Silicon-Based Anode Materials for Lithium-Ion Batteries

Silicon-based anode materials for lithium-ion batteries

Binghui Xu Bachelor of Engineering

A thesis submitted for the degree of Doctor of Philosophy at The University of Queensland in 2015 School of Chemical Engineering

Abstract

While lithium-ion batteries (LIBs) have found wide applications in customer electronics, such as cellular phones and lap-top computers, the performance of the currently LIBs does not meet the market requirements for massive energy storage, such as electric vehicles and smart grids. One of the critical issues is the low energy capacity of the electrode materials. Graphite has been a common anode for LIBs, but the maximum capacity (372 mAh·g-1) of graphite has hindered its application in advanced LIBs. Substituting graphite with materials of high capacity has become an active research area. Graphene can store more Li+ than graphite because Li+ can not only be stored on both sides of graphene sheets, but also on the edges and covalent sites. Being electric conducting and mechanically strong, graphene is also an attractive support for other high capacity materials, such as silicon (Si). Si possesses the highest known theoretical capacity (4200 mAh·g-1, fully lithiated). However, Si suffers from a critical problem, namely severe volume change during charging/discharging, resulting in poor reversibility and rapid capacity decay. This PhD project aims to improve the stability and suitability of Si nanoparticles (SiNPs) for LIBs.

The first aspect in this thesis is to demonstrate a novel approach to wrapping SiNPs in a reduced graphene oxide (RGO) aerogel framework. The aerogel-typed RGO architecture not only provided a porous network to accommodate the volume change of entrapping SiNPs, but also facilitated electrolyte transport and electron transfer.

The second aspect in this thesis is to report a simple, green and scalable method to prepare RGO using gallic acid (GA) as a chemical reducing reagent. RGO samples reduced by GA showed a superior Li+ storage capacity compared with graphite and other reported graphene materials as anode for LIBs.

The third aspect in this thesis is to report on the preparation of RGO-stabilized SiNPs, a sandwich-typed composite material (Si@RGO). As a result, good battery performance of Si@RGO was obtained, 1074 mAh·g-1 at the 5th cycle and 900 mAh·g-1 at the 100th cycle with a capacity retention of about 84% as well as excellent rate capability.

Declaration by author

This thesis is composed of my original work, and contains no material previously published or written by another person except where due reference has been made in the text. I have clearly stated the contribution by others to jointly-authored works that I have included in my thesis.

I have clearly stated the contribution of others to my thesis as a whole, including statistical assistance, survey design, data analysis, significant technical procedures, professional editorial advice, and any other original research work used or reported in my thesis. The content of my thesis is the result of work I have carried out since the commencement of my research higher degree candidature and does not include a substantial part of work that has been submitted to qualify for the award of any other degree or diploma in any university or other tertiary institution. I have clearly stated which parts of my thesis, if any, have been submitted to qualify for another award

I acknowledge that an electronic copy of my thesis must be lodged with the University Library and, subject to the policy and procedures of The University of Queensland, the thesis be made available for research and study in accordance with the Copyright Act 1968 unless a period of embargo has been approved by the Dean of the Graduate School.

I acknowledge that copyright of all material contained in my thesis resides with the copyright holder(s) of that material. Where appropriate I have obtained copyright permission from the copyright holder to reproduce material in this thesis.

Publications during candidature

Journal Papers:

1. Graphene-based electrodes for electrochemical energy storage, Chaohe Xu, † Binghui Xu, † Yi Gu, † Zhigang Xiong, Jing Sun and X. S. Zhao, Energy and Environmental Science, 2013, 6, 1388-1414. († equally contributed first co-authors). 2. Stabilization of Silicon Nanoparticles in Graphene Aerogel Framework for Lithium Ion Storage, Binghui Xu, Hao Wu, C. X. (Cynthia) Lin, Bo Wang, Zhi Zhang, X. S. Zhao, RSC Advances, 2015, 5, 30624-30630. 3. Lithium-storage Properties of Reduced Graphene Oxide and its Silicon Based Composites, Binghui Xu, Jintao Zhang, Yi Gu, Zhi Zhang, Wael Al Abdulla, Nanjundan Ashok Kumar and X. S. Zhao*, Submitted to Scientific Reports. 2+ 4. Mesoporous carbon-coated LiFePO4 nanocrystals co-modified with graphene and Mg doping as superior cathode materials for lithium ion batteries, Bo Wang, Binghui Xu, Tiefeng Liu, Peng Liu, Chenfeng Guo, Shuo Wang, Qiuming Wang, Zhigang Xiong, Dianlong Wang and X. S. Zhao, Nanoscale, 2014, 6, 986-995.

5. Growth of LiFePO4 nanoplatelets with orientated (010) facets on graphene for fast lithium storage, Bo Wang, Shuo Wang, Peng Liu, Jie Deng, Binghui Xu, Tiefeng Liu, Dianlong Wang, X. S. Zhao, Materials Letters, 2014, 118, 137-141.

6. The synergy effect on Li storage of LiFePO4 with activated carbon modifications, Bo Wang, Qiuming Wang, Binghui Xu, Tiefeng Liu, Dianlong Wang and George Zhao, RSC Advances, 2013, 3, 20024-20033.

Conference Proceedings:

1. Self-assembly of Graphene Silicon Nanoparticles into 3D Hierarchical Aerogel Nanocomposite for Lithium Ion Storage, Binghui Xu, Hao Wu, Zhigang Xiong, C. X. (Cynthia) Lin, Bo Wang, Peng Liu, and X. S. Zhao. Poster presentation at Asia-Pacific Conference On Electrochemical Energy Storage & Conversion (APEnergy 2014), 5–8 February 2014, Brisbane Convention & Exhibition Centre, Australia.

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Publications included in this thesis

1. “Chaohe Xu†, Binghui Xu†, Yi Gu†, Zhigang Xiong, Jing Sun and X. S. Zhao*, Graphene-based Electrodes for Electrochemical Energy Storage, Energy and Environmental Science, 2013, 6, 1388-1414.” – incorporated as Chapter 2.

Contributor Statement of contribution Co-first author Author Binghui Xu (Candidate) Wrote and edited paper (25%) Co-first author Author Dr Chaohe Xu Wrote and edited paper (25%) Co-first author Author Yi Gu Wrote and edited paper (25%) Author Dr Zhigang Xiong Wrote and edited paper (5%)

Author Prof. Jing Sun Wrote and edited paper (5%)

Author Prof. X. S. Zhao Wrote and edited paper (15%)

2. “Binghui Xu, Hao Wu, C. X. (Cynthia) Lin, Bo Wang, Zhi Zhang, X. S. Zhao*, Stabilization of Silicon Nanoparticles in Graphene Aerogel Framework for Lithium Ion Storage, RSC Advances, 2015, 5, 30624-30630.” – incorporated as Chapter 4.

Contributor Statement of contribution Experimental design and performance (80%) Author Binghui Xu (Candidate) Wrote and edited paper (85%) Author Dr Hao Wu Experimental design and performance (20%)

Author Dr C. X. (Cynthia) Lin Helped to run freeze drying process

Author Bo Wang Helped to take EDS images

Author Zhi Zhang Helped to take HR-TEM images

Author Prof. X. S. Zhao Wrote and edited paper (15%)

3. “Binghui Xu, Jintao Zhang, Yi Gu, Zhi Zhang, Wael Al Abdulla, Nanjundan Ashok Kumar and X. S. Zhao*, Lithium-storage Properties of Reduced Graphene Oxide and its Silicon Based Composites, Submitted to Scientific Reports. – incorporated as Chapter 5 and Chapter 6.

Contributor Statement of contribution Experimental design and performance (100%) Author Binghui Xu (Candidate) Wrote and edited paper (80%) Author Prof. Jintao Zhang Wrote and edited paper (5%)

Author Yi Gu Helped to take FE-SEM images

Author Zhi Zhang Helped to take HR-TEM images

Author Dr Wael Al Abdulla Helped to record Raman spectroscopy

Author Dr Nanjundan Ashok Kumar Wrote and edited paper (5%)

Author Prof. X. S. Zhao Wrote and edited paper (15%)

Contributions by others to the thesis

No contributions by others.

Statement of parts of the thesis submitted to qualify for the award of another degree

None

Acknowledgements

First of all, I would like to thank my parents for their sensible suggestions, directions and unselfish support in my study career.

Then I would like to express my sincere gratitude to my respectable supervisors, Prof. X. S (George) Zhao and Prof. Lianzhou Wang for giving me guidance, encouragement and support throughout this work.

Many thanks go to Prof. Joe Diniz da Costa and Prof. Chengzhong (Michael) Yu for being in my thesis committee, and for their constructive suggestions throughout my Ph. D candidature career.

I am also grateful for the staffs in School of Chemical Engineering who helped to solve the financial as well as administrative issues. This project would not have been possible without the kind assistance from my dear colleagues in Prof. George Zhao’s research group. The facilities and technical assistance of the Australian Microscopy and Microanalysis Research Facility at the Centre for Microscopy and Microanalysis at the University of Queensland are acknowledged.

I will take this opportunity to express my appreciation to the China Scholarship Council and the Australian Research Council for financial support over the PhD study.

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Keywords

Lithium ion batteries, anode, stabilization, silicon nanoparticles, graphene, aerogel, composite

Australian and New Zealand Standard Research Classifications (ANZSRC)

ANZSRC code: 091205, Functional Materials, 60%

ANZSRC code: 090406, Powder and Particle Technology, 20%

ANZSRC code: 090403, Chemical Engineering Design, 20%

Fields of Research (FoR) Classification

FoR code: 0912, Materials Engineering, 50%

FoR code: 0904, Chemical Engineering, 50%

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Table of Contents

Abstract ...... I

Acknowledgements ...... VII

Keywords ...... VIII

List of figures ...... XII

List of schemes...... XVI

List of tables ...... XVI

List of abbreviations...... XVII

Chapter 1 Introduction ...... 1

1.1 Background ...... 1

1.2 Development history of battery ...... 2

1.3 Working principles and terms for LIBs ...... 3

1.4 Significance and research objective of the project ...... 6

1.5 Thesis outline ...... 7

Chapter 2 Literature review ...... 9

2.1 Cathode ...... 10 2.1.1 Layered structured cathode ...... 10 2.1.2 Spinel structured cathode ...... 12 2.1.3 Olivine structured cathode ...... 12 2.1.4 Strategies to enhance the performance of cathode ...... 12

2.2 Anode ...... 13 2.2.1 Metallic Li anode ...... 13 2.2.2 Carbonaceous anode ...... 14

2.2.3 Sn and SnO2 based anode...... 14 2.2.4 Lithium titanate anode ...... 15 2.2.5 Other metal oxide anode ...... 15

2.3 Graphene-based anode ...... 16 2.3.1 Preparation of graphene ...... 16 2.3.2 Graphene-based electrode for LIBs ...... 24

2.4 Silicon-based anode ...... 29 2.4.1 Carbon coated Si electrode ...... 30 2.4.2 Metal oxide coated Si electrode ...... 32 2.4.3 Polymer modified Si electrode ...... 32 2.4.4 Si electrode with designed morphology ...... 34

2.5 Si/Graphene composite anode ...... 38

2.6 Separator ...... 45

2.7 Electrolyte ...... 46

2.8 Summary and perspectives ...... 47

Chapter 3 Research methodology ...... 48

3.1 Preparation of GO ...... 48

3.2 Anode materials preparation ...... 49

3.3 Fabrication of coin cell batteries ...... 49

3.4 Electrochemical measurements ...... 50

3.5 Materials Characterization ...... 52

Chapter 4 Stabilization of Silicon Nanoparticles in Graphene Aerogel Framework for Lithium Ion Storage ...... 54

4.1. Introduction ...... 55

4.2 Experimental ...... 57

4.3 Results and discussion ...... 59

4.4 Conclusions ...... 70

Chapter 5 Lithium-storage Properties of Reduced Graphene Oxide Using Gallic Acid as Reducing Agent ...... 71

5.1 Introduction ...... 72

5.2 Experimental ...... 73

5.3 Results and Discussion ...... 74

5.4 Conclusions ...... 82

Chapter 6 Silicon Nanoparticles Encapsulated in Reduced Graphene Oxide Framework as a Stable Anode for Lithium Ion Storage ...... 83

6.1 Introduction ...... 84

6.2 Experimental ...... 85

6.3 Results and Discussion ...... 86

6.4 Conclusions ...... 95

Chapter 7 Conclusions and Recommendations for Future Work ...... 96

7.1 Conclusions ...... 97

7.2 Recommendations for future work ...... 98

References ...... 99

List of figures

Figure 2.1 (A) Intramolecular and (B) intermolecular dehydration of GO at hydrothermal conditions. Procedure (B) gives rise to aggregated products.

Figure 2.2 Illustration of the preparation of graphene based on Fe reduction.

Figure 2.3 Schematic diagrams of the possible growth mechanism for graphene by CVD methods. (A) Carbon segregation/precipitation process; (B) Surface growth mechanism.

Figure 2.4 Illustration of the bubbling transfer process of graphene from a Pt substrate.

Figure 2.5 Illustration of edge-carboxylation of graphite for graphene in the presence of dry ice. (A)

Pristine graphite; (B) dry ice (solid phase CO2); (C) ECG prepared by ball milling; (D) schematic representation of physical cracking and edge-carboxylation of graphite by ball milling in the presence of dry ice.

Figure 2.6 Bottom-up fabrication of atomically precise graphene nanoribbons.

Figuer 2.7 (A) Cyclic voltammograms of the FGNs electrode, (B) Initial discharge/charge profiles of the FGNs and the RGNs at a current density of 100 mA·g-1, (C) The cycling stabilities of the FGNs and RGNs at a high current density of 1000 mA·g-1, (D) The rate capabilities and cycle performances of the FGNs and the RGNs at various current densities, Nyquist plots of the FGNs and RGNs (E) before cycling and (F) after 3 cycles.

Figure 2.8 Schematic illustration of the electrospinning and carbonization steps. (A) The electrospinning process using a dual nozzle. The PMMA solutions containing SiNPs and the PAN solution were injected into the core and shell channels of the nozzle, respectively. (B) After electrospinning, stabilization, and carbonization steps were employed to complete the core–shell 1D fibers, respectively.

Figure 2.9 (A) Schematic illustration of 3D porous SiNP/conductive polymer hydrogel composite electrodes. Each SiNP is encapsulated within a conductive polymer surface coating and is further connected to the highly porous hydrogel framework. SiNPs has been conformally coated with a polymer layer either through interactions between surface -OH groups and the phosphonic acids in the crosslinker phytic acid molecules (right column), or the electrostatic interaction between negatively charged -OH groups and positively charge PANi due to phytic acid doping. (B-D) Photographs showing the key steps of the electrode fabrication process. (B) First, SiNPs were dispersed in the

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hydrogel precursor solution containing the crosslinker (phytic acid), the monomer aniline and the initiator ammonium persulphate. (C) After several minutes of chemical reaction, the aniline monomer was crosslinked, forming a viscous gel with a dark green colour. (D) The viscous gel was then bladed onto a 5 × 20 cm2 copper foil current collector and dried.

Figure 2.10 Synthesis and characterization of an interconnected Si hollow nanosphere electrode. (A) Schematic of hollow sphere synthesis. Silica particles (R ~175 nm) are first coated onto a stainless steel substrate, followed by CVD deposition of Si. The SiO2 core is then removed by HF etching. (B) Typical cross-sectional SEM image of hollow Si nanospheres showing the underlying substrate and the electrode thickness of ∼12 μm. (C) SEM side view (45° tilt) of the same sample of hollow Si nanospheres. (D) SEM image of hollow Si nanospheres scraped using a sharp razor blade, revealing the interior empty space in the spheres. (E) TEM image of interconnected hollow Si spheres (Rin ~175 nm, Rout ~200 nm). The area outlined by red dotted line indicates a shared shell between two interconnected spheres. Inset, the selected area electron diffraction pattern shows the amorphous nature of the sample. (F) Energy dispersive X-ray spectroscopy (EDS) of a hollow Si sphere. The carbon and copper signals come from the holey carbon TEM grids.

Figure 2.11 Preparation of bulk 3DNP-Si by dealloying in a metallic melt. (A) Schematic of producing 3DNP-Si by dealloying in a metallic melt. (B) Schematic of the evolution of the 3DNP structure by dealloying. (C) SEM image and corresponding EDX mapping of Si–Mg precursor. Color of the mapping represents atomic concentration defined by the color bar in the right side of image. (D) SEM image and corresponding EDX maps of sample after immersion in Bi. (E) SEM image and corresponding EDX map of sample after etching. (F) Photograph of the bulk 3DNP-Si powder. (G) TEM image and corresponding selected area diffraction pattern of single grain from the powder.

Figure 2.12 Fabrication of naturally rolled-up C/Si/C multilayer microtubes.

Figure 2.13 Cross-sectional schematic drawing (not to scale) of a high-capacity, stable electrode, made of a continuous, conducting 3-D network of graphite (red) anchoring regions of graphene–Si composite. Blue circles: SiNPs, black lines: graphene sheets.

Figure 2.14 Schematic of C-Si-graphene composite formation: (A) natural graphite is transformed to (B) graphene and then (C) coated by SiNPs and (D) a thin C layer.

Figure 2.15 Schematic of the fabrication (upper panel) and adapting (lower panel) of

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SiNW@G@RGO. The fabrication process mainly includes (A) chemical vapor deposition (CVD) growth of overlapped graphene sheets on as-synthesized SiNWs to form SiNW@G nanocables, and (B) vacuum filtration of an aqueous SiNW@G-GO dispersion followed by thermal reduction. The resulting SiNW@G@RGO can transform between an expanded state and a contracted state during lithiation–delithiation cycles, thus enabling the stabilization of the Si material.

Figure 2.16 Fabrication process of the nanocomposite: (A) self-assembly between SiNPs and PDDA to render the SiNP charged positively (hereafter abbreviated as Si-PDDA nanoparticles); and, (B) self-assembly between positively charged Si-PDDA nanoparticles and negatively charged GO followed by freeze-drying, thermal reduction, and HF treatment. The golden balls and the black coatings represent SiNPs and graphene sheets, respectively.

Figure 2.17 Schematic of the fabrication process for hybrid nanostructures. (A) A nickel foam substrate. (B) CVD growth of CNTs on nickel foams. (C) Mixture of Si NPs and GO suspension in ethanol. (D) Addition of PMMA into the mixed solution. (E) Deposition of GO/Si/PMMA mixtures onto the CNT-coated nickel foam and (F) thermal annealing at 700 °C for 4 h.

Figure 3.1 Glove box system MBraun-unilab (1200/780)

Figure 3.2 BTS-5V1mA battery testing equipment

Figure 3.3 CHI660D electrochemical workstation

Figure 4.1 HRTEM image of SiNPs used in this work.

Figure 4.2 XPS survey (A) and Si 2p spectrum of SiNPs (B).

Figure 4.3 SEM (A, B), TEM (C), and HRTEM images (D) of Si/RGO-AG composite.

Figure 4.4 SEM images of SiNPs (A) sample Si/RGO (B), which was prepared by using physical mixing method and sample Si/RGO-SWAG (C, D), which was prepared by single RGO wrapping.

Figure 4.5 EDS mapping of Si/RGO-AG.

Figure 4.6 XRD patterns of SiNPs and Si/RGO-AG.

Figure 4.7 Raman spectra of SiNPs and Si/RGO-AG.

Figure 4.8 C 1s XPS spectra of GO (A) and Si/RGO-AG (B).

Figure 4.9 N2 adsorption/desorption isotherms (A) and BJH pore size distribution based on

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desorption isotherm of Si/RGO-AG (B). N2 adsorption/desorption isotherm of SiNPs (C); TGA curves for SiNPs, RGO aerogel and Si/RGO-AG conducted in air (D).

Figure 4.10 (A) Cyclic voltammetry curves for Si/RGO-AG for the first two cycles; (B) Galvanostatic discharge/charge profiles of the first three cycles; (C) cycling performance of electrode Si/RGO-AG, Si/RGO-SWAG, and Si/RGO; (D) rate capability of electrode Si/RGO-AG.

Figure 4.11 SEM image of Si/RGO-AG after 40 cycles.

Figure 4.12 Electrochemical impedance spectroscopy (EIS) curves of Si/RGO-AG.

Figure 5.1 (A-B) FESEM images of RGO under different magnifications.

Figure 5.2 (A) Survey XPS patterns of GO and RGO; (B) C1s XPS patterns of RGO and GO.

Figure 5.3 XRD patterns of RGO, GO and graphite.

Figure 5.4 Raman spectra of RGO, GO, and graphite.

Figure 5.5 (A) Cyclic voltammetry curves for RGO; (B) Galvanostatic discharge-charge profiles of the first three cycles for RGO; (C) cycling performances of RGO and graphite at the current rate of 200 mA·g-1; (D) rate capability of RGO composite.

Figure 6.1 (A and B) Top-view FESEM images of Si@RGO under different magnifications (The white arrows and red circles in Figure 6.1B show RGO wrapping layers and encapsulated SiNPs, respectively.); (C) Cross-sectional view FESEM image of Si@RGO; (D) TEM image of Si@RGO composite.

Figure 6.2 Survey XPS patterns of sample Si@RGO 1:1 (A) and N2 adsorption/desorption isotherms of sample Si@RGO 1:1. Figure 6.3 FESEM images of RGO (A) and Si@RGO composites with different Si/GO ratios (1:2 B, 1:1 C, 2:1 D).

Figure 6.4 (A) XRD patterns of Si@RGO, SiNPs and RGO; (B) Raman spectra of Si@RGO, SiNPs and RGO.

Figure 6.5 (A) Cyclic voltammetry curves for Si@RGO 1:2 composite; (B) Galvanostatic discharge-charge profiles of the first three cycles.

Figure 6.6 (A) cycling performances of Si@RGO composites with different Si/GO mass ratios at a current density 200 mA·g-1; (B) rate capabilities of Si@RGO composites with different Si/GO mass ratios.

Figure 6.7 (A) cycling performances of Si@RGO 2:1 composite using different binders; (B) cycling performances of Si@RGO 1:1 composite using different binders.

List of schemes

Scheme 1.1 The scheme of a common LIB: (A) aluminium current collector; (B) metal oxide cathode active material; (C) porous separator in organic electrolyte; (D) solid electrolyte interface layer; (E) graphite anode active material and (F) copper current collector. Scheme 3.1 Illustration of coin cell battery fabrication.

Scheme 4.1 Schematic illustration of the preparation of Si/RGO-AG: (A) surface modification of

SiNPs with PDDA; (B) formation of Si@GO composite via electrostatic interactions; (C)

reduction of Si@GO using GA to form Si@RGO hydrogel; (D) dispersing Si@RGO in a

GO suspension for further reduction using GA.

Scheme 6.1 Schematic of the preparation procedure of Si@RGO composite: (A) surface modification of SiNPs with poly(diallydimethylammonium chloride, PDDA); (B) adding GO to form a Si@GO composite via electrostatic interactions; (C) reduction of the Si@GO to obtain a Si@RGO composite; (D) freeze drying and thermal treatment to obtain the final product.

List of tables

Table 1.1 History of electrochemical cell development

Table 5.1 Comparison of graphene anode materials prepared by reducing GO using various methods

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List of abbreviations

3DNP three-dimensional nanoporous

AG artificial graphite

ALD atomic layer deposition

APTES (3–aminopropyl)triethoxysilane

BET Brunauer-Emmett-Teller

CMC sodium carboxymethyl cellulose

CNT carbon nanotube

CV cyclic voltammetry

CVD chemical vapor deposition

DEC diethyl carbonate

DI deionized

DMC dimethyl carbonate

EC ethylene carbonate

ECG edge-selectively carboxylated graphite

EDA Ethylenediamine

EDS energy dispersive X-ray spectroscopy

EI electrified interface

EIS electrochemical impedance spectroscopy

EMC ethyl methyl carbonate

GA gallic acid

GNS graphene nanosheet

GO graphene oxide

HEV hybrid electric vehicle

HOPG highly oriented pyrolitic graphite

HRTEM High-resolution transmission electron microscopy

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LIB lithium-ion battery

MPECVD microwave plasma enhanced chemical vapor deposition

NMR nuclear magnetic resonance

PAA poly (acrylic acid)

PAN polyacrylonitrile

PANI polyaniline

PDDA poly(diallydimethylammonium chloride)

PE polyethylene

PHEV plug-in hybrid electric vehicle

PP polypropylene

PPy polypyrrole

PVDF polyvinylidene fluoride

RGO reduced graphene oxide

SA sodium alginate

SEI solid electrolyte interphase

SEM scanning electron microscopy

SHE standard hydrogen electrode

SiNP Si nanoparticle

SiNW Si nanowire

SNWA Si nanowire arrays

SWCNT single-walled carbon nanotubes

TEM transmission electron microscopy

TGA thermogravimetric analysis

XPS X-ray photoelectron spectroscopy

XRD X-ray diffraction

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Chapter 1 Introduction

1.1 Background

Environmental pollutions, global warming, increasing demand of energy supply as well as the accompanying surging price of the traditional energy sources have been attracting worldwide concerns. People all over the globe are trying their utmost to explore an alternative way to substitute the exhausting fossil energy sources. Renewable energy such as: solar radiation, wind and waves have already entered people’s horizon. On the other hand, these clean energy sources require energy storage devices. The common energy carriers are the electricity grid, electromagnetic waves, and chemical energy. The most convenient one for energy storage is chemical battery.

Battery possesses the advantages such as: the portability of stored chemical energy and deliver this chemical energy as electrical energy with high conversion efficiency, no gaseous exhaust, etc.1 The great success of LIBs in portable electronic divisions has been achieved since it was first 1

commercialized by Sony at early 1990s. The higher volumetric and gravimetric energy density of a LIB has already made it popularly used in small electrical devices, such as cellular phones and lap-top computers.2 Furthermore, LIBs are also regarded as the most promising green generation of energy storage technologies applied in hybrid electric vehicles (HEVs), plug-in hybrid electric vehicles (PHEVs). However, several unfavorable issues including cost, safety, stored energy density, charge/discharge rates, and lifetime continue to plague the further utilization of the LIBs for the potential mass market of electric vehicles.3

1.2 Development history of battery

The first electrochemical cell is discovered by an Italian physicist Alessandro Volta (1800) because he demonstrated that two different metals, zinc and copper, could generate electric current by decomposing water and generating hydrogen when they were immersed in acidic electrolyte. This discovery was followed a few decades later by Michael Faraday’s major advances in developing the laws of electrochemistry, and the subsequent development of rechargeable batteries with aqueous electrolytes, notably lead–acid (Gaston Planté, 1859), nickel–cadmium (Waldemar Jungner, 1899), and nickel–iron (Thomas Edison, 1901) systems. Nickel–cadmium and nickel–iron batteries were the forerunners of modern-day nickel–metal hydride batteries, introduced into the market in 1989.4

Li-based rechargeable batteries were first proposed in the early 1960s and since then the battery has undergone several transformations. Initially, metallic Li was selected as the anode but it induced serious safety hazards because of dendritic Li grows and pierces the separator membrane during cycling and finally triggers short circuit.5 The most inspiring breakthrough in LIB technology happened in 1991 owing to Sony Corporation’s introduction of a high-voltage (~3.7 V) and high-energy LixC6/non-aqueous liquid electrolyte/Li1-xCoO2 cell for portable electronic applications. Furthermore, the discovery of stable, liquid organic carbonate solvents allowed the reversible operation of these LIBs at high voltage, at least up to 4.2 V. And since 1991, graphite has remained the anode material of choice.4 Table 1.14, 6 lists the landmarks of the history of batteries.

Table 1.1 History of electrochemical cell development

Battery Invented Battery inventor Battery name type year Primary 1800 Alessandro Volta Voltaic pile battery 1836 John Frederic Daniel Daniel cell 1844 William Robert Grove Grove cell 1860 Callaud Gravity cell 1866 Georges-Lionel Leclanché Leclanché wet cell 1888 Carl Gassner Zinc-carbon dry cell 1955 Lewis Urry Alkaline battery 1970 No information Zinc-air battery 1975 Sanyo Electric Co. Lithium-manganese cell 2004 Panasonic Corporation Oxyride battery

Secondar 1859 Raymond Gaston Planté Planté lead-acid cell y battery 1881 Camille Alphonse Faure Improved lead-acid cell 1899 Waldmar Jungner Nickel-cadmium cell 1901 Thomas Edison Nickel-iron cell 1946 Union Carbide Company Alkaline manganese secondary cell 1970 Exxon laboratory Lithium-titanium disulfide 1980 Moli Energy Lithium-molybdenum disulfide 1990 Samsung Nickel-metal hydride 1991 Sony Lithium-ion 1999 Sony Lithium polymer

1.3 Working principles and terms for LIBs

A typical LIB can be represented below:

( )Cn |1 mol/L LiPF 6 -EC+DEC| LiMO x ( ) (1)

Where C stands for a carbonaceous material and M indicates a metal. On the cathode, the following reaction occurs:

-+charge LiMO -ye Li MO +yLi (2) xdischarge 1-y x

Whereas on the anode, the following reaction takes place:

+-charge C +yLi +ye Li C (3) ndischarge y n

The overall reaction becomes:

charge LiMO +C Li MO +Li C (4) x ndischarge 1-y x y n

Scheme 1.1 illustrates the working mechanism of a typical LIB. During charging process, lithium ions are deintercalated from the layer-structure cathode, then transported through the porous separator in the electrolyte and intercalated into the carbonaceous anode. During discharging process, the lithium ions are deintercalated from the anode and intercalated back to the empty sites between the layers in the cathode. Therefore, such a LIB is called “rocking chair” battery from the view point of lithium ion transfer mechanism.

Scheme 1.1 The scheme of a common LIB: (a) aluminium current collector; (b) metal oxide cathode active material; (c) porous separator in organic electrolyte; (d) solid electrolyte interface layer; (e) graphite anode active material and (f) copper current collector.

The battery capacity is defined as the amount of charge that can be obtained via the charging and discharging processes under specific conditions. The theoretical capacity of a battery is determined by the amount of an active electrode material given by:

mo 1 Coo26.8 n m (Ah) Mq (5)

where Co in Ah is the theoretical capacity, mo in g is the mass of the active material participating in the electrochemical reaction, M in g/mol is the molar weight of the active material, n is the number of electron involved in the reaction, and q is the electrochemical equivalence.

For example, the theoretical capacity of the commonly used graphite anode LiC6 is calculated as below:

LiC Li+ +C +e 66 (6)

1 LiC6 =26.8  1000  339.50 (mAh/g) 78.94 (7)

The actual capacity of a battery at constant current is given by:

C i t (8) or

1 t2 C Vdt t R 1 (9) at a constant resistance.

The theoretical energy value of a battery is the maximum value that can be delivered:

WCV o o o (10)

Where Vo is the standard potential.

The maximum power that can be produced by the battery is given by:

P iV oo (11)

The actual energy and power are determined by substituting the working voltage into Equations (10) and (11), respectively. 5

1.4 Significance and research objective of the project

Graphite, the conventional and commercialized anode material for LIBs suffers from low theoretical capacity (372 mAh·g-1), hampering its practical application in advanced LIBs. Si is a promising anode material for new-generation LIBs because of its high theoretical Li ion storage capacity (~4200 mAh·g-1), significantly higher than that of graphite. Si also features low discharge potential (~0.4 V vs Li/Li+), rich natural abundance, and environmental benignity.7, 8, 9 However, the realization of Si as a LIB anode has been hindered because Si suffers from low intrinsic electronic conductivity and large specific volume changes (>400%) during lithiation and delithiation, resulting in poor rate capability, pulverization of Si particles and electrical disconnection from the current collector, finally leading to a rapid capacity fading.10, 11, 12 Two approaches are usually taken to improve its electrode performance. One is to shorten the electronic and ionic transport length by fabricating nanostructured Si. The other is to increase the electron transport by conductive carbon coating or adding conductive additives such as carbon nanotubes, graphene.13 Although using nano-size Si particles or wires can efficiently decrease Li+ diffusion length and accommodate large stress and strain without cracking, nano-sized Si is easy to form aggregations, thus leading to poor cycling performance.14

Graphene, a monolayer of SP2 carbon atoms arranged in a 2D honeycomb network, has been shown to be effective for stabilizing Si nanoparticles (SiNPs) to prepare composite anode materials for LIB.15, 16, 17, 18, 19, 20, 21, 22, 23 Due to its high electric conductivity, superior mechanical strength and flexibility, graphene can enhance overall electric conductivity and facilitate Li+ diffusion when preparing composite material with Si, thus improving the electrochemical performance of the materials.24, 25, 26, 27, 28 Apart from stabilizing SiNPs, graphene as an anode candidate for LIB can store more lithium ions than graphite because lithium ions can not only be stored on both sides of graphene sheets, but also on the edges and covalent sites.15 Chemical reduction of GO is regarded as an economical and scalable method, which has been widely employed for graphene preparation. Generally, the RGO product tends to aggregate because of π–π stacking. Besides, some reducing agents (i.e. sodium borohydride29, hydroxylamine30 and hydrazine31) are toxic or explosive. Therefore, sustained efforts have been paid to explore green and safe approaches to reducing GO. Moreover, it is important to design a unique mechanically strong and conductive RGO framework to confine nano-sized Si and form stable solid electrolyte interphase (SEI) layers.

The objective of this project is set to design and prepare Si-based composite electrode with RGO for

advanced LIBs taking the followings into consideration:

(1) Explore a facile, green, scalable and cost-efficient way to prepare RGO architecture by chemically reducing GO; (2) Prohibit the severe aggregation of SiNPs during material synthesis by surface modification and GO sheets confining; (3) Alleviate the heavy volume changes of SiNPs, and improve the overall electric conductivity and ionic conductivity of the electrode using designed RGO architecture; (4) Prepare mechanical strong electrode and simplify the electrode preparation process.

1.5 Thesis outline

This thesis contains seven chapters presented in the following sequences:

Chapter 1 Introduction – This chapter first gives a brief introduction of research and development history of battery as well as LIB working principles, then introduces the significance of the project and summarizes the research objectives.

Chapter 2 Literature Review – Different components for LIBs are overviewed in this chapter first. Then Si and RGO based anodes are analyzed. Finally the latest research achievements of Si/Graphene composite anode are emphasized and summarized.

Chapter 3 Research Methodology – This chapter describes configuration and assembly of LIBs in lab, the battery performance measurement methods, and the characterization techniques.

Chapter 4 Stabilization of Silicon Nanoparticles in Graphene Aerogel Framework for Lithium Ion Storage – This chapter demonstrates a novel approach to wrapping SiNPs in three-dimensional RGO aerogel. The graphene-wrapped SiNPs aerogel showed improved cycling performance with excellent rate capacity than the samples prepared using simple mixing method.

Chapter 5 Lithium-storage Properties of Reduced Graphene Oxide Using Gallic Acid as Reducing Agent – This chapter introduces a facile and scalable chemical reducing approach to preparing highly-reduced porous RGO architecture using gallic acid as reductant. The RGO particles showed inspiring lithium ion storage capability as anode for LIB.

Chapter 6 Silicon Nanoparticles Encapsulated in Reduced Graphene Oxide Framework as a Stable Anode for Lithium Ion Storage – This chapter describes a new approach to stabilize SiNPs by encapsulating SiNPs in reduced RGO framework to form a micron-sized sandwich-type composite with different starting Si/GO mass ratios. Good cycling performance and rate capability were achieved.

Chapter 7 Conclusions and recommendations for future work – This chapter is a summary of the thesis with recommendations for the future work.

Chapter 2 Literature review

A LIB mainly consists of a cathode, an anode, electrolyte and a separator. The cathode materials are commonly Li-containing transition metal oxides with a layered structure (such as lithium cobalt oxide) or tunnel-structured materials (such as lithium manganese oxide). The anode materials include insertion-type materials (such as graphite, TiO2, etc.), conversion-type materials (such as iron oxides, nickel oxides, cobalt oxides, etc.), and alloying-type materials (such as Si, Sn, etc.). The electrolyte is supposed to be a good ionic conductor and electronic insulator. Most of the electrolytes are based on the solution of inorganic lithium salts dissolved in a mixture of two or more organic solvents. The function of the separator is to prevent short circuit between the two electrodes and to provide abundant channels for transportation of Li ion during the process of charging/discharging.

2.1 Cathode

Cathode materials are typically Li-containing oxides of transition metals with layer structure, which undergo oxidation reaction to higher valences when lithium is removed. While oxidation of the transition metal can maintain charge neutrality in the compound, large compositional changes often lead to phase changes, so crystal structures that are stable over wide ranges of composition must be used. This structural stability is a particular challenge during charging when most (ideally all) of the lithium is removed from the cathode. During discharge lithium is inserted into the cathode material and electrons from the anode reduce the transition metal ions in the cathode to a lower valence.32

The key requirements for a successful cathode material in a rechargeable LIB are as follows:33

(1) The material contains a readily reducible/oxidizable ion, for example a transition metal; (2) The material reacts with lithium in a reversible manner; (3) The material reacts with lithium with a high free energy of reaction; (4) The material reacts with lithium very rapidly during both insertion and extraction processes; (5) The material is supposed to be a good electronic conductor, preferably a metal; (6) The material should be stable. For example, there won’t be structure changing or degrading if over-discharging or over-changing takes place; (7) The material should be less expensive; (8) The material is better to be environmentally friendly;

During the early age of LIBs, metal chalcogenides (e.g., TiS2 and MoS2) and manganese or vanadium oxides were chosen as the cathode materials to pair with the metallic Li or the graphite anode. After several decades’ development and the increasing requirement of high-performance electrode material from large-scale energy storage, the main focus on the cathode materials for LIBs has been devoted to layered, structured hexagonal LiCoO2, spinel LiMn2O4, olivine LiFePO4, and their derivatives.

2.1.1 Layered structured cathode

There are three basic layered cathode materials: LiCoO2, LiNiO2 and LiMnO2.

LiCoO2 has been and is still widely used in commercialized LIBs since its first patented as a lithium intercalation cathode material in 1980 by Goodenough and its first appearance in the commercial LIBs by SONY in 1991.34 It possesses the advantages of large theoretical capacity (274 mAh·g-1, + 35 calculated based on extraction of 1 mole of Li from LiCoO2) and high working voltage (> 3.4 V). The cathode undergoes serious phase transformation between hexagonal and monoclinic when the cathode charged above 4.2 V. Since this phase transformation leads deterioration of cycling 36 performance, practical capacity is restricted as half of theoretical capacity. Although the LiCoO2 cathode dominates the LIB market, there is a limited availability of cobalt, which makes the cost of a single battery expensive. This price limits its application only in the area of small cells of portable electrical devices, such as those used in computers, cell phones, and cameras.33

3+ 6 1 Lithium nickel oxide, LiNiO2, due to the configuration of low spin Ni :t2g eg , electrons are 2- removed only from the eg band that barely touches the top of the O :2p band, thus most of the lithium can be extracted from LiNiO2 prior to the oxygen loss and a high capacity of around 200 -1 37 mAh·g is generated. LiNiO2 has the same structure with lithium cobalt oxide but has not been pursued in the pure state as a battery cathode for a variety of reasons, even though nickel is less expensive than cobalt. First, it is not clear that stoichiometric LiNiO2 exists. Most reports suggest excess nickel as in Li1-yNi1+yO2; thus, nickel is always found in the lithium layer, which pins the

NiO2 layers together, thereby reducing the lithium diffusion coefficient and the power capability of the electrode. Second, compounds with low lithium contents appear to be unstable due to the high effective equilibrium oxygen partial pressure, so that such cells are inherently unstable and therefore dangerous in contact with organic solvents.33

There are two structures of LiMnO2, the layered structure and the orthorhombic phase. The layered structure is thermodynamically more stable. The layered LiMnO2 can be synthesized using the soft 38 chemical method such as ion-exchange from the sodium analogue NaMnO2. Generally speaking, the manganese based material has advantages of low price and low toxicity and high initial charge capacity. However the main shortcoming of the layered LiMnO2 is the crystallographic transformation to more stable spinel structure during electrochemical cycling. Mn3+ is a Jahn-Teller ion and the redox reaction between Mn3+ and Mn4+ will result in local distortion and quick capacity decay.39

2.1.2 Spinel structured cathode

LiMn2O4 is an extensively studied cathode material for LIBs. Compared with LiCoO2, it possesses the advantages of cheaper price, lower toxicity, and higher rate capability. However, the main drawback of spinel LiMn2O4 cathode is drastic capacity fading, especially at high temperatures, and shows a low capacity of around 120 mAh·g-1.40 The dissolution of manganese from the lattice into the electrolyte due to a disproportion of Mn3+ ions into Mn2+ and Mn4+ plays a vital role for the problem, which is facilitated by trace amount of HF generated in the electrolyte.41

2.1.3 Olivine structured cathode

In 1997, Goodenough and co-workers42 firstly reported reversible lithium intercalation in the phosphate polyanionic compound of LiFePO4. The Olivine-type of LiFePO4 with a theoretical capacity of 170 mAh·g-1 has been considered as a promising cathode material for LIBs.43,44 The main drawbacks of LiFePO4 are the poor conductivity resulting from the low Li ion diffusion rate and low electronic conductivity.45

2.1.4 Strategies to enhance the performance of cathode

One can never neglect the fact that the nano-particle forms of lithium transition metal oxides such as LiCoO2, LiNiO2, or their solid solutions can react with the electrolyte and lead to safety problems.

In the case of LiMn2O4, the use of nanoparticles causes undesirable dissolution of Mn as mentioned before. Significant efforts have been made to increase the stability of these nanocrystalline lithium metal oxides. Better stability can be achieved by coating the electrode materials with a nano-sized stabilizing surface layer.

To improve the electrochemical kinetics, the cathode materials need to be embedded within an electronically conducting network, for example, some thin coating of conductive material. The coatings must be thin enough, on the nanoscale, so that ions can penetrate through them without

appreciable polarization. Furthermore, the internal electrical field generated by electrons may enhance the ionic motions. Such surface modifications alleviate the problem of low electronic conductivity, at the same time, reducing the size of active material would shorten the diffusion length of Li+. The realization of such nanostructured composites consisting of cathode materials and conductive additives makes it possible to utilize theoretical capacities at intermediate or even higher current rates.43

2.2 Anode

Anode is the negative electrode of a primary cell and is always associated with oxidation reaction, accompanying with release of electrons to external circuit. In a secondary cell, anode is the negative pole during discharge but the positive pole during charge. The basic requirements for a LIB anode material are that the material should have minimal volume expansion and resulting stress during charge/discharge process, high electronic conductivity, low irreversible capacity during initial charging or intercalation process, stable under wide operating temperature window in a highly reducing environment, and low specific surface area (typically < 2 m2/g) for optimal performance and safety.34

2.2.1 Metallic Li anode

The most elementary anode material for lithium-based batteries is metallic lithium, which has been used for primary batteries since the early 1960s. By possessing the lowest standard potential (–3.05 V vs. a standard hydrogen electrode (SHE)) and the lowest atomic weight (6.94 g·mol–1; specific gravity: ρ= 0.53 g·cm–3) among all metals, the utilization of metallic lithium as an anode offers the realization of galvanostatic cells having an extremely high capacity (3820 mAh·g-1) and energy density (1470Wh·kg-1).46 Consequently, metallic lithium was also considered as the candidate of choice for secondary LIBs. However, lithium metal cells have one severe drawback, namely, inhomogeneous lithium plating, which hinders their commercial development. This uneven deposition of lithium onto the anode surface upon charge results in the formation of so-called dendrites. And the disordered metallic deposit gives rise to a poor Coulombic efficiency, resulting

from such a fine Li metal often acts as an active site inducing reductive decomposition of electrolyte components. Part of the deposit may become electrically isolated and shedding may also occur. These dendrites consist of high surface area, highly branched lithium metal structures, which continuously grow, may easily penetrate into the separator and eventually cause internal short, resulting in heat generation and contingent ignition. Hence the hazardous nature of Li has paved way to explore some other safer anode materials, possessing comparably the same electrochemical features as that of metallic lithium.

2.2.2 Carbonaceous anode

The most commonly used anode material is carbonaceous materials as they possess the advantages of low cost, excellent reversible capacity and long cycling life.47 Carbon as an insertion anode material was first reported in 1973. Carbonaceous materials have large variations in crystallinity, chemical composition and microtexture depending on their preparation, processing, precursor, thermal and chemical activation treatments. All carbonaceous materials are able to storing lithium ions.48

Graphitic carbons are characterized by layers of hexagonally arranged carbon atoms in a planar graphene layers. Graphite shows a rather flat potential profile when reversibly hosting Li+ at potentials below 0.5 V vs. Li/Li+. Additionally, it offers a significantly higher specific capacity of –1 32 372 mAh·g (corresponding to one lithium per hexagonal carbon ring, i.e. LiC6). The volume expansion during lithium intercalation and extraction between the graphene layers is just over 10%, resulting in high reversibility and stable capacity over repeated cycling.49 However, compared with other alternative anode materials, the theoretical capacity of graphite is far behind to meet the requirement from batteries with high energy density.

2.2.3 Sn and SnO2 based anode

Sn and SnO2 are promising anode materials with high theoretical capacity. The reversible capacity -1 of Sn is as high as 994 mAh·g when alloy Li4.4Sn forms during lithiation and the theoretical

-1 32,50 capacity of SnO2 is 782 mAh·g . However both Sn and SnO2 suffer from severe volume change during cycling.51,52

One of the most efficient approaches to alleviate the severe volume change and stabilizing the structure is to design nanostructured materials, for example, in the form of nanotubes, nanorods, nanowires and hollow structures.53, 54, 55 Although most of the attempts have been made to synthesize new nanostructures for Sn and SnO2, up to now most efforts have been characterized as having technical difficulties or high synthesis costs, and most of them have proved not to be very 56 successful in improving the cycling stability of Sn and SnO2-based electrodes. Adding coating materials (such as carbon or other conductive materials) is another method to accommodate volume changes. In addition, the electrical conductivity of the whole electrode could also be enhanced. 57,58

2.2.4 Lithium titanate anode

Lithium titanate (Li4Ti5O12) is another promising anode material with spinel structure for certain applications that require fast charging capability, first reported in 1994.46 It has superior high rate performance with very long cycle life. It is being developed for automotive and energy storage applications. Some of the challenges for the Li4Ti5O12-based batteries include lower voltage (2.5 V 34 vs. LiCoO2), lower capacity, and excellent high-temperature stability.

2.2.5 Other metal oxide anode

Apart from SnO2, other metal oxides, such as Co3O4, TiO2, NiO and Fe3O4 have been studied. These metal oxides also suffer from severe aggregation and rapid capacity fading due to dramatic volume expansion and low electrical conductivity.32,59,60 The most commonly studied methods to solve the problems are coating carbon layers on the surface of the metal oxides or dispersing them into carbon matrix.61, 62, 63

2.3 Graphene-based anode

Graphene is a single layer carbon nanosheet consisting of two equivalent sub-lattices of sp2 carbon atoms connected by σ bonds.64 An theoretical calculation65 showed that graphene possesses a maximum capacity of 740 mAh·g-1 on the basis of a double-layer adsorption configuration, which is much larger than that of layered graphite. Sato and co-workers66 used the 7Li nuclear magnetic resonance (NMR) technique to study lithium ion transport mechanism in a disordered carbon, and proposed a model of Li2 covalent molecule configuration that a Li2 molecule is loosely trapped by two adjacent benzene rings in the carbon. The most saturated Li storage state can be LiC2 to give a capacity of up to 1116 mAh·g-1. Both of the two configurations suggest a higher capacity of graphene anode. Another study showed that graphene and graphite have a similar interaction mode with Li ions after analysing in situ Raman spectra of lithiation processes in both single- and few-layer graphene anodes. However, because of the repulsion forces between Li+ at both surfaces of graphene, the amount of adsorbed Li on single-layered graphene is low.67 Uthaisar and Barone68 studied the edge effects of Li diffusion in graphene by means of density functional theory and showed that narrower graphene nano-ribbons, especially with a zigzag morphology, can provide a faster discharge rate owing to the lowering of energy barriers and diffusion lengths. The theories mentioned above can explain the Li storage mechanisms to some extent; however, more work is still needed in the coming future to erase the debate.

2.3.1 Preparation of graphene

Preparation of graphene with large scale, controlled sizes is an important step to realize its applications. The first successful method to prepare graphene was by pealing (use scotch-tape) from highly oriented pyrolytic graphite (HOPG), then deposited on a Si substrate assisted by organic solvents.69 The highest quality graphene obtained was 10 μm in size. But the yield and efficiency are very low.70 A prerequisite for graphene as anode for LIBs is the availability of processable graphene in large quantities.

2.3.1.1 Chemical reduction of graphene oxide

Oxidative exfoliation of natural graphite by thermal or oxidation technique and subsequent

chemical reduction has been considered as one of the most efficient methods for low-cost and large-scale production of graphene.71 Ruoff et al. 72 demonstrated the solution-based route involving chemical oxidation of graphite to graphite oxide, which can be readily exfoliated as individual GO by ultrasonication. The GO can be converted back to graphene by chemical reduction, for instance, using hydrazine or NaBH4. However, graphene obtained from GO by chemical reduction often suffers from limited dispersion or even irreversible agglomeration, and the reducing agents are also toxic and dangerously unstable, making the further processing and application very difficult.

Surface modification and green chemistry route for the reduction of GO are the most desirable to produce high concentration and well dispersed graphene colloids. Li et al. 70 demonstrated that reduced graphene colloids could be prepared through electrostatic stabilization without polymeric or surfactant stabilizers. And also, the amphiphilic molecules, which had been widely used to improve the solubility of carbon nanotubes,73, 74, 75, 76, 77, 78 have been shown to act as agents for enhancing the solubility of graphene.79, 80, 81, 82, 83, 84, 85, 86, 87, 88 This is mainly because the hydrophobic parts could form strong π-π interaction with graphene, while the hydrophilic sides prevented graphene aggregation via electrostatic repulsions.

Zhang et al. 89 introduced a method producing dispersed graphene solution that employs L-ascorbic acid as the reductants and L-tryptophan as the stabilizer. Other green reductants have been widely applied to produce stable graphene colloids, such as Vitamin C and reducing sugars (glucose, sucrose, fructose etc.).90, 91, 92 The chemicals containing a similar electron-rich aromatic group function as an electron donors and can be absorbed onto graphene through π-π interaction.92

Loh et al.93 presented a green reduction of GO to stable graphene solution using hydrothermal dehydration. Compared to reduction using hydrazine, the water-only route has combined advantages on removing the oxygen functional groups and repairing the aromatic structures (Figure 2.1). Based on this, a lot of solvothermal methods have been developed to prepare high quality and high concentration graphene colloids.94, 95, 96, 97, 98, 99

Wei et al.100 reported a green and facile approach to the synthesis of graphene based on Fe reduction of GO, resulting in a substantial removal of oxygen functionalities (Figure 2.2). The C/O atomic ratio and conductivity can reach as high as 7.9 and 2300 S·m-1, respectively. Ouyang et al. 101 reported that with Zn powder as reductants for GO, the C/O atomic ratio and conductivity can reach 33.5 and 15000 S·m-1, respectively.

Figure 2.1 (a) Intramolecular and (b) intermolecular dehydration of GO at hydrothermal conditions. Procedure (b) gives rise to aggregated products.

Figure 2.2 Illustration of the preparation of graphene based on Fe reduction.

2.3.1.2 Chemical vapor deposition

Another promising and widely applied way for production of graphene is chemical vapor deposition (CVD). The CVD method for graphene mainly makes use of the pyrolysis of hydrocarbon gases on the surface of catalysts at high temperature.102 18

There are two types of growth mechanisms based on the difference of the carbon solubility in different metals.103 The first one is the carbon segregation/precipitation process (Figure 2.3). For catalysts with high solubility of carbon (such as Ni), carbon sources crack at high temperature and carbon atoms dissolve into the surface of catalysts; when cooling, the carbon atoms precipitate from the surface of catalysts followed by nucleation and growth into continuous graphene.103 Therefore, the numbers of graphene layer are strongly dependent on the cooling rate. Kong et al. 104 and Hong et al. 105 firstly utilize the polycrystalline nickel as catalyst to prepare graphene through CVD. However, it is difficult to prepare graphene with large size and controllable morphology. The number of layers is unevenly distributed. In order to resolve this problem, Cheng et al.106 adopted nickel foam as the catalyst and synthesized scale-up and macroscopic three-dimensional graphene foams by CVD method. These opened the door to various applications of graphene such as chemical sensor and energy storage devices.107, 108, 109, 110 The second mechanism is via surface growth, in which the carbon atoms do not diffuse into the matrix of catalysts (such as Cu).103 Carbon atoms are firstly adsorbed on the surface of catalysts, and then nucleate and grow into graphene island, and finally into graphene by continuous growth. Ruoff et al. 102 used copper foil to grow graphene in large scale. Ahn et al.111 reported the roll-to-roll production and wet-chemical doping of predominantly monolayer 30-inch graphene films grown by CVD onto flexible copper substrates.

Figure 2.3 Schematic diagrams of the possible growth mechanism for graphene by CVD methods. (a) Carbon segregation/precipitation process; (b) Surface growth mechanism.

Although large scale graphene can be prepared by CVD method in the presence of metal catalysts, it remains a great challenge to transfer graphene from these metals to other substrates. Currently, the transfer processes are mostly based on the etching of metal substrates in suitable etchants,105, 112 which leads to inevitable damage to graphene, loss of metal catalysts, increasing of costs, production of residues, and serious pollution. Therefore, developing a green and high efficient transfer technology becomes one of the key factors for graphene applications. Wang et al.113 demonstrated an electrochemical delamination method to repeatedly transfer the graphene from copper to another substrate. The growth and transfer processes on copper catalyst have been repeated hundreds of times. However, the substrate was partially etched for each transfer. Cheng et al. 114 reported the bubbling method to transfer single graphene grains and films to arbitrary substrate, which is non-destructive to graphene and the Pt substrates (Figure 2.4). The Pt substrates can be repeatedly used for graphene growth and the quality of graphene grown on the recycled Pt is almost the same as that on the Pt substrate of the initial condition. Further studies should be carried out to achieve scathe less substrate transfer of graphene.

Figure 2.4 Illustration of the bubbling transfer process of graphene from a Pt substrate.

2.3.1.3 The arc discharge method

The arc discharge method has been extensively used for the production of fullerenes, single- and multi-walled carbon nanotubes.115 The temperature of arc discharge method can be instantaneously increased to more than 2000 °C during arc discharge process.116 Therefore, it is naturally expected that the arc discharge can be used for efficient exfoliation and deoxygenation of GO and healing of

the resultant exfoliated graphite. Rao et al.117 reported that the synthesis of graphene by arc discharge between graphite electrodes under a relatively high pressure of hydrogen yields 2-4 layered graphene flakes. Shi et al.118 achieved low-cost and large-scale production of graphene by arc discharge in an air atmosphere instead of mixed H2 and He. They found that the yield of graphene is greatly dependent on the pressure of the atmosphere, i.e., high pressure facilitates the formation of graphene, while low pressure favours the growth of carbon nanohorns and nanospheres. Besides, they found that the nitrogen doped graphene can be produced in large scale by direct 119 current arc discharge between pure graphite rods by use of NH3 as one of the buffer gas. Moreover, hydrogen arc discharges as a rapid heating method is also used to produce good quality graphene from GO combined with solution-phase dispersion and centrifugation.116, 120 The obtained graphene exhibited a high electrical conductivity and high thermal stability compared with conventional thermal exfoliation method. These demonstrate that this hydrogen arc discharge exfoliation method is a good approach for the preparation of graphene with high quality.

2.3.1.4 The top-down approach

Graphite is a crystal consisting of stacked carbon monolayers. Each individual monolayer has the same structure as graphene. Therefore, exfoliated pristine graphite provides enormous potential for obtaining high quality graphene with high electrical conductivity and high yield because the graphene can keep the basic structure of graphite and has fewer defects. This top-down fabrication method mainly contains mechanical cleavage, high-energy ball milling, ultrasonication and electrochemical exfoliation. As discussed above, the mechanical cleavage can obtain the highest quality graphene, but the yield and efficiency are very low.

High-energy ball milling is an effective method to prepare few-layer graphene and its composites.121, 122, 123, 124 For instance, Fan et al. 123 developed a ball milling route to prepare graphene/Al2O3 composites from expanded graphite with thickness of 2.5-20 nm, which were homogeneously dispersed in the ceramic matrix. Dai’s group125 have achieved high yield of edge-selectively carboxylated graphite (ECG) by a simple ball-milling of graphite in the presence of dry ice (Figure 2.5). Then, the ECG is highly dispersible in various solvents to self-exfoliate into single- and few-layer (≤ 5 layers) graphene. This progress provides simple, but efficient and versatile, and eco-friendly approaches to low-cost mass production of high-quality graphene. In addition to mechanical cleavage and ball-milling methods, liquid exfoliated by ultrasonication is 126, 127, another effective method to prepare two-dimensional materials, such as MoS2, BN, WS2, etc.

128, 129 Coleman et al. 128, 129, 130 demonstrated that when graphite powder is exposed to ultrasonication in the presence of a suitable solvent, the powder fragments into nanosheets, which are stabilized against aggregation by the solvent. However, liquid exfoliation by ultrasonication method still has some critical disadvantages, such as low concentrations of colloidal suspensions of graphene, difficulty in obtaining single-layer graphene with high yield, and the high cost of some of the solvents. Compared with ultrasonication, electrochemical exfoliated method is an effective tool to synthesize high quality graphene.131, 132, 133 For instance, Luo et al. 134 developed electrochemical exfoliation method for the preparation of ionic liquid functionalized graphene with the assistance of ionic liquid and water. Loh et al. 133 demonstrated a solution route inspired by the LIB for the high yield (>70%, thickness <5 layers) exfoliation of graphite into highly conductive few-layer graphene. Although the above strategies provided a green route to prepare high-quality graphene, it is still very difficult to prepare graphene in a large scale by this method due to the limited volume of electrochemical chamber or devices.

Figure 2.5 Illustration of edge-carboxylation of graphite for graphene in the presence of dry ice. (A)

2.3.1.5 The bottom-up approach

The bottom-up method has been proven to be a potential procedure for preparing nanomaterials with different size, shape, composition and microstructures. As known, graphene is composed of 22

interconnected polycyclic hydrocarbons. Therefore, synthesis of graphene with different size from small molecules of organics could provide an effective route for controllable synthesis of graphene.135, 136, 137

Cai et al. 137 reported a simple method for the production of atomically precise graphene nanoribbons with different topologies and widths. They proposed a surface-assisted coupling of molecular precursors into linear polyphenylenes and their subsequent cyclodehydrogenation (Figure 2.6). They found that the topology, width and edge periphery of graphene products were defined by the structure of the monomers. Another bottom-up method is based on the high temperature solvothermal reaction. Stride et al. 138 demonstrated that single-layer graphene can be synthesized by flash pyrolysis of a solvothermal product of sodium and ethanol, followed by gentle sonication of the nanoporous carbon products. This process mainly used the high reductive ability and chemical activity of alkali metals such as sodium and potassium to produce carbon radicals, which then coupled together to form graphene. The ability to produce bulk graphene from non-graphitic precursors with scalable and low-cost is a step closer to real applications of graphene.

Figure 2.6 Bottom-up fabrication of atomically precise graphene nanoribbons.

2.3.2 Graphene-based electrode for LIBs

The electrochemical characteristics of graphene can be different when graphene is prepared in different ways.

Guo et al. 139 prepared crumpled-paper-shaped graphene by the means of oxidation, rapid thermal expansion, ultrasonic treatment from artificial graphite (AG) and compared the electrochemical performances of graphene and AG as active anode materials. The results showed that improved reversible capacity of 672 mAh·g-1 and fine cycle performance for the graphene nanosheets. Lian and co-workers140 used thermal expansion method in nitrogen atmosphere to prepare curled, few-layered graphene and the reversible specific capacity of the first cycle was as high as 1264 mAh·g-1 at a current density of 100 mA·g-1, it could still maintain 848 mAh·g-1 after 40 cycles at the same current density. The high capacity might be attributed to the large surface area, curled

morphology as well as the hydrogen atom on the prepared graphene. However, there was no clear voltage plateau on the charge and discharge curves, which would be one of main blocks for graphene to be commercially utilized. Besides, graphene sheets derived from graphite oxides with different oxidation degree was also reported.141 In another report, RGO nanosheets were prepared by oxidation and rapid expansion of graphite using microwave radiation. The prepared RGO delivered a reversible capacity around 400 mAh·g-1 for 50 cycles.142 Recently, Few-layer graphene sheets have been successfully prepared by a microwave digestion RGO method using expanded graphite as raw material. The first discharge/charge capacities are 2113.8/1069.3 mAh·g-1 at 100 mA·g-1, respectively. The reversible capacity remains 733.3 mAh·g-1 after 100 cycles.143

Thermal reduction is also a widely used method to prepare RGO. A flexible RGO paper was produced using thermal reduction method and the electrochemical performances were tested as a binder-free anode. It delivered a reversible capacity of 576 mAh·g-1 at 50 mA·g-1 for 50 cycles.144 Freeze-dried GO sample was thermal reduced under Ar atmosphere was synthesized (FGNs) and compared with the RGO sample prepared using rotary evaporation process (RGNs) instead of freeze-drying. Results showed that FGNs had wrinkled paper-like structure and larger interlayered distance, while RGNs were seriously stacked and had smaller interlayer distance, indicating freeze-drying may significantly minimize the restacking of graphene nanosheets. The battery performances of the FGNs and RGNs are shown in Figuer 2.7. The FGNs exhibit superior cycle stability (556.9 mAh·g-1 after 300 cycles at a high current density of 1000 mA·g-1) and excellent rate capability (257.3 mAh·g-1 at a superhigh current density of 5000 mA·g-1).145 A microwave-induced thermal reduction process combined post-process treatment (ultra-sonication) were used to reduce GO to prepared morphology-controlled and wrinkle free RGO nanosheets. Graphene nanosheets prepared using the modified-improved method exhibits discharge capacities of 1079 mAh·g-1 (at a constant current of 40 mA·g-1) and 1002 mAh·g-1 after 50 cycles. The capacity retention of the synthesized graphene nanosheets is 1070 mAh·g-1 at a current of 40 mA·g-1after the rate capability test, and their rate capability is 463 mAh·g-1 at a current of 400 mA·g-1.146 Liu and co-workers147 prepared 3D porous RGO using a two-step process: hydrothermal reaction and thermal reduction. Highly porous RGO material with a specific surface area of ~1,066 ± 2 m2·g−1 used as an active electrode material for LIBs, the RGO exhibited an initial discharge capacity of ~1,538 mAh·g-1. Even cycled at higher discharge/charge rates, the RGO-based electrodes still delivered large capacities and excellent cycling performance.

Chemical reduction process is regarded as a scalable and efficient approach to prepare RGO from GO apart from thermal reduction method. Flower-petal shaped RGO were prepared, using hydrazine monohydrate to reduce GO. When cycling at 1 C rate, the prepared RGO anode exhibited 25

discharging capacity of 945 mAh·g-1 and charging capacity of 650 mAh·g-1. It could still maintain a specific capacity of 460 mAh·g-1 after 100 cycles, which showed an excellent cycling stability. They attributed it to the double-layer adsorption model and believed that the nano-cavities between RGO layers also contributed to the Li storage.148 Some researchers prepared non-annealed RGO paper by reducing GO using hydrazine monohydrate and employed the fabricated RGO as single-component binder-free anode for LIB. The cycling performance of the prepared anodes was less impressive, the highest reversible capacity was about 200 mAh·g-1.149 In another work, a facile and green method to produce RGO nanosheets based on supercritical alcohols was reported. High lithium storage capacity with good cyclability was achieved (652 mAh·g−1 at a 50 mA·g−1 after 40 cycles) when tested as an anode material.150 Li and co-workers151 synthesized few-layer RGO using Ethylenediamine (EDA) as reducing agent in GO solution and the prepared RGO sample delivered a reversible capacity of 346 mAh·g−1 at 350 mA·g−1 after 60 cycles. A novel porous RGO paper was prepared via freeze drying a wet GO gel, followed by thermal and ascorbic reduction. It can deliver a high discharge capacity of 420 mAh·g−1 at a current density of 2000 mA·g−1 for 100 cycle when used as binder-free LIB anode.152

Figuer 2.7 (a) Cyclic voltammograms of the FGNs electrode, (b) Initial discharge/charge profiles of the FGNs and the RGNs at a current density of 100 mA·g-1, (c) The cycling stabilities of the FGNs and RGNs at a high current density of 1000 mA·g-1, (d) The rate capabilities and cycle performances of the FGNs and the RGNs at various current densities, Nyquist plots of the FGNs and RGNs (e) before cycling and (f) after 3 cycles.

In another work, microwave plasma enhanced CVD was employed to prepare vertically aligned graphene electrode for LIBs using Ni foil as catalyst and current collector. Stable reversible capacity was achieved around 300 mAh·g-1 for about 175 cycles.153 Wei and co-workers reported porous graphene obtained by CVD using porous MgO sheets as template is demonstrated to exhibit

a high reversible capacity (1723 mAh·g-1), excellent high-rate capability and cycling stability (926 mAh·g-1 after 100 cycles at 1 C) for LIBs.154

The morphology and structure of graphene have also become an interesting research topic. Yoo and co-workers28 believe that the space between the graphene layers and the electronic structure play an important role during the process of Li accommodation. The specific capacity of single-component graphene anode in their report was 540 mAh·g-1. When they selected carbon nanotubes (CNT) and fullerene (C60) to act as interacting molecules, the specific capacity of graphene could reach up to 730 mAh·g-1 and 784 mAh·g-1, respectively, which confirmed their assumption. Nonetheless, the stability of the prepared anode was not improved significantly. The anodes they prepared didn’t show obvious potential plateau and the reason might be the disordered stacking of the graphene. In another research, one-dimensional graphene nanoribbons were used as anode material of LIBs, which were prepared using a solution-based oxidative process to unzip pristine multiwalled carbon nanotubes. Interestingly, the reduced graphene nanoribbons didn’t present remarkable improvement of reversible capacity, whereas the oxidized graphene nanoribbons exhibited much higher specific capacity with reversible capacity of 800 mAh·g-1. They proposed that the residual oxygen enhanced the formation of a more stable Li-rich SEI layer. There was still capacity loss of about 3% per cycle during cycling.155 It is hold that the large surface area of graphene can induce severe inter-sheet aggregation during the process such as drying and thermal annealing, which could result in limiting permeation of electrolyte between graphene as well as reducing the ion mobility. Therefore, Zhao et al.156 prepared flexible holey graphene by the method of mechanical cavitation-chemical oxidation. The cycling performance showed that the highest reversible capacity was below 300 mAh·g-1, although it could maintain stable high-rate capacity for more than 400 cycles. Metal etching method was reported for large-scale preparation of porous graphene with controlled pore size.157 Nanoporous RGO sheets were prepared by Guo et al.158 through a simple thermal annealing procedure using composites of ferrocene nanoparticles and GO sheets as precursors in a large scale. Specific discharge and charge capacities were observed at 1036.2 and 936.3 mAh·g-1 at the 2nd cycle for the preprared RGO, respectively. After 100 cycles of discharge/charge, the specific capacity of the retained at 800 mAh·g-1, and the Coulombic efficiency was up to 95.2%. In another work, a series of metals (Ni, Fe, Cu and Co) have been applied to prepare porous graphene. The size of the pores and pore density in the basal plane of graphene sheets is controllable by the reaction time and the ratio of the reactants. The porous graphene as anode materials exhibited first charge specific capacity of 933 mAh·g-1 at 50 mA·g-1 with good cycling stability after 50 cycles. Xiao et al. reported a simple approach to grow nanolayered graphene sheets vertically aligned on the current collector as anode for LIBs in a microwave plasma enhanced chemical vapour deposition system.

Good rated capability was obtained. Recently, a scalable co-assembly strategy for the preparation of graphene-encapsulated porous carbon (CMK-3 and CMK-8) composites was reported.159 The composites show superior electrochemical performances compared with pure porous carbon or pure graphene. Reversible capacities were achieved of ~680 mAh·g-1 and ~620 mAh·g-1 after 100 cycles at a current density of 100 mA·g-1, for graphene/CMK-3 and graphene/CMK-8, respectively. The core–shell structure of these novel composites forbid the aggregation of graphene nanosheets, resulting in improved lithium ion storage performance.

2.4 Silicon-based anode

Si has drawn attention as a promising anode material since the discovery of its superior theoretical Li storage capacity (∼4200 mAh·g-1) in the 1980s, significantly higher than that of commercialized graphite (372 mAh·g-1). The packing density in Li-Si alloys is very close to, or slightly higher than, −1 that in pure Li metal, e.g., Li22Si5: 0.0851 vs. Li: 0.0769 mol· ml . Si also possesses low discharge potential (~0.4 V vs Li/Li+), rich natural abundance, and environmental benignity.7, 8, 9 However, the realization of Si as a LIB anode has been hindered because Si suffers from not only a low intrinsic electronic conductivity, resulting in poor rate capability, but also a large specific volume changes (>400%) during lithiation and delithiation. Because the intermetallics of Li/Si have much greater molar volume than the nanostructure Si phase, large stresses will be generated from the repetitive volume expansion and contraction, leading to pulverization of Si particles and electrical disconnection from the current collector, then the battery will show a rapid capacity fading10, 11, 12. And the major reasons for the poor cycling properties have been established as shown below160:

(1) Mechanical/electrical separation in the Si electrode; (2) Capacity loss by reductive reaction in the components of the native layer (silicon oxide and silanol) during lithiation; (3) Formation of SEI by reductive decomposition of the electrolyte solution; (4) Continuous SEI-filming process on fresh Si surfaces following crack formation.

At current stage, one effective protocol is to shorten the electronic and ionic transport length by fabricating nanostructured Si. Morphological design of Si aims to mitigate the effect of volume change during repetitive cycling by exploiting Si structures with at least one dimension below ~100 nm to accommodate Li without leading to the cracking or pulverization.14, 160, 161 Although using 29

nano-size Si particles or wires can efficiently accommodate large stress and strain without cracking14, and decrease diffusion length, nano-sized Si is easy to form aggregations leading to poor cycling performance. Other protocols include surface modification of nano-size Si, designing porous structures or adding conductive additives.13

2.4.1 Carbon coated Si electrode

Carbon coating has been widely studied to prepare Si-based anode materials for LIBs, being one of few materials that have both electrical and ionic conductivity. Additionally, carbon coating is also used to stabilize the Si-electrolyte interface, promote stable SEI formation, and extend cycle life.

Cui’s group engineered empty space between Si nanoparticles (SiNPs) by encapsulating them in hollow carbon tubes using electrospinning methods. The prepared anode demonstrated a high gravimetric capacity (~1000 mAh·g-1 based on the total mass) and long cycle life (200 cycles with 90% capacity retention).162 Almost at the same time, Choi and co-workers163 also developed an electrospinning process to produce core−shell fiber electrodes using a dual nozzle method shown in Figure 2.8. The unique core−shell structure efficiently resolved various issues of Si anode, such as pulverization, vulnerable contacts between Si and carbon conductors, and an unstable SEI. Outstanding cell performances were achieved: a gravimetric capacity as high as 1384 mAh·g-1, a 5 min discharging rate capability retaining 721 mAh·g-1, and cycle life of 300 cycles with almost no capacity loss. In addition to electrospinning process only, a flexible 3D Si/C fiber paper electrode synthesized by simultaneously electrospraying nano-Si-PAN (polyacrylonitrile) clusters and electrospinning PAN fibers followed by carbonization was reported. The combined techniques allowed uniform incorporation of SiNPs into a carbon textile matrix to form a nano-Si/carbon composite fiber paper. The flexible 3D Si/C fiber paper electrode showed an overall capacity of ~ 1600 mAh·g-1 with capacity loss less than 0.079% per cycle for 600 cycles and excellent rate capability.164 In another work from Cui’s group9, they prepared Si@void@C powders used a sol–gel method to conformally coat amorphous SiO2 sacrificing layer onto the SiNPs first. The obtained

Si@SiO2 then mixed with dopamine hydrochloride. Finally, carbonized the polydopamine coating, and remove the SiO2 sacrificial layer. The well-defined void space allowed the Si particles to expand freely without breaking the outer carbon shell, therefore stabilizing the SEI on the shell surface. High capacity (~2800 mAh·g-1 at C/10), long cycle life (1000 cycles with 74% capacity retention), and high Coulombic efficiency (99.84%) have been realized in this yolk-shell structured 30

165 Si electrode. Yang and co-workers reported a novel Si/void/SiO2/void/C dual yolk-shell nanostructure included a hard SiO2-coated layer, a conductive carbon-coated layer, and two internal void spaces. In this structure, the carbon shell could enhance conductivity, the SiO2 shell had mechanically strong qualities, and the two internal void spaces could confine and accommodate volume expansion of Si during lithiation. Therefore, these dual yolk-shell structures exhibited a stable capacity of 956 mAh·g-1 after 430 cycles with capacity retention of 83%. Recently, a yolk–shell structured Si–C nanocomposite was synthesized using a facile NaOH etching method, high specific capacity, good cycling stability and rate performance were observed.166 Different from the above methods: etching SiO2@C to form voids using NaOH or HF aqueous solution, a straightforward approach to prepare yolk–shell mesoporous Si/C nanocomposites was developed using magnesiothermic-reduction method to obtain mesoporous Si and citric acid as the precursor of C shell. The prepared composite showed good retention of specific capacity (1264.7 mAh·g-1 after 150 cycles at 100 mA·g-1).167

Apart from coating Si particles, Si nanowires (SiNWs) can also be modified. Si-C nanohybrid anode with SiNWs dwelling in hollow graphitic tubes was developed comprising four steps: synthesis of SiNWs, formation of SiO2 sacrificial layer, growth of graphitic carbon sheaths and removal of the SiO2 sacrificial layer. The Si-C nanohybrid exhibited good rate capability and remarkable cycling stability (~1100 mAh·g-1 at 4200 mA·g-1 over 1000 cycles).168

2.4.2 Metal oxide coated Si electrode

Cui et al.169 reported SiNPs with atomic layer deposited zinc oxide (ZnO) coating showed significantly improved electrochemical performance mainly attributed from enhanced mechanical integrity and stablized interface. In-situ structural characterization confirmed that the coating is mechanically robust to remain intact even though Si expanded inside, and a battery electrode made of this Si/ZnO structure underwent minimal structural damage upon cycling. These ZnO coated SiNPs demonstrated stable cycling with a reversible capacity of 1500 mAh·g-1 over 260 cycles with ZnO protected interface.

Besides ZnO, Al2O3 deposited by the atomic layer deposition (ALD) method has also been studied as a protection layer for Si. Significantly improved cycling performance was achieved due to the 170, 171, 172 suppression of the side reaction between Si and the electrolyte by the thin Al2O3 layer. Recently, nanostructured micron-sized Al–Si particles were synthesized via a facile and cost-effective selective etching process of Al–Si alloy powder. Subsequently thin Al2O3 layers were introduced on the Si foam surface via a selective thermal wet oxidation process of etched Al–Si particles. The resulting Si/Al2O3 foam anodes exhibited outstanding cycling stability (a capacity retention of 78% after 300 cycles at the C/5 rate) and excellent rate capability.173

2.4.3 Polymer modified Si electrode

Polymer binders are components used in preparation of electrodes to keep the integrity of the electrodes. Previous studies show that binders play a critical role in cell performance. A number of 32

polymers have been investigated as binders in LIBs.174, 175, 176, 177 Polyvinylidene fluoride (PVDF), one conventional binder, exhibits a good ductility and has been examined for Si anodes. But it is not a good binder agent. The Si anodes using PVDF always suffered rapid capacity fading. Therefore, the impact of binder on the battery performance of Si-based anodes indicates the need for a strong connection between the swelling Si particles in order to keep electrode integrity. A chemical reaction between the electrochemical active sites and the binder must occur, forming a covalent chemical bond between the binder and the active mass and improving cycling stability.178

In addition to the impact of volume changes during repetitive cycling, the properties and chemistry of the Si interface bring an additional hurdle to its utilization in advanced batteries. Different from carbonaceous materials, Si is covered by an intrinsic oxide layer. This native oxide layer terminated with a hydroxyl group always exists in commercially available Si materials. It is readily formed during synthesis, storage, and transportation due to the presence of trace amounts of oxygen and water.179 This native oxide increases the irreversible capacity loss, and is involved in the formation of a complicated SEI layer. This surface SiO2 layer was shown to contribute to a higher first-cycle capacity loss compared to that of the Si samples etched by hydrofluoric acid.180 However, from the binder point of view, existing functional groups on the Si surface provide a reaction point with the polymer binders. This has been regarded as the origin of some useful binders, such as poly (acrylic acid) (PAA), sodium carboxymethyl cellulose (CMC), and sodium alginate (SA), for Si-based anodes.179 Respectable LIB performances were obtained with PAA, CMC, and SA.176, 181, 182, 183, 184 They all have carboxylic functional groups and are soluble in water with environment-friendly characteristic.

In addition to the polymer binder modification, Cui’s group185 reported an anode by incorporating SiNPs into a nanostructured 3D porous conducting polymer hydrogel shown in Figure 2.9. The in-situ polymerized hydrogel created a well-connected 3D network structure consisting of SiNPs conformally coated by the conducting polymer. The hierarchical hydrogel framework possessed several advantages, including a continuous electrically conductive polyaniline network, binding with the Si surface through either the crosslinker hydrogen bonding with phytic acid or electrostatic interaction with the positively charged polymer, and porous space for volume expansion of Si. The as-prepared anode delivered a cycle life of 5,000 cycles with over 90% capacity retention at current density of 6000 mA·g-1. In another report, a ternary hybrid anode composed of SiNPs wrapped in a 3D hierarchically porous nanostructured conductive polypyrrole (PPy) framework with single-walled carbon nanotubes (SWCNTs) as the electronic fortifier delivered reversible discharge capacity over 1600 mAh·g-1 and 86% capacity retention over 1000 cycles.186

Figure 2.9 (a) Schematic illustration of 3D porous SiNP/conductive polymer hydrogel composite electrodes. Each SiNP is encapsulated within a conductive polymer surface coating and is further connected to the highly porous hydrogel framework. SiNPs has been conformally coated with a polymer layer either through interactions between surface -OH groups and the phosphonic acids in the crosslinker phytic acid molecules (right column), or the electrostatic interaction between negatively charged -OH groups and positively charge polyaniline (PANI) due to phytic acid doping. (b–d) Photographs showing the key steps of the electrode fabrication process. (b) First, SiNPs were dispersed in the hydrogel precursor solution containing the crosslinker (phytic acid), the monomer aniline and the initiator ammonium persulphate. (c) After several minutes of chemical reaction, the aniline monomer was crosslinked, forming a viscous gel with a dark green colour. (d) The viscous gel was then bladed onto a 5 × 20 cm2 copper foil current collector and dried.

2.4.4 Si electrode with designed morphology

SiNWs have recently attracted intense attention for application as LIB anodes for the following attributes: (1) SiNWs can be directly connected to the current collector without additional binders or conducting additives; (2) SiNWs can offer direct one-dimensional electronic pathways for rapid 34

charge transport; (3) SiNWs can provide rapid charge/discharge rates due to high surface-to-volume ratio; and (4) SiNWs with small diameters can accommodate large volume changes stemming from the alloying reaction and can prevent fragmentation.187 Si nanowire arrays (n-SNWAs) with a coral-like surface on Cu foam were synthesized via a one-step CVD method, in which the Cu foam simultaneously acted as a catalyst and current collector. The as-prepared n-SNWAs on Cu foam directly applied as the anode for LIBs, exhibiting a reversible discharge capacity of 2745 mAh·g-1 at 200 mA·g-1, 2178 mAh·g-1 at 400 mA·g-1 after 50 cycles, and fast charge and discharge capability of 884 mAh·g-1 at 3200 mA·g-1.188

In addition to design SiNWs, Cui et al.189 reported an interconnected Si hollow nanosphere electrode by a template approach using solid silica nanospheres as shown by the schematic in Figure 2.10. The as-prepared Si hollow nanosphere electrode is capable of accommodating large volume changes without pulverization during cycling wit high initial discharge capacity of 2725 mAh·g-1 and less than 8% capacity degradation every hundred cycles for 700 total cycles as well as a Coulombic efficiency of 99.5% in later cycles. Superior rate capability is demonstrated and attributed to fast lithium diffusion in the interconnected Si hollow structure.

(B) Typical cross-sectional SEM image of hollow Si nanospheres showing the underlying substrate and the electrode thickness of ~12 μm. (C) SEM side view (45° tilt) of the same sample of hollow Si nanospheres. (D) SEM image of hollow Si nanospheres scraped using a sharp razor blade, revealing the interior empty space in the spheres. (E) TEM image of interconnected hollow Si spheres (Rin ~175 nm, Rout ~200 nm). The area outlined by red dotted line indicates a shared shell between two interconnected spheres. Inset, the selected area electron diffraction pattern shows the amorphous nature of the sample. (F) Energy dispersive X-ray spectroscopy (EDS) of a hollow Si sphere. The carbon and copper signals come from the holey carbon TEM grids.

Porous structured Si composite anode is also a promising choice. Liu et al.190 synthesized large (>20 μm) mesoporous Si sponge with thin crystalline Si walls surrounding large pores that are up to ~50 nm in diameter using an electrochemical etching method. In-situ transmission electron microscopy showed that the volume change of the Si walls during the charge/discharge processes was primarily accommodated by the inner pores in the mesoporous Si sponge, leading to only ~30% expansion in the full particle size rather than >300% expansion as in the case of conventional Si particles. Furthermore, the particles didn’t pulverize after 1000 charge/discharge cycles. The mesoporous Si sponge anode delivered a capacity of up to ~750 mAh·g-1 based on the whole electrode weight with 80% capacity retention. In another report, hierarchical micro/nano porous Si powders were prepared through a wet-chemical etching method. which showed no capacity fading at 1500 mAh·g-1 in 50 charge/discharge cycles.191 Recently, a large-area preparation of mesoporous SiNW as LIB anode from metallurgical Si (purity ~99.74%) with adjustable porosity by metal-assisted chemical etching was reported. It exhibited a reversible capacity of about 2111 mAh·g-1 at a current rate of 0.2 C for over 50 cycles as well as a good rate capability.192 Du et al.193 demonstrated that a Si electrode with a large amount of nanopores fabricated by an in situ thermal generating approach using triethanolamine as a sacrificing template achieved high reversible capacities (2500 mAh·g-1) with excellent cycling stability (∼90% capacity retention after 100 charge–discharge cycles). To simplify the preparation process of three-dimensional nanoporous (3DNP) Si, a top-down process to dealloy a metallic melt was designed to prepare bulk Si nanoligaments with widths of several hundred nanometers directly interconnected without need for a template shown in Figure 2.11. Basically, Si–Mg–Bi system was used, with Si acting as the porous-structure-forming element, Mg acting as sacrificial element, and Bi acting as the dealloying melt medium. By immersing a Si–Mg precursor in a Bi melt, Mg should selectively dissolve from the Si–Mg alloy into the Bi melt; the Si atoms left at the solid/liquid interface cohere, growing Si islands in the melt. After solidification by cooling, the Bi elements filling the interspace of Si 36

islands were etched away; after drying in air, the 3DNP-Si was produced. The capacity of the nanoporous Si (∼1000 mAh·g-1) could be retained over 1500 cycles.194

Figure 2.11 Preparation of bulk 3DNP-Si by dealloying in a metallic melt. (a) Schematic of producing 3DNP-Si by dealloying in a metallic melt. (b) Schematic of the evolution of the 3DNP structure by dealloying. (c) SEM image and corresponding EDX mapping of Si–Mg precursor. Color of the mapping represents atomic concentration defined by the color bar in the right side of image. (d) SEM image and corresponding EDX maps of sample after immersion in Bi. (e) SEM image and corresponding EDX map of sample after etching. (f) Photograph of the bulk 3DNP-Si powder. (g) TEM image and corresponding selected area diffraction pattern of single grain from the powder.

A naturally rolled-up C/Si/C trilayer nanomembrane using rolled-up nanotechnology shown in Figure 2.12 was designed by Schmidt et al.195 First, a sacrificial layer (red color, photoresist ARP 3510) was deposited on Si substrate by spin-coating, then trilayer C/Si/C (10/40/10 nm, respectively) nanomembranes were sequentially deposited by radio frequency sputtering, during which the intrinsic strain was generated from thermal expansion effects. The sacrificial layer was then 37

selectively under-etched. Finally, the composite nanomembranes naturally peeled off from the Si substrates and broke into micron-sized pieces because of the pulling strain. These pieces further self-rolled into multilayer tubular structures. The nanomembrane delivered a highly reversible capacity of ~2000 mAh·g-1 at current density of 50 mA·g-1 with nearly 100 % capacity retention at 500 mA·g-1 after 300 cycles. Based on the same mechanism, a sandwich nanoarchitecture of rolled-up Si/RGO bilayer nanomembranes was designed by the same group and showed long cycling life of 2000 cycles at 3000 mA·g-1 with a capacity degradation of only 3.3% per 100 cycles.21

Figure 2.12 Fabrication of naturally rolled-up C/Si/C multilayer microtubes.

2.5 Si/Graphene composite anode

Graphene can be utilized to enhance the cycling performance of Si anode and become a hot research topic.17, 18, 19, 20, 21, 22, 23, 196, 197, 198, 199, 200 Graphene is electrically conducting, mechanically strong, and rather easy to prepare from exfoliated graphite. Which means graphene is an attractive support for other high-storage capacity materials.201 The excellent properties of graphene, as mentioned above, can be taken advantage of to enhance the performance of Si anode.

Sandwich-type Si/graphene composites have been widely investigated as anode for LIBs. In 2009, Dou and co-workers202 prepared Si/graphene composite anode from commercially available SiNPs and solvothermal fabricated graphene by the method of simply mixing in mortar. The prepared electrode exhibited a specific capacity of 1168 mAh·g-1 after 30 cycles with a Coulombic efficiency of 93% on average. The impedance spectroscopy showed that the charge-transfer resistance of the composite anode was less than half of SiNPs. Almost at the same time, SiNPs well dispersed in 38

graphene were prepared by thermal reducing Si/GO composite where a portion of graphene reconstituted to form graphite structure (shown in Figure 2.13) reported by Kung et al.203 The electrode showed highly conducting 3-D network and the greatly improved cycling performance, more than 2200 mAh·g-1 after 50 cycles and more than 1500 mAh·g-1 after 200 cycles. The reconstitution of graphene ensured better electron conductivity of the whole anode as well as better connection with surrounding SiNPs. It was also pointed out that excessive constitution of graphene should be avoided so that SiNPs could disperse homogeneously in the composite. Tao et al.201 studied a self-supporting Si/RGO nanocomposites film following the process of preparing GO by Hummers method, simply filtrating Si/GO suspension and thermal reducing GO under Ar atmosphere. The first time reversible specific capacity of the prepared anode was 1040 mAh·g-1 and it could still maintain 786 mAh·g-1 after 300 cycles at a current rate of 50 mA·g-1. However, the rate capacities decreased significantly as the current rates rose from 50 mA·g-1 to 5000 mA·g-1. Similar work was done by Liu and co-workers204 ,the free-standing, paper-like Si/graphene composite materials exhibited a discharge capacity of 708 mAh·g-1 after 100 cycles.

In another work, SiNPs have been successfully inserted into graphene sheets via a novel method combining freeze-drying and thermal reduction. Freeze-drying is defined as direct sublimation of solvents under vacuum to dry samples and maintain their microstructures. The as-obtained Si/graphene nanocomposite exhibits remarkably enhanced cycling performance (ca. 1153 mAh·g-1 after 100 cycles) and rate performance compared with bare SiNPs for LIBs.200 Recently, Kung et al.20 reported ordered sandwich-structured magnesiothermo-reduced SiNPs-thermally RGO

nanocomposites prepared by combined magnesiothermic reduction, freeze-drying, and thermal reduction. Good cycling performance (746 mAh·g-1, over 160 cycles) was achieved over the electrode fabricated with the 80:10:10 weight ratio of active material: acetylene black: sodium carboxymethyl cellulose (CMC) and the 5 v% vinylene carbonate in electrolyte. Chang et al.205 reported a 3D multilayered Si/RGO nanohybrid anode was prepared by assembly of alternating Si/RGO layers on porous Ni foams based on a dip-coating method. The 3D Si/RGO electrode exhibited high reversible specific capacity (2300 mAh·g-1 at 0.05 C, 700 mAh·g-1 at 10 C), fast rate capability (up to 10 C), and good capacity retention (87% capacity retention with a rate of 10 C after 152 cycles up to 630 mAh·g-1, 780 mAh·g-1 at 3 C after 300 cycles).

Different from Si/graphene sandwich structure with Si particles wrapped by graphene wrapping layers, Yushin et al.206 reported that uniform Si and C coatings can be deposited on high surface area graphene via vapor decomposition routes shown in Figure 2.14. The produced Si and C-coated graphene granules exhibit specific surface area of ~5 m2·g−1 and an average size of ~25 μm. The anode composed of the nanocomposite particles exhibited specific capacity in excess of 2000 mAh·g-1 at the current density of 140 mA·g-1 and excellent stability for 150 cycles. In another report, in-situ growth binder-free Si-based anode was prepared. Adaptable SiNPs were encapsulated in graphene nanosheets (SiNPs@GNS), simultaneously, the SiNPs@GNS composites anchored on vertically aligned graphene trees with loose intersecting leaves in a microwave plasma enhanced chemical vapor deposition system (MPECVD). The composite material exhibited a capacity (1528 mAh·g-1 at 150 mA·g-1), good cycle stability (88.6% after 50 cycles), and fast charge/discharge rate (412 mAh·g-1 at 8000 mA·g-1).17

Figure 2.14 Schematic of C-Si-graphene composite formation: a) natural graphite is transformed to b) graphene and then c) coated by SiNPs and d) a thin C layer.

Researchers also showed great interest in 1-D SiNWs in addition to SiNPs and studied the composite of SiNW/graphene197. Zhi’s group207 reported a self-supporting binder-free anode (SiNW@G@RGO) via the encapsulation of SiNWs with dual adaptable apparels (overlapped graphene (G) sheaths and RGO overcoats) shown in Figure 2.15. The overlapped graphene sheets prevented direct exposure of Si to electrolyte and the flexible conductive RGO overcoats accommodated the volume change of embedded SiNW@G nanocables. The SiNW@G@RGO electrodes exhibited high reversible specific capacity of 1600 mAh·g-1 at 2100 mA·g-1, 80% capacity retention after 100 cycles, and superior rate capability (500 mAh·g-1 at 8400 mA·g-1) calculated on total electrode weight.

Figure 2.15 Schematic of the fabrication (upper panel) and adapting (lower panel) of SiNW@G@RGO. The fabrication process mainly includes (I) chemical vapor deposition (CVD) growth of overlapped graphene sheets on as-synthesized SiNWs to form SiNW@G nanocables, and (II) vacuum filtration of an aqueous SiNW@G-GO dispersion followed by thermal reduction. The resulting SiNW@G@RGO can transform between an expanded state and a contracted state during lithiation–delithiation cycles, thus enabling the stabilization of the Si material.

The oxide layer on the surface of Si endows it with a negative charge. GO also shows a negative charge owing to ionization of the carboxylic acid and phenolic hydroxyl groups. Based on surface modification and graphene coating methods, Zhou and co-workers24 intentionally used a positively charged polyelectrolyte poly(diallydimethylammonium chloride) (PDDA) to make surface modification of Si and then realize the self-assembly of SiNPs and GO. After thermal reduction, a self-assembled nanocomposite of well-dispersed SiNPs encapsulated in graphene was obtained shown in Figure 2.16. The Si/graphene nanocomposite possessed a flexible graphene shell that effectively encapsulated the SiNPs, as a result, the composite delivered stable cycling performance (approximately 1205 mAh·g-1 after 150 cycles) and excellent rate capability. In another work, monolayer SiNPs were firmly anchored on GOs by covalent immobilization using (3–aminopropyl)triethoxysilane (APTES) and subsequent thermal reduction. The most stable SiNP–graphene hybrid sample exhibited an initial Li+ insertion capacity of 1297 mAh·g-1, and the 50th Li+ insertion capacity of 1203 mAh·g-1 with 92.7% capacity retention.199 Apart from using 208 APTES surface modification method to prepare Si@NH2/GO composite, Chen et al. further combined Si@NH2/GO composite with sodium alginate, acting as a binder, to yield a stable

composite negative electrode. These two chemical cross-linked/hydrogen bonding interactions–one between NH2-modified SiNPs and GO, the other between the GO and sodium alginate–along with highly mechanically flexible GO, produced a robust network in the negative electrode system to stabilize the electrode during discharge and charge cycles. The as-prepared Si@NH2/GO electrode exhibited good capacity retention capability and rate performance, delivering a reversible capacity of 1000 mAh·g-1 after 400 cycles at a current of 420 mA·g-1 with almost 100% capacity retention. Wu and co-workers209 designed a 3D porous interconnected network of graphene-wrapped porous Si spheres micro/nanostructure (3D Si@G network) constructed through layer-by-layer assembly and in situ magnesiothermic-reduction method. First, SiO2 sphere templates were synthesized via a modified Stober method. Then, the 3D interconnected network of GO sheets-wrapped SiO2 spheres

(SiO2@GO network) was prepared through a layer-by-layer assembly approach based on electrostatic attraction using PDDA and poly(sodium 4-styrenesulfonate) (PSS). Finally, the as-synthesized SiO2@GO network was reduced by Mg powder then annealed. The Si@G network was obtained by etching the formed MgO and possible Mg2Si by-product with HCl solution. The Si@G network exhibited enhanced performance in terms of specific capacity, cycling stability, and rate capability.

Figure 2.16 Fabrication process of the nanocomposite: i) self-assembly between SiNPs and PDDA to render the SiNP charged positively (hereafter abbreviated as Si-PDDA nanoparticles); and, ii) self-assembly between positively charged Si-PDDA nanoparticles and negatively charged GO followed by freeze-drying, thermal reduction, and HF treatment. The golden balls and the black coatings represent SiNPs and graphene sheets, respectively.

Instead of using SiNPs with a native oxide layer, in another work, a simple method to fabricate graphene-encapsulated pyrolyzed polyaniline-grafted SiNPs has been developed using HF-treated SiNPs. The composite materials exhibited good cycling stability and Coulombic efficiency as anodes in LIBs. After 300 cycles at a current density of 2000 mA·g-1, the composite electrode 43

delivered a specific capacity of ~900 mAh·g-1 with ~76% capacity retention.210

In another interesting work, a multilayered structural Si-RGO electrode was synthesized from bulk Si particles through inexpensive electroless chemical etching to obtain porous macro Si particles and graphene self-encapsulating approach governed by electrostatic interaction. The prepared composite electrode delivered a reversible capacity of 2787 mAh·g-1 at a charging rate of 100 mA·g-1, and a stable capacity over 1000 mAh·g-1 at 1000 mA·g-1 after 50 cycles.19 Chang et al.211 developed a binder-free Si-based anode through encapsulation of SiNPs with carbon coating and RGO coating, where CNT rooted from a nickel foam resulted in a strong connection mechanically and electrically between active materials and current collectors. In the resulting architecture shown in Figure 2.17, a dense cellular carbon coating from carbonization of poly(methyl methacrylate) (PMMA) on Si surfaces improved the electrical conductivity and accommodated the volume change, whereas RGO networks provide additional mechanical strength to maintain the integrity of electrodes. The nanostructure exhibited a reversible capacity up to 2700 mAh·g-1 at 0.05 C (130 mA·g-1) and 70% of capacity retention (up to 1311 mAh·g-1) at 2600 mA·g-1after 900 cycles.

Figure 2.17 Schematic of the fabrication process for hybrid nanostructures. (a) A nickel foam substrate. (b) CVD growth of CNTs on nickel foams. (c) Mixture of Si NPs and GO suspension in

ethanol. (d) Addition of PMMA into the mixed solution. (e) Deposition of GO/Si/PMMA mixtures onto the CNT-coated nickel foam and (f) thermal annealing at 700 °C for 4 h.

2.6 Separator

A separator is a porous membrane placed between electrodes and soaked in electrolyte, which is permeable to ionic ﬂow but preventing electronic contact between the electrodes. Another important function of separator is a shutter action to stop the mass transfer in the case of accidental heat generation, which cause it melt down, resulting in shut down of a cell. Separators should be very good electronic insulators and have the capability of conducting ions by either intrinsically being an ionic conductor or by soaking an electrolyte. They should minimize any processes that adversely affect the electrochemical energy efﬁciency of the batteries. The material should be soft and flexible enough to act as buffer between the both sheets of cathode and anode. The material should also be sufficiently stable for a long time because it is directly in contact with the electrolyte all the time.212

A variety of separators have been used in batteries over the years. Starting with cedar shingles and sausage casings, separators have been manufactured from cellulosic papers and cellophane to nonwoven fabrics, foams, ion exchange membranes, and microporous flat sheet membranes made from polymeric materials. Currently, the most commonly used separators for LIBs are polyolefin membranes, which are made of polyethylene (PE), polypropylene (PP), or laminates of PE and PP. They are made up of polyolefin materials because they provide excellent mechanical properties, chemical stability, and acceptable cost. They have been found to be compatible with the cell chemistry and can be cycled for several hundred cycles without significant degradation in chemical or physical properties.34 As batteries have become more sophisticated, highly-performance separator has also become more demanding. A wide variety of properties are required of separators used in batteries,212, 213 including:

(1) Electronic insulator; (2) Minimal electrolyte (ionic) resistance; (3) Mechanical and dimensional stability; (4) Sufficient physical strength to allow easy handling; (5) Chemical resistance to degradation by electrolyte, impurities, and electrode reactants and products; (6) Effective in preventing migration of particles or colloidal or soluble species between the two electrodes; 45

(7) Readily wetted by electrolyte; (8) Uniform in thickness and other properties.

2.7 Electrolyte

A mixture of aprotic organic solvents is typically used in commercial LIBs, including cyclic and linear alkyl carbonates such as ethylene carbonate (EC), dimethyl carbonate (DMC), diethyl carbonate (DEC), ethyl methyl carbonate (EMC) and propylene carbonate (PC). The solvent composition is chosen to achieve a balance between viscosity, the ability of the solvent to form a stable passivation layer, and sufficiently high dielectric constant to dissolve lithium salts such as lithium hexafluorophosphate (LiPF6) or lithium perchlorate (LiClO4). The operating potential of the anode in lithium cells is close to that of metallic lithium, which lies outside of the thermo-dynamic stability region of organic electrolytes. Therefore, the electrolyte solution undergoes electrochemical decomposition, resulting in the formation of a SEI layer, which prevents electron travelling between the electrolyte and electrode, and maintains kinetic stability outside of the solvent potential window.32

Electrolyte is actually the main component determining the current (power) density, the time stability, and the safety of a battery. Taking it into account that the electrolyte is in intimate contact with the two electrodes, a good chemical and physical compatibility with both the electrodes is required in order to obtain a stable electrified interface (EI) and, consequently, energy and power density values stabilized with time. In spite of a wide variety of liquid electrolytes developed during the 1990s, many LIBs preferentially chose some optimized formulations, which consist of proper 214 concentrations of LiPF6 in a mixture of EC and a linear alkyl carbonate.

An ideal electrolyte solvent for LIBs shall meet the following minimal criteria, namely, high dielectric constant, ability to dissolve salts of sufficient concentration, low viscosity for facile ion transport, inert to all cell components, especially the charged surfaces of the electrodes, low melting point, and high boiling point to remain in liquid phase in a wide temperature range.

An ideal electrolyte solute in LIBs completely dissolves and dissociate, in the nonaqueous media, and the solvated ions should be able to move in the media with high mobility. The electrolyte solute should be stable against oxidative decomposition at the positive electrode, should be inert to electrolyte solvents and other cell components, and should be nontoxic and remain stable against thermally induced reactions with electrolyte solvents and other cell components.212

2.8 Summary and perspectives

There is no doubt that great markets exist for higher energy density batteries and increasing research efforts to develop high-power and high-capacity LIBs. Current commercial LIBs are limited by the capacity of their electrode materials. Although Si possesses the highest specific capacity as anode material for LIBs, development of Si to replace graphite faces several hurdles; most notably, the volume changes of Si during Li+ intercalation and de-intercalation causes cracking and pulverization of the active material. Nano-sized Si has proven to be successful in overcoming mechanical degradation, providing a good substrate adhesion and allowing for shorter lithium diffusion distances. But Nano-sized Si is easy to aggregate and large surface area leads to massive SEI growth, causing heavy irreversible capacities.

Graphene has shown growing research attentions since it was given birth. Graphene as an anode for LIB can store more lithium ions than graphite because lithium ions can not only be stored on both sides of graphene sheets, but also on the edges and covalent sites. Graphene as coating layer for Si is also attractive because graphene is electrically conducting and mechanically strong. Various fabrication methods have been developed for graphene and its derivatives preparation, and approaches for the synthesis of high-quality graphene in large scale at low price are required.

Recent works have shown surface coating, surface modification and good structure design of Si-based anode are effective protocols to improve the electrode performance. As development of Si anodes continues to extend the cycle life and improve high-rate performance, it may not be long before these high-capacity Si anodes find their way into practical applications.

2.8 References

Chapter 3 Research methodology

This chapter describes the experimental setup and technical theories. The designed materials were synthesized, characterized and electrochemically tested in the laboratory. Electrode preparation technique and material characterization protocols are introduced in this chapter. The battery assembly procedure and electrochemical testing details are also presented.

3.1 Preparation of GO

Natural graphite flake (325 mesh, 5.0 g) and NaNO3 (2.5g) are mixed with 120 ml H2SO4 (95 %) in a 500 ml flask. Within ice bath, stir the mixture for 30 min. Under vigorous stirring, potassium 48

permanganate (15.0 g) is carefully added little by little to the suspension. The temperature of the reactants should be controlled below 20 °C and keep stirring overnight. Then 150 ml H2O is slowly added under vigorous stirring and the temperature quickly rises to 98 °C, meanwhile the color of the suspension changes to yellow. Keep stirring for one day and add 50 ml of 30 % H2O2 slowly. To purify the product, the mixture is washed by rinsing and centrifugation with 5 % HCl then deionized (DI) water for several times. Dry the purified GO and store it in a dry tower for further research.

3.2 Anode materials preparation

See the following chapters for details.

3.3 Fabrication of coin cell batteries

A manual cutter is used to cut the anode disks with a diameter of 15 mm. The CR2032 coin cells are assembled in glovebox (MBraun Company, Germany, Figure 3.1) under ultra-pure Ar atmosphere. Scheme 3.1 demonstrates the schematic diagram of a typical used CR2032 coin cell in this project. A metallic Li disk (99.9% purity, 200um thick, 15mm diameter) is chosen as the counter electrode and 1 M LiPF6 in ethylene carbonate (EC) : Diethyl carbonate (DEC) : dimethyl carbonate (DMC) 1:1:1 w/w will be used as electrolyte. A sheet of circular polyethylene (PE) membrane, whose diameter is a little larger than the electrodes, will be placed between the two electrodes as separator. A steel disk as spacer and a wave spring are employed between the anode and the negative cell shell to make sure a close contact of each component. After assembly using a crimping machine (MTI company, China), the prepared coin cells are placed aside for 24 hours and then undergo electrochemical testing.

Figure 3.1 Glove box system MBraun-unilab (1200/780)

Scheme 3.1 Illustration of coin cell battery fabrication.

3.4 Electrochemical measurements

To test the electrochemical performances of the prepared anode disks, below measuring protocols

are chosen to be conducted including the galvanostatic charge-discharge test, Cyclic Voltammetry (CV) and Electrochemical Impedance Spectroscopy (EIS) tests.

Charge-discharge Test

Cycling testers from Neware Company (Figure 3.2) are used to test the cycling performances of the coin cells at constant current in a setting voltage range for RGO and Si@RGO composite materials. Charging and discharging capacities are tested at a constant current density between the voltages from 0.005 to 3.500 V and from 1.20 to 0.01 V for carbon materials and Si@RGO composites, respectively.

Figure 3.2 BTS-5V1mA battery testing equipment

Cyclic Voltammetry

CV curves are recorded with a CHI660D electrochemistry workstation (Figure 3.3). CV measurement is an electrochemical technique applied to record the current development in an electrochemical cell when the potential at the working electrode is ramped linearly with time in the set potential range. The same type of coin cells are used for the CV analysis as that for the cycle and rate capability tests with lithium metal as the counter and reference electrodes.

Electrochemical Impedance Spectroscopy 51

EIS measurements of coin cells with metallic lithium as both counter and reference electrode are taken at room temperature with CHI660D electrochemistry workstation over the frequency range 100 kHz to 10 mHz, and the signal amplitude is 5 mV. EIS measurement is used to study the reaction mechanism of electrochemical process. The rate determining step can be revealed by the frequency response shown by EIS test because different reaction steps will dominate at certain frequencies. The variation of resistance can be examined by monitoring the current response. The obtained impedance Nyquist curve of the testing battery is made up by a depressed semicircle in the high-frequency region representing the resistance of the SEI film and the charge-transfer resistance, a straight line in the low-frequency region corresponding to the diffusion of Li+.

Figure 3.3 CHI660D electrochemical workstation

3.5 Materials Characterization

In order to achieve comprehensive study of the prepared anode materials, the widely used morphology characterization method is electron microscopy.

A scanning electron microscope (SEM) is a type of electron microscope that produces images of a sample by scanning it with a focused beam of electrons. The electrons interact with atoms in the sample, producing various signals that contain information about the sample's surface topography and composition. The electron beam is generally scanned in a raster scan pattern, and the beam's position is combined with the detected signal to produce an image. SEM can achieve resolution better than 1 nanometer.

The most common SEM mode is detection of secondary electrons emitted by atoms excited by the electron beam. The number of secondary electrons that can be detected depends, among other things, on the angle at which beam meets surface of specimen, i.e. on specimen topography. By scanning the sample and collecting the secondary electrons that are emitted using a special detector, an image displaying the topography of the surface is created.

Apart from SEM, below are necessary strategies and instruments utilized in the research:

(1) Scanning electron microscopy (SEM, JEOL 6610, JEOL 7001 and JEOL 7800F at 5.0 kV) (2) Transmission electron microscopy (TEM, Philips Tecnai F20 and Philips Tecnai F30 at 200 kV) (3) High-resolution transmission electron microscopy (HRTEM, Philips Tecnai F30 at 200kV) (4) Energy-dispersive X-ray spectroscopy analysis (EDX, JSM6610 EDS system) (5) X-ray photoelectron spectroscopy (XPS, Kratos Axis ULTRA, using a monochromatic Al Ka (1486.6 eV) X-ray source and a 165 mm hemispherical electron energy analyzer) (6) X-ray diffraction (XRD, German Bruker D8 Advanced X-Ray Diffractometer with Ni filtered Cu Kα radiation (40 kV, 30 mA)) (7) Raman Spectroscopy (Thermo-Fischer Almega dispersive Raman instrument fitted with both 633 and 785 nm lasers)

(8) N2 adsorption/desorption isotherms (Tristar II 3020) (9) Thermogravimetric analysis (TGA, DTG-60A system under air flow)

Chapter 4 Stabilization of Silicon Nanoparticles in Graphene Aerogel Framework for Lithium Ion Storage

4.1. Introduction

With the fast growing demands for energy storage devices, the development of new-generation LIBs is becoming increasingly important. Achieving high capacity, excellent cycling performance and rate capability with innovative electrode materials has been one of the main challenges.32 Silicon (Si) is a promising anode material for new-generation LIBs because of its high theoretical Li ion storage capacity (∼4200 mAh·g-1), significantly higher than that of commercialized graphite (372 mAh·g-1). Si also features low discharge potential (~0.4 V vs Li/Li+), rich natural abundance, 7, 8, 9 and environmental benignity. However, the realization of Si as a LIB anode has been hindered because Si suffers from not only a low intrinsic electronic conductivity but also a large specific volume changes (>300%) during lithiation and delithiation, resulting in pulverization of Si particles and electrical disconnection from the current collector, leading to a rapid capacity fading. 10, 11, 12 To solve these problems, nanostructured Si electrodes such as nanoparticles, 9, 200, 215, 216 nanowires, 217, 218, 219, 220, 221 nanotubes and hollow spheres 189, 222 have been explored. Combining SiNPs with carbon materials has also been studied. 24, 200, 216, 223, 224, 225, 226, 227

Graphene, a monolayer of carbon atoms arranged in a two-dimensional (2D) honeycomb network, has been used to improve the stability and electric conductivity of nanostructured Si electrodes for LIBs. 15 Due to its high electronic conductivity, superior mechanical strength and flexibility, graphene can improve electron transport and Li+ diffusion, thus enhancing the electrochemical performance of SiNPs for lithium ion storage. 24, 25, 26, 27 In spite of the observed improvement of the electrochemical performance of SiNPs by graphene, SiNPs still tend to aggregate. As a result, the performance of the graphene-SiNPs composites inevitably degrades during charge/discharge. Recently, three-dimensional (3D) graphene materials, including hydrogels and aerogels, have been shown to possess advantages, such as high surface area and good electrical conductivity. 228, 229, 230 These 3D graphene materials may be favourable for stabilizing SiNPs.

This work describes a method for preparing graphene-stabilized SiNPs on the basis of electrostatic interactions. Because both graphene oxide (GO) and SiNPs (there is a thin layer of SiO2 on the surface of the SiNPs) are negatively charged in a wide pH range, the SiNPs used in this work were firstly modified using poly(diallydimethylammonium chloride) (PDDA, a positively charged polyelectrolyte) to change the surface charge nature from being negative to being positive following a protocol described elsewhere.24 As schematically illustrated in Scheme 4.1, the PDDA-modified SiNPs interacted strongly with negatively charged GO sheets to form a Si-GO composite 55

(hereinafter designated as Si@GO). Because of the flexibility of GO sheets, SiNPs were wrapped in by the GO sheets. The GO in the Si@GO suspension was then reduced using gallic acid (GA, a natural plant phenolic acid) in an oil bath at 95 °C for 4 h. It has been reported that in the presence of natural phenolic acid, GO can be reduced to assemble into RGO hydrogels driven by the enhancing hydrophobicity and π–π interactions among the nanosheets during the reduction. 228, 231 During this GA reduction process, yellow-brown-coloured Si@GO suspension gradually turned transparent to form a dark grey gel separating from the suspension, indicating that the GO had been reduced to reduced graphene oxide (RGO) and the SiNPs had been wrapped by the RGO architecture. In order to understand the stabilization effect of RGO sheets on SiNPs, the obtained Si@RGO hydrogel was crushed and redispersed in an aqueous GO suspension, which was further reduced with GA at 95 °C for 12 h. After freeze-drying and thermal reduction at 700 °C for 2 h under H2/Ar atmosphere, a black 3D RGO aerogel with confined SiNPs, designated as Si/RGO-AG, was obtained. For comparison purpose, another two samples were prepared. One, designated as Si/RGO, was prepared by mixing SiNPs with aqueous GO suspension, followed by filtration separation and thermal reduction at 700 °C for 2 h under H2/Ar atmosphere. The other one, designated as Si/RGO-SWAG with SiNPs entrapped in RGO single-step wrapping aerogel framework, was prepared according to the procedure of preparing sample Si/RGO-AG without Step 3 (see Experimental Section for details).

Scheme 4.1 Schematic illustration of the preparation of Si/RGO-AG: 1) surface modification of 56

SiNPs with PDDA; 2) formation of Si@GO composite via electrostatic interactions; 3) reduction of Si@GO using GA to form Si@RGO hydrogel; 4) dispersing Si@RGO in a GO suspension for further reduction using GA.

4.2 Experimental

Preparations of graphene oxide (GO)

The GO used in this work was prepared from natural graphite flake (Sigma Aldrich, 325 mesh) by using a modified Hummers method. 232

Preparation of RGO aerogel wrapped SiNPs (Si/RGO-AG)

SiNPs (120 mg, 99% purity, from Nanostructured & Amorphous Materials, Inc.) and 20wt% PDDA (3.0 g, Sigma Aldrich, MW = 10000-20000) aqueous solution were dispersed in water (120 mL) under sonication in a water bath (KQ3200DE, 40 kHz). Excess PDDA was washed away by centrifugation (10000 rpm) four times, followed by vacuum drying. An aqueous GO suspension (15 mL, 2 mg·mL-1) was centrifuged at 5000 rpm for 30 min to remove any graphite particles and then mixed with the above PDDA-modified SiNPs (120 mg) under sonication for 1 h. After the addition of GA (13 mg), the mixed suspension was transferred to a 20 mL capped vial and placed into an oil bath at 95 °C for 4 h to form a Si@GO hydrogel. The hydrogel was collected and redispersed in another GO aqueous suspension (15 mL, 2 mg·mL-1) and GA (13 mg) under sonication. The suspension was then transferred into a 20 mL vial to prepare hydrogel at 95 °C for 12 h. The resulting hydrogel monolith was soaked in distilled water, followed by freeze-drying. Reduction was conducted in a quartz tube at 700 °C for 2 h under a H2/Ar (5:95 v/v) atmosphere with a heating rate of 2 °C min-1 to obtain sample Si/RGO-AG.

Preparation of other samples

For comparison purpose, another two samples were prepared. One of them, designated as Si/RGO, was prepared as follows. An aqueous GO suspension (30 mL, 2 mg·mL-1) was mixed with PDDA-modified SiNPs (120 mg) under sonication for 1 h. After the addition of GA (26 mg), the suspension was reduced at 95 °C in an oil bath for 12 h under magnetic stirring. Then the solid was collected by filtration, followed by vacuum drying and reduction in a quartz tube at 700 °C for 2 h

-1 under a H2/Ar (5:95 v/v) atmosphere with a heating rate of 2 °C min . The other sample, designated as Si/RGO-SWAG, was prepared as follows. An aqueous GO suspension (15 mL, 2 mg·mL-1) was mixed with PDDA-modified SiNPs (60 mg) under sonication for 1 h. After the addition of GA (13 mg), the suspension was transferred into a 20 mL capped vial and placed in an oil bath at 95 °C for 12 h. The resulting hydrogel monolith was soaked in distilled water, followed by freeze-drying.

Reduction was conducted in a quartz tube at 700 °C for 2 h under an H2/Ar (5:95 v/v) atmosphere with a heating rate of 2 °C min-1.

Characterization

The TEM and HRTEM measurement was performed on a Philips Tecnai F20 and Philips Tecnai F30 field emission transmission electron microscopes operated at 200 kV. SEM images were obtained from a JEOL 7800F scanning electron microscope operated at 5.0 kV. XRD patterns were collected on a German Bruker D8 Advanced X-Ray Diffractometer with Ni filtered Cu Kα radiation (40 kV, 30 mA). Thermo gravimetric analysis (TGA) was carried out on a TGA/DSC1 STARe System under air flow (25-900 ºC, 5 ºC/min). X-ray photoelectron spectra (XPS) were collected on a Kratos Axis ULTRA X-ray photoelectron spectrometer using a monochromatic Al Ka (1486.6 eV) X-ray source and a 165 mm hemispherical electron energy analyzer. Nitrogen physisorption isotherms were measured at 77 K on Tristar II 3020. All samples were degassed at 200 °C for 12 h prior to the measurements. The specific surface areas of the samples were calculated using the Brunauer–Emmett–Teller (BET) method and the total pore volumes were estimated from the nitrogen volumes adsorbed at the relative pressure of 0.99. The pore size distribution curves were derived from the Barrett-Joyner-Halenda (BJH) method of isotherm analysis. Raman spectra were collected on a Thermo-Fischer Almega dispersive Raman instrument. The instrument was fitted with both 633 and 785 nm lasers. Energy-dispersive X-ray spectroscopy analysis (EDX) was conducted on a JSM6610 EDS system.

Electrochemical measurements

The electrochemical properties of the prepared samples were evaluated with CR2032 coin cells. The electrodes were prepared by mixing the active materials with conductive carbon black (Super-P) and poly(vinylidene fluoride) (PVDF) dissolved in N-methyl-2-pyrrolidone (NMP) with a mass ratio of 80:10:10 to form a homogeneous slurry under magnetic stirring. The slurry was then spread onto a pure Cu foil. Pure lithium metal discs were selected as the counter electrode. The electrolyte used was 1 M LiPF6 in ethylene carbonate/dimethyl carbonate (1:1 v/v). A Celgard 2400 microporous polypropylene membrane was used as the separator. The CR2032 cells were

assembled in an argon-filled glove box (Mbraun Unilab) with water and oxygen contents less than 1 ppm. Cyclic voltammetry was carried out on a CHI 660D electrochemical workstation at a scan rate of 0.1 mV·s-1. The discharge and charge measurements of the batteries were performed on a Neware system in the fixed voltage window between 0.02 and 1.20 V at room temperature.

4.3 Results and discussion

The morphology and microstructure of the Si/RGO-AG composite were characterized using scanning electron microscopy (SEM), transmission electron microscopy (TEM) and high-resolution TEM (HRTEM) techniques. The HRTEM image in Figure 4.1 shows that there was a shell around the SiNPs with a thickness of about 5 nm. The X-ray photoelectron spectroscopy (XPS) results of 233 the SiNPs in Figure 4.2 shows a peak at 103 eV, which is attributed to Si 2p electrons in Si-O4, confirming the shell was silicon oxide (SiO2). This SiO2 layer plays a vital role in modifying the 24 surface chemistry of SiNPs. The silanol groups on the surface of the SiO2 shell can be ionized under the experimental conditions to carry a negative charge, thus enabling strong electrostatic interactions between the SiNPs and PDDA to change the SiNPs from a negatively charged surface 234 to a positively charged surface. On the other hand, according to a recent study, the SiO2 layer can suppress the volume expansion of SiNPs although it consumes some lithium ions during the first discharge cycle.

Figure 4.1 HRTEM image of SiNPs used in this work.

Figure 4.2 XPS survey (A) and Si 2p spectrum of SiNPs (B).

Figures 4.3A and 4.3B are the SEM images of Si/RGO-AG under different magnifications. SiNPs with an average diameter of 100 nm dispersed the 3D RGO aerogel framework can be clearly seen. On contrast, severe aggregation of SiNPs can be observed from pure SiNPs (Figure 4.4A), sample Si/RGO (Figure 4.4B) and sample Si/RGO-SWAG (Figure 4.4C, 4.4D). In particular, the SiNPs in sample Si/RGO formed micro-sized aggregates. After a single-step wrapping by RGO, the aggregation of SiNPs was not significant. However, some SiNPs were not fully wrapped by RGO and aggregates can be still observed. This explains the capacity fading of sample Si/RGO-SWAG. The elemental analysis results of Si/RGO-AG (Figure 4.5) confirmed the existence of major 60

elements Si and C and their homogeneous distribution in the 3D aerogel. In this Si/RGO-AG composite, the graphene aerogel provides a porous network for the entrapped SiNPs, thus is beneficial for accommodating the volume change of these SiNPs during electrochemical reactions. The porous network also facilitates electrolyte transport, potentially enhancing the rate capacity of LIBs. Besides, the continuous interconnected graphene network creates favourable electron pathways against cycling processes. The HRTEM image shown in Figure 4.1C also clearly demonstrates the existence of a continuous RGO network with SiNPs entrapped. The HRTEM image in Figure 4.1D shows that SiNPs were well-encapsulated within the RGO aerogel network

A B

C D RGO

SiNP

RGO Figure 4.3 SEM (A, B), TEM (C), and HRTEM images (D) of Si/RGO-AG composite.

Figure 4.4 SEM images of SiNPs (A) sample Si/RGO (B), which was prepared by using physical mixing method and sample Si/RGO-SWAG (C, D), which was prepared by single RGO wrapping.

Figure 4.5 EDS mapping of Si/RGO-AG.

Figure 4.6 shows the X-ray diffraction (XRD) patterns of Si/RGO-AG and SiNPs. All diffraction peaks due to SiNPs can be seen from sample Si/RGO-AG, indicating that the silicon crystalline structure in the Si/RGO-AG composite retained after the freeze-drying and thermal reduction treatments.

Figure 4.6 XRD patterns of SiNPs and Si/RGO-AG.

Figure 4.7 presents the Raman spectra of Si/RGO-AG and pristine SiNPs. The absorption bands due to SiNPs can be seen from Si/RGO-AG, confirming the presence of crystalline Si particles in the composite. There are two more absorption bands at 1350 and 1596 cm-1, respectively. These two peaks are assigned to the D band and G band of graphene, respectively, 235 confirming the presence of RGO in the composite.

Figure 4.8 presents the C1s XPS spectra of samples. As can be seen, GO showed three peaks at 284.9, 286.8, and 289.0 eV, corresponding to C=C in aromatic rings, C-O-C in epoxy and alkoxy, and C=O in carbonyl and carboxyl groups, respectively. It is clearly seen that the peak intensity due to C-O-C and C=O bonds of sample Si/RGO-AG is significantly lower than that of sample GO. The intensity of the peak ascribed to C=C however increased after reduction. These data suggest that the oxygen-containing groups on GO were largely removed, and most of the conjugated bonds were restored by GA reduction and thermal treatment.

Figure 4.7 Raman spectra of SiNPs and Si/RGO-AG.

The N2 adsorption/desorption isotherms of the Si/RGO-AG composite (see Figure 4.9A) showed a type IV isotherm with an H3 hysteresis loop, indicating a mesoporous structure of the material. The Brunauer-Emmett-Teller (BET) surface area and pore volume of the composite was measured to be 125 m2/g and 0.37 cm3/g, respectively, much higher than that of pure SiNPs (25 m2/g and 0.055 cm3/g, respectively, see Figure 4.9B). The Barrett-Joyner-Halenda (BJH) pore size distribution curve of Si/RGO-AG (see Figure 4.9C) demonstrates that Si/RGO-AG composite had a mesoporous structure with two pore systems. The intensive peak around 5.0 nm is due to the presence of small gaps among randomly stacked graphene sheets, while the broad peak centred at about 26 nm may be due to the large voids between inter-twisted graphene sheets. 236 The high surface area along with the existence of mesopores in Si/RGO-AG should offer a large material-electrolyte contact area and promote the diffusion of Li+ ions if Si/RGO-AG is used as electrode materials for lithium storage.

Thermal gravimetric analysis (TGA) was carried out in air and the results are shown in Figure 4.9D. It can be seen that the mass of the pristine SiNPs increased slightly with a mass gain of about 0.150 wt% at 700 °C. This gain was due to the oxidation of Si by O2 to form SiO2. At 700 °C, the residual of the RGO aerogel was about 3.02 wt%, which was due to the impurities in the GO sample. The weight loss of the Si/RGO-AG composite at 700 °C was about 19.3 wt%. Based on the TGA data, the Si/RGO-AG composite was calculated to contain about 80.0 wt% SiNPs and 20.0 wt % RGO aerogel.

Figure 4.8 C 1s XPS spectra of GO (A) and Si/RGO-AG (B).

Figure 4.9 N2 adsorption/desorption isotherms (A) and BJH pore size distribution based on desorption isotherm of Si/RGO-AG (B). N2 adsorption/desorption isotherm of SiNPs (C); TGA curves for SiNPs, RGO aerogel and Si/RGO-AG conducted in air (D).

The electrochemical performance of the Si/RGO-AG composite as an anode was assembled and tested in a CR2032 coin cell where lithium foil was used as a counter electrode. For comparison, the cycling performance of electrode Si/RGO was also tested under the same experimental conditions. Figure 4.10A shows typical cyclic voltammetry (CV) curves of electrode Si/RGO-AG in the potential range of 0.02-1.20 V (vs Li+/Li) at a scan rate of 0.1 mV·s-1 of the first two cycles, starting at the open circuit potential of 1.59 V. A broad cathodic peak in the first cycle appeared at 0.69 V, indicating the formation of solid electrolyte interphase (SEI).200 This cathodic peak disappeared in the second cycle and correlated to an initial capacity loss. The main cathodic part of the second cycle displayed a peak at 0.19 V, corresponding to the formation of Li-Si alloy phases. 237 The anodic part showed two peaks at 0.34 and 0.52 V, corresponding to the phase transition from Li-Si alloys to amorphous Si. 24, 200

Figure 4.10B displays the discharge/charge profiles of the initial three cycles of electrode Si/RGO-AG under a current density of 150 mA·g-1 in the voltage window of 0.02 to 1.2 V (vs. Li+/Li). The onset slope at about 0.7 V in the initial discharging curve, which disappeared in the following cycles, corresponds to the SEI formation 200. Besides, the main discharge plateau is around 0.2 V and the charge plateau is around 0.5 V. All these features are in good agreement with the CV results discussed above. The specific capacity was calculated based on the total mass of Si/RGO-AG. The initial discharge/charge capacities were 3446 and 2535 mAh·g-1, respectively, to give an initial Coulombic efficiency of 73.6 %. The initial irreversible capacity of electrode Si/RGO-AG can be attributed to the formation SEI, the unexpected reaction of the remaining oxygen-containing groups in the RGO aerogel and SiO2 layer on the surface of SiNPs with Li ions. 27, 195 After the second cycle, the Coulombic efficiency tended to increase and stabilize. It is interesting to note that the curves of the second and third cycles almost overlapped each other, which indicates a good cycling stability of the electrode. The reversible capacities of the second and third cycles compared with the first cycle were slightly increased, which can be attributed to the activation of the SiNPs in the Si/RGO-AG composite.

Figure 4.10C shows the cycling performance of Si/RGO-AG at a current density of 150 mA·g−1, together with electrodes Si/RGO-SWAG and Si/RGO. It can be seen that the initial charge and discharge capacities of electrode Si/RGO were the lowest among the three electrodes studied. This poor performance of electrode Si/RGO was probably due to severe aggregation of the SiNPs, indicating the RGO sheets did not stabilize the SiNPs well. After about 20 cycles, the reversible capacity dropped drastically to about 450 mAh·g-1, confirming that the SiNPs were not stabilized well by the RGO sheets. These results suggest that the physical mixing method is not a good approach to stabilizing SiNPs. While the initial discharge/charge capacities of electrode Si/RGO-SWAG reached to 3360 and 2450 mAh·g-1, respectively, its reversible capacity decreased to about 615 mAh·g-1 after 40 cycles. The improved performance of electrode Si/RGO-SWAG indicates the single-step wrapping method illustrated in Scheme 4.1 is advantageous over the physical mixing method. This fast capacity fading however indicates that the single-step wrapping method still could not afford efficient stabilization of SiNPs. On contrast, the Si/RGO-AG electrode showed a significantly improved cycling performance with the highest initial charge/discharge capacities and delivered a reversible capacity of 1984 mAh·g-1 after 40 cycles.

Figure 4.10D demonstrates the rate capability of the Si/RGO-AG at current densities ranging from 170 mA·g-1 to 2500 mA·g-1. The battery delivered a reversible capacity of about 1000 and 600 mAh·g-1 at the current densities of 2000 and 2500 mA·g-1, respectively. Furthermore, the capacity

reached around 2000 mAh·g-1 when the current density was decreased to 170 mA·g-1 after having been cycled at higher current densities, indicating a good cycling stability of Si/RGO-AG.

A Delithiathion B

1st cycle

2nd cycle 1st cycle

Lithiathion 2nd and 3rd cycles

C D

Si/RGO-AG

Si/RGO-SWAG

Si/RGO

Figure 4.11 shows the SEM image of electrode Si/RGO-AG after 40 cycles. As can be seen, the Si/RGO-AG composite maintained its integrity and porous structure. In addition, the SiNPs were entrapped by the RGO framework, contributing to the significantly improved cycling performance of Si/RGO-AG.

Figure 4.11 SEM image of Si/RGO-AG after 40 cycles.

The Nyquist plots of the Si/RGO-AG electrode are presented in Figure 4.12. The depressed semicircle in the high-frequency region represents the resistance of the SEI film and the charge-transfer resistance, while the straight lines in the low-frequency region corresponds to the diffusion kinetics of lithium ions. No obvious impedance increase is observed after cycling due to the stable SEI. The impedance decrease may be attributed to the gradual electrolyte transport into the electrode and the increasing conductivity of the SiNPs after lithiation.

The improved cycle performance and enhanced rate capability of Si/RGO-AG can be attributed to the following reasons: i) the Si/RGO-AG composite created sufficient space and efficiently accommodated the drastic volume change of the entrapped SiNPs during cycling; ii) the interconnected 3D RGO aerogel network maintained the integrity of the electrode structure, prohibited the detachment of the SiNPs from the current collector and improved the electrical conductivity of the electrode; iii) the existence of meso- and macro- pores provided an efficient pathway for electrolyte transport and facilitated Li+ diffusion, thus enhancing the rate performance.

Figure 4.12 Electrochemical impedance spectroscopy (EIS) curves of Si/RGO-AG.

4.4 Conclusions

In conclusion, we have demonstrated an approach to preparing a high-performance LIB negative electrode material by encapsulating SiNPs in 3D RGO aerogel. The RGO aerogel with a porous network offers space for accommodating SiNPs, as well as facilitating electron and electrolyte transport. The composite electrode delivered a stable cycling performance with a capacity of about 2000 mA h g-1 after 40 cycles, together with an excellent rate capability.

Chapter 5 Lithium-storage Properties of Reduced Graphene Oxide Using Gallic Acid as Reducing Agent

5.1 Introduction

Lithium-ion batteries (LIBs) are one of the most promising energy-storage systems. Graphite, a conventional and commercialized anode material for LIBs has a theoretical capacity of about 372 mAh·g-1, limiting the development of advanced LIBs for emerging technologies, such as smart grids, electric vehicles, and solar energy. Graphene-based materials have been shown to have tremendous potentials for a myriad of applications, including energy storage.28 In particular, graphene as an anode for LIB can store more lithium ions than graphite because lithium ions can not only be stored on both sides of graphene sheets, but also on the edges and covalent sites.15 To move the graphene-based energy storage technology further, a scalable and ecofriendly technique to produce large-quantity graphene is paramountly important.

Chemical reduction of graphene oxide (GO) is regarded as a feasible method for mass-production of reduced graphene oxide (RGO). The most commonly used reductants are sodium borohydride and hydrazine,29, 31, 238 which are toxic. Thus, alternative reducing methods are desirable. Zhao and co-workers239 have demonstrated that urea is an effective reducing agent in removing oxygen-containing groups from graphite oxide for restoring the conjugated electronic structure of graphene. Other ecofriendly reducing agents, such as L-ascorbic acid and tea polyphenol, have also been proved to reduce GO successfully.240, 241, 242

Gallic acid (GA), a natural organic acid with three adjacent hydroxyl groups located at its only benzene ring, has recently been shown to possess considerable ability to reduce GO.243 It has also been observed that phenolic compounds including GA can function as a stabilizer to prevent RGO sheets from aggregation.243, 244, 245 In the present study, the electrochemical properties of GO samples reduced by GA as anode material for LIB were investigated for the first time, and compared with previously reported RGO-based anodes. The RGO anode material prepared in this work showed a significantly improved electrochemical performance with a reversible capacity of around 1280 mAh·g-1 after 350 cycles at a current density of 200 mA·g-1 and good rate capability. The improvements are attributed to the fairly complete reduction of GO using GA, the stabilization function of GA prohibited RGO sheets from aggregation thus created micron-sized disordered porous RGO architecture, the nano-cavities between RGO sheets and edges of RGO sheets contributed superior lithium storage.

5.2 Experimental

Synthesis of Graphene Oxide (GO)

GO was synthesized from natural graphite flake (Sigma Aldrich, 325 mesh) using a modified Hummers method.232.

Preparation of RGO

GA (13 mg) was dissolved in 20 mL of the above-prepared GO suspension (2 mg·mL-1) under sonication. The mixed suspension was transferred into a 20 mL capped vial and placed into an oil bath at 95 °C under magnetic stirring for 12 h. The resulting black solids were collected after removing the upper liquid phase. The collected granules were soaked in distilled water 7 times for 7 days and then frozen quickly with liquid nitrogen and lyophilized, followed by heat treatment in a -1 quartz tube at 700 °C for 2 h under a H2/Ar (5/95, v/v) with a heating rate of 5 °C min to eventually obtain Si@RGO composite.

Characterization

TEM measurement was performed on a Philips Tecnai F30 field emission transmission electron microscope operated at 200 kV. SEM images were obtained from a JEOL 7001 scanning electron microscope operated at 5.0 kV. XRD patterns were collected on a German Bruker D8 Advanced X-Ray Diffractometer with Ni filtered Cu Kα radiation (40 kV, 30 mA). X-ray photoelectron spectra (XPS) were collected on a Kratos Axis ULTRA X-ray photoelectron spectrometer using a monochromatic Al Ka (1486.6 eV) X-ray source and a 165 mm hemispherical electron energy analyzer.

Electrochemical measurements

The electrochemical properties of the prepared samples were evaluated using CR2032 coin cells. The RGO electrodes were prepared by mixing the RGO with conductive carbon black (Super-P) and poly(vinylidene fluoride) (PVDF) dissolved in N-methyl-2-pyrrolidone (NMP) with a mass ratio of 80:10:10 to form a homogeneous slurry under magnetic stirring. The slurry was then spread onto a pure Cu foil. Pure lithium metal foil was used as the counter electrode. The electrolyte was 1

M LiPF6 in ethylene carbonate/dimethyl carbonate (1:1, v/v). A Celgard 2400 microporous polypropylene membrane was used as the separator. The CR2032 cell was assembled in an

argon-filled glove box (Mbraun Unilab) with water and oxygen contents less than 0.1 ppm. Cyclic voltammetry (CV) measurements were carried out on a CHI 660D electrochemical workstation at a scan rate of 0.1 mV·s-1. Electrochemical impedance spectral measurements were also performed using the CHI 660D electrochemical workstation over the frequency range 100 kHz to 10 mHz, and the signal amplitude was 5 mV. The discharge and charge measurements of the batteries were performed on a Neware system in the fixed voltage window between 0.005 and 3.5 V for RGO anode.

5.3 Results and Discussion

Oxidative breakup of graphite was done to produce GO and GO dispersions were deoxygenated using GA in a fairly straightforward process. The morphology and microstructures of the RGO sample were obtained using field-emission scanning electron microscopy (FESEM) and transmission electron microscopy (TEM).

Figures 5.1A and 5.1B show the FESEM images of sample RGO at different magnifications. Porous micron-sized RGO particles with similar sizes can be seen. Form Figure 5.1B corrugated RGO architecture can be seen. No significant stacking of RGO sheets was observed.

Figure 5.2A shows that XPS results of both the GO and RGO samples. A drastic decrease in the oxygen content after reduction can be seen. It was estimated that the atom percentage of oxygen in the RGO sample was about 2.0%, less than that in the GO sample (31.1%). Apart from the surface adsorbed oxygen, it is interesting to note that the RGO prepared using GA as the reducing agent contained minor oxygen content in comparison with other reported reducing agents, such as sodium borohydride (15.9%),29 hydrazine hydrate (6.2%),246 and other reduction methods namely microwave digestion (10.5%),143 microwave radiation then hydrazine hydrate reduction (4.2%),142 and mild thermal reduction (17.5%).247 The high-resolution C1s XPS spectra of GO and RGO are shown in Figure 5.2B. The C1s spectrum of GO was deconvoluted with five peak components with binding energies of 284.6, 285.1, 286.9 287.7 and 288.9 eV, corresponding to C=C (sp2), C–C (sp3), C–O (hydroxyl and epoxy), C=O (carbonyl), O–C=O (carboxyl), respectively.243, 244, 248 Meanwhile, for the RGO sample, no obvious peaks belonging to oxygen containing groups were observed, suggesting a fairly complete reduction of GO. Only one prominent C1s peak was observed at 284.8 suggesting the sp2 graphitic nature of the sample.

Figure 5.1 (A-B) FESEM images of RGO under different magnifications.

Figure 5.2 (A) Survey XPS patterns of GO and RGO; (B) C1s XPS patterns of RGO and GO.

Figure 5.3 shows the X-ray diffraction (XRD) patterns of samples. For the GO sample, a wide diffraction peak catered at 9.8° corresponding to an interlayer spacing of 0.901 nm can be seen. The graphite sample displayed a sharp peak at 26.5° (corresponding to an interlayer spacing of 0.335 nm). The XRD results indicated the presence of functional groups and the disordered layers after oxidization and exfoliation. After reduction, a new broad peak appeared at around 26.1° for RGO, corresponding to an interlayer spacing of 0.340 nm fairly close to that of graphite. The new broad peak suggests the removal of oxygen containing functionalities and the partially restoration of the conjugate sp2 networks in the porous RGO network during the reduction of GO.152, 245

Figure 5.3 XRD patterns of RGO, GO and graphite.

Figure 5.4 shows the Raman spectra of samples. The pristine graphite displayed a prominent G peak (corresponding to sp2 hybridizing carbon) at 1581 cm−1, as well as a very small D peak (originating from disordered carbon) at 1338 cm−1. The G band of sample GO was broadened and shifted to 1598 cm−1. The D band at 1342 cm−1 became more prominent, indicating the reduction in size of the in-plane sp2 domains with more defects resulting from the extensive oxidation, thus generating a large number of oxygen-containing groups. While the Raman spectrum of the RGO sample also contained both G and D bands, the peak intensity ratio of the D band over the G band (ID/IG=1.12) is higher than that of GO (ID/IG=0.993). This change in intensity ratio is due to the decrease in the dimension of sp2 domains during reduction and increased graphitization.248, 249

Figure 5.4 Raman spectra of RGO, GO, and graphite.

The electrochemical performances of the prepared RGO were tested in CR2032 coin cells. Lithium foil was used as both counter and reference electrode. For comparison, graphite was tested under the same experimental conditions as RGO. The specific capacity was calculated based on the mass of RGO and graphite only, respectively.

Figure 5.5A shows the first three cyclic voltammetry (CV) curves of the RGO electrode at room temperature in the potential range of 0.005 – 3.500 V (versus Li+/Li) at a scan rate of 0.1 mV·s-1. The CV curve of the first cycle was different from those of the subsequent cycles, especially for the cathodic curve, a strong peak is observed at 0.7 V, which was resulted from the irreversible side reactions on the electrode surfaces and interfaces due to the solid electrolyte interphase (SEI) film. From the second cycle on, it is important to note that the CV curves almost overlapped, indicating the stable and superior reversibility of RGO anode.

The discharge-charge profiles of RGO anode in the first 3 cycles are shown in Figure 5.5B. The RGO anode delivered a specific capacity of 1762 mAh·g-1 in the initial discharging and a reversible capacity of 844 mAh·g-1 in the first charging. The irreversible capacity could be associated with the formation of the SEI layer in the first cycle. From the second cycle, the curves of RGO anode showed no distinguishable plateaus, which can be attributed to the smaller crystallite structure, high specific surface area and disorganized RGO architecture.139 It has been proposed that lithium ions 78

can be adsorbed on both sides of the graphene sheet that arranged like a ‘‘house of cards’’ in hard carbons, leading to two layers of lithium for each graphene sheet, with a theoretical capacity of 744 -1 mAh·g through the formation of Li2C6. On the other hand, nano-cavities between RGO sheets due to scrolling and crumpling could also contribute to the lithium storage.15

The cycling performances of RGO and graphite electrodes were examined at a current density of 200 mA·g-1. In Figure 5.5C, the RGO electrode delivered a reversible initial discharging capacity of 844 mAh·g-1 and its reversible capacity maintained a stable rising trend from the 15th cycle for RGO. After 200 cycles, the reversible capacity still kept a rising trend overall but more fluctuant. At the 350th cycle, the RGO anode delivered a reversible capacity of 1280 mAh·g-1, which represented a much superior performance than that of graphite anode. The possible reason for the fluctuant rising trend of its reversible capacity is that the gradual activation process of lithium insertion active sites, such as edge-type sites, nanopores due to the permeation of the electrolyte during repetitive charging and discharging process. The stable SEI layer formed during the first cycle also prohibited further side reactions.

Figure 5.5D exhibits the rate capability of RGO anode at the current densities from 100 to 2000 mA·g-1. It could deliver reversible capacities of about 900, 950, 500, 450, 400 mAh·g-1 at the current densities of 100, 200, 500, 1000 and 1500 mA·g-1, respectively. Even at 2000 mA·g-1, it could still deliver a capacity of more than 350 mAh·g-1. After cycled at higher current densities, it could still maintain the reversible capacities when the current densities were lower down. It is reasonable to believe that more edge sites on the RGO are beneficial to the improved capacity at high current rates, and the porous structure is able to reduce diffusion length of lithium ions, resulting in better rate capability and cycling performance. In addition, the strongly crumpled and scrolled RGO sheet could also contribute to the structural buffer for the volume expansion during the Li insertion/extraction.

Summarizing the above results of RGO anode, it is concluded that micron-sized RGO particles reduced by GA delivered a reversible capacity of around 1280 mAh·g-1 after 350 cycles at the current density of 200 mA·g-1 as well as good rate capability. The RGO showed superior battery performances than the recent reported ones: highest specific capacity of 817 mAh·g-1 for over 100 cycles using hydriodic acid (HI) as reductant250; reversible capacity of 800 mAh·g-1 after 100 cycles at a current density of 200 mA·g-1 using thermal annealing procedure158; reversible capacity of -1 -1 about 600 mAh·g after 100 cycles at a current density of 149 mA·g using hydrazine (N2H4) as reductant251; first charge specific capacity of 933 mAh·g-1 at 50 mA·g-1 by a metal etching method157; reversible capacity of about 620 mAh·g-1 at 100 mA·g-1 after 100 cycles for a core–shell structured porous carbon–graphene composite159; first discharge capacity of 1069.3 mAh·g-1 at 100 mA·g-1, 733.3 mAh·g-1 after 100 cycles by a microwave digestion reduced graphene oxide method143.

Table 5.1 compares the electrochemical performance of RGO materials prepared by using different reducing methods. It can be seen that micron-sized RGO reduced by GA delivered not only high reversible capacity but also exhibited long cycle life. It is worth pointing out that the RGO prepared using GA as a reductant is among the few ecofriendly chemical reducing methods to prepare RGO anode for lithium ion storage without doping or adding extra additives.

Table 5.1 Comparison of graphene anode materials prepared by reducing GO using various methods

Reversible Capacity Current rate Reducing method capacity Cycle life retention (mA·g-1) (mAh·g-1) (%) Solvothermal treatment and 800 100 81.6 200 thermal reduction158

Hydrothermal treatment and 1053 130  186 thermal reduction147

Thermal reduction252 760 35  100

Thermal reduction145 ~634 300  1000

Thermal reduction144 576 50  50

Carbothermal reduction253 ~576 20  100

Microwave digestion ~733 100  100 reduction143

Microwave-induced thermal 1002 50 93 40 reduction146

Infrared irradiation induced 755 50 ~60 100 photothermal reduction254

Supercritical isopropanol 1331 100  50 reduction255

Supercritical alcohol 652 40  50 reduction150

Ascorbic reduction152 420 100 97 2000

Ethylenediamine (EDA) 346 60  350 reduction151

Hydrazine reduction251 ~600 50 ~48 149

Hydrazine reduction148 460 100 71 

GA reduction (This work) 1280 350  200

5.4 Conclusions

In summary, micron-sized RGO particles prepared by reducing GO using GA displayed a good electrochemical performance when used as anode for LIBs with a reversible capacity of around 1280 mAh·g-1 after 350 cycles at a current density of 200 mA·g-1 as well as a good rate capability, attributing to the fairly complete reduction of GO by GA, existence of the edge sites from the defects of RGO, and the porous network structure formed.

Chapter 6 Silicon Nanoparticles Encapsulated in Reduced Graphene Oxide Framework as a Stable Anode for Lithium Ion Storage

6.1 Introduction

The development of lithium ion batteries (LIB) with both high energy density and power density is of great importance to meet the ever-growing demands for high power sources. Achieving high capacity, excellent cycling performance and rate capability for both anode and cathode of a LIB is the main challenge for the improvement of the electrochemical performance of this energy storage device.32 Compared with carbon materials, especially graphite as anode for LIBs, silicon (Si) is regarded as a promising anode material because of its large theoretical capacity for Li ion storage (4200 mAh·g-1) and low discharge potential. However, the severe volume change (> 400%) of bulk Si leads to rapid capacity fading due to pulverization during repetitive lithiation and delithiation processes. The rate capability of Si is limited by its poor intrinsic electric conductivity.10, 11 Two approaches are usually taken to improve its electrode performance. One approach is to shorten the electronic and ionic transport length by fabricating nanostructured Si, but SiNPs trend to aggregate during treatment. The other approach is to increase the electron transport by conductive carbon coating or adding conductive additives such as carbon nanotubes, graphene.13 Especially graphene or RGO, with a few-layer two-dimensional honeycomb carbon structure, has great advantages for its high electron mobility, huge specific surface area and superior electrical conductivity when prepared meticulously.15 Although improvements have been achieved using graphene sheets to stabilize SiNPs16, 17, 19, 20, 21, 22, 23, 209, 211, 256, 257, it is still highly needed to explore a cost-efficient and green way to prepare Si/graphene composite.

Recently, it has been reported that sodium alginate (SA) can produce an improved Si anode performance because SA contains a large amount of carboxylic groups, leading to a great number of possible binder-bonds for electrode materials. It also has excellent mechanical properties and no interaction in electrolyte solvents, showing 50 times higher stiffness than PVDF.258, 259

In this work, a new approach to encapsulating SiNPs within reduced graphene oxide (RGO) for stabilizing the SiNPs is described. The RGO was used to stabilize SiNPs to prepare Si@RGO composite materials with different Si/GO mass ratios. It was found that the Si@RGO composite with Si/GO mass ratio of 1:2 (Si@RGO 1:2) delivered reversible capacities of 1074 mAh·g-1 at the 5th cycle and 900 mAh·g-1 at the 100th cycle with a capacity retention of about 84% as well as excellent rate capability. Well-dispersed SiNPs in numerous cluster forms were stabilized between the RGO architecture. The continuous RGO sheets not only enhanced the electronic conductivity of

the electrode, but also created enough space to facilitate electrolyte transport. Results also showed that sodium alginate was a good binder for the Si@RGO composite.

6.2 Experimental

Synthesis of Graphene Oxide (GO)

GO was synthesized from natural graphite flake (Sigma Aldrich, 325 mesh) using a modified Hummers method.232.

Preparation of Si@RGO composite

SiNPs were first modified by using PDDA. Briefly, 120 mg of SiNPs (Aldrich, 40-100 nm) was added to 120 mL de-ionised water, followed by adding 3 g of a 20 wt% PDDA (Sigma Aldrich, MW = 10000-20000) solution under sonication in a water bath (KQ3200DE, 40 kHz). Excess PDDA was washed away using deionized water four times. The PDDA-modified SiNPs (20, 40, 80 mg) was added to 20 mL of the GO suspension (2 mg·mL-1) under sonication for 1 h. After the addition of GA (13 mg), the mixed suspension was transferred into a 20 mL capped vial and placed into an oil bath at 95 °C under magnetic stirring for 12 h. The following treatments were the same with those of RGO above.

Characterization

TEM measurement was performed on a Philips Tecnai F30 field emission transmission electron microscope operated at 200 kV. FESEM images were obtained from a JEOL 7001 scanning electron microscope operated at 5.0 kV. XRD patterns were collected on a German Bruker D8 Advanced X-Ray Diffractometer with Ni filtered Cu Kα radiation (40 kV, 30 mA). X-ray photoelectron spectra (XPS) were collected on a Kratos Axis ULTRA X-ray photoelectron spectrometer using a monochromatic Al Ka (1486.6 eV) X-ray source and a 165 mm hemispherical electron energy analyzer. Raman spectra were collected on a Thermo-Fischer Almega dispersive Raman instrument. The instrument was fitted with both 633 and 785 nm lasers.

Electrochemical measurements

The electrochemical properties of the prepared samples were evaluated using CR2032 coin cells. For PVDF binder, the Si@RGO electrodes were prepared by mixing the active materials with conductive carbon black (Super-P) and PVDF dissolved in N-methyl-2-pyrrolidone (NMP) with a mass ratio of 80:10:10 to form a homogeneous slurry under magnetic stirring; For CMC and SA binders, the Si@RGO electrodes were prepared by directly mixing Si@RGO composites with carbon black (Super-P) and CMC or SA dissolved in de-ionised water with a mass ratio of 80:10:10 to form a slurry under magnetic stirring. The above slurries were then spread onto a Cu foil using doctor’s blade method. Pure lithium metal foil was used as both counter and reference electrode.

The electrolyte was 1 M LiPF6 in ethylene carbonate/dimethyl carbonate (1:1, v/v). A Celgard 2400 microporous polypropylene membrane was used as the separator. The CR2032 cell was assembled in an argon-filled glove box (Mbraun Unilab) with water and oxygen contents less than 0.1 ppm. Cyclic voltammetry measurements were carried out on a CHI 660D electrochemical workstation at a scan rate of 0.1 mV·s-1. The discharge and charge measurements of the batteries were performed on a Neware system in the fixed voltage window between 0.02 and 1.20 V for Si@RGO composites at room temperature.

6.3 Results and Discussion

Thus the porous RGO architecture was also utilized to stabilize SiNPs as shown in Scheme 6.1. To prepare Si@RGO composite, SiNPs were firstly modified using poly(diallydimethylammonium chloride) (PDDA, a positively charged polyelectrolyte) to change the surface charge nature from being negative to being positive. Note, both GO and SiNPs (there is a thin layer of SiO2 about 5 nm on the surface of the SiNPs in our reported research.257) are negatively charged in a wide pH range. The positively charged PDDA-modified SiNPs interacted strongly with negatively charged GO sheets by electrostatic force. After 12 h reaction at 95 °C under magnetic stirring, homogeneous yellow suspension turned black. Upon sedimentation for a short while, black solids and colourless transparent liquid phase were seen, indicating that the SiNPs had been combined with RGO sheets to form a composite. After freeze-drying and thermal treatment, Si@RGO composite was obtained.

Scheme 6.1 Schematic of the preparation procedure of Si@RGO composite: (1) surface modification of SiNPs with poly(diallydimethylammonium chloride, PDDA), (2) adding GO to form a Si@GO composite via electrostatic interactions, (3) reduction of the Si@GO to obtain a Si@RGO composite, (4) freeze drying and thermal treatment to obtain the final product.

Characterization results of Si@RGO 1:1 composite (Mass ratio of Si/GO = 1:1) are shown in Figure 6.1 It can be seen (Figure 6.1A) that SiNP clusters with different sizes were well dispersed among the graphene layers without obvious aggregation, leaving enough space between each cluster. The survey XPS patterns of sample Si@RGO in Figure 6.2A confirmed that the sample contained only

Si, C and O elements. The N2 adsorption/desorption isotherms of the Si@RGO sample are shown in Figure 6.2B. The Brunauer-Emmett-Teller (BET) surface area and pore volume of the composite are measured to be 75 m2/g and 0.46 cm3/g, respectively. The RGO wrapping layers with crumpled and rough textures on the surface were associated with the flexible and corrugated nature of GO sheets as well as the encapsulated SiNPs. Particularly, in Figure 6.1B, it could be clearly seen SiNP clusters were perfectly wrapped by the continuous corrugated RGO layer without any exposure to outside. The continuous RGO wrapping layer is beneficial for accommodating the volume change of the inside wrapped SiNPs during electrochemical reactions. The space between SiNP clusters also facilitates electrolyte transport, potentially enhancing the rate capacity of LIBs. Besides, the continuous interconnected graphene network creates favourable electron pathways during cycling processes. Figure 6.1C is the cross-sectional view of the sample Si@RGO. It could be clearly

observed the sample has a rough surface and SiNP clusters were encapsulated inside the RGO wrapping layers homogeneously, indicating the success of prohibiting the aggregation of SiNPs. The thickness of the Si@RGO particles was about 1 μm. Such thin particle morphology is favourable for electrolyte ions to go through, thus enhancing rate capacity. The top-view FESEM images of other samples are shown in Figure 6.3

Figure 6.1D shows the TEM image of sample Si@RGO. SiNPs wrapped by thin RGO sheets can be clearly seen. The similar size of SiNPs (~100 nm) before and after wrapping by RGO sheets indicated that the Si particles did not aggregate during the sample processing process. In addition, the formation of porous RGO network is of importance for the electron transport. The porous structure is also ideal for accommodating the volume change of the SiNPs during electrochemical reactions.

Figure 6.2 Survey XPS patterns of sample Si@RGO 1:1 (A) and N2 adsorption/desorption isotherms of sample Si@RGO 1:1.

Figure 6.3 FESEM images of RGO (A) and Si@RGO composites with different Si/GO ratios (1:2 B, 1:1 C, 2:1 D).

Figure 6.4A and 6.4B shows the XRD patterns and Raman spectra of Si@RGO, SiNPs and RGO, respectively. All the major XRD and Raman peaks of Si@RGO are overlapped with those of the pristine SiNPs and RGO, indicating that the silicon crystalline structure in the Si@RGO composite are still retained after the freeze-drying and thermal reduction treatment.

Figure 6.4 (A) XRD patterns of Si@RGO, SiNPs and RGO; (B) Raman spectra of Si@RGO, SiNPs and RGO.

The electrochemical performances of Si@RGO composites using traditional poly(vinylidene fluoride) (PVDF) binder to fabricate electrode are shown in Figure 6.5. As is widely understood, SiNP has the highest Li+ storage capacity but suffers from rapid capacity fading. On the other hand, porous RGO architecture could efficiently stabilize SiNPs and improve the cycling stability, but the overall reversible capacity of Si@RGO would drop with the increasing RGO mass ratio. Here, three

samples were prepared and tested with different Si/GO mass ratios, 1:2, 1:1 and 2:1 under the same testing protocols. The specific capacity was calculated based on the total mass of Si and RGO.

Figure 6.5A shows the CV curves of the Si@RGO 1:2 in the potential range of 0.02 – 1.20 V (versus Li+/Li) at a scan rate of 0.1 mV·s-1. In the first cycle, a broad cathodic peak due to the formation of SEI was observed at 0.69 V. The absence of the cathodic peak in the subsequent cycles indicates a complete formation of a stable SEI layer in the first cycle, which is beneficial for the Si@RGO anode in the following cycles. The anodic part showed two peaks at 0.34 and 0.52 V, corresponding to the phase transition of Li–Si alloys to amorphous Si, in agreement with previous works.237, 260

Figure 6.5B displays the discharge–charge profiles of the Si@RGO 1:2 composite at a current density of 200 mA·g-1 in the voltage window of 0.02 to 1.20 V vs. Li+/Li. The onset slope at about 0.7 V in the initial discharging curve, which disappeared in the following cycles, corresponded to the SEI formation. Besides, the main discharge plateau was around 0.2 V and charging plateau was around 0.5 V. All these features were in good agreement with the CV curve. The initial discharge/charge capacities were 898 and 1608 mAh·g-1, respectively. As a result, the initial Coulombic efficiency was about 55.8 %. The observed initial irreversible capacity of the Si@RGO anode is attributed to the SEI formation and the existence of silicon dioxide on the surface of the SiNPs, both of which consume Li+ ions.27, 195 From the second cycle, the Coulombic efficiency gradually increased and became stable and stable.

Figure 6.5 (A) Cyclic voltammetry curves for Si@RGO 1:2 composite; (B) Galvanostatic discharge-charge profiles of the first three cycles.

Figure 6.6A shows the cycling performance of samples Si@RGO 1:2, 1:1 and 2:1 at a current density 200 mA·g-1 for 100 cycles. Overall, the reversible capacities of all the three samples kept a descending trend. Obviously the Si@RGO 1:2 delivered the most stable capacity maintenance throughout the 100 cycles, 1074 mAh·g-1 at the 5th cycle and 900 mAh·g-1 at the 100th cycle with a capacity retention of about 84%. For sample Si@RGO 1:1, the highest reversible capacity was 1270 mAh·g-1 observed at the 5th cycle. At the 40th cycle, it could still deliver a reversible capacity of 1094 mAh·g-1 with a capacity retention of about 86 %. After that, fast capacity fading was observed and the reversible capacity could only maintain 526 mAh·g-1 at the 100th cycle with a capacity retention of about 41 %. For sample Si@RGO 2:1, the highest reversible capacity was 1700 mAh·g-1 observed at the 5th cycle. At the 30th cycle, it delivered a reversible capacity of 1477 mAh·g-1 with a capacity retention of about 87 %. More severe capacity fading was observed since then and the reversible capacity could only maintain 466 mAh·g-1 at the 100th cycle with a capacity retention of about 27 %.

Figure 6.6B demonstrates the rate capabilities of samples Si@RGO 1:2, 1:1 and 2:1 with the current density ranging from 100 to 2000 mA·g-1. The Si@RGO 1:2 battery delivered a reversible capacity of about 800 and 650 mAh·g-1, when the current density was increased to 1000 and 2000 mA·g-1, respectively. On the contrast, although higher starting reversible capacities were seen for samples Si@RGO 1:1 and 2:1 batteries, they delivered only about 520 and 420, 450 and 270 mAh·g-1, respectively when the current density was increased to 1000 and 2000 mA·g-1. Particularly, when the current density was lowered to 100 mA·g-1 after cycled at high ones, for sample Si@RGO 1:2 the reversible capacity rose up to around 900 mAh·g-1 compared with the highest capacity of 1019 mAh·g-1 at the 6th cycle with a capacity retention of 88 %, indicating a good stability of the electrode in a wide range of discharge/charge currents. On the contrast, for Si@RGO 1:1 and 2:1 the values were 1000/1313, 76 % and 1250/1779, 70 %, respectively.

Summarizing the above results of Si@RGO composite, it is concluded that higher initial capacity can be obtained with more SiNP mass ratio; more stable capacity maintenance as well as better rate capability can be achieved with more RGO mass ratio. Among the three prepared Si@RGO composites, Si@RGO 1:2 delivered both stable cycling performance and excellent rate capability. Composite Si@RGO 1:2 showed comparable battery performance with the latest reported ones18, 19, 20, 23, 261, but the preparation rout is more efficient and environmental-friendly.

It is reported binder play an important role for the electrochemical performance of Si based anode materials.183, 259, 262 To improve the electrode performance of the samples Si@RGO 1:1 and 2:1, carboxyl methyl cellulose (CMC) and sodium alginate (SA) were used to prepare electrode and the electrode performances were tested for the three binders, PVDF, CMC and SA.

Figure 6.7A and 6.7B show the cycling performances of Si@RGO 2:1 and 1:1 using different binders for 100 cycles at the current rate of 200 mA·g-1, respectively. In Figure 5A, relatively stable cycling performance could maintain for about 30 cycles for the electrodes using PVDF and CMC binders, then rapid capacity decay were observed. The capacity decay rate of the electrode using PVDF was more rapid than the one using CMC. For the electrode using SA as binder, relatively stable cycling performance could maintain for more than 50 cycles with a capacity retention about 86 % (1420/1652). Although capacity fading could be observed after that, it could deliver the highest reversible capacity among the three electrodes after 100 cycles. In Figure 5B, relatively stable cycling performance could maintain for about 40 cycles for the electrodes using PVDF and CMC binders, then rapid capacity decay were observed. Still the capacity decay rate of the electrode using PVDF was more rapid than the one using CMC. For the electrode using SA as binder, relatively stable cycling performance could maintain for about 70 cycles with a capacity retention about 86 % (1024/1186).

Figure 6.7 (A) cycling performances of Si@RGO 2:1 composite using different binders; (B) cycling performances of Si@RGO 1:1 composite using different binders.

6.4 Conclusions

In conclusion, SiNPs stabilised by GA-reduced RGO sheets exhibited capacities of 1074 mAh·g-1 at the 5th cycle and 900 mAh·g-1 at the 100th cycle, as well as excellent rate capability. A good coverage of SiNPs by RGO sheets can not only buffer the volume change of SiNPs during charge/discharge and prevent the SiNPs from aggregating, but also improve the overall ionic and electrical conductivities of the electrode. Sodium alginate was found to be an ideal binder for the Si@RGO based composite materials.

Chapter 7 Conclusions and Recommendations for Future Work

7.1 Conclusions

The objective of this project is set to design and prepare Si-based composite electrode for LIBs with reduced graphene oxide (RGO), by accommodating the volume changes, improving the electronic conductivity and preventing aggregation of Si nanoparticles (SiNPs). Highlights of the thesis include: a novel approach to wrap SiNPs in a dual RGO aerogel framework as anode for LIBs; a simple, ecofriendly and scalable chemical reducing method to prepare RGO anode for LIBs using gallic acid as reducing agent; a sandwich-typed micron-sized RGO-stabilized SiNPs composite anode for LIBs prepared in a simple way under mild condition. These studies introduced potential RGO and Si-based anodes for high-performance LIBs.

Stabilization of SiNPs in a 3D Aerogel-typed RGO Framework: A high-performance LIB negative electrode material was prepared by encapsulating SiNPs in 3D porous RGO aerogel. The RGO aerogel with a porous network (125 m2/g, BET surface area) offered space for accommodating SiNPs, as well as facilitating electron and electrolyte transport. The composite electrode delivered a stable cycling performance with a high capacity of about 2000 mA h g-1 after 40 cycles, together with an excellent rate capability.

A Porous RGO Architecture Prepared by a Chemical Reducing Method: Micron-sized porous RGO particles prepared by reducing GO using gallic acid displayed a good electrochemical performance when used as anode for LIBs. The RGO sample delivered a reversible capacity of around 1280 mAh·g-1 after 350 cycles at a current density of 200 mA·g-1 as well as a good rate capability over 800 cycles, superior than most of the reported results. The inspiring electrochemical performance of the RGO sample was attributed to the fairly complete reduction of GO by GA, existence of the edge sites from the defects of RGO, and the porous network structure formed.

Stabilization of SiNPs in a Sandwich-typed RGO Framework: SiNPs were stabilized by RGO sheets under mild condition. Results of the cycling performance of the sample displayed a high reversible capacity of 1074 mAh·g-1 at the 5th cycle and 900 mAh·g-1 at the 100th cycle at a current density 200 mA·g-1 with a capacity retention of about 84%; Results of the rate capability of the sample showed a capacity around 900 mAh·g-1 compared with the highest capacity of 1019 mAh·g-1 (88 %, capacity retention) after 65 cycles with current densities ranging from 100 to 2000 mA·g-1. A good coverage of SiNPs by RGO sheets can not only buffer the volume change of SiNPs during charge/discharge and prevent the SiNPs from aggregating, but also improve the overall ionic

and electrical conductivities of the electrode. Micron-sized composite particles with an average thickness of about 1 μm were favourable for electrolyte ions to go through, thus enhancing rate capacity. The sandwich-typed composite anode showed comparable battery performance with the latest reported results, but the preparation rout was more efficient and environmental-friendly.

7.2 Recommendations for future work

Based on the research results collected in this thesis work, a few recommendations for future research work are listed below:

1. 3D porous RGO aerogel can confine SiNPs and facilitate electrolyte transport, but the cycling stability with high reversible capacities is poor. To stabilize the entrapped SiNPs in the RGO aerogel, carbon materials such as carbon nanotubes can be added into the RGO aerogel, especially by designing vertically grown CNTs between RGO layers.

2. Preparing sandwich-type Si/RGO composite has proven to be a facile and effective method for improving SiNPs as anode for LIBs, by balancing well the weight ratio of RGO and SiNPs as shown in Chapter 6. Future research in this area could try other Si nanomaterilas such as porous Si prepared by using the electrochemical etching method, SiNPs prepared by magnesiothermic reduction of SiO2. And the topic of synthesis route of material preparation will be studied.

3. The observation that a significant improvement on battery performances by replacing PVDF with CMC and SA indicates that binders used in battery fabrication deserve future studies, and this is a research area that has not been widely exploited. Especially certain binders can form chemical bonds with silicon to enable new properties or improve some properties of silicon anode materials.

4. The observation that Si-based anode materials for LIBs detached from the current collector after long-time cycling indicates that copper foil as a current collector can be replaced by other 3D metal foams. This is another method to maintain the integrity of the fabricated electrode over long-time battery cycling in addition to finding alternative binders to substitute PVDF.

5. The initial CEs of the Si@RGO samples are relatively low, originating from SEI formation. The mechanism of such SEI formation will be further studied.

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