Advanced Battery Materials 9781119407553, 1211211231, 3563573573, 1119407559

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Advanced Battery Materials

Scrivener Publishing 100 Cummings Center, Suite 541J Beverly, MA 01915-6106 Publishers at Scrivener Martin Scrivener ([email protected]) Phillip Carmical ([email protected]) Managing Editors: Sachin Mishra, S. Patra and Anshuman Mishra

Advanced Battery Materials

Edited by

Chunwen Sun Beijing Institute of Nanoenergy and Nanosystems, China

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 © 2019 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 978-1-119-40755-3 Cover image: Pixabay.Com Cover design Russell Richardson Set in size of 11pt and Minion Pro by Exeter Premedia Services Private Ltd., Chennai, India Printed in the USA 10 9 8 7 6 5 4 3 2 1

Contents Preface 1 Carbon Anode Materials for Sodium-Ion Batteries Hongshuai Hou and Xiaobo Ji 1.1 Introduction 1.2 Sodium Storage Mechanism of Carbon Materials 1.2.1 Graphite Materials 1.2.2 Hard Carbon Materials 1.2.2.1 Insertion-Absorption Mechanism 1.2.2.2 Absorption-Insertion Mechanism 1.2.2.3 Absorption-Filling Mechanism 1.3 Carbon Anode Materials for Advanced Sodium-Ion Batteries 1.3.1 Graphite Materials 1.3.2 Hard Carbon Materials 1.3.3 Heteroatom-Doped Carbon Materials 1.3.4 Biomass Derived Carbon Materials 1.4 Conclusion and Prospects Acknowledgements References 2 Lithium Titanate Based Lithium-Ion Batteries Jiehua Liu, Xiangfeng Wei and Fancheng Meng 2.1 Introduction 2.2 Benefits of Lithium Titanate 2.3 Geometrical Structures and Fabrication of Lithium Titanate 2.3.1 Zero-Dimensional Nano-Architectures 2.3.2 One-Dimensional Nanoarchitectures 2.3.3 Two-Dimensional Nanostructures 2.3.4 Three-Dimensional Nanostructures 2.3.5 Other Nanostructures

xv 1 2 6 6 8 9 11 19 21 21 24 31 47 55 70 70 87 88 91 93 94 98 105 109 117 v

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2.4 Modification of Lithium Titanate 2.4.1 Surface Coating 2.4.2 Doping 2.4.3 Hybrids 2.5 LTO Full Cells 2.6 Commercial LTO Batteries 2.7 Other Applications 2.8 Summary and Outlook Acknowledgements References 3 Rational Material Design and Performance Optimization of Transition Metal Oxide-Based Lithium Ion Battery Anodes Qingshui Xie and Dong-Liang Peng 3.1 Introduction 3.2 Transition Metal Oxide-Based Anodes with Intercalation Mechanism 3.2.1 The Lithium Storage Properties of Titanium Dioxide-Based Anodes 3.3 Transition Metal Oxide-Based Anodes with Conversion Mechanism 3.3.1 The Lithium Storage Properties of Manganese Oxide-Based Anodes 3.3.2 The Lithium Storage Properties of Iron Oxide-Based Anodes 3.3.3 The Lithium Storage Properties of Cobalt Oxide-Based Anodes 3.3.4 The Lithium Storage Properties of Nickel Oxide-Based Anodes 3.3.5 The Lithium Storage Properties of Copper Oxide-Based Anodes 3.3.6 The Lithium Storage Properties of Molybdenum Oxide-Based Anodes 3.4 Transition Metal Oxide-Based Anodes with Alloying Mechanism 3.4.1 The Lithium Storage Properties of Zinc Oxide-Based Anodes 3.4.2 The Lithium Storage Properties of Tin Dioxide-Based Anodes 3.5 Summary and Outlook References

121 121 123 131 132 137 140 142 143 143 159 160 162 162 166 166 171 173 178 182 185 189 189 192 195 196

Contents

4 Effects of Graphene on the Electrochemical Properties of the Electrodes of Lithium Ion Batteries Wenzhuo Shen and Shouwu Guo 4.1 Introduction 4.2 Effects of Graphene on the Electrochemical Properties of Cathode Materials 4.2.1 LiFePO4 4.2.1.1 Doping Modification 4.2.1.2 Surface Modification 4.2.2 LiMn2O4 4.2.3 LiCoO2 or LiNiO2 4.2.4 Li[Ni1-X-yCoxMny]O2 4.3 Effects of Graphene on the Electrochemical Properties of Anode Materials 4.3.1 Graphene and Its Derivatives 4.3.2 Metal Oxides 4.3.3 Carbon Nanotube or Fullerene Hybrid 4.3.4 Li4Ti5O12 4.3.5 Si or Ge 4.3.6 Sulfide 4.3.7 Alloy 4.4 Conclusions and Perspectives References 5 Practically Relevant Research on Silicon Based Lithium-Ion Battery Anodes Qing Ji, Yonggao Xia, Jin Zhu, Binjie Hu, Peter Müller-Buschbaum and Ya-Jun Cheng 5.1 Introduction 5.2 Lifetime of Batteries 5.3 High Energy Density 5.4 Low Cost Silicon 5.5 Large-Sized Silicon 5.5.1 Mechanical Milling (Physical Mixing): Carbon, Metal 5.5.2 Chemical Coating (Chemical Mixing): Metal Deposition, Carbon Coating 5.5.3 Chemical Etching 5.5.3.1 Silicon-Metal Alloy 5.5.4 Silicon From Silica Through Magnesiothermic Reduction

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209 209 210 210 210 214 225 228 229 235 235 236 239 239 241 242 243 243 244 261

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5.5.5 Synthesized Silica 5.5.6 Commercial Silica 5.5.7 Natural Silica Sources 5.5.8 Perspective of Magnesiothermic Reduction 5.5.9 SiOx (x 1):Disproportionation to Produce Si Convertible Oxide 5.6 Industrial-Related Perspective 5.7 Conclusion References 6

Mo-Based Anode Materials for Alkali Metal Ion Batteries Zhanwei Xu, Weiwei Guan, Yixing Zhao, Kai Yao, Juju He, Xinyue Liu and Xintong Duan 6.1 Mo-Based Anode Materials for Alkali Metal Ion Batteries 6.1.1 MoO2 6.1.1.1 Storage Mechanism 6.1.1.2 Pure MoO2 as Anode Materials 6.1.1.3 MoO2-Based Composite as Anode Materials 6.2 MoO3 6.2.1 Storage Mechanism 6.2.1.1 Storage Mechanism for LIBs 6.2.1.2 Storage Mechanism for SIBs 6.2.2 Pure MoO3 as Anode Materials 6.2.2.1 Pure MoO3 for LIBs 6.2.2.2 Pure MoO3 for SIBs 6.2.3 MoO3-Based Composite as Anode Materials 6.2.3.1 MoO3-Carbon for LIBs 6.2.3.2 MoO3-Carbon for SIBs 6.2.3.3 Heterogeneous MoO3-Based for LIBs 6.2.3.4 Other MoO3-Based for LIBs 6.2.4 MoO3-X 6.3 MoS2 6.3.1 Storage Mechanism 6.3.1.1 Storage Mechanism for LIBs 6.3.1.2 Storage Mechanism for SIBs 6.3.2 Pure MoS2 as Anode Materials 6.3.3 MoS2-Based Composite as Anode Materials 6.3.3.1 MoS2-Graphene for LIBs 6.3.3.2 MoS2-N-Carbon for LIBs 6.3.3.3 MoS2-Carbon for SIBs

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6.3.3.4

MoS2-PEO for SIBs

6.4 MoSe2 6.5 Oxysalts (MMoO4 (M = Fe, Co, Ni, Ca)) 6.5.1 Cobalt Molybdate (CoMoO4) 6.5.1.1 CoMoO4 as Anode Materials for LIBs 6.5.2 Iron Molybdate (FeMoO4) 6.5.2.1 FeMoO4 as Anode Materials for LIBs 6.5.3 Nickel Molybdate (NiMoO4) 6.5.3.1 NiMoO4 as Anode Materials for LIBs 6.5.4 Calcium Molybdate (CaMoO4) 6.5.4.1 CaMoO4 as Anode Materials for LIBs 6.6 Summery and Lookout References 7 Comprehensively Understanding Lithium-Sulfur Batteries: Current Status and Outlook Xue Bai, Tao Li, Umair Gulzar, Remo Proietti Zaccaria, Claudio Capiglia and Yu-Jun Bai 7.1 Introduction 7.2 Fundamental Li-S Electrochemistry 7.2.1 Electrochemical Principles of Li-S Batteries 7.2.2 Challenges and Problems 7.3 Sulfur Cathodes 7.3.1 Sulfur/Carbon Composites 7.3.2 Porous Carbon 7.3.2.1 Graphene 7.3.2.2 Carbon Fibers and Carbon Nanotubes (CNTs) 7.3.2.3 Hollow Structured Carbon 7.3.3 Li2S Cathode 7.4 Anode 7.4.1 Protection of Metal Lithium 7.4.2 Electrolyte Additives 7.4.3 Silicon Anode Materials 7.4.4 Tin Anode Materials 7.4.5 Carbon-Based Anode Materials 7.5 Separator 7.5.1 Carbon 7.5.2 Polymer 7.5.3 Inorganic Fillers

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334 334 337 337 337 339 340 342 343 343 344 347 347 355

356 357 357 358 360 360 360 362 362 364 366 369 369 370 371 373 374 376 376 378 379

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7.6

Electrolyte 7.6.1 Liquid Electrolyte 7.6.2 Inorganic Solid Electrolytes 7.6.3 Polymer Electrolyte 7.6.3.1 Gel Polymer Electrolytes 7.6.3.2 Solid Polymer Electrolytes 7.7 Application and Prospects 7.8 Summary References

8 Graphene in Lithium-Ion/Lithium-Sulfur Batteries Guillermina L. LuqueMaría Laura Para, Emiliano N. Primo, M. Victoria Bracamonte, Manuel Otero, María del Carmen Rojas, J. Francisco, García Soriano, German Lener and Andrea Calderón 8.1 Introduction 8.2 Graphene in Lithium-Ion Batteries 8.2.1 Graphene in Anodes of Lithium-Ion Batteries 8.2.2 Graphene Applied to Cathode Materials for Lithium-Ion Batteries 8.3 Graphene in Lithium-Sulfur Batteries 8.3.1 Graphene in Lithium-Sulfur Batteries Cathodes 8.3.2 Graphene in Lithium-Sulfur Batteries Separators 8.3.3 Graphene for Lithium Anode Protection in Lithium-Sulfur Batteries 8.4 Conclusions and Outlooks References 9 Graphene-Ionic Liquid Supercapacitors: Design, Fabrication and Applications Minjie Shi, Xuefeng Song, Cheng Yang, Yuyu Tian, Kai Tao, Jijin Xu, Peng Zhang and Lian Gao 9.1 Introduction 9.2 Investigation of the Storage Mechanism of Graphene-ILs SCs 9.2.1 Influence of the Composition of ILs Electrolytes on the Electrochemical Characteristics of Graphene-ILs SCs 9.2.1.1 Research of Interaction Between Graphene Electrode and IL Electrolyte in SCs

380 380 381 382 382 384 385 388 388 399

400 404 405 410 418 419 426 432 433 435 451

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Design of the Graphene-ILs SCs with Favorable Electrochemical Performance 9.2.2.1 Graphene with Various Reasonable Structures 9.2.2.2 Graphene Integrated with Electro-Active Materials 9.3 Development of Flexible Graphene-ILs SCs 9.4 Fabrication of Planar Graphene-ILs Micro-SCs 9.5 Challenges and Prospect Acknowledgements References

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10 Development of Battery Electrodes Based on Carbon Species and Conducting Polymers Shahbaz Khan and Abdul Majid 10.1 Introduction 10.2 Battery Operation 10.2.1 Rechargeable Batteries 10.2.2 Lithium Ions Batteries (LIBs) 10.2.3 History of Lithium Ions Battery Electrodes 10.2.4 Lithium Ions Battery Operation 10.3 Cathode Materials for LIBs 10.4 Anode Materials for LIBs 10.4.1 Carbon-Based Anode Materials 10.4.2 Graphene Nanosheets (GNS) 10.4.3 Carbon Nanotubes (CNTs) 10.4.4 Single Walled Carbon Nanotubes (SWCNTs) 10.4.5 Multi Walled Carbon Nanotubes (MWCNTs) 10.4.6 Free-Standing Carbon Nanotubes Paper Anodes for LIBs 10.5 Conducting Polymer-Based Anodes 10.6 Lithium Oxygen Batteries (LOBs) 10.6.1 General Principles 10.6.2 Cathode Materials for LOBs 10.6.3 Graphene Nanosheets (GNS) 10.6.4 Carbon Nanotubes (CNTs) 10.6.5 Anode Materials for LOBs 10.7 Synthesis and Characterization of Carbon-Based Cathode Catalytic Materials for LOBs 10.7.1 Free Standing Paper Anode of Ge/SWCNT Composite

457 457 459 465 467 469 472 472 477 478 479 479 480 480 480 482 483 483 483 483 484 484 491 491 493 493 494 495 495 496 496 496

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10.7.2 GNS/PbGeO3 Composite 10.7.3 Pd/PNCNF-2 Catalytic Cathode Composite Acknowledgements References 11 Doped Graphene for Electrochemical Energy Storage Systems Jie Yang, Zhenghui Pan, Yuegang Zhang and Yongcai Qiu 11.1 Introduction 11.2 Properties of Graphene 11.3 Brief Introduction to Undoped Graphene for Electrochemical Energy Storage Systems 11.4 Preparation Methods of Doped Graphene 11.4.1 Nitrogen (N)-Doped Graphene (NG) 11.4.1.1 Chemical Vapor Deposition 11.4.1.2 Hydrothermal Methods 11.4.1.3 Thermal Treatment 11.4.1.4 Plasma Treatment 11.4.1.5 Pyrolysis 11.4.1.6 Ball Milling 11.4.1.7 Supercritical 11.4.1.8 Electrochemical 11.4.2 Boron (B)-Doped Graphene (BG) 11.4.2.1 Thermal Decomposition 11.4.2.2 Gas-Solid Reactions 11.4.2.3 Direct Solid State Reactions 11.4.2.4 Microwave Plasma 11.4.2.5 Hydrothermal Methods 11.4.2.6 Reflux Processes 11.4.3 Sulfur (S)-Doped Graphene (SG) 11.4.3.1 Hydrothermal Methods 11.4.3.2 Thermal Treatment 11.4.3.3 Exfoliation of Graphite 11.4.3.4 Ion Exchange Resin 11.4.3.5 Ball Milling 11.4.4 Phosphorus (P)-Doped Graphene (PG) 11.4.4.1 Pyrolysis 11.4.4.2 Post-Thermal Annealing Synthesis 11.4.4.3 Liquid Process 11.4.5 Co-Doped Graphene 11.4.5.1 B and N-Doped Graphene (BNG) 11.4.5.2 S and N-Doped Graphene (SNG)

500 501 504 504 511 511 512 514 516 516 516 520 521 521 522 523 524 525 525 527 527 527 528 528 529 529 530 531 532 534 534 534 535 536 536 537 537 539

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11.4.5.3 P and N-Doped Graphene (PNG) 11.4.6 Halogen-Doped Graphene 11.5 Doped Graphene for Electrochemical Energy Storage Systems 11.5.1 Supercapacitors(SCs) 11.5.1.1 NG for Supercapacitors 11.5.1.2 Other Heteroatom-Doped Graphene for SCs 11.5.1.3 Co-Doped Graphene for Supercapacitors 11.5.2 Lithium-Ion Batteries (LIBs) 11.5.2.1 NG as Anode Materials for LIBs 11.5.2.2 B-Doped Graphene as Anode Materials for LIBs 11.5.2.3 Other Heteroatom-Doped Graphene as Anode Materials for LIBs 11.5.3 Lithium-Sulfur Batteries (LSBs) 11.5.3.1 O-Doped Graphene as Host Materials for LSBs 11.5.3.2 NG as Host Materials for LSBs 11.5.3.3 BG as Host Materials for LSBs 11.5.3.4 Co-Doped Graphene as Host Materials for LSBs 11.5.4 Metal-Air Batteries 11.5.4.1 Doped Graphene as Air Cathodes for ZABs 11.5.4.2 Doped Graphene as Air Cathodes for LABs 11.5.5 Sodium-Ion Batteries (SIBs) 11.6 Prospective Outlook Acknowledgement References 12 Processing of Graphene Oxide for Enhanced Electrical Properties Bhargav Raval, S. K. Mahapatra and Indrani Banerjee 12.1 Introduction 12.2 Processing of Graphene Oxide 12.2.1 Low Temperature Plasma Processing 12.2.1.1 Hydrogen Based Glow Discharge Plasma

543 544 545 545 546 551 557 559 561 565 566 567 569 572 575 576 578 579 582 585 587 588 588 613 614 617 617 618

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12.2.1.2 Ammonia Based Glow Discharge Plasma 12.2.1.3 Methane Based Glow Discharge Plasma 12.2.1.4 Argon, Oxygen and Other Gas-Based Glow Discharge Plasma 12.2.2 Electron Beam Irradiation Processing 12.3 Properties of Graphene Oxide 12.3.1 Hydrogen Plasma Treated Graphene Oxide 12.3.2 Ammonia Plasma Treated Graphene Oxide 12.3.3 Methane Plasma Treated Graphene Oxide 12.3.4 Electron Beam Irradiated Graphene Oxide References Index

619 619 619 620 621 622 625 629 631 635 645

Preface Electrochemical energy storage has played important roles in energy storage technologies for portable electronics and electric vehicle applications. During the past three decades, great progress has been made in research and development of various batteries, in terms of energy density increase and cost reduction. However, the energy density has to be further increased to achieve long endurance time. In this book, recent research and development in advanced electrode materials for electrochemical energy storage devices are showcased, including lithium ion batteries, lithium-sulfur batteries and metal-air batteries, sodium ion batteries and supercapacitors. The materials involve transition metal oxides, sulfides, Si-based material as well as graphene and graphene composites. The contributors to the volume are battery scientists and engineers with excellent academic records and expertise. Each chapter is relatively independent of the others, with a structure which is easy for readers quickly find topics of interest. I hope that this book will be helpful for scientists and engineers working in the field of energy storage, especially the graduate students. This book mainly addresses a primary discussion, latest research & developments, industry and the future of battery materials. The book begins with the discussion on the recent progress of the carbonaceous anode materials including the sodium storage performances of amormphous carbons, graphite/graphene-based carbons, heteroatoms-doped carbons, biomass derived carbons, and the corresponding sodium storage mechanism. In addition, the current critical issues, challenges and perspectives of carbon anode materials for sodium ion batteries are also discussed in this chapter. Chapter 2 summarizes the performance of lithium titanate-based lithium-ion batteries with three classified themes including organic half Li-ion cells, organic full Li-ion cells and Na-ion batteries. The outlook and perspective on lithium titanate-based lithium-ion batteries have also been concisely provided in this chapter. Recent research advance in the controllable fabrication and the future of various transition metal oxide-based electrode materials and their lithium storage properties are xv

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presented in chapter 3. In chapter 4, there is a discussion on the recent progresses on the effects of the graphene on the electrochemical performances of cathode materials. Additionally, the preparation and applications of the composites of carbonaceous materials with graphene in the anodes of lithium ion batteries has also been incorporated in this chapter. Chapter 5 summarizes the practically relevant studies on silicon anodes for Li-ion batteries. Chapter 6 provides a systematic summary of the synthesis techniques, modification methods, as well as electrochemical property and performance of Mo-based compounds in lithium/sodium-ion batteries. The electrochemical performances and the related charge/discharge mechanism is also discussed in this chapter. The current application situations have been described to introduce the state-of-art of Li-S battery in chapter 7. Chapter 8 presents a critical overview of the state-of-art in the optimization and application of graphene derived materials for anodes, cathodes and separators in lithium batteries. In chapter 9, the recent achievements in the design and fabrication of flexible graphene-ionic liquids supercapacitors, and their application in portable electronics has been discussed. Chapter 10 provides an overview on composites of conducting polymers and activated carbon. Along with a detailed perspective on the reported experimental techniques and theoretical strategies to tune the properties of carbon and conducting copolymers composites based electrode materials. The preparation and application of doped graphene in the electrochemical energy storage systems has been summarized in chapter 11. Chapter 12 emphasizes the techniques of low temperature plasma processing and electron beam irradiation techniques to enhance the electrical properties of graphene oxide. I would like to express my gratitude to all the contributors for their collective and fruitful work. It is their efforts and expertise that have made this book comprehensive, valuable and unique. Grateful thanks to Sachin Mishra, S. Patra and Anshuman Mishra for managing the chapters and their help and useful suggestions in preparing Advanced Battery Materials. Finally, I would like to thank the International Association of Advanced Materials for all their help and direction. Chunwen Sun Beijing November 2018

1 Carbon Anode Materials for SodiumIon Batteries Hongshuai Hou1,2 and Xiaobo Ji1,2,* 1

College of Chemistry and Chemical Engineering, Central South University, Changsha, China 2 State Key Laboratory of Powder Metallurgy, Central South University, Changsha, China

Abstract A rechargeable ion battery is a kind of high-efficiency energy storage and conversion system. Lithium-ion battery (LIB) has been widely applied to lots of fields since it was commercialized in the 1990s, including various electronic products and electric vehicles. With the fast development of electronic products and electric vehicles, the market demand of LIB has dramatically expanded, while the lithium resource is not rich on the earth, which will inevitably lead to the rise of LIB cost, limiting the large-scale application of LIBs. Although the sodium-ion battery (SIB) with the similar working principle as LIBs was ignored after the commercialization of LIBs, it has attracted attention again due to the abundant sodium resource, and now it is often considered as the promising alternative for LIBs. One of the main limiting factors for the development of SIBs is the absence of proper anode materials, but in recent years, important progress on the anode materials have been made. Among lots of reported anode materials for SIBs, carbon material may be one of the most attractive candidates owing to abundant resource, low cost, good stability, nontoxicity, and high safety. A variety of carbonaceous materials have been evaluted as sodium storage anodes, involving graphite, graphene and amorphous carbon materials. In comparison with the graphitized carbon materials, the amorphous carbon materials exhibited better electrochemical performances due to the multiple sodium storage modes, including adsorption, intercalation nanopores filling. To further improve the electrochemical sodium storage properties, micro/nanostructure design and heteroatoms-doping was conducted. In consideration of the resource and sustainability, a large number of biomass derived carbon anode materials were developed. For the

*Corresponding author: [email protected] Chunwen Sun (ed.) Advanced Battery Materials, (1–86) © 2019 Scrivener Publishing LLC

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Advanced Battery Materials

sodium storage mechanism of hard carbons, there are conflicting opinions regarding the assignment of Na+ storage mode at different voltage regions. Although some important achievements have been made, the disadvantageous points are still to be solved, like low initial Coulomb efficiency, poor rate capability and relatively high manufacturing cost. In this chapter, the recent progress of the carbonaceous anode materials is summarized and discussed, including the sodium storage performances of graphite/graphene-based carbons, amormphous carbons, heteroatoms-doped carbons, biomass derived carbons, and the corresponding sodium storage mechanism. In addition, the current critical issues, challenges and perspectives of carbon anode materials for SIBs are discussed as well. We really wish that this chapter can help readers to understand the carbonaceous anode materials in SIBs. Keywords: Energy storage, sodium-ion battery, sodium storage, anode, carbon material, sodium storage mechanism

1.1

Introduction

Energy is the base of human existence and the development of human society is highly dependent on the emergence of high-quality energy and the application of advanced energy technology. For thousands of years, people obtained energy from natural world to survive and reproduce. Nowadays, the progress of energy, especially clean energy, is one of the hottest topics concerned by the people all over the world. In the past several decades, the over-exploitation and utilization of fossil fuel, leading to rapid exhaustion of fossil fuel resources and emerging environmental problems. Employing novel renewable and clean energy sources to substitute for the fossil fuel is highly desired. The wind energy, tidal energy, geothermal energy, hydroenergy, and solar energy are growing rapidly; nevertheless, they are all intermittent. To realize the integration of these renewable energies into the electrical grid, building the large-scale energy storage system (ESS) is vital to the operation of peak shift [1]. Among a variety of energy storage technologies, electrochemical secondary battery is a promising large-scale electricity storage device due to high energy conversion efficiency, flexibility, and simple maintenance [1–3]. The rechargeable alkali metal-ion battery, like lithium-ion battery (LIB) and sodium-ion battery (SIB), were proposed in 1970–1980s. And the LIB was successfully commercialized in 1991 by Sony, and then lots of studies were focused on the LIBs. On the contrary, SIBs were left out and bare related investigations were reported for three decades. In recent years, the flourishing development of various electric-equipment, including multifarious consumer electronics, electric tools and electric vehicles,

Carbon Anode Materials for Sodium-Ion Batteries Cathode

3

Anode

Layered oxide

Carbonaceous

Na-ion battery

O3-type Metal oxide

P2-type Polyanion Metal alloy

Additive

Binder Solvent Salt

Electrolyte & Binder Anode

Working potential / V (vs. Na+/Na)

2.0 Carbonaceous

Metal alloy

TMO, TMS

Organic

FeSx Na2DBQ

1.5 Schiff

Na2C10H2O4

bases

Li4Ti5O12 NaTiO2

1.0

CoSx, MoS2

CuxO, SnSx

CoS3O4 SnOx

Fe3O4 Fe2O3

TiO2

Natural

Na2C8H4O4

Sb

Sn

Graphite

SnSb

0.5

Ge ZnS Na4Ti5O12 Na2Ti3O7

P

Non-graphitic carbon (hard carbon)

0.0 0

SnxP3

S, N-droped carbon, Graphene

100

200

300

400

500

1000

300

Specific capacity / mAh g–1

Figure 1.1 Schematic illustrations of SIBs (a) and recent research progress in anode for SIBs (b). Reproduced with permission [1]. Copyright 2017, Royal Society of Chemistry.

largely increased the market demand of LIBs. Unfortunately, the lithium is not an abundant element, the lithium resource reserve is only 20 ppm in the earth cluster [2, 3], and this would seriously limit the supply of LIBs. On account of this, now the SIB attracts more and more attentions owing to the abundant sodium resource (23600 ppm in the earth cluster) [4, 5]. Since the elemental Na is adjacent to Li in the group of alkaline metal on the periodic table of elements, their physiochemical properties are quite similar. Thus, the electrochemical work principle of SIBs (Figure  1.1a)

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Advanced Battery Materials

resembles that of LIBs. When charging, the Na+ ions are released from the cathode materials and transferred to the anode and then inserted into the anode material. While in the discharging process, the reverse process is performed, Na+ ions are extracted from the sodiated anode material and move to the cathode to restore to the original state. The discharge process is accompanied with the electronic migration in the external circuit, providing the electric energy [3, 5]. In the past decades, the absence of suitable anode materials is one of the key factors which lead to the shelved research of SIBs [6]. In 1980s, carbon materials were initially investigated as anode materials for LIBs. And it was found that graphite has a high theoretical specific capacity of 372 mAh g-1, corresponding to the formation of Li-graphite intercalation compounds (LiC6), and the voltage-capacity curve of graphite shows a flat plateau at a low potential region of 0.1–0.2 V (vs. Li/Li+) [7]. Subsequently, the graphite was successfully applied in the commercialized LIBs. Regrettably, it was discovered that the Na+ ions cannot be easily inserted into the graphite interlayer. Before 2000, only several anode materials for SIBs were reported, and their energy densities were much lower than that of graphite for LIBs. In 1988, Ge and Fouletier [8] investigated the electrochemical intercalation process of Na+ into graphite, they proposed that NaC64 was formed during the intercalation of Na+, and the corresponding theoretical specific capacity of graphite was 35 mAh g-1. In 1993, based on the calculated results from the time and amount of current passed, Doeff and coworkers proposed that the insertion of Na+ ions into graphite, petroleum coke and Shawinigan black gave rise to the formation of NaC70 (31 mAh g-1), NaC30 (70 mAh g-1) and NaC15 (132 mAh g-1), respectively [9]. In 2000, Dahn et al., demonstrated that the hard carbon materials can exhibit high reversible specific capacity of ≈300 mAh g-1, which is close to the lithium storage capacity of graphite (372 mAh g-1) [10]. Although its cycle ability was poor, it still brought new hope to the development of SIBs. After 2010, the research regarding SIBs increased sharply, a large quantity of sodium storage anode materials are exploited (Figure 1.1b), such as, alloys [11] (Sn [12–17], Sb [18–27], Ge [28–31], Bi [32–34], P [35–39] and their alloys [11], metal oxides40/sulfides41/selenides42/phosphides [43] (e.g., SnO2 [44–47], Sb2O3 [48–51], Sb2O4 [52], CuO [53, 54], Fe2O3 [55, 56], Co3O4 [57, 58], TiO2 [59–61], NiCo2O4 [62, 63], MnFe2O4 [64], MnCo2O4 [65], MoS2 [66–69], SnS2 [70, 71], Sb2S3 [72–75], Sb2S5 [76], Bi2S3 [77–79], CoS [80], CoS2 [81, 82], NiS2 [83], Ni3S2 [84], NiSe2 [85], CoSe2 [86], Sb2Se3 [87, 88], FeP4 [89], Co2P [90], NiP3 [91], CuP2 [92], MoP [93], and Sn4P3 [94–97]), and carbonaceous materials [98, 99] (including

Carbon Anode Materials for Sodium-Ion Batteries

5

expanded graphite, graphene and amorphous carbon). Benefit from the unique structure and composition, some emerging anode materials showed greatly improved sodium storage performances, which is even comparable to LIBs. Similar to LIBs, the alloy-type anode materials often have high theoretical capacities, but they usually suffer from the large volume change in the charge-discharge process, resulting in the poor cycle stability. Wang and coworkers constructed a phosphorus/graphene nanosheet composite anode, which exhibited a high reversible capacity of 2077 mAh g-1 and it can be remained at 1700 mAh g-1 after 60 cycles [100]. Kovalenko and coworkers reported that the Sb nanocrystals had a high capacity of >500 mAh g-1 at the high current density of 13200 mA g-1 [101]. Based on the multi-electron conversion reactions, the metal oxides/sulfides/selenides/ phosphides exhibit high capacities and energy densities as well. However, these materials are commonly burdened with poor cycle stability, low initial coulombic efficiency and large hysteresis. Chen et al., constructed a porous Fe2O3/C nanocomposite to improve the sodium storage of Fe2O3, the capacity of the nanocomposite electrode could reach ≈1100 mAh g-1 at the second cycle and 740 mAh g-1 was kept after 200 cycles [55]. Yu and coworkers also designed the yolk-shell Sn4P3@C nanosphere composite electrode to enhance the battery properties of Sn4P3. It delivered a high reversible capacity (790 mAh g-1) and good cycle stability [97]. Ji and coworkers reported a Sb2S3@C nanorod composite electrode, which delivered a high specific capacity of 699.1 mAh g-1 after 100 cycles, and the related capacity retention can reach 95.7% [72]. The sodium storage way of Ti-based oxides is the insertion reaction, and they usually presented excellent cycle ability, but the capacity is relatively low and the rate capability is poor. Ji et al., designed the petal-like TiO2 wrapped by graphene to enhance the rate capability [102], and Yu et al., constructed the S-doped TiO2 to promote the sodium storage properties [103]. Among lots of the reported anode materials, the carbonaceous materials are most attractive due to the abundant resource, stability and low cost, which involve expanded graphite, graphene, carbon spheres, carbon fibers, carbon tubes, carbon sheets, porous carbons, and other amorphous carbons [98, 99, 104, 105]. In general, the capacities of carbonaceous anode materials are between 200 and 500 mAh g-1. As is well-known, the large-scale commercialized batteries should have some crucial features, including high safety, low cost and long cycle life. Therefore, the selection of electrode material should be also based on these standards, the abundance, nontoxicity, stability and durability of electrode materials should be preferentially considered. It is nice to find that carbon materials can

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Advanced Battery Materials

meet almost all of the above-mentioned requirements. In this chapter, we mainly summarize the sodium storage mechanism of different carbons and the progress of various carbon anode materials in SIBs, including graphite anode, amorphous carbon anode, heteroatom-doped carbon anode, and biomass derived carbon anode.

1.2

Sodium Storage Mechanism of Carbon Materials

1.2.1 Graphite Materials For commercialized LIBs, graphite materials, including natural graphite and synthetic graphite, are the most common anode materials. As shown in Figure 1.2a, during the charge-discharge process, the graphite intercalation compound (LiC6) is formed due to the reaction of Li+ + 6C + 6eLiC6, corresponding to a theoretical specific capacity of 372 mAh g-1. SIBs and LIBs have the similar work principle, in the same way, the storage of sodium in graphite also derives from the generation of graphite intercalation compound (NaCx). But a crucial difference of Na+ and Li+ has a dramatic influence on the specific capacity, the effective intercalation of Na+ in to graphite is limited, the related electrochemical reaction is Na+ + 70C + 70e- NaC70, a quite low specific capacity of 31 mAh g-1 is obtained [105]. Apparently, the pristine graphite with such a low capacity is not suitable for practical application in SIBs. Of late, graphite has been confirmed to be able to store K+, Rb+ and + Cs through the intercalation of these alkali metal ions into graphite interlayer [107–109]. The diverse intercalation behaviors of above-mentioned alkali ions into graphite lead to various electrochemical performances of graphite electrode in alkali ion batteries, and the fundamental origin of the performance diversity is desired to be explored. Early, researchers often attribute the poor sodium storage properties of graphite to the larger size of Na+ ions. Since SIBs have rapidly grown into the research focus in the past several years and been considered the promising recharge ion batteries, some researchers have been devoted to unfolding the mystery of the poor sodium storage properties of graphite, and hoping to propose an effective solution. A number of theoretical studies were carried out to simulate and compute the Na+ insertion process. Nobuhara and co-workers [109] investigated the alkali metal-graphite intercalation compounds (AM-GICs) utilizing the density functional theory (DFT). The calculated investigation regarding the atomic structure of the AM-GICs indicated that the lengths of C-C bond in the Na-GICs are

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Figure 1.2 (a) The graphite intercalation compounds of Li+ and Na+ ions. Reproduced with permission.[105] Copyright 2015, Wiley-VCH. (b) The formation energy (Ef ) values of MC6 and MC8 structures (M = Li, Na, K, Rb, and Cs). (c) Contribution factors to Ef values of AM-GICs. Reproduced with permission.[106] Copyright 2016, Wiley-VCH.

stretched more than that in the K-GICs, meaning that graphite is stressed after the intercalation of Na+ into it, bring about the result that sodiumgraphite intercalation compounds (Na-GICs), such as NaC16, NaC12, NaC8 and NaC6, suffer from energetic instability, which is the main reason for poor sodium storage properties of graphite. Based on the calculated results, the K-GICs are stable at least until KC8, additionally, according to the calculated energy barriers of the Li, Na and K ions jumps between the sites

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Advanced Battery Materials

in the graphite, it is found that the larger radius ions more smoothly diffuse in the graphite. Thus, it is obvious that the size of the alkali ion is not a factor, which directly limits the intercalation of alkali ion into graphite. In addition, Okamoto [110] examined the energetics of alkali metal intercalation into graphite through first-principles calculations, stating that the formation of NaC6 and NaC8 is difficult, which results from the high redox potential of Na/Na+. Wang et al., [111] Liu et al., [112] and Kang [106] further analyzed the instability of Na+ intercalation into graphite from another angle, as shown in Figure 1.2b, c, they divided the formation energy (Ef) of Na-GICs into three parts base on Hess’s law: (1) the energy cost of forming isolated atoms from the bulk metal (Ed, i. e., the cohesive energy); (2) the reconfiguration energy of the graphite resulting from the formation of the GICs (Es); (3) the energy drop (Eb) deriving from the insertion of Na+ into graphite. Because the weak Na binding exceeds the decrease of Es +Ed, the Ef of Na-GICs is higher than that of other alkali metal-GICs. As a consequence, the low sodium storage capacity of graphite should be directly associated with Eb.

1.2.2 Hard Carbon Materials Different from the graphite, amorphous carbons have few long-range order structures in the plane and ordered stacked structures in the c direction. Because of the high degree of disorder in c direction and perfect hexagonal network of carbon atoms in the plane, the amorphous carbons are difficult to be graphitized. Typically, amorphous carbons consist of three components, namely, graphitized microdomains with random distribution, distorted graphene sheets, and nanovoids among these two regions. The randomly dispersed and cross-linked graphite microdomains would be inclined to suppress the graphitized process and maintain the state of amorphous structure. Generally, two wide peaks located at around 24.8° and 43.8° can be observed in the XRD patterns of amorphous carbons, which are related to the crystal plane of (002) and (100). Additionally, there are two characteristic peaks at ~1350 cm-1 (D band) and ~1580 cm-1 (G band) can be found in Raman spectrum of amorphous carbon materials, which are associated with defects/disorder and ordered graphitic structures, respectively. And the intensity ratio of D band to G band (ID/ IG) can be utilized to evaluate the disordered degree of amorphous carbon materials. According to the possibility of further graphitizing, amorphous carbons can be divided into soft carbons and hard carbons, the one which can be converted into graphite over 2000 °C is classified as soft carbon, and the one which always holds the disordered structure is categorized as

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hard carbon. The nature of carbon precursor has a significant influence on the graphitization, the aromatic hydrocarbon precursors often lead to high graphitized degree. Commonly, soft carbons can be achieved from pyrolytic aromatic compounds and polymers, and hard carbons can be obtained from pyrolytic biomass with less aromatic structures [99, 105]. The sodium storage properties are tightly determined by the microstructure of carbon materials. The narrow interlayer distance of graphite and soft carbon lead to the fact that the main sodium storage way of them is the absorption of Na+ on their surface. Because there is no layer stacking structure in graphene, the absorption of Na+ is the main way of sodium storage as well. In the sodiation-desodiation profiles (specific capacity vs. voltage), graphite and graphene with little micropores and functional groups show the analogous behaviour of capacitive oblique lines, and soft carbons with micropores and functional groups present relatively curving lines [99]. For hard carbon, the electrochemical sodium storage behavior is remarkably different from all the above-mentioned carbon materials [99, 113]. Its sodiation curve is composed of a sloping line in the high potential region and a flat line in the low potential region. Nowadays, there are divergent opinions regarding the sodium storage mechanism at different potential regions. Some researchers advocate that the sloping line in the high potential region is related to the intercalation of Na+ into the expanded carbon layers, and the plateau region corresponds to the adsorption/deposition of Na+ in the micropores, which is similar to the lithium storage mechanism. In contrast, other researchers think that the sloping region is associated with the Na+ absorption of defective sites, active sites and functional groups on the surface of carbon materials, and the plateau region is ascribed to the insertion behavior of Na+. In addition, for some hard carbon materials, the intercalation behavior of Na+ was not found in both sloping region and flat region, so some researchers propose the third mechanism without sodium insertion process.

1.2.2.1

Insertion-Absorption Mechanism

In 2000, Dahn and Stevens firstly proposed the insertion-absorption mechanism for the first time according to the investigation of lithium and sodium storage behavior of hard carbon materials from the pyrolytic glucose [10]. It was discovered that the sodiation-desodiation profile of hard carbon is very similar to the lithiation-delithiation profile, (Figure 1.3a,b). To be convenient for understanding this electrochemical behavior, they built a “house of cards” model (Figure 1.3c), namely, the hard carbon is

Advanced Battery Materials 0.2 0.1 0.0

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Figure 1.3 The charge-discharge (voltage vs. capacity) curves of glucose derived hard carbon in (a) SIBs and (b) LIBs. The insets present the magnified low-potential regions. (c) “House of cards” model for the sodium/lithium storage in hard carbon. Reproduced with permission.[10] Copyright 2000, The Electrochemical Society. (d) Schematic illustration of “insertion-adsorption” mechanism for sodium storage in hard carbon. Reproduced with permission. [114] Copyright 2017, Wiley-VCH.

comprised of numerous microcrystal carbon layers (cards), a portion of carbon layers (cards) arrange in parallel to constitute the graphitic microcrystal domains and other carbon layers (cards) disperse in disordering to create the nanosized microporous domains. They attributed the sloping line at the high potential range and the flatting line at the low potential range to the insertion of Na+ or Li+ into carbon interlayer the adsorption of the Na+ or Li+ in the micropores among the randomly stacked layers, respectively (Figure 1.3d). Later, they further explored the sodium/lithium storage in hard carbon using the in-situ X-ray scattering, confirming that the intercalation of Na+ and Li+ into the carbon interlayer of hard carbon is truly feasible, which is accompanied with the expansion interlayer. Meanwhile, the proposed adsorption of Na+ and Li+ ions and in micropores was also verified [115]. Although this mechanism has explained the corresponding experimental phenomenon and been demonstrated by the related experimental results, with the large increase of the investigation regarding hard carbons in SIBs, researchers find that this mechanism cannot be suitable for all hard carbons. For instance, it was discovered that low capacity in the plateau region is observed for some hard carbons with abundant micropores, and with the decrease of the micropores’ volume and the increase of graphitized degree, the plateau capacity gradually increases, which is

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in contradiction with the “insertion-absorption” mechanism. Thus, a new sodium storage mechanism of hard carbons is desired to understand the electrochemical process.

1.2.2.2 Absorption- Insertion Mechanism In 2012, Cao and co-workers firstly proposed a new mechanism which is opposed to the insertion-absorption mechanism based on experimental and theoretical approaches [116]. They pointed out that the sloping region should be ascribed to the adsorption of Na+ on the surface and nanovoids while the plateau region is attributed to the insertion of Na+ into carbon interlayers. Their study was based on the evaluation of sodium and lithium storage behaviors in hollow carbon nanowires (HCNWs) obtained from pyrolyzing polyaniline, suggesting that the electrochemical behavior of HCNWs in SIBs (Figure 1.4a) is distinctly different from that in LIBs (Figure  1.4b). The electrochemical sodium storage behavior of HCNWs in the low potential region is analogous to the lithium storage behavior of graphite (Figure  1.4c), corresponding to the insertion of metal ions into the carbon interlayers. Based on the lithium storage behavior of hard carbon and graphite, it was inferred that the electrochemical reaction in the high potential region should be associated with the charge transfer on the surface of graphitic microdomains. To further discover the Na+ intercalation/extraction process of the HCNWs, theoretical simulations were carried out base on the balance of attractive van der Waals interactions between carbon layers and repulsive interactions between Na+ and carbon, suggesting that the equilibrium interplanar spacings for NaC6 is 0.37 nm (Figure 1.4d), respectively, and the energy cost of Na+ insertion into graphite interlayers is 0.12 eV, which is much larger than the thermal fluctuations energy (0.0257 eV) at room temperature, leading to the difficult insertion of Na+ into graphite. Note that the energy barrier would decrease with the increase of interlayer spacing, when the interlayer spacing reaches 0.37 nm, the energy barrier could reduce to 0.053 eV, which is low enough to be overcome, agreeing well with the experimental results. Cao et al.,[114] have measured the diffusion variation of Na+ throughout the intercalation-extraction process in the hard carbon nanoparticles prepared by pyrolyzing polyaniline, demonstrating the phase conversion from graphitized domains with enlarged interlayer distance to intercalation compounds (NaCx). They gave an important insight into the stepwise insertion/extraction of Na+, phases and rates in hard carbon materials. In 2017, to further understand the sodium storage process, Cao and coworkers [117] synthesized a series of well-controlled nano hard carbon

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Figure 1.4 Initial two charge-discharge curves of the HCNWs in (a) SIBs and (b) LIBs. (c) Initial two charge-discharge curves of the graphite in LIBs. (d) The relationship between the theoretical energy cost for the insertion of Na+/Li+ into carbon and the interlayer distance of carbon. The schematic of the insertion of Na+/Li+ into carbon interlayer is presented in the inset. Reproduced with permission [116]. Copyright 2012, American Chemical Society. (e) The correlation between interaction of Na+ ions with two graphitic surfaces and the distance to the surfaces normalized by the inter surface spacings. The related spacings are depicted in the legend. The shaded rectangles on the bottom show the characteristic energy regions related to three sodium storage modes in graphitic carbon electrode. (f) Schematic illustration of “adsorption-insertion” mechanism for sodium storage in hard carbon. Reproduced with permission [117] Copyright 2017, Wiley-VCH.

materials without heteroatomic doping through carbonizing the cellulose precursor at temperatures of 900, 1100, 1300, and 1500 °C. And they systematically evaluated the correlation between the electrochemical characteristics and structural evolution of hard carbon, employing the combination of in-situ X-ray diffraction mapping, ex-situ nuclear magnetic resonance,

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electron paramagnetic resonance, electrochemical techniques and computational simulations. Based on the systematic experimental results, eventually they established the “adsorption-insertion” mechanism for Na+ ions storage in hard carbons. During the initial sodiation stage, Na+ ions adsorb on the defective sites of hard carbon with a wide adsorption energy distribution, resulting in the sloping voltage profile. In the following stage, Na+ ions are intercalated into carbon interlayers with right spacing to generate a NaCx compound, which is similar to the intercalation process of Li+ ions into graphite, giving rise to the flat plateau at low voltage region. Mitlin and co-workers also confirmed this mechanism [118]. They investigated the electrochemical sodium behaviors of hard carbons through carbonizing peat moss at different temperatures. When the carbonizing temperature rises, the graphitic domains would grow up and the graphitization degree would be improved, meanwhile, the carbon interlayer distance would be decreased (Figure  1.5a). And it was discovered that the amount of micropores was gradually decreased and the number of mesopores/macropores was increased with elevating the carbonizing temperature (Figure 1.5c, d). From the sodiation curves (Figure 1.5e, f), it was found that the hard carbon obtained at higher temperature had a larger capacity in low potential region, which indicated that the low potential region should be associated with the intercalation of Na+ into carbon interlayers, instead of storage in micropore or nanovoid. Moreover, the sodiation curve of the activated carbon with lots of micropores presented no plateau in the low potential region (Figure 1.5e, f), further confirming that the low potential plateau should be not related to the storage of Na+ in the micropores. In addition, they utilized the ex-situ XRD (Figure 1.5g) to investigate the structure evolution of the pyrolytic hard carbon and activated carbon at the sodiation potential of 0.0–0.2 V. With the dropping of the sodiation potential, it was observed that the peak (002) of the pyrolytic hard carbon moved towards the low diffraction angle, suggesting the expansion of carbon interlayer distance (Figure 1.5h). Nonetheless, the shift of peak (002) was not seen in the XRD patterns of the activated carbon. Thus, it can be deduced that storage of Na+ in the low potential plateau region might be ascribed to the intercalation of Na+ into the carbon interlayers, rather than the micropores. Ghimbeu et al., [113] also reported the Na+ ions storage mechanism in hard carbon and the discrimination between the porosity, surface functional groups and defects (Figure 1.6). In order to systematically investigate the correlation between porosity, surface chemistry/defects and the Na storage mechanism of hard carbons, they synthesized a series of hard carbon materials by pyrolyzing cellulose at different temperatures. N2, Kr

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Advanced Battery Materials

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Figure 1.5 (a) XRD patterns of the activated samples (CPM-A) obtained at different temperatures and commercial activated carbon (CAC), the activated samples prepared at 600, 900, 1100, and 1400 °C are denoted as CPM-600-A, CPM-900-A, CPM-1100-A, and CPM-1400-A, respectively. (b) Raman spectra of several CPM-A samples and CAC samples. (c) Nitrogen adsorption-desorption isotherms and (d) pore size distribution of several CPM-A samples. (e) The charge-discharge profiles of several CPM-A and CAC electrodes. (f) The capacity distribution in sloping and plateau region of CPM-A and CAC electrodes. Data are obtained from 10th cycle at a current density of 50 mA g-1 in both e) and f). (g) XRD patterns of the CPM-1400 and CAC electrodes at different discharged potential of A) 0.2 V, B) 0.1 V, C) 0.05 V, and D) 0.001 V. (h) The relationship between the average interlayer spacing and discharge potential, which are attained from g). Reproduced with permission [118] Copyright 2013, American Chemical Society.

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Carbon Anode Materials for Sodium-Ion Batteries

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T1.0 V), resulting from the high content of heteroatoms (N and O). Increasing the temperature from 800 °C to 950 °C, the heteroatoms were gradually removed, so the capacities of high potential region were declined. Note that, a short plateau was observed at the low potential region (0.1 V) of CNFs-950, which should be associated with the formation of some tiny graphitic microdoamins. Obviously, the potential-specific capacity curve of stage II CNFs (Figure 1.16b) is composed of a sloping region (1.0–0.1V) and a plateau region (≈ 0.1 V). For stage III CNFs with high graphitized degree (Figure 1.9c), two long plateau regions at around 0.5 and 0.1V are seen in the initial sodiation curve, and only a long plateau region is found in the subsequent sodiation process, no sloping region like that of the stage II CNFs is found in this stage.

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Figure 1.9 (a,b,c) A series of charge-discharge curves of hard carbons carbonized at various temperatures and the corresponding schematic illustrations of sodium storage mechanisms (bottom). In-situ XRD patterns of CNF-2200 in d) SIBs and (e) LIBs during the first discharge process. Reproduced with permission.121 Copyright 2016, Wiley-VCH. Ex-situ TEM images of HCT1300 electrodes before (f) and after (g) sodiation. (h) The schematic of sodium storage mechanism for cotton derived hard microtubes. Reproduced with permission. [122] Copyright 2016, Wiley-VCH.

The plateau at 0.5 V should be ascribed to the generation of solid electrolyte interface (SEI) film. To further unfold the nature of plateau region at 0.1 V, the in-situ XRD analysis was also utilized to monitor the change of CNF-2200 during the first sodiation process (Figure 1.9d), it was discovered that location of the (002) peak (2θ = 25.5°) is not changed in the entire sodiation process, revealing that Na+ ions are not intercalated into the carbon interlayers.

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Differently, for LIBs, the (002) peak was obviously shifted to the direction of low diffraction angle in the low potential plateau region, and the (002) peak can be recovered to the original location after the desodiation, suggesting that the reversible intercalation/extraction between the graphite interlayers. Through evaluating electrochemical behaviors of a series of hard carbons with tailored microstructures and compositions, a reasonable sodium storage mechanism was proposed. The sodium storage ways can be divided into three types: (1) Na+ uptake in defects derived from the introduction of heteroatoms, which happens over 1.0 V, (2) Na+ adsorption on the randomly dispersed graphene sheets, which is related to the sloping region between 1.0 and 0.1 V, (3) Na+ filling into pores, which is associated with the plateaus region at ≈ 0.1 V. The study of Hu and co-workers confirmed that Na+ ions were not inserted into carbon interlayers as well [122]. They prepared hard carbon microtubes (HCT) through carbonizing cottons and utilized the ex-situ TEM to investigate the microstructure change of the HCT before (Figure 1.9f) and after (Figure 1.9g) sodiation. It was observed that the interlayer distance of sodiated HCT (0.404 nm) is similar to that of HCT, which indicates that Na+ ions are not inserted into the interlayer. Interestingly, it was discovered that the edges of graphitic nanodomains and nanovoids of the HCT become illdefined after sodiation, deducing that some Na+ ions are stored in these sites. They attributed the sloping region to the adsorption of Na+ on defects, edges, and surface of graphitic nanodomains, and ascribed the plateau region to the filling of Na+ into nanovoids. The sodium storage mechanism of HCT proposed by Hu et al., is shown in Figure 1.9h.

1.3 1.3.1

Carbon Anode Materials for Advanced SodiumIon Batteries Graphite Materials

Although the graphite cannot be directly employed as anode material for SIBs in the common way, researchers have not given up investigating the sodium storage of graphite, and a series of solutions were proposed to enhance the sodium storage performance of graphite. Increasing the interlayer distance is an effective approach to promote the intercalation of Na+, it has been found that graphite oxide and expanded graphite show improved sodium storage activity,123 Figure 1.10a displays the schematic illustration of Na+ storage in modified graphite-based materials. Chou and co-workers [124] have demonstrated that the reduced graphene oxide with an enlarged interlayer spacing of 0.371 nm delivered a reversible specific capacity of 174.3 mAh g-1 at the

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Figure 1.10 (a) The schematic illustration of Na+ ions storage in graphite-based materials, involving graphite, graphite oxide and expanded graphite. Reproduced with permission.123 Copyright 2014, Nature Publishing Group. (b) Cyclic voltammetry curves of several Na-solvent systems, 1 mol L-1 NaPF6 in these solvents are utilized as electrolyte. (c) The molecular structure models for several categories of solvents. (e) The schematic illustration of the influence of solvents on the co-intercalation behavior of Na-solvent into graphite. PC, EC, DEC, DMC, THF, DOL, TEGDME, DEGDME, and DME are the abbreviations of propylene carbonate, ethylene carbonate, diethyl carbonate, dimethyl carbonate, tetrahydrofuran, dioxolane, tetraethylene glycol dimethyl ether, diethylene glycol dimethyl ether, and ethylene glycol dimethylether, respectively. Reproduced with permission. [106] Copyright 2016, Wiley-VCH.

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current density of 40 mAh g-1, which is much higher than that of graphite (31 mAh g-1). And Wang et al., [123] found that the expanded graphite with a larger interlayer distance of 0.43 nm exhibited good sodium storage performances, a high reversible specific capacity of 284 mAh g-1 was attained at a current density of 20 mA g-1, and this material could run stably for 2000 cycles. This expanded graphite was fabricated from reducing graphene oxide (GO), and the GO was prepared through a special treating process, which is different from the traditional Hummer's method, the ultrasonic treatment was deliberately omitted when the graphene oxide suspension was obtained from the exfoliation of graphite oxide, promoting the conversion from ordered graphite to expanded one. Furthermore, it has been also verified that Na+ ions and some special solvent molecules are able to be co-intercalated into the graphite interlayer. With the help of specific solvents, more Na+ ions can be inserted into graphite, it seems to be an attractive idea to realize the effective sodium storage of conventional graphite materials [125–129]. It is verified that the diglyme is one of these specific solvents [125], a ternary intercalation compound of Na(diglyme)2C20 would be formed during the sodiation process of commercialized graphite, and a reversible specific capacity of ≈100 mAh g-1 was observed after 1000 cycles at a current density of 37 mA g-1, the corresponding initial coulombic efficiency can reach 90%. In LIBs, the co-intercalated solvent is often unstable and tends to be decomposed, which would cause the exfoliation of graphite, weakening the cycle performances. As a consequence, the co-intercalation of ions and solvent is typically considered as a terrible problem. Nevertheless, this phenomenon has not been observed in SIBs, indicating that the co-intercalation may become an available way to realize the sodium storage of graphite. Kang et al., evaluated the sodium storage behavior in natural graphite utilizing ether-based electrolytes [128]. The electrochemical and ex-situ analyses confirmed that the sodium storage mechanism of natural graphite is the co-intercalation of Na+-solvent combined with pseudocapacitive behavior. When cycled at 100 mA g-1, it exhibited an initial reversible specific capacity of 150 mAh g-1, and a capacity of 127 mAh g-1 can also be retained after 300 cycles. Increasing the current density to 500 mA g-1, a reversible capacity of 125 mAh g-1 was still achieved, and the corresponding capacity retention can reach 80% after 2500 cycles. Notably, the natural graphite still displayed a reversible capacity of 75 mAh g-1 at the high current density of 10 A g-1. Kang and co-workers also investigated the interaction among alkali metal ions, solvent molecules and graphite to further unfold the internal chemistry of the Na+ intercalation into graphite [106]. It was discovered that the largely repulsive local interaction between Na+

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and the graphene layer would lead to the unstable formation of Na-GICs, since the co-intercalation of Na+ ions and solvent molecules can avoid the direct interaction of Na+ and graphene, thus the repulsive interaction can be remarkably suppressed by the co-intercalation behavior. They systematically examined the electrochemical behaviors (Figure 1.10b) of a series of solvents (Figure  1.10c), involving linear/cyclic ethers and carbonates, finding that linear ethers and their derivatives can be co-intercalated into graphite, and no intercalation behaviors of cyclic ethers and carbonates were observed, the related electrochemical behaviors were presented by cyclic voltammetry curves (Figure  1.10b). It was also discovered that the reversible co-intercalation ability depends on the solvation energy of Na and chemical stability of Na-solvent complexes. The co-intercalation behavior of Na-solvent into graphite is schematically illustrated in Figure 1.10d, benefiting from the high energy of Na and good chemical stability of the corresponding co-intercalated complexes, Na+ ions and linear ethers show superior reversible co-intercalation property. Accordingly, the low solvent energy would result in a fact that the related Na-solvent complex is unstable and would easily desolvate, leading to the failed cointercalation. In addition, if the chemical stability of the Na/solvent-GICs is poor, the GICs would be inclined to decompose with the release of gas, giving rise to the exfoliation of graphite. Therefore, to realize the reversible co-intercalation, both of the high solvent energy and good chemical stability of the Na/solvent-GICs are indispensable. Although the co-intercalation seems to be an available way to improve the sodium storage performances of graphite, challenges and problems still remain. First of all, the co-intercalation of solvent molecules would inevitably consume electrolyte solvent, triggering the increase of resistance and demand of excess electrolyte. Secondly, the specific capacity is still low and the intercalation voltage is relatively high, which would weaken the energy density of the full battery system. Moreover, the co-intercalation behavior would give rise to large volume expansion (≈350%), bringing about the pulverization of graphite particles in the repetitive intercalation-extraction cycles, leading to the degradation of cycling performance. [127]

1.3.2 Hard Carbon Materials The conventional graphite anode shows poor electrochemical sodium storage performances in the common way, however, numerous amorphous have presented praiseworthy electrochemical properties in SIBs. According to the successful experiences of carbon materials in LIBs, a variety of carbon materials with diversiform micro/nanostructures and functional groups

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have been designed and constructed to improve the specific capacity, rate capability and cycle ability, including carbon spheres, carbon fibers, and carbon sheets, etc. In 2012, Maier and co-workers [130] investigated the sodium storage performances of hollow carbon nanospheres (HCNSs, Figure  1.11a,b), which are achieved by employing poly(styrene) latexes and D-Glucose as the carbon precursor. The related electrochemical properties are shown in Figure 1.11c , d. When cycled at the current density of 50 mA g-1, the HCNSs electrode delivered initial sodiation and desodiation capacities of 537 and 223 mAh g-1, respectively, corresponding to a coulombic efficiency of ≈ 41.5%, and the reversible specific capacity could be remained at 160 mAh g-1 after 100 cycles. The good rate capability was also observed (Figure  1.2d), the reversible capacities of 75 and 50 mAh g-1 could be attained at high current densities of 5 and 10 A g-1, respectively. In the same year, Cao and co-workers [116] prepared hollow nanowire (HCNWs, Figure 1.11e–g) by pyrolyzing polyaniline and evaluated its electrochemical behaviors as anode material for SIBs, exhibiting good electrochemical performances (Figure  1.11h,i). At the current density of 50 mA g-1, the HCNWs electrode delivered an initial reversible capacity of 251 mAh g-1, and the capacity retention can reach 82.2% after 400 cycles. Increasing the current density to 500 mA g-1, a capacity of 149.9 mAh g-1 could be still obtained. In comparison with cotemporaneous carbon anode materials, the above-mentioned hollow nanostructures presented preeminent electrochemical properties, which may depend on their unique structures: (1) the hollow interior can generate more active sites and increase the electrode/electrolyte contact area, and the nanosized thin wall can mostly shorten the Na+ ions diffusion path; (2) the enlarged interlayer distance of these two materials (0.40 and 0.37 nm) can be favor of the effective intercalation and transfer of Na+ ions. Guo et al., constructed a sandwich-like hierarchical porous carbon/ graphene composite (G@HPC, Figure  1.12a–d) to further improve the sodium storage performances of amorphous carbon materials [131]. In the obtained composite, porous carbon with an enlarged interlayer spacing of ≈ 0.42 nm was uniformly dispersed on both sides of the graphene. The sandwiched graphene can serve as the fast electron transfer path, the hierarchical porous structure can facilitate the fast diffusion of Na+ ions, and the enlarged interlayer distance can be beneficial to the Na+ insertion, the corresponding schematic is shown in Figure 1.12e , f. Owing to the synergistic effect of these desirable features, this composite electrode exhibited outstanding electrochemical sodium storage performances (Figure 1.12g–i). A high reversible capacity of 400 mAh g-1 was observed

Advanced Battery Materials

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Figure 1.11 (a) TEM image, (b) HRTEM image, (c) cycle performance, and (d) rate property of HCNSs. Reproduced with permission.130 Copyright 2012, Wiley-VCH. (e) SEM image, (f) TEM image, (g) HRTEM image, (h) cycle performance, and (i) rate property of HCNWs. Reproduced with permission. [116] Copyright 2012, American Chemical Society.

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Figure 1.12 (a,b) SEM) images, (c) TEM image, and (d) HRTEM image of the G@HPC. (e) The schematic illustration of the detailed structure of G@HPC. (f) The schematic illustration of Na+ storage in the G@HPC. (g) The charge-discharge curves of G@HPC at the current density of 50 mA g-1. (h, i) Cycle properties of G@HPC at the current densities of 50 mA g-1 and 1000 mA g-1. Reproduced with permission. [131] Copyright 2014, Wiley-VCH.

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Figure 1.13 (a) SEM image, (b) TEM image, and (c) HRTEM image of the 3D PCFs. (d, f, g) Cycle performances at different current densities and (e) rate property of 3D PCFs electrodes. Reproduced with permission. [132] Copyright 2015, Wiley-VCH.

at a current density of 50 mA g-1. And at a higher current density of 1000 mA g-1, the composite electrode can still deliver a high reversible capacity of 250 mAh g-1 after 1000 cycles, no obvious fading was observed. Ji et al., [132] reported a new method to fabricate 3D porous carbon frameworks (PCFs, Figure 1.13a–c) by calcining the carbon quantum dots

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with abundant oxygenic functional groups at high temperature. As anode material for SIBs, 3D PCFs presented high specific capacity and ultralong cycle life (Figure 1.13d–g). The 3D PCFs electrode delivered a reversible capacity of 356.1 mAh g-1 at a current density of 100 mA g-1, even at high current densities of 10000 and 20000 mA g-1, it can still have considerable capacities of 104 and 90 mAh g-1, respectively. Excitingly, the 3D PCFs electrode presented extremely superior cycle stability, and it can be stably cycled at a high current density of 5000 mA g-1 for 10000 cycles. Such excellent electrochemical performances are closed linked to its unique structural features: 3D structure with high specific surface area is favor of the electrolyte permeation and Na+ ions transportation, the thin nanosheets and expanded interlayer distance of 0.42 nm can offer more active sites. Additionally, the porous frameworks can effectively accommodate the volume changes during the repeated sodiation-desodiation process and improve the capacitive storage of Na+ ions. Li and coworkers [133] reported the preparation of reduced graphene nanowires on three-dimensional graphene foam (3DGNW) through a facile template strategy including self-assembly of graphene sheets, elimination of polystyrene spheres template and heating activation. Owing to the extensively lateral exposed edges/pores, expanded graphene interlayer spacing and low sodiation plateau between graphene and graphite, the 3DGNW showed superior electrochemical performances. As anode material for SIBs, the 3DGNW electrode delivered a reversible capacity of over 301 mAh g-1without obvious capacity decay after 1000 cycles at the current density 372 mA g-1. Even at a high current density of rate of 7440 mA g-1, a high reversible capacity of 200 mAh g-1 can be still maintained. Zhang et al., [134] synthesized porous carbon nanotubes (CNTs) by eliminating MoO2 nanoparticles of the MoO2@C nanofibers obtained from a single-needle electrospinning method (Figure 1.14a–d). As anode for SIBs (Figure 1.14e–h), the obtained porous carbon nanotubes exhibited superior cycle performance and rate capability. At 50 mA g-1, it displayed a high capacity of 503 mAh g-1. And increasing the current densities to 1000 and 5000 mA g-1, it delivered capacities of ≈130 mAh g-1 after 1000 cycles and 110 mA h g-1 after 1200 cycles, respectively. Li et al., utilized a combined experimental and theoretical approach to investigate the solid electrolyte interface (SEI) films formed on the surface of hard carbon anode LIBs and SIBs [135]. It was presented that the SEI properties could be tuned by the electrolyte additive or preforming an SEI layer on the surface of the electrode, and the stable SEI film could be constructed by precycling the electrode in the suitable electrolytes for Na+ or Li+. And it was demonstrated that the in-situ formed SEI films are

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Figure 1.14 (a) The SEM image of MoO2@C nanofibers annealing at 850 °C. (b) The SEM image of porous CNTs after thermal treatment with nitric acid. (c) TEM and (d) HRTEM images of porous CNTs. (e) Rate capability of porous CNT electrode. (f-h) Cycle performances of porous CNTs electrodes at different current densities. Reproduced with permission. [134] Copyright 2017, Wiley-VCH.

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stable and highly selective, they produced the selective Li- and Na-based SEI films employing proper Li- or Na-based electrolytes. Note that Li+ ions can easily pass through Na-based SEI film, while Na+ ions cannot transit the Li-based SEI film. Na+ ions storage can be manipulated by modifying the SEI film with film-forming electrolyte additives or preforming an SEI film on the electrode surface. It was demonstrated that the additive of fluoroethylene carbonate (FEC) can control the intercalation of Na+ in hard carbon by tuning the SEI film, but it cannot affect the intercalation of Li+.

1.3.3 Heteroatom-Doped Carbon Materials In addition to the construction of morphology, altering the electronic and chemical structure of the carbon materials by heteroatom-doping is considered as an effective approach to boost the electrochemical performances of carbon materials. For carbon anode materials of SIBs, introducing heteroatoms (like B, O, N, S, and P) has become a quite promising method to tailor the surface wettability, charge transfer, electronic conductivity, electrode/electrolyte interactions and specific capacity [136–138]. A large number of heteroatom-doped carbon materials have been used as anode materials for SIBs, and lots of important achievement have been made. Hitherto, nitrogen is the most common heteroatom to modify the carbon materials, which can generate extrinsic defects and alter the electron distribution around carbon atoms to enhance the reaction activity and electronic conductivity. In 2013, Huang co-workers reported the sodium storage behaviors of N-doped interconnected carbon nanofibers from carbonizing polypyrrole for the first time [139]. At the current density of 200 mA g-1, it delivered a reversible capacity of 134.2 mAh g-1 after 200 cycles, and further increasing the current density to 20000 mA g-1, the reversible capacity can still be remained at 73 mA h g-1. In the following several years, more and more N-doped carbon anode materials for SIBs were developed. In 2015, Dai et al., [140] designed a 3D N-doped graphene foam with a high N content of 6.8 at% (Figure 1.15a–d), when used as anode material for SIBs, it delivered a quite high reversible capacity of 1057.1 mAh g-1 at a current density of 100 mA g-1 (Figure 1.15f). Even at 500 mA g-1 (Figure 1.15h), it can still exhibit a high reversible capacity of 594 mAh g-1 after 150 cycles, corresponding to capacity retention of 69.7%. Compared with sodium storage behaviors of graphene oxide foams, reduced graphene and N-doped graphene, it can be deduced that both N-doping and foamlike morphological structure can effectively enhance the electrochemical performances, and the synergistic effect of these two points gave rise to the splendid sodium storage performances of the 3D N-doped graphene foam.

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Figure 1.15 (a,b) SEM images, (c) TEM image, and (d) HRTEM image of the N-doped graphene foams (N-GF), the selected area electron diffraction (SAED) pattern is shown in the inset of d). (e) Initial charge-discharge profiles and (f) rate properties of the N-doped graphene (N-G) electrode, reduced graphene foams (rGF) electrode and the N-GF electrode at the rates of 0.2 C (1 C = 500 mA g-1), Scapacity = specific capacity. (g) Capacity retention of N-G, rGF, and N-GF electrodes at different C-rates to the initial capacity at 0.2 C. h) Cycle performances and Coulombic efficiencies of the N-G, rGF, and N-GF electrodes at the rates of 1 C. Reproduced with permission. [140] Copyright 2015, Wiley-VCH.

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Figure 1.16 (a,b) SEM images, (c) TEM image, (d) HRTEM image, (e) high-angle annular dark-field (HAADF) image and corresponding elemental mapping of N-CNFs, the inset of d) depicts the corresponding SAED pattern. (f, g) SEM images, (h) TEM image, (i) HRTEM image (the inset shows the related SAED pattern), (j) HAADF image and the corresponding elemental mapping of N-CNFs after 3800 cycles. (k) Cycle performance of the N-CNFs electrode at a current density of 0.1 A g-1, (l) rate capability of the N-CNFs electrode at various current densities and (m) long-term cycle performance of the N-CNFs electrode at a high current density of 5 A g-1. Reproduced with permission. [141] Copyright 2016, Wiley-VCH.

In 2016, Lou’s group reported a free-standing N-doped carbon nanofiber films (N-CNFs) anode for SIBs, which were obtained through electrospinning and carbonizing [141]. As shown in Figure 1.16a–e, the N-CNFs

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exhibit a morphology of 3D carbon fiber network and have multiscale pores and high N content (7.15 wt%). The remarkable electrochemical performances of the freestanding N-CNFs electrode are shown in Figure 1.16k–m. At the high current density of 5000 mA g-1, a reversible capacity of 210 mAh g-1 is kept after 7000 cycles and the corresponding capacity retention is as high as 99%. Note that the designed freestanding electrode can still present a high specific capacity of 154 mAh g-1 at a very large current density of 15 A g-1, which is nearly the highest level of carbon anode materials in SIBs. The superb electrochemical performances must result from the unique internal structure, a series of characterization of the tested electrode after 3800 cycles (Figure 1.16f–j) confirmed the high structural durability and mechanical flexibility of the free-standing N-doped carbon nanofiber films, which ensure the superior performances. In 2018, Zhao and coworkers reported the preparation of N-doped hard carbon nanoshells (N-GCNs, Figure 1.17) with homogeneous defective nano graphitized domains through the pre-chelation between Ni2+ and chitosan and with simultaneous N-doping and KOH-activation [142]. The as-prepared N-GCNs present highly ordered structure (ID/IG = 0.85), expanded interlayer spacing (0.347–0.400 nm) of (002) plane, and successful N-doping (1.5 at%). The obtained N-GCNs show 3D interconnected nanoshell architecture with diameter sizes from 200 to 600 nm and open hollow characteristic of the nanoshell with the wall thickness of ≈20–50 nm, and the specific surface area of N-GCNs reaches a high value fo1490 m2 g-1. The turbostratic carbon structure was confirmed by the HRTEM images, it can be seen that the graphitized microdomains are separated by some nanovoids, leading to the discontinuous graphitic structure, and some carbon layers are criss-crossed with each other, generating some defects (edges and pores). As anode material for SIBs, the N-GCNs electrode exhibited a reversible specific capacity of 325 mA h g-1, which was kept at 174 mA h g-1 after 200 cycles. When used as anode material for LIBs, the N-GCNs electrode deliver a quite high reversible specific capacity of 1253 mAh g-1 and outstand rate capacity (175 mAh g-1 at the high current density of 20000 mA g-1). The density functional theory (DFT) calculations demonstrated the synergistic effect between N-doping and nanopore defects for ions storage. And the in-situ Raman analysis confirmed that the sodium storage process of N-GCNs is adsorption and the lithium storage process of N-GCNs is adsorption-intercalation mechanism. Sulphur is also the common doping heteroatom, different from the N doping, the sulfur doped in the carbon has electrochemical activity for sodium storage and can provide more sodium storage sites, leading to the improved specific capacity [143]. Additionally, introducing the S atoms

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Figure 1.17 (a) Raman spectra, (b) XPS spectra (N1s), (c) porosity distribution, (d, e, g) SEM images, (f, h, i) TEM images, and (j-l) HRTEM images of N-GCNs. (m) Cycle performance of the N-GCNs electrode at a current density of 0.1 A g-1. (n, o) Chargedischarge profiles of N-GCNs electrodes at different cycles and current densities. (p) rate performance of N-GCNs electrode. Reproduced with permission. [142] Copyright 2018, Wiley-VCH.

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Advanced Battery Materials

with large radius can enlarge the carbon interlayer spacing, which can facilitate the intercalation of Na+ ions and enhance the electrode kinetics. Therefore, the S-doping has been considered be an effective way to promote the sodium storage performances of carbon by lots of researchers. Chen and co-workers [144] investigated the electrochemical sodium storage behaviors of sulfur covalently bonded graphene (SG), the SG electrode exhibited a reversible specific capacity of 291 mAh g-1 at a current density of 50 mA g-1 and good cycle stability over 200 cycles at 1000 mA g-1, deriving from the alteration of electronic structure in graphene sheets. Afterwards, Jiang and co-workers [145] evaluated the sodium storage properties of S-doped disordered carbon with a quite high S content of 26.91 wt%, which was prepared by calcining the mixture of 1,4,5,8-Naphthalenetetracarboxylic dianhydride and sulfur. As anode material for SIBs, this S-doped disordered carbon presented a high reversible capacity of 516 mAh g-1 at the current density of 20 mA g-1, and increasing the current density to 1000 mA g-1, its specific capacity can be still maintained at 271 mAh g-1 after 1000 cycles, the related capacity retention is 85.9%. Simultaneously, Huang and co-workers [137] also explored the effects of S-doping on the sodium storage performances of carbon (Figure 1.18). The S-doped carbon with S content of 15.17 wt% and an enlarged carbon interlayer distance of 0.386 nm was obtained from carbonizing the poly(3,4-ethylenedioxythiophene) (PEDOT). When used as anode material for SIBs, at the current density of 500 mA g-1, it delivered a specific capacity of 303.2 mAh g-1 after 700 cycles, and the initial Coulombic efficiency could reach 73.6%. Recently, phosphorus has been also doped into carbon materials to tailor the electrochemical sodium storage behaviors. Ji and co-workers [146] have obtained P-doped carbon nanosheets (P-CNSs, Figure 1.19) with large area and enlarged interlayer distance of 0.42 nm through calcining the mixture of carbon dots and NaH2PO4 under Ar atmosphere. The electrochemical tests indicated that the P-CNSs electrode had excellent sodium storage properties with high specific capacity, outstanding rate capability and cycle stability. It delivered a high reversible specific capacity of 328 mAh g-1 at a current density of 100 mA g-1, and at higher current density of 20000 mA g-1, the reversible specific capacity can be still kept at 108 mAh g-1. Better yet, at the current density of 5 A g-1, the P-CNSs electrode exhibited a stable capacity of 149 mAh g-1 after 5000 cycles. The enhanced electrochemical properties should derive from the collective effect of following aspects. a) The ultrathin nanosheet with large surface area can offer more active sites and shorten Na+ diffusion distance. b) The enlarged interlayer spacing can effectively facilitate the Na+ of insertion/extraction. c) The modified electronic state of carbon

Carbon Anode Materials for Sodium-Ion Batteries

37

2 nm

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(b)

(a)

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Figure 1.18 (a) SEM image, (b) HRTEM image, and (c) high-resolution XPS spectra of S2p of the S-doped carbon. Inset of c) is the schematic model of the S-doped carbon. (d) Rate property and (e) cycle performance of the S-doped carbon electrode. Reproduced with permission. [137] Copyright 2015, Wiley-VCH.

resulting from the P-doping may be favor of the adsorption of electrolyte ions. Cui and coworkers [138] constructed the graphene-based phosphorusdoped carbon (GPC) through annealing the mixture of triphenylphosphine and graphite oxide under an Ar atmosphere (Figure 1.20). And it was demonstrated that the P atoms are successfully introduced into fewlayer graphene with the formation of P-O and P-C bonds. When used as anode material for SIBs, at the current density of 50 mA g-1, the GPC exhibited a charge capacity of 284.8 mAh g-1 after 60 cycles. At the current

38

Advanced Battery Materials

Figure 1.19 (a) TEM image, (b) HRTEM image, (c) SEM image, and (d) elemental mappings of the P-CNSs, (e-g, i) cycle performances and (h) rate capability of the P-CNSs electrode. Reproduced with permission. [146] Copyright 2017, Wiley-VCH.

Carbon Anode Materials for Sodium-Ion Batteries

39

Hummer’s

Graphite

GO

TPP:triphenyl phosphine

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GPC

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Figure 1.20 (a) The schematic illustration of the preparation process of GPC by annealing the mixture of triphenylphosphine (TPP) and GO. (b) SEM image, (c) TEM image, and (d) HRTEM image of GPC, the inset of (c) is the corresponding SAED pattern. (c) Cycle property of the GPC electrode at the current density of 50 mA g-1. (d) Rate capability of the GPC electrode at various current densities. (e) Long-term cycle capability of GPC electrode at the high current density of 500 mA g-1. Reproduced with permission. [138] Copyright 2017, Wiley-VCH.

40

Advanced Battery Materials

density of 500 mA g-1, a charge capacity of 145.6 mAh g-1 was obtained after 600 cycles. They attributed the enhanced electrochemical sodium performances to the enlarged interlayer distance, abundance of defects, vacancies, and active sites induced by the doping of P atoms. Boron with the lower electronegativity and analogous atomic size to carbon could be served as an effective dopant as well. Wang et al., [147] found that the interlayer distance of reduced graphene oxide (rGO) can be broadened by the “pillars” originated from the strong chemical bonds between boron and oxygen, and introducing boron could generate abundant defects, giving rise to increasing sodium storage sites. As a consequence, at a current density of 20 mA g-1, the boron functionalized rGO presented a capacity of 280 mAh g-1, much higher than that of pure rGO, and in good accordance with the results of first-principle calculations offered by Ling and Mizuno [148]. Besides, codoped carbons with several different heteroatoms have been also developed to enhance the electrochemical storage properties. Considering that N doping can improve the electronic conductivity, the S or P-doping can induce the expansion of carbon interlayer distance and create additional active sites, it seems that the N, S-codoping and N, P-codoping could be a nice approach to further promote the sodium storage performances of carbon materials [149, 150]. Zhang and co-workers [151] designed the hierarchical N/S-codoped carbon microspheres (NSC-SP) through carbonizing cellulose/polyaniline, the designed NSC-SP electrode exhibited an excellent cycle stability with a stable capacity of ≈150 mAh g-1 for 3400 cycles, it is much better than that of undoped/N-doped/S-doped carbon microspheres. Firstprinciple calculations were employed to reveal the effect of the doping, interlayer spacing and defects on the electronic conductivity, adsorption energy and diffusion barrier of Na+ ions, which suggested that the N/Scodoping, enlarged interlayer spacing and defects have positive effects on the electronic conductivity of the carbon materials and adsorption stability and mobility of Na+ ions. Zhou et al., [152] synthesized an S/N-codoped carbon (S-N/C, Figure 1.21b–e) through utilizing S atoms to replace N atoms of N-rich carbon nanosheets (N-C) via the gas-solid reaction (Figure 1.21a). Displacing the pyrrolic-N and reserving the pyridinic-N can improve the utilization of S owing to the formation of strong C-S-C bonds, which can enlarge the interlayer distance and generate more active sites. And it was confirmed that the S was mainly introduced through replacing the pyrrolic-N with S atoms. It was discovered that the introduction of S atoms could increase the specific surface area, enlarge the interlayer distance and distort the

Carbon Anode Materials for Sodium-Ion Batteries

N N

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Figure 1.21 (a) The schematic of synthetic process of S/N-codoped carbon (S-N/C). (b) SEM image, (c) TEM image, (d) HRTEM image, and (e) elemental mappings of the obtained S-N/C. (f) The calculated interlayer distance of graphite, N-rich carbon (N-C) and S-N/C. The cyclic voltammetry curves of (g) S-N/C electrode and (h) N-C electrode. (i) Long-term cycle properties of S-N/C electrode and N-C electrode. Reproduced with permission. [152] Copyright 2016, Wiley-VCH.

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Advanced Battery Materials

carbon structure. The effect of S-doping was evaluated by using the density functional theory (DFT) computations, which indicated that the calculated d values of N-C and S-N/C are 3.47 and 3.73 Å (Figure 1.21f), respectively, agreeing well with the TEM and XRD results. Compared with the N-doped carbon, a couple of additional cathodic/anodic peaks (1.05/1.85 V) was observed in the cyclic voltammetry curves (Figure 1.21g,h) of S/Ncodoped carbon, and it is also quite different from cathodic/anodic peaks in the typical Na-S battery, suggesting that the bonds of C-S-C are not broken and reconstructed in the insertion/extraction process Na+ ions. Note that the Faradaic reaction between Na+ and S atoms bonded to the carbon can further improve the capacity to some extent. In consequence, the S-N/C electrode presented excellent electrochemical properties. When tested at current densities of 50 and 10000 mA g-1, the S-N/C electrode can show reversible capacities of 350 and 110 mAh g-1, respectively. At the current density of 1000 mA g-1 (Figure 1.21i), a stable capacity of 211 mAh g-1 can be kept for 1000 cycles. Liu and coworkers [153] prepared a 3D hierarchical N/S co-doped carbon by simply calcining the mixture of thiourea and citric acid at 700 °C. When used as anode material for SIBs, at the current density of 500 mA g-1, the reversible capacity can be kept at 260 mAh g-1 after 1000 cycles. And even increasing the current density to 10000 mA g-1, it can still deliver a capacity of 172 mAh g-1. In comparison with the mono-doped carbon (contrast samples of single N-doped and S-doped carbon), the co-doped carbon showed larger capacity, better cycle stability and higher rate capability. And they calculated the physicochemical properties of the doped carbon according to first-principle calculation based on dispersion-corrected density functional theory, indicating that the co-doped model with a lower diffusion barrier and proper adsorption energy can improve the electronegativity and enhance the spin density, which are advantageous for the absorption and transmission of Na+ ions. In the similar way, Li et al., [154] constructed N/P-codoped carbon microspheres through hydrothermal synthesis and thermal treatment. When tested as anode material, the N/P-codoped carbon showed excellent sodium storage properties. It delivered a reversible capacity of 305 and 136 mAh g-1 at current densities of 100 and 5000 mA g-1, respectively. Yu et al., [155] explored the electrochemical sodium storage performances of N/O-codoped carbon (NOC) network (Figure 1.22a,b), which was obtained from carbonizing the mixture of bacterial cellulose and polyaniline and activating with KOH. The schematic is shown in Figure 1.22c, it can be seen that the interconnected carbon nanofibers network can serve as the fast electron transportation channels and the dense nanopores in

Carbon Anode Materials for Sodium-Ion Batteries

43

100 nm

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Figure 1.22 (a) SEM and (b) TEM images of the N/O-codoped carbon (NOC) network. (c) The schematic illustration of the Na+ storage and electron transfer process in the NOC network. (d) Cycle performance of the NOC electrode at the current density of 100 mA g-1. (e) Long-term cycle performance of the NOC electrode at a high current density of 2000 mA g-1 with an activation at a low current density of 100 mA g-1 for the first five cycles. (f) Rate property of the NOC electrode. Reproduced with permission. [55] Copyright 2016, Wiley-VCH.

carbon nanofibers can generate abundant sodium storage active sites. And the N/O-codoping could enhance the electrical conductivity and surface wettability of carbon material. Moreover, the 3D network structure with high specific surface area can be favor of the electrode/electrolyte reaction and shorten the diffusion distance of Na+ ions. Thus, the NOC network electrode presented remarkable sodium storage properties with high

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Advanced Battery Materials

capacity and cycle stability. At a current density of 100 mA g-1, a high capacity of 545 mAh g-1 can be attained after 100 cycles (Figure  1.22d), and at a higher current density of 2000 mA g-1, it could still exhibit a capacity of 240 mAh g-1 after 2000 cycles (Figure 1.22e). Nonetheless, its initial coulombic efficiency was relatively low (only ~30%) (Figure 1.22d), which is mainly caused by its high specific surface area (1426.1 m2 g-1). Yu and coworkers [156] also designed a self-supported B/N-codoped 3D interconnected carbon nanofibers thin film (BN-CNFs, Figure  1.23) by fully infiltrating NH4HB4O7·H2O into the bacterial cellulose pellicle and carbonizing. The prepared BN-CNFs present a porous cross-linked 3D network structure, an expanded interlayer spacing of 0.44 nm and a large specific surface of 1585 m2 g-1. Since the porous and 3D interconnected structure can facilitate the Na+ migration, the dual-atom doping can not only create defective sites but also enhance the electrochemical activity and electronic conductivity, and the robust 3D interconnected network structure can accommodate the volume change during Na+ repeated insertion/extraction, the obtained BN-CNFs electrode exhibited superior electrochemical performances with high capacity, excellent rate capacity and long cycle life. At the current density of 100 mA g-1, it delivered a high specific reversible capacity of 581 mAh g-1 after 120 cycles, and increasing the current density to 10000 mA g-1, a high capacity of 277 mAh g-1 can be still attained after 1000 cycles. In addition, they also further investigated the synergistic effect of N and B codoping utilizing the first-principles DFT calculations. Huang et al., [157] designed four kinds of oxygen-rich carbon (ORC) materials through two-step synthetic methods: pyrolysis of potassium citrate followed by hydrothermal treatment (Figure 1.24). The precursor of ORC was obtained by calcining potassium citrate at 750 °C for 1 hr under an Ar atmosphere (marked as PC750). And then the PC750 was modified by dilute nitric acid to introduce oxygenous functional groups on the carbon surface at 90 °C for various times of 1.5, 3, 6, and 12 hr, the corresponding samples were named as ORC-1.5, ORC-3, ORC-6, and ORC12, respectively. The PC750 presents a porous structure with amorphous texture, which is comprised of cross-linked and curly carbon nanosheet with a thickness of 10–20 nm and an interlayer distance of 0.38 nm. The ORC materials inherited the architecture characteristics of PC750, no obvious morphology change was observed after the functionalization of nitric acid. But the interlayer distances of these samples were discriminating, the functionalization of nitric acid resulted in the expanded interlayer distances of ORC-1.5, ORC-3, ORC-6, and ORC-12, which were 0.38 nm 0.42 nm, 0.40 nm, and 0.39 nm, respectively. It can be seen that the

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Carbon Anode Materials for Sodium-Ion Batteries

Figure 1.23 (a) The synthesis process of BN-CNFs. SEM images of (b) bacterial cellulose (BC) treated with NH4HB4O7·H2O solution and (c) the BN-CNFs. (d) TEM and (e) HRTEM images of BN-CNFs. (f) Cycle property of the BN-CNFs electrode at a current density of 100 mA g-1. (g) Rate capability of the BN-CNFs electrode at different current densities. (h) Long-term cycle performance of the BN-CNFs electrode at a high current density of 10 A g-1. Reproduced with permission. [156] Copyright 2017, Wiley-VCH.

interlayer distance decreases with the extension of reaction time when the reaction time is over 3 hr, which maybe result from the fact that the hydrothermal reaction could facilitate the graphitization of carbon and cause the shrink of carbon layers. And the oxygen content presents an increasing trend with the prolonged hydrothermal reaction time, for PC750, ORC1.5, ORC-3, ORC-6, and ORC-12, the oxygen contents are 8.4, 16.2, 17.5,

46

Advanced Battery Materials + o– oK

o o –o K+

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Figure 1.24 (a) The schematic illustration of the preparation process of oxygen-rich carbon (ORC). (b) SEM and element mapping images (inset) of PC750, (c) TEM and (d) HRTEM images of PC750. (e) SEM and element mapping images (inset) of ORC-3, (f) TEM and g) HRTEM images of ORC-3. h) The charge-discharge curves of the different samples at the current density of 200 mA g-1, (i) rate performances of PC750 and ORC samples. Long-term cycle property of the ORC-3 electrodes at (j) 2 A g-1 and (k) 5 A g-1. (l) The schematic diagrams of sodium storage mechanism in PC750, ORC-3, and ORC-12 electrode. Reproduced with permission. [157] Copyright 2017, Wiley-VCH.

Carbon Anode Materials for Sodium-Ion Batteries

47

19.8, and 20.2 at%, respectively. It was discovered that the ORC-3 exhibited the highest sodium storage capacity (447 mAh g-1 at 200 mA g-1) and best rate capability (172 mAh g-1 at 20 A g-1) among the five samples, and with the formation and increase of oxygen functional groups on the carbon surface, the initial coulomb efficiencies gradually increase from 45.6% (PC750) to 81.5% (ORC-12), suggesting that the redox reactions between Na+ and oxygenous functional groups are highly reversible. According to the equation of i = k1ν + k2ν1/2 and cyclic voltammograms at 1 mV s-1, the calculated results indicated that the surface capacitive contributions for the capacities of PC750, ORC-1.5, ORC-3, ORC-6, and ORC-12 are 39.6%, 41.0%, 42.6%, 43.0%, and 43.8% at 1 mV s-1, respectively, which demonstrated that the surface capacitive contribution would be enhanced with the increase of oxygen functional groups content. Although the surface capacitive behavior is fast and reversible, in comparison with the ORC-3, the ORC-12 presents a lower capacity, which should be due to the higher graphitizing degree of ORC-12 with smaller interlayer distance, limiting the electrochemical intercalation of Na+ ions. Finally, they proposed that an excellent sodium storage property cannot be obtained if the capacity only comes from the contribution diffusion-controlled intercalation reaction or surface-induced capacitive reaction, the ideal electrode material should combine the fast surface reaction with the reversible intercalation reaction.

1.3.4 Biomass Derived Carbon Materials As is well known, the elemental carbon, hydrogen and oxygen are the major composition of biomass materials, thus, utilizing them as precursors to produce carbon materials is practicable. In general, carbonization is the main way to obtain carbon materials from biomass materials, during the carbonizing process, several gases (CO, CH4, H2 and H2O, etc.) would be released and then the residual would be transformed into the amorphous carbons [158]. There is no doubt that the biomass materials are abundant, green, renewable and eco-friendly, and hence exploiting these sustainable materials meets the requirement of social development. Commonly, the main procedure of producing carbon materials from biomass materials consists of the following steps: (1) washing with water, (2) drying, and (3) pyrolyzing at high temperature under inert atmosphere. In the past decades, a large number of biomass derived carbon materials have been achieved and widely applied into energy storage systems [159].

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Advanced Battery Materials

Recently, with the rapid development of SIBs, more and more biomass derived carbon materials are utilized as anode materials for SIBs [158]. In 2013, Mitlin and coworkers successfully obtained the disordered carbon nanosheet frameworks with abundant micro/mesopores by pyrolyzating peat moss [118], and the corresponding interlayer distance of the attained carbon is 0.388 nm, which is larger than that of conventional graphite (0.335 nm). When tested at the current density of 50 mA g-1, its initial sodiation and desodiation capacities were 532 and 306 mAh g-1, respectively, and the reversible capacity was 298 mAh g-1 after 10 cycles, then increasing the current density to 100 mA g-1, the reversible capacity can be kept at 255 mAh g-1 after 200 cycles. Whereafter, Mitlin et al., also prepared amorphous carbon with low specific surface area by carbonizing banana peels [160], which delivered a reversible capacity of 336 mAh g-1 at the current density of 100 mA g-1, and only a capacity loss 11% was observed after 300 cycles. At a higher current density of 500 mA g-1, the reversible capacity is maintained at 221 mAh g-1, and only 7% of the capacity was lost after 600 cycles. Hu and coworkers investigated the sodium storage performances of a series of biomass derived hard carbon materials. In 2014, Hu et al., fabricated the monodisperse hard carbon spheres (HCS) utilizing sucrose as carbon sources [161]. It was discovered that the HCS coated with a soft carbon layer delivered an initial coulombic efficiency of 83%, which was much higher than that of HCS (54%). Successively, they evaluated the sodium storage behaviors of hard carbon microtubes (HCT, Figure 1.25b) attained by pyrolyzating natural cotton (Figure  1.25a) at different temperatures of 1000 °C (HCT1000), 1300 °C (HCT1300), 1600 °C (HCT1600) [122]. It was found that the carbonization temperature has a significant influence on the electrochemical properties of the hard carbon microtubes. The HCT1000 electrode shows a low reversible specific capacity of 88 mAh g-1, and the sodiation curve only consists of a sloping region (Figure 1.25c), and the initial Coulombic efficiency is as low as 26%, which may be due to the lacking active sites in the sample carbonized at low temperature. The cycle performances of HCT1000, HCT1300 and HCT1600 are presented in Figure 1.25d. In comparison with the HCT1000, both HCT1300 and HCT1600 have higher reversible capacities (≈300 mAh g-1) and higher initial Coulombic efficiencies (>80%) at the current density of 30 mA g-1. The HCT1300 electrode presents the best comprehensive electrochemical performances with the highest reversible capacity (315 mAh g-1), high initial Coulombic efficiency (83%) and excellent cycle stability. In addition, they also examined the sodium storage properties hard carbon from calcining corn cobs at various temperatures of 1000 °C (HCC1000), 1300 °C (HCC1300) and

Carbon Anode Materials for Sodium-Ion Batteries

49

5 μm

20 μm

100 μm (b) 600 HCT1000 HCT1300 HCT1600

2.0

Specific capacity (mAh g–1)

Voltage (V vs. Na+/Na)

2.5

1.5 1.0 0.5 0.0 0

(c)

100

200

300 –1

Specific capacity (mAh g )

400

(d)

120

500

100 HCT1000 HCT1300 HCT1600

400

80

300

60

200

40

100

20

0

0

20

40 60 Cycle number

80

Coulombic efficiency (%)

(a)

0 100

Figure 1.25 SEM images of (a) cotton and (b) carbonized cotton, the inset is the digital image of cotton. (c) Charge-discharge profiles and (d) cycle properties of three HCT samples carbonized at different temperatures. Reproduced with permission. [122] Copyright 2016, Wiley-VCH.

1600 °C (HCC1600) [162]. Similar to the HCT, the HCC1300 showed the best comprehensive performances, and it exhibited a reversible capacity of 298 mAh g-1 and good cycle stability with high capacity retention of 97% after 100 cycles. Wu and coworkers employed the discarded bio-waste apple to produce hard carbon material and investigated its sodium storage performances [163]. Figure 1.26 gives the main procedure, and this is also a typical process for producing biomass derived carbon materials. Interestingly, the elemental nitrogen and sulphur are found in the obtained hard carbon, which should be from the natural proteins in the apple. This hard carbon electrode showed capacities of 245 mAh g-1 at 20 mA g-1 and 112 mAh g-1 at 1000 mA g-1, and it presented a good long-term cycling stability (1000 cycles). Zhao et al., [164] prepared hard carbon nanoparticles through a flame deposition method, utilizing coconut oil as a biomass precursor. When tested as anode material for SIBs, it delivered a sodiation capacity of 277 mAh g-1 in the second cycle at the current density of 100 mA g-1, and

50

Advanced Battery Materials 72 h 80 ºC

75 h RT

H3PO4

1. Washing 2. Drying 1100 ºC Ar

Ball milling

Figure 1.26 The typical synthetic process of hard carbon from biowaste apple. Reproduced with permission. [163] Copyright 2016, Wiley-VCH.

the sodiation capacity was kept at 217 mAh g-1 after 20 cycles. The oatmeal has been also used to prepare N-doped carbon microspheres (NCSs) electrode for SIBs. When tested at 50 mA g-1, the NCSs electrode exhibited a high capacity of 336 mAh g-1 after 50 cycles. When increasing the current density to 10000 mA g-1, and the capacity can be still maintained at 104 mAh g-1 after 12500 cycles, and no obvious capacity fading was seen [165]. Yan et al., [166] found that biomass byproduct okara can be utilized to prepare N-doped carbon sheets (NDCS), which has a high nitrogen content of 9.89 at%, the byproduct okara served as both carbon and nitrogen source. As anode material for SIBs, the NDCS exhibited a reversible capacity of 292.2 mAh g-1 at a current density of 56.25 mA g-1. When activating at a low current density of 112.5 mA g-1 for initial 100 cycles, the NDCSs electrode presented stable cycle ability for 2000 cycles at the current density of 1.687.5 mA g-1. Note that the NDCS exhibited a high energy density of 146.1 Wh kg-1. Huang and coworkers achieved a porous hard carbon material through carbonizing H3PO4-treated pomelo peels [167], the obtained hard carbon has a linked porous structure and high specific surface area (1272 m2 g-1). It was discovered that this hard carbon surface was modified with oxygen-/ phosphorus-containing functional groups, which can trigger the pseudocapacitive effect from the surface redox reaction. At the current density

Carbon Anode Materials for Sodium-Ion Batteries

(a)

(b)

(c)

Na+

Na+

e+ e+

e+

51

Na+

25 μm

(d)

5 μm

Na+

e+

25 μm

(e)

25 μm

(f) 1st 2nd 5th

2.0 1.5 1.0 0.5

Capacity retention (%)

+ Potential (V vs. Na/Na )

3.0 2.5

0

(h)

100 200 300 400 500 Specific capacity (mAh/g)

(i)

5 nm

150

100

120

80

90

60

60

40

30

20

0

0.0

(g)

0

20

40

60

0 80 100 120 140 160 180 200

Coulombic efficiency (%)

e+ e+

Cycle number

Figure 1.27 Digital images of the leaf (a) before and (b) after carbonization. (c) The SEM image of the cross-section of the carbonized leaf, (d) the SEM image of the carbon nanosheet in the sponge layer. SEM images of (e) the stomata on the back surface and (f) the bricklike grains on the upper surface. (g) The HRTEM image of the carbonized leaf. (h) Charge-discharge profiles for 1st, 2nd, and 5th cycles. (i) The capacity retentions and corresponding Coulombic efficiencies in 200 cycles. Reproduced with permission. [168] Copyright 2016, American Chemical Society.

of 200 mA g-1, the reversible capacity can be maintained at 181 mAh g-1 after 220 cycles. At a higher current density of 5000 mA g-1, a reversible capacity of 71 mAh g-1 was observed. Unfortunately, the initial coulombic efficiency is quite low, only 27%, which may be caused by the large specific surface area and abundant surface functional groups. Hu et al., [168] attained the carbon membrane through a facile one-step pyrolysis of natural oak leaves (Figure 1.27a,b). The obtained carbon membrane has anisotropic surfaces and internal hierarchical pores (Figure 1.27c–g), which can be directly used as a binder-free and current-collector-free anode for SIBs (Figure 1.27h,i). And it showed superior electrochemical properties with a high specific capacity of 360 mAh g-1 and a large initial Coulombic efficiency of 74.8%. Li and coworkers [169] reported an ultrathin mesoporous carbon with low tortuosity prepared from directly carbonizing natural wood, and it was discovered that the mesoporous carbon has abundant

52

Advanced Battery Materials

mesopores in the horizontal direction and straight channels in the vertical direction, originating from the inherent microstructure of natural wood. More remarkable, this mesoporous carbon could be utilized as binder-free and current collector-free anode for SIBs, exhibiting a high capacity of 13.6 mAh cm-2, much higher than that of conventional anode for LIBs (~3.5 mAh cm-2 for graphite). Zheng et al., [170] prepared lamellar hard carbon through hydrothermal treatment and pyrolysis of holly leaves, and the obtained hard carbon material has large pores of tiny isolated graphitized domains and mesopores. It was found that the hydrothermal treatment is the key step to generate the lamellar morphology and enlarged pores, and the product obtained from directly pyrolyzing the holly leaves presented the bulk morphology and nanopores of graphitized domains. Moreover, the temperature of hydrothermal treatment also has a crucial influence on the pore structure. Used for sodium storage, the lamellar carbon anode showed excellent rate property and a high capacity of 318 mAh g-1 at 20 mA g-1. Interestingly, it was demonstrated that this approach can be expanded to produce lamellar hard carbons from betula platy phylla and sophora japonica leaves. Ji and coworkers [171] obtained a porous carbon material through directly carbonizing the dried camphor wood without any activating agent (Figure 1.28). The prepared carbon material has an interconnected carbon network structure with a specific surface area of 678 m2 g-1, and enlarged interlayer distance of 0.39 nm is measured, leading to excellent electrochemical sodium storage performances. At a current density of 100 mA g-1, the prepared carbon electrode delivered a reversible capacity of ≈310 mAh g-1 after 100 cycles. Note that a desirable capacity of 120.6 mAh g-1 was retained after 5000 cycles at a high current density of 5000 mA g-1. Wang et al., [172] reported an effective route involving a pyrolysis process and a reductive strategy to prepare high-performance hard carbon anode materials for SIBs from waste apricot shell. The synthesized hard carbon materials inherited the original architecture of the apricot shells, presenting a well-connected structure and a larger interlayer distance (Figure 1.29a–c), which can facilitate the insertion and transport of Na+ ions. The H2 reduction was employed to enhance the initial Coulombic efficiency (Figure 1.29d). The prepared hard carbon material delivered a desirable reversible capacity of ≈400 mAh g-1 and a high initial Coulombic efficiency of 79%. At the current density of 250 mA g-1, a reversible specific capacity of 250 mAh g-1 was obtained after 500 cycles, corresponding to a high capacity retention of 91.7%. Han and coworkers [173] obtained the hard carbon materials from biowaste mangosteen shell through a simple one-step carbonizing process (Figure 1.29e–h). They investigated the

Carbon Anode Materials for Sodium-Ion Batteries

53

Figure 1.28 (a) The digital image of camphor tree. (b, c) SEM, d, e) TEM, and (f) HRTEM images of the camphor tree derived hard carbon. Cycle performances of the camphor tree derived hard carbon electrodes at (g) 100 and (h) 5000 mA g-1. Reproduced with permission. [171] Copyright 2017, The Electrochemical Society.

influence of carbonizing temperatures and times on the microstructure and electrochemical properties. It was found that the optimal carbonization condition was at the temperature of 1500 °C for 2 hr. The prepared hard carbon material delivered a reversible capacity of ≈330 mAh g-1 at the current density of 20 mA g-1, corresponding to a high initial Coulombic efficiency of ≈330%, and the capacity retention can reach ≈98% after 100 cycles. Besides, it has been reported that peanut shell [174], pinecone [175], alfalfa leaf [176], walnut shell [177], wood cellulose fiber [178], corn stalk [179], sorghum stalk waste [180], grass [181], dandelion [182], cherry

54

Advanced Battery Materials

(a)

(b)

(c)

Dangling bond

OR

O

H2O+RH

H2 reduction

H2O

O

O

O

OR

H2O H2O

H2O+RH

Potential (V vs. Na+/Na)

(d)

(e)

50μm (g)

1st 10th 50th 100th

1.5 1.0 0.5 0.0

10 nm (f)

2.0

0 (h)

100 200 300 400 Specific capacity (mAh g–1)

Figure 1.29 (a) SEM image of the cross-section of the raw apricot shell, (b) SEM and (c) TEM images of the carbonization samples. (d) Possible H2 reduction mechanism in the pyrolysis process. Reproduced with permission.172 Copyright 2018, Elsevier. (e) The digital image of mangosteen, (f) TEM image, (g) SEM image, and (h) charge-discharge curves of the mangosteen shell derived hard carbon. Reproduced with permission. [173] Copyright 2017, American Chemical Society.

petal [183], wheat straw [184], lotus stem [185], kelp [186], pistachio shell [187], durian shell [188], rice husk [189], and rape seed shuck [190], could be employed to produce hard carbon anode for SIBs as well. In consideration of the cost and environment, if the sodium storage properties can reach the considerable level, the biomass (especially the bio-waste) derived carbon materials should be the ideal electrode materials for SIBs. In general, in the large-scale application, the selection of biomass precursors should follow the principles below. First of all, the

Carbon Anode Materials for Sodium-Ion Batteries

55

biomass should not have high value in other fields. For instance, some biomass precursors which are food for humans should not be selected. By comparison, the bio-waste, like banana peels, corn stalk, rice husk, pomelo peels, peanut shells, and natural leaves, are most suitable biomass precursors. Secondly, biomass often has the unique microstructure and chemical composition, which can have significant influence on the structure/ morphology and composition of the final derived carbon materials, and then determining the sodium storage performances. For example, most of biomass precursors consist of various biopolymers, including lignins, crystalline cellulose, hemicelluloses, free sugars, pectins and proteins. The lignin and hemicelluloses with the features of highly cross-linking and nanocrystalline are adverse to the graphitization, thus they are usually used to prepare the hard carbon anode materials for SIBs. And porous carbon materials are often obtained by pyrolyzing lignin, especially pyrolyzing the lignin with impurities [191, 192]. As we know, the elements of O, N, S and P are the important compositions of creatures, thus the biomass materials can be directly used to obtain heteroatoms-doping carbon materials by thermal treatment. The structure of biomass derived carbon materials usually depends on the structural features of the biomass precursors. It is practicable to efficiently construct high-performance carbon electrode materials with scheduled structure and composition through selecting the proper biomass precursor. However, now the relatively low initial Coulombic efficiency and specific capacity still hamper the practical application of biomass derived carbons. Further systematic studies to better unveil the relationship between carbon precursors, preparation/ pyrolysis parameters, microstructure and electrochemical properties are desired, which can offer the available guidelines for the current bottlenecks of the battery performances. Optimizing the biomass precursors and pre-/post-pyrolysis treatments, designing carbon materials with enlarged interlayer distance, low graphitization degree, suitable porosity and functional groups to achieve the optimum performances is the most significant research direction.

1.4

Conclusion and Prospects

Owing to great market demands of LIBs and the limited lithium resource, searching the feasible substitute for LIBs become imperative. Under this background, SIBs have attracted wide interest again due to the abundant sodium resource, and the related researches have increased rapidly. In the past several years, a great deal of researchers have made a

56

Advanced Battery Materials

lot of efforts for exploring suitable electrode materials for SIBs, and a large quantity of materials have been demonstrated to have sodium storage activity, Table 1.1 gives the summary of the sodium storage performances of reported carbon materials. Thereinto, some significant achievements of carbonaceous anode materials with unique structure/texture have been reported, and it was demonstrated that the carbonaceous materials would be a kind of promising sodium storage anode materials. For graphite, it was found that the intercalation of Na+ ions into the graphite interlayers can be realized through the enlarging the interlayer spacing or Na+-solvent molecules co-intercalation behavior. And it was discovered that a large number of amorphous carbon materials have higher specific capacity than that of graphite due to the different sodium storage mechanism. The heteroatom-doping was employed to improve the sodium storage properties of carbon materials. The hard carbon materials derived from the biomass are considered as the promising anode materials for large-scale application due to the abundant resource and low cost. And a series of biomass derived materials were evaluated as anode for SIBs, which exhibit good electrochemical properties. For the sodium storage mechanism of amorphous carbon anode materials, researchers have different opinions, and three mechanisms were proposed based on respective experimental results Recently, there are some significant progresses in the study of SIBs, but the problems and challenges still exist in the practical application of SIBs. In the field of carbon anode materials, the initial coulombic efficiency, rate capability and manufacturing cost still leave much to be desired. The initial coulombic efficiency of most of the reported carbon materials is less than 70% (Table 1.1), which is much lower than the required value of 90% for the practical application. Although several carbon anode materials exhibited the satisfactory battery performance, the complex preparation process and high cost restrict their practical large-scale application. Further exploration of carbonaceous anode materials with better battery performances is still necessary. The material designing should focus on the construction of novel micro/nanostructured materials with abundant reversible sodium storage sites and reduced Na+ diffusion distance. The electrolyte has an important influence on the formation of SEI film, which can directly affect the initial Coulombic efficiency. The optimization of electrolyte composition and selection of the suitable additive are desired to promote the initial Coulombic efficiency. Preparing highly ordered layered structure carbons with appropriate interlayer distance (0.37–0.38 nm) to enhance the sodium storage capacity in the low potential plateau region and fabricating nonporous carbon materials to improve the Coulombic efficiency are needed to meet the practical application. In addition, develop novel carbon anode

Initial C.E.

78%

20%

41.5%

50.5%

58.8%

57.5%

/

Carbon materials

commercially available hard carbon

templated carbon

hollow carbon nanospheres

hollow carbon nanowires

cellulose derived carbon nanofibers

carbonized peat moss

nanocellular carbon foams

137 (300th cycle, 100 mA g-1)

255 (210th cycle, 100 mA g-1)

176 (600th cycle, 200 mA g-1)

≈ 220 (200th cycle, 50 mA g-1)

160 (100th cycle, 100 mA g-1)

120 (400th cycle, 74 mA g-1)

≈ 225 (100th cycle, 25 mA g-1) 2011/ [194] 2012/ [130]

2012/ [116] 2013/ [195] 2013/ [118]

2013/ [196]

≈ 140 (74 mA g-1); ≈ 100 (1.85 A g-1) 168 (200 mA g-1); 100 (2 A g-1); 75 (5 A g-1) 210 (250 mA g-1); 149 (500 mA g-1) 255 (40 mA g-1); 85 (2 A g-1) 250 (200 mA g-1); 106 (2 A g-1); 66 (5 A g-1) 140 (200 mA g-1); 100 (1 A g-1); 50 (5 A g-1)

(Continued)

2011/ [193]

Year/Ref.

/

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Table 1.1 Summary of the sodium storage performances of reported carbon materials.

Carbon Anode Materials for Sodium-Ion Batteries 57

46%

N-doped porous carbon nanofibers

2014/ [198]

259 (200 mA g-1); 157 (1 A g-1); 98 (2 A g-1)

36%

(Continued)

2014/ [199]

2014/ [197]

210 (200 mA g-1); 134 (2 A g-1); 101 (5 A g-1)

243 (100th cycle, 50 mA g-1)

/

2014/ [123]

284 (20 mA g-1); 91 (200 mA g-1)

150 (2000th cycle, 100 mA g-1)

≈ 260 (280th cycle, 50 mA g-1)

≈ 49.53%

expanded graphite

2013/ [136]

2013/ [139]

≈ 190 (200 mA g-1); ≈ 50 (2 A g-1); ≈ 45 (5 A g-1)

Year/Ref.

134.2 (200th cycle, 200 mA 150 (200 mA g-1); g-1) 121 (2 A g-1); 100 (5 A g-1)

carbon nanofibers

41.8%

N-doped interconnected carbon nanofibers

155 (200th cycle, 50 mA g-1)

327.6 (45th cycle, 100 mA g-1)

34.9%

N-doped carbon nanosheets

Cycle performances (mAh Rate capability g-1) (mAh g-1)

N-doped ordered mesoporous carbon /

Initial C.E.

Carbon materials

Table 1.1 Cont.

58 Advanced Battery Materials

Initial C.E.

67.8%

53.5%

/

≈ 38%

≈ 77%

≈ 52%

Carbon materials

banana peel derived pseudographite

free standing porous carbon nanofibers

porous carbon/graphene composite

carbon microspheres

nanoporous hard carbon

natural graphite

Table 1.1 Cont.

2014/ [202] 2015/ [203] 2015/ [128]

/ 149 (100 mA g-1); 80 (1 A g-1) 307 (20 mA g-1); 95 (500 mA g-1) ≈ 145 (200 mA g-1); ≈ 103 (5 A g-1); ≈ 78 (10 A g-1)

250 (1000th cycle, 1 A g-1) 183 (50th cycle, 30 mA g-1)

127 (300th cycle, 100 mA g-1); ≈ 100 (2500th cycle, 500 mA g-1)

289 (100th cycle, 20 mA g-1)

(Continued)

2014/ [201]

2014/ [200]

≈ 300 (50 mA g-1); 164 (2 A g-1); 60 (10 A g-1)

266 (100th cycle, 50 mA g-1); 140 (1000th cycle, 500 mA g-1)

298 (300th cycle, 100 mA g-1)

2014/ [160]

Year/Ref.

290 (200 mA g-1); 100 (2 A g-1); 70 (5 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Carbon Anode Materials for Sodium-Ion Batteries 59

Initial C.E.

≈ 93%

/

42.6%

≈ 30%

72%

52%

Carbon materials

graphite (solvent co-intercalation)

amorphous carbon/graphene composite

N-doped graphene foams

N-doped bamboo-like carbon nanotubes

wood fibre derived carbon

honeycomb carbon bubbles

Table 1.1 Cont.

2015/ [204]

230 (100 mA g-1); ≈ 130 (5 A g-1); 120 (10 A g-1)

209 (400th cycle, 100 mA g-1); 122 (700th cycle, 1 A g-1)

196 (200th cycle, 100 mA g-1)

100 (160th cycle, 500 mA g-1)

2015/ [206]

359 (100 mA g-1); 112 (5 A g-1)

(Continued)

2015/ [178]

2015/ [205] /

270.6 (50 mA g-1); 81 (1 A g-1)

2015/ [140]

2015/ [126]

Year/Ref.

110 (100 mA g-1); 102 (10 A g-1)

594.0 (150th cycle, 500 mA 1057.1 (100 mA g-1) g-1); 244.7 (2 A g-1); 137.7 (5 A g-1)

142.0 (2500th cycle, 500 mA g-1)

110 (6000th cycle, 200 mA g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

60 Advanced Battery Materials

Initial C.E.

34.8%

69.9%

63.2%

73.6%

35.5%

57.4%

3D porous carbon frameworks

hard carbon from polyvinyl chloride nanofibers

S-doped disordered carbon

S-doped carbon

N-doped carbon nanofiber films

S-covalently bonded graphene

Cont.

Carbon materials

Table 1.1

2015/ [145]

2015/ [137] 2015/ [141]

516 (20 mA g-1); 275 (1 A g-1); 211 (2 A g-1) 192.5 (2 A g-1); 119.5 (5 A g-1) 315 (500 mA g-1); 154 (15 A g-1) 262 (100 mA g-1); 161 (1 A g-1); 83 (5 A g-1)

271 (1000th cycle, 1 A g-1)

377 (100th cycle, 100 mA g-1); 210 (7000th cycle, 5 A g-1) 150 (200th cycle, 1 A g-1)

(Continued)

2015/ [144]

2015/ [207]

271 (12 mA g-1); 147 (240 A g-1)

215 (120th cycle, 12 mA g-1)

302.2 (700th cycle, 0.5 A g-1)

2015/ [132]

Year/Ref.

303.2 (100th cycle, 100 mA 290 (200 mA g-1); g-1); 104 (10 A g-1); 147.3 (2500th cycle, 2.5 A 90 (20 A g-1) -1 g ); 99.8 (10000th cycle, 5 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Carbon Anode Materials for Sodium-Ion Batteries 61

Initial C.E.

46.7%

26.7%

83%

51.6%

≈ 58%

48.8%

P-doped carbon nanosheets

hierarchical N/S-codoped carbon

hard carbon microtubes

hard carbon nanoparticles

few-layered graphene (solvent-assisted intercalation)

self-standing S-doped flexible graphene films

Cont.

Carbon materials

Table 1.1

2015/ [122] 2015/ [114] 2016/ [129]

2016/ [209]

275 (150 mA g-1); 180 (300 mA g-1) 270 (50 mA g-1); 45 (2.5 A g-1) 150 (200 mA g-1); 125 (10 A g-1); 100 (30 A g-1) 377 (100 mA g-1); 89 (1 A g-1)

244 (300th cycle, 100 mA g-1)

≈ 115 (8000th cycle, 12 A g-1)

207 (500th cycle, 50 mA g-1)

305 (100th cycle, 30 mA g-1)

(Continued)

2015/ [151]

280 (30 mA g-1); 131 (5 A g-1); ≈ 130 (10 A g-1)

150 (3400th cycle, 500 mA g-1)

2016/ [208]

Year/Ref.

321.2 (100th cycle, 100 mA 269 (200 mA g-1); g-1); 117 (10 A g-1); 159.9 (2500th cycle, 2.5 A 108 (20 A g-1) -1 g ); 108.8 (5000th cycle, 5 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

62 Advanced Battery Materials

Initial C.E.

27.8%

67.3%

43.9%

≈ 48%

≈ 30%

86%

S/N-codoped hollow carbon spheres

microporous spherical carbon

S-doped N-rich carbon nanosheets

P/N-codoped carbon

N/O-codoped carbon networks

corn cobs derived carbon

Cont.

Carbon materials

Table 1.1

≈ 250 (100 mA g-1); 140 (5 A g-1); 110 (10 A g-1)

275 (100th cycle, 60 mA g-1)

545 (100th cycle, 100 mA g-1); 240 (2000th cycle, 2 A g-1)

275 (90th cycle, 100 mA g-1)

≈ 350 (100th cycle, 50 mA g-1); 211 (1000th cycle, 1 A g-1)

2016/ [154]

2016/ [155]

2016/ [162]

305 (100 mA g-1); 189 (1 A g-1); 136 (5 A g-1) 650 (100 mA g-1); 235 (1 A g-1); 161 (5 A g-1) 360 (30 mA g-1); 275 (300 mA g-1); 211 (600 mA g-1)

(Continued)

2016/ [212]

2016/ [211]

2016/ [210]

Year/Ref.

350 (50 mA g-1); 150 (5 A g-1); 110 (10 A g-1)

232 (40th cycle, 20 mA g-1); 223 (20 A g-1); 115 (1000th cycle, 200 mA 67 (1 A g-1) g-1)

180 (200th cycle, 500 mA g-1); 169 (2000th cycle, 500 mA g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Carbon Anode Materials for Sodium-Ion Batteries 63

Initial C.E.

61%

≈ 38%

74.8%

60%

70.4%

58.6%

Carbon materials

apple biowaste derived hard carbon

Oatmeal derived N-doped carbon microspheres

carbonized-leaf membrane

holly leaf derived lamellar carbon

N-doped holey carbon nanosheets

disordered 3D multi-layer graphene

Table 1.1 Cont.

2016/ [214] 2016/ [215]

230 (500 mA g-1); ≈ 105 (1 A g-1)

190 (15th cycle, 15 mA g-1); 190 (15 mA g-1); 100 (500th cycle, 750 mA 118 (750 mA g-1) -1 g )

268 (200th cycle, 250 mA g-1)

(Continued)

2016/ [170]

103 (200 mA g-1)

253 (1000th cycle, 20 mA g-1)

/

2016/ [213]

2016/ [165]

336 (50th cycle, 50 mA g-1); ≈ 330 (50 mA g-1); 104 (12500th cycle, 10 A 102 (10 A g-1) -1 g ) 360 (10 mA g-1); 270 (40 mA g-1)

2016/ [163]

Year/Ref.

230 (80th cycle, 20 mA g-1); 235 (20 mA g-1); 85 (1000th cycle, 1 A g-1) 112 (2 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

64 Advanced Battery Materials

2016/ [218]

2016/ [219]

2016/ [220] 2017/ [190]

159.3 (100th cycle, 100 mA ≈ 230 (50 mA g-1) 120 (1 A g-1) g-1); ≈ 100 (1000th cycle, 500 mA g-1) ≈ 430 (100th cycle, 100 mA 422 (100 mA g-1); g-1); 263 (1 A g-1); -1 223 (1200th cycle, 1 A g ) 210 (2 A g-1) 433 (100 mA g-1); 122 (3.2 A g-1) 196 (25 mA g-1); 62 (2 A g-1); 32 (5 A g-1)

rod-like ordered mesoporous carbons 71.26%

63.1%

53.7%

≈ 80%

O-doped 3D interdigital carbon

polydopamine derived carbon

rape seed shuck derived hard carbon

143 (200th cycle, 100 mA g-1)

508 (1000th cycle, 50 mA g-1)

(Continued)

2016/ [217]

226 (150th cycle, 30 mA g-1)

≈ 200 (200th cycle, 50 mA g-1); 103 (3000th cycle, 500 mA g-1) 254 (30 mA g-1); 212 (150 mA g-1); 162 (300 mA g-1)

82%

pitch and lignin derived carbon

Year/Ref. 2016/ [216]

45%

mesoporous soft carbon

Cycle performances (mAh Rate capability g-1) (mAh g-1) 331 (30 mA g-1); 90 (2 A g-1); 62 (5 A g-1)

Initial C.E.

Cont.

Carbon materials

Table 1.1

Carbon Anode Materials for Sodium-Ion Batteries 65

Initial C.E.

79.6%

67.99%

41%

41.9%

≈ 40%

Carbon materials

carbon nanosheets

porous carbon nanotubes

3D neat porous carbon aerogels

3D hollow porous carbon microspheres

orange peel derived carbon

Table 1.1 Cont.

2017/ [222]

362 (40 mA g-1); 172 (400 mA g-1); 139 (1 A g-1)

156 (100th cycle, 500 mA g-1); 117 (1000th cycle, 1 A g-1)

224 (50 mA g-1); 125 (1 A g-1)

313.8 (100th cycle, 100 mA 303.5 (100 mA g-1); 142.6 (3.2 A g-1); g-1); 115.8 (5 A g-1)

(Continued)

2017/ [224]

2017/ [223]

2017/ [134]

183.3 (750th cycle, 500 mA 503 (50 mA g-1); g-1); 187 (1 A g-1); 131.9 (1000th cycle, 1 A 132(2 A g-1) g-1); 110 (1200th cycle, 5 A g-1) 287 (100th cycle, 50 mA g-1); 154 (1000th cycle, 500 mA g-1)

2017/ [221]

Year/Ref.

367 (60th cycle, 50 mA g-1); 305 (100 mA g-1); 190 (500th cycle, 1 A g-1) 133 (10 A g-1); 112 (20 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

66 Advanced Battery Materials

Initial C.E.

≈ 40%

38.2%

60.3%

56.02%

42.8%

60.9%

B/N dual-doping carbon nanofibers

nitrogen-rich hard carbon

sulfur-doped graphitic carbon nanosheets

sulfur-doped carbon

rose-like N-doped porous carbon

S/N codoped carbon nanosheets

Cont.

Carbon materials

Table 1.1

2017/ [226] 2017/ [149]

253.8 (50 mA g-1); 104 (3.2 A g-1)

383.1 (100th cycle, 100 mA 393.1 (100 mA g-1); g-1); 212.3 (3.2 A g-1) -1 ≈ 178 (5000th cycle, 5 A g ) 184.2 (5 A g-1)

224 (100th cycle, 100 mA g-1)

(Continued)

2017/ [188]

264 (200th cycle, 100 mA g-1); 100 (4500th cycle, 5 A g-1)

345 (50 mA g-1); 155.1 (3.2 A g-1)

321.5 (100th cycle, 100 mA 322.5 (100 mA g-1); g-1); 182.4 (3.2 A g-1) -1 161.8 (5000th cycle, 5 A g )

2017/ [143]

2017/ [225]

427 (20 mA g-1); 277 (100 mA g-1); 162 (1 A g-1)

≈ 204 (1000th cycle, 1 A g-1)

2017/ [156]

Year/Ref.

644 (200 mA g-1); 581(120th cycle, 100 mA -1 g ); 323 (5 A g-1); -1 277 (1000th cycle, 10 A g ) 314 (10 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Carbon Anode Materials for Sodium-Ion Batteries 67

Initial C.E.

≈ 62.5%

≈ 33%

40.9%

≈ 30.6%

Carbon materials

oxygen-rich carbon

nitrogen-doped carbon nanosheets

graphene-based phosphorus-doped carbon

ball-in-ball nitrogen-doped carbon microspheres

Table 1.1 Cont.

2017/ [138]

2017/ [228]

305.5 (50 mA g-1); 284.8 (60th cycle, 50 mA -1 136.8 (2 A g-1); g ); 145.6 (600th cycle, 500 mA 101.4 (5 A g-1) g-1) 104 (1000th cycle, 5 A g-1)

(Continued)

2017/ [227]

291.4 (200th cycle, 100 mA 388 (100 mA g-1); g-1); 209.8 (1 A g-1); 217.1 (200th cycle, 500 mA 162.5 (10 A g-1) g-1); 198.6 (10000th cycle, 10 A g-1)

324 (50 mA g-1); 106 (5 A g-1); 90 (10 A g-1)

2017/ [157]

Year/Ref.

447 (200 mA g-1); 194 (10 A g-1); 172 (20 A g-1)

225 (1000th cycle, 2 A g-1); 164 (1000th cycle, 5 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

68 Advanced Battery Materials

Initial C.E.

39%

79%

70%

≈ 30%

≈ 25%

intertwined nitrogen-doped carbon nanotubes

apricot shell derived carbon

lotus stem derived hard carbon

N-doped hard carbon nanoshells

N/S codoped carbon nanorods

Cont.

Carbon materials

Table 1.1

2018/ [142]

63 (5 A g-1) 350 (100 mA g-1); 204 (5 A g-1)

174 (200th cycle, 100 mA g-1) 230 (3000th cycle, 1 A g-1)

2018/ [230]

2018/ [185]

351 (40 mA g-1); 150 (500 mA g-1)

330 (450th cycle, 100 mA g-1)

338 (300th cycle, 100 mA g-1); 250 (500th cycle, 250 mA g-1)

2018/ [172]

2017/ [229]

Year/Ref.

400 (25 mA g-1); 100 (1 A g-1)

179.1 (100th cycle, 200 mA 185.3 (200 mA g-1); g-1); 115.3 (5 A g-1); 125.5 (1000th cycle, 1 A 91.7 (10 A g-1) -1 g ); 104.2 (1000th cycle, 5 A g-1)

Cycle performances (mAh Rate capability g-1) (mAh g-1)

Carbon Anode Materials for Sodium-Ion Batteries 69

70

Advanced Battery Materials

materials from cheap raw materials through simplified synthetic routes to reduce the processing cost. Constructing the carbon materials with excellent sodium storage performances through an effective synthetic approach with cheap precursor and high yield is the ultimate goal. In conclusion, we believe that carbon materials could be the practical anode materials for commercialized SIBs in the future.

Acknowledgements This work was financially supported by National Postdoctoral Program for Innovative Talents (BX00192), Young Elite Scientists Sponsorship Program by CAST (2017QNRC001), China Postdoctoral Science Foundation (2017M6203552), the National Natural Science Foundation of China (21473258, 21673298 and 51622406), Innovation Mover Program of Central South University (2017CX004), National Key Research and Development Program of China (2017YFB0102000), and Hunan Provincial Science and Technology Plan (2017TP1001).

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192. Titirici, M.M., Thomas, A., Yu, S.-H., Müller, J.-O., Antonietti, M., A direct synthesis of mesoporous carbons with bicontinuous pore morphology from crude plant material by hydrothermal carbonization. Chem. Mater., 19(17), 4205–4212, 2007. 193. Komaba, S., Murata, W., Ishikawa, T., Yabuuchi, N., Ozeki, T., Nakayama, T., et al., Electrochemical Na insertion and solid electrolyte interphase for hard‐ carbon electrodes and application to Na-Ion batteries. Adv. Funct. Mater., 21(20), 3859–3867, 2011. 194. Wenzel, S., Hara, T., Janek, J., Adelhelm, P., Room-temperature sodium-ion batteries: Improving the rate capability of carbon anode materials by templating strategies. Energy Environ. Sci., 4(9), 3342, 2011. 195. Luo, W., Schardt, J., Bommier, C., Wang, B., Razink, J., Simonsen, J., et al., Carbon nanofibers derived from cellulose nanofibers as a long-life anode material for rechargeable sodium-ion batteries. J. Mater. Chem. A, 1(36), 10662, 2013. 196. Shao, Y., Xiao, J., Wang, W., Engelhard, M., Chen, X., Nie, Z., et al., Surfacedriven sodium ion energy storage in nanocellular carbon foams. Nano Lett., 13(8), 3909–3914, 2013. 197. Fu, L., Tang, K., Song, K., van Aken, P.A., Yu, Y., Maier, J., Nitrogen doped porous carbon fibres as anode materials for sodium ion batteries with excellent rate performance. Nanoscale, 6(3), 1384–1389, 2014. 198. Wang, Z., Li, Y., Lv, X.-J., N-doped ordered mesoporous carbon as a high performance anode material in sodium ion batteries at room temperature. RSC Adv., 4(107), 62673–62677, 2014. 199. Liu, Y., Fan, F., Wang, J., Liu, Y., Chen, H., Jungjohann, K.L., et al., In situ transmission electron microscopy study of electrochemical sodiation and potassiation of carbon nanofibers. Nano Lett., 14(6), 3445–3452, 2014. 200. Li, W., Zeng, L., Yang, Z., Gu, L., Wang, J., Liu, X., et al., Free-standing and binder-free sodium-ion electrodes with ultralong cycle life and high rate performance based on porous carbon nanofibers. Nanoscale, 6(2), 693–698, 2014. 201. Yan, Y., Yin, Y.-X., Guo, Y.-G., Wan, L.-J., A sandwich-like hierarchically porous carbon/graphene composite as a high-performance anode material for sodium-ion batteries. Adv. Energy Mater., 4(8), 1301584, 2014. 202. Chen, T., Pan, L., Lu, T., Fu, C., Chua, D.H.C., Sun, Z., Fast synthesis of carbon microspheres via a microwave-assisted reaction for sodium ion batteries. J. Mater. Chem. A, 2(5), 1263–1267, 2014. 203. Prabakar, S.J.R., Jeong, J., Pyo, M., Nanoporous hard carbon anodes for improved electrochemical performance in sodium ion batteries. Electrochim. Acta, 161, 23–31, 2015. 204. Li, S., Qiu, J., Lai, C., Ling, M., Zhao, H., Zhang, S., Surface capacitive contributions: Towards high rate anode materials for sodium ion batteries. Nano Energy, 12, 224–230, 2015.

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205. Li, D., Zhang, L., Chen, H., Ding, L.X., Wang, S., Wang, H., Nitrogen-doped bamboo-like carbon nanotubes: promising anode materials for sodium-ion batteries. Chem. Commun. (Camb.)., 51(89), 16045–16048, 2015. 206. Yang, G., Song, H., Cui, H., Wang, C., Honeycomb in honeycomb carbon bubbles: excellent Li- and Na-storage performances. J. Mater. Chem. A, 3(40), 20065–20072, 2015. 207. Bai, Y., Wang, Z., Wu, C., Xu, R., Wu, F., Liu, Y., et al., Hard carbon originated from polyvinyl chloride nanofibers as high-performance anode material for Na-ion battery. ACS Appl. Mater. Interfaces, 7(9), 5598–5604, 2015. 208. Hou, H., Shao, L., Zhang, Y., Zou, G., Chen, J., Ji, X., Large‐area carbon nanosheets doped with phosphorus: a high‐performance anode material for sodium‐ion batteries. Adv. Sci., 2016. 209. Deng, X., Xie, K., Li, L., Zhou, W., Sunarso, J., Shao, Z., Scalable synthesis of self-standing sulfur-doped flexible graphene films as recyclable anode materials for low-cost sodium-ion batteries. Carbon N Y, 107, 67–73, 2016. 210. Ye, J., Zang, J., Tian, Z., Zheng, M., Dong, Q., Sulfur and nitrogen co-doped hollow carbon spheres for sodium-ion batteries with superior cyclic and rate performance. J. Mater. Chem. A, 4(34), 13223–13227, 2016. 211. Zhou, D., Peer, M., Yang, Z., Pol, V.G., Key, F.D., Jorne, J., et al., Long cycle life microporous spherical carbon anodes for sodium-ion batteries derived from furfuryl alcohol. J. Mater. Chem. A, 4(17), 6271–6275, 2016. 212. Yang, J., Zhou, X., Wu, D., Zhao, X., Zhou, Z., S‐doped N‐rich carbon nanosheets with expanded interlayer distance as anode materials for sodium‐ ion batteries. Adv. Mater. Weinheim., 2016. 213. Li, H., Shen, F., Luo, W., Dai, J., Han, X., Chen, Y., et  al., Carbonized-leaf membrane with anisotropic surfaces for sodium-ion battery. ACS Appl. Mater. Interfaces, 8(3), 2204–2210, 2016. 214. Vadahanambi, S., Chun, H.-H., Jung, K.H., Park, H., Nitrogen doped holey carbon nano-sheets as anodes in sodium ion battery. RSC Adv., 6(44), 38112– 38116, 2016. 215. Smith, K., Parrish, R., Wei, W., Liu, Y., Li, T., Hu, Y.H., et  al., Disordered 3 D multi-layer graphene anode material from CO2 for sodium-ion batteries. ChemSusChem, 9(12), 1397–1402, 2016. 216. Cao, B., Liu, H., Xu, B., Lei, Y., Chen, X., Song, H., Mesoporous soft carbon as an anode material for sodium ion batteries with superior rate and cycling performance. J. Mater. Chem. A, 4(17), 6472–6478, 2016. 217. Li, Y., Hu, Y.-S., Li, H., Chen, L., Huang, X., A superior low-cost amorphous carbon anode made from pitch and lignin for sodium-ion batteries. J. Mater. Chem. A, 4(1), 96–104, 2016. 218. Yu, L., Song, H., Li, Y., Chen, Y., Chen, X., Zhou, J., et al., Rod-like ordered mesoporous carbons with various lengths as anode materials for sodium ion battery. Electrochim. Acta, 218, 285–293, 2016. 219. Fan, L., Lu, B., Reactive oxygen-doped 3D interdigital carbonaceous materials for Li and Na ion batteries. Small, 12(20), 2783–2791, 2016.

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220. Sun, T., Li, Z.-jun., Wang, H.-guo., Bao, D., Meng, F.-lu., Zhang, X.-bo., A biodegradable polydopamine‐erived electrode material for high‐apacity and long‐ife lithium‐on and sodium‐on batteries. Angew. Chem., 128(36), 10820–10824, 2016. 221. Chen, Y., Shi, L., Guo, S., Yuan, Q., Chen, X., Zhou, J., et al., A general strategy towards carbon nanosheets from triblock polymers as high-rate anode materials for lithium and sodium ion batteries. J. Mater. Chem. A, 5(37), 19866–19874, 2017. 222. Chen, Y., Zhang, Z., Lai, Y., Shi, X., Li, J., Chen, X., et al., Self-assembly of 3D neat porous carbon aerogels with NaCl as template and flux for sodium-ion batteries. J. Power Sources, 359, 529–538, 2017. 223. Zou, G., Hou, H., Cao, X., Ge, P., Zhao, G., Yin, D., et al., 3D hollow porous carbon microspheres derived from Mn-MOFs and their electrochemical behavior for sodium storage. J. Mater. Chem. A, 5(45), 23550–23558, 2017. 224. Xiang, J., Lv, W., Mu, C., Zhao, J., Wang, B., Activated hard carbon from orange peel for lithium/sodium ion battery anode with long cycle life. J. Alloys Compd., 701, 870–874, 2017. 225. Gaddam, R.R., Farokh Niaei, A.H., Hankel, M., Searles, D.J., Kumar, N.A., Zhao, X.S., Capacitance-enhanced sodium-ion storage in nitrogen-rich hard carbon. J. Mater. Chem. A, 5(42), 22186–22192, 2017. 226. Zhao, G., Zou, G., Qiu, X., Li, S., Guo, T., Hou, H., et al., Rose-like N-doped porous carbon for advanced sodium storage. Electrochim. Acta, 240, 24–30, 2017. 227. Luo, Z., Zhou, J., Cao, X., Liu, S., Cai, Y., Wang, L., et al., Graphene oxide templated nitrogen-doped carbon nanosheets with superior rate capability for sodium ion batteries. Carbon N Y, 122, 82–91, 2017. 228. Xiong, W., Wang, Z., Zhang, J., Shang, C., Yang, M., He, L., et al., Hierarchical ball-in-ball structured nitrogen-doped carbon microspheres as high performance anode for sodium-ion batteries. Energy Storage Materials, 7, 229–235, 2017. 229. Ding, K., Gao, B., Fu, J., An, W., Song, H., Li, X., et al., Intertwined nitrogendoped carbon nanotubes for high-rate and long-life sodium-ion battery anodes. ChemElectroChem, 4(10), 2542–2546, 2017. 230. Hu, A., Jin, S., Du, Z., Jin, H., Ji, H., NS codoped carbon nanorods as anode materials for high-performance lithium and sodium ion batteries. Journal of Energy Chemistry, 27(1), 203–208, 2018.

2 Lithium Titanate-Based Lithium-Ion Batteries Jiehua Liu*, Xiangfeng Wei and Fancheng Meng Future Energy Laboratory, School of Materials Science and Engineering, Hefei University of Technology, Anhui, China

Abstract Safe fast-charging lithium-ion batteries have huge potential market size on demand according to their shortened charging time for high-power devices. Safety and fastcharging are also two key features for next-generation lithium-ion batteries, which have higher requirements than conventional anode materials (e.g., graphite). In principle, the charging-rate capability of safe lithium-ion batteries depends largely on the performance of anode for lithium storage. Graphite-based anode materials for lithium storage are also accompanied by volume swell as well as high-capacity anode materials, such as micro/nano-silicon, tin dioxide and cobalt-based oxide. However, highcapacity anode materials often exhibit poor initial Coulombic efficiency, unsafety and poor high-rate performance. Lithium titanate, Spinel Li4Ti5O12 (LTO), as an alternative to carbon-based anode candidates, has aroused much attention because of its inherent characteristics, such as the zero volume change during lithiation/delithiation, no lithium dendrites and solid-electrolyte interphase (SEI) layer. These properties make it become one of ideal candidates in the field of safe electric vehicles and electronic devices. However, its poor electronic conductivity and low Li-ion diffusion coefficient lead to serious polarization and poor capability at high charge/discharge rates. In order to improve above properties, there are three developed strategies, which include constructing novel microstructures with shortened transport lengths of electrons and lithium ions, doping with foreign atoms and surface modification. Numerous nanostructured LTO materials, such as nanosheets, nanowires, mesoporous nanocrystalline, hollow spheres, nanotube and nanofiber were demonstrated superior rate performance and long cycle life. This chapter is devoted to the recent progress in the fabrication of LTO with various synthetic strategies and their applications for lithium storage in lithium-ion batteries. These architectures are categorized

*Corresponding author: [email protected] Chunwen Sun (ed.) Advanced Battery Materials, (87–158) © 2019 Scrivener Publishing LLC

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into five styles, i.e., zero-dimensional LTO nanospheres, one-dimensional LTO nanowires, two-dimensional LTO nanosheets, three-dimensional hierarchical LTO and their hybrids. Based on the role in batteries, we further summarised their electrochemical performance of cells with three classified themes including organic half Li-ion cells, organic full Li-ion cells and Na-ion batteries. Finally, the outlook and perspective is concisely provided on LTO based lithium-ion batteries. Keywords: Lithium-ion batteries, anode, lithium titanate, fast charging, safety, full cells

2.1

Introduction

In principle, safe fast-charging lithium-ion batteries (LIBs) depend largely on the performance of anode for lithium storage. Fast charging produces sudden heating of LIBs because of large polarization potential of the anode or/and internal short circuit, which often brings the fire hazard for highpower batteries with a high-cost metallic lithium anode. Compared with the progress of the high-performance cathode materials [1–3], an urgent task is to develop high-performance anode materials. Graphite-based anode materials [4–7] for lithium storage are also accompanied by volume swell as well as high-capacity anode materials, such as micro/nano-silicon [8–12] SiOx,13, 14 cobalt-based oxide/sulfides [15–20] and others [21–24], in which case researchers are desired to synthesize safe stable anode materials for high-power LIBs. Spinel Li4Ti5O12 is an ideal host owing not only to its ‘zero-strain insertion’ structural characteristics, but also to its low cost, abundance and environmental benignity [25–29]. Spinel Li1.33Ti1.67O4 (Li4Ti5O12, LTO) was first reported in 1956 [30], but a detailed structural analysis of single-crystal Li4Ti5O12 was reported by ¯ Deschanvres et al. in 1971 [31]. Li4Ti5O12 is a cubic spinel with a space group and a unit cell parameter (a = 8.352 and 8.370 Å) [32–35]. The tetrahedral 8a sites are occupied by Li+ ions alone, and the octahedral 16d sites are randomly occupied by 1/6 Li+ ions and 5/6 Ti4+ ions. O atoms occupied all of the 32e sites. The other half of the octahedral cation sites in the ccp structure (16c sites) and both the tetrahedral 8b and 48f cationic sites are empty, which are suitable for lithium insertion and extraction. When increasing temperature, the lithium ions in the 8a tetrahedral sites can shift to the 16c octahedral sites, demonstrated by using infrared spectroscopy [36]. The results were supported by Leonidov et al., with the finding that the tetrahedral 8a sites in Li4Ti5O12 were occupied solely by lithium ions based on X-ray and neutron powder diffraction data [34]. Single crystal data of Li4Ti5O12 was obtained by a flux method using LiCl

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Table 2.1 Single crystal data of Li4Ti5O12 Was obtained by a flux method [34]. Structural formula

Li4Ti5O12

Temperature (K)

295

Crystal system

Cubic

Space group

Fd3¯m

Lattice parameters A (Å)

8.352(4)

V (Å [3]

583.6(8)

Z

8

Dx (g/cm3)

3.48

Crystal size (mm)

0.05 × 0.05 × 0.20

Maximum 2θ (deg)

135

Absorption correction

Gaussian integration

Transmission factors: min and max

0.660 and 0.804

Measured reflections

1144

Rint

0.067

Independent reflections

209

Observed reflections (>3σ)

297

Number of variables

8

R

0.036

wR [w = 1/σ2F]

0.033

flux (Table  2.1). Li4Ti5O12 crystallizes in the cubic spinel-type structure, space group ¯ , a = 8.352 (4) Å, V = 583.6 (8) Å3, and Z = 8. There are other many advanced characterization techniques to characterize the phase structure including neutron diffraction, Mossbauer spectrometry, Raman spectroscopy, and nuclear magnetic resonance [37–40]. Ziebarth and co-workers investigate the diffusion barriers for lithium ions in two different crystal structures of LTO using the density functional

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[111]

[111] – [110]

– [110] – [101]

– [101] (a)

(b)

Figure 2.1 (a) Hexagonal supercell of LTO in spinel phase with correct stoichiometry (Li8Ti10O24). The remaining lithium ions occupy the 8a sites. (b) Hexagonal supercell of LTO in rocksalt phase with correct stoichiometry (Li14Ti10O24). Red spheres are oxygen ions (32e), blue spheres are titanium ions (16d), and green spheres are lithium ions. Lithium ions on 16d sites are encircled; the remaining lithium ions occupy the 16c sites. The tripod indicates the directions in the cubic LTO structure by Miller indices [hkl] [41].

theory. Their calculations show that the activation barriers vary between 0.30–0.48 eV for the spinel phase Li4Ti5O12 (Figure  2.1a) and between 0.20–0.51 eV in the lithiated rocksalt phase Li7Ti5O12 (Figure  2.1b). The origins of the rather broad ranges of activation energies are related to different chemical environments of the diffusion channels due to mixed occupancies of some sites in LTO. They suggested that the determination of lithium diffusion constants in LTO cannot be carried out by using a single activation barrier. Instead, the local environment of the diffusion paths must be considered to correctly capture the variety of activation barriers. They also found that the sites have mixed occupation in LTO to trap lithium vacancies in the spinel phase, but this effect is not observed in the rocksalt phase, which explains the low lithium diffusivity found in experiments for lithium concentrations in the vicinity of the spinel phase [41]. In 1989, Colbow et al., first recorded the galvanostatic charge-discharge plateaus and stable cycling performance for Li||Li4Ti5O12 half cells [33]. In 1995, Ohzuku et al., reported the “zero-strain” characteristic based on the no volume change between Li4Ti5O12 and Li7Ti5O12 [32]. The calculated capacity is 175 mA h g−1 when discharged to 1 V using the mass of Li4Ti5O12, which is generally authorized as its theoretical value. Li4Ti5O12 could maximally store 5 Lithium to form Li9Ti5O12 when discharging to 0.01 V with a capacity of 293 mA h g−1, which was limited by the number of tetravalent titanium ions rather than the number of octahedral or

Lithium Titanate-Based Lithium-Ion Batteries 5

LiMn1.5Ni0.5O4 LMO

LCO

4

LNO

L333

1.9 V

2.3 V

2.2 V

2.4 V

3

3.2 V

LFP 2.6 V

Voltage / V

91

2 LTO 1

Figure 2.2 The voltage of a spinel lithium titanium oxide anode and various cathode materials. Reproduced from Ref. 45 with permission from the Royal Society of Chemistry.

tetrahedral sites [42]. Li4Ti5O12 was neglected due to its Li insertion/extraction potential of 1.55 V in lithium batteries, but it can serve as an anode electrode when matches with high-voltage cathodes such as LiMn2O4, LiCoO2, and LiNi0.5Mn1.5O4 to yield safe 2.5–3.2 V lithium-ion cells. 2.5 V LTO batteries were reported by Ferg et al., with different cathodes, such as Li1.03Mnl.97O4 and LiZn0.025Mn1.95O4, and layered LiCoO2 [43]. 5 V LiNi0.5Mn1.5O4 was employed to match Li4Ti5O12 and gave a 3.2 V safe Li-ion cells [44]. Li4Ti5O12-based anode materials can offer a more significant safety than the graphite and lithium anodes, which can be used to develop extremely safe power lithium-ion batteries for HEV or EV applications and storage power station. Various possible candidates have been reported based on different cathodes (Figure 2.2), such as LiFePO4, LiNi0.8Co0.15Al0.05O2, LiFe0.2Mn0.8PO4, LiCoO2, Li1+x(Ni1/3Co1/3Mn1/3)1−xO2, LiMn2O4 (LMO), Li4Ti5O12/LiNi0.5Mn1.5O4, LiCoPO4 (LCP), LiCoMnO4 and so on. [45]

2.2

Benefits of Lithium Titanate

Graphite, which has maximum theoretical specific charge of LiC6 is 372 mA h g−1, was the first thing widely used to replace Li-metal in lithium-ion batteries. Commercial graphite has a low cost, low potential, and general cycling performance [16]. However, its low reversible capacity, large polarization at high rates make the batteries instability and unsafety. Compared with commercial graphite materials, LTO has a high Li+ diffusion coefficient of 2 × 10−8 cm2 s−1, two orders of magnitude larger than graphite

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Table 2.2 Different features of commercial anodes. Anode

LTO

Graphite

Lithium

Voltage /V

1.55

0.05–0.2

0

SEI layer

No

Yes

Yes

Safety

Yes

No

No

Capacity /mA h g−1 175

372

3870

Cost

Low

Low

Expensive

Fast charging

Excellent

General

General

Lithium dendrites

No

Yes

Yes

Cycling performance

>10000

200–1000

200–1000

High power

Excellent

General

General

Volume change

0

10%

Large

Li+ diffusion coefficient /cm2 s−1

2 × 10−8

10−10

-

of 10−10 cm2 s−1. Li4Ti5O12 exhibits unique features including no SEI film, safety, fast charging, no Li dendrites, no volume swelling and high ion conductivity as shown in Table 2.2. However, there are still several challenges for LTO materials and batteries: 1) The inherent insulating characteristic of LTO seriously limits its high-rate capability, which is one of the key parameters to obtain the highpower density of batteries [10, 11]. 2) The associated gas (CO2, H2, CO) evolution of Li4Ti5O12 batteries during the charge/discharge cycles, which leads to serious expansion and further safety issues. 3) Beside rate performance and safety, the energy density of Li4Ti5O12 batteries is are urgently needed to achieve a higher level, because the energy density directly is related to the practical specific capacities of the electrode materials and the potential difference between the cathode and anode of full cells. To solve the above problems, one way is to shorten Li-ion transfer path by nano-architectures form 0D to 3D for increasing the active material/electrolyte interface and shortening the time of Li-ion insertion/extraction. Moreover, the addition of conductive additives could

Lithium Titanate-Based Lithium-Ion Batteries

(a)

(b)

(c)

(d)

93

Figure 2.3 (a) 0D , (b) 1D, (c) 2D and (d) 3D nano-architectures.

be used to improve its surface electronic conductivity for achieving a high-rate capability. The doping method in LTO lattices is also another way to improve the performance of LTO materials and batteries. The strategies will be addressed in details in the following two sections, that is geometrical structures and fabrication of lithium titanate, and modification of lithium titanate

2.3

Geometrical Structures and Fabrication of Lithium Titanate

Different geometrical structures exhibit their unique performance based on surface and structural properties respectively shown in Figure 2.3. 0D structures (ultrafine nanoparticle) possess the large surface area and exhibit excellent shortened pathway in X, Y, Z directions. 1D structures(nanofibers, nanotubes or nanowires), which obtained by oriented growth provide the opportunity to fabricate stable devices, also have the shortened path in Y and Z direction. 2D structures (nanosheets) often have large exposed surfaces and specific facets, but only have a shortened path for lithium-ion transfer in one direction. Most of the 3D structures could be obtained with the aid of structure directing agents (SDAs) or assembled by 0D, 1D and/or 2D structures, which also have the performance of secondary structures.

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We will address in details on the effect of geometrical structures and fabrication methods of LTO materials.

2.3.1

Zero-Dimensional Nano-Architectures

0D nano-architectures have an irresistible attraction for lithium storage due to their exposed surface and the path for Li-ion transfer. The most attractive feature of 0D nano-architectures is the shortened path for lithium insertion and open meso-nanopores holding nanodrops of electrolyte which may reduce electrode polarization when fast lithium storage. Compared with 1 um LTO microsphere, the LTO nanosphere with a diameter of 100 nm only needs the Li storage time as same as 1/1000 of the time for LTO microspheres. Ultrafine nanoparticles, nanosphere, hollow structures are common LTO materials. The cubic Li4Ti5O12 nanoparticles were continuously synthesized in supercritical water using a flow hydrothermal reaction system. [46] Kim and coworkers synthesized Li4Ti5O12 nanoparticles by the precipitation method from ethylene glycol solution of titanium tetra-isopropoxide and Li2O2 by refluxing at 197 °C for 12 hrs [47]. The obtained particles were filtered and dried at 100 °C for 12 hrs, and the dried powder samples were heated at 320, 500 °C and 800 °C for 3 hrs. The X-ray diffraction patterns of the LTO samples exhibited a good fit with the spinel phase. The FESEM images of the dried powder sample and the samples heated at 320, (b) H2O CTAB

TTIP Oleic n-hexane

CTAB Oleic

mix, stir

TiO2 Li2CO3 Li4Ti5O12

100nm

(c)

Carbon

Carbon Hydrolysis Li4Ti5O12 Filter Washing

(a)

+Li2CO3

Argon Annealing 50nm

Figure 2.4 (a) Schematic formation process of microemulsion-assisted ultrafine Li4Ti5O12/C with oleic acid as carbon precursor, (b) and (c) TEM images of ultrafine Li4Ti5O12/C. Reproduced from Ref. 48 with permission from the ScienceDirect.

Lithium Titanate-Based Lithium-Ion Batteries (a)

(e)

(b)

(c)

95

(d)

(f) I

SiO2

II

SiO2@a-TiO2

Hollow Li4Ti5O12

Figure 2.5 FESEM (a) and TEM (b) images of SiO2@a-TiO2. FESEM (c) and TEM (d) images of Li4Ti5O12 hollow spheres. TEM images show an individual Li4Ti5O12 hollow sphere (e). (f) Schematic illustration of the formation of mesoporous Li4Ti5O12 hollow spheres through a templating approach: I) uniform coating of an a-TiO2 shell on the silica sphere through a sol-gel process; II) chemical lithiation in a LiOH solution and subsequent annealing treatment to convert the core-shell sphere into the mesoporous Li4Ti5O12 hollow sphere. Reproduced from Ref. 53 with permission from the Wiley.

500 °C and 800 °C showed a uniform spherical morphology with a particle size of 5, 8, 10 and 400 nm, respectively. The dried powder sample and the samples heated at 320, 500, and 800 °C for 3 hrs showed initial capacities of 200, 310, 320, and 260 mA h g-1, respectively, at a current density of 0.05 mA cm−2. The sample heated at 500 °C for 3 hrs exhibited a high capacity and an excellent rate capability over 60 cycles. The ultrafine Li4Ti5O12/C composite was synthesized by microemulsion with oleic acid as carbon precursor and particle size controller as shown in Figure 2.4 [48]. The as-prepared sample was composed of ultrafine Li4Ti5O12 particles with average size of 25 nm, which are uniformly dispersed in carbon matrix with a carbon content as low as 2.69 wt%. At 10C, the highest discharge capacity reaches 136.3 mA h g−1 and the capacity retention is 96.4% after 100 cycles. The improved performance of the Li4Ti5O12/C composite is attributed to the ultrafine particle size and the uniform composite of Li4Ti5O12 with carbon. To obtain spinel Li4Ti5O12 with smallest possible grain size and highest possible phase purity, Abe et al., synthesized 81–88% phase purity of Li4Ti5O12 with its average grain size ca 600 nm by a solid-state route. After optimizing, They obtained 130 nm Li4Ti5O12 with its 90% pure phase by milling the mixture preliminarily calcined at 500 °C and heating subsequently at 700 °C [49]. Homogeneous Li4Ti5O12 nanoparticles

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were prepared via a sonochemical method, operated at 20 kHz and 220 W for 20 min [50]. These LTO nanoparticles had an average grain size of about 30–40 nm with excellent phase purity in good stoichiometric ratios of Li4Ti5O12. Phosphidated-Li4Ti5O12 was reported by Jo and co-workers, which shows high capacity with a significantly enhanced kinetics opening new possibilities for ultra-fast charge/discharge of lithium rechargeable batteries. [51] A mesoporous Li4Ti5O12/C nanocomposite was fabricated by a nanocasting technique using the porous carbon CMK-3 as a hard template [52]. The modified CMK-3 template is impregnated with the Li4Ti5O12 precursor, followed by heat treatment at 750 °C for 6 hrs under N2. Li4Ti5O12 nanocrystals are successfully synthesized in micrometer-sized porous carbon foam to form a highly conductive network. It exhibited greatly improved electrochemical performance compared with bulk Li4Ti5O12, and shows an excellent rate capability of 73.4 mA h g−1 at 80 °C with significantly enhanced cycling performance (only 5.6% capacity loss after 1000 cycles at a high rate of 20C). The enhanced lithium storage properties of the mesoporous Li4Ti5O12/C nanocomposite was attributed to the interpenetrating conductive carbon network, ordered mesoporous structure, and the small Li4Ti5O12 nanocrystallites, which increased the ionic and electronic conduction throughout the electrode. Lou et al., have developed an efficient templating approach for the preparation of mesoporous Li4Ti5O12 hollow spheres with high quality [53]. FETEM and TEM images show uniform SiO2@a-TiO2 and LTO nanospheres in Figure 2.5a–e. The Li4Ti5O12 hollow spheres possess high uniformity and mesoporous shells with tunable thickness. This approach involves the formation of SiO2@a-TiO2 core–shell particles, and the subsequent chemical lithiation and simultaneous removal of the SiO2 template, which is provided in Figure 2.5f. The as-prepared Li4Ti5O12 hollow structures exhibit remarkable rate capability up to 20C and stable long-term capacity retention for over 300 cycles. More 0D LTO materials and their performances are provided in Table 2.3.

2.3.2

One-Dimensional Nanoarchitectures

There are few reports to obtain 1D Li4Ti5O12 directly. There are three main ways to synthesize 1D LTO materials including ion exchange [61], 1D template [62] and electrospinning methods [63]. Novel spinel Li4Ti5O12 with nanotubes/nanowires morphology and the high surface area has been prepared by a low-temperature hydrothermal lithium-ion exchange processing from hydrogen titanate nanotubes/

Structures

Li4Ti5O12 nanoparticles

Nanoparticlesconstructed LTO

Mg-doped LTO(20–30 nm)

LTO/Cu nanocomposites

Li4Ti5O12/n-Ag composite

No

1

2

3

4

5

Three-step solid-state synthesis

Hydrothermal method

Hydrothermal method

Solid-state method

Solvothermal (polyol) techniques

Method

Table 2.3 The methods and performance of 0d lto based materials.

(Continued)

[58]

[57]

167 mA h g−1 at 1C and 117 mA h g−1 at 20C Less than 6% of capacity loss after 50 charge/ discharge cycles at 1C

[56]

[55]

[54]

Ref.

The discharge capacity remains at 170 mA h g−1 over 100 cycles at 2C; 115 mA h g-1 at 50 C

High reversible capacity of 173 mA h g−1 at 0.5C; 145 mA h g−1 at 30C between 1.0 and 2.5 V

Initial discharge/charge specific capacities of 230 and 179 mA h g−1

Performance

Lithium Titanate-Based Lithium-Ion Batteries 97

Structures

Microspherical Li4Ti5O12/C

MoO2-modified Li4Ti5O12

6

7

Cont.

No

Table 2.3

[60]

116 mA h g−1 at 10C, and 124 mA h g−1 after 450 cycles at 5C.

Solution-based method

[59]

The specific capacity of the electrode at 20C rate is 131 mA h g−1 and the loss of capacity is less than 2% after 60 cycles

Ball-milling, and spraydrying

Ref.

Performance

Method

98 Advanced Battery Materials

Lithium Titanate-Based Lithium-Ion Batteries

99

nanowires precursors [61]. The shape and morphology of spinel Li4Ti5O12 are controllable by varying the hydrogen titanate precursors (nanotube, nanowire, nanorod, and nanobelt) from the alkaline-hydrothermal approach. The formation temperature of spinel Li4Ti5O12 nanotubes/ nanowires is lower than that of bulk materials counterpart prepared by solid-state reaction or by sol-gel processing. The cyclic voltammetric results of both electrodes indicated the enhanced electrochemical kinetics for lithium insertion. Self-supported Li4Ti5O12 nanowire arrays grew directly on Ti foil by a facile solution-based method, further enhancing Li-ion storage properties by creating Ti3+ sites through hydrogenation (Figure 2.6a) [64]. This architecture ensures that every Li4Ti5O12 nanowire participates in the fast electrochemical reaction, enabling remarkable rate performance and a long cycle life. Mesoporous Li4Ti5O12/carbon nanofibers are prepared by a facile electrospinning method combined with soft-template self-assembly (Figure  2.6b) [65]. The diameter of as-prepared LTO/C NFs is approximately 400 nm with highly crystallinity LTO nanoparticle completely embedded in carbon framework. The mesoporous LTO/C NFs exhibited a high specific surface area of 212.1 m2 g−1 and large pore volume. At a current rate of 5C, the reversible capacity of the mesoporous LTO/C NFs electrode is up to 127.4 mA h g−1 and still remains at 122.7 mA h g−1 after 100 cycles. The excellent electrochemical performances are related to welldefined 1D mesoporous nanostructure with LTO nanoparticles embedded in the carbon framework, which could efficiently shorten the path for Li+ diffusion, and enhance electrolyte-active material contact area and facilitate rapid electron transfer. A novel in situ nickel doped 1-D lithium titanate nanofibers (Li4Ti5−xNixO12, where x = 0, 0.05 and 0.1) have been synthesized using a facile electrospinning process as shown in Figure 2.7. Physical characterizations revealed that nickel is homogeneously incorporated into the lattice of lithium titanate nanofibers (LTONFs) which significantly improves their properties yielding outstanding electrochemical performance in a lithium-ion battery at high power rates and a significant reduction in the voltage gap between the oxidation and reduction peaks. A capacity of 190 mA h g−1 has been obtained at 0.2C for the 10% nickel doped nanofibers (Ni-LTONF10), which is higher than the theoretical capacity of lithium titanate bulk (175 mA h g−1). Their results also showed superior rate capability resulting in 63 mA h g−1 obtained at 50C, which is 20 times higher than that of un-doped pristine LTO nanofibers and lithium titanate nanoparticles [66].

100 Advanced Battery Materials

(a) 220 ºC (b) 0.5 M HCI Nano-sized

Large open space

e–1 e–1 e–1

H2TiO5

2O

LiOH solution

e–1 e–1 e–1 Surface conductivity Annealing in Ar/H2

(a)

L-T-O NWAs

H-LTO NWAs Gases

C LTO pores e–

Air

800 ºC, 10 h

350 ºC, 2 h

e–

Li+

As-spun fibers PVP

LTO/C fibers

LTO precursor

LTO

PLTO/C fibers C

Li+

12 nm

Ar/H2

e–

LTO/C

(b)

Figure 2.6 (a) Schematic illustration for the fabrication of H-LTO nanowires[64]; (b) Schematic illustration of the formation of the PLTO/C nanofibers through electrospinning and a subsequent two-step heat-treatment [65].

Park et al., synthesized surface functionalized 1D Li4Ti5O12 nanofibers using simple electro-spinning and subsequent nitridation process [67]. Uniformly coated conducting TiN/TiOxNy layer on the surface of Li4Ti5O12 nanofiber enables fast electron transport along its one-dimensional geometry, which leads to significant improvement in rate capability. Nitridated Li4Ti5O12 nanofibers electrode delivers about 1.35 times larger discharge capacity than that of pristine nanofibers electrode at a current density of 10C. Song et al., designed the kinetically favored Li4Ti5O12 by modifying its crystal structure to improve intrinsic Li diffusivity for high power density in Figure 2.8 [68]. The first-principles calculations revealed

Lithium Titanate-Based Lithium-Ion Batteries 101

720 nm 5 μm

500 nm

(a)

(b)

300 nm

400 nm

200 nm

(c)

(d)

400 nm 790 nm 750 nm 500 nm (e)

300 nm

350 nm

(f)

Figure 2.7 (a and b) SEM images of electrospun LTONFs before heat treatment at low and high magnifications, (c and d) SEM images of the LTONF after heat treatment at low and high magnifications, (e and f) high magnification SEM images of Ni-LTONF10 before and after heat treatment. Reproduced from Ref. 66 with permission from the Royal Society of Chemistry.

that the substituted Na expanded the oxygen framework of Li4Ti5O12 and facilitated Li-ion diffusion in Li4Ti5O12 through 3D high-rate diffusion pathway secured by Na ions. Accordingly, we synthesized sodium-substituted Li4Ti5O12 nanorods having not only a morphological merit from 1D nanostructure engineering but also sodium substitution-induced open framework to attain ultrafast Li diffusion. The new material exhibited an outstanding cycling stability and capacity retention even at a higher current density of 20 C. One-dimensional Ce3+-doped Li4Ti5O12 (Li4Ti5-xCexO12, x = 0, 0.01, 0.02, and 0.05) submicrobelts with the width of approximately 500 nm and

102 Advanced Battery Materials Ion exchange Li+

Na

+ i-O) Na (T

Li+

Li+

Starting material (a)

1D structural formation

Li+

Ti-O + Li+ Li

+ Li+ Li+ Li

Li+

12

-Ti -O

Na+ TiO2

Calcination + Li+

Li+ Li

Li T 4 i O 5

Hydrothermal growth

Synthesis of Li4Ti5O12 Nanorod

Li substitutes for Na Ti

Intensity/ arb. units.

(c)

C Ti

Na Pt 0

2

Li4Ti5O12

20 (b)

2μm

O Ti

40

60

80

100

4

400nm 6

TiO2

120

2 θ/O (d)

Figure 2.8 (a) Schematic representation of the fabrication sequence of NaxLi4–xTi5O12 nanorods. Structural characterization of NaxLi4–xTi5O12nanorods. (b) Rietveldcalculated X-ray diffractograms for NaxLi4–xTi5O12nanorods (inset: EDS results of NaxLi4–xTi5O12 nanorods (Pt signal come from the conducting agent)). (c) Low-resolution TEM micrograph of NaxLi4–xTi5O12nanorods. The inset shows SEM micrographs of NaxLi4–xTi5O12 nanorods. (d) High-resolution TEM image of NaxLi4–xTi5O12 nanorods. Reproduced from Ref. 68 with permission from ACS.

thickness of about 200 nm have been synthesized via the facile electrospinning method. Importantly, 1D Li4Ti5O12 submicrobelts could be well preserved with the introduction of Ce3+ ions, while CeO2 impurity is obtained when x is greater than or equal to 0.02. The comparative experiments prove that Ce3+-doped Li4Ti5O12 electrodes exhibit the brilliant electrochemical performance than the undoped counterpart. Particularly, the reversible capacity of Li4Ti4.98Ce0.02O12 electrode reaches up to 139.9 mA h g-1 and still maintains at 132.6 mA h g-1 even after 100 cycles under the current rate of 4C. The superior lithium storage properties of the Li4Ti4.98Ce0.02O12 electrode could be attributed to their intrinsic structure advantage as well as enhanced overall conductivity [69]. A detail comparisons of methods and performances are also provided in Table 2.4.

Structures

Li4Ti5O12@C hierarchical nanofibers

CNT/ Li4Ti5O12/C nanofibers

LTO/nitrogen-doped porous graphene fibers

Li2CoTi3O8 fibers

No

1

2

3

4

Electrospinning method

High-temperature treatment

Electrospinning and thermal treatment

Electrospinning and annealing

Methods

Table 2.4 The methods and performance of 1d lto based materials.

[72]

[62]

[71]

[70]

Ref.

(Continued)

A specific capacity of 388 mA h g–1 for the first cycle; 237 mA h g-1 after 30 cycles at 50 mA g−1

A high reversible capacity of 164 mA h g–1 at 0.3C); 102 mA h g–1 at 10C, and long cycling stability

Discharge capacity of 119 mA h g–1 after 100 cycles at a current density of 100 mA g(−1)

Discharge capacities of about 145.5 mA h g–1 after 1000 cycles at 10C

Performance

Lithium Titanate-Based Lithium-Ion Batteries 103

Solid-state method

Hydrothermal and calcination method

Li4Ti5O12/carbon/carbon nano-tubes

Li4Ti5O12/TiO2 tubes

6

7

Electrospinning method

Li4Ti5O12 fibers

5

Methods

Structures

Cont.

No

Table 2.4

Initial discharge capacities of 420, 225, and 160 mA h g–1 at 0.01, 0.1, and 1.0 A g−1 respectively.

Discharge capacities of 163 mA h g–1, 148 mA h g–1 and 143 mA h g–1 at rate of 0.5C, 5.0C and 10.0C respectively

Its discharge capacities are 172.4, 168.2, 163.3, 155.9, 138.7, 123.4, 108.8, and 90.4 mA h g0–1 at rates of 0.2, 0.5, 1, 2, 10, 20, 40, and 60C

Performance

[75]

[74]

[73]

Ref.

104 Advanced Battery Materials

Lithium Titanate-Based Lithium-Ion Batteries 105

2.3.3 Two-Dimensional Nanostructures Hydrothermal methods are the welcome routes to synthesize 2D LTO materials. Sheets of Li4Ti5O12 with high crystallinity are coated with nitrogen-doped carbon using a controlled process comprising hydrothermal reaction followed by chemical vapor deposition (CVD) [76]. Acetonitrile vapor is used as carbon and nitrogen source to obtain a thin coating layer of nitrogen-doped carbon. The layer enabled the NC-LTO material to maintain its sheet structure during the high-temperature CVD process and to achieve high crystallinity. Doping with nitrogen introduced defects into the carbon coating layer, and this increased degree of disorder allows fast transportation of lithium ions in the layer. An electrode of NC-LTO synthesized at 700 °C exhibited greatly improved the rate and cycling performance due to a markedly decreased total cell resistance and enhanced Li-ion diffusion coefficient (DLi). Specific capacities of 159.2 and 145.8 mA h g−1 are obtained using the NC-LTO sheets at rates of 1C and 10C, respectively. These values are much higher than values for LTO particles did not undergo the acetonitrile CVD treatment. A capacity retention value as high as 94.7% is achieved for the NC-LTO sheets after 400 cycles in a halfcell at 5C. In our previous work, we successfully synthesized novel wavelike spinel LTO nanosheets using a facile 'co-hydrolysis' method, which is superior to molten-salt approach and traditional solvothermal method in some respects [77]. The unique 2D structures have a single-crystal framework with a shortened path for Li-ion transport. As a result, the N-doped 2D wavelike LTO with 0.6 wt.% of 'carbon joint' not only exhibits exciting capacity of similar to 180 and similar to 150 mA h g−1 for fast lithium storage at high discharge/charge rates of 1.7 and mA h g−1 (10C and 50C) respectively, but also shows excellent low-temperature performance at −20 °C (Figure 2.9). In addition, the cost may be further decreased due to recycled functional reagents. This novel nanostructured 2D LTO anode material makes it possible to develop safe fast-charging high-power lithium-ion batteries. Zero-strain and long-term stability of nanoscale lithium titanate (LTO) anode materials make possible the fabrication of exceptionally stable lithium-ion batteries. But one issue must be considered that of nanostructureinduced relaxation in 2D LTO nanosheets which profoundly modifies their Li storage properties and structural stability. Excessively intercalated Li ions at both 8a and 16c sites trigger nucleation of the relaxed LTO structure in the near-surface region, which impedes Li-ion diffusion and causes the increasing polarization of LTO nanosheet electrodes. Nuclei

106 Advanced Battery Materials

20nm

20nm (a)

(b) Open channel

Electrolyte/LTO interface

Carbon joint (c) 250

120 100

200

80

150

60 100

Charging 40 Discharging 20

50 0

0

30

60 90 Cycle number

(d)

120

Coulombic efficiency /%

Capacity /mA h g–1

50C/8.5 A g–1

0 150

Capacity /mA h g–1

250 200 20 ºC

0 ºC

20 ºC

–10 ºC –20 ºC

150 100 50 0 0

(e)

50

100 Cycle number

150

200

Figure 2.9 TEM images and scheme of single-layer (a) and overlapped LTO nanosheets (b); (c) Schematic formation of ‘carbon joint’ between the overlapped wavelike LTO nanosheets. (d) Cycling performance of the CLTO at a constant current drain of 50C and the corresponding Coulombic efficiency; (e) Cycling performance at a constant current drain of 10C at different temperatures from −20 °C to 20 °C [77].

Lithium Titanate-Based Lithium-Ion Batteries 107 of relaxed LTO then undergo isotropic growth along the 3D Li-ion pathways in LTO to completely convert near-surface regions into relaxed LTO. With increasing population of trapped Li ions, the enhanced conductivity due to Ti4+/Ti3+ reduction gradually eliminates the raised polarization. In the meantime, spontaneous electrolyte/LTO reduction to form the solid electrolyte interphase starts playing a major role in a capacity loss once the transformation of the near-surface region into relaxed LTO becomes saturated. Elucidation of these fundamental intercalation-induced surface structure transformations contributes greatly to the design of highly performing 2D nanoscaled LTO and other electrode materials [78]. Zero-strain and long-term stability of nanoscale lithium titanate (LTO) anode materials make possible the fabrication of exceptionally stable lithium-ion batteries. But one issue must be considered that of nanostructure-induced relaxation in 2D LTO nanosheets which profoundly modifies their Li storage properties and structural stability. Excessively intercalated Li ions at both 8a and 16c sites trigger nucleation of the relaxed LTO structure in the near-surface region, which impedes Li-ion diffusion and causes the increasing polarization of LTO nanosheet electrodes. Nuclei of relaxed LTO then undergo isotropic growth along the 3D Li-ion pathways in LTO to completely convert near-surface regions into relaxed LTO. With increasing population of trapped Li ions, the enhanced conductivity due to Ti4+/Ti3+ reduction gradually eliminates the raised polarization. In the meantime, spontaneous electrolyte/LTO reduction to form the solid electrolyte interphase starts playing a major role in a capacity loss once the transformation of the near-surface region into relaxed LTO becomes saturated. Elucidation of these fundamental intercalation-induced surface structure transformations contributes greatly to the design of highly performing 2D nanoscaled LTO and other electrode materials [78]. Cu-doped Li4Ti5O12–TiO2 nanosheets were synthesized by a facile, cheap, and environmentally friendly solution-based method [79]. These nanostructures were investigated as an anode material for lithium-ion batteries. Cu doping was found to enhance the electron conductivity of the materials, and the amount of Cu doped controlled the crystal structure and content of TiO2. TEM images of as-prepared samples were provided in Figure  2.10. In addition, the samples, which benefit from multiphase and doping, exhibited much-improved capacity, cycle performance, and high rate capability over Cu-free Li4Ti5O12–TiO2. The discharge capacity of the 0.05 Cu-doped sample was 190 mA h g−1 at 1C, and even 144 mA h g−1 was obtained at 30C after 100 cycles. Moreover, after 500 cycles at 30C, the discharge capacity remained at approximately 130 mA h g−1.

108 Advanced Battery Materials (a)

(b)

(c)

(d)

(e)

(f)

(g)

Figure 2.10 TEM images of as-prepared samples: (a) precursor of LTO and (b) precursor of LTO + 0.05 Cu. HRTEM images and fast Fourier transform image patterns of (c) LTO and (d) LTO + 0.05 Cu. (e) Comparison of the surfaces of LTO and LTO + 0.05 Cu. (f) SAED pattern of LTO + 0.05 Cu. (g) Nanosheet model. Reproduced from Ref.79 with permission from Wiley.

Various LTO nanosheet crystals with designated degrees of residual strain were prepared by varying the annealing temperature. The cyclability and rate capability of the prepared 2D LTO nanocrystals were evaluated. During annealing, dehydration of layered lithium titanate hydrate precursors (LTH, (Li2-xHx) Ti2O5-yH2O), leads to the topotactic transformation from a C-centered orthorhombic to a body-centered orthorhombic system. Dehydrated LTH has then converted into spinel LTO nanosheets. Residual strain induced during transformation causes LTO structure distortion rendering Li-O bonding more covalent. Strongly covalent Li-O bonds screen the Li-Li Coulomb repulsion when 8a-16c intercalation occurs, leading to largely irreversible capacity loss and deteriorated cyclability. Meanwhile, rich irreversible Li intercalation degrades the power capability of LTO nanocrystals despite their high surface area. In addition to mechanism elucidation, conditions are identified to eliminate residual strain-induced structural relaxation giving rise to nanosheet anodes with superior power and highly reversible cycling performance. [80] Chen et al., reported a facile synthesis of vertically aligned LTO nanosheet arrays grown directly on Ti foil by hydrothermal growth in LiOH solution. SEM and TEM images of Li4Ti5O12 nanosheet arrays standing on Ti foil were provided in Figure 2.11a–d. When used as a binderfree anode for LIBs, the self-supported LTO nanosheet arrays exhibited an excellent rate capability (a reversible capacity of 163 mA h g−1 and 78 mA h g−1 at 20C and 200C, respectively) and an outstanding cycling performance (a capacity retention of 124 mA h g−1 after 3000 cycles at 50C)

Lithium Titanate-Based Lithium-Ion Batteries 109 (g) 250

(f)

(c)

Capacity/(mAh/g)

5μm

(b)

Current (mA)

0.08 0.04 0.00 −0.04 −0.08 1.00

2μm

200nm

(e) 0.48nm

Capacity/(mAh/g)

(h) 14nm

(111) (111) (200)

1.75

2.00

20C

150

20C 50C

100C

100

200C

50

2.25 2.50

0

0

20

40

100

250 Charge capacity Discharge capacity

200

80 60

150

40

100

20

50 0

5nm

80

300

0

100nm

60

500

100

Cycle Number

Voltage (V) vs Li*/Li

200nm

(d)

1.25 1.50

LTO-NSA-1 LTO-NSA-2 LTO-NS

200

1000

1500

2000

2500

0 3000

Coulombic Efficiency %

(a)

Cycle Number

Figure 2.11 Cross-sectional (a) and top-view (b) SEM images of Li4Ti5O12 nanosheet arrays standing on Ti foil. TEM (c and d) and HRTEM (e) images of large Li4Ti5O12 nanosheets scratched from Ti foil. The inset in (b) is a high-magnification image showing the co-existence of large nanosheets and small nanosheets, and the inset in (d) is the related SAED pattern. Electrochemical properties of Li4Ti5O12 nanosheet arrays (LTONSA) and randomly dispersed Li4Ti5O12 nanosheets (LTO-NS): (f) cyclic voltammetry curve at a scan rate of 0.5 mV s−1 for LTO-NSA-1. (g) Rate performance of different LTO nanosheet structures. (g) Specific capacity and Coulombic efficiency for 3000 cycles at 50C for LTO-NSA-1. Reproduced from Ref. 81 with permission from the Royal Society of Chemistry.

as shown in Figure 2.11f–h. Furthermore, a flexible lithium-ion battery, which could be fully recharged within 30 s and was able to light an LED, was assembled by using the self-supported LTO nanosheet arrays as the anode, demonstrating their potential applications in flexible fast charging electronics [81]. The methods and performance of 2D LTO based materials are summarized in Table 2.5.

2.3.4

Three-Dimensional Nanostructures

Nature inspired researchers to synthesize the flower-like structure assembled by 0D, 1D, 2D and their hybrids with the assistance of SDAs, which is one of the most effective solutions to avoid overlapping. Hierarchical spheres exhibit high surface area and larger open pore for fast lithium diffusion. With the use of ammonium chloride as the pore-forming agent, 3D “fishnet-like” lithium titanate/reduced graphene oxide (LTO/G) composites with hierarchical porous structure are prepared via a gas-foaming method (Figure 2.12a). The nitrogen-sorption analyses reveal the existence of micro-/mesopores, which is attributed to the introduction of NH4Cl into the gap between the graphene sheets that further decomposed into gases

An excellent rate capability due to the thin carbon coating and porous nanosheet structures, The initial discharge capacities of 185.5, 177.2, 167.3, 145.8, 137.7, 127.5 and 112.5 mA h g−1 at the 0.1, 0.5, 1, 2, 5, 10 and 20C rates

LTO nanosheets with an N-doped Hydrothermal process carbon coating

Li4Ti5O12-TiO2 composite

4

Hydrothermal method

[85]

[84]

[83]

(Continued)

The initial discharge capacity of 183 mA h g−1 together with a discharge capacity of 160 mA h g−1 after 100 cycles at 1C; mA h g−1 even after 300 cycles at 10C.

3

Hydrothermal treatment

Li1.81H0.19Ti2O5·xH2O nanosheets

2

[82]

A reversible capacity of 180 mA h g−1 after 200 cycles at 200 mA h g−1; Up to 132 mA h g−1 after 200 cycles at 10,000 mA h g−1 (57C)

Li4Ti5O12 microspheres assembled Hydrothermal process by nanosheets and subsequent thermal treatment

Ref.

Performance

1

Methods

Structures

No

Table 2.5 The methods and performance of 2d lto based materials.

110 Advanced Battery Materials

Hydrothermal method

Hydrothermal process and calcination

Li4Ti5O12 -TiO2 nanosheet arrays

Carbon-coated Li4Ti5O12 nanosheets

6

7

Solution-based method

Cu-doped Li4Ti5O12-TiO2 nanosheets

5

Methods

Structures

No

Table 2.5 Cont.

Li4Ti5O12@C exhibits higher specific capacity, better rate capability and capacity retention than the pristine Li4Ti5O12.

[87]

[86]

[79]

Ref.

(Continued)

The initial discharge capacity of 184.6 mA h g−1 at 200 mA g−1 and possessing excellent electrochemical stability with only 8.3% loss of specific capacity at 1 A g−1 in a prolonged charge-discharge process (1000 cycles).

The discharge capacity of the 0.05 Cu-doped sample was 190 mA h g−1 at 1C, and even 144 mA h g−1 was obtained at 30C after 100 cycles. Moreover, after 500 cycles at 30C, the discharge capacity remained at approximately 130 mA h g−1.

Performance

Lithium Titanate-Based Lithium-Ion Batteries 111

Structures

Ultrathin dual phase nanosheets

8

Cont.

No

Table 2.5

Hydrothermal method

Methods

Discharge capacities of 178.5, 154.9, 148.4, 142.3, 138.2, and 131.4 mA h g−1 at rates of 1, 10, 20, 30, 40, and 50C, respectively; A capacity retention of 93.1% even after 500 cycles at 50C.

Performance

[88]

Ref.

112 Advanced Battery Materials

Lithium Titanate-Based Lithium-Ion Batteries 113

Added GO

Microwave

Dropped NH4CI

sol

Hydrothermal

Solution

Ti(OH)4

Li+

CI+

NH4+

Dried

2O

Thermal treatment

Precursor LTO

in N2

LTO

NH4CI

(a)

Li2CO3

g-C3N4

Annealing

Annealing

Porous CNx Amorphous titanate (b)

N-doped LTO/C hybrid

Figure 2.12 (a) Scheme illustration of the preparation of LTO/G materials89; (b) The synthetic strategy of 3D NCLTO [90].

and produced hierarchical pores during the thermal treatment process. The loose and porous structure of 1-LTO/G composites enabled the better penetration of electrolytes, providing more rapid diffusion channels for lithium ion. As a result, the 1-LTO/G electrode delivered an ultrahigh specific capacity of 176.6 mA h g–1 at a rate of 1C. Even at 3C and 10C, the specific capacity can reach 167.5 and 142.9 mA h g–1, respectively. Moreover,

114 Advanced Battery Materials the 1-LTO/G electrode showed excellent cycle performance with 95.4% capacity retention at 10C after 100 cycles. [89] Nitrogen-doped LTO/C (NCLTO) is firstly synthesized by thermal decomposition of an amorphous titanate cross-linking g-C3N4 inorganic polymer and then reaction with lithium salts utilizing a simple solid-state reaction (Figure 2.12b). The doping nitrogen has a fair distribution which allows fast transportation of lithium ions, and the content of nitrogen is tunable through control the decomposition time of g-C3N4. It exhibits a better high-rate capability and cycling performance than that of LTO, which delivers a high capacity of 122 mA h g−1 after 500 cycles at a current density of 10C with 102% capacity retention while the capacity retention of only 50.5% for pure LTO. Furthermore, the NCLTO also exhibits an overwhelming advantage of high-rate performance than LTO owing to the introduced conductive nitrogen-doped carbon and TiN. [90] “Flower-like” Li4Ti5O12 were synthesized by using a facile and largescale hydrothermal process involving unique Ti foil precursors followed by a short, relatively low-temperature calcination in air. Moreover, a detailed time-dependent growth mechanism and a reasonable reaction scheme were proposed to clearly illustrate and highlight the structural evolution and subsequent formation of this material. Specifically, the resulting “flower-like” Li4Ti5O12 microspheres consisting of thin nanosheets provide for an enhanced surface area and a reduced lithium-ion diffusion distance. The high surface areas of the exposed roughened, thin petal-like component nanosheets are beneficial for the interaction of the electrolyte with Li4Ti5O12, which thereby ultimately provides for improved high-rate performance and favorable charge/discharge dynamics. Electrochemical studies of the as-prepared nanostructured Li4Ti5O12 clearly revealed their promising potential as an enhanced anode material for lithium-ion batteries, as they present both excellent rate capabilities (delivering 148, 141, 137, 123, and 60 mA h g−1 under discharge rates of 0.2, 10, 20, 50, and 100C, at cycles of 50, 55, 60, 65, and 70, respectively) and stable cycling performance (exhibiting a capacity retention of ≈97 % from cycles 10–100, under a discharge rate of 0.2C, and an impressive capacity retention of ≈87 % by using a more rigorous discharge rate of 20C from cycles 101–300) [91]. Hierarchical hollow Li4Ti5O12 microspheres (HLTOMs) assembled by zigzag-like nanosheets were synthesized by hydrothermal treatment of scalable lithium peroxotitanate complex solution using low-cost commercial H2TiO3 particles as titanium sources, followed by a calcination treatment in Figure  2.13. Precursor solution concentration, Li/Ti ratio, hydrothermal temperature, and duration are found correlative and should

Lithium Titanate-Based Lithium-Ion Batteries 115 Hydrothermal treatment

Calcination Li O

Probable Ti-complex TiO6 octahedron

Intermediate hollow spheres

Ti

Li4Ti5O12 hollow spheres

H2O

+ H2O2, LiOH

H2TiO3 particles

Layered hydrous lithium titanate

Li4Ti5O12

Figure 2.13 Schematic illustration of the formation process of HLTOMs [92].

be optimized to obtain pure Li4Ti5O12 products. A high yield of HLTOMs up to 120 g L−1 was achieved. Due to the unique morphology, the HLTOMs deliver an outstanding rate capability of 139, 125 and 108 mA h g−1 at 10, 20 and 30C, respectively, and exhibit 94% capacity retention after 1000 cycles at 30C indicating excellent stability. These values are much superior to those of commercial Li4Ti5O12 particles [92]. Exploring advanced high-rate anodes is of great importance for the development of next-generation high-power lithium-ion batteries (LIBs). Therefore, the novel carbon nanotubes (CNTs)/Li4Ti5O12 core/shell arrays on carbon cloth (CC) as an integrated high-quality anode are constructed via a facile combined chemical vapor deposition–atomic layer deposition (ALD) method (Figure 2.14). ALD-synthesized LTO is strongly anchored on the CNTs' skeleton forming core/shell structures with diameters of 70–80 nm the combined advantages including highly conductive network, large surface area, and strong adhesion are obtained in the CC-LTO@CNTs core/shell arrays. The electrochemical performance of the CC-CNTs/LTO electrode is completely studied as the anode of LIBs and it shows noticeable high-rate capability (a capacity of 169 mA h g−1 at 1C and 112 mA h g−1 at 20C), as well as a stable cycle life with a capacity retention of 86% after 5000 cycles at 10C, which is much better than the CC-LTO counterpart. Meanwhile, excellent cycling stability is also demonstrated for the full cell with the LiFePO4 cathode and CC-CNTs/LTO anode (87% capacity retention after 1500 cycles at 10C) [93].

116 Advanced Battery Materials ALD

(a)

(b)

Lithiation

CC-CNTs/TiO2

CC-CNTs

200 μm

5 μm

(c)

200 nm

(d)

5 μm

(e) C

(h)

CC-CNTs/LTO

(f)

(g) Ti

200 nm

5 μm O

2 μm

Figure 2.14 (a) Fabrication schematics of CC-CNTs/LTO core/shell arrays; SEM images of (b,c) CC-CNTs arrays, (d,e) CC-CNTs/TiO2 arrays, and (f,g) CC-CNTs/LTO arrays; (h) EDX mapping images of C, Ti, and O elements in CC-CNTs/LTO electrode. Reproduced from Ref. 93 with permission from Wiley.

3D ordered macroporous (3DOM) Li4Ti5O12 membrane was prepared by a colloidal crystal templating process. A colloidal crystal consisting of monodisperse polystyrene particles (1 μm diameter) was used as the template for the preparation of macroporous Li4Ti5O12. A precursor sol consisting of titanium isopropoxide and lithium acetate was impregnated into the void space of template, and it was calcined at various temperatures. A microporous membrane of Li4Ti5O12 with inverse-opal structure was successfully prepared at 800 °C. The interconnected pores with uniform size (0.8 μm) were clearly observed on the entire part of the membrane. The electrochemical properties of the 3D ordered Li4Ti5O12 were characterized by cyclic voltammetry and galvanostatic charge and discharge in an organic electrolyte containing a lithium salt. The 3DOM Li4Ti5O12 exhibited a discharge capacity of 160 mA h g−1 at the electrode potential of 1.55 V versus Li/Li+ due to the solid state redox of Ti3+/4+ accompanying with

Lithium Titanate-Based Lithium-Ion Batteries 117 Li+ ion insertion and extraction. The discharge capacity was close to the theoretical capacity (167 mA h g−1), which suggested that the Li+ ion insertion and extraction took place at the entire part of the 3DOM Li4Ti5O12 membrane [94]. The poor electrical conductivity of lithium titanate (Li4Ti5O12) has limited its practical application as high-rate performance anode for lithiumion batteries. Liu et al., utilized the oxygen vacancy strategy to improve the electrical conductivity of Li4Ti5O12 nanospheres by synthesizing 3D oxygen-deficient Li4Ti5O12 on a titanium mesh from Ti3+-doped anatase TiO2. When the oxygen-deficient Li4Ti5O12 nanospheres were used as anode material for high-rate lithium-ion batteries, the electrode exhibited excellent lithium storage performance delivering a capacity of 143 mA h g−1 at a current density of 100C. The outstanding performance is due to the 3D architecture from the 3D substrates with large space and increased surface area that facilitate the rapid transfer of lithium ion and small diameter of nanospheres, which provides shorter diffusion path-length [95]. Flower-like Li4Ti5O12 hollow microspheres consisting of nanosheets are prepared via a hydrothermal process and subsequently wrapped by graphene through electrostatic interactions. In comparison with pristine Li4Ti5O12, Li4Ti5O12@G exhibited higher capacities and improved rate capability in the 0.01–3.0 V or 1.0–3.0 V potential range. The Li4Ti5O12@G composite shows the specific capacity of 272.7 mA h g−1 at 750 mA g−1 after 200 cycles in the potential range from 0.01 to 3.0 V, while the pristine Li4Ti5O12 only delivered a discharge capacity of 235.6 mA h g−1. The improved electrochemical performances of Li4Ti5O12@G should be attributed to lower charge-transfer resistance, larger lithium-ion diffusion coefficient, and lower activation energy. [96]

2.3.5

Other Nanostructures

The commercial Li4Ti5O12 is modified with SiO2 directly via a simple solgel method [97]. Compared with pure Li4Ti5O12, the structure of SiO2 modified Li4Ti5O12 has no change, and there is a SiO2 coating layer over the Li4Ti5O12 particles. An appropriate amount of SiO2 could effectively reduce the electrochemical polarization of Li4Ti5O12 and enhance electrochemical reaction kinetics of Li+ insertion/deinsertion. Li4Ti5O12 modified with 2.5 mol% SiO2 exhibited higher specific capacity and better rate capability. Moreover, the SiO2 coating layer is likely to cover the catalytic sites on the Li4Ti5O12 surface for the decomposition of the electrolyte, thereby restraining the formation of solid electrolyte interphase (SEI), which is very favorable for improving the cycle performance of Li4Ti5O12.

118 Advanced Battery Materials Pawlitzek, et al., prepared a nanocomposite consisting of vertically aligned carbon nanotube arrays decorated with in-situ grown necklace type Li4Ti5O12 nanoparticles. The morphology of the composites was examined by SEM investigations. In Figure 2.15, it can be clearly seen that the vertical alignment of the CNT was conserved during the composite synthesis. Furthermore, LTO nanoparticles are homogeneously distributed throughout the whole volume of the VACNT array. The particle size and distribution strongly depend on the LTO load within the composites. Rather narrow size distributions around 25 nm and 45 nm are obtained for 45 wt% and 60 wt% composites, respectively. Whereas much more polydisperse particles in the range of 30–200 nm and 40–250 nm are maintained for the higher LTO loads with 70 wt% and 78 wt%, respectively. Owing to this structure the electrodes exhibited outstanding rate performances with specific capacities of 110 mA h g−1 up to 300C and cycling performance with high-capacity retention of 97% after 500 cycles at 10C. Thus, the combination of short Li+ diffusion distances within Li4Ti5O12 particles, remarkable electronic conductivity by carbon nanotubes directly grown on the current collector as well as a high contact surface area due to an open pore geometry was essential in achieving high-power Li4Ti5O12 anodes [98]. Nanostructured Li4Ti5O12/CNT composite particles were synthesized using an aerosol spray drying process followed by thermal annealing in Figure 2.16. After assembling into CNT networks, lightweight, flexible and binder-free Li4Ti5O12/CNT anodes were fabricated for lithium-ion batteries. Electrodes fabricated from these nanocomposites showed effective electron and ion transport and, more importantly, more robust structure compared to traditional binder-bonded electrodes. The electrodes delivered a high reversible capacity and superior rate performance up to an extremely high rate of 100C. Moreover, ultra-long cycling stability was attained through 8000 rapid charge-discharge cycles with 89% capacity retention. [99] Hierarchical Li4Ti5O12-TiO2 microspheres are synthesized via an oil/ water interface method with the assistance of cetyltrimethylammonium bromide (CTAB) and following calcination in Figure  2.17. The as-prepared Li4Ti5O12-TiO2 microspheres consist of nanoflakes (ca. 15 nm in thickness); moreover, Li4Ti5O12 and anatase TiO2 in the nanoflake building blocks are bounded by (011) and (001) facets, respectively. More especially, it is found that CTAB played a very important role in the phase and morphology of the Li4Ti5O12-TiO2 composite. When employed as the anode materials for lithium-ion batteries, the Li4Ti5O12-TiO2 microspheres showed initial discharge capacities of 234, 170, 151 and 140 mA h g−1 at charge/discharge rates of 1, 5, 10, and 20C,

Lithium Titanate-Based Lithium-Ion Batteries 119

(a)

500 nm

(b)

250 nm

(c)

250 nm

(d)

250 nm

(e)

250 nm

(f)

250 nm

2.50 45 wt% 60 wt% 70 wt% 78 wt%

0.5

0.0

–0.5

2.25 Potential / V vs Li/Li+

Current / mA mg–1

1.0

2.00 1.75 1.50 1.25

–1.0 1.35 1.40 1.45 1.50 1.55 1.60 1.65 1.70 1.75 (e) Potential / V vs Li/Li+

25

50

75

150 125 100 75 50 25

100 125

150 175

200

Specific capacity / mAh g–1

100

175

0

(g)

0

(f)

80 60 40 20

Coulombic efficiency / %

Specific capacity / mAh g–1

200

1.00

0 0 50 100 150 200 250 300 350 400 450 500 Cycle number

Figure 2.15 SEM images of LTO/VACNT composites: (a) side view, (b) higher magnification SEM image illustrating the necklace-type structure of the LTO/VACNT composite, (c) 45 wt% LTO, (d) 60 wt% LTO, (e) 70 wt% LTO, (f) 78 wt% LTO. Electrochemical characterizations of LTO/VACNT composites (45 wt% black, 60 wt% red, 70 wt% blue, 78 wt% purple): (e) cyclic voltammograms at a sweep rate of 0.1 mV s−1, (f) voltage profiles at 2C, (g) long-term stability test at 10C. Reproduced from Ref. 98 with permission from the ScienceDirect.

respectively. Furthermore, after 100 cycles at 5C, the discharge capacity remained at 158 mA h g−1 with a capacity retention of 92.9%. The excellent electrochemical properties of the product can be ascribed to the

120 Advanced Battery Materials Drying zone

Heating zone High temperature (450 ºC)

Atomizer

Filter Carrier gas (N2)

CNT Sucrose Hydrated lithium acetate Isopropyl titanate

Exhaust

Aerosol process

Mixture solution Aerosol droplet

Aerosol particle

(a)

CNT

Channel Li4Ti5O12

(b)

Figure 2.16 Schematic of (a) aerosol spray drying process and (b) the prepared nanostructured Li4Ti5O12/CNT composites [99].

synergistic effect of the exposed Li4Ti5O12 (011) and anatase TiO2 (001) facets, abundant phase interfaces, hierarchical structures, and thin flakelike building blocks [100]. Large-scale functionalized carbon-anchored Li4Ti5O12 ultrathin nanosquares were synthesized via a combination of ball milling and solid state routes. The transformation from microspheres to nanosquares was achieved through a "mechanical activation and Ostwald ripening" mechanism via examining the intermediate product structure at different reaction stages. This research indicates that C-LTO nanosquares with high crystallinity are a promising anode material for high-specific-energy in rechargeable lithium batteries [101].

Lithium Titanate-Based Lithium-Ion Batteries 121

Ti(OC4H9)4 cyclohexane Nucleation and LiOH CTAB

adsorption of CTA+

Calcination

CTAB

Anisotropic growth

Oriented attachment

Ostwald ripening

Li1.81H0.19Ti2O5 xH2O

Assembly of nanoflakes

Li4Ti5O12-TiO2

Figure 2.17 Schematic illustration of the formation process of LTO-TO microspheres [100].

2.4

Modification of Lithium Titanate

To improve electronic conductivity of LTO, surface coating and lattice doping are two key routes for fabricating the high-performance LTO materials.

2.4.1 Surface Coating The carbon layer is one of the most attractive surfaces modified materials due to its low cost, designability and high electronic conductivity. Zhu et al., reported a facile process for preparing a carbon-coated nanosized Li4Ti5O12 nanoporous micro-sphere by a carbon pre-coating process combined with the spray drying method as shown in Figure 2.18 [102]. The obtained material consisted of a micron-size secondary sphere (10–20 μm) accumulated by carbon-coated nanosized primary particles with 200 nm. The nanosized primary particles and nano-thickness carbon layer uniformly coated over the particles as well as the interconnected nanopores greatly improve its rate capability. As a consequence, the resulting sample delivered a reversible capacity of 160 mA h g−1 at 0.2C, and showed remarkable rate capability by maintaining 79% of the capacity at 20C (vs. 0.2C), as well as excellent cycling stability with a capacity retention of 95% after 1000 cycles at 1C rate (vs. 0.2C). Microscale C-Li4Ti5O12 particles with high tap density were synthesized by a simple solid-state reaction using TiO2, Li2CO3, and pitch. On calcination of the particles at a high temperature in an inert atmosphere, the uniformly coated carbon layer from pitch inhibited the growth of primary

122 Advanced Battery Materials

Li salt

High temp.

Melting

Treating Micron-size LTO

Conventional solid-state process Mixing with sugar

Ball milled

600 ºC - 5h

with Li2CO3

Coated carbon layer

Carbon-coated nano-sized LTO Spray drying 800 ºC - 10h

Nanoporous micro-sphere LTO/C with high tap density

Interconnected open pores TiO2 Sugar Li2Co3

Our new strategy

Figure 2.18 Schematic presentation of the conventional solid-state process for micronsize Li4Ti5O12 and our new strategy for carbon-coated nano-sized Li4Ti5O12 nanoporous microsphere (CN-LTO-NMS) [102].

Mixture (Pitch, +Li2CO3+TiO2)

Anatase-TiO2 precursor Infiltration of pitch into voids

C-coated Li4Ti5O12

Carbonization in Ar at 900 ºC

Figure 2.19 Schematic illustration of the synthetic process of microscale spherical, carbon-coated Li4Ti5O12 and photographs of material formed during synthesis [103].

particles, maintaining the spherical morphology, similar to the TiO2 precursor in size and shape, and also enabling partial reduction of the starting Ti4+ to Ti3+ (Figure 2.19). The excellent electronic conductivity of the C-coated Li4Ti5O12 resulted from the presence of the highly conducting carbon coating layer and the mixed valence state of Ti3+ and Ti4+. Both the nanoporous morphology and highly conducting carbon coating layer in Li4Ti5O12 particles gave rise to high-rate capability [103].

Lithium Titanate-Based Lithium-Ion Batteries 123 A novel synthesis of highly oxygen-deficient “defective-LTO” anode material with high-rate performance is reported [104]. It is synthesized using conventional precursors via a one-pot thermal reduction process. A high level of oxygen vacancies of ≈6.5 at% and conformal amorphous carbon coating are achieved simultaneously, resulting in a compounding effect for a high discharge capacity of 123 mA h g−1 with a Coulombic efficiency of 99.8% at 10C. Ab initio calculations show that oxygen vacancies increased the electron donor density of the neighboring titanium atoms, while conformal amorphous carbon significantly reduced interfacial charge-transfer resistance. The core-shell Li4Ti5O12@polyaniline composites (LP) have been synthesized via an in situ synthesis with different mole ratios of aniline: LTO (25:1, 50:1 and 100:1). LP-2 electrodes (50:1) exhibited a high initial reversible capacity of 205 mA h g−1 with an initial Coulombic efficiency of 97.6% at 0.1C. Even at 10C, the reversible capacity of the LP-2 electrodes still remains at 102 mA h g−1. Moreover, the LP-2 electrodes retained an impressive high capacity of 161 mA h g–1 after 100 cycles at 1C, with 0.11% capacity fading per cycle, attributing to the significantly decreasing chargetransfer impedance of the LP composite and reductive polarity difference between the cathode and the electrolyte. [105] TiN modified micron-sized Li4Ti5O12 has been synthesized via a combination of solid-state reaction and surface thermal nitridation. The results demonstrated that a thin amorphous TiN coated on the surface of the Li4Ti5O12 particles. The TiN coating layer enhanced surface electronic conductivity and electrical contact between particles. The Li4Ti5O12 electrode showed the good electrochemical performance of 144.5 mA h g−1 at 5C after 500 cycles with a capacity retention of 91.6%, much higher than that of pristine Li4Ti5O12 and tantalum doped or TiN coated Li4Ti5O12, only Ta-doped Li4Ti5O12. This design by exploring both surface modification and bulk doping is highly attractive for high-performance Li4Ti5O12 manufacturing and may be applicative to other micron-sized electrode materials with inferior conductivity [106]. More coating layers and their performance are provided in Table 2.6.

2.4.2

Doping

Doping is one of effective methods for altering the structural properties and thus the performance of lithium titanate based anode materials. A small amount of ion doping will not change their phase structure and may significantly affect the bond strength, the lattice, the valence state of cations, and lattice defects. To mitigate the drawbacks of Li4Ti5O12, especially

Agents

acetate

Citric acid

Polyacrylonitrile

Polyvinyl pyrrolidone

Sugar

NH3

Layer

Carbon

Carbon

N-doped carbon

Carbon

Carbon

TiN

1.1–3.4

\

\

1–10

2–10

4–6

Thickness/nm

Table 2.6 Effect of different coating layers on the li4Ti5O12.

Solid state

Solid state

Spray drying

Solid state

Sol gel

Solid-state

Synthesized method

1.0–2.5

1.0–2.5

1.0–2.5

1.0–2.5

1.0–3.0

1.0–2.0 0.0–2.0

Voltage range/V

[110] [111] [106]

107.2 mA h g−1 at 10C 155.7 mA h g−1 at 0.1C 144.5 mA h g−1 at 5C after 500 cycles

(Continued)

[109]

[108]

133.5 mA h g−1 after 50 cycles at 1C Initial capability of 158 mA h g−1 at 0.2C

[107]

Ref.

155.0 mA h g discharged at 20 mA g−1 220.2 mA h g−1 discharged at 20 mA g−1

−1

Electrochemical performance

124 Advanced Battery Materials

Agents

Graphene

Aniline

AgNO3

GO

NH4F

Graphene

Polyaniline

Ag

RGO

LiF

Cont.

Layer

Table 2.6

\

\

3–4

20, 35

\

Thickness/nm

Solid state

Microwave

Hydrothermal method

In situ polymerization

Sol gel

Synthesized method

1.0–2.5

1.0–3.0

1.0–2.5

1.0–2.5

1.0–3.0

Voltage range/V

[112] [105] [113] [89] [114]

161 mA h g–1 after 100 cycles at 1C 186.3 mA h g−1 at 0.5C 176.6 mA h g–1 at a rate of 1C 167 mA h g–1 at 1C

Ref.

100 mA h g at 40C

−1

Electrochemical performance

Lithium Titanate-Based Lithium-Ion Batteries 125

126 Advanced Battery Materials its poor electronic conductivity, extensive research has been conducted to increase the transport properties. The LTO performances have been improved through doping monovalent (Na+, K+, Ag+) [115, 116], divalent (Mg2+, Ca2+, Cu2+, Sr2+, Zn2+, Pb2+, Ni2+, Co2+) [56, 117–124], trivalent cations (Fe3+, Al3+, Cr3+, La3+, Sc3+, Nd3+) [125–130], anions (F−, Br−) [131, 132] and mixed ions (W6+, Br−; Al3+, F−; La3+, F−; Sr2+, Na+; Sr2+, Pb2+; Ba2+, La3+; Pb2+, Ba2+) [132–138] at the 8a tetrahedral site, the octahedral 16d site, the oxygen anion site, or a combination of these sites. Table 2.7 provided different dopants, structural formula, doping site and electrochemical performance after doping. A series of NaxLi4−xTi6O14 are prepared by the solid-state method and their structures, morphologies and electrochemical properties are compared by changing the Na/Li atomic ratio. It can be found that two single-phase samples of Na2Li2Ti6O14 and Na4Ti6O14 and three multi-phase samples of Li4Ti6O14, NaLi3Ti6O14 and Na3LiTi6O14 can be formed by using a stoichiometric ratio of starting materials based on the formula of NaxLi4− Ti6O14 [140]. In Figure 2.20, electrochemical tests showed that Li4Ti6O14 x delivered the highest charge capacities of 116 mA h g−1 at 100 mA g−1, 91.5 mA h g−1 at 200 mA g−1, 77.5 mA h g−1 at 300 mA g−1 and 66.9 mA h g−1 at 400 mA g−1. It suggested that Li4Ti6O14 was the most potential lithium storage material among all the NaxLi4−xTi6O14 samples. Li4-xMgxTi5O12 anode material was synthesized by a solid-state method using Li2CO3, MgO and anatase TiO2. The effects of Mg element doping on the crystal structure, phase composition, morphology and electrochemical properties of Li4-xMgxTi5O12 were investigated by many characteristics. The electrochemical results indicated that when x = 0.10, the resultants showed better electrochemical performance in higher rates cycle stability and the polarization degree. The initial discharge capacity reached 122.5 mA h g–1 at 2C, and the discharge capacity still remained 119.3 mA h g–1 after 100 cycles [118]. Trivalent cations also were investigated in the recent years. Li et al., firstly presented the Al-doping-induced memory effect in LIBs. Pristine Li4Ti5O12 was free from the memory effect, while a distinct memory effect could be induced by Al-doping. After being discharged to a lower cut-off potential, Al-doped Li4Ti5O12 exhibited poorer electrochemical kinetics, delivering a larger overpotential during the charging process. This dependence of the overpotential on the discharging cutoff led to the memory effect in Al-doped Li4Ti5O12 [126]. Chromium (Cr)-modified LTO with a synergistic effect of bulk doping, surface coating, and size reducing was synthesized by a facile sol-gel method in Figure  2.21a. The size of LTO particles can be significantly

Structure

Li4-xNaxTi5O12

Li4-xNaxTi5O12

Ag doped LTO

Na+

Na+

Ag+

Dopant

Monovalent cations

Table 2.7 The performance of doped lto materials.

/

8a

/

Doping sites

/1.53

1.6/1.52

1.68/1.51

Charge/discharge plateau V/V

130/5C at 100th

Capacity/mA h g-1

193/0.2C; 196.9/1C

197/1C at 10th

155/0.1C; 152/1C; 150/0.1C at 50th 130/10C

139/2C

Rate

(Continued)

[139]

[68]

[115]

Ref.

Lithium Titanate-Based Lithium-Ion Batteries 127

Li3.9Mg0.1Ti5O12

Li4-xCaxTi5O12

Li3.8Cu0.3Ti4.9O12

Sr-doped LTO

Li2ZnTi3O8

PbLi2Ti6O14

Li(Ni1/6Li2/9Ti1/9)Ti3/2O4

Mg2+

Ca2+

Cu2+

Sr2+

Zn2+

Pb2+

Ni2+

Co

Co-doped LTO

Li27MgTi27O64

Mg2+

2+

Li4-xMgxTi5O12

Structure

Mg2+

Cont.

Divalent cation

Dopant

Table 2.7

\

16d

\

\

\

16d

8a

8a

8a

8a,16d

Doping sites

164/152

1.74/1.42

1.43/1.51; 1.40/1.46; 1.22/1.33

1.41/1.17

1.7/1.52

1.65/1.55

1.71/1.51

1.58/1.53

1.6/1.5

1.6/1.48

Charge/discharge plateau V/V

112.6/50C

165/0.2C

266.9/100 mA g-1

526.6/100 mA g-1

160/0.5C

146/1C; 114/10 C

174/1C; 147/10C

122.5/2 C

211/2C

169/0.25 mA cm-2

Rate

150.8/5C at 1000th

140/0.2C at 50th

[124]

[122]

[123]

[122]

[121]

[120]

[119]

[118]

[56]

[117]

Ref.

(Continued)

102/1A g−1 at 1000th

124.4/1A g−1 at 1000th

150/0.5C at 100th

112/10C at 100th

162/1C,139/10C at 100th

119.3/2C at 100th

170/2C, 160/5C at 100th

130 at 8th /0.25 mA cm-1

Capacity/mA h g-1

128 Advanced Battery Materials

Pb1-xBaxLi2Ti6O14

Pb2+,Ba2+

Ba ,La

Ba0.9M0.1Li2Ti6O14

Sr1-xPbxLi2Ti6O14

3+

2+

Sr2+,Pb2+

Sr , Na

Sr1−xNa2xLi2Ti6O14

Li3.95La0.05Ti5O11.7F0.3

La3+, F-

2+

Li4AlxTi5−xFyO12−y

W,Br- doped LTO

F-doped LTO

Al3+,F-

W ,Br

Mixed ions

6+

F-

Anion

Nd

+

Sc-doped LTO

Sc3+

−

La-doped LTO

3+

Li4Ti4.95Nd0.05O12

Cr-doped LTO

Cr3+

3+

Al-doped LTO

Al3+

La

Fe-doped LTO

Structure

Fe3+

Cont.

Trivalent cation

Dopant

Table 2.7

\

\

\

\

\

\

32e(Br)

32e

\

16d

\

\

\

\

Doping sites