Investigation of Ca-Doped Yttrium Iron Garnet for Solid Oxide Fuel Cells

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2020-04-28 Investigation of Ca-doped Yttrium Iron Garnet for Solid Oxide Fuel Cells

Zhang, Zheyu

Zhang, Z. (2020). Investigation of Ca-doped Yttrium Iron Garnet for Solid Oxide Fuel Cells (Unpublished master's thesis). University of Calgary, Calgary, AB. http://hdl.handle.net/1880/111927 master thesis

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Investigation of Ca-doped Yttrium Iron Garnet for Solid Oxide Fuel Cells

Zheyu Zhang

A THESIS

SUBMITTED TO THE FACULTY OF GRADUATE STUDIES

IN PARTIAL FULFILMENT OF THE REQUIREMENTS FOR THE

DEGREE OF MASTER OF SCIENCE

GRADUATE PROGRAM IN CHEMISTRY

CALGARY, ALBERTA

APRIL, 2020

Solid oxide fuel cells (SOFCs) are among the next-generation electrochemical energy conversion devices with higher energy efficiency. A major drawback hindering the mass commercialization of SOFCs is the slow oxygen reduction reaction kinetics at the cathode when the operating temperature is lowered from a high temperature (~1000 °C) to an intermediate temperature (500–750 °C). While many perovskite materials have been considered for use in novel

SOFC cathodes, there remains a number of issues, such as large thermal expansion coefficient of

Ba0.5Sr0.5Co0.8Fe0.2O3-δ due to the presence of Co and chemical instability under humidity and CO2- containing atmosphere of various perovskites. Therefore, the search, characterization and optimization of new cathode materials is of vital importance to the development of SOFCs.

In this work, garnet-type Y3-xCaxFe5O12-δ was investigated as a promising candidate for a novel SOFC cathode. Structural variations from Ca-substitution in the parent phase Y3Fe5O12 was studied using powder X-ray diffraction. Different charge compensation mechanisms because of aliovalent doping of Ca were studied and validated through iodometric titration and X-ray absorption spectroscopy. It was suggested that formation of oxide ion vacancy and electron-hole co-exist in Ca-doped garnet samples. The maximum electrical conductivity was found in the x =

0.1 garnet composition, whose ionic transport number was revealed to be about 1.82×10-5 at 750

°C. Symmetrical cells with garnet and La0.8Sr0.2Ga0.8Mg0.2O3-δ composite electrode were fabricated and evaluated for their electrochemical properties at 600–900 °C in air and at 750 °C in various oxygen partial pressures. At 750 °C, the lowest area-specific resistance in air was obtained by x =

0.3 garnet composition, with dissociation and partial reduction of adsorbed oxygen identified as rate-limiting steps.

Acknowledgements

First and foremost, I would like to express my sincere appreciation and deepest gratitude to Dr. Venkataraman Thangadurai, my supervisor, for his excellent and constant support, guidance and encouragement throughout my M.Sc. study at University of Calgary. It was a privilege to be in his lab and pursue this degree under his supervision. His expertise, personality and the attitude towards scientific research will always inspire me.

I would like to extend my thanks to my supervisory committee, Drs. Todd Sutherland and

Simon Trudel for their helpful comments and advices on my research. I would also like to thank

Dr. Michelle Dolgos for joining my examination committee. I would like to thank Dr. Viola Birss for her insightful electrochemistry course.

I am grateful to every former and present member of our research group and members in

Dr. Birss’ group whom we share office with in comradery. I want to thank Dr. Kalpana Singh for her tremendous help – this work would not have been possible without her input. I would like to have a special thank to Jialang Li. I would also like to thank my colleagues and friends: Dr. Xia

Tong, Chengtian Zhou, Miao Wang, Bin Pan, Taozhe Wu and Lei Wang. There will always be many things that I need to learn from each of you. I want to extend my gratitude to Drs. Yan Jiang,

Guangwei Wang, Tao Wang and Xiaoan Li. I am thankful to Sanoop Kammampata and Kyle

Hofstetter for their generous help with lab instruments and a lot of constructive discussions. I would like to thank Drs. Senthil Venkatesan, Arpita Nandy and Anand Singh for their useful suggestions on experimental designs. I am grateful to Drs. Sourav Bag, Alfred Samson, Scott

Paulson and Jason Young for their valuable advices and assistance. In addition, I thank Marwa

Atwa, Chengying Ai, Haris Ansari, Orrsam Abubaker and Samantha Luong. It is you all mentioned above who have made this learning environment so friendly and supportive to me.

iii

I would like to thank Janice Crawford and all other chemistry department staffs who helped me during the course of this degree program. I extend my thanks to Drs. Jigang Zhou and James

Dynes at Canadian Light Source for the collaboration on synchrotron-based X-ray absorption spectroscopy experiments. I would like to thank Dr. Yoed Tsur at Technion – Israel Institute of

Technology for his support of using impedance spectroscopy genetic programming (ISGP). I want to thank Dr. Wang Hay Kan at China Spallation Neutron Source for sharing his view of performing

Rietveld refinement.

Finally, I would like to thank the Natural Sciences and Engineering Research Council of

Canada (NSERC) and the Chemistry Department, Faculty of Science at University of Calgary for their financial support.

Dedication

I dedicate this thesis to my parents.

Table of Contents

Abstract ...... ii Acknowledgements ...... iii Dedication ...... v Table of Contents ...... vi List of Tables ...... viii List of Figures ...... ix List of Symbols ...... xii List of Abbreviations ...... xiii

Chapter One: Introduction ...... 1 1.1 Project background ...... 1 1.2 Project objectives ...... 3 1.3 Thesis organization ...... 4

Chapter Two: Background ...... 5 2.1 Solid oxide fuel cell (SOFC)...... 5 2.1.1 Basic principle of SOFCs...... 6 2.1.2 Reversible potential and cell efficiency ...... 7 2.1.3 Electrochemical performance ...... 9 2.2 Materials of SOFCs...... 12 2.2.1 Electrolytes ...... 12 2.2.1.1 Oxide ion conducting electrolytes for O-SOFCs ...... 13 2.2.1.2 Proton conducting electrolytes for H-SOFCs ...... 16 2.2.2 Anodes ...... 20 2.2.3 Cathodes ...... 22 2.2.4 Interconnects and sealants...... 25

Chapter Three: Experimental Methods...... 27 3.1 Material preparation ...... 27 3.1.1 Synthesis of garnet-type Y3-xCaxFe5O12-δ ...... 27 3.2 Symmetrical cell fabrication ...... 28 3.3 Characterization methods...... 28 3.3.1 Powder X-ray diffraction (PXRD) ...... 28 3.3.2 X-ray absorption spectroscopy (XAS) ...... 30 3.3.3 Iodometric titration ...... 31 3.3.4 Thermogravimetric analysis (TGA) ...... 32 3.3.5 Density measurement ...... 33 3.3.6 4-Probe DC measurement ...... 34 3.3.7 Hebb-Wagner polarization method ...... 35 3.3.8 Scanning electron microscopy (SEM) ...... 37

3.3.9 Electrochemical impedance spectroscopy (EIS) ...... 38 3.3.10 Impedance spectroscopy genetic programming (ISGP) ...... 40 3.4 Error considerations ...... 41

Chapter Four: Synthesis, Rietveld Refinement of Crystal Structure, Oxygen Non- Stoichiometry, Electronic Structure and Electrical Transport Properties of Garnet- Type Y3-xCaxFe5O12-δ ...... 42 4.1 Introduction ...... 42 4.2 Results and discussion ...... 44 4.2.1 Phase analysis ...... 44 4.2.2 Investigation of oxygen non-stoichiometry and electronic structure ...... 52 4.2.2.1 Oxygen non-stoichiometry ...... 52 4.2.2.2 Electronic structure at O and Fe sites ...... 56 4.2.3 Characterization of electrical properties ...... 59 4.2.3.1 Electrical conductivity and conduction mechanism ...... 59 4.2.3.2 Ionic conductivity and ionic transport number ...... 64 4.3 Summary ...... 66

Chapter Five: Electrocatalytic Properties of Oxygen Reduction Reaction (ORR) Studied by ISGP of Garnet-Type Y3-xCaxFe5O12-δ ...... 68 5.1 Introduction ...... 68 5.2 Results and discussion ...... 69 5.2.1 Chemical reactivity with LSGM ...... 69 5.2.2 Electrocatalytic properties of ORR for SOFC cathode application ...... 72 5.2.2.1 Evaluation of ORR performance ...... 72 5.2.2.2 Determination of rate-limiting step (RLS) ...... 79 5.3 Summary ...... 86

Chapter Six: Conclusions and Future Work ...... 87 6.1 Conclusions ...... 87 6.2 Future work ...... 89

References ...... 91

vii

List of Tables

Table 4.1 Structural parameters of Y3-xCaxFe5O12-δ obtained by Rietveld refinement*...... 51

Table 4.2 Structural parameters of reaction products after reduction of Y2.7Ca0.3Fe5O12-δ under 2%H2-98%N2 atmosphere in TGA obtained by Rietveld refinement*...... 56

Table 4.3 Activation energy (Ea) of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) determined by 4-probe DC method, and density of pellets used in the measurements...... 64

Table 5.1 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air...... 77

Table 5.2 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 600-900 °C in air...... 78

Table 5.3 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in various oxygen partial pressures...... 85

viii

List of Figures

Figure 2.1 Schematic illustration of working principles for O-SOFC and H-SOFC...... 6

Figure 2.2 Generalized current-voltage (I-V) characteristics for SOFC...... 11

Figure 2.3 Electrical conductivity of different electrolytes43,47–52...... 16

Figure 3.1 Optical image of a garnet-LSGM/LSGM symmetrical cell...... 28

Figure 3.2 Schematic diagram of Bragg’s law for PXRD...... 29

Figure 3.3 Schematic illustration of arrangement for 4-probe DC method...... 35

Figure 3.4 Illustration of Hebb-Wagner polarization method used for ionic conductivity measurement of Y2.9Ca0.1Fe5O12-δ...... 37

Figure 4.1 PXRD patterns of as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3, 0.5 and 0.7) powders at room temperature, and the selected PXRD pattern between 32.0 to 32.8 two- theta degrees for [024] reflection...... 46

Figure 4.2 PXRD Rietveld refinement of Y3-xCaxFe5O12-δ with (a) x = 0, (b) x = 0.05, (c) x = 0.1, (d) x = 0.3, and (e) x = 0.5. The initial structure models of Ca2Fe2O5 and CaFe2O4 are COD #9013469 and ICSD #16695, respectively...... 49

Figure 4.3 Illustration of crystal structure of garnet-type Y3-xCaxFe5O12-δ...... 50

Figure 4.4 Variation of bond characteristics of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) determined by PXRD Rietveld refinement...... 50

Figure 4.5 Evaluation of oxygen non-stoichiometry (δ) for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) sample powder by iodometric titration...... 54

Figure 4.6 Experimental TGA curve for Y2.7Ca0.3Fe5O12-δ powder under 2%H2-98%N2 atmosphere in order to achieve complete reduction at 800 °C...... 54

Figure 4.7 PXRD Rietveld refinement of reaction products in thermogravimetric hydrogen reduction method...... 55

Figure 4.8 Weight change and oxygen content variation in air as a function of temperature between 30-1000 °C for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) powders...... 55

Figure 4.9 Fe L-edge XANES spectra for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.1 and 0.5) powders...... 58

Figure 4.10 O K-edge XANES spectra for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.1 and 0.5) powders...... 58

Figure 4.11 Temperature dependence of electrical conductivity measured by 4-probe DC method for Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) sample pellets in air...... 61

Figure 4.12 Variation of electrical conductivity at 750 °C, with the inset showing p-type conductivity for x = 0.1 and 0.3 in different oxygen partial pressures...... 62

Figure 4.13 Surface SEM images of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) pellets used in the 4-probe DC conductivity measurements...... 63

Figure 4.14 Current responses as a function of time under various DC polarizations for Y2.9Ca0.1Fe5O12-δ at 750 °C in air. The inset shows the steady-state current as a function of applied potential...... 65

Figure 4.15 Comparison of ionic conductivity and electrical conductivity as a function of temperature for Y2.9Ca0.1Fe5O12-δ in air...... 66

Figure 5.1 Room temperature PXRD patterns for heated mixture of Y3-xCaxFe5O12-δ for (a) x = 0.1, (b) x = 0.3, and (c) x = 0.5) with LSGM at 1100 °C for 2 hours. “♠” indicates peaks of secondary phases seen after heating...... 71

Figure 5.2 Schematic diagram and cross-sectional SEM images of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells...... 74

Figure 5.3 Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air...... 75

Figure 5.4 DFRT plots (normalized by Rpol) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air...... 75

Figure 5.5 Temperature variations of Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell in air...... 76

Figure 5.6 Temperature variations of DFRT plots (normalized by Rpol) of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell in air...... 76

Figure 5.7 Arrhenius plot of total ASR with respect to temperature of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells in air...... 77

Figure 5.8 Oxygen partial pressure variations of Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y2.9Ca0.1Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C...... 81

Figure 5.9 Oxygen partial pressure variations of Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C...... 81

Figure 5.10 Oxygen partial pressure variations of Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y2.5Ca0.5Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C...... 82

Figure 5.11 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.9Ca0.1Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C. .... 82

Figure 5.12 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C. .... 83

Figure 5.13 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.5Ca0.5Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C. .... 83

Figure 5.14 Log (ASR) as a function of log (pO2) obtained from ISGP of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C...... 84

List of Symbols

Symbol Unit Definition

Ea eV Activation energy θ ° Angle ω Rad/s Angular frequency C F Capacitance I A Current ρ g/cm3 Density Γ - Distribution function of relaxation times

σe S/cm Electronic conductivity / - Electrode-electrolyte boundary F C/mol Faraday constant (≈ 96485.3 C/mol) f Hz Frequency R J/(mol·K) Ideal gas constant (≈ 8.3145 J/(mol·K)) Z Ω Impedance

σi S/cm Ionic conductivity

ti - Ionic transport number a Å Lattice constant δ - Oxygen non-stoichiometric value

pO2 atm Partial pressure of oxygen

Rpol Ω Polarization resistance p Pa Pressure Dependency factor of area-specific n - resistance on the partial pressure of oxygen τ s Relaxation time

Rs Ω Series resistance T K or °C Temperature t s Time V V Voltage λ m Wavelength

xii

List of Abbreviations

Abbreviation Definition AC Alternating current APU Auxiliary power unit ASR Area specific resistance BCY Yttrium-doped barium cerate BSCF Barium strontium iron cobaltite COD Crystal Open Database DC Direct current DFRT Distribution function of relaxation times ECM Equivalent circuit model EIS Electrochemical impedance spectroscopy FC Fuel cell GDC Gadolinium-doped ceria HT High temperature ICSD Inorganic Crystal Structure Database ISGP Impedance spectroscopy genetic programming IT Intermediate temperature LSC Lanthanum strontium cobaltite LSCF Lanthanum strontium cobalt ferrite LSCM Lanthanum strontium manganese chromite LSGM Strontium‐ and manganese‐doped lanthanum gallate LSM Lanthanum strontium manganite MIEC Mixed ionic-electronic conductor OCP Open-circuit potential O-SOFC Oxide-ion conducting solid oxide fuel cell ORR Oxygen reduction reaction PXRD Powder X-ray diffraction H-SOFC Proton-conducting solid oxide fuel cell RLS Rate-limiting step SEM Scanning electron microscopy SOFC Solid oxide fuel cell SSZ Scandia-stabilized zirconia TEC Thermal expansion coefficient TGA Thermogravimetric analysis TPB Triple-phase boundary XAS X-ray absorption spectroscopy YIG Yttrium iron garnet YSZ Yttria-stabilized zirconia

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Chapter One: Introduction

1.1 Project background

Energy is critical to economic development and social advancement of human society. In the world today, growing populations, rapid industrialization in developing countries and a desire for higher living standards for all individuals are driving an increasing global demand for energy1–

3. The majority of the current global energy demand has been fulfilled by the consumption of fossil fuel reserves, resulting in many environmental issues such as air and water pollution, global warming, and concerns of unsustainability due to the fact that they are finite2,4,5. As a result, cleaner and renewable energy sources need to be found, and more efficient energy conversion devices must be developed, in order to mitigate the impacts of human activities on the environment and optimize the utilization of available energy sources6.

Fuel cells (FCs) are electrochemical energy conversion devices that directly transform chemical energy of various fuels into electricity without combustion as an intermediate step3,7–9.

The efficiencies of fuel cells are not limited by Carnot engine constraints, and are generally higher

(~40%–60%) than conventional thermomechanical devices, such as an internal combustion engine

(~30–40%)4,7,9–14. Amongst all types of FCs, solid oxide fuel cells (SOFCs) stand out for their unique properties, including fuel flexibility and all-solid-state design, and attract attention due to these distinctive merits15.

The fuels that power SOFCs are hydrogen, carbon monoxide or methane16. Other common hydrocarbons can also be supplied after an internal or external reforming process17. Hydrogen is considered an ideal energy carrier. By combining it with SOFCs, the only reaction product is water.

This provides a promising clean solution for energy generation at the point of use. The fuel flexibility for the use of hydrocarbons makes SOFCs compatible with the current energy

1 infrastructures, where commercialized hydrogen supply has not been widely accessible and can be cost-prohibitive3,16. The high operating temperature of SOFCs increases the rate of electrode processes and reduces the harsh requirements for electrode catalysts1,3,16. Therefore, the use of expensive precious metals such as platinum is avoided, lowering the material costs for the device1,3,16. In addition, the unique all-solid-state feature of SOFCs provides high stability and low-noise operation benefits3,15.

Despite many advantages of SOFCs, several factors still hinder their mass commercialization. The high thermal stress caused by a potential mismatch of thermal expansion coefficients (TECs) between the electrode and electrolyte, long start-up and shut down time and high manufacturing cost of various cell components, especially sealing and interconnect materials due to high working temperature, challenge the ease of implementation3,15,16,18. Therefore, there is a collective effort in the academic community to reduce the operating temperature of SOFCs from high temperature (HT) range (750-1000 °C) to intermediate temperature (IT) range (500-750 °C)19.

However, a reduction in operating temperature is found to result in subsequent reductions in electrochemical performance of electrodes and ionic conductivity of the electrolyte, which in turn decreases the overall efficiency of the cell8. The challenges need to be addressed for IT-SOFCs to be widely accepted.

Consequently, researchers have tried many approaches to ensure that IT-SOFCs maintain the same or exceed the performance of the HT counterpart. For example, the electrolyte-supported structure has been replaced by most developers with the anode-supported configuration, in an effort to decrease the thickness of the electrolyte while minimizing the electrolyte ohmic losses8.

Besides, alternative cathode materials have also been proposed and investigated for IT-SOFCs.

Novel cathodes, such as perovskite-type Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF), are promising

8,20 candidates to replace the traditional material La1-xSrxMnO3-δ (LSM) used in HT-SOFCs . At 650

°C, the polarization resistance of BSCF (~ 0.1 Ωcm2) is around 550 times smaller than LSM

2 21 (La0.8Sr0.2MnO3-δ; ~ 55 Ωcm ) . Other double-perovskite and Ruddlesden-Popper types of compounds have also been actively studied in the field of IT-SOFC cathode22–30. Nevertheless, one type of structure – garnet, which has a wide range of applications across many disciplines, has rarely been investigated for SOFCs. It is meaningful to explore the potential of garnets, especially calcium-doped yttrium iron garnet (Y3-xCaxFe5O12-δ), which exhibits interesting electrochemical features from limited records, for IT-SOFC cathode application31,32.

1.2 Project objectives

The primary objective of this thesis was to systematically explore the effect of Ca substitution in yttrium iron garnet (Y3Fe5O12; YIG) on the phase, oxygen non-stoichiometry, electrical transport properties and electrocatalytic properties of oxygen reduction reaction (ORR).

Additionally, to overcome the disadvantages of conventional equivalent circuit model (ECM) technique, impedance spectra in the present work were analyzed by impedance spectroscopy genetic programming (ISGP), which is a MATLAB-based software developed for accurate separation of electrochemical processes based on difference in relaxation times, in an effort to gain a better understanding of the electrochemical properties of Ca-doped YIGs towards potential

SOFC cathode application.

1.3 Thesis organization

This thesis contains six chapters. Chapter 1 provides a brief introduction to the project overview, motivation and organization of the current thesis. Chapter 2 introduces the background on SOFC operating principle and electrochemical fundamentals. A short literature review regarding the development and functionality of each SOFC component is included in this chapter.

Chapter 3 presents detailed material synthesis and cell fabrication methods, as well as the characterization techniques employed in the present work.

Chapters 4 and 5 describe the results in this thesis, with each chapter having an overview, results and discussion, and summary of the analysis conducted. Chapter 4 analyzes structure variations of synthesized Ca-doped YIG (Y3-xCaxFe5O12-δ) powder. It reports the impact on oxygen non-stoichiometry and electronic structure at the O and Fe sites of Ca-doped YIG. Electrical transport properties, including the conduction mechanism and ionic transport number, are also discussed in Chapter 4. Chapter 5 is focused on electrocatalytic properties of Ca-doped YIG and the potential application as an SOFC cathode material. Evaluation of performance on oxygen reduction reaction (ORR) and the corresponding rate-limiting steps are discussed in Chapter 5.

Chapter 6 summarizes the main conclusions and outlines future work.

Chapter Two: Background

2.1 Solid oxide fuel cell (SOFC)

Solid oxide fuel cell is a high-temperature electrochemical device that enables direct conversion of chemical energy into electricity. It is comprised of all-solid-state materials, such as ceramics, metals and ceramic-metal composites for the use as electrolyte, cathode, anode, sealant and interconnect. The world first SOFC, which operated at around 1000 °C and was shaped in a tubular geometry, was demonstrated by Baur and Preis in 19379,35. The electrolyte of the cell was made from a ceramic mixture of zirconium oxide with 15% (wt.) of yttrium oxide9,33,34. Magnetite

9,33,34 (Fe3O4) was chosen as the cathode, and iron and coal were used as the anode . Eight SOFC cells were connected together to construct an SOFC stack, with a net volume of 250 cm3 and an open circuit potential of 0.83 V35. Under polarization at 0.65 V, the cell stack was able to achieve a current of 0.07 A and a power density of 0.045 W35. The performance of this cell stack stimulated a though that, with proper further developments, SOFC would be able to compete with other power generation technologies such as the steam turbine9,35. Nowadays, with years of continuous efforts,

SOFC is one of the most promising next-generation energy devices that is efficient and pollution- free at the point of use. It offers the unique advantage of fuel flexibility, avoids expensive precious metal catalysts, and has the ability to internally reform hydrocarbon fuels with heat recovery that yields an overall efficiency up to 90%7,16,35. The current targeted applications of SOFC are to distribute stationary heat and power sources and/or function as an auxiliary power unit (APU) in commercial transportations such as heavy-duty trucks and high-end passenger vehicles36. The prospects have generated great interest from academia and industry to advance SOFC technology and its potential benefits.

2.1.1 Basic principle of SOFCs

SOFCs consist of three main components: electrolyte, anode (fuel electrode) and cathode

(air electrode)37. The electrolyte is a dense ceramic ion-conducting membrane that separates the two porous electrodes. Depending on the type of ion the electrolyte conducts, SOFC can be categorized into oxide-ion conducting SOFC (O-SOFC) and proton-conducting SOFC (H-SOFC).

At the cathode, oxygen is fed as the oxidant. Whereas at the anode, a number of fuels may be chosen as the reductant, most common are hydrogen, carbon monoxide and methane. In a hydrogen

O-SOFC, when the cell is in operation, oxygen molecules at the cathode are reduced into oxide ions, these are transported through the electrolyte and oxidized by hydrogen gas into water at the anode. The electrons flow through an external electrical circuit and therefore, generate electricity.

The scenario is quite similar, but also slightly different in the case of an H-SOFC, where hydrogen molecules at the anode are reduced into protons, transported through the electrolyte and oxidized into water by oxygen molecules at the cathode. A schematic illustration containing both hydrogen

O-SOFC and H-SOFC is found in Figure 2.1.

Figure 2.1 Schematic illustration of working principles for O-SOFC and H-SOFC. 6

2.1.2 Reversible potential and cell efficiency

Assuming that one SOFC cell under isothermal and isobaric conditions, it will undergo a quasi-static reversible redox reaction at thermodynamic equilibrium. The theoretical reversible potential (ERev) of this cell can be calculated, based on the difference of Gibbs free energy (ΔG) between the reactants and the products as34:

∆퐺 = −푛퐹퐸푅푒푣 (2.1) where n is the number of electrons participating in the electrochemical reaction and F is Faraday constant (F = 96485 C/mol)34. Considering a hydrogen SOFC, where hydrogen and oxygen are the reactants and water is the product, the overall chemical reaction equation can be given as:

1 퐻 + 푂 → 퐻 푂 (2.2) 2 2 2 2

In this reaction, the number of electrons transferred is two. The Gibbs free energies (G) of hydrogen, oxygen and water, at the standard state (25 °C and 1 bar) are 0, 0 and -237.17 kJ/mol

(hydrogen and oxygen are in the gas form, while water is in the liquid form) respectively38. The theoretical cell potential is calculated, from equation 2.1, as follows:

푘퐽 푘퐽 (−237.17 )−(0 ) ∆퐺 푚표푙 푚표푙 −3 푘퐽 −3 퐸푅푒푣/ 25°퐶,1 푏푎푟 = − = − 퐶 = 1.23 × 10 ( ) = 1.23 × 10 × 푛퐹 2×96485 퐶 푚표푙

퐽 103 ( ) = 1.23 (푉) (2.3) 퐶

In reality, SOFCs operate at elevated temperatures, up to 1000 °C for conventional cells.

The condition can no longer be taken as standard state in the above calculation, since Gibbs free energy is temperature dependant and will vary when temperature changes. Furthermore, it is likely that the produced water is in the form of vapour at this operating temperature, whose Gibbs free energy also differs from that of the liquid state. Taking 1000 °C and one bar as an example, the

Gibbs free energy of water (g) at this condition is -185.33 kJ/mol, yielding a theoretical potential of 0.96 V37:

푘퐽 푘퐽 (−185.33 )−(0 ) ∆퐺 푚표푙 푚표푙 −2 푘퐽 −2 퐸푅푒푣/ 1000°퐶,1 푏푎푟 = − = − 퐶 = 9.60 × 10 ( ) = 9.60 × 10 × 푛퐹 2×96485 퐶 푚표푙

퐽 103 ( ) = 0.96 (푉) (2.4) 퐶

In addition, the partial pressure of the gases in the two electrode chambers may change from the standard state of one bar, depending on different experimental settings, which can cause cell reversible potential to vary further. Nernst equation can be applied in this case to theoretically predict the potential when gases are at non-standard state (partial) pressures (and/or temperature):

푅푇 푝 퐸 = 퐸0 − ln ( 퐻2푂 ) (2.5) 푅푒푣 푛퐹 1 푝 푝2 퐻2 푂2 where E0 is the reversible potential at standard state, R is the ideal gas constant, T is temperature,

and 푝퐻2푂, 푝퐻2, 푝푂2 are the partial pressure of water, hydrogen gas and oxygen gas, respectively.

In all, it can be seen from above that, both operating temperature and gas partial pressure are influential to the theoretical reversible cell potential, and should be carefully considered when making predictions.

On the other hand, besides the cell potential, efficiency is also an important measure when evaluating an energy conversion device. The theoretical efficiency (ε) of an SOFC is defined as11:

∆퐺 휀 = × 100% (2.6) ∆퐻 where ΔG is the change of Gibbs free energy in the reaction, indicating the amount of energy available to generate electricity. And ΔH is the change of enthalpy, meaning the total energy that can be released by the chemicals in the reaction. For example, at 1000 °C and one bar, the change of enthalpy ΔH of the reaction of interest (equation 2.2), is approximately -241.82 kJ/mol37. The

8 change of Gibbs free energy ΔG is -185.33 kJ/mol. Therefore, the theoretical efficiency of this hydrogen SOFC operating at 1000 °C and one bar is calculated as:

−185.33(푘퐽/푚표푙) 휀 = × 100% = 77% (2.7) −241.82(푘퐽/푚표푙)

Comparison can be drawn to an internal combustion engine, whose theoretical efficiency is governed by the principle of Carnot cycle39:

푇 휀 = (1 − 퐶 ) × 100% (2.8) 푇퐻 where TC is the absolute temperature of the cold reservoir and TH is the absolute temperature of the hot reservoir. For a heat engine that operates at 400 °C (673.15 K), TH is considered the temperature of this engine, and TC is the outside environment temperature (25 °C; 293.15 K), where the exhaust is ejected11. The theoretical efficiency of internal combustion engine at these given conditions is determined by:

293.15(퐾) 휀 = (1 − ) × 100% = 56% (2.9) 673.15(퐾)

From the calculations above, it can be seen that the hydrogen SOFC usually has a higher theoretical efficiency than an internal combustion engine for converting fuels to energy. It should also be noted that, although the conditions chosen for either a hydrogen SOFC operating at 1000 °C, or an internal combustion engine operating at 400 °C, may be extreme, and could lead to variations of numbers used in the comparison. However, the benefits obtained from the advanced operating principle of SOFC in terms of theoretical efficiency are still conveyed.

2.1.3 Electrochemical performance

Practically, the actual efficiency of a hydrogen SOFC hardly matches the theoretical value.

The deviation may be due to parasitic side reactions, where a portion of the total current is

9 consumed, resulting in a drop of overall efficiency below the thermodynamic value1. Besides, several kinds of irreversible losses present during cell operation, indicated by a difference between theoretical and actual cell potentials, also contribute to the decrease of overall cell efficiency1.

For an ideal SOFC, theoretical cell potential would be maintained, regardless of the current generated. However, given all irreversible losses existed in the non-ideal case, performance of this cell cannot be simplified and is usually evaluated by current-voltage (I-V) characteristics, as shown in a generalized illustration in Figure 2.2.

At the open circuit potential (OCP), the actual measured value for the cell is often not the same as the theoretical thermodynamic one, represented by the two purple arrows in Figure 2.2.

The potential difference is denoted by UL. It indicates that the electrolyte material used is not a perfect ionic conductor, and possesses residual electronic conductivity, which causes internal short circuits and the subsequent loss of potential37,40,41. Another possible scenario is that the electrolyte may have micro cracks and fissures, causing the gases on each sides to cross over37.

Upon drawing current from the cell, the actual potential is further reduced, mainly because of three types of losses: 1) activation polarization (ηact) from the slow charge transfer processes,

2) ohmic losses (ηohmic) due to the electronic and ionic resistance and 3) concentration polarization

11,37 (ηconc) attributed to the slow mass transfer processes . Therefore, the actual cell potential is expressed as:

퐸퐶푒푙푙 = 퐸푅푒푣 − 푈퐿 − 휂푎푐푡 − 휂표ℎ푚𝑖푐 − 휂푐표푛푐 (2.10)

The activation losses are the extra potential required to overcome the energy barrier of the rate- liming step in the electrode reactions. The slow chemical reaction kinetics could be from, for example, the reduction of adsorbed oxygen atom into oxide ion, which involves a limited rate of charge transfer between the reactants and products. Since the reaction kinetics is associated with

10 many factors such as temperature, pressure and the electrode material. Improvements can be made by increasing the reaction temperature and pressure and by enhancing the catalytic activity of the electrode through strategic material selection9. The region in I-V curve where activation losses are dominating is shown in the blue background in Figure 2.2.

At a higher current density, ohmic losses become dominant, represented by the light orange background region in Figure 2.2. The ohmic losses arise from resistance in the electronic and ionic conduction of the cell and are proportional to the current density value.

When current density is extremely high, meaning that chemical reactions at the electrodes are occurring very rapidly, it is likely that mass transport of both bringing in reactants and removing products will be left behind. In this instance, the system is dominated by the concentration losses, shown in the green background region in Figure 2.2, which are mainly attributed to the reactant depletion and product clogging at the electrodes9.

Theoretical Reversible Potential

Actual Reversible Potential

Activation Ohmic Concentration Cell Potential (V) Potential Cell Polarization Polarization Polarization

2 Current Density (A/cm )

Figure 2.2 Generalized current-voltage (I-V) characteristics for SOFC. 11

2.2 Materials of SOFCs

In order for a single SOFC cell to achieve high performance and good stability during long- term operation, it is critical to have its materials selected, based on a number of strict criteria, and with sufficient foresight.

For example, the thermal expansion coefficient (TEC) of each cell component needs to be as close as possible8. A mismatch of TECs at high working temperature risks mechanical fracture and material delamination, which are likely to induce high contact resistance and malfunction of the cell8. Other desired properties include chemical stability between different components: the materials of electrodes are not supposed to react with the electrolyte to form insulating layers, preventing the conductions of ions and decreasing the current density8. The total cost of manufacturing such a cell also must be kept low to allow mass production and the potential commercialization. In addition, there are also several component-specific requirements besides these general rules and they are discussed in the following sections.

2.2.1 Electrolytes

The primary role of a SOFC electrolyte is to conduct ions between the two electrodes, balancing the charge from the electron flow to complete the electrical circuit of the cell. Since the electrolyte works as a separator between the fuel and oxidant, it needs to be stable in both reducing and oxidizing environments3. Mechanical properties such as being dense and gas-tight are also necessary to avoid the cross-over and recombination of the gases3. Additionally, the electrolyte is required to possess high ionic conductivity, with ideally no electronic conductivity in the dual working atmospheres and at the elevated temperature3. There are generally two kinds of SOFC electrolyte, categorized by the ions it conducts: oxide ion conducting electrolyte and proton

12 conducting electrolyte. In some cases, the electrolyte materials can be found to conduct both oxide ion and proton.

2.2.1.1 Oxide ion conducting electrolytes for O-SOFCs

Currently, the most widely used oxide ion conducting electrolyte is yttria-stabilized

42 zirconia (YSZ), with general chemical formula of (ZrO2)1-x(Y2O3)x (commonly, 0.08 ≤ x ≤ 0.1) .

Zirconium dioxide (ZrO2), in its pure form, is monoclinic (space group: P21/c) at room

43 temperature . As temperature rises, the structure changes into tetragonal (space group: P42/nmc) above 1170 °C, and further into cubic (space group: Fm3m) over 2370 °C43. It is then found that with the addition of certain aliovalent metal oxide, the phase transitions, which often involve abrupt crystal volume changes and inhibit the long-term mechanical stability in an integrated system, can be avoided. As a result, the cubic fluorite phase is stabilized regardless of temperature43. Furthermore, by substituting the aliovalent cation at the zirconium site, a large concentration of oxide ion vacancies is created, which in-turn enhances the oxide ion mobility in the lattice and improves the material’s ionic conductivity34.

When the dopant is yttrium oxide (Y2O3), the formation of YSZ is expressed by the defect equation using Kröger-Vink notation as17:

(푍푟푂 ) 푌 푂 2 2푌′ + 3푂× + 푉∙∙ (2.11) 2 3 → 푍푟 푂 푂

′ 3+ 4+ × where 푌푍푟 indicates Y occupying the Zr site and possessing one negative net charge, 푂푂

2+ ∙∙ indicates O at its original site with neutral charge and 푉푂 represents an oxygen vacancy with two positive net charges.

YSZ is considered the traditional electrolyte for HT-SOFC. It exhibits a high oxide ion conductivity at around 1000 °C (~ 0.1 S/cm) and is relatively stable in both oxidizing and reducing

13 atmospheres, with negligible electronic leakage present in a wide range of oxygen partial pressures

11,19,34 (pO2) . Since zirconium is one of the abundant elements in the earth crust, the availability of the raw material also helps to keep the overall manufacturing cost low7,43.

However, the high working temperature required for the use of YSZ, despite of its popularity, have led to many practical problems, including the fast degradation rate of the cell and the limited choice of materials for cell fabrication44,45. Potentially, these issues could be eliminated by switching the electrolyte to the one that has an adequate ionic conductivity at a relatively lower temperature, for example, in the IT range44.

Among the candidates investigated, scandia-stabilized zirconia ((ZrO2)1-x(Sc2O3)x; SSZ) is

44 a possible alternative to YSZ . It originates from the same parent ZrO2 phase as YSZ, but uses scandium oxide (Sc2O3) as a stabilizer. Although scandium may be too expensive for widespread application, the conductivity of SSZ does show a markedly higher value than YSZ, as summarized in Figure 2.3 along with a variety of other kinds of electrolyte materials44.

Besides, another electrolyte material with the same cubic fluorite phase as stabilized-

43 zirconia, is cerium oxide (CeO2) . By adopting a similar aliovalent doping strategy, its ionic

43 conductivity can be greatly improved . Gadolinium oxide (Gd2O3) is one of the most commonly used dopants, forming the solid solution of gadolinium-doped ceria ((CeO2)1-x(Gd2O3)x; GDC), where the optimum mole fraction of gadolinium is found to be 0.1 in the compound46. Other doping

43 options include samarium oxide (Sm2O3), lanthanum oxide (La2O3) and yttrium oxide (Y2O3) .

Overall, doped-ceria has high conductivity, as shown in Figure 2.3 for the selected doping compositions. However, one critical drawback of doped-ceria is that it tends to undergo a reduction

4+ 3+ of Ce to Ce in the low pO2 environment and at high temperatures (> 600 °C), which induces

14 an electronic current by allowing electrons to hop between cerium ions of different valences and therefore limits its performance as electrolyte35,43.

In addition to the aforementioned fluorite-structured electrolytes, perovskite-based oxides such as lanthanum gallate – doped with Sr and Mg (La1-xSrxGa1-yMgyO3-δ; LSGM) also exhibit strong potential for the role of a structured electrolyte16,44,46. LSGM has a good compatibility with many of the electrode materials in SOFC, given the fact that these electrodes are commonly of the

16 same perovskite structure . It shows high ionic conductivity and is also stable in low pO2 environment, without experiencing significant electronic leakage44. The primary challenges with

LSGM are the complicated preparation steps, including harsh conditions for densification of the electrolyte layer and possible Ga evaporation in sintering, such that making a desired single-phase

16 compound can be difficult . Other materials such as LaSrGaO4 of Ruddlesden-Popper phase,

Ba2In2O5 of brownmillerite phase, and Gd2Ti2O7 and Y2Zr2O7 of pyrochlore phase, with their corresponding benefits and limitations, have also been explored as SOFC electrolytes by researchers34.

Figure 2.3 Electrical conductivity of different electrolytes43,47–52.

2.2.1.2 Proton conducting electrolytes for H-SOFCs

While the oxide ion conductors are of the traditional focus by the researchers, there are other efforts in the development of proton conducting ceramics, which typically have lower activation energies (0.4 – 0.7 eV) due to the unique properties of the proton, and are expected to exhibit higher ionic conductivity at reduced (IT or even lower) operating temperature46,53. When the proton conductor is exposed to the humidified environment, water from the gas phase is combined with an oxide ion vacancy in the lattice, and dissociates into a hydroxide ion and a proton46,54. The hydroxide ion fills the oxide ion vacancy site, whereas the proton forms a covalent

16 bond with a normal lattice oxygen46,54. This dissociative adsorption process of water is expressed as46,54:

∙∙ × ∙ 퐻2푂 + 푉표 + 푂푂 ↔ 2(푂퐻)푂 (2.12)

∙ where (푂퐻)푂 represents a hydroxide ion occupying an oxide ion site with one net positive charge.

The proton incorporated in the ceramic, in the form of hydroxide ion, can migrate through the lattice by hopping between the adjacent oxide ion via the Grotthuss-type mechanism, where the steps of rotational diffusion and the transfer of the proton are involved46.

On the other hand, proton conduction can also take place in a dry hydrogen atmosphere. In this case, the proton uptake process in the electrolyte is written as46:

∙ × ∙ 퐻2 + 2ℎ + 2푂푂 ↔ 2(푂퐻)푂 (2.13) where ℎ∙ indicates a hole with one positive net charge. The hole in the material is formed during the high-temperature sintering in an oxygen-rich environment by the reaction between gaseous oxygen molecule and the lattice oxygen vacancy, shown as43:

1 푂 + 푉∙∙ ↔ 푂× + 2ℎ∙ (2.14) 2 2 표 푂

It can be seen, from the discussion above, that the degree of proton incorporation in the electrolyte material, which is influential to the overall proton conductivity, is closely related to the concentration of oxide ion vacancy present in the lattice17. The acceptor doping route, where the electrolyte elements are partially or fully substituted with the alternatives of lower valences, is a common strategy adopted by scientists for the creation of additional oxide ion vacancies55. Based on charge neutrality, this substitution process, for example, where a tetravalent cation (B4+) is replaced by a trivalent cation (M3+), can be written as46:

× × ′ ∙∙ 2퐵퐵 + 푂푂 + 푀2푂3 → 2푀퐵 + 푉푂 + 2퐵푂2 (2.15)

Therefore, the relationship between the oxide ion vacancy concentration and the dopant concentration can be expressed, according to equation 2.15, as follows46:

1 [푉∙∙] = [푀(퐼퐼퐼)] (2.16) 푂 2

The high temperature proton conductivity in ceramic oxides was first discovered by Iwahara et al.

56 in doped SrCeO3 in 1981 . Afterwards, a variety of materials have been explored, with several groups of perovskites being the most extensively studied, including SrCeO3, BaCeO3 and SrZrO3- based chemicals55.

Doped strontium cerates, such as SrCe0.95Yb0.05O3-δ, have been reported to exhibit

43 - significant proton conductivity . At 600 °C, SrCe0.95Yb0.05O3-δ showed ionic conductivity of 2×10

3 43,57 S/cm in 5% H2 (balanced N2) with an activation energy of 0.59 eV . Other families of the

43 doping materials include SrCe1-xScxO3-δ and SrCe1-xYxO3-δ (x = 0.05 or 0.1) . The conductivities of selected doped strontium cerates are shown in Figure 2.3.

Alternatively, materials from the BaCeO3 family have been revealed to have proton conductivity often even higher than doped SrCeO3, possibly due to a larger lattice parameter of

58 the parent phase BaCeO3, which provides additional free volume for the protons to migrate .

However, barium cerates are found to be unstable in the humidity and CO2-containing atmosphere, given the fact that a larger lattice constant also gives arise to a longer and weaker metal-oxygen bond, and causes unwanted reactions between the cerium and water or CO2 to form the barium hydroxide or barium carbonate, shown as follows58:

퐵푎퐶푒푂3 + 퐻2푂 → 퐵푎(푂퐻)2 + 퐶푒푂2 (2.17)

퐵푎퐶푒푂3 + 퐶푂2 → 퐵푎퐶푂3 + 퐶푒푂2 (2.18)

The partial substitution of Zr at Ce site yields a solid solution of BaCe1-xZrxO3 and improves the

54 stability of the barium cerates . Un-doped BaZrO3 is found to hardly react with water and CO2,

18 but usually exhibits an overall proton conductivity lower than BaCeO3, owing to the presence of high grain boundary resistance from unqualified sintering condition54. The un-doped barium zirconate also requires a high sintering temperature over 1700 °C to obtain a dense ceramic layer54.

Therefore, the combined solid solution of barium cerate and zirconate is expected to own complementary properties with a balance of chemical stability and proton conductivity.

By further introducing the aliovalent doping of trivalent cation at Zr/Ce site, materials such as BaCe0.8Zr0.1Y0.1O3-δ, BaCe0.8Zr0.1Nd0.1O3-δ and BaCe0.8Zr0.1Gd0.1O3-δ have been demonstrated with high performance and are also water vapour- and CO2- tolerant to some extent.

Besides perovskites, ortho-niobates and ortho-tantalates, described as Re1-xAxMO4 (Re =

La, Nd, Gd, Te, Er or Y; A = Ca, Sr or Ba; M = Ta or Nb), have been studied by R. Haugsrud et al. and are discovered to show dominant proton conductivity up to 1000 °C46,59,60. For example,

-3 the ortho-niobate La0.99Ca0.01NbO4 is found to have a maximum conductivity of 10 S/cm at 800

°C whereas the ortho-tantalate La0.99Ca0.01TaO4 at the same temperature has a conductivity of

5×10-4 S/cm58. However, ortho-niobates undergo phase transformation from monoclinic to tetragonal at around 500 °C, which is prone to cause abrupt thermal expansion and severely prohibits the practical application of the materials46.

K. Liang et al. have reported the high proton conductivity of the non-stoichiometric double

61 perovskites such as Ba2Ca0.67Nb1.33O6 below 600 °C . Bhella et al. made a partial substitution of

Ta for Nb in the same compound and successfully enhance the chemical stability of the material

62 in the water and CO2 – containing atmosphere .

2.2.2 Anodes

The anode of an SOFC is subject to exposure to extremely reducing environment with a

-20 pO2 as low as around 10 . Therefore, the chemical stability in this reducing atmosphere has a vital role in the selection of anode materials. In addition, the material chosen must be an excellent combination of high catalytic activity towards fuel oxidation reaction, high ionic conductivity and high electronic conductivity17. The catalytic property is necessary to ensure the fast reaction kinetics, while the ionic and electronic conductivities are essential to allow ions to spread over a large area of anode/electrolyte interface, and electrons to convey rapidly through the external circuit after being produced17. A high degree of porosity (20% - 40%) and a large area of triple- phase boundary (TPB) are also required, given the fact that an enhanced fuel supply and product release rate and an increased electrochemical active surface area are known to improve the overall reaction yield63.

In the very early stage of SOFC development, single-phase materials, such as platinum, iron, graphite and nickel were proposed34. However, it was found that platinum, besides being expensive, tends to spall off from the electrode surface, possibly due to the evolution of water vapor34. Graphite is seen to corrode by electrochemical reactions while iron is prone to be oxidized when the partial pressure of oxidants is above a certain level34. The pure nickel metal experiences a large TEC mismatch with common SOFC electrolytes and is likely to aggregate at high temperature through grain growth, which decreases the TPB area and the electrode porosity34.

In 1970, a patent filed by General Electric Co. with inventor H.S. Spacil was published, introducing an electrode made from Ni-YSZ cermet64. The combination of nickel and YSZ brings many benefits and has made the mixture the conventional state-of-the-art SOFC anode material.

The presence of YSZ in the anode is able to support nickel metal particles and inhibit them

20 coarsening together, to some extent, at the operating temperature of SOFC43. It can also mitigate the difference in TECs between the anode and electrolyte43. The other advantages of using Ni-YSZ cermet include the additional ionic conductivity from YSZ part to the anode and the resulting extended electrochemically active areas of TPB43.

Apart from the aforementioned merits, there are still a few drawbacks for the Ni-YSZ anode, which hinder further improvements of cell performance. One of them is the mechanical instability due to dimensional changes of Ni during the alternating oxidation and reduction processes, termed “redox cycling” (equation 2.19). In the manufacturing of Ni-YSZ cermet, Ni is usually put as NiO, which is to be reduced to its metallic form in the actual reducing operating environment of anode65. On the other hand, the oxidation of Ni back to NiO occurs when there is a shortage of fuel supply at the anode such that the atmosphere becomes less reducing, or the SOFC system is shut down by the operator with its anode exposed to the ambient air65. A large volume change associated with this redox reaction in equation 2.19 is as high as ~65%, which can result in severe mechanical and the corresponding electrochemical performance degradation of the cell65–

67. Other issues of a Ni-YSZ anode involves potential poisoning of sulphur and coking.

1 푁𝑖 + 푂 ↔ 푁𝑖푂 (2.19) 2 2

Alternative anode materials to Ni-YSZ cermet have been explored by scientists, for example, perovskite-based La0.75Sr0.25Cr0.5Mn0.5O3 (LSCM), which is shown to be stable in both fuel and air atmospheres34,46. The electronic conductivity of LSCM is revealed to be p-type and

46 decreases with decreasing pO2 from ~ 38 S/cm in air to ~ 1.5 S/cm in 5% H2/Ar at 900 °C . It has good electrochemical activity towards the oxidation of methane in the humidified environment, and a quite small volume change ~ 1% during the redox cycling34. Research has suggested that the coking issue with LSCM can be minor when hydrocarbons are used as the fuels. However, it is

21 reported that the material may have less stability when in contact with high concentration of

34,46 H2S . Other promising candidates of novel SOFC anode include doped-SrTiO3 (La-doped

SrTiO3 and Y-doped SrTiO3), double perovskite-based Sr2FeNbO6, Sr2MgMoO6 and Sr2MnMoO6,

46 and tungsten-bronze-structured niobates (Sr0.2Ba0.4Ti0.2Nb0.8O3) .

2.2.3 Cathodes

The cathode material for SOFC needs to possess high catalytic activity for oxygen reduction reaction (ORR), as well as high ionic and electronic conductivity (preferably mixed ionic-electronic conductor (MIEC))46. Similar to the requirement of an excellent anode, it is important for the SOFC cathode to have sufficient porosity for gaseous oxygen diffusion and good stability in an oxidizing environment46.

The conventional state-of-the-art SOFC cathode material is the perovskite-based

58 strontium-substituted lanthanum manganite La1−xSrxMnO3−δ (LSM) . The aliovalent substitution

2+ 3+ of Sr for La (20%-30%) in the parent phase lanthanum manganite LaMnO3 gives rise to the formation of a certain concentration of Mn4+ in the lattice, which in turn improves p-type electrical conductivity and electrocatalytic properties of oxygen reduction reaction (ORR)46. Given the limitation that LSM only offers acceptable performance at temperatures over 900 °C, in order for

SOFCs to operate at lower temperatures, for example, in the IT range, one strategy is to mix LSM with another highly ionic conducting material (typically the electrolyte material that is used in the same cell) to make a composite cathode46. The active surface area is increased in this way, such that the drawbacks of being mostly electronic conducting and lacking adequate ionic conductivity for LSM is mitigated46. Examples of a composite LSM cathode include the work on an x: y wt.%

LSM:YSZ electrode as reported by Kim et al., where the lowest ASR of 0.04 Ωcm2 at 950 °C and

68 0.2 atm of pO2 was found when x: y = 60:40 .

An alternative strategy is to directly use MIEC as the SOFC cathode, instead of making a

46 composite . Lanthanum cobaltite-based perovskites, such as La1-xSrxCoO3-δ (LSC), are known to

46 show mixed electronic and ionic properties . By further substituting Co with Fe to form La1- xSrxCo1-yFeyO3-δ (LSCF), the phase stability can be enhanced and the mixed conductivity is kept

58 at a reasonable level . For instance, the composition La0.6Sr0.4Co0.8Fe0.2O3-δ has been investigated, which is revealed to possess an electronic conductivity of 269 S/cm and an ionic conductivity of

0.058 S/cm69. Other studies in the literature look into replacing La with Ba in LSCF, yielding the compound Ba1-xSrxCo1-yFeyO3-δ (BSCF). The work done by Shao et al. using the composition

2 Ba0.5Sr0.5Co0.8Fe0.2O3-δ demonstrates an ASR as low as 0.055-0.071 Ωcm at a reduced operating temperature of 600 °C in air under open-circuit potential and symmetrical cell configuration20.

However, although the aforementioned materials have been studied extensively, with some even being available in the market at present, yet several disadvantages still exist and are hindering further applications46. One of the main concerns is the large mismatch of TECs between the Co-

46 based perovskite cathode and the common SOFC electrolyte . La0.6Sr0.4Co0.8Fe0.2O3-δ is reported

-6 to have a TEC of 21.4 (×10 /K), while Ba0.5Sr0.5Co0.8Fe0.2O3-δ is found to exhibit a similar value of 20 (×10-6/K)69. On the other hand, most of the common electrolytes, such as YSZ, GDC and

LSGM, are recorded to have TECs ~ 10 (×10-6/K), which differ nearly half from those of Co-based

-6 -6 perovskite cathodes (Zr0.92Y0.08O2-δ: 10.5 (×10 /K); Ce0.8Gd0.2O1.9: 12.5 (×10 /K);

-6 69 La0.9Sr0.1Ga0.8Mg0.2O3-δ: 10.7 (×10 /K)) . The large difference in TECs tends to cause delamination issue during cell operation at elevated temperature and eventually leads to cell

23 failure. Other problems include reacting with YSZ to form insulating layers of La2Zr2O7 and

46 SrZrO3 as secondary phases .

Among the Co-free perovskite candidates, ferro-nickelate LaNixFe1-xO3-δ is an attractive alternative material, with TEC value matching that of GDC electrolyte and the ability to resist Cr

58 2 poisoning . Rajendra et al. showed that for the composition LaNi0.6Fe0.4O3, an ASR of 0.06 Ωcm was achieved at 800 °C in air, and a high electrical conductivity of 305 S/cm was measured70. The

-6 70 corresponding TEC at 800 °C was determined to be 11.8 (×10 /K) . Moreover, SrFeO3-δ and

58 3+ Bi0.5Sr0.5FeO3-δ are also promising Co-free cathodes for IT-SOFC . With the substitution of Bi , an extra lone pair is present and is able to improve the electrocatalytic activity of the material

58 significantly . The activation energy of Bi0.5Sr0.5FeO3-δ is found to be smaller than other typical

Co-free electrodes58.

Besides simple perovskite-type oxides, double perovskite is another group of novel cathode material. It has a general formula of AA’M2O5+δ, where A is usually a lanthanide cation and A’ an alkaline-earth cation46. A 3-d element often occupies the M site46. The compound has an ordering of A and A’ cations in the lattice, such that peculiar structural and electrical properties, which are not seen in normally A-site disordered doped-perovskite, can be observed46. In particular, layered cobaltites LnBaCo2O5+δ (Ln = Gd, Pr, Y, and La) have been extensively studied in the recent years46. With the transformation from cubic perovskite phase into layered structure, the oxygen bonding strength is undermined in the layer of A’Oδ, providing more free space for the ion to migrate through the channel and enhancing the material’s ionic conductivity46. At 600 °C, an ASR

2 46 of 0.213 Ωcm was obtained using PrBaCo2O5+δ, mostly due to its fast surface exchange kinetics .

2 Brownmillerite-type oxides such as Ca2Fe2-xCoxO5 (ASR of 0.23 Ωcm at 700 °C in air) and

Ruddlesden-Popper-type oxides such as La2CuO4 and La2NiO4 are all considered as promising candidates for the next generation SOFC cathode46.

2.2.4 Interconnects and sealants

Since each individual SOFC cell is able to provide only an open circuit potential of ~ 1 V, an interconnect is needed to connect multiple cells in series (to construct the SOFC stack) for higher voltage and greater power output71. The interconnect links the anode of a previous cell to the cathode of a subsequent cell within an SOFC stack, and is therefore required to be stable in reducing and oxidizing environments43. It also functions as a physical separator of the fuel and oxygen gases between the cells and thus must be dense and impermeable43. In order to minimize

Ohmic losses, it is important for the material to have a high electrical conductivity, which does not

43 vary significantly in a large pO2 range . In addition, good chemical and thermomechanical stability with other cell components at elevated temperature is necessary43.

From the 1970s, a lanthanum chromite (LaCrO3) – based ceramic interconnect has been

43 widely used for HT-SOFCs . The parent phase LaCrO3 is a p-type semiconductor, with ~ 1 S/cm electrical conductivity at 1000 °C in air71. It has the advantage of being chemically stable both at

71 high temperature (melting point ~ 2510 °C) and in low pO2 atmosphere . The electrical conductivity can be tuned by substituting the A-site cation of La3+ with Ca2+ or Sr2+, and B-site

3+ 2+ cation of Cr with Mg . For example, the Ca-doped compound La0.7Ca0.3CrO3 is reported to have an electrical conductivity of 43 S/cm at 1000 °C, with a TEC value of 10.2 (×10-6 K-1) compatible

71 with many common SOFC electrolytes and electrodes . One of main drawbacks of LaCrO3 – based ceramic oxides is the high temperature (~ 1600 °C) required to sinter a dense interconnect

25 layer71. Furthermore, poor thermal conductivity of the ceramics could lead to high thermal stresses and eventually cause mechanical failure of the device71.

When the operating temperature of SOFC is reduced to the IT range, metallic interconnect becomes another option58. Cr-based alloys, such as Cr-Fe and Cr-Ni are often used in this case58.

The metal or alloy is able to provide excellent thermal and electrical conductivities, and is easy to fabricate into an integrated device58. However, there remain further concerns regarding the mismatch of TECs, potential oxidation of metals and ion diffusion between the interconnect and electrode58.

For the planar design of SOFC, sealing is also required to keep the different gases separate, in addition to the use of interconnect71. Glass and glass-ceramic sealants are generally chosen, for their electrical-insulating nature and low-cost58. Moreover, they are easy to be processed and are stable in a variety of environments58. One major problem associated with glass and glass-ceramic- based materials is the brittleness to cracking when there is high tensile stress58. On the other hand, the design of a tubular SOFC, which is also referred to as the “seal-less design”, offers the benefit of not requiring any sealants in the current instance71.

Chapter Three: Experimental Methods

3.1 Material preparation

3.1.1 Synthesis of garnet-type Y3-xCaxFe5O12-δ

Garnet-type metal oxides with the nominal chemical formula Y3-xCaxFe5O12-δ (x = 0, 0.05,

0.1, 0.3, 0.5 and 0.7) were synthesized through sol-gel method72–74. Starting materials with stoichiometric amount (0.007 mole of the targeting synthesized garnet product) of Ca(NO3)2·4H2O

(99%, Alfa Aesar), Y(NO3)3·6H2O (99.9%, Alfa Aesar) and Fe(NO3)3·9H2O (98%-101%, Alfa

Aesar) were dissolved in distilled water and mixed at room temperature by stirring. Required amount of citric acid (C6H8O7, Fisher Scientific), with citrate to total metal ion molar ratio of 1:1, was dissolved in another container and later added to the nitrate solution as a chelating agent. The pH was monitored using pH paper and adjusted to around 8 with addition of ammonium hydroxide

o (NH4OH 28-30%, Sigma-Aldrich). The solution was heated to ~ 80 C under constant stirring until it turned into gel. The gel was then aged at room temperature for at least 12 hours and subsequently calcined in the oven at 300 oC for one hour. Powder was collected and grounded using a mortar and pestle and sintered in air at 1100 oC for two hours in order to form the desired garnet phase.

About four grams of the sintered powder was cold pressed into a cylindrical disk (~1cm diameter, ~2cm thickness) rubber mould under an isostatic pressure of 200 kN for 5-10 minutes75.

Due to the fine porous nature of the synthesized powder, the pressing process was repeated several times until a dense pellet was successfully made. The cylindrical disks were further sintered at

1150 °C for five hours in air at a heating and cooling rate of 5 °C/min. It was also found that using a higher sintering temperature, for example, at 1250 °C for five hours for the composition with x

= 0.5, the pellet is likely to experience a partial meltdown.

3.2 Symmetrical cell fabrication

The symmetrical cell with composite electrode of either Y3-xCaxFe5O12-δ: LSGM

(La0.8Sr0.2Ga0.8Mg0.2O3-δ, Fuel Cell Materials) = 6:4 wt. or Y3-xCaxFe5O12-δ: YSZ (Zr0.92Y0.08O2-δ,

TOSOH) = 6:4 wt. (x = 0.1, 0.3 and 0.5) was fabricated by screen-printing a mixed slurry of garnet powder, LSGM/YSZ powder and an organic binder on each side of a dense LSGM/YSZ electrolyte pellet. The LSGM pellet was obtained by die-pressing LSGM powder under a pressure of 10 kN for 5-10 minutes and sintered at 1450 °C for ten hours in air, while the YSZ pellet was purchased from Fuel Cell Materials. The screen-printed cell was sintered at 1100 °C for two hours in air and brush-painted with gold paste on each side to act as current collectors. The gold paste was cured at 800 °C for 30 minutes in air.

A picture of symmetrical cell with the garnet-LSGM composite electrode on an LSGM electrolyte is shown as an example in Figure 3.1.

Figure 3.1 Optical image of a garnet-LSGM/LSGM symmetrical cell.

3.3 Characterization methods

3.3.1 Powder X-ray diffraction (PXRD)

PXRD was used to characterize unit cell dimensions, bond angles and phase purity of the polycrystalline powders. It is a powerful technique that relies on the resolution of X-ray diffraction pattern of crystalline (and/or amorphous) materials for the investigation of phase, lattice structure

28 and crystallinity. One common laboratory X-ray source is the X-ray tube, or so called “hot cathode tube”76. It involves a piece of positively charged metal target, usually made of Cu or Mo, placed on one side, and a heated filament on the other side (with cooling system)76. Through thermionic effect, negatively charged electrons are generated from the hot filament, which are then accelerated by strong electric field (thousands of volts) in the tube and eventually hit metal target to produce a group of certain wavelengths of X-rays76. In the case of Cu target, a Ni foil is often employed as the filter to create an almost monochromatic incident beam, except for the separation of Cu Kα1 (λ

76 = 1.5405 Å) and Kα2 (λ = 1.5443 Å), whose difference is less than 1% . The X-ray beam is then focused and directed towards the powder sample, and scanned at different angles with the help of a goniometer76. Constructive interference occurs when Bragg’s law is satisfied (Figure 3.2):

푛휆 = 2푑푠𝑖푛휃 (3.1) where n is the order of diffraction, λ is the wavelength of the incident beam, d is the lattice spacing

(d-spacing) and θ is the angle of incidence76. After all designated angles are covered, the diffraction pattern is obtained and subject to comparison with standard reference patterns from either Crystal

Open Database (COD) or Inorganic Crystal Structure Database (ICSD). In addition, quantitative analysis can be made through Rietveld refinement of the PXRD pattern.

Figure 3.2 Schematic diagram of Bragg’s law for PXRD. 29

PXRD measurements in the current thesis were carried out using a Bruker D8 Advance

o o (ECO) powder X-ray diffractometers at 2θ = 10-80 , with a step size of 0.02 and 12 second per step at room temperature. The X-ray diffraction data was refined using a conventional Rietveld method using General Structure Analysis System – II (GSAS-II) package77. The background, scale factor, zero, cell parameter, atomic positions, thermal parameters and profile coefficients including

U, V, W, X and Y were refined until convergence.

3.3.2 X-ray absorption spectroscopy (XAS)

Besides diffraction, materials can also absorb X-rays when the energy level of an incident beam is equal to the electron binding energy of a core level for the element of interest, such that core electrons are excited to the continuum, and a sharp increase in X-ray absorbance is recorded, namely the absorption edge78. Depending on difference in electronic structure of the investigated element, the binding energy and corresponding absorption edge vary78. The region of spectrum near the absorption edge (± 10 eV) is called X-ray absorption near edge structure (XANES)79. In addition, the ejected electron (or “photoelectron”) further interacts with the surrounding atoms in the lattice and produce constructive or destructive interference, based on whether the two electromagnetic waves are completely in phase or out of phase78. The relevant region (50 – 1000 eV after the absorption edge) is named extended X-ray absorption fine structure (EXAFS), which provides useful information on local chemical environment of the studied element78. Electrons from higher energy shells are also likely to fill in the “holes” created by the ejection of core electrons, which will then release energy in the form of photons in order for relaxation to happen78.

XAS was employed to investigate the electronic structure at oxygen and iron sites for as- prepared Y3-xCaxFe5O12-δ (x = 0, 0.1 and 0.5) garnet samples. The XANES measurements were carried out on the Spherical Grating Monochromator (SGM) beamline 11ID-1 at Canadian Light

Source in Saskatoon80. Powder samples were pressed into double-sided carbon tape and attached to a brass sample plate, which was then put in the X-ray absorption chamber under 5 x 10-6 Torr.

The O K-edge and Fe L-edge XANES spectra were recorded in both partial fluorescence yield

(PFY) using a four-energy resolved Amptek Silicon Drift Detectors and in total electron yield

(TEY) using the drain current from the sample. For each sample 15 spectra were recorded from a new spot on sample using slew scanning for 60 seconds and averaged81. Spectra were normalized using drain current from an Au mesh before the sample81,82.

3.3.3 Iodometric titration

Iodometry is an analytical technique and can be applied to the determination of cation average oxidation state in metal oxides. In the present thesis, the average oxidation state of Fe (and the corresponding oxygen non-stoichiometry value, calculated based on charge neutrality) for as- prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3, 0.5) powders at room temperature were analyzed using iodometric titration83–86. Ca. 100 mg of the garnet powder was added to 10 ml hydrochloric acid (37%, Sigma-Aldrich) and boiled until it dissolved. The solution was cooled down and diluted to 100 ml using distilled water in a volumetric flask. 15 ml of the analyte was taken from the volumetric flask each time to be titrated against a pre-standardized Na2S2O3 solution in a flowing

N2 atmosphere. 10 ml of 6.6 (wt.%) KI (>99%, Sigma-Aldrich) solution was added to the analyte in advance to reduce any Fe3+ or Fe4+ into Fe2+. 0.2 ml of freshly prepared 1 (wt.%) starch (soluble, for iodometry, Alfa Aesar) solution was used as the indicator and was added just before the end-

31 point where the colour of solution was pale yellow. The titration of each sample was performed at least 3 times until consistent results were achieved. The reactions occurred in the titration were

- 2- oxidation of I to I2 by the higher-valent Fe ions, and the reduction of generated I2 by the S2O3 in titrant, as shown below:

푥−2 퐹푒푥+ + (푥 − 2)퐼− → 퐹푒2+ + 퐼 (3.2) 2 2

2− − 2− 퐼2 + 2푆2푂3 → 2퐼 + 푆4푂6 (3.3)

The Na2S2O3 solution was standardized in advance using potassium dichromate. Ca. 12 g of sodium thiosulfate anhydrous (99%, Alfa Aesar) and 0.15 g of sodium carbonate anhydrous

(ACS grade, Anachemia) was dissolved in 750 ml distilled water and boiled on the hot plate for

15 mins. The solution was subject to settling at room temperature for a week to obtain a stabilized concentration of Na2S2O3. Ca. 0.2 g of potassium dichromate (K2Cr2O7, ACS grade, Fisher

Scientific) was accurately weighed and dissolved in H2SO4 solution and mixed with an excess of

KI solution. The mixing process was performed in a dark container. The container was sealed tight right afterwards and was allowed to sit for 10 minutes, in order for the reaction to occur while avoiding decomposition of the generated I2. 25 ml of the mixture was taken each time to be titrated against the Na2S2O3 solution, using 0.2 ml 1% starch solution as the indicator.

3.3.4 Thermogravimetric analysis (TGA)

TGA is used to monitor the weight change of a sample as a function of temperature under designated atmospheric condition. It is a versatile thermal analysis tool and provides information on the phase transition, redox reaction, material decomposition and gas uptake/release of the chemical during single or multiple heating and cooling (and isothermal) cycles. The atmosphere can be controlled using pure gas, such as N2, O2 and H2, or a combination of several gases.

The garnet sample with the highest oxygen non-stoichiometry (determined by iodometric titration) was studied by TGA for thermogravimetric hydrogen reduction87,88. Ca. 20 mg powder was placed on a microbalance under 2% H2 in N2 environment in a thermogravimetric analyzer

(TGA, Mettler Toledo TGA/DSC/HT1600). The powder underwent reduction reaction and was converted into metals and/or metal oxides, whose phases were later confirmed using PXRD and

Rietveld refinement. The weight change during the reduction reaction was recorded and used for the calculation of oxygen non-stoichiometry of the original sample.

The oxygen content variation of the garnet samples in air at elevated temperature were also determined by TGA under a flowing air environment89.

3.3.5 Density measurement

Archimedes method is used to measure the experimental density of sample pellet75,90. It is based on the principle that the buoyant force exerted on a submerged object equals to the weight of liquid it displaces. During the measurement, the pellet is weighed in dry condition first (denoted as Mdry), and then suspended in a liquid for measurement again (denoted as Msus). Afterwards, it is taken out on balance for a third measurement, where the weight of pellet with saturated liquid in it (denoted as Msat) is recorded. Upon knowing the density of liquid (denoted as ρsolv), the pellet experimental density can be expressed as:

휌푠표푙푣푀푑푟푦 휌푒푥푝푒푟𝑖푚푒푛푡푎푙 = (3.4) 푀푠푎푡−푀푠푢푠

Mettler Toledo Density Kit was used for the experimental density determination, whereas the theoretical density of the pellet was obtained from Rietveld refinement. The relative density was calculated based on the equation:

휌푒푥푝푒푟𝑖푚푒푛푡푎푙 휌푟푒푙푎푡𝑖푣푒 = × 100% (3.5) 휌푡ℎ푒표푟푒푡𝑖푐푎푙

3.3.6 4-Probe DC measurement

4-probe DC method is a commonly used technique for accurate measurement of electrical impedance. It has the advantages of eliminating contact and lead resistance from the measured result, which is critical for the evaluation of object with very low resistance value. The method employs two separate pairs of current and voltage probes. During the measurement, a direct current is passed through the pair of current probes and creates a voltage drop across the tested sample, which is detected by the pair of voltage probes. In the case of probe-configuration shown in Figure

3.3, the conductivity is calculated as:

4퐿 휎 = (3.6) 푅퐷2휋 where L is the length between two voltage probes, D is the diameter of cylindrical pellet and R is the resistance of the pellet between two voltage probes. Depending on the applied current (A) and detected voltage (V), R is determined as:

푉 푅 = (3.7) 퐴

The electrical conductivity of sintered Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3, 0.5) garnet pellet was measured using 4-probe DC method with a BioLogic potentiostat (Model VSP-300,

Seyssinet-Pariset, France) from room temperature to 900 °C in air. All pellets were cut into two grooves along the rod and brush-painted with gold paste as current collectors for voltage-probes.

The top and bottom surfaces were gold-pasted and attached with two current-probes. A constant current in the range of 0.01-50 mA was applied to the current-probes, and the corresponding voltage response was recorded. Five constant currents and corresponding voltages were obtained for each measurement to collectively calculate resistance through linear regression. Conductivity is then determined based on equation 3.6.

Figure 3.3 Schematic illustration of arrangement for 4-probe DC method.

3.3.7 Hebb-Wagner polarization method

Hebb-Wagner DC polarization is a useful method to separate the ionic conductivity (σi)

91–94 from electronic conductivity (σe) (or vice versa) in MIECs . By using this technique, the ionic transport number (ti) of the material (Equation 3.8) can be determined, which functions as an important indicator for its potential usage and performance. For example, if the studied compound is intended to be employed as an oxide-ion conducting electrolyte for SOFC, it is expected that its ionic transport number is close to 1, such that there would be no electronic leakage during the device operation. On the other hand, if the material is to be used as an SOFC electrode, a balance of possessing both ionic and electronic conductivity is preferred. Common SOFC cathodes of

- La0.6Sr0.4CoO3 and La0.6Sr0.4Co0.2Fe0.8O3 are reported to have ionic transport numbers of 1.37×10

4 and 3.62 ×10-5 respectively at 800 °C in air69.

휎𝑖 푡𝑖 = (3.8) 휎𝑖+휎푒

There are variations in arrangements of Hebb-Wagner polarization method in the literature.

In the case of measuring ionic conductivity (and σi is expected to be relatively smaller than σe), one configuration is to have the MIEC of interest sandwiched between the two different electrodes, with one chosen to be electronic-blocking and the other electronic-conducting. An DC-bias is applied to this system and the current response as a function of time is recorded. The driving force for all charge carriers within the system, at the very beginning, is the applied external DC-potential.

However, due to the presence of electronic-blocking electrode, where only ions are able to pass through, the received current will decay as there is an internal electrical field building up inside

MIEC, since ionic current is lower than electronic current and the extra electrons will accumulate at the interface of MIEC and electronic-blocking electrode. Eventually, an equilibrium of electrical field is achieved (electrostatic potential of internal = external) inside MIEC, such that the driving force of ion migration is merely diffusion. The final equilibrium current measured is used to calculate the ionic conductivity of MIEC.

The ionic (oxide-ion) conductivity of garnet-type Y2.9Ca0.1Fe5O12-δ pellet was determined by Hebb-Wagner polarization method (Figure 3.4). A dense YSZ pellet was made through pressing commercial (TOSOH) YSZ powder under pressure of 200 kN for 5-10 minutes in rubber mould and sintering at 1400 oC for 10 hours in air. The obtained pellet was cut into thin slices by a diamond-saw. One slice of YSZ pellet was then attached to a thin garnet pellet (both ~ 1 mm in thickness) using gold-paste in between and cured at 800 °C for 30 minutes. The surroundings were covered with alumina seal (Ceramabond 552, Aremco) in order to prevent any oxygen leakage and potential gas-phase reaction. The top and bottom surfaces were also pasted with gold. The DC polarization measurements were conducted between 550 – 900 °C, with an interval of 50 °C, under various constant biases in air. Each polarization was applied for at least one hour to obtain a

36 stabilized current. A 30 minutes open-circuit potential (OCP) period was placed in between the two consecutive polarizations for the system to restore the initial state of charge distribution.

Figure 3.4 Illustration of Hebb-Wagner polarization method used for ionic conductivity measurement of Y2.9Ca0.1Fe5O12-δ.

3.3.8 Scanning electron microscopy (SEM)

SEM is a powerful imaging tool used heavily in many fields, including solid-state chemistry. It relies on scanning sample surface with a focused beam of high-energy electrons and analyzing the received signals from the electron-sample interaction to provide information on the morphology of the studied material95. Various signals can be generated when electrons interact with sample, such as Auger electrons, secondary electrons, backscattered electrons, characteristic

X-rays and fluorescence X-rays95. Among them, secondary electrons and backscattered electrons are two of the most commonly used imaging sources95. The former has a strength in reflecting morphology and topology of the sample, whereas the latter is competent of drawing contrast in a multi-phase specimen95. By coupling with an energy-dispersive detector, the emitted X-rays from

37 different elements of a sample can be separated, and the results can be used for identification of chemical composition in a certain selected area, both qualitatively and quantitively95.

The microstructure analysis of sintered Y3-xCaxFe5O12-δ pellets and the cross-sectional area of garnet-LSGM / LSGM symmetrical cells after electrochemical measurements was performed using a Zeiss Sigma VP Series scanning electron microscope at University of Calgary.

3.3.9 Electrochemical impedance spectroscopy (EIS)

EIS can be considered as a testing or diagnostic tool for the investigation of processes or mechanisms within an unknown electrochemical system. The system tested is sometimes referred as a “black box”. A small sinusoidal AC perturbation, usually 10-100 mV in amplitude and 10-2-

107 Hz in frequency range, is applied as the input for the examination of this black box. The output of the current response is analyzed to obtain information on system properties through reasonable hypothesis and fitting experimental data with appropriate models.

Given that the probing signal is a sinusoidal wave (potential), it is written as:

푉(푡) = 푉푚sin (2휋푓푡) (3.9) where V(t) is the instant potential at a given time t, Vm is the peak potential, f is the frequency, and t is the time96. Accordingly, the received current response is expressed as:

퐼(푡) = 퐼푚sin (2휋푓푡 + 휃) (3.10) where I(t) is the instant current at the given time t, Im is the peak current and θ is the phase shift of the present sinusoidal wave compared to the original probing signal in equation 3.996.

The impedance (Z) of the tested electrochemical system, therefore, is calculated as (Z0 is

96 defined as Vm/Im) :

푉(푡) 푉푚sin (2휋푓푡) sin (2휋푓푡) 푍(푡) = = = 푍0 (3.11) 퐼 (푡) 퐼푚sin (2휋푓푡+휃) sin (2휋푓푡+휃)

It is then acknowledged that, with a fixed frequency f, the impedance Z can be represented

96 in terms of a magnitude parameter Z0 and a phase shift θ, based on equation 3.11 .

In an experimental setting, the frequency is usually swept in a large range, in order to obtain a set of impedance results, represented by a set of Z0 and θ, as a function of f. There are a few ways of presenting the experimental data. The phase shift θ can be plotted against frequency f, which is called a Bode plot97. In addition, by using phasor representation and considering impedance data as vectors of Z0 in length and θ in angle under a complex plane, a Nyquist plot can be drawn with

97 the real part of Z0 (Z’) as the x-axis and imaginary part of Z0 (Z’’) as the y axis .

In order to gain understanding of the presented data, equivalent circuit model (ECM) is a conventional method for the analysis of these impedance results. It draws direct analog between the behavior of the probed system and an idealized electrical circuit consisting of one or multiple electrical components97. With careful selection of electrical components and proper fitting of the component parameters, it is claimed that the properties shown by the fitted electrical circuit are able to reflect the real chemical and physical properties of the studied electrochemical system97.

However, there are limitations to the ECM technique, despite of it being well-established over the past decades97. One of the main concerns is that the choice of equivalent circuit for the experimental data can seldom be unique, which leads to unambiguity of resolution and potential errors due to improper fitting97. The problem may rise due to the indiscrimination of all the selected electrical components during fitting process, and can be alleviated to some extent by switching to a novel EIS analysis protocol, named “impedance spectroscopy genetic programming (ISGP)”.

The EIS was performed in the current thesis on a BioLogic potentiostat (Model VSP-300,

Seyssinet-Pariset, France) or a Solartron AC impedance analyzer (Model 1260). An AC potential of 100 mV and 0.1-106 Hz was applied at open-circuit condition in a 2-probe configuration. A

39 minimum of one hour of waiting time was established to ensure the system reaches thermal and atmospheric equilibria during the measurements.

3.3.10 Impedance spectroscopy genetic programming (ISGP)

ISGP is a computer software that utilizes genetic algorithm to find the distribution function of relaxation times (DFRT)98,99. It applies discrepancy-complexity approach to avoid overfitting, and does not use filter or Lagrange coefficients for regularization, which further helps to avoid the appearance of artificial peaks or too smooth solutions 98,99. Electrochemical processes are easily deconvoluted in this method, as it finds an analytic function that follows the behaviour of each peak separately. The correlation between the measured impedance and DFRT is shown as98,99:

∞ 훤(푙표푔(휏)) 푍(휔) = 푅 + 푅 ∫ 푑(푙표푔( 휏)) (3.12) ∞ 푝표푙 −∞ 1+𝑖휔휏 where Z is the impedance, R∞ is the series resistance, Rpol is the polarization resistance, Γ is the

DFRT, τ is the relaxation time, and ω is the angular frequency.

Advanced numerical tools such as MATLB were used to run ISGP, since the nature of this ill-posed inverse problem hinders the direct extraction of DFRT from measured impedance data in equation 3.1298,99. Each peak in the DFRT vs. log (τ) or frequency plot is characterised by its relaxation time, height and width. By multiplying Rpol with the area under the peak, the polarisation resistance of individual process is determined. Rs (or R∞) was subtracted from the real part of the impedance data, since the Rpol for the electrochemical process was smaller than Rs and would be obscured by relatively large value of Rs. Thus, two iterations of ISGP were performed on AC impedance data, where in the first iteration Rs was calculated and then subtracted for the analysis of the second ISGP iteration.

3.4 Error considerations

• The starting chemicals were used as-purchased for material synthesis, without additional

purification. Small amounts of impurity elements presented in the precursors would be

directly transferred into the synthesized product.

• The dimension of sintered pellet was not perfectly cylindrical and the shape of gold current

collector was not perfectly circular. Multiple measurements of diameter and thickness were

performed using a digital caliper, and the average of respective readings were recorded.

• The relative density of sintered pellet was below the theoretical value of 100%, which could

lead to potential inaccuracy of electrical and ionic conductivity measurement.

• The X-ray diffractometer has a detection limit ~ 5%, which could lead to small amounts of

secondary phases being missed in the PXRD analysis.

• For TGA measurement, Mettler Toledo TGA/DSC1/1600HT thermogravimetric analyzer

is reported to have an error of ± 0.5 °C at the controlled temperature and an error of ± 1 μg

for the measured weight.

• The Alicat Scientific mass flow controller used to create designated atmospheric condition

has an error of ± 0.2% under full scale gas flow (100 sccm).

• A Thermolyne F21100 tubular furnace was employed for electrical and electrochemical

characterizations at elevated temperature. A thermocouple (Omega Engineering Inc.) was

used for temperature readings and has an error of ± 1 °C.

• The author takes responsibility for any oversight in data interpretation.

Chapter Four: Synthesis, Rietveld Refinement of Crystal Structure, Oxygen Non-

Stoichiometry, Electronic Structure and Electrical Transport Properties of Garnet-

Type Y3-xCaxFe5O12-δ

4.1 Introduction

Electrochemical devices play an important role in the sustainable energy production and storage with minimum environmental impacts100,101. Specifically, mixed ionic-electronic conductors (MIECs), which can function as SOFC electrodes and membranes for the use of oxygen separation and hydrocarbon conversion, are of vital significance in the field102–109.

Ceramics based on perovskite structures exhibit high performance in these electrochemical devices, owing to their high mixed ionic-and-electronic conductivities and high permeability of oxygen103,107,110. However, many groups of the perovskite-type oxides are found to be structurally unstable under large pO2 gradients, and reactive in the presence of CO2, H2O and SO2, as well as possessing high TECs which can lead to mechanical instability in an integrated system7,19,111–115.

As a result, alternative MIECs have been investigated by researchers, including materials of the double perovskite, fluorite, Ruddlesden-popper and Brownmillerite – type structures.

YIG is a well-known garnet-type metal oxide with interesting ferrimagnetic properties and has applications in many areas, such as microwave devices, bio-antennas, optical oscillators, circulators and nanodevice phase shifters116–122. Electrical conduction and defect property relations have been studied in polycrystalline and single crystals of YIG123–128. Similar to perovskites, oxide

.. ion conductivity occurs in YIG through 푉푂, in addition to its polaron-based electronic conductivity, both of which can be enhanced by adopting an acceptor-doping route32,128,129. Moreover, YIG is known to possess a low TEC, and is expected to be stable under CO2 and H2O since it does not

42 contain any alkali or alkaline earth elements126. Therefore, it would be beneficial to explore the potential of YIG as a promising MIEC in the field of electrochemical energy conversion.

Kharton et. al. studied the partial substitution of Y-site in YIG by acceptor-doping of Ca with a unit cell formula of Y2.5Ca0.5Fe5O12 (YCFO5), which exhibited an oxygen permeability of

7.9×10-11 mol/(s cm2) and an electrical conductivity of 2.1 S/cm at 900 °C32. Subsequently, Zhong et al. employed a composite of YCFO5 and SDC as novel SOFC cathode material, and obtained an ASR of 0.55 Ωcm2 at 650 ºC in air31. However, the current literatures have only focused on the investigation of YCFO5 and a systematic study regarding the effect of Ca-doping in YIG on the structural and electrical properties is absent.

A better understanding of the phase, oxygen non-stoichiometry, electronic structure and electrical transport properties of Ca-doped YIG (Y3-xCaxFe5O12-δ) is believed to not only help later optimizations of its performance as an SOFC cathode, but also extend applications beyond SOFC for other electrochemical devices, once its MIEC features are known.

In the current chapter, PXRD and Rietveld refinements were used for phase analysis on as- prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3, 0.5 and 0.7). Oxygen non-stoichiometry of the bulk sample powder was quantified by iodometric titration and TGA, whereas electronic structure at O and Fe sites was characterized by XAS on the sample surface. 4-Probe DC measurements were carried out for the electrical conductivity at elevated temperatures. The ionic conductivity of the composition x = 0.1 was separated from its electrical conductivity using Hebb-Wagner polarization method. Based on the defect model and experimental results of electrical conductivity variation in varying pO2 atmosphere, the conduction mechanism in Ca-doped YIG was discussed.

4.2 Results and discussion

4.2.1 Phase analysis

The PXRD patterns for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3, 0.5 and 0.7) powders at room temperature are shown in Figure 4.1. Compositions with x = 0-0.3 were single- phase and were indexed with cubic symmetry (Ia-3d, space group # 230). A further increase of Ca content to x = 0.5 and 0.7 led to the segregation of CaFe2O4 (Pnma; #62) and Ca2Fe2O5 (Pnma;

#62) phases. Selected area of diffraction patterns between 32.0 and 32.8 two thetas shows the slight shifting of peaks to higher degrees for x = 0, 0.05 and 0.1, and subsequent back-shifting towards the original two theta peak position (compared to x = 0) for x = 0.3, 0.5 and 0.7.

Rietveld refinements (Figure 4.2) were performed on all PXRD patterns except for x = 0.7, where impurity phases are significant and may contribute to a large degree of error in the actual composition of the garnet phase. The values of wRp in all refinements were below 10%, indicating good agreement of experimental data with the calculations. The detailed structural parameters from refinements were summarized in Table 4.1. The illustration of crystal structure of Ca-doped YIG is shown in Figure 4.3 and the variation of bond characteristics determined from refinements are shown in Figure 4.4.

Looking into the results obtained from Rietveld refinements shown in Table 4.1, it can be seen that the lattice constants and volumes do not show any significant variations due to Ca doping.

In theory, since Ca2+ (1.00 Å) has a larger ionic radius than Y3+ (0.90 Å) at dodecahedral site in the garnet unit cell, its substitution should have led to an increase in the metal-oxygen bond length and therefore an increase in the lattice constant (or shifting of peaks towards lower two thetas).

However, this supposed trend was not observed in our results, and we suspect it is because the combination of the following factors, which also contributes to the lattice changes in addition to

44 the Ca substitution: i) The tilting of metal-oxygen polyhedra in the unit cell into other directions, which would affect the lattice spacing32. ii) The aliovalent doping of Ca could result in two possible charge compensation mechanisms, listed using Kröger-Vink notation as130:

3 ( ) 3퐶푎푂 + 푂 푌3퐹푒5푂12 3퐶푎′ + 3퐹푒∙ + 2퐹푒× + 12푂× (4.1) 2 2 → 푌 퐹푒 퐹푒 푂

( ) 3 21 3퐶푎푂 푌3퐹푒5푂12 3퐶푎′ + 5퐹푒× + 푉∙∙ + 푂× (4.2) → 푌 퐹푒 2 푂 2 푂

′ 2+ 3+ . 4+ where, 퐶푎푌 indicates Ca in the Y site with one effective negative charge, 퐹푒퐹푒 indicates Fe

3+ 푥 3+ 3+ in the Fe site with one effective positive charge, 퐹푒퐹푒 indicates Fe in the Fe site with effective

.. 푥 neutral charge, 푉푂 indicates oxide ion vacancy with two effective positive charges, and 푂푂 indicates O2- in the O2- site with effective neutral charge.

The changes due to either or both of these charge compensation mechanisms would further

. impact the lattice spacing, and in-turn the peak position. In the case of electron-hole (퐹푒퐹푒 ) formation (Equation 4.1), as Fe4+ (0.585 Å) is smaller than Fe3+ (0.645 Å), the appearance of a certain concentration of Fe4+ could result in a smaller Fe-site mean ionic radius131–134. Similarly,

∙∙ with the formation of oxide ion vacancies ( 푉푂 ) (Equation 4.2), the lattice constant would experience either reduction or increment depending on the circumstances135.

In all, the subtle variations in the lattice parameters and peak positions are inferred to be the collective effects from all the factors stated above, which consequently results in a non- monotonic change of unit cell with respect to Ca-doping level in the present report.

O Y/Ca Fe

Figure 4.3 Illustration of crystal structure of garnet-type Y3-xCaxFe5O12-δ.

Figure 4.4 Variation of bond characteristics of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) determined by PXRD Rietveld refinement.

Table 4.1 Structural parameters of Y3-xCaxFe5O12-δ obtained by Rietveld refinement*.

x = 0 x = 0.05 x = 0.1 x = 0.3 x = 0.5 Unit cell a (Å) 12.3782(2) 12.3802(6) 12.3786(8) 12.379(35) 12.3784(6) V (Å3) 1896.59(9) 1897.53(4) 1896.80(8) 1897.1(19) 1896.70(7) Y/Ca (24c) x 0.125 0.125 0.125 0.125 0.125 y 0 0 0 0 0 z 0.25 0.25 0.25 0.25 0.25 Fe (16a) x 0 0 0 0 0 y 0 0 0 0 0 z 0 0 0 0 0 Fe (24d) x 0.375 0.375 0.375 0.375 0.375 y 0 0 0 0 0 z 0.25 0.25 0.25 0.25 0.25 O (96h) x 0.151418 0.15053 0.151014 0.150208 0.151942 y -0.028075 -0.02919 -0.028142 -0.027963 -0.026264 z 0.055811 0.05673 0.055109 0.056817 0.057565 Occupancy Y 1 0.983 0.967 0.9 0.833 Ca 0 0.017 0.033 0.1 0.167 Fe (16a) 1 1 1 1 1 Fe (24d) 1 1 1 1 1 O 1 1 1 1 1

Uiso Y 0.0316 0.0285 0.0307 0.0181 0.0284 Ca - 0.0199 0.0136 0.0339 0.024 Fe (16a) 0.0244 0.0261 0.0219 0.0241 0.0408 Fe (24d) 0.0146 0.0241 0.011 0.025 0.0432 O 0.0265 0.0327 0.0239 0.0343 0.0584 Bond length Fe(16a)-O 2.02759 2.0242 2.02015 2.018 2.03736 Fe(24d)-O 1.84541 1.84842 1.84513 1.86116 1.86417 Y/Ca-O 2.45065 2.44042 2.45867 2.43643 2.42708 Bond angle Fe(16a)-O-Fe(24d) 126.5231 126.591 126.9927 126.1988 124.8893 wRp (%) 7.379 6.569 7.123 9.257 7.178 * Background, scale factor, zero, cell parameter, atomic positions, thermal parameters and profile coefficients of U, V, W, X and Y were refined.

4.2.2 Investigation of oxygen non-stoichiometry and electronic structure

4.2.2.1 Oxygen non-stoichiometry

The oxygen non-stoichiometry (δ), calculated based on the iron average oxidation states by iodometric titration is presented in Figure 4.5. It can be seen that, all Ca-doped samples are oxygen-deficient at room temperature, with δ ranging from 0.06 (x = 0.05) to 0.19 (x = 0.3). From

Figure 4.5, it can also be seen that, with the increase of Ca substitution level (x), δ also increases, except for the end member x = 0.5, where two impurity phases are found through PXRD and may be the reason to break this trend. The presence of oxygen deficiency in all Ca-doped samples indicates the existence of charge compensation mechanism proposed in equation 4.2 (formation of oxide ion vacancies).

Moreover, the most oxygen-deficient sample composition, x = 0.3, was further studied by thermogravimetric hydrogen reduction. In this method, as-prepared powder undergoes reduction in 2% H2 atmosphere at elevated temperature, described as:

15.3−2훿 2.7 15.3−2훿 푌 퐶푎 퐹푒 푂 + ( ) 퐻 (푔) → 푌 푂 + 0.3퐶푎푂 + 5퐹푒 + ( ) 퐻 푂(푔) (4.3) 2.7 0.3 5 12−훿 2 2 2 2 3 2 2 where the powder Y2.7Ca0.3Fe5O12-δ is reduced by H2 into Y2O3, CaO and Fe, and generates water

(H2O) as the by-product.

Figure 4.6 shows the experimental curve, reflecting weight change of the powder during heating, isothermal and cooling segments. It can be seen that the reduction reaction starts around

650 ºC and is completed shortly after temperature reaches 800 ºC, where the curve becomes flat and the total weight remains constant. The oxygen non-stoichiometry (δ) was calculated, based on the weight change before and after the reaction according to equation 4.3. The obtained result (δ

= 0.18) is listed in Figure 4.6 in comparison with the value (δ = 0.19) obtained by iodometric

52 titration method. Since there is only small difference between the values from two methods, it can be concluded that the two methods tell the same story and both the results can be trusted.

Besides, the proposed decomposed phases (Y2O3 (Ia-3); CaO (Fm-3m); Fe (Im-3m)) were confirmed by PXRD and Rietveld refinement, as shown in Figure 4.7 and Table 4.2. An wRp value of 8.231% was achieved, and no other extra peaks were observed, proving that equation 4.3 describes the actual reduction reaction occurring to the tested sample.

The oxygen non-stoichiometry at elevated temperatures was further investigated by TGA under flowing air (Figure 4.8). Where it can be seen that all compositions experience a minimal amount of oxygen gain/loss (less than 0.1 % in overall weight) from room temperature up to 1000

°C. A slight dip between 300-500 °C is observed and may be ascribed to the Fe valance change due to reduction and loss of oxygen136. In the present case, the normally (relative) significant mass loss associated with the oxygen release from lattice and facile reduction of Fe is not seen, which is different from those TGA results of typical SOFC cathode materials, such as BSCF137,138.

Figure 4.5 Evaluation of oxygen non-stoichiometry (δ) for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) sample powder by iodometric titration.

Figure 4.6 Experimental TGA curve for Y2.7Ca0.3Fe5O12-δ powder under 2%H2-98%N2 atmosphere in order to achieve complete reduction at 800 °C.

Figure 4.7 PXRD Rietveld refinement of reaction products in thermogravimetric hydrogen reduction method.

Figure 4.8 Weight change and oxygen content variation in air as a function of temperature between 30-1000 °C for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) powders.

Table 4.2 Structural parameters of reaction products after reduction of Y2.7Ca0.3Fe5O12-δ under 2%H2-98%N2 atmosphere in TGA obtained by Rietveld refinement*.

Y2O3 CaO Fe Space group Ia-3 Space group Fm-3m Space group Im-3m Unit cell Unit cell Unit cell a (Å) 10.606(4) a (Å) 4.8131(7) a (Å) 2.8686(6) V (Å3) 1193.1(7) V (Å3) 111.50(5) V (Å3) 23.60(7) Y (24d) Ca (4b) Fe (2a) x 0 x 0 x 0 y 0.25 y 0 y 0 z 0.28209 z 0.5 z 0 Y (8a) O (4a) Occupancy x 0 x 0 Fe 1

y 0 y 0 Uiso z 0 z 0 Fe 0.0219 O (48e) Occupancy x 0.09619 Ca 1 y 0.14091 O 1

z 0.62839 Uiso Occupancy Ca 0.0196 Y (24d) 1 O 0.0391 Y (8a) 1 O 1

Uiso Y (24d) 0.0207 Y (8a) 0.024 O 0.0202 wRp (%) 8.231 * Background, scale factor, zero, cell parameter, atomic positions, thermal parameters and profile coefficients of U, V, W, X and Y were refined.

4.2.2.2 Electronic structure at O and Fe sites

The electronic structures at O and Fe sites were studied by XAS. The Fe L-edge spectra

(Figure 4.9) show that compositions with x = 0 and x = 0.5 have similar Fe L-edge spectra, which is an indication of a similar iron oxidation state. Whereas the extra spectroscopic features in the composition with x = 0.1, labelled by arrow in Figure 4.9, shows a higher iron oxidation state than

56 x = 0 and x = 0.5139–141. While it might be imprudent to make an interpretation on the result of x =

0.5, given the impurities present, the increase in the iron oxidation state between x = 0 and x = 0.1 suggests the formation of a small concentration of Fe4+ with the increase of Ca doping level through charge balance reaction (Equation 4.1). The O K-edge spectra shown in Figure 4.10, on the other hand, reveal the direct presence of the electron-hole state, with its feature marked by arrow, for compositions with x = 0.1 and x = 0.5142,143. This result supports the formation of Fe4+

· seen in Fe L-edge spectrum for x = 0.1, since the electron-hole is nothing but Fe Fe. Furthermore, this observation validates the charge compensation mechanism proposed in Equation 4.1

(formation of electron-hole) for the Ca-doped samples.

Overall, by combining the analysis of data obtained from several methods discussed above, it may be concluded that both the formation of oxide ion vacancy (Equation 4.2) and the electron- hole (Equation 4.1) are true and co-exist when the aliovalent doping of Ca is introduced into the garnet lattice. It is also worth noticing that, given the detection limit of XAS in the present study to be ca. 100 nm on the powder surface, where the environment is almost identical to instrumental settings. While iodometry is a characterization technique that is performed in ambient air to the bulk samples. Possible deviations between the results of the two methods could occur, and future research should be planned to address these potential sources of error.

Figure 4.9 Fe L-edge XANES spectra for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.1 and 0.5) powders.

Figure 4.10 O K-edge XANES spectra for as-prepared Y3-xCaxFe5O12-δ (x = 0, 0.1 and 0.5) powders.

4.2.3 Characterization of electrical properties

4.2.3.1 Electrical conductivity and conduction mechanism

The electrical conductivity (σ) of sintered Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) pellets was measured by 4-probe DC method in air. Figure 4.11 summarizes the conductivity as a function of temperature for the different Ca-doped samples, with the bottom-left inset showing result for the parent un-doped phase and the upper right inset reflecting linear regression lines in the Arrhenius plot of conductivity at 500-900 °C. In general, electrical conductivity increases with increasing temperature for all the compositions. Through a closer look over various compositions at the same temperature, shown as an example at 750 °C in Figure 4.12, it is found that with increase of Ca doping level, the conductivity experiences a sharp increase from the parent phase x

= 0 (3.09×10-3 S/cm at 750 °C) to x = 0.1 (1.58 S/cm at 750 °C). And after reaching the maximum for x = 0.1, the electrical conductivity then slowly decreases to 1.26 S/cm at 750 °C for x = 0.5.

The sharp increase in the electrical conductivity between x = 0 and x = 0.1 can be attributed to an

. increase in the concentration of p-type charge carriers (퐹푒퐹푒) due to increasing Ca substitution level (Equation 4.1). Whereas the slow decrease between x = 0.1 and x = 0.5 of the electrical

′ . conductivity is speculated because of the formation of defect associates [퐶푎푌 − 퐹푒퐹푒], which can lead to a lowered mobility of p-type electronic charge carriers. Additionally, for x = 0.5, the extra

Ca substitution has led to the formation of CaFe2O4 and Ca2Fe2O5, which are also known to show predominantly p-type electronic conductivity, but with values lower than garnet-type oxides144.

The overall low electronic conductivity of the studied garnets, compared to typical MIEC cathodes in SOFC (BSCF: 31.5 S/cm at 750 °C in air; LSCF (La0.8Sr0.2Co0.8Fe0.2O3-δ): 1038 S/cm at 700 °C in air145,146), can be attributed to the absence of linear Fe−O−Fe chains (Figure 4.4 and Table 4.1),

59 which hampers the migration of p-type charge carriers, as normally seen in perovskite-type oxides147.

The inset of Figure 4.12 shows the effect of pO2 on the electrical conductivity for x = 0.1 and x = 0.3 at 750 °C. The slopes of both compositions are very small (~0.02), indicating an almost

128,129 pO2 independent conductivity region at this temperature and pO2 range . Neutrality in the pO2

. ′ independent region might follow the relation: 퐹푒퐹푒 = [퐶푎푌], where pO2 independent p-type

129 -15 conduction mechanism is observed . Increasing the measured pO2 range down into ~ 10 should

1/4 allow the sample to exhibit both the pO2 independent (p-type) and dependent (푝푂2 ) regions, as seen in other garnet phases128,129.

The activation energies (Ea) was calculated from the Arrhenius equation and was obtained from linear regressions in the Arrhenius plot of electrical conductivity. Ea and densities of pellets used in the measurements are listed in Table 4.3. Where it can be seen that with the increase in Ca,

. Ea decreases from 1.08 eV (x = 0) to 0.19-0.24 eV (x = 0.1–0.5), indicating small-polaron (퐹푒퐹푒) conduction mechanism with an induced p-type conductivity, as shown in equation 4.132,126,127,129.

The relative densities of the pellets for the Ca-substituted compositions (x = 0.05-0.5) are between 78 % to 85 %, while the un-doped phase (x = 0) has a much lower density of 54 %. In order to analyze the effect of pellet porosity on the measured conductivity: in the present study, for x = 0.5, the electrical conductivity determined at 750 °C is 1.26 S/cm, with an 84.8 % relative density. Whereas in the literature study performed by V. V. Kharton et al., the electrical conductivity obtained with a 97.7 % dense pellet at 750 °C is 1.07 S/cm32. On comparing these values, it can be suggested that the electrical conductivity values are not much affected by the pellet porosity in the current experimental setting.

Furthermore, from SEM images shown in Figure 4.13, it can be seen that, although pellet surfaces are indeed porous, in agreement with Archimedes measurements, however, the grains are still well-connected to each other, providing continuous conduction pathways for the charge carriers, which explains the reason for the relatively un-influenced electrical conductivities. In addition, the grain size in Figure 4.13 can be seen to increase with increasing Ca substitution level, probably due to the fact that Ca can act as a sintering aid for the garnet pellet.

Figure 4.11 Temperature dependence of electrical conductivity measured by 4-probe DC method for Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) sample pellets in air.

Figure 4.12 Variation of electrical conductivity at 750 °C, with the inset showing p-type conductivity for x = 0.1 and 0.3 in different oxygen partial pressures.

Figure 4.13 Surface SEM images of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) pellets used in the 4-probe DC conductivity measurements.

Table 4.3 Activation energy (Ea) of Y3-xCaxFe5O12-δ (x = 0, 0.05, 0.1, 0.3 and 0.5) determined by 4-probe DC method, and density of pellets used in the measurements.

3 3 Y3-xCaxFe5O12-δ Ea (eV) ρExperimental (g/cm ) ρTheoretical (g/cm ) ρRelative (%)

x = 0 1.08 2.798 5.169 54.13

x = 0.05 0.24 4.311 5.149 83.73

x = 0.10 0.19 4.385 5.134 85.41

x = 0.30 0.20 3.959 5.065 78.17

x = 0.50 0.21 4.236 4.997 84.78

4.2.3.2 Ionic conductivity and ionic transport number

The ionic conductivity of Y2.9Ca0.1Fe5O12-δ, which shows the highest electrical conductivity among the series discussed previously, was determined by Hebb-Wagner polarization method. A typical long-term current response as a function of time under various DC biases is shown in Figure

4.14, where the initial current experiences a gradual decline until it stabilizes to a steady-state value. The steady-state current is considered to be only due to oxide ion conduction since electronic current has been blocked by the ionic-conducting YSZ. The linear portion of the I-V relationship, shown in the inset of Figure 4.14, was used to calculate ionic conductivity148. Therefore, the ionic

-5 conductivity of Y2.9Ca0.1Fe5O12-δ at 750 °C is determined to be 2.89×10 S/cm.

Furthermore, the ionic transport number of Y2.9Ca0.1Fe5O12-δ at 750 °C can be calculated, based on equation 3.8, to be around 1.82×10-5. A comparison of ionic and electrical conductivities

64 at elevated temperatures for this garnet composition is shown in Figure 4.15. Since an ionic transport number in the order of 10-5 can be considered as relatively low, it may be concluded that the studied MIEC material is mainly an electronic conductor.

It should also be noticed that, given the results obtained from Archimedes measurements in Table 4.3 and SEM images in Figure 4.13 indicating certain porosity of the pellets, the condition of an accurate Hebb-Wagner polarization experiment is not strictly fulfilled, since the pellet used is not entirely gas-tight. There could be oxygen molecules moving through the pores and contribute errors to the measured ionic conductivity value. Even though, when we look at the literature, where a dense pellet of 97.7% relative density for Y2.5Ca0.5Fe5O12-δ at 900 °C was determined to have an

-5 -4 ionic transport number of 4.6×10 , our experimental result (3.5×10 for Y2.9Ca0.1Fe5O12-δ at 900

°C from Figure 4.15) is reasonable for the purpose of estimation32.

Figure 4.15 Comparison of ionic conductivity and electrical conductivity as a function of temperature for Y2.9Ca0.1Fe5O12-δ in air.

4.3 Summary

To summarize, single-phase cubic (Ia-3d) powders of Y3-xCaxFe5O12-δ have been obtained for compositions with x = 0, 0.05, 0.1, and 0.3. Increasing x to 0.5 and 0.7 leads to the appearance of secondary phases of CaFe2O4 (Pnma) and Ca2Fe2O5 (Pnma). Oxygen non-stoichiometry and electronic structure of Ca-doped YIG have been investigated through iodometric titration, TGA and XAS. Highest oxygen non-stoichiometry was discovered in x = 0.3 garnet, and p-type charge

. carriers (퐹푒퐹푒) were confirmed by XAS, due to charge compensation of aliovalent Ca-substitution.

Electrical conductivity was found to increase with an increase in Ca content till x = 0.1, and then decrease because of the decrease in concentration of free charge (hole) carriers. Highest electrical

66 conductivity of 1.58 S/cm at 750 °C was shown by the composition with x = 0.1. Hebb-Wagner polarization method was employed for the separation of ionic from electrical conductivity, which yields an ionic conductivity of 2.89 × 10-5 S/cm and an ionic transport number of 1.82 × 10-5 at

750 °C in air for the x = 0.1 garnet composition.

Chapter Five: Electrocatalytic Properties of Oxygen Reduction Reaction (ORR) Studied by

ISGP of Garnet-Type Y3-xCaxFe5O12-δ

5.1 Introduction

SOFC is a highly efficient, quiet and pollution-free energy converter, and has long been one of the major subjects in the research of electrochemical devices18,149–153. Perovskites such as

SrCo0.8Nb0.1Ta0.1O3-δ, (La0.7Sr0.3)0.95Co0.2Fe0.8O3-δ (LSCF), Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF), La1- xSrxCoO3-δ (LSC), NdBa0.5Sr0.5Co2O5+δ, and Ca3Co2O6 have been studied by scientists as IT-SOFC cathodes20,103,107,154–157. Among them, BSCF has shown the lowest area specific resistance (ASR) of 0.03 Ωcm2 for ORR at 650 ºC in air20. Nevertheless, Co-based cathodes suffer from a large TEC mismatch with most of the SOFC electrolytes, and the application of BSCF is hindered to some extent by the decomposition from cubic into hexagonal phase in long-term operation and the

19,113–115,158 carbonate formation when exposed to CO2 .

As discussed in Chapter 4.1, YCFO5 has been investigated in the literature for novel SOFC cathode application, with an achieved ASR of 0.55 Ωcm2 at 650 ºC in air31. In addition, the parent phase YIG is reported to have a TEC value of ~ 10.6 × 10-6 K-1, which is very close to those of conventional SOFC electrolytes (YSZ = 10.5 × 10-6 K-1, SDC = 11.4 × 10-6 K-1)31,159,160. This closeness in TEC values can help avoid delamination issues and ensures the mechanical stability of the device. However, to date, there lacks any complete study on the electrocatalytic properties of ORR of Ca-doped YIG system, except the research work on composition YCFO5. Therefore, it is unclear whether the doping optimization of Ca on SOFC cathode performance has been reached or not. Furthermore, the analysis of impedance spectra in the previous publication was made using equivalent circuit model (ECM), where error could exist when difference between relaxation times of similar electrochemical processes is small.

In the present chapter, electrocatalytic properties of ORR were explored for garnet-type

Y3-xCaxFe5O12-δ (x = 0.1, 0.3 and 0.5). Chemical reactivity between the garnet and LSGM powders was examined. Symmetrical cell with the composite electrode of LSGM and Ca-doped YIG was tested under OCP using EIS in air and various oxygen partial pressures. The obtained impedance spectra were analyzed by ISGP for the separation of electrochemical processes. Rate-limiting steps were determined and discussed.

5.2 Results and discussion

5.2.1 Chemical reactivity with LSGM

The chemical reactivity of as-prepared Y3-xCaxFe5O12-δ (x = 0.1, 0.3 and 0.5) garnet powder with LSGM powder was determined, by mixing them in a 1:1 weight ratio and die-pressing into a thin pellet. The pellet was heated at 1100 °C for 2 hours in air and then grounded into powder for

PXRD analysis. Figure 5.1 shows the PXRD patterns of the mixtures after thermal treatment. It can be seen that, there are small amounts of secondary phases present after heating both powders together at high temperature, evidenced by the extra peaks marked in the plots, and indicate a small extent of chemical reactivity between the two. In order to avoid any potential chemical reactivity,

YSZ was tried as the alternative electrolyte and part of the composite electrode material. Although it was found that YSZ did not react with Ca-doped YIG, however, the composite garnet cathode with YSZ yielded a much larger ASR value than that with LSGM. Therefore, the symmetrical cell studies on garnet-LSGM composite electrode were further pursued, and the impact of all secondary phases were rationalized in the following paragraphs.

The secondary phases in Figure 5.1 can be assigned to un-doped or Ca-doped YFeO3

(Pnma), CaGa4O7 (C2/c) and Ca0.7Sr0.3O3 (Fm-3m) or Fe-doped Ca0.7Sr0.3O3-δ (Fm3m – cubic perovskite or Ibm2 – brownmillerite). The electrical conductivity of un-doped YFeO3 is reported

푥 to be ~ 0.07 S/cm in air at 750 ºC, arising from the disproportionation reaction of iron: 퐹푒퐹푒 =

′ . 161 퐹푒퐹푒 + 퐹푒퐹푒 . Doping Ca at Fe site in YFeO3 is revealed to increase its electrical conductivity to about 8 S/cm at 750 ºC, with additional induced p-type conductivity owing to the charge compensation mechanism shown as follows161:

1 1 퐶푎푂 + 푂 + 푌푥 → 퐶푎′ + ℎ. + 푌 푂 (5.1) 4 2 푌 푌 2 2 3

162 In the literature, Ca-doped YFeO3 has been demonstrated as SOFC cathode material .

Since the un-doped or Ca-doped YFeO3 is identified as the major impurity phase, its presence is expected to bring a relatively large impact on the electrochemical performance of symmetrical cells, possibly promoting ORR due to its high electrical conductivity and the existence of p-type charge carriers. On the other hand, CaGa4O7 might exhibit a very low electrical conductivity because of the structural restrictions and the absence of redox cation. However, its true conductivity value has not been reported, except one reference where it is shown that the formation of CaGa4O7 impurity phase can lead to a decrease in the electrical conductivity of Ca-doped

163 PrGaO3 . In addition, Ca0.7Sr0.3O3 or Fe-doped Ca0.7Sr0.3O3-δ is known to possess p-type electronic charge carriers with a total conductivity around 0.1 S/cm at 750 °C, caused by an

164 increased oxygen non-stoichiometry at elevated temperature . Similar to Ca-doped YFeO3, the presence of this phase is also anticipated to promote the electrochemical performance of ORR.

5.2.2 Electrocatalytic properties of ORR for SOFC cathode application

5.2.2.1 Evaluation of ORR performance

Symmetrical cell studies of garnet powder with LSGM as the composite electrode (Y3- xCaxFe5O12-δ: LSGM = 6:4 wt., x = 0.1, 0.3 and 0.5) were performed. The schematic diagram and cross-sectional SEM images of symmetrical cells after electrochemical measurements are shown in Figure 5.2. It can be seen that the electrode layers are all porous and adhere well to the LSGM electrolytes, with a uniform thickness of ~ 8 μm. There are no additional reaction layers observed, although a small extent of reactivity was found and discussed previously.

Figure 5.3 shows Nyquist plots of AC impedance data and ISGP fittings for symmetrical cells at 750 °C in air, where the garnet composition x = 0.3 exhibits the lowest ASR of 1 Ωcm2.

Figure 5.4 shows DFRT (г vs frequency) as a function of Ca content at 750 °C, where in general two peaks are seen, with a higher frequency peak labelled P1 and a lower frequency peak labelled

P2. The difference between two peak frequencies can be observed to be small at 750 °C from the plot. The ASR, capacitances and relaxation times associated with P1 and P2 at 750 °C in air are listed in Table 5.1. The capacitance values of the two peaks are found to be similar and both in the order of ~ 10-3 F/cm2. Typically, the electrochemical process with the capacitance value between

10-1 – 10-4 F/cm2 can be assigned to the dissociative/non-dissociative adsorption of oxygen, which is in agreement with our findings of ORR rate-limiting steps discussed later88,165,166. The difference in the relaxation times of P1 and P2 is also seen to be small, all within one order of magnitude and could be easily missed when using the method of ECM.

Similarly, in the temperature range of 600 – 800 °C, it is found that two peaks exist for all the three compositions and consistently exhibit thermal activation effect, where peaks shift towards higher frequencies with increasing temperature (Figures 5.5-5.6 for x = 0.3 symmetrical cell as a representative). The separation of the identified two peaks are visibly distinct at 600 °C, shown as grey line in Figure 5.6, while at higher temperatures they gradually converge on each other. On the other hand, at 850 °C and 900 °C, an additional third peak (P3) appears at lower frequencies.

The capacitance of P3 process lies in the range of 1-10 F/cm2, associated with the gas phase oxygen diffusion. The contribution of P3 to the total ASR is discovered to be minor, between 10-2 – 10-3

Ωcm2 at 900 °C in air, and can be further minimized through proper modification of the electrode microstructure, such as adding pore formers in the fabrication step165,167. The ASR, capacitances and relaxation times between 600-900 °C for x = 0.3 symmetrical cell are shown in Table 5.2.

The Arrhenius plot of total polarization resistance (Rp or ASRTotal) for x = 0.1, 0.3, and 0.5 is shown in Figure 5.7, where activation energies are obtained from linear regression, and are in

73 the range of 1.0 – 1.5 eV168. The activation energy values illustrate a decreasing trend from x = 0.1

(1.36 eV) to x = 0.5 (1.07 eV), and are relatively lower than those reported in the literature with similar composite cathode configuration using LSGM (For example, the composite cathode of

LSCF+LSGM (1:1 wt.) shows an activation energy of 1.7 eV)169.

Figure 5.2 Schematic diagram and cross-sectional SEM images of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells.

Figure 5.3 Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air.

Figure 5.4 DFRT plots (normalized by Rpol) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air.

Figure 5.5 Temperature variations of Nyquist plots, where symbols represent experimental data and solid lines indicate fitted data, of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell in air.

Figure 5.6 Temperature variations of DFRT plots (normalized by Rpol) of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell in air. 76

Figure 5.7 Arrhenius plot of total ASR with respect to temperature of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells in air.

Table 5.1 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in air.

-3 -3 R1 C1 (×10 R2 C2 (×10 -3 -3 Y3-xCaxFe5O12-δ τ1 (×10 s) τ2 (×10 s) (Ωcm2) F/cm2) (Ωcm2) F/cm2)

x = 0.1 0.41 6.90 2.86 1.16 3.97 4.60

x = 0.3 0.33 18.27 6.08 0.67 11.12 7.45

x = 0.5 0.32 15.73 4.99 0.87 6.00 5.20

Table 5.2 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 600-900 °C in air.

-2 τ1 -2 -2 τ3 2 C1 (×10 2 C2 (×10 -2 R3 (×10 2 T (°C) R1 (Ωcm ) 2 R2 (Ωcm ) 2 τ2 (×10 s) 2 C3 (F/cm ) F/cm ) (×10-2 s) F/cm ) Ωcm ) (×10-2 s)

600 3.20 2.95 9.44 16.55 2.20 36.39 - - -

650 1.21 2.61 3.16 3.93 1.38 5.41 - - -

700 0.59 2.07 1.22 1.39 1.17 1.63 - - -

750 0.33 1.83 0.61 0.67 1.11 0.75 - - -

800 0.19 2.03 0.38 0.46 1.01 0.46 - - -

850 - - - 0.39 0.64 0.25 1.05 7.35 7.69

900 - - - 0.27 0.59 0.16 0.66 10.05 6.65

5.2.2.2 Determination of rate-limiting step (RLS)

In order to assign the physical process to each peak, pO2 dependency of ASR was studied for all three compositions. Electrode behavior as a function of pO2 was measured using a mixture of O2 and N2 gas from 100% O2 to 5% O2. The Nyquist plot for x = 0.1, 0.3 and 0.5 symmetrical cells at 750 °C under different pO2 are shown in Figures 5.8-5.10, where the lowest ASR is seen

2 in 100% O2 atmosphere (for example, 0.49 Ωcm for x = 0.3), and the highest ASR is seen in 5%

2 O2 atmosphere (for example, 2.04 Ωcm for x = 0.3). The corresponding DFRT (г vs frequency) as a function of pO2 is shown in Figures 5.11-5.13, where two processes are present indicated by

P1 and P2, as also seen in the air atmosphere discussed previously. It can be observed that, with the increase in pO2, DFRT peaks shift to higher frequencies, demonstrating an enhanced catalytic activity of ORR with increasing pO2, as typically found in SOFC cathode materials. However, the changes of peak positions for P1 and P2, as well as their peak shapes, seen in Figures 5.11-5.13, are relatively small. The corresponding ASR, capacitances and relaxation times at 750 °C in different pO2 are shown in Table 5.3.

170–173 In addition, the ASR with respect to pO2 is known to follow the relationship as :

−푛 ASR = 퐴푆푅0(푝푂2) (5.2) where ASR0 is a constant and n is a parameter related to the rate-limiting step (RLS) of ORR.

Under current experimental condition of open-circuit potential (OCP), the change of ASR and the value of corresponding n are in fact associated with an average effect of both ORR and oxygen evolution reaction (OER).

Figure 5.14 summaries relationships between pO2 and polarization resistances attributed to total, low frequency peak and high frequency peak, respectively, for x = 0.3 symmetrical cell. It is

79 found that n for high frequency peak is 0.40, a value close to 0.375, which can be ascribed to the electrochemical process of partial reduction of adsorbed atomic oxygen into oxide ion (Equation

5.3; n = 0.375)170,174. On the other hand, the low frequency peak has a n value of 0.51, which is near 0.5 and can be assigned to the electrochemical process of dissociation of adsorbed molecular oxygen into atomic oxygen (Equation 5.4; n = 0.5)170,171,174.

− − 푂푎푑푠 + 푒 ↔ 푂푎푑 (5.3)

푂2,푎푑푠 ↔ 2푂푎푑푠 (5.4)

Therefore, it is believed that the symmetrical cell with x = 0.3 garnet-LSGM composite electrode at 750 °C is limited by a combination of dissociation and partial reduction of the adsorbed oxygen molecules. A previous study by Zhong et al. on a symmetrical cell with x = 0.5 garnet-

SDC composite electrode at 650 °C showed that ASRHF was independent of pO2 and was associated with the process of oxygen ion diffusion, while ASRLF with a slope of 0.63 indicated contribution of both dissociative adsorption of oxygen (n = 0.5) and oxygen gas diffusion (n = 1)31.

The respective capacitances for high and low frequency processes were reported to be 2.10×10-4

(F/cm2) and 5.76×10-3 (F/cm2)31. Comparing it to the current research, we do not see gas diffusion issue, since the cathode layer is very porous and is also thinner (evidenced by SEM images in

Figure 5.2, the cathode layer is ~ 8 µm in the present work vs. ~ 50 µm as reported in the literature).

There is also no presence of oxygen ion diffusion as an RLS in our current study, which could be due to the difference in the electronic-ionic properties of electrolytes used and mixed in the composite electrode, as well as temperature difference that influences thermally activated ionic conductivity. SDC is known to show some electronic conductivity above 600 °C, while LSGM remains mostly ionic even at 750 °C. In addition, an increase of Ca substitution level from 0.3 to

0.5 may also lead to property variations in oxide ion conductivity. On the other hand, similar to

Zhong et al., we have found that the dissociation of adsorbed molecular oxygen into atomic oxygen limits the ORR, which is likely the material property of these Ca-doped yttrium iron garnets.

Figure 5.11 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.9Ca0.1Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C.

Figure 5.12 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C.

Figure 5.13 Oxygen partial pressure variations of DFRT plots (normalized by Rpol) obtained from ISGP fitting of Y2.5Ca0.5Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C.

Figure 5.14 Log (ASR) as a function of log (pO2) obtained from ISGP of Y2.7Ca0.3Fe5O12-δ + LSGM / LSGM symmetrical cell at 750 °C.

Table 5.3 Calculated ASR (R), capacitance (C) and relaxation time (τ) of Y3-xCaxFe5O12-δ + LSGM / LSGM (x = 0.1, 0.3 and 0.5) symmetrical cells at 750 °C in various oxygen partial pressures.

Log [pO C (×10-3 C (×10-3 Y Ca Fe O 2 R (Ωcm2) 1 τ (×10-3 s) R (Ωcm2) 2 τ (×10-3 s) 3-x x 5 12-δ (atm)] 1 F/cm2) 1 2 F/cm2) 2 0 0.22 4.43 0.97 0.58 1.91 1.10

-0.68 0.53 3.63 1.91 1.06 2.27 2.40 x = 0.1 -1 0.81 3.45 2.79 1.49 2.52 3.75

-1.3 1.07 3.68 3.93 2.11 2.65 5.55

0 0.16 3.86 0.62 0.33 2.20 0.72

-0.68 0.34 4.33 1.46 0.66 2.63 1.75 x = 0.3 -1 0.44 5.49 2.40 1.05 2.75 2.88

-1.3 0.52 7.14 3.75 1.52 2.84 4.33

0 0.25 4.78 1.19 0.39 3.61 1.39

-0.68 0.58 4.47 2.79 0.61 5.20 2.92 x = 0.5 -1 1.24 3.60 4.45 0.43 10.99 4.69

-1.3 1.79 3.75 6.69 0.47 16.70 8.60

5.3 Summary

In conclusion, LSGM was examined to have a small extend of chemical reactivity with Ca- doped YIG. Symmetrical cell studies revealed the lowest ASR of 1.0 Ωcm2 at 750 °C in air for

Y2.7Ca0.3Fe5O12-δ garnet-LSGM composite electrode. Impedance spectra were deconvoluted by

ISGP into two electrochemical processes with similar relaxation times, which would be difficult to achieve with the conventional ECM method. By measuring ASR in varying pO2 environments, oxygen dissociation and the partial reduction of adsorbed oxygen molecules have been suggested to be the rate-limiting steps of ORR at 750 °C for x = 0.3 garnet-LSGM composite electrode.

Chapter Six: Conclusions and Future Work

6.1 Conclusions

In Chapter 4, garnet-type metal oxide Y3-xCaxFe5O12-δ (0, 0.05, 0.1, 0.3, 0.5 and 0.7) was synthesized using the sol-gel method. PXRD pattern of the sintered powder revealed that a single cubic-phase (Ia-3d) compound formed when Ca substitution level was below 0.5. For the compositions where there were 0.5 or more Ca substitution per unit formula, two secondary phases of CaFe2O4 (Pnma) and Ca2Fe2O5 (Pnma) were observed, along with the major garnet cubic-phase

(Ia-3d). The garnet structural variations from the Ca substitution were studied in detail by performing Rietveld refinements on the obtained PXRD patterns and analyzing changes in lattice parameter, metal-oxygen bond length and metal-oxygen-metal bond angle. It was found that although doping element of Ca2+ (1.00 Å) has a larger ionic radius than Y3+ (0.90 Å), the resulting unit cell length did not demonstrate a monotonic increasing trend, which could be owing to charge compensation mechanism of forming either oxide ion vacancies or higher valency Fe (Fe4+). In order to validate the existence of presumed charge compensation mechanisms, iodometric titration was performed on compositions of x = 0, 0.05, 0.1, 0.3 and 0.5. The results showed all Ca-doped garnets possess oxygen deficiency, with oxygen non-stoichiometry (δ) ranging from 0.06 (x =

0.05) to 0.19 (x = 0.3). On the other hand, electronic structure at the Fe and O sites were investigated by XAS for garnet compositions of x = 0, 0.1 and 0.5, where Fe was seen to have a higher valency in x = 0.1 than x = 0, and an electron-hole-related spectroscopic feature was found in O K-edge spectra for both Ca-doped samples. These discoveries could be concluded that the formation of both oxygen vacancy and Fe4+ might co-exist in Ca-doped garnets. Furthermore, the electrical transport properties were explored using synthesized garnet pellets. Electrical conductivity was measured by 4-probe DC method and a maximum was found in the composition

87 of x = 0.1. The initial increasing tendency of electrical conductivity from x = 0 to x = 0.1 is suspected to be due to an increase in concentration of electron-hole charge carriers from aliovalent

Ca substitution. A later decreasing trend from x = 0.1 to x = 0.5 is deduced owing to possible

′ ∙ formation of defect associates (퐶푎푌 − 퐹푒퐹푒). An attempt of measuring ionic conductivity and calculating ionic transport number was made through Hebb-Wagner polarization method for the composition of x = 0.1. As a result, a relatively low ionic conductivity of 2.89×10-5 S/cm at 750

°C was determined and a corresponding ionic transport number was obtained as 1.82×10-5. A comparison was made of this result to a similar literature report to crosscheck the validity of the data in the current work, where reasonable agreement was found.

In Chapter 5, synthesized Y3-xCaxFe5O12-δ (0.1, 0.3 and 0.5) powder was mixed with either

YSZ or LSGM and tested as an electrode material under 2-probe symmetrical cell configuration for the potential application of SOFC cathode. Although LSGM and Ca-doped garnets showed a small extent of chemical reactivity when heated together at 1100°C in air, forming secondary phases such as undoped or Ca-doped YFeO3 (Pnma), CaGa4O7 (C2/c) and undoped or Fe-doped

Ca0.7Sr0.3O3 (Fm3m or Ibm2), however, the area-specific resistance (ASR) of symmetrical cell using garnet-LSGM composite electrode was still proven to be much smaller than with garnet-

YSZ mixture. In addition, some of these impurities, from literature search, were analyzed to be able to promote oxygen reduction reaction (ORR) instead of supressing the process. Therefore, garnet-LSGM (3:2 wt.) composite electrode was chosen to be studied further. Electrochemical impedance spectroscopy was used in the current work for the evaluation of ORR performance at

600–900 °C and different oxygen partial pressures (pO2). It was found that the composite electrode with garnet composition of x = 0.3 showed the lowest ASR (1.00 Ωcm2) at 750 °C in air compared to the other two compositions of x = 0.1 and x = 0.5. The obtained EIS spectra were deconvoluted

88 through impedance spectroscopy genetic programming (ISGP), where two electrochemical sub- processes were identified for all the three compositions at 600–800 °C in air and at 750 °C in pO2 range of -1.5 – 0. The capacitance values of these two deconvoluted peaks were found to be between 10-1–10-4 F/cm2, which could be assigned to the dissociative/non-dissociative adsorption of oxygen. When the temperature was above 800 °C in air, a third sub-process was found with a capacitance value between 1–10 F/cm2, which can be associated with gas phase oxygen diffusion.

−푛 By relating changes in pO2 to ASR using the equation ASR = 퐴푆푅0(푝푂2) for symmetrical cell with x = 0.3 garnet-LSGM composite electrode at 750 °C, the parameter n can be calculated and the corresponding rate-limiting step (RLS) can be inferred. It was shown that n for high frequency peak is 0.4, indicating the process of partial reduction of adsorbed atomic oxygen into oxide ion, while n for low frequency peak is 0.51, suggesting the process of dissociation of adsorbed molecular oxygen into atomic oxygen. As a result, these two processes were deduced as RLS for the x = 0.3 symmetrical cell at 750 °C. Finally, comparison with a similar literature report using x

= 0.5 garnet-SDC composite electrode was performed.

6.2 Future work

Besides the methods of iodometric titration and TGA employed in the current work, the oxygen non-stoichiometry of the synthesized garnet powder can also be determined by neutron diffraction. It is possible that with a joint neutron/x-ray diffraction data sets refinement, more information on the structural parameters and oxide ion vacancies can be obtained, which could lead to a more consistent explanation with the existing results of electronic structure from XAS measurements.

Hebb-Wagner polarization measurement was performed to separate ionic from electrical conductivity. However, a rigorous experimental design requires the sample pellet used to be dense and completely gas-tight. Using higher pressure to make the garnet pellet (for example, isostatic pressure > 200 kN) and/or sintering the pellet at a higher temperature would help to improve the density of the pellet, and therefore offer greater validity of the obtained results.

Potential theoretical calculations using density functional theory (DFT) on the charge distribution and oxide ion vacancy site in the garnet lattice can be carried out for a better understanding of the phase and ionic conduction pathway of the material. In addition, comparison can be made between samples of different Ca-substitution level to provide insight in the variations of their electrical transport properties.

The determination of rate-limiting step of garnet-LSGM composite electrode in the current work was conducted only at 750 °C and in a pO2 range of -1.5 – 0. Increasing the tested pO2 range and performing the experiment at multiple temperatures would broaden the understanding of electrochemical processes that limit the electrode ORR performance.

It is worth exploring the substitution of yttrium with other elements besides calcium in the parent compound yttrium iron garnet, and the corresponding electrochemical properties of ORR.

Finally, it is known that beyond electrochemical properties, the parent phase yttrium iron garnet also possesses interesting magnetic properties. Further studies should be conducted to probe the variations in magnetic properties of Ca-doped garnets and explore their potential applications.

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