Novel Nanomaterials for Lithium-Ion Batteries and Lithium-Sulfur Batteries

Novel Nanomaterials for Lithium-ion Batteries and Lithium-sulfur Batteries

A thesis presented for the award of the degree of

Doctor of Philosophy

from

University of Technology Sydney

Jing Xu, B. Sc., M.Sc.

September, 2017

CERTIFICATE OF ORIGINAL AUTHORSHIP

I, Jing Xu, certify that the work presented in this thesis has not previously been

submitted for a degree nor has been submitted as part of requirements for a degree

except as fully acknowledged within the text.

I also certify that the thesis has been written by me. Any help that I have received in

my research work and the preparation of the thesis itself has been acknowledged. In

addition, I certify that all information sources and literature used are indicated in the

thesis.

Jing Xu

Production Note: Signature removed prior to publication.

Sydney, Australia

September, 2017

I DEDICATION

This thesis is dedicated to my parents. Especially thanks for their deep love and selfless dedication.

II ACKNOWLEDGEMENTS

Firstly, I would like to express my great appreciation to my supervisor, Professor

Guoxiu Wang, for his invaluable guidance, kind understanding, positive encouragement that enable me to fulfill my full potential in my Ph.D study.

I also would like to acknowledge my co-supervisor Dr. Dawei Su. His professional dedication, sincere assistance and constant support shrew a promising light for me to successfully complete my research.

Special thanks are expressed to my colleagues at CCET for their kind help and excellent cooperation in my research activities. Mr. Weizhai Bao, Dr Wenxue Zhang,

Miss Yufei Zhao, Dr Shuangqiang Chen, Dr Xiuqiang Xie, Dr Kefei Li, Mr. Jinqiang

Zhang, Dr. Hao Liu, Dr. Anjon Kumar Mondal, Mr. Xin Guo. I would like to especially acknowledge Dr. Jane Yao. Her professional management of laboratory resources makes me can fully devote to my research.

In addition, I appreciate the administrative and technical support I received from Dr.

Ronald Shimmon, Dr. Linda Xiao, Miss Sarah King, and Miss Emaly Black. I take this chance to mention my cordial thanks to China Scholarship Council (CSC) and

ARENA project (ARENA 2014/RND106) for financing my Ph.D research.

Last but not the least, I would like to convey my deeply grateful to my parents for the warmest love and immense understanding they give me, which inspire me to

III overcome tough difficulties and make breakthrough progress in my research. Many thanks also express to Xiaolong Shi for his motivation and kind support during my four years overseas study.

Jing Xu

IV RESEARCH PUBLICATIONS

1) Jing Xu, Qiujiang Liu, Caiping Zhang, Research on the internal resistance cycle performance of lithium-ion batteries echelon use, ITEC Asia-Pacific, 2014. 2˅Jing Xu, Dawei Su, Weizhai Bao, Yufei Zhao, Xiuqiang Xie and Guoxiu Wang,

Rose Flower-Like NiCo2O4 with Hierarchically Porous Structures for Highly Reversible Lithium Storage, Journal of Alloys and Compounds, 2016, 684, 691-698. 3˅Jing Xu, Dawei Su, Wenxue Zhang, Weizhai Bao, Guoxiu Wang, Nitrogen-Sulfur Co-doped Porous Graphene Matrix as a Sulfur Immobilizer for High Performance Lithium Sulfur Battery, Journal of Materials Chemistry A, 2016, 4, 17381-17393.

4˅Jing Xu, Dawei Su, Guoxiu Wangˈ3D Co3O4-CFCs freestanding cathode for lithium sulfur battery, Materials Science and Engineering, 2017, 222, 012013. 5˅Jing Xu, Dawei Su, Guoxiu Wang, Entrapment of Polysulfides by Hierarchically Mesoporous Yolk-Shell Carbon Spheres for High Performance Lithium-Sulfur Batteries, Carbon (Under Review) 6˅Jing Xu, Dawei Su, Guoxiu Wang, Novel Cathode Materials for Lithium-Sulfur Batteries, Advanced Energy Materials, 2018, 1702607. 7˅JingXu, Wenxue Zhang, Dawei Su, Guoxiu Wang, Nitrogen-doped Metal Organic

Framework Derived Co3O4/N-C-rGO Hollow Dodecahedrons with Strong Polysulfide Affinity for Lithium-Sulfur Batteries, Journal of Materials Chemistry A, 2017. DOI: 10.1039/C7TA10272K. 8˅Weizhai Bao, Jing Xu, Dawei Su, Sinho Choi, Xin Guo, Xiuqiang Xie, Jinqiang Zhang, Changhao Liang, Guowen Meng, Guoxiu Wang, A New Sulfur Host based on Graphene-like Porous Mxene Nanosheets for High performance Lithium-Sulfur Batteries, Energy Storage Materials, 2017. 10) Weizhai Bao, Xiuqiang Xie, Jing Xu, Xin Guo, Jianjun Song, Dawei Su, Guoxiu

Wang, Wenjian Wu Confine sulfur in 3D flexible hybrid MXene/reduced

graphene oxide nanosheets for lithium sulfur battery, Chemistry-A European

Journal, 2017, 10.1002/chem.201702387

11) Weizhai Bao, Anjon Kumar Mondal, Jing Xu, Chenying Wang, Dawei Su,

Guoxiu Wang, 3D hybrid–porous carbon derived from carbonization of metal

V organic frameworks for high performance supercapacitors, Journal of Power

Source, 2016, 325, 286-291.

12) Xin Guo, Kefei Li, Weizhai Bao, Yufei Zhao, Jing Xu, Hao Liu, Guoxiu Wang,

Highly Reversible Lithium-polysulfide Semi-liquid Battery with Nitrogen-rich

Carbon Fiber Electrodes, Energy Technology 2017, 8, 1001

VI TABLE OF CONTENTS

Novel Nanomaterials for Lithium-ion Batteries and Lithium-sulfur Batteries ...... I

CERTIFICATE OF ORIGINAL AUTHORSHIP ...... I

DEDICATION ...... II

ACKNOWLEDGEMENTS ...... III

RESEARCH PUBLICATIONS ...... V

TABLE OF CONTENTS ...... VII

LIST OF TABLES ...... XIII

LIST OF FIGURES ...... XIV

ABSTRACT ...... XXVII

INTRODUCTION ...... XXXI

Chapter 1 Literature Review ...... 1

1.1 Lithium-ion Batteries ...... 1

1.1.1 Electrochemistry of lithium-ion batteries ...... 1

1.1.2 Anode materials for lithium-ion batteries ...... 4

1.1.3 Cathode materials for lithium-ion batteries ...... 25

1.2 Lithium-sulfur Batteries ...... 33

1.2.1 Electrochemistry of lithium-sulfur batteries ...... 34

VII 1.2.2 Categories of lithium-sulfur batteries ...... 35

1.2.3 Cathode materials for lithium sulfur batteries ...... 41

1.2.4 Anode materials for lithium-sulfur batteries ...... 82

Chapter 2 Experimental Method and Characterization ...... 84

Overview ...... 84

Material Preparations ...... 87

2.2.1 Solid-state reaction...... 87

2.2.2 Hydrothermal method ...... 89

Material Characterizations ...... 90

2.3.1 X-ray Diffraction (XRD) ...... 90

2.3.2 N2 sorption/desorption measurement ...... 91

2.3.3 Raman spectroscopy ...... 92

2.3.4 X-ray photoelectron spectroscopy (XPS) ...... 94

2.3.5 Thermogravimetric analysis (TGA) ...... 94

2.3.6 Fourier transform infrared spectroscopy (FTIR) ...... 95

2.3.7 Ultraviolet-visible spectroscopy (UV) ...... 96

2.3.8 Scanning electron microscopy (SEM) ...... 97

2.3.9 Transmission Electron Microscopy (TEM) ...... 98

VIII Electrode Preparation and Batteries Assembly ...... 99

2.4.1 Lithium-ion batteries ...... 99

2.4.2 Lithium–sulfur batteries ...... 100

Electrochemical Performance Characterization ...... 101

2.5.1 Cyclic voltammetry ...... 101

2.5.2 Galanostatic charge and discharge ...... 102

2.5.3 Electrochemical impedance spectroscopy ...... 104

2.5.4 Computational methods ...... 104

Chapter 3 Rose Flower-Like NiCo2O4 with Hierarchically Porous Structures for

Highly Reversible Lithium Storage ...... 106

3.1 Introduction ...... 106

3.2 Experimental Section ...... 109

3.2.1 Synthesis of hierarchical porous rose flower-like NiCo2O4 ...... 109

3.2.2 Characterization of materials ...... 109

3.2.3 Electrochemical measurement ...... 110

3.3 Results and Discussion ...... 111

3.4 Conclusions ...... 126

Chapter 4 Nitrogen-Sulfur Co-doped Porous Graphene Matrix as a Sulfur

Immobilizer for High Performance Lithium Sulfur Battery ...... 128 IX 4.1 Introduction ...... 128

4.2 Experimental Section ...... 131

4.2.1 Synthesis of NSG (Nitrogen-Sulfur Co-doped Graphene) ...... 131

4.2.2 Synthesis of A-NSG (Activated Nitrogen-Sulfur Co-doped Graphene)

132

4.2.3 Synthesis of A-NSG@S (Activated Nitrogen-Sulfur Co-doped

Graphene@S) ...... 132

4.2.4 Characterization of materials ...... 133

4.2.5 Electrochemical measurement ...... 134

4.2.6 Computational methods ...... 135

4.3 Results and Discussion ...... 136

4.4 Conclusions ...... 164

Chapter 5 Nitrogen Doped Mesoporous Yolk-Shell Carbon Spheres for High

Performance Lithium-Sulfur Batteries ...... 166

5.1 Introduction ...... 166

5.2 Experimental section ...... 169

5.2.1 Synthesis of NYSC(Nitrogen Doped Yolk-Shell Carbon Sphere) ...... 169

5.2.2 Synthesis of NYSC@S (Nitrogen Doped Yolk-Shell Carbon Sphere@S)

170

X 5.2.3 Characterization of materials ...... 170

5.2.4 Electrochemical measurement ...... 171

5.3 Results and Discussion ...... 172

5.4 Conclusions ...... 197

Chapter 6 Nitrogen-doped Hollow Co3O4 Nanoparticles Coated With Reduced

Graphene as High-Capacity Cathodes for Lithium-Sulfur Batteries ...... 199

6.1 Introduction ...... 199

6.2 Experimental Section ...... 202

6.2.1 Material synthesis ...... 202

6.2.2 Characterization of materials ...... 203

6.2.3 Electrochemical measurement ...... 204

6.2.4 Computational methods ...... 205

6.3 Results and Discussion ...... 206

6.4 Conclusions ...... 230

Chapter 7 Co3O4-Carbon Cloth Free Standing Cathode for Lithium Sulfur Battery

232

7.1 Introduction ...... 232

7.2 Experimenl Section...... 232

7.3 Results and Discussion ...... 233 XI 7.4 Conclusions ...... 240

Chapter 8 Conclusions and Future Perspective ...... 241

8.1 Conclusions ...... 241

8.2 Future Perspective ...... 245

APPENDIX: NOMENCLATURE ...... 248

REFERENCES ...... 251

XII LIST OF TABLES

Table 1.1 Comparison of metal compounds (-S, -N, -OH, -C, MOFs) used as cathode materials for lithium-sulfur batteries...... 79

Table 2.1 Chemicals used in the research project ...... 85

Table 3.1 Comparison of the electrochemical performances of the as-prepared

NiCo2O4 with the reported ones ...... 123

st th Table 3.2 Electrochemical parameters of the NiCo2O4 for the 1 ˈ100 cycles ...... 124

Table 4.1 Percentage of Carbon atoms and Full width at Half-Maximum Values of C

1s Peaks in Graphene, A-NSG...... 144

Table 4.2 Carbon, oxygen, nitrogen, sulfur atomic percent of GO, A-NSG materials.

...... 146

Table 4.3 Comparisons of the adsorption energies of graphene, N-graphene and

N,S-graphene with Li2S, and L2S4...... 152

Table 4.4 Comparisons of the atomic charge transfers (QT) of graphene, N-graphene and N,S-graphene with Li2S, and Li2S4...... 152

Table 4.5 Parameters identification by modeling the impedance spectra in Figure 18b

...... 161

Table 4.6 Recent advance in the dual doped carbon framework to host sulfur for Li-S batteries ...... 164

Table 5.1 Comparison of the electrochemical performance of the as-prepared

NYSC@S with the reported...... 196

XIII LIST OF FIGURES

Figure 1.1 Schematic of the conﬁguration of rechargeable Li-ion batteries...... 2

Figure 1.2 Electrode materials and corresponding electrochemical performances in the current LIB technologies ...... 3

Figure 1.3 (a) Schematic representation of the fabrication of hydrogenated-LTO

(H-LTO); (b) Electrochemical performance of H-LTO and LTO nanowires. (From Ref.

[134].)...... 9

Figure 1.4 Schematic of morphological changes that occur in Si during electrochemical cycling a, The volume of silicon anodes changes by about 400% during cycling. As a result, Si films and particles tend to pulverize during cycling.

Much of the material loses contact with the current collector, resulting in poor transport of electrons, as indicated by the arrow. b, NWs grown directly on the current collector do not pulverize or break into smaller particles after cycling. Rather, facile strain relaxation in the NWs allows them to increase in diameter and length without breaking. This NW anode design has each NW connecting with the current collector, allowing for efficient 1D electron transport down the length of every NW...... 12

Figure 1.5 Electrochemical data for Si NW electrodes. a, Cyclic voltammogram for

Si NWs from 2.0 V to 0.01 V versus Li/Li+ at 1 mV s-1 scan rate. The first seven cycles are shown. b, Voltage profiles for the first and second galvanostatic cycles of the Si NWs at the C/20 rate. The first charge achieved the theoretical capacity of

-1 4,200 mAh g for Li4.4Si. c, The voltage profiles for the Si NWs cycled at different power rates. The C/20 profile is from the second cycle. d, Capacity versus cycle number for the Si NWs at the C/20 rate showing the charge (squares) and discharge capacity (circles). (From Ref. [19].) ...... 13

Figure 1.6 (a) TEM image of SiOx nanoconifer on NiSix nanowires (NWs) (inset:

SEM image of SiOx/NiSix NW); (b) scanning TEM-HAADF images and EDX elemental colour map of single SiOx/NiSix NW; (c) electrochemical cycling stability along with coulombic efficiency and specific capacity vs. charge/discharge cycle XIV numbers of SiOx/NiSix NWs (inset: galvanostatic chargeedischarge curves in the range of 0.01 and 1.2 V); (d) cycling performance of SiOx/NiSix NWs at different current rates. (From Ref. [22].) ...... 15

Figure 1.7 (a) Schematic representation of the lithium reaction mechanism in Ge,

GeO2/C and GeO2/Ge/C; (b) Cycling performance of all samples for 50 cycles at 1C rate; (c)rate performance of all samples from 0.1 to 10 C. (From Ref. [26].) ...... 18

Figure 1.8 (a, b) TEM images of Fe2O3 nanotubes (FNTs) from MONTs template;

Li-ion battery performance of FNTs, (c) discharge capacity vs. number of cycles of

FNTs (red)and commercial Fe2O3 nanoparticles (black). The coulombic efficiency of

FNT is also included (violet); (d) galvanostatic chargeedischarge curves; (e) cycling performance of FNT at different current rates. (From Ref. [30].) ...... 21

Figure 1.9 (a) Galvanostatic first chargeedischarge curves and (b) cycling performance for three CoO samples at the current rate of 0.2 C. (From Ref. [32].) ... 23

Figure 1.10 Lamellar structure of LiCoO2...... 26

Figure 1.11 Crystal structure of spinel LiMn2O4...... 28

Figure 1.12 (a) Crystal structure of LiFePO4. (b) Schematic representation of the processes during charge/discharge of LiFePO4...... 31

Figure 1.13 (a) Schematic comparison of theoretical gravimetric/volumetric energy densities of Li-ion (graphite/LiCoO2) and Li–S batteries. (b) The discharge reaction mechanism of a lithium sulfur battery. Copy right Nature communications, 2013, 4,

2985 (c) Energy density of various electrochemical storage systems. Copyright Adv.

Energy Mater. 2015, 5, 1401986 ...... 34

Figure 1.14 Structures of (a)elemental sulfur, (b)small sulfur cathode materials and their schematic discharge/charge profiles Copyright Adv. mater, 2017, 1606823...... 37

Figure 1.15 Structures of (a)elemental sulfur, (b)small sulfur, (c)Li2S, (d)catholyte cathode materials and their schematic discharge/charge profiles Copyright Adv. mater,

2017, 1606823...... 38

Figure 1.16 a) TEM image of the Ti4O7 particle; b) enlarged zone in (a); c) TEM

XV image of the Ti4O7 particle coated with a thin layer of carbon; d) enlarged zone in (c); e) Rate capabilities of the Ti4O7 based cathode and the carbon-coated Ti4O7 based cathode at different current rates (0.1, 0.2, 0.5, 1, and 0.1 C). f) Cycling performance of the Ti4O7, carbon-coated Ti4O7 and TiO2 based cathodes with ≈50 μL of electrolyte over 300 cycles at a charge/discharge rate of 0.1 C. The solid squares represent capacity and hollow squares represent Coulombic efficiency...... 44

Figure 1.17 Morphological and structural characterizations of MMNC and

CeO2/MMNC nanospheres. (a) SEM image of MMNC nanospheres. (b,c) SEM and

(d−f) TEM images of CeO2/MMNC nanospheres. (f) HRTEM image of CeO2 nanocrystals embedded in the pores of MMNC nanospheres. The lattice distance in the inset of (f) is 0.31 nm, corresponding to the (111) planes of CeO2 nanocrystals. . 47

Figure 1.18 Unit cell of different types of MOFs that applied in lithium sulfur batteries ...... 50

Figure 1.19 (a) Cycling performace of S@MIL-100(Cr)ˈCopyright Journal of the

American Chemical Society, 2011, 133, 16154-16160. (b) Cycling performane of

S@HKUST-1, Copyright Crystal Growth & Design, 2013, 13, 5116-5120. (c) Cycling performane of S@ZIF-8,S@MIL-53, S/NH2-MIL-53, S@HKUST-1, and Schematic of the largest apertures of the four MOFs. Copyright Energy & Environmental

Science, 2014, 7, 2715. (d) Comparision of binding energies of lithium polysulfides to

Ni-MOF and Co-MOF, and cycling performance of Ni(II)-based and Co(II)-based

MOF/S composites at 0.2 C within a voltage range of 1.5−3.0 V. Copyright Nano letters, 2014, 14, 2345-2352. (e) XPS S 2p spectra of S@MOF-525(Cu), Cycle performance of S@MOF-525(2H), S@MOF-525(FeCl), and S@MOF-525(Cu) with the Coulombic efficiency of S@MOF-525(Cu). Copyright ACS applied materials & interfaces, 2015, 7, 20999-21004. (f) V 2p spectrum of S@MIL-100(V), cycling performance at a current of 0.1 C, Copyright Nano Research, 2016, 10, 344-353. .... 51

Figure 1.20 (a) XPS spectra of the Na2Fe[Fe(CN)6] and S@Na2Fe[Fe(CN)6] and high resolution XPS spectra of Fe, N, C, and S of the Na2Fe[Fe(CN)6] and

XVI S@Na2Fe[Fe(CN)6]. (b) Photo of 5 mM pristine Li2S6 solution and upper solution after soaking the Na2Fe[Fe(CN)6] nanocrystals. (b) UV-Vis spectra of 5 mM pristine

Li2S6 solution and upper solution after soaking the Na2Fe[Fe(CN)6] nanocrystals. (c)

Cycling performances of S@Na2Fe[Fe(CN)6]@PEDOT composite at 5 C current rate.

Atomic model configurations showing the interactions between Na2Fe[Fe(CN)6] and polysulfide Li2Sx (x = 8, 6, 4, and 2). The optimized structure and the electron density of the PEDOT with the S8 and polysulfide Li2Sx (x = 8, 6, 4, and 2). Copyright

Advanced materials, 2017, 1700587 ...... 55

Figure 1.21 (a) Schematic illustration for the preparation of WS2 vertically aligned on the CNFs. (b) Schematics of various polysulfides conformations on C@WS2/S. (c)

Binding energies (Eb) for the Li−S composites at six different lithiation stages (S8,

Li2S8, Li2S6, Li2S4, Li2S2, and Li2S) on C@WS2/S, as given by first-principles calculations designed to study the interaction between the lithium sulfide species and

WS2. UV−vis spectra of the Li2S6 solution with C@WS2/S and C/S (inset photograph of the Li2S6 solution with different active materials). (d) Long cycles with various

C-rates. (e) Long-term cycling stability test showing an unprecedented high capacity retention with an excellent Coulombic efficiency over 1500 cycles at 2 C...... 60

Figure 1.22 (a) Schematic of the synthesis of the MoS2-x/RGO composite and the conversion of Li2Sx on the MoS2-x/RGO surface. (b) Cyclic voltammograms of symmetric cells with identical electrodes of MoS2-x/RGO, MoS2/RGO and RGO in

-1 electrolytes with and without 0.2 M Li2S6 at 3 mV s . (c) XPS spectra of the RGO,

MoS2-x/RGO counter electrodes of symmetric cells after scanning to 1.4 V, or after scanning to 1.4 V and returning to 0 V. Copyright Energy Environ. Sci., 2017, 10,

1476-1486...... 63

Figure 1.23 In situ TEM study of MoS2-encapsulated hollow sulfur spheres. (a)

Photographs of a flexible film of MoS2-encapsulated hollow sulfur spheres. (b, c)

SEM images of MoS2-encapsulated hollow sulfur sphere with typical wrinkles generated by the stacking of 2-D flakes marked by violet arrows in (c). (d) Schematic

XVII of in situ TEM setup. (e−i) Time-lapse images of the continuous lithiation and delithiation of MoS2-encapsulated hollow sulfur spheres to demonstrate the high reversibility. Copyright Journal of the American Chemical Society, 2017, 139,

10133-10141...... 64

Figure 1.24 (a) Schematic illustration of the synthesis of the CH@LDH/S composite.

SEM and TEM images of ii,vi) ZIF-67, iii,vii) single-shelled ZIF-67@LDH, iv,viii) double-shelled CH@LDH nanocages, v, ix) CH@LDH/S is given. Copyright

Angewandte Chemie, 2016, 55, 3982-3986. (b) i) XRD patterns and ii) TGA curve of the S@Ni(OH)2 composite are shown. iii) FESEM image of S@Ni(OH)2 and corresponding EDX elemental mappings of iv) O, v) S and vi) Ni are provided.

Figure 1. 25 (a) i)A low-magnification SEM image, ii)high-angle annular dark-field

(HAADF) STEM image, iii)TEM images of the as-prepared porous VN/G composite are shown. (b) i) A STEM image of a VN nanoribbon after cycling with the corresponding elemental maps of ii)vanadium, iii) nitrogen and iv)sulfur is provided.

Scale bars indicate 100 nm. (c) Cycling stability of the VN/G cathode at 1C for 200 cycles is shown. (d) i) A side view of a Li2S6 molecule on a nitrogen-doped graphene surface is shown; the binding energy between Li2S6 and pyridinic N-doped graphene was calculated at 1.07 eV. ii)A side view of a Li2S6 molecule on a VN (200) surface is shown; the binding energy between Li2S6 and VN was calculated at 3.75 eV.Copyright Nature communications, 2017, 8, 14627...... 71

Figure 1. 26 (a) SEM images of the Co3O4 phase, Co4N phase, Co4N @S phase. (b)

Co 2p3/2 X-ray photoelectron spectroscopy of the Co4N phase and Co4N/Li2S6, respectively. (c) Sealed vials of a lithium polysulfide solution (Li2S6 dissolved in

DOL/DME solvents) containing 5, 10, 20, and 30 mg of Co4N phase a) and Co3O4 phase b) after 12 h, respectively. (d) Charge and discharge capacity of Co4N/90S versus cycle number at current densities of 2 and 5 C and Co4N/95S at 2 C. Copyright

Nano letters, 2015, 15, 5137-5142...... 73

XVIII Figure 1.27 (a) Replacement of the Ti-OH bond on the Mxene surface with a S-Ti-C bond on heat treatment or by contact with polysulfides. (b) Schematic demonstrating the two-step interaction between a representative hydroxyl-decorated Mxene phase and polysulfide. (c) First-principles calculations for the interaction between Ti3C2

Mxene and polysulfide Li2S4, showing the most stable Li2S4 binding geometry configuration after full relaxation on: i) Ti3C2(OH)2 (used to represent the pristine

Mxene), ii) Ti3C2(OH)2 with one hydroxyl vacancy, iii) Ti3C2 without any surface functional groups. Iv) The variations of the binding energies of polysulfide molecules

(Li2S, Li2S2, and Li2S4) bonding to the respective substrates. Blue, brown, red, pink, yellow, and green spheres represent Ti, C, O, H, S, and Li. (d) long term cycling performance. Cells with 3.6 and 5.5 mg cm–2 sulfur loadings were examined at C/5.

Figure 2.1 Framework of the experiments ...... 84

Figure 2.2 Autoclaves for hydrothermal synthesis ...... 90

Figure 2.3 The Bruker D8 Discover XRD instrument...... 91

Figure 2.4 The 3 Flex surface characterization analyser instrument produced by

Micromeritics...... 92

Figure 2.5 The Renishaw inVia Raman microscope equipped with a Leica DMLB microscope (Wetzlar, Germany) and a 17 mW at 633 nm Renishaw helium neon laser.

...... 93

Figure 2.6 TGA Analyzer (SDT 2960 model, TA Instrument) ...... 95

Figure 2.7 FTIR Equipment ...... 96

Figure 2.8 Carry 300 UV/vis spectrophotometer ...... 96

Figure 2.9 The field emission scanning electron microscopy in a mode of Supera 55

VP produced by Zeiss and equipped with EDS detector...... 98

Figure 2.10 TEM instrument (JEM-2010FS) equipped with EDX detector...... 99

Figure 2.11 The argon-filled glove box (Mbraun, Unilab, Germany)...... 101

Figure 2.12 The CH instruments (CHI 660D) for CV and EIS testing...... 102

XIX Figure 2.13 The computer-controlled Neware battery test system...... 103

Figure 3.1 XRD patterns of NiCo-precusor ...... 111

Figure 3.2 (a, b and c) FSEM images of the NiCo-precusors at different magnifications. (d and e) TEM images for NiCo-precusors. (f) SAED pattern for

NiCo-precusors...... 112

o Figure 3.3 (a, b, c and d) FESEM images of the NiCo2O4 after calcination at 450 C for 1h...... 113

Figure 3.4 (a and b) TEM images for the hierarchical porous flower-like NiCo2O4. (c)

Corresponding SAED pattern. (d) HRTEM image for NiCo2O4...... 114

Figure 3.5 (a) XRD patterns of China rose-flower-like NiCo2O4. (b) TGA analysis of

NiCo-precursor. (c) FTIR spectra of NiCo-precusors and China rose-flower-like

NiCo2O4. (d) Nitrogen adsorption/desorption isotherms of China rose-flower-like

NiCo2O4, inset is pore size distribution...... 115

Figure 3.6 XRD patterns of the products calcined at 500Ԩ ...... 116

Figure 3.7 FSEM images of NiCo2O4 with different amount of PVP. (a) 0g. (b) 0.05g.

Figure 3.8 Illustration for the formation of hierarchical, porous flower-like NiCo2O4

...... 119

Figure 3.9 Electrochemical performance of the flower-like NiCo2O4. (a) The first four consecutive CV curves. (b) Discharge-charge capacity vs cycle number at current densities of 100, 200, 500, 1000, 2000 mA g-1. (c) Discharge and charge profiles for the 1st, 2nd, 100th cycles at 1000mA g-1. (d) Discharge/charge capacity and coulombic efficiency vs. cycle number at a 1000mA g-1 current density...... 120

Figure 3.10 EIS spectra of flower-like NiCo2O4 and corresponding equivalent circuit.

...... 124

Figure 4.1 Illustration for the formation of A-NSG ...... 136

Figure 4.2 SEM image of NSG...... 137

Figure 4.3 (a) SEM image of A-NSG. (b) EDX elemental mapping for A-NSG. (c)

XX SEM image of the A-NSG@S. (d) AFM image of A-NSG...... 138

Figure 4.4 (a, b) TEM images of A-NSG. (c, d) High-magnification TEM images of the A-NSG@S ...... 139

Figure 4.5 Thermogravimetric curves of pure sulfur powder and A-NSG@S in the N2 with a heating rate of 10 Ԩ min-1...... 140

Figure 4.6 XRD patterns and Raman spectra of the pure sulfur powder, A-NSG,

A-NSG@S-155, A-NSG@S-230 ...... 141

Figure 4.7 (a, b, c) High-resolution C 1s, N 1s and S 2p XPS spectrum of A-NSG and corresponding XPS survey spectra of A-NSG. (d) Schematic structure of A-NSG .. 142

Figure 4.9 (a) Raman spectra of GO, NSG, A-G and A-NSG. (b) Pore characterization of the NSG, A-G and A-NSG materials ...... 147

Figure 4.10 (a) Li2S4 solution (b) Initially mixed (c) Aging for 3h (d) Aging for 24h

...... 148

Figure 4.11 UV-vis spectra of 5 mM pristine Li2S4 solution and solution after soaking the N, S-graphene...... 148

Figure 4.12 High-resolution XPS (a) Li 1s (b) S 2p spectra of N,S graphene-Li2Sx 149

Figure 4.13 High-resolution S 2p spectra of N,S graphene@ S ...... 150

Figure 4.14 Optimized configurations for the adsorption of Li2S and Li2S4 on pristine graphene (a, b) and N,S codoped graphene (c, d) with corresponding adsorption energies in eV and atom Mulliken charge. Gray, blue, yellow, and purple balls represent C, N, S, and Li atoms, respectively...... 151

Figure 4.15 Electrochemical characterization of A-NSG material as the cathode of a

Li-S battery. (a) Cyclic voltammetry (CV) measured between 1.7 and 2.7 V at a sweep rate of 0.1 mV s-1 for the first, second, third, and fifth cycles. (b) Galvanostatic charge

XXI and discharge profiles for different cycles at 0.2 C. (c) The discharge capacities of high and low voltage plateaus. The onset voltage of the low plateau was defined around 2.0 V. (d) Long term cycling performance test of the A-NSG@S electrode at

0.2 C discharge rate and corresponding Coulombic efficiency...... 153

Figure 4.16 Long-term cycling performance test of the A-NSG@S electrode at 1C discharge rate and the corresponding Coulombic efficiency...... 155

Figure 4.17 Electrochemical characterization of A-NSG material as the cathode of a

Li-S battery. (a) Galvanostatic charge and discharge profiles at 0.1 C, 0.5 C, 1 C, 2 C,

5C. (b) Voltage plateaus for charge and discharge processes over 200 cycles at 0.1 and

1 C. (c) Discharge/Charge capacity cycled at various rates from 0.1 C, 0.5 C, 1 C, 2 C,

5 C...... 156

Figure 4.18 (a) dQ/dV plots of A-NSG@S electrode. (b) Electrochemical impedance spectra of A-NSG @S lithium-sulfur battery at different cycles...... 158

Figure 4.19 Cycling performance of A-NSG@S material cycled at 0.2 C, in comparison with NSG@S, A-G@S, rGO@S material. Specific capacity values were all calculated based on the mass of sulfur...... 163

Figure 5.1 Synthesis procedure of the mesoporous NYSC@S...... 172

Figure 5.2 (a) SEM image of the precusors obtained from hydrothermal process. (b)

SEM image of carbon-silica after calcination. (c, d) SEM image of the NYSC. ... 173

Figure 5.3 (a-d) EDX maps of the precursor obtained from hydrothermal process. (e)

EDX spectrum of the precursor. (f) Line scanning profiles of a single precursor. .... 175

Figure 5.4 (a, b) SEM image of carbon-silica after calcination...... 176

Figure 5.5 (a-d) EDX maps of carbon-silica after calcination...... 176

Figure 5.6 (a) SEM image of NYSC. (b) High-magnification SEM images of the

NYSC. (c) Particle size distribution. (d) TEM image of the NYSC...... 177

Figure 5.7 (a) N2 isothermal curves and pore size distributions of NYSC and

NYSC@S. (b) XPS spectrum of NYSC, (c) XRD patterns and (d) Raman spectra of the NYSC@S, NYSC, Bare Sulfur...... 178

XXII Figure 5.8 N2 isothermal curves and pore size distributions of (a)carbon/silica composite , (b)NYSC@S...... 179

Figure 5.9 (a) SEM image of the NYSC@S. The inset is the magnified SEM image of

NYSC@S. (b) TEM image of NYSC. (c) HRTEM image of NYSC@S. (d) TEM image of NYSC@S. (e, f, g, h) EDX spectra of NYSC@S...... 181

Figure 5.10 TGA curve of the (a)NYSC@S, (b)NCSC@S ...... 183

Figure 5.11 High resolution SEM and TEM images of the NYSC (a, b, c) and NCSC

(d, e, f). In all figures, the scale bar is 300 nm ...... 185

Figure 5.12 Electrochemical characterization of materials as the cathodes for Li-S batteries. (a, b) Cyclic voltammetry (CV) tested between 1.7 and 2.7 V at a sweep rate of 0.1 mV s-1 for NYSC@S, NCSC@S, respectively. (c, d) Galvanostatic charge and discharge profiles for 1st and 300th cycle at 0.2 C, respectively. (e) Long cycling performance of the NYSC@S, NCSC@S electrode at 0.2 C discharge rate and corresponding Coulombic efficiency...... 186

Figure 5.13 (a) Discharge capacity at different rates of 0.5 C, 1 C, 2 C, 3 C and 5 C for NYSC@S,NCSC@S electrodes. (b) Galvanostatic charge and discharge profiles at

0.5 C, 1 C, 2 C, 3 C and 5 C for NYSC@S electrodes...... 189

Figure 5.14 Electrochemical impedance spectra of NYSC@S, NCSC@S batteries for the (a) first cycle and (b) final cycle (500 th)...... 190

Figure 5.15 (a,b) FESEM images of NYSC@S composites before and after 100 cycles. (c,d) FESEM images of NCSC@S composites before and after 100 cycles. 192

Figure 5.16 (a-c) EDX maps of NCSC@S after 100 cycles. (d) EDX spectrum of

NCSC@S after 100 cycles ...... 193

Figure 5.17 Schematic of the lithiation process in core shell and yolk shell sulfur based morphologies...... 195

Figure 5.18 (a-c) EDX maps of NYSC@S after 100 cycles. (d) EDX spectrum of

NYSC@S after 100 cycles...... 194

Figure 5.19 Digital images of NCSC(a), NYSC(b) nanoparticles immersed in

XXIII Li2S6/DOL/DME (1:1, v-v) solution...... 196

Figure 6.1 Schematic illustration for the formation of rGO/N-C-Co3O4@S...... 206

Figure 6.2 (a, b) SEM image of ZIF-67 ...... 207

Figure 6.3 (a) SEM image of N-C-Co3O4, (b) TEM image of N-C-Co3O4, (c) FESEM image of N-C-Co3O4 composite. (d) HRTEM image of N-C-Co3O4...... 209

Figure 6.4 (a) XRD patterns of N-C-Co3O4, rGO/N-C-Co3O4 @S and pure S. (b)

FTIR spectra of ZIF-67, N-C-Co3O4, rGO/N-C-Co3O4@S and pure S. (c) Nitrogen adsorption-desorption isotherm curves of rGO/N-C-Co3O4. (d) Thermogravimetric

-1 curves of rGO/N-C-Co3O4 @S in the N2 with a heating rate of 10 Ԩ min ...... 210

Figure 6.5 SEM images of (a, b) N-C-Co3O4, (c) rGO/N-C-Co3O4 and (d) rGO/N-C-Co3O4@S ...... 211

Figure 6.6 (a)EDX mapping results of rGO/N-C-Co3O4, (b)Overall, (c) Carbon, (d)

Cobalt, (e) Nitrogen, (f) Oxygen...... 211

Figure 6.7 (a) FSEM image of rGO/N-C-Co3O4@S, (b) TEM image of rGO/N-C-Co3O4@S, (c) FSEM image of rGO/N-C-Co3O4@S, (d) HRTEM image of rGO/N-C-Co3O4@S, corresponding elemental mapping images of (e)overall image.

(f)cobalt, (g) oxygen, (h) sulfur, (i) carbon and (j) nitrogen...... 212

Figure 6.8 Nitrogen adsorption-desorption isotherm curves of rGO/N-C-Co3O4@S

...... 213

Figure 6.9 (a) XPS survey spectra of rGO/N-C-Co3O4@S ...... 214

Figure 6.10 XPS spectra of (a) Co 2p, (b) S 2p, (c) C 1s and (d) N 1s in the rGO

/N-C-Co3O4 @S composite ...... 215

Figure 6.11 (a) Cyclic voltammetry (CV) tested between 1.5 and 3 V at a sweep rate

-1 of 0.1 mV s for rGO/N-C-Co3O4@S (b) The first-cycle galvanostatic charge/discharge voltage profiles of rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S cathodes at 0.1 C. (c) Nyquist plots of rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S cathodes. (d) The discharge capacities along high and low voltage plateaus of rGO/N-C-Co3O4@S cathodes. The onset voltage of the low plateau was defined

XXIV around 2.0 V. (e) Prolonged cycle life and Coulombic efficiency of the rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S electrodes at 0.1C...... 217

Figure 6.12 The discharge capacities of high and low voltage plateaus of (a)

N-C-Co3O4@S, (b) rGO@S. The onset voltage of the low plateau was defined around

2.0 V...... 218

Figure 6.13 The 100th and 600th cycle galvanostatic charge/discharge voltage profiles at 0.1 C. (a) rGO/N-C-Co3O4@S, (b) N-C-Co3O4@S, (c) rGO@S...... 220

Figure 6.14 (a) The potential differences changes of charge and discharge plateaus for rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S cathodes between the charge discharge plateaus at various current densities. (b) Voltage profiles of rGO

N-C-Co3O4@S cathodes at various current densities from 0.1 C to 3 C. (c) Discharge capacity of rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S electrodes cycled at various rates spaning 0.1 C, 0.5C, 1C, 2C, 3C. (d) Cycling performance test of the rGO/N-C-Co3O4@S electrode at 1C discharge rate and corresponding Coulombic efficiency...... 222

Figure 6.15 (a) Dissembled electrode of rGO/S after 50 cycles. (b) Dissembled electrode of rGO/N-C-Co3O4@S after 50 cycles...... 224

Figure 6.16 (a)UV-vis absorption spectra of the solution obtained by immersing the cycled rGO@S and rGO/N-C-Co3O4@S cathodes in a mixture of DOL/DME electrolyte. The inset images are visualized colour changes after N-C-Co3O4@S and rGO@S cathodes were immersed in (1), (2) solution for 6h, respectively, (3)

DOL/DME solvent only as reference. (b) Ex situ Raman spectra of bare rGO/N-C-Co3O4-S and the electrode after 100 cycles (discharged to 2.1V). (c) XPS spectra of Li 1s for an electrode after discharging to 2.1V...... 225

Figure 6.17 (a) Adsorption energies and (b) Electrons transfered for Li2S and Li2S4 compounds on rGO and Co3O4 surfaces. Schematic diagram showing the crystal structure of a Co3O4 with top and side views and after adsorption (c) Li2S and (d)

Li2S4 with corresponding atom Mulliken charge. Blue, red, yellow, and purple balls

XXV represent Co, O, S, and Li atoms, respectively...... 227

Figure 6.18 Optimized configurations for the adsorption of Li2S and Li2S4 on rGO (a and b) with corresponding adsorption energies in eV and atom Mulliken charge. Gray, blue, yellow, and purple balls represent C, N, S, and Li atoms, respectively...... 228

Figure 7.1 Schematic illustration for the formation of Co3O4-CC-S ...... 233

Figure 7.2 (a, b)SEM image of Co3O4-CC. (c)Mapping results of Co3O4-CC-S.

(d)SEM image of Co3O4-CC-S. (e)TEM image of Co3O4. (f) SAED pattern for Co3O4.

...... 235

Figure 7.3 TGA curve(a), XRD patterns (b) and Raman spectra (c) of Co3O4-CC-S

...... 236

Figure 7.4 (a) Galvanostatic charge and discharge profiles of Co3O4-CC-S for different cycles at 0.2C. (b) Galvanostatic charge and discharge profiles of CC-S for different cycles at 0.2C. (c) Discharge/Charge capacity cycled at various rates from

0.1 C, 0.5 C, 1 C, 2 C, 3 C. (d) Long term cycling performance test of the

Co3O4-CC-S, CC-S electrodes at 0.2C discharge rate and corresponding Coulombic efficiency...... 238

Figure 7.5 (a) Initially mixed. (b) Aging for 3h. ((1)L2S6 soultion. (2) L2S6 soultion immersed with CC. (3) L2S6 solution immersed with Co3O4-CC...... 239

XXVI ABSTRACT

Advanced energy storage is an intrinsic driving force for modern life. There is a large spectrum of storage technologies with wide variations in terms of energy and power density, service life, efficiency, cost, etc. Batteries have achieved great triumphs in this landscape, as they can be utilized conveniently at low cost. The properties of electrode materials are of great significance for the electrochemical performances of batteries. In this doctoral work, a series of electrode materials were designed and fabricated and their electrochemical properties for lithium-ion batteries and lithium-sulfur batteries were investigated.

Novel porous NiCo2O4 nanoparticles are synthesized by a solvothermal method using poly (vinylpyrolidone) (PVP) as the structure-directing agent followed by a simple

thermal annealing treatment. Through the XRD, FESEM, TEM, and N2 sorption

analyses, it has been found that the as-prepared NiCo2O4 nanoparticles show hierarchical rose flower-like architecture constituted by 2D hierarchically porous nanosheets. The 2D porous nanosheets provide sufficient void space generated during thermal annealing treatment, benefiting electrolyte penetration and fast electron transfer. The porous structure also can tolerate the volume variation upon prolonged

charge/discharge cycling. Therefore, when the as-prepared NiCo2O4 nanoparticles are used as anode materials for the Li-ion batteries, they exhibit high capacity, remarkable capacity retention at increased current densities, and outstanding cycling

XXVII stability.

Lithium-sulfur batteries have attracted extensive attentions because of their high theoretical capacity and high energy density compared with lithium-ion batteries.

Highly ordered mesoporous nitrogen doped yolk-shell carbon spheres were synthesized via a facile sol-gel method, this sulfur host possesses large pore volume and interconnected mesopore structure, which can effectively prevent polysulfide diffusion and stabilize the dissolved polysulfides. The inner mesoporous “yolk” acts as a sulfur reservoir to entrap polysulfide species; Meanwhile, the outer “shell” serves as a physical barrier to confine the dissolution of polysulfides and also enhances the cycling stability of the cathode. When applied as cathode material for lithium-sulfur

(Li-S) batteries, these mesoporous nitrogen doped yolk-shell carbon spheres exhibit a high specific capacity of 1329 mAh g-1 at 0.2C and an extended cycle life, demonstrating a promising cathode host material for lithium-sulfur batteries.

3D nitrogen–sulfur co-doped porous graphene matrix was synthesized via chemical

activation of polypyrrole (PPy) functionalized graphene sheets using K2CO3. The dopant N and S atoms act as electron attracting atoms, leading to the nearby C atoms and causing oxygen functional groups to be polarized and more active for anchoring sulfur and polysulfides. Meanwhile, highly developed defects and edges, as well as porous structure derived from graphene chemical activation, not only achieve a high sulfur loading in a well dispersed amorphous state, but also serve as polysulfide

XXVIII reservoirs to alleviate the shuttle effect. When applied as cathode hosts for lithium-sulfur batteries, the nitrogen-sulfur co-doped porous graphene architecture exhibited a high specific capacity of 1178 mAh g-1 at 0.2C, 1103 mAh g-1 at 0.5C, 984 mAh g-1 at 1C rate, and excellent cycling stability for 600 cycles with a retained capacity of 780 mAh g-1 (0.2C).

Nitrogen-doped hollow Co3O4 nanoparticles coated with reduced graphene oxide

(rGO) were synthesized by a facile solid-state pyrolysis process that using metal

organic framework (ZIF-67) as precursor. The obtained rGO/N-C-Co3O4 architecture offer different types of polar interactions to suppress polysulfide shuttle effect. The

open metal centre in the obtained rGO/N-C-Co3O4 architecture serving as the Lewis acid sites show high affinity to the polysulfide, the doped nitrogen introduces more defects, active sites and can additionally immobilize lithium polysulfide within the cathode. Moreover, the well-defined porous structure and the rGO simultaneously contribute to the electron transfer and remarkably buffer the volume expansion/contraction of active materials upon cycling. Owing to these synergistic

interactions between rGO/N-C-Co3O4 and sulfur species, the rGO/N-C-Co3O4@S composite generated a high reversible capacity (1205 mAh g-1 at 0.2C) and excellent stability (865 mAh g-1 at 1C after 300 cycles). Ex situ Raman, Ex situ X-ray

Photoelectron Spectroscopy, UV-vis absorption spectra and first-principle

calculations further confirmed that rGO/N-C-Co3O4 nanoparticles can effectively bind polysulfides in the electrode over cycles and exhibit high binding energies. XXIX A binder-free cathode was developed by chemisorption of Co3O4 to activated carbon cloth (CC), which was used as a 3D current collector to accommodate a large amount of sulfur, multiwall carbon nanofiber (MWCNF) and carbon black (CB) hybrids within the conductive scaffold, enabling the fabrication of ultrahigh sulfur loaded electrodes. The interconnected carbon fibers established a long-range conductive matrix for an efficient electron transport, the multiple conductive pathways guarantee

high sulfur utilization. More importantly, the polar Co3O4 could also effectively entrapped the intermediated polysulfides preventing their free diffusion to the lithium

anode, guaranteeing good cycling stability. Consequently, the Co3O4-CC-S electrodes exhibit excellent electrochemical performance with sulfur loading of 4.3 mg cm-2.

XXX INTRODUCTION

Availability of energy at an affordable cost without adverse environmental consequences is one of the major challenges to modern society. The increasing consumption and limited availability of fossil fuels, along with the environmental impact caused by burning fossil fuels, have prompted the development of alternate, sustainable, clean energy technologies. Renewable energy sources, such as solar and wind, are appealing in this regard, but the effective utilization of these intermittent energy sources requires efficient and economical electrical energy storage (EES) systems. Meanwhile, the energy storage system market is supposed to experience booming development in the following decades due to the wide range application demands spaning from smart power grids, electric vehicles to portable electronic devices. It is reported that the energy storage system market is expected to increase from 27 GW h in 2017 to 90.9 GW h in 2020, keeping an annual growth rate of

21.7%. These new demands require the energy storage systems with a much higher energy density than is currently available.

Rechargeable batteries are the most viable option for energy storage system. Among the existing energy storage solutions, lithium-ion batteries are considered to be by far the most promising owing to their relatively high capacity, rate capability, high voltage and long cycle life. Despite the rapid advancement in the development of high performance Li-ion batteries in the past decades, there are still some technical

XXXI obstacles that must be overcame to enable a broader range of applications. For instance, the energy density of conventional lithium-ion batteries with

insertion-compound cathodes (e.g., LiCoO2, LiMn2O4, and LiFePO4) and anodes

(graphite) is limited to fulfill the demands of electric vehicles and smart grids, although graphite is the most common commercialized anode material for Li-ion batteries, it can only deliver a limited capacity (372 mAh/g, theoretical capacity). In addition, the safety of Li-ion batteries must be further improved.

Significant advancements in battery performance and reductions in cost are expected to come from new battery chemistries, based on different storage mechanisms at the materials level, and different configurations at the cell and system level. Based on this condition, lithium sulfur battery is the most promising candidate for next-generation rechargeable batteries due to its overwhelming theoretical capacity (1672 mAh g-1) and high energy density (2567 W h kg 1) at low cost. It has different working principle compared with the conventional intercalation/de-intercalation principle of lithium-ion battery. The primary challenges in the development of Li–S batteries include the suppression of the dissolution of polysulfides, improving the conductivity of the sulfur host while maintaining high sulfur loading. Extensive researches have been devoted to improve the electrochemical performance of lithium sulfur battery from the aspects of cathode materials, electrolytes, separators, and lithium-metal anodes.

XXXII This thesis mainly focuses on the synthesis of anode material for Li-ion batteries and a series of new nanocomposites as cathode materials for Li-S batteries, with the aim of preparing suitable nanomaterials and improving their electrochemical performances. The aforementioned problems about Li-ion batteries and Li-S batteries will be carefully addressed via the synthesis novel nanocomposites. Meanwhile, the cost to manufacture the Li-ion batteries and Li-S batteries will also be reduced by the use cheap raw materials.

Each chapter of this doctoral thesis is outlined as follows:

(1) Chapter 1 briefly introduces the background of Li-ion batteries and Li-S batteries.

The basic reaction mechanism and working principles are also presented. For

Li-ion batteries, the anode and cathode materials are both discussed, but this thesis

mainly focuses on anode materials for lithium-ion batteries. For Li-S batteries, the

categories of lithium sulfur batteries also mentioned, this thesis mainly discusses

the latest development of cathode materials for lithium sulfur batteries.

(2) Chapter 2 is the experimental sections, it can be divided into three parts including

material preparations, physiochemical characterizations as well as the

electrochemical measurement. Solid state reaction and hydrothermal synthesis

were mainly applied to prepare different electrode materials in this doctoral work.

The structure and morphology of the prepared materials were characterized by

X-ray diffraction (XRD), field emission scanning electron microscopy (FESEM),

XXXIII transmission electron microscopy (TEM), nitrogen adsorption-desorption, Raman

Spectroscopy, X-ray photoelectron spectroscopy (XPS), and thermogravimetric

analysis. Cell assembly and electrochemical testing techniques were also

presented.

(3) In chapter 3, 3D hierarchical porous rose flower-like NiCo2O4 was successfully

synthesized via a solvothermal method using PVP as the structure-directing agent

followed by a simple thermal annealing treatment. The influences of solvothermal

reaction time and the different amount of PVP on their morphologies and

electrochemical performances are discussed. The mechanism about the formation

of rose flower-like morphology are also explained. When applied as anode

material for lithium-ion battery, it exhibits excellent electrochemical performance.

(4) In chapter 4, activated N–S co-doped porous graphene (A-NSG) was obtained by

using polypyrrole (PPy) as nitrogen source, the PPy monomers polymerize on the

surfaces of graphene oxide (GO) during the polymerization process. The doping

of sulfur was achieved by ammonium persulfate (APS). After carbonization,

nitrogen and sulfur atoms are doped in the graphene nanosheets. Both the

physiochemical characterizations and electrochemical performance of graphene,

N-doped graphene, activated graphene and activated N–S co-doped graphene are

comparatively evaluated. In addition, optimized configurations for the adsorption

of Li2S and Li2S4 on none doped graphene, single doped graphene, N-S co-doped

XXXIV graphene are conducted by quantum mechanical calculations.When applied these

hosts as cathode materials for lithium sulfur batteries, the activated N–S co-doped

graphene exhibited superior specific capacities, good cycling stabilities and high

rate capabilities compared with others. The reason for the enhanced

electrochemical performances is also discussed.

(5) Uniform yolk-shelled carbon spheres (YS-CSs) with a hierarchically porous

nanoarchitecture have successfully been synthesized in Chapter 5 through a new

gradient sol–gel process with surfactant-directing co-assembly by using cationic

surfactant cetyltrimethylammonium bromide (CTAB) as a template,

resorcinol-formaldehyde (RF) as a carbon source and tetraethoxysilane (TEOS) as

an assistant pore-forming agent. The effects of CTAB and TEOS on the yolk-shell

structure and pore properties of the carbon spheres are systematically investigated.

At the same time, nitrogen doped core shell carbon spheres were also prepared in

order to contrastively evaluate the influence of hollow cavity on the performance

of sulfur host. Assisted by the analyses of SEM, TEM and impedance spectra, the

reasons for excellent cycling stabilities are discussed.

(6) Research findings in Chapter 5 inspired the preparation of mesoporous hollow

structure for lithium sulfur batteries in order to alleviate the volume expansion

during sulfur reduction. In chapter 6, nitrogen-doped rGO/N-C-Co3O4 hollow

dodecahedra was obtained through direct pyrolysis of Co-based MOFs (ZIF-67)

by two step calcination, during which N-C-Co3O4 hollow dodecahedra with XXXV meso-sized hierarchical pores were formed and rGO effectively wrapped on the

surface of N-C-Co3O4 polyhedra through electrostatic interactions, UV-vis

absorption spectra and Ex situ Raman spectra were performed to demonstrate the

strong affinity for lithium polysulfides, adsorption energies and electrons

transfered for Li2S and Li2S4 on rGO/N-C-Co3O4 surfaces were carried out by

DFT calculation, the rGO/N-C-Co3O4@S composites exhibited high capacity and

long cycling performance as cathode material for lithium-sulfur battery.

(7) Aiming to improve the energy density of lithium sulfur batteries, we significantly

stabilized cycle life of high sulfur loading binder-free cathode by chemisorption of

Co3O4 to activated carbon fiber cloth, which was used as a 3D current collector to

accommodate a large amount of sulfur, MWCNF and CB hybrids within the

conductive scaffold, enabling the fabrication of ultrahigh sulfur loaded electrodes.

This special nanoarchitecture combines the advantage of strong chemisorption of

lithium polysulfides as well as excellent electrical conductivity, enabling high

sulfur utilization and effective trap of lithium polysulfides. When applied as

cathode materials for lithium sulfur batteries, the cathodes exhibit a reversible

capacity of 1007 mAh g-1 after 300 cycles with a high percent of sulfur loading

(81%).

(8) The last chapter of this doctoral thesis (Chapter 8) briefly summarizes the research

outcomes of this doctoral work and future scope of related research is also

XXXVI presented.

XXXVII Chapter 1 Literature Review

1.1 Lithium-ion Batteries

Among all metals, lithium is the lightest and has the greatest electrochemical potential and provides the largest energy density for weight. However, attempts to develop rechargeable lithium batteries failed because of poor lithium metal recharge ability due to the formation of dendrites and the consequent safety concerns. Research then shifted to a non-metallic lithium battery using lithium ions. In 1991, the first commercialized Lithium-ion battery appeared. Besides the improved safety, lithium-ion batteries have merits such as high energy density, low maintenance, long cycle life, no memory effect and low self-discharge. Lithium-ion batteries have been attracting R&D investment from all corners of the world, with wide applications for portable devices, power tools, and electric vehicles.

1.1.1 Electrochemistry of lithium-ion batteries

The common concept of present lithium-ion batteries relies, at the cathode side, on

lithium transition metals oxides or lithium phosphates active material (LiCoO2,

LiMn2O4, LiCo1/3Mn1/3Ni1/3O2, LiFePO4, etc.), while graphite is commonly used as anode active material. Cathode and anode are separated by a membrane made of polypropylene/polyethylene ﬁlled with electrolyte which contains lithium salts (i.e.

LiPF6) in alkyl organic carbonates such as ethylene, propylene, dimethyl carbonates at

1 different ratio. The separator prevents the electrical contact between the electrodes, at the same time, it allows the diffusion of lithium ions from cathode to anode during the charging and the reverse discharging process. It also need to be noted that the

Figure 1.1 Schematic of the conﬁguration of rechargeable Li-ion batteries. Copyright

Nano Energy, 2013, 2, 425-434. electrolyte solvent should be able to withstand the potentials of electrodes without decomposing. A typical operation of LIBs is shown in Figure 1.1. During the charge-discharge cycles, lithium ions are shuttled between the cathode and anode

electrodes. Taking the typical LiCoO2//C cell as example, during the charging process, electrons transfer via the external circuit from the cathode to the anode. Meanwhile,

Li-ions are extracted from the layered LiCoO2 cathode, move within the electrolyte and intercalate into the anode electrode, storing electrochemical energy within the battery in the form of chemical energy. When the cell is discharged, the reverse occurs

2 and the electrons flow in the opposite direction through the external circuit powering electrical devices.The corresponding chemical reactions are as follows:

+ - Anode: 6C + xLi + xe o LixC6

+ - Cathode: LiCoO2 o Li1-xCoO2 + xLi + xe

In order to be considered suitable candidates for LIBs, active materials should fulfill the requirements of reversible capacity, good ionic and electrical conductivity, long cycle life, high rate of lithium diffusion into active material and conclusively low cost and eco-compatibility. Tremendous efforts have been devoted to developing cathode and anode materials for LIBs.1-11 Figure 1.2 summarizes some of these electrode materials and it can be seen that their potential and specific capacity varies with the chemical compositions. In the following part, recent advancements of electrode materials for LIBs will be reviewed.

Figure 1.2 Electrode materials and corresponding electrochemical performances in the current LIB technologies. Copyright Electrochemistry Communications, 2010, 12,

1302-1306. 3 1.1.2 Anode materials for lithium-ion batteries

Graphite allows the intercalation of only one Li-ion with six carbon atoms, with a

resulting stoichiometry of LiC6 and thus an equivalent reversible capacity of 372 mAh

-1 g . The diffusion rate of lithium into carbon materials is between 10-12 and 10-6 cm2 s-1 (for graphite it is between 10-9 and 10-7 cm2 s-1), which results in batteries with low power density. Hence, there is an urgency to replace graphite anodes to materials with higher capacity, energy and power density. Even though lithium metal holds one of the highest capacity among anode materials (3860 mAh g-1), safety issues prevent the use of lithium as anode material in secondary batteries. In fact, dendrites formation on the lithium metal can cause short circuit between anode and cathode. Therefore, as major challenge, it is very important to select suitable anode materials which can provide high capacity and ease diffusion of Li-ions into the anode, along with good cycling life and free from safety concerns. In this section we will discuss the state of the art of anode materials for LIBs, with particular emphasis on the recent nanotechnology research outcomes and outstanding results. For sake of simplicity, we will classify the discussed innovative anode materials in three main groups, depending on their Li-ion battery performances and reaction mechanism:

1) Intercalation/de-intercalation materials, such as carbon based materials, porous

carbon, carbon nanotubes, graphene, TiO2, Li4Ti5O12, etc;

2) Alloy/de-alloy materials such as Si, Ge, Sn, Al, Bi, SnO2, etc;

4 3) Conversion materials like transition metal oxides (MnxOy, NiO, FexOy, CuO, Cu2O,

MoO2 etc.), metal sulphides, metal phosphides and metal nitrides (MxXy; here X = S,

P, N).

Anode materials based on intercalation reaction

Carbon based materials

The variety of carbon based materials used as active anode in LIBs are classiﬁed into two categories, according to the degree of crystallinity and carbon atoms stacking12-14: i) soft carbon (graphitizable carbons) where crystallites are stacked almost in the same direction and ii) hard carbon (non-graphitizable carbons) where crystallites have disordered orientation. In particular, the former is quite popular in the battery community. In fact, it shows an appropriate reversible capacity (i.e. 350-370 mAh g-1), long cycling life and good columbic efﬁciency (more than 90%). The reaction mechanism between lithium and graphite, following an intercalation/de-intercalation process, has been extensively studied with various analytical techniques.15 Among the types of commercially available graphite is worth mentioning Mesocarbon Microbead

(MCMB), Mesophase-pitchbased carbon ﬁbre (MCF), vapour grown carbon ﬁbre

(VGCF) and Massive Artificial Graphite (MAG). Despite their massive production and the relative low cost of the industrial processes, these classes of carbon materials have, as major issue, a low specific capacity (i.e. 372 mAh g-1), especially for applications such as HEV, PHEV or PEV. Hence, the use of graphitic carbon as anode is still limited to low power devices like mobile phones and laptops. Therefore, 5 further advances in anode materials for lithium batteries are necessary to the development of competitive performance electric vehicles, smart electric grid systems and, in general, of portable high power consuming devices. One of the possible scenarios involves the use of carbon base materials showing high capacity. Presently, the research activity is strongly focused on porous carbon, carbon nanotubes (CNTs), nanofibers and graphene, as the most promising carbon based anode materials. The size reduction and the unique shape of these structures introduce novel properties that can substantially improve the energy storage capacity in LIBs.16, 17 For example, carbon nanorings (CNRs) with 20 nm outer diameters and 3.5 nm wall thickness, have shown outstanding performances as anode active materials: high lithium uptake and larger capacity, i.e. more than 1200 mAh g -1, and over a hundreds of cycles at the current density of 0.4 A g -1. Even at the higher current rates of 45 A g -1, it has been observed a capacity as high as 500 mAh g-1. The larger specific capacity and the high rate capability have been rationalized due to the reduction of the diffusion distance and the increase of the number of storage sites for Li-ions.18 These advantages are paradigmatic results of the developing of nanosized carbon based materials.

Titanium based oxides

Titanium based oxides have drawn signiﬁcant attention in the lithium battery community because they allow the designing of operational devices with minor safety concerns. Moreover, this class of active materials shows other interesting features

6 such as inexpensiveness, low toxicity, low volume change (2-3%) on both lithium insertion and de-insertion, along with an excellent cycling life. However, it shows low theoretical capacity, in the range 175-330 mAh g -1, and low electronic conductivity.

The electrochemical performance and the lithium insertion/extraction capacity of titanium based oxides mainly depend on their structure, morphology and size. In particular, it has been found that nanostructured titanium oxides lead to better capacity, longer cycling life and higher rate capability than the bulk materials. To date, titanium

dioxide with various allotropic forms and spinel Li4Ti5O12 have been extensively studied for anode purposes.

Spinel Li4Ti5O12 (LTO) is considered the most appropriate titanium based oxide material for lithium storage purposes because it exhibits excellent Li-ion reversibility at the high operating potential of 1.55 V vs. Li/Li+. Lithium insertion/extraction in

LTO occurs by the lithiation of spinel Li4Ti5O12 yielding rock salt type Li7Ti5O12.

During the insertion process, the spinel symmetry and its structure remain almost unaltered.19 The high operating potential guarantees safety conditions, in fact the formation of the solid electrolyte interface (SEI) is mitigated and the development of dendrites, typical issue in carbon based anodes, is avoided. However, the low theoretical capacity of 175 mAh g -1 and the low electronic conductivity 10-13 S cm-1 limit the full capacity at high charge discharge rates and reduce the Li-ion diffusion.20

To overcome these issues two approaches have been followed: the ﬁrst one is to enhance the electronic conductivity of LTO by surface treatments, the second one 7 consists in enhancing the Li-ion diffusion by downsizing the LTO to the nanoscale.

Nanocrystalline LTO, with particle sizes between 20 and 50 nm, has been synthesized by simple combustion method in short time period (less than 1 min).21 An experimental discharge capacity, i.e.170 mAh g-1, very close to the maximum theoretical value, was measured at 0.5C rate, while stable capacities of 140 and 70 mAh g-1 were observed at higher charge discharge rates of 10 C and 100 C, respectively. The doping of low conductivity materials, such as LTO, is one of the prominent techniques to enhance their electrical conductivity. In this regard, Shen et al. implemented a methodology for direct growing of LTO nanowires on titanium foil and they further showed an improvement in the conductivity of LTO nanowires by introducing Ti3+ ions through hydrogenation. The results were conﬁrmed through XPS technique.22 These nanowires containing Ti foil were directly used as electrodes without any conductive additives and binders, and they exhibited brilliant rate performance by reaching a capacity value close to the theoretical one, i.e. 173 mAh g

-1 at 0.2 C rate with good cycle life. Moreover, this value became 121 mAh g-1 at 30 C

(Figure 1.3). These nice results are related to the enhancement of the electron conductivity in hydrogenated LTO with respect to the pristine one.20

Anode materials based on alloying reaction

Materials which can satisfy the requirement of high capacity are, for example, silicon, germanium, silicon monoxide and tin oxide which react with lithium according to an

Figure 1.3 (a) Schematic representation of the fabrication of hydrogenated-LTO

(H-LTO); (b) Electrochemical performance of H-LTO and LTO nanowires. Copyright

Advanced Materials, 2012, 24, 6502-6506.

9 alloy/de-alloy mechanism. Their speciﬁc theoretical capacity ranges from 783 mAh g-1 for tin oxide up to 4211 mAh g-1 for silicon.23, 24 Even though these alloy based materials can provide larger speciﬁc capacity than graphite (372 mAh g-1) and LTO

(175 mAh g-1), the major drawbacks are the poor cycling life due to the high volume expansion/contraction and the larger irreversible capacity at the initial cycles. In order to overcome these issues, various approaches have been followed: the downsizing from micro to nanoscale particle and the fabrication of composites with both lithium active and inactive material, are the most promising. In the latter case the lithium active/inactive material serves as a conductive buffer matrix between the alloy materials and lithium source.25 Nanostructured alloy materials with different morphologies like nanowires and nanotubes have been considered as an implementable path to achieve high capacity with a good rate capability and long cycling life.26

-1 Silicon has both the highest gravimetric capacity (4200 mAh g , Li22Si5) and volumetric capacity (9786 mAh cm-3) among the anode material candidates.27, 28 In addition, the discharge (lithiation) potential of silicon is almost close to graphite i.e.

0.4 V vs. Li/Li+. Finally, Si is the 2nd most abundant element on earth, hence inexpensive and eco-friendly. The electrochemical lithiation of Si electrodes has been deeply investigated by many groups. It has been clariﬁed that the high speciﬁc capacity value is due to the formation of intermetallic Li-Si binary compounds such as

10 28 Li12Si7, Li7Si3, Li13Si4, Li22Si5. However, some issues prevent the employment of Si as anode in LIBs. Firstly, the large volume modiﬁcation (~400%) during the charge/discharge process causes poor cycling life and irreversible capacity. Secondly, the formation of Si compounds at the solid electrolyte interface inhibits the alloy/de-alloy process. It was evidenced that the electrical contact between the active material and both the conductive carbon and current collector undergoes a drop due to the large volume expansion/contraction of the Si anode, leading to an irreversibility of lithium insertion/extraction.29 Eventually, these volume changes result in shorter cycling life and capacity fading. To overcome these issues, a lot of efforts were focused on nanostructured Si, especially on their morphology aspect. For example, nanowires, nanotubes and nanospheres were considered due to their capability of providing the necessary free volume for accommodating the Si expansion during the alloying/de-alloying process. In particular, Si nanowires (Si NWs) and nanotubes have shown a reversible capacity over 2000 mAh g -1 with good cycling stability. The best results were obtained by the direct growth of Si NWs on the metallic current collectors, following a vapour-liquide solid approach (Figure 1.4). Such technique gives better accommodation to the large volume changes and all the nanowires participate to the lithium alloy process because of the direct contact with the current collector.30 This anode fabrication approach led to remarkable electrochemical performances.This anode fabrication approach led to remarkable electrochemical performances. In fact, the observed charge-discharge capacity was 3500 mAh g-1 over

11 20 cycles at 0.2 C rate, while at the higher current rate of 1 C the capacity value was set at 2100 mAh g-1 (Figure 1.5). This impressive lithium storage results were possible due to the good electronic contact between the current collector and the Si

NWs, to an efficient charge transport mechanism and to the small diameter of Si

12 breaking. This NW anode design has each NW connecting with the current collector, allowing for efficient 1D electron transport down the length of every NW. Copyright

Nature nanotechnology, 2008, 3, 31-35.

Figure 1.5 Electrochemical data for Si NW electrodes. a, Cyclic voltammogram for

4,200mAh g21 for Li4.4Si. c, The voltage profiles for the Si NWs cycled at different power rates. The C/20 profile is from the second cycle. d, Capacity versus cycle number for the Si NWs at the C/20 rate showing the charge (squares) and discharge capacity (circles).Copyright Nature nanotechnology, 2008, 3, 31-35.

NWs, which allows better accommodation of the volume change without any initiation of fractures.

13 Silicon monoxide (SiO)

SiO is considered an alternative choice to silicon, as anode candidate, due to its high theoretical capacity (>1600 mAh g-1). In addition, lithium oxygen co-ordination implies minimal volume change and, at the same time, lower activation energy.31 The mical reactions happening during the charge discharge process determine the conversion of SiO into Si and lithium oxides and, possibly, the formation of Si-Li or, alternatively, the direct formation of Si-Li alloy and lithium silicates. The mechanism can be schematized as follows:

SiO + xLi l Si + LixO

xLi + Si l LixSi

(or)

xLi + SiO l Lix-ySiz + LixSiOz

Solid SiO is thermodynamically unstable at all temperatures. Therefore, it can be

transformed into Si and SiO2 in a temperature triggered disproportionation reaction.

As previously pointed out about silicon, also SiO undergoes consistent changes in the volume during the lithiation and de-lithiation processes. Moreover, both the electrical conductivity and the lithium insertion/de-insertion rates remain quite poor. To solve these problems and to improve the reversible capacity as well as the cycle retention, different approaches have been tested. The most promising are carbon coating, electrochemical reduction of SiO with lithium and, ﬁnally, the size reduction of SiO.

14 Among these solutions, particle size reduction combined with carbon coating is an effective way to overcome the above issues by shortening the diffusion distance of

Li-ions and thus to enhance the ionic and the electrical conductivities.32 Carbon

Figure 1.6 (a) TEM image of SiOx nanoconifer on NiSix nanowires (NWs) (inset:

SEM image of SiOx/NiSix NW); (b) scanning TEM-HAADF images and EDX elemental colour map of single SiOx/NiSix NW; (c) electrochemical cycling stability along with coulombic efficiency and specific capacity vs. charge/discharge cycle numbers of SiOx/NiSix NWs (inset: galvanostatic chargeedischarge curves in the range of 0.01 and 1.2 V); (d) cycling performance of SiOx/NiSix NWs at different current rates.Copyright Journal of Power Sources, 2013, 229, 229-233.

15 coated 3-D porous SiO represents an effective way to achieve improved electrochemical conductivity and low volume expansion during the cycling process.

In fact, carbon coated porous SiO have shown an excellent capacity around 1490 mAh g-1 over 50 charge-discharge cycles at 0.1 C. A good response was also observed at the larger current rates of 3 C and 5 C with stable capacities of 1100 and 920 mAh

-1 g , respectively. Recently, structurally and kinetically stable hierarchical SiOx (0.9 < x < 1) nanoconifers were reported by Song’s group for applications as anode material

33 in LIBs. Columnar-shape SiOx nanoconifers were directly self-assembled on

metallic NiSix nanowires by means of chemical vapour deposition technique (Figure

1.6). These SiOx nanoconifers revealed stable charge-discharge performance over a hundred of charge-discharge cycles with speciﬁc capacity close to the theoretical one,

along with good coulombic efﬁciency and high rate capability compared to bulk SiOx.

This exciting result is due to the high conductivity of the metallic NiSix core, which

assists the efﬁcient charge transport pathway from/to SiOx phase.

Germanium(Ge)

Germanium is an extensively studied active anode material owing to its high lithium

-1 storage capability (1623 mAh g ) with Li22Ge5 as equivalent stoichiometry and reversible alloy/de-alloy reactions.23, 34 Even though Ge is considerably more expensive and shows lower capacity than silicon, it has desirable advantages such as high intrinsic electrical conductivity (104 times higher than silicon), higher capacity

16 than graphite anode and a narrow band gap (0.67 eV). Furthermore, it has been reported that the lithium diffusion into Ge is 15 and 400 times faster than in Si at 360 qC and at room temperature, respectively. This ensures higher rate capability and more efﬁcient charge transport than in silicon and graphite as well.35 Ge high power capability is then extremely important in advanced high power density applications such as electric vehicles. However, as discussed for silicon, the practical usage of Ge as active electrode in LIBs is hindered by the dramatic volume change (~300%) during lithium insertion/de-insertion.36 Ge nanostructures, such as nanoparticles, nanowires and nanotubes can effectively sustain the volume change with better efﬁciency than bulk and microstructures. Noticeably, improvements have been observed with hybrid composite of Ge nanoparticles using conductive matrices, obtained through simple preparation routes, for example solid state pyrolysis.12 For example, Ge nanoparticles, with diameter between 5 nm and 20 nm, were encapsulated inside carbon nanospheres with diameters in the range from 50 to 70 nm.

The role of the carbon nanospheres is to act as structural buffer and electro-active material during the lithium insertion and de-insertion process and to avoid direct contact with the electrolyte. This last aspect prevents Ge from the formation of SEI.

These obtained composites exhibited highly reversible lithium storage along with high rate capability. Nanoscale germanium oxides were also studied as novel anode

material for LIBs. Very recently, Seng et al. improved germanium dioxide (GeO2)

Figure 1.7 (a) Schematic representation of the lithium reaction mechanism in Ge,

GeO2/C and GeO2/Ge/C; (b) Cycling performance of all samples for 50 cycles at 1C rate; (c)rate performance of all samples from 0.1 to 10 C. Copyright Nano letters,

2011, 11, 3991-3997.

GeO2 from hydrolysis of GeCl4 and followed by carbon coating with acetylene gas,

18 resulting in GeO2/C. The last step was the reduction of GeO2/C to GeO2/Ge/C through high temperature was 650 qC. The analysis of the lithium battery performance was

carried out for materials such as GeO2/Ge/C, GeO2/C, nano GeO2 and bulk GeO2. The

-1 GeO2/Ge/C composite showed high capacities of 1860 and 1680 mAh g , with good cycling stability, at the current rates of 1 C and 10 C, respectively (see Figure 1.7).

Such a notable Li battery performance was related to the carbon coating of Ge, in particular it was found that carbon enhances the reversibility of the conversion

reaction of GeO2 with lithium.

Anode materials based on conversion reaction

In this section we will provide an overview on the transition metal compounds such as

oxides, phosphides, sulphides and nitrides (MxNy; M=Fe, Co, Cu, Mn, Ni and N=O, P,

S and N) when utilized as anodes in LIBs. The electrochemical reaction mechanism involving these compounds together with lithium, implies the reduction (oxidation) of the transition metal along with the composition (decomposition) of lithium

compounds (LixNy; N=O, P, S and N). Anodes based on these compounds exhibit high reversible capacities (500-1000 mAh g-1) owing to the participation of a high number of electrons in the conversion reactions. The electrochemical conversions reactions

+ - can be described as follows:MxNy +zLi +ze l4 LizNy + xM (Here M=Fe, Co, Cu,

Mn, Ni & N= O, P, S and N).

Iron based oxides have been extensively studied for rechargeable lithium batteries 19 because of their low cost, non-toxicity and high natural abundance. Iron oxides, both

haematite (D-Fe2O3) and magnetite (Fe3O4), are capable of participating in reversible conversion reactions with lithium, providing a theoretical capacity of 1007 and 926 mAh g-1, respectively.37 However, iron oxides tend to exhibit poor cycling performance due to low electrical conductivity, low diffusion of Li-ions, high volume expansion and iron aggregation during charging and discharging. Therefore, to overcome the above identiﬁed limits, many recent investigations have been focused on developing new methods for the preparation of iron oxide nanomaterials as well as for modifying their size, shape and porosity.38, 39 Other studies have been focused on methods to stabilize their structure and to improve the electrochemical kinetics and power capability, mainly by carbon coating or by carbon based composites of

39 D-Fe2O3 and Fe3O4. For example, Sohn et al., through an aerosol-assisted process,

followed by vapour coating, synthesized nanocomposites of Fe3O4 cores and porous carbon-silicate layers. Carbon was chosen for its characteristic of increasing the electrical conductivity, hence to mitigate the aggregation of iron oxides particles and to reduce the mechanical stress during cycling. These nanocomposites showed quite stable capacity around 900 mAh g-1 with a coulombic efﬁciency close to 100% over

50 charge-discharge cycles. Very recently, nanoparticulate Fe2O3 tubes have been obtained from microporous organic nanotubes (MONT) used as template.40(Figure

1.8) The prepared porous Fe2O3 nanotubes exhibited excellent electrochemical performances with large capacities such as 918 and 882 mAh g-1 at current densities

Figure 1.8 (a, b) TEM images of Fe2O3 nanotubes (FNTs) from MONTs template;

Li-ion battery performance of FNTs, (c) discharge capacity vs. number of cycles of

FNTs (red) and commercial Fe2O3 nanoparticles (black). The coulombic efficiency of

FNT is also included (violet); (d) galvanostatic chargeedischarge curves; (e) cycling performance of FNT at different current rates. Angewandte Chemie, 2012, 51,

6626-6630.

21 of 500 and 1000 mA g-1, respectively. These results indicate that low cost iron based oxides with highly conductive carbon composites can be a valid alternative to graphite anodes.

-1 41 Co3O4 and CoO show theoretical capacities of 890 and 715 mAh g , respectively.

Similarly to other kinds of materials, a number of different forms of cobalt oxides have been studied. Porous nanostructures, nanosheets, nanocubes, nanowires and nanotubes have been prepared by various synthetic routes such as wet chemical, solid-state, hydrothermal and microwave, therefore allowing the tailoring of their electrochemical performance. For example, Guan et al. developed a template free method for the preparation of pure phase CoO octahedral nanocages by simply

42 utilizing NH3 as a coordination etching agent. Herein, uniform size of the octahedral nanocages, with edge length of 100-200 nm, were obtained. In the test as anode material, they exhibited excellent cycling performance along with a good rate capability and enhanced lithium storage capacity, as seen in Figure 1.9. Even at the high current rate of 5C, these octahedral nanocages were able to deliver a reversible capacity of 474 mAh g-1. Their large capacity and the high rate capability were possible owing to the existing huge voids which can accommodate the large volume changes. Various other metal oxides, showing a conversion reaction mechanism with

lithium, have also been studied. For example NiO, MnOx, CuOx, MoOx, CrOx were extensively investigated as anode materials for LIBs and they a showed reversible capacity around 500 mAh g-1.43, 44 22

Figure 1.9 (a) Galvanostatic first chargeedischarge curves and (b) cycling performance for three CoO samples at the current rate of 0.2C. Copyright Chemical communications, 2012, 48, 4878-4880.

Metal phosphides have also been deeply studied as anodes for rechargeable lithium batteries.45-47 They can react with lithium, both in a conversion reaction schema and in an intercalation/de-intercalation reaction, depending on the nature of the transition metal and on phosphorous bonding stability upon electrochemical environment. It is

possible to divide MPx into two groups. The ﬁrst one involves the lithium insertion/extraction without breaking the metal-phosphorous bond, known as

23 + - insertion/de-insertion mechanism: MxPy + zLi + ze l LizMx-zPy, The second group involves a conversion reaction mechanism. In such electrochemical reactions the bonds between metal and phosphorous are broken, resulting in nanosized metal

+ - particles and Li-phosphides: MxPy + zLi + ze l LizPy + xM, copper, cobalt, iron, nickel and tin based phosphides are usually considered to belong to the second group,

i.e. conversion mechanism. Nevertheless, some reports showed that MPx could participate to the insertion as well as the conversion mechanism with respective to potential vs. Li/Li+. Metal phosphides can deliver high capacities between 500 and

-1 1800 mAh g . In addition, MPx as anode, show a high degree of electrons delocalization, which leads to low formal oxidation state and strong covalent bond of

M-Phosphorous. Another advantage is the lower insertion potential of MPx than their

oxide counterparts. However, MPx usually have low electrical conductivity and high

volume changes upon charge-discharge cycling. The use of MPx as anodes deserves

further exploration to overcome these drawbacks. For example, cubic Ni2P crystals

and lithiated NiP2, were prepared through the milling of a mixture of nickel powder and red phosphorus. Furthermore, Li-Ni-P ternary active composites were also

synthesized. Both Ni2P and Li-Ni-P showed reversible lithium reactions in electrochemical environment. Especially the ternary composites exhibited a high reversible capacity of 600 mAh g-1 by conversion process.48

Other kinds of materials, which have been thoroughly investigated for anode

applications in the LIBs ﬁeld, are the transition metal sulphides (MSx) and nitrides 24 (MNx). Iron, molybdenum, tin, antimony, nickel, cobalt and tungsten have attracted major interest due to their high lithium storage capability and structural advantages during the lithium insertion/de-insertion process.49, 50 As anticipated in the previous

sections, MSx and MNx reaction mechanisms with lithium involve the reduction to metal and, respectively, the formation of lithium sulphur and lithium nitride, through conversion reaction.For instance, Wang et al. reported phase controlled polyhedron

CoSx prepared by solid state reaction. The synthesis procedure was based on a high temperature Co-sulphur powders reaction in presence of potassium halide (KX).

51 Therein, KX served as reaction agent. In particular, CoS2 and CoS were the most promising composites exhibiting superior reversible lithium insertion and extraction reaction, hence resulting in capacities of 929 and 835 mAh g-1, respectively. In

another approach, Paolella et al. prepared nanoplates of Fe3S4 using sulphur and thiosulfate, octadecylamine, iron salt and 3-methyl catechol.52 Therein, 3-methyl catechol acted as phase control agent and growth moderator as well. Furthermore,

cyclic voltammograms of Fe3S4 electrode showed good reversible lithiation/delithiation process.53

1.1.3 Cathode materials for lithium-ion batteries

Although a wide range of cathode materials have been developed, the most promising

ones for practical applications include LiCoO2, Li(Ni1/3Co1/3Mn1/3)O2, olivine

LiFePO4, and spinel LiMn2O4.

LiCoO2 25

Figure 1.10 Lamellar structure of LiCoO2.

LiCoO2 (LCO) introduced by Goodenough is the ﬁrst and the most commercially

-1 + theoretical capacity of LiCoO2 is 270 mAh g when all Li ions are extracted from

the crystal. LiCoO2 materials are typically cycled between the fully lithiated discharge

state LiCoO2 and a half-delithiated charge state LixCoO2 (x=0.5 - 0.6, up to 4.3 V) to yield a useable speciﬁc capacity about 140 mAh g-1, above which the capacity reaches approximately 180 mAh g-1 up to 4.5 V charge cut-off. It has high theoretical volumetric capacity of 1363 mAh cm-3, low self-discharge, high discharge voltage, and good cycling performance. The major limitations are high cost, low thermal stability, and fast capacity fade at high current rates or during deep cycling. LCO cathodes are expensive because of the high cost of cobalt. Low thermal stability refers to exothermic release of oxygen when a lithium metal oxide cathode is heated above a certain point, resulting in a runaway reaction in which the cell can burst into ﬂames.

While this issue is general to transition metal oxide intercalation cathodes, LCO has the lowest thermal stability of any commercial cathode material. Although thermal

26 stability is also largely dependent on non-materials factors such as cell design and cell size, LCO typically experiences thermal runaway past ~200 qC due to an exothermic reaction between the released oxygen and organic materials. Deep cycling

(delithiation above 4.2 V, meaning approximately 50% or more Li extraction) induces lattice distortion from hexagonal to monoclinic symmetry and this change deteriorates cycling performance. Many different types of metals (Mn, Al, Fe, Cr) were studied as dopants/partial substitutes for Co and demonstrated promising but limited

performance. Coatings of various metal oxides (Al2O3, B2O3, TiO2, ZrO2) were more effective in enhancing LCO stability and performance characteristics even during deep cycling, because mechanically and chemically stable oxide material could reduce structural change of LCO and side reactions with electrolyte. Besides these,

nanostructured LiCoO2 including 1D nanowires, nanotubes, and 3D mesoporous structures have been prepared to improve the electrode/electrolyte interaction, and thus, increase the power density of the battery.

Spinel LiMn2O4 (4 V)

The spinel cathode LiMn2O4 originally proposed by Thackeray et al. has been extensively developed by the Bellcore laboratory. The anion lattice contains cubic

close-packed oxygen ions in the structure. The spinel LiMn2O4 is composed of Li, Mn, and O ions which are located at the tetrahedral 8a, octahedral 16d and 32e sites, respectively, as shown in Figure 1.11. In detail, the average oxidation state of Mn is ideally 3.5+, which means the molar ratio of Mn3+ and Mn4+ is 1:1. Hence, the product

27 would be stabilized in its stoichiometric form, namely, LiMn2O4. In fact, in view of their properties, that include material availability and environmental compatibility, these oxides appear as ideal cathode substitutes for the common, high-cost and

partially toxic lithium cobalt oxide, LiCoO2. Unfortunately, the wide practical use of the lithium manganese spinel electrode has so far been hindered by their severe

Figure 1.11 Crystal structure of spinel LiMn2O4. capacity fading on cycling, especially when operating at temperatures above ambient, e.g., above 40 qC.54 This limiting phenomenon is due to the manganese dissolution, in turn associated with the disproportionate reaction of Mn3+, i.e.: 2Mn3+ l Mn2+ +

Mn4+ . As well known, Mn2+ dissolves into the electrolyte solution while Mn4+ remains into the bulk of the electrode. These events are deleterious for the electrode performances. The dissolved Mn2+ ion deposits as metal Mn on the surface of the negative electrode, thus inhibiting its performance.55 The Mn4+ ion promotes ﬁlm

formation of Li2MnO3 that is electro-inactive since it induces discontinuity in the networks for electron transfer. As a consequence, various strategies have been 28 attempted to overcome this issue. A common one has involved the partial substitution of manganese ions with a series of foreign ions, such as Li, B, Mg, Al, Ni, Co, Fe, Ti, or Zn. The goal was to reduce the amount of Mn3+ with the ﬁnal aim of preventing

Mn2+ dissolution. However, the doping by these ions may result in capacity decrease and in environmental concern. An alternative approach consists in the surface modiﬁcation of the lithium manganese spinel by an oxide coating, including ZnO,

56, 57 Al2O3, Co3O4, NiO and BiOF. It has also been demonstrated that the capacity of

LiMn2O4 can be enhanced by reducing the particle size. Therefore, LiMn2O4 materials are able to lead faster diffusion kinetics through realization of nanostructure.58

Li(Co1/3Ni1/3Mn1/3)O2

The LiNi1-y-zMnyCozO2 compound, in which the metals Co, Ni, and Mn all are accommodated in the layered metal oxide structure, was been extensively investigated in the last few years and found to have properties that qualify them as possible

candidates, for the replacement of LiCoO2, especially LiCo1/3 Ni1/3Mn1/3O2 which was reported in 2001 by Ohzuku et al. In addition to their high lithiation capacities and reversibility, these compounds show higher thermal stabilities, promising electrochemistry and intriguing structural behaviour compared to the Co-free

compounds. LiCo1/3Ni1/3Mn1/3O2 possesses the same α-NaFeO2 type layer structure,

which can be considered as a 1:1:1 solid solution of LiCoO2, LiNiO2, and LiMnO2 or

a 1:2 solid solution between LiCoO2 and LiNi0.5Mn0.5O2 (i.e., LiCo1-2xNix MnxO2),

29 with Ni, Co, and Mn adopting valence states of 2+, 3+, and 4+, respectively. In

addition, Li 1-xCo1/3Ni1/3Mn1/3O2 exhibits a very tiny volume change (1–2 %) within 0 and 0.7 range for x. It also shows a low level of cation disorder. Both the stable structure and regular cation order contribute to its excellent electrochemical performance. The material shows high rate capability (200 mAh g-1 at 18.3 mA g-1

-1 -1 and 150 mAh g at 1600 mA g ). Moreover, LiCo1/3Ni1/3Mn1/3O2 has excellent safty

at a high state of charge compared to LiNiO2, LiCoO2, and LiNi0.8Co0.15Al0.05O2, which is another attractive property suitable for the large scale practical applications, especially as a potential cathode candidate for electric vehicles. 59

Olivine LiFePO4

Many studies had been directed towards iron oxides as potential cathode materials for

+ LIBs. However, layered LiFeO2 has shown poor Li insertion ability. In 1997, there

was a discovery of electrochemical-active olivine phase, speciﬁcally, LiFePO4 by

Padhi and Goodenough. This is the ﬁrst cathode material emphasizing substantially

low cost and earth-abundant element of iron, Fe. For LiFePO4, the crystal structure is shown in Figure 1.12a in which P occupies tetrahedral sites, Fe occupies octahedral

sites and Li forms one-dimensional chains along the [010] direction. LiFePO4

2+ 3+ delithiates to FePO4 as Fe is oxidized to Fe and a miscibility gap exists between

FePO4 and LiFePO4, as schematically shown in Figure 1.12b. During the lithiation

process, FePO4 is reversed back to LiFePO4 upon lithiation. Electrochemically,

30 + LiFePO4 has a flat discharge voltage of 3.4 V vs. Li /Li and a theoretical capacity of

-1 -1 170 mA h g , which is higher than available capacity of LiCoO2 (140-150 mAh g ) with good cyclability due to its stable structure supported by the P-O covalent bond.

Figure 1.12 (a) Crystal structure of LiFePO4. (b) Schematic representation of the processes during charge/discharge of LiFePO4.

Since then, many researchers have improved its signiﬁcant drawback, in particular

poor rate capability stemming from low electric conductivity because LiFePO4 is intrinsically an electrical insulator. In previous attempts to enhance the electronic conductivity, coating a conductivity layer like carbon, conductivity polymers seem to be a very promising way to overcome the limited rate capability, because the conductivity layer provides a pathway for electron transport, resulting in improvement of the electric conductivity. To alleviate the low ionic conductivity, constructing

nanostructured LiFePO4 to shorten the diffusion distances and increase the

electrode/electrolyte interaction has been widely investigated. For example, LiFePO4 with an average size of 140 nm has been prepared, which delivered a specific capacity 31 of 147 mA h g-1 at 5 C rate and a good cycle stability with no significant capacity loss after more than 400 cycles at 0.5 C rate.

32 1.2 Lithium-sulfur Batteries

The energy storage system market experiences booming development in the past few decades due to the wide range application demands spaning from smart power grids, electric vehicles to portable electronic devices. It is reported that the energy storage system market is expected to increase from 27 GW h in 2017 to 90.9 GW h in 2020, keeping an annual growth rate of 21.7%.60 These new demands require the energy storage systems with a much higher energy density than is currently available, which will inevitably reach the toplimit at which conventional lithium ion battery system cannot meet. Compared with the established batteries systems, lithium–sulfur (Li–S) batteries with high theoretical energy density of 2567 W h kg -1, which is estimated to be 5 times higher than that of state-of-the-art lithium-ion batteries based on metal oxide cathodes and graphite anodes. (Figure 1.13a) The practical energy density of

−1 lithium-sulfur batteries in the future is estimated to be 500–600 Wh kg , more than twice that of state-of-the-art LIBs, showing great potential for electric vehicles to be driven 500 km after fully charging.60-62 Coupling with cost effectiveness, nature abundance, and nontoxicity, making lithium sulfur batteries promise great potential to be the next generation large scale energy storage system.6, 7, 63-75(Figure 1.13b)

2985 (c) Energy density of various electrochemical storage systems. Copyright Adv.

Energy Mater. 2015, 5, 1401986

1.2.1 Electrochemistry of lithium-sulfur batteries

The working principle of a Li–S battery is fundamentally different from the intercalation process that occurs in conventional Li-ion batteries, it involves a

multi-electron transfer electrochemistry based on the conversion reaction of S8 + 16

34 + − Li + 16 e → 8 Li2S, where simplified reaction sequence of S8 → Li2S8 →

Li2S6/Li2S4 → Li2S2/Li2S has been generated as shown in Figure 1.13c, long chain

polysulfides (Li2Sx (4 ≤ x ≤ 8)) are easily dissolve in the organic electrolyte, while

60, 76-78 lithium sulfides (Li2S2, Li2S) do not. The main obstacles impede the commercialization of lithium sulfur batteries are:(I) The insulating nature of sulfur

and lithium sulfides. (II)The dissolution of intermediate polysulfides (Li2Sx (4 ≤ x ≤

8)) in the electrolyte. (III) Large volumetric expansion (80%) of sulfur upon lithiation.

(IV) The use of a metallic lithium anode.64, 79-86 To tackle these challenges, extensive researches have been conducted on cathodes, electrolytes, separators, and lithium-metal anodes for preventing the dissolution of polysulfides and facilitating the utilization of sulfur. In this thesis, we mainly focused on the research of cathode materials for lithium sulfur batteries.

1.2.2 Categories of lithium-sulfur batteries

As for the cathodes of lithium sulfur batteries, the starting exist states of sulfur are

generally divided into four types: elemental sulfur (S8), small sulfur (S2-4), lithium

sulfide (Li2S), catholyte (Li2Sx, 4≤x≤8). There are various type of cathodes that using

87 elemental sulfur(S8) as the starting materials, such as carbon based materials, metal

88 89 90 91 92 93 oxides, metal sulfides, metal nitrides, MOFs, MXene, C3N4, et al.(Figure

1.14a) In most conditions, sulfur exists in the form of cyclo-S8 molecule with the dimension of 0.7 nm,80 exhibits multi-step reduction sequences of

35 S8՜Li2S8՜Li2S6/Li2S4՜Li2S2/Li2S with two typical plateaus, corresponding to the

transformation of cyclo-S8 to high-order lithium polusylfides (2.4 mV), high-order lithium polysulfides to lithium sulfides (2.1 mV), respectively. During this process the theoretical capacity generated is around 1672 mAh g-1, this type of sulfur cathode possesses the advantage of relatively high discharge voltage plateaus and potentially high sulfur contents in the cathode.94-96 However, this system still bothers by the concern of polysulfides dissolution, electrodes mechanical damage caused by the large volume expansion of sulfur in the lithiation process. Apart from that, metallic lithium used as an anode also gives rise to the dendrite formation and security problems.72, 75, 88, 97, 98

Small sulfur molecule refers to the short-chain sulfur (S2–4) that are distributed in the

conductive matrix at molecular levelˈwhich directly transition from S2–4 to the Li2S

2- phase avoiding the unfavorable transition between S8 and S4 .(Figure 1.14b) The short chain sulfur in this regard is compatible with carbonate electrolyte, thus no special electrolyte is required in this respect, which could effectively alleviate the shuttle effect, contribute to the high Coulombic efficiency and stable capacity retention. There are two main types of short-chain sulfur-based composite cathode materials, microporous-carbon/small-sulfur99 and sulfurized carbon.100, 101 Sulfurized carbon is a compound in which short sulfur chains are covalently bound to the surface of carbon particles.101 For instance, Guo and co-workers proposed small sulfur

102 allotropes S2–4 (S2, S3, and S4 ) as the reactant phases in cathode materials. An 36 microporous carbon matrix was prepared via the pyrolysis of D-glucose with pore size of 0.5 nm was adopted to trap the metastable small sulfur molecules with a theoretical sulfur loading of 49 wt%. Wang’s group applied the elemental sulfur to dehydrogenate polyacrylonitrile (PAN),103 the polar side chain -CN groups in the

PAN cyclize to form a heterocyclic compound in which sulfur is embedded and incorporated, resulting in the formation of the sulfurized PAN. However, it should be noted that this small sulfur-based cathodes have obvious problem that the sulfur content is normally

Figure 1.14 Structures of (a)elemental sulfur, (b)small sulfur cathode materials and their schematic discharge/charge profiles Copyright Adv. mater, 2017, 1606823. 37

Figure 1.15 Structures of (a) Li2S, (b)catholyte cathode materials and their schematic discharge/charge profiles Copyright Adv. mater, 2017, 1606823. less than 50% in the cathode, and the relatively low discharge voltage plateau (1.9 V) also results in low energy density.

Recently, the lithium sulfide (Li2S) cathode, with a theoretical capacity of 1166 mAh g−1, has attracted much attention due to it is compatible with safer Li-free anodes (e.g., graphite or Si/Sn-based composites)(Figure 1.15a).104-110 In addition, since Li

incorporates into the structure of Li2S, the electrode shrinks in delithiation and provides sufficient space to endure volume expansion of sulfur in lithiation process, makes it easier to form a stable electrically insulated SEI, which is permeable to Li ions but impermeable to polysulfides, thus not only eliminating the irreversible side

38 reactions between polysulfide anions and Li metal, but also avoiding the gradual

111-114 cathode dissolution. However, the high melting point of commercial Li2S

(940 Ԩ), its low solubility in organic solvents, and very sensitive to moisture, making it unstable in the air and difficult to infiltrate it into conductive carbon hosts via

115 conventional approaches such as ball milling method or disperse Li2S into the anhydrous ethanol solution.109, 115, 116 Apart from that, the high electronic resistivity

-13 -1 and poor ionic conductivity (ca. 10 S cm ) of Li2S, also makes this cathode

electrochemically inactive. Cui et al. have found that, for micrometer-sized Li2S particles, a large potential barrier (ca. 1 V) should be overcome at the beginning of

104 charging process. In 2014, Cui’ group innovatively encapsulated Li2S with TiS2 which exhibits high conductive and have affinity for polysulfides, but the prepare processes are too complicated that requiring high annealing temperature. To enhance

the Li storage kinetics, a combination of Li2S with electronic conductors is highly desired in preparing the cathode materials. Recently, Khalil Amine’s group invented a

117 novel method by burning lithium foils in a CS2 vapour. The obtained structure

features crystalline Li2S nanoparticles wrapped by few-layer graphene, this electrode exhibits a high reversible capacity of 1,160 mAh g-1 at a loading of 10mg cm-2, this work overcomes the major constraints associated with traditional sulfur electrodes

and previous reported Li2S composites, throwing new light for our research on Li2S based cathode. More recently, Arumugam Manthiram’s group proposed a novel

strategy that forming encapsulation layer comprised of MnS on the surface of Li2S

39 bulk particles via in situ surface chemical reactions between Li2S and electrolyte

118 additives containing Mn(acac)2, It also confirmed that the electronic band structure

of this encapsulation layer is vital to the initial charging resistance of the Li2S bulk particles, which leads to improved material encapsulation and capacity retention. This work opens up new prospects for tuning the properties of electrode materials by selective modifications of the surface electronic and chemical structures desired for energy-storage applications.

Another type of lithium sulfur batteries are based on lithium polysulfide cathodes, instead of applying solid sulfur as starting active material followed by melt-diffussion or liquid infiltration methods, these dissolved lithium polysulfides have been proved to have better kinetics comparing to the sluggish reaction of insulating solid phases, and can also disperse into the host matrix uniformly and be reached by both Li+ ions and electrons readily, thus resulting in improved sulfur utilization.119, 120 121, 122(Figure

1.15b) However, although the utilization of soluble polysulfides as cathode materials can achieve a high reaction activity, it may decrease the volumetric energy density of

Li-S batteries because the increased mass of electrolyte contributes to the inactive mass of the cell.123 In addition, some severely practical problems also arise during synthesis and battery assembly processes due to the extraordinarily hygroscopic feature of polysulfides. In recent research, some workers use surface modification to address the polysulfide deposition issues. For instance, Cui’s group prepared carbon nanofibers with tin-doped indium oxide nanoparticles (exhibits a hydrophilic surface) 40 decorating the surface as hybrid three-dimensional electrodes which using Li2S8 as the catholyte 124. A reversible specific capacity of 710 mAhg-1 could achieve over 500 cycles. Arumugam’s group utilized a lightweight 3D nitrogen sulfur codoped graphene current collector as the additive/binder-free electrode structure to

accommodate large amounts of Li2S8 catholyte, the maximum sulfur loading is as much as 8.5 mg cm-2,120 capacity stabilizes at around 670mAh g-1 after 200 cycles.

Huiming Chen’s group reported a 3D conductive porous vanadium nitride

nanoribbon/graphene composite accommodating the Li2S6 catholyte as the cathode of a lithium–sulfur battery.125 The capacity could be retained at 1252 mAh g-1 after 100 cycles at 0.2C, offering a potential for use in high energy lithium–sulfur batteries.125

These pioneer work intrigue us to do more exploration on lithium sulfur batteries.

1.2.3 Cathode materials for lithium sulfur batteries

Carbon based materials as cathode materials for lithium sulfur batteries

Enormous researchers have been devoted to the development of cathode materials for high performance lithium-sulfur batteries, considerable groups used to focus on encapsulating sulfur particles into various carbon matrix materials due to their good compatibility with sulfur. For instance, Nazar’s research group firstly confirmed that by encapsulating sulfur into nanostructured mesoporous carbon (CMK-3), high reversible capacities and good rates can be obtained due to the enhanced conductivity and physical confinement of the polysulfides via the mesopores. Since then,

41 significant advances have been achieved by using multiple carbon materials as hosts for sulfur cathodes, including carbon nanotube/fibers, graphene coated hybrid structures, hollow carbon structures and composite carbon structures. However, due to limited sites to strongly anchor polar molecules like lithium polysulfides and lithium sulfides, non-polar carbon materials do not exhibit specific bonding configurations with LiPSs: very low binding energies (0.1–0.7 eV) are observed and are attributed to van der Waals forces. To add anchoring sites and enhance anchoring ability of the carbon, tunable polar sites that can chemically confine the polysulfides were introduced. Nanostructured carbon doped with heteroatoms (N, S, P, B), and carbon modified with functional groups (amino-functionized, carboxyl-functioned, sulfonated, fluorinated, hydroxylated) as well as combined with conductive polymers, have all been widely investigated for obtaining high performance Li–S batteries.

Nevertheless, the N-, O- or S- containing polymers/molecules exhibits a relatively low binding energy of 0.5–1.3 eV in the form of Li-X (X=N, O, S). Although carbon materials doped with heteroatoms (N, O, S) show increased binding energies in the range of 1.3–2.6 eV, this value still far from satisfactoriness. Resulting in undesired long term cycle performance as well as high-rate capacity.

Metal oxides as cathode materials for lithium sulfur batteries

To further modulate the binding energy with polysulfides and increase the tap density of electrodes, nanostructured polar inorganic compounds like transitional-metal

42 oxides have appeared as polar host materials for lithium sulfur batteries. Metal oxides that typically contain an anion of oxygen in the oxidation state of O2− are always with a strong polar surface. These metal oxides are directly added to the cathode with the role of an additive which is usually less than 10 wt%. Metal oxides afford abundant polar active sites for absorption of polysulfides. The volumetric energy density of Li–

S cells is significantly improved through the engineering fabrication of nanostructured metal oxides and sulfur composites due to the high intrinsic density of oxides in comparison with nanocarbon materials.

Ti4O7 is a member of the TinO2n−1 Magnéli phases, substoichiometric compositions of metallic titanium oxides that form a homologous series between the end members of

−1 TiO2 and Ti2O3, due to the high electrical conductivity (>103 S cm at room temperature) and high affinity for polysulfide derived from the polar O-Ti-O units, making it is promising to be applied as cathode material for Li–S batteries. Nazar and

co-workers have reported a pioneering electrochemical study on the Ti4O7 particles without interconnected pores structure. Their particles exhibited a surface area of 290 m2g-1 and the electrodes delivered a capacity of 1000 mAh g-1 at 0.2 C rate. Another

work reported by Caruso and co-workers used porous Ti4O7 particles with a surface area of 197.2 m2 g-1 as cathode materials. They present a high specific capacity of

1190 mAh g-1 at 0.2 C rate in the beginning. Recently, Yan Lu’s group designed

Figure 1.16 (a) TEM image of the Ti4O7 particle; (b) enlarged zone in (a); (c) TEM image of the Ti4O7 particle coated with a thin layer of carbon; (d) enlarged zone in (c);

(e) Rate capabilities of the Ti4O7 based cathode and the carbon-coated Ti4O7 based cathode at different current rates (0.1, 0.2, 0.5, 1, and 0.1 C). (f) Cycling performance of the Ti4O7, carbon-coated Ti4O7 and TiO2 based cathodes with ≈50 μL of electrolyte over 300 cycles at a charge/discharge rate of 0.1 C. The solid squares represent capacity and hollow squares represent Coulombic efficiency. Copyright Advanced

Functional Materials 2017, 27, 1701176.

44 Ti4O7 particles with interconnected-pore structure (Figure 1.16), this Ti4O7 have the advantage of mesopores for encapsulating of sulfur and provides a polar surface for chemical binding with polysulfides to suppress their dissolution. Moreover, a thin

layer of carbon is coated on the Ti4O7 surface without destroying its porous structure.

−1 This carbon coated Ti4O7 cathode displays the highest capacity of 1411 mAh g at

0.1 C, and a good cycling stability with a capacity decay of 0.099% per cycle over

300 charge/discharge cycles.

Figure 1.17 The use of MnO2 for Li–S batteries. (a) Visual confirmation of polysulfide entrapment of MnO2 at specific discharge depths (75S/KB, top; 75S/

MnO2, down. (b) The ex situ XPS of S/MnO2 electrodes after discharge to specific states: from top to bottom: dis-charged to 2.15 V, discharged to 2.15 V and then aged in the cell for 20 h, discharged to 800 mA h g−1 and discharged to 1.8 V. Copyright

Nature.Communication. 2015, 6, 5682-5687.

45 Using MnO2 as a host material has been proposed recently. MnO2 is always characteristically nonstoichiometric and is deficient in oxygen atoms. Liang et al. firstly designed a highly efficient polysulfide mediator monoclinic potassium

birnessite G-MnO2 for Li–S batteries. This as-obtained 75S/MnO2 nanocomposite

−1 (containing 75% sulfur) cathode displayed an initial capacity of ≈1300 mA h g at

−1 −1 C/20, 1120mAhg at C/5, and 950mAhg at1.0 C. An in-situ visual–electrochemical

investigation with 75S/MnO2 and a control sample 75S/KB (75% sulfur and 25%

Ketjen Black) electrode was conducted. Figure 1.17a illustrates that the electrolyte in the 75S/KB cell changes from colorless to bright yellow-green on partial discharge of the cell over 4.0 h. At the end of the discharge (12.0 h), the electrolyte is still yellow,

indicating the polysulfides remain in solution. In contrast, in the 75S/MnO2 cell, the electrolyte exhibits only a faint yellow color at 4.0 h. On full discharge, the electrolyte

is rendered completely colorless. The comparative experiment reveals that MnO2 is

effective in conversion of polysulfides into insoluble reduced species of Li2S2/Li2S.

XPS study of the interaction of lithium (poly-) sulfides and MnO2 nanosheets is shown in Figure 1.17b. Two are terminal and bridging S environments, which are the

same as in Li2S4. The S 2p3/2 peak at 167.2 eV corresponding to the S = O sulfur in

2− thiosulfate ([SSO3] ) arises from a surface redox reaction between Li2S4 and G-MnO2.

3+ The peak at 641.4 eV is from the Mn contribution in the Mn 2p3/2 XPS spectrum and two additional Mn 2p3/2 peaks arising at lower energy (640.4 and 639.4 eV) are

46 2+ readily attributable to Mn . The existence of thiosulfate and the reduction of Mn ions

in G-MnO2 confirm the reactions between the polysulfides and the G-MnO2.

Wang and co-workers demonstrated a ternary hybrid material consisting of highly conductive CNTs as an electron-conduction framework, nonconductive spinel

NiFe2O4 as a polysulfide absorber, and sulfur as the electrochemical active material.

The 2D metal oxide nanosheets afford strong binding sites for polysulfides and thus

restrict shuttling. At 0.1C, the cell with the CNT/NiFe2O4–S ternary material as the

−1 cathode exhibited a high reversible specific capacity of 1350 mA h g . Reversible specific capacities at different current densities of 0.1, 0.2, 0.5, 1, and 2C were 1350,

Figure 1.18 Morphological and structural characterizations of MMNC and

CeO2/MMNC nanospheres. (a) SEM image of MMNC nanospheres. (b,c) SEM and

47 −1 1200, 1050, 900, and 700 mA h g , respectively. Long-term cycling performance of

−1 the ternary material was tested. A specific capacity higher than 850 mA h g was retained after 500 recharging cycles, corresponding to a capacity decay as low as

0.009% per cycle with excellent Coulombic efficiency higher than 99.2%.

Zhong Jin demonstrated a design of sulfur host by implanting CeO2 nanocrystals homogeneously in the hierarchical pores of bimodal micro-mesoporous nitrogen-rich

carbon nanospheres (CeO2/MMNC) for the effective confinement of sulfur species.

This hybrid structure benefits from both the physical confinement of the polysulfides

by the mesoporous carbon microspheres and their chemical bindings to the CeO2

nanocrystals and N-doped carbon species (Figure 1.18). More importantly, the CeO2 nanocrystals can promote the chemical redox reactions of polysulfides, thus

significantly enhancing their retentions upon cycling. The CeO2/MMNC cathodes with 1.4 mg cm-2 sulfur exhibit high reversible capacities (836 mAh g-1 at 1.0 C after

500 cycles) and good rate capability (737 mAh g-1 at 2.0 C), and high cycle stability

(721 mAh g-1 at 2.0 C after 1,000 cycles with a low capacity decay of 0.024% per cycle). Furthermore, a high and stable reversible capacity of 611 mAh g -1 is achieved after cycling for 200 cycles with higher sulfur loading of 3.4 mg cm-2.

On the other hand, other metal oxides with polar surfaces such as SiO2, Al2O3, La2O3,

MoO2, V2O5, SnO2, Fe2O3, MgO, Co3O4, Nb2O5 have been widely introduced as cathode materials for lithium sulfur batteries, but the application of meal oxides is still

48 restricted by the poor conductivity and low energy density. Based on this condition, most oxide nanostructures are coupled with conductive polymers or carbon materials to enhance the overall conductivity of the cathode rather than relying on the intrinsic

conductivity to attain the best service performances in the Li–S batteries. Al2O3 was coated on carbon/sulfur cathode materials through the ALD method, the casting method, and the mixing method. These studies provide new materials for consideration to effectively anchor polysulfides in the cathode of a Li–S battery.

MOFs as cathode materials for lithium sulfur batteries.

Metal–organic frameworks (MOFs) are a class of porous materials assembled by connecting metal ions and organic linkers with tremendous extensiveness in multiplicity, their tunable chemical composition is favorable and can be designed at the molecular level, thus there is great opportunity to rationally design and systematically adjust for effective MOFs as sulfur host and fully connected the mobile di-electron redox centers (Figure 1.19).126-129 Moreover, their highly porous framework allows fast mass/ionic transportation of the relevant species. When compared with the carbonaceous cathode materials, the pores of MOFs can be decorated with chemically active sites, such as Lewis acidic sites and functional organic groups. These active

49 Figure 1.19 Unit cell of different types of MOFs that applied in lithium sulfur batteries sites can provide chemical affinities to sulfur and polysulfides in MOFs that are superior to those of porous carbon materials.

For instance, MIL-100 (MIL: Materiaux Institut Lavoisier) solids, are inorganic– organic hybrid materials composed of trimesic acid ligands and metal(III) octahedral, exhibit a zeotype architecture with two types of mesoporous cages, high surface areas and a significant number of accessible Lewis metal sites, making this unique porous structure of MIL-100 solids are favourable for sulfur impregnation. In 2011, MIL-100

50 Figure 1.20 (a) Cycling performace of S@MIL-100(Cr)ˈCopyright Journal of the

American Chemical Society, 2011, 133, 16154-16160. (b) Cycling performane of

Science, 2014, 7, 2715. (d) Comparision of binding energies of lithium polysulfides to

Ni-MOF and Co-MOF, and cycling performance of Ni(II)-based and Co(II)-based

(Cr) was the first to be used as host for Li-S batteries because it was evidenced a weak bonding between polysulfide anions and MIL-100 (Cr) framework,130 coupling with

51 the two types of mesoporous cages (׽ 25- 29 Å) connected through microporous

pentagonal (׽ 5Å) and hexagonal windows (׽ 9 Å) that could reversibly capture and release polysulfides during cycling, making it a suitable candidate for lithium sulfur battery. But the sulfur content in the sulfur cathode was relatively low (48% of the

MIL-100- (Cr)/S@155), and the capacity could only retained at 461 mAh g-1 after 50 cycles under the current rate of 0.1C (Figure 1.20a) making it difficult for evaluating the effective function of MOF. Afterwards, HKUST-1 was reported to play an important role in slowing the release of sulfur into the electrolyte by the strong confinement effect,131 that derived from suitable pore space and the open Cu2+ sites.

However, due to the low electronic conductivity, the capacity was limited to 500 mAh g-1 after 170 cycles (Figure 1.20b). In the meantime, Li’group thoroughly proposed that both the aperture window and particles size play important roles on the rational design for MOFs hosts in high performance lithium sulfur batteries.132 As demonstrated, the particle sizes of MOF host mainly affect the internal Li+/e- transport, thus in turn determine the utilization of sulfur and initial discharge capacity. While the aperture window that associated the functionalities in the open framework dominates the escape diffusion of polysulfides, can affect the cycling stability. In this paper the author synthesized ZIF-8 with particle sizes of 150nm, 1um, 3um and

compared ZIF-8 (3.4 Å) with other HKUST-1, NH2-MIL-53 and MIL-53 MOFs hosts with the aperture windows of 6.9 Å, 7.5 Å, 8.5 Å, respectively. Conluding that ZIF-8 with the aperture windows of 3.4 Å and particle size of 100-200 nm exhibits

52 prominent rate capabilities with exceptional capacity retention (Figure 1.20c). This work threw light on preparing small nanoparticle of MOFs with giant cage-type pores that have small apertures to increase sulfur loading, in order to achieve high cycling performance. These former literatures are the first stage development of MOFs mainly focused on the physical confinement for lithium sulfur batteries.

In the second stage, Xiao et al started to explore the electrochemical reaction of sulfur and MOFs from the perspective of “Lewis acid-base interaction” concept.128 As reported in the work, the Lewis acidic metallic center is inclined to coordinate with soluble polysulfide anions (soft Lewis base) as axial ligand, which could effectively entrap the soluble polysulfides within the cathode. Based on this mechanism, the author also put forward an order to form stable complexes with metallic centers: Mn

(II)˘Fe(II) ˘ Co(II) ˘Ni(II) ˘Cu(II), providing new insight on the choice of

MOFs hosts for Li–S batteries. This conclusion can also be confirmed by the cycling

performance of Ni6(BTB)4(BP)3 and Co(BTB)4(BP)3 as hosts for lithium sulfur batteries (Figure 1.20d). Although Co-based MOF has improved electronic conductivity than Ni-based MOF, it shows inferior cycling performance due to the weaker coordination between center Co(II) and polysulfides. After that, Guodong

Qian and Co-workers proposed that the strong interaction between sulfur species and

MOFs are closely associated with the number of Lewis acidic sites in the MOFs.133

MOF-525(2H), MOF-525 (FeCl), MOF-525(Cu) can respectively provide zero, one and two Lewis acidic sites for the binding and inclusion of sulfur, these also 53 consistent with the shifts in the corresponding XPS binding energies, no obvious chemical shift of the S 2p spectra was observed in S@MOF-525(2H) compared with pristine sulfur, while large S 2p chemical shifts of about 0.7 eV to a lower energy were observed in S@MOF-525(FeCl) and S@MOF-525(Cu). Meanwhile, the corresponding cycling stability of these three electrodes also confirms the difference between the various interaction (Figure 1.20e). Recently, some researchers also demonstrated that MIL-100(V) is attractive for lithium sulfur batteries as it is different from other members in MIL-100 solid,134 it contains V3+, V4+ in its vanadyl groups. Because different valence states of vanadium ions can offer various Lewis acid sites and form multiple strength chemical interactions with sulfur and lithium polysulfides compared with metal–organic frameworks with a single valence state of metal ions.135 In fact, the cycling stability of S@MIL-100(V) is evidenced superior to that of S@MOF-525, indicates the potential applications of S@MIL-100(V) as hosts for Li-S batteries (Figure 1.20f).

More Recently, Dawei Su et al demonstrated sodium iron cyanide

91 (Na2Fe[Fe(CN)6]), can confine polysulfides via Lewis acid–base bonding derived from the Fe(II)-S interactions this strong interaction was evidenced by the shifts of bonding energy in XPS of Fe spectrum (Figure 1.21a). The corresponding

UV-Spectrum (Figure 1.21b) and the Ex-situ Raman further proved that polysulfide anions formed during the discharge process and were effectively entrapped on the

cathode. The author also proposed the affinity between Na2Fe[Fe(CN)6] framework 54 and polysulfide species becomes stronger with the increase of the polysulfide chain

length, because the coordinated Fe in the Na2Fe[Fe(CN)6] is a soft Lewis acid and prone to coordinated with soft Lewis base. The longer chain polysulfides exhibits a

Figure 1.21 (a) XPS spectra of the Na2Fe[Fe(CN)6] and S@Na2Fe[Fe(CN)6] and high resolution XPS spectra of Fe, N, C, and S of the Na2Fe[Fe(CN)6] and

S@Na2Fe[Fe(CN)6]. (b) Photo of 5 Mm pristine Li2S6 solution and upper solution after soaking the Na2Fe[Fe(CN)6] nanocrystals. (b) UV-Vis spectra of 5 mM pristine

Li2S6 solution and upper solution after soaking the Na2Fe[Fe(CN)6] nanocrystals. (c)

Cycling performances of S@Na2Fe[Fe(CN)6]@PEDOT composite at 5 C current rate.

Advanced materials, 2017, 1700587

55 stronger Lewis base feature, resulting in a higher binding energy with

Na2Fe[Fe(CN)6]. Meanwhile, it also worth noting that the large interstitial sites of the

5.27*5.27 Å (S8 has 4.65*4.72 Å) in the Na2Fe[Fe(CN)6] framework benefits to

tolerate the expansion of the pore content on the full lithiation to Li2S as only 66.7

vol.% large interstitial sites will be occupied by the Li2S, implying that this pore structure is ideal for alleviating the volume expansion during sulfur reduction. Next, the author further calculated out that only ~59.4% of sulfur can be stored within the

large interstitial sites of Na2Fe[Fe(CN)6] framework, and around 22.6% of sulfur covered on the surface of the nanocubes that could not be efficiently utilized. Based on this condition, the author did PEDOT coating to accelerate electron transport and prevent the dissolution of polysulfides. As a result, both the 82 wt.%

S@Na2Fe[Fe(CN)6] and S@Na2Fe[Fe(CN)6]@PEDOT electrodes can maintain the high capacity of 763 and 1101 mAh g-1 at 0.1C after 100 cycles (Figure 1.21c), respectively.

In conclusion, MOFs exhibit great potential in the field of lithium sulfur batteries due to its controllable pore structure and tunable chemical composition. The irreplaceable advantages of MOFs is that both the open metal centers serve as Lewis acid sites and the heteroatom dopant sites (N, S and P) in MOFs show strong affinity to polysulfide anions. Especially, the different valence states of metal ions in the certain MOFs can offer various Lewis acid sites and form multiple strength chemical interactions with sulfur and lithium polysulfides, this throws new lights for our future research. 56 Moreover, the electrochemically catalytic character of MOFs for accelerating the conversion of polysulfides are expected to be developed in the lithium sulfur batteries.

However, by comparison recent reported MOFs as sulfur host literatures, the average sulfur loading level in each electrode is less than 40% and the cycling stability are the main barriers that restricted MOFs further development in lithium sulfur batteries.

Specially, the existence of organic linkers leads to the limited electronic conductivity, restricting the kinetic reactions and effective conversion of polysulfides. Apart from that, manufacture of controllable porous structure and morphology that are suitable for high amount sulfur accommodation are still challenging for MOFs or

MOFs-derived composites, in which the uniform distribution between metal particles and host frameworks is not easy to achieve. At the same time, the bonding effect, component proportional of metal sites in MOFs still need comprehensive understanding in order to balance optimization between polysulfide adsorption and diffusion on the surface of the frameworks.

Metal sulfides as cathodes materials for lithium-sulfur batteries

136 137, 138 Although some metal oxides such as Ti4O7, MnO2, and tin-doped indium oxide124 have proved to effectively retain polysulfides without sacrificing the conductivity of the whole electrode. However, the cycling stability of these electrodes with high sulfur loading (> 5 mg cm−2) has not yet been reported due to the high density of the host materials and lack of enough active sites for ion/electron transport.

57 Therefore, exploring conductive polar host materials with adequate active sites is highly desirable for high-energy Li-S batteries. The use of nickel139, 140 and cobalt129 as the ‘‘catalysts’’ for polysulfide conversion is a relatively recent development, as such the catalysis of polysulfide conversion is still in an early phase of development.139, 141, 142

Firstly, Co-S binary systems consist of five intermediate phases that are Co4S3 ±x,

Co9S8, Co1−xS, Co3S4, and CoS2, respectively. Co4S3 ±x and Co1−xS are only stable at

5 −1 5 −1 high temperatures. CoS2 (6.7*10 S m ) and Co3S4 (3.3*10 S m ) exhibits much

−1 143 higher conductivity compared with Co9S8 (1.36 S m ). In the early stage, Linda

Nazar’s group first proposed high-surface-area Co9S8 with hierarchical porosity to afford superior LiPS adsorptivity for long-life and high-loading Li−S batteries.144

Although it is noted that the heteropolar feature of conductive surface favors the charge transfer from host material to polysulfides, the mechanism in which the polar surface accelerates polysulfide redox does not clearly revealed yet. In this

contribution, Qiang Zhang’s group employed the pyrite-type CoS2 that possess high conductivity to power Li-S battery performance by propelling polysulfide redox.145 It

demonstrated that CoS2-polysulfide interactions not only statically exist but also dynamically accelerate the electrochemical reactions of lithium polysulfides, in which

CoS2 working as an effective electrocatalyst to tune the redox reaction of polysulfides.

Similarily, Jun Pu ect al proved that this concept is also appropriate for Co3S4

nanotube which has 2–3 times electrocatalytic capability than that of CoS2 for oxygen 58 reduction reactions.146 As proposed in this work, the author further proved the enhanced kinetics of the lithiation/delithiation reaction of polysulfides induced by the

Co3S4 nanotube, and concluded that high catalytic capability of this Co3S4 nanotubes with high volume fraction combining sulfur are able to form the percolation network, which is ideal for improving the high performance of lithium sulfur batteries.

In 2014, Yi Cui et al demonstrated the use of 2D layered transition metal disulphides

for effective encapsulation of Li2S cathodes due to their high conductivity (9–10

times higher than that between Li2S and carbon-based graphene) and strong binding

147 with Li2S/Li2Sn species. As proposed in the literature, the results of ab initio

simulations show strong binding between Li2S and TiS2, ZrS2, VS2, as evidenced by

the strong Li–S interaction (between the Li atoms in Li2S and S atoms in TiS2) and

strong S–S interaction (between the S atoms in Li2S and S atoms in TiS2), and the

binding energy between Li2S and these transition metal disulphides are 2.99, 2.7, 2.94

eV, respectively. These values are 10 times higher than that between Li2S and a single

layer of carbon-based graphene (0.29 eV). The Li2S@TiS2 cathode achieved an

-1 unprecedented specific capacity(503 mAh gLi2S ) under high C-rate (4C) conditions

and unprecedented areal capacity under high mass loading conditions (5.3 mgLi2S cm-2), opening up the new prospect of using transition metal disulphides instead of conventional carbon-based materials for effective encapsulation of high-capacity electrode materials.

Figure 1.22 (a) Schematic illustration for the preparation of WS2 vertically aligned on the CNFs. (b) Schematics of various polysulfides conformations on C@WS2/S. (c)

Binding energies (Eb) for the Li−S composites at six different lithiation stages (S8,

Li2S8, Li2S6, Li2S4, Li2S2, and Li2S) on C@WS2/S, as given by first-principles calculations designed to study the interaction between the lithium sulfide species and

WS2. UV−vis spectra of the Li2S6 solution with C@WS2/S and C/S (inset photograph of the Li2S6 solution with different active materials). (d) Long cycles with various

C-rates. (e) Long-term cycling stability test showing an unprecedented high capacity retention with an excellent Coulombic efficiency over 1500 cycles at 2 C.

60 Recently, Jie Xiong’s group first proposed polar WS2 nanosheets deposited on carbon nanofibers as freestanding electrodes for lithium sulfur batteries.89 In this

flaky-structured C@WS2 composite electrode, dense WS2 nanosheets are wrapped around and anchored on the CNFs (Figure 1.22a), which was supported by a

systematic simulation study, confirming that WS2 shows different binding strengths

(0.8 eV˘Eb˘2 Ev) on the Li2Sn (n=2, 3, 4, 6, 8) species at different lithiation stages

due to the WS2 polar functional groups (Figure 1.22b, c). Based on this condition, the

-1 -1 C@WS2 composites exhibits high rate capability (1501 mAh g at 0.1C, 450 mAh g at 3C), outstanding long-term cycling and excellent stability, a specific capacity of more than 502 mAh g-1 (initial capacity: 563 mAh g-1) could be retained even after

1500 cycles at 2C rate (Figure 1.22d, e), which is to the best of our knowledge. This work also opens up an effective way of applying nonpolar/polar composite materials as 3D current collectors to produce long-cycle-life Li–S batteries.

Jim Yang Lee’s group recently reports an electrocatalyst MoS2-x/reduced graphene

148 oxide (MoS2-x/rGO), that can accelerate the kinetics of polysulfide conversion reactions to insoluble products (Figure 1.23a). The amount of sulfur deficiencies can be varied by changing the time and temperature in a heat treatment in hydrogen. This

paper demonstrated that the sulfur deficiencies in the MoS2 nanoflakes were the

catalytic centers that shown good electrochemical activity for Li2S deposition. This

catalytic effect of MoS2-x on the polysulfides redox reactions was further confirmed

by CV of symmetric cells in 0.2 M Li2S6 electrolyte (Figure 1.23b). It is shown that 61 6262 Figure 1.23 (a) Schematic of the synthesis of the MoS2-x/RGO composite and the conversion of Li2Sx on the MoS2-x/RGO surface. (b) Cyclic voltammograms of symmetric cells with identical electrodes of MoS2-x/RGO, MoS2/RGO and RGO in

-1 electrolytes with and without 0.2 M Li2S6 at 3 mV s . (c) XPS spectra of the RGO,

MoS2-x/RGO counter electrodes of symmetric cells after scanning to 1.4 V, or after scanning to 1.4 V and returning to 0 V. Copyright Energy Environ. Sci., 2017, 10,

1476-1486.

MoS2-x/rGO with sulfur deficiencies exhibit sharp peaks and narrow peak separation in each redox pair (-0.39V/0.047V, -0.047V/0.39V), indicating superior electrochemical reversibility and facile polysulfide conversion compared with

MoS2/rGO and rGO electrodes. Apart from that, XPS of the counter electrodes of symmetric cells after scanning from 0 to -0.14 V, and then from -0.14 V to 0 V also provide direct proof that sulfur deficiencies as the origin of enhanced catalytic activity in polysulfide electrochemical reactions (Figure 1.23c). Because the sulfur-deficient

MoS2 component (blue curve) in MoS2-x/rGO was significantly diminished in intensity when scan to -1.4V, while sulfur-deficient were restored once scan back to

0V. It was these changes that establish the correspondence between sulfur deficiency and the extent and reversibility of polysulfide conversion. These types of electrodes exhibit high rate capacity (826.5 mA h g-1 at an 8 C rate) and good cycle stability at around 628.2 mAh g-1 after 600 cycles at 0.5C, making it as one of the best polysulfide conversion catalysts reported to date, more importantly, the mass loading

63 of sulfur in each electrode was around 60%, which is promising for lithium sulfur batteries’s advanced development.

More recently, Guangyuan Zheng et.cl applied the in-situ TEM to probe the detailed

149 sulfur lithiation/delithiation processes of MoS2-encapsulated hollow sulfur spheres,

as shown in Figure 1.24, it proved that the presence of the MoS2 encapsulation layer can limit the volume expansion of sulfur to 48% compared with the theoretical value of 80%, and the

Figure 1.24 In situ TEM study of MoS2-encapsulated hollow sulfur spheres. (a)

Photographs of a flexible film of MoS2-encapsulated hollow sulfur spheres. (b, c)

SEM images of MoS2-encapsulated hollow sulfur sphere with typical wrinkles generated by the stacking of 2-D flakes marked by violet arrows in (c). (d) Schematic of in situ TEM setup. (e−i) Time-lapse images of the continuous lithiation and delithiation of MoS2-encapsulated hollow sulfur spheres to demonstrate the high reversibility. Copyright Journal of the American Chemical Society, 2017, 139,

10133-10141.

64 reversible lithiation and delithiation of sulfur is possible with the conductive MoS2 coating due to the high flexibility and strong van der Waals force. Additionally, it also

confirmed that the hermetical encapsulation of sulfur particles by MoS2 cages is also effective in prevent soluble lithium polysulfides dissolution from the perspective of physical confinement or chemisorption. Thus resulting in a highly reversible specific capacity 585 mAh g-1 after 1000 cycles at 1C.

On the other hand, some workers also intend to explore other metal sulfides with catalysts property in lithium sulfur batteries. For instance, Zhang also investigate

pyrite FeS2 as an efficient adsorbent of lithium polysulfide for improved lithium–

150 sulfur batteries, and confirms the interactions between FeS2 and Li2Sx derived from

the S–S covalent bonds; Li utilizes the idea that employ small amount of SnS2 nanoparticles to immobilize S/polysulfides in the hollow carbon sphere;151 γ-MnS was also applied into the lithium sulfur batteries due to its its highly activity at the surface of porous carbon;152 CuS as a capacity contributing additive was also introducted into lithium sulfur batteries,153 but the dissolution of CuS in low current density still be a

main concern that restricting its application; NiS2 with superior electrocatalytic activity was also utilized in lithium sulfur batteries;139 Shizhang Qiao’s group recently also designed an 3D carbon hollow spheres in which nanosized NiS uniformly distributed on,154 the capacity retention of this NiS@C-HS was 96% after 300 cycles at 0.5 C.

65 By way of conclusion, metal sulfides exhibit higher electrical conductivity and mechanical stability than their corresponding metal oxides. The interaction between stoichiometric metal sulfides and polysulfides exist mainly in the form of Li–S bonding, the binding energy is at around 2.6-3.5 eV. While as for the non-stoichiometric metal sulfides, depending on the metal exposed facets, they are able to bind LiPSs via both polar–polar Li–S interaction and Lewis acid–base bonding, exhibits much higher binding energy at about 4.0–6.0 eV. The development of metal sulfides still in the first stage, the deep mechanism about electrocatalytic property of metal sulfides till need further exploration. On the other hand, the maximum areal sulfur loading of the metal sulfides reported in this work is 4.5 mg cm-2, showing great potential for the practical application of lithium sulfur batteries with high energy density, it seems to be significantly promising to integrate non-stoichiometric metal sulfides that have electrocatalytic function for polysulfides conversion with 3D ultralight conductive freestanding electrodes.

Metal hydroxides as cathodes materials for lithium sulfur batteries

Providing that the high volume ratio of sulfur and the metal oxides/sulfides, it also has a problem that the anchoring powder, even with nanoscale particle sizes, would not be able to provide sufficient interfaces to entrap all polysulfide species in the electrode.

A more practical way is to apply the limited anchoring materials as the coating layers.

155 Recently, thin layered metal hydroxides, including Co(OH)2 and

66 Figure 1.25 (a) Schematic illustration of the synthesis of the CH@LDH/S composite.

SEM and TEM images of ii,vi) ZIF-67, iii,vii) single-shelled ZIF-67@LDH, iv,viii) double-shelled CH@LDH nanocages, v, ix) CH@LDH/S is given. Copyright

67 156 Ni3(NO3)2(OH)4, have been used as effective encapsulation materials for the sulfur cathodes benefitting from the abundant hydrophilic and hydroxy groups. Yu’s

156 research group demonstrated that Ni3(NO3)2(OH)4 could turn into layered

(Li,Ni)-mixed hydroxide compounds by irreversibly reacting with Li+ ions during the early cycles,156 but restricted by the undesired polarizations of the electrode and the

inevitable process of the slow transformation from Ni3(NO3)2(OH)4 to layered (Li,

Ni)-mixed hydroxides. However, these researches imply the possibility of using layered transition-metal hydroxides as promising encapsulation materials to establish better Li–S cells.

David Lou’s group innovatively demonstrated a new concept using double-shelled cobalt hydroxide and layered double hydroxides (CH@LDH) as sulfur host for Li–S batteries (Figure 1.25a).157 This MOF derived Ni–Co LDH polyhedral preserve the morphology and dimension of ZIF-67, not only provides sufficient void space to accommodate a high content of sulfur (75%), but also exhibits strong binding affinities toward polysulfides due to the hydroxyl-functionlized polar surfaces, which are supposed to provide sufficient interfaces to fix all of the LiPSs in the electrode.

This idea of introducing layered double hydroxides in the Li–S battery system opens a new avenue for future development of high performance Li–S batteries.

Inspired by this ideas, Maowen Xu’s group developed uniform hollow sulfur

nanospheres decorated by ultrathin α-Ni(OH)2 nanosheets for Li-S batteries (Figure

68 1.25b).158 This design maximized the intimate contact between the inner sulfur core

and outer α-Ni(OH)2 shell while guarantee the electron-transfer properties of the electrode, the hollow nanospheres in this work provides sufficient void space to load a

large amount of sulfur materials (82%), it should be noted that the outer α-Ni(OH)2 shell anchors the movable polysulfide both by physical confinement and strong chemical interactions. The author also indicates that metal ions with different valence states can offer various Lewis acid sites and form stronger chemical interactions with sulfur and lithium polysulfides, especially for the low valence. As evidenced in the

2+ 3+ XPS spectrum of Ni, Ni and Ni are detected to present in the S@Ni(OH)2, the variation of chemical valence of Ni2+ imply a strong binding interaction between

sulfur and α-Ni(OH)2, that can effectively suppress the dissolution of polysulfides and thus leading to high rate performance in high mass loading of sulfur.

The publications of metal hydroxides as cathode materials for lithium sulfur battery are rare, and the cycling stability cannot compare with other materials, as for reasons, the low electronic conductivity of the metal hydroxides greatly restricted the effective

Li+ transports and sulfur utilization, it may be meaningful to do modification by inserting some metal particles at the interface of metal hydroxides in this regard. In addition, fully anchoring of metal hydroxides on the sulfur particles surface has not been effectively realized, it is necessary to do further exploration on the amount of metal salt, surfactant as well as reaction time in order to utilize the hydroxy polar surfaces of metal hydroxides. Finally, Computational calculation about the binding 69 energy between the metal hydroxides and different state polysulfides is also important to evaluation the adsorption mechanism of metal hydroxides but this field is blank up to now.

Metal nitride as cathode materials for lithium-sulfur batteries

Most metal oxides/ hydroxides are not conductive and ultimately impede electron transport pathways, leading to low sulfur utilization and rate capability. Therefore, it is of great importance to explore other conductive substrate for lithium sulfur batteries.

Transition metal nitrides often crystallize in rock–salt structure and exhibit mixed metallic, ionic, and covalent bonding. Depending on the exact nature of bonding, these nitrides XN (X = Sc, Ti, V, Cr, Zr, Nb) represent a unique combination of high hardness, high melting point and metallic conductivity, which can withstand the volume changes of sulfur/polysulfides during lithiation, and be an ideal anchoring materials. Most importantly, metal nitrides can efficiently suppress the dissolution of polysulfide and its diffusion into the electrolyte by the chemical bonding between metal atoms or non-metallic heteroatoms and lithium polysulfide, while it is also found that the interface of such metal compounds is more favorable for the deposition of the discharge product lithium polysulfide, which can greatly enhance the sulfur cathode stability.

Titanium nitride (TiN) has a number of desirable properties as host materials for sulfur batteries such as high electrical conductivity (46 S cm-1) and excellent chemical

70 Figure 1.26 (a) i)A low-magnification SEM image, ii)high-angle annular dark-field

71 stability due to the formation of an oxide passivation layer.159 John B. Goodenough’s group reported a mesoporous TiN with a high surface area that was synthesized by a solid–solid phase separation method,160 it is confirmed that the strong interaction between TiN and polysulfides derived from N–S surface bonding, which helps to trap

soluble polysulfide intermediates within the agglomerations. In addition, the TiO2 passivation layer on the surface of TiN possesses hydrophilic Ti-O groups and may provide a polar surface for strong binding with a polysulfide. Under this conditions, the TiN is stably remained a sintered porous structure in the organic electrolyte after

500 cycles, but the problem of serious capacity fade still needs to be solved.

Vanadium nitride (VN) has a number of desirable properties for a potential host materials for sulfur including the following: (1) a strong chemical adsorption for polysulfides that can effectively inhibit the shuttle effect, (2) a high electrical conductivity (1.17*106 S m-1 at room temperature) that is conducive to the electrochemical conversion of adsorbed sulfur species on the surface and (3) catalytic properties similar to the precious metals that may facilitate redox reaction kinetics.

Recently, Feng Li’s group designed a highly conductive porous VN

nanoribbon/graphene (VN/G) composite accommodating a suitable amount of Li2S6 catholyte as the cathode of Li–S batteries (Figure 1.26a,b).125 The VN not only shows strong chemical anchoring of the polysulfides, but also accelerates the redox

reaction kinetics. It was confirmed that the strong adsorption of polar VN to Li2S6 due to ionic bonding 72

Figure 1. 27 (a) SEM images of the Co3O4 phase, Co4N phase, Co4N @S phase. (b)

Co 2p3/2 X-ray photoelectron spectroscopy of the Co4N phase and Co4N/Li2S6, respectively. (c) Sealed vials of a lithium polysulfide solution (Li2S6 dissolved in

Nano letters, 2015, 15, 5137-5142

73 of V-S, N-Li, the bonding energy is around 3.75 eV, 1.07 eV, respectively. In detail,

the strong polar–polar interaction between Li2S6 and VN results in an obvious

deformation of the Li2S6 molecule, forming three S–V and one Li–N bonds. The bond lengths of these S–V (2.49–2.61Å) and Li–N (2.08Å) bonds are very close to the

corresponding bond lengths in bulk VS (2.42 Å) and LiNH2 (2.06 Å), respectively

(Figure 1.26d). These results clearly show the good affinity and strong chemical anchoring of polar VN for polysulfides. Moreover, the surface of the VN also contain small amounts of V–N–O and V–O bonds, which exhibits a high affinity for polysulfides. Based on these outsanding characteristics, the author synthesized a 3D free standing lithium/polysulfide battery generated a high specific capacity of

1,461mAhg-1 at 0.2 C, a Coulombic efficiency approaching 100%, and a high rate performance of 956mAhg-1 at 2 C (Figure 1.26c).

Recently, Dong ’s group developed Co4N mesoporous sphere that composed of nanosheets for high performance lithium sulfur (Figure 1.27a),90 it was demonstrated

that Co and N atoms in Co4N have strong adsorption capacity for both S and Li in the

LiPSs, respectively. Specifically, the Co4N mesoporous sphere shows bifunctional

catalytic activities for the Li−S battery, the Co element with low valence in Co4N not at the full oxidation valence of +3 are highly capable of donating electrons, leading to a strong affinity for sulfur atoms/ions in the LiPSs (promote the discharge process), as proved by the formation of additional peak at 778.9 eV and the decreased peak at

around 778.5 eV in the XPS spectrum of Co 2p3/2 in Co4N/Li2S6 (Figure 1.27b). At

74 the same time, one electron lone pair of the N atom in Co4N can serve as a conductive

Lewis base “catalyst” matrix to enhance the adsorption energy of Li in the LiPSs, thus improve the charge process of lithium sulfur battery (Figure 1.27c). Most importantly, it need to be pointed out that even when the sulfur loading reaching 94.88%, this cathode can still generated an extremely high specific discharge capacity of 1259

−1 mAh g with good cycle performance (Figure 1.27d), which is really high value at present reported.

Apart from that, K.Y. Simon Ng et al also investigated WN, Mo2N, VN as cathode materials for lithium polysulfides conversion,161 although the electrochemical performance of these three cathodes cannot compared with the above works, it gave us a new insight that WN is a highly promising candidate for high performance Li-S batteries due to the existence of S-W-N bonding on the electrode surface, some structural modification were expected to made to further improve its cycling stability.

Other host materials for lithium-sulfur batteries

To overcome the insulating nature of MOFs, metal oxides and metals sulfides while combines the advantage of entrapping LiPSs via Lewis acid-base bonding, the exploration cathode materials for lithium sulfur batteries have been extended to

MXene, which belongs to a large family of semiconducting or metallic early-transition-metal carbides or carbonitrides. The Mxenes are produced by

selectively etching an atom from layered MAX phases (Mn+1AXn, A = IIIA/IVA elements), and then delaminating the sheets in a polar solvent. They have a general 75 formula of Mn+1CnTx, where T is assigned to the surface-terminated functional groups

162-165 (OH, F, O, etc.). A completely delaminated MXene phase, Ti2C, possesses metallic conductivity with 2D surface that has ample Lewis acid Ti sites and hydroxyl groups.166-168 Linda Nazar’s group firstly reported that the use of 2D conductive

MXene nanosheets as cathode hosts to improve the performance of Li–S batteries due

to the fact that the self-functionalized surface of the delaminated HF-etched Ti2CTx promotes the chemisportion of polysulfides on the “acidic” Ti sites (Figure 1.28a).92

More recently, this group promoted the understanding about MXene, innovatively proposed the two steps mechanism about MXene entraps polysulfides (Figure

1.28b),169 it demonstrated that the hydroxyl terminal groups on MXene first undergo redox reaction with polysulfides forming the surface thiosulfate groups, then followed by the Wackenroder reaction exposing the Ti atoms. Second, the exposed and metastable Ti atoms readily accept electrons from additional polysulfides in the electrolyte and form Ti-S bonds by Lewis acid–base interactions. The author then confirmed this concept by the DFT calculation (Figure 1.28c), as the binding energy on pristine MXene was only 3.42 eV due to the polar-polar interaction. After gradually removing the hyfroxyl groups, the polysulfides readily absorbs the vacancy

and bonds to Ti, resulting in the remarkable distortion of Li2S4 molecule. Specially when all the hyfroxyl groups removed, the binding energy reached the maximum value. Moreover, the variation of binding energies of different polysulfides molecules

(Li2S2, Li2S4, Li2S6) also further confirmed that the binding energy increases

76 Figure 1.28(a) Replacement of the Ti-OH bond on the Mxene surface with a S-Ti-C bond on heat treatment or by contact with polysulfides. (b) Schematic demonstrating the two-step interaction between a representative hydroxyl-decorated Mxene phase and polysulfide. (c) First-principles calculations for the interaction between Ti3C2

Mxene and polysulfide Li2S4, showing the most stable Li2S4 binding geometry configuration after full relaxation on: i) Ti3C2(OH)2 (used to represent the pristine

Mxene), ii) Ti3C2(OH)2 with one hydroxyl vacancy, iii) Ti3C2 without any surface functional groups. iv) The variations of the binding energies of polysulfide molecules

77 with a decrease in the number of the hydroxyl groups on the surface, with the fully exposed MXene exhibiting the highest binding energy, all these results first demonstrates that the consist of hydroxyl on the functionalized MXene are subject to redox activity with polysulfide. Based on this high polysulfide adsorption, the author then integrated the CNTs into the MXene layers to achieve a porous, electronically conductive network with stable performance, the fading rates as low as 0.043% per cycle for up to 1200 cycles. The highest sulfur mass loading in this work is around 5.5 mg cm−2 which is the best level for MXene up to date. Guoxiu Wang’s group also successfully developed metal carbide@mesoporous carbon hybrid architecture

(HF-etched Ti3C2Tx@Meso-C) by a facile step pyrolysis of Ti3C2Tx/MOF-5 at 900 °C

170 for 3 h in argon and subsequent HF etching to remove ZnO. This Ti3C2Tx@

Meso-C framework not only combines the advantage of MXene, but also generates hybrid porous architecture that is favourable for sulfur accommodation. When applied as cathode material for lithium sulfur batteries, it could maintain a capacity of 704.6

−1 mAh g after 300 cycles at 0.5C. This work also opens a new direction for integrating other unique material to enrich the properties of MXene.

78 Table 1.1 Comparison of metal compounds (-S, -N, -OH, -C, MOFs) used as cathode

materials for lithium-sulfur batteries.

Year Sulfur loading in Maximum discharge Retained capacity (mAh g-1) or Voltag R

electrode by weight capacity(mAh g-1) Capacity retentions rate (%) e ef

(%)&( mg cm-2) Range er

(V) en

Ti4O7 2017 50–55 wt%&1.0–1.2 1411 (0.1C) 988, 70% (300 cycles, 0.1C) 1.8-3

G-MnO2 2015 56% & 0.7-1.0 1300 (0.05C) 380(1200 cycles, 0.2C) 1.7-3 1120(0.2C)

CNT/NiFe2O4 2015 55% & 1.0–1.2 1350(0.1C) 850, 95%(500 cycles, 1C) 1.7-2.6 1200(0.2C)

1050(0.5C)

900(1C)

700(2C)

CeO2/MMNC 2017 56% & 1.4 1368(0.2C) 1066, 78%(200 cycles, 0.2C) 1.7-2.8 1352(1C) 836, 62%(500 cycles, 1C)

MIL-100(Cr) 2011 19% & N/A 1103 (0.1C) 461 , 42% (50 cycles, 0.1C) 1-3 130

MIL-100(V) 2016 35% & ~0.9-1 849 (0.1C) 600 , 70% (70 cycles, 0.1C) 1.5-3 134

HKUST-1(Cu) 2013 16% & ~0.5 1498 (0.1C) 500, 33% (170 cycles, 0.1C) 1-3 131

ZIF-8 2014 30% & N/A 1055 (0.1C) 553, 67% (300 cyclesˈ0.5C) 1.8-2.8 132

738(0.5C)

MIL-53 (Al) 1215 (0.1C) 347, 44% (300 cyclesˈ0.5C)

793(0.5C)

NH2-MIL-53 (Al) 1055 (0.1C) 332, 58%(300 cyclesˈ0.5C)

568(0.5C)

HKUST-1 526 (0.1C) 286, 66%(300 cyclesˈ0.5C)

431(0.5C)

MOF-525(2H) 2015 35% & ~0.7 1184(0.5C) 402, 34% (200 cycles, 0.5C) 1.5-3 133

MOF-525(FeCl) 1190(0.5C) 616, 52% (200 cycles, 0.5C)

MOF-525(Cu) 1197(0.5C) 704, 59% (200 cycles, 0.5C)

128 Ni6(BTB)4(BP)3 2014 48% & N/A 689(0.1C) 611, 89% (100 cycles, 0.1C) 1.5-3 617(0.2C) 500, 84% (200 cycles,0.2C)

Co(BTB)4(BP)3 287(2C) 585 (0.2C) 372, 64% (200 cycles, 0.2C)

91 Na2Fe[Fe(CN)6] 2017 65% & N/A 1129(0.1C) 763, 68% (100 cycles, 0.1C) 1.7-2.7 1147(0.2C) 770, (100 cycles,1C)

1001(1C) 697, (100 cycles, 2C)

79 819(2C) 544, (200 cycles, 5C)

683(5C)

Na2Fe[Fe(CN)6]@P 1291(0.1C) 1101, 85% (100 cycles, 0.1C) EDOT 913(0.5C)

838(1C)

770(2C)

654(5C)

Co(OH)@CB 2015 42% & N/A 993(0.1C) 779, 78% (100 cycles, 0.1C) 1.6-2.7 155

860(0.2C) 576,79% (200 cycles, 1C)

780(0.5C)

730(1C)

156 Ni3(NO3)2(OH)4 2015 63% & ~1.8-2.5 1326(0.1C) 1250, 94% (500 cycles, 0.2C) 1.5-3

158 α-Ni(OH)2 2017 56% & ~2.2 708(1C) 422, 60% (1000 cycles, 1C) 1.5-3 1020(0.2C)

842(0.5C)

618(1.5C)

523(2C)

CH@LDH 2016 52% & 3 1014(0.1C) 653, 64% (100 cycles, 0.1C) 1.7-2.8 157

747(0.5C) 491, 66%(100 cycles, 0.5C)

144 Co9S8 2016 60% & 2.5-4.5 1130(0.05C) 643,75% (400 cycles,2C) 1.6-3 890(0.5C)

895(1C)

863(2C)

145 CoS2 2016 60%& 2.9 1003(2C) 321, 32%(2000 cycles,2C) 1.7-2.8

146 Co3S4 2017 59%& 2 517(5C) 305, 59% (1000 cycles, 5C) 1.6-2.6 N/A& 4 810(0.1C) 825,100%(50 cycles, 0.1C)

89 WS2@C 2017 11%&1-1.2 1180(0.5C) 995,84%(500 cycles,0.5C) 1.7-2.7 (binder free 3current 954(1C) 843,88%(500 cycles, 1C) collector) 762(1.5C) 719, 94%(500 cycles, 1.5C)

563(2C) 502,90%(1500 cycles, 2C)

151 SnS2/Hollow carbon 2015 30%& N/A 1237(0.2C) 924,75%(200 cycles,0.2C) 1.8-3 sphere 930(0.5C)

827(1C)

718(2C)

570(5C)

148 MoS2-x (4 %)/RGO 2017 60% &2-4 1156(0.5C) 628, 54% (600 cycles, 0.5C) 1.8-2.6 826(8C)

149 MoS2 2017 52% & 1.5 1660(0.1C) 585, 55% (1000 cycles, 1C) 1.8-2.6 1089(1C)

NiS@C-HS 2017 50% & 2.3 1002 (0.2C) 718, 72% (200 cycles, 0.2C) 1.6-2.8 154

723 (0.5C) 695, 96% (300 cycles, 0.5C)

-2 -2 150 FeS2(10%) 2016 60% & 2 1129(0.5mA cm ) 700, 62%(200 cycles,0.5mA cm ) 1.7-2.8 80 N/A & 1 -2 -1 -1 147 TiS2(78%) 2014 mgLi2S cm 806mAh g Li2S(0.2 C) 718 mAh g Li2S, 89%(100 cycles, 0.2C) 1.8-2.6

-1 -1 650 mAh g Li2S (0.5C) 501 mAh g Li2S,77%(400 cycles, 0.5C)

-1 608 mAh g Li2S (1C)

-1 595 mAh g Li2S (2C)

-1 560 mAh g Li2S (3C) 5.3 -2 -1 mgLi2S cm 503 mAh g Li2S (4C) 3mAh cm-2 (0.1mA

cm-2)

2.2 mAh cm-2(0.3 mA

cm-2)

1.6 mAh cm-2(0.6 mA

cm-2)

N/A & 1 -2 147 ZrS2(72%) 2014 mgLi2S cm 777(0.2C) 660, 85%(100 cycles, 0.2C) 1.8-2.6 2.7mAh cm-2 (0.1mA 4.8 -2 -2 mgLi2S cm cm ) 2 mAh cm-2(0.3 mA

cm-2)

1.7 mAh cm-2(0.6 mA

cm-2)

N/A & 1 -2 147 VS2(75%) 2014 mgLi2S cm 747(0.2C) 642, 86% (100 cycles, 0.2C) 1.8-2.6 5 -2 -2 mgLi2S cm 2.7mAh cm (0.1mA cm-2)

1.9 mAh cm-2(0.3 mA

cm-2)

1.6 mAh cm-2(0.6 mA

cm-2)

90 Co4N(28%) 2017 49%& 1.5-2 1659(0.1C) 1000, 60%(100 cycles, 0.1C) 1.7-2.7 63%& 2.4-2.8 1428(2C) 690, 48%(800 cycles, 2C)

67%&2.4-2.8 1259(2C) 540, 43%(800 cycles, 2C)

WN 2016 N/A & 8 1768(0.1C) 697, 39%(100 cycles, 0.1C) 1.5-3 161

665(1C)

N/A &9.5 980(100 cycles, 0.1C)

N/A &12.5 1283(100 cycles, 0.1C)

161 Mo2N 2016 N/A & 8 1001(0.1C) 569, 57%(100 cycles, 0.1C) 1.5-3 347(1C)

VN 2016 N/A & 8 1068(0.1C) 264, 25% (100 cycles,0.1C) 1.5-3 161

145(1C)

VN(30%)/GO 2017 N/A &3 1471(0.2C) 1252, 85% (100 cycles, 0.2C) 1.7-2.8 125

(L2S6) 1241(0.5C) 1131(1C) 917, 81% (200 cycles,1C) 81 953(2C)

701(3C) d-Ti2C 2015 56% & 1 1090(0.5C) 723,66% (650 cycles, 0.5C) 1.8-3 92

1000(1C)

169 CNT(10%)-Ti2C, 2017 66.4% & 1.5 1240(0.05C) 450,48.4% (1200 cycles, 0.5C) 1.7-3

CNT(10%)-Ti3CN 66.4% & 1.5 1263(0.05C) 450,48.4% (1200 cycles, 0.5C)

CNT(10%)-Ti3C2 63.2% & 1.5 1216(0.05C) 450,48.4% (1200 cycles, 0.5C) 64%& 3.6 1093(0.2C) 607, 55.5% (250 cycles, 0.2C)

64%& 5.5 910(0.2C) 512, 56.3% (250 cycles, 0.2C)

170 Ti3C2Tx@ 2016 58% & 2 1226(0.5C) 705, 56% (300 cycles, 0.5C) 1.7-2.8 Meso-C

1.2.4 Anode materials for lithium-sulfur batteries

With the significant improvements in the development of sulfur cathodes, the anode,

which is a major limiting factor in the use of Li-S batteries as practical energy storage

devices, has recently attracted much more attention. Metallic lithium is the most

promising anode candidate due to its high specific capacity and low electrochemical

potential. However, challenges including lithium dendrite formation, unstable solid

electrolyte interfaces and volume changes lead to safety concerns, low Coulombic

efficiency, and short cycle life. Therefore, the stability of the lithium metal anode may

determine the fate of Li-S batteries. Three aspects of the stabilization of the lithium

anode are very important :1) stabilization of the interface between the anode and the

electrolyte. Because metallic lithium is highly reactive in organic electrolytes,

strategies for improving the stability and uniformity of the interfaces between the

lithium anode and electrolytes have been explored to suppress dendrite formation and

stabilize the solid electrolyte interphase (SEI). Most approaches focus on a

modification of liquid electrolytes by either adjusting the electrolyte components or

82 optimizing additives to control SEI formation. 2) a 3D lithium host. a planar lithium metal foil is directly used as the anode of lithium batteries, and the initial plating of lithium on the planar substrate is prone to the inhomogeneous deposition of lithium particles on which lithium dendrites may then grow. In this regard, constructing a 3D host to accommodate lithium deposition shows promise in preventing dendrite growth. and 3)a “lithiophilic” lithium host. The use of alternative anodes such as carbon and silicon will also be a trend..

83 Chapter 2 Experimental Method and Characterization

Overview

Figure 2.1 Framework of the experiments

Figure 2.1 summarized the main procedures and techniques for the application of nanomaterials for energy storage batteries system, it includes three main steps: a) Design and synthesis of the nanomaterials, by the means of chemical vapor

deposition (CVD), hydrothermal method and solid state reaction. All chemicals

used in this thesis, along with their formula, purity and supplier, are list below in

Table 2.1.

84 b) Characterisation of the physical properties of the as-made materials, in this step

it alternatively involves different characteristic approaches such as the X-ray

diffraction (XRD), scanning electron microscopy (SEM), transmission electron

microscopy (TEM), high-resolution TEM (HR-TEM), Brunauer Emmett Teller

(BET), thermogravimetric analysis (TGA), Raman spectroscopy, Energy

dispersive spectroscopy (EDS), Ultraviolet-visible spectroscopy (UV-vis),

Atomic force microscopy (AFM). c) Electrochemical performance evaluation for lithium ion batteries or lithium

sulfur batteries. It basically includes cyclic voltammetry (CV), galvanostatic

charge-discharge testing, and electrochemical impedance spectroscopy. To

further explore the mechanisms and reasons for different cycling performance of

the materials, in-situ Raman spectroscopy, ex-situ SEM are also used to

investigate the electrodes after/during cycling.

Table 2.1 Chemicals used in the research project

Chemicals Formular Purity Supplier

Carbon black C 100 % Sigma-Aldrich

Poly(vinylidene (CH2CF2)n - Sigma-Aldrich difluoride) (PVDF)

N-methyl pyrrolidinone (NMP, C5H9NO 99.5 % Sigma-Aldrich anhydrous)

Hohsen Lithium foil Li 99.999 % Corporation Japan

85 Lithium-ion battery 1 M LiPF6 in ethylene - Guotai-Huarong electrolyte (LB-303) carbonate(EC) and New Chemical dimethyl carbonate Materials (DMC) (1:1 w/w) Co.Ltd, China

Sufur S 99.5% Sigma-Aldrich

1,2-Dimethoxyethane CH3OCH2CH2OCH3 99.5% Sigma-Aldrich

(DME, anhydrous)

1,3-Dioxolane C3H6O2 99.8% Sigma-Aldrich (DOX, anhydrous)

Lithium CF3SO2NLiSO2CF3 99.95% Sigma-Aldrich bis(trifluoromethane sulfon)imide (LiTFSI) poly(vinylpyrrolidon PVP 99% Sigma-Aldrich e) (Mw = 360000 g mol-1)

Nickel(II) nitrate Ni(NO3)2x6H2O 98% Sigma-Aldrich hexahydrate

Cobalt(II) nitrate Co(NO3)2x6H2O 98% Sigma-Aldrich hexahydrate

Ethanol CH3CH2OH 95 % Chem Supply

Graphite (natural C 75 % Sigma-Aldrich flakes)

Cetyltrimethylammo CH3(CH2)15N(Br)(CH3)3 99% Sigma-Aldrich nium bromide

Ammonium (NH4)2S2O8 98% Sigma-Aldrich persulfate

Formaldehyde HCHO 36.5%-38 Sigma-Aldrich %

Pyrrole monomer C6H10N2 99% Sigma-Aldrich

Potassium carbonate K2CO3 99% Sigma-Aldrich

86 Tris(hydroxymethyl) NH2C(CH2OH)3 99.8% Sigma-Aldrich aminomethane

Dopamine (HO)2C6H3CH2CH2NH2·HCl 98.5% Sigma-Aldrich hydrochloride

Sodium thiosulfate Na2S2O3 98% Sigma-Aldrich

Sodium sulfide Na2S - Sigma-Aldrich

Methanol CH3OH 99.9% Sigma-Aldrich

2-Methylimidazole C4H6N2 99% Sigma-Aldrich

Sodium formate HCOONa 99% Sigma-Aldrich

1-Methylimidazole C4H6N2 99% Sigma-Aldrich

Carbon disulfide CS2 99% Sigma-Aldrich

Material Preparations

2.2.1 Solid-state reaction

Solid state reaction is the most widely used method for the preparation of polycrystalline solids by simply mixing of the solid starting materials. In order to react at an appreciate rate, it is necessary to heat treatment at 500 to 1500 ć.

Because both thermodynamic and kinetic factors are important in the solid reactions, thermodynamic factor reveals whether or not a particular reaction should occur by considering the changes in free energy that are involved, kinetic factor determine the rate at which the reaction occurs. The nucleation is rather difficult because of (a) the considerable differences in structure between reactants and product, (b) the large amount of structural reorganization is involved in forming the product: bonds must be

87 broken and reformed, atoms must migrate, perhaps over considerable distances. The subsequent stage, growth of the product layer is even more difficult considering the fact that counter diffusion of precursors must occur right through the existing product layer to the new reaction interfaces. The above discussion shows that three of the important factors that influence the rate of reaction between solids are (a) the area of contact between the reacting solids and hence their surface areas, (b) the area of contact between the reacting and (c) the rates of diffusion of ions through the various phases and especially through the product phase. Clearly, it is necessary to maximize all of these factors in order to reduce the time taken for solids to react together. A typical procedure of solid state reaction is outlined as follows:

Reagent: The selection of reactant chemicals depends on the reaction conditions and expected nature of the product. The reactants are dried thoroughly prior to weighing.

As increase in surface area enhances the reaction rate, fine grained materials should be used if possible.

Mixing: After the reactants have been weighed out in the required amounts, they are mixed. For manual mixing of small quantities, usually an agate mortar and pestle are employed. Sufficient amount of some volatile organic liquid – preferably acetone or alcohol – is added to the mixture to aid homogenization. This forms a paste which is mixed thoroughly. During the process of grinding and mixing, the organic liquid gradually volatilizes and has usually evaporated completely after 10 to 15 minutes.

88 For quantities much larger than ~20 g, mechanical mixing is usually adopted using a ball mill and the process may take several hours.

Heat treatment: The heating programme to be used depends on the form and reactivity of the reactants. In the control of either temperature or atmosphere (argon, nitrogen, helium, hydrogen, etc.), nature of the reactant chemicals are considered in detail.

In this thesis, the solid-state reaction route was employed to prepare nitrogen-sulfur

co-doped graphene, nitrogen doped yolk shell carbon spheres, rGO/N-C-Co3O4 cathode materials for Lithium-sulfur battteries.

2.2.2 Hydrothermal method

Hydrothermal synthesis can be defined as a method of synthesis of single crystals that depends on the solubility of nanomaterials in hot solvent under high pressure. The crystal growth is performed in an apparatus consisting of a steel pressure vessel called an autoclave that can withstand high temperatures and pressure for a long time, in which material is supplied along with solvent. A temperature gradient is maintained between the opposite ends of the growth chamber. At the hotter end the material solute dissolves, while at the cooler end it is deposited on a seed crystal, growing the desired crystal. Advantages of the hydrothermal synthesis method include the ability to synthesis crystals of substances which are unstable near the melting point, and the ability to synthesis large crystals of high quality while maintaining control over their

89 composition. Disadvantages of the method include the need of expensive autoclaves, and the impossibility of observing the crystal as it grows.

In this doctoral study, the hydrothermal methods were employed to synthesize

NiCo2O4 anode material for Lithium-ion battery.

Figure 2.2 Autoclaves for hydrothermal synthesis

Material Characterizations

2.3.1 X-ray Diffraction (XRD)

X-Ray Diffraction (XRD) is a rapid analytical technique primarily used in the determination of crystal structure, qualitative phase identification, quantitative phase analysis, particle size and strain measurements, and the study of preferred orientation in crystals. X-ray diffraction is based on constructive interference of monochromatic

X-rays and a crystalline sample. These X-rays are generated by a cathode ray tube, filtered to produce monochromatic radiation, collimated to concentrate, and directed

90 toward the sample. The interaction of the incident rays with the sample produces constructive interference (and a diffracted ray) when conditions satisfy Bragg's Law ˖

nλ=2d sin θ where d is the inter planar spacing, θ is the Bragg angle, n is the order of reflection and λ is the wavelength of the X-rays irradiation. The crystal phase can be identified by using the standard database (JCPDS cards).

Figure 2.3 The Bruker D8 Discover XRD instrument.

2.3.2 N2 sorption/desorption measurement

The BET instrument applied by particle analysis (Micromeritics 3 Flex analyzer at 77

K) determines the specific surface area, pore size and pore volume of desired samples. 91 The samples are dried with nitrogen purging or in a vacuum applying elevated

temperatures. Unless instructed we use P/P0 = 0.05-0.25 as standard measurement points. The volume of gas adsorbed to the surface of the particles is measured at the boiling point of nitrogen (-196 °C). The amount of adsorbed gas is correlated to the total surface area of the particles including pores in the surface. The calculation is based on the Barret-Joyner-Halenda (BJH) method. Traditionally nitrogen is used as adsorbate gas.

Figure 2.4 The 3 Flex surface characterization analyser instrument produced by

Micromeritics.

2.3.3 Raman spectroscopy

Raman spectroscopy is a vibrational spectroscopic technique used to investigate molecular vibrations and crystal structures. This technique uses a laser light source to irradiate a sample, and generates an infinitesimal amount of Raman scattered light,

92 which is detected as a Raman spectrum using a CCD camera. The characteristic fingerprinting pattern in a Raman spectrum makes it possible to identify substances including polymorphs and evaluate crystallinity, orientation and stress. In comparison to other vibrational spectroscopy methods like FTIR, Raman has several major advantages. These advantages stem from the fact that the Raman manifests itself in the light scattered off a sample as opposed to the light absorbed by a sample. As a result, Raman spectroscopy requires little to no sample preparation and is insensitive to aqueous absorption bands. This property of Raman facilitates the measurement of solids, liquids, and gases not only directly, but also through transparent containers such as glass, quartz, and plastic.

Figure 2.5 The Renishaw inVia Raman microscope equipped with a Leica DMLB microscope (Wetzlar, Germany) and a 17 mW at 633 nm Renishaw helium neon laser.

In this thesis, Raman spectra were collected on an inVia Renishaw Raman spectrometer system using a 632.8 nm wavelength laser.

93 2.3.4 X-ray photoelectron spectroscopy (XPS)

X-ray photoelectron spectroscopy (XPS) is a surface-sensitive quantitative spectroscopic technique that measures elemental composition, empirical formula, chemical state and electronic state of the elements that exist within a material. XPS spectra are obtained by irradiating a material with a beam of X-rays while simultaneously measuring the kinetic energy and number of electrons. A photoelectron spectrum is recorded by counting ejected electrons over a range of electron kinetic energies. Peaks appear in the spectrum from atoms emitting electrons of a particular characteristic energy. The energies and intensities of the photoelectron peaks enable identification and quantification of all surface elements (except hydrogen).

In this thesis, X-ray photoelectron spectroscopy (XPS) measurement was performed on an ESCALAB250Xi (Thermo Scientific, UK) equipped with monochromated Al K alpha (energy 1486.68 eV).

2.3.5 Thermogravimetric analysis (TGA)

Thermogravimetric analysis is a method of thermal analysis in which changes in physical and chemical properties of materials are measured as a function of increasing temperature (with constant heating rate). For different purposes, it usually performed in air or nitrogen atmosphere. In this thesis, TGA was conducted by simultaneous

TG-DTA (SDT 2960), mainly applied for evaluating the weight ratio of sulfur in the

94 host materials, normally under nitrogen atmosphere with the temperature increasing to

800 °C at a speed of 5-10 °C min-1.

Figure 2.6 TGA Analyzer (SDT 2960 model, TA Instrument)

2.3.6 Fourier transform infrared spectroscopy (FTIR)

Fourier Transform-Infrared Spectroscopy (FTIR) is an analytical technique used to identify organic/inorganic materials. This technique measures the absorption of infrared radiation by the sample material versus wavelength. The infrared absorption bands identify molecular components and structures. When a material is irradiated with infrared radiation, absorbed IR radiation usually excites molecules into a higher vibrational state. The wavelength of light absorbed by a particular molecule is a function of the energy difference between the at-rest and excited vibrational states.

The wavelengths that are absorbed by the sample are characteristic of its molecular structure.

95 Figure 2.7 FTIR Equipment

2.3.7 Ultraviolet-visible spectroscopy (UV)

Ultraviolet-visible spectroscopy is a technique that acquires the absorption or reflectance spectroscopy by using electromagnetic radiation in the range of ultraviolet to visible and adjacent light. Molecules in specimens containing π-electrons or non-bonding electrons excite the electrons to higher anti-bonding molecular orbitals after absorbing the energy from the ultraviolet or visible lights. This technique is widely applied in analytical chemistry for qualitative measurements. In this thesis, the

UV-vis absorption spectra according to the Beer Lambert Law, are conducted by a

Carry 300 UV/vis spectrophotometer.

Figure 2.8 Carry 300 UV/vis spectrophotometer

96 2.3.8 Scanning electron microscopy (SEM)

Scanning electron microscopy (SEM) is a method for high-resolution imaging of nanomaterials’ morphology. The SEM uses electrons for imaging, much as a light microscope uses visible light. The advantages of SEM over light microscopy include much higher magnification (>100,000X) and greater depth of field up to 100 times that of light microscopy. Qualitative and quantitative chemical analysis information is also obtained using an energy dispersive x-ray spectrometer (EDS) with the SEM.

The SEM generates a beam of incident electrons in an electron column above the sample chamber. The electrons are produced by a thermal emission source, such as a heated tungsten filament, or by a field emission cathode. The energy of the incident electrons can be as low as 100 eV or as high as 30 keV depending on the evaluation objectives. The electrons are focused into a small beam by a series of electromagnetic lenses in the SEM column. Scanning coils near the end of the column direct and position the focused beam onto the sample surface. The electron beam is scanned in a raster pattern over the surface for imaging. The beam can also be focused at a single point or scanned along a line for x-ray analysis.

In this thesis, field-emission electron microscopy (FESEM, Zeiss Supra 55VP) and electron energy dispersive spectroscopy (Zeiss Evo SEM) were conducted to characterize the morphology and surface structure of as-prepared materials.

Figure 2.9 The field emission scanning electron microscopy in a mode of Supera 55

VP produced by Zeiss and equipped with EDS detector.

2.3.9 Transmission Electron Microscopy (TEM)

Transmission electron microscopy is a microscopy technique in which a beam of electrons is transmitted through a specimen to form an image. The specimen is most often an ultrathin section less than 100 nm thick or a suspension on a grid. An image is formed from the interaction of the electrons with the sample as the beam is transmitted through the specimen. The image is then magnified and focused onto an imaging device. Transmission electron microscopes are capable of imaging at a significantly higher resolution than light microscopes, owing to the smaller de Broglie wavelength of electrons. This enables the instrument to capture fine detail—even as small as a single column of atoms, which is thousands of times smaller than a resolvable object seen in a light microscope.

98 In this work, the morphology and chemical composition of the as-prepared samples were conducted by transmission electron microscopy (TEM, Model JEM2011, JEOL) with a normal operation accelerating voltage of 200 kV.

Figure 2.2 TEM instrument (JEM-2011FS) equipped with EDX detector.

Electrode Preparation and Batteries Assembly

2.4.1 Lithium-ion batteries

The working electrodes were made from 80 wt. % of active materials, 10 wt. % of conductive agent (carbon black) and 10 wt. % of binder (polyvinylidene difluoride) in

N-methyl-2-pyrrolidone (NMP). CR2032 coin cells were assembled in an argon-filled glove box (Mbraun, Unilab, Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm. Cu foil was used as anode current

99 collector. The electrolyte was 1 M LiPF4, dissolved in a 1:1 (weight ratio) mixture of ethylene carbonate and diethyl carbonate. The electrodes were dried at 80 oC under vacuum for 12 h. Electrochemical measurements were conducted using a

LAND-CT2001C battery test system. The cells were discharged and charged galvanostatically in the fixed voltage range of 0.01−3 V with a current density of 100,

200, 500, 1000, 2000 mA g-1. Cyclic voltammetry (CV), and electrochemical impedance spectroscopy (EIS) were conducted with a CHI660C electrochemistry workstation in the voltage range of 0.01 to 3.0 V(vs. Li/Li+).

2.4.2 Lithium–sulfur batteries

Working electrodes were made from 80 wt. % of active materials, 10 wt. % of conductive agent (carbon black) and 10 wt. % of binder (polyvinylidene difluoride) in

N-methyl-2-pyrrolidone (NMP). The mass loading of sulfur on the electrode is around

2 mg/cm2. CR2032 coin cells were assembled in an argon-filled glove box (Mbraun,

Unilab, Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm. Al foil was used as a cathode current collector. The electrolyte was 1 M lithium bis-(trifluoromethanesulfonyl) imide (LiTFSI) and 1 wt% lithium

nitrate (LiNO3) in 1,3-dioxolane and 1,2-dimethoxy-ethane (volume ratio 1:1), The electrodes were dried at 70 ć under vacuum for 12 h. The electrolyte used for each coin cell is around 30 μL. Electrochemical measurements were conducted using a

LAND-CT2001C battery test system. The cells were discharged and charged galvanostatically in the fixed voltage range 1.7−2.7 V with a current rate of 0.1 C, 0.5 100 C, 1C, 2 C and 3 C, respectively. Cyclic voltammetry (CV), and lectrochemical impedance spectroscopy (EIS) were conducted with a CHI660C Electrochemistry

Workstation in the voltage range 1.6 to 2.7 V (vs. Li/Li+).

Figure 2.3 The argon-filled glove box (Mbraun, Unilab, Germany).

Electrochemical Performance Characterization

The electrochemical measurements for evaluating the electrochemical performance of

Li-ion batteries and Li-S batteries included cyclic voltammetry (CV), galvanostatic charge-discharge testing, and electrochemical impedance spectroscopy. The details of these electrochemical measurements are discussed as follows.

2.5.1 Cyclic voltammetry

Cyclic voltammetry (CV) is a potentiodynamic electrochemical measurement that used to traces the relation of current vs. voltage. In a cyclic voltammetry experiment, the working electrode potential is ramped linearly versus time, which applies a potential between a reference electrode and a working electrode while monitors the

101 current between a working electrode and a counter electrode. These cycles of ramps in potential may be repeated as many times as required. When reduction or oxidation reactions occur on the working electrode at a certain potential, current peaks will appear. For lithium sulfur batteries, CV provides detailed information on the transformation phase of sulfur.

In this thesis, CHI 660C or CHI 660D electrochemical workstation (CH Instrument,

Cordova, TN) was applied in CV test (Figure 12).

Figure 2.4 The CH instruments (CHI 660D) for CV and EIS testing.

2.5.2 Galanostatic charge and discharge

Galvanostatic charge-discharge testing plays an important role in evaluating the electrochemical performance of lithium-ion batteries and lithium-sulfur batteries in the aspects of specific capacity, cycling performance, coulombic efficiency, rate performance. Based on different materials, each cell is tested at constant current under given voltage range. For instance, lithium ion batteries (anode material) are testing at the voltage range of 0.01-3 V, while the cut-off voltage of sulfur cathode for lithium

102 sulfur batteries is 1.7-2.6 V. The specific charge/discharge capacities (Q) of electrode materials can be calculated by the following formula:

Q = I * t where I is the current density and t is the charge/discharge time. It should be noted that the specific capacity of each material is calculated based on the specific mass of each electrode.

The cycling performance and coulombic efficiency can be evaluated after certain cycles. In order to investigate the rate capability of batteries, step-wise current densities testing strategy is applied to electrodes.

In this thesis, the galvanostatic charge/discharge measurements are performed on

Neware battery test systems and Land battery test machines, as shown in Figure 2-13.

Figure 2.5 The computer-controlled Neware battery test system.

103 2.5.3 Electrochemical impedance spectroscopy

Impedance spectroscopy represents a powerful method for investigation of electrical properties of materials and interfaces of conducting electrodes. The impedance and phase angles of the materials is measured by a multiple frequency impedance analyser

(impedance meter) that is able to scan each sample at different frequencies. On an electrochemical system in equilibrium a small signal acts (time-dependent potential or current) over a range of frequencies from 0.01 Hz to 100,000 Hz, the linear response of the system is measured then. Data obtained by EIS is expressed graphically in a

Bode plot or a Nyquist plot. By applying physically-sound equivalent circuit models wherein physiochemical processes occurring within the cell are represented by a network of resistors, capacitors and inductors, meaningful qualitative and quantitative information regarding the sources of impedance within the cell can be extracted. EIS is useful for research and development of new materials and electrode structures.

In this thesis, CHI 660C or CHI 660D electrochemical workstation (CH Instrument,

Cordova, TN) was used to conducted this experiment.

2.5.4 Computational methods

104 properties. Ultrasoft pseudopotentials and a plane-wave expansion of the wave functions were chosen for computations. The Brillouin zone is sampled by 2 × 1 × 1 k-points and the energy cutoff of 340 eV is chosen in the geometry optimization calculations. In order to take into account the contributions of the van der Waals

(vdW) interactions between different layers, the DFT-D (D stands for dispersion) approach within the Grimme scheme is adopted for the vdW correction. These setups are proven to be accurate enough for describing the results after careful test calculations.

The adsorption energies (Eads), are defined as: Eads = Etotal - Especies - Esubstrate, where

Etotal is the total energy of the adsorbed system, E species is the energy of the

adsorbate in vacuum and Esubstrate is the energy of the surface. According to this definition, a more negative value indicates a more energetically favorable (exothermic) reaction between the adsorbate and materials surface.

105 Chapter 3 Rose Flower-Like NiCo2O4 with Hierarchically Porous Structures

for Highly Reversible Lithium Storage

3.1 Introduction

The rapid depletion of fossil fuels and increasing environmental pollution issues are major stimulating forces for the development of renewable and clean energy sources.

Among various electrochemical energy storage systems available to date, lithium-ion batteries (LIBs) have attracted extensive attention due to high specific power and energy densities, long cycle lifespan as well as being environmental friendly. The ever-growing demands for LIBs with high power / energy density and perfect cycle life have promoted numerous research efforts devoted to the development of novel and high performance electrode materials as the overall performance of LIBs primarily depends on the electrochemical properties of the electrode materials171-174.

For commercial LIBs, graphite with a theoretical capacity of 372 mA h g-1 cannot meet the current requirements of high capacity10, long cycling life, as well as excellent tolerance of fast charge / discharge. Recently, exploration of anode materials with good electrochemical performances has been carried out, including carbon-based materials, Si, Sn, and various transition metal oxides175. Compared with other counterparts, transition metal oxides with nanostructure characteristics attracted wide attention due to their high surface-to-volume ratio and short path length for Li-ion diffusion176-178. Among transition metal oxides, cobalt oxide exhibited capacity 3

106 times higher than that of conventional graphite, but considering the high cost and toxic nature of cobalt, it has not been regarded as a suitable alternative anode material179-182. Many efforts are made toward replacing cobalt oxide partially by eco-friendly and cheaper alternative metals. Some binary metal oxides, such as

183-187 CuCo2O4, MnCo2O4, ZnCo2O4, NiCo2O4, have also been reported . With two different metal cations, these binary metal oxides exhibit high electrochemical characteristics due to their complicate chemical composition and synergic effects of

multiple metal species. For instance, spinel nickel cobaltite (NiCo2O4), in which one

Co atom is replaced by Ni, possesses much better electrical conductivity and higher electrochemical characteristics compared with plain nickel oxides or cobalt oxides188.

As a result, NiCo2O4 exhibits exceptionally high specific capacity, which is typically

2-3 times higher than that of corresponding monometal oxides189-191.

It is well known that the electrochemical performance of electrode materials is generally determined by the unique structural properties, such as morphologies,

192-195 particle sizes, as well as surface texture, etc . Therefore, synthesizing NiCo2O4 with a suitably designed nanostructure is imperative if employed as an anode material in high-performance LIBs. For instance, David Lou’s group has recently reported

highly uniform NiCo2O4 hollow spheres with a core-in-double-shell interior structure, which exhibit superior electrochemical performances as advanced electrode materials for LIBs189. Jin Liang et al. have successfully prepared one–dimensional hierarchical

nanostructures of NiCo2O4 nanosheets@halloysite nanotubes through a facile 107 coprecipitation method. The as-obtained NiCo2O4 nanosheets@halloysite nanotubes reveal remarkable cycling stability by virtue of the ultrathin and hierarchical

196 nanosheets . Jingfa Li et al. have also reported monodisperse NiCo2O4 mesoporous microspheres by a facile solvothermal method with subsequent annealing of the precursor for achieving excellent cycling performance and rate capability197.

Obviously, electrode materials with hierarchical porous structure possess excellent electrochemical performances due to the high specific surface area which efficiently facilitate the interface contact between electrode and electrolyte198. Furthermore, a porous structure enables the liquid electrolyte to easily diffuse into the electrode material, which facilitates a high flux of lithium ions across the interface and provides space for volume expansion during charge and discharge process198.

Herein, we report 3D hierarchical porous flower-like NiCo2O4, which consist of a 2D nanoplate architecture synthesized by a simple template-free hydrothermal method.

The as-prepared NiCo2O4 nanostructures manifest outstanding cycling stability (4% loss after 100 cycles at 1000 mA g-1), and remarkable capacity retention at high current densities, which can be widely used in high-performance energy-storage devices.

108 3.2 Experimental Section

3.2.1 Synthesis of hierarchical porous rose flower-like NiCo2O4

All chemicals were purchased from Sigma-Aldrich and used without further

purification. For the synthesis of flower-like structure NiCo2O4, a typical process is:

-1 0.1 g of poly(vinylpyrrolidone) (PVP, Mw = 360000 g mol , 99 %) was well dispersed in 7.5 mL of deionized water and 7.5 ml of ethanol (95 %) by stirring

treatment, then 0.25 mmol of Ni(NO3)26H2O (> 98%), 0.5 mmol of Co(NO3)36H2O

(> 98%) were dissolved into the above dispersion to form a light pink solution. After being stirred for a further 30 min, the resulting solution was transferred into a 25 mL

Teflon-Lined stainless steel autoclave and reacted at 180 oC for 12 h. When the reaction finished, the solution was cooled down to room temperature naturally and centrifuged to obtain the green NiCo-precursor. Subsequently, the precursor was washed with deionized water and dried in a vacuum oven at 80 oC overnight. The

China rose flower-like NiCo2O4 was obtained through annealing the dried precursor at

450 oC for 1 h with a heating rate of 1 oC min-1 in normal air atmosphere.

3.2.2 Characterization of materials

The morphology of the obtained materials was characterized by field emission scanning electron microscopy (FESEM, Zeiss Supra 55VP) and transmission electron microscopy (TEM, Model JEM-2011, JEOL). The crystallographic information for the samples was collected on Siemens D5000 diffractometer using Cu-KD radiation with a scanning step of 0.02° per second. Thermogravimetric analysis (TGA) was 109 used for analysis of weight loss from precursor to the final product. Nitrogen adsorption−desorption measurements were conducted on a 3 Flex surface characterization analyzer to determine the Brunauer−Emmett−Teller (BET) specific surface areas using a Quadrasorb SI analyzer at 77 K. The BET surface area was

calculated using experimental points at a relative pressure of P/P0=0.05-0.25.

3.2.3 Electrochemical measurement

The working electrodes were made from 80 wt. % of active materials, 10 wt. % of conductive agent (carbon black) and 10 wt. % of binder (polyvinylidene difluoride).

CR2032 coin cells were assembled in an argon-filled glove box (Mbraun, Unilab,

Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm. Lithium foil was used as a counter electrode. The electrolyte was 1 M

LiPF4, dissolved in a 1:1 (weight ratio) mixture of ethylene carbonate and diethyl carbonate. The electrodes were dried at 80 oC under vacuum for 12 h. Electrochemical measurements were conducted using a LAND-CT2001C battery test system. The cells were discharged and charged galvanostatically in the fixed voltage range of 0.01−3 V with a current density of 100, 200, 500, 1000, 2000 mA g-1. Cyclic voltammetry

(CV), and electrochemical impedance spectroscopy (EIS) were conducted with a

CHI660C electrochemistry workstation in the voltage range of 0.01 to 3.0 V (vs.

Li/Li+).

110 3.3 Results and Discussion

Two steps were employed to prepare NiCo2O4 nanoparticles. Firstly, mixed-metal

(Ni, Co) hydroxide precursors were formed through the solvothermal approach. XRD patterns were indexed as the hexagonal symmetry structure. Typical XRD patterns are shown in Figure 3.1. The morphology of as-prepared (Ni, Co) hydroxide precursors was characterized by the scanning electron microscopy (FESEM) (Figure 3.2). From the low magnification FESEM image (Figure 3.2a), it can be seen that the (Ni, Co)

Figure 3.1 XRD patterns of NiCo-precusor hydroxide precursors have a uniform sphere structure with a size of ~ 1.6 μm.

Through the high magnification FESEM images (Figure 3.2b and c), it can be identified that the spheres are constituted by nanosheets with a thickness of ~ 20 nm.

The transmission electron microcopy (TEM) images (Figure 3.2d and e) demonstrate the assembled nanosheet architecture of the (Ni, Co) hydroxide precursors, which are thin enough to be transparent. Furthermore, each nanosheet has the single crystal

111 Figure 3.2 (a, b and c) FSEM images of the NiCo-precusors at different magnifications. (d and e) TEM images for NiCo-precusors. (f) SAED pattern for

NiCo-precusors. characteristics, as confirmed by the selected area electron diffraction patterns (SAED,

Figure 3.2f) taken from the monolayer nanosheet in Figure 3.2e. All the diffraction patterns can be indexed as [001] zone axis of the hexagonal phase of (Ni, Co) hydroxide.

112

o Figure 3.3 (a, b, c and d) FESEM images of the NiCo2O4 after calcination at 450 C for 1h.

When the hydroxide precursors were thermally treated, they converted into black

hierarchical porous China rose flower-like NiCo2O4 as shown in Figure 3.3. From the

SEM image shown in Figure 3.3a and b, it can be seen that the as-prepared NiCo2O4 preserved the morphology of the NiCo-precursor, presenting uniform and discrete 3D hierarchical flower-like nanostructures with sizes around 1.6 um. Furthermore, these flower-like structures are also constructed from 2D nanosheets which interconnect with each other as shown in the high magnification FESEM images (Figure 3.3b and inset in Figure 3.3b). The typical thickness of a nanosheet is about 20 nm (Figure

3.3c). In Figure 3.3d, it needs to be noted that the 2D nanosheets of NiCo2O4 exhibit a porous structure after thermal treatment of the (Ni, Co) precursor in air for 1 hour. 113

Figure 3.4 (a and b) TEM images for the hierarchical porous flower-like NiCo2O4. (c)

Corresponding SAED pattern. (d) HRTEM image for NiCo2O4.

This is mainly due to the large weight loss accompanying removal of PVP and gases.

Moreover, these numerous exit pores on the “leaves” or “petals” of the nanoparticles would be expected to definitely improve the quantity and speed of electrochemical reactions, resulting in greatly enhanced lithium storage properties.

To further investigate the 3D flower-like architecture and the crystalline phase of the

as-prepared NiCo2O4, the TEM and high-resolution TEM (HRTEM) images associated with SAED were employed as shown in Figure 3.4. Consistent with the

114

Figure 3.5 (a) XRD patterns of China rose-flower-like NiCo2O4. (b) TGA analysis of

NiCo-precursor. (c) FTIR spectra of NiCo-precusors and China rose-flower-like

NiCo2O4. (d) Nitrogen adsorption/desorption isotherms of China rose-flower-like

NiCo2O4, inset is pore size distribution. above FESEM analysis, a low-magnification TEM image (Figure 3.4a) shows that

this flower-like NiCo2O4 is assembled by radially oriented flake-like nanosheets with a thickness of approximately 20 nm. The details of the nanosheets are shown in

Figure 3.4b, from which it can be readily seen that lots of pores exit in the nanosheets, consistent with the FESEM results. The SAED patterns of monolayer nanosheets (Figure 3.4c) shows ring diffraction patterns, illustrating a polycrystalline nature, which can be indexed as the face centred cubic (fcc) spinel phase (Fd3m) of

115 NiCo2O4. The structural characteristic of a typical nickel cobalt oxide particle with visible lattice fringes is also distinctly observed in a HRTEM image (Figure 3.4d).

The inter-planar distance is measured to be 0.287 nm, corresponding to (220) crystal

plane of the spinel NiCo2O4 phase.

Figure 3.6 XRD patterns of the products calcined at 500Ԩ

The crystalline structure and phase purity of the NiCo2O4 materials were further characterized by XRD. Figure 3.5a shows their XRD patterns. The diffraction peaks observed at 2θ values of 31.1, 36.6, 44.6, 55.3, 58.9, 64.7, 76.6 and 81.9°, can be indexed to (220), (311), (400), (422), (511), (440), (533), and (444) crystal planes

respectively, and assigned to cubic NiCo2O4 with a spinel crystalline structure

(JCPDS card no. 02-1704). No other impurity peaks were detected, which manifests

that the pure follow-like NiCo2O4 was formed after thermal treatment. The conversion

process from (Ni, Co) precursor to final NiCo2O4 product was also investigated by thermogravimetric analysis as shown in Figure 3.5b. From this it can be seen that

116

Figure 3.7 FSEM images of NiCo2O4 with different amount of PVP. (a) 0g. (b) 0.05g.

(c) 0.1g. (d) 0.4g. there are two major weight loss steps: the first ~ 7 % weight loss below 200oC is attributed to the loss of absorbed moisture, and the evaporation of residual solvent.

While the second sharp slope at around 250 – 320 °C corresponds to the removal of poly(vinylpyrolidone) (PVP)190, 193. It is also worth noting that there is no obvious weight loss in the temperature range between 320 and 600 °C, which indicates the complete decomposition of the metal precursor and PVP polymer, suggesting that the product retains structural integrity above 320°C. Additionally, the XRD pattern

(Figure 3.6) confirms that Co2.74O4 can appear in the products after calcination at

500oC. Thus, a temperature of 450oC is chosen as the calcination temperature in order

to obtain a single phase of NiCo2O4 and thoroughly remove PVP. To confirm that

117 PVP has been completely removed from the product, FTIR spectroscopy was conducted (Figure 3.5c). In the NiCo-precusor FTIR spectra, the peaks at around

3469 and 1603 cm-1 represent the vibration of –OH groups from absorbed water or residual ethanol. The band at around 1384 cm-1 can be ascribed to C-H stretching and bending modes due to the presence of PVP193. However, after calcination at 450 oC for 1h, the bands corresponding to PVP disappeared and two peaks at 668 and 580 cm-1 formed instead, which can be ascribed to the metal-oxygen vibrations of

193 o NiCo2O4 . It demonstrates that calcinating NiCo2O4 nanoparticles at 450 C results

in the simultaneous removal of PVP and crystallization of NiCo2O4 nanocrystallites.

It should be noted that the addition of PVP is crucial to achieve the hierarchical spherical shape. When no PVP is used, the primary nanoplates suffer from random aggregation, resulting in the formation of aggregate nanostructure (Figure 3.7a).

When increase the amount of PVP to 0.05 g (Figure 3.7b), the problem of aggregation was to some extent alleviated while the flower-like morphology did not appear, and the size of the nanoparticles is not uniform. Nevertheless, when the amount of PVP is 0.1 g (Figure 3.7c), the flower-like structure achieved. Continue increase the amount of PVP to 0.4 g (Figure 3.7d), the flower-like morphology could not preserve, and generate large and uneven hierarchical flower-like structure.

Obviously, excessive PVP in the system can not only provide many high-energy sites for further growth as many free PVP molecules absorbed on the surface of the nanocrystals, but also increases the viscosity of the solution, resulting in an uneven

118 distribution of the products. Based on this condition, it is a specific amount of PVP that contribute to the formation of the delicate hierarchical flower-like structures. As illustrated in Figure 3.8, the shape-inducing effects of PVP could be ascribed to the groups of hydrophobic vinyl and hydrophilic carbonyl, resulting into the formation of polarized micelles (Figure 3.8)2, 199 . The small primary particles were absorbed into the carbonyl groups on the vinyl group, as it is hydrophobic (Figure 3.8b). PVP also functions as a surfactant in the solution, which aggregates the primary particles. With further processing, Co3+ and Ni2+ will oligomerize via the Lamer scheme200. The particles thus formed are flat on the surface with relatively low abundance, but gradually turn and become more tightly packed (Figure 3.8c). The spheres retain their geometry after the PVP is removed, owing to the spherical blocking effect.

Furthermore, after heat treatment the final hierarchical, porous, flower-like NiCo2O4 can be achieved (Figure 3.8d).

Figure 3.8 Illustration for the formation of hierarchical, porous flower-like NiCo2O4

119 The porosity of the NiCo2O4 hierarchitectures was further investigated by BET

analysis, which is shown in Figure 3.5d. The N2 adsorption-desorption isotherm can

be classified as a type IV isotherm with H2 hysteresis loops observed in the range

0.8-1.0 P/P0, indicating that the as-prepared NiCo2O4 typically has a porous structure.

Moreover, the porous structure is further confirmed by the Barrett-Joyner-Halenda

(BJH) pore size distribution data shown in the inset of Figure 3.5d. Two sizes of

pores exist in the as-prepared NiCo2O4, (~ 2 and ~ 20 nm) consistent with the TEM

120 observations of the porous structured nanosheets. The unique flower-like structure has a high Brunauer-Emmett-Teller (BET) specific surface area of 60.3 cm2 g-1 and a pore volume of 0.188 cm3 g-1. It is evident that the 3D flower-like porous structure of the

as-prepared NiCo2O4 demonstrates large surface area and porous architecture, which not only provides more active sites within the pores, but also favours the diffusion of

Li+ ions and electrode-electrolyte contaction during electrochemical reaction, and can tolerate the volume expansion and extraction during the discharge and charge processes, which could significantly improve the electrochemical performance of

LIBs. Therefore, the as-prepared NiCo2O4 is a good candidate as an anode material for LIBs.

The electrochemical performances of flower-like NiCo2O4 nanocomposites were first evaluated by cyclic voltammetry (CV) measurement as shown in Figure 3.9a. In the first cathodic scan, there are two intense peaks located at around ~0.9 and 0.65 V, which can be ascribed to the reduction of Ni2+ and Co3+ to metallic Ni and Co, respectively201. During the following anodic sweep, the peak is observed at around

2.27 V, which could be attributed to the oxidation of metallic Ni and Co to nickel oxide and cobalt oxide. In the subsequent cycles, the reduction peak becomes weak and shifts to 1.23 V, while there is no significant change in the potential of the oxidation peak at 2.27 V. This demonstrated the good electrochemical reversibility of

193 as-prepared NiCo2O4. Based on previous reports , the redox reactions can be described as follows: 121 + - NiCo2O4 + 8Li + 8e → Ni + 2Co + 4Li2O (1)

+ - Ni+Li2O↔NiO+2Li +2e (2)

+ - Co + Li2O ↔ CoO + 2Li + 2e (3)

We further tested the NiCo2O4 electrode cycled at varied current densities ranging

-1 from 100 to 2000 mA g as shown in Figure 3.9b, The flower-like NiCo2O4 electrode shows a good rate capability, with discharge capacities of 1560, 1397, 1301, and 1205 mA g-1 at current densities of 100, 200, 500, and 1000 mA g-1, respectively.

It can still deliver a capacity of 1100 mA h g-1 even at a high current density of 2000 mA g-1, which demonstrate the excellent rate performance of the electrode for high-power LIBs. When gradually decreasing the current rates to 100 mA g-1, the

discharge capacities of NiCo2O4 electrodes recovered the previous values, confirming

the promising reversibility of the as-prepared NiCo2O4.

We also provided the cycling performance data for the China rose flower-like

-1 NiCo2O4 electrodes at a high current density of 1000 mA g , As shown in Figure

3.9c the discharge and charge capacities of the first cycle is 1282 and 1204 mA h g-1, respectively, corresponding to a moderate irreversible loss (about 6%), probably caused by the formation of SEI layer. Clearly, it achieved the high initial columbic efficiency of 93.9%, which suggests the good electrochemical performance of

as-prepared NiCo2O4. In the subsequent charge/discharge curves, a large deviation of voltage is observed. This may be the result of polarization related to ion transfer

122 during the cycling, which is often observed in many transition-metal oxides193. After continuous cycling for 100 cycles at the current density of 1000 mA g-1(Figure 3.9d),

Table 3.1 Comparison of the electrochemical performances of the as-prepared

NiCo2O4 with the reported ones Material Synthesis method Specific capacity Capacity retention Refer (mAh g-1) ence

-1 NiCo2O4 Co-precipitation 750 (1A g ) 86.4% (1000 cycles) 182

-1 NiCo2O4 Self-assembly and thermal 706 (200mA g ) 78% (100 cycles) 181 nanospheres treatment

-1 NiCo2O4 Hydrothermal and 705 (400mA g ) 65% (500 cycles) 165 nanospheres annealing

-1 NiCo2O4 Chemical bath deposition 656 (500mA g ) 74.6% (100 cycles) 180 nanosheets

-1 NiCo2O4 Electrochemically 990 (400mA g ) 86% (300 cycles) 167 nanosheets synthesized

-1 NiCo2O4 Microwave synthesis 891 (100mA g ) 86% (50 cycles) 164 nanosheets Urchin-like Electrochemical deposition 835 (1A g-1) 93% (100 cycles) 166

NiCo2O4

-1 NiCo2O4 Hydrothermal and 900 (500mA g ) 83% (100 cycles) 136 nanoflakes annealing Flower-like Hydrothermal and 915 (1A g-1) 94% (100 cycles) This

NiCo2O4 annealing a reversible discharge capacity of 915 mA h g-1 is retained, indicating the excellent

reversibility and high-rate performances of the China rose flower-like NiCo2O4 electrodes. The lithium storage properties achieved in this present study are

190 remarkably superior to that of many other different NiCo2O4 forms: flower-like ,

189 174 188 NiCo2O4 microspheres , nanocomposites , NiCo2O4/C and NiCo2O4/rGO composite175 as compared in Table 3.1. The good electrochemical properties of

123

Figure 3.10 EIS spectra of flower-like NiCo2O4 and corresponding equivalent circuit.

as-prepared China rose flower-like NiCo2O4 electrodes could be ascribed to the unique 3D hierarchical flower-like nanostructure, which is beneficial for enhanced lithium storage properties. Specifically, the presence of small primary nanoparticles and pores can boost Li+ ion transport, resulting in high capacity and excellent rate capability10. More importantly, the unique 2D nanosheet interior structure could to some extent buffer the large volume change linked to the repeated Li+ insertion/extraction processes during cycling and preserve the structural integrity, thus effectively alleviating the pulverization issue and improving the cycling stability. In these two respects above, this hierarchical porous flower-like architecture is useful for application in LIBs.

st th Table 3.2 Electrochemical parameters of the NiCo2O4 for the 1 ˈ100 cycles

-1 -2 -n Cycle number L (H) Rs (ohm) Rct (ohm) Y0(Ω cm s ) n˄0˘n˘1˅ 1 2.226¯10-6 2.438 64 3.521¯10-5 0.7517 100 2.286¯10-6 3.094 82.29 3.351¯10-5 0.741

124 Electrochemical impedance spectra were also measured before and after 100 cycles.

Nyquist impedance plots were obtained in the frequency range 100 kHz to 0.01 Hz at different charge/discharge cycles (as shown in Figure 3.10). To analyze the

electrochemical parameters of the China rose flower-like NiCo2O4 in detail, we

established the electrochemical model LR (Qdl Rct) (Qw Rw) based on the data obtained from electrochemical impedance spectra. From parameter identification results (Table

3.2), it can be seen that the inductance L has little effect on the electrochemical

performance of the NiCo2O4 electrode before and after 100 cycles due to the low magnitude of inductance L. The depressed semicircle in the medium frequency region is closed, associated with the charge transfer process through the electrode/electrolyte

interface, which could be expressed by Rct/Qdl in the equivalent circuit. Equivalent component Q contains two parameters which are Y and n. the dimension for Y is

Ω-1·cm-2·s-n or S·cm-2·s-n, parameter n is a dimensionless index. Admittance response expression for Q is:

n n Y=Y0ω cos(nπ/2)+jY0ω sin(nπ/2) (5)

Where w is the angular frequency, j is the imaginary unit. When n is 0, the equivalent becomes a resistor. Conversely, when n is 1, the equivalent represents a capacitor,

Which indicating that the smaller n is, the more the electric double layer capacitance deviates from the pure capacitance, and the more purely resistive the equivalent

circuit becomes. However, neither the parameters Y0 nor n experience any remarkable

125 change during repeated cycling, and the charge-transfer resistance Rct showed only a

slight increase, which further confirmed the stability of NiCo2O4 nanocomposites during lithiation and de-lithiation. Meanwhile, the inclined line which is linked with lithium diffusion process within the electrode materials represents the typical

Warburg resistance Rw/Qw, however, these two parameters are difficult to identify, mainly being due to insoluble or barely soluble oxide film covering on the electrode surfaces.

3.4 Conclusions

In conclusion, 3D hierarchical porous China rose flower-like NiCo2O4 was successfully synthesized via a solvothermal method using PVP as the structure-directing agent followed by a simple thermal annealing treatment. Through

the XRD, FESEM, TEM, and N2 sorption analyses, it was revealed that each 3D

flower-like NiCo2O4 sphere is composed of nanoplates, on which numerous pores are distributed due to the thermal treatment from the precursor. Such a novel porous hierarchical architecture not only offers more active sites within the pores for fast electrochemical reaction, but also alleviates the volume expansion/contraction during lithiation and de-lithiation processes. When evaluated as an anode material for lithium-ion batteries, high reversible capacities of 915 mA h g-1 could be retained after 100 cycles at a high current density of 1000 mA g-1, corresponding to about 96% of the second discharge capacity (953 mA h g-1), and remarkable capacity retention was also obtained at increased current densities. The improved electrochemical 126 performance enables the as-prepared 3D hierarchical porous China rose flower-like

NiCo2O4 to be a promising anode material for next-generation, high-power lithium-ion batteries. Furthermore, the currently reported simple synthetic approach could be extended to other transition metal oxides, potentially yielding facile preparation and high yields of electrochemically outstanding products.

127 Chapter 4 Nitrogen-Sulfur Co-doped Porous Graphene Matrix as a Sulfur

Immobilizer for High Performance Lithium Sulfur Battery

4.1 Introduction

The rapid depletion of fossil fuels and increasing environmental pollution issues are major stimulating forces for the development of clean energy technologies.

Renewable energy sources, such as solar, wind, and tide, are attractive in this regard.

However, efficient utilization of these intermittent resources requires efficient and economical energy storage (EES) systems. Advanced rechargeable batteries are the most suitable option for EES. However, confined by the limited theoretical energy density, as well as high cost, the state-of-the-art lithium-ion batteries (LIBs) are not capable to fulfill the specific energy requirements for future electric vehicles (EVs) and large-scale smart grid applications. Therefore, it is of great importance to develop advanced rechargeable batteries to satisfy future needs. Invented in the 1960s, rechargeable lithium-sulfur batteries have now attracted extensive interests, mainly due to their high theoretical energy density (≈ 2567 Wh kg−1), which is estimated to be 5 times higher than that of state-of-the-art lithium-ion batteries. Combined with natural abundance and low cost as well as environmental friendliness, Li-S batteries are one of the most promising rechargeable battery technologies for a wide range of applications.64, 94, 95, 202, 203

128 Despite considerable advantages of lithium-sulfur batteries, their commercialization still suffers from some tough obstacles, such as the insulating nature of sulfur, large volume expansion/contraction (~80%) during cycling, dissolution of polysulfides

(Li2Sx, 4 ≤ x ≤ 8) in organic electrolytes, and shuttle effect of polysulfides, which inevitably result in the low electrochemical utilization of sulfur and rapid capacity degradation.98, 120, 204, 205 Many approaches have been devoted to address these issues.

These include: (1) Encapsulating sulfur particles into various carbon matrix materials and/or conducting polymers, such as mesoporous carbon, carbon nanofiber interlayers, and reduced graphene oxide, polyaniline (PAni), polypyrrole (PPy), and their derivatives.206, 207(2) Strengthening sulfur’s chemical bonds via chemical

interactions or surface-mediated redox reactions with metal oxides (e.g. SiO2, TiO2, metal-organic frameworks) and metal carbides (MXene phases).92, 208-212 (3) Adding

+ 207, novel electrolyte additives such as P2S5 or Cs to suppress the polysulfide shuttle.

213, 214 215-220(4) Modifying the surface chemistry of the hosts to prevent the shuttle effect of the soluble polysulfides between the cathode and the anode, enabling better cycling performance.214, 221, 222 The chemical properties of sp2 carbon materials can be further modified by introducing heteroatom doping, such as phosphorus, sulfur, boron and nitrogen, which demonstrated significant potentials for lithium polysulfide confinement. Song et al reported that nitrogen doping in mesoporous carbon effectively promoted the chemical adsorption of sulfur atoms on oxygen-containing functional groups because surface modification with nitrogen atoms greatly enhances

129 the electrochemical reactivity and electronic conductivity of carbon matrices.223

Additionally, sulfur-doped porous carbon also attracted wide attention for battery applications due to the fact that sulfur-doping can change the charge state of the neighbouring carbon atoms. Therefore, the sulfur-doped carbon materials exhibit enhanced adsorption ability. Most importantly, when sulfur and nitrogen are simultaneously doped into the carbon matrix, synergistic effects can occur due to the highly active multiple doping elements. The co-doping not only strengthens chemical bonding between carbon host and sulfur chains in the heat treatment during sulfur loading process, ensuring the uniform distribution of sulfur, but also generates more active sites and increases their activity, which greatly improve the adsorption of soluble lithium polysulfide intermediates, thus increasing cycle ability and Coulombic efficiency. Therefore, the co-doping strategy can achieve superior electrochemical activity beyond that of mono-element doping.

Graphene, a free-standing atom-thick layer of sp2 carbon atoms, has emerged as the most significant carbon material in the last decade due to its unique physical and chemical properties.224-228 Nitrogen-sulfur co-doped graphene is a promising nanoscale immobilizer as a cathode host material for lithium-sulfur batteries.

However, graphene nanosheets tend to form agglomerates or even restack through van der Waals interactions during the preparation processes and subsequent electrode fabrication, resulting in the loss of specific surface area.229-232 This effect will also lead to lower polysulfide adsorption due to the decrease of the accessible active sites 130 on the graphene surface. Therefore, it is very important to prevent graphene aggregation by creating porous structure of nitrogen-sulfur co-doped graphene nanosheets.

Herein, we have developed a highly crumpled nitrogen-sulfur co-doped 3D graphene nanosheets matrix as a sulfur immobilizer for lithium-sulfur batteries. These highly porous nitrogen-sulfur co-doped graphene nanosheets were synthesized through a facile synthesis approach, in which polypyrrole (PPy) serves as a nitrogen source and ammonium persulfate (APS) as sulfur source. Highly developed defects and edges, as well as hierarchical pore structures derived from graphene chemical activation endow nitrogen-sulfur co-doped graphene nanosheets with large surface area (1012 m2 g-1) and a wide pore size distribution of 2 nm and 25~40 nm. The as-prepared materials have strong chemisorption capabilities for polysulfides and high sulfur loading. When applied as cathode materials in lithium-sulfur batteries, the cathodes exhibit a high initial capacity of 1178 mAh g-1 at 0.2 C and excellent long life cycling performance with a retained capacity of 780 mAh g-1 after 600 cycles.

4.2 Experimental Section

4.2.1 Synthesis of NSG (Nitrogen-Sulfur Co-doped Graphene)

Graphene oxide (GO) was synthesized using natural graphite flakes by a modified

Hummers’ method. 50mg of the as-made graphene oxide was well suspended in 50ml

DI water after ultrasonically stirring for 2h, then 0.73g cetyltrimethylammonium

131 bromide (CTAB) and 1.37g ammonium persulfate (APS) were dissolved into the prepared graphene solution magnetic stirring for 1 h. After that, 0.83 mL of 99% pyrrole monomer was added and the mixture was stirred in an ice water bath for about

24 h. When the reaction finished, the resultant products were washed with DI water and dried in a vacuum at 80 qC for 24 h to remove any residual solvents, this final product is denoted as NSG.

4.2.2 Synthesis of A-NSG (Activated Nitrogen-Sulfur Co-doped Graphene)

Typically, 400 mg NSG powder was dispersed in 20 ml 3M K2CO3 solution and

stirred for 24 hours. The extra K2CO3 solution was removed by briefly filtering the mixture through a polycarbonate membrane (Whatman, 0.2 um); then the mixture was

dried in the lab environment at 70 qC for 24 hours. The NSG/ K2CO3 mixture was put in a tube furnace under flowing nitrogen at 800 qC for 2 h (heating rate 5 qC min-1).

Subsequently, the final product (denoted as A-NSG) was thoroughly washed with

0.1M HCl and DI water until the PH value reached 7 and finally dried in a vacuum at

80 qC for 20h.

4.2.3 Synthesis of A-NSG@S (Activated Nitrogen-Sulfur Co-doped Graphene@S)

A-NSG was homogeneously dispersed in DI water by ultrasonication, then transfered

to an appropriate amount of elemental sulfur/CS2/ethanol solution with a mass ratio of

mA-NSG : mS=1:3, the mixed solution was magnetically stirred for 24 h over an ice

water bath . During this time period, the CS2 and ethanol were allowed to completely 132 evaporate while stirring, then the A-NSG@S composite was filtered and dried at 50 qC in a vacuum oven for 12h. Finally, the as-synthesized A-NSG@S was heated to

155qC at a heating rate of 2 qC min-1 in a tube furnace under flowing argon at 100 sccm before re-heating at 230 qC for another 2 h under same condition.

4.2.4 Characterization of materials

Renishaw in Via Raman spectrometer system (Gloucestershire, UK) equipped with a

Leica DMLB microscope (Wetzlar, Germany) and a 17 mW 633 nm Renishaw helium neon laser at 50% power. XPS analysis was performed on an ESCALAB MK

II X-ray photoelectron spectrometer with a JEOL JSM-6700F electron microscope with an accelerating voltage of 10 kV. Thermogravimetric analysis (TGA) was used for analysis of weight loss from precursor to the final product. Nitrogen adsorption−desorption measurements were conducted on a 3 Flex surface characterization analyzer to determine the Brunauer−Emmett−Teller (BET) specific surface areas using a Quadrasorb SI analyzer at 77 K. The BET surface area was calculated using experimental points at a relative pressure of P/P0=0.05-0.25. The

UV−Vis spectra were measured in the spectra range of 300−1000 nm by Cary 60 133 UV-Vis variable wavelength spectrophotometer. All the samples for UV−Vis measurement were prepared and sealed in an argon-filled glove box.

4.2.5 Electrochemical measurement

Working electrodes were made from 90 wt. % of active materials, 2 wt. % of conductive agent (carbon black) and 8 wt. % of binder (polyvinylidene difluoride).

The mass loading of sulfur on the electrode is around 1.5 mg/cm2. CR2032 coin cells were assembled in an argon-filled glove box (Mbraun, Unilab, Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm.

Lithium foil was used as a counter electrode. The electrolyte was 1 M lithium bis-(trifluoromethanesulfonyl) imide (LiTFSI) and 1 wt% lithium nitrate (LiNO3) in

1,3-dioxolane and 1,2-dimethoxy-ethane (volume ratio 1:1), for each electrode, around 30 uL electrolyte was added in the coin cell. The electrodes were dried at 60 Ԩ under vacuum for 12 h. Electrochemical measurements were conducted using a

LAND-CT2001C battery test system. The cells were discharged and charged galvanostatically in the fixed voltage range 1.7−2.7V with a current rate of 0.1C,

0.5C, 1C, 2C, 5C. Cyclic voltammetry (CV), and electrochemical impedance spectroscopy (EIS) were conducted with a CHI660C electrochemistry workstation in the voltage range 1.7 to 2.7 V (vs. Li/Li+).

134 4.2.6 Computational methods

The simulations are based on density functional theory (DFT), which is provided by

DMOL3. The generalized gradient approximation (GGA) with the

Perdew-Burke-Ernzerhof scheme (PBE) is adopted for the exchange-correlation potential to optimize geometrical structures and calculate properties. The All-Electron

Relativistic Kohn-Sham wave functions (AER) and double numeric plus polarization

(DNP) basis set are adopted in the local atomic orbital basis set with the global orbital cutoff set to 4.4 Å. A single-layer of graphene was modelled using a hexagonal 4×4 supercell. The nearest distance between nanosheets in neighboring cells is greater than 15 Å to ensure no interactions between different layers. For geometric optimization, both the cell and the atomic positions are allowed to fully relax. The

Brillouin zone is sampled at 6 × 6 × 1 k-points for all structures in the geometric optimization calculations, which brings out the convergence tolerance of energy of

1.0 × 10-5 Ha (1 Ha = 27.2114 eV), maximum force of 0.002 Ha/Å, and maximum

displacement of 0.005 Å. The adsorption energies (Ea), is defined as:

Ea = Etotal - Eads - Esubstrate where Etotal is the total energy of the adsorbed system, Eads is the energy of the adsorbate in vacuum and Esubstrate is the energy of the doped and undoped graphene substrate. According to this definition, a more negative value indicates a more energetically favorable (exothermic) reaction between polysulfides and graphene surface.

135 4.3 Results and Discussion

Figure 4.1 Illustration for the formation of A-NSG

The 3D nitrogen-sulfur co-doped porous graphene matrix was synthesized via

chemical activation of polypyrrole (PPy) functionalized graphene sheets using K2CO3 as shown in Figure 4.1. PPy was chosen as nitrogen source due to its high nitrogen content. More importantly, because of the π - π interaction between PPy and graphene oxide (GO), the PPy monomers will polymerize on the surfaces of graphene oxide

(GO) during the polymerization process. The doping of sulfur was achieved by ammonium persulfate (APS). After carbonization, nitrogen and sulfur atoms are expected to dope in the graphene nanosheets. Scanning electron microscopy (SEM) images of nitrogen and sulfur co-doped graphene (denoted as NSG) composites are shown in Figure 4.2, it is clear that the NSG composites show quite similar morphologies to the rGO, displaying a thin smooth surface. The as-prepared NSG was

chemically activated by K2CO3 to obtain final porous nitrogen-sulfur co-doped

136 Figure 4.2 SEM image of NSG. graphene (denoted as A-NSG). Figure 4.3a and Figure 4.4a show the scanning electron microscopy (SEM) images and transmission electron microscopy (TEM) images of A-NSG. The corrugated morphology with wrinkles and folded regions is clearly visible, which is markedly different from rGO and NSG. Meanwhile, most of the A-NSG layers are discontinuous, giving rise to a large number of defects such as edges and pores. Figure 4.3d shows the typical AFM image of the A-NSG, which further reveals that the thin wall typically consists of only a few graphene layers as the A-NSG nanosheets with an average thickness of about 0.396 nm. The crumpling of the A-NSG can effectively prevent them from agglomerating and restacking, thus increasing the electrolyte accessible surface area. The corresponding energy-dispersive X-ray spectroscopy (EDX) illustrates uniform distribution of carbon, nitrogen and sulfur in A-NSG (Figure 4.3b).

137 Figure 4.3 (a) SEM image of A-NSG. (b) EDX elemental mapping for A-NSG. (c)

SEM image of the A-NSG@S. (d) AFM image of A-NSG.

Applying the 3D A-NSG hierarchical network structure as sulfur host, we loaded sulfur to the A-NSG matrix by liquid-phase infiltration and melt vulcanization processes. The initial liquid phase infiltration process ensures sulfur uniformly distribute in the A-NSG. The melt vulcanization processes include two major steps:

(i) a typical step at 155 Ԩ for 10 hours is aimed to encapsulate sulfur into the porous graphene 3D frameworks (denoted as A-NSG@S-155), and (ii) heating at 230 qC is to

trigger ring-opening polymerization of elemental sulfur (S8) into a linear polysulfide along porous graphene (denoted as A-NSG@S-230). In the A-NSG@S composites, a uniform layer of sulfur is homogeneously anchored on the surface of A-NSG without

138 any obviously aggregated sulfur particles being observed in the TEM image (Figure

4.4c) or FSEM image (Figure 4.3c). The sulfur content in the A-NSG@S nanocomposites was determined by thermogravimetric analysis (TGA) under nitrogen atmosphere with a heating rate of 10 qC min-1, as shown in Figure 4.5. The TGA analyses reveal that the sulfur loading in A-NSG@S was 72.4 wt%, thus the mass

Figure 4.4(a, b) TEM images of A-NSG. (c, d) High-magnification TEM images of the A-NSG@S loading of sulfur in each electrode is around 65%. To verify the structural characteristics of A-NSG@S, X-ray diffraction (XRD) was conducted on pure sulfur,

139

Figure 4.5 Thermogravimetric curves of pure sulfur powder and A-NSG@S in the N2 with a heating rate of 10 Ԩ min-1.

A-NSG@S-155, A-NSG@S-230 (Figure 4.6a). It should be noted that the overlapped broad peak at about 26° appear in both A-NSG@S-155 and

A-NSG@S-230, which can be ascribed to the A-NSG from thermal restoration of sp2 carbon. Compared with the intact crystalline peaks of pure sulfur and A-NSG@S-155, the absence of sulfur crystal peaks for A-NSG@S-230 is noteworthy (Figure 4.6a).

This indicated that sulfur now exists in a highly dispersed amorphous state.

Meanwhile, as shown in the Raman Spectra (Figure 4.6b), the pure sulfur powder exhibits characteristic peaks in the range 100-500 cm-1, associated with vibration of

the S-S bond in S8 species while both the A-NSG@S-155 and A-NSG@S-230 did not show noticeable characteristic sulfur peaks, which is also consistent with the

conclusion from XRD patterns. This implies that S8 molecules were loaded into the

A-NSG matrix without long-rang ordering, similar to small sulfur molecule dispersion. Other groups have reported that the amorphous state of sulfur could result

140

Figure 4.6 XRD patterns and Raman spectra of the pure sulfur powder, A-NSG,

A-NSG@S-155, A-NSG@S-230 in the high ultilization of sulfur in electrochemical lithiations.

To investigate the elemental composition and functional groups in the A-NSG, X-ray photoelectron spectroscopy (XPS) measurements were performed. As shown in

Figure 4.7a, the XPS survey spectrum of A-NSG exhibits four characteristic peaks located at ca. 164, 400.23, 285.07 and 532.25 eV, corresponding to S 2p, N 1s, C 1s, and O 1s. The elemental content of S and N in A-NSG is 0.85 at% and 4.18 at%,

141 respectively. This suggests the effective doping of N and S into the graphene sheets.

In the N 1s spectrum of A-NSG (Figure 4.7b), three components including pyridinic

N (21.68 wt%), pyrrolic N (50.16 wt%), and graphitic N (28.16 wt%), can be identified at 398.4, 400.2 and 402.1 eV, respectively. It is also reported that the

Figure 4.7 (a, b, c) High-resolution C 1s, N 1s and S 2p XPS spectrum of A-NSG and corresponding XPS survey spectra of A-NSG. (d) Schematic structure of A-NSG dominant pyridinic and pyrrolic N had been proven to contribute to the improvement of the affinity and binding energy of non-polar carbon atoms with polar and rate capability of Li–S batteries. We also displayed the molecular geometries and the

adsorption energies of Li2S and Li2S4 on pyridinic and pyrrolic N sites in Figure

142 4.8a-d. In the most stable configuration, terminal Li atoms in Li2S and Li2S4 tended to directly bind to the pyridinic N atom. There was distinctive electron concentration

Figure 4.8 Optimized configurations for the adsorption of Li2S on pyridinic and pyrrolic N sites (a, b), Li2S4 on pyridinic and pyrrolic N sites (c, d) with corresponding adsorption energies in eV and atom Mulliken charge. Gray, blue, yellow, and purple balls represent C, N, S, and Li atoms, respectively. between N and Li atoms, suggesting a strong Li–N electrostatic interaction. The electron migration can be well explained by the Lewis acid-base theory. The pyridinic

N with an extra pair of electrons was considered as an electron-rich donor that naturally acted as a Lewis-base site to interact with the strong Lewis acid of terminal 143 Li atom in lithium (poly)sulfides. Similarly, as indicated by Figures 4.8b and d, pyrrolic N played an analogous role as pyridinic N, which were consistent with Hou et al’s theoretical results. In addition, the dependence of adsorption strength between

LiPSs and N doping configuration can be seen from Li-N distances. The Li-N

distances of Li2S on pyridinic and pyrrolic N sites were 1.99 and 2.02 Å, which are

consistent with the large Ea obtained from Li2S on pyridinic N site. In the case of

Li2S4 on pyridinic and pyrrolic N site, the Li-N distances were calculated to be about

2.04 and 2.05 Å, which further confirmed the strong anchoring effect between

pyridinic, pyrrolic N and acid terminated Li in Li2Sx. On the other hand, in the high

Table 4.1Percentage of Carbon atoms and Full width at Half-Maximum Values of C

1s Peaks in Graphene, A-NSG.

carbon % (atom %) fwhm (eV)

Graphene 93.4% 1.13

A-NSG 89.81 % 1.21 resolution C 1s spectra (Figure 4.7a), due to the removal of most oxygen groups as well as the partial reconstruction of graphitic carbon network, peaks corresponding to

C-O (286.2 eV), C=O (288.2 eV) and O-C=O (289.6 eV), decrease considerably for

A-NSG, compared with that of GO, two particular new peaks located at 286.1 and

287.2 eV suggest the bond formation of graphene, the sharp peak in the C 1s spectrum of A-NSG, still attributed to the sp2-hybridized graphitic carbon atoms, shifts to higher binding energy and its full width of half-maximum (fwhm) at 284.8 eV increases with the presence of nitrogen content (Table 4.1). All of these results 144

Figure 4.8 Optimized configurations for the adsorption of Li2S on pyridinic and pyrrolic N sites (a, b), Li2S4 on pyridinic and pyrrolic N sites (c, d) with corresponding adsorption energies in eV and atom Mulliken charge. Gray, blue, yellow, and purple balls represent C, N, S, and Li atoms, respectively. confirm the formation of C-N bonding configurations in the A-NSG. Moreover, in the

X-ray photoelectron spectroscopy (XPS) analysis, the S 2p spectrum for the A-NSG nanocomposites in Figure 4.7c has 2p3/2 and 2p1/2 spin-orbit levels with an energy separation of 1.27eV and intensity ratio of 2:1, confirming the effective doping of sulfur atom into the carbon lattice. Two sulfur species are assigned to carbon-bonding: thiophenic S (163.6 eV) and sulfonic S (167.9 eV),233 the thiophenic 145 Table 4.2 Carbon, oxygen, nitrogen, sulfur atomic percent of GO, A-NSG materials.

Sample C (at%) O (at%) N (at%) S (at%) C/O

GO 70.2 29.8 _ _ 2.4

A-NSG 89.81 5.16 4.18 0.85 17.25

S (sulfur atom bonded with carbon directly by a C-S bond) is believed to make the carbon matrix positively charged, thereby increasing its affinity to absorb polysulfides. At the same time, it is also interesting to find that the C/O ratio of GO is

2.4 and this value increases to 17.52 for A-NSG (Table 4.2), indicating that nitrogen-sulfur co-doping increases the reduction efficiency of GO. This seems to be further supported by the Raman spectroscopy analysis. In Figure 4.9a. the Raman spectra for A-NSG exhibit two remarkable peaks at around 1350 and 1580 cm-1. The

G band located at 1580 cm-1 is related to the E 2g vibration mode of sp2 carbon atoms while the D band at 1350 cm-1 is attributed to the defects and disorder hybridized

vibrational mode of graphene. Therefore, the higher intensity ratio (ID/IG) of A-NSG than GO (1.41 vs 1.06), should be ascribed to more defect sites generated by N and S co-doping, resulting the high disorder of graphene nanosheets. More importantly, these defects can provide considerable active sites for lithium polysulfides adsorption.234

To further prove the interaction between polysulfides and A-NSG, the Li2S4 solution

was prepared by dissolving Li2S and sublimed sulfur (with a stoichiometric ratio of

1:3) in the mixed solvents of 1,3-dioxolane (DOL) and 1,2-dimethoxyethane (DME)

146 (1:1 v/v). Then soaking 15 mg of N-S co-doped graphene in 15 mL of 5 mM Li2S4 solutions at 30 °C for 24 h, respectively, the upper solutions were carefully collected and sealed for UV-Vis spectra measurement. Figure 4.10a–d display color of pristine

Li2S4 solution and the color changes of the Li2S4 solution at various periods after

Figure 4.9 (a) Raman spectra of GO, NSG, A-G and A-NSG. (b) Pore characterization of the NSG, A-G and A-NSG materials adding N-S co-doped graphene composite. After adding N-S co-doped graphene into

the Li2S4 solution, a homogeneous black suspension was initially formed, and it took

147

Figure 4.10 (a) Li2S4 solution (b) Initially mixed (c) Aging for 3h (d) Aging for 24h several minutes for N-S co-doped graphene to precipitate. After aging for 3 and 24 h,

the color of Li2S4 turns from the initial yellow to light yellow, indicating that much of

the Li2S4 has been adsorbed by the N, S graphene. The result indicates the strong adsorption ability of N-S co-doped graphene for lithium polysulfides. The corresponding UV-Vis spectra of Figure 4.11 shows that the reflect index at 513 nm

of the Li2S4 solution was increased after soaking the N, S graphene composite for

24h, confirming that large amount of Li2S4 was absorbed by N-S co-doped graphene.

Figure 4.11 UV-vis spectra of 5 mM pristine Li2S4 solution and solution after soaking the N, S-graphene.

148 We conducted X-ray photoelectron spectroscopy (XPS) analysis to identify the interaction between polysulfides and N, S-graphene. As shown in Figure 4.12, the Li

1s XPS spectrum of electrode after discharging to 2.1 V, it exhibits a single

Figure 4.12 High-resolution XPS (a) Li 1s (b) S 2p spectra of N,S graphene-Li2Sx asymmetric peak at around 56.0 eV, therefore, an additional peak with a +1.1 eV shift was fitted and attributed to the Li in the polysulfides interacting with doped N (Li–N) because normally, the Li 1s XPS spectrum only show the symmetric peak around 55.5 eV, this result is consistent with the observation for nitrogen/sulfur-doped cellulose bonded LiPSs, suggesting that N, S doped graphene can trap the polysulphides. The S

149 2p spectrum also sheds light on the interaction between the lithium polysulfide and

N-S co-doped graphene, the new peak at around 164.6 could be ascribed to the –C=S– bonding which derived from the interaction between polysulfide and the N, S doped graphene because this peak was not detected in the S 2p spectrum of N-S co-doped graphene@S (Figure 4.13). Which further confirming that interaction between polysulfide and the N-S co-doped graphene. We also note that the peaks at 170.64 and

169.4 eV in N, S co-doped graphene electrode which discharged to 2.1V (Figure

4.13) originates from 1) surface oxy-group of graphene and 2) oxidization due to exposing to the atmosphere during the transferring to the XPS measurement.

Figure 4.13 High-resolution S 2p spectra of N,S graphene@ S

To further confirm and quantitatively characterize the polysulfide adsorption ability of N-S co-doped graphene, N doped graphene and none doped graphene, we carried out quantum mechanical calculations based on density functional theory, the results are shown in Table 3, and corresponding models for N-S co-doped graphene, none doped graphene are in Figure 4.14, for N-doped graphene is added in Figure 4.8. The

150 adsorption energies of Li2S and Li2S4 on pristine graphene are -0.94 and -0.49 eV,

respectively. The adsorption energies of Li2S and Li2S4 on N doped graphene are

around -1.41 and -1.01 eV, respectively. While for Li2S and Li2S4 on N, S co-doped graphene, our DFT calculations predict that thiophene-like sulfur can exist adjacent to pyridinic N, as shown in Figure 4.14c and d. The adsorption energies are -1.85 and

151 -1.25 eV, respectively (Table 4.3), which exhibits much stronger bonding effects than those on pristine graphene and N doped graphene. Furthermore, there are 0.227 and

0.019 electrons transferred from Li2S and Li2S4 to pristine graphene while 0.555 and

Table 4.3 Comparisons of the adsorption energies of graphene, N-graphene and

N,S-graphene with Li2S, and L2S4.

Pyridinic like Pyrrolic like N, S doped Eads (eV) Graphene graphene graphene Graphene

Li2S -0.94 -1.41 -1.15 -1.85

Li2S4 -0.49 -1.01 -0.68 -1.25

Table 4.4 Comparisons of the atomic charge transfers (QT) of graphene, N-graphene and N,S-graphene with Li2S, and Li2S4.

QT (e) Graphene Pyridinic N Pyrrolic N N, S doped Gr

Li2S 0.227 0.787 0.69 0.555

Li2S4 0.019 0.12 1.15 0.317

0.317 electrons are transferred from Li2S and Li2S4 to N, S co-doped graphene

(Figure 4.4), suggesting that more charge migrates into co-doped graphene as lithiation proceeds. Therefore, the co-doping of N and S in graphene can significantly enhance the adsorption of lithium polysulphides as well as lithium sulfides.

The Brunauer–Emmett–Teller (BET) results also elucidate the function of K2CO3 chemical activation. As shown in Figure 4.9b, the optimized A-NSG showed a high

BET surface area of 1012 m2 g−1 and a large pore volume of 1.8 cm3 g−1. A

152

Figure 4.15 Electrochemical characterization of A-NSG material as the cathode of a

Li-S battery. (a) Cyclic voltammetry (CV) measured between 1.7 and 2.7 V at a sweep rate of 0.1 mV s-1 for the first, second, third, and fifth cycles. (b) Galvanostatic charge and discharge profiles for different cycles at 0.2C. (c) The discharge capacities of high and low voltage plateaus. The onset voltage of the low plateau was defined around 2.0 V. (d) Long term cycling performance test of the A-NSG@S electrode at

0.2C discharge rate and corresponding Coulombic efficiency.

153 hierarchical nanoporosity including both micropores (2 nm) and mesopores (25–40

nm) resulted from the gas evolution upon pyrolysis of the polymers and K2CO3 chemical activation. In contrast, the NSG sample, which only experienced a

nitrogen-sulfur co-doping process without K2CO3 chemical activation, exhibits a low

2 −1 3 −1 surface area of 230 m g and a pore volume of 0.3 cm g . Therefore, the K2CO3 chemical activation greatly contributes to the nanoporosity of the material. Moreover, the large pore volume plays a vital role in embedding sulfur efficiently and providing enough void space for volume expansion, while the high surface area is important for

interfacial chemisorption of lithium polysulfides. On the other hand, the ID/IG intensity ratio of A-G of Raman spectrum analysis is clearly higher than that of GO or

NSG (1.12 vs 1.06, 1.1), which indicates that K2CO3 activation could also induce GO

to generate more defects for effective reactions. In summary, the activation by K2CO3 yields a high BET surface area, creats porous structure and continuous 3D networks.

The electrochemical performance of A-NSG@S nanocomposite electrodes were tested by cyclic voltammetry (CV) and charge/discharge cycling behaviors of sulfur cathodes. Two cathodic peaks at 2.35 V and 2.05 V (Figure 4.15a) are consistent with the voltage plateaus in Figure 4.15b. The upper plateau at 2.35 V contributes to the transformation of cyclo-octasulfur to long-chain soluble lithium polysulfides. The lower plateau at 2.05 V is associated with the decomposition of those polysulfides to insoluble short-chain lithium sulfides, which contributes to the major capacity of

A-NSG@S electrode. Notably, there are no shape or position changes of the redox 154 peaks, implying excellent reversibility and stability of the A-NSG@S nanocomposite electrodes.6 Figure 4.15d shows the long-term cycling performance of the A-NSG@S at 0.2C up to 600 cycles. The electrode delivered an initial capacity of 1178 mAh g-1 and retained a high capacity of 780 mAh g-1 after 600 cycles. It should be noted that the capacity gradually increases in the first 10 cycles upon cycling at 0.2C current rate, which has been reported as a common phenomenon for electrodes applied in lithium sulfur battery.235, as shown in Figure 4.15c and d, the total capacity increased is mainly derived from the capacity accumulated in the low voltage plateau. There is almost no polysulfide lost in the high voltage plateau during the first few cycles.

Figure 4.16 Long-term cycling performance test of the A-NSG@S electrode at 1C discharge rate and the corresponding Coulombic efficiency.

Normally, for sulfur battery, the capacity is gradually decreased due to dissolution of polysulfides at the low voltage plateau. However, due to the strong polysulfide absorption effect of N-S co-doping A-NSG, the polysulfides could be effectively

155 trapped by the N-S co-doping graphene. Additionally, we suppose that not all sulfur takes part in the reaction with lithium ion due to poor conductivity in the initial

Figure 4.17 Electrochemical characterization of A-NSG material as the cathode of a

Li-S battery. (a) Galvanostatic charge and discharge profiles at 0.1 C, 0.5C, 1C, 2C,

5C. (b) Voltage plateaus for charge and discharge processes over 200 cycles at 0.1 and 1 C. (c) Discharge/Charge capacity cycled at various rates from 0.1 C, 0.5C, 1C,

2C, 5C. several cycles, and because of the “sulfur activation” in the first initial cycles, more and more sulfur participated in the discharge reaction step by step, thus the capacity generated in the low voltage plateau gradually improvedˈand resulting in the increase of total capacity. Furthermore, the capacity experienced only a slight

156 decrease during the 100th to 600th cycles, which further demonstrates long term cycling stability. When increasing the current rates to 1C (Figure 4.16), the

A-NSG@S electrodes achieved a capacity of 994 mAh g-1 in the initial cycle, and showed superior cycling performance with capacity degradations rate of 0.81% per cycle up to 300 cycles. The cycling performance is much better than that of the previously reported data with high-loaded sulfur cathodes.

The high conductivity and porous structure enables the A-NSG@S composites to be kinetically sustainable for a wide range of current rates. Generally, the rate performance of a battery can be reflected from two characteristics i.e. the capacity maintenance at stepped rates and polarization during high-current charge and discharge processes. To explore the rate performance of A-NSG@S cathodes, we tested batteries at different current rates. The voltage gap between charge and discharge plateaus was slightly enlarged when the current rate was increased tenfold, and was stable during 200 cycles (Figure 4.17b). When increasing current rates, reversible capacities of 1203, 1120, 987, 855 mAh g-1 were obtained at current rates of 0.1C, 0.5C, 1C, 2C (Figure 4.17c), respectively. Even at a high current rate of 5C, a reversible capacity of 651 mAh g-1 was still retained. When the current rate was gradually returned back to 0.1C, a reversible capacity of 1021 mAh g-1 was recovered, showing excellent high rate capability.

157 The Incremental Capacity Analysis (ICA) allows us to probe any gradual changes in the electrochemical behavior of rechargeable batteries during a cycle-life test with greater sensitivity than those based on the conventional charge and discharge curves.

Figure 4.18 (a) dQ/dV plots of A-NSG@S electrode. (b) Electrochemical impedance spectra of A-NSG @S lithium-sulfur battery at different cycles.

The advantage of ICA is achieved by transforming either the voltage or inflection points on V vs. Q curves into clearly identifiable DQ/DV peaks on the ICA curves,

158 depicting the solid solution or phase transformation characteristics. By monitoring the evolution of these dQ/dV peaks upon cycling, we can access key information on the cell chemistry behavior. Figure 4.18a shows the incremental capacity peaks derived from the discharge curves (0.2C) for different cycles. The ICA peak located at 2.05 V is slightly broadened upon cycling, suggesting that the characteristics of charge transfer might have been altered, which may be associated with an increased diffusion path for lithium polysulfides due to volume expansion during cycling. However, this voltage vibration caused by volume expansion is almost negligible compared with other reports, because our 3D porous graphene networks and graphene layers can buffer the volume expansion/contraction of activated materials upon cycling.

Meanwhile, both positions of the two peaks remain almost unchanged, implying that polarization is almost unchanged during 600 cycles. This result is in consistent with the outcomes shown in Figure 4.17b and Figure 4.18b. Moreover, although a slight intensity reduction appears in both peaks due to the inevitable active material loss, the position as well as sharpness of both peaks is maintained unchanged during 600 cycles, which further confirms that electrode materials preserve excellent electrochemical stability. In particular, the chemically activated A-NSG@S functions as a 3D immobilizer to accommodate a large amount of sulfur active materials, and the internal graphene layer greatly increases the polysulfide utilization due to the improved overall electrical conductivity of the structure.

159 To further understand the improved electrochemical performance of A-NSG@S composites, electrochemical impedance spectroscopy (EIS) was employed to characterize the cathode before cycling, after the first cycle and 600 cycles. Nyquist plots are shown in Figure 4.18b. For the cathode before cycling, the impedance curve is composed of one depressed semicircle at high frequency and a short inclined line in the low frequency regions. The semicircle corresponds to the internal resistance of the cathode including bulk impedance and interfacial impedance. In addition, the inclined line at low frequency reflects the lithium ion diffusion into the active mass.86, 236 In contrast, the cathode after cycling exhibited two depressed semicircles followed by a long sloping line. The depressed semicircle in the high frequency region represents the charge transfer process at the carbon matrix interface, which also dominates the reduction reaction along the upper voltage plateau. While the semicircle at middle

frequencies could be ascribed to the formation of insoluble Li2S, Li2S2, which controls the lower voltage plateau.236 Based on this analysis, the two features of impedance spectra in Figure 4.18b can be used to deduce the equivalent circuits 1 and 2 respectively (shown as the inset). In the equivalent circuits, Re represents the resistance of the electrolyte, Rct is the charge transfer resistance at the conductive agent interface, and CPE is a constant phase element, which is associated with the roughness of the particle surface. CPE2 describes the space charge capacitance of the

Li2S (or Li2S2) film and Rg is the resistance in the Li2S (or Li2S2) film. ZW is the

160 Warburg impedance due to the diffusion of the polysulfides within the cathode. The fitting results are shown in Figure 4.18b. Parameter identification results are

Table 4.5 Parameters identification by modeling the impedance spectra in Figure 18b

-1 -2 -1 -2 Cycle Re Rct Rg (ohm) Y1(Ω cm n 1 Y2(Ω cm n 2 number (ohm) (ohm) s-n) s-n)

Pristine 7.23 34.71 - 0.000027 0.757 - -

1 7.98 36.23 18.23 0.00003 0.84 0.087 0.93

600 8.329 41.56 23.23 0.0009 0.6523 0.1094 0.865 summarized in Table 4.5, according to the equivalent circuit. It can be seen that Re experienced a slight increase during cycling due to the inevitably dissolution of polysulfides. In contrast, Rg increased remarkably during the first cycle, compared with that of a pristine cell, while it shows a modest rise in the following cycles. We

can draw a conclusion that the solid Li2S (or Li2S2) appeared at the very beginning of the second voltage plateau, but due to the dominant pyridinic and pyrrolic N in the carbon lattice of A-NSG@S, the electron distribution is modified and the affinity for

insoluble Li2S is improved, and the reversible conversion of Li2S/polysulfide/S is correspondingly promoted. Similiarly, Rct exhibits the same trend as Rg during cycling. As for the equivalent component CPE2, which contains two parameters Y and n, the dimension for Y is Ω-1·cm-2·s-n or S·cm-2·s-n, the parameter n is a dimensionless index. The admittance response expression for CPE is:

Y=Y0ωncos(nπ/2)+jY0ωnsin(nπ/2) (1)

161 However, after 600 cycles, both Y and n mostly remained stable as shown in Table

4.15, showing that the electrochemical stability of CPE2 is excellent after 600 cycles.

Most importantly, this electrochemical stability also contributes greatly to improve the electrochemical performance of lithium-sulfur batteries.

To demonstrate whether the co-doped heteroatoms of N and S and the graphene

chemically activated by K2CO3 have positive functions on the electrochemical performance of Li-S batteries, several other comparison mateials were prepared and tested. These include sulfur composites with about 70 wt% of sulfur prepared via a two step melt-diffusion method to produce nitrogen-sulfur co-doped graphene (NSG),

K2CO3 activated graphene (A-G), nitrogen-sulfur co-doped and K2CO3 activated graphene (A-NSG), and graphene(rGO), denoted as NSG@S, A-G@S, A-NSG@S, and rGO@S, respectively. For comparision, the cycle performances of these electrodes were tested at the same current density 0.2C rate within a voltage window of 1.7-2.7 V, as shown in Figure 4.19. The capacity retention for rGO@S, NSG@S and A-G@S electrodes up to 100 cycles is 38%, 55.2%, 56.8%, respectively, which is not comparable with that of A-NSG@S. It also further confirmed the effective

function of co-doping as well as K2CO3 activation. Furthermore, we compared the stability of A-NSG@S with other co-doped hosts in Table 4.6 120, 208, 221, 237, 238

Apparently, the sulfur electrode made of A-NSG exhibits a decay rate of 0.056% per cycle up to 600 cycles, which it superior to that of other co-doped hosts.

162 Figure 4.19 Cycling performance of A-NSG@S material cycled at 0.2C, in comparison with NSG@S, A-G@S, rGO@S material. Specific capacity values were all calculated based on the mass of sulfur.

Additionally, we can see that the cycle performance of N,S doped carbon framework is better than that of N,O doped. This is because that N and S co-doped carbon exhibits asymmetric charge density distribution, namely, N bears a negative charge and S bears a positive charge, which benefits charge transfer. Therefore, the excellent electrochemical performance of the A-NSG@S electrodes is clearly associated with their unique structure: (i) The 3D porous graphene network and graphene layers facilitate electron transfer, and remarkably buffer the volume expansion/contraction of active materials upon cycling. (ii) The inner defects, edges and porous structure not only permit a high sulfur loading in a homogeneously dispersed amorphous state, but also are beneficial for electrolyte access to the active sulfur component, leading to efficient reactions with Li+. (iii) Combined physical adsorption of lithium polysulfides onto porous graphene and the chemical binding of polysulfides to N and S sites in the

163 A-NSG, promotes reversible Li2S/polysulfide/S conversion, realizing high performance Li-S batteries with long cycle life and high-energy density.

Table 4.6 Recent advance in the dual doped carbon framework to host sulfur for Li-S batteries Carbon Composites Cycle Decay Rate(per Ref life(cycles) cycle) A-NSG(N,S doped 600 0.056% This work porous graphene) N,S-graphene 500 0.078% Nature Commun. 2015, 6, 7760 N,S doped carbon 500 0.065% Adv Mater. 2015, 27, 6021-6028 N,O doped CNTs 200 0.15% Adv Mater Interfaces. 2014, 1, 1400227 N,O-graphene 350 0.057% Nature Commun. 2014, 5, 5002 N,O-doped porous 200 0.05% Angew Chem Int Ed. 2015, 54, carbon 4325–4329

4.4 Conclusions

In summary, we designed an efficient strategy for preparing crumpled N-S co-doped porous graphene (A-NSG) via chemical activation of polypyrrole (PPy)

functionalized graphene sheets with K2CO3. The A-NSG @S electrodes with a high sulfur loading of 72.4 wt% exhibits low polarization, stable cycling performance, and excellent rate capability, compared with rGO@S, NSG@S and A-G@S electrodes.

The excellent performance could attributed to synergic effects, including the high surface area and high conductivity of the porous matrix, as well as the unique lithium polysulfides binding capability of the N, S functional groups in the A-NSG sheets.

Furthermore, pyridinic nitrogen and thiophenic sulfur in carbon lattice of A-NSG@S

164 can modify the electron distribution and improve the affinity for Li2S as well as

Li2S4, leading to the improvement in reversibility of Li2S/polysulfide/S conversion.

This work not only presents a simple method to synthesize N-S co-doped hierarchical porous graphene, but also demonstrates the excellent performance of A-NSG cathode materials for lithium-sulfur batteries.

165 Chapter 5 Nitrogen Doped Mesoporous Yolk-Shell Carbon Spheres for High

Performance Lithium-Sulfur Batteries

5.1 Introduction

As energy storage systems, rechargeable lithium-ion batteries (LIBs) play an important role in portable electronic devices because of their high performance compared with other batteries such as nickel-cadmium batteries and nickel-metal hydride batteries. However, the increased demands of fast-developing electrical vehicles, large-scale energy storage systems, and smart portable electronics require batteries with high-energy density, low cost and long cycle life. The tough issue for further increasing the performance of LIBs is the capacity mismatch between

cathodes and anodes. The commercial cathode materials such as LiCoO2, LiMn2O4,

LiFePO4 and LiNixMnyCozO2 can only deliver capacities in the range of 140-200 mAh g-1. However, as cathodes for lithium-sulfur batteries, the sulfur cathode has a high theoretical capacity of 1672 mAh g-1.120, 125 Additionally, some valuable features of sulfur, such as the natural abundance, low cost, and nontoxicity, making lithium-sulfur batteries an attractive system for energy storage and conversion.61, 72, 98,

239-250

In lithium-sulfur batteries, the large molecules cyclo-S8 present in S cathodes and undergo multi-step open-ring reduction reactions with lithium, through breaking S–S bonds, resulting in long-chain lithium polysulfides during the discharge process.

166 Eventually those lithium polysulfides are reduced to Li2S2 and Li2S. However, the practical application of Li-S batteries still suffer from some tough obstacles, such as

the insulating nature of sulfur and discharge products of Li2S2 and Li2S, large volume

expansion/contraction (~80%) during cycling, dissolution of polysulfides (Li2Sx, 4 ≤ x

≤ 8) in electrolytes and shuttle effects, which inevitably lead to the low sulfur utilization as well as rapid capacity degradation.66, 86, 107, 210, 211, 251 To tackle these challenges, extensive research have been conducted on cathodes, electrolytes, separators, and lithium-metal anodes for preventing the dissolution of polysulfides and facilitating the utilization of sulfur. Among these approaches, confining sulfur in a conductive host has been proven promising. In the pioneer work of Nazar’s research group, sulfur was confined into highly ordered mesoporous carbon (CMK-3) and showed a reversible capacity over 1000 mAh g-1 for 20 cycles.87 Since then, significant advances have been achieved using multiple carbon materials as hosts for sulfur cathodes, including carbon nanotube/fibers, graphene coated hybrid structures, hollow carbon structures and composite carbon structures. These sulfur hosts serve not only as conducting agents to facilitate the electron transport of the active sulfur materials, but also as reservoirs to capture the soluble lithium polysulfides and to improve the utilization of the active materials.73, 145, 202, 214, 252, 253 Although these sulfur composites exhibit high specific capacities in the initial cycle, they usually experience a rapid decrease of capacities in the subsequent cycles. Hence, cycle stability remains a tough problem for practical applications. Conventionally, sulfur is

167 introduced into the pores within a carbon matrix via a melt-diffusion or vapor impregnation process. It is widely accepted that sulfur can diffuse out of the small pores of the porous carbon and then be accessed by the electrolyte for extended durations.254, 255 Thus lithium polysulfides will eventually be dissolved due to the weak physical adsorption between sulfur and porous carbon matrics. Additionally,

significant volumetric expansion from sulfur to Li2S during lithiation renders the protective host material layers susceptible to crack or break, resulting in the loss of

159, 237, 256, 257 polysulfides. Moreover, the isolated Li2S particles formed in the discharge process will not participate in the subsequent charge process, leading to capacity degradation.107, 239, 258, 259 Therefore, it is necessary to develop well-defined architectures that confining suitable hollow cavity around sulfur particles, which not only alleviate volume expansion during cycling but also prevent the dissolution of polysulfides by the strong chemisorption.113, 209, 211, 260-262

The chemical properties of sp2 carbon materials can be further modified by introducing heteroatom doping, such as phosphorus, sulfur, boron and nitrogen, which demonstrated significant potential for lithium polysulfide confinement. Some groups reported that nitrogen doping in mesoporous carbon effectively facilitates the chemical adsorption of sulfur atoms on oxygen-containing functional groups because surface modification with nitrogen atoms greatly enhances the electrochemical reactivity and electronic conductivity of carbon matrices.205

168 Herein, we report a strategy to dually confine sulfur and polysulfides. The nitrogen doped yolk-shell carbon composites (denoted as NYSC) with a hierarchically porous architecture synthesize through a facile approach, in which polydopamine serves as a nitrogen source, the interpenetrating mesoporous in the NYSC is specific for sulfur anchor and accommodation, and this host shows strong affinity for sulfur and polysulfides, the appropriate cavity in the NYSC is effectively designed for alleviating the sulfur volume changes during cycling. Nitrogen doped core-shell carbon spheres (NCSC) were also synthesized and applied as cathode hosts for Li-S batteries. The electrochemical testing demonstrated that the dissolution/diffusion of polysulfides in NYSC@S cathodes can be more effectively suppressed by the yolk-shell architecture, compared with core-shell structure cathodes. The nitrogen doped yolk-shell sulfur cathodes showed a high reversible capacity of 909 mAh g-1 at the current rate of 0.2C and excellent cycling stability.

5.2 Experimental section

5.2.1 Synthesis of NYSC(Nitrogen Doped Yolk-Shell Carbon Sphere)

A mixed solution was made of 0.15 mL ammonia solution, 6.0 mL absolute ethanol and 15 mL deionized water (DI). 0.15 g resorcinol and 0.15 g CTAB were dissolved in the solution, and then stirred for 1 h. 0.75 mL TEOS and 0.21 mL 37 wt.% formaldehyde solution were added to the reaction solution and vigorously stirred for

24 h at room temperature, and then heated for 24 h at 100 ć in a Teflon-lined

169 autoclave. The as-prepared precusors were filtered and dried at 70 ć overnight.

Then dissolving 100 mg precusors and 20 mg Tris buffer (PH=8.5) into the 100 mL deionized water (DI), following by the adding of 30 mg dopamine, and stirring overnight before filtered and dried. After that, the products were heated at 700 ć

for 3 h under a N2 flow (2ć/min) in a tube furnace to obtain carbon/silica spheres.

Finally, the carbon/silica spheres were washed by 10 wt.% HF aqueous solution to remove silica. Followed by subsequent washing with deionized water, the nitrogen doped yolk-shell carbon spheres were obtained.

5.2.2 Synthesis of NYSC@S (Nitrogen Doped Yolk-Shell Carbon Sphere@S)

0.806 g sodium thiosulfate and 1g sodium sulfide were added into a mixture solution containing 100 ml deionized water and 20 ml ethanol. Then 100 mg NYSC were added into the solution and vigorously stirred for 1h before adding in Hydrochloric acid (1 M, 40 ml) solution. Afterward, the dark suspension was washed with DI water and dried in a vacuum oven at 60ć for 12h. Finally, the composite was heated to

155°C at a heating rate of 2°C min-1 in a tube furnace under flowing argon at 100 sccm.

5.2.3 Characterization of materials

The morphology and crystal structures of the materials were characterized by field emission scanning electron microscopy (FESEM, Zeiss Supra 55VP) and transmission electron microscopy (TEM, JEM-2011, JEOL). The crystal structure and

170 phase of materials were characterized on a Siemens D5000 X-ray diffractometer using

Cu-KD radiation at a scanning step of 0.02° per second. Raman spectra were measured by a Renishaw in Via Raman spectrometer system (Gloucestershire, UK) equipped with a Leica DMLB microscope (Wetzlar, Germany) and a 17 mW 633 nm Renishaw helium neon laser at 50% power. Thermogravimetric analysis (TGA) was used for

evaluating the sulfur amount in the NYSC@S composites. N2 adsorption–desorption measurements were conducted using a 3 Flex surface characterization analyser at 77

5.2.4 Electrochemical measurement

Commercially available carbon paper (SIGRACET CDL 10 BA, Ion Power, Inc.) was used as a current collector in this work. An acid treatment were performed to functionalize CP with oxygen containing groups. Pristine CPs were cut into circular

disks with a diameter of 1.4 cm and immersed in a mixture of HNO3 for 3 h. The acid-treated CPs were then washed by water and dried at 60 °C for 24 h. 0.5 M

NYSC@S slurry was prepared by adding 50 mg NYSC@S into 5 mL 1,

2-dimethoxyethane (DME). To fabricate CP-NYSC@S electrodes, 85 PL NYSC@S slurry were slowly dropped onto CPs disk and then dried in argon. The actual sulfur content in the whole electrodes was around 1.8 mg/cm2, which was further confirmed by the TGA analysis. CR2032 coin cells were assembled in an argon-filled glove box

(Mbraun, Unilab, Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm. Lithium foil was used as a counter electrode. The 171 electrolyte was made of 1 M lithium bis-(trifluoromethanesulfonyl) imide (LiTFSI)

and 1 wt% lithium nitrate (LiNO3) in 1,3-dioxolane and 1,2-dimethoxy-ethane

(volume ratio 1:1). For each electrode, around 30 uL electrolyte was added in the coin cell. The electrodes were dried overnight at 60 Ԩ. Electrochemical performances were tested by a LAND-CT2001C battery system. The cells were tested in the fixed voltage range 1.7−2.7V at current rates of 0.5C, 1C, 2C, 3C, 5C. Cyclic voltammetry

(CV) was measured using a CHI660C electrochemistry workstation in the voltage range 1.7 to 2.7 V (vs. Li/Li+).

5.3 Results and Discussion

Figure 5.1 Synthesis procedure of the mesoporous NYSC@S. 172 The synthesis procedure of NYSC is shown in Figure 5.1. The synthesis procedure starts with the Stöber sol-gel method using cetyltrimethyl ammonium bromide

(denoted as CTAB) as templates, resorcinol-formaldehyde (denoted as RF) as carbon source, tetraethyl orthosilicate (denoted as TEOS) as a pore forming agent. In the beginning, RF precursor forms negatively charged emulsion droplets through the hydrogen bonding of ammonia, water, alcohol, resorcinol, and formaldehyde. Silicate oligomers hydrolysed from TEOS are also negatively charged; while CTAB with

Figure 5.2 (a) SEM image of the precusors obtained from hydrothermal process. (b)

SEM image of carbon-silica after calcination. (c, d) SEM image of the NYSC. positive charge can bind to the surface of the formed RF emulsion droplets and silicate oligomers through electrostatic interactions.263 Under the catalysis of ammonia molecules, the cross-linkage of RF droplets and silicate oligomers occur but 173 at different rates. The hydrolysis polymerization of RF is much faster than that of

TEOS at the beginning of the sol–gel process. Thus, only a few silicate oligomers participate in the co-assembly of CTAB and RF through the electrostatic interactions, resulting in the formation of yolk. With the sol–gel process prolonged, the concentration of RF emulsion was gradually decreased due to the RF being consumed by the growth of yolk, leading to a slower polymerization rate, but matches well with

174

Figure 5.3 (a-d) EDX maps of the precursor obtained from hydrothermal process. (e)

EDX spectrum of the precursor. (f)Line scanning profiles of a single precursor. the hydrolysis polymerization rate of TEOS. Consequently, the gradient hydrolysis and condensation of silicates start. The silicate oligomers from the hydrolysis of

TEOS together with RF emulsion droplets interact with CTAB to co-assemble at the surface of the yolk, thus creating the hybrid shells and then forming a yolk-shell architecture.

The precursor obtained through the solvothermal approach, exhibited a uniform

distribution of spherical morphology with the particle size of around ׽815 nm as shown in Figure 5.2. The energy-dispersive X-ray (EDX) mapping images clearly reveal the existence of C, Si, O elements, and the line scaning results in Figure 5.3 show that C is mainly concentrated in the centre and sparsely distributed elsewhere

175 throughout the sphere; While Si is primarily found near the surface of the sphere.

After carbonization, CTAB were removed and the RF

Figure 5.4 (a, b) SEM image of carbon-silica after calcination.

Figure 5.5 EDX maps of carbon-silica after calcination. (a) overall image. (b) carbon,

176 resin polymers were converted into a porous carbon-silica framework. At the same time, the de-hydrogenation of silica and polymerization of resin polymer occurs, leading to the shrinkages of frameworks. The obtained carbon-silica composites preserved the same morphology as the precursor, exhibiting homogeneous carbon spheres with sizes around 815 nm (Figure 5.2b, 5.4). The silica and carbon

Figure 5.6 (a) SEM image of NYSC. (b) High-magnification SEM images of the

NYSC. (c) Particle size distribution. (d) TEM image of the NYSC. distributions in the EDX mapping shown in Figure 5.5 are also consistent with the precursor, which demonstrates the robustness of this material. The silica is primally distributed in the shell, partially occupies the inner void, and only a small amount is left in the centre area. The silica distribution is according to a radially increasing trend,

177 while the carbon element follows an opposite distribution. This suggests that the shell of the NYSC sphere consists of carbon and silica, indicating

Figure 5.7 (a) N2 isothermal curves and pore size distributions of NYSC and

NYSC@S. (b) XPS spectrum of NYSC, (c) XRD patterns and (d) Raman spectra of the NYSC@S, NYSC, Bare Sulfur. that these carbon/silica spheres have a structure based on the mutual inter-penetration of these two elements. The inner core is totally composed of carbon. Finally, porous carbon spheres were obtained by etching silica away. It is clearly observed from the

SEM images that the final porous carbon spheres remained the spherical shape with an average diameter about 815 nm (Figure 5.6a, b, Figure 5.2c-d). Moreover, the

178 TEM image (Figure 5.6d) also exhibits distinct yolk-shell structure, the spacious void between the exterior shell and the interior yolk is clearly identified due to low contrast

Figure 5.8 N2 isothermal curves and pore size distributions of (a)carbon/silica composite , (b)NYSC@S. in the TEM result. The thickness of the shell is about 140 nm and the average diameter of the spherical yolk is around 410 nm. It is also worth noting that by removing the silica template from this “inter-penetration mesostructure”, these

179 inter-penetrating interfaces are converted into interconnected meso-size tunnels, showing a large pore volume and acting as a crucial channel for small sulfur molecules to easily penetrate or diffuse within the hollow tunnels. Thus, it can be confirmed that the mesopores formed throughout the shell of NYSC is generated by the removal of silica. Nitrogen sorption analyses of both NYSC-precusor and NYSC

show type-IV curves with H2 hysteresis (Figure 5.7a, Figure 5.8a), demonstrating the presence of mesopores with a uniform pore size. The BET surface area for the

NYSC-precursor is 290.22 m2/g and the pore volume is 0.34 cm3/g (Figure 5.8a); while NYSC has BET surface area of 780 m2/g and the pore volume is 1.38 cm3/g

(Figure 5.7a), respectively. This confirms that the removal of silica helps to increase the mesopore surface area and generate more mesopores in the carbon hosts, particularly in the shell. The pore size distributions are associated with: (i) mesopores

(2-5 nm) yolk from the phenolic resin pyrolysis, (ii) inter-penetration mesoporous

(2-10 nm) shell by removing the surfactant and the silica from the inter-penetration mesostructure of the precusor. This porous structure design is beneficial for increasing the effective contact area with conducting host. In addition, the effective doping of N atom has been confirmed by the X-ray photoelectron spectroscopy (XPS) results, in the N 1s spectrum of NYSC (Figure 5.7b), three components including pyridinic N (28.68 wt%), pyrrolic N (50.13 wt%), and graphitic N (21.19 wt%), can

180

Figure 5.9 (a) SEM image of the NYSC@S. The inset is the magnified SEM image of

NYSC@S. (b) TEM image of NYSC. (c) HRTEM image of NYSC@S. (d) TEM image of NYSC@S. (e, f, g, h) EDX spectra of NYSC@S. be attributed to 398.4, 400.2 and 402.1 eV, respectively. The pyridinic and pyrrolic N had been reported to improve the affinity and binding energy of non-polar carbon atoms with polysulfides.264

However, for the unique inter-penetration mesoporous yolk-shell architecture, it is relatively difficult to homogeneously load sulfur into the inner yolk of NYSC by the

181 melting-diffusion method via simple capillary force, especially for those inner and narrow mesopore in the NYSC nanospheres. Therefore, we developed a “two-step” approach (wet-precipitation and melting-diffusion) to load sulfur. Sulfur, prepared by wet-precipitation, is uniformly pre-loaded into the inner and narrow carbon chambers, which ensures the diffusion of sulfur into the yolk of NYSC. The subsequent melting-diffusion at 155 ć can facilitate the homogenous distribution of sulfur into the interconnected mesoporous NYSC host and increase the intimate contact between the active material and the conductive carbon spheres. This “two-step” approach is particularly effective for hierarchically hollow hosts. Sulfur was firstly pre-impregnated into NYSC via liquid-phase infiltration. This process can be described as:

2Na2S + Na2SO3 + 6HCl Æ 3S↓ + 6NaCl + 3H2O.

During the heat treatment at 155 ć for 10 h, melted sulfur with low viscosity was then absorbed into the mesopores in shells through strong capillary forces, forming

NYSC@S composites. Compared with the semi-transparent pristine NYSC, the obtained NYSC@S composites still preserve the spherical morphology (Figure

5.9a), no agglomerated sulfur particles are detected on the surface of NYSC@S in

Figure 5.9a. This further indicates that sulfur homogeneously diffused into the porous of NYSC@S. The SEM and TEM observation in Figure 5.9 are consistent with the

XRD patterns as shown in Figure 5.7c. Compared with the crystalline peaks of pure

182 sulfur and NYSC@S, the absence of sulfur crystal peaks for NYSC@S is noteworthy.

This indicates that amorphous sulfur is highly dispersed in NYSC spheres. Meanwhile, as shown in the Raman Spectra (Figure 5.7d), the pure

Figure 5.10 TGA curve of the (a)NYSC@S, (b)NCSC@S

ulfur powders exhibit characteristic peaks in the range 100-500 cm-1, associated with

vibrations of the S-S bond in S8 species. However, NYSC@S did not show

183 noticeable peaks of sulfur, which is also consistent with the conclusion from XRD

patterns. This implies that S8 molecules were loaded into the NYSC mesoporous without long-chain format, similar to small sulfur molecule dispersion. It is reported that the amorphous state of sulfur also contributes to high ultilization of sulfur in electrochemical lithiations.259 In order to identify the sulfur distribution in the NYSC.

High-resolution (HRTEM) were conducted to identify the sulfur distribution. As can be seen in Figure 5.9b and d, the dark area in the carbon shell and yolk implys a heavy element, which can be attributed to sulfur (heavier than carbon). The elemental mapping results in Figure 5.9e-h also illustrate the spread of sulfur (blue)ˈcarbon

(red)ˈnitrogen(yellow) in NYSC@S. It is clear that sulfur is effectively distributed in the the outer shell. Normally, sulfur prefers to accommodate in the inner core and outer shell instead of the cavity area, owing to the strong capillary forces between sulfur and yolk-shell carbon structure. The sulfur impregnation procedure begins with pores in the shell and gradually diffuses inward, finally entrapped by the yolk. The sulfur amount in the NYSC@S nanocomposites was determined by thermogravimetric analysis (TGA) under nitrogen atmosphere with a heating rate of

10 Ԩ min-1, as shown in Figure 5.10. The sulfur loading in NYSC@S was 75.4 wt%, thus the mass loading of sulfur in each electrode is around 1.8 mg/cm2. Nitrogen adsorption-desorption results show the change of inter structure after sulfur loading.

The specific surface area and the number of mesopores dramatically decreased. The

BET results (Figure 5.8b) confirm the homogenenous sulfur occupation of the

184 mesopores, which reduces the surface area to 152 m2/g, the pore volume to 0.17 cm3/g-1 in the NYSC@S. Obvious peak does not appear in the pore size distribution plot, suggesting that the sulfur occupied the mesoporous of NYSC@S. Meanwhile, the hollow cavity still existed after sulfur loading, contributing to alleviate the volume

Figure 5.11 High resolution SEM and TEM images of the NYSC (a, b, c) and NCSC

(d, e, f). In all figures, the scale bar is 300 nm expansion of sulfur during cycling. Additionally, the homogenous distribution of nitrogen in the element mapping result (Figure 5.9h) also confirms the sucessfully doping of nitrogen atom.

Interestingly, when adjusting the amount of CTAB in the synthesis procress, the final product still keeps the spherical morphology but the structural changes to the 185

Figure 5.12 Electrochemical characterization of materials as the cathodes for Li-S batteries. (a, b) Cyclic voltammetry (CV) tested between 1.7 and 2.7 V at a sweep rate of 0.1 mV s-1 for NYSC@S, NCSC@S, respectively. (c, d) Galvanostatic charge and discharge profiles for 1st and 300th cycle at 0.2C, respectively. (e) Long cycling performance of the NYSC@S, NCSC@S electrode at 0.2C discharge rate and corresponding Coulombic efficiency. core-shell structure in which the shell thickness shrinks to 93 nm from the original

NYSC thickness of 140 nm. Meanwhile, the core diameter expands to around 550 nm

186 compared with 410 nm in the NYSC (Figure 5.1c, f). Consequently, the void space decreased significantly. To compare the advantages of the yolk-shell structure in alleviating the dissolution of polysulfide and stabilizing the sulfur cathode, we tested the electrochemical performances of both NYSC@S and NCSC@S. The sulfur loading in NCSC@S was around 68.32 wt% (Figure 5.10b). Cyclic voltammetry (CV) results of the NCSC@S and NYSC@S as cathode materials at a scan rate of 0.1 mV s-1 are provided in Figure 5.12a and b, respectively. Apparently, two cathodic peaks at 2.35 V and 2.05 V (Figure 5.12a) are consistent with the voltage plateaus in

Figure 5.12c. The upper plateau at around 2.35 V is associated with the conversion of

cyclo-octasulfur to soluble intermediate polysulfides (Li2Sn, 4˘n˘8). The lower plateau at 2.05 V corresponds to the conversion of intermediate polysulfides to

insoluble short-chain lithium sulphides (Li2Sn, nİ2), which contributes to the major capacity of sulfur electrodes.7 Compared with the sharp peak changes in the first five cycles of the NCSC@S electrode, there are rarely shape or position changes of the redox peaks for NYSC@S electrodes, implying excellent reversibility and stability of the NYSC@S nanocomposite electrodes.

Figure 5.12c and d show the typical galvanostatic charge/discharge profiles within a potential window of 1.7–2.7 V at 0.2C in the first and 300th cycles of two sulfur electrodes. The NYSC@S electrode exhits two obvious discharge/charge plateaus,

representing the conversion of elemental sulfur (S8) to intermidate polysulfides and

finally to the short-chain Li2S2/Li2S. As for charging process, the anodic plateau 187 represents the reverse reaction from sulfides to polysulfides and then to sulfur.

Apparently, both the initial capacity and voltage gap between anodic and cathode plateaus of the NYSC@S and NCSC@S are almost same in the first cycle. However, after 300 cycles, the capacity of NCSC@S significantly decreased to 433 mAh g-1, while the capacity of NYSC@S is stable at 961 mAh g-1 (Figure 5.12d). Furthermore, these plateaus for NYSC@S electrode maintains well with low over potentials

(denoted as ∆E, voltage gap between anodic and cathodic plateaus) even after 300 cycles (e.g., 221 mV at 1st cycle and 301 mV at 300 cycles), suggesting a rapid kinetic reaction almost have no barrier. However, the NCSC@S electrode shows sloping discharge/charge plateaus and these plateaus even disappeared after 300 cycles. The voltage gap (∆E) is much higher (221 mV at 1st cycle and 489 mV at 300 cycles), compared to NYSC@S electrode, implying the relatively slow redox reaction kinetics.

Therefore, the dual confinement system by a yolk-shell structure contributes to electrochemical performance as well as cycling stability of the electrode. The long term cycling performance and corresponding Coulombic efficiencies of the NYSC@S and NCSC@S electrodes (0.2C) are shown in Figure 5.12e. The NYSC@S and

NCSC@S electrodes delivered an initial capacity of 1329 and 1052 mAh g-1 (Figure

5.12c), respectively. After 500 cycles, the capacity of NYSC@S was stable at 909 mAh g-1, while the capacity of NCSC@S dramatically decreased to 323 mAh g-1(Figure 5.12e), and the Coulombic efficiency of NYSC@S is also superior to that of NCSC@S. The NYSC@S and NCSC@S electrodes were also subjected to cycle at

188

Figure 5.13 (a) Discharge capacity at different rates of 0.5C, 1C, 2C, 3C and 5C for

NYSC@S,NCSC@S electrodes. (b) Galvanostatic charge and discharge profiles at

0.5C, 1C, 2C, 3C and 5C for NYSC@S electrodes. increasing current densities. When increasing current rates to 0.5C, 1C, 2C, 3C

(Figure 5.13a), reversible capacities were stable for NYSC@S, around 1280, 1050,

810, 670 mAh g-1, respectively. Even at a high current rate of 5C, a reversible capacity of 510 mAh g-1 was still retained. The corresponding charge/discharge profiles are provided in Figure 5.13b. Moreover, when the current density was gradually reversed back to 0.1C rate, a reversible capacity of 930 mAh g-1 was

189 recovered, showing excellent stability of this sulfur cathode structure. In sharp contrast, the NCSC@S electrode can only generate a capacity of 46 mA h g−1 at 5 C rate and then recovered to 277 mA h g−1 at a 0.1C rate. This further elucidated that the enlarged hollow cavity in the yolk-shell structure is

Figure 5.14 Electrochemical impedance spectra of NYSC@S, NCSC@S batteries for the (a) first cycle and (b) final cycle (500 th).

190 adequate to accommodate the volume expansion of sulfur during cycling, thus preserving the structural integrity of the shells to minimize polysulfide dissolution.

The robustness of the intact carbon shell and its mesoporous interpenetration tunnels could also limit polysulfide dissolution. Furthermore, the effective doping of nitrogen atom could also improve the affinity of polysulfides, facilitate the utilization of active material, which is contributing to the stability of long term cycle performance.

To evaluate the electrochemical characteristics of the NYSC@S and NCSC@S composite, Figure 5.14 presents the electrochemical impedance spectroscopy (EIS) analyses of NYSC@S cathode and NCSC@S cathode before cycling and after 500 cycles, respectively. Before cycling, the impedance spectra consist of a depressed semicircle and an inclined line from high frequency to low frequency. The depressed semicircle represents the internal resistance that composed of bulk impedance and interfacial impedance. The inclined line at low frequency reflects the Li ion diffusion into the active mass. As identified by EIS, the internal resistance of NYSC@S cathode is nearly the same as that of NCSC@S cathode as shown in Figure 5.14a.

However, after 500 cycles, the impedance plots of both cathodes consist of two depressed semicircles from high frequency to middle frequency (Figure 5.14b), respectively, and an inclined line at low frequency.236 The high frequency semicircle corresponds to the resistance of a solid-state layer of accumulated lithium sulfide, and the middle frequency semicircle is considered to be the charge transfer resistance. The charge transfer resistances of the NYSC@S cathode is much smaller than that of the 191 NCSC@S cathode,265 suggesting the superior electrochemical performance of

NYSC@S composite.

To exploit the integrity of the NYSC@S structure, the cells of NYSC@S and

NCSC@S were dissembled after 100 cycles and the corresponding cathodes were

Figure 5.15 (a,b) FESEM images of NYSC@S composites before and after 100 cycles. (c,d) FESEM images of NCSC@S composites before and after 100 cycles. observed by SEM. As shown in Figure 5.15b, a few broken spheres can be found in the electrode, but most of the NYSC@S particles show no morphology change after

100 cycles. On the contrary, the SEM image of the NCSC@S electrodes after 100 cycles gave a distinct evidence, in which more than half of the NCSC@S carbon spheres were cracked or broken(Figure 5.15d), suggesting that the integrity of core

192 shell structure not well remained during the lithiation process as shown in Figure

5.18. The elemental mapping result also confirms this as shown in Figure 5.16. Large amount of active materials leaked out of the shell and most spheres cracked after cycling for the NCSC@S electrode. In comparision, sufur is still well confined in the

Figure 5.16 (a-c) EDX maps of NCSC@S after 100 cycles. (d) EDX spectrum of

NCSC@S after 100 cycles yolk-shell carbon spheres as shown in Figure 5.17. Almost no agglomerated sulfur can be seen on the surface of the NYSC@S, revealing that this robust yolk-shell structure can maintain its integrity, effectively entrap lithium polysulfides derived from inner yolk and prevent them to diffuse into the electrolyte (Figure 5.18).

193 To experimentally prove the adsorption ability for polysulfides, an adsorption test was

conducted by soaking the as-synthesized NYSC and NCSC composites in 5 mM Li2S6

catholyte solutions. The Li2S6 solution was prepared by stoichiometrically dissolving

Figure 5.17 (a-c) EDX maps of NYSC@S after 100 cycles. (d) EDX spectrum of

NYSC@S after 100 cycles.

Li2S and sublimed S in a molar ratio of 1:5 in the mixed solvents of 1,3-dioxolane

(DOL) and 1,2-dimethoxyethane (DME) (1:1 v/v). Then, 15 mg of NYSC, NCSC

were soaked in 15 mL of 5 mM Li2S6 solutions at 30 °C for 24 h, respectively. The upper solution was carefully collected and sealed. Figure 5.19 display the digital

photos of pristine Li2S6 solution and the color changes of the Li2S6 solution after adding NYSC and NCSC spheres, respectively. After aging for 24 h, the color change

194 of Li2S6 in the dispersed suspension of NYSC is more obvious than that of NCSC,

indicating that much of the Li2S6 has been adsorbed by the NYSC spheres. This

Figure 5.18 Schematic of the lithiation process in core shell and yolk shell sulfur based morphologies. outcome proved that the yolk-shell structure coupling with the dopant nitrogen atom is more efficient in preventing polysulfide diffusion and stabilizing the dissolved polysulfides within the cathode compared with the core-shell morphology.

When compared with other carbon nanospheres mophologies, like double shell hollow carbon sphere,266 hollow-in-hollow carbon sphere,211 double-layered core−shell carbon sphere,254 hollow carbon sphere,267, 268 porous carbon sphere,99, 269 the nitrogen doped yolk shell sphere in this work accommodates high amount of sulfur content, and exhibits excellent stability compared with that of others provided 195 in Table 5.1. The excellent electrochemical performance of the NYSC@S electrodes could be ascribed to the rationally designed nitrogen doped yolk-shell architecture: (a)

The dopant N

Figure 5.19 Digital images of NCSC(a), NYSC(b) nanoparticles immersed in the

Li2S6/DOL/DME (1:1, v-v) solution. actoms serves as electron attracting actoms, leading to the nearby C atoms to be polarized and more active for anchoring sulfur and polysulfides. (b) The inner mesoporous “yolk” acts as a sulfur reservoir to entrap polysulfide species, enabling an effective utilization of the active material;212 (c) The inner void between the shell and yolk provides sufficient cavity to cushion the volume expansion during cycling;270, 271 which contributs to the stability of the cycling performance. (d) The robust mesoporous shell confers fast ion channels to facilitate ion transport and minimizes polarization effects, and the multimode mesoporous NYSC can successfully prevent polysulfide diffusion and stabilize dissolved polysulfides within cathodes.

Table 5.1 Comparison of the electrochemical performance of the as-prepared

NYSC@S with the reported.

Material Sulfur Specific Capacity Capacity Retention Reference

196 Content Nitrogen doped yolk shell 75.4% 909 mAh/g (0.2C) 68.3% (500 cycles) This work carbon sphere

Double shell hollow carbon 64% 690 mAh/g (0.1C) 68% (100 cycles) 241

sphere

Hollow-in-hollow carbon 70% 780 mAh/g (1 A/g) 72% (300 cycles) 233

sphere

Double-Layered Core−Shell 65% 900 mAh/g (0.2C) 84% (150 cycles) 227

carbon sphere

Hollow carbon sphere 70% 974 mAh/g (0.5C) 91% (100 cycles) 243

Hollow carbon sphere 50.2% 1357 mAh/g (0.05C) 86.6% (50 cycles) 244

Hollow carbon sphere 70% 902 mAh/g (1C) 90% (100 cycles) 242

Porous carbon sphere 70% 830 mAh/g (1C) 83% (100 cycles) 90

Porous carbon sphere 42% 650 mAh/g (400 mA/g ) 76%(500 cycles) 244

5.4 Conclusions

In summary, we have successfully synthesized highly ordered mesoporous nitrogen

doped yolk-shell carbon spheres using a sol-gel method. This nitrogen doped

yolk-shell carbon spheres possess a hierarchical mesoporous nanoarchitecture,

comprising an immobilizer yolk, a robust porous shell, and a cavity tunnel between

them. This unique architecture is beneficial for high sulfur loading and facilitates

complete redox reactions of active materials. The robust carbon shell of NYSC@S

can successfully entrap the polysulfides derived from the inner components and

greatly alleviate the “shuttle effect”. Furthermore, the dopant nitrogen atom in carbon

lattice of NYSC@S can modify the electron distribution and improve the affinity for

197 polysulfides, leading to the improvement in reversibility of Li2S/polysulfide/S conversion. Compared with the NCSC@S cathode, the NYSC@S cathode showed higher capacity, better rate capability, and superior stability. The NYSC@S cathode exhibits a high reversible capacity of 909 mAh/g after 500 cycles. The excellent cycling performance of the NYSC@S cathode is closely associated with its unique nitrogen doped yolk shell structure, such a cathode structure is promising for practical applications in high-performance Li−S batteries.

198 Chapter 6 Nitrogen-doped Hollow Co3O4 Nanoparticles Coated With Reduced

Graphene as High-Capacity Cathodes for Lithium-Sulfur Batteries

6.1 Introduction

Booming developments in electric vehicles, grid energy storage systems and portable electronic devices have accelerated the demand for more powerful, durable batteries with high energy density and long cycle life. Lithium-ion batteries, which dominate today’s portable electronic markets, have obstructed by their limited theoretical energy density. In this regard, Lithium-sulfur (Li-S) system is one of the promising candidates due to its high theoretical capacity (1675 mA h g-1) and energy density

(2567 Wh kg-1), which is at least 3 to 5 times that of conventional lithium-ion batteries.64, 72 Additionally, Sulfur is abundant, inexpensive, and environmentally benign, making Li–S batteries even more commercially attractive than LIBs. Despite their promising prospects, Li–S batteries are facing tough challenges such as the insulating nature of sulfur, large volume expansion/contraction (~80%) during cycling,

dissolution of polysulfides intermediates (Li2Sx, 4İxİ8) in organic electrolytes and the shuttle effect of polysulfides,61, 66, 67, 73, 145, 241-243, 256, 272, 273 among which restrict the shuttle effect of polysulfide is the main task for the advanced development of lithium-sulfur batteries.

It is generally acknowledged that an ideal sulfur host are expected to endow with highly porous interconnected architecture to accommodate sulfur, strong capability

199 for confining soluble polysulfide, sufficient electronic conductivity to ensure high sulfur utilization as well as robust framework to withstand the volume expansion of sulfur. From this perspective, infiltrating molten sulfur into well-designed porous conductive carbon materials like carbon nanotubes, graphene oxide and mixtures thereof seem to be meaningful in the beginning of cycling, 68, 85, 92, 106, 203, 223, 225, 232, 234,

238, 274-276 but rather weak intermolecular interactions between non-polar hydrophobic carbonaceous material and polar hydrophobic lithium polysulfide species is not sufficient to prevent polysulfide diffusion and shuttle effect over long-term cycling, resulting in serious capacity degradation 88 Therefore, polar host materials with well-designed porous structure, have been developed providing the fact that lithium polysulfides are intrinsically polar species with the terminal sulfur bearing most of the negative charge. In this regard, some workers by introduction of electronegative heteroatom atoms (N, S, P, O) into the carbon based materials, affect the nearby net polarity and creat sites for binding polusulfides.127, 155, 240, 277-280 In addition, hosts endow with ample Lewis acid site are proved to strongly interact with polysulfides and entrap them within or on surface of the host, considering that polysulfides are soft

Lewis base. 281 Such good examples are Metal-organic frameworks (MOFs), polar mental oxides, 137 metal carbide (MXene), which exhibit much better electrochemical performance due to the enhanced chemical interactions. 108, 133, 157, 277, 282

Recently, polar metal compounds derived from MOFs have attracted wide attention in the application of lithium sulfur batteries. As for reasons, the irreplaceable advantages 200 of these metal compounds are that both the abundant open metal centers serve as

Lewis acid sites and the heteroatom dopant sites (N, S and P) that derived from MOFs show strong affinity to polysulfide anions, which can form multiple strength chemical interactions with sulfur and lithium polysulfides.157 Meanwhile, MOFs assembled by connecting metal ions and organic linkers with tremendous extensiveness in multiplicity, 283 their tunable chemical composition can be rationally designed at the molecular level, making the appealing characteristics like controllable porous structure and high-surface area possible to be inherited by the follow up metalcompounds, 127, 264, 284, 285 thereby not only prevent the diffusion of polysulfides via both physical and chemical confinement, but also provide sufficient space to accommodate high sulfur loading and alleviate volumetric expansion of sulfur when applied as cathode host. 128, 282, 286-289

Herein, nitrogen-doped hollow Co3O4 nanoparticles coated with reduced graphene were prepared and used as the cathode host in lithium-sulfur battery. The as-prepared

rGO/N-C-Co3O4 nanoparticles were synthesized by a solid-state pyrolysis process that using metal organic framework (ZIF-67) as precursor, owing to its advantages of abundant Co–N moieties and unique dodecahedral morphology. In the as-prepared

RGO/N-C-Co3O4 framework structure, the well-designed porous structure coupling with RGO can effectively accommodate sulfur molecules and alleviate the volume expansion during sulfur lithiation. Moreover, both open metal centers serving as

Lewis acid sites and the nitrogen dopant sites in the rGO/N-C-Co3O4 nanoparticles 201 show strong affinity to polysulfide anions, which can form multiple strength chemical interactions with sulfur and lithium polysulfides. Ex situ Raman, Ex situ X-ray

Photoelectron Spectroscopy, UV-vis absorption spectra results and first-principle

calculations further confirmed that rGO/N-C-Co3O4 nanoparticles can effectively bind polysulfides in the electrode over cycles and exhibit strong binding energy. Therefore,

-1 the rGO/N-C-Co3O4@S cathodes generated a high reversible capacity (1205 mAh g at 0.2C) and excellent stability (865 mAh g-1 at 1C after 300 cycles).

6.2 Experimental Section

6.2.1 Material synthesis

Preparation of nanocrystal ZIF-67: 1.455 g of Co(NO3)2·6H2O was dissolved in the binary mixture of 50 ml methanol and 50 ml ethanol. 1.642 g of 2-methylimidazole,

0.5g sodium formate were dissolved in another mixture of 50 ml methanol and 50 ml ethanol, then 800ul of 1-Methylimidazole was added. The above two solutions were then mixed vigorously for 30 s, then the obtaining solution was incubated at room temperature for 12 h. The resulting purple precipitates were collected by centrifugation and washing with ethanol several times and finally vacuum-dried at

80 eC, thus ZIF-67 precursor was obtained.

Preparation of porous GO/N-C-Co3O4: ZIF-67 was pyrolyzed at 450 °C in flowing air

−1 with a ramp rate of 1 °C min to give N-C-Co3O4. The graphene oxide (GO) was synthesized through chemical exfoliation of graphite powders using a modified

202 290 Hummers’ method. Then, N-C-Co3O4 was dispersed into an aqueous solution of 0.5 wt% PDDA that contained 20 mM NaCl and 20 mM Tris solution, and mixed with the GO solution. After stirring for 3 hours, the product was washed with DI water several times and dried in a vacuum oven at room temperature.

Preparation of rGO/N-C-Co3O4@S: GO/N-C-Co3O4@S was homogeneously dispersed in DI water by ultrasonication, then transferred to an appropriate amount of

elemental sulfur/CS2/ethanol solution with a mass ratio of mCo3O4 / N-C-rGO : mS=1:4, the mixed solution was magnetically stirred for 24 h in an ice water bath. During this,

the CS2 and ethanol were allowed to completely evaporate while stirring, then the

rGO/N-C-Co3O4@S composite was filtered and dried at 50 ć in a vacuum oven for

12 h. Finally, the as-synthesized rGO/N-C-Co3O4@S was heated to 230 ć at a heating rate of 2 Ԩ min-1 for 2 h in a tube furnace under flowing argon at 80 sccm.

The GO nanosheets are transformed into reduced graphene oxide (rGO) during this

heating treatment, hence rGO/N-C-Co3O4@S was obtained.

6.2.2 Characterization of materials

203 Renishaw in-Via Raman spectrometer system (Gloucestershire, UK) equipped with a

Leica DMLB microscope (Wetzlar, Germany) and a 17 mW 633 nm Renishaw helium neon laser at 50% power. XPS analysis was performed on an ESCALAB MK

II X-ray photoelectron spectrometer in a JEOL JSM-6700F electron microscope with an accelerating voltage of 10 kV. FTIR spectra were recorded on a NEXUS 670FTIR spectrometer using KBr disks. Thermogravimetric analysis (TGA) was used for analysis of weight loss from precursors to the final products. Nitrogen adsorption−desorption measurements were conducted on a 3 Flex Surface

Characterization Analyzer to determine the Brunauer−Emmett−Teller (BET) specific surface areas using a Quadrasorb SI Analyzer at 77 K. The BET surface area was calculated using experimental points at a relative pressure of P/P0=0.05-0.25.

6.2.3 Electrochemical measurement

Working electrodes were made from 80 wt. % of active materials, 10 wt. % of conductive agent (carbon black) and 10 wt. % of binder (polyvinylidene difluoride).

The mass loading of sulfur on the electrode is around 2 mg/cm2. CR2032 coin cells were assembled in an argon-filled glove box (Mbraun, Unilab, Germany), in which both the moisture and oxygen contents were controlled to be less than 0.1 ppm. Al foil was used as cathode current collector. The electrolyte was 1 M lithium

bis-(trifluoromethanesulfonyl) imide (LiTFSI) and 1 wt% lithium nitrate (LiNO3) in

1,3-dioxolane and 1,2-dimethoxy-ethane (volume ratio 1:1), The electrodes were dried at 80 Ԩ under vacuum for 12 h. The electrolyte used for each coin cell is around 204 30 μL. Electrochemical measurements were conducted using a LAND-CT2001C battery test system. The cells were discharged and charged galvanostatically in the fixed voltage range 1.7−2.7 V with a current rate of 0.1C, 0.5C, 1C, 2C and 3C, respectively. Cyclic voltammetry (CV), and lectrochemical impedance spectroscopy

(EIS) were conducted with a CHI660C Electrochemistry Workstation in the voltage range 1.7 to 2.7 V (vs. Li+/Li).

6.2.4 Computational methods

The simulations are based on first-principles density functional theory (DFT), which is provided by the CASTEP package. The generalized gradient approximation (GGA) and Perdew-Burke-Ernzerhof scheme (PBE) was adopted for the exchange-correlation potential to optimize geometrical structures and calculate properties. Ultrasoft pseudopotentials and a plane-wave expansion of the wave functions were chosen for computations. The Brillouin zone is sampled by 2 × 1 × 1 k-points and the energy cutoff of 340 eV is chosen in the geometry optimization calculations. In order to take into account the contributions of the van der Waals

(vdW) interactions between different layers, the DFT-D (D stands for dispersion) approach within the Grimme scheme is adopted for the vdW correction.5 These setups are proven to be accurate enough for describing the results after careful test

calculations. A three-layer Co3O4 (110) surface model is constructed, in which the bottom layer of atoms are fixed, while the other two layers of atoms are fully relaxed.

The nearest distance between nanosheets in neighboring cells is greater than 18 Å to 205 ensure no interactions between different layers. For geometric optimization, the atomic positions of all structures are allowed to relax until the convergence tolerances of energy, maximum force, and displacement of 1 × 10−5 eV, 3 × 10−2 eV Å-1, and 1 ×

10−3 Å are reached, respectively. Periodic boundary conditions were adopted for all models utilized models in this work.

The adsorption energies (Eads), are defined as: Eads = Etotal - E species - Esubstrate, where

Etotal is the total energy of the adsorbed system, E species is the energy of the adsorbate

in vacuum and Esubstrate is the energy of the Co3O4 (110) surface. According to this definition, a more negative value indicates a more energetically favorable (exothermic)

reaction between the adsorbate and Co3O4 (110) surface.

6.3 Results and Discussion

Figure 6.1 shows the synthesis process of the rGO/N-C-Co3O4 polyhedra nanohybrids. An one step prolysis process was introduced to transform ZIF-67 into

N-C-Co3O4 hollow polyhedra. Then the negative charged graphene oxide (GO)

nanosheets can tightly coated on the surface of positive charge decorated N-C-Co3O4 polyhedra through electrostatic interactions. The as-prepared ZIF-67 and the

N-C-Co3O4 nanoparticles were characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). As shown in Figure 6.1, the as-made

206 Figure 6.2 Schematic illustration for the formation of rGO/N-C-Co3O4@S

ZIF-67 exhibits regular dodecahedral morphology and has two typical kinds of vertices in each dodecahedron which share three or four edges, aggregating to a

uniform particle size of approximate 2 μm. SEM images of N-C-Co3O4 in Figure 6.3a,

b clearly illustrate that N-C-Co3O4 nanoparticles preserve the polyhedral morphology of ZIF-67 after calcination. Actually, uniform dodecahedra also exhibit two typical types of vertex that share three or four edges, which can be clearly observed, and the

surface of the N-C-Co3O4 consist of numerous nano-sized particles and pores (Figure

6.3c), which is consistent with the TEM result in Figure 6.3b. The high-resolution

TEM image in Figure 6.3d further discloses that the N-C-Co3O4 sample is composed of closely packed nanoparticles with a size of about 15-20 nm and they are interconnected, forming a 3D porous structure. It also reveals that the interstices of

10-20 nm in size are observed among the adjacent nanoparticles, further confirming

207

Figure 6.2(a, b) SEM image of ZIF-67

that the N-C-Co3O4 is highly porous. The formation mechanism of the porous structure may be attributed to the release of gases from the decomposition and carbonization of organic component (2-methylimidazole). X-ray diffraction analysis

(Figure 6.4a) shows high crystallinity, indicating that Co3O4 (JCPDS card number

01-074-2120)is the primary crystalline phase of the as-obtained product. N-C-Co3O4 polyhedra were functionalized with the cationic polyelectrolyte PDDA to get a positive charged surface, then by mixing with GO solution, the positive charged

N-C-Co3O4 can be wrapped by the negative charged GO nanosheets vis electrostatic

interactions. After calcination rGO/N-C-Co3O4 nanohybrid is formed. Figure 6.5c,d

208

Figure 6.3 (a) SEM image of N-C-Co3O4, (b) TEM image of N-C-Co3O4, (c) FESEM image of N-C-Co3O4 composite. (d) HRTEM image of N-C-Co3O4.

shows the morphology of N-C-Co3O4 after rGO coating. It is apparent that rGO is

quite chiffon-like and closely wrapped on the surface of N-C-Co3O4 polyhedra. These tightly wrapped rGO nanosheets are supposed to act as a physical barrier to prevent

the diffusion of polysulfide anions from the porous rGO/N-C-Co3O4 polyhedra. To

further confirm elements in the rGO/N-C-Co3O4 nanoparticles, energy-dispersive

X-ray spectra (EDS) with elemental mapping images of C, N, Co and O are displayed in Figure 6.6. The EDS mapping results not only clearly reveal the homogenous distribution of Co3O4 nanoparticles, but also implies the effective incorporation of N atom into the carbon matrix. These results are also consistent with the FTIR results in

209 Figure 6.4b. It is obvious that all peaks of ZIF-67 disappear in the N-C-Co3O4 sample except for the absorption peak at 978 cm-1, which comes from the stretching

291 (vibration) and bending modes of C-N bonds. N2 adsorption-desorption analysis

was performed to further investigate the porous structure of the rGO/N-C-Co3O4

composite. The Brunauer-Emmett-Teller (BET) surface area of rGO/N-C-Co3O4 was measured to be 377 m2g-1 (Figure 6.4c), and the typical type I Langmuir isotherm indicates the wide range of pore size distribution from 0 to 50 nm,

Figure 6.4 (a) XRD patterns of N-C-Co3O4, rGO/N-C-Co3O4 @S and pure S. (b)

FTIR spectra of ZIF-67, N-C-Co3O4, rGO/N-C-Co3O4@S and pure S. (c) Nitrogen adsorption-desorption isotherm curves of rGO/N-C-Co3O4. (d) Thermogravimetric

-1 curves of rGO/N-C-Co3O4 @S in the N2 with a heating rate of 10 Ԩ min .

210

Figure 6.5 SEM images of (a, b) N-C-Co3O4, (c) rGO/N-C-Co3O4 and (d) rGO/N-C-Co3O4@S

Figure 6.6 (a)EDX mapping results of rGO/N-C-Co3O4, (b)Overall, (c) Carbon, (d)

Cobalt, (e) Nitrogen, (f) Oxygen. with an average pore size of 25nm, and a mesoporous and microporous volume of

0.49 and 0.09 cm3 g-1, respectively. This high surface area and abundant mesoporous

volume of porous rGO/N-C-Co3O4 are not only beneficial to accommodate high 211 amount of sulfur, but also provide sufficient surfaces area for effectively entrapping polysulfides via physical and chemical adsorptions during electrochemical cycles.

For this unique porous nanopolyhedral architecture, it is relatively difficult to monodisperse sulfur nanoparticles into the metal oxide by the traditional melting-diffusion method. In this paper, we adopt the wet-precipitation method to

(f)cobalt, (g) oxygen, (h) sulfur, (i) carbon and (j) nitrogen. achieve 75% sulfur content loading as identified by the TGA result in Figure 6.4d,

212 and the monodisperse sulfur was uniformly permeated into the mesopores of the

rGO/N-C-Co3O4 (denoted as rGO/N-C-Co3O4@S). SEM observations reveal that the

rGO/N-C-Co3O4@S composite maintains the original dodecahedral shape (Figure

6.7a, c), further implying that no extra sulfur exists outside the

rGO/N-C-Co3O4@Sstructure. TEM image (Figure 6.7b) shows that the inner space of

rGO/N-C-Co3O4@S is much darker after sulfur impregnation, and the crystalline

nanoparticles of Co3O4 on the shells cannot be easily identified, confirming that a

high content of sulfur is present inside the rGO/N-C-Co3O4@S host. In addition, energy-dispersive X-ray spectra (EDS) (Figure 6.7e-j) also clearly illustrate the

homogeneous distribution of sulfur in the rGO/N-C-Co3O4 matrix. From the X-ray

Figure 6.8 Nitrogen adsorption-desorption isotherm curves of rGO/N-C-Co3O4@S

diffraction pattern of the rGO/N-C-Co3O4 shown in Figure 6.4a, it can be seen that all the diffraction peaks are well indexed as sulfur structure, confirming that sulfur was successfully encapsulated into the porous structure. After sulfur is impregnated

into pores of rGO/N-C-Co3O4, the textural properties show a tremendous change as

213 shown in Figure 6.8. The specific surface area decreases to 14.3 m2 g-1, the mesoporous volume also decreases to 0.05 cm3 g-1, indicating that the majority pores are occupied by sulfur species. Obvious peak does not appear in the pore size distribution plot, confirming that the sulfur occupied the mesopores of rGO/

N-C-Co3O4. The surface chemical composition and the valence states of

rGO/N-C-Co3O4 @S were revealed by X-ray photoelectron spectroscopy (XPS). In the survey spectrum (Figure 6.9a), the characteristic peaks of Co, C, O, N, S can be observed, demonstrating the coexistence of these elements. From the high resolution

Figure 6.9 (a) XPS survey spectra of rGO/N-C-Co3O4@S

XPS spectrum of Co 2p (Figure 6.10a), the binding energy of Co 2p1/2 and Co 2p3/2 was detected to be approximately 795.6 and 781.5 e V, and the distance between the

two peaks was about 14.1 e V. Compared with pure Co3O4 reported in the literature, the slight shift in the binding energy could be attributed to the transfer of electrons between nitrogen doped carbon and cobalt oxide, which eventually leads to strong

277 electronic interaction between Co3O4 and the doped nitrogen bonded to carbon.

214 This seems to be further supported by the peak at 780.3 e V, which could be attributed to the incorporation of cobalt into a network consisting of a carbon matrix and the

292 nitrogen species Co-N. Two peaks at 796.9 eV and 781.0 eV for Co 2p1/2 and 2p3/2 respectively, confirm the existence of Co–S bonds while indicating the entrapping

effect of the rGO/N-C-Co3O4 framework. It has been reported that Co (II) centres

2– with free d-orbitals in rGO/N-C-Co3O4 are able to coordinate to nucleophilic Sx anion clusters.127 In addition, some authors also proposed that the Co may not only

Figure 6.10 XPS spectra of (a) Co 2p, (b) S 2p, (c) C 1s and (d) N 1s in the rGO

/N-C-Co3O4 @S composite contribute to convert polysulfide deposits back to soluble long-chain polysulfides, but

also catalyse long-chain polysulfides to Li2S2 and even to Li2S and thus enhance the reaction kinetics and result in a high specific capacity.292 The two small extra peaks at 215 2+ 293 804.2 and 787.9 e V are probably Co statellite peaks of Co3O4. In the X-ray photoelectron spectroscopy (XPS) analysis, the S 2p spectrum for the rGO/

N-C-Co3O4@S has 2p3/2 and 2p1/2 spin-orbit levels at 163.5 and 164.6 e V (Figure

6.10b), respectively. The peak located at 163.2 e V can be observed, which is possibly assigned to the chemical bond formation between cobalt and sulfur after melting with

S,292 this is consistent with the result in Figure 6.10a. The other peaks at 165.4, 168.3,

2− and 169.5 eV in the S 2p spectrum should be ascribed to the SO4 , which originates

from the surface oxy-groups on rGO/N-C-Co3O4. The fitted C 1s spectrum (Figure

6.10c) shows a primary peak ataround 284.8 eV, corresponding tosp2-hybridized graphitic carbon atoms, and peak at 287.1 e V is assigned to C-N bonds. All of these

results confirm the formation of C-N bonding configurations in the rGO/N-C-Co3O4.

The N-doped carbon surface is reported to serve as a conductive Lewis base “catalyst”

matrix to improve the adsorption ability of Li2Sn (n=4-8), which promotes the

oxidization of Li2S6, Li2S8, and S8 and thus improves the S utilization and cycle stability.259 The N 1s spectrum has four main characteristic peaks at binding energies of 398.17, 399.64, 400.23 and 402.12 eV in Figure 6.10d, which correspond to pyridinic N, Co-N, pyrrolic N, and graphitic N, respectively.294, 295 The presence of a

Co-N bond peak at 399.6 eV is consistent with the corresponding peak detected in the

Co 2p spectrum due to Co-N species. The synergistic Co-N plays a significant role in the improvement of electrochemical performance for lithium-sulfur batteries.292

216

Figure 6.11 (a) Cyclic voltammetry (CV) tested between 1.5 and 3 V at a sweep rate

217 rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S electrodes at 0.1C.

To evaluate structural effects of host materials on the electrochemical performance, both the sulfur mass loading and the amount of electrolyte added to the coin cells

were kept at the same level for rGO/N-C-Co3O4@S, Co3O4/S, and rGO@S composites. Cyclic voltammetry (CV) curves were tested to illustrate the electrochemical processes during discharging and charging performance. Figure

Figure 6.12 The discharge capacities of high and low voltage plateaus of (a)

N-C-Co3O4@S, (b) rGO@S. The onset voltage of the low plateau was defined around

2.0 V.

218 6.11a shows typical CV curves of rGO/N-C-Co3O4@S cathodes, including two reduction peaks and two oxidation peaks. Two cathodic peaks at 2.35 V and 2.05 V are consistent with the voltage plateaus in Figure 6.11b. The upper plateau at 2.35 V

contributes to the transformation of cyclo-S8 to long-chain soluble lithium polysulfides. The lower plateau at 2.05 V is associated with the decomposition of

those polysulfides to insoluble short-chain lithium sulfides (Li2S2, Li2S), which

contributes to the major capacity of the rGO/N-C-Co3O4@S electrodes. It can be seen that the capacity is not significantly decreased during the first 5 cycles, which suggests excellent electrochemical kinetics.

From the Nyquist plots (Figure 6.11c), it can be observed that the rGO/ N-C-Co3O4 electrode has the smallest semicircle in the high-frequency region, indicating the

lowest charge transfer resistance compared with Co3O4@S and rGO@S. Since all the test cathodes were given the same amount of sulfur loading, the different charge transfer resistances can be attributed to the conductivity of the host materials.

Benefiting from the metallic nature of rGO/N-C-Co3O4 and the good confinement of

sulfur, rGO/N-C-Co3O4 exhibited better conductivity and lower electrochemical resistance to facilitate the charge transfer for surface reactions than the other two electrodes. Figure 6.11d, Figure 6.12a-b compare the discharge capacity accumulated along the high and low discharge voltage plateau within 600 cycles,

respectively. As for the rGO/N-C-Co3O4 electrode, the low plateau capacity corresponding to the conversion from polysulfides to lithium sulphide experienced 219 slightly capacity degradation during the cycling, implying a marginally irreversible loss of active materials. Meanwhile, the high plateaus capacity attributed to the formation of polysulfides, was also well maintained since the first few round trip

Figure 6.13 The 100th and 600th cycle galvanostatic charge/discharge voltage profiles at 0.1 C. (a) rGO/N-C-Co3O4@S, (b) N-C-Co3O4@S, (c) rGO@S.

220 during cycling, suggesting that enormous polysulfides were effectively entrapped by

the polar sites in rGO/N-C-Co3O4, which can be verified by the overlapping voltage profiles of the 100th and 600th cycles (Figure 6.13). In contrast, without the polar sites for chemisorption of polysulfides, the pure rGO/S electrodes showed severe capacity degradation in the both voltage plateaus particular on the high plateau as shown in Figure 6.13, implying the poor confinement of polysulfides during cycling.

Figure 6.12e shows the cycling performance of these three electrodes at a 0.1 C rate, benefitting from highly-ordered hollow porous structure for sulfur accommodation.

rGO/N-C-Co3O4@S and N-C-Co3O4@S electrodes delivered a high initial capacity of

1275 mAh g-1, while rGO@S only generated about 870 mAh g-1. After 600 cycles, the capacity of rGO@S dramatically dropped to only 53 mAh g-1. In sharp contrast, with

strong affinity for lithium polysulfides, the rGO/N-C-Co3O4@S and N-C-Co3O4@S electrodes showed much better cycling stability with capacities of 845 and 632 mAh

-1 g respectively after 600 cycles, respectively. This implies that rGO/N-C-Co3O4@S electrodes make up for the shortages of carbon materials as they yield stronger chemical reactions with LiPSs (lithium polysulfides) and can also form a porous conductive structure for useful ion-shuttling while significantly reducing the influence of the unwanted LiPSs “shuttle effect”. At the same time, the rGO wrapping also

improves the electrical conductivity of Co3O4 and addresses the insulating nature of the sulfur cathode, resulting in an enhancement of cycling performance.

221 Figure 6.14 (a) The potential differences changes of charge and discharge plateaus for rGO/N-C-Co3O4@S, N-C-Co3O4@S and rGO@S cathodes between the charge discharge plateaus at various current densities. (b) Voltage profiles of rGO

222 efficiency.

Next, the rate capabilities and electrode kinetics of these cathode materials were evaluated at various current densities. The calculated potential differences between the charge/discharge voltage plateaus at current rates 0.1 C, 0.5 C, 1 C, 2 C, 3 C are

shown in Figure 6.14a. The result here implies that rGO/N-C-Co3O4@S possesses much less polarization and better reaction kinetics than other samples, which is attributed to its improved conductivity and efficiency for LiPSs adsorption. When the

current rate is increased from 0.1 C to 0.5, 1, 2 and 3C, the rGO/N-C-Co3O4@S electrodes delivered high rate capacities of 1301, 1199, 927, 756 and 652 mAh g-1,

respectively. The corresponding charge/discharge profiles of rGO/N-C-Co3O4@S at different rates are provided in Figure 6.14b. Even when the current density is increased to 3 C, it still preserves a similar discharge voltage plateau. In contrast, because of the ineffective LiPS confinement of bare carbon particles, the rGO@S electrode shows the worst rate capability, especially when the current rate is increased to 3 C, where upon the capacity is abruptly decreased to merely 130 mAh g-1. When the current rate was returned to 0.1 C, a specific capacity of 400 mAh g-1 was

recovered. The rate capability of N-C-Co3O4@S was slightly inferior to that of

rGO/N-C-Co3O4@S electrode due to poor conductivity (Figure 6.14c). In addition,

when cycling the rGO/N-C-Co3O4@S electrode at the current rate of 1C as shown in

Figure 6.14d, it still exhibits stable cycling performance. Its capacity was maintained at 865 mAh g-1 after 300 cycles. This could be attributed to the synergy between

223 N-C-Co3O4 and rGO that can fully satisfy the demands of both strong surface binding

and effective charge transfer between lithium polysulfides and rGO/N-C-Co3O4 electrode, resulting in enhanced electrochemical kinetics. It is generally acknowledged that the interfacial electrochemical kinetics are primarily dominated by the binding affinity and the efficient charge transfer between the polysulfides and the whole electrode. The strong interaction with polysulfides allows sufficient surface adsorption, the highly conductive electrode can facilitate the transport of electrons generated during the redox reactions of the lithium polysulfides.

Figure 6.15 (a) Dissembled electrode of rGO/S after 50 cycles. (b) Dissembled electrode of rGO/N-C-Co3O4@S after 50 cycles.

To investigate the enhancement mechanism for improving the cycling performance of

Li-S batteries, we dissembled the cells with rGO/N-C-Co3O4@S and rGO@S electrodes in the same discharge state after 50 cycles in a glovebox as shown in

Figure 6.15. Then both cathodes were immersed into a mixture of

1,3-dioxolane/1,2-dimethoxyethane (DOL/DME, 1:1 vol:vol) for 6 hours. It is clear

that the solution colour of rGO/N-C-Co3O4@S became slightly light yellow which was difficult to observe, suggesting only a negligible amount of soluble polysulfides 224

in the rGO/N-C-Co3O4@S electrode. In contrast, the solution of rGO@S electrode exhibits a distinct colour change from colourless to bright yellow, indicating the accumulation of soluble polysulfides in the electrolyte. Meanwhile, UV-vis

225 absorption spectra were also consistent with these results, as shown in Figure 6.16a.

2- The peak observed around 280 nm is attributed to S8 and S6 species, the small peak

2- 2- at 310 nm should be assigned to S6 or S4 species, while the noticeable peak located

2- 296-298 at 420 nm is associated with S4 species, indicating that there are serious detachment of polysulfides from the rGO@S cathode. However, the electrode with

Co3O4 is proved to effectively confine soluble polysulfides during cycling. Hence, the results from the visible experiment are highly consistent with the cycling performance

of the batteries and further imply the strong polysulfide affinity of rGO/N-C-Co3O4.

Meanwhile, the ex-situ Raman spectrum measurement of the rGO/N-C-Co3O4@S after 300 cycles (discharged to 2.1V, which is an appropriate voltage to generate the polysulfides) also consistent with the result of UV-spectrum. As shown in Figure

6.16b, D and G bonds derive from rGO while Eg and F2g, A1g bond corresponds to

-1 - Co3O4. It can be seen that the peaks at 274 and 549 cm can be assigned to S6 , while

-1 - the peaks located at around 154, 194, and 394 cm should be S8 . These results

- - readily confirm that the polysulfides (S8 , S6 ) were effectively entrapped by this

rGO/N-C-Co3O4 cathode material.

2– Normally, polysulfide anions (Sx ) are soft Lewis bases due to the sulfur lone electron pairs, the abundant open coordination metal sites in the MOF are soft Lewis acids and thus can readily coordinate to polysulfide anions.128, 133 Therefore, the

rGO/N-C-Co3O4@S which exhibits Lewis acid characteristics is able to chemically

226

Li2S4 with corresponding atom Mulliken charge. Blue, red, yellow, and purple balls represent Co, O, S, and Li atoms, respectively. interact strongly with LiPSs and thus entrap them within or on their surface. Figure

6.16c shows a Li 1s XPS spectrum of a rGO/N-C-Co3O4@S electrode after discharging to 2.1 V. It exhibits a single asymmetric peak at around 55.9 eV.

227 Therefore, an additional peak with a +0.5 eV shift was fitted and attributed to the Li in the polysulfides interacting with N (Li–N) because normally the Li 1s XPS spectrum only shows the symmetric peak around 55.5 eV. The introduction of electronegative N atoms into the carbon lattice, induces asymmetric charge distribution, which affects the net polarity, creating sites for binding LiPSs.221

Additionally, metal oxides are also able to bind LiPSs via polar–polar Li–O (57 eV) interaction depending on its exposed facets.142, 299

The enhancement of adsorption strength can be further understood by comparing the

adsorption energy between the Co3O4 nanocrystals and polysulfides. Density

adsorption energies (Figure 6.17, Figure 6.18). For simplicity, Li2S and Li2S4 were 228 employed as the representative polysulfides. The lowest-energy adsorption

configurations of Li2S and Li2S4 on Co3O4 are presented in Figure 6.9 along with

their Mulliken populations. It is clear that there is specific bonding between Li2S or

Li2S4 with Co3O4, resulting in much higher adsorption energies of -3.551 and -1.398 eV, respectively. This value is obviously lower than the adsorption energy (-0.94 eV)

between Li2S and graphene, and it is also lower than the adsorption energy (-0.49 eV)

259 between Li2S4 and graphene. These interaction exist mainly in the formation of

Li-O and Co-S bonds. For Li2S and Li2S4 with Co3O4 systems, the closest distances between Li and O are approximately 1.967 Å and 1.916 Å, respectively, implying strong ionic bonds between Li and O atoms. The strong Li-O bonding interaction is

the main cause of the strong binding energy between lithium polysulfides and Co3O4 surface. And the closest contact between Co and S is approxiamately 2.13~2.3 nm, implying weak ionic bonds and van der Waals attractions, which can also partially

enhance the Li2Sx attractions. Moreover, there is one Co-S and four Li-O pairs (per

Li2S) for the Li2S -Co3O4 system, two Co-S and four Li-O pairs (per Li2S4) for the

Li2S4 - Co3O4 system in contrast to no pairs for pure graphene. Furthermore, there are

1.38 e and 1.48 e transferred from Li2S and Li2S4 to a Co3O4 surface, respectively,

while only 0.227 and 0.019 electrons transfer from Li2S and Li2S4 to pristine graphene,

suggesting that more charge migrates into Co3O4 substrates. Therefore, stronger

attraction between Li2Sx - Co3O4 is expected, demonstrating that the presence of

Co3O4 nanocrystals enhance the retention of lithium polysulfides.

229 Therefore, the excellent electrochemical performance of MOF-derived

rGO/N-C-Co3O4@S composites could be attributed to the following aspects: (I) The abundant mesopores, micropores and sufficient hollow cavity in the MOF-derived polar hosts are favourable to accommodate sulfur and lithium polysulfides, rGO serves as a physical sulfur barrier, which contributes to the internal transport of Li+/e-, maximizing the utilization of sulfur. (II) The strong adsorption capability enriches the concentration of soluble polysulfides in the interior or on the surface of the conductive host, facilitating chemical transformation of polysulfide intermediates. (III)

rGO/N-C-Co3O4 nanocrystals probably participate in multiple polysulfide transformation, affecting the redox reaction environment favourably. This point inspired us to rationally introduce electrocatalytic-active nanocrystals to accelerate the transfer kinetics of sulfur redox reaction, achieving a high rate and long life Li-S battery.

6.4 Conclusions

In summary, we synthesized nitrogen-doped hollow Co3O4 nanoparticles coated with reduced graphene (rGO) by a facile solid-state pyrolysis methode that using metal organic framework (ZIF-67) as precursor. Comprehensive physical characterisations

confirmed that this nitrogen-doped hollow Co3O4 particles had a well-defined porous structure, excellent conductivity and unique chemisorptive nature, which is an ideal

cathode host for lithium sulfur battery. The porous rGO/N-C-Co3O4@S composite electrodes exhibited excellent cycling stability (865 mAh g-1 at 1C after 300 cycles), 230 good rate capability (756 and 652 mAh g-1 at 2 and 3 C rates, respectively). The excellent electrochemical performance can be attributed to the different types of chemical interactions (Lewis acid-based/polar-polar interaction) existed in the

rGO/N-C-Co3O4 host, which induced by the open metal centre and nitrogen doping.

This host can maximize the effectiveness of moderating lithium polysulfides diffusion and enhance the redox reaction kinetics of sulfur species at the same time. Ex situ

Raman, Ex situ X-ray Photoelectron Spectroscopy, UV-vis absorption spectra and

first-principle calculations further confirmed that rGO/N-C-Co3O4 nanoparticles can effectively bind polysulfides in the electrode over cycles and exhibit strong affinity, thus rendering good electronic contact for sustainable utilization of sulfur and favorable cycle stability.

231 Chapter 7 Co3O4-Carbon Cloth Free Standing Cathode for Lithium Sulfur

Battery

7.1 Introduction

We significantly stabilized cycle life of high sulfur loading binder-free cathode by

chemisorption of Co3O4 to carbon fiber cloth, which was used as a 3D current collector to accommodate a large amount of sulfur, MWCNF and CB hybrids within the conductive scaffold, enabling the fabrication of ultrahigh sulfur loaded electrodes.1,2 This special nanoarchitecture combines the advantage of strong chemisorption of lithium polysulfides as well as excellent electrical conductivity, enabling high sulfur utilization and effective trap of lithium polysulfides. When applied as cathode materials for lithium sulfur batteries, the cathodes exhibit a reversible capacity of 1007 mAh g-1 after 300 cycles.

7.2 Experimenl Section

Firstly, Co3O4 nanocages were constructed on carbon cloth (CC) by a solvothermal

method. In a typical procedure, 0.1 of poly(vinylpyrrolidone) (PVP, Mw = 360000 g mol-1, 99 %) was well dispersed in 7.5 ml of deionized water and 7.5 ml of ethanol

(95 %) by stirring treatment, then 0.5 mmol of Co(NO3)2•6H2O (> 98%) were dissolved into the above dispersion to form a light pink solution by continuous stirring for 30 min. CC which was pre-treated by an acid treatment were performed to functionalize CC with oxygen containing groups. Pristine CC were cut into circular

232 disks with a diameter of 1.4 cm and soaked in HNO3 solution at 60 °C for 3 h.

Acid-treated CC were then washed by water and dried at 60 °C for 24 h. Then transform the as-made solution and CC into a 25 ml Teflon-Lined stainless steel autoclave and reacted at 180 °C for 12 h. When the reaction finished, the solution was

cooled down to room temperature naturally and washed to obtain the black Co3O4-CC.

Then heated at 450 °C for 1 h under N2 atmosphere. Homogeneous sulfur-containing slurry was fabricated by mixing 90 wt% sulfur, 5 wt% carbon black, and 5 wt%

MWCNFs in N-methyl-2-pyrrolidone (NMP) followed by high power ultrasonication

for 0.5 h. The as-prepared Co3O4-CC was immersed into the slurry for 10 min, then

the Co3O4-CC was removed from the slurry and placed in a vacuum oven at 60 °C

overnight to obtain the Co3O4-CC-S electrode.

7.3 Results and Discussion

Figure 7.1 Schematic illustration for the formation of Co3O4-CC-S

By using our simple method, an activated carbon cloth was used as a template for the

preparation of Co3O4-CC-S composite (Figure 7.1), Co3O4 particles were

successfully coated on the surface of CFCs. The as-prepared Co3O4 characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM).

233 As shown in Figure 7.2b, the as-made Co3O4 exhibits regular nanocubic morphology

and has a uniform particle size of approximate 400 nm, also the surface of Co3O4 was smooth, which was consistent with the TEM result in Figure 7.2f, the surface of which shows high crystallinity without any detectable by-products as revealed by the

X-ray diffraction pattern (Figure 7.3). The structure of the Co3O4-CC-S was characterized by scanning electron microscopy (SEM) in Figure 7.2, it can be seen that sufficient void space, tens of micrometers in size, was generated by the randomly interconnected carbon fibers, which is capable of holding a large amount of active material, maintaining high electrolyte absorbability and effectively accommodating the volume expansion of sulfur during discharge.

The Co3O4-CC-S electrodes were prepared by immersing the Co3O4-CC in a premixed sulfur slurry containing 70 wt% commercial pure sulfur powder, 10 wt% carbon black(CB), and 20 wt% MWCNFs after high-power ultrasonication, it needs to be pointed out that MWCNFs were used in the sulfur slurry instead of conventional polymetric binders in integrated binder-free electrodes. From the X-ray diffraction

pattern of the Co3O4-CC-S shown in Figure 7.3b, it can be seen that all the diffraction peaks are well indexed as sulfur structure, indicating that sulfur was successfully encapsulated into the porous structure. This seems to be further supported by the

Raman spectroscopy analysis. In Figure 7.3c, we can detect distinct sulfur peaks at

234

Figure 7.2 (a, b)SEM image of Co3O4-CC. (c)Mapping results of Co3O4-CC-S.

(d)SEM image of Co3O4-CC-S. (e)TEM image of Co3O4. (f) SAED pattern for Co3O4.

-1 around 200 cm , confirming that sulfur effectively diffuse into the Co3O4-CC

composites. The 3D Co3O4-CC skelton is capable of holding a high, uniformly distributed concentration of the active material within its interconnected pores, enabling a high sulfur loading. We also characterized the structure of the

sulfur-MWCNF-CB clusters loaded on the Co3O4-CC, considering the distribution of sulfur is closedly associated with the electrochemical performance of the electrode. It was seen that the MWCNFs and CB are not simply surrounded around the sulfur

235

Figure 7.3 TGA curve(a), XRD patterns (b) and Raman spectra (c) of Co3O4-CC-S

236 particles, but homogenously dispersed through the sulfur and formed an interconnected and embedded conductive network, the corresponding elemental maps of carbon and sulfur revealing that sulfur is uniformly distributed within the carbon conductive network constructed by MWCNFs and CB. In this work, we demonstrate that an embedded conductive scaffold formed by MWCNFs and CB can be constructed through the insulated sulfur, which is highly desirable for high sulfur utilization during the electrochemical redox process.

Electrochemical measurements were conducted to compare the performance of

Co3O4-CC-S and CC/S. Both the sulfur mass loading and the amount of electrolyte

added in the coin cells were kept at the same level for Co3O4-CC-S and CC-S

composites. Figure 7.4a, b illustrate the charge/discharge profiles of Co3O4-CC-S and

CC/S in the first, second and 300th cycles, it is noteworthy that Co3O4-CC-S exhibits higher first-discharge plateau at around 2.37 V and wider second-discharge plateau at around 2.1 V compared with pure CNF/S electrodes, and even after 300 cycles, the discharge plateau still remain well, indicating more sulfur utilization in the

Co3O4-CC-S. The cycling performance and Coulombic efficiency of the Co3O4-CC-S composite and pure CC-S electrodes over long-term cycles are shown in Figure 7.4d.

-1 After 300 cycles, the reversible capacity of Co3O4-CC-S is around 1007 mAh g , the capacity retained 81% of the initial capacity after 300 cycles at 0.5C, while the stability of pure CNF/S electrode is significantly lower than that of Co3O4-CC-S, the capacity degraded to merely 434 mAh g-1 by 300th cycles. We also 237

0.1 C, 0.5 C, 1 C, 2 C, 3 C. (d) Long term cycling performance test of the

Co3O4-CC-S, CC-S electrodes at 0.2C discharge rate and corresponding Coulombic efficiency.

cycled the Co3O4-CC-S electrodes at different current rates. Figure 7.4c shows the

results of Co3O4-CC-S and CC-S electrode cycled at step-wise current rates. During

238 the discharge-charge process, the electrodes were consecutively cycled at 0.1C, 0.5C,

1C, 2C, 3C and then reversed back to low rates. It should be noted that when the

current reversed back to 0.1C, the Co3O4-CC-S electrode could retain a capacity of

950 mAh g-1, while the figure for CC-S electrodes was only 370 mAh g-1.

The interaction of different sulfur host materials (CC, Co3O4-CC) with lithium

polysulfides were probed by visual discrimination, taking Li2S6 as the representative

polysulfide. An equivalent amount of host materials were first added to Li2S6 solution.

The superior intrinsic capability of Co3O4-CC to absorb Li2S6 was clearly obvious, as

shown in Figure 7.5. The adding of Co3O4-CC rendered the Li2S6 solution light yellow after resting for 3h, implying strong adsorption while the carbon black solution did not exhibit obvious change, indicative of no interaction. Hence, the results from the visible experiment are highly consistent with the cycling performance of the

batteries and further confirm the strong capability of Co3O4-CC-S in confining the

LiPS species.

Figure 7.5 (a) Initially mixed. (b) Aging for 3h. ((1)L2S6 soultion. (2) L2S6 soultion immersed with CC. (3) L2S6 solution immersed with Co3O4-CC.

239 7.4 Conclusions

In summary, we designed a binder-free cathode by chemisorption of Co3O4 to carbon fiber cloth, and applied this as a 3D current collector to accommodate a large amount of sulfur, multiwall carbon nanofiber (MWCNF) and carbon black (CB) hybrids within the conductive scaffold. The electrode with a high sulfur loading exhibits low polarization, stable cycling performance, and excellent rate capability compared with pure CC-S cathodes. The excellent performance could be attributed to synergic effects, including the high surface area and high conductivity of the CC-S matrix, as well as

the strong polysulfides binding capability of Co3O4 particles. Moreover, Co3O4 nanocrystals probably participate in multiple polysulfide transformation, affecting the redox reaction environment favourably, leading to the improvement in the

reversibility of Li2S/polysulfides/S conversion. When applied as cathode material for lithium sulfur batteries, it exhibits superior cycling performance and excellent rate capability. This work opens a new opportunity for the realization of high-energy and commercially viable Li-S batteries.

240 Chapter 8 Conclusions and Future Perspective

8.1 Conclusions

The aim of this thesis was to explore the relationship between designed nanostructure materials and their electrochemical performance for Lithium-ion batteries and

Lithium-sulfur batteries, by means of theoretical calculation, structural, physical, and electrochemical characterization techniques. In this thesis, we have designed flower-like transition metal oxide as anode material for lithium-ion battery, and also developed carbon based materials and hybrid metal nanocomposites as cathode hosts for lithium sulfur batteries from the perspective of surface chemistry modification and nanostructure optimization. These as-prepared materials show several advantages as electrode materials for Lithium-ion battery and Lithium-sulfur battery, which benefit from the high surface area, improving the accessibility of electrolyte, shorting the diffusion length of ions (Li+), or the strong chemisorption for polysulfides, highly porous architecture to accommodate sulfur and alleviate the volume expansion during sulfur reduction. Following are the detailed summary of the outcomes.

Transition metal oxide as anode material for high performance lithium-ion battery

3D hierarchical porous China rose flower-like NiCo2O4 was successfully synthesized via a solvothermal method using PVP as the structure-directing agent followed by a

simple thermal annealing treatment. Each 3D flower-like NiCo2O4 sphere is

241 composed of nanoplates, on which numerous pores are distributed due to the thermal treatment from the precursor. Such a novel porous hierarchical architecture not only offers more active sites within the pores for fast electrochemical reaction, but also alleviates the volume expansion/contraction during lithiation and de-lithiation processes. When evaluated as an anode material for lithium-ion batteries, high reversible capacities of 915 mA h g-1 could be retained after 100 cycles at a high current density of 1000 mA g-1, corresponding to about 96% of the second discharge capacity (953 mA h g-1), and remarkable capacity retention was also obtained at increased current densities. The improved electrochemical performance enables the

as-prepared 3D hierarchical porous rose flower-like NiCo2O4 to be a promising anode material for next-generation, high-power lithium-ion batteries. Furthermore, the currently reported simple synthetic approach could be extended to other transition metal oxides, potentially yielding facile preparation and high yields of electrochemically outstanding products.

Graphene oxide as sulfur immobilizer for high performance lithium sulfur battery

We designed an efficient strategy for preparing crumpled N-S co-doped porous graphene (A-NSG) via chemical activation of polypyrrole (PPy) functionalized

graphene sheets with K2CO3. The A-NSG @S electrodes with a high sulfur loading of

72.4 wt% exhibits low polarization, stable cycling performance, and excellent rate

242 capability, compared with rGO@S, NSG@S and A-G@S electrodes. The excellent performance could attributed to synergic effects, including the high surface area and high conductivity of the porous matrix, as well as the unique lithium polysulfides binding capability of the N, S functional groups in the A-NSG sheets. Furthermore, pyridinic nitrogen and thiophenic sulfur in carbon lattice of A-NSG@S can modify the electron distribution and improve the affinity for Li2S as well as Li2S4, leading to the improvement in reversibility of Li2S/polysulfide/S conversion. This work not only presents a simple method to synthesize N-S co-doped hierarchical porous graphene, but also demonstrates the excellent performance of A-NSG cathode materials for lithium-sulfur batteries.

Carbon spheres with yolk-shell architecture for high performance lithium-sulfur batteries

We have successfully synthesized highly ordered mesoporous nitrogen doped yolk-shell carbon spheres using a sol-gel method. This nitrogen doped yolk-shell carbon spheres possess a hierarchical mesoporous nanoarchitecture, comprising an immobilizer yolk, a robust porous shell, and a cavity tunnel between them. This unique architecture is beneficial for high sulfur loading and facilitates complete redox reactions of active materials. The robust carbon shell of NYSC@S can successfully entrap the polysulfides derived from the inner components and greatly alleviate the

“shuttle effect”. Furthermore, the dopant nitrogen atom in carbon lattice of NYSC@S

243 can modify the electron distribution and improve the affinity for polysulfides, leading

to the improvement in reversibility of Li2S/polysulfide/S conversion. Compared with the NCSC@S cathode, the NYSC@S cathode showed higher capacity, better rate capability, and superior stability. The NYSC@S cathode exhibits a high reversible capacity of 909 mAh/g after 500 cycles. The excellent cycling performance of the

NYSC@S cathode is closely associated with its unique nitrogen doped yolk shell structure, such a cathode structure is promising for practical applications in high-performance Li−S batteries

Metal oxide as cathode material for high performance lithium-sulfur batteries

We synthesized nitrogen-doped hollow Co3O4 nanoparticles coated with reduced graphene (rGO) by a facile solid-state pyrolysis methode that using metal organic framework (ZIF-67) as precursor. Comprehensive physical characterisations

confirmed that this nitrogen-doped hollow Co3O4 particles had a well-defined porous structure, excellent conductivity and unique chemisorptive nature, which is an ideal

cathode host for lithium sulfur battery. The porous rGO/N-C-Co3O4@S composite electrodes exhibited excellent cycling stability (865 mAh g-1 at 1C after 300 cycles), good rate capability (756 and 652 mAh g-1 at 2 and 3 C rates, respectively). The excellent electrochemical performance can be attributed to the different types of chemical interactions (Lewis acid-based/polar-polar interaction) existed in the

rGO/N-C-Co3O4 host, which induced by the open metal centre and nitrogen doping.

244 This host can maximize the effectiveness of moderating lithium polysulfides diffusion and enhance the redox reaction kinetics of sulfur species at the same time. Ex situ

Raman, Ex situ X-ray Photoelectron Spectroscopy, UV-vis absorption spectra and

8.2 Future Perspective

The main method and design strategy of synthesising nanostructure electrode material presented here can be extended to the preparation of other analogue cathode or anode materials for both Lithium-ion batteries and Lithium-sulfur batteries. At the same time, many aspects which might affect the electrochemical properties of asprepared active materials were not evaluated due to the limitation of doctoral duration.

Therefore, further research is needed to deeply characterise the relationship between the materials structure and the electrochemcial performance, particularly in relation to factors related to the improvement of capacity (the majority electrode materials cannot achieve their theoretical capacity) and cycliablity.

Moreover, as discussed in this thesis, Lithium-sulfur batteries are emerging in importance due to high energy density and high theoretical capacity at low cost when compared with its counterpart Lithium-ion batteries, of course, some disadvantages

245 relating to polysulfides shuttle effect limit the practical development of

Lithium-sulfur batteries.Consequently, finding and optimising suitable electrode materials for the application of Lithium-sulfur batteries is a meaningful pursuit. Based on the understanding of the mechanisms of lithium sulfur battery, the future perspective of nanostructured cathode materials for lithium-sulfur batteries should focus on these key factors: (I) excellent electronic conductivity of the sulfur host, (II) strong chemisorption to polysulfides (III) highly porous or hollow architecture to alleviate the volume expansion of sulfur. (IV) the possible catalytic property for accelerating the mutual conversions of lithium polysulfides and lithium sulfides.

Among these factors, both the first-principle theory calculation and the visual polysulfide absorption experiments in this thesis implies that cathodes with polar surface showing strong affinity for polysulfides, this factor greatly contributes to the improvement of cycling stability. Especially, hosts endow with Lewis acid site can strongly interact with polysulfides and entrap them within or on surface of the host, this finding throws new light on our future research. Meanwhile, the non-stoichiometric metal oxides/sulfides with catalystic property which reported to contribute to mutual polysulfide conversion are also appealing for lithium sulfur batteries. In addition, to achieve long term cycling stability and high capacity particular in the high sulfur loading, it is important to value the electronic conductivity of the hosts as important as the surface chemistry. Based on this condition, MXene and metal nitrides combining the advantages of strong Lewis acid–

246 base bonding and high electronic conductivity with robustness structure seem to be promising for high performance lithium sulfur batteries.

On the other hand, it is generally acknowledged that the areal sulfur loadings of lithium-sulfur batteries need to be higher than 7 mg cm-2 for the sake of practical applications. The general levels demonstrated by laboratory experiments are around 2 mg cm-2, thus significant efforts need to be devoted into explore the high areal sulfur loading, long life cathode that can sustain low electrolyte volumes especially under high current rate. From this perspective, establishing ultra-light conductive networks by some flexible materials seem to be a positive trend. Moreover, the development of electrolyte is also meaningful for the design of high areal sulfur loading cathodes, which can effectively alleviate the polysulfide dissolution and suppress the shuttling effect. One solution to date is to employ high donor number solvent to maximal extent solubilize the polysulfide in a minimum volume of electrolyte. On the other hand, sulfur redox based flow batteries are proposed, in this system non-solvent ionic liquids are adopted as alternative solution forcing redox chemistry to take place at the interface. In addition, much attention should be also paid to mediate the redox

conversion from Li2S2 to Li2S as this process negatively affects the extended cycling performance by blocking access to the electrode surface. Considering the merits and drawbacks of lithium sulfur batteries, much progresses are expected to be achieved in the near future, tomorrow looks positive!

247 APPENDIX: NOMENCLATURE

Abbreviations/Symbols Full name a.u. Arbitrary unit

Ar Argon

BET Brunauer-Emmett-Teller

BJH Barrett-Joyner-Halenda

CB Carbon black

CNT carbon nanotube

CV Cyclic Voltammetry

C-rate Current rate

DI de-ionized

DMC Dimethyl carbonate

EC Ethylene carbonate

EIS Electrochemical Impedance Spectroscopy

EVs Electric vehicles

FESEM Field-Emission Scanning Electron Microscopy

FTIR Fourier transform infrared spectroscopy g Gram

GO Graphene oxide h Hour

Hz Hertz

248

I Intensity

HEVs Hybrid electric vehicles

HRTEM High-resolution transmission electronic spectroscopy

JCPDS Joint Committee on Powder Diffraction Standards

Li Lithium

LIBs Lithium-ion Batteries

M Molar concentration mA h g-1 Milli ampere hour per gram min Minute mm Millimeter nm Nanometer

NMP 1-methyl-2-pyrrolidinone

PC Propylene carbonate

PVDF Poly(vinylidene difluoride)

Rct Charge transfer resistance

RΩ Ohmic resistance

SAED Selected area electron diffraction

SEI Solid Electrolyte Interface

SEM Scanning electron Microscopy

SIBs Sodium-ion Batteries

249 TEM Transmission electron microscopy

TGA Thermogravimetric analysis

XRD X-ray diffraction

° Degree

Ω Ohm

°C Degree Celsius

Zw Warburg impedance

250 REFERENCES

1. S. Chen, Z. Ao, B. Sun, X. Xie and G. Wang, Energy Storage Materials, 2016,

5, 180-190.

2. S. Chen, P. Bao, X. Huang, B. Sun and G. Wang, Nano Research, 2013, 7,

85-94.

3. S. Chen, P. Bao and G. Wang, Nano Energy, 2013, 2, 425-434.

4. S. Chen, P. Bao, L. Xiao and G. Wang, Carbon, 2013, 64, 158-169.

5. S. Chen, P. Chen, M. Wu, D. Pan and Y. Wang, Electrochemistry

Communications, 2010, 12, 1302-1306.

6. S. Chen, X. Huang, H. Liu, B. Sun, W. Yeoh, K. Li, J. Zhang and G. Wang,

Advanced Energy Materials, 2014, 4, 190-197.

7. S. Chen, B. Sun, X. Xie, A. K. Mondal, X. Huang and G. Wang, Nano Energy,

2015, 16, 268-280.

8. S. Chen, Y. Wang, H. Ahn and G. Wang, Journal of Power Sources, 2012, 216,

22-27.

9. S. Chen, W. Yeoh, Q. Liu and G. Wang, Carbon, 2012, 50, 4557-4565.

10. S. Chen, Y. Zhao, B. Sun, Z. Ao, X. Xie, Y. Wei and G. Wang, ACS applied

materials & interfaces, 2015, 7, 3306-3313.

11. Y. Wei, S. Chen, D. Su, B. Sun, J. Zhu and G. Wang, Journal of Materials

Chemistry A, 2014, 2, 8103.

12. G. Cui, L. Gu, L. Zhi, N. Kaskhedikar, P. A. van Aken, K. Müllen and J. Maier,

251 Advanced materials, 2008, 20, 3079-3083.

13. H. Wang, M. Yoshio, T. Abe and Z. Ogumi, Journal of The Electrochemical

Society, 2002, 149, A499.

14. H. Nakagawa, Y. Domi, T. Doi, M. Ochida, S. Tsubouchi, T. Yamanaka, T. Abe

and Z. Ogumi, Journal of Power Sources, 2013, 236, 138-144.

15. H. Zheng, Q. Qu, L. Zhang, G. Liu and V. S. Battaglia, RSC Advances, 2012, 2,

4904.

16. B. J. Landi, M. J. Ganter, C. D. Cress, R. A. DiLeo and R. P. Raffaelle, Energy

& Environmental Science, 2009, 2, 638.

17. J. Hou, Y. Shao, M. W. Ellis, R. B. Moore and B. Yi, Physical chemistry

chemical physics : PCCP, 2011, 13, 15384-15402.

18. J. Sun, H. Liu, X. Chen, D. G. Evans, W. Yang and X. Duan, Advanced

materials, 2013, 25, 1125-1130, 1124.

19. X. Li and C. Wang, J. Mater. Chem. A, 2013, 1, 165-182.

20. G.-N. Zhu, L. Chen, Y.-G. Wang, C.-X. Wang, R.-C. Che and Y.-Y. X i a,

Advanced Functional Materials, 2013, 23, 640-647.

21. A. S. Prakash, P. Manikandan, K. Ramesha, M. Sathiya, J. M. Tarascon and A.

K. Shukla, Chemistry of Materials, 2010, 22, 2857-2863.

22. L. Shen, E. Uchaker, X. Zhang and G. Cao, Advanced materials, 2012, 24,

6502-6506.

23. C. M. Park, J. H. Kim, H. Kim and H. J. Sohn, Chem Soc Rev, 2010, 39,

252 3115-3141.

24. W.-J. Zhang, Journal of Power Sources, 2011, 196, 13-24.

25. U. Kasavajjula, C. Wang and A. J. Appleby, Journal of Power Sources, 2007,

163, 1003-1039.

26. H. Zhang and P. V. Braun, Nano letters, 2012, 12, 2778-2783.

27. H. Wu and Y. Cui, Nano Today, 2012, 7, 414-429.

28. J. R. Szczech and S. Jin, Energy Environ. Sci., 2011, 4, 56-72.

29. M. T. McDowell, S. W. Lee, J. T. Harris, B. A. Korgel, C. Wang, W. D. Nix

and Y. Cui, Nano letters, 2013, 13, 758-764.

30. C. K. Chan, H. Peng, G. Liu, K. McIlwrath, X. F. Zhang, R. A. Huggins and Y.

Cui, Nature nanotechnology, 2008, 3, 31-35.

31. S. W. Hwang, J. K. Lee and W. Y. Yoon, Journal of Power Sources, 2013, 244,

620-624.

32. J.-I. Lee and S. Park, Nano Energy, 2013, 2, 146-152.

33. K. Song, S. Yoo, K. Kang, H. Heo, Y.-M. Kang and M.-H. Jo, Journal of

Power Sources, 2013, 229, 229-233.

34. N. G. Rudawski, B. R. Yates, M. R. Holzworth, K. S. Jones, R. G. Elliman and

A. A. Volinsky, Journal of Power Sources, 2013, 223, 336-340.

35. A. M. Chockla, K. C. Klavetter, C. B. Mullins and B. A. Korgel, ACS applied

materials & interfaces, 2012, 4, 4658-4664.

36. X. H. Liu, S. Huang, S. T. Picraux, J. Li, T. Zhu and J. Y. Huang, Nano letters,

253 2011, 11, 3991-3997.

37. C. C. Li, Q. H. Li, L. B. Chen and T. H. Wang, Journal of Materials Chemistry,

2011, 21, 11867.

38. B. Koo, H. Xiong, M. D. Slater, V. B. Prakapenka, M. Balasubramanian, P.

Podsiadlo, C. S. Johnson, T. Rajh and E. V. Shevchenko, Nano letters, 2012,

12, 2429-2435.

39. I. T. Kim, A. Magasinski, K. Jacob, G. Yushin and R. Tannenbaum, Carbon,

2013, 52, 56-64.

40. N. Kang, J. H. Park, J. Choi, J. Jin, J. Chun, I. G. Jung, J. Jeong, J. G. Park, S.

M. Lee, H. J. Kim and S. U. Son, Angewandte Chemie, 2012, 51, 6626-6630.

41. L. Zhang, P. Hu, X. Zhao, R. Tian, R. Zou and D. Xia, Journal of Materials

Chemistry, 2011, 21, 18279.

42. H. Guan, X. Wang, H. Li, C. Zhi, T. Zhai, Y. Bando and D. Golberg, Chemical

communications, 2012, 48, 4878-4880.

43. A. Paolella, R. Brescia, M. Prato, M. Povia, S. Marras, L. De Trizio, A. Falqui,

L. Manna and C. George, ACS applied materials & interfaces, 2013, 5,

2745-2751.

44. H. Liu, X. Du, X. Xing, G. Wang and S. Z. Qiao, Chemical communications,

2012, 48, 865-867.

45. C. Villevieille, F. Robert, P. L. Taberna, L. Bazin, P. Simon and L. Monconduit,

Journal of Materials Chemistry, 2008, 18, 5956.

254 46. M. C. Stan, R. Klöpsch, A. Bhaskar, J. Li, S. Passerini and M. Winter,

Advanced Energy Materials, 2013, 3, 231-238.

47. M. G. Kim and J. Cho, Advanced Functional Materials, 2009, 19, 1497-1514.

48. A. Hayashi, A. Inoue and M. Tatsumisago, Journal of Power Sources, 2009,

189, 669-671.

49. K. Aso, A. Hayashi and M. Tatsumisago, Electrochimica Acta, 2012, 83,

448-453.

50. J. Cabana, C. M. Ionica-Bousquet, C. P. Grey and M. R. Palacín,

Electrochemistry Communications, 2010, 12, 315-318.

51. Y. Wang, J. Wu, Y. Tang, X. Lu, C. Yang, M. Qin, F. Huang, X. Li and X.

Zhang, ACS applied materials & interfaces, 2012, 4, 4246-4250.

52. A. Paolella, C. George, M. Povia, Y. Zhang, R. Krahne, M. Gich, A. Genovese,

A. Falqui, M. Longobardi, P. Guardia, T. Pellegrino and L. Manna, Chemistry

of Materials, 2011, 23, 3762-3768.

53. N. H. Idris, M. M. Rahman, S.-L. Chou, J.-Z. Wang, D. Wexler and H.-K. Liu,

Electrochimica Acta, 2011, 58, 456-462.

54. J. Arai, T. Yamaki, S. Yamauchi, T. Yuasa, T. Maeshima, T. Sakai, M. Koseki

and T. Horiba, Journal of Power Sources, 2005, 146, 788-792.

55. S.-T. Myung, S. Komaba and N. Kumagai, Journal of The Electrochemical

Society, 2001, 148, A482.

56. S.-T. Myung, K. Hosoya, S. Komaba, H. Yashiro, Y.-K. Sun and N. Kumagai,

255 Electrochimica Acta, 2006, 51, 5912-5919.

57. K.-S. Lee, S.-T. Myung, K. Amine, H. Yashiro and Y.-K. Sun, Journal of

Materials Chemistry, 2009, 19, 1995.

58. S.-M. Bak, K.-W. Nam, C.-W. Lee, K.-H. Kim, H.-C. Jung, X.-Q. Yang and

K.-B. Kim, Journal of Materials Chemistry, 2011, 21, 17309.

59. K. Shaju, G. S. Rao and B. Chowdari, Electrochim. Acta, 2002, 48, 145-151.

60. H.-J. Peng, J.-Q. Huang, X.-B. Cheng and Q. Zhang, Advanced Energy

Materials, 2017, 2, 1700260.

61. Q. Pang, X. Liang, C. Y. Kwok and L. F. Nazar, Nature Energy, 2016, 1,

16132.

62. A. Rosenman, E. Markevich, G. Salitra, D. Aurbach, A. Garsuch and F. F.

Chesneau, Advanced Energy Materials, 2015, 5, 1500212.

63. M. Hagen, D. Hanselmann, K. Ahlbrecht, R. Maça, D. Gerber and J. Tübke,

Advanced Energy Materials, 2015, 5, 1401986.

64. A. Manthiram, S. H. Chung and C. Zu, Advanced materials, 2015, 27,

1980-2006.

65. M. Wild, L. O'Neill, T. Zhang, R. Purkayastha, G. Minton, M. Marinescu and

G. J. Offer, Energy Environ. Sci., 2015, 8, 3477-3494.

66. H. J. Peng, Z. W. Zhang, J. Q. Huang, G. Zhang, J. Xie, W. T. Xu, J. L. Shi, X.

Chen, X. B. Cheng and Q. Zhang, Advanced materials, 2016, 4, 1603401.

67. G. Li, J. Sun, W. Hou, S. Jiang, Y. Huang and J. Geng, Nature communications,

256 2016, 7, 10601.

68. R. Fang, S. Zhao, P. Hou, M. Cheng, S. Wang, H. M. Cheng, C. Liu and F. Li,

Advanced materials, 2016, 28, 3374-3382.

69. M. Yu, J. Ma, H. Song, A. Wang, F. Tian, Y. Wang, H. Qiu and R. Wang,

Energy Environ. Sci., 2016, 9, 1495-1503.

70. P. G. Bruce, S. A. Freunberger, L. J. Hardwick and J. M. Tarascon, Nature

materials, 2012, 11, 19-29.

71. J. Gao and H. D. Abruña, The Journal of Physical Chemistry Letters, 2014, 5,

882-885.

72. J.-G. Wang, K. Xie and B. Wei, Nano Energy, 2015, 15, 413-444.

73. Q. Zhang, Y. Wang, Z. W. Seh, Z. Fu, R. Zhang and Y. Cui, Nano letters, 2015,

15, 3780-3786.

74. S. Li, T. Mou, G. Ren, J. Warzywoda, B. Wang and Z. Fan, ACS Energy

Letters, 2016, 5, 0182.

75. H. J. Peng and Q. Zhang, Angewandte Chemie, 2015, 54, 11018-11020.

76. D. Su, K. Kretschmer and G. Wang, Advanced Energy Materials, 2016, 6,

1501785.

77. S. Urbonaite, T. Poux and P. Novák, Advanced Energy Materials, 2015, 5,

1500118.

78. D. Zheng, X. Zhang, J. Wang, D. Qu, X. Yang and D. Qu, Journal of Power

Sources, 2016, 301, 312-316.

257 79. Y. Xiang, J. Li, J. Lei, D. Liu, Z. Xie, D. Qu, K. Li, T. Deng and H. Tang,

ChemSusChem, 2016, 9, 3023-3039.

80. R. Fang, S. Zhao, Z. Sun, D. W. Wang, H. M. Cheng and F. Li, Advanced

materials, 2017, 34, 1606823.

81. T. Z. Hou, W. T. Xu, X. Chen, H. J. Peng, J. Q. Huang and Q. Zhang,

Angewandte Chemie, 2017, 23, 1704324.

82. X. Liang, C. Y. Kwok, F. Lodi-Marzano, Q. Pang, M. Cuisinier, H. Huang, C.

J. Hart, D. Houtarde, K. Kaup, H. Sommer, T. Brezesinski, J. Janek and L. F.

Nazar, Advanced Energy Materials, 2016, 6, 1501636.

83. Y. S. Su, Y. Fu, T. Cochell and A. Manthiram, Nature communications, 2013, 4,

2985.

84. X. B. Cheng, R. Zhang, C. Z. Zhao and Q. Zhang, Chem Rev, 2017, 23, 115.

85. R. Cao, W. Xu, D. Lv, J. Xiao and J.-G. Zhang, Advanced Energy Materials,

2015, 5, 145-154.

86. Z. Li, J. T. Zhang, Y. M. Chen, J. Li and X. W. Lou, Nature communications,

2015, 6, 8850.

87. X. Ji, K. T. Lee and L. F. Nazar, Nature materials, 2009, 8, 500-506.

88. Z. Li, J. Zhang, B. Guan, D. Wang, L. M. Liu and X. W. Lou, Nature

communications, 2016, 7, 13065.

89. T. Lei, W. Chen, J. Huang, C. Yan, H. Sun, C. Wang, W. Zhang, Y. Li and J.

Xiong, Advanced Energy Materials, 2017, 7, 1601843.

258 90. D. R. Deng, F. Xue, Y. J. Jia, J. C. Ye, C. D. Bai, M. S. Zheng and Q. F. Dong,

ACS nano, 2017, 11, 6031-6039.

91. D. Su, M. Cortie, H. Fan and G. Wang, Advanced materials, 2017, 56,

1700587.

92. X. Liang, A. Garsuch and L. F. Nazar, Angewandte Chemie, 2015, 54,

3907-3911.

93. Q. Pang and L. F. Nazar, ACS nano, 2016, 10, 4111-4118.

94. H. Chen, Q. Zou, Z. Liang, H. Liu, Q. Li and Y. C. Lu, Nature

communications, 2015, 6, 5877.

95. H. Kim, J. Lee, H. Ahn, O. Kim and M. J. Park, Nature communications, 2015,

6, 7278.

96. G. Hu, C. Xu, Z. Sun, S. Wang, H. M. Cheng, F. Li and W. Ren, Advanced

materials, 2015, 68,1504765.

97. Z. Li, H. B. Wu and X. W. Lou, Energy Environ. Sci., 2016, 9, 3061-3070.

98. W. Li, H. Yao, K. Yan, G. Zheng, Z. Liang, Y. M. Chiang and Y. Cui, Nature

communications, 2015, 6, 7436.

99. B. Zhang, X. Qin, G. R. Li and X. P. Gao, Energy & Environmental Science,

2010, 3, 1531.

100. J. Wang, J. Yang, C. Wan, K. Du, J. Xie and N. Xu, Advanced Functional

Materials, 2003, 13, 487-492.

101. S. S. Zhang, Frontiers in Energy Research, 2013, 1, 123-130.

259 102. S. Xin, L. Gu, N. H. Zhao, Y. X. Yin, L. J. Zhou, Y. G. Guo and L. J. Wan,

Journal of the American Chemical Society, 2012, 134, 18510-18513.

103. F. Zhan, B. Geng and Y. Guo, Chem.–Eur. J., 2009, 15, 6169–6174.

104. Y. Yang, G. Zheng, S. Misra, J. Nelson, M. F. Toney and Y. Cui, Journal of the

American Chemical Society, 2012, 134, 15387-15394.

105. Y. Son, J.-S. Lee, Y. Son, J.-H. Jang and J. Cho, Advanced Energy Materials,

2015, 5, 1500110.

106. J. He, Y. Chen, W. Lv, K. Wen, C. Xu, W. Zhang, W. Qin and W. He, ACS

Energy Letters, 2016, 272, 820-826.

107. F. Han, J. Yue, X. Fan, T. Gao, C. Luo, Z. Ma, L. Suo and C. Wang, Nano

letters, 2016, 16, 4521-4527.

108. J. He, Y. Chen, W. Lv, K. Wen, C. Xu, W. Zhang, Y. Li, W. Qin and W. He,

ACS nano, 2016, 23,5696.

109. F. Wu, A. Magasinski and G. Yushin, J. Mater. Chem. A, 2014, 2, 6064-6070.

110. C. Zu, M. Klein and A. Manthiram, The Journal of Physical Chemistry Letters,

2014, 5, 3986-3991.

111. D. Sun, Y. Hwa, Y. Shen, Y. Huang and E. J. Cairns, Nano Energy, 2016, 26,

524-532.

112. F. Wu, J. T. Lee, E. Zhao, B. Zhang and G. Yushin, ACS nano, 2016, 10,

1333-1340.

113. C. Nan, Z. Lin, H. Liao, M. K. Song, Y. Li and E. J. Cairns, Journal of the

260 American Chemical Society, 2014, 136, 4659-4663.

114. H. Noh, J. Song, J.-K. Park and H.-T. Kim, Journal of Power Sources, 2015,

293, 329-335.

115. K. Cai, M. K. Song, E. J. Cairns and Y. Zhang, Nano letters, 2012, 12,

6474-6479.

116. Y. Yang, M. T. McDowell, A. Jackson, J. J. Cha, S. S. Hong and Y. Cui, Nano

letters, 2010, 10, 1486-1491.

117. G. Tan, R. Xu, Z. Xing, Y. Yuan, J. Lu, J. Wen, C. Liu, L. Ma, C. Zhan, Q. Liu,

T. Wu, Z. Jian, R. Shahbazian-Yassar, Y. Ren, D. J. Miller, L. A. Curtiss, X. Ji

and K. Amine, Nature Energy, 2017, 2, 17090.

118. M. J. Klein, A. Dolocan, C. Zu and A. Manthiram, Advanced Energy Materials,

2017, 23, 1701122.

119. J. Conder, R. Bouchet, S. Trabesinger, C. Marino, L. Gubler and C. Villevieille,

Nature Energy, 2017, 2, 17069.

120. G. Zhou, E. Paek, G. S. Hwang and A. Manthiram, Nature communications,

2015, 6, 7760.

121. S. Liatard, K. Benhamouda, A. Fournier, R. Ramos, C. Barchasz and J. Dijon,

Chemical communications, 2015, 51, 7749-7752.

122. Y. Fu, Y. S. Su and A. Manthiram, Angewandte Chemie, 2013, 52, 6930-6935.

123. R. Fang, S. Zhao, S. Pei, Y. Cheng, P. Hou, M. Liu, H.-M. Cheng, C. Liu and F.

Li, Carbon, 2016, 109, 719-726.

261 124. H. Yao, G. Zheng, P. C. Hsu, D. Kong, J. J. Cha, W. Li, Z. W. Seh, M. T.

McDowell, K. Yan, Z. Liang, V. K. Narasimhan and Y. Cui, Nature

communications, 2014, 5, 3943.

125. Z. Sun, J. Zhang, L. Yin, G. Hu, R. Fang, H. M. Cheng and F. Li, Nature

communications, 2017, 8, 14627.

126. X. Cui, Z. Xie and Y. Wang, Nanoscale, 2016, 8, 11984-11992.

127. Z. Li, C. Li, X. Ge, J. Ma, Z. Zhang, Q. Li, C. Wang and L. Yin, Nano Energy,

2016, 23, 15-26.

128. J. Zheng, J. Tian, D. Wu, M. Gu, W. Xu, C. Wang, F. Gao, M. H. Engelhard, J.

G. Zhang, J. Liu and J. Xiao, Nano letters, 2014, 14, 2345-2352.

129. H. Hu, L. Han, M. Yu, Z. Wang and X. W. Lou, Energy Environ. Sci., 2016, 9,

107-111.

130. R. Demir-Cakan, M. Morcrette, F. Nouar, C. Davoisne, T. Devic, D. Gonbeau,

R. Dominko, C. Serre, G. Ferey and J. M. Tarascon, Journal of the American

Chemical Society, 2011, 133, 16154-16160.

131. Z. Wang, X. Li, Y. Cui, Y. Yang, H. Pan, Z. Wang, C. Wu, B. Chen and G.

Qian, Crystal Growth & Design, 2013, 13, 5116-5120.

132. J. Zhou, R. Li, X. Fan, Y. Chen, R. Han, W. Li, J. Zheng, B. Wang and X. Li,

Energy & Environmental Science, 2014, 7, 2715.

133. Z. Wang, B. Wang, Y. Yang, Y. Cui, Z. Wang, B. Chen and G. Qian, ACS

applied materials & interfaces, 2015, 7, 20999-21004.

262 134. Y. Hou, H. Mao and L. Xu, Nano Research, 2016, 10, 344-353.

135. A. Lieb, H. Leclerc, T. Devic, C. Serre, I. Margiolaki, F. Mahjoubi, J. S. Lee,

A. Vimont, M. Daturi and J.-S. Chang, Microporous and Mesoporous

Materials, 2012, 157, 18-23.

136. H. Wei, E. F. Rodriguez, A. S. Best, A. F. Hollenkamp, D. Chen and R. A.

Caruso, Advanced Energy Materials, 2016, 23, 1601616.

137. R. Wu, X. Qian, X. Rui, H. Liu, B. Yadian, K. Zhou, J. Wei, Q. Yan, X. Q.

Feng, Y. Long, L. Wang and Y. Huang, Small, 2014, 10, 1932–1938..

138. Y. Li, D. Ye, W. Liu, B. Shi, R. Guo, H. Zhao, H. Pei, J. Xu and J. Y. Xie, ACS

applied materials & interfaces, 2016, 23, 4270.

139. G. Zhou, J. Sun, Y. Jin, W. Chen, C. Zu, R. Zhang, Y. Qiu, J. Zhao, D. Zhuo, Y.

Liu, X. Tao, W. Liu, K. Yan, H. R. Lee and Y. Cui, Advanced materials, 2017,

29, 121-130.

140. y. Lu, X. Li, J. Liang, L. Hu, Y. Zhu and Y. Qian, Nanoscale, 2016, DOI:

10.1039/c6nr05626a.

141. X.-Y. Yu, L. Yu and X. W. D. Lou, Advanced Energy Materials, 2016, 6,

12-18.

142. L. Ma, H. Zhuang, Y. Lu, S. S. Moganty, R. G. Hennig and L. A. Archer,

Advanced Energy Materials, 2014, 4, 1400390.

143. H. Xu and A. Manthiram, Nano Energy, 2017, 1.007.

144. Q. Pang, D. Kundu and L. F. Nazar, Mater. Horiz., 2016, 3, 130-136.

263 145. Z. Yuan, H. J. Peng, T. Z. Hou, J. Q. Huang, C. M. Chen, D. W. Wang, X. B.

Cheng, F. Wei and Q. Zhang, Nano letters, 2016, 16, 519-527.

146. J. Pu, Z. Shen, J. Zheng, W. Wu, C. Zhu, Q. Zhou, H. Zhang and F. Pan, Nano

Energy, 2017, 37, 7-14.

147. Z. W. Seh, J. H. Yu, W. Li, P. C. Hsu, H. Wang, Y. Sun, H. Yao, Q. Zhang and

Y. C u i , Nature communications, 2014, 5, 5017.

148. H. Lin, L. Yang, X. Jiang, G. Li, T. Zhang, Q. Yao, G. W. Zheng and J. Y. Lee,

Energy Environ. Sci., 2017, 10, 1476-1486.

149. W. Tang, Z. Chen, B. Tian, H. W. Lee, X. Zhao, X. Fan, Y. Fan, K. Leng, C.

Peng, M. H. Kim, M. Li, M. Lin, J. Su, J. Chen, H. Y. Jeong, X. Yin, Q. Zhang,

W. Zhou, K. P. Loh and G. W. Zheng, Journal of the American Chemical

Society, 2017, 139, 10133-10141.

150. S. S. Zhang and D. T. Tran, J. Mater. Chem. A, 2016, 4, 4371-4374.

151. X. Li, L. Chu, Y. Wang and L. Pan, Materials Science and Engineering: B,

2016, 205, 46-54.

152. J. D. Liu, X. S. Zheng, Z. F. Shi and S. Q. Zhang, Ionics, 2013, 20, 659-664.

153. K. Sun, D. Su, Q. Zhang, D. C. Bock, A. C. Marschilok, K. J. Takeuchi, E. S.

Takeuchi and H. Gan, Journal of The Electrochemical Society, 2015, 162,

A2834-A2839.

154. C. Ye, L. Zhang, C. Guo, D. Li, A. Vasileff, H. Wang and S.-Z. Qiao,

Advanced Functional Materials, 2017, 70, 1702524.

264 155. X.-q. Niu, X.-l. Wang, D.-h. Wang, Y. Li, Y.-j. Zhang, Y.-d. Zhang, T. Yang, T.

Yu and J.-p. Tu, J. Mater. Chem. A, 2015, 3, 17106-17112.

156. J. Jiang, J. Zhu, W. Ai, X. Wang, Y. Wang, C. Zou, W. Huang and T. Yu,

Nature communications, 2015, 6, 8622.

157. J. Zhang, H. Hu, Z. Li and X. W. Lou, Angewandte Chemie, 2016, 55,

3982-3986.

158. C. Dai, L. Hu, M.-Q. Wang, Y. Chen, J. Han, J. Jiang, Y. Zhang, B. Shen, Y.

Niu, S.-J. Bao and M. Xu, Energy Storage Materials, 2017, 4, 003.

159. D. Deng, T. An, Y. Li, Q. Wu, M.-s. Zheng and Q.-F. Dong, J. Mater. Chem. A,

2016, 12, 07221.

160. Z. Cui, C. Zu, W. Zhou, A. Manthiram and J. B. Goodenough, Advanced

materials, 2016, 28, 6926-6931.

161. N. Mosavati, S. O. Salley and K. Y. S. Ng, Journal of Power Sources, 2017,

340, 210-216.

162. X. Wang, S. Kajiyama, H. Iinuma, E. Hosono, S. Oro, I. Moriguchi, M. Okubo

and A. Yamada, Nature communications, 2015, 6, 6544.

163. M.-Q. Zhao, C. E. Ren, Z. Ling, M. R. Lukatskaya, C. Zhang, K. L. Van Aken,

M. W. Barsoum and Y. Gogotsi, Advanced materials, 2015, 27, 339-345.

164. V. M. Hong Ng, H. Huang, K. Zhou, P. S. Lee, W. Que, J. Z. Xu and L. B.

Kong, J. Mater. Chem. A, 2017, 5, 3039-3068.

165. Z. Ling, C. E. Ren, M. Q. Zhao, J. Yang, J. M. Giammarco, J. Qiu, M. W.

265 Barsoum and Y. Gogotsi, Proceedings of the National Academy of Sciences of

the United States of America, 2014, 111, 16676-16681.

166. X. Xie, M.-Q. Zhao, B. Anasori, K. Maleski, C. E. Ren, J. Li, B. W. Byles, E.

Pomerantseva, G. Wang and Y. Gogotsi, Nano Energy, 2016, 26, 513-523.

167. M. Q. Zhao, C. E. Ren, Z. Ling, M. R. Lukatskaya, C. Zhang, K. L. Van Aken,

M. W. Barsoum and Y. Gogotsi, Advanced materials, 2015, 27, 339-345.

168. J. Luo, X. Tao, J. Zhang, Y. Xia, H. Huang, L. Zhang, Y. Gan, C. Liang and W.

Zhang, ACS nano, 2016, 10, 2491-2499.

169. X. Liang, Y. Rangom, C. Y. Kwok, Q. Pang and L. F. Nazar, Advanced

materials, 2017, 29, 120-130.

170. W. Bao, D. Su, W. Zhang, X. Guo and G. Wang, Advanced Functional

Materials, 2016, 26, 8746-8756.

171. B. Sun, J. Jiang, F. Zheng, W. Zhao, B. Y. Liaw, H. Ruan, Z. Han and W.

Zhang, Journal of Power Sources, 2015, 282, 146-157.

172. Y. Hua, A. Cordoba-Arenas, N. Warner and G. Rizzoni, Journal of Power

Sources, 2015, 280, 293-312.

173. M. Dubarry, N. Vuillaume and B. Y. Liaw, International Journal of Energy

Research, 2010, 34, 216-231.

174. G. Gao, H. B. Wu, S. Ding and X. W. Lou, Small, 2015, 11, 432-436.

175. E. Mitchell, A. Jimenez, R. K. Gupta, B. K. Gupta, K. Ramasamy, M.

Shahabuddin and S. R. Mishra, New J. Chem., 2015, 39, 2181-2187.

266 176. G. Zhang, B. Y. Xia, X. Wang and X. W. David Lou, Advanced materials,

2014, 26, 2408-2412.

177. W. Xiong, Y. Gao, X. Wu, X. Hu, D. Lan, Y. Chen, X. Pu, Y. Zeng, J. Su and Z.

Zhu, ACS applied materials & interfaces, 2014, 6, 19416-19423.

178. Y. Xu, J. Wei, L. Tan, J. Yu and Y. Chen, J. Mater. Chem. A, 2015, 3,

7121-7131.

179. Q. Wang, B. Liu, X. Wang, S. Ran, L. Wang, D. Chen and G. Shen, Journal of

Materials Chemistry, 2012, 22, 21647.

180. Y. D. Mo, Q. Ru, X. Song, S. J. Hu, L. Y. Guo and X. Q. Chen, Electrochimica

Acta, 2015, 176, 575-585.

181. M. Li, J. P. Cheng, F. Liu and X. B. Zhang, Electrochimica Acta, 2015, 178,

439-446.

182. M. U. Anu Prathap, C. Wei, S. Sun and Z. J. Xu, Nano Research, 2015, 8,

2636-2645.

183. A. K. Mondal, D. Su, S. Chen, K. Kretschmer, X. Xie, H. J. Ahn and G. Wang,

Chemphyschem : a European journal of chemical physics and physical

chemistry, 2015, 16, 169-175.

184. X. Xu, B. Dong, S. Ding, C. Xiao and D. Yu, Journal of Materials Chemistry

A, 2014, 2, 13069.

185. C. An, Y. Wang, Y. Huang, Y. Xu, C. Xu, L. Jiao and H. Yuan, CrystEngComm,

2014, 16, 385-392.

267 186. D. Chen, Q. Wang, R. Wang and G. Shen, J. Mater. Chem. A, 2015, 3,

10158-10173.

187. D. Li, Y. Gong, Y. Zhang, C. Luo, W. Li, Q. Fu and C. Pan, Scientific reports,

2015, 5, 12903.

188. W.-w. Liu, C. Lu, K. Liang and B. K. Tay, J. Mater. Chem. A, 2014, 2,

5100-5107.

189. L. Shen, L. Yu, X. Y. Yu, X. Zhang and X. W. Lou, Angewandte Chemie, 2015,

54, 1868-1872.

190. J. Zhang, F. Liu, J. P. Cheng and X. B. Zhang, ACS applied materials &

interfaces, 2015, 7, 17630-17640.

191. L. S. Zhong, J. S. Hu, H. P. Liang, A. M. Cao, W. G. Song and L. J. Wan,

Advanced materials, 2006, 18, 2426-2431.

192. X. Li, L. Jiang, C. Zhou, J. Liu and H. Zeng, NPG Asia Materials, 2015, 7,

e165.

193. L. Li, S. Peng, Y. Cheah, P. Teh, J. Wang, G. Wee, Y. Ko, C. Wong and M.

Srinivasan, Chemistry, 2013, 19, 5892-5898.

194. D. Cheng, Y. Yang, J. Xie, C. Fang, G. Zhang and J. Xiong, J. Mater. Chem. A,

2015, 2, 3455.

195. T. Li, X. Li, Z. Wang, H. Guo and Y. Li, J. Mater. Chem. A, 2015, 3,

11970-11975.

196. J. Liang, Z. Fan, S. Chen, S. Ding and G. Yang, Chemistry of Materials, 2014,

268 26, 4354-4360.

197. J. Li, S. Xiong, Y. Liu, Z. Ju and Y. Qian, ACS applied materials & interfaces,

2013, 5, 981-988.

198. Y. He, L. Xu, Y. Zhai, A. Li and X. Chen, Chemical communications, 2015, 51,

14768-14771.

199. E. Uchaker, N. Zhou, Y. Li and G. Cao, J. Phys. Chem. C, 2013, 117,

1621-1626.

200. V. K. LaMer and R. H. Dinegar, J. Am. Chem. Soc., 1950, 72, 4847-4854.

201. A. K. Mondal, D. Su, S. Chen, X. Xie and G. Wang, ACS applied materials &

interfaces, 2014, 6, 14827-14835.

202. X. Fang, W. Weng, J. Ren and H. Peng, Advanced materials, 2016, 28,

491-496.

203. L. Qie, C. Zu and A. Manthiram, Advanced Energy Materials, 2016, 459,

12-18.

204. Z. Ma, S. Dou, A. Shen, L. Tao, L. Dai and S. Wang, Angewandte Chemie,

2015, 54, 1888-1892.

205. M. Wang, H. Zhang, Q. Wang, C. Qu, X. Li and H. Zhang, ACS applied

materials & interfaces, 2015, 7, 3590-3599.

206. Q. Sun, X. Fang, W. Weng, J. Deng, P. Chen, J. Ren, G. Guan, M. Wang and H.

Peng, Angewandte Chemie, 2015, 54, 10539-10544.

207. X. G. Sun, Y. Fang, X. Jiang, K. Yoshii, T. Tsuda and S. Dai, Chemical

269 communications, 2016, 52, 292-295.

208. J. Song, M. L. Gordin, T. Xu, S. Chen, Z. Yu, H. Sohn, J. Lu, Y. Ren, Y. Duan

and D. Wang, Angewandte Chemie, 2015, 54, 4325-4329.

209. F. Wu, J. Li, Y. Su, J. Wang, W. Yang, N. Li, L. Chen, S. Chen, R. Chen and L.

Bao, Nano letters, 2016, 23, 1981.

210. J. Yang, Y.-X. Wang, S.-L. Chou, R. Zhang, Y. Xu, J. Fan, W.-x. Zhang, H.

Kun Liu, D. Zhao and S. Xue Dou, Nano Energy, 2015, 18, 133-142.

211. J. Zang, T. An, Y. Dong, X. Fang, M. Zheng, Q. Dong and N. Zheng, Nano

Research, 2015, 8, 2663-2675.

212. S. Xu, C. M. Hessel, H. Ren, R. Yu, Q. Jin, M. Yang, H. Zhao and D. Wang,

Energy Environ. Sci., 2014, 7, 632-637.

213. M. Rao, X. Geng, X. Li, S. Hu and W. Li, Journal of Power Sources, 2012,

212, 179-185.

214. Y. -L. Ding, P. Kopold, K. Hahn, P. A. van Aken, J. Maier and Y. Yu, Advanced

Functional Materials, 2015, 94, 12-19.

215. J. Chen, K. S. Han, W. A. Henderson, K. C. Lau, M. Vijayakumar, T. Dzwiniel,

H. Pan, L. A. Curtiss, J. Xiao, K. T. Mueller, Y. Shao and J. Liu, Advanced

Energy Materials, 2016, 12, 160.

216. S. Chen, F. Dai, M. L. Gordin, Z. Yu, Y. Gao, J. Song and D. Wang,

Angewandte Chemie, 2016, 55, 4231-4235.

217. S. Liu, G. R. Li and X. P. Gao, ACS applied materials & interfaces, 2016, 8,

270 7783-7789.

218. J. R. Nair, F. Bella, N. Angulakshmi, A. M. Stephan and C. Gerbaldi, Energy

Storage Materials, 2016, 3, 69-76.

219. T. Nakazawa, A. Ikoma, R. Kido, K. Ueno, K. Dokko and M. Watanabe,

Journal of Power Sources, 2016, 307, 746-752.

220. R. Ponraj, A. G. Kannan, J. H. Ahn and D. W. Kim, ACS applied materials &

interfaces, 2016, 8, 4000-4006.

221. Q. Pang, J. Tang, H. Huang, X. Liang, C. Hart, K. C. Tam and L. F. Nazar,

Advanced materials, 2015, 27, 6021-6028.

222. Y. Qiu, W. Li, W. Zhao, G. Li, Y. Hou, M. Liu, L. Zhou, F. Ye, H. Li, Z. Wei, S.

Yang, W. Duan, Y. Ye, J. Guo and Y. Zhang, Nano letters, 2014, 14,

4821-4827.

223. M. Agostini, B. Scrosati and J. Hassoun, Advanced Energy Materials, 2015,

81, 121-130.

224. M. Chen, J. Liu, D. Chao, J. Wang, J. Yin, J. Lin, H. Jin Fan and Z. Xiang

Shen, Nano Energy, 2014, 9, 364-372.

225. Z. Chen, W. Ren, L. Gao, B. Liu, S. Pei and H. M. Cheng, Nature materials,

2011, 10, 424-428.

226. B. Li, S. Li, J. Liu, B. Wang and S. Yang, Nano letters, 2015, 15, 3073-3079.

227. H. Yu, D. Ye, T. Butburee, L. Wang and M. Dargusch, ACS applied materials

& interfaces, 2016, 8, 2505-2510.

271 228. G. Zhou, L. Li, C. Ma, S. Wang, Y. Shi, N. Koratkar, W. Ren, F. Li and H.-M.

Cheng, Nano Energy, 2015, 11, 356-365.

229. L. Fei, X. Li, W. Bi, Z. Zhuo, W. Wei, L. Sun, W. Lu, X. Wu, K. Xie, C. Wu, H.

L. Chan and Y. Wang, Advanced materials, 2015, 27, 5936-5942.

230. L. Ji, M. Rao, H. Zheng, L. Zhang, Y. Li, W. Duan, J. Guo, E. J. Cairns and Y.

Zhang, Journal of the American Chemical Society, 2011, 133, 18522-18525.

231. H. W. Shim, Y. H. Jin, S. D. Seo, S. H. Lee and D. W. Kim, ACS Nano, 2011,

5, 443–449..

232. J. Wang, Y. S. He and J. Yang, Advanced materials, 2015, 27, 569-575.

233. C. Zu and A. Manthiram, Advanced Energy Materials, 2014, 4, 23-30.

234. T. Yang, T. Qian, M. Wang, X. Shen, N. Xu, Z. Sun and C. Yan, Advanced

materials, 2016, 28, 539-545.

235. L. Miao, W. Wang, K. Yuan, Y. Yang and A. Wang, Chemical communications,

2014, 50, 13231-13234.

236. L. Yuan, X. Qiu, L. Chen and W. Zhu, Journal of Power Sources, 2009, 189,

127-132.

237. Z. Wang, Y. Dong, H. Li, Z. Zhao, H. B. Wu, C. Hao, S. Liu, J. Qiu and X. W.

Lou, Nature communications, 2014, 5, 5002.

238. H.-J. Peng, T.-Z. Hou, Q. Zhang, J.-Q. Huang, X.-B. Cheng, M.-Q. Guo, Z.

Yuan, L.-Y. He and F. Wei, Advanced Materials Interfaces, 2014, 1, 1400227.

239. C. J. Hart, M. Cuisinier, X. Liang, D. Kundu, A. Garsuch and L. F. Nazar,

272 Chemical communications, 2015, 51, 2308-2311.

240. Q. Fan, W. Liu, Z. Weng, Y. Sun and H. Wang, Journal of the American

Chemical Society, 2015, 137, 12946-12953.

241. J. Song, Z. Yu, M. L. Gordin and D. Wang, Nano letters, 2016, 16, 864-870.

242. B. Cao, Li, B. Hou, Y. Mo, L. Yin and Y. Chen, ACS applied materials &

interfaces, 2016, 12, 109-114.

243. J. Cao, C. Chen, Q. Zhao, N. Zhang, Q. Lu, X. Wang, Z. Niu and J. Chen,

Advanced materials, 2016, 23, 262.

244. m. chen, X. Wang, s. cai, z. ma, P. Song and A. C. Fisher, J. Mater. Chem. A,

2016, 45, 863-872.

245. X.-B. Cheng, C. Yan, J.-Q. Huang, P. Li, L. Zhu, L. Zhao, Y. Zhang, W. Zhu,

S.-T. Yang and Q. Zhang, Energy Storage Materials, 2017, 6, 18-25.

246. X. Liu, J. Q. Huang, Q. Zhang and L. Mai, Advanced materials, 2017, 4,

1601759.

247. S. Choi, T. Bok, J. Ryu, J.-I. Lee, J. Cho and S. Park, Nano Energy, 2015, 12,

161-168.

248. H. Park, S. Choi, S. Lee, G. Hwang, N.-S. Choi and S. Park, J. Mater. Chem. A,

2015, 3, 1325-1332.

249. J. Ryu, D. Hong, S. Choi and S. Park, ACS nano, 2016, 10, 2843-2851.

250. S. H. Yu, B. Lee, S. Choi, S. Park, B. H. Hong and Y. E. Sung, Chemical

communications, 2016, 52, 3203-3206.

273 251. H.-S. Kang and Y.-K. Sun, Advanced Functional Materials, 2016, 26,

1225-1232.

252. L. Hu, Y. Huang, F. Zhang and Q. Chen, Nanoscale, 2013, 5,4186.

253. G. Hu, Z. Sun, C. Shi, R. Fang, J. Chen, P. Hou, C. Liu, H. M. Cheng and F. Li,

Advanced materials, 2016, 23, 1603835.

254. W. Zhou, X. Xiao, M. Cai and L. Yang, Nano letters, 2014, 14, 5250-5256.

255. W. Zhou, Y. Yu, H. Chen, F. J. DiSalvo and H. D. Abruna, Journal of the

American Chemical Society, 2013, 135, 16736-16743.

256. J. Sun, Y. Sun, M. Pasta, G. Zhou, Y. Li, W. Liu, F. Xiong and Y. Cui,

Advanced materials, 2016, 23, 1602172.

257. H. Xu, L. Qie and A. Manthiram, Nano Energy, 2016, 26, 224-232.

258. X. Liang, C. Hart, Q. Pang, A. Garsuch, T. Weiss and L. F. Nazar, Nature

communications, 2015, 6, 5682.

259. J. Xu, D. Su, W. Zhang, W. Bao and G. Wang, J. Mater. Chem. A, 2016, 12,

878.

260. S.-H. Chung, C.-H. Chang and A. Manthiram, Energy Environ. Sci., 2016, 9,

3188-3200.

261. H. Gao, Q. Lu, N. Liu, X. Wang and F. Wang, J. Mater. Chem. A, 2015, 3,

7215-7218.

262. G. Zhou, Y. Zhao and A. Manthiram, Advanced Energy Materials, 2015, 5,

11-18.

274 263. J. Hou, T. Cao, F. Idrees and C. Cao, Nanoscale, 2016, 8, 451-457.

264. Z. Li and L. Yin, ACS applied materials & interfaces, 2015, 7, 4029-4038.

265. W. Bao, Z. Zhang, C. Zhou, Y. Lai and J. Li, Journal of Power Sources, 2014,

248, 570-576.

266. E. Ghosh, L. Shmuylovich and S. J. Kovacs, Conf Proc IEEE Eng Med Biol

Soc, 2009, 2009, 2883-2886.

267. N. Jayaprakash, J. Shen, S. S. Moganty, A. Corona and L. A. Archer,

Angewandte Chemie, 2011, 50, 5904-5908.

268. K. Zhang, Q. Zhao, Z. Tao and J. Chen, Nano Research, 2012, 6, 38-46.

269. J. Schuster, G. He, B. Mandlmeier, T. Yim, K. T. Lee, T. Bein and L. F. Nazar,

Angewandte Chemie, 2012, 51, 3591-3595.

270. H. Ren, R. Yu, J. Wang, Q. Jin, M. Yang, D. Mao, D. Kisailus, H. Zhao and D.

Wang, Nano letters, 2014, 14, 6679-6684.

271. Y. Yang, Q. Jin, D. Mao, J. Qi, Y. Wei, R. Yu, A. Li, S. Li, H. Zhao, Y. Ma, L.

Wang, W. Hu and D. Wang, Advanced materials, 2017, 29, 31-39.

272. J. Tang, R. R. Salunkhe, J. Liu, N. L. Torad, M. Imura, S. Furukawa and Y.

Yamauchi, Journal of the American Chemical Society, 2015, 137, 1572-1580.

273. J. Zhou, N. Lin, W. l. Cai, C. Guo, K. Zhang, J. Zhou, Y. Zhu and Y. Qian,

Electrochimica Acta, 2016, 218, 243-251.

274. J. Yan, X. Liu, M. Yao, X. Wang, T. K. Wafle and B. Li, Chemistry of

Materials, 2015, 78, 127-130.

275 275. J. Liang, L. C. Yin, X. Tang, H. Yang, W. Yan, L. Song, H. M. Cheng and F. Li,

ACS applied materials & interfaces, 2016, 47, 129-136.

276. X. Gu, C.-j. Tong, C. Lai, J. Qiu, X. Huang, W. Yang, B. Wen, L.-m. Liu, Y.

Hou and S. Zhang, J. Mater. Chem. A, 2015, 3, 16670-16678.

277. M. Zhang, C. Yu, C. Zhao, X. Song, X. Han, S. Liu, C. Hao and J. Qiu, Energy

Storage Materials, 2016, 5, 223-229.

278. H. Wang, T. Zhou, D. Li, H. Gao, G. Gao, A. Du, H. Liu and Z. Guo, ACS

applied materials & interfaces, 2016, 8, 13-20.

279. L. Jia, T. Wu, J. Lu, L. Ma, W. Zhu and X. Qiu, ACS applied materials &

interfaces, 2016, 23, 10366.

280. L. Zhang, L. Ji, P. A. Glans, Y. Zhang, J. Zhu and J. Guo, Physical chemistry

chemical physics : PCCP, 2012, 14, 13670-13675.

281. Q. Pang, D. Kundu, M. Cuisinier and L. F. Nazar, Nature communications,

2014, 5, 4759.

282. X. Li, Q. Sun, J. Liu, B. Xiao, R. Li and X. Sun, Journal of Power Sources,

2016, 302, 174-179.

283. S. Zheng, F. Yi, Z. Li, Y. Zhu, Y. Xu, C. Luo, J. Yang and C. Wang, Advanced

Functional Materials, 2014, 24, 4156-4163.

284. J. Liu, C. Wu, D. Xiao, P. Kopold, L. Gu, P. A. van Aken, J. Maier and Y. Yu,

Small, 2016, 12, 2354-2364.

285. F. Zou, X. Hu, Z. Li, L. Qie, C. Hu, R. Zeng, Y. Jiang and Y. Huang, Advanced

276 materials, 2014, 26, 6622-6628.

286. G. Xu, B. Ding, L. Shen, P. Nie, J. Han and X. Zhang, Journal of Materials

Chemistry A, 2013, 1, 4490.

287. J. Zhou, X. Yu, X. Fan, X. Wang, H. Li, Y. Zhang, W. Li, J. Zheng, B. Wang

and X. Li, J. Mater. Chem. A, 2015, 3, 8272-8275.

288. H. Hu, B. Guan, B. Xia and X. W. Lou, Journal of the American Chemical

Society, 2015, 137, 5590-5595.

289. R. Wu, X. Qian, X. Rui, H. Liu, B. Yadian, K. Zhou, J. Wei, Q. Yan, X. Q.

Feng, Y. Long, L. Wang and Y. Huang, Small, 2014, 10, 1932-1938.

290. P. Poizot, S. Laruelle, S. Grugeon, L. Dupon andn J. M. Tarascon, Nature,

2000, 407, 496–499..

291. Z. Wu, L.-P. Sun, M. Yang, L.-H. Huo, H. Zhao and J.-C. Grenier, J. Mater.

Chem. A, 2016, 4, 13534-13542.

292. Y. -J. Li, J.-M. Fan, M.-S. Zheng and Q.-F. Dong, Energy Environ. Sci., 2016,

9, 1998-2004.

293. V. Finsy, H. Verelst, L. Alaerts, D. D. Vos, P. A. Jacobs, G. V. Baron and J. F.

M. Denayer, J. Am. Chem. Soc., 2008,130, 7110–7118..

294. Y. Hou, Z. Wen, S. Cui, S. Ci, S. Mao and J. Chen, Advanced Functional

Materials, 2015, 25, 872-882.

295. X. Zeng, C. You, L. Leng, D. Dang, X. Qiao, X. Li, Y. Li, S. Liao and R. R.

Adzic, J. Mater. Chem. A, 2015, 3, 11224-11231.

277 296. Z. Xiao, Z. Yang, L. Wang, H. Nie, M. Zhong, Q. Lai, X. Xu, L. Zhang and S.

Huang, Advanced materials, 2015, 27, 2891-2898.

297. C. Barchasz, F. Molton, C. Duboc, J. C. Lepretre, S. Patoux and F. Alloin,

Anal Chem, 2012, 84, 3973-3980.

298. Y. Li, H. Zhan, S. Liu, K. Huang and Y. Zhou, Journal of Power Sources, 2010,

195, 2945-2949.

299. H. Chen, C. Wang, Y. Dai, S. Qiu, J. Yang, W. Lu and L. Chen, Nano letters,

2015, 15, 5443-5448.

278