Vanadium-Based Nanomaterials for Electrochemical Energy Storage 9783031447952

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Liqiang Mai Lin Xu Wei Chen

Vanadium-Based Nanomaterials for Electrochemical Energy Storage

Vanadium-Based Nanomaterials for Electrochemical Energy Storage

Liqiang Mai • Lin Xu • Wei Chen

Vanadium-Based Nanomaterials for Electrochemical Energy Storage

Liqiang Mai State Key Laboratory of Advanced Technology for Materials Synthesis and Processing Wuhan University of Technology Wuhan, China

Lin Xu School of Materials Science and Engineering Wuhan University of Technology Wuhan, China

Wei Chen School of Materials Science and Engineering Wuhan University of Technology Wuhan, China

ISBN 978-3-031-44795-2 ISBN 978-3-031-44796-9 https://doi.org/10.1007/978-3-031-44796-9

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Preface

The energy issue is the most important topic of the twenty-first century. An excessive reliance on the combustion of nonrenewable fossil fuels brings not only ecological problems but also harsh impacts on the global economy and society. There is a pressing requirement for cost-effective, efficient, and environmentally benign energy storage devices that can power energy-demanding areas, ranging from portable electronics (e.g., cell phones, camcorders, laptops) to transportation (e.g., electric vehicles, hybrid electric vehicles). Typical energy storage devices include fuel cells, solar cells, photoelectrochemical water splitting cells, batteries (especially Li-ion batteries), and supercapacitors. The performance of these energy devices relies strongly on the properties of their nanostructured materials. New development in the field of nanomaterial chemistry is believed to hold the key to further breakthroughs in energy storage systems. Nanostructured materials are advantageous in offering huge surface to volume ratios, favorable transport properties, altered physical properties, and confinement effects resulting from the nanoscale dimensions, and have been extensively studied for energy-related applications such as solar cells, catalysts, thermoelectrics, lithium-ion batteries, supercapacitors, and hydrogen storage systems. Vanadium element from VB group has a valence electron layer of 3d34s2. All the five valence electrons of V atom can take part in bonding, thereby giving rise to a multivalent V from V2+ to V5+, and contributing to various vanadium-based nanomaterials, such as vanadium oxides, vanadates, and vanadium phosphates. Apart from the rich electrochemical properties of V (V2+ to V5+), the low price and abundant source also make V-based nanomaterials as promising electrode candidates for energy storage field. This book provides an overview of this vibrant area of vanadium-based nanomaterials research, summarizes the characteristics of various vanadium-based nanomaterials and their applications in energy storage, and introduces some advanced characterization techniques to further analyze the mechanism of materials during energy storage reactions. In Chap. 1, basic theory and general information of vanadium are introduced to show the characteristics of vanadium material, and major vanadium-based electrode is classified. v

vi

Preface

Chapter 2 introduces the fundamentals of electrochemical energy storage, including the concept and development in the early years of this field, and the advantages and importance of nanomaterials in the electrochemical energy storage field are highlighted. In Chap. 3, we overview main strategies for the synthesis of vanadium-based nanomaterials, including hydro/solvo-thermal method, electrospinning strategy, template synthesis, sol-gel route, and vapor deposition techniques, where the increased modification in the nanomaterials can enable unique functional properties. In Chap. 4, to investigate the mechanism of capacity decay of electrochemical energy storage devices, we focus on the advanced in situ characterization techniques, aiming to reveal the fundamental reasons for performance change during energy storage reactions and find an effective way to improve the properties of nanomaterials. In Chap. 5, various optimization approaches are introduced to meet the elevated demands of energy storage systems, specifically highlighting the electric transport performance and structural stability field. The optimization strategies are listed and discussed in detail to solve some defects of original vanadium-based materials. Benefiting from the multiple values of vanadium, various vanadium-based nanomaterials were developed and used in different energy storage systems (e.g., vanadium oxides, vanadates, vanadium phosphates, and oxygen-free vanadiumbased nanomaterials), and every kind of nanomaterial is introduced in detail in terms of advantages, applications, and outlook in Chaps. 6, 7, 8, and 9. Due to the layered structure of vanadium-based materials, one-dimensional nanowires and other flexible nonstructural materials can be easily fabricated. In Chap. 10, the application of vanadium-based flexible nanomaterials in micro/nano and flexible energy storage devices is introduced emphatically, representing their unique functional properties. Finally, in Chap. 11, we conclude this book and look into the future of the exciting opportunities of vanadium-based nanomaterials. We thank the support of the National Key Research and Development Program of China (2020YFA0715000), the National Natural Science Foundation of China (51832004, 52127816), the Foshan Xianhu Laboratory of the Advanced Energy Science and Technology Guangdong Laboratory (XHT2020-003), and the Key Research and Development Program of Hubei Province (2021BAA070). Wuhan, China

Liqiang Mai Lin Xu Wei Chen

Contents

1

Fundamentals of Vanadium-Based Nanomaterials . . . . . . . . . . . . . 1.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2 General Information on Vanadium . . . . . . . . . . . . . . . . . . . . . . 1.3 History of Vanadium-Based Electrode Materials . . . . . . . . . . . . 1.4 Classification of Vanadium-Based Electrode Materials . . . . . . . . 1.5 Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1 1 1 4 7 10 12

2

Basic Information of Electrochemical Energy Storage . . . . . . . . . . . 2.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2 Electrochemical Energy Storage Technology . . . . . . . . . . . . . . . 2.2.1 Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.1.1 Lithium-Ion Batteries . . . . . . . . . . . . . . . . . . 2.2.1.2 Sodium-Ion Batteries and Potassium-Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.1.3 Rechargeable Multivalent Batteries . . . . . . . . 2.2.1.4 Flow Batteries . . . . . . . . . . . . . . . . . . . . . . . 2.2.2 Supercapacitors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.2.1 Electrical Double Layer Capacitors . . . . . . . . 2.2.2.2 Pseudocapacitors . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

17 17 18 19 20

3

Synthesis of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . 3.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2 Hydro-/Solvothermal Method . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1 Template-Free Hydro-/Solvothermal Synthesis . . . . . . . 3.2.2 Soft Template-Assisted Hydro-/Solvothermal Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.3 Hard Template-Assisted Hydro-/Solvothermal Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3 Electrospinning Strategy . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

26 31 35 37 39 40 41 49 49 50 52 56 59 61 vii

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3.4 Template Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5 Sol-Gel Route . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.6 Vapor Deposition Techniques . . . . . . . . . . . . . . . . . . . . . . . . . 3.7 Other Methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.8 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

68 71 75 79 80 80

4

In Situ Characterizations of Vanadium-Based Nanomaterials . . . . . 87 4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87 4.2 In Situ Spectroscopic Characterizations of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88 4.2.1 In Situ X-Ray Diffraction of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89 4.2.2 In Situ Raman Spectroscopic Characterization of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . 92 4.2.3 In Situ XANES Characterization of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 94 4.3 In Situ Microscopic Characterizations . . . . . . . . . . . . . . . . . . . . 95 4.4 Other In Situ Characterization . . . . . . . . . . . . . . . . . . . . . . . . . 98 4.5 Summary and Outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102

5

Performance Optimization of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2 Electric Transport Performance Optimization . . . . . . . . . . . . . . 5.2.1 Band Structure Optimization . . . . . . . . . . . . . . . . . . . . 5.2.1.1 Ionic Pre-intercalation . . . . . . . . . . . . . . . . . . 5.2.1.2 Elemental Doping . . . . . . . . . . . . . . . . . . . . . 5.2.1.3 Extra Field Application . . . . . . . . . . . . . . . . . 5.2.2 Surface/Interface Optimization . . . . . . . . . . . . . . . . . . 5.2.2.1 Compositing with Conductive Material . . . . . 5.2.2.2 Design of Nanostructure . . . . . . . . . . . . . . . . 5.2.3 Ion Diffusion Channel Optimization . . . . . . . . . . . . . . 5.2.3.1 Metal Ion Pre-intercalation . . . . . . . . . . . . . . 5.2.3.2 Inorganic Molecule/Nonmetal Ion Preintercalation . . . . . . . . . . . . . . . . . . . . . . . . . 5.2.3.3 Organic Molecular Pre-intercalation . . . . . . . . 5.2.4 Electron/Ion Bi-continuous Optimization . . . . . . . . . . . 5.2.4.1 Coaxial Semi-Hollow Structures . . . . . . . . . . 5.2.4.2 Nanostructure Array . . . . . . . . . . . . . . . . . . . 5.3 Regulation and Control of Structural Stability . . . . . . . . . . . . . . 5.3.1 Internal Stress Buffering . . . . . . . . . . . . . . . . . . . . . . . 5.3.1.1 Hybrid Nanostructures . . . . . . . . . . . . . . . . . 5.3.1.2 Hierarchical Nanostructures . . . . . . . . . . . . . .

105 105 105 106 106 107 108 109 110 111 112 113 115 116 117 117 119 121 121 121 122

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5.3.2 Volume Expansion Suppression . . . . . . . . . . . . . . . . . 123 5.4 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . 124 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 125 6

7

Vanadium Oxide Nanomaterials for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2 Orthorhombic V2O5 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.1 V2O5 for LIB . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.1.1 V2O5 Nanocomposites for Lithium-Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.2 V2O5 for Sodium-Ion Batteries . . . . . . . . . . . . . . . . . . 6.2.3 V2O5 for Other Emerging Rechargeable Batteries . . . . . 6.3 Bilayered V2O5 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3.1 Bilayered V2O5 for Lithium-Ion Batteries . . . . . . . . . . 6.3.2 Bilayered V2O5 for Sodium-Ion Batteries . . . . . . . . . . . 6.3.3 Bilayered V2O5 for Other Emerging Rechargeable Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.4 VO2 (B) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.4.1 VO2 (B) for Lithium-Ion Batteries . . . . . . . . . . . . . . . . 6.4.2 VO2 (B) for Sodium-Ion Batteries . . . . . . . . . . . . . . . . 6.4.3 VO2 (B) for Other Emerging Rechargeable Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5 V6O13 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.6 V2O3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.7 V3O7
H2O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.8 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Vanadate Nanomaterials for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2 Alkali Metal Vanadates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2.1 LiV3O8 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2.2 Na1 + xV3O8 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . AxV2O5 (A = Li, Na, K) . . . . . . . . . . . . . . . . . . . . . . 7.2.3 7.2.4 LixVO2 and NaxVO2 . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2.5 Li3VO4 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2.6 Other Alkali Metal Vanadates . . . . . . . . . . . . . . . . . . . 7.3 Alkali-Earth Metal Vanadates . . . . . . . . . . . . . . . . . . . . . . . . . 7.4 Transition Metal Vanadates . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.1 Ag-V-O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.2 Cu-V-O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.3 Co-V-O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.4 Fe-V-O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

129 129 130 130 138 141 146 149 149 152 154 155 156 158 159 160 161 164 166 167 177 177 177 177 180 182 185 187 191 191 194 194 197 198 199

x

8

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10

Contents

7.4.5 Zn-V-O . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.6 Other Transition Metal Vanadates . . . . . . . . . . . . . . . . 7.5 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

201 204 208 208

Vanadium Phosphate Nanomaterials for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.2 Li3V2(PO4)3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.3 Na3V2(PO4)3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.4 Vanadium Fluorophosphates . . . . . . . . . . . . . . . . . . . . . . . . . . 8.4.1 AVPO4F (A = Li, Na, K) . . . . . . . . . . . . . . . . . . . . . . 8.4.2 Na3V2O2x(PO4)2F3 - 2x . . . . . . . . . . . . . . . . . . . . . . . 8.5 Vanadium Pyrophosphate . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.6 Vanadium-Based Mixed Polyanion Materials . . . . . . . . . . . . . . 8.7 VOPO4 and AOPO4 (A = Li, Na, K) . . . . . . . . . . . . . . . . . . . . 8.8 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

221 221 222 231 236 236 238 240 244 247 249 250

Oxygen-Free Vanadium-Based Nanomaterials for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . 9.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9.2 Vanadium Sulfides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9.2.1 Vanadium Disulfide . . . . . . . . . . . . . . . . . . . . . . . . . . 9.2.2 Vanadium Tetrasulfide . . . . . . . . . . . . . . . . . . . . . . . . 9.3 Vanadium Nitride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9.3.1 Vanadium Nitride for Supercapacitors . . . . . . . . . . . . . 9.3.2 Vanadium Nitride for Batteries . . . . . . . . . . . . . . . . . . 9.4 Vanadium Carbide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9.5 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

259 259 260 260 266 270 271 273 277 281 282

Vanadium-Based Nanomaterials for Micro-Nano and Flexible Energy Storage Device . . . . . . . . . . . . . . . . . . . . . . . . 10.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10.2 Micro-Nano Energy Storage Device . . . . . . . . . . . . . . . . . . . . . 10.2.1 Single Nanowire Energy Storage Device . . . . . . . . . . . 10.2.2 Planar Micro-Nano Energy Storage Device . . . . . . . . . 10.2.3 Thin-Film Micro-Nano Energy Storage Device . . . . . . . 10.3 Flexible Energy Storage Device . . . . . . . . . . . . . . . . . . . . . . . . 10.3.1 Linear Flexible Energy Storage Device . . . . . . . . . . . . 10.3.2 Stacked Flexible Energy Storage Device . . . . . . . . . . . 10.4 Summary and Future Directions . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

287 287 287 288 290 291 294 296 299 300 301

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11

Conclusions and Outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11.1 Characteristic of Vanadium-Based Materials . . . . . . . . . . . . . . 11.2 Classification of Vanadium-Based Materials . . . . . . . . . . . . . . 11.3 Synthesis and Characteristic of Vanadium-Based Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11.4 Application of Vanadium-Based Materials . . . . . . . . . . . . . . . 11.5 New Research Approaches . . . . . . . . . . . . . . . . . . . . . . . . . . 11.6 Outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

xi

. 303 . 303 . 304 . . . . .

304 305 306 306 308

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 311

Abbreviations

1D 2D 3D AHD AIBs ALD CEs CIBs CV CVD DFT EDLCs EIS EIS GITT HRTEM LIBs LVP MIBs NASICON NMR NVP PIBs PVD SAED SEI SEM SHE SIBs TEM VOs

One-dimensional Two-dimensional Three-dimensional Ambient hydrolysis deposition Aluminum-ion batteries Atomic layer deposition Coulombic efficiencies Calcium-ion batteries Cyclic voltammetry Chemical vapor deposition Density functional theory Electric double layer capacitors Electrochemical impedance spectroscopy Electrical impedance spectroscopy Galvanostatic intermittent titration technique High-resolution transmission electron microscopy Lithium-ion batteries Li3V2(PO4)3 Magnesium-ion batteries Na super ionic conductor Nuclear magnetic resonance Na3V2(PO4)3 Potassium-ion batteries Physical vapor deposition Selected-area electron diffraction Solid electrolyte interphase Scanning electron microscopy Standard hydrogen electrode Sodium-ion batteries Transmission electron microscopy Vanadium oxides xiii

xiv

VRFBs XAFS XANES XAS XRD ZIBs

Abbreviations

Vanadium redox flow batteries X-ray absorption fine structure X-ray absorption near edge structure X-ray absorption spectroscopy X-ray diffraction Zinc-ion batteries

Chapter 1

Fundamentals of Vanadium-Based Nanomaterials

1.1

Introduction

Vanadium-based electrode materials, like V2O5, have been researched for more than 40 years [1, 2]. The valence state of vanadium can vary from +5 to +1 when used as battery electrodes, which indicates that multi-electrons reaction with high capacity can be achieved. For example, V2O5, as a lithium-ion battery (LIB) cathode, has a ~440 mAh g-1 theoretical specific capacity, which is about three times of that of commercial LiCoO2 and, basically, the highest value among all metal oxide cathodes for LIBs. The potential for high capacity is the main reason why vanadium-based electrode materials receive a continuous attention for next-generation batteries. Besides, ascribe to the rich valence state of vanadium, vanadium-based materials show various electrochemical properties, compositions, and structures [3]. As shown in Fig. 1.1, vanadium-based materials show diverse crystal structures with various valence states and chemical composition of vanadium. For example, V2O5, LiV3O8, and VS2 show a typical layered structure, which endows them the ability to accommodate the intercalation/deintercalation of guest ions (such as Li+, Na+, and Zn2+). Besides, other kinds of structures which are promising in energy storage fields, like quasi-layered structure (such as VO2 and K0.25V2O5), 3D tunnel structure (such as Na3V2(PO4)3), chain-like structure (such as VS4), and rock salt structure (such as VN), can be also found in vanadium-based materials family. The different structures lead to versatile properties of vanadium-based compounds, thus providing considerable choices/possibilities for breakthrough in the investigation of electrode materials, not only for LIBs but also for other energy storage devices.

© Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_1

1

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1

Fundamentals of Vanadium-Based Nanomaterials

Fig. 1.1 Crystalline structures of various vanadium-based materials. (Reproduced from [4]. Copyright 2020 Wiley-VCH) Table 1.1 Comparison of elemental abundance in Earth’s crust and resource price of typical transition metal elements used for rechargeable battery electrodes Element Elemental abundance (ppm) Resource price ($/kg)

Ti 5600 0.72

V 120 6.83

Mn 950 –

Fe 56,300 0.082

Co 25 26.23

Ni 84 9.30

Data for elemental abundance in Earth’s Crust from [5] and data for resource price derived from [6]

1.2

General Information on Vanadium

Vanadium is an early transition metal that belongs to the fourth period and the VB group in the periodic table. Among transition metals, vanadium is relatively abundant; its elemental abundance is about five times of that of cobalt (Table 1.1). Based on the data in Mineral Commodity Summaries 2017 from the US Geological Survey, the world vanadium resources outdo 63 million tons [6]. China, Russia, and South Africa are the main countries for the world vanadium mine production and reserves. Vanadium production from China takes up ~50% of the world total

1.2

General Information on Vanadium

3

Table 1.2 World vanadium mine production and reserves in 2016

China Russia South Africa Brazil Others World total

Mine production (thousand metric tons) 42 16 12 6 – 76

Reserves (thousand metric tons) 9000 5000 3500 – 1800 19,300

Data from [6]

(Table 1.2). Note that based on the production rate and reserves at present, vanadium resources can be mined for about 300 years (Table 1.2). Vanadium is an important alloy element. At present, the main industrial consumption of vanadium centers on steel manufacture. A small amount of vanadium doping can greatly increase the strength, toughness, and wear resistance of steel. The strengthening steels are widely used in the construction, infrastructure, and transportation industries. Also when alloyed with titanium and aluminum to form Ti-AlV alloys, a handful of vanadium helps to construct extremely lightweight and strong materials [7]. Outside of the steel industry, vanadium-based compounds also have wide applications in many other fields, for example, as catalysts for sulfuric acid industry, as colorants for glass and ceramic industry, and as electrolytes for vanadium redox flow batteries (VRFBs) for large-scale energy storage [6, 8]. The outer electron configuration of vanadium is 3d34s2, which leads to the rich valence state from V3+, V4+, V5+ to V2+. Due to the change of the valence state of V, vanadium-based compounds can display a variety of colors, mainly including red, yellow, green, blue, black, and purple [8]. That is why vanadium-based compounds can be used as colorants for glass or ceramic. Also the rich valence state of V is the fundamental basis for the all-vanadium redox flow batteries, in which V2+/V3+ redox peak as anolyte and the V4+/V5+ redox couple as catholyte [9, 10]. Especially, since the increased market demand for grid and renewable energy storage technologies, VRFBs installations have begun to grow dramatically, and as a result, electrolyte for VRFBs is now the second largest specialty market for vanadium. The electrochemical property of vanadium-based compounds remains a continual area for research and has sparked renewed interests in energy conversion and storage in recent years. Since vanadium has abundant resource, vanadium-based electrodes have a distinct advantage in cost compared to the commercialized cobalt-based electrodes. The average price of vanadium resource in 2016 was estimated to be ~US$6.83/kg, about a quarter of that of cobalt (~US$26.23/kg), and was also lower than that of nickel (~US$9.30/kg) (Table 1.1). Apart from the abovementioned points, the effect of vanadium on human health also needs to be paid attention to. Vanadium is a necessary trace element for the human body, which has important impacts on promoting the growth of the body. The

4

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Fundamentals of Vanadium-Based Nanomaterials

daily meal can provide several dozen micrograms of vanadium, which can meet the needs of the body and generally does not require additional supplements. It should be noted that additional intake of vanadium will result in poisoning [11]. Vanadium metal has no toxicity, but vanadium oxides or other compounds are toxic for the human body. In general, the higher the valence state, the greater the toxicity. Given the highly toxic nature of V2O5, it is essential to exercise caution and take appropriate protective measures while conducting experiments involving vanadium oxides, including VO2, in both laboratory and industrial settings [12, 13].

1.3

History of Vanadium-Based Electrode Materials

Vanadium-based materials are one of the groups which were paid attention to research on LIBs in the earliest period. The Li+ intercalation properties of V2O5 have been studied by Whittingham since 1976 [1]. After that, research works about vanadium-based materials used in lithium storage devices were successively reported. In 1981 [14], Nassau et al. were the first to report the electrochemical behavior of LiV3O8, which is another important phase for Li storage. It should be noted that the investigations about V2O5 and LiV3O8 for energy storage are undertaken even today. The identification of these two phases for Li storage opens up the developing history of vanadium-based electrode materials (Fig. 1.2). On the other hand, the successful commercialization of vanadium-based electrodes in lithium primary (not rechargeable) batteries at that period of time also plays an important role for the further development of vanadium-based electrode materials afterward. Silver vanadium oxides (SVOs), especially Ag2V4O11, have to be emphasized when it comes to the lithium primary batteries because of their lucrative application in implantable cardiac defibrillators in the early 1980s [15] (Fig. 1.2). The first implantable grade commercial cell was introduced by Keister et al. in 1986, in which Ag2V4O11, lithium metal, and an organic electrolyte worked as the cathode, anode, and electrolyte, respectively [15]. It was demonstrated that Ag2V4O11

Fig. 1.2 A history line of the exploitation of the vanadium-based electrodes materials for LIBs

1.3

History of Vanadium-Based Electrode Materials

5

exhibits high current and high stability to power the ICDs. Further studies indicated that Ag2V4O11 shows distinct superiority in capacity and volumetric energy density compared to pristine other SVO or V2O5 cathodes with different V/Ag ratios [16]. Since then, studies about primary batteries mainly focused on reducing size and increasing energy density. To achieve these goals, several new cathodes were explored, such as Ag4V2O6F2 [17, 18], Ag3Fe(VO4)2, AgFeV2O7 [19], AgNa (VO2F2)2 [20], CuV2O6 [21], etc. A detailed review about the development of cathode materials was summarized by Chen and Cheng in 2011 [22]. In spite of the commercialization of lithium primary batteries, the limitation in its application determines that lithium rechargeable batteries are more desired. Through the 1980s, many effects are dedicated to commercializing the Li metal anode systems, for example the Li//V2O5 battery [23]. However, the safety conundrum root in the Li dendrites block the further development of Li metal batteries. The LiCoO2//graphite rechargeable batteries were put forward by Sony in 1990. This device, called “lithium-ion battery,” gained significant prosperity afterward and predominated the rechargeable energy storage market. In spite of this, the research about vanadium-based electrodes (such as LiV3O8, VO2, Li3V2(PO4)3, and V2O5) has never stopped (as shown in Fig. 1.2), which could be ascribed to their highcapacity merit. The investigations about V2O5, VO2, and LiV3O8 for lithium batteries have been continuing until today, and their Li storage performance has been greatly improved in recent years due to the development of nanostructured materials. Various kinds of nanostructured vanadium oxides have been fabricated in the past few years. But unlike LiCoO2 with mobile Li+ ions in the structure, the Li-poor properties of these vanadium oxides restrict their application as cathodes in lithiumion batteries coupled with graphite anode. Recently, a new direction appeared for these high-capacity vanadium-based materials interests in Li metal rechargeable batteries for higher energy density revived since 2010 [24]. After investigation about vanadium oxides, vanadium phosphates attracted much research interests. The typical two phases are Li3V2(PO4)3 and VOPO4. Because of the inductive effect produced by the [PO4] group, these phases exhibit much higher output voltage compared to vanadium oxides when used as LIB cathodes, thus showing the potential to achieve higher energy density. Besides, both Li3V2(PO4)3 and VOPO4 can realize multi-electron reactions to exhibit a higher specific capacity than other phosphate cathodes (such as LiFePO4). Therefore, both Li3V2(PO4)3 and VOPO4 are very promising candidates for next-generation high-energy LIB cathodes. Recently considerable works are devoted to clarify the structure evolution and improve the cycling/rate performance during multi-lithium intercalation [25, 26]. Apart from cathode materials, high-capacity anode materials are also an important direction and have received much attention. After 2008, several kinds of vanadium-based high-capacity anodes were identified, such as VN, Li3VO4, and VS4. Especially, Li3VO4, which was first reported in 2013 [27], is a new anode based on insertion reaction. It was found that this new anode holds the potential to achieve higher capacity and better safety when compared to graphite. Therefore, a lot of works have been reported in the past 5 years about Li3VO4 regarding its synthesis,

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Fundamentals of Vanadium-Based Nanomaterials

intrinsic properties, lithium insertion mechanism, performance, and optimization strategies. These progresses have been summarized in a recent comprehensive review [28]. From the above, it is known that the investigations about vanadium-based electrode material for lithium batteries are successive from 1970s to present. On the other sides, since the emergence of energy storage systems around 2010, the application of vanadium-based electrode materials for other metal-ion batteries drew great attention. These energy storage systems mainly include SIBs, PIBs, magnesium-ion batteries (MIBs), calcium-ion batteries (CIBs), aqueous zinc-ion batteries (ZIBs), and aluminum-ion batteries (AIBs). The growth of these emerging battery systems results in a further boom in the research of vanadium-based nanomaterials. For one thing, the existing phases were successively investigated in emerging battery systems, and many new interesting fundamentals/mechanisms/ reactions were observed with exciting electrochemical performances, such as V2O5 and VO2 for ZIB [29, 30]. Additionally, many new vanadium-based materials were identified for emerging battery systems in recent years, such as Na3V2(PO4)3 and Na3(VO)2(PO4)2F for Na-ion battery cathodes [31, 32], CaV4O9 for Na-ion battery anode [33], Zn0.25V2O5
nH2O for ZIB cathode [34], and Mg0.25V2O5
H2O for CIB cathode [35]. Such a history trend suggests that vanadium-based electrode materials will sequentially receive great research interests in energy storage fields. Moreover, the number of reported scientific publications on vanadium-based nanomaterials in electrochemical energy storage has increased obviously in recent years (Fig. 1.3).

Fig. 1.3 Number of publications dealing with “Vanadium-based nanomaterials + electrochemical energy storage” in Web of Science (Mar 31, 2022)

1.4

1.4

Classification of Vanadium-Based Electrode Materials

7

Classification of Vanadium-Based Electrode Materials

As mentioned above, a lot of different vanadium-based electrode materials have been identified in the past 40 years. The rich valence state of vanadium results in numerous vanadium-based compounds. When composite with different anions and cations (Fig. 1.4), a variety of vanadium-based phases can be obtained, including vanadium nitrides, vanadium phosphates, vanadium oxides, vanadium carbides, metal vanadates, and vanadium sulfides. In the past decades, studies have been undertaken mainly on metal vanadates, vanadium oxides, and vanadium phosphates for electrochemical energy storage. By comparison, studies on vanadium sulfides, carbides, and nitrides in energy storage are relatively less because their synthesis is generally complex and difficult. The reports about these materials for energy storage mainly occurred in recent years. Because of the big family of vanadium-based electrode materials, a suitable classification is needed to organize the chapters and discuss their properties and progresses. In this book, vanadium-based electrode materials are divided into four kinds based on their structure features and electrochemical characteristics: oxygenfree vanadium-based compounds, vanadium phosphates, vanadates, and vanadium oxides (Fig. 1.5). Vanadium oxides are the most accessible species in this family, which includes V2O5, bilayered V2O5
nH2O, VO2, V3O7
H2O, V6O13, and V2O3. These vanadium oxides show different crystal structures and valence state of vanadium, which results in their diverse electrochemical properties. Among them, investigations about V2O5 are most in the past years. The open structure between layers of V2O5 allows 3 Li+ intercalation per unit, resulting in a ~440 mAh g-1 high theoretical capacity [2, 3]. However, the fast capacity fading, resulting from the decreased conductivity and structural degradation during electrochemical processes, is a big issue that needs to be resolved [2, 36]. Besides, since V2O5 is a Li-poor cathode, the application of V2O5 when matched with metallic Li anode will suffer from safety issue. Therefore,

Fig. 1.4 The periodic table of the elements. (Reproduced from [4]. Copyright 2020 Wiley-VCH)

8

1

Vanadium oxides

V2O5 Bilayered

LiV3O8

Li3VO4

LiVO3

NaV3O8

KV3O8

NH4V3O8

Ag2V4O11

AgVO3

Co3V2O8

Co2V2O7

Zn2(OH)3VO3

V6O13

VO2

VS2 Na3(VO)2(PO4)2F

CuV2O6 CoV2O6

Li3V2(PO4)3

Zn3V2O8

Na3V2(PO4)3

FeVO4

MoV2O8

Li0.3V2O5

Na0.33V2O5

K0.25V2O5

Na1.25V3O8

Zn0.25V2O5∙ nH2O

Ag0.33V2O5

Cu2.33V4O11 NH4V4O10

LiVPO4F

V5+/V4+

VSe2

Na3V2(PO4)2F3 Na7V3(P2O7)4

VN

Na7V4(P2O7)4(PO4)

NaVO2

V5+

VS4

NaVPO4F

CaV4O9

V3O3

Oxygen-free vanadium-based compounds

VOPO4

AIV3O9

Zn3(OH)2V2O7 ∙ 2H2O V3O7∙H2O

Vanadium phosphates

Vanadates

CaV6O16 ∙ 3H2O

V2O5∙nH2O

Fundamentals of Vanadium-Based Nanomaterials

V2 C

Na3V(PO3)3N

V4+

V3+

V2+

Fig. 1.5 Classification of vanadium-based electrode materials. (Reproduced from [4]. Copyright 2020 Wiley-VCH)

the application prospect of V2O5 has been queried for a long time. With the development of nanomaterials for energy storage and the metallic Li anode technology recent years, both of the two issues are potential to be overcome [24, 37]. Besides, the open structure with high capacity and moderate voltage makes V2O5 a promising candidate as cathode in emerging battery systems, such as MIBs or ZIBs, where metallic anode can be applied without safety issue. Therefore, recently, the research interests in V2O5 have revived, both for LIBs and emerging metal-ion batteries. Except for V2O5, nanostructured bilayered V2O5
nH2O also attracted great attentions in recent years [38]. Its special structure with large interlayer spacing makes it a good host material for large-sized metal cations (such as K+ and Na+) or multivalent ions (such as Zn2+ and Mg2+), thus a promising candidate for other energy devices. For other vanadium oxides (VO2, V3O7
H2O, V6O13, and V2O3) with low valence state or mixed valence state, their synthesis is relatively more complex. But in the past years, considerable research works were also carried out about the synthesis of these materials with nanostructures and electrochemical performance. The detailed progress about nanostructured vanadium oxides will be discussed in Chap. 6. For vanadium oxides, an important feature that needs to take special attention is the type of V-O bonds, especially the single-connected V-O bonds in the layer surface (Fig. 1.1, the crystalline configuration of V2O5, LiV3O8, and K0.25V2O5), which generally exists in vanadium oxides but rarely in other material systems. This kind of single-connected bond is flexible and results in the deformation of V-O polyhedrons [39]. Besides, the single-connected terminal oxygen atoms are

1.4

Classification of Vanadium-Based Electrode Materials

9

negatively charged, which together result in the powerful capacity of V-O structure to accommodate other ions. When composite with different cations, a lot of binary A-V-O (A represents NH4+ or metal ions) derivatives could be obtained and, thus, form a new group called vanadates. As shown in Figs. 1.4 and 1.5, most of the metal elements in periodic table could form compounds with vanadium oxides and then generate plenty of binary vanadates with diversity in electrochemical performance, composition, and structure. Compared with the pristine vanadium oxides, vanadates exhibit improved rate performance and cycling stability because of the benefits produced by introduced cations. For examples, in comparison to V2O5, alkali metal vanadates such as LiV3O8 and K0.25V2O5 have been reported to show better cycling performance during Li+ intercalation/deintercalation. It is believed that the incorporated alkali metal ions can act as pillars in the interlayers and, thus, result in better structure stability during cycling [39–41]. Therefore, the incorporated metal ions have significant impacts on the properties of vanadates. When composite with different metal ions, the achieved materials will show diversity in both crystal structures and electrochemical properties. The detailed classification of vanadates, the comparison between the structure and properties after different metal ions introduction, and their application potential are fully discussed in Chap. 7. When the oxygen anion is replaced by the [PO4] group in the structure, vanadium phosphates could be obtained. The incorporation of [PO4] tetrahedrons will result in a significant change in the structure and properties. In the structure aspect, unlike vanadium oxides or vanadates, majority of vanadium phosphates exhibit 3D framework constructed by [VO6] octahedrons and [PO4] tetrahedrons [25, 42, 43], such as Li3V2(PO4)3 and Na3V2(PO4)3. But α1-VOPO4 shows a layered structure [44–47] (as shown in Fig. 1.1). About the electrochemical properties, vanadium phosphates generally show much higher redox potential than that of vanadates. The higher redox potential originates from the inductive effect produced by the [PO4] tetrahedrons just as other phosphate electrodes [48–50]. Take Li3V2(PO4)3 as an example; its three discharge voltage plateaus at about 4.05, 3.66, and 3.58 V are all much higher than that of V2O5 or LiV3O8 [25]. The higher voltage together with the good structural stability makes Li3V2(PO4)3 a promising candidate for power batteries and, thus, attracts considerable attentions since its emergence [25]. Similarly, Na3V2(PO4)3 is also a hopeful cathode material for SIBs and has been widely studied recently [42, 43, 51, 52]. Na3V2(PO4)3 has a stable framework with NASICON (Na super ionic conductor) structure; therefore, facile Na+ diffusion with fast diffusion kinetics in the host structure and good structure stability during Na+ insertion/extraction is the distinct advantage. Besides, vanadium fluorophosphates (such as Na3V2(PO4)2F3) and vanadium pyrophosphate (such as LiVP2O7) have also attracted much attention in the past years [53], to realize higher voltage and thus higher energy density. Low electric conductivity is a general drawback for vanadium phosphates, vanadium fluorophosphates, and vanadium pyrophosphate, which is ascribed to the fact that [VO6] octahedrons are isolated and separated by the [PO4] or [P2O7] groups [49]. The most popular optimizing strategies are based on conductive coating together with nanostructure design, which will be discussed in Chap. 8 in detail.

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Fundamentals of Vanadium-Based Nanomaterials

Vanadium carbides, vanadium sulfides, and vanadium nitrides have also attracted considerable interests in recent years, such as V2C [54–56], V2N [57–60], and VS2 [61–63]. Because the oxygen is replaced by other anions with weaker electronegativity, distinctive electrochemical properties and structure features are found in these oxygen-free vanadium-based materials. However, the facile synthesis for these materials is a big challenge; that’s why only little works were reported in very recent years. Considering the relatively less study about these materials and their similar features and properties, the development of oxygen-free vanadium-based materials for energy storage devices is introduced in Chap. 9.

1.5

Vanadium-Based Nanomaterials

According to the equation t = L2/D, where L is the distance for ion diffusion, D is the ion diffusion coefficient, and t is the time of ion diffusion, decreasing the diffusion distance could efficiently increase the ion diffusion kinetics. Fabricating nanostructured materials is a general and realizable way to decrease the ion diffusion distance. Besides, it is demonstrated that nano-sized materials have better strain-release property compared to the bulk samples, which could avoid the crack or pulverization of the electrodes during ion insertion/extraction [64, 65]. Moreover, the large surface area of nanomaterials provides a higher electrode-electrolyte contact area, which, in general, could increase the use ratio of active materials and further accelerate the electrochemical reaction. All these merits of nanostructured materials could effectively result in better rate performance, cycling stability, and higher specific capacity when used as battery electrodes. Therefore, it is popular to fabricate or design nanostructured materials. For vanadium-based electrodes, as mentioned above, high capacity is one of the most important advantages that interest the researchers. However, fast capacity fading is a general issue that is caused by the structure degradation during multiion insertion together with poor electron/ion conduction. In this case, investigations focused on nanostructure design/fabrication and structure–performance relationships are more necessary for vanadium-based electrodes. In the past decade, the vast majority of the studies about vanadium-based electrodes for electrochemical energy storage are based on “nano.” Based on the dimension, nanomaterials are divided into three-dimensional (3D), two-dimensional (2D), one-dimensional (1D), and zero-dimensional (0D). Figure 1.6 shows the typical nanostructured vanadium-based materials with different dimensions. 0D nanoparticles are the most common structure among nanomaterials, which can be achieved both by top-down approach (mechanical pulverization from bulk materials) and bottom-up fabrication (chemical synthesis from ions/clusters). However, the high surface energy of the small-sized 0D nanoparticles generally results in the great tendency of self-agglomeration during cycling. Besides, the weak connections between nanoparticles will produce a high interfacial resistance. In comparison, 1D nanomaterials, including nanowires, nanobelts, nanorods, nanotubes, and

1.5

Vanadium-Based Nanomaterials

11

Fig. 1.6 Typical nanostructured vanadium-based materials with different dimensions. (a) Zerodimensional Na3V2(PO4)3 [66]. (b) One-dimensional V2O5 nanobelts; inset is the digital photograph of the V2O5 nanobelts solution [67]. (c) Two-dimensional VOPO4 ultrathin nanosheets [44]. (d) Nanosheet-assembled three-dimensional CaV4O9 microflowers; inset is the TEM image of an individual nanosheet [68]. (Reprinted by permission of the Royal Society of Chemistry, Elsevier and Nature Publishing Group)

nanoribbons, generally exhibit alleviated self-agglomeration. In addition, the intrinsic 1D continuous electron transport pathway along the axial direction and short ion diffusion distance along the radial direction endow the 1D nanomaterials with outstanding electron/ion conduction. Moreover, 1D ultra-long nanomaterials have innate superiority to generate free-standing electrode. Therefore, 1D nanostructured vanadium-based electrodes have received great attention in the past years. Recently, 2D nanomaterials, with graphene as a typical representative, are of great interest in many fields due to their unique surface characteristics and physicochemical properties. 2D vanadium-based nanomaterials have been introduced numerously. However, even though 1D and 2D nanomaterials show many advantages in electrochemical performance compared to bulk materials, sacrifice also exists. The low tap density of these materials will pull down the volumetric energy density of the batteries, which is detrimental to the real application. An efficient way to make a compromise between the advantage of 1D/2D nanostructures and drawbacks of low

12

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Fundamentals of Vanadium-Based Nanomaterials

tap density is constructing 3D hierarchical microstructures. The 3D microstructures are usually assembled by 1D/2D nanostructure units (Fig. 1.6d), which could effectively increase the space utilization and offset the sacrifice in tap density. Meanwhile, the 1D/2D nanostructure units could effectively maintain the nanoeffects.

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18. F. Sauvage, V. Bodenez, H. Vezin, T.A. Albrecht, J.M. Tarascon, K.R. Poeppelmeier, Ag4V2O6F2 (SVOF): A high silver density phase and potential new cathode material for implantable cardioverter defibrillators. Inorg. Chem. 47(19), 8464–8472 (2008) 19. T.A. Albrecht, F. Sauvage, V. Bodenez, J.M. Tarascon, K.R. Poeppelmeier, Room temperature synthesis of the larger power, high silver density cathode material Ag4V2O6F2 for implantable cardioverter defibrillators. Chem. Mater. 21(13), 3017–3020 (2009) 20. M.D. Donakowski, A. Görne, J.T. Vaughey, K.R. Poeppelmeier, AgNa(VO2F2)2: A trioxovanadium fluoride with unconventional electrochemical properties. J. Am. Chem. Soc. 135(26), 9898–9906 (2013) 21. H. Ma, S. Zhang, W. Ji, Z. Tao, J. Chen, α-CuV2O6 nanowires: Hydrothermal synthesis and primary lithium battery application. J. Am. Chem. Soc. 39(30), 5361–5367 (2008) 22. F. Cheng, J. Chen, Transition metal vanadium oxides and vanadate materials for lithium batteries. J. Mater. Chem. 21(27), 9841–9848 (2011) 23. M. Winter, B. Barnett, K. Xu, Before Li ion batteries. Chem. Rev. 118(23), 11433–11456 (2018) 24. D. Lin, Y. Liu, Y. Cui, Reviving the lithium metal anode for high-energy batteries. Nat. Nanotechnol. 12(3), 194–206 (2017) 25. C. Liu, R. Massé, X. Nan, G. Cao, A promising cathode for Li-ion batteries: Li3V2(PO4)3. Energy Stor. Mater. 4, 15–58 (2016) 26. Z. Chen, Q. Chen, L. Chen, R. Zhang, H. Zhou, N.A. Chernova, M.S. Whittingham, Electrochemical behavior of nanostructured ɛ-VOPO4 over two redox plateaus. J. Electrochem. Soc. 160(10), A1777–A1780 (2013) 27. H. Li, X. Liu, T. Zhai, D. Li, H. Zhou, Li3VO4: A promising insertion anode material for lithium-ion batteries. Adv. Energy Mater. 3(4), 428–432 (2013) 28. S. Ni, J. Liu, D. Chao, L. Mai, Vanadate-based materials for Li-ion batteries: The search for anodes for practical applications. Adv. Energy Mater. 9(14), 1803324 (2019) 29. N. Zhang, Y. Dong, M. Jia, X. Bian, Y. Wang, M. Qiu, J. Xu, Y. Liu, L. Jiao, F. Cheng, Rechargeable aqueous Zn–V2O5 battery with high energy density and long cycle life. ACS Energy Lett. 3(6), 1366–1372 (2018) 30. J. Ding, Z. Du, L. Gu, B. Li, L. Wang, S. Wang, Y. Gong, S. Yang, Ultrafast Zn2+ intercalation and deintercalation in vanadium dioxide. Adv. Mater. 30(26), 1800762 (2018) 31. Z. Jian, L. Zhao, H. Pan, Y.S. Hu, H. Li, W. Chen, L. Chen, Carbon coated Na3V2(PO4)3 as novel electrode material for sodium ion batteries. Electrochem. Commun. 14(1), 86–89 (2012) 32. M. Xu, L. Wang, X. Zhao, J. Song, H. Xie, Y. Lu, J.B. Goodenough, Na3V2O2(PO4)2F/ graphene sandwich structure for high-performance cathode of a sodium-ion battery. Phys. Chem. Chem. Phys. 15(31), 13032–13037 (2013) 33. X. Xu, C. Niu, M. Duan, X. Wang, L. Huang, J. Wang, L. Pu, W. Ren, C. Shi, J. Meng, Alkaline earth metal vanadates as sodium-ion battery anodes. Nat. Commun. 8(1), 460 (2017) 34. D. Kundu, B.D. Adams, V. Duffort, S.H. Vajargah, L.F. Nazar, A high-capacity and long-life aqueous rechargeable zinc battery using a metal oxide intercalation cathode. Nat. Energy 1(10), 16119 (2016) 35. X. Xu, M. Duan, Y. Yue, Q. Li, X. Zhang, L. Wu, P. Wu, B. Song, L. Mai, Bilayered Mg0.25V2O5•H2O as a stable cathode for rechargeable Ca-ion batteries. ACS Energy Lett. 4, 1328–1335 (2019) 36. L. Mai, Y. Dong, L. Xu, C. Han, Single nanowire electrochemical devices. Nano Lett. 10(10), 4273–4278 (2010) 37. J. Wang, H. Tang, L. Zhang, H. Ren, R. Yu, Q. Jin, J. Qi, D. Mao, M. Yang, Y. Wang, Multishelled metal oxides prepared via an anion-adsorption mechanism for lithium-ion batteries. Nat. Energy 1(5), 16050 (2016) 38. A. Moretti, S. Passerini, Bilayered nanostructured V2O5•nH2O for metal batteries. Adv. Energy Mater. 6(23), 1600868 (2016)

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39. Y. Zhao, C. Han, J. Yang, J. Su, X. Xu, S. Li, L. Xu, R. Fang, H. Jiang, X. Zou, Stable alkali metal ion intercalation compounds as optimized metal oxide nanowire cathodes for lithium batteries. Nano Lett. 15(3), 2180–2185 (2015) 40. H. Liu, Y. Wang, L. Li, K. Wang, E. Hosono, H. Zhou, Facile synthesis of NaV6O15 nanorods and its electrochemical behavior as cathode material in rechargeable lithium batteries. J. Mater. Chem. 19(42), 7885–7891 (2009) 41. X. Xu, Y.Z. Luo, L.Q. Mai, Y.L. Zhao, Q.Y. An, L. Xu, F. Hu, L. Zhang, Q.J. Zhang, Topotactically synthesized ultralong LiV3O8 nanowire cathode materials for high-rate and long-life rechargeable lithium batteries. NPG Asia Mater. 4(6), e20 (2012) 42. W. Ren, Z. Zheng, C. Xu, C. Niu, Q. Wei, Q. An, K. Zhao, M. Yan, M. Qin, L. Mai, Selfsacrificed synthesis of three-dimensional Na3V2(PO4)3 nanofiber network for high-rate sodiumion full batteries. Nano Energy 25, 145–153 (2016) 43. Y. Fang, L. Xiao, X. Ai, Y. Cao, H. Yang, Hierarchical carbon framework wrapped Na3V2(PO4)3 as a superior high-rate and extended lifespan cathode for sodium-ion batteries. Adv. Mater. 27(39), 5895–5900 (2015) 44. C. Wu, X. Lu, L. Peng, K. Xu, X. Peng, J. Huang, G. Yu, Y. Xie, Two-dimensional vanadyl phosphate ultrathin nanosheets for high energy density and flexible pseudocapacitors. Nat. Commun. 4(9), 2431 (2013) 45. Y. Zhu, L. Peng, D. Chen, G. Yu, Intercalation pseudocapacitance in ultrathin VOPO4 nanosheets: Toward high-rate alkali-ion-based electrochemical energy storage. Nano Lett. 16(1), 742–747 (2016) 46. H. Guang, K. Wang Hay, M. Arumugam, A 3.4 V layered VOPO4 cathode for Na-Ion batteries. Chem. Mater. 28, 682–688 (2016) 47. H. Li, L. Peng, Y. Zhu, D. Chen, X. Zhang, G. Yu, An advanced high-energy sodium ion full battery based on nanostructured Na2Ti3O7/VOPO4 layered materials. Energ. Environ. Sci. 9(11), 3399–3405 (2016) 48. A.K. Padhi, J.B. Goodenough, K.S. Nanjundaswamy, Phospho-olivines as positive-electrode materials for rechargeable lithium batteries. J. Electrochem. Soc. 144(4), 1188–1194 (1997) 49. A. Kraytsberg, Y. Ein-Eli, Higher, stronger, better: A review of 5 volt cathode materials for advanced lithium-ion batteries. Adv. Energy Mater. 2(8), 922–939 (2012) 50. K.S. Nanjundaswamy, A.K. Padhi, J.B. Goodenough, S. Okada, H. Ohtsuka, H. Arai, J. Yamaki, Synthesis, redox potential evaluation and electrochemical characteristics of NASICON-related-3D framework compounds. Solid State Ion. 92(1–2), 1–10 (1996) 51. Y. Xu, Q. Wei, C. Xu, Q. Li, Q. An, P. Zhang, J. Sheng, L. Zhou, L. Mai, Layer-by-layer Na3V2(PO4)3 embedded in reduced graphene oxide as superior rate and ultralong-life sodiumion battery cathode. Adv. Energy Mater. 6(14), 1600389 (2016) 52. S. Li, Y. Dong, L. Xu, X. Xu, L. He, L. Mai, Effect of carbon matrix dimensions on the electrochemical properties of Na3V2(PO4)3 nanograins for high-performance symmetric sodium-ion batteries. Adv. Mater. 26(21), 3545–3553 (2014) 53. V. Palomares, P. Serras, I. Villaluenga, K.B. Hueso, J. Carreterogonzález, T. Rojo, Na-ion batteries, recent advances and present challenges to become low cost energy storage systems. Energ. Environ. Sci. 5(3), 5884–5901 (2012) 54. M. Naguib, J. Halim, J. Lu, K.M. Cook, L. Hultman, Y. Gogotsi, M.W. Barsoum, New two-dimensional niobium and vanadium carbides as promising materials for Li-ion batteries. J. Am. Chem. Soc. 135(43), 15966–15969 (2013) 55. Y. Dall’Agnese, P.L. Taberna, Y. Gogotsi, P. Simon, Two-dimensional vanadium carbide (MXene) as positive electrode for sodium-ion capacitors. J. Phys. Chem. Lett. 6(12), 2305–2309 (2015) 56. S. Bak, R. Qiao, W. Yang, S. Lee, X. Yu, B. Anasori, H. Lee, Y. Gogotsi, X.Q. Yang, Na-ion intercalation and charge storage mechanism in 2D vanadium carbide. Adv. Energy Mater. 7, 1700959 (2017)

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Chapter 2

Basic Information of Electrochemical Energy Storage

2.1

Introduction

There is no doubt that energy is one of the key factors in modern society [1, 2]. Energy conversion and storage are huge challenges for economic development and social prosperity [3]. However, since the 1900s, the massive consumption of non-renewable fossil fuels has led to concerns about the energy crisis and corresponding carbon emissions, which have also led to the deterioration of environmental problems. There is a crucial demand for clean, economical, and reliable energy sources to replace fossil fuels. Compared with traditional energy sources, renewable energy sources, such as tidal energy, wind energy, solar energy, and biomass energy, have the characteristics of abundant resources, pollution-free operation, multi-channel utilization, and sustainability [4]. However, the use of these renewable energy sources is often restricted by factors such as time, space, and climate change [5]. For example, solar energy cannot be used continuously and can only be used during the day, and the solar energy is not evenly distributed [6, 7]. On the other hand, as the demand for consumer electronics, portable devices, and electric vehicles continues to grow, people are eager to realize low-cost, highefficiency, multifunctional energy, and environmentally friendly energy conversion and storage equipment (Fig. 2.1) [8]. Therefore, energy conversion and storage technology are the keys to achieving sustainable energy development [9, 10]. Typical energy storage systems can be separated into chemical energy storage, mechanical energy storage, electrochemical energy storage, charge energy storage, thermal energy storage, and mixed storage according to different energy storage methods [12]. The energy storage secondary battery, based on electrochemical storage, is considered to be one of the new energy storage equipment with the greatest potential for large-scale application in energy storage on account of its high energy density, smart power, adaptability and long life.

© Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_2

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18

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Basic Information of Electrochemical Energy Storage

Fig. 2.1 Schematic diagram of energy storage, power transmission, power distribution, end customers, and future smart grid energy storage applications, combining potential renewable energy and hybrid vehicles, through a two-way digital communication grid between power generation and load or power distribution. (Reproduced from Yang et al. [11]. Copyright 2011, American Chemical Society)

2.2

Electrochemical Energy Storage Technology

According to different working mechanisms, electrochemical energy conversion and storage devices can be divided into batteries and electrochemical capacitors [13]. Ion/electron transfer is utilized in the procedure of discharging/charging to realize the release/storage of energy in the process of working electrochemical energy release and storage [11]. Batteries are mainly divided into primary battery (such as Zn-Mn dry batteries), secondary battery (such as nickel-chromium batteries

2.2 Electrochemical Energy Storage Technology

19

or LIBs), and emerging rechargeable batteries (such as ZIBs and SIBs). Supercapacitors include two patterns: pseudo-capacitive charge-storage capacitors and electric double layer capacitors (EDLCs) [14]. The pseudocapacitors are based on a fast Faraday process, which involves electrochemical oxidation-reduction reactions, while EDLC is based on an electrolyte and an electrode charge adsorption/desorption at the interface [15]. In particular, batteries represent a viable energy storage technology for integrating renewable resources and providing stable energy to the grid (Fig. 2.1) [16]. The compact size of the battery system makes it an ideal choice for applications. In distributed locations (microgrid or off-grid), they can achieve energy storage for solar output and then applied for electric vehicles. Storing electricity is used for residential use or to electrify the entire rural community. Grid-connected batteries are used to manage and store peak load energy generated by fluctuations in load level output in wind farms and large photovoltaic power stations. When the consumption exceeds the normal output, the storage library can be used to supply power to the grid and/or for off-peak use, such as for EV charging.

2.2.1

Batteries

Battery systems that have been extensively studied and recognized by the industry mainly include flow batteries, lead-acid batteries, alkali metal-ion batteries, nickelchromium batteries, sodium-sulfur batteries, and zinc-air batteries. Their characteristics are shown in the Table 2.1. These energy storage batteries are based on electrochemical energy storage systems. Energy is reversibly converted between electrical energy and chemical energy, and this process is accompanied by a certain energy conversion efficiency and some physical changes. Lead-acid batteries are currently one of the most extensively used and most mature energy storage batteries [17]. Since they were discovered in 1860, they have been used as commercial electrical energy storage equipment and internal combustion engine starting power sources. Its basic working principle is that

Table 2.1 Characteristics of typical energy storage batteries [17] Battery type Lead-acid batteries Nickel-chromium batteries Flow batteries Sodium-sulfur batteries Zinc-air batteries LIBs

Energy density (Wh L-1) 50 ~ 90 15 ~ 150

Power density (W L-1) 10 ~ 400 80 ~ 600

Lifetime (years) 5 ~ 15 10 ~ 20

Capital cost ($ kWh-1) 100 ~ 400 800 ~ 2400

20 ~ 30 150 ~ 300

300%) causes the material to pulverize and fall off from the current collector, thus causing a decrease in riding ability.

2.2.1.2

Sodium-Ion Batteries and Potassium-Ion Batteries

Since Sony Company first developed LIBs in the 1990s, they have become the most common portable electronic product energy storage device and the most potential large-scale energy storage device. The introduction of LIBs into the automobile market as a power device for hybrid electric vehicles and pure electric vehicles can greatly reduce the dependence on fossil fuel energy. As an important component of LIBs, the distribution of lithium in the earth’s crust is extremely uneven, and the reserves are extremely low (crust enrichment content Li/Li+ (2.79 V) > K/K+ (-2.88 V); low potential is conducive to increasing the energy density of the battery [66, 67]. 3. Good rate performance: Compared with Na+ ion and Li+ ion, although K+ ion has the largest ion radius, it exhibits the smallest Stokes’s radius in PC and aqueous solutions, the diffusion rate of K+ ions in the electrolyte is the fastest, and the molar conductivity is the highest. Density functional theory (DFT) calculation results show that in PC solutions, potassium ions have low desolvation energy. In terms of cation conductivity and solvent removal performance, potassium-ion batteries are expected to have better rate performance than LIBs and SIBs [68]. 4. Good safety: Potassium metal has a very low melting point, only 63.4 °C [69]. When the potassium dendrite grows and produces a short circuit between the anode and the cathode, the potassium dendrite melts with joule heat, cutting off the short circuit before the cell gets seriously out of control. In addition,

2.2

Electrochemical Energy Storage Technology

31

potassium and sodium metal can form potassium-sodium alloy, which is liquid at room temperature, so as to avoid the generation of dendrite [70]. More importantly, compared with SIBs, K+ ions can be efficiently embedded in the graphite anode, which is unmatched by SIBs, and it turn reduces the difficulty of industrialization of PIBs to a certain extent. Due to the large radius of K+ ions and the high atomic weight, the cathode material inevitably encounters the problems of low reversible capacity, large volume changes during operation, and poor cycle stability. Researchers have done numerous studies on the selection of cathode materials, currently mainly focusing on Prussian blue and its analogs, organic materials, layered transition metal oxides, and polyanionic compounds [71–73]. Due to the relatively high chemical activity of metallic potassium, like metallic lithium, it will react with the electrolyte to form a solid electrolyte interphase (SEI) film. During the discharging and charging procedure of the battery, the dissolution and growth of potassium will lead to the growth of potassium dendrites and the rupture of the SEI film. The rupture will result in the continuous production of new SEI layer on the electrode surface and consumption of electrolyte; the electrode polarization will increase and eventually lead to capacity degradation. Therefore, metallic potassium is not suitable for use as an anode material. It is necessary to find suitable negative electrode materials. At present, research mainly focuses on alloybased materials, titanium-based materials, carbon materials, and so on [74–76].

2.2.1.3

Rechargeable Multivalent Batteries

In addition to the aforementioned alkali metal-ion batteries, multivalent metal-ion batteries represented by MIBs and aqueous ZIBs have also received widespread attention in recent years [77]. This type of emerging secondary battery is characterized by replacing the monovalent alkali metal ions with divalent or trivalent metal ions [78]. Therefore, when the same amount of metal ions is embedded in the electrode material, the specific capacity of the multivalent ion battery will be doubled [79]. For example, when 1 mol Li+ ion is inserted into the electrode material, it corresponds to only 1 mol of electron transfer, while when 1 mol of Mg2+ ion is inserted into the electrode material, the corresponding electron transfer number is 2 mol, and the corresponding specific capacity will be double. In addition, since metallic magnesium is not so active compared to alkali metals (metallic Li, Na, K), it can exist stably in the air [80, 81]. At the same time, related studies have shown that the deposition and separation of Mg2+ ions on the surface of metallic magnesium will not produce branches like lithium. Therefore, the use of metallic magnesium as an anode basically does not cause safety problems [82]. Therefore, MIBs can be directly based on metal magnesium as the anode, which can effectively avoid the problems caused by the matching of cathode and anode in alkali metal-ion batteries, and they also provide more space for the choice of cathode materials. In addition, the metal magnesium itself has a very high capacity (2205 mAh g-1, 3833 mAh mL-1),

Basic Information of Electrochemical Energy Storage 10

8000

8

6000

6

4000

4

2000

2

20

15

10

5

Cost (USD kg–1)

Gravimetric Capacity Volumetric Capacity Abundance Cost

Abundance ( × 1000 ppm)

2

Capacity (mAh g–1)-(mAh mL–1)

32

0.1 0

Li Na

K Mg Ca Zn

0

0

0

0.0

Al

–3.0

20

–2.0

30 40

–1.5 50 –1.0 –0.5

Standard Potential Atomic Weight Cation Radius

0.5

1.0

Cation Radius (A)

–2.5

Atomic Weight

Standard Potential (V)

10

60 70

1.5

Fig. 2.8 Comparison between atomic weight, cation radius of metal anodes, gravimetric and volumetric capacities, cost, heart crust abundance, and standard potential (vs. SHE). (Reproduced from Zhang et al. [91]. Copyright 2018, Wiley-VCH)

in which the volume specific capacity is nearly twice that of lithium [83] (Fig. 2.8). Based on the above advantages, magnesium-ion batteries have become another hot spot for electrochemical energy storage in recent years [78]. However, the challenges faced by MIBs are also very large. The main challenges include two aspects: First, MIBs currently lack a mature electrolyte system. For conventional carbonate electrolyte, the electrolyte will decompose at low potential and form the SEI layer deposited on the surface of magnesium metal [84]. According to related research, due to the bivalent nature of Mg2+ ions, once the SEI layer is formed, Mg2+ ions will not be able to pass through the SEI layer to continue the reaction. Once the electrochemical reaction is stopped, the formation of SEI passivation film must be prevented in the MIBs, which means that a new electrolyte system must be developed, and this in turn also brings greater challenges to the development and commercialization of MIBs [85]. Although a variety of MIB electrolytes have

2.2

Electrochemical Energy Storage Technology

33

been developed, they are not yet fully mature, and related research needs to be further promoted. Second, although the radius of Mg2+ ions is about the same as that of Li+ ions, Mg2+ ions are divalent ions with two positive charges, which allows Mg2+ ions to diffuse in the cathode material, especially in the metal oxide electrode material, due to the strong electrostatic interaction with oxygen [86]. The diffusion barrier of Mg2+ ion is significantly larger than Li+ ion; hence, the diffusion kinetic properties of Mg2+ ion are often very poor, and it is difficult for common electrode materials to achieve rapid diffusion of magnesium ion. The development of a suitable cathode material system is a key part of the development of MIBs [77]. Transition metal sulfides are an important class of anode materials for MIBs, such as TiS2, MoS2, and Mo6S8. Since the electrostatic interaction between Mg2+ ions and sulfur is weaker than the electrostatic interaction with oxygen, the diffusion kinetic properties of Mg2+ ions in transition metal sulfides will be better, and the corresponding electrochemical performance will be better [87]. Another important type of cathode material for MIBs is the layered vanadium oxide aerogel. The characteristic of this type of material is that the layer spacing is very large, reaching more than 10 Å, which is more than twice that of ordinary layered V2O5. At the same time, the layers contain crystal water. The large interlayer spacing provides enough space for the diffusion of Mg2+ ions. In addition, the crystal water between the layers can surround the Mg2+ ions after the Mg2+ ions are intercalated to produce a charge shielding effect on the Mg2+ ions, thereby effectively weakening the Mg2+ ions and the structure. The electrostatic effect of oxygen improves magnesium storage performance [88–90]. ZIB is another kind of divalent metal-ion battery. Due to the relatively high metal zinc reduction potential (0.8 V vs. SHE), Zn2+ ions in the general electrode materials tend to take off the potential of the embedded lower than the level of zinc anode. Therefore, ZIB in organic electrolytic liquid is limited by the output voltage, without a too big advantage compared with other systems (Fig. 2.8) [92–94]. Utilizing an aqueous system electrolyte for a battery is a judicious selection owing to several advantages such as high ionic conductivity and low flammability. Among the anode options for an aqueous system battery, the metal zinc stands out due to its remarkable stability in this environment. This characteristic combined with the unique advantages of aqueous battery, such as environmental friendliness, low price, and simple preparation process, makes the aqueous ZIB to have a very large commercial prospect and a wide application prospect in large-scale energy storage and other fields [95]. In the past 2 years, aqueous ZIBs have also attracted wide attention and became a new upsurge of research [96]. In recent years, research on aqueous ZIBs have mainly focused on the development of suitable high-performance cathode materials and revealing the storage mechanism of Zn2+ ions in aqueous environments. At present, the research on anode materials for aqueous ZIBs have made significant progress, and many materials have been reported to have excellent zinc storage properties [96–99]. AIB is a kind of trivalent metal-ion battery, which has also received some attention in recent years. Aluminum is the third most abundant element in the earth’s crust, after silicon and oxygen, and is the most abundant metal element. Metal aluminum is stable in the air and has no safety hazards. It can be directly used as

34

2 Basic Information of Electrochemical Energy Storage

the anode of AIBs. Al3+ ion is a +3 metal ion, which means that an Al atom becomes an Al3+ ion corresponding to the transfer of three electrons. In addition, aluminum has a lighter atomic weight (26.98 g mol-1) and a high density (2.7 g cm-3). These properties make metal aluminum have a very high volumetric specific capacity (~8040 mAh cm-3) when used as a negative electrode, which is about four times that of metallic lithium and more than twice that of metallic magnesium [100–103] (Fig. 2.8). However, when metal aluminum is used as an anode material, its reduction potential is higher, which is -2.7 V vs. SHE, which is 2.3 V higher than metal lithium and 0.7 V higher than metal magnesium, which means that the same cathode material and the same anode polarization are used [104] (Fig. 2.8). Under the circumstances, the output voltage of batteries based on metal aluminum as the negative electrode will be relatively low, and most of the reported discharge voltages of AIBs are generally less than 2.0 V [105]. In addition, the research and development of AIB cathode materials is a big challenge. Due to the high surface charge density caused by the small radius and three charges of Al3+ ions, a noticeable polarization effect takes place. Consequently, the diffusion of Al3+ ions within the electrode material becomes quite difficult. Moreover, the lattice of certain electrode materials is subject to a potent electrostatic effect of anions. This effect brings a significant challenge to the diffusion of Al3+ ions in the anode material [106, 107]. There are very few cathode materials reported to achieve reversible Al3+ ion insertion and extraction. V2O5 and VS2 are one of the few reported cathode materials that can realize reversible Al3+ ion storage, but the cycle performance is still difficult to meet the requirements [108, 109]. In addition to the battery system mentioned above, the CIBs are also a new type of battery system that has gradually attracted people’s attention [110]. The potential advantages of CIBs are as follows: First, the calcium content is also very abundant in the earth’s crust, and it is higher than sodium, potassium, and magnesium. Second, similar to MIBs, it also belongs to the divalent metal-ion battery. The insertion of one Ca2+ ion corresponds to the transfer of two electrons, which has potential highcapacity characteristics. Third, compared to metallic magnesium, the reduction potential of Ca (-2.87 V vs. SHE) is closer to metallic lithium, which means that CIBs can achieve potentially high output voltages, resulting in higher energy density than MIBs (Fig. 2.8) [111]. Fourth, because the radius of Ca2+ ions is larger than that of Mg2+ ions, and the charges are the same, the surface charge density of Ca2+ ions is smaller than that of Mg2+ ions, and the polarization intensity is smaller. Therefore, when Ca2+ ions diffuse in the electrode material, they are subjected to static electricity. It is smaller than Mg2+ ions, which means that for the same electrode material, when the diffusion channel is large enough, the diffusion barrier for the diffusion of Ca2+ ions is smaller than that of Mg2+ ions, and the diffusion kinetic properties of Ca2+ ions are more important. Although CIBs have many advantages as described above, the research progress of CIBs is very slow compared to other battery systems, which is mainly due to the difficulty of metal calcium as the anode of CIBs [112]. In other battery systems, such as alkali metal-ion batteries, MIBs, ZIBs, and AIBs, the corresponding elemental

2.2

Electrochemical Energy Storage Technology

35

metals can be used as counter electrodes, which greatly facilitate the development of basic research. However, metallic calcium has been proven unable to play the important role of counter electrode [113–115]. The main reasons are as follows: (1) According to previous studies, in most organic electrolytes, the deposition of Ca2+ ions on the surface of metal calcium cannot be carried out normally, which is mainly due to the SEI passivation layer that will hinder the passage of Ca2+ ions. (2) Metal calcium is hard, unlike alkali metals, metallic magnesium, and metallic zinc, which have excellent ductility, and it is difficult to perform simple processing on its shape; hence, this brings about the assembly of the battery in the process of basic research. It is extremely difficult, which is not conducive to the development of basic research. (3) Due to the two reasons mentioned above, the research on the electrolyte and electrode materials of CIBs has been progressing slowly, and there is still no more suitable and mature electrolyte and electrode material system.

2.2.1.4

Flow Batteries

In addition to the several battery systems described above, there are many battery technologies that are still being developed and perfected, such as flow batteries. Flow battery is a redox cell with a circulating liquid state active material (Fig. 2.9a) [116]. Different systems such as Cr series, V series, and Br series have been developed successively. Flow batteries have many unique features: (1) The reactive materials of flow batteries are all located in the electrolyte and only appear in liquid form. There is no solid phase change when the reaction occurs; hence, the active materials will have a relatively long theoretical life span, while the degree of polarization is small. (2) The design of flow battery can be very flexible, because its power is only affected by the size of the stack. In addition, the capacity is only associated to the total amount of reactive materials; hence, enhancing the capacity of the flow battery can be achieved by promoting the liquid storage tank or increasing the electrolyte concentration. (3) There is no irreversible reaction in the liquid flow battery, even if it is 100% discharge, there will be no damage, and the response speed

a

End Plate

Membrane

Graphite Plate

b

V

Inlet Inlet

V4+/V5+

Outlet

Outlet

Silicon gasket

Copper Current Colector

Current collector

+

Pump

–

Charge

Charge

V4

+

V5

+

V2+

V3+

V3+/V2+

Discharge Discharge Membrane

Electrode

Carbon/graphite felt

Fig. 2.9 (a) Exploited view of the single flow cell. (b) Schematic configuration of the working principle of an all-vanadium flow battery. (Reproduced from Ulaganathan et al. [120]. Copyright 2015, Wiley-VCH)

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of the system is very fast. (4) Flow battery has the advantages of green and environmental protection. Most of its constituent materials are carbon and plastics. Such raw materials have the merits of long working time, low cost, and convenient handling, and will not cause environmental pollution. (5) The operation of the flow battery energy storage system is carried out under closed conditions, will not cause leakage and pollution, and has a high degree of freedom in design and selection. It is one of the first choices for large-scale energy storage technology [117–119]. The all-vanadium redox flow battery realizes electric energy-chemical energy through the conversion of the valence state of vanadium ions between the anode and the cathode, specifically the mutual conversion between the two sets of different electric pairs of V2+/V3+ and V5+/V4+ in the sulfuric acid solution of the two electrodes [11]. The working mechanism of the all-vanadium flow battery is presented in Fig. 2.9b. The battery unit is mainly composed of current collector (electrode plate), electrode material (carbon felt), and ion exchange membrane in sequence and symmetrically [121]. The ion exchange membrane divides the battery into two half-cells, and each half-cell is injected with vanadium electrolytes of different valences, in which V5+/V4+ sulfuric acid solution is injected into the positive electrode, and V2+/V3+ sulfuric acid solution is injected into the negative electrode of the battery [122]. When the battery is charged, the external electrolyte is circulated between the battery and the external storage tank under the push of the circulating pump, and the V4+ ions at the cathode lose electrons and become V5+ ions. Electrons flow from the cathode to the anode from the external circuit, and the V3+ ions of the negative electrode of the battery become V2+ ions. The H+ ions in the electrolyte move from the cathode to the anode through the membrane, forming a complete closed circuit. When discharged, the ions move in the opposite direction, the V2+ of the negative electrode loses electrons and turns into V3+, and the V5+ of the positive electrode turns into V4+, which completes the battery reaction and realizes energy storage and release [123]. The electrodes and battery reactions of the all-vanadium flow battery are as follows: Negative electrode: V3þ þ e $ V2þ

ð2:7Þ

VO2þ þ H2 O - e $ VO2 þ þ 2Hþ

ð2:8Þ

Positive electrode:

Overall battery reaction: VO2þ þ V2þ þ H2 O $ VO2 þ þ V3þ þ 2Hþ

ð2:9Þ

The standard potential of the cathode is 2.004 V, and the standard potential of the negative electrode is -0.255 V; hence, the total reaction standard potential of the all-vanadium flow battery is 2.259 V [124]. Equivalent to other types of flow

2.2

Electrochemical Energy Storage Technology

37

batteries, all-vanadium energy storage batteries have many advantages: (1) The power and the capacity depend on the size of the battery stack and the concentration and volume of the negative ion electrolyte. (2) There is only one effective chemical substance in the electrolyte of positive and negative electrodes—vanadium, and there is no cross-contamination problem in the electrolyte caused by ion penetration in the battery work. (3) The concentration polarization is small, and the active material exists in liquid form. In the storage tank, the active material flows continuously between the battery and the storage tank through the circulating pump, so that the concentration difference of the electrolyte is minimized. (4) The electrolyte can be recycled and reused, saving resources; vanadium ions are safe and environmentally friendly; and vanadium compounds are inexpensive and easy to popularize and use [125–127].

2.2.2

Supercapacitors

Supercapacitor is an energy storage apparatus between the traditional capacitor and the battery [128]. It has high power density (102 ~ 106 W kg-1), high energy density (~10 Wh kg-1), high charge-discharge efficiency (>99%), long life (>100,000 times), and other advantages, which are widely applied in electronic equipment, transportation and military equipment, and other fields [129–131]. The concept of supercapacitor originated in 1957 and was proposed in a patent by H.I. Becker. In view of the large specific surface area of carbon materials, H.I. Becker used carbon materials as electrode materials. The capacitance of the capacitors made by H.I. Becker is three to four orders of magnitude higher than that of ordinary capacitors, reaching the Farad level. Therefore, people call this kind of capacitors “super capacitors.” In 1962, the American Ohio Mobil Petroleum Company used sulfuric acid aqueous solution as the electrolyte and activated carbon as the electrode material to obtain a 6 V supercapacitor [132, 133]. It was the first to commercialize it in 1969, but it did not get the attention of researchers at that time. It was not until the 1990s that supercapacitors emerged in hybrid vehicles [134]. Energy storage devices include traditional capacitors, super capacitors, batteries, and fuel cells. The comparison between them is shown in Fig. 2.10. It is shown in the figure that the power density of super capacitors is lower than that of traditional capacitors, and the energy density is higher than that of batteries and fuel cells [135]. The overall performance is between traditional capacitors and batteries, which can make up for the gap between the two [136]. In view of the characteristics of supercapacitors, in terms of application, it can be used as an independent energy storage device to replace traditional batteries or fuel cells, utilized in situations requiring high power density, and can be combined with traditional batteries and fuel cells to form a composition hybrid power system [137]. So far, after decades of development, supercapacitors have been widely used in many fields and have great development potential.

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Power density (Wh/kg)

10000000 1000000

Capacitors

100000 10000 1000 100

Supercapacitors

Batteries

10 1 0.01

0.1

1

10

100

Fuel cells 1000

10000

Energy density (Wh/kg) Fig. 2.10 Ragone plot for the various electrochemical energy conversion systems. (Reproduced from Inganäs and Admassie [138]. Copyright 2013, Wiley-VCH) Table 2.4 Performance comparison between supercapacitors and batteries Performance Charging time (s) Discharging time (s) Power density (W kg-1) Energy density (Wh kg-1) Cycle life (cycles)

Supercapacitors 1 ~ 30 1 ~ 30 1000 ~ 2000 1 ~ 10 >1,000,000

Batteries 3600 ~ 18,000 18 ~ 180 50 ~ 300 20 ~ 180 500 ~ 2000

By linking the performance comparison between the super capacitor and the battery (Table 2.4), it can be enunciated that the super capacitor has the following advantages: 1. High capacity: The capacity of supercapacitors can reach thousands of farads, which is thousands of times higher than that of tantalum electrolytic capacitors and aluminum electrolytic capacitors of the same volume. 2. Long cycle life: The capacitance of supercapacitors has little attenuation, and the cycle life can reach hundreds of thousands of times, which is 5–20 times the number of cycles of storage batteries. 3. Short charging time: Supercapacitors are charged with high current and can be quickly charged in a few seconds to a few minutes. Even if the battery is charged quickly, it takes tens of minutes, and frequent fast charging will affect the service life. 4. High power/energy density: While supercapacitors can provide a 1000 ~ 2000 W kg-1 power density, they can also output a 1 ~ 10 Wh kg-1 energy density at the same time. 5. Wide working temperature range: The operating temperature range of supercapacitors is -40 ~ 70 °C, while the temperature range of general batteries is between -10 and 50 °C.

2.2

Electrochemical Energy Storage Technology

39

Fig. 2.11 Schematic of (a) EDLC and (b) pseudocapacitor. (Reproduced from Chen et al. [149]. Copyright 2017, Oxford University Press.) Schematic diagram illustrating the operation principles of (c) EDLC and (d) pseudocapacitor. (Reproduced from Meng et al. [150]. Copyright 2017, Elsevier)

6. Reliable operation, maintenance-free, and environment-friendly: The super capacitor has a certain degree of anti-overcharge ability, which will not have much influence on its work in a short time, and in turn can ensure the reliability of system operation.

Supercapacitors are mainly composed of electrodes, current collectors, electrolytes, and diaphragms [139]. According to the different energy storage mechanism of the device, supercapacitors are divided into pseudocapacitors and electric double layer capacitors (Fig. 2.11a, b).

2.2.2.1

Electrical Double Layer Capacitors

EDLCs use the electric double layer interface between the electrolyte and the electrode to accumulate static charge to store energy [140–143]. Its working principle is shown in Fig. 2.11c. During the discharging/charging procedure, electrons

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only enter and exit on the surface of the electrode through the external circuit. The cations and anions in the electrolyte move in the solution [144]. There is no chemical reaction on the electrode, that is, no Faraday process is involved. Therefore, EDLC is highly reversible. The discharging and charging processes of EDLC describe by formulas as follows: One electrode: ES1 þ A - $ ES1 þ ==A - þ e -

ð2:10Þ

ES2 þ Cþ þ e - $ ES2 - ==Cþ

ð2:11Þ

ES1 þ ES2 þ A - þ Cþ $ ES1 þ ==A - þ ES2 - ==Cþ

ð2:12Þ

The other electrode:

Overall capacitors:

Among them, ES1, ES2 represent the surface of the two electrodes; A-, C+ represents the cation and the anion in the electrolyte; and // represents the interface between the electrode and the electrolyte. It can be known from the working mechanism of EDLCs that the electrodes are mainly composed of carbon materials with large specific surface area and good chemical stability and conductive nanomaterials, such as metal oxide nanotubes, graphene, activated carbon, and carbon nanotubes [145–148].

2.2.2.2

Pseudocapacitors

The working principle of a pseudocapacitor is shown in Fig. 2.11d. It stores energy through the oxidation-reduction reaction on or near the electrode material surface, and it is related to the Faraday process; hence, it is also called a Faraday capacitor [151]. Conway divides pseudocapacitors into three categories according to the Faraday process that produces the electrochemical characteristics of capacitance: one is underpotential deposition pseudocapacitance, the other is redox pseudocapacitance, and the third is embedded pseudocapacitance [152]. Compared with EDLCs, the electrochemical energy storage process of pseudocapacitors can be extended to the inside of the electrode material, that is, the amount of substances involved in energy storage is greatly increased, making the pseudocapacitors have higher specific capacitance and energy density [153]. On the other hand, under the control of solid-state bulk diffusion, the energy storage and release rate of pseudocapacitors is greatly reduced, resulting in lower power density, and accompanied by the volume increase of the electrode material, which reduces the cycle stability [154].

References

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The electrode materials of pseudocapacitors mainly include metal oxides and conductive polymers. Typical metal oxide electrode materials include RuO2, MnO2, NiO, Co3O4, VO2, V2O5, and SnO2. Conductive polymers include polyaniline, polypyrrol, and polythiophene [139, 155–157].

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Chapter 3

Synthesis of Vanadium-Based Nanomaterials

3.1

Introduction

Since there are diverse crystalline phases (e.g., V2O5, VO2, LiV3O8, and Li3V2(PO4)3) and oxidation states (e.g., V3+, V4+, and V5+), the controllable synthesis of vanadium-based materials with specific phase, morphology, and desirable properties is essential for the research and applications of this material family [1, 2], especially for the nanostructured materials with unique chemical and physical properties [3]. Two important nanoscale parameters, namely, shape and size, greatly influence their properties, including the electron-transport process, the density of energy states, and band-gap energy [4]. The ability to manipulate the size and shape of materials offers researchers the facility to regulate their properties and thereby offer possibilities for better and newer applications in many frontier fields [5]. With the advancement in materials science, the creation and evolution of nanostructured materials range from simple to complex [6, 7]. During these multilevel architectures, the interior spaces and multiphase interfaces provide the microscopic chemical environment and stimulate the intrinsic reaction activity, bringing many potential applications [8–10]. Therefore, various synthetic strategies were developed to design and construct vanadium-based nanomaterials with delicate morphologies and abundant compositions. Generally, these strategies can be mainly categorized into bottom-up and top-down methods. The top-down approach is based on high-energy physical tools (e.g., milling and sputtering) to reduce bulk materials into tiny nanomaterials [11, 12]. However, obtaining the nanomaterials with morphology and desired size is a great challenge due to their high-energy conditions. On the contrary, the bottomup approach involves the self-assembled chemical processes from single-molecule components to delicate nanostructured materials. In the previous report, the reaction processes of the bottom-up approach are determined and controlled by thermodynamics and kinetics during the reaction systems, including nucleation and growth. The bottom-up approach has become more popular because of energy savings, © Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_3

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3 Synthesis of Vanadium-Based Nanomaterials

convenient manipulation, and easy control over size and shape. Therefore, this part focuses on the bottom-up approach and chemistry for synthesizing vanadium-based nanomaterials, mainly including the hydro-/solvothermal method, electrospinning, sol-gel route, template synthesis, and vapor deposition techniques. A comparison of the current main synthetic methods for nanostructured vanadium-based materials is shown in Table 3.1. The corresponding synthetic approaches are discussed in detail as follows.

3.2

Hydro-/Solvothermal Method

Hydro-/solvothermal methods are common and promising bottom-up approaches to designing and constructing well-defined nanostructured materials because of its low cost, low air pollution, effortless control, uniform products, and simple manipulation [34]. The hydrothermal method refers to preparing materials in sealed pressure vessels with water as solvent and powder dissolved and recrystallized. Compared with other powder preparation methods, the powder obtained by the hydrothermal method has the advantages of complete grain development, small particle size, uniform distribution, lighter particle agglomeration, cheaper raw materials, easyto-obtain appropriate stoichiometry, and crystal shape. In particular, the ceramic powder prepared by the hydrothermal method does not need high-temperature calcination, which avoids the grain growth, defect formation, and impurity introduction caused by the calcination process; hence, the prepared powder has high sintering activity. The hydro-/solvothermal processes mainly undergo three steps: (1) the rapid increase of monomer concentration to super-saturation levels in the solution and then the aggregation of monomers into seeds; (2) crystalline growth of monomers by continuous aggregation on the seeds, along with the reduction of the monomers concentration; and (3) surface steadiness of the products in the reaction condition (Fig. 3.1) [4]. These processes are driven by the decrement in the Gibbs free energy. In addition, the processes can be readily adjusted by managing the interior reaction conditions (such as templates, organic additives, monomer concentration, time, pH value, and pressure) [34]. For example, under different pH values, the corresponding solution states of V (V) ions varied, including VO43- (pH: >12.6), V2O6(OH)3- (pH: 12 ~ 9), V3O93- (pH: 9 ~ 7), V10O286- (pH: 7 ~ 6.5), V2O5
xH2O (pH: 6.5 ~ 2.2), and VO2+ (pH: 12.6 VO43-

12 ~ 9 V2O6(OH)3-

9~7 V3O93-

7 ~ 6.5 V10O286-

6.5 ~ 2.2 V2O5
xH2O

2.2 ~ 1 V10O286- $ VO2+

x ≥ 0.04), C2 (0.25 > x ≥ 0.09), and C3 (0.54 > x ≥ 0.25) in charging. Figure 4.2c quantitatively analyzes the variation rate of the crystal parameters at each stage and gives three characteristics as follows: (1) The entire charging and discharging process is reversible. At discharging, the evolution processes experience through VO2, VO2-Zn0.07, VO2Zn0.29, and VO2-Zn0.54. In contrast, the reactions of VO2-Zn0.54, VO2-Zn0.25, VO2Zn0.09, and VO2-Zn0.04 were observed upon charging. Especially the variation of lattice parameters in D3 is completely opposite to C3; in addition, D2 and D1 are the same as C2 and C1, respectively. (2) The variation rate of c and a roughly follows the

92

4

In Situ Characterizations of Vanadium-Based Nanomaterials

same laws at each stage. For example, state D3 expands by 0.5% in c and 0.1% in a; state D2 expands by 1.6% in c and 2.3% in a. (3) The most significant variations in the crystalline structure occur at states D2 and C2. The volume shrinkage at stage C2 is 3.6%, and the volume expansion at stage D2 is 4.5%. Meanwhile, the shrinkage and expansion proceed along directions of c and a, and the direction of b is hardly varied in stages C2 and D2. Therefore, the transfer of Zn2+ in VO2 is along the [010] crystalline direction. The theoretical model of Zn2+-inserted sites in VO2 crystalline is displayed in Fig. 4.2d. It is the first report that when Zn2+ extracted/inserted in ZIB, the crystal parameters of the electrode material can be systematically researched through in situ measurements and corresponding quantitative calculations.

4.2.2

In Situ Raman Spectroscopic Characterization of Vanadium-Based Nanomaterials

Raman spectrum is a scattering spectrum; by analyzing the frequency difference between incident light and scattered light, the molecular vibration and rotation could be studied [31]. The composition of in situ Raman mold battery is shown in Fig. 4.3a, almost the same as the in situ XRD battery mentioned before, but the sapphire window replaces the Be (Al Cu) window for the penetration of the light source (the wavelength of the light source is 532 nm or 633 nm) [33]. Besides, for better encapsulation, Vaseline is evenly applied on the copula of the mold battery. Figure 4.3b shows the mechanism of Raman spectra. After the monochromatic light Fig. 4.3 (a) Scheme of the in situ cell. (1) Stainless steel body contacting the working electrode; (2) holder for the graphite working electrode (the graphite is deposited on top of the holder); (3) current collector for the lithium counter electrode; (4) counter-electrode (lithium foil); (5) sapphire window; (6) retainer ring; (7) cell body made of (poly)propylene. (b) Scheme of the Raman spectra. (Reproduced from [32]. Copyright 2016, American Chemical Society)

4.2 In Situ Spectroscopic Characterizations of Vanadium-Based Nanomaterials

93

passes through light filter to interact with the sample, the incident light can be reflected, absorbed, or scattered. The scattered light can reveal the molecular structure of the sample. By analyzing the frequency (wavelength) of the scattered light, it can be found that there are not only components with the same wavelength as the incident light (Rayleigh scattering) but also scattered light with a small number of wavelengths changed (stokes and anti-stokes Raman scattering), and the Raman scattering light intensity is about 10-1 of the total scattering light intensity. It is Raman scattering light whose changed wavelength gives us the information of structure and chemical composition for the sample. The in situ Raman spectroscopy is a kind of powerful test to investigate the structure change during the electrochemical process for vanadium-based materials. Zhang et al. studied the chemical bond distance between V-O and V-V of V2O5 in the Li/V2O5 battery after the charge/discharge process by in situ Raman spectra [34]. Figure 4.4a shows the molecular structure of the V2O5 and lists three different V-O distances. Figure 4.4b shows the Raman spectrum of the V2O5; all the vibration bands are well attributed to V2O5. The Raman band at 980 cm-1 is corresponds to the vibration of the shortest V-O bond along the directions of b. In situ Raman spectra were carried out for four charging/discharging cycles within 2.0–4.0 V at

Fig. 4.4 (a) The schematic illustration of V2O5. (b) Raman spectrum of V2O5 nanowire. (c) In situ Raman spectrums of V2O5 during voltage window of 2.0–4.0 V for 4 cycles at a current density of 100 mA g-1. (d) The fitting curves of three different Raman intensity decay bands. The exponential decay function is applied to fit the curve, while the electrochemical process follows the exponential decay trends. (Reproduced from [34]. Copyright 2019, Springer Nature)

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100 mA g-1 (Fig. 4.4c). Three significant Raman vibration band variations of V2O5 are corresponded to the ν (d1), ν (d3), and б (O3-V-O2). Figure 4.4c shows the peak intensity changes at three selected important band regions upon electrochemical process. As the cycle progresses, the intensity of the peak decreases rapidly. It suggests that the degree of crystallinity at local structure is also decreased, which may be related to the capacity decay fading. In order to investigate the fading, Zhang et al. fitted the three profiles as shown in Fig. 4.4d. The calculated results of the left and right band at 980 cm-1 fit well with the trend of capacity decay and could be regarded as a symbol of capacity decay. In addition, the Raman shift represents the shortest V-O bond of 1.58 Å, which means the V-O bond becomes longer and more dispersed during the cycle.

4.2.3

In Situ XANES Characterization of Vanadium-Based Nanomaterials

In order to investigate the variation of the vanadate during cycling, XANES were performed after 10 cycles. Figure 4.5a is the V K-edge XANES spectra of the initial electrode material and over 10-times cycling. The V symmetry environment shows a

Fig. 4.5 (a) Vanadium K-edge XANES spectra of CoV2O6, after 10th discharge and charge. (b) Cobalt XANES spectra of CoV2O6, after 10th discharge and charge. (c) Moduli of the vanadium K-edge Fourier transform of CoV2O6 after 10th discharge and charge. Distances have not been corrected from phase shift. (d) Moduli of the cobalt K-edge Fourier transform of CoV2O6 after 10th discharge and charge. Distances have not been corrected from phase shift. (Reproduced from [35]. Copyright 2001, Elsevier)

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In Situ Microscopic Characterizations

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pyramidal-shaped with the lower energy oxidation at 3.5 V, which suggests that the V is gradually reduced to a valence close to +2. The EXAFS spectra (Fig. 4.5b) present that in both cases, the distances of V-Co or/and V-V located in the 2–3 Å region (EXAFS is difficult to differentiate atoms with similar atomic number) is not influenced by electrochemical cycling. As the amount of reactive Li+ increases, only the contribution of the first neighbor of V-O moves to a shorter length. Except in situ XRD and Raman spectra, which are mentioned previously, there are also other powerful characterizations, such as in situ XAFS and XANES, used when studying vanadium-based nanomaterials. For example, Laruelle et al. used the XANES tool to investigate the electronic and structural modifications of CoV2O6 during the cycle process in a CoV2O6/Li battery [35]. The XANES spectra of the Co K-edge (Fig. 4.5b) indicates that the valence of Co changes between 0 and 2 reversibly. Figure 4.5d shows that the peaks near 1.6 Å undoubtedly correspond to the Co-O length. The second peaks at 2.6 Å are derived from the second nearest neighbors of Co-Co and ultimately to the Co-V length. Then the spectrum consists of a single peak at 2 Å at 0.02 V. This length in Co metal cluster is consistent with the XANES spectrum, implying the valence of 0 for Co. Therefore, the V local environment is composed of V and O atoms independently of the charging state of CoV2O6/Li battery.

4.3

In Situ Microscopic Characterizations

To investigate the process and mechanism of intercalation/deintercalation of lithium ion in cathode or anode active materials, in situ characterizations play an important role to achieve real-time monitoring. As described before, in situ spectroscopy characterizations such as in situ Raman spectrum could demonstrate the valence bond effect or element status information. Also, there are other in situ characterizations such as in situ X-ray diffraction that expound the phase change during lithiation and delithiation. During the exploration of the change on the phase, morphology and mechanical characteristics, in situ microscopy characterizations especially for in situ transmission electron microscope (in situ TEM) could achieve real-time visualization of electrode materials during electrochemical reaction including the morphology [32]. Based on the more complex micro-fabrication process of in situ TEM system with high resolution, the phase change and the status of interface layers like SEI could be investigated deeply [36]. However, even though it has great potential in in situ research of LIBs, there are still some disadvantages. For example, lithium metal is sensitive to electron beams, so the electron beam intensity is not too high, which means some pictures are not very good in quality; electrochemical lithiation reaction is too fast, which often brings some difficulties to real-time monitoring. Furthermore, it is difficult to control the degree of lithiation of active materials by the change of potential. Huang et al. were the first to apply in situ TEM to observe directly the electrochemical behavior of lithiation of SnO2 in 2010 [37]. As expected, the phase change

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Fig. 4.6 (a) Schematic of the experimental setup. (b) (i) to (iv) are sequential high-magnification images showing the progressive migration of the reaction front, swelling, and the twisted morphology of the nanowire after the reaction front passed by. (c) TEM images revealed a high density of dislocations emerging from the reaction front (marked by chevron-shaped dotted lines). (d) Schematic drawing showing the high Li diffusion flux in Li2O. (Reproduced from [37]. Copyright 2010, Science)

process of SnO2 during lithiation was first researched. In Fig. 4.6a, the schematic of the experimental setup is shown and infers the requirement of one dimensional active materials. Interestingly, during lithiation, the SnO2 began to flex rapidly and formed a bend and a coil of a spiral later, which needs to be further analyzed deeply with the force and mechanical reaction (Fig. 4.6b). Due to the structure of the experimental setup, the electrolyte could be contact with the SnO2 nanowire, and the other side of SnO2 could apply the reaction potential. Thus, the phase change could be observed directly (Fig. 4.6c). The transition between the crystalline and amorphous states of SnO2 during lithiation has been proved. Combined with in situ TEM, the crystal, dislocation and amorphization were shown that there are some other deep mechanisms on the transition of them. The shape of dislocation stay is six rhombus, which means that there is reaction orientation between inner and outer of SnO2 nanowire. Herein, in situ TEM reconstructed the lithiation process of SnO2 (Fig. 4.6d). Also, this work is just the beginning in electrochemical energy storage combining with in situ TEM.

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Through the above description, it is not difficult to find that the in situ TEM can directly observe the change in volume and interface state of the electrode material during the charging and discharging process, providing theoretical guidance and technical support for understanding the properties of electrode materials [36]. Due to the environmental limitations of in situ TEM, the instability of the entire microbattery device is caused during high resolution, thus making it difficult to directly observe the specific phase change. Thus, the researchers combined the quasi-in situ TEM to observe the crystal state of electrode materials under different charge and discharge conditions. Therefore, using in situ TEM or quasi-in situ TEM, the researchers have achieved an accurate analysis of the material behavior changes of the electrode materials under different charge and discharge conditions. For vanadium-based electrode materials, in situ TEM still exhibits unique analytical advantage. By utilizing the in situ characterization of vanadium-based electrode materials, it is possible to directly observe the properties of the electrode materials, which are closely related to electrochemical properties such as charge and discharge morphology changes and volume expansion. Xu et al. synthesized alkaline-earth metal vanadates as SIBs anodes and studied the electrochemical behavior during sodium ion intercalation [38]. The research results show that Ca2V4O9 nanowires, as anode electrode material for sodium ion batteries, exhibit high conductivity and excellent specific capacity exceeding 300 mAh g-1. As a negative electrode material for metal ion batteries, in the process of deintercalation ions, the volume expansion problem of the electrode material seriously affects the structural stability, Coulombic efficiency, and the cycling stability. Therefore, Xu et al. directly observed the volume expansion of calcium vanadate nanowires during the process of intercalating sodium ions by in situ TEM, which was calculated to be less than 10%, proving the correspondence between the cycle stability and volume change of the electrode material. As we have described before, in situ TEM can not only directly observe the volume change of the electrode material during charge and discharge but also use the other functions of the transmissive electrode to achieve more observations. Xu et al. further verified the diffusion properties of sodium ions in nanowires by using electron energy loss spectroscopy (EELS), further demonstrating the results of in situ TEM (Fig. 4.7). Some of the above examples are based on single-component materials for in situ TEM studies, direct or indirect observation of volume expansion, interface shift, and ion diffusion during charge and discharge to explore the various materials in the ion battery. Tan et al. combined the carbon nanotubes and reduced graphene oxide and vanadium pentoxide to form a multidimensional synergistic nanoarchitecture [39]. The authors compare the V2O3/CNTs with V2O3/CNTs/rGO to analyze the volume change during the sodium ion intercalation process and summarize the strength and elasticity of the composite electrode material during charge and discharge, which is the relevance of structural stability. Comparing with the in situ spectrum characterization, it is hard for in situ TEM to provide the detailed valencebond structure. In situ TEM is demanding for the sample, while it could visualize the changes in the structure of the sample.

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Fig. 4.7 In situ TEM results of calcium vanadate. (a) Configuration of the in situ TEM device. Scale bar, 500 nm. (b) Time-lapse TEM images after applying a voltage bias. Scale bar, 500 nm. (c) EELS data collected from the red circle area at 0 and 35 min. (Reproduced from [38]. Copyright 2017, Nature Publishing Group)

4.4

Other In Situ Characterization

Besides in situ spectroscopy and microscopy characterization, other in situ characterization techniques are also reported, such as in situ electrical impedance spectroscopy (EIS) testing and in situ electrical conductivity testing. In order to explore the fading mechanism of electrode materials performance during the charge and discharge process, the Mai group [40] designed and constructed single nanowire devices and in situ electric conductivity testing was conducted. Figures 4.8a and 4.9a show the two configurations of single nanowire electrochemical devices. Both the configurations are designed with four contacts. Contact 1 was used to perform the electrochemical test to achieve the intercalation and deintercalation of ions, and others were employed to in situ test the electric conductivity. In Configuration 1 (Fig. 4.8a), the H2V3O8 nanowire was all exposed to the electrolyte, whereas in Configuration 2 (Fig. 4.9a), only one end was exposed to the electrolyte and the other area was covered by photoresist. The in situ electric

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Fig. 4.8 Single nanowire device in the first configuration. (a) The first configuration, in which the nanowire is fully immersed in the electrolyte. (b) The dark field optical microscopic image of the nanowire electrode. (c) The ion transport path model of the first configuration. The V3O8 layer consisting of VO6 octahedra and VO5 distorted trigonal bipyramids interconnected with each other. The pink and red are V atoms and O atoms. The ion can transport through b or c axis. (d-f) I-V curves of section 1–2, 2–3, and 3–4 before and after the electrochemical process, respectively. (Reproduced from [40]. Copyright 2015, American Chemical Society)

conductivity tests were conducted to evaluate section 1–2, 2–3, and 3–4 in Configuration 1 and 2. In their work, the conductivity of the nanowires before being immersed into the electrolyte was tested, proving the high quality of the nanowire and metal-nanowire contacts. After immersion into the electrolyte, no conductance change was observed without potential applied, indicating the electrolyte would not affect the nanowire. What’s more, the conductivity of the current collector (Cr/Au) was much higher than the nanowire and the average conductivity of the whole system (1709 S m-1) was much higher than the electrolyte (10-1–10-3 S m-1); thus, the change of tested conductivity can reflect the degeneration of electric conductivity. Figure 4.8b is the dark-field optical microscopic image of the nanowire electrode of Configuration 1, and Fig. 4.8c shows the interconnected distorted VO5 trigonal bipyramids and VO6 octahedra in V3O8 layers. As shown in Fig. 4.8d–f, the degradation of electro-conductivity for section 1–2 was more severe than section 2–3 and 3–4. So it is clear that the shorter length from Contact 1 is, the more seriously the conductivity degrades, implying that the longer the distance, the more difficult the electron transfer is, and the shortest path is selected for Li+ inserting (Fig. 4.6d–f). Figure 4.7b shows the SEM image of electrode in Configuration 2, and Fig. 4.7c shows ions transport direction in H2V3O8 nanowire. For Configuration 2, as shown in Fig. 4.7d–i, the conductivity reduces sharply as the distance from Contact 1 is farther, which suggests that the electrochemical reaction can only occur when both ions and electrons (including Li+ and Na+) are accessible. This study greatly explored the Na+ diffusion mechanism in the nanostructure electrode.

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Fig. 4.9 Single nanowire device in the second configuration. (a) The second configuration: in which the nanowire is covered by a photoresist layer with only one end exposed to the electrolyte. (b) The SEM image of the device with the second configuration. Only one end of the nanowire is exposed to the electrolyte (inset). (c) The ion transport path model of the second configuration. The ions can only transport through the a axis of the H2V3O8 nanowire. (d-f) I-V curves of section 1–2, 2–3, and 3–4 before and after the electrochemical processes for Li ions, respectively. (g-i) I– V curves of section 1–2, 2–3, and 3–4 before and after the electrochemical processes for Li and Na ions, respectively. (Reproduced from [40]. Copyright 2015, American Chemical Society)

Sun et al. performed the in situ EIS experiment to study interface properties of Li+ during the initial charging/discharging process [41]. The set point of alternating voltage amplitude was set as 5 mV, and the range of frequency was from 100 kHz to 10 mHz in their experiment. The initial two cycles in situ EIS testing results of Li3VO4 are shown in Fig. 4.10a, b. Figure 4.10c shows the equivalent circuit fitted by Z-View software. The semicircle in the high-frequency region and the semicircle in the middle-frequency region, respectively, refer to the diffusion of Li+ in the surface coat (Rs) and charge transfer resistance (Rct) (Fig. 4.10a, b). In the low frequency region, the Warburg impedance is associated to the straight line which is related to the diffusion of Li+ in the bulk phase material. Figure 4.11a, b, respectively, shows the simulated variation tendencies of Rs and Rct. In their process, the respective element had been canceled from the equivalent circuit to fit the data if the plots didn’t have the Warburg impedance. Bonded with the voltage data of PITT, Rs kept rising from 2.00 V rapidly during the first discharging process, which was in accord with the reduction peak at 1.95 V during the first negative scan in the CV experiment. The formation of SEI results in the rising of Rs.

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Fig. 4.10 (a) In situ EIS profiles of Li3VO4 at the initial cycle, (b) at the second cycle, (c) the equivalent circuit used to fit the experimental data. (Reproduced with permission [41]. Copyright 2016, American Chemical Society)

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It is worthy to point out that this is the first time the existence of SEI is proven by the EIS experiment of Li3VO4. Due to the formation of the two new phases, the minimum Rs was obtained at about 0.60 V. Rs decreased quickly and reached a minimum at 1.30 V after charging to 0.90 V, which meant the SEI film decomposed. While Rs was still higher than the original 19 Ω, which clarified that this was not a reversible course. As for the low initial Coulombic efficiency, this was an important reason. During the discharging process, the second period’s change trend of Rs was smaller than that of the first period, which meant that there was no SEI produced in the second period. The CV curve proved the decomposition of SEI well, too. During the positive scan, the oxidation peak was around 1.12 V, which could be assigned to the decomposition of the SEI film. During the first positive scan, this peak was strong, but it was much weaker during the second process, and during the third process, it vanished. Figure 4.11b shows the variation tendency of Rct. Rct became

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much smaller during the first discharging process when it came to the voltage plateau 0.60 V, which meant Li+ inserted into Li3VO4. The maximum was 1.15 V In the first charging course, this was corresponding to the low DLi+. After 1.60 V, the Rct raised sharply because the deinsertion of Li+ ended. In the second discharging course, because of the increment of voltage plateau, Rct became much smaller after 1.35 V. In the range of the second charging course, there was not a maximum, which signified that the irreversible phase transformation did not occur at the second cycle.

4.5

Summary and Outlook

In summary, we have reviewed the recent advanced characterization techniques and their applications in understanding and developing high performance vanadiumbased nanomaterials used in Li-ion battery. Various characterization techniques based on in situ and operational design have recently been developed. It has been clearly demonstrated that these advanced characterization techniques are important for recent battery research. The new insights they provided help to mechanically understand device failures, and the real-time monitoring of electrode phase and SEI interface evolution guides future direction in optimization and innovation of highperformance electrode materials. In particular, advanced electron microscopy techniques such as TEM have proven to be an unparalleled approach that provides a unique combination of time and spatial resolution of nanostructured electrode materials. Clearly, the information and insights obtained from in situ TEM cannot be achieved by other batch-level characterizations. In the pursuit of new electrode materials, TEM provides valuable real-time information and insights into the interaction of electrode machinery.

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26. Y. Terada, K. Yasaka, F. Nishikawa, T. Konishi, M. Yoshio, I. Nakai, In situ XAFS analysis of Li(Mn, M)2O4 (M=Cr, Co, Ni) 5V cathode materials for lithium-ion secondary batteries. J. Solid State Chem. 156(2), 286–291 (2001) 27. M.H. Groothaert, J.A. van Bokhoven, A.A. Battiston, B.M. Weckhuysen, R.A. Schoonheydt, Bis(μ-oxo)dicopper in Cu-ZSM-5 and its role in the decomposition of NO: A combined in situ XAFS, UV-Vis-near-IR, and kinetic study. J. Am. Chem. Soc. 125(25), 7629–7640 (2003) 28. J. Polte, T.T. Ahner, F. Delissen, S. Sokolov, F. Emmerling, A.F. Thünemann, R. Kraehnert, Mechanism of gold nanoparticle formation in the classical citrate synthesis method derived from coupled in situ XANES and SAXS evaluation. J. Am. Chem. Soc. 132(4), 1296–1301 (2010) 29. L. Shao, S. Wang, K. Wu, M. Shui, R. Ma, D. Wang, N. Long, Y. Ren, J. Shu, Comparison of (BiO)2CO3 to CdCO3 as anode materials for lithium-ion batteries. Ceram. Int. 40(3), 4623–4630 (2014) 30. L. Chen, Y. Ruan, G. Zhang, Q. Wei, Y. Jiang, T. Xiong, P. He, W. Yang, M. Yan, Q. An, L. Mai, Ultrastable and high-performance Zn/VO2 battery based on a reversible single-phase reaction. Chem. Mater. 31(3), 699–706 (2019) 31. L. Xing, C. Zhang, M. Li, P. Hu, X. Zhang, Y. Dai, X. Pan, W. Sun, S. Li, J. Xue, Q. An, L. Mai, Revealing excess Al3+ preinsertion on altering diffusion paths of aluminum vanadate for zinc-ion batteries. Energy Stor. Mater. 52, 291–298 (2022) 32. X.H. Liu, J.Y. Huang, In situ TEM electrochemistry of anode materials in lithium ion batteries. Energy Environ. Sci. 4(10), 3844 (2011) 33. Y. Lu, X. Li, J. Liang, L. Hu, Y. Zhu, Y. Qian, A simple melting-diffusing-reacting strategy to fabricate S/NiS2-C for lithium-sulfur batteries. Nanoscale 8(40), 17616–17622 (2016) 34. G. Zhang, T. Xiong, X. Pan, Y. Zhao, M. Yan, H. Zhang, B. Wu, K. Zhao, L. Mai, Illumining phase transformation dynamics of vanadium oxide cathode by multimodal techniques under operando conditions. Nano Res. 12, 905–910 (2019) 35. S. Laruelle, P. Poizot, E. Baudrin, V. Briois, M. Touboul, J.M. Tarascon, X-ray absorption study of cobalt vanadates during cycling usable as negative electrode in lithium battery. J. Power Sources 97-98, 251–253 (2001) 36. Y. Yuan, K. Amine, J. Lu, R. Shahbazian-Yassar, Understanding materials challenges for rechargeable ion batteries with in situ transmission electron microscopy. Nat. Commun. 8(1), 15806 (2017) 37. J.Y. Huang, L. Zhong, C.M. Wang, J.P. Sullivan, W. Xu, L.Q. Zhang, S.X. Mao, N.S. Hudak, X.H. Liu, A. Subramanian, H. Fan, L. Qi, A. Kushima, J. Li, In situ observation of the electrochemical lithiation of a single SnO2 nanowire electrode. Science 330(6010), 1515–1520 (2010) 38. X. Xu, C. Niu, M. Duan, X. Wang, L. Huang, J. Wang, L. Pu, W. Ren, C. Shi, J. Meng, B. Song, L. Mai, Alkaline earth metal vanadates as sodium-ion battery anodes. Nat. Commun. 8(1), 460 (2017) 39. S. Tan, Y. Jiang, Q. Wei, Q. Huang, Y. Dai, F. Xiong, Q. Li, Q. An, X. Xu, Z. Zhu, X. Bai, L. Mai, Multidimensional synergistic nanoarchitecture exhibiting highly stable and ultrafast sodium-ion storage. Adv. Mater. 30(18), e1707122 (2018) 40. X. Xu, M. Yan, X. Tian, C. Yang, M. Shi, Q. Wei, L. Xu, L. Mai, In situ investigation of Li and Na ion transport with single nanowire electrochemical devices. Nano Lett. 15(6), 3879–3884 (2015) 41. L.L. Zhou, S.-Y. Shen, X.-X. Peng, L.N. Wu, Q. Wang, C.-H. Shen, T.-T. Tu, L. Huang, J.-T. Li, S.-G. Sun, New insights into the structure changes and interface properties of Li3VO4 anode for lithium-ion batteries during the initial cycle by in-situ techniques. ACS Appl. Mater. Interfaces 8(36), 23739–23745 (2016)

Chapter 5

Performance Optimization of Vanadium-Based Nanomaterials

5.1

Introduction

Secondary batteries, which can reversibly store electrical energy multiple times, are currently the main power supply devices used in electric vehicles and portable electronic equipment. With the continuous development of nanoscience and nanotechnology, the preparation of electrode materials in secondary batteries into various nanostructures to improve battery energy density, power density, cycle life, and safety has achieved remarkable results. Among many nanomaterials, nanowire materials have received extensive attention in the fields of energy storage and other fields, due to the advantages of high aspect ratio, large specific surface area, many active sites, and easy formation of three-dimensional networks. However, during the cycle of the battery, the continuous insertion and extraction of ions will destroy the structure of the electrode material, causing problems such as decreased conductivity and self-aggregation. Therefore, it is essential to control the performance of nanowire electrode materials to solve the problems of structural deterioration and poor stability during the cycle. This section mainly discusses the regulation and control methods of the electrical transport performance of nanowire electrode materials from four aspects: band structure regulation, surface interface regulation, ion diffusion channel regulation, and ion/electron dual continuous regulation.

5.2

Electric Transport Performance Optimization

The electric transport of nanowire materials includes electronic and ionic transport. Methods to adjust and control the electric transport performance will be discussed in following parts from band structure optimization, surface/interface optimization, ion diffusion channel optimization, and electron/ion bi-continuous optimization. © Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_5

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Band Structure Optimization

The energy band theory is often used to study the state of the electronic and its characteristics of motion in crystals (including metal, insulator, and semiconductor). It is often used to improve the electronic conductivity and enhance the electric transport performance of the nanowire electrode materials through ionic pre-intercalating, elemental doping, and extra field applying.

5.2.1.1

Ionic Pre-intercalation

Ionic pre-intercalation usually pre-intercalates ions with electrochemical transport performance, such as lithium ions, sodium ions, and calcium ions, into interlayers of layered crystals (α-MoO3, σ-MnO2, V2O5) to optimize the band structure of crystals, enhancing its electric transport performance. The Mai group reported δ-Na0.33V2O5 (NVO) ZIBs anode materials synthesized by chemically pre-intercalating sodium ions into V2O5 [1], with sodium ions existing between [V4O12]n layers (Fig. 5.1a). Electrical conductivity measurements are carried out on the single nanowires before and after pre-intercalation of sodium ions, which shows that compared to α-V2O5 nanowires, the electronic conductivity of the sodium pre-intercalated δ-Na0.33V2O5 nanowires improves from original 7.3 Sm-1 to 5.9 * 104 Sm-1 for four orders of magnitude (Fig. 5.1b), thereby optimizing greatly the electric transport performance. As cathode materials for

Fig. 5.1 (a) The crystal structure of NVO nanowire. (b) The I-V curves of NVO and V2O5. (c) Rate performance of NVO. (d) Cycling performance of NVO at 1.0 A g-1. (Reproduced from [1]. Copyright 2018, Wiley-VCH)

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Fig. 5.2 Schematic illustration of the layered vanadium oxide xerogel structure. (a) Normal vanadium oxide xerogel structure with large lattice breathing during sodiation/desodiation shows unstable structure and poor cycling stability. (b) Iron pre-intercalated vanadium oxide xerogel with inhibited lattice breathing during sodiation/desodiation displays enhanced structural stability and cycling performance. (c) Schematic formation process of the iron pre-intercalated vanadium oxide xerogel ultrathin nanobelts. (Reproduced from [2]. Copyright 2015, American Chemical Society)

ZIBs, it exhibits excellent rate performance (Fig. 5.1c) and stable cycling performance, as it achieves a capacity retention over 93% after 1000 cycles (Fig. 5.1d). The Mai group also reported a novel and facile strategy to inhibit the lattice breathing of vanadium oxide xerogel during sodiation/desodiation through iron pre-intercalation approach (Fig. 5.2) [2]. With the iron pre-intercalation, the lattice breathing along c-axis of VOx is largely reduced from 3.74 to 0.49 Å. The pre-intercalation of iron also leads to larger stabilized interlayer spacing for reversible Na+ insertion/extraction. Benefiting from the inhibited lattice breathing and stabilized interlayer spacing, the Fe-VOx exhibits enhanced cycling and rate performance.

5.2.1.2

Elemental Doping

Elemental doping usually dopes a little of other elements in the material purposively, aiming to improve the performance of some material or matter. Doping can make materials generate some particular electrical, magnetic, and optical performance, which equips them with specific values or uses. It is reported that doping can introduce impurity levels to materials lattices, effectively adjusting the electronic structure of the lattice, which increases the concentration of the charge carriers to improve its electric transport performance without damaging its structural stability [3, 4]. This principle has been applied to the multivalent ion batteries system up to now. Damboumet et al. introduced a large number of titanium vacancies through

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Fig. 5.3 The electronic structure calculations of the Ti-MnO2. EIS spectrum of (a) MnO2 NWs and (b) Ti-MnO2 NWs. Charge density distribution of (c) Ti-MnO2 and (d) MnO2. (e) Charge density differences of Ti-MnO2 and MnO2. Yellow and blue regions indicate electron accumulation and depletion zones, respectively. (f) Schematic illustration of corresponding charge transfer behavior upon charging/discharging. (Reproduced from [8]. Copyright 2019, Elsevier)

doping, achieving more Mg2+ and Al3+ storage in TiO2 [5, 6]. It is also reported that doping Ag+ in α-MnO2 can induce vacancies, further proving that doping metallic elements in the material is beneficial to introduce vacancies and adjust its electrochemical properties [7]. The Mai’s group studied the surface gradient titanium doping MnO2 nanowire (Ti-MnO2) as cathode materials for aqueous zinc-ion batteries [8]. As Fig. 5.3 shows, compared to original MnO2, Ti-MnO2 after doping presents a smaller charge transfer resistance. DFT calculations predict that the Ti substitution and its derived oxygen vacancy create a charge depletion zone and finally form a built-in electric field. The unbalanced charge and the spin distribution in the tunnel around the Ti substitution site induce an interfacial electric field and derived oxygen vacancies, which could open the [MnO6] octahedral walls, and thus promoting ion diffusion/ electron transport kinetics.

5.2.1.3

Extra Field Application

Electrostatic field is a universal force between charged particles. In energy storage devices, electrical double layer on the surface of the electrode material contributes a lot to the energy storage reaction, especially in electrochemical devices based on pseudo capacity.

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Fig. 5.4 (a) Schematic of the HER device. 0.5 M H2SO4 was used as the electrolyte. Heavily doped silicon (gray) is used as the back gate, and 300 nm SiO2 (ocher) is used as the gate dielectric. Such devices enable easy control of different back-gate voltages. (b) Statistic-based influence of the back gate on the onset overpotential and Tafel slope. The square dots are the average values, and the error bar represents standard error. (Reproduced from [9]. Copyright 2017, American Chemical Society)

If it is possible to adjust and control ions on the electrochemical interface and charge carries inside electrode materials through an extra electrostatic field applied on the interface of the electrochemical reaction, we can optimize the kinetics and thermodynamics of the electrochemical reaction. The Mai group reported a method to tune the dynamics of the adsorption process in HER by applying a back-gate voltage to the pristine VSe2 nanosheet (Fig. 5.4) [9]. The back-gate voltage induces the redistribution of the ions at the electrolyte-VSe2 nanosheet interface, which realizes the enhanced electron transport process and facilitates the rate-limiting step (discharge process) under HER conditions. A considerable low-onset overpotential of 70 mV is achieved in VSe2 nanosheets without any chemical treatment. Such unexpected improvement is attributed to the field-tuned adsorption dynamics of VSe2 nanosheet, which is demonstrated by the greatly optimized charge transfer resistance (from 1.03 to 0.15 MΩ) and time constant of the adsorption process (from 2.5 × 10-3 to 5.0 × 10-4 s).

5.2.2

Surface/Interface Optimization

The surface/interface structure has important effects on charge transfer and ion storage of vanadium-based electrode materials; therefore, vanadium-based electrode materials with different surface/interface structures often exhibit different electrochemical performance. The surface/interface structure of vanadium-based electrode materials is mainly adjusted by compounding conductive materials and structural designing.

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Fig. 5.5 Schematic illustration of the mesoporous nanowire composites with bi-continuous electron/ion transport pathways, larger electrode-electrolyte contact area, and facile strain relaxation during Li+ extraction/insertion. (Reproduced from [15]. Copyright 2014, American Chemical Society)

5.2.2.1

Compositing with Conductive Material

Compositing electrode material with conductive material can effectively overcome the problem of unstable interface structure of nanowire and buffer volume expansion to a certain extent. At the same time, the introduction of conductive material can significantly improve the electrical conductivity of the material to ensure the rapid transmission of electrons. Commonly used conductive materials include conductive carbon materials (graphene [10], carbon nanotubes [11], amorphous carbon [12]), conductive polymers (polypyrrole [13], polyaniline [14]), which are synthesized with vanadium-based electrode materials to optimize the electrochemical properties. The Mai group successfully constructed the mesoporous carbon-coated Li3V2(PO4)3 nanowire structure and studied its lithium storage performance (Fig. 5.5) [15]. Li3V2(PO4)3 nanowire structure can effectively shorten the migration path of Li+, and Li3V2(PO4)3 nanocrystals are uniformly coated in mesoporous carbon. This carbon-coated mesoporous structure can provide rapid electron transport for the electrode material, alleviate the local volume expansion caused by Li+ inserting during the cycling process, and stabilize the structure of Li3V2(PO4)3. In addition, interfacial nanocavities created by heterobonds or van der Waals interactions can provide extra ion storage sites. The Mai group constructed a VOx sub-nanometer cluster/reduced graphene oxide (rGO) cathode material composed of interfacial V-O-C bonds (Fig. 5.6) [16]. Zn2+ are proved to mainly store at the interface between VOx and rGO, which results in abnormal valence state changes compared with traditional mechanisms, and exploits the storage capacity of highly conductive rGO with non-energy storage activity. Furthermore, this interfacial storage triggers decoupled electron/Zn2+ transport, and reversible destruction/reconstruction allows the interface to store more ions than the bulk phase. The results show that the cathode material has obtained high specific capacity (443 mAh g-1 at 100 mA g-1, which exceeds the theoretical capacity of various surface components). Such interface led storage is a new way to build high energy and high power density devices.

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Electric Transport Performance Optimization

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Fig. 5.6 (a) Schematic diagram of the VOx-G heterostructure. (b) Schematic of the interface pseudocapacitance. (Reproduced from [16]. Copyright 2021, Wiley-VCH)

5.2.2.2

Design of Nanostructure

Three-dimensional, porous, or ultra-thin electrode materials with high specific surface area can provide channels for electron transport and ion migration and promote the infiltration of electrolyte, so as to improve the electrochemical performance of vanadium-based electrode materials. The Mai’s group designed a threedimensional H2V3O8 hydrogel structure with ultrathin nanoribbons and self-curling nanoriols by a universal liquid phase stripping method [17]. The hydrogel structure formed by cross-linking ultrathin nanoribbons and self-crimping nanoriols shortens the ion diffusion distance greatly. The porous hydrogel structure increases the contact area between electrode and electrolyte, and provides an effective interconnecting channel, which is conducive to ion diffusion. In addition, ultrathin pre-lithiated V6O13 nanosheet cathodes [18] and three-dimensional porous V2O5 hierarchical microplates [19] are also constructed. In particular, poor solid-solid interfacial contact and interfacial chemical reactions around the cathode in solid-state lithium batteries (SSLBs) hinder their practical

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Fig. 5.7 (a) The distribution and function of the PEO gradient interfaces in the gradient NW cathode, PEO gradient interfaces only occupy a small area, most of PEO is evenly distributed. (b) Internal electron/ion transport characteristics of the gradient NW cathode. (c) Cycling performance of gradient NW cathode at 100 mA g-1 based on SSE at room temperature. (Reproduced from [20]. Copyright 2021, Wiley-VCH)

application. To address this problem, the Mai group designed a gradient H2V3O8 nanowire (NW) cathode [20]. In this unique gradient anode, two polyvinyl oxide (PEO) gradient distribution interfaces are constructed as interfacial buffer layers between the two surfaces (Fig. 5.7). One side of the surface of the more ionic conductive polymer provides smooth contact with SSE, while the other side of the surface of the more electron conductive H2V3O8 NWs/rGO provides fast electron transport as a collector fluid. In addition, the internal NW cathode material is uniformly coated with rGO and PEO-based solid polymer electrolyte (SPE). This strong bonding enables point-to-point contact between the positive electrode and the SSE to become large-area contact, thus providing a continuous channel for fast electron/ion transport and improving structural stability [21–27].

5.2.3

Ion Diffusion Channel Optimization

Ion diffusion coefficient is one of the key parameters reflecting the ion transport characteristics of materials. The Ceder group systematically studied the migration characteristics of lithium ions in layered transition metal oxides by means of DFT theoretical calculation, and they concluded that the spacing between lithium layers and electrostatic repulsion are two decisive influencing factors [28]. Because the layer spacing is also affected by the electrostatic interaction force, the two influencing factors are also coupled to a certain extent. According to Coulomb’s law, the electrostatic interaction force satisfies the following relationship:

5.2

Electric Transport Performance Optimization

u=

k e q1 q2 r2

113

ð5:1Þ

In the formula, ke represents coulomb coefficient (ke ≈ 9.0 × 109 N m2/C2); q stands for charge; r is the distance between the two charges. In the pre-intercalated material, due to the expansion of layer spacing along the c-axis, the electrostatic force u decreases with the increase of distance r, thus reducing the migration resistance of current-carrying ions. Adjusting the diffusion channel of electrochemical transport ions can optimize the ion diffusion coefficient of materials and improve the electrical transport properties of materials. This section will outline how to regulate ion diffusion channels through metal ion pre embedding, inorganic/nonmetallic ion pre-intercalated, and organic molecule pre-intercalated, so as to enhance the ion transport capacity of nanowire electrode materials [2, 29–38].

5.2.3.1

Metal Ion Pre-intercalation

The Mai group studied the typical nanowire cathode material A-M-O (A = Li, Na, K, Rb; M = V, Mo, Co, Mn, Fe) with various alkali metal ions pre-intercalated and quantitatively described the process of alkali metal ions pre-intercalated to increase ion diffusion channels to improve the electrochemical properties of the materials [31]. The results show that vanadium oxide nanowire electrode material with the alkali metal ions pre-intercalated shows considerable rate performance; Fig. 5.8a shows that the pre-intercalation of large alkali metal ions can increase the diffusion channel of lithium ions in lithium-ion batteries. Combined with experimental characterization (rotational electron diffraction, X-ray diffraction) and DFT theoretical calculation, the size of the ion diffusion channel after pre-intercalation of a variety of different ions was calculated. In addition, the following conclusions were drawn: Potassium ion pre-intercalated vanadium oxide has the largest interlayer distance, and its multiplier performance is the best in the cathode material of lithium-ion battery, that is, increasing the ion diffusion channel is helpful to improve the diffusion performance of lithium ion (Fig. 5.8b). Pomerantseva et al. [37] from Drexel University constructed a pre-intercalated vanadium oxide (δ-KxV2O5nH2O, the pre-intercalated layer spacing is 9.56 Å), and the specific capacity of δ-KxV2O5nH2O is as high as 268 mAh g-1 when δ-KxV2O5nH2O is used as the cathode material of potassium-ion battery due to the large layer spacing, which is favorable for potassium ion transport (Fig. 5.8c). In addition to the pre-intercalation of alkali metal ions, the research of alkalineearth metal ion pre-intercalation has also received a lot of attention. As shown in Fig. 5.8d, Pomerantseva et al. compared vanadium oxides (δ-MxV2O5, M = Li, Na, K, Ma, Ca) pre-intercalated with different alkali metal ions and alkaline-earth metal ions, and they obtained the law that the layer spacing is proportional to the hydration radius of pre-intercalated ions [33]. The Mai group prepared the doublelayer structure Mg0.3V2O51.1H2O with the synergistic action of Mg2+ and lattice

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Fig. 5.8 (a) Schematic representation of large alkali metal ion intercalation. (b) Rate performance of A-V-O nanowires. A-V-O nanowires are cycled at various rates from 0.05 to 4.0 A g-1. Here nC denotes the rate at which a full charge or discharge takes 1/n hours. (c) First charge cycle and second discharge/charge cycle curves at a current rate of C/50 of δ-K0.42V2O5nH2O. (d) X-ray diffraction results of different ions pre-intercalated in vanadium oxide and the relationship between layer spacing and hydration ion radius. (Reproduced from [31, 33, 37]. Copyright 2015 and 2018, Elsevier and American Chemical Society)

water, in which the pre-intercalation of Mg2+ can give the electrode material high conductivity and excellent cycle stability [35]. Because of the charge shielding effect, the lattice water can accelerate the migration rate of Mg2+. When used as the cathode material of magnesium-ion battery, it shows excellent rate performance and outstanding cycle stability (the specific capacity retention rate is 80.0% after 10,000 cycles). In addition, other metal ions (Mn2+, Fe3+, Al3+) can also be used for pre-intercalation materials. The Mai group regulated the ion channel of vanadium oxide by pre-intercalation iron ions [2] and synthesized vanadium oxide nanoribbons

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Fig. 5.9 (a) Crystal structure diagram of typical vanadium oxide gel. (b) Schematic diagram of crystal structure of vanadium oxide after pre-intercalated iron ion. (c) Electrochemical impedance spectroscopy of vanadium oxide before and after iron ion pre-intercalation. (d) Cyclic properties of vanadium oxides before and after iron ion pre-intercalation. (Reproduced from [2]. Copyright 2015, American Chemical Society)

pre-intercalated iron ions (Fe-VOx) by simple hydrothermal method. As shown in Fig. 5.9a, b, the pre-intercalation of iron ions can inhibit the “lattice breathing” generated in the process of sodium ion de-embedding, resulting in the instability of crystal structure. This strategy can alleviate the collapse of the diffusion channel of sodium ions, thus improving the cyclic stability of sodium ion storage. Nanocharged transport performance of vanadium oxide pre-intercalated with sodium ions is improved as shown in Fig. 5.9c, charge transfer impedance is reduced, that is, the diameter of half circle is reduced. When used as anode material of sodium-ion batteries, specific capacity and cycle stability performance are improved (Fig. 5.9d).

5.2.3.2

Inorganic Molecule/Nonmetal Ion Pre-intercalation

In addition to metal ions, inorganic molecules/nonmetal ions (H2O, NH4+, etc.) can also be pre-intercalated to increase and optimize the ion diffusion channels of electrode materials. Yang et al. and the Mai group designed a vanadium pentoxide gel (V2O5 nH2O) pre-intercalated with water molecules [34], as shown in Fig. 5.10a, which was used as a cathode material of zinc-ion battery. During the discharge process, zinc ions were embedded between V2O5nH2O layers. The solvation of the pre-intercalated water molecules with zinc ions increases the layer spacing, reduces the effective charge of zinc ions, and reduces the electrostatic interaction between zinc ions and V2O5. This process can be understood as the pre-intercalated water molecules play a role of “lubricant” in the diffusion process of zinc ions, effectively

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Fig. 5.10 (a) The proposed crystal structures of pristine VOG, VOG after charging to 1.3 V, and discharging to 0.2 V. (b) Comparison of the Ragone plot between the V2O5H2O cell and two other reported materials for aqueous Zn batteries. (Reproduced from [34]. Copyright 2017, Wiley-VCH)

realizing the efficient diffusion of zinc ions, and finally presenting excellent electrochemical zinc storage performance (Fig. 5.10b). Kanatzidis et al. [29] reveals gel structure-single photograph of V2O5 by water molecules through atomic pair distribution function (APDF), which are separated and infinite stack in the Z axis direction to enlarge layer spacing. Therefore, this V2O5 gel achieved better performance compared with the original photograph vanadium pentoxide.

5.2.3.3

Organic Molecular Pre-intercalation

In addition to metal ions, inorganic molecules/nonmetal ions and organic molecules (such as triglycoll, aniline, and pyridine) can also be pre-intercalated into the crystal of nanowire electrode material to optimize the ion diffusion channel. Yu et al. pre-intercalated the organic molecule triethylene glycol into the layered polyanionic VOPO4 electrode material, effectively increasing the layer spacing of VOPO4, and used the VOPO4 pre-intercalated with triethylene glycol as the electrode material of sodium-ion battery. Compared with the VOPO4 without pre-intercalated triethylene glycol, the diffusion energy barrier of sodium ion decreased significantly, effectively improving the diffusion kinetics of sodium ion [32]. The Mai group pre-intercalated aniline molecules into the crystal of VOPO4 to further expand the layer spacing of VOPO4, which is conducive to the transport of polyvalent metal ions with large radius and has a great advantage in polyvalent metal-ion batteries [38]. The Mai group also studied the sodium storage performance of V2O5nH2O nanowires pre-intercalated with pyridine molecules [36]. Due to the large interlayer spacing and the free movement of protons between layers of bilayer V2O5nH2O, pyridine molecules can easily interlayer, combine with interlayer protons, and form hydrogen bonds between layers. This strategy can effectively stabilize the interlayer structure and avoid the collapse of the interlayer structure, thus effectively protecting the interlayer transport channel of sodium ions (Fig. 5.11a). Based on the above

5.2 Electric Transport Performance Optimization

117

a Pyridine

Na+

(proton acceptor)

d=13.36 Å

V2O5.nH2O (proton donor) 4

3

2

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30

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c

Pyridine-V2O5.nH2O V2O5.nH2O

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180

Specific capacity (mAh g-1)

Potential (V vs. Na+/Na)

b

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Discharge d=11.36 Å Na+

d=11.36 Å

Sodiated Pyridine-V2O5.nH2O Pyridine-V2O5.nH2O V2O5.nH2O

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Fig. 5.11 (a) Functions of pyridine intercalated into V2O5nH2O, stabilizing the interlayer structure during rapid Na+ (de)intercalation. (b) Charge-discharge curves of pyridine-V2O5nH2O and V2O5nH2O at 0.5 A g-1 in the potential range of 1–4 V (vs Na+/Na). (c) Long-term cycling performance at 1 A g-1. (Reproduced from [36]. Copyright 2019, Wiley-VCH)

structures, V2O5nH2O pre-intercalated with pyridine exhibits high specific capacity and excellent stability in sodium-ion batteries (Fig. 5.11b, c).

5.2.4

Electron/Ion Bi-continuous Optimization

The electrical transport of materials is limited by the slower transport of ions and electrons, while the optimization of electrical transport of electrode materials usually only starts from ions or electrons unilaterally, even the unoptimized party has a negative impact. In order to realize the electron/ion bi-continuous transmission of the nanowire electrode material, the electrical transport performance can be controlled by designing the coaxial semi-hollow structure and the nanowire array structure.

5.2.4.1

Coaxial Semi-Hollow Structures

The active material can be tightly coated by conductive carbon, graphene, conductive polymer, and other conductive materials to obtain good electronic conduction. However, the tight coating will hinder the diffusion of ions in the electrolyte in the electrode material, which makes it difficult for the electrode material to take into account the efficient transmission of electrons and ions. Compared to the tightly coated nanowires, the coaxial semi-hollow structure can not only provide a fast electron conduction path through the outer conductive material, but also make the electrolyte fully infiltrated through the semi-hollow structure.

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Fig. 5.12 Synthetic schematic diagram and SEM image of coil coaxial semi-hollow V3O7/ graphene nanowires. (Reproduced from [21]. Copyright 2013, American Chemical Society)

The charge transfer characteristic formulas of tightly coated nanowires and coaxial semi-hollow nanowires have been introduced. Based on the formula description, the Mai group designed a V3O7/graphene material with the coaxial semi-hollow nanowire structure, benefiting for the electron/ion bi-continuous transport [21]. They proposed a growth mechanism of “oriented assembly” and “self-curling” by combining experiments and molecular dynamics simulations: providing a certain amount of energy in the synthesis environment to allow graphene self-curl on the surface of V3O7 nanowires. This process makes a certain interspace between graphene and V3O7 nanowires to construct coaxial semi-hollow nanowires for the bi-continuous ions/electrons transport process (Fig. 5.12). This structure not only enhances electronic conduction, but also creates a cavity between the nanowires and the graphene rolls. The electrolyte can fully infiltrate the interior of the semi-hollow structure, which effectively improves the interface contact between the electrode material and the electrolyte, benefiting for the uniform distribution of charge and reducing the polarization in the energy storage process. The charge transport performance test shows that, compared with the conventional V3O7 nanowires, the interface charge transfer resistance of the coil coaxial semi-hollow V3O7/graphene nanowires is reduced from 81 Ω to 32 Ω, while the conductivity is increased by 27 times. The energy density of constructed energy storage device not only maintains high energy density, and but also the power density has been enhanced by 6 times. Moreover, the cavity can accommodate the volume expansion of the material, and the electron/ion bi-continuous transport is maintained during the cycle. In addition, in order to verify the above ideas, The Mai group also constructed MnO2 nanorods with a coaxial semi-hollow structure [22]. The study found that the electron/ion bi-continuous effect is also suitable for controlling the conversion reaction kinetics of the negative electrode material system. The electron/ion bi-continuous transport is realized under the condition of large expansion/shrinkage in volume during the electrochemical process. They created a cavity between the metal oxide nanowire and the carbon material, constructing the MnO2/porous carbon material with a coaxial semi-hollow structure (one-dimensional yolk-shell structure). The carbon can provide a fast electron conduction path, and the hollow structure helps the electrolyte to infiltrate the active material and promote ion transport. In this

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119

case, the interface charge transfer resistance is reduced from 263 Ω to 89 Ω. At the same time, the electrode rate performance is significantly improved (179 mAh g-1 at a high current density of 1 A g-1), while the pure MnO2 has almost no capacity. This structure also enhances the capacity to accommodate volume expansion, effectively hindering SEI. The electrode cycle stability will be improved during cycling (especially under high-rate charging and discharging) to avoid the rupture of membrane. There is a significant improvement at a current density of 200 mA g-1, while the specific capacity of the synthesized sample after 200 cycles is 509 mAh g-1, higher than that of pure MnO2 attenuating to 61 mAh g-1.

5.2.4.2

Nanostructure Array

The formation of nanowire arrays is generally controlled by the conditions of the growth process, in which the nanowire array material with uniform orientation is grown on the surface of the substrate (metal-based or carbon-based material). The prepared material does not need to add a binder, and it also overcomes the shortcomings of the electrochemical inertness of the binder. Compared with the way the active material adheres to the current collector, the nanowire array has the following advantages: 1. The higher specific surface area of the nanowire array material will avoid the agglomeration of the active material. 2. The uniform electron conduction between the nanowire array and the substrate greatly improves the overall conductivity of the material. 3. The gap between the nanowires makes the active material fully contact the electrolyte, which is beneficial to the transport of ions during the reaction (Fig. 5.13) [23, 24]. Based on the advantages of nanowire arrays, The Mai group studied the structure of SnO2 polyaniline nanowire arrays with high-efficiency electron transport characteristics [24]. By a combination of hydrothermal treatment and electrodeposition, a

Fig. 5.13 Schematic diagram of the electron/ion bi-continuous transmission of nanowire array. (Reproduced from [23]. Copyright 2013, American Chemical Society)

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Fig. 5.14 The structure and performance changes of nanowire arrays when used as electrode materials for lithium-ion batteries. (a) Schematic diagram of structure destruction of SnO2 nanowire arrays after cycling and structure retention of SnO2 PANI arrays after cycling. (b) Impedance comparison of nanowire arrays. (c) Cyclic performance comparison of nanowire arrays. (Reproduced from [24]. Copyright 2014, Elsevier)

simple and controllable method for synthesizing heterogeneous multi-branched core-shell SnO2 PANI nanowire arrays was proposed. Compared with the structural damage of the initial SnO2 arrays and SnO2 arrays after cycling, the multi-branched core-shell structure SnO2 PANI array can still maintain a complete mechanical structure after cycling (Fig. 5.14a). Due to its unique composition and structure, the prepared heterogeneous SnO2 PANI array with a multi-branched core-shell structure has strong adhesion between the SnO2 core and the conductive polymer shell, exhibiting excellent structural stability and mechanical integrity. At the same time, it has the advantages of the nanowire array, promoting the dynamic process in the three-dimensional electron transport and ion transport. Therefore, the cycle performance and the rate performance of the SnO2 PANI array are significantly improved as a lithium-ion battery electrode material (Fig. 5.14b, c). The construction

5.3

Regulation and Control of Structural Stability

121

of this multi-branched core-shell nanowire array and its effective strategy in material synthesis design will likely be further applied to other high-performance energy storage devices.

5.3

Regulation and Control of Structural Stability

Nanowire electrode materials will expand or contract to varying degrees during the electrochemical reaction, resulting in structural collapse and performance degradation. Therefore, improving the structural stability of the nanowire electrode materials during the electrochemical reaction is very important for improving the performance of the material. This section will mainly discuss the control methods of structural stability of the nanowire electrode materials from two aspects: internal stress buffering and volume expansion suppression.

5.3.1

Internal Stress Buffering

The internal stress generated by the volume change of the nanowire electrode material during the electrochemical reaction is an important factor leading to material deterioration. By designing hybrid nanowires and hierarchical nanowires, the internal stress generated by the reaction process can be effectively buffered.

5.3.1.1

Hybrid Nanostructures

The Mai group studied the nanowire with a hybrid structure of VO2, as shown in Fig. 5.15a, the hybrid structure is composed of nanowires, nanobelts, and nanovolumes. This hybrid structure not only has the better electrical transport performance (Fig. 5.15b), but also effectively buffer the structural damage caused by stress change during the cycling process. And the cycle stability of energy storage devices is also improved accordingly [39, 40].

Fig. 5.15 (a) SEM image of VO2 hybrid structure nanowire. (b) Impedance spectrum of VO2 hybrid structure nanowire. (c) Cycle performance of VO2 hybrid structure nanowire. (Reproduced from [39]. Copyright 2013, Wiley-VCH)

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Schmidt of Chemnitz University of Technology in Germany also used the method of buffering internal stress to construct a new hybrid structure composed of multilayer Ge and Ti nanofilms [41]. A “high-speed channel” will be formed for charge storage and transfer due to the Ti film in the middle layer, and the self-curling process of the Ge/Ti nanomembrance also can minimize the overall system energy, and effectively release the inner space of the nanomembrance. Therefore, it can play a role in stabilizing the structure while ensuring rapid electron transport, which in turn leads to its high reversible specific capacity and excellent cycle performance. In addition, Schmidt et al. also applied this stress-buffering method to the construction of SiOx/SiOy double-layer nanomembrane materials [42]. Due to this double-layer nanomembrane structure with stress-buffering effect, when it is used as a lithium-ion battery negative electrode material, it can effectively solve the problem of material powdering, and thus exhibit excellent cycle stability.

5.3.1.2

Hierarchical Nanostructures

Due to the interface slip effect of hierarchical nanowires, the stress during the electrochemical reaction of materials can also be effectively released. Based on the above ideas, the Mai group and others constructed MnMoO4/CoMoO4 hierarchical nanowires, as shown in Fig. 5.16a. The material is used as the electrode material of supercapacitors, showing excellent cycle stability and high specific capacity [43]. Hongjin Fan of Nanyang Technological University in Singapore and others have studied a novel dendritic hierarchical nanomaterial with SnO2 nanowires as the backbone and Fe2O3 nanorods as branches. Based on the stress buffer effect of the abovementioned hierarchical structure, this structure can also effectively alleviate the internal stress changes generated during the lithium-ion insertion/extraction process and improve the cycle stability of electrode materials.

Fig. 5.16 (a) The construction of hierarchical MnMoO4/CoMoO4 nanowires. The green rod represents the backbone MnMoO4 nanowire, and the orange rods the CoMoO4 nanorods. Red and blue balls are different ions dispersed in the aqueous solution. (b, c) SEM image of the backbone MnMoO4 and hierarchical MnMoO4/CoMoO4 heterostructurednanowires. (d, e) TEM and high-resolution TEM images at the heterojunction of hierarchical MnMoO4/CoMoO4 heterostructured nanowires. (Reproduced from [43]. Copyright 2011, Nature Publishing Group)

5.3

Regulation and Control of Structural Stability

5.3.2

123

Volume Expansion Suppression

Volume expansion refers to the deformation phenomenon of electrode material structure caused by participating in electrochemical reactions in the charging and discharging processes. Searching for suitable electrode materials to satisfy the longterm stability requirement is an important step to realize the large-scale energy storage [44]. As a one-dimensional material with relatively stable structure, nanowires can inhibit the volume expansion effect by constructing dense coating structure and cavity structure, so as to effectively alleviate the electrode degradation caused by stress and improve the electrode cycle stability. Other materials with excellent mechanical properties (carbon, metal, oxide, etc.) are tightly coated on the nanowires, which can buffer the volume expansion and protect the nanowire electrode materials. The purpose of buffer is to maintain a stable morphology and structure, while the purpose of protection is to block the excessive contact between active substances and electrolyte, so as to prevent the continuous exposure of active substances and the continuous formation of SEI film during volume expansion. The coating layer is divided into homogeneous coating layer and heterogeneous coating layer. The coating layer with the same composition but different crystal phase with nanowires is called homogeneous coating layer. Compared with homogeneous coating, heterogeneous coating (i.e., nanowire materials composed of different substances) has been studied extensively. The Mai group synthesized graphene-modified vanadium oxide nanowire composites, and the structure is shown in Fig. 5.17 [45]. The graphene layer has a certain buffer effect on the volume expansion of V2O5 nanowires during the Mg2+ ion insertion/extraction process. It has excellent cycle performance and magnification performance, which proves the excellent structural stability of the material. Close coating is an effective means to inhibit volume expansion and stabilize the structure. In contrast, giving materials a certain expansion freedom can also effectively stabilize the structure. The structure of the material can be effectively stabilized by designing the cavity structures such as nanotube with coating on the surface and semi-hollow “pipe centerline.” Cavity structure nanowires include tubular nanowires (“nanotube” structure) and semi-hollow nanowires (“tube centerline”

Fig. 5.17 Schematic illustration of the fabrication process and proposed formation mechanism of the V2O51.42H2O@rGO aerogel. (Reproduced from [45]. Copyright 2015, Elsevier)

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structure). The former has an inward expansion space along the radial direction, and the latter has an outward expansion space along the radial direction. Both nanowires can reserve space for the expansion of materials. At the same time, both have shell protection, so they can stabilize the structure. The Mai group synthesized semihollow bi-continuous structure as cathode material for LIBs by nanowire (V3O7) template method [21]. The nanowire templated graphene scroll nanostructure provides internal void space for swelling during lithiation and an effective internal electrolyte channel to enhance the ion diffusion coefficient. In a pure nanowire structure, the strain could not be completely and promptly released.

5.4

Summary and Future Directions

Nanowire electrode materials are prone to problems such as conductivity drop and structural deterioration in actual electrochemical energy storage applications. In order to improve the performance of nanowire electrode materials, this chapter mainly focuses on the regulation of electrical transport performance and structural stability. The optimization of nanowire electrode materials is summarized as follows: (1) The regulation of electrical transport performance mainly includes band structure regulation, surface interface regulation, ion diffusion channel regulation, and ion/electron dual continuous regulation. More carriers will be provided and induceelectric field to drive charge transport to achieve energy band structure regulation via the ion pre-intercalation strategy. And the surface and interface was regulated through composite conductive materials for increasing active sites. And, the regulation of ion diffusion channels can be achieved during the pre-intercalation process of metal ions, inorganic molecules/nonmetal ions and organic molecules. Additionally, the dual continuous regulation of ions/electrons can be achieved through the design of specific struture of coaxial semi-hollow and nanowire array structures. (2) Structural stability control mainly includes internal stress buffering and volume expansion suppression. The internal stress is buffered by designing hybrid nanowires and hierarchical nanowires; the volume expansion can be suppressed by designing tightly wrapped structures and cavity structures. In addition to optimizing electrical transport and stabilizing the structure, the following two points need to be considered in the future: (1) The performance control of nanowire electrode materials still needs to be combined with in situ characterization methods, and further targeted research on the performance degradation mechanism of specific materials is required to achieve the performance of “targeted” control, so as to obtain high-performance nanowire electrode materials. (2) It is still necessary to explore new methods and new ideas for synergistically improving the electrical transport performance and structural stability of nanowire electrode materials to provide support for the development of secondary batteries, supercapacitors, and micro-nano energy storage devices.

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Chapter 6

Vanadium Oxide Nanomaterials for Electrochemical Energy Storage

6.1

Introduction

In 1990, Sony introduced LIB with LiCoO2 as a positive electrode, which promotes the development of portable chemical power supply and brings an innovation to the battery industry. Since then, LiNiO2, LiMn2O4, LiNi1/3Co1/3Mn1/3O2, and LiFePO4 have emerged as positive electrode for LIB. However, these cathode materials have some shortcomings, respectively. For example, the commercialized LiCoO2 is expensive and toxic [1]. The synthesis of LiNiO2 is difficult for mass production, and its thermal stability is poor [2]. LiMn2O4 suffers from the John-Taller effect, which results in rapid capacity fade in the process of deep charge/discharge cycling [3, 4]. Although taking some advantages of the above three materials, the structure and properties of LiNi1/3Co1/3Mn1/3O2 extensively depend on the synthesis method, and the cation mixing will destroy the performance of this ternary system electrode. LiFePO4 has its disadvantages such as low conductivity, low Li ion diffusion coefficient, and poor high/low temperature performance. Therefore, researchers are looking for a compound which is suitable for lithium intercalation to replace the commercial LiCoO2 as a cathode material for LIB [5–9]. Recently, vanadium oxides (VOs) have widely attracted attention from researchers in energy storage field. Vanadium has various oxidation valence states (V5+, V4 +, V3 +) and crystal structures including VO2, V2O5, and V6O13. These compounds have an open layered structure leading a strong covalent bond in layer as well as a weak van der Waals bond or hydrogen bond between the layers which can be embedded in an atom or a molecule. VOs exhibit outstanding interaction with ions or molecules [10], strong electron correlation and excellent catalytic activity, and abundant resources of vanadium in nature, which makes vanadium oxides a widespread application in the field of energy conversion. However, the limited cycling stability, low conductivity, and rapid fading capacity of VOs greatly hinder the practical applications in cathode materials. In the past few decades, a large number of researchers have worked to solve the intrinsic defects caused by VOs © Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_6

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via nanostructured materials strategy. VOs are prepared into a series of nanostructured composite materials, such as 1D nanorods [11–13], nanowires [14–16], nanotubes [17], and nanobelts [18–20]. 2D nanosheets [21, 22] and 3D or hierarchical nano-hollow spheres and nanoflowers, which can enlarge the storage performance of Li ions and decrease the diffusion pathway of ions, reducing the mechanical stress associated with lithium ion deintercalation, thereby improve the rate performance and specific capacity of the LIB [23–26]. At the same time, combining with carbon black, graphene oxide, carbon nanotubes [15, 27, 28], and metal ions doping [29, 30], VOs can enhance the conductivity, which shows higher specific capacity and good cycling stability. Therefore, VOs are expected to be the promising candidate of cathode materials for the next-generation LIB with high performance.

6.2 6.2.1

Orthorhombic V2O5 V2O5 for LIB

In the past decades, VOs have been widely explored as negative electrode, since it was firstly reported by Whittingham in 1976 [31]. Among them, most works are focused on the Li-ion intercalations of orthorhombic V2O5 [32], which belongs to the space group Pmmn (a = 1.1540 nm, b = 0.3571 nm, c = 0.4383 nm). A typical layered structure of orthorhombic V2O5 is shown in Fig. 6.1 [33]. The single-layer V2O5 is shaped by two units of “VO6 octahedra” linked to one another [33]. Among them, one of the terminal V-O bonds is shorter than the others, which makes the V coordination polyhedron form a square pyramid [VO5] (Fig. 6.1a). On the other hand, the interlayer spacing can be readjusted by molecules or ions because each single-layer V2O5 is connected via the weak V-O bond [34– 36]. Theoretically speaking, V2O5 can release the specific capacity of 440 mAh g-1 with 3 lithium ions intercalated per V2O5 which outclasses those commercialized materials (e.g., LiCoO2, LiMn2O4, and LiFePO4). The lithium intercalation mechanism is as follows [38] (Fig. 6.2):

a

b

b

c a

a

Fig. 6.1 Crystal structures of orthorhombic V2O5: (a) single layer, (b) layered structure. (Reproduced from Chernova et al. [33]. Copyright 2009, Royal Society of Chemistry)

6.2

Orthorhombic V2O5

131

Fig. 6.2 Evolution of phases with the degree of lithium intercalation into V2O5 and the cycling of ω-phase: collected at 180 μA cm-2. (Reproduced from Delmas et al. [38]. Copyright 1994, Elsevier B.V.)

Fig. 6.3 Comparison of the arrangement of VOn polyhedral in (a) V2O5, (b) α-LixV2O5, (c) ε-LixV2O5, (d) δ-LixV2O5, and e γ-LixV2O5. (Reproduced from Cocciantelli et al. [39]. Copyright 1991, Elsevier B.V.)

xLi þ V2 O5 = Lix V2 O5

ð6:1Þ

The crystalline V2O5 exhibits multiple platforms in the discharge process [38] and is shown in Fig. 6.2. After several deep charge/discharge cycling, the crystal structure changes and the lengths of apical V-O bond are outstretched as well as the corner-shared oxygen atoms are pushed mildly into the interlayer space (Fig. 6.3), while the specific energy and charge capacity faded. At the same time, there are other problems in the crystalline V2O5 cathode such as low electronic conductivity, irreversible lithium insertion under medium pressure, and the electrolyte oxidation during battery charging, which limits the excellent electrochemical performance of V2O5 in LIB.

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Over the past decades, a number of researchers have found that electrochemical properties (e.g., specific capacity, cycle stability, and rate capacity) extremely rely on nanostructures of cathode materials. Among many nanomaterials based on V2O5, 1D V2O5 is widely studied due to its relatively facile preparation method, larger active specific surface area than that of the traditional cathode material which exhibits high specific capacity and good cycling performance. 1D nanostructures of V2O5 include nanofibers, nanorods, nanobelts, and nanotubes. Via electrospinning combined and high temperature annealing, the Yu group [40] manufactured mesoporous V2O5 nanofibers. The specific surface area is up to 97 m2 g-1 and the porosity is about 67%. The excellent nano-mesoporous structure provides a lot of space for Li-ion storage, which results in superior cycling stability and reversibility, and releases an initial discharge capacity of 370 mAh g-1. The Mai group [14] also synthesized ultra-long nanowires with a hierarchical structure via electrospinning method. Since the nanorods are adhered to the surface of the nanowire from the transmission electron microscopic image (Fig. 6.4), it prevents the agglomeration of nanorods and increases the contact area of the active material for improving the capacity of the nanowires as the negative electrode material of LIB. The initial specific capacity is 390 mAh g-1 (1.75–4 V) and remains 201 mAh g-1 after 50 cycles, showing a good capacity performance. Compared with nanowires, nanotubes have a multi-walled tubular structure that provides a larger active area and more space for the rapid transportation of Li ions. As V2O5 nanotubes, hydrothermal method is the most usually used way of production, which is a very complicated process under hydrothermal synthesis conditions.

Fig. 6.4 (a) FESEM images of electrospun composite nanowires before annealing. (b) FESEM images of the ultralong hierarchical V2O5 nanowires after annealing. (c) TEM image of the ultralong hierarchical V2O5 nanowires after annealing. (d) HRTEM image. (e) Fast Fourier transformation pattern of a single nanorod on the hierarchical V2O5 nanowires. (f) Charge/discharge curves of hierarchical V2O5 nanowires at voltages of 2–4 and 1.75–4 V, respectively. (g) Capacity versus cycle number and Coulombic efficiency versus cycle number of the ultralong hierarchical V2O5 nanowires. (Reproduced from Mai et al. [14]. Copyright 2010, American Chemical Society)

6.2

Orthorhombic V2O5

133

Cui et al. [17] prepared V2O5 multi-wall nanotubes by a sol-gel approach combined with hydrothermal method. The introduction of LiOH can increase the specific capacity of V2O5 and promote the deintercalation of lithium ions. Due to the introduction of LiOH, the initial specific capacity is as high as 457 mAh g-1 at 30 mA g-1 in a voltage range from 1 to 4 V, which is superior than the theoretical specific capacity of V2O5 (442 mAh g-1). However, the capacity retention of V2O5 nanotubes is poor, the specific capacity decreased to 270 mAh g-1 after 10 cycles. In the meantime, they proposed the rolling mechanism of V2O5 nanotubes that in the hydrothermal synthesis process, the decomposition of some organic amine molecules will lead to the conversion of V5+ to V4+, forming a layered V6O17 compound. Because the ionic radius of V4+ is bigger than V5+, a layer of [VO5] tetragonal pyramid rich in V4+ will increase the size of one side of the V6O17 layer, which directly causes the V6O17 layer to curl to reduce the system energy to form a nanotube structure. Qin et al. [18] prepared a porous V2O5 nanobelt with a length of 2 μm and a thickness of 20 nm by hydrothermal method. At 50 mA g-1 in 2.5 to 4 V, the initial specific capacity is 142 mAh g-1, and retains 141 mAh g-1 after 100 cycles, which is very close to the theoretical capacity of V2O5 microcrystalline electrode (147 mAh g-1). Even at 1000 mA g-1, the initial specific capacity is as high as 130 mAh g-1, the retention capacity reaches 128 mAh g-1 after 100 cycles, the electrode material shows superior cycle stability and specific capacity due to the porous structure of V2O5 nanobelt, which shortens the diffusion pathway of Li ions and the electron transport distance, promotes the penetration of the electrolyte, alleviates the volume change of V2O5 during cycling. To some extent, common nanofibers, nanowires, nanotubes, etc., are used as a negative electrode material to improve the performance of V2O5 due to reducing the diffusion distance of electrons and Li ions as well as increasing effective contact area and the storage of Li ions. Although the abovementioned nanomaterials have a high active specific surface area when just formed, the 1D nanostructures are prone to agglomeration due to high surface energy and low mechanical stability, which decrease the active specific surface area. When the charge and discharge tests are performed, their initial specific capacity will be higher; however, there is a significant capacity fading and no further improvement in long-term cycle performance. There are few reports on 2D V2O5 nanosheets. 2D nanostructure (such as nanoplates and nanosheets) can further enhance the electrochemical kinetics and structural stability [22, 41, 42]. Rui et al. [43] obtained an ultrathin V2O5 nanosheets by liquid phase stripping method, and the thickness of nanosheets is 2.1–3.8 nm. When using the ultrathin nanosheets as a cathode material, Fig. 6.5 shows that the initial discharge specific capacity is 290 mAh g-1 at 59 mA g-1 from 2.05 to 4 V, and the specific capacity of V2O5 nanosheet negative electrodes is 40% higher than bulk V2O5 samples, which is largely owing to the much longer diffusion distance for electron and Li-ion transportation. Moreover, nanosheet negative electrodes retain 93.8% capacity retention after 50 cycles in a voltage window between 1.5 and 4 V (vs. Li/Li+), while more than 60% capacity fading is observed in bulk cathode material. Furthermore, compared to bulk V2O5 cathode, the V2O5 nanosheets can charge and discharge at a much higher rate. Even at a rate of 50 C, the specific capacity of ultrathin nanosheet is still up to 117 mAh g-1.

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Fig. 6.5 Characterizations for {001}-oriented few-layer V2O5 nanosheets: (a) XRD patterns, (b) the TEM image and the corresponding colloidal acetone dispersion, (c) HRTEM the corresponding selected-area electron diffraction, (d) cycling performance at 0.2 C, (e) initial galvanostatic chargedischarge voltage profiles at a current density of 59 mA g-1 (0.2 C), (f) rate capability at various charge and discharge rates. (Reproduced from Rui et al. [43]. Copyright 2013, Royal Society of Chemistry)

The 2D nanostructure can enlarge the effective contact area between the electrolyte and the active material; reduce the diffusion distance of Li+ ions, which results in much rapid charge/discharge, so as to improve the storage capacity of Li+ ions. Therefore, it is a promising structure for the fast storage of Li+ ions. Nevertheless, the mechanical stress generated via the nanosheet during the cyclic process tends to cause the nanosheets to overlap and recombine which weakens the contact area of the nanosheets. Thus, the dispersion of the 2D nanosheets is a key factor for an excellent electrochemical performance. Recently, in order to address the problem of low capacity and poor cycle stability caused by low ion diffusivity and poor structural stability of V2O5, 3D nanostructure of V2O5 cathode materials has gradually become a research focus. Wang et al. [37] used an electrolytic deposition method to prepare a V2O5 nanotube array with a length of about 10 μm. Since the tubular structure has a big specific surface area and shortens the ion diffusion distance, the initial discharge capacity can reach 300 mAh g-1, which is twice as large as V2O5 film. Although the capacity decay is obvious during the cycle, it shows good capacity stability during 10 cycles with the capacity of 160 mAh g-1. In contrast with fast fading capacity reported by the Wang group, Chen et al. [25] obtained a V2O5 nanoflower (2–4 μm in size) consisting of nanosheets by hydrothermal method. The thickness of nanosheets is about 9 nm and the specific surface area is 15.9 m2 g-1. V2O5 nanoflower cathode exhibits excellent long-term cycling performance (Fig. 6.6). At 200 mA g-1 from 2.5 to 4 V, the specific capacity is reduced from 145 mAh g-1 to 128 mAh g-1 as well as the capacity fading rates is

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135

Fig. 6.6 (a) XRD pattern, (b and c) FESEM and (d) TEM images of V2O5 micro-flowers (VFs), (e) cycling performances of VFs and bulk V2O5 at a current density of 200 mA g-1, the inset is the cycling performances of VFs and bulk V2O5 at a current density of 500 mA g-1, (f) rate performance of the VFs at various current densities. (Reproduced from Chen et al. [25]. Copyright 2014, Elsevier B.V.)

only 0.08% after 150 cycles. Even at 2000 mA g-1, the capacity fading rates is 0.01% after 3000 cycles, which is much lower than the previously reported fading rates [21, 44]. Hierarchical nanostructures are reported to possess excellent capability to restrain the agglomeration of low-dimensional nanostructures and maintain the capacity [45]. Moreover, hierarchical nanostructures still retain high surface areas as well as much higher rate capability [46, 47]. A lot of hierarchical nanostructures are prepared as active electrode materials for LIB. For example, Dong et al. [23] obtained V2O5 microspheres with hierarchical structure by an easy two-step method. Compared to V2O5 nanorods, nanosheet-assembled V2O5 microspheres have presented higher cycling stability and specific capacity (Fig. 6.7). Compared to the obtained nanorods, V2O5 microspheres exhibit a high specific capacity of 275 mAh g-1 at 1 C and they still retain 243 mAh g-1 after 200 cycles. Particularly, even at 5 C after 500 cycles, the specific capacity of V2O5 microspheres as electrode materials is still up to 200 mAh g-1. Furthermore, Liu et al. [48] summarized the hierarchical V2O5 nanostructures as electrode material for LIBs (Table 6.1 and 6.2). The electrochemical properties (e.g., discharge capacity, rate capacity, and cycling stability) were mainly relied on the nanostructures. Low-dimensional structures were widely researched in the voltage window from 2.5 to 4 V (vs. Li/Li+). Within this voltage window, the phase transitions were regarded reversible. The conclusions show that the capacity

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Fig. 6.7 SEM images of V2O5 (a) nanorods and (b) nanosheet-assembled microspheres. (c) Cycling performance at a current density of 300 mA g-1 of 200 charge-discharge cycles. (Reproduced from Dong et al. [23]. Copyright 2015, Elsevier B.V.)

Table 6.1 Summary the morphologies of low-dimensional V2O5 nanostructures and their corresponding performance for lithium-ion batteries

LD-V2O5 structures Nanocrystal

Nanosphere Hollow sphere Nanorod Nanowire Nanotube Nanosheet Nanoplate Nanobelt

Specific capacity [mA h g1 ] 140 280 350 290 200

Current density [mA g1 ] 100 100 100 29.4 200

Voltage range [V vs. Li/Li+] 2.5–4 2–4 1.5–4 2–4 2–4

Capacity retention 78.5% 25.0% 14.2% 55.3% 92.6%

260 427 350 139.8 131.2 290 125 147 427 295 430

147 50 50 100 2000 58.8 1500 100 7.5 30 43.7

2–4 1.5–4 1.5–4 2.5–4 2.5–4 1.5–4 2.5–4 2.5–4 1.5–4 1.5–4 1.5–4

93.8% 59.7% 50.0% – 85.0% 93.8% 92.6% 98.0% 30.4% 25.4% 62.8%

Cycles 100 100 100 50 50

High-rate capacity [mA h g1 ] 100 – – – 74.5

High current density [mA g-1] 2000 – – – 2000

30 50 20 – 250 50 200 100 20 20 50

144 40 – 114.1 60 117 103 130 – – 119

2352 874 – 6000 15,000 14,700 5000 6000 – – 874

Reproduced from Liu et al. [48]. Copyright 2017, Wiley-VCH Verlag GmbH

retention of hierarchical V2O5 nanostructures has been enhanced to over 90% from 2.5 to 4 V after 50 cycles. On the contrary, the capacity retention was about 75–85% for low-dimensional nanostructures. In wider voltage ranges, it was found that the hierarchical nanostructured V2O5 cathode had an excellent cycling performance. After 50 cycles, around 40–50% and 75–95% specific capacity can be maintained from 1.5 to 4 V and 2 to 4 V, respectively. Furthermore, at 1500–2000 mA g-1, the hierarchical nanostructures present a superior specific capacity of 220–230 mAh g-1.

142 135 266

275.7

275 274 275 275 390 264 288 402 185.6

Porous Porous –

Sheet

Sheet Wire Wire Plate

Sheet Plate

Rod

Nanocage Octahedron Micro spheres

Nanowire

Nanoleaf Nanobelt

Nanosheet

2–4 2–4 2–4 2–4 2–4 1.75–4 2–4 2–4 1.5–4 2–4

– 300 50 30 30 500 50 50 294

2.5–4 2.5–4 2–4

Voltage range [V vs. Li/Li+] 2–4 2.5–4 2–4 2–4 2–4 2–4 2.5–4 2–4

300

73.5 100 300

Current density [mA g-1] 300 300 100 – 300 300 147 300

Capacity retention 74.30% 93.43% 76.68% – 86.64% 89.45% 84.47% 75.8%/ 64.5% 90.00% 104.44% 87.2%/ 75.2% 94.0%/ 88.4% 88.36% 79.93% 87.00% 68.00% 51.54% 78.00% 85.42% 41.79% 96.71% Cycles 100 50 55 – 50 50 350 50/ 100 100 60 50/ 100 100/ 200 50 50 50 50 50 100 50 50 50 – – 188 – – 100 82 – 62.7

230

86.7 96 223

High-rate capacity [mA h g-1] – 125 119 – 160 – 107.2 127

– – 2000 – – 5000 4000 – 2940

1500

8232 2000 2400

High current density [mA g-1] – 1200 2000 – 2400 – 2940 2400

Orthorhombic V2O5

Reproduced from Liu et al. [48]. Copyright 2017, Wiley-VCH Verlag GmbH

Triple shell Rattle-type

Yolk-shell

Specific capacity [mA h g-1] 284 137 283 286.4 262 275 146.8 290

Surface details Sheet Sheet Flake Rod Plates Plates Plates Sheet

Hollow V2O5 structures Hollow sphere

Table 6.2 Summary the morphologies of hierarchical V2O5 nanostructures and their corresponding performance for lithium-ion batteries

6.2 137

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The storage capacity of lithium ions will decrease with the increase of the number of cycles in LD nanomaterials, largely owing to LD nano-scale materials that have high surface energy with large specific surface area, which is facile to agglomerate. In contrast, 3D or hierarchical nanostructure has a good effect in suppressing agglomeration, and it not only has a high active specific surface area, but also provides a large space, which is better for the penetration of the electrolyte. Conversely, the 3D or hierarchical nanostructures are also self-supporting. During cycling process, the nano-components are not inclined to agglomerating, and the pore structure prepared in the electrode material can exist for a long term, with regard to the same conditions. 3D or hierarchical nanostructure of V2O5 electrode material exhibits superior electrochemical performance and better cycling stability compared to the LD V2O5 electrode material.

6.2.1.1

V2O5 Nanocomposites for Lithium-Ion Batteries

The combination of V2O5 and conductive materials (e.g., carbonaceous materials, conductive polymers, and metal oxides) can effectively enhance electron transport and the conductivity in LIB [49]. The structure stability can also be boosted by compositing strategies, including particle supporting and surface coating. Typical conductive materials (such as Ppy [50] and PEDOT [19]) are largely used as coating materials on V2O5 nanostructures, which cause the performance of electrode materials to improve effectively. Via a hydrothermal method, Chao et al. [19] fabricate self-supported, binder-free anode by growing a V2O5/PEDOT core/shell nanobelt array on ultrathin graphite foam. The electrode material shows superior cycling stability at high current density. At 5 C and 60 C, it delivers high discharge capacities of 265 mAh g-1 and 168 mAh g-1 respectively, which can retain 98% capacity retention after 1000 cycles. The whole surface area of V2O5 nanostructures can be practically enhanced by decoration of metal oxide nanoparticles. The metal oxides, such as MnO2 with high theoretical capacity [51], have been decorated to exhibit enhanced cycling stability and specific capacity of V2O5 cathodes. For example, by using MnO2 nanoparticles as the protuberance and PEDOT as the shell, the Mai group [51] designed the heterostructured nanomaterial and fabricated a cucumber-like V2O5/PEDOT coaxial nanowire enriched with MnO2 nanoparticles (Fig. 6.8). Compared with V2O5 nanowires, the heterostructured nanomaterial shows better cycling performance. At 100 mA g-1, the capacity fading decreases from 0.557% to 0.173% after 200 cycles. Combining carbonaceous materials (graphene, carbon nanotube, carbon fiber), it was found that V2O5 had enhanced cycle stability and electrochemical properties for LIB [52, 53]. For instance, Zhang et al. [53] described the fabrication of carboncoated nanocrystalline V2O5 by a special capillary-induced filling method. When used as a negative electrode material for LIB, the nanocrystals exhibit evidently enhanced rate capability and excellent cyclability. At 2.0, 5.0, 10.0 A g-1, the material delivers discharge capacities of 255, 168, and 130 mAh g-1, respectively.

6.2 Orthorhombic V2O5

139

Fig. 6.8 (a) Schematic illustration of the synthesis of V2O5/PEDOT &MnO2 nanowires (NWs). (b) FESEM image and EDS mapping of V, Mn, and S from V2O5/PEDOT&MnO2 NWs. (c) The cycling performance of V2O5 NWs, V2O5/PEDOT NWs, and V2O5/PEDOT&MnO2 NWs, at the current densities of 50 and 100 mA/g, respectively. (Reproduced from Mai et al. [51]. Copyright 2013, American Chemical Society)

Meanwhile, it still shows very high cycling stability. At 1.0 A g-1, a discharge capacity of 288 mAh g-1 is still reached after 50 cycles. Even at 10000 mA g-1, less than 2.3% capacity fading is obtained over 50 cycles. However, the thickness of carbon coatings needs to be controlled carefully because it might hinder Li-ion diffusion [54]. In view of this, mesoporous carbon layers are much more conducive to increase the ion transportation between electrolytes and VOs [15]. Therefore, Wang et al. synthesized a sandwich-like mesoporous composite by an easy solvothermal method [55]. As a positive electrode for LIB, the composite shows a high reversible capacity of 1006 mAh g-1 at 0.5 A g-1 after 300 cycles and excellent rate performance with a specific capacity of 500 mAh g-1 at 3.0 A g-1. It was found that the stacked structures of V2O5 and rGO can enhance capacity and cycling performance [15, 56]. Lee [26] reported an electrode structure where ultrathin V2O5 nanowires are uniformly integrated with graphene sheets to make graphene composite paper by a facile vacuum filtration step. The composite electrodes show excellent cycling performance that a large part of initial discharge capacity can be maintained after 100,000 cycles. These cycle numbers are two or three orders of magnitude larger compared to those of a typical battery. At the same time, a conductive active material such as graphene is compounded in the V2O5 nanosheets. Not only the agglomeration of the nanosheets can be prevented, but also the high conductivity of the electrode material can be fully utilized. Cheng et al. [57] used a facile solvothermal approach to obtain self-assembled V2O5 nanosheets/rGO hierarchical nanocomposite, which has high specific surface area and good

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Fig. 6.9 (a) Schematic illustration of the formation of 3D V2O5 nanosheets/RGO hierarchical nanocomposite. (b) SEM images of V2O5 nanosheets/RGO hierarchical nanocomposite. (c) Rate capability of V2O5 nanosheets/RGO nanocomposite electrode and V2O5 microsphere electrode at different current densities: 0.2 C (60 mA g-1), 0.5 C (150 mA g-1), 1 C (300 mA g-1), 2 C (600 mA g-1), 5 C (1500 mA g-1), 10 C (3000 mA g-1), 20 C (6000 mA g-1), and 50 C (15 A g-1). (d) Cycling performance and Coulombic efficiency of V2O5 nanosheets/RGO nanocomposite electrode and V2O5 microsphere electrode at a current rate of 2 C. (Reproduced from Cheng et al. [57]. Copyright 2013, Royal Society of Chemistry)

electronic/ionic conducting pathway (Fig. 6.9). When used as a cathode material, the nanocomposite exhibits higher reversible specific capacity and better rate performance than the bulk electrode sample. At 3 A g-1, 6 A g-1 and 15 A g-1, the electrode material reaches capacities of 138, 112, and 76 mAh g-1, respectively. At the same time, a discharge capacity of 102 mAh g-1 can be retained at 2 C after 160 cycles, exhibiting a good cycling performance. V2O5 has a unique layered structure and is suitable for lithium-ion storage. Compared with traditional negative electrode materials (e.g., LiCoO2, LiMn2O4, and LiFePO4), V2O5 exhibits high theoretical specific capacity and power density as LIB. However, the unstable structure itself and low conductivity make the actual specific capacity much lower than the theoretical value, resulting in a poor cycle stability. Therefore, a plenty of synthesis methods are used to prepare various nanostructures of V2O5 electrode materials (such as 1D nanowires, 2D nanosheets, and 3D nanoflowers) to improve the inherent morphology of the electrode material and enhance lithium storage capacity and specific capacity, so that V2O5 is a promising material served as the negative electrode for LIB.

6.2

Orthorhombic V2O5

141

Fig. 6.10 Abundance (atom fraction) of the chemical elements in earth’s upper continental crust as a function of atomic number. (Reproduced from Massé et al. [59]. Copyright 2015, Springer Science and Business Media)

6.2.2

V2O5 for Sodium-Ion Batteries

The initial research on room-temperature SIB dates back to the 1980s [58]; however, the researches on SIB were nearly interrupted in the following 30 years, which is largely owing to the lower energy density of SIB compared to LIB, and the safety problems caused by the use of Na in SIB. Recently, SIB has retook interest mainly owing to the abundant resources of sodium (2.83 wt%), along with similar properties of sodium to lithium because of the same group IA and electrochemical mechanisms of SIB to LIB. Figure 6.10 exhibits the abundance of several elements of the periodic table in earth’s crust [59]. Obviously, it is cost-effective to employ abundant and cheap materials. Therefore, SIB has been regarded as one of the most promising materials for inexpensive, high-energy rechargeable batteries [60, 61]. Among all, vanadium-based negative electrode materials have high theoretical energy densities, capacities and working voltages (Fig. 6.11). In addition, their plentiful electrochemical reaction with Na+, rich valent state of vanadium (V2+ ~ V5+) and low price, make vanadium-based cathode materials promising candidates for SIB. Recently, the vanadium-based materials can be largely separated into four groups [60]: VOs (such as VO2and V2O5); vanadium bronzes (such as NaV3O8, NaV6O15, NaxVO2, δ-NH4V4O10); vanadium-based phosphates (such as VOPO4, NaVOPO4, Na2(VO)P2O7, Na3V2(PO4)3, and Na7V3(P2O7)4); and F-containing vanadium-based polyanions (e.g., Na3V2 (PO4)2F3, NaVPO4F, and Na3(VOx)2(PO4)2F3-x).

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Fig. 6.11 Average discharge potential (V vs. Na+/Na) as a function of theoretical capacity (mAh g-1) and energy density (Wh kg-1) of representative V-based materials as cathodes for SIBs. (Reproduced from Wang et al. [60]. Copyright 2018, Royal Society of Chemistry)

The layered V2O5 used as an electrode material for SIB was initially reported in 1988 [62]; however, few of the V2O5-based SIB was studied in the next decades [63]. Recently, due to its high theoretical capacity, and a plenty of applications in SIB, V2O5 has been retook as a negative electrode material for SIB. Despite the layer spacing (4.37 Å) of orthorhombic V2O5 is beneficial to Li-ion intercalation/ deintercalation, it is hard for Na-ion intercalation/deintercalation because of the bigger radius of Na ion than Li ion [62]. Meanwhile, the innate low diffusion coefficient and conductivity of Na ion as well as big volume expansion makes V2O5 show inferior electrochemical performance [19]. To solve these problems, many methods have been exploited for V2O5, such as morphological control, crystallinity, crystal structure modification, and hybridization of V2O5 with other conductive materials. Nanostructures can greatly increase the electrochemical performance of V2O5 electrode materials for SIB. Nguyen et al. [64] prepared V2O5 nanoparticles (200–500 nm) by ball milling, and the material releases an initial discharge capacity of 208.1 mAh g-1. After Na-ion intercalation at the initial discharge, the structure of V2O5 transforms to a NaxV2O5 structure, which keeps stable during cycling, and the specific capacity retained 61.2% of the capacity of the 2nd cycle after 40 cycles. Similarly, Li et al. [65] prepared electrochemically grown nanocrystalline V2O5 with a two-layer structure by electrochemical growth directly on a steel supporting base in a VOSO4-based solution. The optimum specific capacity of 220 mAh g-1 is reached whenV2O5 has grown in the activation controlled voltage range (e.g., 0.8 V vs Ag/AgCl). After 500 cycles, it is only 8% capacity decay for the V2O5 electrode. A few 3D V2O5 micro-/nanostructures were fabricated to improve the performance

6.2 Orthorhombic V2O5

143

Fig. 6.12 (a) Schematic illustration of the evolution of V2O5 hollow nanospheres. (b) High magnification FESEM images of V2O5 nanospheres. (c) Atomic resolution HRTEM image, from which the interlayer structure of V2O5 can be directly observed. Right top inset in (c) is the FFT pattern along the (110) zone axis of V2O5. Inset in the middle of (f) is the simulated V2O5 (110) crystal plane (with V and O atoms marked as yellow and red in color, respectively. (d) 1st, 2nd, 3rd, and 100th cycle discharge and charge profiles of V2O5 hollow nanospheres at 20 mA g-1 current density. (Reproduced from Su et al. [67]. Copyright 2014, Royal Society of Chemistry)

of SIBs [26, 66]. For example, Vadivukarasiet al. [66] demonstrated that when orthorhombic V2O5 was homogeneously coated inside nanoporous carbon by a novel environmental hydrolysis deposition, which can present excellent electrochemical performance in SIB. The encapsulated V2O5 exhibits a high discharge capacity of 276 mAh g-1; meanwhile, the whole nanocomposite releases a capacity of 170 mAh g-1. The special structure of the nanocomposite markedly improves the rate performance and the capacity of V2O5. At 640 mA g-1, the composite releases a capacity of over 90 mAh g-1. Via solvothermal process and sintering step, Su et al. [67] obtained a V2O5 hollow nanospheres fabricated from hierarchical nanocrystals with mainly exposed {110} facets (Fig. 6.12). When used as a negative electrode for SIB, it reached a specific capacity of 150 mAh g-1 for the V2O5 hollow nanospheres. Theoretical modeling showed that the long-term cycling performance could be owing to the porous hollow spherical structure. Moreover, the (110) crystal planes of V2O5 nanocrystals have 2D diffusion pathways for Na-ion insertion, which are conducive to superior cycling performance as well as high rate capacity. From the abovementioned example, it can be shown that the morphology control can realize the performance enhancement of orthorhombic V2O5 via the exposed active plane. Su et al. [68] synthesized single-crystalline bilayered V2O5 nanobelts by a simple solvothermal method. When used as negative electrode in SIB, V2O5 nanobelts showed a high specific capacity of 231.4 mAh g-1 at 80 mA g-1. It is close to the theoretical capacity of Na2V2O5 formed on Na+ ion intercalation. V2O5 nanobelts also exhibited excellent cycle stability and superior high-rate performance.

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Fig. 6.13 (a) FESEM images of bilayered V2O5 nanobelts. (b–d) Bragg positions are represented by light green ticks. Refined structural model of the bipyramidal layered structure of V2O5 viewed along the [010] (b), [69] (c), and [001] (d) directions. The blue base-faced square-pyramidal VO5 unit consists of one vanadium atom and five oxide atoms, which are colored blue and red, respectively. (e) Lattice resolution TEM image, in which the (001) crystal plane can be identified. (f) Charge and discharge capacities versus cycle number at current densities of 80 mA g-1. (Reproduced from Su and Wang [68]. Copyright 2013, American Chemical Society)

These excellent electrochemical performances could be owing to the special bilayered-V2O5 nanobelts in which exposed (001) facets provide big interlayer spacing (11.53 Å) for Na+ ion intercalation/ deintercalation (Fig. 6.13). Furthermore, the exposure of (001) crystal planes of the V2O5 nanosheets is also conducive to charge transportation at the electrode/electrolyte interface [70]. Along the [001] axis, the thin nanosheets provide Na+ ions with short diffusion pathways. Electrochemical tests show that the material has a high discharge capacity of 216 mAh g-1 at 20 mAh g-1 as well as a good capacity retention of 73% at 100 mAh g-1 after 100 cycles. Adjusting the nanostructure of the material provides an unprecedented opportunity to take advantage of its functional properties; several researches [68, 70, 71] show that bilayered-V2O5 presented better Na+ ion storage performance compared to orthorhombic V2O5. For instance, the bilayered-V2O5 [71] shows much bigger layer spacing (13.5 Å of bilayered-V2O5 vs. 4.4 Å of orthorhombic V2O5). Because of the big interlayer spacing of (001) crystal plane, bilayered-V2O5 provides open channels

6.2 Orthorhombic V2O5

a

145

reversible Li+ ion insertion/extraction

c

Li+

d

4.37 Á

c

Na+

a

+

irreversible Na ion insertion/extraction

D-V2O5 300 nm

b +

d

12.3 Å

Na+

c reversible Li+/Na+ ion insertion/extraction

V2O5·nH2O

d Discharge capacity (mAh/g)

Li

a

400

300

200 V2O5·nH2O D-V2O5

100

0

0

10 20 Cycle number

30

Fig. 6.14 The structural models of α-V2O5 (a) and V2O5nH2O (b) viewed along the [010] direction. α-V2O5 is not suitable for reversible Na+ ion insertion/extraction, while V2O5nH2O is beneficial to reversible and rapid Na+ ion insertion/extraction. TEM images (c) of the V2O5nH2O xerogel. (d) Cycling performance of the V2O5nH2O cathode at the current density of 0.1 A g-1. (Reproduced from Wei et al. [72]. Copyright 2015, Royal Society of Chemistry)

for easy Na+ ion intercalation/deintercalation. Hence, the material releases a high initial specific capacity of 250 mAh g-1 at 20 mA g-1 from 1.5 to 3.8 V (vs. Na+/ Na). The layered structure V2O5 provides various inorganic or organic composites with an ability for inserting V2O5. To expand the interlayer spacing of V2O5, the Mai group [72] fabricated a V2O5nH2O xerogel (Fig. 6.14) by a facile freeze-drying process, which exhibited an enhanced electrochemical performance and released a high initial specific capacity of 338 mAh g-1 at 0.05 A g-1. With an interlayer spacing range of 8.8–13.8 Å, the intercalated V2O5 xerogel favors a large amount of Na+ storage [73]. Similarly, when served as a cathode for SIB, the hydrated V2O5 film shows improved Na storage performance [74]. At 200 mA g-1 from 1.5 to 0.5 V, the hydrated V2O5 film still retains a high discharge capacity of 166 mAh g-1 after 100 cycles. Moreover, it shows excellent rate performance with specific capacities of 133 mAh g-1 at 1 A g-1 and 121 mAh g-1 at 2 A g-1, respectively, which are much higher than those of orthorhombic V2O5 film (41 mAh g-1 at 1 A g1 and 34 mAh g-1 at 2 A g-1, respectively).

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The existence of severe “lattice breathing” needs to be given special attention to V2O5 xerogel because it could result in the collapse and shrinkage of the layeredV2O5 structure, which causes rapid capacity fading and the deactivation of V2O5 [75, 76]. Except the V2O5 crystal used as active electrode materials for SIB, amorphous V2O5 also exhibited promising electrochemical performances [76]. Amorphous structure can provide V2O5 with rich-opening ion diffusion pathways. For instance, the amorphous V2O5 based on Ni foam presented high reversible specific capacity of 241 mAh g-1 at the second cycle as well as capacity of 160 mAh g-1 after 100 cycles [77]. Presently, SIB is a very promising candidate for energy storage. Theoretically, SIB presents a complementary alternative to LIB due to the rich Na resources, low cost, and better safety. However, because radius of Na+ ions is much bigger than that of Li+ ions, resulting in much more invasive and destructive on the host material, one of the main problems with electrode materials of SIB is the capability to insert/deinsert Na ions smoothly. Given its high capacity and easy synthesis, V2O5 is a suitable cathode material to address the abovementioned issue for SIB. Nevertheless, there are a lot of research work to do for the development of V2O5-active electrode materials before SIB can be used in practical applications.

6.2.3

V2O5 for Other Emerging Rechargeable Batteries

Recently, V2O5 has attracted much attention as probable host materials for other metal ion (such as Mg2+, Zn2+, Ca2+, Al3+) storage in energy conversion field. Compared to lithium, magnesium (Mg) has several advantages such as a higher reduction, more abundant, and cost-effective. Theoretically, divalent Mg ion can release twice as much charge as Li ion [78–80]. As with SIB, study of MIB is in the exploration stage yet, despite interest is quickly rising [81]. Gregory et al. [82] first studied V2O5 for rechargeable MIB, which has a capacity of 196 mAh g-1 and considerably high open-circuit voltage of 2.66 V (vs. Mg/Mg2+). However, MIB has been studied rarely due to a shortage of Mg-compatible electrolytes, particularly after the application of LIB in the 1990s. Subsequently, Aurbach et al. [83–85] have made some important breakthroughs and showed that a full MIB is capable of thousands of cycles. With raising the material surface area and shrinking the particle size, the diffusion length of Mg ion can be shortened in the active material of V2O5. Amatucci et al. [86] reported Mg insertion into V2O5 nanocrystal, which was fabricated by combustion flame-chemical vapor condensation. The electrode releases a reversible specific capacity of 180 mAh g-1, although it shows a slow reaction kinetics. Similar work was reported for a GO/V2O5 composite [87]. When used as a dichloro-complex electrolyte, Du et al. [85] achieved a high specific capacity of 178 mAh g-1 and maintained 140 mAh g-1 at 0.2 C after 20 cycles The strong influence of H2O on the intercalation of Mg ion into microcrystalline V2O5/carbon black composite was first observed by Novák et al. [88]. They noted that the electrode materials cannot deliver any capacity in dry electrolytes; however,

6.2

Orthorhombic V2O5

147

when water is added to the electrolyte, the electrochemical performance is improved and the initial capacity of V2O5 electrode reaches 170 mAh g-1, but rapidly decays below 50 mAh g-1 after 20 cycles. However, addition of H2O in the electrolyte makes it incompatible with magnesium metal used as a positive electrode. To address this issue, an aerogel composite of rGO and V2O5 nanowires [89] was studied. The electrodes show a high specific capacity of 330 mAh g-1, and retain a good capacity retention of 81% at 1 A g-1 after 200 cycles. The structural water is conducive to retaining the interlayer spacing for Mg ion insertion and decreasing electrostatic effects by dipole interactions. Many researches have shown that V2O5 is a promising cathode for MIB, although poor transport kinetics and limited stability as well as rare Mg-compatible electrolytes remain a challenge. Inspired by the encouraging results of V2O5 electrode materials for SIB and MIB, the possibility of reversible insertion of other cations (e.g., Zn2+, Ca2+, and Al3+) are also studied recently. Presently, rechargeable aqueous ZIB is attracting much attention due to zinc resource abundance and low cost. Chen et al. [90] prepared a porous V2O5 nanofiber cathode via an electrospinning technique followed by calcination. A reaction mechanism is proposed based on phase transition from orthorhombic V2O5 to zinc pyrovanadate on first discharging and reversible Zn2+ (de)insertion in the open-structured hosts during subsequent cycling (Fig. 6.15). The electrode materials release a high reversible specific capacity of 319 mAh g-1 at 20 mA g-1 and maintain a 81% capacity retention after 500 cycles owing to the open and stable architecture. Li et al. [91] reported a 2D V2O5 nano paper consisting of V2O5 nanofibers and carbon nanotubes as cathode. The layered V2O5 with 2D diffusion pathways and the nanofiber morphology endow the material with short Zn-ion diffusion distance and may tolerate high volume expansion. Therefore, this open structure enables a high discharge capacity of 375 mAh g-1 and a long-term life of 500 cycles. Qin et al. [92]

Fig. 6.15 (a) Schematic illustration of the reaction mechanism of vanadium oxide electrode. (b) Galvanostatic discharge/charge profiles at 20 mA g-1. (c) Long-term cycling performance at current density of 2 C. (d) SEM image, and (e) TEM image of as-prepared V2O5 nanofibers. Inset (e) shows the HRTEM image. (Reproduced from Chen et al. [90]. Copyright 2019, Elsevier B.V.)

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synthesized 3D V2O5 hollow spheres for aqueous rechargeable ZIB by a solvothermal method. The as-prepared V2O5 electrode material consists of hollow sphere particles with a larger specific area. When used as negative electrode, V2O5 hollow spheres show better cycling performance and outstanding rate capability than commercial V2O5. The hollow spheres deliver a high specific capacity of 132 mAh g-1 and retain 82.5% capacity retention at 10 A g-1 after 6200 cycles. Bigger calcium ion radius (vs. Mg2+, Zn2+ and Al3+) [93], abundant calcium resources [94], and higher redox potential (vs. Mg2+, Zn2+ and Al3+) make CIB a markedly higher output voltage than that available with MIB, AIB, or ZIB. There are some reports on V2O5 insertion of Ca ion. For instance, the Amatucci group [86] studied the electrochemical reactivity of Ca ion in orthorhombic V2O5, which releases a reversible specific capacity of 200 mAh g-1 from 1 to 0.5 V. Using V2O5 as electrode for CIB, Hayashi et al. [95] reported a discharge capacity of 450 mAh g-1 at 50 μA cm-2. However, there is very slow diffusion kinetics of Ca2+ in V2O5 due to a high obstacle for Ca ion movement in the orthorhombic V2O5 [96, 97]. Compared to the crystalline V2O5, it was found that the amorphous V2O5 improved insertion capacity of Ca ion [95]. Bervas et al. [98] fabricated V2O5 xerogel/propylene carbonate (PC) nanocomposite via a sol-gel approach. The nanocomposite delivers a discharge capacity of 310 mAh g-1 at 6.58 mA g-1, which is 2–10 times better than the sample without PC. The excellent electrochemical performance is ascribed to the PC added in mesoporous V2O5 xerogel network, which ensures a very facile and rapid transportation of calcium ion to the xerogel network. As an abundant element [94], Al can provide 3 electrons for redox reaction; hence, it could deliver much higher capacity compared to Li+, Na+, Zn2+, and Mg2+. However, there are few researches on cathode materials that can reversibly (de)intercalate Al ion due to extremely high charge density of Al ion. In 2011, Jayaprakash et al. [99] initially reported the orthorhombic V2O5 nanowires used as cathode and AlCl3 added in 1-ethyl-3-methylimidazolium chloride used as electrolyte for rechargeable AlB. At 125 mA g-1, the cathode materials release an initial specific capacity of 305 mAh g-1 and maintain good capacity retention of 273 mAh g-1 after 20 cycles. However, the discharge plateau is only around 0.55 V, and the Coulombic efficiency of the battery is poor. Recently, the Gu group [100] reported the reversible Al ion storage in V2O5 nanowire and put forward the following mechanism (Fig. 6.16): In the first discharge stage, Al ion intercalates into crystalized V2O5 nanowires. At the same time, the insertion of Al ion results in the reduction of V5+, thus on the edge of nanowires, where an amorphous layer is formed. In the following process, along the nanowires’ edges, a new phase takes shape, and a two-phase transition reaction takes place. The researches on rechargeable metal ion (Mg2+, Ca2+, Zn2+, Al3+) batteries are still at its exploration stage and much more efforts should be made to push their further development. When used for rechargeable metal-ion batteries, V2O5 shows a high specific capacity; however, its average voltage plateaus, cycling performance, and rate capability appear far from practical utilization. There is more detailed work need to be done for the following issues [69]: (1) the cycling stability of nanostructures during long time operation, (2) improvement of the discharge

6.3 Bilayered V2O5

149

Fig. 6.16 (a) XRD patterns for V2O5 nanowires; inset: SEM image. (b) Galvanostatic charge/ discharge profiles. (c) Schematic diagram of Al3+ electrochemical insertion/extraction in crystallized V2O5 nanowire. (Reproduced from Gu et al. [100]. Copyright 2017, Elsevier B.V.)

voltage, (3) influences of interlayer distance change of V2O5 on the storage capacity and transport properties, and (4) co-intercalation with other species (e.g., water) in layered structure.

6.3

Bilayered V2O5

Different from the orthorhombic V2O5, bilayered V2O5 consists of V2O5 two layers separated by H2O layers and is generally described in the form of V2O5nH2O, lacking crystalline order in the long range. In the V2O5nH2O structure, the distance (in the [001]) of V-V atoms in two layers of the same bilayer is ~2.9 Å, which is much shorter than the value of 4.4 Å in the bulk α-V2O5. The low degree crystalline renders the amorphous V2O5 to be good host for big cations and offers more sites for ion insertion. In recent years, bilayered V2O5 has been greatly researched as monovalent and multivalent intercalation cathode materials for rechargeable batteries.

6.3.1

Bilayered V2O5 for Lithium-Ion Batteries

Similar to the orthorhombic V2O5, bilayered V2O5 was found to be a host for 4 equivalents of Lithium of each V2O5 unit cell. Later, a higher equivalent of lithium composition Li5.8V2O5 was achieved [101]. But the bilayered V2O5 encountered

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poor cyclability when more than 5 Li+/V2O5 were intercalated [102]. Combining atomic PDF, XPS, and DFT material characterization and electrochemical methods, the capacity fading of bilayered V2O5 occurred due to the formation and accumulation of LiOH poisonous species during cycling. Thus, it indicates that the interlayer water is inherently essential for stabilizing the bilayered structure. It has also been demonstrated that the water content impacts the electrochemical properties of bilayered V2O5 [103]. Specifically, the V2O50.3H2O film showed a higher discharge capacity and better cyclability than V2O5nH2O films (n = 0.1, 0.6, 1.6), due to the decreased water content and maintained interlayer distance. When electrochemically lithiated to 0.1 V (vs. Li/Li+), the bilayered V2O5 yielded a high reversible discharge capacity of 1000 mAh g-1 with the valent of vanadium changing from +5 to +2 [104]. Ex situ TEM images exhibited that the fibrous morphology of bilayered V2O5 could be retained when discharge to 0.1 V. This deep cycling ability was severely related to its interconnected porosity, high surface area, and the fibrous network. When V2O5nH2O nanowires were uniformly decorated with graphene paper, the decent capacities were maintained after 100,000 cycles, which are extraordinary excellent [105]. However, V2O5 naturally has moderate electronic conductivity (10-2 ~ 10-3 S cm-1), which hinders its rate performance. Coated on CNTs (as V2O5/CNT), bilayered V2O5 layer showed both faradic and capacitive (nonfaradic) electrochemical characterization [106]. At 1 C rate, V2O5/CNT presented a high specific capacity of 2700 C g-1, of which 2/3 is capacitive and 1/3 is Li+ intercalation capacity. V2O5/CNT composite structure possesses short lithium-ion diffusion path and fast electronic conductivity due to its entangled network structure and conductive additive CNTs (Fig. 6.17a). The V2O5/CNT nanocomposites exhibit ultra-high reversible discharge capacity of 850 mAh g-1 in the potential window of 1.5–4.0 V at 0.5 C, which is much higher than those of pure bilayered V2O5 without CNTs (200 mAh g-1) and commercially available orthorhombic V2O5 (280 mAh g-1). According to i(V ) = k1υ + k2υ1/2, the capacitive contribution in the capacity of V2O5/CNT nanocomposites is as high as 67% of total capacity. The improved capacity is owing to the special nanostructure, which makes short lithium-ion diffusion path and fast electronic conduction network. Based on the ultralong 1D nanostructure, bilayered V2O5 holds a high promise to be fabricated as a freestanding electrode, which is significant for its application in flexible LIBs. Jia et al. [54] fabricated the robust network of bilayered-V2O5 nanowires and CNTs via a hydrothermal reaction and filtration process. The as-synthesized freestanding V2O5/CNTs films display excellent mechanical strength and flexibility (Fig. 6.17b, c). The optimized freestanding V2O5/CNTs electrode (with 25% CNTs) achieves a capacity of 318 mAh g-1 at 1 C (280 mA g-1) in the voltage window of 1.8–4.0 V. Moreover, a specific capacity of 169 mA g-1 is still reached, even at 10 C. The sufficient electrolyte permeation, fast electronic transport network, and short lithium-ion diffusion length of V2O5/CNTs nanocomposites play a role in of the superior rate performance. Carbon coating is an efficient method to enhance the electronic conductivity and structural stability of electrode materials, but it is inappropriate for bilayered V2O5. This is because the effective carbon coating is usually obtained via high-temperature

6.3 Bilayered V2O5

151

Fig. 6.17 (a) Transmission electron microscopy image showing the absence of V2O5 particles or agglomeration over CNT surface. (Reproduced from Sathiya et al. [106]. Copyright 2011American Chemical Society). (b) The SEM image, digital photograph (inset) and (c) cross-sectional SEM image of V2O5xH2O nanowire/CNT freestanding electrodes. (Reproduced from Jia et al. [54]. Copyright 2012, Royal Society of Chemistry). (d) TEM image of V2O5xH2O-graphene composites. (Reproduced from Liu et al. [107]. Copyright 2015, Springer Nature)

treatment, while bilayered V2O5 is usually obtained via a mild method (10, 7, and 3 μm, respectively) and obtained different electrochemical performance. As a result, the VO2 nanobelts with the length of 7 μm and the width of 120 nm demonstrates higher discharge capacity (201.3 mAh g-1) in the initial cycle, and 53.5% capacity

6.4

VO2 (B)

157

Fig. 6.19 (a, b) SEM images of VO2/rGO nanorods and rGO nanosheets. (c) STEM image of the single VO2 rod coated by rGO layers. (d) Elemental mapping of a single VO2/rGO nanorod. (e–g) The charge-discharge profiles in different windows of 3.3–1.5 V, 3.0–0.05 V, and 2.0–0.05 V, respectively. (Reproduced from He et al. [127]. Copyright 2015 Royal Society of Chemistry)

retention after 50 cycles. The poor long-term cyclability may be caused by the volume expansion along with the Li+ insertion, and even the crystal structure pulverization. In a word, the electrochemical performance still needs to be improved. Although the above reported VO2 nanostructures have modest enhancement in the cyclic stability, the issues of rate capability dependent on the efficient ion and electron pathways are still a challenge. To overcome the problem, the 3D structures with enhanced electrical conductivity have been proposed [134, 135]. Yang et al. synthesized the single-crystalline VO2/rGO nano-ribbons via a bottom-up approach and demonstrated ultrafast Li storage performance. Figure 6.19h shows the morphology of the prepared VO2/rGO nano-ribbons with the width of 200–600 nm, the average thickness of around 10 nm, and the length of several tens of micrometers. The fabricated three-dimensional architectures enable ultrafast, supercapacitor-like charge-discharge capability. Niu et al. reported the VO2 nanowires assembled hollow microspheres through a hydrothermal method for high-performance LIBs [130]. Through the control of self-assembly of VO22+ ions and C12H25SO4- spherical micelles, the 3D nanostructures of VO2 nanowires assembled hollow microspheres were successfully synthesized. The VO2 nanowires are uniformly distributed with length of about 2 μm and the diameter of around 50 nm. When used as the electrode for LIBs, it displays high specific capacity of 163 mAh g-1 at 1 A g-1 and 90% of the initial capacity is retained at 100 mA g-1 after 1000 cycles. The superior electrochemical performance is owing to the high surface and efficient self-expansion/shrinkage of VO2 nanowires.

158

6.4.2

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Vanadium Oxide Nanomaterials for Electrochemical Energy Storage

VO2 (B) for Sodium-Ion Batteries

Most of the research papers about VO2 (B) have focused on the application for LIBs, the layered structure and nano-sized VO2 (B) may be promising electrode for SIBs. To explore the electrochemical performance of VO2 (B) as cathode for SIBs, Wang et al. conducted a simple hydrothermal reaction and synthesized the single VO2 (B) parallel ultrathin nanosheets for application in SIBs. The prepared VO2 (B) has an ultrathin nanosheet morphology with the diameter of around 50–60 nm and the length of hundreds of nanometers. When evaluated as the electrode between the voltage window of 1.5–4.0 V (vs. Na+/Na), the VO2 (B) nanosheets display high initial specific capacity of 214 mAh g-1 at 50 mA g-1 as well as the capacity of 108 mAh g-1 is still obtained at 500 mA g-1 after 50 cycles. The constant interlayer distance of NaxVO2 and reversible de-/intercalation between Na0.3VO2 and NaVO2 had been demonstrated through the XRD, XPS, and first-principles simulations analysis. The synthesized VO2/rGO nanorods by He et al. show better Na storage performance [127]. The nanorods exhibit improved capacity retention of about 90% at 60 mA g-1 after 200 cycles and high rate capacity of about 100 mAh g-1 at 800 mA g-1 over 50 cycles. The authors also demonstrate the difference of electrochemical reactions when VO2/rGO used as the electrode of LIBs and SIBs through XRD and TEM analyses. To further enhance the Na storage performance, Balogun et al. reported the carbon quantum dot surface-engineered VO2 interwoven nanowires (C-VOCQD) [136]. Figure 6.20a shows the interwoven nanowires get in touch with carbon cloth.

Fig. 6.20 (a) SEM image of carbon quantum dot surface-engineered VO2 interwoven nanowires. (b) CV curves of C-VOCQD between the voltage of 3.5 and 1.5 V at 0.2 mV s-1. (c) Rate performance from 0.3 to 60 C. (Reproduced from Balogun et al. [136]. Copyright 2016 Royal Society of Chemistry). (d) SEM image of VO2/MXene. (e) Long-term cycling performance at 0.1 A g-1. (f) Rate capability at 50, 100, 200, 400, 800, and 1600 mA g-1. (Reproduced from Wu et al. [137]. Copyright 2016, American Chemical Society)

6.4

VO2 (B)

159

The three-dimensional nanostructures have the advantage of excellent structural stability characteristics and high surface area. With the decreased Na+ diffusion paths and enhanced electron conductivity, the first 10 cycles have little capacity decay with a pair of redox peaks at 2.64/3.09 V (Fig. 6.20b) and high-rate performance of 133 mAh g-1 at 60 C (Fig. 6.20c). Wu et al. also fabricated a 3D flowerlike VO2/MXene hybrid architecture and obtained excellent positive electrode performance for NIBs [137]. Figure 6.20d shows the flower-like morphology assembled with hybrid nanosheets. The unique structure possesses enhanced electric conductivity, a big surface area, and superior structural stability. When employed as an anode for SIBs between the voltage window of 3.0 and 0.0 V, it delivers high initial specific capacity of 229.2 mAh g-1, after 200 cycles, it still exhibits high 141% capacity retention (Fig. 6.20e) and outstanding rate capacity of 206 mAh g-1 at 1.6 A g-1 (Fig.6.20f).

6.4.3

VO2 (B) for Other Emerging Rechargeable Batteries

VO2 (B) nanostructures for other rechargeable battery systems are rarely investigated, and only a few works are reported. Luo et al. prepared the VO2 (B) nanorods and nanosheets through a facile hydrothermal method, which served as promising negative electrode material for magnesium-ion battery [138]. The VO2 (B) nanorods exhibits better Mg-storage capacity of about 0.61 mol Mg for per unit of VO2 with a high discharge plateau of 2.17 V (vs. Mg2+/Mg). It also shows high specific capacity of 394 mAh g-1 and retains a capacity retention of 52.2% at 50 mA g-1 after 60 cycles. However, the electrochemical performance remains to be improved. Pei et al. explored the application of VO2 (B) nanoflakes as negative electrode of hybrid Mg-Li-ion battery. The nanoflakes display high specific capacity of 244.4 mAh g-1 with the stable plateau at about 1.75 V, which corresponds to the highest energy density of 427 Wh kg-1. It also shows excellent long-term cycling stability (75% capacity retention after 100 cycles) and high rate performance (138.8 mAh g-1 at 500 mA g-1). Chen et al. reported the VO2 (B) nanorods as an insertion negative electrode for aqueous ZIBs. The synthesized VO2 (B) are uniform nanorods with a diameter of around 100–150 nm and the length of around 1–2 μm. When evaluated as the cathode of ZIBs, the VO2 nanorods release high initial specific capacity (325.6 mAh g-1), superior cyclic performance (capacity retention of 86% at 3 A g-1 after 5000 cycles), and good rate capability (72 mAh g-1 at 5 A g-1). The highly reversible single-phase reaction zinc storage mechanism has been exhibited through in situ XRD and the analysis of calculated lattice parameters.

160

6.5

6

Vanadium Oxide Nanomaterials for Electrochemical Energy Storage

V6O13

V6O13 is a typical Wadsley phase with a slightly deformed VO6 octahedron in its structure, in which vanadium has valence states of +4 and +5 [139–141] (Fig. 6.21a). They are connected by an angle sharing single slice and an edge sharing double slice parallel to the plane (010). Corners share single sheets to form a “V2O5-layer of octahedral,” by contrast, edges share a double sheet to form the “VO2 (B) layer of the octahedron” [143]. Obviously, the crystal structures of V2O5, VO2 (B), and V6O13 are very similar. In addition, V6O13 phase is usually formed as an intermediate in the synthesis of VO2 under reduction conditions [144]. V6O13 is considered a promising candidate of cathode materials for battery because of its high electrochemical capacity [144–149]. The theoretical capacity and energy density are 417 mAh g-1 and 890 Wh/kg, respectively, equivalent to 8 Li/unit of formula (average formal charge of vanadium has dropped from +4.33 to +3) [139, 142]. However, limited electron sites in V6O13 reduce lithium uptake to a certain extent, leading to a capacity of about 310 mAh g-1 [139, 150]. Zou et al. [139] reported that anoxic non-stoichiometric V6O13-y nanosheets can effectively improve their electrochemical properties in LIBs. Furthermore, Ding et al. [151] reported the preparation of 3D V6O13 nanoplates from continuously interconnected 1D nanogrooves at room temperature using an easily realized method of dissolved oxygen-redox self-assembly. When employed as cathodes in LIBs, the 3D nanotextiles exhibit excellent lithium storage properties. In addition, recombination with the collector fluid is also an effective choice. Wu et al. [152] reported that V6O13 was loaded on etched stainless steel mesh substrate, and they studied the effect of the mass (thickness) of the active material on the rate capability, which was believed to affect lithium ion and electron conduction. Moreover, Tian et al. [153] effectively enhanced the electrical transport and cycling performance of ultrathin V6O13 nanosheet by pre-lithium treatment. As shown in the Fig. 6.21b, Meng et al. [142] clarified the complex lithiation/delithiation conversion mechanism of highcapacity negative electrode material V6O13. These works reveal the high specific

Fig. 6.21 (a) Crystal structure of monoclinic V6O13 (C2/m). (b) Proposed mechanism of the first discharge and charge process of V6O13 determined by in situ XRD. (Reproduced from Meng et al. [142]. Copyright 2017, American Chemical Society)

6.6

V2O3

161

capacity of V6O13 and provide a reference for the development of this material for the negative electrode of secondary lithium batteries. We believe that the research on the utilization of V6O13 in the field of energy storage will be more thorough.

6.6

V2O3

Figure 6.22a shows the crystal structure of V2O3 [154], where rich tunnels are exhibited in the crystal structure of V2O3, which is conducive to the insertion/ desertion of metal ions. Itinerating along the V-V chains in V2O3, the electrons in

Fig. 6.22 (a) The supercell structure of V2O3. (Reproduced from Sun et al. [154]. Copyright 2011, Elsevier B.V.) (b) TEM image and (c) the first four CV curves of uniform V2O3@NC hollow sphere. (Reproduced from Han et al. [155]. Copyright 2018 Royal Society of Chemistry). (d) TEM image and (e) galvanostatic charge-discharge curves for the 2nd, 5th, 10th, 50th, and 100th cycles at 100 mA g-1. (Reproduced from An et al. [156]. Copyright 2017, Science China Press)

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6 Vanadium Oxide Nanomaterials for Electrochemical Energy Storage

the V-3d orbital display metallic behavior, which make them relatively high conductivity [155]. These features endow V2O3 as one ideal candidate of electrode materials. However, different from other vanadium oxides, V2O3 shows a low capacity in the voltage range of over 2 V (vs Li+/Li) with no obvious plateau in cycling curves and has hardly been used as a negative electrode for LIBs [157]. Moreover, V2O3 is considered as a promising candidate of anode for LIBs due to a high theoretical capacity of 1070 mAh g-1 [158]. Pure V2O3 delivers relatively poor electrochemical properties compared with V2O3 composite materials due to volume expansion and inferior conductivity. Hence, many V2O3 composite materials as positive electrode materials for LIB have been investigated recently, which can be divided into low-dimensional and three-dimensional composite materials [159–161]. Low-dimensional V2O3 composite materials include nanowire, nanorod, nanobelts, and nanoparticles. Li et al. synthesized peapod-like V2O3 nanorods encapsulated into carbon used as highperformance positive electrode material [162]. The ultralong peapod-like V2O3@C nanowires as flexible and freestanding electrode has a reversible specific capacity of 210 mAh g-1 at 0.1 C. It displays appreciable rate performance, with 68% capacitance retention at various current densities ranging from 0.1 to 1 C, and superior prolonged cyclic performance with no capacity fading after 125 cycles. This is superior to bare V2O3 nanorods and active carbon mixed V2O3 nanorods. The improved performance can be owing to existence of the outer conductive carbon network, further relieving the volume expansion of V2O3 nanorods. Sun et al. reported highly ordered lamellar V2O3-based hybrid nanorods for aqueous LIBs [154]. When served as positive electrode materials for aqueous LIBs, V2O3-based hybrid nanorods exhibit a reversible capacity of 131 mAh g-1, nearly several times higher compared to V2O3 nanocrystals (90 mAh g-1) and bulk V2O3 (73.9 mAh g1 ). Xiao et al. reported V2O3/rGO composite material as positive electrode material for LIB [158]. In the synthesis material, V2O3 nanoparticles, which show sizes of 5–40 nm, are encapsulated by rGO sheets. When tested at 100 mA g-1, the V2O3/ rGO nanocomposite reaches a high specific capacity of 823.4 mAh g-1. It still retains the capacity of 406.3 mAh g-1, even at 4 A g-1. Recently, some work on the 3D V2O3 composite materials were published, including carbon-supported nanosheet-assembled microspheres, porous V2O3/C, porous V2O3@C hollow spheres, V2O3/3D carbon nanosheet, and porous V2O3/ VO2@C heterostructure, which show excellent electrochemical performance for LIBs [163–165]. For example, Niu et al. synthesized the V2O3/C microspheres as a stable positive electrode for LIBs [166]. The as-synthesized microspheres with diameter of ~2 μm are composed of V2O3/C nanosheets with thickness of ~10 nm, in which V2O3 nanocrystals are encapsulated in carbon matrix. When served as positive electrode materials for LIBs, the V2O3/C microspheres display excellent cycling performance. Even at 2000 mA g-1 after 9000 cycles, V2O3/C microspheres still remains at 98% retention of the discharge capacity at second cycles. The superior cycling stability is owing to the nanosheet-assembled microsphere

6.6

V2O3

163

structure, which hinders the self-agglomeration of nanosheets and buffers the strain resulting from Li-ion insertion/desertion. Wang et al. reported porous V2O3/ VO2@carbon heterostructure structure electrode for LIBs [164]. The sample shows a high specific capacity of 503 mAh g-1 and 453 mAh g-1 at 100 mA g-1, 500 mA g-1, respectively, as well as superior cycling performance with a capacity of 569 mAh g-1 at 100 mA g-1 after 105 cycles, which are attributed to the synergic effects of two-phase hybridization and porous nanostructure. Han et al. reported uniform carbon-confined V2O3 hollow spheres (V2O3@NC) as anode material for stable and rapid Li storage [155]. In the as-synthesized V2O3@NC hollow spheres, the V2O3 spheres, which were composed of many nanoparticles, had thin shells around 30 nm in thickness and an average diameter of around 200 nm. The carbon layer was about 2 nm in thickness confined onto the external V2O3 spheres. At 200 mA g-1, the V2O3@NC hollow spheres release a high initial specific capacity of 915 mAh g-1 and high initial coulombic efficiency (80.3%). The specific capacity still retained 811 mAh g-1 after 120 cycles. Furthermore, the electrode material showed a high reversible specific capacity of 472 mAh g-1 at 2 A g-1 after 700 cycles. The excellent electrochemical performance was owing to the unique carbon-confined hollow sphere structure. Except for V2O3 composite materials as positive electrode materials for LIB, the investigation of V2O3 composite materials as positive electrode materials for SIBs is still rare. An et al. synthesized coherent porous V2O3/C nanocomposites by hydrothermal method for high-performance SIBs [156]. At 100 mA g-1, the specific capacity of the V2O3/C nanocomposites, V2O3 nanoparticles, and pure carbon on the first cycle is 744.6, 281, and 879.6 mAh g-1, respectively. The V2O3/C nanocomposites deliver the specific capacity of 270 mAh g-1 after 150 cycles, which is much higher compared to pure carbon (100 mAh g-1) and V2O3 nanoparticles (98 mAh g-1). The rate capability of V2O3/C nanocomposites was tested at different current densities from 100 to 1000 mAh g-1. The specific capacity of the electrode still remains 270.8 mAh g-1 when the current density returned to 100 mAh g-1. However, the pure V2O3 nanoparticles only show 62.9 mAh g-1. Xia group reported graphene foam-supported CNTs decorated with V2O3 nanoflake arrays (GF + V2O3/CNTs) as an anode material for SIBs [167]. At 100 mA g-1, the GF + V2O3/CNTs can deliver a high reversible specific capacity of ~612 mAh g1 with an initial Coulombic efficiency of 72.5%. Even at 10 A g-1, the specific capacity of GF + V2O3/CNTs anode is still 207 mAh g-1, higher than GF supported V2O3 arrays without CNTs (23 mAh g-1). Moreover, the GF + V2O3/CNTs also exhibit excellent cycling stability. After 6000 cycles at 2000 mA g-1, the specific capacity of GF + V2O3/CNTs anode still reaches 402 mAh g-1, corresponding to the capacity retention of 100%. Besides, the capacity retention of 70% is reached at 10 A g-1 after 10,000 cycles. The improved Na-ion storage performance of GF + V2O3/CNTs is ascribed to the existence of CNTs and GF substrate, which provides high electronic conductivity and hinders the self-aggregation of V2O3 nanoflakes.

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V3O7H2O

Théobald and Cabala first reported V3O7H2O (H2V3O8) synthesized by hydrothermal method in 1970 [168]. The crystalline structure of H2V3O8 is orthorhombic symmetry (S.G. #62 Pnam). Oka et al. [169] first described the crystalline structure model of H2V3O8 in 1990. H2V3O8 is a layered structure of vanadium oxide in which V3O8 layers pile along the a axis. The V3O8 layers are composed of [VO5] square pyramids and [VO6] octahedrons, which are links by sharing corners and edges (Fig. 6.23). The V valence in H2V3O8 is a mixed state with V4+/V5+ ratio of 1: 2, which makes H2V3O8 with higher electronic conductivity. Compared with orthorhombic V2O5, H2V3O8 has a larger layer spacing, which can buffer the structure expanding/shrinking during charging/discharging. In addition, H2V3O8 owns the unique 1D nanostructure, which effectively promotes the electron and ion transmission. Therefore, H2V3O8 has been widely studied in the field of energy storage, including LIBs [170–174], SIBs [174, 175], PIBs [176], MIBs [177, 178], and ZIBs [179, 180]. Zhu et al. [171] studied the Li+ storage performance of H2V3O8/rGO composite in 2014. The H2V3O8/rGO shows a high specific capacity of 256 mAh g-1 at 100 mA g-1. Zhang et al. [172] reported the H2V3O8 nanobelts and rGO sandwich composites (H2V3O8-rGO) as cathode material for LIBs. At 1.0 A g-1, the H2V3O8rGO showed the reversible capacity of 120 mAh g-1 over 400 cycles, which

O(2) O(5) V(1) V(2)

V(3)

O(6) O(3)

O(1)

O(8)

O(7)

b O(4)

a

Fig. 6.23 Crystalline structure of V3O7H2O viewed along the c-axis. (Reproduced from Theobald and Cabala [168]. Copyright 1970)

6.7

V3O7H2O

165

Delithiation

Fissures

Lithiation

H2V3O8 NRs

Li+

I am breathing!

Lithiation

Li+

Breathing fissures

V2O3

H2V3O8

Lithiation

Li+

VO2

Li+

Delithiation

Fig. 6.24 The lithium storage mechanism of H2V3O8. (Reproduced from Wu et al. [170]. Copyright 2019, American Chemical Society)

corresponds to the capacity retention of 87%. Wu et al. [170] studied the lithium storage mechanism for H2V3O8 by in situ transmission electron microscopy (Fig. 6.24). The results indicated the local phase is transformed from H2V3O8 to V2O3 through the intermediate phase of VO2 and large structural cracks are piecemeals generated during discharge process. The large structural cracks were healed upon charge process with the VO2 phase as product. In the subsequent cycles, a highly reversible phase transition between VO2 and V2O3 is observed. Wang et al. [175] developed the sodium storage performance of flexible and binder-free H2V3O8 nanowire membrane. At 10 mA g-1, H2V3O8 nanowire film showed a high capacity of 168 mAh g-1. Moreover, the sodium storage mechanism is investigated by XPS, FTIR, and ex situ XRD measurements. In order to investigate the basic mechanism of Na+ and Li+ transmit in the H2V3O8 nanostructure electrodes. Xu et al. [174] designed a single nanowire device with multi-contacts to in situ explore the transportation performances of Na+ and Li+ at the single nanowire scale (Fig. 6.25). The results indicated that the conductivity of H2V3O8 nanowire decreases after insertion/deinsertion of Li+ or Na+, demonstrating that the similar electrochemical reaction mechanisms for H2V3O8 in LIBs and SIBs. However, compared with Li+, the structure failure and conductivity degeneration of H2V3O8 for sodium are more serious during the charge/discharge processes, which are mainly due to the larger ion radius and higher diffusion barrier of Na+. In addition, Mohadese et al. [176] reported that the H2V3O8 as cathode materials for PIBs. The H2V3O8 exhibited a discharge capacity of 168 mAh g-1 with 75% capacity retention at 5 mA g-1over 100 cycles and the K+ storage mechanism of H2V3O8 was investigated by ex situ XPS and in situ XRD.

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Fig. 6.25 The single nanowire device based on H2V3O8 nanowire electrodes. (Reproduced from Xu et al. [174]. Copyright 2015, American Chemical Society)

The applications of H2V3O8 in multivalent ion batteries have also been extensively studied. Tang et al. [177] and Mohadese et al. [178] reported the H2V3O8 nanowires as cathode materials for MIBs. The H2V3O8 nanowires showed a high discharge capacity of 304 mAh g-1 and long cycle life of 100 cycles. The Mg2+ storage mechanism of H2V3O8 was studied by in situ XRD, ex situ FTIR, and ex situ XPS. Unfortunately, H2V3O8 is incompatible with those electrolytes that are compatible with Mg metal anodes. Therefore, Tang et al. [177] assembled a Mg-based Li+/Mg2+ hybrid battery based on H2V3O8 as a cathode material, which combines the advantages of LIBs and MIBs. He et al. [179] and Pang et al. [180] studied the Zn2+ storage performance of H2V3O8 nanowires. H2V3O8 showed a high discharge capacity of ~400 mAh g-1, outstanding cycle stability (over 1000 cycles), and excellent rate performance (up to 5 A g-1). The Zn2+ storage mechanism of H2V3O8 was researched by Raman, XPS, and ex situ XRD. These results demonstrate that H2V3O8 is an excellent cathode material for energy storage device and worthy of further exploration and development.

6.8

Summary and Future Directions

Owing to its low price and high specific capacity, vanadium oxide cathode materials are receiving widespread attention and are likely to develop into a new generation of LIB, although there are some drawbacks such as low conductivity and poor stability of the structure. Studies have shown that vanadium oxide cathode materials with various nanoscale structures have different metal ion intercalation capabilities. When

References

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the material scale is reduced from micron to nanometer, the energy storage and release properties of the electrode materials are greatly improved, mainly due to the benefit of the unique morphology and large surface area caused by nanostructures of electrode materials. The stability of vanadium oxide in charge/discharge cycling process is still unsatisfactory because the nanostructures are susceptible to cracking, so that the single method is difficult to completely solve the defect of vanadium oxide. In the future, the modification of vanadium oxide will be combined with various methods such as ion doping, surface coating, and strategies for complexing with conductive materials to enhance the conductivity and cycle performance of vanadium oxide. In summary, future research on vanadium oxide will focus on the following aspects: (1) the reversible transformation mechanism of layer structure vanadium oxide during metal ion insertion/deinsertion; (2) the influence of the nanostructured morphology on the properties of vanadium oxide cathode materials; and (3) cost-effective synthetic technologies of vanadium oxide for future practical application.

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Chapter 7

Vanadate Nanomaterials for Electrochemical Energy Storage

7.1

Introduction

Vanadate-based materials are important parts of vanadium-based materials. Owing to combine with other clusters or metal ions, they are thought to be important derivatives of vanadium oxides. At present, numerous vanadate-based materials were researched deeply in energy storage field. The wide vanadium element valence state and flexible V-O bonds in polyhedral lead to a big family of vanadates. From the previous researches on vanadate-based materials, different kinds of clusters or metal ions, such as alkali metal ions (Li+, Na+, K+, Rb+), alkali-earth metal ions (Mg2+, Ca2+, Sr2+), transition metal ions (Ag+, Zn2+, Cu2+, Co2+, Fe3+, etc.), and other clusters (H+, NH4+, In3+, Bi3+, Al3+, etc.), can composite with vanadium oxide. With the introduction of different components, the resulting vanadate-based electrodes exhibit various electrochemical performance, making vanadates possess a promising applications potential in electrochemical energy storage fields. This part introduces the research status and electrochemical performance of the typical vanadates. The different vanadates are discussed in detail in this chapter.

7.2 7.2.1

Alkali Metal Vanadates LiV3O8

LiV3O8 has been researched deeply in energy storage fields since it has appropriate working voltage window and enough guests ions storage sites [1–3]. Moreover, owing to the occupation of lithium ions in layered spacing, the crystalline structure stability of LiV3O8 is supposed to be better than that of orthogonality V2O5 [4]. In 1981, Nassau et al. researched LiV3O8 as cathode material for LIBs at first. According to the previous research [5], the LiV3O8 vanadium oxygen layer is © Springer Nature Switzerland AG 2023 L. Mai et al., Vanadium-Based Nanomaterials for Electrochemical Energy Storage, https://doi.org/10.1007/978-3-031-44796-9_7

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constitutive of quadruple chains of VO6 octahedra and double chains of VO5 square pyramids. Since the vanadium-oxygen polyhedrons connected through terminalcorner oxygen atoms, the layered structure is flexible and the guest lithium ions could occupy all octahedral sites and partial tetrahedral sites. It is worth mentioning that the structural lithium ions in LiV3O8 are immobile in the layered spacing. When LiV3O8 is used as cathode material in LIBs, the capacities are ~280 mAh g-1 and ~372 mAh g-1 when 3 and 4 mole lithium ions inserted, respectively [1, 6]. In the discharge curves of LiV3O8, multiple voltage plateaus appeared, illustrating the multiple-step intercalation of lithium ions. The discharge voltage is around 2.6 V; when it multiplied by the theoretical capacity, LiV3O8 could possess considerable energy density, about 728 ~ 967 Wh kg-1. However, LiV3O8 still suffers from many puzzles, such as poor lithium diffusion rate (~10-13 cm2 s-1) and sluggish electronic conductivity (~10-6 S cm-1), which will lead to unsatisfactory rate property in practical application [6]. Many previous studies have proved that the different synthesis methods will lead to the obvious effect on the electrochemical property of LiV3O7 [7–9]. Diverse synthesis methods result in much influence on the process of crystallization, surface defects or state, particle size, and morphology, which will impact the lithium ions diffusion coefficient and electronic conductivity. In order to optimize the electrochemical properties of LiV3O8 cathode in LIBs, many strategies have been adopted by the different researchers, including morphology design [1, 4, 10], modification surface [11–14], adding conductive coating [15], doping metal ions [13], and introducing oxygen-deficient [6]. When the particle size decreases to nanoscale, the diffusion distance of lithium ions in materials will certainly shorten and the charge/discharge kinetics will be improved. In 2004, Xu et al. were the first to report the LiV3O8 nanoscale [16]. In this work, LiV3O8 nanorods were synthesized by the hydrothermal method, and they exhibit the highest discharge capacity of 302 mAh g-1. Moreover, LiV3O8 nanorods possess the wide working voltage window of 1.8–4.0 V. The cycling performance of LiV3O8 nanorods is not very well, since only 92% capacity remained after 30 cycles. For improving the rate and long-term cycling property, Luo et al. applied topotactic synthesis method to fabricated ultralong LiV3O8 nanowires [10]. Because LiV3O8 possess a similar layered structure with H2V3O8 nanowires, H2V3O8 were selected as the precursor. After topotactic lithium ions intercalation, the precursor could still maintain the ultralong morphology. When ultralong LiV3O8 nanowires were applied in the LIB as cathode materials, it exhibited excellent high-rate electrochemical performance with discharge capacity of 137 mAh g-1 at 2000 mA g-1. Furthermore, a high capacity of 120 mAh g-1 could be maintained even after 600 cycles, which corresponds to 0.022% capacity fading every cycle. Based on the materials design, electrochemical property of LiV3O8 can be further optimized by decorating conducting materials, surface morphology modification, or metal ions preintercalation. For instance, Al2O3-modified LiV3O8 was able to display better cycling performance. The authors thought that the Al2O3 layer could protect the nano-sized electrode surface and prevent the potential side reaction in the cycling process [17–19]. Sun et al. reported another facile and important surface modification strategy, involving a self-transformation process of external LiV3O8 to

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LixV2O5 under a certain reduction atmosphere [20]. When the LixV2O5 layer is regulated with various thickness, the compound effect between LiV3O8 and LixV2O5 layer will be more excellent than other conventional coating method. In addition, the surface LixV2O5 layer possesses a much faster Li+ diffusion rate than pure LiV3O7. From the result, the LixV2O5/LiV3O8 nanoflakes exhibit improved cycling performance (82% initial capacity could be remained after 420 cycles) and rate property. Conductive materials decoration strategy is another method for resolving the problem of inferior electronic conductivity. The incompatibility of V5+ in LiV3O8 with the reducing atmosphere required for carbon coating makes the synthesis of carboncoated LiV3O8 a difficult thing. On the contrary, graphene-coated LiV3O8 nanomaterials were synthesized by different researchers, such as LiV3O8/rGO nanostructure [21], graphene-nanosheet-wrapped LiV3O8 nanoarchitecture [22], and sandwich-shaped LiV3O8/multilayer graphene nanomembranes [23]. Different nanocomposites structure improved rate and cycling properties owing to the obviously improved ion/electron transportation efficiency. Other conducting-materialscoated LiV3O8, such as LiV3O8/PANI [24], LiV3O8/Ag [25], and Cr-coated LiV3O8 [26], were also synthesized successfully. For improving the electronic conductivity coefficient and ion diffusion efficiency, doping metal ions is an effective way. The Cao group [27] synthesized Mo-doped LiV3O8 nanorods, in which 3.5% oxygen vacancies and 25% V4+ were founded by XPS measurements. The introduced oxygen vacancies and reduced vanadium element enhanced the electronic transportation efficiency and optimized the exoteric structure for Li+ insertion, reducing the reaction resistance and enhancing the cycling performance. As to the LiV3O8 practical application, an important fact has to be discussed. Unlike the commercialized cathodes materials (LiCoO2, LiFePO4, and LiMn2O4) in LIBs, the lithium ions in the crystal structure of LiV3O8 are immobile, which determines that LiV3O8 is not compatible with commercial graphite anode. For introducing the lithium ions in the system, the LiV3O8 cathode must combine with the Li-contained anode for practical applications [26, 28]. Recently, the rechargeable batteries with Li metal anodes have attracted much attention and relevant researches have got a great improvement [29, 30]. Due to its moderate voltage and high capacity, LiV3O8 has a considerable theoretical energy density of 728 ~ 967 Wh kg-1. The LiV3O8/Li batteries can achieve a high energy density of 250 ~ 350 Wh kg-1, which is apparently better than the commercial LIBs (130 ~ 180 Wh kg-1). With the optimization of materials, LiV3O8 will be one of the best choices for the cathodes. Except for the application in LIBs with organic electrolytes, LiV3O8 is also considered as one of the best anodic candidates in aqueous LIB, owing to its suitable voltage, high capacity, and good structural reversibility [31–33]. The redox reaction potential of LiV3O8 is slightly higher than the hydrogen evolution potential, which prevents the decomposition of the aqueous electrolyte during electrochemical reaction. In 2000, Kohler et al. were the first to apply LiV3O8 as an aqueous LIB anode with 1 M LiCl (or Li2SO4) as the electrolyte and LiNi0.81Co0.19O2 as the cathode [34]. The assembled aqueous battery shows an operating voltage range of 1–1.2 V, but only 20–40 mAh g-1 (weight of anode + cathode) can be obtained. Moreover,

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the system only exhibited an inferior stability of capacity retention of about 25% after 100 cycles. After that, various aqueous batteries with LiV3O8 as anode were reported, including LiV3O8//LiFePO4 [35], LiV3O8//LiMn2O4 [36], LiV3O8// LiCoO2 [37, 38], and LiV3O8//LiFe0.5Mn0.5PO4 [39]. The early relevant researches generally illustrated that LiV3O8 exhibited the worse electrochemical properties in aqueous electrolytes than in organic systems. The possible reasons were investigated by several researchers from different aspects, and it was speculated that the dissolution of vanadium element was the main cause [40]. In addition, H+ co-intercalation may also lead to the crystal structure degradation and then result in fast capacity fading [34]. In a word, except for electrode optimization, electrolyte modification is also very important for LiV3O8 as aqueous battery electrodes. The relevant details about the optimization strategies for aqueous batteries have been reviewed by the Wang et al. [31], Tang et al. [33], and Kim et al. [32], respectively. In short, for LiV3O8, a neutral electrolyte (such as Li2SO4 and LiNO3) with pH value of 7, is the best selection to reduce the dissolution of the electrode. Moreover, higher electrolyte concentration could bring about better electrochemical performance. Besides, it is also important to match with a suitable cathode material that is adaptive to neutral electrolyte, such as LiMn2O4 and LiFePO4. For example, the LiV3O8/LiFePO4 aqueous battery in a high-concentration LiNO3 electrolyte (9 M) without dissolved oxygen could exhibit excellent electrochemical performance [35].

7.2.2

Na1 + xV3O8

Na1 + xV3O8 is the analogue of LiV3O8 with the substitution of Li by Na [41]. Similar to anchored Li ions in LiV3O8, the existing Na ions are also immobile in the structure. Because of the larger radius of Na ions, Na1 + xV3O8 has a larger interlayer distance than LiV3O8, which indicates faster lithium diffusion kinetics. In addition, Na element is more abundant than Li element, which leads to the lower cost of Na1 + xV3O8 than LiV3O7. Owing to these superiorities, Na1 + xV3O8 has attracted much attention in recent years. Wang et al. synthesized several kinds of Na1 + xV3O8 nanomaterials, including ultrathin Na1.08V3O8 nanosheets [42], NaV3O8 nanoflakes [43], and hydrous Na2V6O16
xH2O nanowires [44]. The Na1.08V3O8 nanosheets with about 10 nm thickness exhibited outstanding high rate property with a capacity of 94.2 mAh g-1 at 30°C and outstanding cycling stability with no capacity fading over 200 cycles [42]. The Na2V6O16
xH2O possess analogical crystal structure with the other two compounds, but have obvious expanded layer spacing, which is attributed to the crystal water in the interlayers. It is noted that the heating treatment caused the reduction of crystal water amount together with shrinkage of the layer spacing, but the heated sample displayed superior cycling stability even though with sacrificed capacity. The authors thought that the better stability is owing to the decreased lithium ions insertion and alleviated degradation degree [44]. Liang et al. explored another two Na1 + xV3O8 nanomaterials, which are Na1.25V3O8 [45] and Na1.1V3O7.9 nanobelts [46]. This work illustrated that the electrochemical

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property of the prepared samples was highly dependent on crystallinity and morphologies. The Na1.25V3O8 nanobelts displayed a considerable capacity of 225 mAh g-1 at 100 mA g-1 and well-cycling performance for 450 cycles with 94% capacity retention. SIBs have been the focus of study in recent years. Searching for satisfied cathode and anode materials is the aim for the development of SIBs. Owing to the larger interlayer spacing, which could accommodate more Na ions, Na1 + xV3O8 was also paid much attention in the field of sodium ion storage. Wang and co-workers applied annealed NaV3O8 nanowires as SIB cathode [47]. This kind of Na1 + xV3O8 exhibits the working voltage around 2.5 V and reversible specific capacity of 145.8 mAh g-1. The authors found a single-phase reaction existing in the storage process. Zhang’s group researched the sodium storage performance of single-crystalline Na1.1V3O7.9 nanobelts [48], which could exhibit a considerable capacity of about 173 mAh g-1 and favorable cycling performance. Dong et al. synthesized a new type of zigzag and hierarchical Na1.25V3O8 nanowires, which exhibited superior electrochemical performance in sodium ions storage [41]. Zigzag Na1.25V3O8 nanowires exhibit a considerable capacity of 172.5 mAh g-1 at 100 mA g-1 and superior cycling and rate property at high current density. The synergistic effect of the favorable structure features of Na1.25V3O8 and zigzag nanowire morphology bring outstanding sodium storage property. Compared with the simple nanowires, zigzag-shaped nanowires will effectively optimize the electrode/electrolyte contact area, restrain the selfaggregation of nanowires, release the strain during the cycling process, and then generate superior structure integrity and ion diffusion kinetic. Because of similar lithium insertion potential and superior structure features compared to LiV3O8, Na1 + xV3O8 is also a potential candidate when applied in aqueous LIB anodes. The Wang group designed an aqueous LIB based on Na2V6O16
0.14H2O nanowires as anode and LiMn2O4 as cathode materials with saturated Li2SO4 solution as electrolyte. This aqueous battery operated between 0 and 1.6 V and displayed an excellent electrochemical performance [49]. Nair et al. designed a symmetric aqueous rechargeable LIB based on Na1.16V3O8 nanobelts as both anode and cathode materials. The wide operating potential of Na1.16V3O8 in aqueous electrolyte is the reason why the symmetric aqueous battery can work well. The original capacity of 150 mAh g-1 was observed in this symmetric aqueous battery, and ~75% initial capacity remained after 100 cycles [50]. Besides aqueous LIB, aqueous SIB based on Na1 + xV3O8 also attracted much attention. The sodium ions insertion/extraction process of Na2V6O16
nH2O in the aqueous battery system was explored by Deng et al. Moreover, they also applied bundled Na2V6O16
nH2O nanobelts as anode and Na0.44MnO2 as cathode in full aqueous SIB. But fast capacity decay at initial cycles was found. From ex situ X-ray diffraction (XRD) results, Deng et al. thought that the capacity decrease was related to the irreversible transformation of Na2V6O16
nH2O with sodium intercalation [51]. Mentus and co-workers constructed Na1.2V3O8 nanobelts and explored the intercalation/deintercalation process of Li+, Na+, and Mg2+ in aqueous nitrates-based electrolyte. The capacities of 101, 55, and 67 mAh g-1 were achieved at 500 mA g-1 for Li, Na, and Mg intercalation, respectively, illustrating the utility insertion capability of the Na1.2V3O8 nanobelts [52].

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AxV2O5 (A = Li, Na, K)

The peculiar property of the V2O5 crystal structure for insertion process was recognized in the 1960s, when the vanadium bronzes were deeply researched. In particular, the LixV2O5 system on both structural chemistry and solid state with its various phases was also researched completely [53]. After that, a number of researches have been devoted to the AxV2O5 (A = Li, Na, K) system, which is thought to be a superior electronic conductor than most oxides, and its 2D lamellar structure realized an excellent ionic mobility and fast insertion of alkaline ions [54, 55]. Due to the above advantages, the vanadium bronzes are likely to be qualified candidates for positive electrode materials in secondary batteries. Related to the amount of A (x) existing in V2O5, numerous structural modifications can be synthesized [56]. Different insertion content of lithium will impact the phase of LixV2O5. When the lithium content is less than 0.01, LixV2O5 is α-phase. Then, with the increase in lithium content, LixV2O5 will change to β-LixV2O5 (0.22 < x < 0.37), ε-phase (0.35 < x < 0.7), δ-phase (0.7 < x < 1), and γ-phase (1 < x < 2). When lithium content is between 2 and 3, the ω-phase LixV2O5 will not reversibly change back to the ω-phase [5, 57–63]. The V2O5 layer structure is similar to the α-LixV2O5 and ε-LixV2O5 (Fig. 7.1a–c). The δ-phase LixV2O5 layers with higher bended degree (Fig. 7.1d) lead to more lithium ions insertion. In γ-phase LixV2O5 (Fig. 7.1e), the puckering V2O5 layers appear more distinct than δ-phase LixV2O5, where the VO5 tetrahedron connects with distortion. Galy et al. discovered the β-LixV2O5 (0.22 < x < 0.37) firstly [58], which existed in the monoclinic system of C2/m space group and consisted of the [VO6] octahedrons network. From Ravnsbæk and the co-workers’ report [64], it was found that

Fig. 7.1 Comparison of the arrangement of VOn polyhedral in (a) V2O5, (b) α-LixV2O5, (c) ε-LixV2O5, (d) δ-LixV2O5, and (e) γ-LixV2O5. (Reproduced from [56]. Copyright 1991, Elsevier)

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the disordered β-LixV2O5 only possesses a domain size of10 Å. AxV2O5 (A = Li, Na, K) were researched deeply as potential anode material for Lithium-ion batteries (LIBs), because these materials possess various chemical and physical properties and high theoretical capacities between 500 and 1000 mAh g-1 [56]. Goodilin et al. [65] reported LixV2O5 (x  0.8) nanobelts via a simple hydrothermal method; the obtained nanobelts were potential cathode material for LIBs with a considerable discharge capacity of 490 mAh g-1. In addition, the δ-type crystal structure of the as-prepared materials convert into the ε and γ mixture phases with the initial nanobelts morphology in the drying condition (200 °C under vacuum). It is found that full discharge with much Li ions intercalation leads to an irreversible structural change to ω-LixV2O5, which was exhibited in 1991 by Delmas et al. [60]. Ravnsbæk et al. [64] make use of operando powder X-ray diffraction and total scattering to research the atomic-scale deep-discharge ω-LixV2O5 phase (x  3). Their researchers found that the discharged ω-LixV2O5 composed of 60 Å rock salt structure on a  15 Å length scale cation ordering (Fig. 7.2). Zhao et al. [66] systematically researched AxV2O5 (A = Li, Na, K) as optimized cathode materials for lithium batteries (Fig. 7.3). They thought that the synergistic effect between the layers and the intercalated ions can relieve the restriction of rate property and cycling stability. In addition, this method optimizes the diffusion channel of guest ions, which is important and promising for further optimization and design of new kinds of intercalation compounds.

Fig. 7.2 (a) Overview plot of the operando PXRD data collected during discharge-charge cycling of an α-V2O5 cathode vs Li and (b) the simultaneously measured galvanostatic potential profile. (Reproduced from [64]. Copyright 2018, American Chemical Society)

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Fig. 7.3 Schematic representation and electrochemical properties of large alkali metal ion intercalation. (a) Schematic representation of large alkali metal ion intercalation. (b) Cycling performance of A-V-O nanowires formed by preintercalating large alkali metal ions into vanadium oxides with a charge/discharge rate of 0.1 A g-1. (c) Rate performance of A-V-O nanowires. A-V-O nanowires are cycled at various rates from 0.05 to 4.0 A g-1. Here n-C denotes the rate at which a full charge or discharge takes 1/n hours. (d) Cycling performance of A-V-O nanowires at the charge/discharge rate of 1.0 A g-1. (Reproduced from [66]. Copyright 2015, American Chemical Society)

Lu et al. [67] synthesized Na0.33V2O5-graphene nanosheet, which obtained 94.8% capacity retention and high capacity of 310 mA h g-1 at 150 mA g-1 after 50 cycles. Besides, the Na0.33V2O5-graphene holds a relative high capacity of about 199 mA h g-1 even at 4.5 A g-1. Liang et al. [68] synthesized potassium vanadate K0.25V2O5 by the sol-gel method, and as-synthesized K0.25V2O5 displayed the special layer-by-layer structure. K0.25V2O5 possessed considerable electrochemical performance at different current densities. In addition, K0.25V2O5 also displayed superior cycling ability. Zhu et al. [69] reported potassium vanadate for potassium ion batteries. In spite of the large radius of K+, the synthesized cathode exhibited a capacity of about 90 mAh g-1 at 10 mA g-1 between 1.5 and 3.8 V (vs. K+/K) and also exhibited good rate performance and cycling property. This work opens up new research direction of potassium vanadates in potassium ion batteries.

7.2

Alkali Metal Vanadates

7.2.4

185

LixVO2 and NaxVO2

There are six crystalline phases of polymorph VO2: tetragonal VO2(R) (P42/mnm), tetragonal VO2(A) (P42/nmc), VO2(C), VO2(D), monoclinic VO2(B) (C2/m), and monoclinic VO2(M) (P21/c), respectively, which construct by connecting distorted VO6 octahedrons in different ways [70, 71]. In particular, several edges share the octahedron of VO2(B) and form two identical layers along the c-axis in the unit lattice. The V-O tunnels stacked by corner sharing connections of the neighboring cells endow VO2(B) a faster ion transport rate and a higher capacity than other types of VO2 [72–74]. The Goodenough group [75] prepared delithiated LiVO2 by chemical and electrochemical methods and utilized powder X-ray diffraction techniques to examine their structural characteristics. In the cubic-close-packed oxygen lattice, the arrangement of vanadium ions accompanied with lithium extraction from the layered LiVO2 structure (space group = R-3 m) by electrochemical data and structural analyses of LixVO2 (x < 0.7) samples. About one-third of the vanadium ions in Li0.22VO2 are located in the vacants of octahedral position in the lithium layer, which gives it a stable structure close to the cubic lattice constants. Through linear sweep voltammetry, chemical insertion, constant current discharge, coulometric titration, and electromotive force measurements, Jacobsen et al. [76] have studied the lithium-ion insertion of LixVO2(B) between 25 °C and 120 °C. Limited by kinetic factors, the insertion x was less than 1. The Li/V insertion can be observed by electrochemical methods: At ambient temperature, 0.5 Li/V can be inserted at 2.55 V vs. Li in the two-phase region; at higher temperature, further insertion can be observed in another two-phase region at 2.1 V vs. Li; and 0.82 Li/V can be obtained at 120 °C. It was confirmed by X-rays that the two phases coexist in the interval of 0.13 < x < 0.45. The typical high-temperature solid-phase chemistry method can prepare only two phases for NaxVO2 system, P2-Na0.7VO2 and O3-NaVO2. Different oxygen packings of these two phases make them not converted into each other without breaking the V-O bond. Afterward, new metastable phases, Na1/2VO2 and Na2/3VO2, have been discovered by Delmas and co-workers [77], which were obtained from the parent O3-NaVO2 compounds by electrochemical deintercalation in a battery under room temperature (Fig. 7.4). Recently, the X-ray diffraction pattern of nanosized-layered rock-salt LiVO2 with a crystallite size about 47 nm, which prepared from peroxo-polyvanadic acid by Isao Tsuyumoto et al. [78], was matched well with the space group R-3 m in the Rietveld analysis. When used LiVO2 as cathodes for lithium-ion batteries (LIBs), nanosizedlayered rock-salt LiVO2 has higher specific capacity (149 mAh g-1 under 149 mA g-1) compared to conventional LiVO2 prepared by the solid-state reactions (below 50 mAh g-1). Rozier and co-workers [79] reported two NaxVO2 polymorphs with either Na located in octahedral or trigonal prismatic sites. Both of them were capable of reversible reactions with 0.5 Na+ per unit formula for sodium ion batteries. Meanwhile, more than 50% capacity concentrated near 1.6 to 1.7 V for the two NaxVO2 polymorphs. Moreover, the two polymorphs can preserve their initial phases of prismatic or octahedral sites during Na+ intercalation/deintercalation processes and show outstanding capacity retention, despite their rocky voltage profiles.

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Fig. 7.4 The electrochemical curve and ex-situ XRD patterns of O3 - NaVO2 during the deintercalation. (Reproduced from [77]. Copyright 2012, American Chemical Society)

Electrochemically activated transformation

P3

270 °C

irreversible reversible

P’3

125 °C

O’3

Thermally activated transformation irreversible

Room temperature

reversible

O’3

O’3

O’3 O3

P’3 1/2

P’3

O’3 with O’3 with stacking faults stacking faults 2/3

x in Nax VO2

1

Fig. 7.5 Phase diagram and schematic diagram of thermally and electrochemically driven transformations for NaxVO2 systems. (Reproduced from [80]. Copyright 2016, American Chemical Society)

Figure 7.5 shows a schematic diagram of the driven transformations of thermal and electrochemical in the sodium layer oxide NaxVO2. The blue, red, and green colors represent the new phases reported in the study, respectively. Because of the larger ion radius of Na+ compared to Li+, the sodium-layered oxides usually experience complex phase transition processes when alkaline ion (de)intercalated

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than lithium layered oxides. This work presents the orderly arrangement of sodium/ vacancy between the MO2 layers or the MO2 plate sliding. The less strength of Na-O bonds than Li-O bonds leads to easy slab gliding for sodium-layered oxides. However, thorough studies of the physical properties and band structure calculations of these NaxVO2 phases will be necessary to fully explain the changes in the voltage distribution during cycling. Such results provide further motivation to research the compounds based on V or Ti to find low-pressure sodium insertion compounds.

7.2.5

Li3VO4

Owing to high second-harmonic generation efficiency, low-temperature β-Li3VO4 single crystals found by Sakata et al. were studied in the material scientific field and applied as nonlinear optical materials [81, 82]. Earlier, high-quality Li3VO4 single crystals were generally grown by using the Czochralski method [83], floating-zone method [84], and heater-in-zone zone-melting method [85]. Characterizing the Li3VO4 crystals through X-ray diffraction, scientists disclose that Li3VO4 belongs to the space group Pmn21 and is an orthorhombic system with a = 5.4471 Å, b = 6.3272 Å, and c = 4.9483 Å. In the Li3VO4 crystals, two different crystallographic sites are occupied by Li+ (Li1 and Li2), and a third cationic site is occupied by V5+ ions. Due to the existance of many empty lattice sites in the intrinsic structure of Li3VO4 crystals, it can be predicted that Li3VO4 has good ionic conductivity. Later, researchers confirmed that the ionic conductivity of Li3VO4 can reach to approximately 10-4 S m-1, indicating the high diffusion activity of Li+ ions in its structure. As an outstanding ionic conductor, Li3VO4 was investigated as an active component of the solid electrolytes, such as Li4SiO4 - Li3VO4, Li4GeO4 - Li3VO4, Li2SO4 - Li3VO4, Li3VO4 - xLi3PO4, and Li3PO4 - Li3VO4 - Li4GeO4 [86– 92]. Meanwhile, Li3VO4, with high mobility of Li+, could be used as coating material to improve the cathode’s electrochemical performances, including LiMnPO4, LiNi0.6Co0.2Mn0.2O2, LiNi0.8Co0.1Mn0.1O2, and LiCoO2, Li1.18Co0.15Ni0.15Mn0.52O2 [93–97]. For example, Pu et al. improved the cyclability and rate capability of LiCoO2 by coating Li+-conductive Li3VO4 [93]. Since Sony Corporation released the first commercial LIB in 1991, LIBs have been widely applied in various fields ranging from portable electronic devices to large-scale energy storage system. The demand of high-energy-density LIBs is stronger with the development of modern society, especially when air pollution becomes more terrible. Generally, the energy density of LIBs is decided by two critical factors, the working voltage and the cell capacity. In terms of the anode materials, the working voltage should be rather low and the capacity should be high as much as possible to increase the energy density of full cells. Although with the extremely high theoretical capacities, conversion/alloy-reaction electrode materials, like transition metal oxides, Si-/Sn-based alloys and their composites, suffer from severe capacity fading caused by the huge volume expansion and low Coulombic efficiency owing to the irreversible side reaction during electrochemical cycling. In

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the intercalation-type anode materials (graphite, layered V-based oxides, and Ti-based oxides), the above defects have been alleviated in large degree. And graphite and Li4Ti5O12 as the typical insertion anodes have been successfully marketed during the past decades. But theoretical capacities of 372 mAh g-1 for graphite and 175 mAh g-1 for Li4Ti5O12 are limited, which makes them not relieve the anxiety of energy density for LIBs. Though the low lithiation potential of graphite is about 0.1 V, LIBs always have the security risks due to forming the lithium dendrite. Moreover, the high Li+ insertion potential of Li4Ti5O12 (1.5 V) severely weakens the energy density of the full cell when coupled with the cathode. Therefore, it is urgent to find the proper anodes to meet the next-generation highenergy-density LIBs. The insertion type anode material Li3VO4, first reported by Li et al., has aroused extensive attention in energy storage field and exhibited a new direction [98]. In the past 6 years, nearly 110 new papers about Li3VO4 have been added to the Web of Science online database. The reason for Li3VO4 being studied extensively is that it has higher theoretical intercalation capacity (394 mAh g-1, corresponding to Li5VO4) and more suitable insertion potential (mainly 0.5–1.0 V) simultaneously than graphite and Li4Ti5O12. Then Ni et al. proposed that the maximum theoretical capacity of LixVO4 can reach at 592 mAh g-1 (x = 3), and the existence of V2+ in the discharging production has been confirmed by the XPS results. Further, the authors constructed a highly competitive anode Li3VO4/C NC-on-NF with excellent initial reversible capacity (594 mAh g-1) for next-generation LIBs [99]. However, Li3VO4 still has some barriers in practical application, including the bad initial Coulombic efficiency and the terrible electronic conductivity ( 2 in Li3 + xVO4 will distort the structure of Li3VO4 and render the irreversible structure change, which will damage the cycling properties of Li3VO4 [143]. Therefore, more experimental results need to be explored to reveal the structural evolution mechanism of Li3VO4. Apart from serving as the LIB anode, Li3VO4 has also been investigated in many other energy storage systems, including lithium-ion capacitors (LICs) [138–141], magnesium ion batteries [142], and aluminum-ion batteries [143]. In these fields, Li3VO4 also exhibited a satisfactory property because of its unique crystal structure with fast ions diffusion channel. Taking LICs as an example, Shen et al. reported a peapod-like nanowire of Li3VO4/N-doped carbon with good energy density (136.4 Wh kg-1) and high power density (532 W kg-1) for LICs [138]. The research results well proved the potential application value of Li3VO4 in ion storage fields.

7.3

Alkali-Earth Metal Vanadates

191

Fig. 7.8 (a) Composing of Na2V6O16
1.63H2O nanowire cathode and Zn foil anode, as well as (b) its cycling performance. (Reproduced from [144]. Copyright 2018, American Chemical Society)

7.2.6

Other Alkali Metal Vanadates

In the energy storage fields, vanadium-based compounds are extensively investigated due to the layered structure and multi-electron reaction mechanism. Alkali metal vanadates with some unique strength are used as electrode materials in various devices, from LIBs to zinc ion batteries (ZIBs). ZIBs attracted wide concern originating from their high-safety and low-cost. Recently, Hu et al. synthesized a Na2V6O16
1.63H2O nanowire via hydrothermal method [144]. The nanowire cathode delivered a considerable capacity of 352 mAh g-1 at 50 mA g-1 in aqueous Zn (CF3SO3)2. Meanwhile, the aqueous ZIB system also exhibited excellent long lifespan (90% capacity retention at 5 A g-1 after 6000 cycles) (Fig. 7.8), which conformed that the Na2V6O16
1.63H2O is a valuable and meaningful cathode for ZIBs. Analogous to Na2V6O16
1.63H2O, Sambandam et al. reported a K2V6O16
2.7H2O nanorod intercalated cathode material for ZIBs [145]. The authors first used inexpensive electrolyte ZnSO4 and layered structural K2V6O16
2.7H2O cathode to assemble ZIB, which showed a capacity of 296 mAh g-1 after 100 cycles. And the specific energy can reach to 128 Wh kg-1, when the power density was 5760 W kg-1. The results make this kind of materials (alkali metal vanadates) to be a new attractive electrode for energy storage.

7.3

Alkali-Earth Metal Vanadates

Compared to alkaline metal elements at the same period, alkaline-earth metal elements are more abundant [146] and have a higher application potential. Alkaline-earth metal contained in vanadates possess some features that are distinguished from that of alkaline metal. First, elements of the second group have a large atomic radius than the first group. When inserted in vanadates lattice, the guest

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Vanadate Nanomaterials for Electrochemical Energy Storage

Fig. 7.9 Fraction of earth’s crust and other properties of elements. (Reproduced from [146]. Copyright 2017, American Chemical Society)

cations are able to form larger layer spacing, which provides smoother access for other metal ion to insert and detach. Second, the alkaline-earth metal elements have two valence electrons in the outermost shell; divalent ions tend to form stronger bonds than monovalent ions, due to which the crystal structure remains more stable. Based on these characters, the applications of alkali-earth metal vanadates nanomaterials for novel electrode materials mostly start from the following two starting points, the buffer of volume change by the nonelectrochemical activity alkaline-earth oxides and stabling the structure by guest cations. Besides, alkaliearth metal vanadates also possess abundant resources (compared to alkaline metal at the same group), higher conductivity (compared to some vanadium oxide), and other advantages (Fig. 7.9). For the former starting point, previous reports have shown that alkaline-earthmetal-containing compounds mostly exhibit excellent long-term cycling stability when used as metal-ion battery anodes, such as Ca-Fe-O [147], Ca-Co-O [148], CaSn-O [149]. Because the Ca ion can form nano-sized CaO, which is electrochemically inactive and buffers the electrodes volume expansion and inhibits the agglomeration of active materials, researchers call it “spectator effect” [147]. In this case, Xu et al. synthesized CaV4O9 and Sr-V-O nanowires [150], through the hydrothermal method and then thermal treatment, with ~100 nm in diameter and ~10 μm in length, and investigated the electrochemical properties for used as SIBs anodes. The heat treatment proceeded at two different temperatures, 450 °C and 550 °C. The sample prepared at 550 °C exhibits much stronger peaks than the material synthesized at 450 °C, and the morphology of nanowires at higher annealing temperature become wrinkle and irregular. For the nanowires prepared at 450 °C, numerous cavities were distributed homogeneously on the whole nanowires, which is owing to the removal of crystal water. The electric conductivity of CaV4O9 nanowires at 450 ° C exceeds 100 S cm-1, and the volume variation in the charge and discharge process is less than 10%, which also exhibits an average capacity of 363.9 mA h g-1 at 100 mA g-1, with an average voltage of about 1.0 V, and when measured at

7.3

Alkali-Earth Metal Vanadates

193

5000 mA g-1, the average specific capacity still remained 56.3% of that at 100 mA g-1. The CaV4O9 nanowires prepared at 550 °C remained 149.8 mA h g-1 at 1000 mA g-1 over 1600 cycles, which corresponds to a 102.4% capacity retention for the 12th discharge capacity. As for Sr-V-O nanowires, it displayed a better cycling stability up to 2000 cycles. The outstanding long-term cycling property of CaV4O9 is due to the uniformly distributed and in situ-generated CaO nanograins produced in the first discharge process. The CaO nanograins generated a self-preserving effect, which restrains the agglomeration of the electrochemistry active parts and preserves the high reversibility of the electrochemical reaction in the desodiation/sodiation process. Nanowire morphology is accessible to research on volume change and mechanism, but it lacks of practical application due to low tap density. Thereafter, Xu et al. synthesized CaV4O9 microflowers [151] with enhanced tap density of 0.74 g cm-3. The size of CaV4O9 miocroflowers is about 2–3 μm, it was synthesized by hydrothermal method and followed heat treatment. The microflower morphologies are assembled via plentiful ultrathin nanosheets. And a mass of mesopores is observed in the whole nanosheets; researchers speculated that the formation of mesopores is derived from the decomposition of glycerol in the reactant or its derivative during the calcining process. Caused by the excellent electrochemical performance of CaV4O9 and the increased tap density of microflowers, when used as LIBs anode, it has reached a high areal capacity of ~2.5 mAh cm-2 at 4.4 mg cm-2 and a stable cycling performance with the areal capacity of 1.5 mAh cm-2 after 400 cycles. It also exhibited an excellent electrochemical property at the high mass loading for Na storage, the initial areal capacity can reach up to ~1.0 mAh cm-2 with a high mass loading of 3.65 mg cm-2, then the reversible capacity retained 82% after 100 cycles. For the another starting point, compared with vanadium oxide, the guest cations deposited in the interlayer are anchored by strong ionic bonds bridging two adjacent layers and can act as pillars to stabilize the V-O interlayer with extended interlayer spacing. The presences of cations do not prevent entering metal ion from occupying available intercalation sites and can prevent destructive structural changes upon metal ion intercalation/de-intercalation. This strategy was mostly concentrated on alkali metal (Li, Na, K, etc.) or conductive metal (Ag, Cu, etc.), as for alkali-metal ions; it also can be expanded. The alkaline-earth metal ions have larger ion radius than alkali metal ions; appropriate ion radius can support more unimpeded layer spacing for other incoming metal ions. Alkali-metal ion offers another advantage of the divalent charge; the divalent cations bonded with oxygen atoms can lead to the formation of stronger M-O ionic bonds compared with monovalent cations, better pinning the layers together, which can prevent the structure from devastative collapse more efficiently, and stable long-term cyclability is highly expected to be achieved. In this case, Zhang et al. fabricated ultralong CaV6O16
3H2O nanoribbons [152] up to millimeter scale in length and 100 to 500 nm in width without any surfactants or templates; when used as LIB cathode materials, by the intercalated water molecules and Ca ion in the layered framework of vanadium oxide, it exhibited excellent high-rate kinetics and long-term cyclability. Liu et al. synthesized CaV6O16
3H2O [153] via microwave reaction and used it as an aqueous ZIBs cathode material. The as-prepared CaV6O16
3H2O was composed of uniform fibers.

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The length of fibers is more than 100 μm, and the diameter is about hundreds of nanometers. The sub-micrometric diameter is of benefit to shortening the Zn ion diffusion length. The calculated interlayer distance is 7.08 Å, larger than V2O5 (4.38 Å, PDF 41–1426), which facilitates diffusion of Zn ion. Moreover, the large interlayer distance in CaV6O16
3H2O contributes to high specific capacity. Liu et al. also carried out ex situ XRD measurement to confirm the reversibility of the Zn ion intercalation/deintercalation, which proved the high structural stability of CaV6O16
3H2O. Ma et al. used a facile hydrothermal method to fabricate δ-Ca0.24V2O5
H2O nanorods [154], which can form the tunnel β-Ca0.24V2O5 by a vacuum annealing process; the length of the nanorods is dozens of micrometers and the width ranging from around 100 to 500 nm. The morphology of β-Ca0.24V2O5 is wrinkled because of the dehydration treatment compared to δ-Ca0.24V2O5
H2O. When used as ZIB cathode materials, due to the phase transition at the 1st cycle for δ-Ca0.24V2O5
H2O, β-Ca0.24V2O5 nanorods showed better cycling performance, with an average decay of 0.035% per cycle at 500 mA g-1. Xia et al. synthesized freestanding Ca0.24V2O5
0.83H2O paper cathode [155] through a one-step hydrothermal method, which not only exhibits high capacity and stable cycle performance but also shows a very high energy density of 267 Wh kg-1 at 53.4 W kg-1 in Zn cell. In conclusion, alkali-earth metal vanadates nanomaterials applying for electrochemical energy storage mainly possess at least one following superiority compared to other vanadium-base nanomaterials, especially alkali metal vanadates. First, the abundance of alkali-earth metal and vanadium resources indicates the potential of them as industrialized electrode material. Second, due to the “spectator effect” of nano-sized alkaline-earth oxides, the structure change of electrode materials can be buffered. Third, the divalent cations can form strong bond to stable lattice and possess the appropriate interlayer spacing. In this case, alkali-earth metal vanadate nanomaterials have a bright future in electrochemical energy storage.

7.4 7.4.1

Transition Metal Vanadates Ag-V-O

The silver vanadium oxide has many allotropes, such as AgVO3, Ag2V4O11, Ag1.2V3O8, Ag0.33V2O5, etc. Most of these materials can exhibit high energy and power density and, thus, have been researched as cathode materials for lithium battery for a long time. In the 1980s, a kind of primary lithium battery based on Ag2V4O11 cathode was wildly used as power source for commercial implantable cardiac defibrillators (ICDs) [156]. The ICD is a device for electric shock treatment, whose mechanism is to conduct pulsing current to the heart to eliminate cardiac arrhythmia and restore sinus rhythm. So it has high demand for power source with excellent rate capability. A series of electrochemical tests show that this Ag2V4O11 lithium battery can not only work steadily in constant small current, but also can

7.4

Transition Metal Vanadates

195 molybdenum feedthrough pin (positive polarity) insulative glass to metal seal multiplate cell stack cathode lead bridge electrolyte fill hole and final close weld

stainless steel case and lid (negative polarity microporus polypropylene separator lithium anode silver vanadium oxide cathode

Fig. 7.10 Cutaway view of primary lithium battery for commercial ICDs. (Reproduced from [156]. Copyright 2001, Elsevier)

tolerate large pulsing current up to 2 A. It is a perfect match for the demand of ICDs. The cutaway view of this primary lithium battery is shown in Fig. 7.10. A metallic lithium anode and liquid organic electrolyte are utilized in this system, and the whole cell is sealed by stainless steel case. The multi-plate cell design can largely increase the effective surface area of cathode, which could be one reason for such rate capability. Another reason might be the reduction of Ag+ and formation of Ag0 during discharge, leading to higher conductivity [157]. Compared to Ag2V4O11, AgVO3 has a higher Ag: V ratio and is regarded as a more potential cathode material for primary lithium battery. In 2013, Zhou et al. reported a β-AgVO3 nanobelt with anchored Ag nanoparticles [158]. The existence of Ag nanoparticles can increase conductivity of material, leading to better electrochemical performance. According to authors, during the discharging process, the V5+ first turns into V4+, and then V3+, corresponding to the second and third plateaus of discharging curves (shown in Fig. 5.11). And the reduction of Ag+ to Ag0 exist in all plateaus. In 2016, David McNulty et al. also reported a β-AgVO3 nanowire as cathode for lithium-ion battery, and they proposed different analysis for the electrochemical reaction [159]. The CV curves are shown in Fig. 7.11b, in which we can

Potential (V vs Li/Li+)

a

7

3.5

Vanadate Nanomaterials for Electrochemical Energy Storage

20 mA g–1 50 mA g–1 100 mA g–1 500 mA g–1 1000 mA g–1

3.0

2.5

2.0

b 0.02 Current (mA)

196

0.00 –0.02

1st 2nd 5th 10th

–0.04

1.5 0

50

100 150 200 250 Specific capacity (mAh g–1)

300

1.5

2.0 2.5 3.0 Voltage (V)

3.5

Fig. 7.11 (a) Initial discharge curves of Ag/β-AgVO3 nanobelts at different current densities. (Reproduced from [158]. Copyright 2013, Royal Society of Chemistry). (b) Cyclic voltammograms for β-AgVO3 NWs, acquired at a scan rate of 0.1 mV s-1. (Reproduced from [159]. Copyright 2016, Royal Society of Chemistry)

observe three strong cathodic peaks at 2.85 V, 2.28 V, and 1.98 V in first scan. According to the authors, the strong peak at 2.85 V is correspond to reduction of V5+ and Ag+, and there is a weak peak at 2.74 V caused by the overlap of two reduction peaks of V5+ and Ag+. The cathodic peaks at 2.70 V and 1.98 V correspond to the reduction from V5+ to V4+ and from V4+ to V3+, respectively. The anodic peaks at 2.70 V and 3.08 V correspond to the oxidation from V3+ to V4+ and partly from V4+ to V5+, respectively. Starting from the second cycle, the strong peak at 2.85 V does not exist, demonstrating that there are irreversible reactions in first cycle. Though AgVO3 and Ag2V4O11 are good cathode materials for primary lithium battery, during discharge and charge cycling, there will be an irreversible phase change in first cycle, forming an amorphous phase and resulting to sudden decrease of capacity [160]. So they are not suitable for rechargeable lithium ion battery. As their allotropes, Ag0.33V2O5 and Ag1.2V3O8 can achieve reversible deintercalation/ intercalation of lithium ions, thanks to the layered structure [161–163]. For example, in 2013, Liang et al. reported a Ag1.2V3O8 nanobelt as cathode material for LIBs [162]. The crystal structure of Ag1 + xV3O8 is similar to Li1 + xV3O8 and both are stacked by (V3O8)n- layers. This layered structure can remain unbroken during the charge and discharge process. From the ex situ XRD results of cathode (shown in Fig. 7.12), the Ag1.2V3O8 turns into LiV3O8 and Ag after first cycle. After the fifth cycle, the LiV3O8 and Ag phase still exists and Ag1.2V3O8 disappears. So the real active material in electrochemical reaction after the first cycle might be LiV3O8, and the XRD results demonstrate good reversibility of lithium ion storage in LiV3O7. The lithium-ion battery assembled in this work exhibits capacities of 215, 198 and 174 mAh g-1 at 50, 100, and 200 mA g-1, respectively. The capacity is lower than AgVO3 primary lithium battery, but it can run for 50 cycles.

7.4

Transition Metal Vanadates

197

Fig. 7.12 XRD patterns of electrodes before cycling, after 1 and 5 cycles. (Reproduced from [162]. Copyright 2013, Elsevier)

7.4.2

Cu-V-O

Copper vanadium oxides (CVOs) or vanadates are lithium active through a multistep reduction by the valence change of the transition metal cations. As copper is relatively light with a high valence (2e- Cu2+/Cu0 redox process) among transition metal elements, CVOs can offer high theoretical gravimetric capacity and energy density. Also, another advantage of CVOs is its low cost with abundant storage in earth. The crystal structure of CVOs with only V5+ typically have three types, namely Cu2V2O7, CuV2O6 and Cu5V2O10. Though with the same valence state, the V ions in these compounds are in different sites of tetrahedral square and pyramid coordination. The vanadium coordination will change with the V5+/V4+ reduction process, for V4+ cations are unstable in a tetrahedral coordination environment. Therefore, Cu2V2O7 and Cu5V2O10 are converted to the amorphous phase during the discharge process [164]. CuV2O6 is a bit different and has edge-sharing VO5 square pyramid chains linked by Cu ions structure. During the discharge process, the exchange of Cu and Li will break the linkage mentioned above, leading to structure failure [165]. Considering these structural, CVOs with only V5+ generally exhibit rapid capacity fade. However, Chen and co-workers have successfully prepared α-CuV2O6 nanowires, microrods and mesowires via a hydrothermal reaction parameter through “ripening-splitting” mechanism [166]. This work suggests that the electrochemical performance of CuV2O6 electrode are related to the sizes and morphologies, and the CuV2O6 nanostructures have the promising application as alternative cathode for energy storage devices. The average plateau voltages of mesowires and nanowires are higher than that of bulk particles and microrods. This phenomenon can be prevented by size effect and reduced electrode polarization because of the contribution of excess surface free energy, so the discharge capacities of the CuV2O6

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7 Vanadate Nanomaterials for Electrochemical Energy Storage

electrodes are in the order of bulk particles < microrods < mesowires < nanowires. Also, the lithium intercalation activation energy of the nanowires is lower than that in bulk materials according to the calculations. Because V4+ cations could elevate the CVO electronic conductivity comparing to that which only has V5+. Those Cu vanadates with both V(V) and V(IV) have been studied as promising cathode materials [167]. A layered structure of Cu2.33V4O11 has been developed by Tarascon and the co-workers [164]. The structure is made up of [V4O11]n layers linked by the Cu cations with no obvious sites to insert Li+; however, this compound is electrochemically active with Li in a special way instead of the classic intercalation process. A new Cu vanadate phase, Cu1.1V4O11, could be produced by the removal of Cu from Cu2.33V4O11 [168]. This material will give ~260 mAh g-1 capacity provided by 5 Li through a displacement reaction. In the discharge process of Cu1.1V4O11, whole copper is not removed and the final structure is crystal, which is considered to be related to the strengthening of the Cu–O bond in Cu1.1V4O11. In addition, ε-Cu0.95V2O5 hollow microspheres fabricated by single-crystal nanoribbons are synthesized via a hydrothermal reaction between Cu(NO3)2 and V (IV)O(acac)2 in a solution containing the polyvinyl pyrrolidone [169]. The obtained ε-Cu0.95V2O5 superstructures exhibits a high capacity of 304 mAh g-1 and an outstanding rate capability. Although many studies have proved the potential applications of CVO as cathode material for energy storage devices, they suffer from copper dissolution, in which the Cu cations may dissolve out from the cathode material to the electrolyte and be lost, seriously limiting the cycling stability. And some strategies including particle coating, electrode configuration, and so no are implemented to reduce such Cu movement, thus increasing structural stability.

7.4.3

Co-V-O

Cobalt oxides, such as Co3O4 and CoO, have attracted great interest for the high theoretical capacity; the replacement of toxic and expensive Co can reduce the cost. V is used to coupling with Co, forming binary cobalt vanadates with good electrochemical property. Baudrin et al. were the first to synthesize a series of cobalt-based vanadates, among which the Co(VO3)2 shows the best performance as the anode material for LIBs that can react with 9.5 Li+ reversibly with just 17% irreversible capacity, which could deliver a high reversible specific capacity of 600 mA h g-1 after five cycles [170]. Wang et al. developed a Co3V2O8 multilayered nanosheets synthesized by a hydrothermal method, followed by the calcination process as a promising anode for LIBs. The reaction mechanism of Co3V2O8 is the reversible conversion reaction between Co and CoO take place on the amorphous LixV2O5 matrix. Also, the ultrathin Co3V2O8 nanosheets deliver a high capacity of 1114 mAh g-1 at 1 A g-1 over 100 cycles [171]. The Mai group reported hollow Co3V2O8 microspheres with

7.4

Transition Metal Vanadates

199

a high surface area as LIBs anode, which can keep a high reversible capacity of 320 mAh g-1 at a high current density of 20 A g-1. Also, a stable capacity of 424 mAh g-1 is measured at 10 A g-1 over 300 cycles [172]. Moreover, the replacement of toxic and expensive Co can reduce the cost and also make materials more environmental friendly [173]. Co2V2O7 can store up to 13.6 Li+ in the potential window of 0.02–3.5 V upon the first lithiation, which is a little less than 15.4 Li+ for Co3V2O8.[3] Therefore, even with one more Co atom, the capacity of Co3V2O8 is still higher than that of Co2V2O7. Wu and co-workers have fabricated the monodisperse mesoporous hexagonal Co2V2O7 nanoplatelets, which shows a high capacity of 866 mAh g-1 at 500 mA g-1 over 150 cycles [174]. The Mai group report a macroporous CoV2O6 nanosheets via using a one-pot approach by acetylene black induced heterogeneous growth as the anode material for LIBs. The CoV2O6 nanosheets present a high capacity of 702 mAh g-1 at 200 mA g-1 and a 89% capacity retention at 500 mA g-1 over 500 cycles, and excellent rate performance of 453 mAh g-1 at 5 A g-1 [175]. Also, cobalt vanadates such as CoV3O8, which has a crystallized tunnel structures, enables extraction and insertion of lithium ions and makes it an ideal candidate for LIB anode material. A CoV3O8 material prepared by a solid-state method exhibits a high reversible capacity of 120 mA h g-1 between the potential window of 1.5 and 4.5 V for 100 cycles [176]. It is believed that the small difference in unit cell volume between Li+ extraction and insertion state contribute to the cycling performance. It is also worth mentioning that the irreversible charge–discharge characters in the initial cycle are originated from a crystal material and an amorphous one induced by Li+ intercalation [177]. CoV2O5 has amorphous distinction synthesized by mild combustion of Co(NO3)2, V2O5, and glycine, which exhibits a high initial capacity of 275 mA h g-1 between the potential window of 2 to 4 V and, thus, is a promising candidate for cathode material [178]. For another case, the cobalt vanadate in spinel-type, LiCoVO4, while performed as anode material for LIBs, shows a high capacity and good rate capability, due to the formation of nanometer scale cobalt particles [179].

7.4.4

Fe-V-O

Iron is the most common metal in our life, and also the cheapest. This is one important reasons why lithium iron phosphate (LiFePO4) still is a hit in the lithium-ion battery market. The electrochemical performance of LiFePO4 is no match to lithium nickel cobalt manganese oxide (NCM, another popular commercial cathode material for lithium-ion battery), but LiFePO4 has much a lower price. This advantage of iron also makes iron vanadium oxide a promising material for battery. Vanadium oxide xerogel (V2O5
nH2O) has large lattice distance and can even adjust the distance between 7.8–13.8 Å by controlling the H2O content [180]. So it is not hard to intercalate metal cations. However, during intercalation and

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7 Vanadate Nanomaterials for Electrochemical Energy Storage

de-intercalation of cations, the change of lattice distance is too intense and might lead to crack of structure [181]. To stabilize the lattice, Wei et al. synthesized an iron pre-intercalated vanadium oxide ultrathin nanobelts (Fe-VOx) with constricted interlayer spacing [182]. The XRD patterns of VOx and Fe-VOx (Fig. 7.13a) show that the location of (001) peak in Fe-VOx is higher than in VOx, which means the lattice distance along c-axis has shrunk after intercalation of iron cations. It can be calculated that the distances in between (001) plane of VOx and Fe-VOx are 14.0 Å and 10.6 Å, respectively. The TGA results of Fe-VOx and VOx (Fig. 7.13b) reveal the same mass loss, which demonstrates that the intercalation of iron has no effect on H2O content. So the shrink of lattice distance is mainly caused by intercalation of iron. Then the two materials were tested in sodium ion battery. Figure 7.13c shows the electrochemical curves in 1.0–4.0 V at 0.1 A g-1. Both VOx and Fe-VOx have a short discharging plateau at 1.70 V. And the CV curves (Fig. 7.13f) also show obvious reduction peaks at ~1.70 V and oxidation peaks at ~2.20 V, perfectly matched with charge and discharge curves. Figure 7.13d shows that Fe-VOx has higher capacity at 0.1 A g-1, which demonstrates that the intercalation of iron cations has not occupied the active site of Na+ and decrease capacity. It is clear that Fe-VOx has better cycle stability. Fe-VOx can maintain 80.0% capacity after 50 cycles, while VOx can only maintain 62.1% capacity. What’s more, Fe-VOx also exhibits a much better rate performance than VOx in Fig. 7.13e. All these tests demonstrate that iron preintercalation does have optimized electrochemical performance of vanadium oxide xerogel. In 2018, Peng et al. reported an iron vanadium oxide Fe5V15O39(OH)9
9H2O nanosheet as aqueous ZIBs cathode material [183]. The synthesis process of Fe5V15O39(OH)9
9H2O nanosheet is simple and quick, with cheap ingredients. According to the authors, it was synthesized by water bath in atmospheric pressure at 90 °C for only 1 h. This material has a high reversible specific capacity of 385 mAh g-1 at 0.1 A g-1 and shows a superior cycling stability even at 5 A g-1. In Fig. 7.14, the XPS results of electrodes under three different conditions (original, charged and discharged) were analyzed. A relatively weaker Zn 2p intensity of charged electrode than that under discharging conditions indicates that the Zn2+ has successfully inserted into the cathode material [184]. As was mentioned in the last paragraph, the V and Fe elements were reduced from the results of fitted curves of V 2p3/2 and Fe 2p3/2 in Fig. 7.14b, c. In detail, V5+ was partially reduced to V4+ and even V3+ from the original state to the discharged state [185]. Fe3+ was completely transferred into the Fe2+ signal. The valence states of Fe and V are reversed into the original states after the following charging process. The results indicate that both V and Fe participate in the electrochemical reaction and together contribute to the high capacity.

7.4

Transition Metal Vanadates

201

Fig. 7.13 Characterizations and electrochemical tests of VOx and Fe-VOx in sodium ion battery: (a) XRD patterns, (b) TG-DSC, (c) charge-discharge curves, (d) cycling performance at a current density of 0.1 A g-1, (e) rate performance and (f) CV curves. (Reproduced from [182]. Copyright 2015, American Chemical Society)

7.4.5

Zn-V-O

Many reported zinc vanadium oxides are composed of vanadium oxide layers, with embedded zinc cations. So they often have good compatibility with zinc ion battery. In 2016, Nazar et al. developed a layered Zn0.25V2O5
nH2O nanobelt as aqueous Zn cathode material [185]. As shown in Fig. 7.15, the structure of Zn0.25V2O5
nH2O is composed of two-dimensional double-sheet structure of V2O5 framework with, similar to the δ-V2O5 layers. The ZnO6 octahedra located in the interspace between

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Vanadate Nanomaterials for Electrochemical Energy Storage

Fig. 7.14 XPS results of the electrodes at different states: (a) Zn 2p region, (b) fitted curves of V 2p3/ in which V3+ is magenta, V4+ is blue, and V5+ is red; (c) fitted curves of Fe 2p3/2, in which Fe2+ is blue and Fe3+ is red. (Reproduced from [183]. Copyright 2018, Royal Society of Chemistry)

2,

Fig. 7.15 (a) Schematic of the Zn metal/Zn0.25V2O5 cell on discharge in aqueous 1 M ZnSO4. (b) Long-term stability test (galvanostatic cycling) at a current density of 1200 mA g-1. (c) Scheme showing reversible water intercalation into Zn0.25V2O5
nH2O immersed in electrolyte/H2O and the water deintercalation accompanying Zn2+ intercalation upon electrochemical discharge. The red and blue spheres represent O and H, respectively; the H2O molecules interact with the oxygen layers through hydrogen bonding. Here y > z > n, because a fraction of intercalated H2O remains in the discharged material. (Reproduced from [185]. Copyright 2016, Nature Publishing Group)

7.4

Transition Metal Vanadates

203

V2O5 layers act as pillars. These pillars fix the V2O5 layers in a framework and stabilize the paths for inserting additional Zn ion during discharge process. The Zn/Zn0.25V2O5 cell exhibit a high capacity (250–300 mAh g-1) and a long cycle life: over 80% of the capacity retention at 2400 mA g-1 over 1000 cycles. During Zn2+ intercalation/deintercalation, the indigenous Zn2+ and crystal water act as pillars and contribute to the stabilization of layered structure and mobilize ions diffusion. The authors also proposed that when the electrode is immersed in electrolyte, water can be reversibly inserted into Zn0.25V2O5
nH2O, and extracted, accompanied by Zn2+ insertion, during the discharge process. The reaction is provided as the following equation (Fig. 7.15c): Zn0:25 V2 O5
nH2 O þ 1:1Zn2þ þ 2:2e - $ Zn1:35 V2 O5
nH2 O

ð7:1Þ

In 2017, Xia et al. successfully synthesized an allotrope Zn3V2O7(OH)2
2H2O [186]. Its structure is constructed by brucite type layers, separated by V2O7 polyhedra (Fig. 7.16a). In the close-packed O layer, Zn atoms occupy three or four

Fig. 7.16 (a) Crystal structure viewed along the b-axis of Zn3V2O7(OH)2
2H2O, which shows a layered structure with porous framework. The Zn atoms in ZnO6 and V atoms in V2O7 polyhedra are depicted in blue and yellow, respectively. The gray atoms in the crystal cavities represent the lattice water. (b) Galvanostatic charge–discharge profiles at a current density of 50 mA g-1. (c) Cycle performance at a current rate of 200 mA g-1. (d) Rate capability. (Reproduced from [186]. Copyright 2018, Wiley-VCH)

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Vanadate Nanomaterials for Electrochemical Energy Storage

Fig. 7.17 Crystal structure of ZnV2O6. (Reproduced from [187]. Copyright 2012, Royal Society of Chemistry)

octahedral sites. The large interspace among layers is full with water molecules. The Zn3V2O7(OH)2
2H2O cathode exhibits the high capacities of 76 and 213 mAh g-1 at 3000 and 50 mA g-1, respectively. A high reversible capacity of 101 mAh g-1 with 68% capacity retention can be obtained at 200 mA g-1 over 300 cycles. The electrochemical reaction in cathode can be described by the following equation: Zn3 V2 O7 ðOHÞ2
2H2 O þ 1:9Zn2þ þ 3:8e - $ Zn4:9 V2 O7 ðOHÞ2
2H2 O

ð7:2Þ

In addition to zinc ion battery, some zinc vanadium oxides also have been reported as electrode material for LIBs. In 2012, Sun et al. synthesized a monoclinic ZnV2O6 nanowire via hydrothermal method [187]. The schematic diagram of the crystal structure is shown in Fig. 7.17. The double-layered vanadium oxide layers are stacked along b-axis, and each layer is composed of [VO4] tetrahedral by corner sharing. Meanwhile, the [ZnO6] octahedral coordinations also array along b-axis. This material can work as anode in LIB. Over a voltage window between 0.025–3.0 V, it achieves a high initial capacity of 1555.5 mAh g-1 at 100 mA g-1 and maintains 972.9 mAh g-1 over 10 cycles.

7.4.6

Other Transition Metal Vanadates

Except for those main metal vanadates have been mentioned, there are other types of transition metal vanadates that also have been studied in recent years. MnxVyOz is also potential LIBs anode material, mainly referring to MnV2O6 and Mn2V2O7. Although the application of MnV2O6 as the LIBs anode was early reported in 2001, there also obscure for the Li+ ion storage mechanism until now. Kim et al. studied the physical and electrochemical properties of MnV2O6 with its phase transformation via the nuclear magnetic resonance (NMR) spectrum and ex situ XRD [188]. He also investigated the electrochemical performance of MnV2O6 as anode for LIB without clearly explaining the charge/discharge mechanism [189]. In order to bring up its electrochemical performance, the authors have successfully synthesized lithiated MnV2O6 with Li5MnV2O6 composition and evaluated the martial as anode in a full-cell system against Mn0.2V2O5
0.9H2O cathode. An initial capacity of

7.4

Transition Metal Vanadates

205

100

1200

80

800

60 40

400 20 0

Current Density = 500 mA g–1 0

100 200 Cycle Index

300

Coloumbic Efficiency / %

Specific capacity / mAh g–1

a

Specific Capacity / mAh g–1

Fig. 7.18 Polyhedral view of showing (a) the unit cell down the b-axis; edge-sharing VO5 square pyramidal chains along the b-axis, bridged through corner oxygen atoms by (b) disordered MnO4(H2O)2+/2NH4+and (c) MnO6 octahedral chains. (Reproduced from [191]. Copyright 2016, Royal Society of Chemistry)

b 1500 1000 1 mg cm–2 500 0

2 mg cm–2 100 1000 Current Density / mA g–1

Fig. 7.19 (a) Capacity for 300 cycles cycled at 0.5 A g-1 along with the corresponding CEs with an aerial loading = 2 mg cm-2; (b) capacities as a function of current density with varied areal loading. (Reproduced from [191]. Copyright 2016, Royal Society of Chemistry)

300 Ah kg-1 was measured, and it shows a promising potential for commercial anode martial [190]. At the time, the introduction of coordinating water and ammonia ions forms special structure for Mn1.5(H2O)(NH4)V4O12, which can prove as a method to combat poor electronic conductivity and limited cycle life with manganese vanadates (Fig. 7.18). This structure allows for the initial capacity of 1600 mAh g-1 over hundreds of cycles with stable capacities of 500 mAh g-1 at 500 mA g-1. They also investigated in situ a rare disorder of Mn(H2O)22+/2NH4+ during lithium intercalation (Fig. 7.19) [191]. Recently, β-Mn2V2O7 haa also been prepared by Sambandam et al. by a facile green method. It was found that the structure of Mn2V2O7 was completely destroyed in the initial lithiation process and

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γ-LixV2O5 was retained in the whole electrochemical reaction processes, indicating the stable structure when Li+ extraction/insertion [192]. Nickel vanadium oxide, like NiV3O8, as a new sort of anode, has also been studied in recent years. Ni and co-workers have constructed a NiV3O8 nano-flakes grown on Ni foam composite material through in situ growth method and tested its electrochemical performances as anode material for LIBs [192]. Except for this work, the electrochemical performances of nickel vanadates with different morphologies have been studied, such as NiV3O8 amorphous wire encapsulated in crystalline tube nanostructure [193] and Ni3V2O8 nano-rods anchored onto ordered mesoporous carbon (CMK-3) [194] and its volume expansion relieved to some extent. A recent study by the Lou group suggests that multi-components and intricate architectures may be an essential element for constructing high performance electrode for LIBs. They successfully synthesized the yolk-shelled NiCo2V2O8 spheres from Ni-Co glycerate spheres and have an extra-high specific capacity of 1228 mAh g-1 at 1.0 A g-1 over 500 cycles (Fig. 7.20) [195]. MoV2O8 becomes a promising electrode material for its high valence states of Mo and V, electronegativity as well as high interlayer spacing of 4.20 Å. During lithium insertion, the material goes through six oxidation states and is able to show a high capacity, and the unoccupied nonbonding orbitals of Mo and V atoms would help the material with volume change problem [196]. MoV2O8 material also displays improved ionic conductivity after the alkaline ion insertion [197]. The Guo Group

Fig. 7.20 (a) Formation of N with enhanced lithium storage properties. FESEM images, TEM images, and XRD pattern of the NiCo2V2O8 yolk–double shell spheres. (b) Electrochemical lithium storage properties of the NiCo2V2O8 and Co3V2O8 samples. (Reproduced from [195]. Copyright 2018, Wiley-VCH)

Transition Metal Vanadates

207

10

15

20

Z

(110)

(001)

Intensity (a.u.)

7.4

25

X 30

35

40

Y

O

Mo

V

2T/degree

Hydrothermal method 3 hours VO2 6 hours

VO2

Energy density (Wh kg-1)

Fig. 7.21 In situ XRD and crystal structure of MoV2O8 nanosheets. (Reproduced from [198]. Copyright 2019, Elsevier) 40

WN2O7-G W-doped V2O5 PANI/WP

30 20

10

10

100

1k

10k

Power density (W kg-1)

100k

overnight WV2O7 VO2 graphene 100 nm

100 nm

Fig. 7.22 Synthesis mechanism ragone plot and TEM images of WV2O7 nanosheets are prepared on graphene NSs. (Reproduced from [198]. Copyright 2018, Elsevier)

designed porous MoV2O8 nanosheets and MoV2O8 nanoparticles fabricated through a solvothermal treatment method by tuning the amount of acetic acid and compared their electrochemical performance by in situ X-ray diffraction. The results show that the porous MoV2O8 nanosheets exhibit excellent capacity, prolonged cyclability, and good rate performance than nanoparticles (Fig. 7.21). Moreover, a full cell was assembled using the porous MoV2O8 nanosheets, which presented desirable electrochemical properties, showing the potential application of the MoV2O8 nanosheets as a commercial anode material [198]. Due to their multiple oxidation states, special tunnel structure for fast ion insertion low cost, tungsten vanadates have been seen as a potential material for electrodes. By the cumulative, self-assembled tungsten vanadate (WV2O7) nanosheets are fabricated on graphene nanosheets (Fig. 7.22). It shows a high

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specific capacitance, which was 346.4 F g-1, and high differential capacitance of 1211.4 F g-1 at -0.1 V when applied into a supercapacitor electrode using H2SO4 electrolyte. Moreover, it is observed that the charge retention during cyclic voltammetry is improved by a great deal (68% at scan rate 100 mV s-1) even with a high electrode loading (5 mg cm-2), which was practically important because the commercial electrode fabrication was based on slurry mixing process [199].

7.5

Summary and Future Directions

Due to the various kinds of interlayer metal ions, V-O polyhedrons distortion, and the wide variation of vanadium element, there are numerous members in the vanadates family. In this chapter, different vanadate materials displayed the superior structure stability and considerable electrochemical properties in electrochemical energy storage. In consideration of multiple optimizing effects, such as charge buffering effect, pillar effect, and synergistic effect, vanadate materials often exhibited favorable electrochemical performances. In addition, with the arising of various metal-ion batteries in recent years, the studies for the advanced electrode materials, which is compatible to the new-type battery systems, plays a considerable role in the future development. The diversity of vanadate materials provides extensive development potential and new development direction for the emerging battery technology. We firmly believe that with the intensive study on the different group of vanadates, more and more electrode materials with excellent properties and significant results will be identified.

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