Solid State Electrolytes for Energy Storage and Conversion

University of Calgary

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Graduate Studies The Vault: Electronic Theses and Dissertations

2016

Solid State Electrolytes for Energy Storage and

Conversion Devices

Narayanan, Sumaletha

Narayanan, S. (2016). Solid State Electrolytes for Energy Storage and Conversion Devices

(Unpublished doctoral thesis). University of Calgary, Calgary, AB. doi:10.11575/PRISM/27823

http://hdl.handle.net/11023/3165

doctoral thesis

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UNIVERSITY OF CALGARY

Solid State Electrolytes for Energy Storage and Conversion Devices

by

Sumaletha Narayanan

A THESIS

SUBMITTED TO THE FACULTY OF GRADUATE STUDIES

IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE

DEGREE OF DOCTOR OF PHILOSOPHY

GRADUATE PROGRAM IN CHEMISTRY

CALGARY, ALBERTA

JULY, 2016

© Sumaletha Narayanan 2016

ii

Abstract

Electrochemical energy conversion and storage devices hold significant importance in the

successful implementation of renewable energy systems. Solid-state electrolytes, with

garnet-type crystal structure for lithium ion batteries, have been synthesized using ceramic

method for energy storage application. A systematic study on lithium-stuffed garnet-type

Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) have been carried out for the understanding of the

effect of Y- and Li- doping on the structural, electrical, chemical and electrochemical

properties. Powder X-ray diffraction (PXRD) studies have revealed the cubic garnet-type

structure of the materials. The AC electrochemical impedance spectroscopy (EIS) has

shown that the sample with highest Li and Y content show the best Li+ ion conductivity of

10-4 Scm-1 at room temperature.

Optimization of lithium salt added during the solid-state preparation of garnets to

compensate lithium loss at higher sintering temperatures was also carried out on

Li6La3Ta1.5Y0.5O12. Crystal structure was unaltered with the change in lithium salt addition,

whereas the porosity and conductivity were affected. Evaluation of fundamental transport

properties of Li5+2xLa3M2-xYxO12 (M = Nb, Ta) (x = 0.25, 0.50 and 0.75) has been carried

out by employing AC EIS method. The dielectric analysis performed below room

temperature suggested that the Li+ ion conduction in garnet-type materials takes place

through a hopping mechanism following Path A (low energy route) or Path B (high energy

route).

Hybrid proton conducting materials derived from ionic liquid and

polyoxometalates have been developed for the potential application in energy conversion

devices. The effect of different heteropoly acids such as H3PW11MoO40, H4PMo11VO40 and

iii

H5PMo11V2O40 used along with (3-(pyridin-1-ium-1-yl) propane-1-sulfonate (PyPs)) ionic

liquid on the properties of hybrid materials were studied. A high thermal stability up to ~

300 °C, electrochemical stability of ~ 3V, and ionic conductivity of 0.01 Scm-1 at 90 °C

were observed for the hybrid proton conductors. A soft-chemistry approach was proposed

for the synthesis of another class of proton conductors, from layered perovskites such as

KLaNb2O7 and K2La2Ti3O10, and an imidazolium based ionic liquid. PXRD, scanning

electron microscopy (SEM), and thermogravimetric analysis (TGA) were used to

understand the ion exchange chemistry.

iv

Acknowledgements

First, I extend my deep gratitude to my supervisor, Prof. Venkataraman Thangadurai, for

his endless support and encouragement. He has been a great mentor throughout my years

at University of Calgary, and I am grateful to his valuable guidance and advice on my

research. This work would not have been possible without his continuous moral support

and motivation. His endless patience, perseverance, and time management are only few

examples which I learned from him for my life. I sincerely thank my supervisory

committee, Drs. Todd Sutherland and Yujun Shi for their critical comments and valuable

suggestions on my research. I also thank my external examiners, Dr. Rajamani Nagarajan,

University of Delhi and Dr. Uttandaraman Sundararaj, University of Calgary for their

valuable time.

I would like to thank Drs. Xia Tong, Ashok Kumar Baral and Farshid

Ramezanipour for all their contributions to my work. I also acknowledge the valuable help

from all our collaborators, Prof. Eric D Wachsman and Greory Hits from University of

Maryland, and Prof. Francoise Le Berre from Université du Maine. I thank my summer

student, Shahrukh Shamim for his help with part of the experimental work on ionic liquid

project.

I also thank all the present and past members of the Thangadurai and Birss groups.

Special thanks to all my friends with whom I cherish my school days at the University of

Calgary. I thank Dr. Anand Singh, Dr. Xiaoan Li, Annie Hoang, Robert Mayall and Aaron

Kirkey for their help.

I thank Jianjun Li (NMR analysis) Mark Toonen (glass shop), Mike Siewert, Keith

Collins (electronic shop), Bob Stockwell, Magdi Khalil, and Nelson Cruz (chemstore) for

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all their support. I also thank Bonnie King, Janice Crawford, and other staff members in

the Chemistry office for their administrative support over all these years. I thank my

teaching team Dr. Kal Mahadev, Dr. Julie Lefebvre, Dr. Vivian Mozol, Doris Jo, Jill

Helfer, and Kristin Yaehne for their support in giving me a wonderful teaching experience.

Finally, I extend my gratitude towards my family for their constant support, prayers

and motivation.

vi

Table of Contents

Abstract ............................................................................................................................... ii Acknowledgements ............................................................................................................ iv

Table of Contents ............................................................................................................... vi List of Tables ..................................................................................................................... ix List of Figures and Illustrations ......................................................................................... xi List of Symbols and Abbreviations..................................................................................xxv

CHAPTER ONE: INTRODUCTION ..................................................................................1

1.1 Project motivation ......................................................................................................1 1.2 Research objectives ....................................................................................................5 1.3 Thesis organization ....................................................................................................7

CHAPTER TWO: BACKGROUND ...................................................................................9 2.1 Introduction ................................................................................................................9 2.2 Lithium ion batteries (LIBs) ....................................................................................10

2.2.1 Operating principle of LIBs .............................................................................12 2.2.2 Challenges of current LIB research .................................................................14

2.3 Electrolytes for LIBs ................................................................................................14 2.4 Solid-state Li ion electrolytes ..................................................................................16

2.4.1 LISICON and NASICON-structured ceramic materials .................................17

2.4.2 A-site deficient perovskites .............................................................................20 2.4.3 Li-rich anti-perovskites (LiRAP) ....................................................................22

2.4.4 Garnet-type Li ion conductors .........................................................................24

2.4.4.1 Li5- garnet family ...................................................................................26

2.4.4.2 Li6-garnet family ....................................................................................27 2.4.4.3 Li7-garnet family ....................................................................................28

2.4.4.4 Li ion occupancy and conductivity of Li5-, Li6- and Li7-garnets ...........30 2.5 Proton exchange membrane fuel cells (PEMFCs) ...................................................33

2.5.1 Challenges of PEMFCs ...................................................................................34

2.6 Proton conductors for PEMFCs ...............................................................................35 2.6.1 Polymer membrane electrolytes ......................................................................35

2.6.2 Modifications on Nafion ..................................................................................36 2.7 Polyoxometalates (POMs) and ionic liquid derived hybrid proton conductors .......36

2.8 Tuning layered perovskites as an alternative to Nafion membrane for PEMFCs ...38

CHAPTER THREE: EXPERIMENTAL SECTION .........................................................42

3.1 Materials synthesis ...................................................................................................42 3.1.1 Synthesis of lithium-stuffed garnet-type metal oxides ....................................42 3.1.2 Synthesis of polyoxometalate and ionic liquid based proton conductors ........43

3.1.2.1 Synthesis of solid acid, H3PW11MoO40 (H3PWMo) ..............................43 3.1.2.2 Synthesis of solid acid H4PMo11VO40 (H4PMoV) .................................43

3.1.2.3 Synthesis of solid acid H5PMo10V2O40 (H5PMoV) ...............................43 3.1.2.4 Synthesis of PyPs ionic liquid ...............................................................44 3.1.2.5 Synthesis of hybrid gels .........................................................................44

vii

3.1.3 Synthesis of proton conductors based on layered perovskites and ionic

liquid ................................................................................................................44 3.1.3.1 Synthesis of KLaNb2O7 and K2La2Ti3O10 layered perovskite and its

proton analogues .....................................................................................44

3.1.3.2 Synthesis of hybrid proton conductors from layered perovskite and

ionic liquid ..............................................................................................45 3.2 Material characterization .........................................................................................45

3.2.1 Powder X-ray diffraction (PXRD) ..................................................................45 3.2.2 Scanning electron microscopy (SEM) .............................................................47

3.2.3 Density measurements .....................................................................................47 3.2.4 Electrochemical impedance spectroscopy (EIS) .............................................49 3.2.5 Cyclic voltammetry (CV) ................................................................................53

3.2.6 Thermogravimetric analysis ............................................................................54 3.2.7 Fourier transform infrared (FTIR) spectroscopic studies ................................56 3.2.8 Solid-state magic angle spinning nuclear magnetic resonance (MAS

NMR) studies ...................................................................................................56 3.2.9 Raman spectroscopy ........................................................................................57

3.2.10 Chemical stability tests of Li5+2xLa3Ta2-xYxO12 garnets ................................57 3.3 Error considerations .................................................................................................58

CHAPTER FOUR: DOPANT CONCENTRATION - POROSITY - Li-ION

CONDUCTIVITY RELATIONSHIP IN GARNET-TYPE Li5+2xLa3Ta2-xYxO12

(0.05 ≤ x ≤ 0.75) AND THEIR STABILITY IN WATER AND 1M LiCl ..............60

4.1 General overview125 .................................................................................................60 4.2 Results and discussion125 .........................................................................................61

4.2.1 Phase analysis of Li5+2xLa3Ta2-xYxO12 ............................................................61 4.2.2 Li ion conductivity as a function of Y-doping and temperature. ....................70

4.2.3 Electrochemical stability .................................................................................90 4.3 Summary125 ..............................................................................................................92

CHAPTER FIVE: EFFECT OF EXCESS Li ON THE STRUCTURAL AND

ELECTRICAL PROPERTIES OF GARNET-TYPE Li6La3Ta1.5Y0.5O12 ................93 5.1 General overview145 .................................................................................................93

5.2 Results and discussion145 .........................................................................................95 5.2.1 Structure and density analysis .........................................................................95

5.2.2 Microstructure and chemical composition analysis ........................................97 5.2.3 Electrical characterization .............................................................................104

5.3 Summary145 ............................................................................................................110

CHAPTER SIX: EVALUATION OF FUNDAMENTAL TRANSPORT

PROPERTIES OF Li-EXCESS GARNET-TYPE Li5+2xLa3Ta2-xYxO12 (x = 0.25,

0.5 AND 0.75) ELECTROLYTES USING AC IMPEDANCE AND

DIELECTRIC SPECTROSCOPY ..........................................................................111

6.1 General overview130 ...............................................................................................111 6.2 Results and discussion ...........................................................................................112

6.2.1 Crystallographic sites and migration pathways in Li-excess garnets ............112

6.2.2 Ionic conductivity ..........................................................................................116

viii

6.2.3 Dielectric permittivity and loss tangent .........................................................119

CHAPTER SEVEN: DIELECTRIC CHARACTERISTICS OF FAST Li ION

CONDUCTING GARNET-TYPE Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 AND

0.75) ........................................................................................................................136

7.1 General overview185 ...............................................................................................136 7.2 Results and discussion185 .......................................................................................137

7.2.1 AC Impedance spectroscopy analysis ...........................................................137 7.3 Summary185 ............................................................................................................156

CHAPTER EIGHT: SYNTHESIS AND ELECTRICAL PROPERTIES OF HYBRID

GEL ELECTROLYTES DERIVED FROM KEGGIN-TYPE HETEROPOLY

ACIDS AND 3-(PYRIDIN-1-IUM-1-YL) PROPANE-1-SULFONATE (PyPs) ..158

CHAPTER NINE: POTENTIAL PROTON CONDUCTING HYBRIDS DERIVED

FROM LAYERED PEROVSKITES AND IONIC LIQUID VIA SOFT-

CHEMISTRY SYNTHESIS ...................................................................................184 9.1 General overview ...................................................................................................184

9.2 Results and discussion ...........................................................................................184 9.2.1 Intercalation Chemistry of Dion-Jacobson – DJ phase KLaNb2O7 ...............184

9.2.2 Intercalation chemistry of Ruddlesdon-Popper-RP phase K2La2Ti3O10 .......193 9.2.3 Intercalation chemistry of H2La2Ti3O10 .........................................................202 9.2.4 Extended study of intercalation chemistry of Ruddlesdon-Popper-RP

phase K2La2Ti3O10 .........................................................................................203 9.3 Summary ................................................................................................................210

CHAPTER TEN: CONCLUSIONS AND FUTURE WORK .........................................212 10.1 Conclusions ..........................................................................................................212

10.1.1 Lithium-stuffed garnet-type material research for the application of

energy storage ................................................................................................212

10.1.2 Proton conductors based on metal oxides and ionic liquid hybrid systems

for energy conversion ....................................................................................214

10.2 Suggestions for future work .................................................................................216

REFERENCES ................................................................................................................218

APPENDIX: COPYRIGHTS......................................................................................233

ix

List of Tables

Table 1.1 Comparison of voltage, energy and power densities of various secondary

batteries.1 ..................................................................................................................... 3

Table 1.2 Classification and the state-of-the-art of fuel cells.4 ........................................... 4

Table 2.1 Summary of different types of batteries and their applications.12 .................... 11

Table 4.1 The Rietveld refinement results for Li5+2xLa3Ta2-xYxO12 (x = 0.05 - 0.20).

Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977. Copyright

(2015) American Chemical Society.125 ..................................................................... 62

Table 4.2 The Rietveld refinement results for Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.50 and

0.75). Reproduced by permission of the PCCP Owner Societies.130 ........................ 63

Table 4.3 AC impedance fitting results of Li5+2xLa3Ta2-xYxO12 (x = 0.05 and 0.10).



Table 4.4 The room temperature AC impedance fitting results of Li5+2xLa3Ta2-xYxO12

(x = 0.20, 0.25, and 0.50). Reprinted with permission from Inorg. Chem. 2015,

54, 6968-6977. Copyright (2015) American Chemical Society.125 .......................... 74

Table 4.5 The room temperature (25 °C) conductivity and activation energies of

Li5+2xLa3Ta2-xYxO12 calculated at 23-325 °C. Reprinted with permission from

Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical

Society.125 .................................................................................................................. 75

Table 4.6 Calculated and measured density results of Li5+2xLa3Ta2-xYxO12 from

PXRD and Archimedes methods. Reprinted with permission from Inorg. Chem.

2015, 54, 6968-6977. Copyright (2015) American Chemical Society.125 ................ 81

Table 5.1 Comparison of densities calculated from PXRD data and measured from He

Pycnometer and the lattice parameter (a) of Li6La3Ta1.5Y0.5O12 with excess

amount of lithium (2.5 - 15 wt.% LiNO3) added during solid-state synthesis.

Reproduced by permission of ECS-The Electrochemical Society.145....................... 96

Table 5.2 The AC impedance fitting results measured at 50 °C of Li6La3Ta1.5Y0.5O12

with excess amount of lithium (2.5, 5 and 7 wt.% LiNO3) added during solid-

state synthesis. Reproduced by permission of ECS-The Electrochemical

Society.145 ................................................................................................................ 107

Table 6.1 Values of conductivities (σ), activation energies (Ea) and pre exponential

factor (A) for Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.50 and 0.75). Phys. Chem. Chem.

Phys. 2014, 16, 11356-11365. Reproduced by permission of the PCCP Owner

Societies.130 ............................................................................................................. 118

x

Table 7.1 The bulk (except bulk + grain-boundary for Li6.5-Nb) conductivity of

Li5+2xLa3Nb2-xYxO12 at -22 and 25 °C and the activation energy calculated from

Arrhenius plots (Figure 7.2) at temperature range -50 to 50 °C. Phys. Chem.

Chem. Phys. 2016, 18, 15418-15426. Reproduced by permission of the PCCP

Owner Societies.185 ................................................................................................. 140

Table 7.2 Comparison of ionic conductivity of both Nb and Ta phases of Li5+2xLa3M2-

xYxO12 at specific temperatures.57, 125, 130 ................................................................ 155

Table 8.1 CHN elemental analysis PyPs and its hybrid compounds, PyPs-H3PWMo,

PyPs-H4PMoV, and PyPs-H5PMoV. ...................................................................... 161

Table 8.2 FTIR Bands of IL-H3PWMo, IL-H4PMoV and IL-H5PMoV and their

Precursors. ............................................................................................................... 167

Table 8.3 The impedance fitting results of PyPs-H4PMoV and PyPs-H5PMoV hybrids

at different temperatures. ........................................................................................ 178

Table 9.1 The shift in ((004) plane of PXRD in terms of d-spacing values of as-

prepared, acetonitrile and IL treated KLaNb2O7. .................................................... 186

Table 9.2 The shift in (002) plane of PXRD in terms of d-spacing values of as-

prepared, acetonitrile and IL treated K2La2Ti3O10. ................................................. 196

xi

List of Figures and Illustrations

Figure 1.1 General comparison of discharge rates and power rating of different

electrochemical energy storage technologies. Copyright permission is obtained.1 .... 2

Figure 2.1 Comparison of volumetric and gravimetric energy densities of various

commercially available secondary batteries.1 ........................................................... 12

Figure 2.2 Operating principle of a LIB using LiCoO2 cathode and graphite anode. ...... 13

Figure 2.3 Li ion conductivity of various electrolytes including liquid, polymer and

solid-state electrolytes.12, 16, 19-23 ............................................................................... 16

Figure 2.4 Idealized crystal structure of LISICON with formula Li3Zn0.5GeO4.

Reproduced by permission of The Royal Society of Chemistry.14 ........................... 17

Figure 2.5 Idealized crystal structure of NASICON with formula LiM2(PO4)3.

Reproduced by permission of The Royal Society of Chemistry.14 ........................... 18

Figure 2.6 Arrhenius plots of recently studied NASICONs.32, 36, 39, 40 ............................. 20

Figure 2.7 Crystal structure of A-site deficient perovskite-type La(2/3)-xLi3xTiO3 with

general formula ABO3. Reproduced by permission of The Royal Society of

Chemistry.14 .............................................................................................................. 21

Figure 2.8 Temperature dependence of total conductivity for some A-site deficient

perovskites.21, 27, 41, 44 ................................................................................................. 22

Figure 2.9 Crystal structure of Li-rich anti-perovskites (LiRAP) with nominal formula

of Li3O(Cl,Br). Adapted with permission from Zhao, Y.; Daemen, L. L., J.

Amer. Chem. Soc. 2012, 134, 15042-15047. Copyright (2012) American

Chemical Society.47 ................................................................................................... 24

Figure 2.10 The crystal structure of ideal garnet structure with general formula

A3B2C3O12. Reprinted with permission from Thangadurai, V.; Pinzaru, D.;

Narayanan, S.; Baral, A. K., J. Phys. Chem. Lett. 2015, 6, 292-299. Copyright

(2015) American Chemical Society.15 ...................................................................... 25

Figure 2.11 The crystal structure of garnet-like Li-stuffed Li5La3M2O12. ........................ 26

Figure 2.12 Conductivity comparison of some representative garnet-type Li ion

conductors from Li5-, Li6-, and Li7-families.14, 78 ..................................................... 30

Figure 2.13 Representation of possible Li-occupancy at different crystallographic

sites in Li-stuffed garnets. Reproduced by permission of The Royal Society of

Chemistry.14 .............................................................................................................. 31

xii

Figure 2.14 Correlation of Li-site occupancy and ionic conductivity towards the Li

content in the garnets. Reproduced by permission of The Royal Society of

Chemistry.14 .............................................................................................................. 32

Figure 2.15 The illustration of proton exchange membrane fuel cell (PEMFC). ............. 34

Figure 2.16 The chemical structure of Nafion membrane.83 ............................................ 35

Figure 2.17 The crystal structure of (a) Anderson [XM6O24]n-, (b) Keggin

[XM12O40]n-, and (c) Dawson [X2M18O62]

n- families. Adapted from Ivanova, S.,

ISRN Chemical Engineering 2014, 2014, 13.95 ........................................................ 37

Figure 2.18 The molecular structure of 3-(pyridine-1-ium-1-yl)propane-1-sulfonate

(PyPs). ....................................................................................................................... 38

Figure 2.19 Comparison of proton conductivity of solid- and liquid-state materials.106 .. 39

Figure 2.20 The crystal structure of a) KLaNb2O7 and b) K2La2Ti3O10. .......................... 41

Figure 2.21 The structure of ionic liquid containing imidazolium cation and imide

anion. For simplicity, the structure is represented as IL. .......................................... 41

Figure 3.1 Illustration of Bragg’s law for powder-X-ray diffraction116. ........................... 46

Figure 3.2 Nyquist plot of a polycrystalline material showing bulk, grain-boundary

and electrode contribution. ........................................................................................ 50

Figure 3.3 Schematic representation of the cell set up used for AC impedance

measurements for garnet samples. ............................................................................ 51

Figure 3.4 The cell set up used for AC impedance measurements for polyoxometalate

and ionic liquid based hybrid materials. ................................................................... 53

Figure 3.5 Schematic representation of cyclic voltammetry cell set up used for garnet

sample. ...................................................................................................................... 54

Figure 3.6 Schematic representation of typical TGA curve showing single step weight

loss. ........................................................................................................................... 55

Figure 4.1 The Rietveld refinement profile for x = 0.05 member of Li5+2xLa3Ta2-

xYxO12. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.

Copyright (2015) American Chemical Society.125 .................................................... 64




xiii





xYxO12. Reproduced by permission of the PCCP Owner Societies.130 ..................... 65





Figure 4.7 The garnet-type structure of Li5+2xLa3Ta2-xYxO12. (a) Ta-O sublattice

where Ta is shown in blue and oxygen in red. (b) Ta-O sublattice with La

(purple) in spaces between TaO6 units. (c) TaO6 units shown by gray octahedra.

La atoms are in purple, while Li1 (24d site) and Li2 (48g site) are shown as blue

and green spheres, respectively. Li3 (96h site) is omitted for clarity. (d) The La-

O sublattice showing LaO8 polyhedra. Reprinted with permission from Inorg.

Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical Society.125 ..... 68

Figure 4.8 Three-dimensional connectivity of oxygen octahedra and tetrahedra, which

can accommodate Li atoms, shown as blue and green spheres. In the actual

structure, not all of the neighboring sites can be simultaneously occupied by Li

and it is not possible for a Li atom to exist at an octahedral site if both tetrahedral

sites adjacent to it are occupied. Reprinted with permission from Inorg. Chem.


Figure 4.9 Typical AC impedance plot at 23 °C for x = 0.05 member of Li5+2xLa3Ta2-

xYxO12. The open symbols represent the collected data and the solid lines

represent the fitting. The inset figure indicates the equivalent circuits used for

fitting analysis. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-

6977. Copyright (2015) American Chemical Society.125 .......................................... 71











Figure 4.12 Typical AC impedance plot at 23 °C for x = 0.25 member of

Li5+2xLa3Ta2-xYxO12. The open symbols represent the collected data and the solid

xiv

lines represent the fitting. The inset figure indicates the equivalent circuits used

for fitting analysis. Reprinted with permission from Inorg. Chem. 2015, 54,

6968-6977. Copyright (2015) American Chemical Society.125 ................................ 73






Figure 4.14 Arrhenius plots showing the conductivity variation against temperature of

a) as-prepared Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Reprinted with permission

from Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical

Society.125 .................................................................................................................. 75

Figure 4.15 7 Li and MAS NMR of Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Chemical

shift was measured against solid LiCl. A spinning frequency of 5 kHz was used.



Figure 4.16 27Al MAS NMR of Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Chemical shift

was measured against Al(NO3)3. A spinning frequency of 5 kHz was used.



Figure 4.17 Morphological studies of Li5+2xLa3Ta2-xYxO12 (x = 0.05–0.75) pellets

using SEM microscopy. Reprinted with permission from Inorg. Chem. 2015, 54,


Figure 4.18 In situ PXRD for x = 0.25 member of Li5+2xLa3Ta2-xYxO12 showing the

thermal stability. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-



thermal stability. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-


Figure 4.20 In situ PXRD patterns of water-treated Li6La3Ta1,5Y0.5O12 showing the

structural stability. Reprinted with permission from Inorg. Chem. 2015, 54,


Figure 4.21 Changes in the pH during the treatment of Li5+2xLa3Ta2-xYxO12 (x = 0.50

and 0.75) in 1 M LiCl solution. Reprinted with permission from Inorg. Chem.


Figure 4.22 Impedance spectra of Li5+2xLa3Ta2-xYxO12 (x = 0.50 and 0.75) after

treatment with 1 M LiCl solution measured at 75 °C. The inset magnifies the

data for x = 0.50 sample to show the small semi-circle at the high frequency

xv

side. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


Figure 4.23 Arrhenius plots showing the conductivity variation against temperature of

Li5+2xLa3Ta2-xYxO12 (x = 0.5, and 0.75) after treatment with 1 M LiCl solution.



Figure 4.24 FTIR spectra of water-treated Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and

0.75). Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


Figure 4.25 SEM images and EDX results for Li5+2xLa3Ta2-xYxO12 (x = 0.5 and 0.75)

after treatment with 1 M LiCl solution. Reprinted with permission from Inorg.

Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical Society.125 ..... 89

Figure 4.26 PXRD of Li5+2xLa3Ta2-xYxO12 (x = 0.5 and 0.75) after reaction with 1 M

LiCl solution. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-


Figure 4.27 Cyclic voltammetry of Li6.5La3Ta1.25Y0.75O12 at 50 °C between -0.2 to +6

V vs. Li/Li+ with a scanning rate of 100 mVs-1. ....................................................... 91

Figure 4.28 Cyclic voltammetry of Li6.5La3Ta1.25Y0.75O12 at 100 °C between -0.2 to

+6 V vs. Li/Li+ with a scanning rate of 100 mVs-1. .................................................. 91

Figure 5.1 The powder X-ray diffraction patterns of Li6La3Ta1.5Y0.5O12 prepared via

solid-state synthesis method with varying amounts of excess (a) 2.5, (b) 5, (c) 7,

(d) 10 and (e) 15 wt.% LiNO3 and a calculated pattern for cubic garnet

(Li5La3Ta2O12, ICSD # 01-074-9856 where ICSD stands for Inorganic Crystal

Structure Database). Expected peak for Li2CO3 is also marked with a dotted line

at ~ 31.92° (JCPDS # 22-1141 where JCPDS stands for Joint Committee on

Powder Diffraction Standards). Reproduced by permission of ECS-The

Electrochemical Society.145 ....................................................................................... 96

Figure 5.2 Scanning electron micrographs of Li6La3Ta1.5Y0.5O12 with varying

amounts of excess (a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3

(corresponding higher magnification images are shown in the same row).

Reproduced by permission of ECS-The Electrochemical Society.145....................... 98

Figure 5.3 SEM image and corresponding elemental mapping of Li6La3Ta1.5Y0.5O12

with varying amounts of excess (a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.%

LiNO3. Reproduced by permission of ECS-The Electrochemical Society.145 ........ 100

Figure 5.4 27Al MAS NMR of Li6La3Ta1.5Y0.5O12 with varying amounts of excess (a)

2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3 LiNO3 measured against

Al(NO3)3 as standard. Reproduced by permission of ECS-The Electrochemical

Society.145 ................................................................................................................ 101

xvi

Figure 5.5 7Li MAS NMR of Li6La3Ta1.5Y0.5O12 with varying amounts of excess (a)

2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3. The standard material used was

LiCl. Reproduced by permission of ECS-The Electrochemical Society.145 ........... 102

Figure 5.6 Raman spectroscopy of Li6La3Ta1.5Y0.5O12 with varying amounts of excess

(a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3. For comparison, Raman

spectrum of commercial Li2CO3 is also shown. Reproduced by permission of

ECS-The Electrochemical Society.145 ..................................................................... 103

Figure 5.7 Typical AC impedance spectrum of Li6La3Ta1.5Y0.5O12 with 2.5 wt.%

excess LiNO3 measured in air at 50 °C. The open symbols and solid lines

represent the measured and fitted data, respectively. The equivalent circuit used

to fit the impedance data is shown in the inset figure. Reproduced by permission

of ECS-The Electrochemical Society.145 ................................................................ 105

Figure 5.8 Typical AC impedance spectrum of Li6La3Ta1.5Y0.5O12 with 5 wt.% excess

LiNO3 measured in air at 50 °C. The open symbols and solid lines represent the

measured and fitted data, respectively. The equivalent circuit used to fit the

impedance data is shown in the inset figure. Reproduced by permission of ECS-

The Electrochemical Society.145 ............................................................................. 105




impedance data is shown in the inset figure. Reproduced by permission of ECS-

The Electrochemical Society.145 ............................................................................. 106

Figure 5.10 Typical AC impedance spectrum of Li6La3Ta1.5Y0.5O12 with 10 wt.%










Figure 5.12 The Arrhenius plots of Li6La3Ta1.5Y0.5O12 (2.5 – 15 wt.% excess LiNO3)

showing the bulk Li ion conductivity as a function of temperature. The closed

and open symbols represent the heating and cooling cycles, respectively.

Reproduced by permission of ECS-The Electrochemical Society.145..................... 108

Figure 5.13 Correlation of density and activation energy (calculated at 24-325 °C)

with the excess LiNO3 wt.% (2.5 – 15 wt. %) used for the synthesis of

Li6La3Ta1.5Y0.5O12. It shows that the activation energy decreases with increase in

xvii

density of the ceramics. Reproduced by permission of ECS-The Electrochemical

Society.145 ................................................................................................................ 109

Figure 6.1 (a) The garnet-type structure viewed down the 111 axis. The TaO6

octahedra are shown in purple. Large red spheres are La atoms and small blue

and green spheres represent Li atoms on 24d and 48g sites, respectively. (b) The

red lines highlight the dodecahedral geometry where an La atom is

accommodated, (c) the octahedra, which encompass Ta atoms, (d) the tetrahedral

geometry around Li atoms on 24d positions, and (e) the octahedral geometry

where Li atoms on the 48g site (center of the octahedron) or the 96h site (shifted

from the center) are located. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365.

Reproduced by permission of the PCCP Owner Societies.130 ................................ 114

Figure 6.2 The connectivity of LiO6 and LiO4 polyhedra is shown. These polyhedra

are connected in 3-dimentions throughout the unit cell. If there is a Li atom on a

tetrahedral site, the Li atoms at the centers of neighbouring octahedra (48g site)

will be shifted to the 96h position to avoid close Li-Li contacts. If the two

octahedral sites that are on the two sides of a tetrahedral position are occupied,

then the tetrahedral site between them will be empty. Two possible migration

pathways (A and B) are shown by arrows that are discussed in details in the later

section. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by

permission of the PCCP Owner Societies.130 .......................................................... 115

Figure 6.3 (a) Complex impedance plots of samples Li5+2xLa3Ta2-xYxO12 (x = 0.25,

0.5, and 0.75) at -20 °C. The equivalent circuit used for fitting z-plots is given in

the inset. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by

permission of the PCCP Owner Societies.130 .......................................................... 116

Figure 6.4 The Arrhenius plots of conductivity Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5,

and 0.75). Inset gives the values of activation energy in these samples. Phys.

Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by permission of the

PCCP Owner Societies.130 ...................................................................................... 117

Figure 6.5 Conductivity (σ) vs. frequency spectra for Li5+2xLa3Ta2-xYxO12 (x = 0.25

and 0.5) at -50 °C. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365.


Figure 6.6 Real part of dielectric permittivity (ε/) as a function of frequency for

Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase) at various temperatures. Phys. Chem. Chem.


Societies.130 ............................................................................................................. 120


Li6La3Ta1.5Y0.5O12 (Li6-phase) at various temperatures. Phys. Chem. Chem.


Societies.130 ............................................................................................................. 121

xviii

Figure 6.8 Imaginary part of dielectric permittivity (ε//) as a function of frequency for

Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase) at various temperatures. Phys. Chem. Chem.


Societies.130 ............................................................................................................. 122


Li6La3Ta1.5Y0.5O12 (Li6-phase) at various temperatures. Phys. Chem. Chem.


Societies.130 ............................................................................................................. 122

Figure 6.10 Frequency spectra of dielectric loss tangent (tan δ) at various

temperatures in Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase). Phys. Chem. Chem. Phys.

2014, 16, 11356-11365. Reproduced by permission of the PCCP Owner

Societies.130 ............................................................................................................. 124


temperatures in Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem. Chem. Phys. 2014,

16, 11356-11365. Reproduced by permission of the PCCP Owner Societies.130 ... 124


temperatures in Li6.5La3Ta1.25Y0.75O12 (Li6.5-phase), respectively. Phys. Chem.


Owner Societies.130 ................................................................................................. 125

Figure 6.13 Arrhenius plots of dielectric loss peak frequency in Li5+2xLa3Ta2-xYxO12

(x = 0.25, 0.5, and 0.75). Phys. Chem. Chem. Phys. 2014, 16, 11356-11365.


Figure 6.14 Frequency spectra of imaginary part of electric modulus (M″) at various



Societies.130 ............................................................................................................. 127


temperatures in Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem. Chem. Phys. 2014,

16, 11356-11365. Reproduced by permission of the PCCP Owner Societies.130 ... 128




Societies.130 ............................................................................................................. 128

Figure 6.17 Modulus master curve at various temperatures in Li5.5La3Ta1.75Y0.25O12

(Li5.5-phase). Here M״max is the value of M״ at peak position in modulus spectra

(a-b) and fM-is the peak frequency. Phys. Chem. Chem. Phys. 2014, 16, 11356 ״

11365. Reproduced by permission of the PCCP Owner Societies.130 .................... 129

xix

Figure 6.18 Modulus master curve at various temperatures in Li6La3Ta1.5Y0.5O12 (Li6-

phase). Here M״max is the value of M״ at peak position in modulus spectra (a-b)

and fM .is the peak frequency. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365 ״


Figure 6.19 Arrhenius plots of relaxation time (τ) in the samples

Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase), Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem.


Owner Societies.130 ................................................................................................. 131

Figure 7.1 Typical complex impedance spectra obtained using Li+ ion blocking Au

electrodes for Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb), Li6La3Nb1.5Y0.5O12 (Li6-Nb) and

Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) at about -22 °C. The line spacing through the

data point is fitted using an equivalent circuit consisting of (RbCPEb)(CPEe) for

Li5.5-Nb and Li6-Nb and (Rb CPEb) (Rgb CPEgb) (CPEe) for Li6.5-Nb. Phys.



Figure 7.2 Arrhenius plots of bulk Li ion conductivity of Li5.5La3Nb1.75Y0.25O12

(Li5.5-Nb), Li6La3Nb1.5Y0.5O12 (Li6-Nb), and total (bulk + grain-boundary) Li+

ion conductivity of Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb). The line passing through

the data points is fitted line using Arrhenius equation. Phys. Chem. Chem. Phys.


Societies.185 ............................................................................................................. 139

Figure 7.3 The electrical conductivity of Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) as a

function of frequency at different temperatures obtained using AC impedance

spectroscopy with Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys.


Societies.185 ............................................................................................................. 141

Figure 7.4 The electrical conductivity of Li6La3Nb1.5Y0.5O12 (Li6-Nb) as a function of

frequency at different temperatures obtained using AC impedance spectroscopy

with Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys. 2016, 18, 15418-


Figure 7.5 The electrical conductivity of Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) as a

function of frequency at different temperatures obtained using AC impedance

spectroscopy with Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys.



Figure 7.6 Estimated real part of permittivity (ε/) as a function of frequency of

Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) at different temperatures from AC impedance

spectroscopy data collected using Li+ ion blocking Au electrodes. Phys. Chem.


Owner Societies.185 ................................................................................................. 143

xx


Li6La3Nb1.5Y0.5O12 (Li6-Nb) at different temperatures from AC impedance



Owner Societies. 185 ................................................................................................ 143





Owner Societies.185 ................................................................................................. 144

Figure 7.9 Imaginary part of permittivity (ε//) as a function of frequency of

Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) at different temperatures. Phys. Chem. Chem.


Societies.185 ............................................................................................................. 145


Li6La3Nb1.5Y0.5O12 (Li6-Nb) and at different temperatures. Phys. Chem. Chem.


Societies.185 ............................................................................................................. 146


Li6La3Nb1.5Y0.5O12 (Li6-Nb) and at different temperatures. Phys. Chem. Chem.


Societies.185 ............................................................................................................. 146

Figure 7.12 Dielectric loss tangent (δ) as a function of frequency of



Societies.185 ............................................................................................................. 148


Li6La3Nb1.5Y0.5O12 (Li6-Nb) at different temperatures. Phys. Chem. Chem. Phys.


Societies.185 ............................................................................................................. 148




Societies.185 ............................................................................................................. 149

Figure 7.15 Comparison of dielectric loss tangent (δ) as a function of frequency of

Li5.5-phase of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and

Ta) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


xxi


Li6-phase of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and Ta)

at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.



Li6.5-phases of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and

Ta) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


Figure 7.18 The relative permittivity of Li5.5-, Li6-, and Li6.5-phases Li5+2xLa3M2-

xYxO12 (M = Nb and Ta) at different temperatures. Phys. Chem. Chem. Phys.


Societies.185 ............................................................................................................. 152

Figure 7.19 Electric modulus M// as a function of frequency of Li5.5La3Nb1.75Y0.25O12

(Li5.5-Nb) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-


Figure 7.20 Electric modulus M// as a function of frequency of Li6La3Nb1.5Y0.5O12

(Li6-Nb) and (c) Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) at different temperatures. Phys.




(Li6.5-Nb) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-


Figure 7.22 Arrhenius plots of relaxation time of Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb),

Li6La3Nb1.5Y0.5O12 (Li6-Nb), and Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb). Phys. Chem.


Owner Societies.185 ................................................................................................. 156

Figure 8.1 Photographs of (a) heteropoly acid H4PMo11VO40 and (b) its hybrid form,

PyPs-H4PMoV derived from HPA and PyPs ionic liquid....................................... 160

Figure 8.2 Synthesis of POMs-based IL hybrid gel electrolytes showing various steps

involved. .................................................................................................................. 161

Figure 8.3 1H NMR of PyPs, PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV.

The chemical shift values are represented against TMS dissolved in D2O. ........... 163

Figure 8.4 13C NMR of PyPs, PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV.

The chemical shift values are represented against TMS dissolved in D2O. ........... 164

Figure 8.5 FTIR spectra of heteropoly acids H3PW11MoO40 (H3PWMo),

H4PMo11VO40 (H4PMoV) and H5PMo10V2O40 (H5PMoV) and their IL-hybrids

(PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV. .......................................... 166

xxii

Figure 8.6 PXRD patterns of (a) PyPs, heteropoly acids H3PW11MoO40 (H3PWMo),

H4PMo11VO40 (H4PMoV) and H5PMo10V2O40 (H5PMoV) and their PyPs-POM

hybrids (PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV) and (b) PyPs-

H3PWMo in the small angle region. ....................................................................... 168

Figure 8.7 The schematic illustration of the crystal structure of POM-IL hybrid gel

electrolyte. ............................................................................................................... 169

Figure 8.8 HTXRD patterns of PyPs-H3PWMo under different temperatures. .............. 170

Figure 8.9 HTXRD patterns PyPs-H4PMoV under different temperatures. ................... 171

Figure 8.10 (a) The TGA of heteropoly acids H3PW11MoO40 (H3PWMo) and

H4PMo11VO40 (H4PMoV) and their PyPs-POM hybrids (PyPs-H3PWMo, and

PyPs-H4PMoV); (b) The proposed decomposition steps of IL-PMo11V sample. ... 173

Figure 8.11 Typical AC impedance spectra of PyPs-H3PWMo hybrid at 23 C

measured in air atmosphere. Impedance plots zoomed at high frequency side is

shown as inset for clarity. ....................................................................................... 174


measured in air atmosphere..................................................................................... 175


measured in air atmosphere. Impedance plots zoomed at high frequency side is

shown as inset for clarity. ....................................................................................... 175

Figure 8.14 Typical AC impedance spectra of PyPs-H4PMoV hybrids at 32 C










Figure 8.19 Arrhenius plots for ionic conductivity of PyPs-POM gel electrolytes,

PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV. ........................................... 179

Figure 8.20 Comparison of electrical conductivity of PyPs-H3PWMo with other

known promising proton conductors.99, 200, 201 ........................................................ 181

xxiii

Figure 8.21 Cyclic voltammetry of PyPs-H3PWMo hybrid gel electrolyte at room

temperature between -4 to +4 V vs. Ag/AgCl with a scanning rate of 10 mVs-1. .. 182


temperature between -4 to +4 V vs. Ag/AgCl with a scanning rate of 10 mVs-1. .. 182

Figure 9.1: PXRD of as-prepared, acetonitrile and IL treated KLaNb2O7 (a) complete

patterns and (b) selected area patterns at 2θ = 8-9 ° of the same to show the shift

in first peak. The indexed hkl values for as-prepared KLaNb2O7 corresponding to

orthogonal phase is marked.203................................................................................ 185

Figure 9.2 TGA of as-prepared, acetonitrile and IL treated KLaNb2O7 in air at 5

°C/min. .................................................................................................................... 187

Figure 9.3 Scanning electron micrograph and corresponding EDS of as-prepared

KLaNb2O7. .............................................................................................................. 189

Figure 9.4 Scanning electron micrograph and corresponding EDS of KLaNb2O7

treated with acetonitrile solvent at room temperature for 24 h. .............................. 190


treated with IL. ........................................................................................................ 191


treated with a mixture of IL and acetonitrile solvent. ............................................. 192

Figure 9.7 PXRD patterns of as-prepared, acid, acetonitrile IL and acetonitrile treated

protonated KLaNb2O7. ............................................................................................ 193

Figure 9.8 PXRD (a) complete patterns of as-prepared, acetonitrile and IL treated

K2La2Ti3O10 and (b) zoomed patterns at 2θ = 5-6.5 ° of the same to show the

shift in first peak ((002) plane). The indexed hkl values for as-prepared

K2La2Ti3O10 corresponding to tetragonal phase is marked.115 ................................ 195

Figure 9.9 TGA of as-prepared, acetonitrile and IL treated K2La2Ti3O10 in air at 5

°C/min. .................................................................................................................... 198

Figure 9.10 Scanning electron micrograph and corresponding EDS of K2La2Ti3O10. ... 199

Figure 9.11 Scanning electron micrograph and corresponding EDS of K2La2Ti3O10

treated with acetonitrile. .......................................................................................... 200


treated with IL. ........................................................................................................ 201


treated with acetonitrile and IL. .............................................................................. 202

xxiv

Figure 9.14 PXRD patterns of as-prepared, acid, acetonitrile IL and acetonitrile

treated protonated K2La2Ti3O10. ............................................................................. 203

Figure 9.15 PXRD patterns of as-prepared K2La2Ti3O10 samples both fresh and 48h

exposed to air to check the stability in ambient condition. ..................................... 204

Figure 9.16 PXRD patterns of acetonitrile (AN) treated K2La2Ti3O10 samples both

fresh and 48h exposed to air to check the stability in ambient condition. .............. 205

Figure 9.17 PXRD patterns of IL treated K2La2Ti3O10 samples both fresh and 48h

exposed to air to check the stability in ambient condition. ..................................... 205

Figure 9.18 PXRD patterns of acetonitrile (AN) and IL treated K2La2Ti3O10 samples

both fresh and 48h exposed to air to check the stability in ambient condition. ...... 206

Figure 9.19 PXRD refinement profile of K2La2Ti3O10 considering the presence of

only hydrated , K2La2Ti3O10.xH2O phase. .............................................................. 207

Figure 9.20 PXRD refinement profile of acetonitrile treated K2La2Ti3O10 considering

the presence of two phases, unhydrated (K2La2Ti3O10) and hydrated

(K2La2Ti3O10.xH2O). ............................................................................................... 208

Figure 9.21 PXRD refinement profile of IL treated K2La2Ti3O10 considering the

presence of only hydrated, K2La2Ti3O10.xH2O phase. ............................................ 208

Figure 9.22 PXRD refinement profile of acetonitrile and IL treated K2La2Ti3O10

considering the presence of only hydrated, K2La2Ti3O10.xH2O phase. .................. 209

Figure 9.23 PXRD patterns of as-prepared, acetonitrile treated, pure and diluted IL

treated K2La2Ti3O10 samples after TGA measurements done in Argon up to a

heating temperature of 700 °C. ............................................................................... 210

xxv

List of Symbols and Abbreviations

Symbol Unit Definition

λ Å Wavelength

ρ Ohm cm Resistivity

σ S cm-1 Conductivity

R J K-1 mol-1 Gas constant

R Ohm Resistance

k J K-1 Boltzmann constant

NA Mol-1 Avogadro’s number

χ2 Goodness of fit

Z/ Ohm Real component of impedance

Z// Ohm Imaginary component of impedance

Z Ohm Impedance

F C mol-1 Faraday constant

E V Potential

I A Current

ƒ Hz Frequency

ω Rads s-1 Radial frequency

θ Degree Phase angle

C

εr

F Capacitance

Relative permittivity of medium

ε0

Ea

τ0

T

F cm-1

eV

s

ºC or K

Vacuum permittivity

Activation energy

Relaxation time

Temperature

Voc V Open circuit voltage

l cm Length

A cm2 Area

xxvi

AC Alternating current

CE Counter electrode

CPE Constant phase element

CV Cyclic voltammetry

DC Direct current

DEC Diethyl carbonate

DFT Density functional theory

DJ Dion-Jacobson

EC Ethylene carbonate

EDS Energy dispersive X-ray spectroscopy

EIS Electrochemical impedance spectroscopy

FTIR Fourier transform infrared spectroscopy

FWHM Full width half maximum

H3PWMo H3PW11MoO40

H4PMoV H4PMo11VO40

H5PMoV H5PMo10V2O40

HPA Heteropoly acid

IL Ionic liquid

JCPDS Joint committee on powder diffraction standards

LIB Lithium ion battery

LiRAP Li-rich anti-perovskites

LISICON Li superionic conductor

LLT Lithium lanthanum titanate

LLTF Lithium lanthanum titanate/fluoride

MAS Magic angle spinning

MIMPS N-methyl imidazolium-1-(3-sulfonic group) propyl

NASICON Sodium super ionic conductors

NMR Nuclear magnetic resonance

PC Propylene carbonate

PEG Polyethylene glycol

PEMFC Proton exchange membrane fuel cell

PEO Ploy ethylene oxide

POM Polyoxometallate

PVdF-HFP poly(vinylidene fluoride-hexa fluoropropylene)

PVP Polyvinylpyrrolidone

PXRD Powder X-ray diffraction

PyPs 3-(pyridine-1-ium-1-yl)propane-1-sulfonate

RE Reference electrode

RP Ruddlesden-Popper

SEM Scanning electron microscopy

TEAPS 1-(3-sulfonic group) trimethylamine

TEC Thermal expansion coefficient

TGA Thermo gravimetric analysis

WE Working electrode

1

Chapter One: Introduction

1.1 Project motivation

Coal, natural gas and petroleum are major fossil fuel sources, which serve the world’s

current energy and electricity demands. However, burning fossil fuels releases gases such

as carbon monoxide (CO), carbon dioxide (CO2), sulphur dioxide (SO2), and nitrogen oxide

(NOx), which has severe impact on the environment. In spite of this issue, these non-

renewable energy sources are in high demand as they are cheap compared to the alternate

energy sources leading to their quick depletion. Clean renewable energy sources such as

solar and wind power can be used to decrease carbon emissions and provide clean power.

However, fluctuation in solar and wind outputs make them unreliable and these energy

sources can be implemented by the integration of electrochemical energy storage

technologies such as rechargeable batteries (e.g. nickel-cadmium, nickel-metal hydride,

lead-acid, lithium ion, sodium-sulphur, and redox flow batteries).1 Another promising

alternative energy source is fuel cells, which use non-polluting fuel (e.g. H2) and can

electrochemically convert the fuel into clean energy.

Batteries convert the chemical energy stored in the electrode material into energy,

whereas fuel cells convert the externally fed fuel into energy as long as a fuel supply is

present. An ideal electrochemical energy storage system must provide electricity when it

is needed at proper time dimensions. The compact nature of batteries allows for easy

distribution where energy supply is needed. The discharge rate of most batteries is in

minutes and can provide a wide range of power ratings, making them more desirable among

2

other electrochemical energy storage technologies. A comparison of discharge time and

power ratings of different grid storage systems are shown in Figure 1.1.1

Figure 1.1 General comparison of discharge rates and power rating of different

electrochemical energy storage technologies. Copyright permission is obtained.1

Table 1.1 lists the energy and power characteristics of various known secondary

batteries.1 Among them, lithium ion batteries (LIBs) are better in the list due to their high

output voltage (~ 4V) and high energy densities. The light weight Li has smaller size (0.69

Å) and possesses lower redox potential (Li+/Li) of -3.04 V vs standard hydrogen electrode

SHE), which make LIBs achieve higher voltage and energy densities compared to Na and

K-based systems.2 Due to these promising properties, LIBs are considered for wide range

of applications from portable electronics to hybrid vehicles and grid storage.3 There are

several safety concerns associated with the LIB operation regardless of their promising

3

applications. The current state-of-the art liquid electrolyte, LiPF6 dissolved in organic

solvent (e.g. ethylene carbonate), grounds issues such as leakage, flammability, explosion,

short circuiting and self-discharge.

Table 1.1 Comparison of voltage, energy and power densities of various secondary

batteries.1

Battery name Voltage

(V)

Volumetric

energy density

(Whl-1)

Gravimetric

energy density

(Whkg-1)

Specific power

(Wkg-1)

Lead-acid 2.1-1.8 30-100 20-45 60-110

Nickel-cadmium 1.3-0.8 70-200 25-60 40-100

Nickel-metal hydride 1.3-0.9 75-300 65-100 70-200

Lithium ion-LiCoO2-

graphite

4.2-2.5 250-400 110-175 200-300

Sodium-sulphur 2.1-1.8 70-150 60-120 15-70

Vanadium redox flow 1.6-1.1 10-20 10-20 1-4

Solid-state electrolytes are identified as the safe substitute for organic solvent-based

electrolyte to overcome the global challenges in current LIBs. All-solid-state LIBs can

provide further advantage of easy fabrication in addition to safety. The major challenge of

solid-state electrolytes is to beat the high ionic conductivity of the liquid electrolyte (10-2

Scm-1) at operating temperature. Among the known solid Li ion electrolytes, garnet-type

metal oxides are identified as suitable electrolyte candidate for solid-state LIBs. One of the

focuses of this thesis is to understand the physical, chemical and structural stability of

garnets and their transport properties.

4

There are five major classifications of fuel cells based on the charge carrier and are

summarized in Table 1.2.4 Among the fuel cell list, proton exchange membrane fuel cells

(PEMFCs) are most desired as portable power source due to its structural compactness,

quick start-up time, high output power density, quiet operation and zero emissions.4 The

Nafion membrane electrolyte and carbon supported Pt (catalyst) electrodes are the major

components of PEMFC.

Table 1.2 Classification and the state-of-the-art of fuel cells.4

Fuel cell name Working

temperature (ºC)

Fuels and by-

products

Mobile ions

in

Electrolyte

Oxidant

Solid oxide

fuel cell

(SOFC)

500-1000 Fuel: H2 or CO

Product: H2O, CO2

O2- Input: O2

(air)

Molten-

carbonate fuel

cell (MCFC)

650 Fuel: H2 or CO

Product: H2O, CO2

CO32- Input: O2

(air) and CO2

Phosphoric

acid fuel cell

(PAFC)

200 Fuel: H2

H+ Input: O2

(air)

Proton

exchange

membrane fuel

cell (PEMFC)

80 Fuel: H2

H+ Input: O2

(air)

Alkaline fuel

cell (AFC)

70 Fuel: H2

Product: H2O

OH- Input: O2

(air)

The operating temperature is limited to < 90 °C due to the dependence of proton conduction

of the polymer membrane on hydration content, resulting in membrane degradation if the

temperature and humidity level is not maintained properly. Demanding low temperature

operation, PEMFC requires expensive Pt catalyst for the desired electrochemical catalytic

activity. However, carbon supported Pt layer is prone to CO (fuel impurity) poisoning at

5

low temperature which can even result in the structural collapse of the electrode.5

Increasing the operating temperature can avoid such electrode deterioration, along with

replacing the expensive catalyst with cheap metals such as copper (Cu), as well as replacing

Nafion with other alternatives, which require less water management. Modifications of

Nafion by inorganic additives such as silica and use of ceramic oxides such as TiO2 have

been tried to improve the working temperature of PEMFC.4 Another key focus of this thesis

is to develop metal oxide based hybrid proton conductors from polyoxometalates, layered

perovskites and ionic liquids for PEMFCs.

1.2 Research objectives

There are two primary goals for this thesis work. The first goal is a contribution towards

energy storage devices especially Li ion battery, and the second goal is a contribution

towards energy conversion devices. Solid-state lithium ion batteries are much desired

nowadays especially for the transport application, with garnet-type Li ion conductors

gaining superior interest among the known solid-state electrolytes. Deeper understanding

of the garnet-type material is adequate in order to safely and effectively implement these

materials in the commercial market. Lithium stuffing has been proved to improve the ionic

conductivity of garnets. In this thesis, we have studied the effect of lithium-stuffing and

yttrium-doping on the porosity, morphology, structure, stability, and electrical properties

of Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Apart from the Li ion battery application, we

wanted to check the feasibility of using these materials for Li-air batteries, which is one of

the emerging rechargeable battery technologies. This was done by examining the stability

of these materials in LiCl solution, used as one of the anode protectors in Li-air battery,

6

which is in contact with the electrolyte. Another interest was to understand the role of

excess lithium salt used (0 – 15 wt. %) in the synthesis of garnet-type Li6La3Ta1.5Y0.5O12

on its structural and electrical properties.

Several researchers have tried to understand the mechanism of Li+ ion movement

in garnets using techniques such as solid-state 6Li and 7Li magic angle spinning nuclear

magnetic resonance (MAS NMR), neutron diffraction, and computational techniques.

However, little attempt is made to make use of electrochemical impedance spectroscopy

(EIS) to understand the dielectric properties for exploring the transport mechanism. A

systematic study was carried out to analyze the dielectric properties of two families of

lithium stuffed Li5+2xLa3M2-xYxO12 (M = Nb, Ta) garnets using EIS, while attempting to

define its transport characteristics as a function of M-site doping. Part of this study is an

extension of my MSc thesis where we initially reported the fundamental physical properties

of Li5+2xLa3Nb2-xYxO12.6

As the second goal of this thesis work, novel hybrid proton conductors were

developed based on ionic liquids (ILs) and metal oxides. The idea was to combine the

unique properties of ionic liquids and highly robust structured polyoxometalates

(heteropoly acids) or layered perovskites. The key objectives of this second goal of the

research are (i) to synthesize the 3-(pyridin-1-ium-1-yl)propane-1-sulfonate (PyPs) ionic

liquid and H3PW11MoO40 (H3PWMo), H4PMo11VO40 (H4PMoV) and H5PMo10V2O40

(H5PMoV) heteropoly acids and synthesize hybrid proton conductors derived from PyPs

and H3PWMo, H4PMoV and H5PMoV and understand its structural and electrical

properties; (ii) to synthesize the layered perovskite metal oxides such as KLaNb2O7 and

K2La2Ti3O10 and their proton analogues as the precursors for hybrid proton conductors and

7

(iii) to develop new proton conducting hybrid electrolytes by ion exchange reaction of as-

prepared and proton exchanged KLaNb2O7 and K2La2Ti3O10 with imidazolium based IL to

better understand their intercalation chemistry, and structural properties.

1.3 Thesis organization

The thesis is organized into 10 chapters. Chapter 1 describes the project overview,

motivation and organization of this thesis. Chapter 2 provides the background information

of lithium ion batteries (LIBs) and proton exchange membrane fuel cells (PEMFCs), and a

brief review of electrolyte materials used for LIBs and PEMFCs with the emphasis on

garnet-type Li ion conductors and metal oxide based hybrid proton conductors. Two review

articles were published in Chemical Society Reviews (Thangadurai, V.; Narayanan, S.;

Pinzaru, D. Chem. Soc. Rev. 2014, 43, 4714-4727) and the Journal of Physical Chemistry

Letters (Thangadurai, V.; Pinzaru, D.; Narayanan, S.; Baral, A. K. J. Phys. Chem. Lett.

2015, 6, 292-299) based on recent research progress of the solid-state Li ion conducting

garnet-type electrolytes for LIB application. Chapter 3 discusses the synthesis methods

and characterization techniques used in this thesis work.

There are six results chapters included in this thesis and each chapter has a general

overview, results and discussion, and conclusion section. Chapter 4 reports the systematic

study of the effect of doping of Y and Li on the porosity and conductivity of garnet-type

Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) and their stability in LiCl solution. These results are

published in Inorganic Chemistry (Narayanan, S.; Ramezanipour, F.; Thangadurai, V.

Inorg. Chem. 2015, 54, 6968-6977). Chapter 5 describes the extensive study of the

influence of excess lithium salt used during the synthesis process of Li6La3Ta1.5Y0.5O12

8

garnets on the structural and electrical properties which has been published in the Journal

of Electrochemical Society (Narayanan, S.; Hitz, G. T.; Wachsman, E. D.; Thangadurai,

V. J. Electrochem. Soc. 2015, 162, A1772-A1777). Chapter 6 and 7 report the detailed

analysis of fundamental transport properties of garnet-type Li5+2xLa3Ta2-xYxO12 and

Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75) using impedance spectroscopy and dielectric

spectroscopy. These results are published in Physical Chemistry and Chemical Physics

(Baral, A. K.; Narayanan, S.; Ramezanipour, F.; Thangadurai, V. Phys. Chem. Chem. Phys.

2014, 16, 11356-11365; and Narayanan, S.; Baral, A. K; Thangadurai, V. Phys. Chem.

Chem. Phys. 2016, 18, 15418-15426).

Proton conducting hybrid materials developed from polyoxometalates and ionic

liquid are reported in Chapter 8. The study was focused to understand the structural and

electrical aspects of the hybrid materials and the results are submitted to Journal of

Materials Chemistry A (Narayanan, S.; Tong, X; Thangadurai, V. J. Mater. Chem. A 2016

(under review)). Chapter 9 reports the potential proton conducting hybrids derived from

the layered perovskites, members of Dion-Jacobson – DJ phase (KLaNb2O7) and

Ruddlesden-Popper – RP phase (K2La2Ti3O10) and an imidazolium based commercial ionic

liquid through ion-exchange process. Initial attempts to understand the ion-exchange

chemistry were done using structural tools such as powder X-ray diffraction pattern

(PXRD), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy

(EDS) and thermogravimetric analysis (TGA). Finally, Chapter 10 summarizes the

conclusions of this thesis and describes the future direction of the research.

9

Chapter Two: Background

2.1 Introduction

The world’s energy demand is rapidly increasing because of population growth and

industrialization. Fossil fuels are the main power source which is currently serving the

global energy demand. However, the usage of hydrocarbon fuels such as oil and natural

gas cause greenhouse gas (GHG) emissions, which result in environmental pollution,

global warming and climate change. The renewable energy resources such as solar, wind,

and hydro power are alternate energy sources but their major flaw is unreliability,

especially for solar and wind energy. As a result, there is a need for the development of

highly efficient, pollution free, zero GHG emission, energy conversion and storage devices.

In this regard, electrochemical energy storage and conversion devices such as batteries and

fuel cells are being considered for energy storage and conversion applications, respectively.

Among the known rechargeable batteries such as lead-acid, nickel-cadmium, and

nickel-metal hydride batteries, Li ion batteries (LIBs) possess high energy densities

(volumetric - 650 WhL-1 and gravimetric - 150 Whkg-1) and have versatile applications

ranging from portable electronics to hybrid-electric vehicles and grid storage.3, 7 They are

one of the most accepted candidates for energy storage for renewable energy systems such

as solar power and wind energy, which are quite unreliable. However, the state-of-the-art

organic solvent based electrolyte (e.g., LiPF6 dissolved in ethylene carbonate) in LIBs

cause some safety concerns due to flammability, leakage, and short circuiting issues.8, 9

Significant LIB research is being focused to replace liquid electrolyte with ceramic (solid

state) electrolyte to produce all-solid-state LIBs. The potential advantages of all-solid-state

batteries over conventional models are safety and ease of fabrication.

10

One of the promising energy conversion devices is proton exchange membrane fuel

cell (PEMFC), which works at low temperature and uses proton conducting Nafion

polymer membrane as an electrolyte. In Nafion, the proton conduction is based on the

hydration content in the polymer membrane and hence the current PEMFCs work at lower

temperatures (< 90 °C).5 At low temperature, the activity of Pt catalyst at the electrode is

reduced by CO poisoning, reducing the performance of PEMFC.5 Thus, there is a need for

new metal oxide-based proton conductors which can replace the Nafion membrane and the

use of expensive catalyst, making it possible to facilitate the cell operation at elevated

temperatures, ≥ 100 °C with less water management. In the following sections, overview

of LIBs and PEMFCs are briefly presented.

2.2 Lithium ion batteries (LIBs)

A battery is an electrochemical energy storage device, which converts the chemical energy

into electrical energy through electrochemical redox reactions. A battery includes three key

components, anode (negative electrode), cathode (positive electrode) and electrolyte (ionic

conductor). The major classification of batteries include non-rechargeable (primary) and

rechargeable (secondary). Primary batteries are designed to be used once and are then

disposable of, for e.g., alkaline and zinc-carbon batteries. Secondary batteries are being

designed to be used in multiple charge-discharge cycles and are recharged using electricity.

The examples for secondary batteries are LIBs, lead-acid and nickel-cadmium batteries.

Lithium or sodium-sulphur and redox flow batteries are other rechargeable battery

technologies which have recently gained much attention.10, 11 An overview of different

batteries and their applications is summarized in Table 2.1.12 Usually the term battery is

referred when different units of parallel and/or series connection of electrochemical cells

11

are combined according to the energy demand for application. Each cell comprised of

positive and negative electrode separated by an electrolyte.9 A battery is ranked according

to the amount of energy available in a given volume or mass, namely volumetric (WhL-1)

and gravimetric (Whkg-1) energy density, respectively. A comparison of relative energy

densities of LIBs with different commercially available secondary batteries are

summarized in Figure 2.1.3 It is shown that the higher the volumetric and gravimetric

densities the lower the size and weight of the batteries and Li ion batteries possess the

highest energy densities.

Table 2.1 Summary of different types of batteries and their applications.12

Type Name Cell Voltage (V) Applications

Primary Zinc-carbon Zn/ Zn,NH4Cl/

C,MnO2

1.5 flashlights, toys

Alkaline Zn/KOH/MnO2 1.5 radios, digital

cameras

Lithium Li/LiClO4/MnO

2 1.5 calculators, clocks

Silver-oxide Zn/KOH/Ag2O 1.55 hearing aids, watches

Zinc air Zn/KOH/O2 1.65 photoflash, cameras

Secondary Lead-acid Pb/aq. H2SO

4/

PbO2

2.1 automobiles,

headlamps

Nickel-

cadmium

Cd/KOH/NiO 1.2 portable electronics,

toys

Nickel-metal

hydride Ni alloy/OH

-

/

NiOOH

1.2 flashlights, electric

vehicles

Lithium ion Graphite/ LiPF6/

LiCoO2

3.7 cell phones, laptops,

hybrid cars

The first commercial launch of LIBs was done by Sony Corporation in 1991 and different

types of electrodes and Li ion electrolytes were tried in the battery fabrication thereafter.13

12

State-of-art LIBs use LiPF6 dissolved in organic solvents such as ethylene carbonate

electrolytes, graphite anode, and LiCoO2 cathode and provide a highest working voltage of

about 4 V vs Li/Li+.12

Figure 2.1 Comparison of volumetric and gravimetric energy densities of various

commercially available secondary batteries.1

2.2.1 Operating principle of LIBs

A schematic representation of the working principle of a LIB is represented in Figure 2.2.

Anode and cathode are supported on current collectors which are externally connected by

wire for the passage of electron during battery operation. While charging, an external

power supply is connected and Li ions move from cathode (LiCoO2) to anode (graphite)

Pd-acid Ni-Cd Ni-MH Li ion Na-S V-redox flow0

100

200

300

400 Volumetric

Gravimetric

En

erg

y d

en

sity (

Wh

L-1 o

r W

hkg

-1)

13

through the electrolyte while electrons move through the external circuit (as the oxidation

happens at cathode) to the anode to form LixC6 alloy (reduction). While discharging, Li

ions and electrons move from anode (LixC6) (oxidation) to cathode and the reverse reaction

happens at the cathode (reduction). The corresponding electrochemical reactions at the

electrodes and overall cell reactions are provided in equations 2.1 - 2.3. The electrodes in

LIBs allow the movement of Li ions back and forth during charging and discharging

processes are known as intercalation compounds. LIBs are also known as rocking-chair

batteries as the Li ions are intercalated or de-intercalated from the interstitial space in the

crystal structure of electrodes through the electrolyte during its charge/discharge cycle.

Figure 2.2 Operating principle of a LIB using LiCoO2 cathode and graphite anode.

Anode:Charge+ -

6 x 6Discharge

C xLi xe Li C (2.1)

14

Cathode:Charge + -

2 1-x 2Discharge

LiCoO Li CoO + xLi xe (2.2)

Overall:Charge

2 6 1-x 2 x 6Discharge

LiCoO + C Li CoO + Li C (2.3)

2.2.2 Challenges of current LIB research

The main challenge of LIB research is to improve the volumetric and gravimetric energy

densities for transport and grid scale applications. The appropriate choice of cathode, anode

and electrolyte must be determined in order to achieve a lightweight, small LIB with

maximum energy capacity that meets the requirements for applications such as

transportation. Researchers are focused on optimizing the properties of each battery

components to achieve this goal. One of the particular interests of this thesis is to address

the challenges of LIB electrolytes. There are some safety concerns with the usage of liquid

electrolytes due to flammability, leakage, and short circuiting issues.8, 9 Also, their

electrochemical stability window and range of operation temperature are limited to below

5 V/Li.14, 15 Hence, there is a need to replace liquid electrolyte with ceramic electrolyte to

produce all-solid-state LIBs. The potential advantages of all-solid-state batteries over

conventional ones are safety and ease of fabrication.

2.3 Electrolytes for LIBs

The current issue of solid-electrolyte-interface (SEI) formation with the use of polymer

electrolytes should be eliminated completely, at the same time the material should possess

the basic requirements of an electrolyte. There are several criteria that need to be satisfied

for a Li ion conductor to be used as an electrolyte for batteries. Which include high ionic

conductivity, negligible electronic conductivity, chemical stability under battery operation

15

conditions, inert to battery components, good thermal expansion coefficient match and low

cost.

The electrolytes used in LIBs are classified into liquid, polymer and solid

electrolytes. The liquid electrolyte comprises of a lithium salt such as lithium

hexafluorophosphate (LiPF6) dissolved in organic solvents such as ethylene carbonates

(EC) or propylene carbonate (PC) and they possess the highest ionic conductivity of ~ 10-

2 Scm-1 at ambient conditions. However, they are prone to safety hazards such as explosions

and leakage threats as mentioned earlier. Polymer electrolytes provide safety compared to

liquid electrolytes where Li salt (e.g., LiPF6, LiClO4) is dissolved in heavy polymers such

as polyethylene oxides (PEO), but their conductivity is compromised (10-8 Scm-1 at room

temperature). Incorporation of organic solvents into the polymer matrix has proven to

increase the ionic conductivity to ~ 10-3 Scm-1 at room temperature.12, 16. These electrolytes

are called gel polymer electrolytes. A widely studied composition of this type is

poly(vinylidene fluoride-hexa fluoropropylene) (PVdF-HFP) in combination with LiClO4

and propylene carbonate (PC) and diethyl carbonate (DEC). However, the gel polymers

have the major disadvantages of poor mechanical strength and compatibility issues with

the Li electrode.17, 18

Several solid-state materials such as Li silicates (e.g. Li4SiO4, σ57 °C = 10-6 Scm-1,

Ea = 0.83 eV), Li phosphorous oxy-nitrides (e.g. Li2.88PO3.73N0.14, σ25 °C = 2.3 x 10-6 Scm-1,

Ea = 0.55 eV), Na superionic conducting (NASICON) structured phosphates (e.g.

LiTi2P3O12, σ25 °C = 2 x 10-6 Scm-1, Ea = 0.30 eV), and A-site deficient perovskite-type

oxides (e.g. La(2/3)-xLi3x□(1/3)-2xTiO3 (where □ represents vacancy), σ25 °C = 10-3 Scm-1, Ea =

0.40 eV) were studied as electrolytes for LIBs.19-21 As the focus of the thesis is on solid-

16

state electrolytes, the following section describes some potential solid-state Li ion

conducting electrolytes. Figure 2.3 summarizes the conductivity behavior of representative

members of potential liquid, polymer and solid-state electrolytes.12, 16, 19-23

Figure 2.3 Li ion conductivity of various electrolytes including liquid, polymer and

solid-state electrolytes.12, 16, 19-23

2.4 Solid-state Li ion electrolytes

In order to be applicable to practical applications, a solid-state electrolyte should meet the

criteria of having high total (bulk + grain-boundary) conductivity of 10-2 Scm-1 at working

temperature with Li transference number (tLi+) ~ unity over the range of Li activities (aLi)

from anode to cathode, high chemical decomposition voltage (6V vs Li/Li+) with the anode

1 2 3 4 5 6-8

-7

-6

-5

-4

-3

-2

-1

0

Gel polymer

(1M LiPF6 in EC/DMC 50:50 vol %

+ PVdF/HFP 10 wt %)

Garnet-type

Li6.4

La3Zr

1.4Ta

0.6O

12

LISICON (Li14

ZnGe4O

16)

NASICON

(Li1.3

Ti1.7

Al0.3

(PO4)

3)

Li3N

LIPON (Li2.9

PO3.3

N0.46

)

A-site deficient perovskite

(Li0.34

La0.51

TiO2.94

)

Liquid electrolyte

(LiPF6 in EC/DC

50:50 vol %)

Thio-LISICON (Li10

GeP2S

12)

log

10 (

Scm

-1)

1000/T (K-1)

700

600

500

400

300

200

100

0 -100

T (oC)

17

and cathode during battery manufacturing and operation, negligible solid-state-interface

resistance, low cost and should be ecofriendly.24-27

2.4.1 LISICON and NASICON-structured ceramic materials

LISICON stands for “Li superionic conductor” which has γ-Li3PO4 phase represented by

a general formula, Li2+2xZn1-xGeO4 (x = -0.36 to 0.87). The crystal structure of LISICON

is provided in Figure 2.4. It has a high conductivity of 0.125 Scm-1 at 300 °C but is weakly

conducting at room temperature (~ 10-7 Scm-1).24 However, new class of thio-LISICON

was discovered (Li4-xGe1-xPxS4, where x = 0.75) with great improvements in conductivity

of 2.2 x 10-3 Scm-1 at 25 °C.28 Recently reported Li10GeP2S12 is found to be an extremely

conducting LISICON with a conductivity of 12 mScm-1 at room temperature.22, 29, 30

Despite of their high ionic conductivity, LISICON and its derivative compounds suffer

from high reactivity towards atmospheric CO2 and its ionic conductivity and performance

degrades over time.

Figure 2.4 Idealized crystal structure of LISICON with formula Li3Zn0.5GeO4.

Reproduced by permission of The Royal Society of Chemistry.14

18

The term NASICON means “sodium super ionic conductors” and these materials

are comparatively cheap and safe to handle. Lithium-based NASICON materials attain

quite some attention because of their high Li+ conductivity which is due to their framework

structure. The general structural formula of NASICON-type Li compounds is LiM2(PO4)3

(M = Zr, Ti, Ge) which consists of MO6 octahedra and PO4 tetrahedra, and are connected

three dimensionally (3D) through edge sharing.31 The crystal structure is shown in Figure

2.5. The aliovalent doping at M-site has proven to increase the ionic conductivity of these

type of conductors. Li1.3Al0.3Ti1.7(PO4)3 is the best Li ion conducting NASICON with a

high conductivity of 10-3 Scm-1 at 25 °C due to stabilization of the rhombohedral crystal

structure.32

Figure 2.5 Idealized crystal structure of NASICON with formula LiM2(PO4)3.

Reproduced by permission of The Royal Society of Chemistry.14

19

However, the facile reduction of Ti4+ to Ti3+ in the presence of Li metal is

problematic with the NASICON structured Li electrolytes. In order to solve this issue,

several substitution studies for Ti have been taken place. Zr, Ge and Hf doping in

NASICON has shown a Li+ conductivity of 10-4 Scm-1 at 25 °C but is expensive to

produce.33, 34 The substitution with Sb did not improve the conductivity.35 Doping at M site

of Li1+xTi2-xMx(PO4) (M = Al, Ga, In, Sc) has resulted in good Li ion conducting

NASICON.32, 36 Temperature dependence of the phase morphology of LiM2(PO4)3 (M =Ti,

InNb, Zr) has been studied by Pinus et al.36 and Ti member did not show any morphological

change whereas the InNb member showed phase change at higher temperature as it shows

thermal hysteresis during heating and cooling. The change of morphology with respect to

temperature has also been studied by Salkus et al. for Li2.9Sc1.9-yYyZr0.1(PO4)3 and found

that a phase change occurs from monoclinic to orthorhombic when the temperature changes

from room temperature to 600-900 K.37 The substitution by Ca for Zr in LiZr2(PO4)3 was

found to stabilize the rhombohedral structure and also increased the bulk ionic conductivity

to 10-4 Scm-1.38 Figure 2.6 shows the conductivity data of some important Li+ ion

conducting NASICON structure compounds.32, 36, 39, 40

20

Figure 2.6 Arrhenius plots of recently studied NASICONs.32, 36, 39, 40

2.4.2 A-site deficient perovskites

The perovskites are represented by a general formula, ABO3 where A and B represent

twelve and six-fold coordinated cations, respectively. The crystal structure of A-site

deficient perovskite with composition, La(2/3)-xLi3xTiO3 (LLTO) is shown in Figure 2.7.

The highest conducting perovskite La(2/3)-xLi3x□(1/3)-2xTiO3 (x = 0.1) shows bulk Li+ ion

conductivity of the order of 10-3 Scm-1 at room temperature, which is about two orders of

magnitude higher than that of grain-boundary (σgb) conductivity.26 The addition of 5 %

SiO2 to LLTO has improved the total conductivity by minimizing the grain-boundary effect

1.5 2.0 2.5 3.0 3.5 4.0

-9

-8

-7

-6

-5

-4

-3

-2

-1

LiGe2(PO

4)

3

log

10 (

Scm

-1)

1000/T(K-1)

LiInNb(PO4)

3

Li1.95

Mn0.10

Ti1.95

(PO4)

3

Li1.2

Zr1.9

Ca0.1

(PO4)

3

LiTi2(PO

4)

3

400

300

200

100

0

T (oC)

21

as it forms amorphous LiSiO4 at the grain boundary. This enhances the grain-grain contact

and further helps to improve the conductivity.41, 42

Figure 2.7 Crystal structure of A-site deficient perovskite-type La(2/3)-xLi3xTiO3 with

general formula ABO3. Reproduced by permission of The Royal Society of

Chemistry.14

The Li+ conduction in LLTO occurs via the A-site vacancy mechanism.21 The

square planar bottleneck allows the Li+ movement through a vacancy mechanism. Also,

the tilting of BO6/TiO6 helps to facilitate the Li+ mobility in the perovskite structure.

Several substitution studies have been performed at the A-site to improve the conductivity.

Doping of Na has lowered the conductivity in Li0.5-xNaxLa0.5TiO3 (0 ≤ x ≤ 0.5).43, 44 The

larger Na+ ions reduce the A-site vacancies and an increase in Na concentration blocks the

mobile Li+. The conductivity falls from 10-3 Scm-1 to 10-10 Scm-1 when the composition

reaches the percolation threshold (x = 0.2). Tailoring the concentrations of Li+ vacancies

or changing the local environment around the mobile Li+ has caused an enhancement in

Li+ ion conductivity of perovskites containing fluoride ion over their oxide ion containing

22

counter-parts. A high Li conductivity of 2.3 x 10-3 Scm-1 at 30 °C was observed for La0.56-

yLi0.33TiO3-3yF3y (LLTF) with y = 0.017.45, 46

Figure 2.8 Temperature dependence of total conductivity for some A-site deficient

perovskites.21, 27, 41, 44

A-site deficient LLT materials are not free of flaws. The major issue involves the

requirement of high sintering temperatures. The lithium oxide volatilizes at temperatures

> 900 °C and which in turn causes a variable Li content in the formula. The facile reduction

of Ti4+ in presence of Li metal restricts its use as solid electrolyte.

2.4.3 Li-rich anti-perovskites (LiRAP)

Recently, a new class of perovskites named ‘electronically-inverted’ anti-perovskites has

been developed by Zhao et al.47 The suggested general formula of these compounds is

2 3 4 5 6-10

-8

-6

-4

-2

0

LLTO + 5 % SiO2

Li0.5-x

NaxLa

0.5TiO

3 (x=0.20)

Li0.34

La0.51

TiO2.94

(LLT)

log

10 (

Scm

-1)

1000/T (K-1)

500

400

300

200

100

0 -100

T (oC)

23

X+3B

2−A− where X is the highly electropositive element, and A and B are the monovalent

and divalent anions, respectively. The compounds Li3OCl, Li3OBr and Li3O(Cl,Br) are

highly promising as they can be chemically, electronically, or structurally modified to

create more channels in the structure and can further enhance the Li+ transport. The

conductivity of mixed member (Li3O(Cl,Br)) is higher compared to end members (Li3OBr

and Li3OCl) because of its topological nature as it accommodates both Br- and Cl- in the

structure. Br- anion is large and fills the very small dodecahedral site completely and it

causes insufficient A-site filling and these results in lower conductivity in Li3OBr. In

contrast, the smaller Cl- anion is not large enough to substantially fill the dodecahedral site,

which causes structural distortion and shows low conductivity.

The crystal structure of Li3O(Cl,Br) is shown in Figure 2.9.47 LiRAP shows highest

Li ion conductivity of greater than 10-3 Scm-1 at room temperature. A super ionic

conductivity of 10 Scm-1 is observed for Li3O(Cl,Br) at 266 °C, which is comparable with

that of high conducting electrolytes such as Li7La3Zr2O12, thio-LISICON (Li10GeP2S12)

and some liquid electrolytes.22, 48 More perovskites can be derived by playing around with

the above mentioned factors and can result in the formation of different perovskites such

as Li3OA1−zA′z, Li3−xMx/2OA1−zA′z and Li3−x−δMx/2O(A1−zA′z)1−δ. These perovskites

provide easy and fast Li+ hopping for fast lithium transport. LiRAP has very promising

properties such as high ionic conductivity, poor electronic conductivity, and wide potential

window; these materials can contribute to a new era in the solid state Li electrolyte

research.

24

Figure 2.9 Crystal structure of Li-rich anti-perovskites (LiRAP) with nominal

formula of Li3O(Cl,Br). Adapted with permission from Zhao, Y.; Daemen, L. L., J.

Amer. Chem. Soc. 2012, 134, 15042-15047. Copyright (2012) American Chemical

Society.47

2.4.4 Garnet-type Li ion conductors

Ideal garnet structure has a general formula, A3B2C3O12, where A (Ca2+, Mg2+, or Fe2+), B

(Al3+, Cr3+, or Fe3+) and C (Si4+, Ge4+, or Al4+) cations occupy the eight, six, and four

coordination sites respectively, in the structure, and stabilize the face centered cubic

structure with a space group Ia-3d. The availability of interstitial space and cations with

different coordination numbers allows for various substitutions in the structure. The crystal

structure of ideal garnet is shown in Figure 2.10.

25

Figure 2.10 The crystal structure of ideal garnet structure with general formula

A3B2C3O12. Reprinted with permission from Thangadurai, V.; Pinzaru, D.;

Narayanan, S.; Baral, A. K., J. Phys. Chem. Lett. 2015, 6, 292-299. Copyright (2015)

American Chemical Society.15

Garnet-type lithium containing oxides, Ln3M2Li3O12 (M = Te, W; Ln = lanthanides) were

first reported by Kasper in 1969 whereas the garnet-type Li5La3M2O12 (M = Nb, Ta) as an

electrolyte material was first reported by Thangadurai et al. in 2003.49, 50 The idealized

garnet-type crystal structure of Li5La3M2O12 is shown in Figure 2.11. The space group (Ia-

3d) of garnet-like cubic Li5La3M2O12 was first reported by Mazza et al. in 1988.51 Detailed

neutron diffraction analysis by Cussen et al. confirmed Ia-3d space group and identified

that the Li occupancy at tetrahedral (Td) and distorted octahedral (Oh) sites are about 2/3

and 1/3 filled, respectively.52 The lattice parameters calculated were a = 12.762 and 12.766

Å for the Nb and Ta member, respectively.

26

Figure 2.11 The crystal structure of garnet-like Li-stuffed Li5La3M2O12.

The introduction of Li5La3M2O12 as a promising solid electrolyte material drastically

enhanced the research in the LIB field.50 Several substitution studies have been carried out

to enhance the electrical conductivity of Li-stuffed Li5La3M2O12. Brief discussion on this

research classifies them into three different families, Li5-phase, Li6-phase, and Li7-phase

and is provided in the following sections 2.4.4.1-2.4.4.3.

2.4.4.1 Li5- garnet family

The parent compound of Li5-family is Li5La3M2O12 (M = Nb, Ta) which shows an ionic

conductivity of 10-6 Scm-1 at room temperature.50 Substitution at the M sites have been

performed using cations such as Bi, Sb, Te, Sc, In, Y, Sm and Gd to increase the ionic

conductivity.53-59 An increase in conductivity of one order was observed with Bi doping

due to an increase in cell parameter whereas no improvement in the conductivity was

27

observed for Sb doping. Doping with larger ions like In3+ (rCN6 = 0.79 Å) for Nb (rCN6 =

0.64 Å) has improved the ionic conductivity (to 10-4 Scm-1 at 50 °C) with expansion of

lattice parameter.60 Incorporation of excess lithium in the system as well as doping of Sc

resulted in the fast ion conducting Li5+2xLa3Nb2-xScxO12 garnets with highest bulk ionic

conductivity of 6 x 10-4 Scm-1 at 50 °C for the x = 0.75 member.61 Similarly a conductivity

of 2.7 x 10-4 Scm-1 was observed for the x = 0.75 member of Li5+2xLa3Nb2-xYxO12 at room

temperature.57 No improvement in the conductivity was found except for the substitution

at Nb site of Li5La3Nb2O12 with Y.57 However, the Li- stuffed Li5+2xLa3Nb2-xMxO12 (M =

Sm, Gd) showed a conductivity in the order of 10-5 Scm-1 at room temperature which is

only one order of magnitude higher than that of parent, Li5La3Nb2O12 compound at the

same temperature.58, 59

2.4.4.2 Li6-garnet family

The Li6-family of compounds was designed by the partial substitution at the La site

(trivalent cation) with alkaline earth metals (divalent cations). As a result the lithium

stuffing increased the occupancy of Li at the Oh position and decreased the same in Td

position. The increase in concentration of Oh Li has helped to improve the ionic

conductivity of garnets. Among a series of compounds having nominal composition,

Li6ALa2M2O12 (A = Ca, Sr, Ba, Mg), only Li6BaLa2Ta2O12 showed a high conductivity of

10-5 Scm-1 at 22 °C and a high electrochemical stability of 6V of Li/Li+ at 23 °C.62-64 It is

also shown that the lattice constant of Li6ALa2M2O12 also increased with an increase in

sintering temperature. The possibility of improving the conductivity was very much

decreased with attempts using Sb and Bi dopants.54, 65 Doping at both La and Ta sites has

increased the conductivity of Li6.5La2.5Ba0.5ZrTaO12 10-4 Scm-1 at room temperature.66, 67

28

2.4.4.3 Li7-garnet family

It has been proved that stuffing more Li into the garnet structure can improve the Li+

conductivity significantly.48 A complete substitution with Zr at M of Li5La3M2O12 has

resulted in Li-stuffed Li7La3Zr2O12. Murugan et al. have synthesized cubic Li7La3Zr2O12 at

1230 °C employing solid-state synthesis method which showed the highest conductivity of

5.11 x 10-4 Scm-1 at 25 °C.48 The polymorph of Li7La3Zr2O12 which is stabilized in

tetragonal phase (synthesized at 980 °C) shows a conductivity two orders of magnitude

lower than that of the cubic phase.68 Recent studies show that doping with Al, Si, Ga, In,

and Ge affects the conductivity of Li7La3Zr2O12.69-71 A decrease in Li+ conductivity was

observed with In doping, but Al, Ga, and Ge increased the Li+ conductivity compared to

the un-doped Li7La3Zr2O12. It was proved that a small amount of Al helps to stabilize the

cubic phase and hence enhances the ionic conductivity.72 Addition of Al2O3 as a sintering

aid has also helped decrease the sintering temperature by 230 °C while still maintain the

high conductivity of the order of 10-4 Scm-1 at room temperature.73

Percival et al. have studied the temperature effect on phase transition of tetragonal

Li7La3Sn2O12 which was synthesized at 900 °C using solid-state synthesis.74 It was found

that the tetragonal phase undergoes a reversible transition to cubic phase at 750 °C and

quenching has resulted in the stabilization of original tetragonal phase. The conductivity

was found to be four orders of magnitude lower than that of Li7La3Zr2O12 and the reason

could be due to the smaller ionic radius (rCN6) of Sn4+ (0.69 Å) compared to Zr4+ (0.72 Å).74

Two polymorphs, both tetragonal and cubic phases of Li7La3Hf2O12 were also reported

with distinct ionic conductivities.75, 76 The tetragonal phase formed at a synthesis

temperature of 980 °C and had a room temperature conductivity of 10-7 Scm-1. However,

29

the cubic structure was formed at a higher temperature of 1250 °C and the conductivity

was 2.11 x 10-4 Scm-1 at room temperature. The substitution of Zr in Li7La3Zr2O12 was

carried out using pentavalent cations such as Ta5+ and Nb5+ and resulted in the composition

of Li7-xLa3Zr2-xMxO12 (M = Nb, Ta).23, 77-79 It has significantly improved the ionic

conductivity of garnets among which 0.3 ≤ x ≤ 0.6 of Ta member 23, 79 and x = 0.25 of Nb77

showed a high conductivity of ~ 10-3 Scm-1 at room temperature. The doping of hexavalent

W and Te at Zr site in the Li7 parent phase, Li7La3Zr2O12 resulted in a composition with a

general formula of Li7-xLa3Zr2-xMxO12 (M = W, Te)80, 81. The compositions

Li6.4La3Zr1.7W0.3O12 and Li6.5La3Zr1.75Te0.25O12 showed a very good conductivity of 7.89 x

10-4 and 1.02 x 10-3 Scm-1 at 30 °C. Comparison of conductivity of some representative

garnet-type Li ion conductors from Li5, Li6, and Li7 families are presented in Figure 2.12.14,

78

30

Figure 2.12 Conductivity comparison of some representative garnet-type Li ion

conductors from Li5-, Li6-, and Li7-families.14, 78

2.4.4.4 Li ion occupancy and conductivity of Li5-, Li6- and Li7-garnets

Tetrahedral (24d), octahedral (48g) and distorted octahedral (96h) sites are three different

crystallographic sites in which Li+ ions in Li-stuffed garnets mainly reside, they are shown

in Figure 2.13.14 It is clearly seen that the occupancy at octahedral site as well as Li+

conductivity increases with an increase in Li content of the garnet and several studies have

suggested that the Li+ ion conductivity as well as the conduction mechanism is governed

by the occupancy of Li at these sites.15 There are two routes of Li ion migration, Route A

and Route B suggested by Xu et al. who used computational studies.82 In Route A, the Li+

1.5 2.0 2.5 3.0 3.5-6

-5

-4

-3

-2

-1

0

log

10 (

Scm

-1)

1000/T(K-1)

Li5La

3Nb

2O

12

Li5La

3Ta

2O

12

Li5.5

La3Nb

1.75In

0.25O

12

Li6BaLa

2Ta

2O

12

c-Li7La

3Zr

2O

12

t-Li7La

3Zr

2O

12

Li6.4

La3Zr

1.4Ta

0.6O

12

350

300

250

200

150

100

50

0

T (oC)

31

ion migration occurs between two neighboring octahedral sites bypassing the common

tetrahedral site, this is common in Li5-family of garnets. In route B, the Li+ ions jump from

one octahedral site to the vacant octahedral site through the tetrahedral sites, this is

common in highly Li-stuffed garnets such as Li7- family of compounds. The variation of

Li ion occupancy at Oh and Td sites of different garnets along with their ionic conductivity

at 25 °C is shown in Figure 2.14 and it suggests that the Li+ ion conduction is highly

dependent on the Oh site.

Figure 2.13 Representation of possible Li-occupancy at different crystallographic

sites in Li-stuffed garnets. Reproduced by permission of The Royal Society of

Chemistry.14

32

Figure 2.14 Correlation of Li-site occupancy and ionic conductivity towards the Li

content in the garnets. Reproduced by permission of The Royal Society of

Chemistry.14

There is still a dispute regarding the mechanism of Li-ion conduction in garnets and

is one that needs to be thoroughly understood before these materials are used for practical

applications such as electrolytes for solid-state batteries. One of the focuses of this thesis

is to thoroughly study Li5+2xLa3Ta2-xYxO12 (x = 0.05-0.75) garnets in terms of effect of Li

content and also to understand the Li-ion conduction mechanism using AC impedance

33

spectroscopy. This work is an extension of my MSc thesis where I developed different

garnet-type metal oxide electrolytes for all-solid-state LIBs.6

2.5 Proton exchange membrane fuel cells (PEMFCs)

As mentioned before, fuel cells (FCs) are electrochemical energy conversion devices which

convert chemical energy of externally supplied fuel and oxidant into electrical energy.

Proton exchange membrane fuel cells (PEMFCs) are low temperature FCs which are

typically used for transport applications due to their quick start-up time. A schematic

representation of a PEMFC is shown in Figure 2.15. State-of-the-art PEMFCs utilize

polymer Nafion electrolyte and carbon supported Pt electrodes and are collectively known

as membrane electrode assembly (MEA). The role of Nafion is to allow the H+ transport

and prevent electron transport while maintaining a sufficient separation between the two

electrodes. Hydrogen gas is fed at the anode compartment and is oxidized to release protons

and electrons (equation 2.4). The protons travel through the electrolyte membrane to reach

the cathode. At the cathode compartment oxygen is reduced and it combines with proton

to produce water (equation 2.5). The electrons travel from anode to cathode through

external electrical circuit to produce electricity. These electrochemical reactions are

summarized in equations 2.4 - 2.6.

Anode: + -

2H 2H 2e (2.4)

Cathode: + -

2 2

1O + 2H 2e H O

2 (2.5)

Overall: 2 2 2

1H + O H O

2 (2.6)

34

Figure 2.15 The illustration of proton exchange membrane fuel cell (PEMFC).

2.5.1 Challenges of PEMFCs

In Nafion, the proton conduction is based on the hydration content in polymer membrane

and hence the current PEMFCs work at lower temperatures (< 90 °C). Due to low

temperature operation, PEMFCs require expensive catalyst like Pt for the electrode

reaction to happen and as a result, high cost is the major disadvantage of PEMFCs. At low

temperature, CO poisoning occurs at the expensive Pt electrodes on the carbon support

which slows down the electrocatalytic reaction.5 This can be avoided by maintaining a high

operating temperature for the fuel cell. Hence, there is a need for new high temperature

anhydrous proton conductors which can replace the Nafion membrane and the use of

expensive catalysts and making it possible to facilitate the cell operation at elevated

temperatures, ≥ 100 °C with less water management.

35

2.6 Proton conductors for PEMFCs

2.6.1 Polymer membrane electrolytes

Nafion is a polymeric membrane which is the widely used electrolyte for PEMFCs.83 It has

polytetrafluoroethylene backbone and sulfonic acid functional groups in its structure

(Figure 2.16). The backbone provides the mechanical strength and the acid chains provide

the charge sites to facilitate proton conduction by keeping water molecules around them.

The main challenges associated with the current PEMFCs which use Nafion as an

electrolyte are:4, 84

The H+ transport of Nafion is highly dependent on the water content, and

hence the water management is a big issue in the development of PEMFCs,

as the working temperature is close to the boiling point of water.

The lower operational temperature causes electrode poisoning.

The heat and water produced must be continuously removed in order to

maintain the operation temperature and to avoid over-humidification inside

the cell.

Figure 2.16 The chemical structure of Nafion membrane.83

36

2.6.2 Modifications on Nafion

The modifications on the polymer chain allow the PEMFCs to operate at elevated

temperatures and also at lower relative humidity conditions.85, 86 When the inorganic

moieties are incorporated into the polymer chain, the cell operation temperature could be

increased to 100-130 °C by maintaining the chemical and mechanical stability of the

polymer.87 The oxides such as silica, alumina, and titania are proven to have the proton

conducting ability at different humid conditions. However, their proton conductivity is very

poor (10-4-10-6 Scm-1) compared to Nafion.87, 88 The durability and cost are major

drawbacks in using the inorganic based electrolytes. The composites of Nafion with

different heteropoly acids were also studied as PEM electrolyte which show good proton

conductivity (10-2 Scm-1) at above 100 °C.89, 90

2.7 Polyoxometalates (POMs) and ionic liquid derived hybrid proton conductors

Hybrid polyoxometalates (POMs) and ionic liquids (ILs) have drawn wide attention

nowadays in the field of electrochemical energy conversion and storage devices.91-93

Heteropoly acids are an important class of solid acids which consist of poly nuclear and

oxygen bridges. They show prominent Brönsted acidity and are very good proton

conductors. However, they are quite sensitive to atmospheric moisture and that prevents

their practical usage in fuel cells. POMs are anionic metal oxide clusters of HPAs and they

are found to be more stable than their corresponding acids. They possess a robust structure

that exhibits very useful functional physical and chemical properties, including electrical,

magnetic, electrochemical and photonic properties.94 There are three major classification

of POMs which are Anderson [XM6O24]n-, Keggin [XM12O40]

n-, and Dawson

[X2M18O62]n-. They are clusters of heteroatoms, where M is transition metals such as Mo

37

or W and X is tetrahedral template. The coordination of heteroatom is octahedral in the

Anderson family and tetrahedral in the Keggin and Dawson families.95 The crystal structure

representation of these families are shown in Figure 2.17.95

Figure 2.17 The crystal structure of (a) Anderson [XM6O24]n-, (b) Keggin [XM12O40]n-,

and (c) Dawson [X2M18O62]n- families. Adapted from Ivanova, S., ISRN Chemical

Engineering 2014, 2014, 13.95

ILs are low melting point salts, which are typically comprised of organic cations

(e.g. imidazolium ions) and small localized inorganic (e.g. Cl-, Br-) or weakly coordinating

organic anions (e.g. bis(trifluoromethylsulphonyl)imide ions) possessing very low vapour

pressure, non-flammability, high ionic conductivity, high chemical and thermal stability

and a large electrochemical stability window.96-98 These unique properties of POMs and

ILs are combined in their respective hybrid compounds and ensure the tunability of desired

physical and chemical properties by varying the transition metal cations in former case and

organic cations and anions in the latter case. The commonly used organic cations are

substituted imidazolium, pyridinium, pyrrolidinium, and alkyl ammonium and

phosphonium cations, and the anions include tetrafluoroborate, hexafluorophosphates and

bis(trifluoromethanesulfonyl)imide. Various hybrid gels were developed using POMs and

38

ILs such as N-methyl imidazolium-1-(3-sulfonic group) propyl (MIMPS), 3-(pyridine-1-

ium-1-yl)propane-1-sulfonate (PyPs), 1-(3-sulfonic group) triethylamine (TEAPS) and N-

propyl-N-methylpiperidinium hexafluorophosphate.99-104 Several heteropoly acids (HPAs)

and polyvinylpyrrolidone (PVP) and polyethylene glycol (PEG) hybrids were also recently

reported.105 The molecular structure of IL used for this study is PyPs which is prepared in

our lab and is illustrated in Figure 2.18. It is used for the development of hybrid proton

conductors.

Figure 2.18 The molecular structure of 3-(pyridine-1-ium-1-yl)propane-1-sulfonate

(PyPs).

2.8 Tuning layered perovskites as an alternative to Nafion membrane for PEMFCs

There are different types of solid-state proton conductors reported and their comparison

with liquid state proton conductors is shown in Figure 2.19.106 It shows that there are

inorganic materials such as CsHSO4 and Y (10 mol%) doped BaZrO3 which conduct

protons at < 100 °C but their conductivity is far lower when compared to Nafion. The

perovskite materials such as BaCeO3 and BaZrO3 are widely used proton conductors at

higher temperatures, > 500 °C for electrochemical devices such as sensors and solid oxide

fuel cells.107, 108 There is a need for more research for the development of solid-state proton

conductors without compromising the conductivity and stability at the gap (~ 150 - 700

°C) seen in the graph.106 In this thesis work, we tried to convert the metal oxide perovskites

as intermediate temperature (> 150 °C) electrolytes for PEMFCs. Perovskite structure has

39

a general formula, ABO3 where A is a large s, p or f-block cation and B is a small transition

metal cation. The crystal structure of perovskite has already been shown in Figure 2.7. The

layered perovskites derived from ABO3 consists of corner connected metal-oxygen

octahedra and the alkali metal ions occupied at the interlayer positions. Ruddlesden-Popper

(RP) and Dion-Jacobson (DJ) series compounds with general formula, A2[An-1BnO3n+1] and

A/[An-1BnO3n+1], respectively are famous for their interlayer chemistry between the

perovskite slabs.109, 110 The interlayer cation can easily undergo ion exchange reaction. The

metal oxide octahedra (BO6) form robust frameworks by sharing their corners.

Figure 2.19 Comparison of proton conductivity of solid- and liquid-state materials.106

0.5 1.0 1.5 2.0 2.5 3.0 3.5

-6

-5

-4

-3

-2

-1

0

1

"gap"

Gd:BaPrO3

Y:BaCeO3

Ba:LaErO3

BCN18

CsHSO4

1M HCl

Nafion117

log

10 (

Scm

-1)

1000/T(K-1)

1600

1400

1200

1000

800

600

400

200

T (oC)

40

Soft chemistry is a powerful tool to tailor a variety of new materials with different

structures and properties for many applications. It is possible to produce new phase of

materials from KLaNb2O7 and K2La2Ti3O10 which belong to DJ and RP series, respectively

at lower temperature by interlayer cation exchange and the best example is their H+

analogue after the proton exchange reaction.110, 111 The solid acids formed from DJ

compounds can intercalate a variety of bases and alcohols to form new molecular

composites.112 However, not all protonated RP compounds show the same trend as DJ

series.110 The crystal structures of KLaNb2O7 and K2La2Ti3O10 are shown in Figure 2.20,

they show that the potassium ions exist between the perovskite slabs. These K+ ions are

prone to undergo ion exchange reaction.

As mentioned in Section 2.7, Ionic liquids (ILs) are salts with unique properties.

No studies have been designed to produce a layered perovskite structure-based proton

conductors in combination with ILs. Combining the discrete property of layered

perovskites and ILs can catalyze the development of new class of hybrid highly efficient,

robust frame work structured proton conducting electrolyte which do not need a complex

water management system to provide excellent proton conductivity, is one of the scopes of

this thesis. The structure of IL used for this purpose is shown in Figure 2.21 and was

provided by our collaborator113.

41

Figure 2.20 The crystal structure of a) KLaNb2O7 and b) K2La2Ti3O10.

Figure 2.21 The structure of ionic liquid containing imidazolium cation and imide

anion. For simplicity, the structure is represented as IL.

In summary, the work done for energy conversion device in this thesis is development of

proton conductors based on POMs, ILs and layered perovskites as an alternative to Nafion

which is limited by the low operating temperature (< 90 °C) which can work beyond 90

°C.

42

Chapter Three: Experimental section

3.1 Materials synthesis

3.1.1 Synthesis of lithium-stuffed garnet-type metal oxides

Conventional solid-state method was used for the synthesis of lithium-stuffed garnet-type

metal oxides of the nominal chemical composition Li5+2xLa3Ta2-xYxO12 (x = 0.05, 0.10,

0.25, 0.50 and 0.75) and Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75). The stoichiometric

quantity of the precursors, LiNO3 (99%, Alfa Aesar), La2O3 (99.99%, Alfa Aesar), Ta2O5

(99.5%, Alfa Aesar), Nb2O5 (99.5%, Alfa Aesar), and Y(NO3)3 (99.9%, Alfa Aesar) were

weighed and mixed using a Pulverisette, Fritsch, Germany ball mill at 200 rpm for 12h in

2-propanol to ensure homogeneous mixing. The ball milling jars and balls were made of

zirconia. As La2O3 is hygroscopic in nature, it was pre-heated at 900 °C for 24h to avoid

any moisture uptake. An extra 10 wt.% of LiNO3 was added during the weighing step to

compensate for the lithium loss through volatilization at higher temperature. After solvent

evaporation, the dried powder samples were subjected to initial heating at 700 °C for 6 h

in clean alumina crucibles and further ball milling was conducted. The resultant powders

were dried and pressed into pellets using P. O. Weber isostatic press by applying a pressure

of 183 MPa for 1 min. The pellets were subjected to a two-step heat treatment in air at 900

°C for 24h and 1100 °C for 6h. It was ensured to cover the bottom of alumina crucibles as

well as the pellets with excess mother powder for each heat treatment. The resultant pellets

were used for further characterization either in powder or disc (pellet) form depending on

the requirement.

Same procedure was followed for the preparation of Li6La3Ta1.5Y0.5O12 with

varying the amount of excess LiNO3. Specifically, 2.5, 5, 7, 10 and 15 wt.% excess LiNO3

43

was used to synthesize a series of garnets to understand the role of excess Li on phase

formation, microstructure and Li ion conductivity.

3.1.2 Synthesis of polyoxometalate and ionic liquid based proton conductors

3.1.2.1 Synthesis of solid acid, H3PW11MoO40 (H3PWMo)

Heteropoly acid (HPA) or polyoxometalate (POM) of the nominal chemical composition

H3PW11MoO40 was synthesized according to literature.99 Na2MoO4 (0.01 mol) (BDH,

99.5%) was dissolved in 20 mL of deionized water. Na2WO42H2O (0.1 mol) (Sigma

Aldrich, 99%) was mixed with 0.02 mol of Na3PO4.12H2O (Fisher Scientific, 100%) and

dissolved in 100 ml of deionized water and added to the sodium molybdate solution. Dilute

H2SO4 (1:1) was used to adjust pH of the mixture to 2.5 and then heated at 90 C for 2h. It

was allowed to cool down to room temperature and the solution was extracted using ether

(50 ml) and dilute H2SO4 (1:1) prior to drying in vacuum oven at 50 C for 24 h, to obtain

the solid product.

3.1.2.2 Synthesis of solid acid H4PMo11VO40 (H4PMoV)

0.01 mol of NaVO3 (BDH, 98%) was dissolved in 20 mL of deionized water. 80 mL of

Na2MoO4 (0.1 mol) (BDH, 99.5%) was mixed with 20 mL of Na3PO4.12H2O (0.01 mol)

(Fisher Scientific, 100%) solution. The rest of the steps are the same as described for

H3PWMo synthesis above.

3.1.2.3 Synthesis of solid acid H5PMo10V2O40 (H5PMoV)

0.02 mol of NaVO3 (BDH, 98%) was dissolved in 20 ml of deionized water. Na2WO42H2O

(0.1 mol) (Sigma Aldrich, 99%) was mixed with 0.01 mol of Na3PO4.12H2O (Fisher

Scientific, 100%) and dissolved in 100 mL of deionized water. The rest of the steps are the

same as described for H3PWMo synthesis above.

44

3.1.2.4 Synthesis of PyPs ionic liquid

Synthesis of ionic liquid, 3-(pyridine-1-ium-1-yl) propane-1-sulfonate (PyPs), was carried

out as per the procedure reported in literature.114 As-received 4.4 ml of 1,2-oxathiolane 2,2-

dioxide (Alfa Aesar, 99%) was added to 50 ml of acetone. 4.0 ml pyridine (Alfa Aesar,

99%) was added to this solution and was heated at 50 C for 4 h while stirring. A white

precipitate was formed which was filtered, washed with ether and dried under vacuum.

3.1.2.5 Synthesis of hybrid gels

PyPs [3-(pyridine-1-ium-1-yl)propane-1-sulfonate] and POM hybrid gel electrolytes were

synthesized by taking 3:1, 4:1, or 5:1 mole ratio of PyPs to H3PW11MoO40, H4PMo11VO40,

and H5PMo10V2O40, respectively. PyPs was then added to aqueous solution of the

heteropoly acids, H3PW11MoO40, H4PMo11VO40, and H5PMo11VO40 under constant

stirring overnight at room temperature. Water was removed from the solution by vacuum

drying to obtain the desired hybrid gel products: [PyPs]3PW11MoO40 (PyPs-H3PWMo),

[PyPs]4PMo11VO40 (PyPs-H4PMoV), and [PyPs]5PMo11V2O40 (PyPs-H5PMoV).

3.1.3 Synthesis of proton conductors based on layered perovskites and ionic liquid

3.1.3.1 Synthesis of KLaNb2O7 and K2La2Ti3O10 layered perovskite and its proton

analogues

The layered perovskites, KLaNb2O7 (Dion-Jacobson – DJ phase) and K2La2Ti3O10

(Ruddlesden-Popper – RP phase), were synthesized using previously reported solid-state

techniques.111, 115 Stoichiometric amounts of high purity precursors, K2CO3 (99%, BDH),

La2O3 (99.99%, Alfa Aesar), Nb2O5 (99.5%, Alfa Aesar) and TiO2 (99.8%, Alfa Aesar)

were homogeneously mixed using ball milling at 200 rpm for 6h in 2-propanol. A 15 wt.%

excess K2CO3 were used to account for volatilization loss of potassium during sintering.

45

The resultant mixtures were heated at 700 °C for 6h and ball milled again prior to making

pellets. The sintering was done at a rate of 5 °C per minute at 1150 °C for 12h to ensure

the phase formation. Ion exchange (soft chemistry) was performed on the layered

perovskites, KLaNb2O7 and K2La2Ti3O10, to make proton analogues by mixing the

powdered layered perovskite samples in 6 M and 2 M HNO3 solutions, respectively for 1-

2 weeks using a magnetic stirrer at room temperature followed by washing with water. The

acid was replaced every 2 days during the ion exchange process.

3.1.3.2 Synthesis of hybrid proton conductors from layered perovskite and ionic liquid

Two routes of ion exchange reactions were carried out with both as received (pure) and

diluted (using acetonitrile solvent) commercial ionic liquid, C13H22N4O5F6S2 (IL). Route 1

was directly treating the perovskite with IL and Route 2 was treating the proton analogue

of layered perovskites with IL. The reaction mixtures were made by mixing 1 g perovskite

powder with 10 mL dilute IL for 1-2 weeks of magnetic stirring. In the case of pure IL,

amount of perovskite powder was adjusted in a way that the IL completely wets the powder

and is good enough to undergo magnetic stirring in a hot plate which is set at 200 °C for 3

days. The resultant powder was washed with acetonitrile after the ion exchange process.

Later it was separated using vacuum filtration and was kept in air for room temperature

drying.

3.2 Material characterization

3.2.1 Powder X-ray diffraction (PXRD)

The most important tool employed in solid-state chemistry is the powder X-ray diffraction

(PXRD) technique to determine the purity of crystalline materials. The X-rays are

generated from cathode ray tube and are filtered to obtain monochromatic X-rays (Cu Kα

46

radiation, λ = 1.5418 Å in this thesis) which are directed towards the sample surface. The

crystal planes reflect the X-rays constructively or destructively. Diffraction happens when

the interferences are constructive, i.e. when Braggs law is satisfied as shown in equation

2.1.116

2 sinn d (3.1)

where n, λ d and θ are the integers representing the order of the diffraction, wavelength of

X-rays, inter-planar distance, and Bragg angle (angle of incidence), respectively. Figure

3.1 illustrates Bragg’s law of powder X-ray diffraction.116

Figure 3.1 Illustration of Bragg’s law for powder-X-ray diffraction116.

In this thesis work, a Bruker D8 powder X-ray diffractometer (CuKα, 40 kV, 40

mA) was employed for the material purity analysis. The room temperature diffraction

patterns were recorded at a count rate of 10s per 0.05 ° in the 2θ range of 5 - 80 °. The long

run scans (13h) for refinements were recorded at a slow scan rate of 0.015 ° per 10s. Crystal

structure analysis of the PXRD data was performed using the Rietveld refinement analysis

(with the help of Dr. Farshid Ramezanipour)117 with the GSAS program118 and EXPGUI

47

interface119. Eva and PROSZKI programs were also used for PXRD analysis to index the

peaks and for the calculation of lattice parameters of the crystal structure, respectively.

3.2.2 Scanning electron microscopy (SEM)

The microstructure of samples can be analyzed using scanning electron microscopy (SEM)

which gives information such as particle size and texture.116, 120 In the electron microscope,

high energy electrons bombard the sample surface under high vacuum (10-4 Torr). The

interaction between the electron and sample result in the production of secondary and back

scattered electrons, X-rays, and light. The secondary and back scattered electrons are

captured by the detector and are converted into micrographs. The characteristic X-rays are

detected to quantify specific elements in the sample composition using energy dispersive

X-ray diffraction spectroscopy (EDS). Multiple imaging at different spots was performed

in order to examine morphological changes.

For SEM analysis, powder or pellet samples sputtered with gold-palladium (Au-

Pd) were used to avoid charge build up. This thesis work utilized a scanning electron

microscope (Philips FEI XL30) equipped with an energy dispersive X-ray diffraction

spectrometer in the Microscopy and Imaging Facility, Health Sciences Center, University

of Calgary (Chapter 4, 6, and 8), Field emission SEM (FESEM-Zeiss Zigma Series),

Nanoscience Program, University of Calgary (Chapter 4), and Hitachi SU-70 scanning

electron microscope, University of Maryland (Chapter 5).

3.2.3 Density measurements

High density is required for a garnet-type Li+ ion conductor for the application as lithium

ion battery (LIB) electrolyte. Two methods were employed to measure garnet sample

density in this thesis work, the Archimedes method and the He Pycnometry technique.

48

According to the Archimedes principle, when a solid object is suspended in a fluid, the

buoyant force exerted by the fluid on the suspended solid object is equal to the weight of

the fluid which is being displaced by the solid object. Methanol is used as the solvent in

this thesis (Chapter 4). Methanol was used instead of water because garnets tend to

exchange Li ions with water, which could affect the accuracy of the density measurements.

The potential exchange of Li in methanol has not been considered in this thesis work. A

METTLER TOLEDO density determination kit and weighing balance were used for this

work. Suspended weight (WSus) was taken when the pellet is suspended in methanol.

Saturated weight (WSat) was obtained for the solvent absorbed pellet. Excess solvent was

wiped off from the pellet surface using Kim wipe prior to weighing the saturated weight.

Pellet samples were dried in air oven at 100 °C for 3h to obtain dry weight, WDry. The

Archimedes density was calculated using the equation 3.2.

0.7918

Dry

Archimedes

Sat Susp

WD

W W

(3.2)

The constant 0.7918 in equation 3.3 is the room temperature density of methanol (gcm-3)

which was used as the suspension medium in the Archimedes technique. The porosity of

the samples was calculated using the formula as shown in equation 3.3.121

% x100Sat Dry

Sat Susp

W Wporosity

W W

(3.3)

Helium (He) displacement Pycnometer experiment was the second method used to

measure the density of powder samples (Micrometrics AccuPyc 1340, Gas Pycnometer),

Department of Chemical and Petroleum Engineering, University of Calgary. Sample

49

density was calculated based on mass (m) and volume (Vs) of the powder sample (equation

3.4).

s

md

V (3.4)

The term Vs depends on various factors such as the volumes of reference (Vr) and empty

cell (Vc) and the gas pressure at different stages during the measurement, which can be

represented as follows (equation 3.5).

1

rs c

i r

f r

VV V

P P

P P

(3.5)

where Pr, Pi, and Pf indicate the system reference, initial and final pressure, respectively.

The theoretical density (d) was calculated using the lattice parameter derived from PXRD

analysis by employing equation 3.6.

3

A

ZMd

a N (3.6)

where Z, M, a, and NA represent the number of chemical formula per unit cell (Z = 8 in the

currently studied Li-stuffed garnets), molar mass of the nominal chemical composition,

lattice parameter (obtained from PXRD), and Avogadro’s number (6.022 x 1023),

respectively.

3.2.4 Electrochemical impedance spectroscopy (EIS)

To understand the electrical properties of ionic conductors, electrochemical impedance

spectroscopy (EIS) is a powerful technique. Impedance of the material is measured over a

frequency range (0.01 to 10 MHz) by applying an AC (alternating current) voltage of 100

mV. The advantage of this technique is that conductivity contribution due to bulk, and

50

grain-boundary could be resolved successfully. Impedance is a complex function of

resistance (R) and capacitance (C) which depends on frequency (ω) and the relationship

can be represented as shown in equation 3.7.122

1RC (3.7)

The capacitance of materials in the range 10-12, 10-11-10-8 and 10-7-10-5 F, represent

the contributions due to bulk, grain-boundary, and electrode interface, respectively.123

Typical impedance spectrum (Nyquist plot) of a polycrystalline material is shown in Figure

3.2 representing different contributions and the corresponding RC circuit is also shown as

an inset. A blocking electrode (typically Au) is used for the measurements that appears as

a tail in the impedance plot. Illustration of AC impedance set up used for measuring

electrical properties of garnet samples is represented in Figure 3.3.

Figure 3.2 Nyquist plot of a polycrystalline material showing bulk, grain-boundary

and electrode contribution.

51

Figure 3.3 Schematic representation of the cell set up used for AC impedance

measurements for garnet samples.

The capacitance is represented by the formula (equation 3.8) where εr the relative

permittivity of medium, εo is the vacuum permittivity (8.854 x 10-14 F cm-1), a is the

electrode area and l is the thickness of pellet.

o r

aC

l (3.8)

Due to material roughness, the capacitance measured is not ideal and is represented by a

constant phase element, CPE instead of C. Hence the bulk and grain-boundary capacitance

are calculated using the formula (equation 3.9),

(3.9)

where n is the integer which varies between 0 to 1 and represents the deviation from pure

capacitor (e.g., for pure capacitor, n = 1).

EIS measurements done in this thesis work employed a Solartron SI 1260

impedance and gain-phase/analyzer and VersaSTAT 3 analyzer. The garnet samples for

impedance measurements were prepared as follows; the pellets after sintering at 1100 °C

1 1n

n nC R CPE

52

were cut in to ~ 2 mm thickness and ~ 10 mm diameter using a diamond saw. Both sides

of the pellets were painted with Au paste (blocking electrode). Two batches of coating were

applied to obtain a good electrical conductivity by ensuring to dry the coating in air oven

at 80 °C each time. Finally, it was cured at 600 °C for 1 h to ensure the continuous contact

of the electrode. The impedance was measured as a function of temperature (22 – 325 °C)

using a tubular furnace, giving enough time to stabilize the temperature at each stage.

Different pellets were used to ensure the reproducibility of results. The conductivity was

calculated the formula given in equation 3.10.

1 l

R a

(3.10)

where σ, R, l, and a are the conductivity, resistance, thickness and area of the pellet,

respectively. The temperature dependence of the conductivity was represented using the

Arrhenius equation (equation 3.11), where A is the pre-exponential term, Ea is the

activation energy, T is the temperature, and k is the Boltzmann constant. The slope of logσT

vs 1000/T plot was used for Ea calculation.

expaE

kTT A

(3.11)

The impedance measurement set up used for polyoxometalate and ionic liquid

based hybrid materials are shown in Figure 3.4. Samples were loaded in a 50 mL 3-necked

flask. Here two parallel Pt plates connected to Pt wire were used as current collectors and

are immersed in the hybrid gel and the flask containing whole set up was heated for 48 h

to remove the water content.

53

Figure 3.4 The cell set up used for AC impedance measurements for polyoxometalate

and ionic liquid based hybrid materials.

The conductivity as a function of temperature was measured using AC impedance method

and temperature was controlled by a hot plate and was monitored by a thermometer. In

order to avoid any error contributing from the absorption of moisture by the sample, it was

heated to 100 °C first and the measurements were taken while cooling down.

3.2.5 Cyclic voltammetry (CV)

Cyclic voltammetry studies were carried out to evaluate the electrochemical stability

windows (ESW) of the materials. A two-electrode measurement set up using Li and Au

electrodes in a glove box was used for garnet materials. Lithium granules were melted on

the surface of pellet (pre- coated with Au on the other side) by heating above 180 °C and

repeated this process few times before applying the final melted Li layer for measurement.

A scanning rate of 10, 50 and 100 mV s-1 between -0.5 V to 4 V was applied for measuring

54

5-20 cycles at different temperatures. A PARSTAT 3000 Potentiostat was used for the CV

study of garnets in this thesis work. The schematic representation of experimental set up

used for this purpose is shown in Figure 3.5.

Figure 3.5 Schematic representation of cyclic voltammetry cell set up used for garnet

sample.

For hybrid materials, three-electrode sample set up was used for the cyclic

voltammetry employing VersaSTAT 3 potentiostat. Pt electrodes were used as the working

(WE) and counter electrodes (CE), and Ag/AgCl was used as the reference electrode (RE)

under N2 atmosphere. The scanning was done at 10-100 mVs-1 between -4 to +4 V and

measurements were repeated for 5-20 cycles.

3.2.6 Thermogravimetric analysis

Thermogravimetric analysis (TGA) measures the weight change of samples as a function

of temperature. TGA results can give an insight on the decomposition and redox nature of

materials tested under specific atmosphere. Thermal analysis can be done under different

atmospheres such as N2, Ar, O2, H2, CO2 and humidified gases. Powder samples were used

under N2 atmosphere for this thesis work. A typical TGA curve with a one-step weight loss

55

is represented by Figure 3.6, the normal way of weight loss calculation, is also marked in

the figure. Depending on the sample nature, multiple step weight loss at different

temperature range can be observed. A TGA curve with no weight loss indicates that the

material is stable within the measured temperature range.

Figure 3.6 Schematic representation of typical TGA curve showing single step weight

loss.

A METTLER TOLEDO thermal system; TGA/DSC1 HT 1600 °C with Stare

system was mostly used for TGA analysis in this thesis work. Heating and cooling cycles

were done at a rate of 5-10 °C min-1 up to a maximum temperature of 700 °C. The

maximum temperature was maintained for 1 h prior to the cooling cycle.

Thermogravimetric analyzer (Setaram TAG 16 TGA/DSC dual chamber balance) was used

to study the thermal stability of the hybrid samples (Chapter 8) under N2 gas up to 650 °C

at a heating rate of 10 °C/min.

56

3.2.7 Fourier transform infrared (FTIR) spectroscopic studies

Fourier transform infrared (FTIR) spectroscopy is a technique which analyzes the

stretching or vibration of chemical bonds present in the material on interaction with the

infrared radiation. Specific wavelengths of radiation are absorbed or transmitted by

different functional groups in organic and inorganic materials. Hence the FTIR spectrum

is unique for each molecule and is known as the molecular fingerprint of compounds. A

Varian Model 7000, NEXUS Model 470 FTIR was used for FTIR measurements in this

thesis work. KBr standard was used for the measurements. Overnight drying at above 100

°C was done for all the samples and KBr powder. The samples were dispersed in KBr

powder for the measurements. Purging the instrument chamber with N2 gas was done to

flush out H2O and CO2 content in air. The scanning was done between 400 and 4000 cm-1.

3.2.8 Solid-state magic angle spinning nuclear magnetic resonance (MAS NMR) studies

The nuclear magnetic resonance (NMR) technique is now widely used to understand the

molecular structure and dynamics in the solid-state materials. NMR measures the average

of anisotropic molecular interactions. Molecular orientations are dependent on the nuclear

spin interactions.124 Due to the presence of anisotropic interactions, the NMR spectra for

solid materials are usually broad which lowers the resolution of the spectra. Magic angle

spinning (MAS) is one of the techniques employed to enhance the resolution of NMR

spectra. For this, the sample is rotated at high speed at an angle of 54.74 ° which makes

faster molecular orientation and results in high-resolution spectra. Facilities at the

University of Calgary were used for solid-state MAS NMR studies carried out in this thesis

work. 7Li and 27Al MAS NMR studies for garnet powder samples were carried out

employing AMS 300 Bruker equipment, at a spin rate 5 kHz. The standards used for the

57

representation of chemical shift values were LiCl and Al(NO3)3, respectively for 7Li and

27Al MAS NMR. 1H and 13C NMR MAS NMR studies for polyoxometate and ionic liquid

based hybrids were analyzed using Bruker RDQ 400 NMR spectrometer.

3.2.9 Raman spectroscopy

Raman spectroscopy is an analytical tool for understanding the vibrational and rotational

motion of molecules under the excitation of a ultra violet-visible (UV-Vis) beam. It

measures the shift in vibrational frequency of the initial beam after interaction with the

sample. For this thesis work, Raman spectroscopy was performed with a 532 nm laser using

a Horiba LabRAM ARAMIS Raman microscope at University of Maryland. The

microscope settings used included a 1800 mm−1 grating and a D1 filter (10 x intensity

reduction). Results were measured between wavenumbers from 50 cm−1 to 1400 cm−1.

3.2.10 Chemical stability tests of Li5+2xLa3Ta2-xYxO12 garnets

The structural stability of as-prepared Li5+2xLa3Ta2-xYxO12 at different temperatures (30 –

900 °C) was studied by variable-temperature in situ PXRD using a high-temperature

chamber, Anton Paar XRK 900 in a Bruker D8 diffractometer. The samples were heated at

10 °C min-1 and held for 30 min to reach thermal stability at each heating stage. In order to

analyze the chemical stability of Li5+2xLa3Ta2-xYxO12 in water, ~ 1 g powder was stirred

with 100 mL de-ionized water in Erlenmeyer flask using magnetic stirrer for 2 days. The

solution was vacuum filtered to collect the powder. Stability of Li5+2xLa3Ta2-xYxO12 in

aqueous LiCl solution was determined by mechanical stirring of the pellet in 1 M LiCl for

1 week employing a magnetic stirrer, after which the pellets were vacuum-dried at ~ 100

°C for 6 h prior to the PXRD and conductivity measurements.

58

3.3 Error considerations

The material synthesis itself could be affected by the impurities present in the

precursors used as its purity is ~ 99 %. Volatilization of alkali metals (e.g., Li, K)

at high temperature sintering could have affected the stoichiometric composition

of samples. Excess amount of volatile salts were used to account for this loss.

Alumina crucibles were used for the ceramic sample preparation and its

contamination on the sample was not quantified. Ionic liquid synthesized in lab

was not recrystallized after synthesis for purity confirmation except for the

washing step at the end of synthesis process. This could have an effect on the

properties of hybrid materials which was not counted in this thesis work.

Methanol used for the Archimedes density measurements was not purified to avoid

the presence of water. The tendency of garnet materials to undergo proton

exchange reaction with water as well as the solvent was not considered for the

calculations performed in this thesis work.

PXRD analysis has a detection limit of 5 % and hence minor impurities are not

detected. Possible sample placement errors were avoided by levelling the powder

samples in the sample holder before measurements. Calibration of PXRD

instrument was done periodically using NIST corundum standard SRM 1976a to

avoid any instrument alignment issue.

The flow rate of gas (90 mL min-1) purged in the Mettler Toledo thermal system

TGA/DSC1 (HT 1600 °C) has an error of ~ 5 % due to the use of a conventional

flow meter. The error associated with the TGA furnace is ±0.5 °C and that of TGA

balance is ±1 μg.

59

Calibration of the Solartron SI 1260 impedance spectrometer was done using

dummy cells (1.8-9.3 k Ω between 1-10 kHz). The possible ohmic resistances at

high frequency measurements could also be accounted from cable contact, cable

length, cable connection, and noise.

Error in temperature reading of the tubular furnaces used for the impedance

measurements is ±3 °C and the thermocouple used to measure the temperature of

sample placed inside the impedance cell is ±1 °C.

Author takes responsibility for any oversight in interpreting the data.

60

Chapter Four: Dopant concentration - porosity - Li-ion conductivity relationship in

garnet-type Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) and their stability in water and 1M

LiCl

Contents in this chapter are reproduced from published work, “Dopant Concentration -

Porosity - Li-ion Conductivity Relationship in Garnet-Type Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x

≤ 0.75) and Their Stability in Water and 1M LiCl” by Sumaletha Narayanan, Farshid

Ramezanipour, and Venkataraman Thangadurai, in Inorg. Chem. 2015, 54, 6968-6977.

4.1 General overview125

Recently, the garnet-type Li-stuffed metal oxides have shown excellent potential for

application as solid Li ion electrolytes.14 Several members of the garnet family have shown

high Li+ ion conductivity, chemical stability to elemental Li and large electrochemical

window at room temperature.25 In addition to solid electrolyte applications, another area

where Li ion conducting solids can be exploited is in lithium-air batteries. Aqueous

electrolytes in these batteries need to be separated from the elemental lithium anode to

prevent the vigorous reaction with lithium. To protect the lithium electrode, protective

layers consisting of Li ion conducting solids that are stable in contact with water and

elemental Li are desired.126-128 Recently, we showed that the ionic conductivity of

Li5La3Nb2O12 can be enhanced by partial doping of Y at the Nb-sites, leading to the Li+

conductivity of up to 10-4 Scm-1 at room temperature.57 Given the wider voltage window

and higher stability of the Ta analogue, Li5La3Ta2O12, to metallic lithium, we envisioned

the possibility of obtaining high ionic conductivity in the Ta compounds, where both high

stability and Li conductivity can be present.

61

This chapter presents the detailed study of highly Li+ ion conductive Y-doped

garnet-type Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) to understand the effects of yttrium- and

lithium-doping on chemical composition, crystal structure, porosity and Li+ ion

conductivity using 7Li MAS NMR, electrochemical AC impedance spectroscopy and

scanning electron microscopy (SEM), as well as ex-situ and in-situ powder X-ray

diffraction (PXRD) to further explore the potential application of garnets in all-solid-state

Li ion batteries. Solid-state 7Li MAS NMR studies showed an increase in the Li ion

mobility as a function of Y- and Li-doping in Li5+2xLa3Ta2-xYxO12, which is consistent with

the results from AC impedance spectroscopy. The SEM studies on sintered pellets

indicated a systematic decrease in porosity and an increase in sinterability as the Y- and

Li-doping levels increase in Li5+2xLa3Ta2-xYxO12. These results are consistent with the

calculated porosity and densities using the Archimedes method. Using the variable-

temperature in situ PXRD in the temperature range of 30- 700 °C, a thermal expansion

coefficient of 7.25 x 10-6 K-1 was observed for Li6La3Ta1.5Y0.5O12 in the temperature range

of 30- 700 °C. To further explore the possibility of a new application for the Li-stuffed

garnets, the stability of these materials in aqueous LiCl solution was also studied. A high

degree of structural stability was observed in these materials upon 1 M LiCl treatment,

making them suitable candidates for further studies as protective layers for lithium

electrodes in aqueous lithium batteries.126-128

4.2 Results and discussion125

4.2.1 Phase analysis of Li5+2xLa3Ta2-xYxO12

Rietveld refinements119, 129 with powder X-ray diffraction data were performed to confirm

the formation of the garnet-type structure in Li5+2xLa3Ta2-xYxO12 phases. The results of

62

these analyses for the x = 0.05, 0.1 and 0.2 phases are shown in Table 4.1, and that of x =

0.25, 0.5 and 0.75 are shown in Table 4.2.

Table 4.1 The Rietveld refinement results for Li5+2xLa3Ta2-xYxO12 (x = 0.05 - 0.20).

Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015)


x = 0.05 x = 0.10 x = 0.20

La (24c) occupancy 1 1 1

X 1/8 1/8 1/8

Y 0 0 0

Z 1/4 1/4 1/4

Uiso 0.0230(8) 0.0209(7) 0.0224(7)

Ta/Y (16a) occupancy 0.975/0.025 0.95/0.05 0.90/0.10

X 0 0 0

Y 0 0 0

Z 0 0 0

Uiso 0.0214(7) 0.0198(6) 0.0206(6)

Li1 (24d) occupancy 0.768 0.7453 0.705

X 1/4 1/4 1/4

Y 7/8 7/8 7/8

Z 0 0 0

Uiso 0.025 0.025 0.025

Li2 (48g) occupancy 0.156 0.174 0.207

X 1/8 1/8 1/8

Y 0.6826 0.6826 0.6826

Z 0.5674 0.5674 0.5674

Uiso 0.025 0.025 0.025

Li3 (96h) occupancy 0.155 0.16 0.1702

X 0.0927 0.0927 0.0927

Y 0.684 0.684 0.684

Z 0.5795 0.5795 0.5795

Uiso 0.025 0.025 0.025

O (96h) occupancy 1 1 1

X 0.2896(6) 0.2883(5) 0.2868(5)

Y 0.0990(6) 0.1012(5) 0.1012(5)

Z 0.19846) 0.1984(6) 0.1996(6)

Uiso

Rp (%)

Rwp (%)

0.0207(31)

9.18

11.83

0.0212(29)

8.58

11.17

0.0215(28)

8.25

11.01

a (Å) 12.8160(3) 12.8268(4) 12.8642(3)

63

The PXRD Rietveld refinement profiles of Li5+2xLa3Ta2-xYxO12 (x = 0.05. 0.1. 0.2, 0.25,

0.5 and 0.75) are presented in Figures 4.1-4.6, and show the good fit to the garnet-type

structure.

Table 4.2 The Rietveld refinement results for Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.50 and

0.75). Reproduced by permission of the PCCP Owner Societies.130

x = 0.25 x = 0.50 x = 0.75

La (24c) occupancy 1 1 1

X 1/8 1/8 1/8

Y 0 0 0

Z 1/4 1/4 1/4

Uiso(Å2) 0.0470(11) 0.0267(11) 0.0315(13)

Ta/Y (16a) occupancy 0.842(11)/0.158(11) 0.743(11)/0.257(11) 0.666(12)/0.334(12)

X 0 0 0

Y 0 0 0

Z 0 0 0

Uiso(Å2) 0.0469(10) 0.0234(11) 0.0269(14)

Li1 (24d) occupancy* 0.681 0.674 0.544

X 1/4 1/4 1/4

Y 7/8 7/8 7/8

Z 0 0 0

Uiso(Å2) 0.030 0.030 0.050

Li2 (48g) occupancy* 0.225 0.220 0.261

X 1/8 1/8 1/8

Y 0.6826 0.6773 0.6826

Z 0.5674 0.5727 0.5674

Uiso (Å2) 0.030 0.030 0.050

Li3 (96h) occupancy* 0.1756 0.218 0.2752

X 0.0927 0.0937 0.0927

Y 0.684 0.6888 0.684

Z 0.5795 0.5817 0.5795

Uiso(Å2) 0.030 0.030 0.050

O (96h) occupancy 1 1 1

X 0.2888(6) 0.2884(6) 0.2899(7)

Y 0.0976(6) 0.0985(6) 0.0977(8)

Z 0.2021(7) 0.1975(6) 0.2025(8)

Uiso(Å2)

Rp

Rwp

0.0316(32)

0.0601

0.0770

0.0127(33)

0.0647

0.0812

0.046(4)

0.0663

0.871

a (Å) 12.8977(4) 12.937(1) 12.964(2)

64


xYxO12. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977. Copyright

(2015) American Chemical Society.125

Figure 4.2 The Rietveld refinement profile for x = 0.1 member of Li5+2xLa3Ta2-xYxO12.



65





xYxO12. Reproduced by permission of the PCCP Owner Societies.130

66


Reproduced by permission of the PCCP Owner Societies.130


xYxO12. Reproduced by permission of the PCCP Owner Societies.130

67

For all Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) phases a cubic structure with space group Ia-

3d can be identified. There is a very small peak at ~ 21.5°, which belongs to Li2CO3. For

x ≈ 0.75, small peaks of unknown phases are usually observed, and were removed in the

final cycles of refinements. A systematic increase in the cell constant as a function of Li

and Y contents was observed in Li5+2xLa3Ta2-xYxO12, as expected due to the larger ionic

radius of Y3+ (0.90 Å) compared to Ta5+ (0.64 Å),131 an increase in the unit cell dimension

is expected as the amount of Y3+ increases in Li5+2xLa3Ta2-xYxO12. This increase can be

observed from the data shown in Table 4.1 and 4.2. A detailed illustration of garnet-type

structure is presented in Figure 4.7. In the garnet-type structure, the transition metals (Ta

or Y) are octahedrally coordinated by oxygen atoms (Figure 4.7a). The Ta(Y)O6 octahedra

are not connected to each other and do not share corners or edges, unlike in the perovskite-

type structure.132 The La and Li atoms reside in the spaces in between these octahedra. The

coordination sphere of La consists of eight oxygen atoms, forming a dodecahedral

geometry with triangular faces around the La atom (Figure 4.7d).

68

Figure 4.7 The garnet-type structure of Li5+2xLa3Ta2-xYxO12. (a) Ta-O sublattice

where Ta is shown in blue and oxygen in red. (b) Ta-O sublattice with La (purple) in

spaces between TaO6 units. (c) TaO6 units shown by gray octahedra. La atoms are in

purple, while Li1 (24d site) and Li2 (48g site) are shown as blue and green spheres,

respectively. Li3 (96h site) is omitted for clarity. (d) The La-O sublattice showing

LaO8 polyhedra. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.

Copyright (2015) American Chemical Society.125

The Li ions sites in Li-stuffed garnets are well established in the literature,72, 133,134-

137 There is an essential role for Li distribution in Li-ion conductivity. Three

crystallographically distinct Li positions, Li(1) (24d site), Li(2) (48g site), and Li(3) (96h

site) are known for this structure.135, 136 Note that the octahedra are connected to the

tetrahedra through face-sharing. There is also edge-sharing between neighboring

69

octahedra. The above arrangement leads to an array of Li atoms surrounded by octahedral

or tetrahedral geometry, which form a three-dimensional network throughout the material

(Figure 4.8), leading to the high Li-ion conductivity observed in these systems. It is

important to note that all of the sites cannot be occupied simultaneously due to short Li –

Li contacts.24

Figure 4.8 Three-dimensional connectivity of oxygen octahedra and tetrahedra,

which can accommodate Li atoms, shown as blue and green spheres. In the actual

structure, not all of the neighboring sites can be simultaneously occupied by Li and it

is not possible for a Li atom to exist at an octahedral site if both tetrahedral sites

adjacent to it are occupied. Reprinted with permission from Inorg. Chem. 2015, 54,

6968-6977. Copyright (2015) American Chemical Society.125

An octahedral site will remain vacant if both tetrahedral sites adjacent to it are

occupied.136, 137 While the occupancy of Li positions cannot be determined by laboratory

PXRD, it is possible to obtain an estimate of the Li occupancy on each site based on neutron

diffraction experiment on similar materials. O’Callaghan and Cussen have studied

Li5+xBaxLa3-xTa2O12 (x = 0, 0.5, 1, 1.2, 1.4, 1.6) and have shown a correlation between the

total Li content in the system and the occupancy of Li on each of the three sites, 24d, 48g

70

and 96h.136 On the basis of their neutron diffraction data, there is an increase in the Li

occupancy on the octahedral sites, 48g and especially 96h, as a function of total Li-content.

On the other hand, the Li occupancy on the 24d site (centers of tetrahedra) decreases as the

total Li content increases in Li-stuffed garnets.1,2 The above information was used to

estimate the Li occupancy on each site in our Li5+2xLa3Ta2-xYxO12 phases, as shown in

Table 4.1 and 4.2.

4.2.2 Li ion conductivity as a function of Y-doping and temperature.

Typical AC-impedance data of Li5+2xLa3Ta2-xYxO12 (x = 0.05-0.5) recorded at room

temperature are shown in Figures 4.9-4.13, and they show one or two semicircle at the

high-frequency side and a spike at the low-frequency side. This behavior is typical for Li-

stuffed garnet-type solid Li-ion electrolytes.50, 56, 62 The data are fitted using an equivalent

circuit of resistances and constant phase elements (CPEs), which are shown as insets in

Figures 4.9-4.13. The open symbols and the solid lines represent the measured and fitted

data, respectively. The equivalent circuits represent the contributions from the electrical

bulk, grain-boundary, and electrode responses. The lower doped Li5+2xLa3Ta2-xYxO12 (x =

0.05 and 0.10) were fitted using the equivalent circuits comprising two sets of parallel R

(resistance) and CPE components and a series CPE component compared to the higher

doped Li5+2xLa3Ta2-xYxO12 (x = 0.20-0.5) where only one set of R and CPE parallel circuit

was used. The calculated bulk (Cb) and grain-boundary (Cgb) capacitance are in the range

of 10-11 – 10-12 and 10-8 F, respectively which are in agreement with values expected for

these materials. The fitting results are provided in Tables 4.3 and 4.4).123 The goodness of

fit is indicated by χ2, where lower χ2 indicates greater reliability of the fit. In all cases very

low χ2 values, in the order of 10-4, were obtained, indicating excellent fits.

71


xYxO12. The open symbols represent the collected data and the solid lines represent

the fitting. The inset figure indicates the equivalent circuits used for fitting analysis.



72











73











74

Table 4.3 AC impedance fitting results of Li5+2xLa3Ta2-xYxO12 (x = 0.05 and 0.10).



Table 4.4 The room temperature AC impedance fitting results of Li5+2xLa3Ta2-xYxO12

(x = 0.20, 0.25, and 0.50). Reprinted with permission from Inorg. Chem. 2015, 54,

6968-6977. Copyright (2015) American Chemical Society.125

x Rb (Ω) CPEb (F) Cb (F) CPEel (F) χ2

0.20 2.00 x 104 2.83 x 10

-9 1.12 x 10

-11 6.74 x 10

-6 0.0008

0.25

0.50

8.89 x 104

2.26 x 104

1.77 x 10-9

8.48 x 10-9

8.06 x 10-12

2.10 x 10-12

3.24 x 10-6

3.00 x 10-6

0.0008

0.0001

The electrical conductivity studies at the temperature range of 23 to 325 °C are given in

the Arrhenius plots (Figure 4.14). As the temperature increases, the total (bulk + grain-

boundary) conductivity increases as expected for garnet-type materials. In addition, the

conductivity increases with increase in Li content and Y-doping, confirming once again

that Li stuffing helps to increase the ionic conductivity in garnets.57, 138 A decrease in

activation energy from 0.43 to 0.33 eV (calculated at 23-325 °C) was observed with the

increase in x in Li5+2xLa3Ta2-xYxO12, as shown in Table 4.5 along with the total conductivity

values at 23 °C.

x Rb (Ω) CPEb (F) Cb (F) Rgb (Ω) CPEgb (F) Cgb (F) CPEel (F) χ2

0.05 1.69 x 104 5.70 x 10

-10 1.37 x 10

-11 2.18 x 10

4 1.59 x 10

-7 2.07 x 10

-8 3.33 x 10

-6 0.0002

0.10 1.49 x 104 1.40 x 10

-9 1.55 x 10

-11 4.84 x 10

4 3.34 x 10

-8 1.32 x 10

-8 9.05 x 10

-7 0.0003

75

Figure 4.14 Arrhenius plots showing the conductivity variation against temperature

of a) as-prepared Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Reprinted with permission

from Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical

Society.125

Table 4.5 The room temperature (25 °C) conductivity and activation energies of

Li5+2xLa3Ta2-xYxO12 calculated at 23-325 °C. Reprinted with permission from Inorg.

Chem. 2015, 54, 6968-6977. Copyright (2015) American Chemical Society.125

x in Li5+2xLa3Ta2-xO12 σ23 °C(Scm-1) Ea (eV)

0.05 9.56 x 10-6 0.43

0.10 1.04 x 10-5 0.40

0.20 1.07 x 10-5 0.44

0.25 2.82 x 10-5 0.43

0.50 1.26 x 10-4 0.37

0.75 1.83 x 10-4 0.33

1.5 2.0 2.5 3.0 3.5-6

-5

-4

-3

-2

-1

Li6.5

Li6

Li5.5

Li5.4

Li5.2

Li5.1

log

10 (

Scm

-1)

1000/T (K-1)

350

300

250

200

150

100

50

Temperature (°C)

76

The conductivity trend as a function of Li content is also consistent with the results

of the 7Li magic angle spinning (MAS) NMR studies. The 7Li MAS NMR spectra, which

are expressed against the chemical shift value of solid LiCl are shown in Figure 4.15. A

characteristic single peak near 0 ppm is observed and is the typical behavior of garnet-type

compounds.139, 140 As the Y-doping and the Li content increase, the peak becomes

narrower, indicating faster mobility of Li, and confirming that the Li-ion mobility is

enhanced as a function of lithium content in the materials.

Figure 4.15 7 Li and MAS NMR of Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Chemical

shift was measured against solid LiCl. A spinning frequency of 5 kHz was used.



-100 -80 -60 -40 -20 0 20 40 60 80 100

X = 0.25

X = 0.50

X = 0.75

X = 0.20

X = 0.10

X = 0.05

Chemical shift (ppm)

77

It has been reported that during sintering, the inclusion of Al from the alumina crucible

into the garnet Li7La3Zr2O12, appeared to stabilize the cubic structure versus tetragonal

phase in that material.72, 141 Therefore, 27Al MAS NMR measurements were performed on

all samples and indicated the presence of Al in Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). As

seen in Figure 4.16, there are two peaks in the 27Al MAS NMR spectra. The peak appearing

in the range of 10−15 ppm represents octahedrally coordinated Al and has been proposed

previously to correspond to small amounts of LaAlO3.8 However, recently, it has been

shown that Al originating from the alumina crucible can be included in the garnet structure

by residing on the octahedral 48g site and sharing that site with lithium.15 The 27Al MAS

NMR peaks appearing in the range of 70-80 ppm, correspond to Al in 4-fold coordination

based on the chemical shift.8 It has been suggested previously that these four coordinated

Al atoms are located in some of the 24d tetrahedral sites in the garnet structure, sharing

that site with lithium.72, 141 Therefore, the 27Al MAS data indicate the inclusion of trace

amounts of aluminum in our materials. However, it is unclear if this Al inclusion plays any

role in the stabilization of the cubic structure in the investigated garnet phases.

78

Figure 4.16 27Al MAS NMR of Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). Chemical shift

was measured against Al(NO3)3. A spinning frequency of 5 kHz was used. Reprinted

with permission from Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015) American

Chemical Society.125

4.2.3 Porosity and sinterabiliy trends as a function of Y-doping

Employing scanning electron microscopy (SEM) and porosity/density calculations, a

correlation has been established between the Y (and Li) content and the sinterability of

Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75). The scanning electron micrographs for Li5+2xLa3Ta2-

xYxO12 are shown in Figure 4.17. A high degree of porosity is present in the sintered pellets

79

of garnets that have a smaller level of doping, while the pellets become considerably denser

upon increasing the Y (and Li) content. The SEM results are also qualitatively consistent

with the percent-porosity values, obtained based on Archimedes principle. The porosities

of sintered pellets for x = 0.05, 0.1, 0.2 and 0.25 are close to each other and range between

28 % and 35 %. However, there is a significant decrease in porosity for x = 0.5 and 0.75,

where the calculated porosities are about 0.5 % and 1.6 %, respectively. These results are

found to be consistent with the densities of sintered pellets (Table 4.6) obtained using the

Archimedes method. The theoretical density (d) was computed from the PXRD. As shown

in Table 4.6, the densities of the x = 0.5 and 0.75 are significantly greater than the densities

obtained for x = 0.05, 0.1, 0.2 and 0.25, confirming the direct correlation between the

doping level and sinterability in these materials.

80

Figure 4.17 Morphological studies of Li5+2xLa3Ta2-xYxO12 (x = 0.05–0.75) pellets using

SEM microscopy. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


x = 0.50

x = 0.20

x = 0.05 x = 0.10

x = 0.25

x = 0.75

5 µm 5 µm

5 µm 5 µm

5 µm 5 µm

81

Table 4.6 Calculated and measured density results of Li5+2xLa3Ta2-xYxO12 from

PXRD and Archimedes methods. Reprinted with permission from Inorg. Chem. 2015,

54, 6968-6977. Copyright (2015) American Chemical Society.125

4.2.4 Structural stability and chemical compatibility with aqueous LiCl solution

The high structural stability of Li5+2xLa3Ta2-xYxO12 was studied by in-situ PXRD in the

temperature range of 30-700 °C as shown in Figures 4.18 and 4.19. The thermal expansion

coefficient (TEC) of Li5+2xLa3Ta2-xYxO12 (x = 0.5) with temperature was calculated using

the lattice parameters derived from in-situ PXRD (Figure 4.19). The formula used for TEC

calculation is shown in equation 4.1:

a =a - a

0

a0

æ

èç

ö

ø÷

1

T -T0

æ

èç

ö

ø÷ (4.1)

where α, T and a represent the thermal expansion coefficient, temperature, and lattice

parameter at that temperature.142 The notations a0, and T0 represent the corresponding

values at room temperature. A thermal expansion coefficient of 7.25 x 10-6 K-1 was

observed for Li6La3Ta1,5Y0.5O12.

x in Li5+2xLa3Ta2-xYxO12 Measured

Density (gcm-3)

Theoretical

Density (gcm-3)

Archimedes

Porosity (%)

0.05 3.463 6.320 29

0.10 3.114 6.279 35

0.20 3.353 6.176 30

0.25 3.479 6.104 28

0.50 4.919 5.928 0.5

0.75 4.776 5.772 1.6

82


thermal stability. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


10 20 30 40 50 60 70 80

30 oC after heating

700 oC

600 oC

400 oC

200 oC

30 oC

In

ten

sity (

a.u

.)

Two theta (degrees)

83


thermal stability. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


To determine the stability in aqueous environments, the garnet materials were treated

in water. The water-treated materials were then studied by variable-temperature PXRD to

determine if there were any changes in their structure or stability. Figure 4.20 shows the

10 20 30 40 50 60 70 80

30 oC after heating

In

ten

sity (

a.u

.)

Two theta (degrees)

900 oC

800 oC

700 oC

600 oC

400 oC

200 oC

30 oC

84

corresponding data for the x = 0.5 phase and indicates that the water-treated material

maintains the cubic structure and is also structurally stable up to 700 °C.

Figure 4.20 In situ PXRD patterns of water-treated Li6La3Ta1,5Y0.5O12 showing the

structural stability. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-

6977. Copyright (2015) American Chemical Society.125

Compatibility tests with 1 M aqueous LiCl solution were done in order to explore the

potential application of Li5+2xLa3Ta2-xYxO12 as a protective layer for the Li anode in lithium

10 20 30 40 50 60 70 80

30 oC after heating

700 oC

600 oC

500 oC

400 oC

300 oC

200 oC

100 oC

30 oC

Inte

nsity (

a.u

.)

Two theta (degrees)

85

aqueous batteries. A highly Li+ conductive material that has chemical compatibility with

aqueous lithium salts is desired for such application. The pH change was monitored for 1

week, and results are shown in Figure 4.21. The change in pH can indicate the exchange

of Li+ with proton in Li5+2xLa3Ta2-xYxO12.

Figure 4.21 Changes in the pH during the treatment of Li5+2xLa3Ta2-xYxO12 (x = 0.50

and 0.75) in 1 M LiCl solution. Reprinted with permission from Inorg. Chem. 2015,

54, 6968-6977. Copyright (2015) American Chemical Society.125

An increase in the pH value was observed over the first 2 days, reaching a maximum value

of 11−11.5 from an initial value of 6, and remained almost constant for the rest of the week.

A similar change in pH was reported in other garnets in water.143, 144 A comparative

impedance study of the aqueous LiCl-treated garnets was also performed. Figure 4.22

0 1 2 3 4 5 6 7 8

6

7

8

9

10

11

12

x = 0.50

x = 0.75

pH

Time (Days)

86

shows the AC impedance plots of LiCl-treated Li5+2xLa3Ta2-xYxO12 (x = 0.50 and 0.75)

measured at 75 °C. The grain-boundary impedance was found to be increasing after the

aqueous LiCl treatment in all the cases. The corresponding Arrhenius plots, in Figure 4.23

indicates a conductivity drop compared to the as-prepared garnets (Figure 4.14).

Figure 4.22 Impedance spectra of Li5+2xLa3Ta2-xYxO12 (x = 0.50 and 0.75) after

treatment with 1 M LiCl solution measured at 75 °C. The inset magnifies the data for

x = 0.50 sample to show the small semi-circle at the high frequency side. Reprinted

with permission from Inorg. Chem. 2015, 54, 6968-6977. Copyright (2015) American

Chemical Society.125

87

Figure 4.23 Arrhenius plots showing the conductivity variation against temperature

of Li5+2xLa3Ta2-xYxO12 (x = 0.5, and 0.75) after treatment with 1 M LiCl solution.



This may be partly due to the replacement of Li+ in the garnet structure with H+ from the

aqueous solution. The pH change during the aqueous LiCl treatment and the presence of –

OH peaks in the FTIR spectra (~ 3500 cm-1 for –OH stretching and ~ 1500 cm-1 for –OH

bending) shown in Figure 4.24 indicate that there was an ion exchange between the garnet

and water during aqueous LiCl treatment.

1.5 2.0 2.5 3.0 3.5-6

-5

-4

-3

-2

LiCl treated - Li6.5

LiCl treated - Li6

log

10 (

Scm

-1)

1000/T (K-1)

350

300

250

200

150

100

50

Temperature (°C)

88

Figure 4.24 FTIR spectra of water-treated Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and 0.75).



A cross sectional SEM imaging was also done to understand the microstructure of

the pellets treated in aqueous LiCl as shown in Figure 4.25. The SEM data indicated a

greater degree of porosity in the samples after aqueous LiCl treatment compared to the

untreated materials (Figure 4.17). The EDX patterns were recorded (Figure 4.25) in search

for chlorine to examine the potential entrapment of LiCl. However, no evidence was found

to indicate the presence of significant amount of chlorine in these samples (Figure 4.25).

Nevertheless, even after the drop in conductivity, the measured values are still significantly

greater than the conductivity of the parent undoped compound.50

4000 3500 3000 2500 2000 1500 1000 500

Wave number (cm-1)

% T

ran

sm

itta

nce

(%

)

X = 0.25

x = 0.50

x = 0.75

89

Figure 4.25 SEM images and EDX results for Li5+2xLa3Ta2-xYxO12 (x = 0.5 and 0.75)

after treatment with 1 M LiCl solution. Reprinted with permission from Inorg. Chem.

2015, 54, 6968-6977. Copyright (2015) American Chemical Society.125

The PXRD data of the samples after the aqueous LiCl test shows that there is no phase

change, indicating high structural stability of the garnets, as shown in Figure 4.26. Thus,

the garnet-type Li5+2xLa3Ta2-xYxO12 materials show a high stability in an aqueous LiCl

environment, making them suitable candidates for potential application in Li aqueous

batteries.

90

Figure 4.26 PXRD of Li5+2xLa3Ta2-xYxO12 (x = 0.5 and 0.75) after reaction with 1 M

LiCl solution. Reprinted with permission from Inorg. Chem. 2015, 54, 6968-6977.


4.2.3 Electrochemical stability

Cyclic voltammetry (CV) was carried out to understand the electrochemical stability of

Li6.5La3Ta1.25Y0.75O12 in glovebox. A scanning rate of 100 mVs-1 between -0.2 to +6 V

were used for the cell, (-)Li/Li6.5La3Ta1.25Y0.75O12/Au(+). The CV measurements carried

out at 50 and 100 °C are presented in Figures 4.27 and 4.28. The anodic and cathodic

currents correspond to the lithium dissolution (Li → Li+ + e-) and deposition (Li+ + e- →

Li). No other peaks (high current) were observed corresponding to the decomposition of

electrolytes in the voltage range scanned, indicating the electrochemical stability of

Li6.5La3Ta1.25Y0.75O12 up to 6 V. Fluctuations in the vacuum level caused noise in the CV

spectra.

10 20 30 40 50 60 70 80

x = 0.75

x = 0.50

Two theta (degrees)

I

nte

nsity (

a.u

.)

91


V vs. Li/Li+ with a scanning rate of 100 mVs-1.


V vs. Li/Li+ with a scanning rate of 100 mVs-1.

0 1 2 3 4 5 6-3.0x10

-6

-2.0x10-6

-1.0x10-6

0.0

1.0x10-6

2.0x10-6

3.0x10-6

Cu

rre

nt (A

cm

-2)

Voltage (V vs Li/Li+)

0 1 2 3 4 5 6-1.0x10

-4

-5.0x10-5

0.0

5.0x10-5

1.0x10-4

Cu

rre

nt (A

cm

-2)

Voltage (V vs Li/Li+)

92

4.3 Summary125

The garnet-type Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤ 0.75) was studied to correlate the Y and

Li content with conductivity, porosity and sinterability. 7Li magic angle spinning NMR

shows a direct correlation between the Li-content and lithium mobility in these materials,

which is consistent with the impedance spectroscopy studies that indicate higher

conductivity for materials with greater Li content. The conductivity of these materials over

a wide temperature range, from 23 °C to 325 °C, has also been studied. In addition, with

the aid of scanning electron microscopy and Archimedes calculations of porosity and

density, a correlation has been established between the sinterability and dopant content,

where an increase in sinterability is observed as the Li (and Y) contents increase. An

electrochemical stability window of up to 6V were observed for Li6.5La3Ta1.25Y0.75O12.

Furthermore, stability tests were performed to determine the compatibility of these

materials with aqueous LiCl, indicating some decrease in Li-conductivity, but very high

stability, making Li5+2xLa3Ta2-xYxO12 a potential candidate for further investigation as

protective layer for lithium electrodes in lithium-air batteries.

93

Chapter Five: Effect of excess Li on the structural and electrical properties of garnet-

type Li6La3Ta1.5Y0.5O12

The results discussed in this chapter is a published work, “Effect of excess Li on the

structural and electrical properties of garnet-type Li6La3Ta1.5Y0.5O12” by Sumaletha

Narayanan, Gregory Thomas Hitz, Eric D Wachsman, and Venkataraman Thangadurai, in

J. Electrochem. Soc. 2015, 162, A1772-A1777.


Volatility of lithium during preparation of lithium-stuffed garnet-type metal oxide solid Li

ion electrolytes is a common problem, which affects phase formation, ionic conductivity,

mechanical strength and density. Therefore, several researchers commonly compensate the

lithium oxide loss by adding a 10 wt.% excess lithium precursors.48, 62, 146 Traditional solid-

state techniques are the most commonly used to synthesize garnet-type metal oxides, but

usually require the highest sintering temperature resulting in the production of large

particles.50, 56, 57, 60 It has been shown that Al migrates from the alumina crucible used for

garnet synthesis, causing significant effects in the material including stabilization of the

cubic phase.72, 147-149 Lithium salts such as Li2O, Li3BO4, Li3PO4 and Li4SiO4 have been

shown to significantly affect the electrical and microstructural properties.148, 150, 151 Liu et

al. showed that excess lithium (0 to 50 mol.%) added during synthesis of

Li6.5La3Ta0.5Zr1.5O12 increase the density and Li ion conductivity by the liquid phase

formation during sintering.150 This seems due to reaction between the excess lithium and

aluminum impurity from the crucible, which was clearly visible from the PXRD studies.

The optimization of Li2O addition for Li6.75La3Zr1.75Ta0.25O12 has been performed by Li et

94

al. and found that 6 wt.% excess Li2O is required for best conducting ceramic (6.4 x 10-4

Scm-1 at room temperature), which is about 3 times higher than that of

Li6.75La3Zr1.75Ta0.25O12 prepared without adding excess Li2O.152 They also found that the

glassy-like phase formed at the grain-boundaries increases the density of material as well

as helps to decrease the grain-boundary and improve the total Li+ conductivity. Sakamoto

group has explored high pressure and high temperature method to prepare dense garnet-

type Li ion electrolytes and such a technique can be used to control the loss of Li content.153

This chapter, describes the systematic investigation of effect of excess LiNO3 (2.5

to 15 wt.%) addition during the ceramic synthesis on the structural, microstructural and

electrical properties of fast Li ion conducting garnet-type Li6La3Ta1.5Y0.5O12 to understand

the role of Li content and its optimization for high ionic conductivity. Powder X-ray

diffraction (PXRD) confirmed that cubic phase was formed in all tested cases, and there is

no significant variation in lattice parameter with amount of excess LiNO3 used. However,

increasing amounts of excess lithium decreased inter-particle contact and increased grain

growth during sintering, producing sharply varied microstructures. PXRD showed no

secondary phase and scanning electron microscopy (SEM) analysis showed rather uniform

morphology and absence of “glassy” materials at the grain-boundaries. The presence of

Li2CO3 in all aged Li6La3Ta1.5Y0.5O12 samples prepared using excess LiNO3 was observed

from Raman microscopy. The bulk Li ion conductivity was found to increase with amount

of excess lithium, reaching a maximum room temperature conductivity of 1.62 x 10-4 Scm-

1 for the sample prepared using 10 wt.% excess LiNO3.

95


5.2.1 Structure and density analysis

Figure 5.1 shows the PXRD patterns of Li6La3Ta1.5Y0.5O12 prepared using 2.5 to 15 wt.%

excess LiNO3 along with the calculated pattern for the cubic phase of Li5La3Ta2O12

(Inorganic Crystal Structure Database (ICSD) # 01-074-9856). PXRD show the formation

of cubic phase in all the investigated Li6La3Ta1.5Y0.5O12 samples. There is no secondary

phase observed in the PXRD. It is found that the crystal structure is stable irrespective of

the change in the amount of excess lithium content, which was added to avoid Li loss

during high temperature sintering. The exact amount of Li in the investigated compound

was unable to be quantified in the present study due to lack of available experimental

method. However, there is no significant change in the cubic lattice constant as a function

of Li content (Table 5.1). The theoretical and measured density results of garnet-type

Li6La3Ta1.5Y0.5O12 prepared using 2.5 to 15 wt.% excess LiNO3 obtained using PXRD and

He pycnometer, respectively, are listed in Table 5.1. The He pycnometer density was found

to be in the range of 4.84 - 5.40 gcm-3 for Li6La3Ta1.5Y0.5O12 prepared using 2.5 to 15 wt.%

excess LiNO3 (Table 5.1), while the theoretical density obtained using XRD is ~ 6 gcm-3.

The density of the samples varied between 81 and 91%. The discrepancy in experimental

density values can be attributed to the porosity of garnet oxides.

96

Table 5.1 Comparison of densities calculated from PXRD data and measured from

He Pycnometer and the lattice parameter (a) of Li6La3Ta1.5Y0.5O12 with excess amount

of lithium (2.5 - 15 wt.% LiNO3) added during solid-state synthesis. Reproduced by

permission of ECS-The Electrochemical Society.145

Excess Li-salt

(wt. %)

Theoretical

density

(gcm-3)

Measured

density

(gcm-3)

Lattice

parameter, a (Å) Density (%)

2.5 5.964 4.839(5) 12.911(1) 81

5

7

6.007 5.101(5) 12.880(3) 85

5.964 5.398(2) 12.911(1) 91

10 5.963 5.161(4) 12.912(1) 87

15 5.985 5.095(3) 12.896(1) 85

Figure 5.1 The powder X-ray diffraction patterns of Li6La3Ta1.5Y0.5O12 prepared via

solid-state synthesis method with varying amounts of excess (a) 2.5, (b) 5, (c) 7, (d) 10

and (e) 15 wt.% LiNO3 and a calculated pattern for cubic garnet (Li5La3Ta2O12, ICSD

# 01-074-9856 where ICSD stands for Inorganic Crystal Structure Database).

Expected peak for Li2CO3 is also marked with a dotted line at ~ 31.92° (JCPDS # 22-

1141 where JCPDS stands for Joint Committee on Powder Diffraction Standards).

Reproduced by permission of ECS-The Electrochemical Society.145

97

5.2.2 Microstructure and chemical composition analysis

Lithium dependence of microstructure is demonstrated by scanning electron microscopy

(SEM) as shown in Figure 5.2. The secondary electron images were recorded at different

magnifications. Imaging was performed on slices of isostatically pressed pellets sintered at

1100 °C for 6h. Significant porosity is seen at the sample with 5 wt.% excess LiNO3 and

below, with a microstructure possibly indicating liquid phase sintering. This observation

supports the lower density obtained from He Pycnometer (Table 5.1). Grain growth

generally increased from 7 to 15 wt.% excess LiNO3, though 10 % excess LiNO3 was out

of the trend with somewhat smaller particle size.

98

Figure 5.2 Scanning electron micrographs of Li6La3Ta1.5Y0.5O12 with varying

amounts of excess (a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3 (corresponding

higher magnification images are shown in the same row). Reproduced by permission

of ECS-The Electrochemical Society.145

In general, the inter-particle contact decreased with increased excess LiNO3, while the

extent of grain growth increased. The beneficial impact of excess LiNO3 on grain growth

1000 µm 50 µm 10 µm

a

b

c

d

e

99

may have to be balanced with the lessened inter-particle contact. It seems that 7 wt.%

excess LiNO3 is close to this balancing point. Also, less porosity with inter-particle contact

and maximum density of 5.4 gcm-3 was observed for 7 wt.% LiNO3 excess sample

compared to others (Table 5.1).

The elemental distribution (EDS mapping) of Li6La3Ta1.5Y0.5O12 prepared using

2.5 to 15 wt.% excess LiNO3 is shown in Figure 5.3, indicating the presence of heavy

elements in the composition along with the Al contamination from the crucible used for

the synthesis. The EDS mapping further indicates that there is no significant segregation

of Al or other elements in Li6La3Ta1.5Y0.5O12 prepared using 2.5 to 15 wt. %. The presence

of Al contamination was further confirmed from the 27Al MAS NMR studies, as shown in

Figure 5.4. The peak close to 0 ppm indicates the octahedral coordination of Al, while the

other peak close to 70 ppm indicates the tetrahedral coordination of Al.57, 72, 139, 154, 155

Geiger et al. proposed that the octahedral sites occupied by Al3+ could lead to distortion.156

It is possible that Al3+ substitutes for Li+ sites and creates lithium vacancies in the structure

which affect the structural stabilization.157

100

Figure 5.3 SEM image and corresponding elemental mapping of Li6La3Ta1.5Y0.5O12

with varying amounts of excess (a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3.


101

Figure 5.4 27Al MAS NMR of Li6La3Ta1.5Y0.5O12 with varying amounts of excess (a)

2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3 measured against Al(NO3)3 as standard.


Li7La3Zr2O12 garnet needs a critical concentration of 0.204 moles of Al to stabilize cubic

structure and it exists as tetragonal polymorph below that concentration limit.157 Though

EDS mapping showed no significant segregation of Al at grain boundaries in this study, a

small amount of Al could also be present at grain-boundaries as an amorphous phase, which

acts as sintering aid and prevents Li volatilization during sintering.23 There was no

indication of aluminum containing impurity peaks in PXRD (Figure 5.1), which could be

due to the detection limit of the technique. Chemical shift value of 7Li MAS NMR was

102

reported against LiCl standard (Figure 5.5). In all the cases, a single peak was obtained

close to 0 ppm which is typical for Li-stuffed garnet-type oxides.56, 57, 72, 158

Figure 5.5 7Li MAS NMR of Li6La3Ta1.5Y0.5O12 with varying amounts of excess (a)

2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3. The standard material used was LiCl.


Raman spectroscopy was performed to reveal the presence of any surface phases and

fingerprint the garnet phase. Figure 5.6 shows that the spectrograph matches that of cubic

garnet regardless of excess lithium amount.159, 160 However, it is also apparent that all

samples show some amount of Li2CO3 formation on the surface, as evidenced by the strong

peak at 1090 cm-1. However, the expected carbonate peak (2θ ~ 31. 92°, JCPDS # 22-1141

where JCPDS stands for Joint Committee on Powder Diffraction Standards) is not present

103

in PXRD patterns (Figure 5.1). The detection of Li2CO3 in Raman, but not in PXRD,

signifies that the Li2CO3 is likely a minority phase existing as a surface layer formed in a

reaction between lithium garnet and atmospheric CO2.

Figure 5.6 Raman spectroscopy of Li6La3Ta1.5Y0.5O12 with varying amounts of excess

(a) 2.5, (b) 5, (c) 7, (d) 10 and (e) 15 wt.% LiNO3. For comparison, Raman spectrum

of commercial Li2CO3 is also shown. Reproduced by permission of ECS-The

Electrochemical Society.145

104

The presence of Li2CO3 due to the sample exposure to atmosphere is also reported recently

using the techniques such as X-ray photoelectron spectroscopy (XPS) and soft X-ray

absorption spectroscopy (sXAS) 161. There is no correlation between excess lithium in the

sample and intensity of the Raman peak signifying Li2CO3. The samples may have had

slightly different exposures to the atmosphere. Lithium vibrations are expected between

300 and 600 cm-1, more specifically in Figure 5.6, the internal modes of octahedrally

coordinated lithium occur at 200-300 cm-1 and that of tetrahedrally coordinated lithium

occur at 350-600 cm-1.148, 162 The peak at ~ 750 cm-1 could be due to the stretching modes

of TaO6 octahedra.163

5.2.3 Electrical characterization

Typical AC impedance spectra of Li6La3Ta1.5Y0.5O12 prepared using 2.5 to 15 wt.%

excess LiNO3 recorded in air at 50 °C are shown in Figure 5.7-5.11, and the equivalent

circuit used for fitting are shown as insets. Two semicircles and a spike due to lithium ion

blocking electrodes are seen for 2.5 wt.% excess sample, and two sets of parallel resistor-

capacitor and a series capacitor was used for the fitting (Figure 5.7). Only one semicircle

and a spike is seen in case of Figures 5.8 and 5.9. The fitting results are provided in Table

5.2 and the capacitance values show that the conductivity contributions are from the bulk

and grain-boundaries. The shape of impedance spectra shows the typical behavior of

garnet-type oxides.148, 150

105

Figure 5.7 Typical AC impedance spectrum of Li6La3Ta1.5Y0.5O12 with 2.5 wt.%

excess LiNO3 measured in air at 50 °C. The open symbols and solid lines represent

the measured and fitted data, respectively. The equivalent circuit used to fit the

impedance data is shown in the inset figure. Reproduced by permission of ECS-The







106











0.0 2.0x103

4.0x103

6.0x103

0.0

-2.0x103

-4.0x103

-6.0x103

10 kHz

1 MHz

Z// (

cm

)

Z/ ( cm)

107






Table 5.2 The AC impedance fitting results measured at 50 °C of Li6La3Ta1.5Y0.5O12

with excess amount of lithium (2.5, 5 and 7 wt.% LiNO3) added during solid-state

synthesis. Reproduced by permission of ECS-The Electrochemical Society.145

Excess

Li-salt

(wt.%)

Rb (Ω) CPEb (F) n Cb (F) n CPEgb (F) Cgb (F) CPEel (F) χ2

2.5 4.93 x 104 1.98 x 10

-10 0.81 1.24 x 10

-11 0.81 1.35 x 10

-8 2.02 x 10

-9 6.75 x 10

-7 3 x 10

-4

5 2.97 x 103 1.34 x 10

-8 0.57 6.79 x 10

-12 - - - 6.63 x 10

-7 2 x 10

-4

7 8.18 x 103 5.30 x 10

-10 0.74 6.32 x 10

-12 - - - 2.05 x 10

-6 4 x 10

-4

Figure 5.12 shows the Arrhenius plots for bulk Li ion conductivity of the

Li6La3Ta1.5Y0.5O12 prepared with varying excess of LiNO3. The maximum conductivity

obtained is for the 10 and 15 wt.% LiNO3 excess samples and a drop in conductivity of

0.0 3.0x103

6.0x103

9.0x103

0.0

-3.0x103

-6.0x103

-9.0x103

1 kHzZ// (

cm

)

Z/ ( cm)

100 kHz

108

couple of orders of magnitude was observed for all other cases (2.5 to 7 wt. %). The

conductivity obtained for 10 wt.% Li sample is 1.62 x 10-4 Scm-1 at 24 °C, while a lower

conductivity of 3.78 x 10-6 Scm-1 was found for the least conductive sample, with 2.5 wt.%

LiNO3 excess. The higher conductivity samples show a decrease in activation energy over

the temperature range of 24-325 °C, as shown in Figure 5.12.

Figure 5.12 The Arrhenius plots of Li6La3Ta1.5Y0.5O12 (2.5 – 15 wt.% excess LiNO3)

showing the bulk Li ion conductivity as a function of temperature. The closed and

open symbols represent the heating and cooling cycles, respectively. Reproduced by

permission of ECS-The Electrochemical Society.145

1.5 2.0 2.5 3.0 3.5-8

-7

-6

-5

-4

-3

-2

-1

2.5 % (Ea = 0.42 eV)

5 % (Ea = 0.40 eV)

7 % (Ea = 0.39 eV)

10 % (Ea = 0.32 eV)

15 % (Ea = 0.34 eV)

log

10 (

Scm

-1)

1000/T (K-1)

350

300

250

200

150

100

50

Temperature (°C)

109

The conductivity increase with increase in excess salt can be attributed to the fact that the

excess LiNO3 added during synthesis acts as a sintering aid, which increases the grain-to-

grain contact and also decreases the grain-boundary resistance. A comparison graph of

calculated and measured densities from PXRD data using the PROSZKI program and He

Pycnometry respectively, along with activation energy as a function of different excess Li

content of Li6La3Ta1.5Y0.5O12 (2.5-15 wt.% excess LiNO3) is shown in Figure 5.13. It is

noticeable that with the increase in measured density, activation energy tends to decrease,

indicating that the excess Li salt has a positive role in increasing the density of the

investigated Li-stuffed Li6La3Ta1.5Y0.5O12.

Figure 5.13 Correlation of density and activation energy (calculated at 24-325 °C)

with the excess LiNO3 wt.% (2.5 – 15 wt. %) used for the synthesis of

Li6La3Ta1.5Y0.5O12. It shows that the activation energy decreases with increase in

density of the ceramics. Reproduced by permission of ECS-The Electrochemical

Society.145

2 4 6 8 10 12 14 16

4.8

5.1

5.4

5.7

6.0

6.3

Pycnometer density

Excess Li in Li6La

3Ta

1.5Y

0.5O

12 (wt %)

De

nsity (

gcm

-3)

0.30

0.35

0.40

0.45

0.50

Ea (e

V)

Theoretical density

110

5.3 Summary145

Li6La3Ta1.5Y0.5O12 was produced with varying excess amounts of LiNO3 and characterized

by powder X-ray diffraction, scanning electron microscopy, AC impedance and Raman

spectroscopy. The desired cubic phase was obtained for all the compositions and the lattice

parameter was essentially constant irrespective of the amount of excess lithium in the

composition. The amount of excess lithium had an insignificant effect on the lattice

structure or propensity to develop lithium carbonates in air, but sinterability and

conductivity showed clear maxima in the range tested. Li6La3Ta1.5Y0.5O12 with 10 and 2.5

wt.% excess Li showed the highest and lowest conductivity of 1.62 x 10-4 Scm-1 and 3.77

x 10-6 respectively, at 24 °C among all other members. This study shows that the

optimization of excess LiNO3 addition can dramatically improve the properties of final

sintered Li-stuffed garnets.

111

Chapter Six: Evaluation of fundamental transport properties of Li-excess garnet-

type Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and 0.75) electrolytes using AC impedance

and dielectric spectroscopy

The results discussed in this chapter is a published work, “Evaluation of fundamental

transport properties of Li-excess garnet-type Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and 0.75)

electrolytes using AC impedance and dielectric spectroscopy” by Ashok Kumar Baral,

Sumaletha Narayanan, Farshid Ramezanipour, and Venkataraman Thangadurai, in Phys.

Chem. Chem. Phys. 2014, 16, 11356-11365.


It is important to understand the Li+ ion conduction mechanism in garnets to further

improve Li+ ion conduction for an advanced all-solid-state Li ion battery. It was predicted

that the Li+ ions at octahedral sites in the garnet structure, mainly attribute to the fast ionic

conduction and tetrahedral site Li+ ions seem to be responsible for maintaining the stability

of the garnet framework structure.137 Solid-state 7Li MAS NMR and density functional

theory (DFT) studies further support above proposed Li+ ion conduction mechanism in the

garnet structure.82, 164 So far, several studies have been carried out to understand the Li+ ion

dynamics in these Li-stuffed garnets using solid-state techniques such as neutron

diffraction, nuclear magnetic resonance, electrochemical AC impedance spectroscopy and

DFT calculations.66, 80-82, 130, 137, 165-173 In this chapter, Li-excess garnet-type cubic (space

group Ia-3d) Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and 0.75) have been studied in the

temperature range of -50 to 50 °C using electrochemical AC impedance spectroscopy to

understand the fundamental electrical transport properties including ionic conductivity,

112

dielectric constants, loss tangent, and relaxation time constants. The synthesis of garnet

materials was carried out using conventional ceramic solid-state synthesis method.

A correlation has been established between the excess Li-content and the Li-ion

migration pathways. The dielectric loss tangent (tan δ) for all samples exhibits a relaxation

peak corresponding to the dielectric loss because of dipolar rotations due to the Li+

migration. Comparing the modulus analysis of Li-excess garnets with fluorite-type oxygen

ion conductors, we propose the local migration of Li+ ions between octahedral sites around

the “immobile” Li+ ions in tetrahedral (24d) sites. In the samples with x = 0.25 and 0.5, Li+

ions seem to jump from one octahedral (96h) site to another bypassing the tetrahedral (24d)

site between them (path A), both in local and long-range order migration processes, with

activation energies of ~ 0.69 eV and 0.54 eV, respectively. For the x = 0.75 member, Li+

ions exhibit mainly long-range order migration, with activation energy of 0.34 eV, where

the Li hopping between two octahedral sites occurs through the edge which is shared

between the two LiO6 octahedra and a LiO4 tetrahedron (path B). The present AC

impedance analysis is consistent with the ab initio theoretical analysis of Li-excess garnets

that showed two conduction paths (A and B) for Li ion conduction with different activation

energies.

6.2 Results and discussion

6.2.1 Crystallographic sites and migration pathways in Li-excess garnets

The crystal structure of Li5+2xLa3Ta2-xYxO12 is garnet-type cubic, with a space group Ia-

3d, the detailed structural analysis was described in Chapter 4. A brief summary of garnet-

structure is given here in Figure 6.1 to explain the possible migration pathways of Li+ ions.

113

The crystal structure involves an octahedral arrangement of oxygens around Ta and Y

atoms, which share the same crystallographic position (16a) as shown in Figure 6.1(a). The

La atoms are located in spaces between these octahedra. For each La atom there are eight

oxygens at close distances such that the La-O sublattice involves LaO8 units with

dodecahedral geometry. Figure 6.1(b and c) highlight the Ta(Y)O6 and LaO8 polyhedra in

relation to each other. Also, as shown in Figure 6.1(d and e), there are spaces between

Ta(Y)O6 and LaO8 units that can be occupied by Li atoms. The tetrahedral sites (Figure

6.1(d)) are commonly occupied by cations in an ideal garnet-type material. However, these

sites can only host 3 Li atoms per formula unit, and in Li-excess phases, these sites are not

enough to accommodate all of the Li. Therefore, some Li atoms are located on octahedral

sites (48g or 96h) shown in Figure 6.1(e). Each tetrahedron shares all of its four faces with

four octahedra, while each octahedron shares only two of its faces with two tetrahedra on

the opposite sides.

114

Figure 6.1 (a) The garnet-type structure viewed down the 111 axis. The TaO6

octahedra are shown in purple. Large red spheres are La atoms and small blue and

green spheres represent Li atoms on 24d and 48g sites, respectively. (b) The red lines

highlight the dodecahedral geometry where an La atom is accommodated, (c) the

octahedra, which encompass Ta atoms, (d) the tetrahedral geometry around Li atoms

on 24d positions, and (e) the octahedral geometry where Li atoms on the 48g site

(center of the octahedron) or the 96h site (shifted from the center) are located. Phys.

Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by permission of the PCCP

Owner Societies.130

115

Figure 6.2 shows the connectivity of LiO6 and LiO4 polyhedra. Note that, due to

the repulsion between Li atoms, not all octahedral and tetrahedral sites are occupied at the

same time, and some of these sites remain unoccupied (intrinsic defects). When a

tetrahedral site (24d) and the adjacent octahedral position are occupied, the Li ions in the

octahedral site will shift away from the center (48g site) to reside at the 96h position.52

Also, a tetrahedral site will remain vacant if the two octahedral sites on its both sides are

occupied.

Figure 6.2 The connectivity of LiO6 and LiO4 polyhedra is shown. These polyhedra

are connected in 3-dimentions throughout the unit cell. If there is a Li atom on a

tetrahedral site, the Li atoms at the centers of neighbouring octahedra (48g site) will

be shifted to the 96h position to avoid close Li-Li contacts. If the two octahedral sites

that are on the two sides of a tetrahedral position are occupied, then the tetrahedral

site between them will be empty. Two possible migration pathways (A and B) are

shown by arrows that are discussed in details in the later section. Phys. Chem. Chem.


Societies.130

116

6.2.2 Ionic conductivity

The typical impedance plots of the garnet-type Li5+2xLa3Ta2-xYxO12 with x = 0.25, 0.5, and

0.75 (which are abbreviated as Li5.5-phase, Li6-phase and L6.5-phase, respectively) at -20

°C are shown in Figure 6.3. The impedance plots of Li-excess garnet type materials usually

consist of one semicircle and a spike over the frequency range of 0.01 Hz to 1 MHz at

lower temperatures, while at higher temperatures they exhibit only a spike.56, 69, 157, 174 It is

difficult to separate the bulk and grain-boundary contributions to the total conductivity of

the samples at higher temperatures. In any case, the intercept of spike with the Z/-axis has

been considered to be the total resistance (R) of the sample.

Figure 6.3 (a) Complex impedance plots of samples Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5,

and 0.75) at -20 °C. The equivalent circuit used for fitting z-plots is given in the inset.

Phys. Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by permission of the

PCCP Owner Societies.130

0.0 0.2 0.4 0.6 0.8 1.00.0

0.2

0.4

0.6

0.8

1.0

Li5.5

-phase

Li6-phase

Li6.5

-phase

-Z// (

M

)

Z/ (M)

117

Figure 6.4 shows the Arrhenius plots of electrical conductivity of Li5+2xLa3Ta2-xYxO12 with

x = 0.25, 0.5, and 0.75. The values of conductivity and activation energy (-50 to 50 °C) for

all samples are given in Table 6.1. It can be mentioned that, since only the spike is observed

in the impedance plots at higher temperatures, for better comparison of conductivity with

dielectric analysis, we have focussed only the lower temperature range i.e., -50 to 50 °C.

Figure 6.4 The Arrhenius plots of conductivity Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5, and

0.75). Inset gives the values of activation energy in these samples. Phys. Chem. Chem.


Societies.130

3.0 3.3 3.6 3.9 4.2 4.5

-5

-4

-3

-2

-1

Li6.5

-phase (Ea = 0.36 eV)

Li6hase (E

a = 0.55 eV)

Li5.5

-phase (Ea = 0.65 eV)

log

T

(S

cm

-1K

)

1000/T (K-1)

60 40 20 0 -20 -40 -60

Temperature (oC)

118

Table 6.1 Values of conductivities (σ), activation energies (Ea) and pre exponential

factor (A) for Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.50 and 0.75). Phys. Chem. Chem. Phys.

2014, 16, 11356-11365. Reproduced by permission of the PCCP Owner Societies.130

Garnet σ-20 °C (Scm-1) σ50 °C (Scm-1) Ea (eV) A

Li6.5La3Ta1.25Y0.75O12 9.84 x 10-6 5.84 x 10-4 0.36 1.50 x 105

Li6La3Ta1.5Y0.5O12 3.12 x 10-6 5.83 x 10-4 0.55 5.88 x 107

Li5.5La3Ta1.75Y0.25O12 3.43 x 10-7 1.19 x 10-4 0.65 8.32 x 108

Figure 6.5 presents the conductivity variation with frequency at -50 °C for

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12. The plots of conductivity spectra follow the

Jonscher’s universal power law175 (i.e., n

dc A , where σdc is the DC conductivity,

“ω” is angular frequency, and “n” is the power factor such that 0 ≤ n ≤ 1) as they contain

the DC plateau at lower frequencies followed by frequency dispersion part at higher

frequency region. According to the jump relaxation model, at lower frequencies the ions

hopping from one site to a neighbouring vacant site successfully contribute to electrical

conductivity.176, 177 The observed charge carrier relaxation at higher frequencies is due to

the correlated forward-backward hopping together with relaxation of mobile ions,

suggesting that the conduction of Li+ ions in garnets seem to take place through a hopping

mechanism. The relaxation could be taking place through a displacement of Li-ions from

the 48g site to the 96h site and adjustment of surrounding ions after each successful jump

of Li+ ions from one octahedral site to new positions (Figure 6.2).

119

Figure 6.5 Conductivity (σ) vs. frequency spectra for Li5+2xLa3Ta2-xYxO12 (x = 0.25

and 0.5) at -50 °C. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced by

permission of the PCCP Owner Societies.130

6.2.3 Dielectric permittivity and loss tangent

Figures 6.6 and 6.7 show the real part of dielectric permittivity (ε/) of the

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 as a function of frequency at various

temperatures. The upturn in the data points at lower frequencies corresponds to the

electrode-electrolyte interface polarization (i.e., Maxwell–Wagner–Sillar polarization).

With an increase in frequency, due to the high periodical reversal of field, the oscillating

charge carriers may not contribute significantly to the dielectric constant. Therefore, the

dielectric constant decreases with increasing frequency and is a typical behaviour for ion

conducting materials.178 At elevated temperatures, the dominant dipolar polarization leads

to the appearance of a plateau in log ε/ plots in the lower frequency region (Figures 6.6 and

6.7).179, 180 The details of dipole formation are given in the next section. With increase in

temperatures, due to the enhanced mobility of thermally activated charge carriers, both

-1 0 1 2 3 4 5 6

-8.0

-7.5

-7.0

-6.5

-6.0

-5.5

Li6-phase

Li6.5

-phase

log

(S

cm

-1)

log f(Hz)

120

electrode polarization and dipolar polarization increase and that leads to a higher value of

dielectric permittivity. The shift in plateau towards higher frequencies with increasing

temperature indicates that dipolar polarization is a thermally activated process. A similar

variation of dielectric permittivity (ε/) with frequency was also observed in

Li6.5La3Ta1.25Y0.75O12.


Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase) at various temperatures. Phys. Chem. Chem. Phys.


121


Li6La3Ta1.5Y0.5O12 (Li6-phase) at various temperatures. Phys. Chem. Chem. Phys.


The variation of imaginary part of permittivity ε// (i.e., dielectric loss factor, which

includes dipolar loss and conduction loss) with frequency is shown in Figures 6.8 and 6.9.

At lower temperatures, the data points fall in a straight line. The conduction loss seems to

predominate in the entire frequency range; therefore, at all the temperatures ε// varies as

1/ω, where ω is the angular frequency. At elevated temperatures, the peak/plateau

corresponding to the dipolar polarization appears at lower frequencies, as it is observed in

the real part of permittivity (Figures 6.6-6.7).179, 180 The dielectric strength ∆ε = εs - ε∞,

(where εs and ε∞ are the low and high-frequency limit of real part of permittivity (ε/) under

the investigated condition) is the measure of electric dipoles formed by Li+ ions. The

dielectric strength (∆ε) of 9.4 x 106, 13.5 x 106 and 3.6 x 106 was observed for

Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12 and Li6.5La3Ta1.25Y0.75O12 at 23 °C, respectively.

122


Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase) at various temperatures. Phys. Chem. Chem. Phys.



Li6La3Ta1.5Y0.5O12 (Li6-phase) at various temperatures. Phys. Chem. Chem. Phys.


123

Figures 6.10-6.12 show the frequency vs dielectric loss tangent plots for

Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12 and Li6.5La3Ta1.25Y0.75O12 at various

temperatures. The spectra exhibit a peak corresponding to the dielectric loss due to dipolar

polarizations (dipolar relaxation), whose peaks shift towards higher frequencies with

increasing temperature. The probable peak frequency in tan δ (ftan δ), peak corresponding

to the dipolar polarization (fp) in εs and ε∞ can be related as equation 6.1:179, 180

1/2

tan p sf f (6.1)

where fp can be obtained from log ε/ vs log f plot. The calculated peak frequency (ftanδ)

corresponding to dipolar polarization match well with the experimentally observed peak in

loss tangent plots (Figures 6.10-6.12) at all the temperatures. The shift in ftanδ with

temperature can be attributed to a thermally activated process. The peak frequency is found

to exhibit the Arrhenius-type behaviour, which is represented in equation 6.2, i.e.,

tan 0 exp a Bf f E k T (6.2)

where fo is the pre-exponential factor and Ea is the activation energy associated with the

dipolar relaxation/rotation which leads to the dielectric loss.

124


temperatures in Li5.5La3Ta1.75Y0.25O12 (Li5.5-phase). Phys. Chem. Chem. Phys. 2014,

16, 11356-11365. Reproduced by permission of the PCCP Owner Societies.130


temperatures in Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem. Chem. Phys. 2014, 16,

11356-11365. Reproduced by permission of the PCCP Owner Societies.130

125


temperatures in Li6.5La3Ta1.25Y0.75O12 (Li6.5-phase), respectively. Phys. Chem. Chem.


Societies.130

Figure 6.13 shows the Arrhenius plots of dielectric loss peak frequency. The activation

energies for dipolar rotation (Em) were found to be 0.69, 0.54 and 0.34 eV (-50 to 50 °C)

(with an error of ±0.005 eV) respectively, for Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12

and Li6.5La3Ta1.25Y0.75O12.

126

Figure 6.13 Arrhenius plots of dielectric loss peak frequency in Li5+2xLa3Ta2-xYxO12

(x = 0.25, 0.5, and 0.75). Phys. Chem. Chem. Phys. 2014, 16, 11356-11365. Reproduced

by permission of the PCCP Owner Societies.130

3.4. Dielectric Modulus Analysis

To further understand the ionic conduction process by suppressing the effect of interfacial

polarization, the impedance data can also be interpreted using the modulus formalism. The

modulus function is the inverse of dielectric permittivity ε (i.e., M = ε−1)122 (equation 6.3):

/ / / / /

0 0( )M j C Z j C Z jZ (6.3)

where M' (-ωCo Z") and M" (ωCoZ') are the real and imaginary parts of modulus function

(M) and Z' and Z" are the real and imaginary parts of impedance (Z), respectively. Co is the

capacitance of the empty measuring cell of electrode area “A” and electrode separation

length L. In the present study, the values of Z' and Z'' were obtained directly from the

impedance measurement.

127

The electric modulus M״ as a function of frequency for Li5.5La3Ta1.75Y0.25O12,

Li6La3Ta1.5Y0.5O12 and Li6.5La3Ta1.25Y0.75O12 are shown in Figures 6.14-6.16. At low

frequencies, the values of M״ approach zero, which confirms that the electrode polarization

does not make any significant contribution to the modulus. At lower temperatures, the

modulus spectra of Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 consist of a relaxation

peak. As in the case of cubic fluorite structured oxygen ionic conductors,178, 181-183 the peak

in the modulus spectra of these Li-excess garnets can be attributed to the charge re-

orientation relaxation of Li+ ion from one octahedral site to another octahedral site around

an immobile Li+ ion present in the tetrahedral site. With increasing temperature, the peak

frequency shifts towards higher values indicating that the relaxation re-orientation of Li+

is a thermally activated process.




-1 0 1 2 3 4 5 6

0.000

0.002

0.004

0.006

0.008

0.010

0.012

0.014

23 oC

-17 oC

-20 oC

-35 oC

-40 oC

M//

log f(Hz)

Li5.5

-phase

128


temperatures in Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem. Chem. Phys. 2014, 16,





-1 0 1 2 3 4 5 6

0.000

0.002

0.004

0.006

0.008

0.010

0.012

23 oC

-17 oC

-20 oC

-45 oC

-50 oC

M//

log f(Hz)

Li6-phase

-1 0 1 2 3 4 5 6

0.000

0.002

0.004

0.006

0.008

0.010

0.012

7 oC

-15 oC

-20 oC

-25 oC

-29 oC

M//

log f(Hz)

Li6.5

-phase

129

Figures 6.17 and 6.18 show the modulus master curves (M״/M״max vs. log f/fM

,״

where M״max is the value of M״ at peak position and fM

is the peak frequency in modulus ״

spectra) at two different temperatures for Li5.5La3Ta1.75Y0.25O12, and Li6La3Ta1.5Y0.5O12.

Overlapping of master curves at different temperatures indicates that the relaxation re-

orientation process is unique and temperature independent. From the Gaussian fitting of

modulus peaks, the values of full width half maximum (FWHM) are found to be about 1.5

and 1.51, respectively for Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12. FWHM being greater

than 1.14 (for Debye type process FWHM= 1.14), relaxation process was found to be non-

Debye type in nature as observed in other class of Li+ ion conductors.184

Figure 6.17 Modulus master curve at various temperatures in Li5.5La3Ta1.75Y0.25O12

(Li5.5-phase). Here M״max is the value of M״ at peak position in modulus spectra (a-b)

and fM .is the peak frequency. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365 ״


130

Figure 6.18 Modulus master curve at various temperatures in Li6La3Ta1.5Y0.5O12 (Li6-

phase). Here M״max is the value of M״ at peak position in modulus spectra (a-b) and

fM .is the peak frequency. Phys. Chem. Chem. Phys. 2014, 16, 11356-11365 ״


The relaxation time (τ) for the local motion of Li+ was obtained by using the relation:

τ = 1/(2πfMwhere fM (״

.is the peak frequency in modulus spectra (Figures 6.14 and 6.15) ״

The relaxation time (τ) is found to exhibit the Arrhenius behaviour with variation in

temperature i.e.; (equation 6.4),

0 exp m BE k T (6.4)

where, τo is the relaxation time at infinite temperature and Em is the migration energy

associated with the location motion of Li+. Figure 6.19 shows the Arrhenius plots of

relaxation times in Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12. The energy associated

with the reorientation process, i.e., the migration energy (Em), was found to be 0.68 and

0.51 eV for Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12, respectively. Relaxation time at

infinite temperature (τo) for this reorientation process was found to be in the order of 10-19

131

and 10-17 s, respectively for Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12. As the temperature

increases the peak corresponding to the local motion of Li+ ion vanishes and M" moves

towards M∞, indicating a long-range order migration of Li+ ions, as observed in oxygen ion

conducting materials.177, 178, 181-183 Absence of the peak corresponding to local migration of

Li+ in Li6.5La3Ta1.25Y0.75O12 indicates that Li+ likely performs only long-range order

migration under the investigated temperatures. Figure 6.2 presents one of the potential

paths of Li+ ions in the long-range migration process.

Figure 6.19 Arrhenius plots of relaxation time (τ) in the samples Li5.5La3Ta1.75Y0.25O12

(Li5.5-phase), Li6La3Ta1.5Y0.5O12 (Li6-phase). Phys. Chem. Chem. Phys. 2014, 16,


In general, the dipolar polarization loss occurs due to the presence of permanent

dipoles. When an octahedral site and its adjacent tetrahedral position are both occupied,

3.9 4.0 4.1 4.2 4.3 4.4 4.5

-6.4

-6.0

-5.6

-5.2

-4.8

-4.4

-4.0

Li6-phase

Li5.5

-phase

log

(s

)

1000/T (K-1)

Em= 0.51 eV

o ~10

-17 s

Em = 0.68 eV

o ~ 10

-19 s

Temperature (oC)

-15 -20 -25 -30 -35 -40 -45 -50

132

the Li+ ion on the octahedral site is displaced from the central 48g site to one of the 96h

positions, which is away from the tetrahedral site, due to the repulsive forces between them.

In that case, the displacement of Li+ could cause a local polarization of charge within the

octahedra. There would be a net effective electric dipole moment pointing from negatively

charged species to the positively charged Li+ present in the tetrahedral site. The thermally

activated motion of Li+ from one octahedra to another leads to the orientation of these

electric dipoles. In presence of an externally applied electric field, these arbitrarily oriented

dipoles are aligned in the direction of the electric field leading to a dipolar polarization

loss. As the migration energy for Li+ ions, obtained from modulus spectra in the

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12, is similar to the energy associated with the

dipolar rotations in the respective samples, it indicates that the charge carriers involved in

both dipolar rotation and local orientation process are the same. Even though there is no

local motion (or reorientation) of octahedral Li+ around the tetrahedral Li+ in the Li6.5-

phase, the dielectric loss is still observed (Figure 6.12) due to the dipolar displacement in

the direction of the applied electric field during the long-range migration of Li+ ions.

Similarly, dielectric loss is also observed in the case of Li5.5La3Ta1.75Y0.25O12 and

Li6La3Ta1.5Y0.5O12 at higher temperature ranges (Figure 6.10 and 6.11), while Li+ ions

seem to perform only a long-range order migration (Figure 6.14 and 6.15).

As in the case of oxygen ion conducting materials,177, 181-183 the migration of Li+ in the

local re-orientation process could be considered to be similar to the migration of the free

charge carrier in the long-range motion. As a result, the migration energies of Li+ ions in

long-range order motion are considered to be 0.68, 0.51 and 0.34 eV, respectively, for

Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12 and Li6.5La3Ta1.25Y0.75O12. This shows that the

133

migration energy of Li+ decreases with increase in Li+ ion and Y content in the

Li5+2xLa3Ta2-xYxO12. XRD studies showed that the lattice parameter increases with the Y

content, due to larger ionic radius of the Y3+ compared to that of Ta5+.125 The larger lattice

parameter leads to wider ionic spacing, which may provide the facile path for ions, and the

activation energy for migration is reduced.

A previous NMR study on Li7La3Zr2O12 provides information on Li+ occupancy

in octahedral sites and the diffusion of Li+ ions.169 The peak observed in the 6Li NMR

spectrum has been described as the homo nuclear dipole-dipole interaction. This peak

probably corresponds to the dipoles formed by the Li+ in the octahedral site and the Li+ in

the tetrahedral site, which leads to the dipolar polarization loss. However, the dipolar

responses in NMR spectra and dielectric loss tangent spectra appear at different frequency

ranges.

Comparing the loss tangent peaks for all samples, it is observed that, the peak

intensities for Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 are higher than that for the

Li6.5La3Ta1.25Y0.75O12, at all the temperatures. The higher the peak intensity, higher is the

dielectric loss in the material.180 The increase in the loss peak intensity can be attributed to

a greater number of dipoles present in the materials, when the dipole moments in all

materials are similar. As the Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 have a lower

concentration of Li ions compared to the Li6.5La3Ta1.25Y0.75O12, there would be fewer

number of dipoles in Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12. Hence, the only

possibility is that the dipole moments in these two phases are larger than that of the

Li6.5La3Ta1.25Y0.75O12, which leads to higher dielectric losses in Li5.5 and Li6-phases.

Considering the dielectric strength (∆ε) as a measure of electric dipole, the higher value of

134

∆ε in Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 compared to Li6.5La3Ta1.25Y0.75O12

support that the dipole moment in former samples are larger than that of later compound.

Using ab initio calculations, Ming Xu et al.82 have proposed two possible pathways

of Li+ ion migration in garnets. In one of them (path A for Li5La3Nb2O12) Li+ migrates

between two octahedral sites (Li(2)) via interstice bypassing their common tetrahedral

neighbour (Li(1)) as shown in Figure 6.2. In another pathway (path B for Li7La3Zr2O12)

Li+ moves through the border (the shared triangular face) that separates the polyhedral

around Li(1) and Li(2), sticking around briefly at the corner of the tetrahedron which

surrounds the Li(1) site, then climbing over the other Li(1)–Li(2) border and finally

arriving at the vacant Li(2) site (Figure 6.2).

The Li+ migration in path A is associated with the activation energy of ~ 0.6 eV

and is preferred when Li+ content is less. Activation energy of Li+ in path B is about 0.3

eV and is preferred in Li+ rich samples.82 In path A, the Li+ migration process could be

associated with the larger dipolar moment, as the average Li–Li inter ionic distance is

expected to be larger in this route, compared to path B. Larger dipole moment can lead to

the higher dielectric loss as observed in Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12

compared to that in Li6.5La3Ta1.25Y0.75O12. This result matches well with our observation

of loss tangents (Figures 6.10 and 6.11) in these materials. Therefore, it can be suggested

that, whether it is long-range or short-range local motion, Li+ follows path A in both

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12, whereas in case of the Li6.5La3Ta1.25Y0.75O12,

the Li+ migration takes place through path B. The values of activation energy obtained

from dipolar loss tangent peaks (i.e., 0.69, 0.54 and 0.34 eV respectively, in

Li5.5La3Ta1.75Y0.25O12, Li6La3Ta1.5Y0.5O12 and Li6.5La3Ta1.25Y0.75O12) also match well with

135

the theoretical predicted values of activation energy for different Li+ ion migration paths

in the investigated garnet-type materials.

6.3 Summary130

The Li ion transport mechanisms in Li-stuffed garnet-type Li5+2xLa3Ta2-xYxO12 (x = 0.25,

0.5 and 0.75) have been investigated by correlating the structural aspects of the materials

to electrical and dielectric properties. These materials have a cubic garnet-type structure,

with a space group Ia-3d where Li+ ions are distributed in octahedral and tetrahedral

positions. The Li+ ion conduction in these materials seems to take place through a hopping

mechanism, primarily involving the Li+ ions on the octahedral sites. In the

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12 phases, the Li+ ions perform local motion

around the immobile Li+ ions place at tetrahedral positions at lower temperatures and long-

range order migrations at higher temperatures. However, the Li6.5La3Ta1.25Y0.75O12 appears

to show only long-range migration of Li+ ions. The Li+ ions appear to be associated with

larger dipole moment (as supported by the higher dielectric strength) in

Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12, causing a higher dielectric loss in these

samples compared to Li6.5La3Ta1.25Y0.75O12. Our experimental study supports the ab initio

theoretical prediction184 about the energetically favourable hopping pathways of lithium

ions, from one octahedral position to another, which vary with Li+ ion concentration in the

material.

136

Chapter Seven: Dielectric characteristics of fast Li ion conducting garnet-type

Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75)

The results discussed in this chapter is a published work, “Dielectric characteristics of fast

Li ion conducting garnet-type Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75)” by Sumaletha

Narayanan, Ashok Kumar Baral, and Venkataraman Thangadurai, in Phys. Chem. Chem.

Phys., 2016, 18, 15418-15426.


The dielectric characteristics of Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75) lithium-

stuffed garnets have been investigated and compared with that of Li5+2xLa3Ta2-xYxO12 (x =

0.25, 0.5 and 0.75) (which is presented in Chapter 6) in the temperature range from about

-53 to 50 °C using AC impedance spectroscopy. All the investigated Li-stuffed garnet

compounds were prepared, under the same condition, using conventional solid-state

reaction at elevated temperature in air. The Nyquist plots show manily bulk contribution

to the total Li+ ion conductivity for Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) and

Li6La3Nb1.5Y0.5O12 (Li6-Nb), while both bulk and grain-boundary effects are visible in case

of Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) phase at ~ -22 °C.

Non-Debye relaxation process was observed in the modulus AC impedance plots.

The dielectric loss tangent of Li5+2xLa3Nb2-xYxO12 are compared with that of corresponding

Ta analogue, Li5+2xLa3Ta2-xYxO12 and showed a decrease in peak intensity for the Nb-based

garnet samples which may be attributed to a slight increase in their Li+ ion conductivity.

The relative dielectric constant values also found to be higher for the Ta member (> 60 for

Li5+2xLa3Ta2-xYxO12) than that of the corresponding Nb analogue (~ 50 for Li5+2xLa3Nb2-

xYxO12) at below room temperature. A long-range order Li+ ion migration pathway with

137

relaxation time (τ0) 10-18 - 10-15 s and an activation energy of 0.59 - 0.40 eV was observed

for the investigated Li5+2xLa3Nb2-xYxO12 garnets and is comparable to that of the

corresponding Ta-based Li5+2xLa3Ta2-xYxO12 garnets.


7.2.1 AC Impedance spectroscopy analysis

Typical AC impedance spectra of cubic structured garnet-type Li5+2xLa3Nb2-xYxO1257 (x =

0.25, 0.5 and 0.75 which are represented as Li5.5-Nb, Li6-Nb and Li6.5-Nb, respectively) at

-22 °C are shown in Figure 7.1. The Li5.5-Nb and Li6-Nb phases of Li5+2xLa3Nb2-xYxO12

show a single semicircle, which corresponds to bulk (b) contribution at higher frequency

range and a spike at the low-frequency range due to polarization resistance. The line

spacing through the data points is fitted using an equivalent circuit consisting of

(RbCPEb)(CPEe) for Li5.5-Nb and Li6-Nb (where Rb is the bulk resistance, CPEb is the

constant phase element due to bulk and CPEe is the constant phase element due to

electrode). The Rb is low-frequency intercerpt to real axis and bulk capacitance was found

to be in the order of ~ 10-11 F. The Li6.5-Nb phase shows an additional distorted semi-circle

corresponding to grain-boundary (gb) contribution (capacitance ~ 10-8 F) at low-frequency

and it can be described using an equivalent circuit (RbCPEb) (Rgb CPEgb) (CPEe) (where

Rgb is the gb resistance, CPEgb is the constant phase element due to gb). The electrode

capacitane was found to be in the order of 10-7 F. As reported earlier,57 the density of

Li5+2xLa3Nb2-xYxO12 garnet was found to be increasing with increase in Y and Li content.

Also, there are some impurity peaks observed in PXRD of highly doped

Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) phase. These factors might be contributing to the presence

of grain-boundary impedance visible in Li6.5-Nb compared to Li5.5La3Nb1.75Y0.25O12 (Li5.5-

138

Nb) and Li6La3Nb1.5Y0.5O12 (Li6-Nb). The total resistance (R) (bulk +grain-boundary) was

used to estimate the total Li+ ion conductivity (σ) from the low-frequency intercept point

of Z// (imaginary part of impedance) to the X-axis (Z/, real part of the impedance) (Figure

7.1).

Figure 7.1 Typical complex impedance spectra obtained using Li+ ion blocking Au

electrodes for Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb), Li6La3Nb1.5Y0.5O12 (Li6-Nb) and

Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) at about -22 °C. The line spacing through the data

point is fitted using an equivalent circuit consisting of (RbCPEb)(CPEe) for Li5.5-Nb

and Li6-Nb and (Rb CPEb) (Rgb CPEgb) (CPEe) for Li6.5-Nb. Phys. Chem. Chem. Phys.


The Li+ ion conductivity dependence on temperature was represented by the Arrhenius

plots, in Figure 7.2. The Li+ ion conductivity of samples increases from

Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) to Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) which is consistent

0.0 0.2 0.4 0.6 0.80.0

-0.2

-0.4

-0.6

-0.8

Rb

Rb

Li5.5

-Nb phase

Li6-Nb phase

Li6.5

-Nb phase

Z// (

M

cm

)

Z/ (M cm)

-22 °C

Rb+R

gb

139

with the increase in cell parameter (a) calculated from powder X-ray diffraction (PXRD).

The cell constant was found to be 12.8582(5) Å, 12.9136(4) Å and 12.9488(11) Å) for

(Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb), Li6La3Nb1.5Y0.5O12 (Li6-Nb), and Li6.5La3Nb1.25Y0.75O12

(Li6.5-Nb), respectively.57 The activation energy calculated for Li5.5-Nb, Li6-Nb and Li6.5-

Nb phases are 0.59, 0.58 and 0.42 eV, respectively, in the temperature range of -53 to 50

°C (Table 7.1) which is in agreement with the Li+ ion conductivity trend.

Figure 7.2 Arrhenius plots of bulk Li ion conductivity of Li5.5La3Nb1.75Y0.25O12 (Li5.5-

Nb), Li6La3Nb1.5Y0.5O12 (Li6-Nb), and total (bulk + grain-boundary) Li+ ion

conductivity of Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb). The line passing through the data

points is fitted line using Arrhenius equation. Phys. Chem. Chem. Phys. 2016, 18,


3.0 3.3 3.6 3.9 4.2 4.5

-5

-4

-3

-2

-1

0

1000/T (K-1)

log

10T

(S

cm

-1K

)

Li5.5

-Nb phase (Ea = 0.59 eV)

Li6-Nb phase (E

a = 0.58 eV)

Li6.5

-Nb phase (Ea = 0.42 eV)

60

40

20

0 -20

-40

-60

Temperature (oC)

140

Table 7.1 The bulk (except bulk + grain-boundary for Li6.5-Nb) conductivity of

Li5+2xLa3Nb2-xYxO12 at -22 and 25 °C and the activation energy calculated from

Arrhenius plots (Figure 7.2) at temperature range -50 to 50 °C. Phys. Chem. Chem.


Societies.185

Sample σ-22 °C (Scm-1) σ25 °C (Scm-1) Ea (eV)

Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) 1.50 x 10-6 7.18 x 10-5 0.59

Li6La3Nb1.5Y0.5O12 (Li6-Nb) 4.20 x 10-6 1.87 x 10-4 0.58

Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) 1.27 x 10-5 2.99 x 10-4 0.42

Shown in Figures 7.3-7.5 are the Li+ ion conductivity as a function of frequency for

Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75) in the temperature range at -53 to 50 °C,

which obeys Jonscher universal power law.175 The dispersion region seen at lower

frequency range is due to the polarization due to Li+ ion-blocking electrode. The frequency

independent DC plateau region is seen at intermediate frequency range followed by a

dispersion region. This could be explained by jump relaxation model which indicates that

at intermediate frequencies, the ions hop from one site to the neighboring vacancy site

contribute to the DC conductivity, and at higher frequencies the correlated alternative

hopping along with relaxation of ions contribute to the conductivity relaxation at the

dispersion region.176, 177 For garnet-type oxide, the jumping of Li+ ions from one octahedral

site to other octahedral- site (for e.g., 48g-site to the 96h-site) followed by the readjustment

of surrounding ions causes the conductivity relaxation.130

141

Figure 7.3 The electrical conductivity of Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) as a function

of frequency at different temperatures obtained using AC impedance spectroscopy

with Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


Figure 7.4 The electrical conductivity of Li6La3Nb1.5Y0.5O12 (Li6-Nb) as a function of

frequency at different temperatures obtained using AC impedance spectroscopy with

Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


-2 -1 0 1 2 3 4 5 6 7-8.0

-7.5

-7.0

-6.5

-6.0

-5.5

-5.0

-4.5 - 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

log(

S c

m-1)

log f (Hz)

Li5.5

-phase

-2 -1 0 1 2 3 4 5 6 7

-7

-6

-5

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

log(

S c

m-1)

log f (Hz)

Li6-phase

142

Figure 7.5 The electrical conductivity of Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) as a function

of frequency at different temperatures obtained using AC impedance spectroscopy

with Li+ ion blocking Au electrodes. Phys. Chem. Chem. Phys. Chem. Chem. Phys.


The real (ε/) and imaginary (ε//) part of permittivity (ε) was computed using the

equations 7.1 and 7.2 as shown below:186

/ //

/2 / /2

0 ( )

Z

C Z Z

(7.1)

// /

/2 / /2

0 ( )

Z

C Z Z

(7.2)

where Co is the vacuum capacitance of the cell. Figures 7.6-7.8 show the relation between

the real part of permittivity (ε/) and frequency of Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and

0.75) at different temperatures. It is noticeable that the dielectric constant decreases with

frequency showing a constant minimum value.

-2 -1 0 1 2 3 4 5 6 7

-8

-7

-6

-5

-4

- 53 °C

- 47 °C

- 40 °C

- 25 °C

- 21 °C

log(

S c

m-1)

log f (Hz)

Li6.5

-phase

143



spectroscopy data collected using Li+ ion blocking Au electrodes. Phys. Chem. Chem.


Societies.185


Li6La3Nb1.5Y0.5O12 (Li6-Nb) at different temperatures from AC impedance



Societies. 185

-1 0 1 2 3 4 5 61

2

3

4

5

6

7

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li5.5

-Nb phase

log

'

log f (Hz)

-1 0 1 2 3 4 5 61

2

3

4

5

6

7

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li6-Nb phase

log

'

log f (Hz)

144





Societies.185

The upturn at the lower frequency region may be due to the blocking electrode-electrolyte

polarization at the interface. When the frequency is increased, the periodical reversal of the

field is increased and as a result the oscillating charges do not contribute to the dielectric

constant. This causes a decrease in dielectric constant with increase in frequency. Dipolar

polarization leads to formation of plateau at the low-frequency region at higher

temperature, which is clearly visible in the case of Li6.5-Nb phase and also in the Ta series,

Li5+2xLa3Ta2-xYxO12 (x = 0.25, 0.5 and 0.75).130 Shifting of plateau towards high

frequencies with increasing temperature indicate the increase in frequency of dipolar

rotations, as Li+ ion mobility is enhanced due to the thermal activation.187

-1 0 1 2 3 4 5 61

2

3

4

5

6Li

6.5-Nb phase

- 53 °C

- 47 °C

- 40 °C

- 25 °C

- 21 °C

log

'

log f (Hz)

145

The plot of imaginary part of the dielectric permittivity (ε//) against frequency is

illustrated in Figures 7.9-7.11. ε// is the dielectric loss factor which could be the

combination of dipolar loss and conduction loss. The conduction loss is predominant at

lower temperatures and appears as a straight line and shows inverse relationship to the

angular frequency (ω) especially for Li5.5-Nb and Li6-Nb phases (Figures 7.9 and 7.10).

The peak/plateau at lower frequencies due to dipolar polarization was observed in all three

cases (Figures 7.6-7.8).


Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb) at different temperatures. Phys. Chem. Chem. Phys.


-1 0 1 2 3 4 5 6

1

2

3

4

5

6

7

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li5.5

-Nb phase

log

"

log f (Hz)

146


Li6La3Nb1.5Y0.5O12 (Li6-Nb) and at different temperatures. Phys. Chem. Chem. Phys.



Li6La3Nb1.5Y0.5O12 (Li6-Nb) and at different temperatures. Phys. Chem. Chem. Phys.


-1 0 1 2 3 4 5 6

1

2

3

4

5

6

7

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li6-Nb phase

log

"

log f (Hz)

-1 0 1 2 3 4 5 6

1

2

3

4

5

6

- 53 °C

- 47 °C

- 40 °C

- 25 °C

- 21 °C

Li6.5

-Nb phase

log

"

log f (Hz)

147

The dielectric loss tangent as a function of frequency of Li5+2xLa3Nb2-xYxO12 is

shown in Figures 7.12-7.14. The peak observed in Li5.5-Nb and Li6-Nb phases can be

attributed to the dielectric loss due to Li-Li-dipolar rotations under the applied electrical

field.130 There seem two relaxation peaks for Li6.5-Nb phases at intermediate frequencies,

suggesting that an additional polarization loss occurs when compared to the Li5.5-Nb and

Li6-Nb phases. The higher frequency peak may be attributed to the dielectric loss due to

Li-Li dipolar rotation in the bulk.188 The lower frequency peak in Figure 7.14 seems to be

due to the dielectric loss, as a result of grain-boundary polarization. This is consistent with

the appearance of grain-boundary arc in the AC impedance spectrum of Li6.5-Nb phase

(Figure 7.1). The shift in loss tangent peaks towards higher frequencies with increase in

temperature, similar to real part of permittivity (ε/) as a function of temperature, indicates

that both grain-boundary space charge polarization and Li-Li dipolar interaction in bulk

are thermal activated processes.

148




Figure 7.13 Dielectric loss tangent (δ) as a function of frequency of Li6La3Nb1.5Y0.5O12

(Li6-Nb) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


-2 -1 0 1 2 3 4 5 6 7

0

10

20

30

40

50

60

70

- 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li5.5

-Nb phase

tan

log f (Hz)

-2 -1 0 1 2 3 4 5 6 7

0

20

40

60 - 48 °C

- 40 °C

- 33 °C

- 30 °C

- 22 °C

Li6-Nb phase

tan

log f (Hz)

149




In Figures 7.15-7.17, the loss tangent of both Nb with Ta members of Li5+2xLa3M2-

xYxO12 (M = Nb, Ta) at different temperatures is compared. A similar plot is observed,

except for Li6.5-Nb phase of Nb series Li5+2xLa3Nb2-xYxO12, correspond to bulk and grain-

boundary contribution.57 However, very small values of peak intensity of Li5+2xLa3Nb2-

xYxO12 (Figures 7.16 and 7.17) compared to that of Ta series Li5+2xLa3Ta2-xYxO12 indicated

that either there are lesser number of dipoles in Nb garnets than Ta garnets or the mobility

of ions is higher in the presently investigated Nb garnets leading to a quick response of Li-

Li dipoles to external electric field and resulting in a less dielectric loss. Overall, the

dielectric relaxation can be considered due to the main-body interactions as is common for

the ionic conductor.175

-2 -1 0 1 2 3 4 5 6 7

0

1

2

3

4

5

6

7

Li6.5

-Nb phase

- 53 °C

- 47 °C

- 40 °C

- 25 °C

- 21 °C

tan

log f (Hz)

150


Li5.5-phase of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and Ta) at

different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426. Reproduced



Li6-phase of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and Ta) at



-2 -1 0 1 2 3 4 5 6 7

0

10

20

30

40

50

60

70

- 40 °C Nb

- 40 °C Ta

- 20 °C Nb

- 20 °C Ta

Li5.5

-phase

tan

log f (Hz)

-2 -1 0 1 2 3 4 5 6 7

0

20

40

60

80

100

- 48 °C Nb

- 50 °C Ta

- 22 °C Nb

- 20 °C Ta

Li6-phase

tan

log f (Hz)

151


Li6.5-phases of both Ta57 and Nb members of Li5+2xLa3M2-xYxO12 (M = Nb and Ta) at



The relative permittivity/dielectric constant (εr) of Li5+2xLa3M2-xYxO12 (M = Nb,

Ta57) were calculated using the capacitance (C), as given in equation 3.8 and the results are

summarized in Figure 7.18. Dielectric constant slightly increases with increase in Li

content in Ta series materials, although there is no trend in variation of εr with Li content

in Nb series. With increase in temperature, value of dielectric constant is increased in both

Nb and Ta based garnets. The dielectric constant of Nb garnets is found to be relatively

higher than that of Ta garnets. Generally, the higher the dielectric constant, the greater is

the polarizability of the medium. Thus, higher dielectric constant (Figure 7.18) and lower

dielectric loss (Figures 7.15-7.17) in Nb series materials indicate that its Li+ ion mobility

seems to be higher compared to that of Ta garnets since both Nb and Ta oxides exhibit

similar ionic radius and electronic structure. A comparison of Li+ ionic conductivity of both

-2 -1 0 1 2 3 4 5 6 7-5

0

5

10

15

20

25

30

35

40

Li6.5

-phase - 53 °C Nb

- 50 °C Ta

- 20 °C Nb

- 21 °C Ta

tan

log f (Hz)

152

Nb and Ta phases of Li5+2xLa3M2-xYxO12 at specific temperatures is shown in Table 7.2.57,

125, 130

Figure 7.18 The relative permittivity of Li5.5-, Li6-, and Li6.5-phases Li5+2xLa3M2-

xYxO12 (M = Nb and Ta) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18,


Using modulus analysis, it is possible to understand the bulk relaxation properties

in in detail by suppressing the lower frequency phenomena, especially the electrode

effect.172 The electric modulus accounts for the electric field relaxation in the material at a

constant electric displacement.189 The modulus spectra of Li5.5-Nb, Li6-Nb and Li6.5-Nb

phases of Li5+2xLa3Nb2-xYxO12 samples, measured at a temperature range of -53 to -21 ºC

as a function of frequency is shown in Figures 7.19-7.21. The electric modulus (M//) shows

a frequency independent behavior at the lower frequency region for all the samples. This

indicates that the electrode polarization does not contribute to the electric modulus and the

-50 -40 -30 -200

20

40

60

80

100

Li5.5

-Nb Li5.5

-Ta

Li6-Nb Li

6-Ta

Li6.5

-Nb Li6.5

-Ta

Re

lative

pe

rmittivity

Temperature (oC)

153

long-range migration of Li+ ion by hopping from one site to the neighboring site.172

Presence of peak in the modulus spectra means the relaxation associated with Li+ ion

mobility.130, 172 The appearance of high frequency side peak at lower temperatures could

be due to the Li+ ion re-orientation relaxation as Li+ ion moves from one octahedron to

another octahedron around the immobile tetrahedral site.130 The shift in peak position

towards higher frequency side with increase in temperature indicates that Li+ ion relaxation

re-orientation is thermally activated phenomenon. Also, it is noticeable that the peak tend

to disappear in Li6.5-Nb phase which can be explained in terms of long-range order

migration of Li+ ion instead of local migration as seen in other members.130


(Li5.5-Nb) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


154

Figure 7.20 Electric modulus M// as a function of frequency of Li6La3Nb1.5Y0.5O12 (Li6-

Nb) and (c) Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb) at different temperatures. Phys. Chem.

Chem. Phys. 2016, 18, 15418-15426. Reproduced by permission of the PCCP Owner

Societies.185


(Li6.5-Nb) at different temperatures. Phys. Chem. Chem. Phys. 2016, 18, 15418-15426.


155

Table 7.2 Comparison of ionic conductivity of both Nb and Ta phases of Li5+2xLa3M2-

xYxO12 at specific temperatures.57, 125, 130

Garnet Li5+2xLa3Nb2-xYxO12 Li5+2xLa3Ta2-xYxO12

T (ºC) σ (Scm-1) T (ºC) σ (Scm-1)

Li5.5- phase -33 4.52 x 10-7 -35 5.94 x 10-8

-40 2.10 x 10-7 -40 3.32 x 10-8

Li6 - phase -33 1.21 x 10-6 -35 5.84 x 10-8

-40 5.48 x 10-7 -40 3.34 x 10-8

Li6.5 - phase -21 1.60 x 10-5 -20 9.85 x 10-6

-25 1.18 x 10-5 -25 7.29 x 10-6

The relaxation time, τ is related to the frequency maximum in the modulus plots,

fM//

which can actually reveal the short range and or long range migration of Li+ in the

crystal structure. Figure 7.22 shows the Arrhenius behavior of relaxation time, τ for the

local motion of Li+ ions. The migration energy calculated for Li5.5-Nb, and Li6-Nb and

Li6.5-Nb phases are 0.59, and 0.52, and 0.40 eV, respectively which is close to the

theoretically calculated activation energy. The calculated relaxation time, τ0 was in the

range of 10-15 – 10-18 s which decreases with Li content and is comparable with that of Ta

family, Li5+2xLa3Ta2-xYxO12.130 The AC impedance or NMR analysis of some of the ionic

conductors is mentioned here to get an idea of how the relaxation time is varied in different

electrolytes. In the polymer nano-composites of polyethylene oxide and lithium

perchlorate, Li+ ions show relaxation times in micro seconds (10-6 s).190 The fastest Na+

conducting, Na β-alumina has exhibited a relaxation time, τ0 of 10-12 s.191 Solid-state 7Li

NMR analysis of Li1.3Al0.15Y0.15Ti1.7(PO4)3 showed a relaxation time in the order of 10-12

s.192 For (BiI3)0.4-(Ag3PO4)0.6 electrolyte, the mobile Ag+ ions show the relaxation time of

10-7 - 10-11 s.193

156

Figure 7.22 Arrhenius plots of relaxation time of Li5.5La3Nb1.75Y0.25O12 (Li5.5-Nb),

Li6La3Nb1.5Y0.5O12 (Li6-Nb), and Li6.5La3Nb1.25Y0.75O12 (Li6.5-Nb). Phys. Chem. Chem.


Societies.185

7.3 Summary185

The transport mechanism of Li5+2xLa3Nb2-xYxO12 (x = 0.25, 0.5 and 0.75) was studied in

terms of their dielectric behavior at below room temperature. The bulk Li+ ion conductivity

of samples increases with increase in Li content and follows the trend,

Li5.5La3Nb1.75Y0.25O12 (Li5.5 -Nb) < Li6La3Nb1.5Y0.5O12 (Li6-Nb) < Li6.5La3Nb1.25Y0.75O12

(Li6.5-Nb) and also follows with the cell constant trend. The Li+ ion conductivity as a

function of frequency was found to obey the Jonscher universal power law. The modulus

plots indicate the non-Debye behavior of Li+ ion relaxation. Relatively higher dielectric

constants for the Nb members compared to that of the Ta members of Li5+2xLa3Nb2-xYxO12

3.9 4.0 4.1 4.2 4.3 4.4 4.5 4.6

-7

-6

-5

-4

E a = 0.59 eV;

= 10-18 s

E a = 0.52 eV;

= 10-17 s

Li5.5

-Nb phase

Li6-Nb phase

Li6.5

-Nb phase

log

1000/T (K-1)

Ea = 0.40 eV;

= 10

-15 s

-20

-25

-30

-35

-40

-45

-50

-55

Temperature (°C)

157

(M= Nb, Ta) could be explained in terms of the higher dipole moments and greater Li+ ion

mobility. The calculated relaxation time, τ0 was 10-18, 10-17 and 10-15 s and activation energy

observed was 0.59, 0.52 and 0.40 eV for Li5.5-Nb, Li6-Nb and Li6.5-Nb, respectively, of

Li5+2xLa3Nb2-xYxO12 from the relaxation profile.

158

Chapter Eight: Synthesis and electrical properties of hybrid gel Electrolytes Derived

from Keggin-type heteropoly acids and 3-(pyridin-1-ium-1-yl) propane-1-sulfonate

(PyPs)

The results discussed in this chapter are submitted as a manuscript entitled, “Synthesis and

electrical properties of hybrid gel Electrolytes Derived from Keggin-type heteropoly acids

and 3-(pyridin-1-ium-1-yl) propane-1-sulfonate (PyPs)” by Sumaletha Narayanan, Xia

Tong, and Venkataraman Thangadurai, for consideration in J. Mater. Chem. A 2016 (under

review).

8.1 General overview

Proton conducting metal oxide hybrid materials derived from ionic liquids (ILs) and

polyoxometalates (POMs) gain attention in the field of electrochemical devices.92, 93, 104, 194

This chapter reports the effect of concentration of protons in POMs on hybrid formation

with ILs and their ionic conductivity relationship to optimize the gel electrolytes for higher

ionic conductivity and electrochemical stability window. The hybrid gels derived from the

Keggin-type heteropoly acids containing different proton concentration, but similar

transition metals, such as H3PW11MoO40, H4PMo11VO40 and H5PMo10V2O40 and 3-

(pyridin-1-ium-1-yl) propane-1-sulfonate (PyPs) IL. Elemental C, H, and N analysis was

found to be consistent with theoretical composition within 4 % for C and N and H content

was found to be slightly higher which may due to potential uptake of water during the

sample preparation.

1H and 13C nuclear magnetic resonance and Fourier transform infrared

spectroscopy (FTIR) confirmed the presence functional groups of PyPs in hybrids. In-situ

variable temperature powder X-ray diffraction (PXRD), thermogravimetric analysis

159

(TGA), electrochemical AC impedance spectroscopy and cyclic voltammetry studies

showed excellent thermal stability (up to ~ 300 C) and electrochemical stability (3V at

room temperature) of hybrid [PyPs]3PW11MoO40. The structural characterizations confirm

the interaction between the organic cation and Keggin-type inorganic heteropoly anion in

the hybrid material. The ionic conductivity of 0.1, 0.01 and 0.0003 Scm-1 at ~ 90 C was

obtained for [PyPs]3PW11MoO40, [PyPs]4PMo11VO40 and [PyPs]5PMo10V2O40,

respectively.


8.2.1 Phase analysis

Figure 8.1 shows the photographs of heteropoly acid H4PMo11VO40 and its derivative

hybrid PyPs-H4PMoV compound at room temperature. The freshly made product was a

dark green transparent semisolid at room temperature. PyPs-H4PMoV turned to liquid by

heating it up to nearly 80 C, exhibiting thermotropic liquid crystal behaviour. The melting

point of the hybrid compounds can be related the Coulombic attraction (Ec) between the

constituent ions, and is described by equation 8.1.195

04

c

MZ ZE

r

(8.1)

Where M is the Coulomb’s law constant, Z+ and Z- are cations and anions charges, 0 is

the permittivity of free space, and r is the interionic separation. For larger ions possessing

lower ionic charge, the formation of low melting salts is favoured. This is because the

larger size enables the effective delocalization of the charge as well as enhances the charge

160

separation. Rickert et al. proposed that the ionic salts made from with a polyoxometalate

anion and an organic cation will have lower melting point when the POM has a lower

charge.195

Figure 8.1 Photographs of (a) heteropoly acid H4PMo11VO40 and (b) its hybrid form,

PyPs-H4PMoV derived from HPA and PyPs ionic liquid.

8.2.2 Composition analysis

CHN analysis results of the as-prepared PyPs (IL), and all its hybrid compounds, PyPs-

H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV are summarized in Table 8.1. The C, H, and

N results were found to be consistent with theoretical composition within 4 % error, and

the discrepancy in % H might be due to the uncertainty in quantifying the water content of

the hybrids. This confirms that the chemical structures proposed in Figure 8.2.

161

Table 8.1 CHN elemental analysis PyPs and its hybrid compounds, PyPs-H3PWMo,

PyPs-H4PMoV, and PyPs-H5PMoV.

Hybrid Elemental % Theoretical Experimental

PyPs % C 47.75 47.09

% H 5.51 5.65

% N 6.96 6.67

PyPs-H3PWMo % C 8.23 8.31

% H 1.3 1.32

% N 1.2 1.14

PyPs-H4PMoV % C 14.38 14.45

% H 2.04 1.82

% N 2.10 2.01

PyPs -H5PMoV % C 17.54 18.17

% H 2.02 2.64

% N 2.56 2.64

Figure 8.2 Synthesis of POMs-based IL hybrid gel electrolytes showing various steps

involved.

162

The functional groups of PyPs and the hybrid POMs-ILs are further proved by NMR

analyses (Figures 8.3 and 8.4). 1H NMR results show the chemical shift (ppm) values

corresponding to the structure. PyPs: 2.46 (p, 2H), 2.98 (t, 2H), 4.77 (t, 2H), 8.08 (t, 2H),

8.56 (t, 1H), 8.88 (d, 2H); PyPs-H3PWMo: 2.47 (s, 2H), 2.98 (m, 3H), 4.81(s, 1H), 8.13 (d,

2H), 8.60 (s, 1H), 8.89 (s, 2H); PyPs-H4PMoV: 2.44 (m, 2H), 2.96 (t, 2H), 4.76 (s, 1H),

8.08 (s, 2H), 8.55 (s, 1H), 8.86 (d, 2H); and PyPs-H5PMoV: 2.45 (m, 2H), 2.96 (t, 2H),

4.77 (s, 1H), 8.08 (s, 2H), 8.55 (s, 1H), 8.87 (d, 2H). The corresponding 13C NMR chemical

shift (ppm) results are also consistent with the 1H NMR results. The 13C NMR chemical

shift values are for PyPs: 26.20, 47.12, 59.98, 128.48, 144.47, 145.99; PyPs-H3PWMo:

26.33, 47.22, 60.14, 128.79, 144.42, 146.26; PyPs-H4PMoV: 26.19, 47.09, 59.96, 128.50,

144.42, 146.01; and PyPs-H5PMoV: 26.23, 47.13, 60.01, 128.45, 144.42, and 146.03.

Figures 8.3 and 8.4 summarize the 1H NMR and 13C NMR results as a comparison plot of

both PyPs and the hybrid IL-POMs.

163

Figure 8.3 1H NMR of PyPs, PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV. The

chemical shift values are represented against TMS dissolved in D2O.

164

Figure 8.4 13C NMR of PyPs, PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV.

The chemical shift values are represented against TMS dissolved in D2O.

8.2.3 FTIR analysis

The FTIR spectra are used to identify structural and bonding changes in the polyoxoanion

units present in POM-based materials. Figure 8.5 presents the FTIR spectra of hybrids,

PyPs-H3PWMo, PyPs-H4PMoV and PyPs-H5PMoV and their solid acid precursors. In the

FTIR spectra, there are four characteristic vibrational bands resulting from the -Keggin-

type heteropolyanion (for e.g., PM12O403- where M = W, Mo, V), (P-Oa), (M-Od), (M-

165

Ob-M), and (M-Oc-M) appearing in the region between 700 and 1100 cm-1 (Od-terminal

oxygen, Ob-bridged oxygen of two octahedral sharing a corner, and Oc-bridged oxygen of

two octahedral sharing an edge). The FTIR spectra of POM-ILs are found to be identical

to that of corresponding pristine solid acids, which also shows the splitting of the P-Oa

stretching, M-Od stretching, stretching of M-Ob-M inter bridges between corner-sharing

MO6 octahedral and stretching of M-Oc-M intra bridges between edge-sharing MO6

octahedral. In the IR spectra of the hybrid materials, those four well-known characteristic

bands prove the presence of the core HPA clusters. The decrease in electrostatic anion-

anion interactions leads to an increase in the stretching and bending vibrational frequencies.

Heteropoly acid, H3PW11MoO40 and its hybrid material PyPs-H3PWMo M-Ob-M and M-

Oc-M asymmetrical stretching vibration are blue-shifted and the wave numbers increase

from 887 and 820 for H3PW11MoO40 to 914 and 855cm-1 for PyPs-H3PWMo, respectively.

There are also three characteristic peaks at 1221, 1185cm-1 (S=O) and 1491 cm-1 (pyridine

ring), indicating that [-PW11MoO40]3- had been successfully immobilized on ionic liquid

(PyPs)-modified compound. Table 8.2 shows the major vibration bands of hybrid

compounds.

166

Figure 8.5 FTIR spectra of heteropoly acids H3PW11MoO40 (H3PWMo),

H4PMo11VO40 (H4PMoV) and H5PMo10V2O40 (H5PMoV) and their IL-hybrids (PyPs-

H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV.

1800 1500 1200 900 600

PyPS-H4PMoV

H4PMoV

PyPS-H3PWMo

H3PWMo

Wavenumber (cm-1

)

PyPS-H5PMoV

H5PMoV

167

Table 8.2 FTIR Bands of IL-H3PWMo, IL-H4PMoV and IL-H5PMoV and their

Precursors.

Wavenumber (cm-1) Vibrations

H3PWMo PyPs -

H3PWMo

H4PMoV PyPs-

H4PMoV

H5PMoV PyPs -

H5PMoV

1077 1098 1064 - - - ʋ(P-Oa)

980 1001 959 942 971 967 ʋ(M-Od)

887 914 871 888 859 862 ʋ(M-Ob-M)

820 855 787 825 - - ʋ(M-Oc-M)

8.2.4 Power XRD phase analysis

Figure 8.6a shows the PXRD patterns of POM-ILs and their pristine solid acids, which

reveal their crystalline and amorphous state. The Keggin-type of the polyoxoanions

exhibits typical peaks at the range of 2θ = 7-11 in the PXRD pattern.196 These peaks at

this range are also observed in XRD patterns of all three hybrids materials. From these

observations, it is clear that the hybrid materials have been successfully prepared. The XRD

pattern peaks of these POM-ILs are significantly different from that of the pristine acid

HPMo11V. The PyPs-POM hybrids have two broad diffraction peaks at 2θ =7-10° and 15-

40°, which show a smectic phase and are consistent with the liquid-state nature of the

samples.197 Figure 8.6b presents the PXRD pattern of PyPs-H3PWMo in the small angle

region. Even though the intensities of peaks at 2θ = 3-5° are much weaker than the peak at

2θ ≈ 9°, they still can be observed. The diffraction patterns contain the fundamental

reflection in the small-angle region, which may be due to the regular arrangement of

molecules in layers.99 As a result, we can speculate that at first, the molecules are organized

in spheroidal aggregates which have a core-shell structure, where each polyoxoanion is

168

surrounded by PyPs cations, due to electrostatic interactions between polyoxoanion and

PyPs cation. After that, IL-polyoxoanion aggregates are regularly arranged in layers

through the organic IL moieties which may be disordered on the surface of the

polyoxoanion core. The schematic illustration of the crystal structure of POM-IL hybrid

gel electrolyte is shown in Figure 8.7. Figures 8.8 and 8.9 present the high temperature

XRD (HTXRD) patterns of the PyPs-H3PWMo under different temperatures. Figures 8.8

shows the inter-planar (002) spacing, d (Å) of the IL-PMo11V at 2θ ≈ 9° vs temperature.

The d value is almost stable from room temperature to 70 °C. Increasing of temperature

after 70 °C, the absorbed water loss seems to be responsible for the decrease of inter-planar

d value. A similar trend was observed for PyPs-H4PMoV hybrid as shown in Figure 8.9.

Figure 8.6 PXRD patterns of (a) PyPs, heteropoly acids H3PW11MoO40 (H3PWMo),

H4PMo11VO40 (H4PMoV) and H5PMo10V2O40 (H5PMoV) and their PyPs-POM

hybrids (PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV) and (b) PyPs-

H3PWMo in the small angle region.

(b)

10 20 30 40

PyPs-H5PMoV

H5PMoV

PyPs-H4PMoV

H4PMoV

PyPs-H3PWMo

H3PWMo

PyPS

Two theta (degrees)

Inte

nsity (

a.u

.)

(a)

4 6 8 10 12 14

Two theta(degrees)

4.72

169

Figure 8.7 The schematic illustration of the crystal structure of POM-IL hybrid gel

electrolyte.

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

SO3H

N

+

IL

self-assembly POM

170

Figure 8.8 HTXRD patterns of PyPs-H3PWMo under different temperatures.

6 8 10 12 14

Decre

ase

Incre

ase

Tem

pera

ture

(C

)

PyPs-H3PWMo

30 °C

50 °C

Inte

nsity (

a.u

.)

Two theta (Degrees)

30 °C

90 °C

70 °C

80 °C

120 °C

171

Figure 8.9 HTXRD patterns PyPs-H4PMoV under different temperatures.

8.2.5 Thermal analysis

Figure 8.10a shows the TGA of the HPAs and their derivative hybrid materials (POMs-

IL). As expected, there is significant weight loss observed for the IL-hybrids compared to

the corresponding HPAs. H4PMoV and its IL derivative (PyPs-H4PMoV) show the initial

low temperature weight loss of 3.8 wt.% for the H4PMoV and 3.4 wt.% for the PyPs-

H4PMoV due to 4 and 5 moles of water, respectively. For the PyPs-H4PMoV salt, the

172

weight loss beginning at 250 °C is primarily due to the decomposition of PyPs organic

moieties, indicating that the material is stable below 250 °C. The huge weight loss of PyPs-

H4PMoV, nearly 27%, occur between 250 °C and 650 °C corresponded to the proportion

of PyPs contained in the PyPs-H4PMoV (30 wt.%). The TGA of PyPs provided as inset

also shows the decomposition after 250 °C. The TGA curves showed the presence of water

in the IL-POMs hybrid samples; hence, the potential uptake of water might have occurred

during the sample preparation, which is consistent with the result of C, H, N elemental

analysis. Figure 8.10b shows the proposed decomposition steps of investigated IL-PMo11V

sample, based on TGA analysis. Similar observation was found for H3PWMo and its

hybrid with lower weight loss, as compared to that of H4PMoV.

173

Figure 8.10 (a) The TGA of heteropoly acids H3PW11MoO40 (H3PWMo) and

H4PMo11VO40 (H4PMoV), their PyPs-POM hybrids (PyPs-H3PWMo, and PyPs-

H4PMoV) and PyPs which is shown as inset; (b) The proposed decomposition steps of

IL-PMo11V sample.

(b)

" 1 2 𝑃2𝑂5 +12 𝑉2𝑂5 + 11 𝑀𝑜𝑂3"

- 5.0𝐻2𝑂 (3.4 wt.%) [𝑃𝑦𝑃𝑠]4[𝑃𝑀𝑜11𝑉𝑂40] ∙ 5.0𝐻2𝑂

[𝑃𝑦𝑃𝑠]4[𝑃𝑀𝑜11𝑉𝑂40]

- 𝑃𝑦𝑃𝑠 (27.0 wt.%)

(a)

174

8.2.6 Electrical properties

Figures 8.11-18 show typical impedance spectra of PyPs-POM hybrids, PyPs-H3PWMo,

PyPs-H4PMoV and PyPs-H5PMoV at different temperatures in air. The resistance was

found to decrease with increase in temperature in all the cases. Among all the hybrids,

PyPs-H5PMoV showed higher resistance. The solid line passing through the symbols

(measured data) represent the fitting using an equivalent circuit consisting of a resistance,

parallel resistance-CPE and CPE components which is shown in the inset. The fitted data

is summarized in Table 8.3.


measured in air atmosphere. Impedance plots zoomed at high frequency side is shown

as inset for clarity.

175


measured in air atmosphere.


measured in air atmosphere. Impedance plots zoomed at high frequency side is shown

as inset for clarity.

0 1000 2000 3000 4000

0

-1000

-2000

-3000

-4000

PyPs-H

3PWMo

65 C

Z''

( c

m)

Z' ( cm)

176





177





0 30 60 90 120 150

0

-30

-60

-90

-120

-150

Z''

( c

m)

Z' ( cm)

91 C

PyPs-H4PMoV

178



Table 8.3 The impedance fitting results of PyPs-H4PMoV and PyPs-H5PMoV hybrids

at different temperatures.

Sample T (ºC) R1 (MΩ) Rb (MΩ) CPEb (F) Cb (F) CPEel (F) χ2 (10-4)

PyPs-H4PMoV 32 12.6 93.9 6.68 x 10-11 5.12 x 10-12 6.69 x 10-5 3

67 4.45 0.96 4.63 x 10-9 1.70 x 10-10 4.32 x 10-4 2

PyPs-H5PMoV 60 10.1 518 7.19 x 10-11 3.36 x 10-12 1.56 x 10-5 8

92 9.02 20.9 3.19 x 10-10 1.12 x 10-11 1.26 x 10-4 2

The conductivity, estimated from the low-frequency intercept to real axis, of the materials

were calculated to be ~ 0.1, 0.01 and 0.0003 Scm-1 for PyPs-H3PWMo, PyPs-H4PMoV,

and PyPs-H5PMoV, respectively at ~ 90 °C. The conductivity trend is in agreement with

179

the impedance results. The conductivity of hybrid materials increases with increase in

temperature. The conductivity sequence of the hybrid materials is in the ascending order

i.e., [PyPs]5PMo10V2O40 < [PyPs]4PMo11VO40 < [PyPs]3PW11MoO40. The Arrhenius plots

for the bulk conductivity (σ) of PyPs-POM hybrids are presented in Figure 8.19.

Figure 8.19 Arrhenius plots for ionic conductivity of PyPs-POM gel electrolytes,

PyPs-H3PWMo, PyPs-H4PMoV, and PyPs-H5PMoV.

It has been suggested that the larger radius of W6+ increases the molecular volume of

[PyPs]3PW11MoO40 compared to corresponding V5+ based hybrids such as

[PyPs]4PMo11VO40 and [PyPs]5PMo10V2O40.103 As the lattice potential energy (UPOT) and

molecular volume (Vm) are inversely related, i.e., (equation 8.2)198,

2.6 2.8 3.0 3.2 3.4-7

-6

-5

-4

-3

-2

-1

Ea = 1.16 eV

Ea = 0.73 eV

Ea = 0.45 eV

PyPs-H3PWMo

PyPs-H4PMoV

PyPs-H5PMoV

log

10 (

Scm

-1)

1000/T (K-1)

100

80

60

40

Temperature (°C)

180

1

1/3

117.3/ 2 51.9POT

m

U kJmolV

(8.2)

the coulombic interaction between counter ions in the former case is lower than that in

latter case, which results in higher conductivity for W containing PyPs-H3PWMo than that

of V containing PyPs-H4PMoV, and PyPs-H5PMoV. However, in the case of N-methyl

imidazolium-1-(3-sulfonic group) propyl (MIMPS) IL based hybrids, an increased

conductivity was observed with H4PMoV compared to than that of H3PWMo.99 It could be

due to the variation in structure of ILs. The activation energy was found to 0.45, 0.73 and

1.16 eV for [PyPs]3PW11MoO40, [PyPs]4PMo11VO40 and [PyPs]5PMo10V2O40, respectively

in the investigated temperature range. It is consistent with the observed trend for

conductivity of the hybrid materials. Based on the higher activation energy (> 0.20 eV), it

can be assumed that the proton conduction in these types of hybrids appears to be following

vehicle mechanism.199 Comparison of proton conductivity of materials studied in this

chapter along with few IL based hybrid materials and Nafion 115 is provided in Figure

8.20.99, 200, 201 It shows that the PyPs-H3PWMo hybrid exhibits promising conductivity

which is similar (slightly higher) to that of bench mark polymer membrane, Nafion.

181

Figure 8.20 Comparison of electrical conductivity of PyPs-H3PWMo with other

known promising proton conductors.99, 200, 201

The cyclic voltammetry analysis was used to understand the electrochemical stability

window (ESW) of PyPs-H3PWMo at room temperature and is shown in Figures 8.21 and

8.22. The scanning was done at 10 and 50 mVs-1 between -4 to +4 V using a 3-electrode

set up employing Pt and Ag/AgCl electrodes. The hybrid gel seemed to show an ESW of

~ 3V which is comparable to other IL-POM hybrids reported in the literature.99, 202

2.6 2.8 3.0 3.2 3.4-6

-5

-4

-3

-2

-1

0

present study

Nafion 115

MIMPS-H3PWMo

MIMPS-H4PMoV

PyPs-H4PWMo

PyPs-H3PWMo (present study)

PyPs-H4PMoV (present study)

log

10 (

Scm

-1)

1000/T (K-1)

present study120

100

80

60

40

20

Temperature (°C)

182


temperature between -4 to +4 V vs. Ag/AgCl with a scanning rate of 10 mVs-1.


temperature between -4 to +4 V vs. Ag/AgCl with a scanning rate of 10 mVs-1.

-4 -3 -2 -1 0 1 2 3 4

-4.0x10-4

-2.0x10-4

0.0

2.0x10-4

4.0x10-4

Cu

rre

nt (A

cm

-2)

Voltage (V)

PyPs-H3PWMo

10 mVs-1

-4 -3 -2 -1 0 1 2 3 4

-4.0x10-4

-2.0x10-4

0.0

2.0x10-4

4.0x10-4

50 mVs-1

Cu

rre

nt (A

cm

-2)

Voltage (V)

PyPs-H3PWMo

183

8.3 Summary

In this work, we successfully synthesized three novel POMs-based ionic liquid (IL) gels

by adjusting the packing efficiency of the ions through introducing the ILs into the

polyoxoanions (POMs). IR spectra and XRD patterns confirm interaction between the

organic cation and inorganic heteropoly anion in the material. The basic Keggin structure

is still remained in the hybrid materials, and TG analysis demonstrates this kind of

materials is thermally stable up to about 300 C. The bulk ionic conductivity of 0.1, 0.01

and 0.0003 Scm-1 at ~ 90 C was obtained for [PyPs]3PW11MoO40, [PyPs]4PMo11VO40 and

[PyPs]5PMo10V2O40, with activation energy of 0.45, 0.73 and 1.16 eV, respectively. The

electrochemical stability window obtained was ~ 3V. Because of the promising

electrochemical properties, these HPA and IL derived hybrid electrolytes have potential

applications in the electrochemical devices such as proton exchange membrane fuel cells.

184

Chapter Nine: Potential proton conducting hybrids derived from layered perovskites

and ionic liquid via soft-chemistry synthesis

9.1 General overview

Layered perovskites such as Dion-Jacobson (DJ) and Ruddlesden-Popper (RP) series

compounds, (e.g., KLaNb2O7 and K2La2Ti3O10) are known to undergo ion exchange

reaction with various species including organic cations.109, 110 In our continuous attempt to

develop alternative proton conductors for PEMFCs, new intercalation chemistry of layered

perovskites such as KLaNb2O7 and K2La2Ti3O10 with an imidazolium based ionic liquid

(IL) was employed and the results are discussed in this chapter. PXRD, SEM, and TGA

studies were used for characterizing the materials.


9.2.1 Intercalation Chemistry of Dion-Jacobson – DJ phase KLaNb2O7

Figure 9.1 shows the powder X-ray diffraction (PXRD) patterns of as-prepared DJ layered

perovskite-type, KLaNb2O7 along with solvent (acetonitrile), pure IL and diluted IL (with

acetonitrile) treated KLaNb2O7 powders. A complete diffraction pattern in the 2θ range of

5-80 ° is shown in Figure 9.1a and selected area of PXRD patterns at 2θ = 8 - 9 ° is shown

in Figure 9.1b to understand the peak shift after ion-exchange reaction. The possible hkl

values for KLaNb2O7 phase with space group, C222 is shown in Figure 9.1a.203 The d-

spacing corresponding to the first line (004) plane)203 is listed in Table 9.1. An increase in

d-spacing value was expected if there is any intercalation of bigger organic cations by

exchanging smaller sized K+ ions residing between the perovskite slabs.112 Only the sample

treated with pure IL has shown a marginal increase in d-value, i.e., from 10.47 to 10.77 Å

indicating some sort of displacement due to intercalation. It is to be noted that all the

185

samples retain parent layered perovskite-type structure of KLaNb2O7 after the ion-

exchange reaction.

Figure 9.1: PXRD of as-prepared, acetonitrile and IL treated KLaNb2O7 (a) complete

patterns and (b) selected area patterns at 2θ = 8-9 ° of the same to show the shift in

first peak. The indexed hkl values for as-prepared KLaNb2O7 corresponding to

orthogonal phase is marked.203

186

Table 9.1 The shift in ((004) plane of PXRD in terms of d-spacing values of as-

prepared, acetonitrile and IL treated KLaNb2O7.

Sample name d-spacing (Å)

of (040) plane

KLaNb2O7 as-prepared 10.47

KLaNb2O7 + acetonitrile 10.28

KLaNb2O7 treated with IL 10.77

KLaNb2O7 treated with IL and acetonitrile 10.38

TGA was carried out in N2 atmosphere at a rate of 5 °C/min up to 700 °C to

understand the weight loss corresponding to intercalation and result is shown in Figure 9.2.

The as-prepared and pure IL treated KLaNb2O7 show a weigh loss of 0.74 and 0.82 wt.%

whereas the acetonitrile + IL treated sample shows a 1.5 wt.% loss. The solvent

(acetonitrile) treated KLaNb2O7 sample show the maximum weight loss of 2.6 %.

All the possible reactions that could occur during ion-exchange process are summarized in

equations 9.1-9.6.

1) If IL replaces K+ in KLaNb2O7 the expected weight loss is 31.2 wt. %. The IL,

can be represented as R+X- (where R+ =

C11H22ON3+ and X- = C2S2F6O4N

-) for simple representation in the equation.

+ -

2 7 2 7KLaNb O R X "KLaNb O " KX (9.1)

RLaNb

2O

7

700 oC / air¾ ®¾¾¾¾¾ decomposed products of organic cations + "LaNb

2O

6.5" (9.2)

2) If acetonitrile, CH3CN replaces K+ in KLaNb2O7 the expected weight loss is 3.3 wt. %.

2 7 3 3 2 7KLaNb O 2CH CN "CH LaNb O " KCN (9.3)

187

CH

3LaNb

2O

7

700 oC / air¾ ®¾¾¾¾¾ "LaNb

2O

6.5"+ CO

2+1.5H

2O (9.4)

3) Assume KLaNb2O7 absorbs water and the water molecules are lost during heating. For

KLaNb2O7.1.5H2O, the expected weight loss corresponds to 1.5 mole H2O is 5.4 wt. %.

KLaNb

2O

7.1.5H

2O

700 oC / air¾ ®¾¾¾¾¾ KLaNb

2O

7+ 1.5H

2O (9.5)

3) Assume KLaNb2O7 absorbs the solvent, acetonitrile and is lost during heating. For

KLaNb2O7.1.5CH3CN, the expected weight loss corresponds to 1.5 mole CH3CN is 11.5

wt. %.

2 7 3 2 7 3

700KLaNb O .1.5CH CN KLaNb O + 1.5CH CN

oC (9.6)

Figure 9.2 TGA of as-prepared, acetonitrile and IL treated KLaNb2O7 in air at 5

°C/min.

100 200 300 400 500 600 70097.0

97.5

98.0

98.5

99.0

99.5

100.0

0.74 %

0.82 %

1.5 %

weig

ht (%

)

Temperature (oC)

KLaNb2O

7 as prepared

KLaNb2O

7 + acetonitrile

KLaNb2O

7 + IL

KLaNb2O

7 + acetonitrile + IL

2.6 %

188

The calculated theoretical weight loss in case of IL ion-exchange is 31.2 wt.% (equations

9.1 and 9.2) which indicates that no IL molecules could replace K+ in the layered perovskite

structure, as the comparative weight loss is too low. In case the solvent undergoes

intercalation in KLaNb2O7 the expected weight loss is 3.3 wt.% (eq. 9.3 and 9.4). DJ phase

compounds are known to be sensitive to moisture and there could be 1-1.5 mole water

absorbed by the layered perovskite.204 Theoretical calculation of 1.5 mole water loss

accounts for 5.4 % of weight loss (equation 9.5). If instead of water, 1.5 mole of acetonitrile

is sitting in the inter layers of perovskite, the expected weight loss during TGA is 11.5 %

(equation 9.6). Intercalation with acetonitrile was done in order to understand if the solvent

is intercalating instead of IL. It is difficult to predict the reason for weight loss, as the

experimental values are way lower than the anticipated values. As expected, there was no

weight change in the cooling cycle.

The intercalating species in KLaNb2O7 is not clear at this point, and SEM and EDS

analyses were done in order to further understand this problem. Aim of EDS study was to

understand the potassium content in the samples before and after intercalation, and SEM

was done in order to understand if there is any morphological change due to the intercalated

species. Figure 9.3 shows the SEM image of as-prepared KLaNb2O7 and EDS data

recorded at different spots which are marked in the micrograph.

189

Figure 9.3 Scanning electron micrograph and corresponding EDS of as-prepared

KLaNb2O7.

SEM shows a flaky nature for KLaNb2O7 perovskite (Figure 9.3). EDS shows the elemental

compositions matching with the theoretical formula. A slightly lower value in K-content

was observed in case of acetonitrile treated KLaNb2O7 as compared to as-prepared

KLaNb2O7 is shown in Figure 9.4. The morphology has changed from flaky to granular

nature and it could be due to the effect of acetonitrile used for the exchange reaction.

190


treated with acetonitrile solvent at room temperature for 24 h.

Figures 9.5 and 9.6 show the SEM images of pure IL treated and diluted (with acetonitrile)

IL treated layered perovskite-type KLaNb2O7. The morphology shows a similar trend as

that of acetonitrile treatment. Relatively lower K-content is observed in both cases

compared to as-prepared KLaNb2O7, which indicates that there could be a chance of

intercalation, which suggests potential partial replacement of K+ ions. However, this result

is not matching with TGA analysis, as there was no indication of ion-exchange, and hence,

further study is needed to understand the intercalation chemistry.

191


treated with IL.

192


treated with a mixture of IL and acetonitrile solvent.

In order to check the feasibility of ion-exchange capability of protonated KLaNb2O7 with

IL, the perovskite powder was treated with 2M HNO3 (as mentioned in Chapter 3, section

3.1.2.1). Even the protonated KLaNb2O7 did not undergo exchange reaction with IL

(diluted with acetonitrile), as there is no increase in d-spacing value (corresponds to the

193

first peak) in the PXRD pattern in Figure 9.7. It is also compared with the PXRD patterns

of as-prepared, acetonitrile, and acid-treated KLaNb2O7. Study of proton-exchanged

KLaNb2O7 was not carried out further as the initial PXRD results did not show any

significance of ion-exchange. The proposed reaction of proton exchange reaction of

KLaNb2O7 is as represented in equation 9.7.

2 7 3 2 7 3KLaNb O HNO HLaNb O KNO (9.7)


treated protonated KLaNb2O7.

9.2.2 Intercalation chemistry of Ruddlesdon-Popper-RP phase K2La2Ti3O10

The RP series compound, K2La2Ti3O10 was also studied in parallel using the IL, R+X-

(where R+ = C11H22ON3+ and X- = C2S2F6O4N

-). Structural change of K2La2Ti3O10 after

10 20 30 40 50 60 70 80

Two theta (degrees)

KLaNb2O

7 + acetonirile

Inte

nsity (

a.u

.)

KLaNb2O

7 + 6M HNO

3+ IL

KLaNb2O

7 + 6M HNO

3

as-prepared KLaNb2O

7

113.2

212.1 1

13

012.2

261

241

170 010.2

010.0

040

194

ion-exchange reaction was investigated using PXRD analysis and is shown in Figure 9.8.

The PXRD patterns for as-prepared K2La2Ti3O10 and acetonitrile, pure IL and diluted IL

(with acetonitrile) treated K2La2Ti3O10 were also recorded for comparison. Figure 9.8a

shows complete diffraction patterns in the 2θ range of 5-80 ° and Figure 9.8b shows the

selected area PXRD patterns at 2θ = 5 to 6.5 °. The indexed hkl values are based on the

tetragonal K2La2Ti3O10 structure with space group I4/mmm.115 It shows a shift of (002)

plane towards low-angle side, which indicates the increase in d-spacing value after

intercalation. Table 9.2 summarizes the change in d-spacing value of (002) plane, which

indicate that an increase of ~ 2.5 Å was obtained for pure as well as diluted IL treated

samples. The acetonitrile treated sample also showed a d-value increase of ~ 2 Å. Since

acetonitrile was used for washing after the ion-exchange reaction with pure IL, there is a

chance that acetonitrile might have intercalated in the layered perovskite instead of IL, as

the d-spacing values in all these cases are close to each other. The molecular structure of

IL is bigger (as shown in Chapter 2 Figure 2.16) compared to acetonitrile (CH3CN). The

increase in d-value of K2La2Ti3O10 before and after pure IL treatment was 13.94 and 16.50

Å, respectively. The expansion of inter layer spacing occurs when a bigger species replaces

K+ in K2La2Ti3O10 layered compound.115 It can be assumed that there is an ion-exchange

reaction happened for K+ in the perovskite K2La2Ti3O10. SEM studies were carried out to

support this argument.

195

Figure 9.8 PXRD (a) complete patterns of as-prepared, acetonitrile and IL treated

K2La2Ti3O10 and (b) zoomed patterns at 2θ = 5-6.5 ° of the same to show the shift in

first peak ((002) plane). The indexed hkl values for as-prepared K2La2Ti3O10

corresponding to tetragonal phase is marked.115

196

Table 9.2 The shift in (002) plane of PXRD in terms of d-spacing values of as-

prepared, acetonitrile and IL treated K2La2Ti3O10.

Sample name d-spacing of

(002) plane (Å)

K2La2Ti3O10 as-prepared 13.94

K2La2Ti3O10 treated with acetonitrile 15.80

K2La2Ti3O10 treated with IL 16.50

K2La2Ti3O10 treated with IL and acetonitrile 16.41

Thermal stability tests were carried out using TGA in N2 atmosphere at a rate of 3

°C/min up to 700 °C and subsequently, a cooling cycle was also done (Figure 9.9). No

weight loss was observed for the as-prepared K2La2Ti3O10 and it may be due to the fact

that the sample was not exposed to air for long time. The observed weight losses were used

to interpret the type of intercalation happened during ion-exchange process. The theoretical

calculations were made based on the proposed reactions shown in equations (9.8-9.11).

The pure IL treated sample shows 1.4 % weight loss (equation 9.8) which is very less

compared to the theoretical value of ~ 42 %. The weight loss observed for acetonitrile

treated K2La2Ti3O10 is 4.6 %, which is close to the predicted value (4.9 % as per equation

9.10). The diluted IL treated sample shows a weight loss of 5.3 %, which is also lower than

the theoretically predicted weight loss for IL (~ 42 %). If any potential absorbed water

molecules are expelled from the structure, the calculated weight loss is 3.9 wt. %. TGA

results show that the extent of ion-exchange of IL with K in the perovskite is minimum.

The effect of solvent, acetonitrile on IL was also monitored in this study showing more

chance of its ion-exchange capability than that of IL, which is not desired. It gives an

indication that acetonitrile should not be used for ion-exchange process. However, there is

197

a possibility of presence of water in all these samples which could account for the observed

weight losses as K2La2Ti3O10 is known to uptake water at ambient atmosphere.205

The proposed reactions for the ion-exchange reaction of K2La2Ti3O10 with different

species involved in the reaction are:

1) Assuming IL undergoes ion-exchange reaction with K+ in K2La2Ti3O10 the calculated

weight loss is 42 wt.% according to the proposed reaction below (equation 9.8). IL can be

represented as R+X- (where R+ = C11H22ON3+ and X- = C2S2F6O4N

-) for simple

representation in the equation.

K

2La

2Ti

3O

10+ 2R+X- ® "R

2La

2Ti

3O

10"+ 2KX (9.8)

2 2 3 10 2 3 9

700 /"R La Ti O " decomposed products of R + "La Ti O "

oC air (9.9)

2) If the solvent acetonitrile (CH3CN) undergoes ion-exchange reaction with K+ in

K2La2Ti3O10 the expected weight loss is 4.9 wt. %.

2 2 3 10 3 3 2 2 3 10K La Ti O 2CH CN "(CH ) La Ti O " 2KCN (9.10)

"(CH

3)

2La

2Ti

3O

10"

700 oC¾ ®¾¾¾ "La

2Ti

3O

9"+ 2CO

2+ 3H

2O (9.11)

3) If K2La2Ti3O10 loses water molecules during heating, the proposed reaction is shown in

equation 9.12. For K2La2Ti3O10.1.5H2O, the expected weight loss corresponds to 1.5 mole

H2O is 3.9 wt. %.

2 2 3 10 2 2 2 3 10 2

700K La Ti O .1.5H O K La Ti O + 1.5H O

oC (9.12)

198

Figure 9.9 TGA of as-prepared, acetonitrile and IL treated K2La2Ti3O10 in air at 5

°C/min.

SEM and EDS studies of K2La2Ti3O10 and all its ion-exchanged samples are shown

in Figures 9.10-9.13. The micrograph at high magnification is also given to monitor any

morphology change due to intercalation of organic species and the spots are marked in the

micrographs where EDS scans were recorded. In Figure 9.10 the EDS of K2La2Ti3O10

shows K-content matching with the chemical formula.

100 200 300 400 500 600 70092

93

94

95

96

97

98

99

100

5.3 %

1.4 %w

eig

ht (%

)

Temperature (oC)

K2La

2Ti

3O

10 as prepared

K2La

2Ti

3O

10 + acetonitrile

K2La

2Ti

3O

10 + IL

K2La

2Ti

3O

10 + acetonitrile + IL

4.6 %

199

Figure 9.10 Scanning electron micrograph and corresponding EDS of K2La2Ti3O10.

On treatment with the solvent, acetonitrile the K content is slightly decreased indicating a

small degree of ion-exchange (Figure 9.11). As shown in Figure 9.12, the IL treated

K2La2Ti3O10 shows ~ 50 % decrease in K-content compared to as-prepared K2La2Ti3O10

and which indicates that the K+ was being exchanged with the IL. Figure 9.13 also shows

a decrease in K-content where the intercalation was done with diluted IL. Presence of traces

of S, in EDS might be coming from the anion part of the IL, which is not washed away

200

after ion-exchange reaction. The morphology of the perovskite, which is flaky in nature,

did not show any significant change after the ion-exchange reactions.


treated with acetonitrile.

201


treated with IL.

202


treated with acetonitrile and IL.

9.2.3 Intercalation chemistry of H2La2Ti3O10

Proton-exchange reaction of K2La2Ti3O10 with 6M HNO3 was carried out and

resultant proton-exchanged H2La2Ti3O10 phase was further used for the ion-exchange

reaction with IL. The expected ion-exchange reaction of K+ with H+ is represented by

equation 9.12. The expansion of first d-spacing value was monitored using PXRD as shown

in Figure 9.14 and a slight increase in trend was observed which is similar to the results of

as-prepared K2La2Ti3O10.110, 115, 206 A drastic increase of d-value is expected if the IL is

203

intercalating between the perovskite slabs in the layered perovskite structure,115 however

the results observed in this study are not quite convincing.

2 2 3 10 3 2 2 3 10 3K La Ti O + excess HNO H La Ti O KNO (9.12)


treated protonated K2La2Ti3O10.

9.2.4 Extended study of intercalation chemistry of Ruddlesdon-Popper-RP phase

K2La2Ti3O10

Since K2La2Ti3O10 results are comparatively promising than that of DJ phase

KLaNb2O7, further studies were on K2La2Ti3O10 samples to understand the ion-exchange

chemistry with the help of one of our collaborators.207 Considering the possibility of

moisture sensitivity of the samples PXRD was recorded before and after exposing to

10 20 30 40 50 60 70 80

Inte

nsity (

a.u

.)

Two theta (degrees)

K2La

2Ti

3O

10 + acetonirile

K2La

2Ti

3O

10 + 2M HNO

3+ IL

K2La

2Ti

3O

10 + 2M HNO

3

as-prepared K2La

2Ti

3O

10

220

20.1

4

2192

17

208

200

00.1

4

10.1

3

118

116

112

107

110

109

105

008

101

004

103

006

002

204

ambient air for 48h and the results are shown in Figure 9.15-9.18. Both as-prepared and

acetonitrile treated K2La2Ti3O10 are strongly unstable in ambient air as extra peaks are

visible in Figures 9.15 and 9.16. It is clearly seen that only K2La2Ti3O10 treated with pure

IL shows the stability when exposed to air and is shown in Figure 9. 17. The K2La2Ti3O10

treated with diluted (with acetonitrile) IL shows only slight change in the diffraction pattern

after air exposure. This study indicates that the samples are air sensitive and uptake water

in ambient condition.

Figure 9.15 PXRD patterns of as-prepared K2La2Ti3O10 samples both fresh and 48h

exposed to air to check the stability in ambient condition.

205

Figure 9.16 PXRD patterns of acetonitrile (AN) treated K2La2Ti3O10 samples both

fresh and 48h exposed to air to check the stability in ambient condition.

Figure 9.17 PXRD patterns of IL treated K2La2Ti3O10 samples both fresh and 48h

exposed to air to check the stability in ambient condition.

206

Figure 9.18 PXRD patterns of acetonitrile (AN) and IL treated K2La2Ti3O10 samples

both fresh and 48h exposed to air to check the stability in ambient condition.

It is reported by Sato et al. that K2La2Ti3O10 show different crystal structure in

anhydrous and hydrated form.205 K2La2Ti3O10 crystallizes in I4/mmm space group with the

cell parameters: a = 3.8769 (1) Å and c = 29.824 (1) Å whereas the hydrated,

K2La2Ti3O10.2H2O crystallizes in P4/mmm space group with cell parameters a = 3.8585 (1)

Å and c = 16.814 (1) Å. The amount of intercalated water depends on the humidity

conditions. They reported that the water molecules occupy at the interlayer spacing. Hence,

it is assumed that the change in PXRD patterns on exposure to ambient conditions is

assumed to be due to water intercalation. To analyze the crystal structure of the fresh

samples, PXRD analysis was done and the data were refined using Fullprof software in

collaboration with Berre group.207 The results are summarized in Figures 9.19-9.22.

K2La2Ti3O10 and K2La2Ti3O10.xH2O phases were taken into account while refining the

207

data. Figures 9.19, as-prepared K2La2Ti3O10 shows the hydrated phase

(K2La2Ti3O10.xH2O), I4/mmm, with cell parameters, a = 3.8694(2) Å, c = 29.762(2) Å. The

acetonitrile treated K2La2Ti3O10 (Figures 9.20) showed a combination of K2La2Ti3O10 and

K2La2Ti3O10.xH2O phases and the refined cell parameters are a = 3.8642(5) Å, c =

29.758(4) Å with space group, I4/mmm. Both pure IL and diluted IL treated K2La2Ti3O10

showed the phase corresponds to K2La2Ti3O10.xH2O with the space group P4/mmm

(Figures 9.21-9.22). The refined cell parameters for pure IL and diluted IL treated

perovskite are: a = 3.8572(7) Å, c = 16.820(3) Å and a = 3.8602(2) Å, c = 16.8118(8) Å,

respectively. The results show that neither acetonitrile nor IL has intercalated in the IL

treated perovskites.

Figure 9.19 PXRD refinement profile of K2La2Ti3O10 considering the presence of only

hydrated , K2La2Ti3O10.xH2O phase.

208

Figure 9.20 PXRD refinement profile of acetonitrile treated K2La2Ti3O10 considering

the presence of two phases, unhydrated (K2La2Ti3O10) and hydrated

(K2La2Ti3O10.xH2O).

Figure 9.21 PXRD refinement profile of IL treated K2La2Ti3O10 considering the

presence of only hydrated, K2La2Ti3O10.xH2O phase.

209

Figure 9.22 PXRD refinement profile of acetonitrile and IL treated K2La2Ti3O10

considering the presence of only hydrated, K2La2Ti3O10.xH2O phase.

TGA measurements were repeated under inert atmosphere using Argon and show

similar results that we observed earlier (Figure 9.9). The first weight loss at < 200 °C is

due to loss of water molecules and is supporting with the PXRD data (Figure 9.20-9.22).

The second weight loss at 200-700 °C might be due to the loss of solvent, which could be

chemically bonded to the perovskites. If this is the case, we can assume that the organic

molecules are not intercalated at the interlayer space in the crystal structure. PXRD studies

were carried out for the samples after TGA measurements (i.e. after heating 700 °C) and

are provided in Figure 9.23. If the intercalation has happened during ion-exchange process,

the structure of K2La2Ti3O10 should have collapsed after heating. However, PXRD patterns

of as-prepared, acetonitrile treated, pure and diluted IL treated K2La2Ti3O10 after TGA

measurements show single phase K2La2Ti3O10 structure which indicates that K is still

present in the solvent/IL treated samples.

210

Figure 9.23 PXRD patterns of as-prepared, acetonitrile treated, pure and diluted IL

treated K2La2Ti3O10 samples after TGA measurements done in Argon up to a heating

temperature of 700 °C.

The present study shows that intercalation reaction with IL was not successful with

the layered perovskites, K2La2Ti3O10 and KLaNb2O7 under the investigated condition. The

PXRD results correspond to K2La2Ti3O10 and/or K2La2Ti3O10.xH2O phases retained after

the ion-exchange reaction under IL. The absorption of water by the sample might have

happened from the ambient air during drying step or from the solvent used for interaction

and washing.

9.3 Summary

Ion-exchange reactions of layered perovskites, KLaNb2O7 (Dion-Jacobson – DJ phase) and

K2La2Ti3O10 (Ruddlesden-Popper – RP phase) with ionic liquid, IL was studied in this

211

work to develop hybrid materials for proton conduction. The PXRD, SEM, and TGA

results show that there is a very minimum chance of intercalation of IL in between the

perovskite slabs of the layered perovskite. The PXRD studies did not show a dramatic

increase in d-spacing value corresponds to the c-plane and TGA results show very little

weight loss (~ 1 %) compared to the expected weight loss for IL (e.g. 31 and 42 wt. %).

Systematic study of optimization is required to understand the feasibility of intercalation

of currently used IL and proper solvent must be selected which can not interfere the

expected replacement of IL with K in the layered perovskites.

212

Chapter Ten: Conclusions and future work

10.1 Conclusions

The summary of all results chapters is gathered in separate sections to encapsulate the

conclusions of the current thesis work.

10.1.1 Lithium-stuffed garnet-type material research for the application of energy

storage

In Chapter 4, a systematic study of Li-stuffed garnet-type Li5+2xLa3Ta2-xYxO12 (0.05 ≤ x ≤

0.75) was performed to correlate the Y and Li co-doping with conductivity, porosity, and

sinterability. It was also aimed at understanding the materials’ chemical as well as

electrochemical stability. Powder X-ray diffraction (PXRD) studies showed that

Li5+2xLa3Ta2-xYxO12 materials have a cubic garnet-type crystal structure, with a space

group Ia-3d. Solid-state 7Li magic-angle spinning nuclear magnetic resonance (MAS

NMR) spectroscopy showed a direct correlation between the Li content and Li+ ion

mobility in these materials. This is consistent with the electrochemical impedance

spectroscopy studies that indicated higher conductivity for materials with greater Li

content. Li5+2xLa3Ta2-xYxO12 with x = 0.75 member showed a maximum conductivity of

1.83 x 10-4 Scm-1 at 23 °C. Employing scanning electron microscopy (SEM) and

Archimedes calculations of porosity and density, a correlation has been established

between the sinterability and dopant content, where an increase in sinterability was

observed with an increase in the Li and Y contents. Structural stability of these materials

is promising even after treating with water. Structural stability tests were also performed

to determine their compatibility with aqueous LiCl, and although a decrease in Li

conductivity was observed, the material was highly stability as indicated by PXRD results.

213

This makes Li5+2xLa3Ta2-xYxO12 a potential candidate for future investigation as a lithium

electrode protective layer in lithium-air batteries.

In Chapter 5, the role of excess lithium salt added during the synthesis of lithium-

garnet, specifically Li6La3Ta1.5Y0.5O12 was investigated with varying excess amounts of

LiNO3 (2.5, 5, 7, 10 and 15 wt. %). The structural analysis by PXRD showed that all the

compositions crystallized in cubic garnet-type phase and the lattice parameter was

essentially constant irrespective of the amount of excess LiNO3. SEM analysis illustrated

that the garnets with lower lithium content ≤ 5 wt.% showed significant porosity, which is

in agreement with Pycnometer density measurements. The propensity to develop lithium

carbonates on exposure to air is also independent of the excess Li content as LiCO3 peak

at ~ 1100 cm-1 was seen in all the cases as shown by Raman spectroscopy.

Li6La3Ta1.5Y0.5O12 with 10 and 2.5 wt.% excess Li showed the highest and lowest

conductivity of 1.62 × 10-4 Scm-1 and 3.77 × 10-6 respectively, at 24 °C among all other

members. This study demonstrated that the optimization of excess LiNO3 addition during

solid-state synthesis of Li-stuffed garnets could dramatically improve its properties.

Evaluation of fundamental transport properties of lithium-stuffed garnet-type

Li5+2xLa3M2-xYxO12 (M = Nb, Ta) (x = 0.25, 0.50 and 0.75) using electrochemical

impedance spectroscopy (EIS) was explored in Chapter 6 and 7. From the dielectric

analysis of these materials, it was hypothesized that the Li+ ion conduction in these

materials takes place primarily involving the Li+ ions at the octahedral sites through a

hopping mechanism. Path A and path B are two different types of Li+ ion hopping routes

(long-range migrations) observed in this study. In path A, hopping of Li+ ions takes place

between two edge-shared octahedral sites and needs a lower activation energy. In path B,

214

Li+ ions hop from one octahedral site to the other bypassing the neighboring tetrahedral

site which requires high activation energy. In Li5.5La3Ta1.75Y0.25O12 and Li6La3Ta1.5Y0.5O12

a short-range local motion of mobile Li+ ions around the immobile tetrahedral Li site was

observed at lower temperature whereas at higher temperature they showed long-range order

migrations. Both of these short-range and long-range order migrations appeared to follow

Path B with activation energies of 0.69 and 0.54 eV for Li5.5La3Ta1.75Y0.25O12 and

Li6La3Ta1.5Y0.5O12, respectively. However, Li6.5La3Ta1.25Y0.75O12 seemed to follow only

the long-range migration of Li+ ions through low energy path B with an activation energy

of 0.34 eV. The difference might be due to the larger dipole moment associated with the

lower Li containing phases (x = 0.25 and 0.50) compared to the higher one (x = 0.75).

These experimental findings support the ab initio theoretical prediction which proposed

that the energetically favourable hopping pathways of Li+ ions from one octahedral

position to another depends on the Li content in the Li-stuffed garnets. Similar observations

in terms of hopping pathways were evaluated for Li5+2xLa3Nb2-xYxO12 compared to Ta

family with activation energies 0.59, 0.52 and 0.40 eV, respectively. However, a decrease

in the relative dielectric constant was observed for Nb members compared to that of Ta

members of Li5+2xLa3M2-xYxO12 and this could be due to the high dipole moment and lower

Li+ ionic conductivity of Ta phases.

10.1.2 Proton conductors based on metal oxides and ionic liquid hybrid systems for

energy conversion

In Chapter 8, we reported the development of novel polyoxometalates (POMs), namely

H3PW11MoO40, H4PMo11VO40 and H5PMo11V2O40 and ionic liquid (3-(pyridin-1-ium-1-

yl) propane-1-sulfonate (PyPs)) based hybrid gels. Packing efficiency of ionic liquid in

215

POMs were tuned by varying the proton content in heteropolyacids (HPAs). PXRD showed

the existence of Keggin-type structure and Fourier transform infrared spectroscopy (FTIR)

spectroscopy proved the successful interaction between the organic cation and inorganic

heteropoly anion in the hybrid material. They exhibited promising thermal stability up to

300 °C, as demonstrated by thermogravimetric analysis (TGA). The maximum bulk ionic

conductivity of 0.1 Scm-1 at ~ 90 °C with an activation energy of 0.45 eV was observed for

[PyPs]3PW11MoO40 hybrid and is comparable to that of Nafion membrane at 100 % relative

humidity. [PyPs]4PMo11VO40 and [PyPs]5PMo10V2O40 showed a bulk ionic conductivity

of 0.01 and 0.0003 Scm-1, at ~ 90 °C with activation energies of 0.73 and 1.16 eV,

respectively. The cyclic voltammetry analysis revealed the electrochemical stability

window (ESW) of ~ 3V for [PyPs]3PW11MoO40 hybrid and proves its potential application

in proton exchange membrane fuel cells.

Chapter 9 demonstrated an approach to derive proton conductors from layered

perovskites, KLaNb2O7 (Dion-Jacobson – DJ phase) and K2La2Ti3O10 (Ruddlesden-Popper

– RP phase) and imidazolium based ionic liquid, IL via soft-chemistry synthesis. The extent

of ion exchange was studied by comparing the different possibilities of ion exchange

processes which include the solvent, and IL with and without solvent on K and H analogues

of KLaNb2O7 and K2La2Ti3O10. PXRD analysis showed the presence of metal oxide phases

in all ion exchanged compounds and did not show a dramatic increase in d-spacing value

corresponds to the c-plane. The findings from PXRD, SEM, and TGA studies revealed

that there is a very minimum chance of intercalation of IL in between the perovskite slabs

of the layered perovskite structure. The TGA results of ion exchanged KLaNb2O7 and

K2La2Ti3O10 showed a very minimum total weight loss of 1 wt.% compared to the expected

216

weight loss for IL of 31 and 42 wt. %, respectively after heating up to 700 °C in inert

medium. It was also found that the solvent used seemed to be competing with the IL for

the ion exchange process. Exposure to atmospheric air also proved to show structural

change due to water absorption on these layered perovskite samples.

10.2 Suggestions for future work

Here are summarized suggestions for the extension of current thesis work and future

directions along this line.

The current thesis reported that garnet-type Li5+2xLa3M2-xYxO12 (M = Ta, Nb) metal

oxides are potential Li+ ion conducting materials. All solid-state LIB could be

developed using these garnet materials as electrolyte, and LiCoO2 and graphite as

the cathode and anode, respectively. However, thin film approaches such as tape

casting, sol-gel and sputtering techniques could also be tried to develop thin

membranes for high power density Li ion batteries.

In this thesis, the effect of excess LiNO3 on the properties of garnet-type

Li6La3Ta1.5Y0.5O12 material was reported. This study could be extended with

different Li salts (e.g., Li2SiO4, and Li3PO4) in order to increase the density of the

material as maximum density is desired for garnet materials to be applied in LIB as

an electrolyte component.

Nano-sized materials exhibit unique properties compared to the bulk materials.

Another possible extension of the garnet work is optimization of the transport and

mechanical properties as a function of the particle size. Different techniques (e.g.,

sol-gel method) could be adopted to carry out this study.

217

Promising hybrid proton conductors were developed from polyoxometalates and

ionic liquid in this thesis work. Since the corresponding solid acids are known to

absorb water from the atmosphere, it is important to know the proton conduction

mechanism of its hybrid materials. Testing the ionic conductivity in different

medium such as D2O, N2/Ar, and H2 could explain the transport mechanism and

could optimize the condition for these hybrid materials to work in PEMFC working

atmospheres.

Ion exchange reactions of layered perovskites and ionic liquid reported in this thesis

work has evoked multiple questions to be answered. First of all, a systematic study

must be carried out for the optimization of proper solvent medium needed for the

exchange reaction, which should not affect the desired ion exchange reaction of IL.

Also, it is important to understand the desired size of IL which could successfully

intercalate between the perovskite slabs in the layered perovskite structure.

218

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Dielectric characteristics of fast Li ion conductinggarnettype Li5+2xLa3Nb2âˆ’xYxO12 (x = 0.25, 0.5 and0.75)S. Narayanan, A. K. Baral and V. Thangadurai, Phys. Chem. Chem. Phys., 2016, 18, 15418DOI: 10.1039/C6CP02287A

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Evaluation of fundamental transport properties ofLiexcess garnettype Li5+2xLa3Ta2−xYxO12 (x = 0.25,0.5 and 0.75) electrolytes using AC impedance anddielectric spectroscopyA. K. Baral, S. Narayanan, F. Ramezanipour and V. Thangadurai, Phys. Chem. Chem. Phys., 2014, 16, 11356DOI: 10.1039/C4CP00418C









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Title: Dopant Concentration–Porosity–LiIon Conductivity Relationshipin GarnetType Li5+2xLa3Ta2–xYxO12 (0.05 ≤ x ≤ 0.75) andTheir Stability in Water and 1 MLiCl

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Title: Superionic Conductivity inLithiumRich AntiPerovskites

Author: Yusheng Zhao, Luke L. DaemenPublication: Journal of the American Chemical

SocietyPublisher: American Chemical SocietyDate: Sep 1, 2012Copyright © 2012, American Chemical Society

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THE AMERICAN ASSOCIATION FOR THE ADVANCEMENT OF SCIENCE LICENSETERMS AND CONDITIONS

Jun 29, 2016

This Agreement between Sumaletha Narayanan ("You") and The American Association forthe Advancement of Science ("The American Association for the Advancement of Science")consists of your license details and the terms and conditions provided by The AmericanAssociation for the Advancement of Science and Copyright Clearance Center.

License Number 3898401227382

License date Jun 29, 2016

Licensed Content Publisher The American Association for the Advancement of Science

Licensed Content Publication Science

Licensed Content Title Electrical Energy Storage for the Grid: A Battery of Choices

Licensed Content Author Bruce Dunn,Haresh Kamath,JeanMarie Tarascon

Licensed Content Date Nov 18, 2011

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Solid State Electrolytes for Energy Storage and Conversion Devices

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Garnettype solidstate fast Li ion conductors for Libatteries: critical reviewV. Thangadurai, S. Narayanan and D. Pinzaru, Chem. Soc. Rev., 2014, 43, 4714DOI: 10.1039/C4CS00020J









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Solid State Electrolytes for Energy Storage and Conversion