 Introduction Environmental pollution and the energy crisis are two prominent concerns in the present world, because of the rapid depletion of natural fossil fuels by the use of a fast-growing world population. Moreover, environmental pollution is also caused by the large release of CO2 and soot from natural fossil fuels [1, 2]. Therefore, huge research effort has been devoted to replacing traditional and natural fossil fuels with other renewable energies as primary resources in our regular daily lives. However, while some of the other renewable energies, such as solar and wind energy, have been explored, they are unbalanced and intermittent, which suggests that there is an urgent need for efficient energy storage devices (ESDs) to meet the current demands of daily lives and industrialization [3, 4]. In this context, rechargeable batteries, fuel cells, and supercapacitors are emerging as effective ESDs, and are expected to play an important role in many practical future applications. Among them, supercapacitors (SCs) are considered the best choice, due to their ultrahigh power density (>10,000 W kg−1), considerable energy density (≤10 Wh kg−1), quick charge–discharge capability (within seconds), long life span (>10,000 cycles), and smart capabilities [3, 5]. In addition, current research results indicate that SCs can serve as effective ESDs that can fill the gap between traditional capacitors and conventional batteries. Generally, SCs are broadly classified into three main categories, based on the charge storage mechanism; they are (i) electric double-layer capacitors (EDLC), (ii) pseudocapacitors (PSC), and (iii) asymmetric supercapacitors (ASC) [6, 7]. In the case of EDLC, electrostatic energy is stored by the separation of charged particles in a Helmholtz double layer and ionic species at the conductive electrode–electrolyte interface, utilizing a non-Faradaic process. For example, highly porous materials, such as carbon materials and their derivatives, allow greater electrolyte accessibility via pores, as well as by the feature of higher surface area [6, 8]. In the case of PSC, the electrochemical energy is stored by effective surface redox reactions between the active working materials and electrolyte, and utilizes a reversible Faradaic process [5, 7]. For example, metal hydroxides, metal oxides, and metal sulfides can exhibit higher specific capacitance/capacity with low conductivity and poor cyclability. The ASC contains two or more active working materials with various complementary advantages, leading to a higher operational potential window, as well as higher energy density, and it has been developed to overcome the problems of lower specific capacitance/capacity and lower energy density [5, 9]. Current practical electronic applications mostly rely on the flexibility and twistability of the devices, which can play an important role in hybrid

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electric vehicles, wireless sensors, wearable electronic devices, portable military equipment, electronic skin, and medical monitoring devices [6, 10]. However, the bottleneck of the practical execution of flexible devices is the obstruction of the availability of highly flexible energy-supplying devices with high energy or power density, light weight, high rate capability, good stability, and long cycle life. In the case of wearable energy storage devices, the backbone of flexibility mainly reliant on the type of substrate, because it protects the activity of working electrode material without any loss, and also conducts internal current to the output supply. To date, a large number of flexible substrates, such as porous metal mesh, metal foam, metal sheet, carbon cloth, carbon and graphene paper, textile materials, have been employed [11]. Among these substrates, textile-based materials have attracted great attention for practical devices, because of their high mechanical flexibility, high surface area, light weight, and low cost, as well as ease of preparation in bulk quantity. Conducting polymers (CPs) were discovered in 1976; they are one of the important types of electrode materials for pseudocapacitors, because they possess improved specific capacitance/ capacity, high electrical conductivity, economical, and ease of synthesis [1]. In the case of CPs, the flow of electrons can be passed via the conjugated backbone of a polymer chain. Among the conducting polymers, polyaniline (PANI), polypyrrole (PPy), and poly[3,4-ethylenedioxythiophene] (PEDOT) and its derivatives have been well explored as electrode materials for pseudocapacitors, because they possess multiple redox states, which results in good control over the electrical conductivity ranging from insulator to metal [6, 9]. In recent years, the CPs has been explored as the most assuring electrode materials for FSC applications, owing to their extraordinary flexibility and economical method of fabrication. With respect to bulk, CPs with nanostructure morphologies, such as nanorods, nanowires, nanowalls, and nanosheets, possess good electrochemical performance, because of their peculiar features of surface interactions, conducting pathways, and high surface-to-volume ratios [1, 2]. The one-dimensional CPs with nanodimensions demonstrates ultrahigh pseudocapacitance, as compared to their bulk counterparts. In the current chapter, we provide a brief synthesis and discussion of PANI, PPy, PEDOT, and their composites, as well as their utility as a flexible electrode for practical SCs.

11.2 Pure Conducting Polymers (PCs) CPs display huge theoretical specific capacitances, which allow them to act as better electrode materials for SCs.

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11.2.1

Polyaniline (PANI)

PANI is a well-known flexible electrode material for SCs, which displays a theoretical specific capacitance of about 750 F g−1 [12]. It can be synthesized via chemical and electrochemical polymerization. The morphology of the PANI electrode material plays a crucial role in bettering the electrochemical properties of the FSCs. For example, Horng et al. synthesized the polyaniline on carbon cloth (CC) via electrochemical polymerization [13]. For the deposition of PANI on CC, a certain amount of aniline monomer (e.g., 0.2 M) was taken in optimum concentration of hydrochloric acid (e.g., 0.5 M) and potential applied (e.g., 0.6 V vs. Ag/AgCl) in a linear potential sweep voltammetry for a few minutes to initiate the nucleation of PANI by the oxidation of aniline monomer. PANI nanowires were grown on CC by prolonging the reaction time for a further few minutes (e.g., 10 min). The extra free particles of aniline monomer on the surface of PANI nanowires (NWs)/CC could be washed off by dipping them in hydrochloric acid (e.g., 0.5 M) solution, and stored at 4°C in hydrochloric acid solution to obtain the PANI NWs/CC. The PANI NWs/CC was explored for FSC electrodes, and exhibited a high specific capacitance of about 1,079 F g−1 at a current density of 1.73 A g−1 in a 1 M H2SO4 aqueous electrolyte. In addition, it also displayed about 70% specific capacitance retention by raising the current density from 1.73 to 17.3 A g−1. The flexibility of the PANI NWs/CC was also examined using cyclic voltammetry at a scan rate of 10 mV s−1 under various bending states from 0.0 to 2.5 cm, and showed only about 0.05 % specific capacitance loss, indicating the high robustness of the PANI NWs/CC flexible electrode for SCs. In another example, Wang et al. synthesized the PANI nanotube arrays via electrochemical polymerization, followed by template removal by weak acid or base solution [14]. First, they synthesized the ZnO nanorods, and then coated PANI on the ZnO nanorods surface by electrochemical polymerization of aniline, followed by dissolution of the ZnO nanorods; here, ZnO acts as a sacrificial template. The hollow PANI nanotubes were explored as FSCs, and displayed a specific capacitance of about 846 F g−1 at a scan rate of 5 mV s−1 in a 1 M H2SO4 aqueous electrolyte. The high specific capacitance of PANI nanotube arrays could be attributed to the fast diffusion of electrolyte ions parallel to the PANI nanotube orientation, the same as ion diffusion in the bulk electrolyte. Moreover, the electrolyte ions could reach or leave the surface of PANI nanotubes rapidly, even at high charge–discharge current densities, because of their vertical alignment, as well as their porous structure. The high electrochemical activity could also be due to the maximum usage of

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electroactive PANI nanotube surface via the orderly arrangement of pores in the nanotube arrays. Unfortunately, these PANI nanotubes displayed about 30% loss of their initial specific capacitance during the first 100 cycles at a scan rate of 50 mV s−1 and later sustained constant at up to 400 cycles. In summary, PANI alone could not meet the criteria for practical FSC device requirements, due to its low cyclic stability. To achieve a better cyclic stability, PANI should be mixed with various carbonaceous materials, metal oxides and metal sulfides, etc. The detailed study related to these composites is discussed in the following chapter.

11.2.2

Polypyrrole (PPy)

Polypyrrole is another well-known CP, which also possesses a good theoretical specific capacitance of about 620 F g−1 with good cyclic stability [15]. Generally, PPy can be synthesized by chemical oxidative polymerization, electrochemical polymerization, radiolytic polymerization, interfacial polymerization, layer-by-layer self-assembly, and enzyme-catalyzed polymerization. Among these methods, chemical oxidative polymerization is the best method, because of its suitability and economic feasibility to scale up for commercial synthesis [16]. In this method, pyrrole monomer is mixed with different oxidants (e.g., ferric perchlorate, ferric chloride, ammonium peroxydisulfate, etc.) along with a conductive substrate, and the reaction can be performed at the optimum conditions. For example, Yuan et al. synthesized the polypyrrole-coated paper by a simple soak and polymerization method [17]. During this process, a normal printing paper was soaked in pyrrole monomer for a few minutes, and then transferred into a certain amount of oxidant (e.g., FeCl3), along with an acid (e.g., HCl). Then, the reaction was performed at 4°C for a few hours, and finally washed with HCl, followed by a certain amount of NaCl and deionized water, respectively. They explored the pyrrole coated paper as a flexible electrode material for SC, which exhibited a high specific capacitance of about 370 F g−1 at a current density of 1 mA cm−2 in a 1 M HCl aqueous electrolyte. Moreover, the symmetric device exhibited an areal capacitance of about 0.42 F cm−2 at 1 mA cm−2. In addition, the symmetric device retained about 75.6% of its initial areal capacitance after a longterm cyclic stability for 10,000 cycles at a scan rate of 5 mA cm−2. The pyrrole coated paper displayed similar shaped CV curves at various bending states, and didn’t show any change in their current response, which indicates better flexibility of the electrode material. Shi et al. synthesized the conductive PPy hydrogel via an interfacial polymerization method [18].

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The PPy hydrogel displayed a specific capacitance of about 400 F g−1 at a current density of about 0.2 A g−1 in a 1 M H2SO4 aqueous electrolyte. Moreover, the PPy hydrogel exhibited only 3% loss of its initial specific capacitance under various bending conditions, indicating the good flexibility of the electrode for FSCs. In summary, PPy exhibits better cyclic stability as compared to PANI, but still doesn’t reach the value of theoretical capacitance alone. Various modifications, such as surface functionality, the formation of composites, and core@shell nanoparticles have been carried out to improve its further performance, and they are presented later in the chapter.

11.2.3

Poly(3,4-ethylenedioxythiophene) (PEDOT)

PEDOT is another well-known CP after the PANI and PPy, which shows a very high theoretical conductivity of about 500 S cm−1, and a comparable theoretical specific capacitance of about 210 F g−1 [19]. It can also show better physicochemical properties with better stability, and also extend the working potential window to a large extent. Anothumakkool et al. synthesized the highly conducting and potent PEDOT on a flexible cellulose paper by prompting surfactant-free interfacial polymerization at the junction of two immiscible liquids. They developed an effective and scalable method for the fabrication of highly flexible conducting PEDOT cellulose paper, which exhibited a very good conductivity of 375 S cm−1, with a low sheet resistance of 3 Ω cm−1. Figure 11.1 shows a schematic of the synthetic strategy for PEDOT coated on cellulose paper, the digital images of PEDOT coated paper before and after scratching, a 0.17 mm thin flexible solid-state SC made up of PEDOT paper, and a digital image of a 3.6 V interdigital SC made up of a monolayer PEDOT paper with LED glowing under flexible conditions [19]. The flexible PEDOT paper electrode exhibited a high specific capacitance of about 115 F g−1 at a current density of 0.5 mA cm−2, and a high volumetric capacitance of about 144 mF cm−2 at a current density of 0.5 A cm−2 in a PVA-H2SO4 gel electrolyte. Moreover, the symmetric flexible device exhibited no change in its galvanostatic charge– discharge curve nature at various bending states, which indicates better flexibility and robustness of the PEDOT-paper electrode. In addition, the flexible PEDOT-paper electrode retained almost 80% of its initial specific capacitance after 3,800 cycles at a current density of 2 mA cm−2 under various bending conditions. The loss of specific capacitance after many cycles in this system can be attributed to the loss of water from the gel electrolyte due to the generation of heat. In summary, PEDOT materials are more stable as compared to PANI and PPy, but the observed results are also too

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Figure 11.1 (a) A schematic representation of the synthetic strategy for PEDOT coated on cellulose paper, (b) a digital image of PEDOT-paper with 40 cm × 25 cm dimensions made by the current synthetic strategy, (c) a digital image showing the cleaned surface of scotch tape after peeling off from the PEDOT-paper surface, (d) a flexible 0.17 mm thin solid-state SC made up of from the PEDOT-paper, and (e) a digital image of 3.6 V interdigital SC made up of from a mono layer PEDOT-paper with LED glowing under flexible conditions. Reproduced with permission from Ref. [19], Copyright 2015, The Royal Society of Chemistry.

far from their actual theoretical specific capacitance value. In order to meet the criteria of practical applications, various PEDOT composite materials with carbon derivatives, metal hydroxides, metal oxides, and Mxenes have been examined. The details of these composite materials are discussed further in the following sections.

11.3 Conducting Polymer Composites (CPCs) Pure CPs exhibit many unique properties; however, they are unable to be utilized as effective electrode materials, because of their pure cyclic stability. To improve the electrochemical activity as well as the cyclic stability of the electrode materials, conducting polymer composites (CPCs) were introduced. The CPCs were typically composed of two types of CPs, including electrochemical double-layer capacitive materials and pseudocapacitive

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PEDOT: polyethylenedioxythiophene PPy: polypyrrole PANI: polyaniline CP: conducting polymer CNT: carbon nanotube G: graphene (r)GO:(reduced) graphene oxide CC: carbon cloth LDH: layered double hydroxide

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CP/Metal Oxide composites

Figure 11.2 Comparison of specific capacitances of CPs and CPCs. Reproduced with permission from Ref. [2], Copyright 2015, The Royal Society of Chemistry.

materials, respectively. The appropriate combination of CPCs can effectively improve the efficiency by tuning the redox active specific capacitance, as well as enhancing the cyclic stability, by modulating their surface morphology and inherent elastic polymeric nature [2]. Moreover, CPCs can also open the possibility of improvement of additional physical properties, which strongly depend on their counterparts, as well as the mutual interactions between them. Figure 11.2 shows the experimental values of CPs and CPCs. Figure 11.2 clearly shows the drastically enhanced specific capacitance of CPCs, compared to individual CPs. Moreover, the specific capacitances of CPCs are much higher when combined with either metal oxides or metal hydroxides, as compared to that of the carbon materials. These results suggest the applicability of these CPCs electrode materials for the practical advanced energy storage device applications of the future. In the current section, we discuss the synthesis of various combinations of metal hydroxide, metal oxide, metal sulfide, and carbon with polymer-based flexible electrodes and the electrochemical studies.

11.3.1 11.3.1.1

PANI-Based Binary Composites PANI- and Carbon-Based Binary Composites

For a long time, carbon, as well as its derivatives, has been attractive as SC electrodes, due to its versatility and abundance in nature, but its low

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specific capacitance restricts its usage for practical applications. On the other hand, carbon, as well as its derivatives, possesses many exceptional advantages, which include ultrahigh specific surface area, high electronic conductivity, better mechanical properties, and exceptional cyclic stability [20]. It is demonstrated that the combinations of these carbon materials with PANI exhibit better electrochemical activity toward FSCs, as compared to their counterparts [21]. For example, PANI is mixed with various carbon derivatives, such as carbon nanoparticles, carbon spheres, carbon fibers, carbon nanotubes (CNTs), multiwall carbon nanotubes (MWCNTs), graphene (G), graphene oxide (GO), and reduced graphene oxide (rGO), to improve the overall electrochemical properties of binary composites. Recently, Hashemi et al. grew the PANI rectangular tubes on functionalized carbon cloth (FCC) via in situ redox additive-assisted electrocatalytic method [22]. They employed 1,4-naphthoquinone (NQ) as a redox additive regenerator, and which helps in tuning the redox shuttle in the PVA/H2SO4 gel electrolyte. The symmetric PANI-FCC// PANI-FCC device exhibited a specific capacitance of about 484 F g−1 at a current density of 1 A g−1; the improved performance can be due to the synergistic effects between PANI rectangular tubes and the functional groups present on the surface of the carbon cloth substrate. This combination favors faster electron transport between the counterparts, as well as improving the redox activity of surface functional groups. Afterward, they also examined the AC-FCC//PANI-FCC asymmetric device in the presence of NQ as a redox component in the PVA/H2SO4 gel electrolyte, which surprisingly exhibited an ultrahigh specific capacitance of 4,007 F g−1 at a current density of 1.4 A g−1. They explained the huge increase in specific capacitance by various charge storage mechanisms, which are shown in Figure 11.3. The improved specific capacitance can be due to several reasons, which they explained as follows: (i) the solubility of the NQ being very much higher in the PVA/H2SO4 gel electrolyte than that in the aqueous H2SO4 electrolyte, as well as the wettability of the functionalized carbon cloth also being higher in the PVA gel electrolyte, resulting in improved specific capacitance, and (ii) the NQ actively promoting regeneration of PANI, which can be reused in the next cycles of redox reactions, which is shown as follows:

PANIox + 2e + 2H+ discharge/charge PANIred

(11.1)

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(11.2)

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e–

e–

–

– – + – – – + – + H2NQ – – + NQ + + + + + + – –– ––

e e e

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–

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+

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PANIred

NQ

PANIox

H2NQ

H2NQ

NQ

H++ H+ H+ H+ H

PANIred PANIox

H2NQ

+2H+ NQ

Figure 11.3 Schematic representation of AC-FCC//PANI-FCC asymmetric device exhibiting different charge storage mechanisms. Reproduced with permission from Ref. [22], Copyright 2017, Elsevier Ltd.

H 2 NQ

NQ 2e 2 H

(11.3)

During the discharge process, the PANI is used twice: it is used once for the primary purpose, and the second time for the regeneration of active electrode material. In addition, NQ also undergoes redox reactions on the carbon cloth substrate surfaces. Hence, NQ can act as a tunable redox shuttle as well as a redox additive, and the charge is stored by PANI via a pseudocapacitive mechanism, and at the junction of the electrode–electrolyte via the redox reactions. The asymmetric AC-FCC//PANI-FCC device exhibited about 84% specific capacitance retention after 7,000 cycles at a high current density of about 35 A g−1. Moreover, the asymmetric AC-FCC//PANI-FCC device also displayed a very high energy density of about 1,091 Wh kg−1, with an ultrahigh power density of about 196,000 W kg−1. To verify the flexibility of the device, they also examined the change in resistance at various bending angles ranging from 0 to 180°. It exhibited only