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Novel Ruthenium Pyrochlore Materials for Cathode Application in Intermediate Temperature Solid Oxide Fuel Cells (IT-SOFCs)

Permanent Link: http://ufdc.ufl.edu/UFE0022800/00001

Material Information

Title: Novel Ruthenium Pyrochlore Materials for Cathode Application in Intermediate Temperature Solid Oxide Fuel Cells (IT-SOFCs)
Physical Description: 1 online resource (117 p.)
Language: english
Creator: Abate, Chiara
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2008

Subjects

Subjects / Keywords: bismuth, cathode, ceria, composite, conductivity, electrical, electrochemical, material, pyrochlore, ruthenium, sofc
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: The performance of solid oxide fuel cells, which operate in the temperature range of 773-973 K (IT-SOFCs), strongly depends on the cathode employed because the interfacial polarization increases rapidly with decreasing temperature. Pyrochlore oxides with Ru on the B-site of the crystal lattice have been shown to have excellent electro catalytic behaviour for oxygen reduction reaction and high electrical conductivity. These characteristics make pyrochlore ruthenates good candidates for IT-SOFCs cathodes. In this work, several compositions of Y sub (two minus x) Pr sub x Ru sub two O sub seven (x = 0, 0.2, 0.5, 1, 1.5, 2) pyrochlore powders were prepared by a soft precipitation method. All the synthesized powders were single pyrochlore phase with particles size depending on the material compositions. Praseodymium (Pr) was introduced in the A-site with the intent to improve the material electrical proprieties and consequently the overall cathode performance. In fact, without destabilizing the pyrochlore structure, Pr caused structural changes that allow higher electron mobility. The electrical measurements showed that the electrical conductivity of the material increased with increasing the Pr content. Compositions of Y sub (two minus x) Pr sub x Ru sub two O sub seven were tested as a cathode to compare its electro-catalytic effect with either of two electrolytes, gadolinium doped ceria (GDC) or erbium stabilized bismuth oxide (ESB). Both systems, Y sub (two minus x) Pr sub x Ru sub two O sub seven / ESB and Y sub (two minus x) Pr sub x Ru sub two O sub seven / GDC, showed a similar variation of the electrode area specific resistance (ASR) with Pr content. This trend was shown to be due to a change of the cathode microstructure upon increasing Pr content. The 25 mol % Pr cathode material on ESB electrolyte presented the best performance. A change of ASR as a function of oxygen partial pressure suggested that the oxygen diffusion is the limiting step of the electrode kinetics. Hence, the better cathode performance on ESB resulted from a much lower charge transfer resistance compared to the GDC system. A partial solid diffusion observed using SEM on the Y sub (one point five) Pr sub (zero point five) Ru sub two O sub seven / ESB interface likely contributed to lower the interfacial polarization in this system. These results suggested that the nanocrystalline yttrium praseodimium rutenate powder, with 25 mol % of Pr, is promising for cathode application in ESB-based electrolyte for intermediate temperature solid oxide fuel cells (IT-SOFCs).
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Chiara Abate.
Thesis: Thesis (Ph.D.)--University of Florida, 2008.
Local: Adviser: Wachsman, Eric D.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2010-12-31

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2008
System ID: UFE0022800:00001

Permanent Link: http://ufdc.ufl.edu/UFE0022800/00001

Material Information

Title: Novel Ruthenium Pyrochlore Materials for Cathode Application in Intermediate Temperature Solid Oxide Fuel Cells (IT-SOFCs)
Physical Description: 1 online resource (117 p.)
Language: english
Creator: Abate, Chiara
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2008

Subjects

Subjects / Keywords: bismuth, cathode, ceria, composite, conductivity, electrical, electrochemical, material, pyrochlore, ruthenium, sofc
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: The performance of solid oxide fuel cells, which operate in the temperature range of 773-973 K (IT-SOFCs), strongly depends on the cathode employed because the interfacial polarization increases rapidly with decreasing temperature. Pyrochlore oxides with Ru on the B-site of the crystal lattice have been shown to have excellent electro catalytic behaviour for oxygen reduction reaction and high electrical conductivity. These characteristics make pyrochlore ruthenates good candidates for IT-SOFCs cathodes. In this work, several compositions of Y sub (two minus x) Pr sub x Ru sub two O sub seven (x = 0, 0.2, 0.5, 1, 1.5, 2) pyrochlore powders were prepared by a soft precipitation method. All the synthesized powders were single pyrochlore phase with particles size depending on the material compositions. Praseodymium (Pr) was introduced in the A-site with the intent to improve the material electrical proprieties and consequently the overall cathode performance. In fact, without destabilizing the pyrochlore structure, Pr caused structural changes that allow higher electron mobility. The electrical measurements showed that the electrical conductivity of the material increased with increasing the Pr content. Compositions of Y sub (two minus x) Pr sub x Ru sub two O sub seven were tested as a cathode to compare its electro-catalytic effect with either of two electrolytes, gadolinium doped ceria (GDC) or erbium stabilized bismuth oxide (ESB). Both systems, Y sub (two minus x) Pr sub x Ru sub two O sub seven / ESB and Y sub (two minus x) Pr sub x Ru sub two O sub seven / GDC, showed a similar variation of the electrode area specific resistance (ASR) with Pr content. This trend was shown to be due to a change of the cathode microstructure upon increasing Pr content. The 25 mol % Pr cathode material on ESB electrolyte presented the best performance. A change of ASR as a function of oxygen partial pressure suggested that the oxygen diffusion is the limiting step of the electrode kinetics. Hence, the better cathode performance on ESB resulted from a much lower charge transfer resistance compared to the GDC system. A partial solid diffusion observed using SEM on the Y sub (one point five) Pr sub (zero point five) Ru sub two O sub seven / ESB interface likely contributed to lower the interfacial polarization in this system. These results suggested that the nanocrystalline yttrium praseodimium rutenate powder, with 25 mol % of Pr, is promising for cathode application in ESB-based electrolyte for intermediate temperature solid oxide fuel cells (IT-SOFCs).
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Chiara Abate.
Thesis: Thesis (Ph.D.)--University of Florida, 2008.
Local: Adviser: Wachsman, Eric D.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2010-12-31

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2008
System ID: UFE0022800:00001


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1 NOVEL RUTHENIUM PYROCHLORE MATERIA LS FOR CATHODE APPLICATION IN INTERMEDIATE TEMPERATURE SOLID OXIDE FUEL CELLS (IT-SOFCS) By CHIARA ABATE A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2008

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2 2008 Chiara Abate

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3 To my mother Maria

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4 ACKNOWLEDGMENTS I would like to thanks my advisor Dr. Enrico Traversa and Dr. Eric Wachsman for their support and guidance throughout this experience. I appreciate the oppor tunity that they gave me of taking part in a joint PhD program between th e University of Tor Vergata and the University of Florida. I think the joint la b was a great experience that help ed me to build my knowledge. I would also like to thank Dr. Juan Nino, Dr. Wolfgang Sigmund, Dr. Mark Orazem, and Dr Jacob Jones for their advice and particip ation as part of my committee. I wish to acknowledge Dr. Keith Duncan for the many conversations that we had on the various theoretical aspects of the field, and Dr Vincenzo Esposito for his advice and collaboration with some experimental aspects of this work. Many thanks are owed to Dr. Luisa Dempere as well as Dr Valentin Craciun and Eric Lambers of the Major Analytical Research Center at the University of Florida for their assistance with the technical instruments and the microstructural analysis. Many thanks go to my lab mates and office mates for their helpful discussions and friendship. Special thanks to Rosaria for her supp ort in some difficult situations. Finally I would like to thanks my family for their support and encouragement in all the circumstances. This work was supported in part by MIUR (F ISR Project), by the U.S. Department of Energy, contract DE-FC26-03NT41959 and by the Mini stry of Foreign Affairs (MAE) of Italy under the frame of the Italy-USA Joint Laborator y on Nanostructured Materials for Solid State Ionic Devices.

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5 TABLE OF CONTENTS page ACKNOWLEDGMENTS...............................................................................................................4 LIST OF TABLES................................................................................................................. ..........7 LIST OF FIGURES................................................................................................................ .........8 ABSTRACT....................................................................................................................... ............11 CHAPTER 1 INTRODUCTION................................................................................................................. .13 2 SOLID OXIDE FUEL CELLS...............................................................................................16 2.1 Fuel Cell Efficiency...................................................................................................... 17 2.2 Fuel Cell Components...................................................................................................20 2.2.1 The Electrolyte..................................................................................................20 2.2.2 The Electrodes...................................................................................................22 2.3 Type of Cathodes.......................................................................................................... 24 2.4 Cathode Materials......................................................................................................... 27 2.5 Summary................................................................................................................... ....29 3 PYROCHLORE RUTHENATES..........................................................................................33 3.1 Pyrochlore Structure.....................................................................................................3 3 3.1.1 Existence and Stability Field.............................................................................35 3.1.2 Crystal Data of Y2Ru2O7 by Neutron Powder Diffraction................................36 3.2 Electrical Properties..................................................................................................... .36 3.3 Summary................................................................................................................... ....38 4 SYNTHESIS OF NOVEL YTTRIU M PRASEODYMIUM RUTHENATES......................41 4.1 Introduction.............................................................................................................. .....41 4.2 Experimental.............................................................................................................. ...42 4.2.1 Materials Preparation........................................................................................42 4.2.2 Materials Characterization................................................................................43 4.3 Results and Discussion..................................................................................................44 4.4 Conclusions............................................................................................................... ....48 5 YTTRIUM PRASEODYMIUM RUTH ENATE ELECTRICAL PROPERTIES..................56 5.1 Introduction.............................................................................................................. .....56 5.2 Experimental.............................................................................................................. ...56

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6 5.3 Results and Discussion..................................................................................................57 5.4 Conclusions............................................................................................................... ....59 6 YTTRIUM PRASEODYMIUM RUTHE NATE CATHODES FOR IT-SOFCS..................63 6.1 Introduction.............................................................................................................. .....63 6.2 Experimental Procedure................................................................................................64 6.3 Results and Discussion..................................................................................................66 6.3.1 Applicability of Y1-xPrxRu2O7 Cathodes on GDC Electrolyte..........................66 6.3.1.1 Chemical compatibility w ith the solid electrolyte..............................66 6.3.1.2 Microstructural an alysis and electroche mical performances..............66 6.3.1.3 Summary.............................................................................................70 6.3.2 Applicability of Y1-xPrxRu2O7 Cathodes on ESB Electrolyte...........................71 6.3.2.1 Electrochemical performances............................................................72 6.3.2.2 Microstructural analysis......................................................................74 6.3.2.3 Summary.............................................................................................77 6.4 Conclusions............................................................................................................... ....77 7 COMPOSITE ELECTRODES ON GDC ELECTROLYTE..................................................93 7.2 Experimental Procedure................................................................................................93 7.3 Composite Y1.5Pr0.5Ru2O7-GDC Electrodes.................................................................95 7.4 Composite Y1.5Pr0.5Ru2O7-ESB Electrodes..................................................................97 7.5 Conclusions............................................................................................................... ....98 8 CONCLUSIONS.................................................................................................................. 106 APPENDIX A DATA REPRODUCIBILITY..............................................................................................109 LIST OF REFERENCES............................................................................................................. 112 LIST OF REFERENCES............................................................................................................. 112 BIOGRAPHICAL SKETCH.......................................................................................................117

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7 LIST OF TABLES Table page 3-1 Fractional atomic coordinates and anisotropic displacement parameters (2 x 103).........39 4-1 Lattice parameters for Y1-x PrxRu2O7 (x = 0 2) from X-ray data.....................................51

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8 LIST OF FIGURES Figure page 2-1. Solid oxide fuel cell..................................................................................................... ......30 2-2. An SOFC stack............................................................................................................. .....30 2-3. Typical current-volta ge profile for SOFC..........................................................................31 2-4. Arrhenius plot for common solid electrolyte materials.....................................................31 2-5. Various reaction pathways fo r the oxygen reduction reaction...........................................32 3-1. One octave of the unit cell in the pyrochlore structure......................................................39 3-2. The Y2Ru2O7 crystal structure...........................................................................................40 3-3. Relationship between the ionic radius of the A cation and the Ru-O Ru bond angle.......40 4-1. Simultaneous TG/DTA analysis on Y2Ru2O7 precursor....................................................49 4-2. The XRD pattern of Y2Ru2O7 after different thermal treatments for 5h each...................49 4-3. The XRD pattern of Y1-xPrxRu2O7 (x = 0 2) powders af ter crystallization at 1050 C....50 4-4. Observed, calculated and difference di ffraction patterns for the X-Ray data for Y1.5Pr0.5Ru2O7 powder.......................................................................................................50 4-5. Composition dependence of the lattice parameters for Y1-xPrxRu2O7 (x = 0 2)..............51 4-6. The FE-SEM pictures af ter crystallization at 1050oC........................................................52 4-7. High magnification FE-SEM pictur es after crystallization at 1050oC...............................53 4-8. The XPS Pr 3d core level spectra of Y1.8Pr0.2Ru2O7..........................................................54 4-9. The XPS Ru 3d core level spectra of Y1.8Pr0.2Ru2O7.........................................................55 5-1. Variation of the electrical conductivity of Y2-xPrxRu2O7 with dopant amount..................60 5-2. Temperature dependence of th e electrical conductivity of Y2-xPrxRu2O7.........................61 5-3. Electrical co nductivity of Y1.8Pr0.2Ru2O7 and Y1.5Pr0.5Ru2O7 as function of oxygen partial pressure at 800oC....................................................................................................62 6-1. The SEM-EDS line scan of Y1.5Pr0.5Ru2O7/GDC interface...............................................79 6-2. The XRD patterns of mixture of Y1.8Pr0.2Ru2O7-GDC (molar ratio 1:1)...........................80

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9 6-3. The FE-SEM pictures of electrodes fired at 850oC on GDC.............................................80 6-4. Low magnification FE-SEM pictur es of electrodes fired at 850oC on GDC.....................81 6-5. Impedance plot of Y1.5Pr0.5Ru2O7 electrodes on GDC electrolyte at 700oC in air............81 6-6. Impedance plot of Y1.5Pr0.5Ru2O7/GDC/Y1.5Pr0.5Ru2O7 cell at 700oC...............................82 6-7. Variation of the electrode polarization resistance as a function of the oxygen partial pressure....................................................................................................................... .......82 6-8. Temperature dependence of the electrode ASR on GDC..................................................83 6-9. The ASR values of Y2-xPrxRu2O7 on GDC pellets tested at di fferent temperatures as a function of Pr content......................................................................................................... 84 6-10. The SEM-EDS line scan of Y1.5Pr0.5Ru2O7/ESB interface................................................85 6-11. The XRD patterns of mixture of Y1.5Pr0.5Ru2O7-ESB (molar ratio 1:1)............................86 6-12. The XRD patterns of mixture of YPrRu2O7-ESB (molar ratio 1:1)..................................87 6-13. Impedance plots of Y2-xPrxRu2O7 electrodes on ESB electrolyte at 700oC in air..............87 6-14. Complex impedance plot for Y1.5Pr0.5Ru2O7 /ESB/ Y1.5Pr0.5Ru2O7 cell, measured in air at 700oC........................................................................................................................88 6-15. Arrhenius plot of the electrode ASR on ESB....................................................................88 6-16. The ASR values of Y2-xPrxRu2O7 on ESB pellets tested at different temperatures as function of Pr content......................................................................................................... 89 6-17. The FE-SEM pictures of YPrRu2O7 and Y0.5Pr1.5Ru2O7 electrodes..................................90 6-18. Arrhenius plots. A) Pr 50 electrodes. B) Pr 75 electrodes.................................................90 6-19. The FE-SEM pictures of Y1.5Pr0.5Ru2O7 electrodes fired on ESB at 800oC for 2h...........91 6-20. Arrhenius plot of the Y1.5Pr0.5Ru2O7 electrode ASR.........................................................92 6-21. The ASR values as function of Pr content of Y2-xPrxRu2O7 on ESB and on GDC pellets tested at 700oC........................................................................................................92 7-1. Variation of the ASR of Y1.5Pr0.5Ru2O7/GDC composite electrodes as a function of vol % GDC...................................................................................................................... .100 7-2. The FE-SEM image of the 80/20 Y1.5Pr0.5Ru2O7-GDC composite electrode..................100

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10 7-3. Impedance plots of 80/20 Y1.5Pr0.5Ru2O7-GDC composite electrodes on GDC electrolyte in air............................................................................................................. ..101 7-4. Impedance plots for Y1.5Pr0.5Ru2O7 single phase and composite 80/20 Y1.5Pr0.5Ru2O7-GDC on GDC electrolyte, measured in air at 700oC...............................101 7-5. Arrhenius plots of the ASR for single phase and composite 80/20 Y1.5Pr0.5Ru2O7GDC electrodes................................................................................................................1 02 7-6. Variation of the ASR of Y1.5Pr0.5Ru2O7/ESB composite electrodes as a function of vol % ESB...................................................................................................................... ..102 7-7. The FE-SEM image of 40/60 Y1.5Pr0.5Ru2O7-ESB composite electrodes.......................103 7-8. Impedance plot for the 40/60 Y1.5Pr0.5Ru2O7-ESB composite on GDC electrolyte measured in air at 700oC..................................................................................................103 7-9. Impedance plots for Y1.5Pr0.5Ru2O7 single phase and composite with 60 vol% of ESB on GDC electrolyte, measured in air at 700oC.................................................................104 7-10. Arrhenius plots of the ASR for single phase and 40/60 Y1.5Pr0.5Ru2O7-ESB composite electrodes........................................................................................................104 7-11. Arrhenius plots of the ASR for singl e phase and composites electrodes on GDC electrolyte.................................................................................................................... .....105 A-1. Arrhenius plots of the ASR for Y2Ru2O7 electrode on ESB electrolyte, tested in the laboratories of the University of Ro me and the University of Florida............................111 A-2. Arrhenius plots of the ASR for Y1.8Pr0.5Ru2O7 electrode on ESB electrolyte, tested in the laboratories of the Un iversity of Rome and the University of Florida......................111

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11 Abstract of Dissertation Pres ented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy NOVEL RUTHENIUM PYROCHLORE MATERIA LS FOR CATHODE APPLICATION IN INTERMEDIATE TEMPERATURE SOLID OXIDE FUEL CELLS (IT-SOFCS) By Chiara Abate December 2008 Chair: Eric Wachsman Cochair: Enrico Traversa Major: Materials Science and Engineering The performance of solid oxide fuel cells, wh ich operate in the temperature range of 500700oC (IT-SOFCs), strongly depends on the cath ode employed because the interfacial polarization increases rapidly with decreasing temperature. Pyrochlore oxides with Ru on the Bsite of the crystal lattice have been shown to have excellent electro catalytic behaviour for oxygen reduction reaction and high electrical conductivity. These characteristics make pyrochlore ruthenates good candi dates for IT-SOFCs cathodes. In this work, several compositions of Y2-xPrxRu2O7 (x = 0, 0.2, 0.5, 1, 1.5, 2) pyrochlore powders were prepared by a soft precipitation method. All the s ynthesized powders were single pyrochlore phase with particles size depending on the material compositions. Praseodymium (Pr) was introduced in the A-site with the intent to improve the material elec trical proprieties and consequently the overall cathode performance. In fact, without destab ilizing the pyrochlore structure, Pr caused structural changes that allow higher electron mobility. The electrical measurements showed that the el ectrical conductivity of the material incr eased with increasing the Pr content.

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12 Compositions of Y2-xPrxRu2O7 were tested as a cathode to compare its electro-catalytic effect with either of two el ectrolytes, gadolinium doped ceria (GDC) or erbium stabilized bismuth oxide (ESB). Both systems, Y2-xPrxRu2O7/ESB and Y2-xPrxRu2O7/GDC, showed a similar variation of the electrode area specific resistance (ASR) w ith Pr content. This trend was shown to be due to a change of the cathode microstructure upon increasing Pr content. The 25 mol % Pr cathode material on ESB electrolyte presented the best performance. A change of ASR as a function of oxygen partial pre ssure suggested that th e oxygen diffusion is the limiting step of the electrode kinetics. Hence, the better cathode performance on ESB resulted from a much lower charge transfer resistance compared to the GDC system. A partial solid diffusion observed using SEM on the Y1.5Pr0.5Ru2O7/ESB interface likely contributed to lower the interfacial polarization in this system. Th ese results suggested that the nanocrystalline powders of Y1.5Pr0.5Ru2O7 electrode are promising material s for cathode application in ESBbased electrolyte for intermediate temper ature solid oxide fuel cells (IT-SOFCs).

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13 CHAPTER 1 INTRODUCTION Nowadays there is a real interest in the development of a new source of energy, because there is an inadequate energy supply in re sponse to the increasing world energy demand. Actually, world energy system is dominated by foss il fuels, in which oil is the main source of energy. According to the US-based Energy Info rmation Administration’ s (EIA) annual report, the world demand for oil is project ed to increase 37% over the 2006 by 2030.1 Most of the energy demand is expected to come from the developing countries due to economic and population growth. On the other hand, oil supplie s dwindle and the fact that global oil supply will decline at some point is the main fundamental cause of rising prices. Indeed, especially in the last five years, the price of a barrel of oil drastically incr eased until reaching four times the original price at the be ginning of this year. Therefore, the de velopment of a new power source is becoming increasingly important. In addition to the problems with energy supply, there are many problems related to environmental concerns and globa l warming. Indeed, the high level of the gas emissions causes air pollution, acid precip itation and ozone depletion, which ca n be dangerous for human health. Moreover, the greenhouse gas emissions (mainly CO2) are considered one of the causes of global warming.2 The increase of global surface temperatures ov er the past few years can cause a series of ecological changes that may increase the am ount of natural disasters like heat waves, hurricanes or tornados. According to the US de partment of Energy (DOE), world emission of carbon are expected to increase by 50% above 1990 levels by 2030.3 Therefore, there is a general need of reducing polluting byproducts ge nerated by conventional energy production. Fuel cells are a good alternative source of power because of their high conversion efficiency and low pollutant emissions.4,5 A fuel cell is an electroche mical device that directly

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14 converts chemical energy of a fuel and an oxi dant into electrical work. No combustion is required, and the conversion efficiency is high er than the conventional thermo-mechanical methods. Even though atmospheric emissions are usually released during the fuel processing, less fuel is consumed to produce the same am ount of electricity compared to a combustion engine. Also, a lower net amount of carbon(CO, CO2), nitrogen(NOx), and sulfur-based (SOx) pollutants are emitted. For all these reasons, fuel cells have become a viable substitute to be an environmentally friendly way to generate electricity. There are several types of fuel cells which mainly differ from each other because of the electrolyte material that is us ed. Consequently, they present different operation temperature range and different prospective applications. Al kaline fuel cells (AFCs) direct-methanol fuel cells (DMFC), phosphoric acid fuel cells (PAFCs), sulfuric acid fuel cells (SAFC), and proton exchange membrane (PEM) fuel cells can operate at low te mperatures (50-210oC) with efficiency of 40-50%. Molten carbonate fuel cells (MCFCs), solid oxide fuel cells (SOFCs), and protonic ceramics fuel cells (PCFC) ha ve a high operation temperature (600-100oC) but they can directly use methane into the ce ll. For this reason, they pres ent an high efficiency, around 4560% or 90% with heat recovery. Compare to the other type of fuel cells, SOFCs present several advantages. At high temperature, high reaction rate, high mass transfer rate and lower cell resistance contribute to improve cell efficiency. Besides, the high operating temperature al lows the direct oxidation of the fuel without external reforming, a nd any CO produced is converted in CO2, therefore SOFCs have extremely low emissions. Finally, they produce high quality byproduct heat for cogeneration.

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15 On the other hand, the high temperature conditions place stringent requirements on SOFC materials and puts very high demands on the technology. To overcome these problems, many studies have been oriented toward lo wering the operation temperature to 500–700 C.6,8 Lower the operation temperature will allow to use cheaper interconnects and structural components, reduce thermal stresses in the ceramic structure, reduce the start-up time and lead to a longer expected lifetime of the system. However, a set of fully compatible materials has not yet been developed for operation at reduced temperature. Therefore, new materi als and manufacturing technologies need to be studied in order to reduce the operation temper ature of SOFCs while maintaining their high performances.

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16 CHAPTER 2 SOLID OXIDE FUEL CELLS A typical single SOFC is composed of a de nse ceramic oxide elect rolyte and two porous electrodes that enable gas diffu sion to and from the electrode-electrolyte interface. At the cathode, the electrochemical re duction of oxygen occurs. The oxyge n ions migrate through the electrolyte via a vacancy mechanism, while the el ectrons are forced to flow through an external circuit. At the anode, an elec trochemical oxidation of hydroge n occurs. Figure 2-1 shows the schematic drawing of a typical so lid oxide fuel cell with the elec trochemical reactions that occur at the electrodes. At equilibrium the cell voltage is related to the change in Gibbs free energy resulting from the reaction between hydr ogen and oxygen (Eq. 2-1). E zF G 0 (2-1) The change in Gibbs free energy can be calculated by the oxida tion potential of H2 which corresponds to the ideal standard po tential for the r eaction between H2 and O2. This potential is equal to 1.229 V with liquid water product or 1. 18 V with gaseous water product. The difference between these values represents the latent heat of vaporization of water at standard conditions (25oC and 1 atm). The driving force for the migration of oxyge n ions is the oxygen chemical potential gradient between the cathode and the anode. Air is usually used in the cathode side, which has a partial pressure for oxygen of 0.21 atm. At the a node side, the oxygen partial pressure is much lower due to the consumption of oxygen during the electrochemical reaction of the fuel. In open circuit conditions, where no charge flows across the electrolyte, the theoretical reversible voltage of the cell is described by the Nernst equation (Eq. 2-2).

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17 a cpO pO F RT E, 2 2ln 4 (2-2) R is the gas constant, T is the absolute te mperature, F is Faraday’s constant, and pO2 a and pO2 c are the partial pressures of oxygen at the an ode and cathode, respectively. The coefficient 4 in the denominator represents the number of el ectrons transferred per mole of oxygen molecules reacted in the cell. 2.1 Fuel Cell Efficiency The thermal efficiency of a fuel cell, is defined as the amount of useful energy produced relatively to the change in stored chemical energy that is released when a fuel reacts with an oxidant. This can be expressed in terms of the ra tio of the operating cell voltage to the ideal cell voltage (Eq. 2-3). ideal cellV V (2-3) The actual cell voltage is less than the ideal one because of losses associated with cell polarization and ohmic loss. Moreov er a fuel cell can operate at different current densities. Thus, the corresponding cell voltage determines the fuel cell efficiency. To incr ease the theoretical cell voltage and power output, multiple cells are usually combined in series to form a ‘stack’ (Figure 2-2) with adjoining anodes and cathodes separated by an interconnect material. In order to develop a high performance fuel cell device, cell polar ization and ohmic loss should be minimized so that Vcell approaches Videal. Cell polarization, or overpotential, is the difference between theoretical and operating vo ltages. Multiple phenomena contribute to irreversible losses in an actual fuel cell, and the total cell polarization can be considered to be the sum of three individual cont ributions: ohmic polarization ( ), activation polarization ( A), and concentration polarization ( D).

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18 tot = + A + D (2-4) Ohmic polarization is caused by ohmic resistance in the electrode, in the electrolyte, and at the interface electrode/electrolyte. It varies dir ectly with the current in creasing over the entire range of current density. = R i (2-5) where R is the universal gas constant. The dominant ohmic losses in the cell come from the electrolyte. Therefore, ohimc losses can be reduced by decreasing the thickness and enhancing the ionic conductivity of the electrolyte. Activation polarization occurs whenever re acting chemical (including electrochemical) species are involved. It is a consequence of the activation energy barri er that needs to be overcome in order for the reaction to proceed. The activation polarization is directly related to the rate of the electrochemical reaction and it is dominant at low cu rrent density. The ButlerVolmer equation (Eq. 2-6) correlates the ac tivation polarization to the current density: RT F n esp RT F n i iA e A e 1 exp0 (2-6) where is the electron transfer coefficien t of the reaction at the electrode, i0 is the exchange current density and ne is referred to the number of el ectrons transferred per reaction, F is the Faraday constant and T is the temperature [K]. In the case of an activation polariza tion higher then 50-100 mV, the Butler-Volmer equation can be simplifie d by the Tafel equation: A = 0ln i i nF RT (2-7)

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19 Another form to express the Tafel equati on is described in the following equation (equation 2-8) where a = 0ln i nF RT and b = nF RT. A = a + b ln(i) (2-8) The term b is called the Tafel slope. In order to obtain a lower voltage drop with increasing current density, there is a strong in terest to develop elec trocatalysts that yiel d a lower Tafel slope for the electrochemical reactions. Concentration polarization occurs when reac ting species are supplie d to reaction sites slower than they are consumed, or when reactio n products are not removed fast enough so that they block the reaction sites. This effectively re sults in lower concentrat ions of reactant species at the reaction sites (lower oxygen partial pressures at the cathode or lower partial pressures of fuel at the anode). Gas transport losses occur ov er the entire range of current density, but they become prominent at high limiting currents. The limiting current (iL) is a measure of the maximum rate at which a reactant can be suppl ied to an electrode. Th e correlation between voltage drop due to concentration polarizati on and the current density is described by the following equation: D = Li i nF RT 1 ln (2-9) In this situation, the charge transfer reaction has such a high exchange current density that the activation polarization is negligible in co mparison with the concentration polarization. Figure 2-3 shows the decrease of the actual cel l potential as a function of the current density. As previously stated, th e activation polarization loss is dominant at low current density, while gas transport losses are prominent at hi gh current density. In order to minimize cell polarization losses, factors gove rning appropriate choice of materials (such as optimal

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20 electrocatalysts and more conductiv e electrolytes), and improvement in electrode structures or thinner cell components n eed to be considered. 2.2 Fuel Cell Components 2.2.1 The Electrolyte The electrolyte has the function of conducting oxygen ions from th e air side to the fuel side of the cell through a vacancy m echanism. Thus, a good solid elect rolyte should have high ionic conductivity and very low electr onic conductivity. Sometimes the total conductivity is expressed in terms of the transference number that is defined as: tot i it (2-10) It follows that tot ion tot elec tott t where tion is the ionic transf erence number and telec is the electronic transference number. For any material tion + telec = 1, therefore a good electrolyte should have tion 1. The state-of-art electrolyte for SOFCs is yt tria stabilized zirconia YSZ. Zirconia oxide ZrO2 has three modifications, monoclinic phase at low temperature, tetragonal phase at intermediate temperature and fluorite phase at high temperature. The high temperature fluorite phase shows a high degree of non-stoichiometry and a high concentration of oxygen vacancies. The introduction of 8 mol % of yttria into the material allows the stab ilization of the cubic fluorite structure at lower temperatures. Moreov er, the amount of the oxygen vacancy into the structure increases because of the different valence between the host (Zr+4) and the dopant (Y+3). This allows for obtaining a highe r ionic conductivity in the doped material. The general equation, in Kroger-Vink notation, is the following: O x O Zr ZrOV O Y O Y 2' 3 22 (2-11)

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21 This electrolyte is suitable fo r a high operation temperature (800oC-1000oC).9 It is commonly used in contact with LaCoO3 or LaMnO3-based cathode materials, which at these high operation temperatures react with the electrolyte forming an insulating phase at the cathodeelectrolyte interface. One strate gy to solve this problem and improve cell performance is to reduce the operating temperature. However, the ionic conductivity of YSZ drastically decreases with decreasing temperatures.10 Reducing the thickness of the electrolyte allows operation at lower temperatures, and good results at 700C were achieved using YSZ of 15 m thick.8 On the other hand, it has been recognized that for small SOFC stacks the ope rating temperature shoul d be further lowered without increasing the internal resistance of the cell. Thus, alternative el ectrolyte materials with higher ionic conductivity at lower temperatures ha ve been investigated.7,10 The ionic conductivity of ceria-based electrolytes is almost five times higher then that of zirconia solid solution at low to intermediate temperature.11 Formation of oxygen vacancies has been shown by doping the material with Gd2O3, Y2O3, CaO or Sm2O3.10 The higher ionic conductivity was obtained by dopi ng the material with Sm3+ or Gd3+.12,13 The ionic conductivity increases monotonically with increasing dopant content until it reaches a maximum. Further increase of dopant level forms defect cluster of the dopant ion with the oxygen vacancies which consequently decrease of oxygen vacancies mobility. A possible disadvantage in using ceria-based electrolytes is th e instability of ceria at low oxygen partial pressure, at the anode site. Indeed at low oxygen partial pressure Ce4+ is reduced to Ce3+ and the material exhibits n-type electron ic conductivity. The conductivity takes place by a small polaron transport in whic h the entire defect migrates by a thermally activated hopping.

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22 As a consequence of the reducing ionic conductivity of the electrolyte, the ope n circuit voltage of the cell becomes lower.15,16 Bismuth-based oxides exhibit the highest oxygen ion conductivity, which is almost 20 times higher then that of YSZ and 10 times higher then GDC at 800oC.14,15 Bi2O3 is stable at high temperature in a cubic defected fluorite structure where of the regular fluorite anion sites are vacant ( -Bi2O3). The high temperature stable defected fluorite structure can be stabilized at room temperature by doping with Ln2O3 oxides (Ln = La, Nd, Sm, Dy, Er, Yb, Y). Among them, erbium-stabilized bismuth oxide (ESB ) has the highest ionic conductivity.17,18 However, unlike YSZ, stabilized bismuth oxides are un stable under fuel atmosphere at pO2 10-2110-18 atm. Due to its high ionic conductivity, the ESB electrolyte has recently been proposed as a cathode side component in a new type of fast pure ionic bi-layered electrolyte, used for IT-SOFCs.15-18 A bilayer electrolyte can combine th e advantages of two electrolyte s and provide a solution to the stability issues inherent with exposure of the oxides to the reducing environment of fuel gasses. In particular, using a material with higher ioni c conductivity at th e cathode side provides a block to the electronic current improving open circuit vo ltage and increasing inte rfacial oxygen partial pressure. A comparison of the ioni c conductivity of severa l electrolytes is reported in Figure 2-4. 2.2.2 The Electrodes The cathode has the function of providing an el ectrochemical reaction site for the reduction of the oxygen gas molecules. The following is th e electrochemical reaction that occurs at the cathode site, using Kroger-Vink notation, in which oxygen vacancies are supplied by the electrolyte to allow the oxygen trans port and block electron migration. x O OO V e O 2 2 12 (2-12)

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23 The anode provides the electrochemical reactiv e site for the oxidation of the fuel gas molecule. e O H O H 4 22 2 2 (2-13) e CO O CO 4 22 2 (2-14) For proper function, an electrode material should have high electr o-catalytic activity towards oxygen reduction, for the cathode, or towa rds fuel oxidation for the anode. The electrode should be chemically stable while in contact with the electrolyte a nd the current collector materials, without forming resistive reaction produc ts at the interface. Furthermore, the material should exhibit similar thermo-mechanical properties as the electrolyte in or der to avoid stresses during the heating and the cooling processes. Finally, the electrode needs to have a high electrical conductivity. Noble metals and transition meta ls have been considered for anode. Pure metals have a strong tendency towards agglomerat ion and grain growth at elev ated temperatures. For this reason they are usually combined with a ceramic material, such as zirconia or ceria, forming a three interconnected framework of metal, ceramic and pores. In this pros pective the nickel-YSZ or nickel-GDC composite have been widely investigated as anodes.19-21 However, many studies have been focused on the development of new cathode materials in order to reduce polarization losses. Indeed, the do minant losses in the electrodes are activation polarization and concentration pol arization. Activation polarization is typically a problem in cathodes because the kinetics of the oxygen reducti on reaction is several orders of magnitude slower than the reactions involving fuel oxidation at the anode site.22-27 At the cathode, the dissociative adsorption of oxygen is a thermally ac tivated process, attrib utable to the high bond strength present in oxygen molecules.27 As the operating temperature is lowered, the rates of

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24 chemical reactions drop dramatically, and activat ion polarization at the cathode becomes even more significant. Therefore, to lower the polariza tion losses, there is a considerable interest in the development of new cathode materials th at will catalyze oxygen reduction reaction at moderate temperatures. 2.3 Type of Cathodes A conventional type of cathode is a porous single phase electronic conductor. Traditional studies have dealt with the kinetics of a nobl e-metal electrode (such as platinum) or an electronically conducting ceramic (such as LSM) on zirconia electrolyte. Because of the particular construction, the oxyge n reduction reaction is restricted to a one dimensional region where the three phases (gas, electrode and electr olyte) meet (a triple phase boundary, or TPB). Therefore, the gaseous oxygen must diffuse thr ough the cathode’s pores towards the reaction site for the reaction to occur. A possible mechanism of the oxygen reduction kine tics at an electroni c cathode is showed in Figure 2-5A and can be described in three main steps.28 At first, the gaseous oxygen molecules are dissociated and adsorbed on the cathode surf ace. Then, they diffuse along the pore walls until the interface electrode/electrolyte. Finally, at th e TPB, the charge transfer process occurs and oxygen ions are incorporated into the electrolyte. The charge transfer process that occurs at the electrode/electrolyte interface is usually expressed by the Butler-Volmer relationship. At certain conditions, this equation can be simplified into the Tafel equation, and a further simplification can be done if the current density is sufficiently low. In the last case the depe ndence of overpotential on current density can be treated ohmically, as report ed in equation (Eq. 2-15). i Rct A (2-15)

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25 In equation (Eq. 2-15) Rct is a charge transfer resistan ce at the electrode/electrolyte interface. Since the charge transf er process takes place at the TP B, several theoretical models related the Rct to the TPB geometry (TPB thickn ess, width and length per unit area of the solid electrolyte). Usually this resistance is normalized, resulting in the area spec ific resistance (ASR), which is calculated by multiplying the re sistance term by the electrode area. One method to accelerate the cathode kinetics is to use a catalyst. Examples of catalysts are noble metal and noble metal oxide powders (e.g., Pt, PdO, RuO2), which have been shown to have a beneficial effect in re ducing the cathode polarization in solid oxide fuel cells (SOFC).29,30 Significant enhancement of electrode performan ce and activity can be achieved by increasing the cathodic surface area. This is possible by introduci ng a nanoscale feature that results in a high surface area/volume ratio in the material. The fi ne microstructure permits an accelerated oxygen adsorption on the cathode surface and consequen tly increases the overall cathode kinetics. As a result, the presence of the nano-porous micr ostructure improved electrode performance.31,32 Another way to lower the cathode overpotential is to spread the reactive zone into the electrode. This is possible by using a porous co mposite cathode made of one ionic conductive phase and one electronic conductive phase. This has been studied and, indeed, the length and the density of the TPB increased because of the expa nsion of the region of contact between electrode and electrolyte.33-36 Taking in account this effect, Virkar et al 35 defined an effective charge transfer resistance that depends on microstructu ral parameters of the electrode (electrode thickness, porosity and particle sizes). ) 1 (2p L R RO ct eff ct (2-16)

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26 In equation (Eq. 2-16) L is the section length, p is the composite porosity and 2O is the ionic conductivity of the electr olyte. According to this model, the best performance of the electrode is obtained usi ng fine microstructure and a thick electrode, assuming that the porosity is large enough to prevent concentr ation polarization. In this same limiting step, model of Deseur et al 36 predicted that grading the composition and the reaction site s (linear variation of the volume fraction of the ionic conduc tor along the electrode thickness) is effective in increasing the electrochemical performance. Finally, a new class of material s that has being recently studi ed is mixed ionic electronic conductors (MIEC).24,25,37 The kinetics of the oxygen reduction at the surface of a mixed ionic electronic conductor (MIEC) versus the surface of a purely electroni c cathode is quite different (Figure 2-5B). In a MIEC, oxygen can be reduced at the surface and diffuse through the bulk of the electrode. Surface and bulk pathways are possibly competitive mechanisms for oxygen diffusion, and that one with the slower rate dete rmines the kinetics at th e electrode. Therefore, these materials should have high oxygen exchange capacity and high oxygen diffusivity for high transport rate. However, the mechanism of oxygen incorporation into MIEC electrode is still under question. Adler et al. 38 proposed a model for oxygen reduction on a porous mixed conducting electrode. According to this model the electrode resistivity is dominated by solid state diffusion and oxygen surface exchange, thus the effective chemical resistance Rchem is expressed as a function of the tortuosity, fracti onal porosity and internal surface area/unit volume of the porous mixed conductor. In conclusion, an understanding of the va rious mechanisms involved in the oxygen reduction reaction is crucial in order to im prove cathode performance. Only after the identification of the slowest (rate -determining) step, a viable stra tegy may be developed in order

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27 to accelerate it. Different cathode typologies, single phase electronic conductor, composite electrode and mixed ioni c electronic conductor, have been studied and their electrochemical behaviour has been related to the cathode geom etry and microstructure. Indeed polarization losses can be minimized by a proper choice of materials and by controlling the cathode microstructure. 2.4 Cathode Materials The state-of-the-art cathode material for SOFCs is La1-xSrxMnO3-x/2 (LSM) perovskite that is widely used on YSZ electrolyte. This material fulfills most of the requirements listed above. However, due to the high operating temperature, th e cathode reacts with th e electrolyte forming a new resistive product at the electrode/e lectrolyte interface, identified as La2Zr2O7.39 The electric conductivity of La2Zr2O7 was 4 to 7 orders of magnitude sm aller then that of LSM, and the electrochemical activity to wards oxygen reduction of La2Zr2O7 was negligible. This increases the ohmic losses and decreases th e performance of the cell. Properties of LSM can be tailored by partial substitution of the A and B site of the ABO3 perovskite. Materils with Ln1-xSrxMnO3 (with Ln=La,Pr,Nd,Sm,Gd,Yb,and Y) compositions have been tested on YSZ and on GDC electrolyte.40,41 The formation of the reaction products between the YSZ and the cathode can be suppressed for Ln1-xSrxMnO3 (Ln=Pr,Nd).40 In all cases, Pr substitution exhibits the best performa nce, improving the electri cal conductivity and the catalytic activity of the material. Materials with the following composition Ln1-xSrxMnO3 (Ln= Gd, Pr, Nd) are also compatible with GDC because they have a good chemi cal stability with the electrolyte and a good thermal expansion coefficient match.41 However, the best performances on ceria based electrolyt e were obtained using La1-xSrxCoyFe1-yO3cathodes.42-44 These materials are mixed ionic and electronic conducto rs, thus the activity of the cathode increases

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28 due to an additional pathway for the oxygen i ons through the bulk of the cathode material. Depending on the composition, the conductivity of La1-xSrxCoyFe1-yO3can vary of about one order of magnitude.42 Performance of La0.6Sr0.4Co0.2Fe0.8O3 symmetrical cell showed a cathodic electrochemical resistance of roughly 0.6 cm2 at 650oC, which suggests that the system is limited by the kinetics of the LSCF cathode. Diffe rent strategies were used to further improve the performance, such incorporation of precious metal into the compound and microstructura l optimization of the electrode.43 Finally, fabrication of a composite made of LSCF and GDC was a good strategy to improve cathode performance at in termediate temperature (500-700oC), because it allowed an enhancement of the ionic conductivity of the electrode at low temperature.44 Again, cathode performance can be improved by substituting one or more of the elements. Recently a new cathode material with the following chemical composition Ba0.5Sr0.5Co0.8Fe0.2O3 (BSCF) has been studied for IT-SOFCs. Shao and Haile45 tested this material on thin film of SDC electrolyte, and it exhibits high power density when ope rated with high humidified hydrogen as a fuel and air as the cathode gas (1010 mW/cm2 at 600oC and 402 mW/cm2 at 500oC). The ASR of BSCF cathode, determined by sy mmetrical cells in air, was substantially lower then the other single phase perovskite cathode. Better results we re obtained on thinner GDC electrolyte (10 m) by Liu and Chan.46 The high performance of this cathode was ascribed to the high rate of oxygen diffusion through the material. To further improve the cell performance at reduced operating temp eratures, cobaltite electrodes have been applied on bismuth oxide based electrolytes, which exhibit an higher oxygen ion conductivity compared to ceria based electrolytes.14,15 However, diffusion of bismuth and cobalt cations occurred at the electrode/e lectrolyte interface, causing changes of the

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29 chemical and phase composition of the materials used.47 The interaction between these compounds was very fast, and a drop of el ectrical conductivity and oxygen permeability was revealed in the reaction phase.47 Therefore, alternatives cathode materials need be studied for application on bismuth based electrolytes for IT-SOFCs. 2.5 Summary This general review provides an essential framework from which this study could be built. To develop a high performance fuel cell, cell pol arization and ohmic loss es should be minimized. Ohmic losses can be decreased by a proper choice of electrolyte material and thickness. For instance, bismuth based electrolyt es presented the highest ionic conductivity at low temperature. On the other hand, the main contribution to the cell polarization co mes from the cathode component, because the oxygen reduction kinetics is not thermodynamically favored and the interfacial polarization increases rapidly with decreas ing temperature. In this prospective, there is a considerable interest in finding alternative materials that can be used as cathode for IT-SOFC.

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30 Figure 2-1. Solid oxide fuel cell. (Source: http://www.thirdorbitpower.com/SOFC_mech.html Last accessed September, 2005) Figure 2-2. An SOFC stack.

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31 Figure 2-3. Typical current-voltage profile for SOFC. (reprinted fr om Solid State Ionics Vol 131, Virkar, J. Chen, C.W. Tanner, and J.W. Kim, 189-198, (2000)). Figure 2-4. Arrhenius plot for common solid electr olyte materials. (reprinted from Ceramic in Advanced Energy Technology, B.C.H. Steele, Ceramic materials for electroceramical conversion devices, D. Reidel, Dordrecht, (1984), pp386-412).

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32 Figure 2-5.Various reaction pathways for the ox ygen reduction reaction. A) In a single phase electrode. B) In a MIEC electrode. (Figures adapted from J. Electrochem. Soc., Vol. 153, B. Kenney and K. Karan, 6, A1172 (2006) and Solid State Ionics. Vol. 135, S.B. Adler, 603 (2000)). Electrode Electrolyte TPB Electrode TPB Electrolyte B A

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33 CHAPTER 3 PYROCHLORE RUTHENATES Physical properties of a crysta lline material, such as electr ical conductivity, are related to its chemical composition and crystallographi c structure. A working knowledge of the crystallographic structure needs to be obtained in order to design a material for a specific property. The following section reviews the pyr ochlore structure of ruthenates and the correlation between the pyrochlore stru cture and its elec tronic conductivity. 3.1 Pyrochlore Structure The general formula of pyroch lores can be written as A2B2O6O`, which includes four crystallographic nonequivalent kinds of atoms. The space group of the ideal pyrochlore structure is Fd 3m (No 227) and there are eight molecules per unit cell (Z=8) with a typically lattice value of ~10 The A cation usually has an ionic radius of ~1 and it is eight fold coordinated. The B cation is smaller (~0.6 ionic radius) and it is six fold coordinated. Arbitrarily decided, the origin can be chosen to be either th e A or the B cation. Typically for ruthenate pyrochlores, the origin is chosen such that the larger A cations occupy the 16(d) sites at (, ) and are loca ted within a scalenohedra (disto rted cube) that contained six equally space O anions and two slightly shorter distanced O` anions. The smaller B cations occupy the 16 (c) sites at (0, 0, 0), and they are bonded to six O anions at equal distance, forming a trigonal antiprism. The oxygen anions, O and O` occupy the 48(f) sites (x, 1/8, 1/8) and the 8(b) sites (3/8, 3/8, 3/8) respectively.48 In order to completely describe the three-dimensional arrangement of ions in the structure, only one positional parameter must be determin ed. In fact, the coordination polyhedra can change shape with the oxygen parameter, x. The limiting values of x are 0.3125 and 0.375. At x = 0.3125, the B ion has a perfect octahedral coordi nation and the A cations are arranged in a way

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34 to form a distorted hexagon of six oxygen (48f). At x = 0.375, the A cations are situated in a regular cube while the B cations are in the centr e of highly distorted octa hedral. Practically, the A-O distance decrease rapidly as x increase a nd when x=0.375 it becomes equal to A-O` distance that is, independent of oxygen x parameter. On the other hand, the B-O-B angle, which is 109o28` for the fluorite, increases to 120-140o for the pyrochlore structure.48 A generalized description of the pyroc hlore structure usually considers two interpenetrating networks, one corner sharing B2O6 octahedra and the other anticrystobalite A2O` chains. The distorted BO6 octahedra share corners to form a tetrahedral lattice of the formula B2O6. The A cations are present at the center of the hexagonal rings of oxygen (formed from six BO6 octahedra) with two more oxyge ns (O`) located above and below the ring. Each O` atom is tetrahedrally coordinated with A cations, and the A and O` atoms form a tetrahedral lattice of the formula A2O`. The B2O6 and A2O` lattices interpenetrate, and the B and A cations interact through B–O–A linkages. Each O is surrounded with two B and two A cations in a distorted tetrahedron (local C 2v symmetry). The pyrochlore structure is clos ely related to the fluorite stru cture and can be described as an anion-deficient, vacancy-ordered fluorite. The fluorite structure is cubic (Fm-3m) with 4 molecules per unit cell (Z=4) and a typically lattice value of ~5 where cations are located in a face centered cubic arrangement and anions are located in the tetrah edral sites. In the pyrochlore structure the cations A and B form a face centered cubic array and the ani ons are located in the tetrahedral interstices of the cationic array. He nce, we can divided the pyrochlore unit cell in eight smaller cubes and consider th e pyrochlore structure in terms of defect fluorite lattice, where the 8a positions are vacant. Figur e 3-1 showed an octave of the unit cell of the pyrochlore

PAGE 35

35 structure, in which there are indicated the f our occupied crystallographic positions and the unoccupied 8a position. The A2O` sublattice is not essential for the stab ility of the structur e and vacancy of both the A-type cations and O` anions are common. If the distincti on between the two type of oxygen atoms O and O` is not made, the formula can be written as A2B2O7-y. The structure of a defected pyrochlore can be obtained in terms of def ected-fluorite lattice by removing 8b oxygens (O`) from the unit cell of the stoichiometric pyrochlore structure (A2B2O7). After the removal of 8b oxygens, the A cations are exposed to each ot her across the 8b vacancy and the resulting electrostatic repulsion would tend to destabilize the structure. Hen ce, the stability of defected pyrochlore is mainly due to the bonding between A cations through the ox ygen. The polarization of the A cations by the oxygen vacancy can result in the stabilization of the A-A bond. When this happens, a defect pyrochlore struct ure can be stabilized in preference to the perovskite. A and B cations with large and small ionic radius, respec tively, may stabilize the pyrochlore structure as they give an x value nearer to the ideal value of 0.3125 (0.30-0.32). 3.1.1 Existence and Stability Field Subramanian et al 48 have shown that the existence and stability field of pyrochlores is governed by the relative ionic radius rA 3+/rB 4+ and the oxygen parameter x. A3+ 2B4+ 2O7 pyrochlores are typically stabilized for 1.46< rA 3+/rB 4+<1.80 whereas the range for A2+ 2B5+ 2O7 pyrochlores is 1.4< rA/rB<2.2. Since a large number of A3+ and B4+ cations have suitable ionic radius for the formation of the pyrochlore struct ure, many of the pyrochlore oxides known in the literature are the (3+,4+) type A3+ 2B4+ 2O7. The A3+ ion can be a rare earth, such as Sc, Y, Bi, Tl, or In, or lanthanide elements whereas B4+ ion can be a transition me tal or any of the group IVa elements (Si, Ge, Sn, Pb). Based on these considerations, Y3+ 2Ru4+ 2O7 Pr3+ 2Ru4+ 2O7 should be stable compounds; therefore their propertie s have been studied in this work.

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36 3.1.2 Crystal Data of Y2Ru2O7 by Neutron Powder Diffraction In this specific thesis, yttr ium ruthenate and doped yttrium ru thenate have been synthesized and their electrical and electrochemical properti es have been studied for application in ITSOFCs. Thus, information on the crystal structure of Y2Ru2O7 were used as initial data for this study. The structure of the pyrochlore yttriu m ruthenium oxide was determined by Neutron Powder Diffraction, by J. Kennedy in the 1995.49 The material adopts a regular pyrochlore structure and no vacancy ordering on the O` site was revealed by the X-ray diffraction data. The fractional atomic coordinates and the anisotr opic displacement parameters are shown in the following table (Table 3-1). The lattice value of the cubic cell is a = 10.1429 (2) and the value of the bond distance are as follows: Y-O = 2.4503 (2) Y-O` = 2.19601 (3) Ru-O = 1.9911 (1) The value of the O-Ru-O angle is equal to 128.45 (2)o. All these data are used in the PowderCell 2.4 software in order to analyze the Y2Ru2O7 crystal structure, as shown in Figure 3-2. 3.2 Electrical Properties A good understanding of the pyrochlore structure is very important becau se it is related to the electronic properties of the material. Pyroch lores electrical conductivit y can be explained in terms of the Mott-Hubbard mechanism of electron localization. The width of the t2g-block bands of A2Ru2O7-y increases with increasing Ru-O-Ru bond angle and decreasing Ru-O distance. When Ru-O-Ru angle is larger than 133o, and the bond distance Ru-O is less than 2, the bandwidth becomes greater than a critical value and A2Ru2O7-y shows metallic behaviour.50-53

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37 The correlation between the structure and th e electronic property b ecomes clear studying the metal-versus-semiconductor behavior. Indeed, the conduction change from semiconductor to metal is consistent with the incr ease in the overlap between the Ru4d t2g and O2p orbitals. K.S. Lee et al 53 have shown that there is a good linear relationship between the ionic radius of the A cation and the Ru-O-Ru bond angle (Figure 3.3). Increasing the size of the A cation, the Ru-O bond length shortens and the Ru-O-Ru angle increases. In some cases pyrochlores exhibit ionic conducti vity. As it was previously mentioned, the oxygen vacancies in the A2O` network influence the ionic properties of pyrochlores. The conduction of the oxygen ion in pyrochlores proceeds via an oxygen vacancy mechanism similarly to fluorites. Pirzada et al 54 used an atomic scale computer simulation in order to predict activation energi es for oxygen migration. They repor ted that pyrochlores with low B cation radii (Ti, Ru) have the highest activation energy for migr ation of oxygen ions in the 48f48f pathway. Therefore, ionic conduction in ruthenates can be difficult to observe. Doping has an effect of charge carrying capac ity in the pyrochlore system. Yoshii et al,55 reported electronic and magnetic properties of A2-xCaxRu2O7 (A = Y,Sm). Calcium (Ca2+) had the function of increasin g the main charge carries, which are holes. Therefore, upon increasing Ca2+ concentration, x, the system exhibited a tran sition to a metallic stat e (Mott transition). For an R atom with larger ionic radius, the cr itical concentration of the transition, xc, was smaller. Extrinsic doping can also enhance ionic conducti vity. It has been shown that incorporating cations such as Ca2+, Mg2+, Sr2+ into the A site of titanate pyrochlores create compensating defects and likely results in a change of the stoichiometry of the pyrochlore.56,57 Moreover, the increase in ionic conductivity can be enhanced by substitution on the B site of pyrochlore structure. Pirzada et al 58 proposed different doping mechan isms depending if the dopant Me2+

PAGE 38

38 was distributed evenly in the A and B sites or favored one site over the other. In each case, oxygen vacancies compensated the solution m echanism and improved oxygen ion conduction. 3.3 Summary Electrical properties of ruth enates are related to their chemical composition and crystallographic structure. Incorp oration of aliovalent cations in fact increases the charge carrying capacity of the material and creates comp ensating defects that likely result in a change of the stoichiometry of the pyrochlore. Moreover, there is a linear relationship between the ionic radius of the A cation and the Ru-O-Ru bond angle, which determines the electrical properties of ruthenates. Thus, doping the A site is a good strate gy to improve the electrical behavior of these materials.

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39 Figure 3-1. One octave of the unit ce ll in the pyrochlore structure. Table 3-1. Fractional atomic coordinates and anisotropic displacement parameters (2 x 103). (reprinted from Acta Cryst. C 51, B. J. Kennedy, Structure Refinement of Y2Ru2O7 by Neutron Powder Df fraction, 790 (1995). Site x y = z U11 U22 = U33 U12 = UI3 U23 Y 16(d) 1/2 1/2 4.51 (6) = U11 -0.97 (7) = U12 Ru 16(c) 0 0 2.33 (7) = U11 0.08 (7) = U12 O’ 48(f) 0.33536(3) 1/8 5.35 (9) = U11 0.0 1.64(10) O 8(b) 3/8 3/8 3.40(19) = U11 0.0 0.0

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40 Figure 3-2. The Y2Ru2O7 crystal structure. Figure 3-3. Relationship between the ionic radius of the A cation and the Ru-O-Ru bond angle ( reprinted from J. Solid State Chem, Vol 131, K.S. Lee, D.K. Seo and M-H.Whangbo, Structural and Electronic Factors Governi ng the Metallic and No nmetallic Properties of the Pyrochlores A2Ru2O7-y, 2 (1997)).

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41 CHAPTER 4 SYNTHESIS OF NOVEL YTTRIUM PRASEODYMIUM RUTHENATES 4.1 Introduction Pyrochlore ruthenates have been widely studied for applications as catalysts,59 electrocatalysts 60,61 and conductive components in thick film resistors.62,63 The catalytic activity is believed to be related to the variable forms of oxidation states available for ruthenium. These include Ru4+/Ru5+, Ru5+/Ru6+, and Ru4+/Ru6+ transitions which cause modifications of electron transfer rates in the oxygen electron reaction.60 These compounds have been recently investigated for cathodic applicati on in solid oxide fuel cells (SOFCs) because they exhibit high electrocatalytic beha vior for oxygen reduction.59,60,64-69 Among them, yttium ruthenium oxide Y2Ru2O7 was evaluated as a possible candidate for cathodic application in IT-SOFC because it showed to be chemical stable in contact with yttrium stabilized zirconia (YSZ) and gadoliniumdoped ceria (GDC) electrolytes.65 However the cathodic performance of Y2Ru2O7 was limited by a low electrical conductivity related to the n-type semic onductive behavior. Bae & Steel65 were able to considerately enhance the performance of this pyrochlore on GDC electrolyte by doping the material with SrO. Those results suggested that the behaviour of the Y2Ru2O7 cathode can be further improved by a proper choice of dopant. In this work, novel yttrium-ruthenates compositions doped with Pr (Y2-xPrxRu2O7, 0 x 2) were synthesized by a soft precipitation method. This chemi cal method was used to control the chemical composition and the final morphology of the materials. N,N,N’,N’ Tetramethyllendiamine (TMEDA) was used as th e precipitating agent to form the metal hydroxide precursor. Subsequent he ating treatments were used to purify and crystallize the precursor in a nanosized shape. Pr element wa s introduced as a dopant in the A-site of the

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42 pyrochlore structure in several contents. The dopan t modified the crystalli ne structure and it may consequently allow enhanced ma terial electrical proprieties. 4.2 Experimental 4.2.1 Materials Preparation All the starting chemicals were reagent-grade and were used with no further purification. Y(NO3)3•6H2O (Alfa Aesar), Pr(NO3)3•6H2O (Alfa Aesar) and RuCl3•xH2O (Alfa Aesar) were used as starting materials. A co-precipitation method was used to prepare the pyrochlore powder Y2-xPrxRu2O7 (0 x 2). In a typical reac tion synthesis 0.77 g Y(NO3)3•6H2O were dissolved in de-ionized water in 0.1 M c oncentration, 0.87 g of Pr(NO3)3•6H2O were dissolved in de-ionized water in 0.1 M concentration and 4.45 mmol of RuCl3•xH2O was dissolved in 40 ml of deionized water. The three prepared solutions we re mixed together and 7 ml of N N N’ N’ Tetramethylethylendiamine (TMEDA, Aldrich 99.5%) was added. The co-precipitation of Y, Pr, and Ru hydroxides was instantaneous at room temp erature. The obtained pr ecipitate was left to settle and age for one day, then filtered and washed several time in de-ionized water. Finally the resulting dark gel was dried at 90oC for two hours, and then calcined at 500oC for two hours and at 1050oC for five hours. A final yield of 100% product (2 mmol) of YPrRu2O7 was obtained. The process was performed va rying the quantities of Y(NO3)3•6H2O and Pr(NO3)3•6H2O to obtain pyrochlore powders with different compos ition. In all chemical processes, the same quantity of TMEDA (7 ml in 80 ml of solu tion) was used to induce the instantaneous precipitation of the metal hydroxide at room temperature. At the end Y2-xPrxRu2O7 with six different compositions (x = 0, 0.2, 0.5, 1, 1.5, 2) were obtained.

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43 4.2.2 Materials Characterization Thermogravimetric and differential therma l analysis (TG/DTA) were performed on precursors using a STA 409 Netzsch instru ment. All samples were heated at 5 oC/min from 100oC to 1200oC in air. The crystal structure of the powders was id entified by X-ray diffraction (XRD) analysis, using X-ray diffractometers (APD-3720 a nd a X-Pert 1900, Philips) with a Cu K source radiation ( 1 = 1.54056 and 2 = 1.54439, Ratio 21 = 0.50000) and a graphite monochromator on the diffracted beam. The inst rumental parameters were: divergence and scatter slits = 1o receiving slit = 0.1 mm, voltage applie d to the X-ray tube = 40 kV, electric current = 40 mA. A -2 step-scan technique was used at a 0.015o, internal in 2 and with a fixed counting time of 1 second in the range 20o 2 80o. X-Pert Data Collector program was used for data collection. Microstructural parameters were obtained from a whole profile XRD Rietveld fitting 70. Powder morphology was determined by field emission scanning electron microscopy (FESEM, LEO SUPRA 1250). The FE-SEM images were obtained in secondary and in lens electron modes. X-ray photoelectron spectroscopy (XPS, Perk in-Elmer PHI 5100) measurements were performed using Al source in order to determ ine the oxidation state of Pr and Ru. The XPS measurements were carried out on the sample as prepared and after sputtering the surface with Ar+ ions for ten minutes. The XPS data were tr eated by a Shirley-type background subtraction and fitted with Gaussian –Lorentzian function.

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44 4.3 Results and Discussion In this synthetic route, TMEDA plays a crucial role in allowing the simultaneous precipitation of yttrium, praseodymium and ruth enium at room temperature and stabilizing the precipitate itself. TMEDA is in fact a chelating-type diamine that forms stable five-membered rings with metal ions due to its two coordinating positions.71,72 The precipitation leads to the formation of stable gel precursor which contains oxo/hydro metals in a prope r stochimetric ratio. Subsequent heating treatments allow the purifi cation from the organics by-products and the formation of a single pyrochlore cristalline phase. TG/DTA technique can be used to determin e the crystallization temperature of the ruthenate. Figure 4-1 shows TG/DTA analysis on Y2Ru2O7 precursor. Two different exothermic peaks and a broad exothermic shoulder, together with a continuous weight loss can be observed. The first intense peak with its maximum at around 170oC, associated with a weight loss of about 20 wt%, can be attributed to the oxidative d ecomposition of the precurs ors, while the broad shoulder centered at about 500C can be associated with combustion of the organic residues. No detectable weight loss was associated with the second smaller exothermic peak with its maximum at about 980oC, which thus can be attributed to the crystallization of the pyrochlore phase. It is worth noting that no weight loss due to element vol atilization was observed with the temperature, confirming the stability of the pyr ochlore phase up to 1200oC. Figure 4-2 shows XRD spectra of Y2Ru2O7 precursor after treatments at different temperatures for 5h each. The formation of the pyrochlore crystalline phase was obtained after thermal treatment at 1050oC. The crystalline phase did not change after thermal treatment at 1200oC, but changes were revealed after thermal treatment at 1400oC. Y2-xPrxRu2O7 with six different compositions (x = 0, 0.2, 0.5, 1, 1.5, 2) were synthesized by this soft precipitation method. In all chemical processes, the same quantity of TMEDA (7 ml

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45 in 80 ml of solution) was sufficient to induce the instantaneous precipitation of the metal hydroxide at room temperature. Figure 4-3 shows the XRD patterns and (h k l) Miller’s indices of Y2-xPrxRu2O7 (x = 0, 0.2, 0.5, 1, 1.5, 2) after calcinations at 1050 C. All samples showed a single pyrochlore crystalline phase (cubic phase, space group Fd 3m). Increasing the Pr amount, the XRD peaks shifted towards lower angles. A change in the lattice size as a consequence of Pr introduction into the structure was expected because of the larger ionic size of Pr compared to Y.73 Rietveld analysis70 of the full XRD patterns provided additional information about the crystal structures of the ruthenates with di fferent compositions. Figur e 4-4 shows the X-Ray diffraction refinement profile of Y1.5Pr0.5Ru2O7. No important variances from the expected stoichiometry were observed after refining occupanc ies of the Y, Pr, Ru and O sites. The lattice parameters of the va rious solid solution Y2-xPrxRu2O7 (x = 0 2) were obtained from the powder X-ray diffraction data, as listed in table 4-1. Figure 4-5 shows the variation of the lattice parameter with the Pr content. According to the Vegard’s law, the lattice size linearly increased as the replacement of the Y with Pr occurs in Y2-xPrxRu2O7 compositions. The crystallite size also increased with increasing the Pr content. Figure 4-6 shows the SEM micrographs of the Y1.5Pr0.5Ru2O7 and YPrRu2O7 powders calcined at 1050oC for 5 hours. FE-SEM observations showed that Y2Ru2O7 powder was made of nanocrystalline particles of about 100 nm average size. The average particle size was kept at the nanometer scale also for Y1.8Pr0.2Ru2O7 and Y1.5Pr0.5Ru2O7 (Figure 4-6A) compositions. However, for Pr content larger than 25 mol% average particle size became larger and the particle size distribution became wider. The average particle size for the YPrRu2O7 sample (50 mol% of Pr) was evaluated to be around 250 nm (Figure 4-6B), while Pr2Ru2O7 presented a

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46 particle size distributi on between 500 nm and 2 m. Figure 4-7 shows SEM images of Y1.8Pr0.2Ru2O7, YPrRu2O7 and Y0.5Pr1.5Ru2O7 and Pr2Ru2O7 at high magnification. All samples presented octahedral crystals, where the extern al faces are only partially developed in { 1 1 1} planes with subordinated {011} planes. This morphol ogy was expected since th e crystal habits of the pyrochlore structure typically include octahe dral crystals eventually modified by other isometric forms. XPS measurements were used to investigate the oxidation state of Pr and Ru. Figure 4-8 showed the Pr 3d core spectra of Y1.8Pr0.2Ru2O7 before and after the Ar+ ions sputtering for ten minutes. A spin-orbit-splitting energy of 20.8 eV was found between the Pr 3d5/2 and Pr 3d3/2 lines. The intensity ratio betw een the two spin-orbit component s is in agreement with the expected statistical ra tio of 2/3. The Pr 3d5/2 spectra, at ~ 933 eV, presented one main peak and one shoulder separated by a binding energy of 4.4 eV. In Figure 4-8A the analysis of praseodymium oxide is unresolved because diff erent oxidation states are difficult to be distinguished. In order to better separate the single c ontribution, the sample was sputtered with Ar+ ions for ten minutes. Figure 48B shows XPS spectra after Ar+ sputtering where the pattern shows a well resolved shoulder at 928. 75 eV, which is characteristic of Pr2O3.74 Previous authors investigated Pr 3d and 4d core level by XPS technique.74-76 Based on what is reported by Lutkehoff,76 the shape and position of XPS peaks in Figure 4.8A are characteristic of Pr6O11 composition. The best fit was obtai ned considering the contribution of four components, in which the line pairs at hi gher binding energy can be associated to Pr4+ (~ 935, ~ 931 eV), while the line pairs at lower bi nding energy (~ 933, ~ 928 eV) can be associated to Pr3+. In conclusion, the Y1.8Pr0.2Ru2O7 sample exhibited mixed Pr oxidation state where the

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47 Pr3+ signal is almost twice the Pr4+ signal. Similar results were observed for the other compositions. Figure 4-9 shows the Ru 3d core spectrum of Y1.8Pr0.2Ru2O7 sample as prepared and after sputtering with Ar+. The curve fitting of the as prepar ed sample (figure 4-9A) showed two different Ru 3d doublets, with Ru 3d5/2 lines at 281.75 and 282.82 eV. In the literature there are different interpretations for these results: th e second doublet may indi cate the presence of a second oxidation state, either Ru(V) or Ru(VI) as reported for Bi2Ru2O7+y,77 or it may arise from final state effects.78 In this second case, the interacti on of a positive hole created in the photoemission process with mobile conductive electrons gives rise to asymmetry in the core-line of X-ray photoemission spectra. This has been proved for several metals.79 Sputtering with Ar+ ions is a common technique in order to eliminate impurities on the surface. However, in this case, the Ru peak is still partially obscure d by the C1s line, probably because of surface roughness of the sample. Therefore, in a comp arison between spectra A and B, the carbon line needs to be considered. Sputtering with Ar+ ions induced a partial reduction of the element, as it is shown in Figure 4-9B. The Ru 3d spectra show ed a shift towards lower binding energy and an increased broadness of the peaks. This curve can be fitted using an additional doublet with Ru 3d5/2 lines at 280.25 eV, which corresponds to Ru (III). In summary, XPS data showed the mixed oxida tion state of Pr. Theoretically, the different valences of Pr may yield to two different product s, where Pr may occupy the A site or the B site of the pyrochlore structure. However, Subramamian et al.48 have shown that the formation and stability of the pyrochlore structure is governed by the relative ionic radius rA 3+/rB 4+. The ionic radius of Pr+4 in sixfold coordinated state (B site) would be too large to allow the formation of a stable pyrochlore structure. Ther efore, Pr must occupy the A-site of the pyrochlore structure,

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48 according to the Rietveld analysis results previous ly shown. In addition, XPS analysis showed a quick change of the elements mixed oxidation stat es in response to the e nvironmental conditions. This effect is confirmed by the multivalent characte r of Ru which enforce the ruthenates catalytic activity towards oxygen reduction.59,60 In this prospective, the mi xed oxidation state of both Pr and Ru can be an advantage because it may enha nce the catalytic activity of the material. 4.4 Conclusions Y1-xPrxRu2O7 (0 x 2) pyrochlore oxide powders we re obtained in a variety of compositions using a soft chemistry method. This chemical route is effective in controlling the composition and the morphology of the final products. Each composition is a single pyrochlore phase stable up to 1200oC. The lattice size of the pyrochlores increased linearly with the Pr content, following the Vegard’s law. Nanoc rystalline particles were obtained for x 0.5, while larger concentrations of Pr led to micrometer particle sizes. Finally, both Pr and Ru showed mixed oxidation states and can easily change thei r oxidation state in response to environmental changes.

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49 Figure 4-1. Simultaneous TG/DTA analysis on Y2Ru2O7 precursor. Figure 4-2. The XRD pattern of Y2Ru2O7 after different thermal treatments for 5h each.

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50 Figure 4-3. The XRD pattern of Y1-xPrxRu2O7 (x = 0 2) powders af ter crystallization at 1050 C. Figure 4-4. Observed, calculat ed and difference diffraction pa tterns for the X-Ray data for Y1.5Pr0.5Ru2O7 powder.

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51 Table 4-1: Lattice parameters for Y1-x PrxRu2O7 (x = 0 2) from X-ray data. Pr content a ( ) error 0 10.1587 0.0005 0.2 10.2115 0.0005 0.5 10.2218 0.0003 1 10.2945 0.0004 1.5 10.3359 0.0004 2 10.3753 0.0005 Figure 4-5. Composition dependence of the lattice parameters for Y1-xPrxRu2O7 (x = 0 2).

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52 Figure 4-6. The FE-SEM pictures after crystallization at 1050oC. A) Y1.5Pr0.5Ru2O7, B) YPrRu2O7.

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53 Figure 4-7. High magnification FE-SEM pi ctures after crystallization at 1050oC. A) Y1.8Pr0.2Ru2O7, B) YPrRu2O7, C) Y0.5Pr1.5Ru2O7 and D) Pr2Ru2O7 particles.

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54 Figure 4-8. The XPS Pr 3d core level spectra of Y1.8Pr0.2Ru2O7. A) As prepared. B) After sputtering with Ar+ for 10 min.

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55 Figure 4-9. The XPS Ru 3d core level spectra of Y1.8Pr0.2Ru2O7. A) As prepared. B) After sputtering with Ar for 10 min.

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56 CHAPTER 5 YTTRIUM PRASEODYMIUM RUTHE NATE ELECTRICAL PROPERTIES 5.1 Introduction As it was reported in chapter 3, ruthenate electrical conductivity can vary from metallic (A = Bi, Pb, Tl, etc.) to semiconducting (A = Y, Nd, Pr, etc.) depending on the A cations. The electrical properties can be e xplained in terms of the Mott-Hubbard mechanism of electron localization. The width of the t2g-block bands of A2Ru2O7-y increases with increasing Ru-O-Ru bond angle and decreasing Ru-O distance. When Ru-O-Ru angle is larger than 133 and the bond distance Ru-O is less than 2 the bandwidth becomes greater than a critical value and A2Ru2O7y shows metallic behavior. 50-53 In this study the elect rical properties of Y2-xPrxRu2O7 (x = 0, 0.2, 0.5, 1, 1.5, 2) were investigated as a function of Pr content. The introduction of Pr into the A-site increased the Asite ionic radius withou t destabilizing the pyrochlore structure. As consequence of the structural change, the overlap between the Ru 4d t2g and O 2p orbitals should increase the material electrical conductivity.50-53 Moreover, the multivalent character of Pr would allow the formation of additional holes, which are the majo r charge carrier of this material.55-58 Therefore, the electrical co nductivity of Y2-xPrxRu2O7 was investigated at differe nt oxygen partial pressures and temperatures. 5.2 Experimental Y2-xPrxRu2O7 (x = 0, 0.1, 0.5, 1, 1.5, 2) powders were isostatically pressed at 200 MPa in a bar shape and sintered for 15 hours at 1100 C. The sintered bars were approximately 14 x 3.13 x 1.6 mm. D.c. 4-probe method was used to meas ure the electrical conductivity of the sintered bars. The data were collected in the temperature range of 300 800 C using a multimeter (2400 series, Keithley) and using plat inum-paste and platinum wires for the contacts. Oxygen partial

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57 pressure varies in the range of pO2 = 10-5 1 atm, controlled by mixtures of oxygen, air and nitrogen gas. The el ectrical conductivity was measured as function of Pr content, oxygen partial pressure and temperature. The relative density of the rectangular bar sa mples increased with increasing Pr content, the values being estimated to be 71, 76, 75, 81, 93 % respectively. Therefore, the raw conductivity must be corrected for a meaningful comparison with data for dense materials. This correction was estimated using the Bruggeman me dium theory, which evaluated the electrical conductivity in biphasic materials having high conductive and lo w conductive phases interconnected in 3D.80 In this specific case, pores were considered the low conductive phase, and the electrical conductivity was co rrected using Equation 5-1, where m is the conductivity accounting for the porosity of the material, h is the conductivity of the high conductive phase and f is the volume fraction of the low conductive pha se, which in this case corresponds to the porosity of the material. 2 / 3) 1 ( fh m (5-1) The corrected values results to be increas ed of almost 1.5 times for the 75% dense sample. It should be stressed that temperature dependence, pO2 dependence and activation energies are not affected by this procedure. 5.3 Results and Discussion The electrical conductivity of Y2-xPrxRu2O7 (0 x 2) was measured as a function of the Pr content in the temperature range of 250-800 C using the dc 4-probe method. Figure 5-1 shows the variation of the electrical c onductivity with Pr content. At 700 C, the electrical conductivity ( ) increased from 37 S/cm to 58 S/cm to 74 S/cm for samples with x = 0, 0.2 and 0.5 respectively, reaching the value of 171 S/cm for Pr2Ru2O7. Experimental data were linearly fitted

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58 using equation 5-2. The conductivity dependence on Pr concentration has a slope of 1.30 at 750oC, 1.29 at 700oC, 1.16 at 600oC up to 1.12 and 0.95 at 500oC and 400oC respectively. Therefore, the variation of with Pr increased with increasing temperatures. [Pr] B A (5-2) Figure 5-2 shows the temperature de pendence of the conductivity of Y2-xPrxRu2O7 (0 x 2). All the samples showed a semiconducti ng behavior described by an increase of with increasing temperature. The Y2Ru2O7 and Pr2Ru2O7 Arrhenius plots are in agreement with those reported in the literature. 51,81As the Pr amount increased, the activation energy for conductivity decreased. Therefore, the introduction of Pr into the pyrochlore structure enhanced the electrically conductive mechanism. The electrical conductivity increase may be due to a risi ng charge–carrying capacity of the material as a consequence of the variable oxi dation state introduced by the dopant. In order to investigate this, the Y2-xPrxRu2O7 (0 x 2) electrical conductivity was tested as a function of the oxygen partial pressure, as repor ted in Figure 5-3. This experi ment was used to evaluate a possible variation of the majority defects concentration with external conditions (pO2). However, these materials were found to be unsta ble at low oxygen partial pressure (pO2 = 10-5 atm) and decomposed. For this reason, a narrow range of oxygen partial pressure s was chosen for the measurements. As result, no a ppreciable conductivity vs pO2 dependence was observed in this oxygen partial pressure range. Therefore, the el ectrical conductivity is dominated by p-type behaviour and its electron defects (holes) resulted to be independent of pO2 within this oxygen partial pressure range. The same result wa s obtained with different samples composition (Y1.8Pr0.2Ru2O7, Y1.5Pr0.5Ru2O7), indicating no change in the defects concentration with Pr

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59 content. Therefore, the electrical conductivity increase was not due to an increased chargecarrying capacity of the doped material. The variation of the elec trical conductivity with Pr was consistent with the results reported by Lee et al 53 and it can be explained by the Mott-Hubba rd mechanism of el ectron localization. In fact, the introduc tion of Pr into Y2Ru2O7 caused structural changes of the material as indicated by the Rietveld analysis in Figure 4-5, where it was reported the increase of the lattice parameter as Y were substituted by Pr. Since the ionic radius of the A-site cation depends on the Pr content, the vs Pr trend confirmed the dependence of from the A-site cation. 53This structural modification allowed an increased bandwidth wh ich favored the electric al conduction. Further analysis using Neutron diffraction technique can be used to conf irm the crystalline changes of Y1-xPrxRu2O7 powders. 5.4 Conclusions Temperature dependence of the electrical conductivity of Y1-xPrxRu2O7 (0 x 2) showed a typical semiconducting behavior for all the sa mples. The electrical conductivity increased linearly with increasing the Pr amount. At 700oC, 25 mol % of Pr is enough to double the oxide conductivity. The observed change in conduction behaviour can be e xplained by pyrochlore structural changes. Therefore, intr oducing Pr in the A-site of the Y2Ru2O7 pyrochlore structure is an effective strategy to improve its electrical performance.

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60 Figure 5-1. Variation of the electrical conductivity of Y2-xPrxRu2O7 with dopant amount.

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61 Figure 5-2. Temperature dependence of the electrical conductivity of Y2-xPrxRu2O7.

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62 Figure 5-3. Electrica l conductivity of Y1.8Pr0.2Ru2O7 and Y1.5Pr0.5Ru2O7 as function of oxygen partial pressure at 800oC.

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63 CHAPTER 6 YTTRIUM PRASEODYMIUM RUTHENATE CATHODES FOR IT-SOFCS 6.1 Introduction In the previous chapters, the synthesis of a nanoscaled powder with the chemical composition of Y1-xPrxRu2O7 (0 x 2) that exhibits a pyrochlor e structure was discussed and the material electrical properties were investig ated as a function of Pr content at different temperatures and oxygen partial pressures. As a result, the introduction of Pr in the A-site enhanced the material electrical conductivity. In addition, the Y1-xPrxRu2O7 (0 x 2) pyrochlore resulted to be stable through 1100 oC, thus a negligible loss of RuO is expected over the range of temperatures for a IT-SOFC operation. Herein, Y1-xPrxRu2O7 (0 x 2) was evaluated for application as an IT-SOFC cathode. Pyrochlore oxides, with Ru on the B-site, have been investigated as cathode for IT-SOFCs because their high electro-catalytic activity towards oxygen reduction reaction. Among them, bismuth ruthenate (Bi2Ru2O7) and lead ruthenate (Pb2Ru2O6.5) have been studied because they exhibit metallic conductivity.64-69 However, issues were raised concerning their thermal stability due to the high volatility of ruthenium oxide (RuO). This study investigates the Y1-xPrxRu2O7 cathode on gadolinium doped ceria (GDC) and erbium stabilized bismuth (ESB) electrolytes. Reactivity tests between the pyrochlores and the electrolytes were performed to verify the chem ical compatibility between the two phases. If a new resistive phase would be formed at the el ectrode/electrolyte inte rface, ohmic losses and cathode polarization could increase and the overa ll cell performance could be degenerated. In addition, microstructural characterization was performed on the electrode and on the interface between the electrode and the el ectrolyte. Cathode microstructure plays an important role because an enhancement of elect rode performance and activity can be achieved by increasing the

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64 cathodic surface area.31,32 Finally, electrical performances of a symmetric cell were evaluated by electrochemical impedance spectr oscopy (EIS), which allowed to separate the contributions of the electrolyte, the in terface and the cathode po larization resistances. 6.2 Experimental Procedure A soft coprecipitation method was used to prepare the Y2-xPrxRu2O7 pyrochlore powder with the following chemical composition x = 0, 0.2, 0.5, 1, 1.5, 2. The chemical synthesis of the powders is widely described in th e previous chapters. For further details on the chemical process the reader can refer to the prev ious chapter and previous works.71,72 Ceria doped with 10% of gadolinium (GDC ) powder was purchased by Rhodia. Dense pellets of GDC were obtained by uniaxially pressi ng and sintering at 1450C for 6 h. The pellets showed a relative density of approximately 97%, a nd they were 0.785 cm in diameter and 0.3 cm in thickness. Erbia-stabilized bismuth oxide (ESB), with (Er2O3)0.2(Bi2O3)0.8 composition, was prepared by solid state reaction. Er2O3 and Bi2O3 (Alfa Aesar) were weighted and mixed in the stoichiometric ratio. The powders were ball milled in ethanol for 24 h using YSZ grinding media. Then the precursors were dried at 80oC overnight and calcined at 800 C for 10 hours in air. The final product was then crushed by mortar and pestle and sieved using a 325 mesh. Green ESB pellets were prepared by uniaxial pressing followed by isostatic pressing at 200 MPa. Then, they were fired at 890 C for 15 h. The pellets showed a rela tive density of approximately 96%, and they were 1.12 cm in diameter and 0.3 cm in thickness. Symmetrical cells were fabricated by paint br ushing the ruthenate slurry on both sides of electrolyte pellets. The electrode slurry was prepared by mixing ruthenate powders with an organic binder and a plasticizer in ethanol. Once an appropriate viscosity wa s obtained, the slurry was applied to both sides of the electrolyte pellets by painting. Fi nally, the sample was dried at

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65 90oC for 1 h. The electrodes on ES B electrolyte were fired at 800oC for 2 h. Firing temperature was increased to 850oC and to 950oC for electrodes on GDC elect rolyte. The el ectrodes had thickness of about 15 m and geometric surface area of about 0.8 1 cm2. To check the chemical compatibility of the el ectrodes with the electr olyte, mixtures of Y2xPrxRu2O7 (x = 0 2) and the electro lyte powders (weight ratio 1: 1) were thoroughly pressed and the resulting pellets were treated at 800oC for 20 h. The materials phase analysis was performed by X-ray diffraction (XRD) (Philips, X’Pert). X-Pe rt Data Collector program was used for data collection. Microstructural characterization and chemical an alysis of the electrode /electrolyte interface was performed by scanning electron microscopy with X-ray microanalysis (SEM/EDS, JEOL JSM 6400). Electrode microstruc ture was analyzed by fiel d emission scanning electron microscopy (FE-SEM, LEO SUPRA 1250). The SEM im ages were obtained in secondary and in lens electron modes. Electrochemical performance of a symmetric cell was estimated by Electrochemical Impedance Spectroscopy (EIS). Samples were m echanically contacted with platinum wires, placed into a quartz chamber and positioned in side a tubular furnace. A frequency response analyzer (FRA Solartron SI 1260) coupled with a Solartron 1296 dielectr ic interface with a 50 mV signal was used for the measurements, which were performed in a synthetic air flux at temperatures from 250 C to 750 C in the frequency range 10 mHz 32 MHz. The EIS data were analyzed using Zview software. To verify the data reproducibility, EIS measurements were repeated on 3 different samples each time. The e rror bars that are reported in some figures indicated the maximum and minimum displacement from an average value.

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66 6.3 Results and Discussion 6.3.1 Applicability of Y1-xPrxRu2O7 Cathodes on GDC Electrolyte 6.3.1.1 Chemical compatibility with the solid electrolyte The chemical compatibility of the cathode material with the solid electrolyte was evaluated using SEM-EDS line scan measurements at th e electrode/electrolyte interface and X-ray diffraction analysis of the powder mixtures. Figure 6-1 shows the SEM-EDS line scan of Y1.5Pr0.5Ru2O7/GDC interface. This measurement s howed a discontinuity of the Y L line scan as well as Ru L line scan and Ce L line scan at the Y1.5Pr0.5Ru2O7/GDC interface. Therefore, within the instrument resolution, no el emental solid diffusion occurred at the electrode/electro lyte interface. This result was confirmed by XRD measuremen ts. Figure 6-2 shows the XRD patterns of Y1.8Pr0.2Ru2O7-GDC powder mixture (weight ratio 1:1) as prepared a nd after being pressed in a pellet and treated at 850C for 20 hours. No addi tional peaks were revealed and the single phases were maintained after pressing and firing. The chemical compatibility between the two phases was confirmed also after thermal treatment at 1100oC for 2 hours. Comparable results were obtained for materials with higher amounts of Pr. 6.3.1.2 Microstructural analysis and electrochemical performances Several compositions of Y2-xPrxRu2O7 (x = 0, 0.2, 0.5, 1, 1.5, 2) electrodes were applied on both sides of GDC pellets and fired at 850oC for 2 h. Figure 6-3 shows SEM images of Y1.5Pr0.5Ru2O7 (Pr 25) and YPrRu2O7 (Pr 50) cathodes. Increasing Pr content in the cathode composition, electrode grain size drastically incr eased. This micro structural change is less evident when comparing Pr-rich electrodes. Figu re 6-4 shows that electrode morphology of Pr 50 and Pr 75 (Y0.5Pr1.5Ru2O7), which presented comparable grain size.

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67 The electrochemical properties of Y2-xPrxRu2O7/GDC symmetrical cells were tested using EIS in air from 300 to 750oC. Figure 6-5 shows EIS data from the Y1.5Pr0.5Ru2O7 / GDC / Y1.5Pr0.5Ru2O7 symmetrical cell tested at 700oC in air. The high frequency intercept of the EIS plot can be attributed to the electrolyte impedance (Rb). Therefore, the low frequency data reported in the Nyquist plot in Figure 6-5 gives information a bout the electrode polarization resistance. All the electrodes presented a similar Nyquist plot, in which two overlapped arcs are presented, one big arc at intermediate frequencie s (characteristic frequency of 500 Hz from fig. 6-5) and a small arc at higher frequencies (charact eristic frequency of 8 kHz from fig. 6-5). In some cases an additional small arc is presented at very low frequencies (below 1 Hz).The data at low frequency are usually correlated to the chem ical-physical processes such as gas adsorptiondesorption, gas ionizatio n and surface diffusion 82. Physical models that correlate impedance data to the rate limiting step of th e electrode oxygen reduction are curren tly under discussion. In some cases, the polarization, governed by diffusion and adsorption processes, can be fitted by a Gerischer model. On the other hand, the data at higher frequency can be correlated to interfacial charge-transfer resistance and they can be fitted with a Voigt element or a resistance in parallel with a constant phase element (CPE). In order to better understand the mechanism of oxygen reduction at the electrodes, EIS measurements were carried out as a function of oxygen partial pressure Figure 6-6 shows the impedance plot for Y1.5Pr0.5Ru2O7/GDC/Y1.5Pr0.5Ru2O7 cell as a function of the oxygen partial pressure. It can be observed th at by decreasing the oxygen partial pressure, the low-frequency arc showed a large variation, while the highfreque ncy arc did not vary sign ificantly. Moreover, the characteristic frequency slightly changed towa rds lower values by decr easing the oxygen partial pressure. In this specific case, data were anal yzed using an equivalent circuit which included a

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68 serial combination of a resistance Rb and two parallel combinations a resistance and a constant phase element (CPE). The equivalent circuit is symbolized as Rb(R1Q1)(R2Q2), where Rb represents the electrolyte resistance, and R a nd Q represent the resistance and the CPE element respectively. Specifically R1 represents the resistance associated with the high-frequency arc, and R2 represents the resistance associated with the low-frequency arc. Figure 6-7 shows the polarizati on resistance variation with th e oxygen partial pressure, at different temperatures. The dependence of R1 on pO2 yielded a slope m of about 0.20 (Fig. 6-7A) while the dependence of R2 on pO2 yielded a slope m of about 0.5 (Fig. 6-7B). This confirmed that the resistance associated with the low frequency arc show ed a larger variation with pO2. Recent studies suggested that th e polarization resistance of an electrode varies with the oxygen partial pressure according to the following equation. 83,84 R=R0(pO2)-m (6-1) The magnitude of m provides an insight into the ratelimiting step in the oxygen reduction reaction at the electrodes. A value of 0.25 has be en associated with the charge-transfer reaction at the triple phase boundary (TPB). A value of 0.5 has been associ ated to the surface diffusion of the adsorbed oxygen at the el ectrodes to the TPB. When m = 1, it has been associated to the gaseous diffusion of oxygen molecule s in the electrode structure. 83,84 Based on these considerations, the pO2 dependence of R1 can be related to a limiti ng step involving the charge transfer reaction at the TPB, and the pO2 dependence of R2 can be associated to the limiting step involving oxygen surface diffusion. Therefor e, the electrochemical process of Y1.5Pr0.5Ru2O7/ GDC/ Y1.5Pr0.5Ru2O7 cell is co-limited by the oxygen dissociative adsorption and surface diffusion and charge transfer resistance at the TPB. It is interesting to notice that at very low fr equencies range (lower then 1 Hz) an additional contribution is shown in the impedance plot for Y1.5Pr0.5Ru2O7/GDC/Y1.5Pr0.5Ru2O7 at low

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69 oxygen partial pressure (Fig. 6-6). In these cases the equivalent circuit needs to be adjusted considering a (R3Q3) parallel combination for further data analyses. The electrode Area Specific Resistance (ASR) was calculated by multiplying the total electrode polarization resistance (R1+R2 (+R3)) by the electrode area and dividing by two to account for the symmetric cell. Figure 6-8 shows the variation of the electrode ASR with temper ature for symmetrical cells in air. All the electrodes (Y2-xPrxRu2O7; x = 0, 0.1, 0.5, 1, 1.5, 2) showed a similar ASR temperature dependence, with activation energy of about 1.5 eV. This indicates that the electrochemical performances of Y2-xPrxRu2O7 on GDC electrolyte were governed by the same rate-limiting step. A similar ASR vs temper ature dependence was observed studing other pyrochlore electrodes on GDC, and it was attri buted to the oxygen species at the TPB. 66 Figure 6-9 shows the ASR values as a function of the Pr content into the electrodes, in the temperature range of 650-750oC. The ASR values decreased with increasing Pr amount until a minimum at 25% Pr doped Y2Ru2O7, and then increased again. In some cases the electrodes exfoliated during the experimental measurements indicating a reduced mechanical adhesion at the electrode/electrolyte interfaces. Electrode pr eservation required special care and the EIS measurements were repeated on different samples to obtain reproducible results. The best performance was obtained using Y1.5Pr0.5Ru2O7 cathode, which showed an ASR of 4.23 cm2 at 700oC and activation energy of 1.56 eV. The Y1.5Pr0.5Ru2O7 on GDC electrolyte showed lower ASR than that reported in th e literature for Stron tium (Sr) doped pyrochlore compounds.65 Steele et al.65 reported an ASR value of 47 cm2 for 5 mol% Sr doped Y2Ru2O7 cathode on GDC at 627oC; in contrast, 25% Pr doped Y2Ru2O7 showed an ASR value of 10.72 cm2 at 650oC and 31.98 cm2 at 600oC.

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70 Increasing the Pr amount into the cat hode composition resulted in decreased electrochemical performances (fig. 6-9). This trend was unexpected be cause the electrical conductivity of these pyrochlore ma terials increased lin early with increasing Pr amount. These results probably depend on a varia tion of the electrode mi crostructure with in creasing Pr content. As previously reported, a large microstructural change was observed comparing Pr 25 and Pr 50. Therefore, the increased electrochemical resistan ce of Pr 50 electrodes with respect to Pr 25 can be attributed to the increased particle size. Ind eed, the larger grain size is indicative of a reduced surface area over volume ratio, which affected the oxygen adsorption mechanism on the cathode surface and consequently decreased the electroch emical performance. On the other hand, no large variations were observed between Pr 50 and Pr 75 microstructures, thus the decreased electrochemical resistance of Pr 75 compared to Pr 50 can be attribut ed to the enhanced electrical properties of the Pr 75 material. 6.3.1.3 Summary Application of single phase Y1-xPrxRu2O7 (0 x 2) electrodes on GDC electrolyte showed a change of the electrochemical perfor mances depending on the amount of Pr into the material. The variation of ASR with Pr content depends on the material electrical properties and cathode microstructure. The best performance was obtained for the electrode containing 25 mol% of Pr (Y1.5Pr0.5Ru2O7), which showed an ASR of 4.23 cm2 at 700oC and activation energy of 1.56 eV. This result is very promis ing if compared with a previous study reported about Sr doped yttrium ru thenate on GDC electrolyte.65 However the electrode ASR is still high if compared to the resistance of other common perovskite ma terials used as cathode on GDC electrolyte, such as LSCF.42-44

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71 6.3.2 Applicability of Y1-xPrxRu2O7 Cathodes on ESB Electrolyte Chemical compatibility of the cathode materi al with the solid electrolyte was evaluated using SEM-EDS line scan measurements at th e electrode/electrolyte interface and X-ray diffraction analysis of the powder mixtures. Figure 6-10 shows the line scan at the cross section of a Y1.5Pr0.5Ru2O7/ESB inteface. The Bi M line is close to the Ru L line and their overlapping may distort the chemical analysis resu lts. For this reason Bi analysis was performed using Bi L line and Ru analysis was performed using the Ru K line although this last one was very noisy. A discontinuity of Y L line scan and Ru K line scan was observed at the Y1.5Pr0.5Ru2O7/ESB interface, while the Bi L line scan did not show ed a sharp change in concentration at the interface. Pr obably a partial solid diffusion of the Bi cations took place at the electrode/electro lyte interface. Figure 6-11 shows the XRD patterns of Y1.5Pr0.5Ru2O7-ESB powder mixtur e (weight ratio 1:1) as prepared and after being pressed in a pe llet and treated at 800C for 20 hours. For a better comparison, XRD data were normalized with resp ect to the most intense peak. No additional peaks were observed and the pure phases were main tained after pressing and firing. Comparable results were obtained for materials wi th lower amounts of Pr, such as Y1.8Pr0.2Ru2O7. On the opposite, materials with Pr amount higher then 50 mol % partially reacted in contact with the ESB powder. Figure 6-12 shows the XRD patterns of YPrRu2O7-ESB mixture (weight ratio 1:1) as prepared (a) and after being pr essed in a pellet and fired at 700oC (b) and at 800oC (c) for 20 hours. Pure phases were maintained after the thermal treatment at 700oC for 20 hours. However some extra peaks were revealed after the thermal treatment at 800oC for 20 hours. As the amount of Pr into the pyrochlore structure increased, the intensity of the extra peaks increased.

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72 It is reasonable to think that a new bism uth praseodymium oxide phase formed at the electrode/electrolyte interface. In the literature the (Bi2O3)1-x(Pr6O11)x systems have been studied several times.85,87 Compounds with less than 35 mol % of Pr2O11/3 exhibited ionic conductivity while compounds with more than 40 mol % of Pr2O11/3 are mixed ionic and electronic conductors. In all cases long thermal treatments at 800 oC are required in order to obtain the phase.85,87 Therefore, no degradation of the electrochemical cell is expected during operation at intermediate temperature (500-700 oC) even for the Pr-rich compositions. 6.3.2.1 Electrochemical performances The electrochemical properties of Y2-xPrxRu2O7 (x = 0, 0.1, 0.5, 1, 1.5, 2) pyrochlores were tested on ESB electrolyte, using EIS in air from 300 to 750oC. In this case, electrodes were sintered on ESB electrolyte at 800oC for 2h. Typical Nyquist plot s at low temperatures (200400oC) presented one arc at high frequency (~1MHz) that known to be related to the electrolyte ionic conduction mechanism 82. EIS analysis at high temperatures (T > 500oC) is used to obtain information about the electrode polarization. Fi gure 6-13 shows EIS data from symmetrical cells tested at 700oC in air. The high frequency intercept can be attributed to the electrolyte bulk impedance because of the high testing temperature (Rs).82 To allow a comparison between the different samples, the data reported in Figure 6-13A are Rs corrected, that is Rs had been subtracted from the real component of each data point in the spectrum. Therefore, the low frequency real-axis intercept co rrespond to the effective resist ance for the electrode reaction (Rchem).82 Figure 6-13B shows the imaginary impe dance vs frequency curve in a semilogarithmic scale. The data reported in the gr aph are normalized with respect to the maximum value of Im(Z). Figure 6-14 shows the EIS data from the Y1.5Pr0.5Ru2O7/ESB/Y1.5Pr0.5Ru2O7 symmetrical cell measured at 700oC in air. For all the electrodes the Nyquist plots presented one elongated arc at low frequency (Figures 6-13A and 614; the characteristic fr equency is less then

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73 10 Hz). As previously mentioned, at such low frequency the elect rochemical process is typically dominated by non-charge transfer processes, which include oxygen adsorption-desorption and surface diffusion.82 Figure 6-15 showed the temperature dependence of the electrode ASR for the symmetrical cells in air. Increasing the temperature the ASR decreased, as e xpected since the oxygen reduction process is thermally activated. The ac tivation energy varied depending on the cathode composition, but for Pr doped electrodes (Y2-xPrxRu2O7; x = 0.1, 0.5, 1, 1.5) the activation energy was estimated to be around 1 eV. A simila r value was observed for other pure pyrochlore electrodes on ESB and was attributed to a rate limiting step of the oxygen species at the TPB.66-69 It is interesting to notice th at the activation energy values resulted to be lower for Y2Ru2O7 (Ea = 0.52 eV) and Pr2Ru2O7 (Ea = 0.35 eV) electrodes (Fig. 6-15). This can be related to a different rate-limiting step that governed the electrochemical process of Y2Ru2O7/ESB and Pr2Ru2O7/ESB symmetric cells. Further studies would be n ecessary to determine it. Previously, the electrochemical performance of Y2-xPrxRu2O7 single phase electrodes wa s studied on ceria-based electrolyte, and the same activation energy va lue was observed for each electrode composition. Therefore, the role of the ESB electrolyte may be determining for explaining the electrochemical performance of Y2Ru2O7/ESB and Pr2Ru2O7/ESB symmetric cells. Figure 6-16 shows the ASR values as a function of the Pr content into the electrodes, in the temperature range of 650-750oC. At 700oC, the ASR of the Y2Ru2O7/ESB/Y2Ru2O7 symmetric cell was 4.4 cm2. Small amount of Pr in the electr ode composition allowed a consistent improvement of the performance. Indeed, at 700oC the ASR value of Y1.8Pr0.2Ru2O7 and Y1.5Pr0.5Ru2O7 electrodes on ESB were 0.32 and 0.34 cm2 respectively. A possible explanation is that the increased electrical conductivity of Pr doped materials permitted to lower the electrode

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74 resistance and promoted faster kinetics of the oxygen reduction reacti on at the electrodes. However, the ASR reached a minimum value for Y1.5Pr0.5Ru2O7 electrodes and then it increased again (fig. 6-16). This trend is apparently in contrast to the elect rical properties of Y2-xPrxRu2O7 materials. The reasons of this behaviour ma y depend on two contemporaneous effects: the presence of a new phase at the electrode/electrolyte interface and the change of the cathode microstructure with increasing Pr content. The best electrochemical performance was obtained for Y1.5Pr0.5Ru2O7 electrodes, which showed an electrochemical response comparable with those of Pb2Ru2O6.5 and Bi2Ru2O7 cathode materials.67,68 In the literature Pb2Ru2O6.5 and Bi2Ru2O7 have been widely studied as cathode materials for bismuth based electrolyte.67,68 Impedance studies on Pb2Ru2O6.5/ESB system have shown ASR of 0.41 cm2 at 750oC and activation energy of 1.18 eV, while impedance studies on Bi2Ru2O7/ESB system have shown ASR of 0.43 cm2 at 700oC and activation energy of 1.07 eV. Results obtained in this st udy have shown an ASR of 0.34 cm2 at 700oC and activation energy of 1.08 eV for Y1.5Pr0.5Ru2O7/ESB system, which are similar to the results obtained for Pb2Ru2O6.5 and Bi2Ru2O7 cathodes. It is worth noting that Y1.5Pr0.5Ru2O7 is a semiconducting material, while Bi2Ru2O7 and Pb2Ru2O6.5 are metallic conductors. Therefore, the enhanced electrical conductivity of Pr doped yttrium ruthenate compared to Y2Ru2O7 cannot be the only reason of the high performance of Y1.5Pr0.5Ru2O7 cathode. Electrode polarization must be related to microstructural feature and a study on the cathode microstructu re is reported in the following section. 6.3.2.2 Microstructural analysis The electrode microstructure plays an important role for the electrochemical performance. For proper function, cathode should be porous in order to allow the oxygen gas molecule to

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75 reach the electrochemical active z one at the electrode/e lectrolyte interface. In addition, reduced particle sizes are preferred becau se a larger surface area per volume can be available for the electro-catalytic process. In th is way the number of active sites at the TPB can be increased and the cathodic electrochemical performance improved. In order to obtain a fine elec trode microstructure, the sintering temperature of the electrode on the electrolyte was lowered to 750oC for 2 hours. In addition, at reduced sintering temperatures the formation of new phases at th e electrode/electrolyte interface should be avoid or reduced. This study wa s carried out using Y2-xPrxRu2O7 pyrochlores with 50 mol % of Pr (Pr 50) and 75 mol % of Pr (Pr 75). Figure 6-17 shows FE-SEM pictur es of Pr 50 sintered at 800oC (fig. 6-17A) and 750oC (fig. 6-17B) and Pr75 sintered at 800oC (fig. 6-17C) and 750oC (fig. 6-17D). Microstructure of the sample with high amount of Pr (Pr 75) drasti cally changed with the sintering temperature. The Pr 75 electrode sintered at 800oC presented big agglomerates and a mostly dense microstructure. Decreasing the sintering temperatur e, the electrode still presented a low porosity but the average particles size was reduced up to 600 nm. On the other hand, microstructural changes in Pr 50 were less evident: the big agglomerates presented after firing at 800oC were not observed after firing at reduced temperature, but the average particle size was almost unchanged. Their electrochemical performances were estimated by EIS technique. Figure 6-18 shows the ASR temperature dependence of the Pr 50/ES B symmetric cells (fig 6-18A) and Pr 75/ESB symmetric cells (fig. 6-18B). If compared the sa me electrode composition prepared at different sintering temperatures, the el ectrochemical performance impr oved decreasing the sintering temperature. On the other hand the electrochemical performance of Pr 50 and Pr 75 sintered at 750oC were comparable, despite the different electri cal properties of these materials. This can be

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76 related to the change of the ca thode microstructure, as it was observed by the FE-SEM analysis (Figures 6-17). Indeed, the porous nanostructure of Pr 50 electrodes can be advantageous respect to the large and dense microstr ucture of Pr 75 electrodes. To clarify the role of the electrode microstructure on th e cathodic electrochemical performance, a specific study was carried out usin g a composition with low amount of Pr. In this way the electrochemical results would no be aff ected by an interfacial phase. This study was carried out using the powder with th e optimized composition that is Y1.5Pr0.5Ru2O7 electrode, having 25 mol % of Pr (Pr 25). Two differe nt processes were used to calcine Y1.5Pr0.5Ru2O7 powders: in the first, a one step process, the powders was calcined at 1050oC for 5h, while in the second, a two steps process, the powders was calcined at 600oC for 3h and at 1050oC for 10h. For simplicity we will call batch1 the first powder and batch2 the second. As a consequence of the different calcination processes, the batch2 showed an increased average particle size and an increased size distribution. Both powders we re applied on ESB pellets and fired at 800oC for 2 hours. Figure 6-19 shows the electrodes microstr ucture of the cells tested by impedance spectroscopy technique. The electrodes prepared with batch2 showed a larger particle size distribution then the electrodes prepared with batch1, as expected. Figure 6-20 shows the temperature dependence of the el ectrodes ASR for the symmetrical cells. The increased ASR of the electrodes prepared with ba tch2 confirmed the influence of the electrode microstructure on the electrochemical performance. Due to the larg er particle size, the co ntact resistance between the electrode and the electrolyte increased and the active sites for the electrochemical reaction (TPB) were reduced. Moreover, the catalytic ac tivity of the ruthenate materials is a surface property, thus the kinetics at the cathode with reduced surface/volume ratio (formed with batch2)

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77 would be slower. All these aspects contributed to increase the resistance of the cathode with a large microstructure. 6.3.2.3 Summary The electrochemical performances of Y1-xPrxRu2O7 (0 x 2) pyrochlores on ESB electrolyte drastically changed de pending on the Pr content into th e material. Small amount of Pr into the pyrochlore structure sign ificantly improved the electrochemical performance of the cell. The best performance were obtained using pyrochl ore with 25% mol of Pr, which showed an ASR value of 0.34 cm2 at 700oC and activation energy of 1.08 eV. The presence of Pr into the pyrochlore structure enhanced the electrical conduc tivity of the material and reduced the charge transfer resistance at the electr ode/electrolyte interface. Moreover, the limited amount of Pr into the material provided the proper ch emical compatibility between th e electrode and the electrolyte phases. Finally, the overall cathode kinetics was a ccelerated by the fine cathode microstructure. 6.4 Conclusions Compositions of Y2-xPrxRu2O7 (0 x 2) were tested as a cat hode to compare its electrocatalytic effect with either two electrolytes, GDC or ESB. Figure 6-21 compares the electrochemical resistance of Y2-xPrxRu2O7 single phase electrode on GDC and on ESB electrolytes. Both systems, Y2-xPrxRu2O7/GDC and Y2-xPrxRu2O7/ESB, showed a similar variation of electrode ASR depending on the amount of praseodymium in the cathode material. The ASR values decreased with increasing Pr amount until a minimum at 25% Pr doped Y2Ru2O7, and then increased again. This trend was due to two simultaneous effects, the increase of the electrical conductivity a nd the change of the cathode micr ostructure with increasing Pr content. Pr-rich compositions showed a large mi crostructure which infl uenced the number of reaction sites (TPBs) and the cathodic catalytic activity, decreasing the overall electrochemical

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78 performance. The optimum cathode composition wa s that one that exhibited an increased electrical conductiv ity and a nanosized microstructure. Y1.5Pr0.5Ru2O7 cathode material presented the best perf ormance, with an ASR value of 0.19 cm2 on ESB and 4.23 cm2 on GDC at 700oC. Since the oxygen mass diffusion is the limiting step of the electrode kinetics, the low value of resistivity of the Y1.5Pr0.5Ru2O7 in contact with ESB resulted from a much lower charge tran sfer resistance if compared to the Y2xPrxRu2O7/GDC system. The enhanced performance on ESB electrolyte was also favored by a partial solid diffusion at the interf ace electrode/electrolyte that lowe rs the interfacial polarization. The low value of resistivity of the Y1.5Pr0.5Ru2O7/ESB/Y1.5Pr0.5Ru2O7 symmetric cell indicated that the nanocrysta lline powders of Y1.5Pr0.5Ru2O7 electrode are promising materials for cathode application in ESB-based electrolyte for intermediate temperature solid oxide fuel cells (ITSOFCs).

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79 Figure 6-1. The SEM-EDS line scan of Y1.5Pr0.5Ru2O7/GDC interface.

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80 Figure 6-2. The XRD patterns of mixture of Y1.8Pr0.2Ru2O7-GDC (molar ratio 1:1). A) As prepared. B) Fired at 850oC for 20h. Figure 6-3. The FE-SEM pictur es of electrodes fired at 850oC on GDC. A) Pr 25. B) Pr 50.

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81 Figure 6-4. Low magnification FE-SEM pi ctures of electrodes fired at 850oC on GDC. A) Pr 50. B) Pr 75. Figure 6-5. Impedance plot of Y1.5Pr0.5Ru2O7 electrodes on GDC electrolyte at 700oC in air.

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82 Figure 6-6. Impedance plot of Y1.5Pr0.5Ru2O7/GDC/Y1.5Pr0.5Ru2O7 cell at 700oC. A) Imaginary vs real impedance. B) Imaginary impe dance vs frequency semi-log scale. Figure 6-7. Variation of the elect rode polarization resistance as a function of the oxygen partial pressure.

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83 Figure 6-8. Temperature dependenc e of the electrode ASR on GDC.

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84 Figure 6-9. The ASR values of Y2-xPrxRu2O7 on GDC pellets tested at different temperatures as a function of Pr content. Error bars indicat ed a displacement from an average value, calculated on 3 samples. Lines are guide to eye.

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85 Figure 6-10. The SEM-EDS line scan of Y1.5Pr0.5Ru2O7/ESB interface.

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86 Figure 6-11. The XRD patterns of mixture of Y1.5Pr0.5Ru2O7-ESB (molar ratio 1:1). A) As prepared. B) After thermal treatment at 800oC for 20h.

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87 Figure 6-12. The XRD pattern s of mixture of YPrRu2O7-ESB (molar ratio 1: 1). A) As prepared. B) Fired at 700oC for 20h. C) Fired at 800oC for 20h. Figure 6-13. Impedance plots of Y2-xPrxRu2O7 electrodes on ESB electrolyte at 700oC in air. A) Imaginary vs real impedance. B) Imagin ary impedance vs frequency semi-log scale.

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88 Figure 6-14. Complex impedance plot for Y1.5Pr0.5Ru2O7 /ESB/ Y1.5Pr0.5Ru2O7 cell, measured in air at 700oC. Figure 6-15. Arrhenius plot of the electrode ASR on ESB.

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89 Figure 6-16. The ASR values of Y2-xPrxRu2O7 on ESB pellets tested at different temperatures as function of Pr content. Error bars indicat ed a displacement from an average value, calculated on 3 samples. Lines are guide to eye.

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90 Figure 6-17. The FE-SEM pictures of YPrRu2O7 and Y0.5Pr1.5Ru2O7 electrodes. A) YPrRu2O7 fired at 800oC. B) YPrRu2O7 fired at 750oC. C) Y0.5Pr1.5Ru2O7 fired at 800oC. D) Y0.5Pr1.5Ru2O7 fired at 750oC. Figure 6-18. Arrhenius plots. A) Pr 50 electrodes. B) Pr 75 electrodes.

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91 Figure 6-19. The FE-SEM pictures of Y1.5Pr0.5Ru2O7 electrodes fired on ESB at 800oC for 2h. A) Batch1. B) Batch2.

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92 Figure 6-20. Arrhenius plot of the Y1.5Pr0.5Ru2O7 electrode ASR. Figure 6-21. The ASR values as function of Pr content of Y2-xPrxRu2O7 on ESB and on GDC pellets tested at 700oC. Error bars indicated a displ acement from an average value, calculated on 3 samples. Lines are guide to eye.

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93 CHAPTER 7 COMPOSITE ELECTRODES ON GDC ELECTROLYTE A possible way to lower the cathode polarization resistance is to increase the TPB by mixing one ionic and one electronic phase to form a porous composite cath ode. In this way, the electrochemically active zone is di stributed through the th ickness of the electrode. In contrast, in a single-phase electrode the electro chemical reactive zone is limite d to the electrode-electrolyte interface. In other words, the le ngth and the density of the TPB of a composite electrode can be increased, and as a result of the expansion of the reactive zone, the solid oxide fuel cell performance improved.33-36 Usually, the ionic conductor material used in the composite electrode is the same material used as the electrolyte. In this way, the inte rfacial cohesion between composite electrode and electrolyte could be enhanced and the ther mal expansion mismatch would be minimized.36 However, many parameters influence the polariza tion resistance, and the major ones include the ratio of electronic to ionic mate rial, their respective microstruc ture, and overall thickness and porosity of the electrode. This study investigates the appl ication of a composite electr ode on GDC electrolyte. Best electrochemical performances were obtained using the Y1.5Pr0.5Ru2O7 electrode (chapter 6). Herein, two type of composite have been ev aluated, at first a composite made of Y1.5Pr0.5Ru2O7 and GDC phases, and secondly a composite made of Y1.5Pr0.5Ru2O7 and ESB phases. Electrical performances of a symmetric cell were evalua ted by electrochemical impedance spectroscopy (EIS), and a comparison between the two di fferent composite electrodes was studied. 7.2 Experimental Procedure A soft coprecipitation method was used to prepare the Y1.5Pr0.5Ru2O7 pyrochlore powder. The chemical synthesis of the powders is widely described in the previous chapters and previous

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94 works.71,72 Ceria doped with 10% of gadolinium (GDC) powder was purchased by Rhodia. Dense pellets of GDC were obtained by uniaxiall y pressing and sintering at 1450C for 6 h. The pellets showed a relative density of approximately 97%, and they were 0.785 cm in diameter and 0.3 cm in thickness. Erbia-stabilized bismuth oxide (ESB), with (Er2O3)0.2(Bi2O3)0.8 composition, was prepared by solid state reaction. The chemical synthesis of the powder is well described in the previous chapter (chapter 6). Symmetrical cells were fabricated by paint br ushing the electrode slur ry on both sides of electrolyte pellets. The electrode slurry wa s prepared by mixing ruthenate and the ionic conductive phase powders in an agata mortar in different volumetric ratio, which varies between 0 and 60 vol % of the ionic conductive phase. An organic binder and a plas ticizer were added to the mixture with ethanol solvent, and, once an ap propriate viscosity was obtained, the slurry was applied to both sides of GDC pellets by pa inting. Finally, the sample was dried at 90oC for 1 h. The electrodes on ESB electrolyte were fired at 850oC for 2 h. Firing temperature was increased to 950oC for electrodes on GDC electrolyte. Th e electrodes had thickness of about 15 m and geometric surface area of about 0.8 1 cm2. Electrochemical performance of a symmetric cell was estimated by Electrochemical Impedance Spectroscopy (EIS). Samples were m echanically contacted with platinum wires, placed into a quartz chamber and positioned in side a tubular furnace. A frequency response analyzer (FRA Solartron SI 1260) coupled with a Solartron 1296 dielectr ic interface with a 50 mV signal was used for the measurements, which were performed in a synthetic air flux at temperatures from 250 C to 750 C in the frequency range 10 mHz 32 MHz. EIS data were analyzed using Zview software.

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95 7.3 Composite Y1.5Pr0.5Ru2O7-GDC Electrodes To evaluate the optimal composite compositi on, the ratio between ionic and electrical phases was varied and electrodes with different composite com positions were tested by EIS. Figure 7-1 shows the variation of ASR as a func tion of the vol % of GDC into the electrodes. The introduction of the GDC into the electrode re sulted in significant impedance reduction. The curve presented a minimum value of ASR ar ound 20 vol% of GDC, which was equal to 1.12 cm2 at 700 oC. This resulted to be four times less then the ASR value of the single phase electrode. The electrical conductivity of a composite depends not only to the amount of each phase but also to the degree to which each phases are connected. According to the effective medium percolation theory (EMPT) when the two phases ar e comparable in size, particle contacts should be maximum if the two phases are present in equal fraction of the overall composite volume.45,80,88 In particular, the volume fraction of a randomly distri bution solid phase should be at least more than 1/3 of the total volume of the composite in order for this phase to became continuous. In this way, the ambipolar conductivity of a com posite can reach the maximum.45,80,88 In this specific analysis, a porous materi al rather then a dense composite should be considered. Indeed, it has been demonstrated that the percolation threshold considerable changes with changing porosity.45 Therefore pores should be consid ered as a third phase in order to examine the effective conductivity of a composite material using the EMPT. In this study, the two phases that constitute th e composite are very different in size. Figure 7-2 shows FE-SEM image of the 80/20 Y1.5Pr0.5Ru2O7-GDC composite electrode. From the SEM picture (Fig. 7-2), the GDC powder presente d a nanosized morphology, while the average particle size of Y1.5Pr0.5Ru2O7 was around 200 nm. Since the GDC particles were much smaller

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96 then the Y1.5Pr0.5Ru2O7 particles, they were forced into interstitial regions between the Y1.5Pr0.5Ru2O7 particles and into contact with one anothe r. This resulted in a lower percolation threshold, and low amount of GDC was sufficient to obtain good nu mber of contacts between the two phases. Figure 7-3 shows EIS data at 500oC and 700oC of the composite made with 20 vol% of GDC and 80 vol% of Y1.5Pr0.5Ru2O7. At 500oC (Fig. 7-1A) the electrode impedance consisted of two arcs, one main arc at low frequency with ch aracteristic frequency of 20 Hz and one at high frequency with characteristic frequency of 2 kHz. At 700oC (Fig. 7.1B), the impedance plot consisted of one main arc with characteristic frequency of 1.5 kHz. This suggested that interfacial charge-transfer resist ance was the rate limiting proce ss, which became very slow at reduced temperatures. Figure 7-4 shows the EIS data from the com posite 80/20 and the single phase electrode, and Figure 7-5 compares the electrodes ASR temp erature variation. Th e relative increased characteristic frequency and decreased activation energy of the composite electrode compared to the single phase electrode clearly show a change of the rate limiting step in the electrochemical process. Using a composite electrode is an e ffective strategy to impr ove the cathodic kinetics because it extends the TPB through the thickness of the cathode. In this way, the oxygen reduction reaction is not localized only at the electrode /electrolyte interface but it takes place in an extended zone. Therefore, the oxygen diffusi on through the electrode is no longer the rate limiting step. Although the low ASR values, impedance data for composite electr odes still showed a remarkable interfacial contribution. The high interfacial resistance of Y1.5Pr0.5Ru2O7-GDC composite can be due to the low electrical prope rties of these materials. Increasing the ionic

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97 conductivity of the ion conducting phase should im prove cathode performance, because it allows increasing the width of the reaction zone. 7.4 Composite Y1.5Pr0.5Ru2O7-ESB Electrodes Bismuth based materials, such as ESB, have an ionic conductivity of ten times higher then GDC.14,15 For this reason, a composite made of Y1.5Pr0.5Ru2O7 and ESB was tested on GDC electrolyte. In this case the sintering temperat ure between the electrode and the electrolyte was lowered to 850oC, to avoid solid diffusion between the two phases that formed the composite. Sintering temperatures lower then 850oC did not allow a good electrode adhesion onto the electrolyte surface. In order to evaluate the optimal composite composition, the ratio between ionic and electronic phases was varied and electrodes with different composite co mpositions were tested by EIS. Figure 7-6 shows the variation of ASR as a function of the vol % of ESB into the electrodes. The lowest ASR value was 0.713 cm2 at 700 oC for 60 vol% of ESB. The electrode microstructure was examined by FE-SEM techni que. Figure 7-7 showed the SEM picture of the composite with 60 vol%. The morphology of the two phases used in the composites was rather dissimilar because the pyrochlore particle size was around 100 nm while the fluorite particle size was around 5 m (Fig. 7-7). As a consequence, a good di stribution and connectivity between the electronic and the ionic material s and the maximization of the TP B length occurred at non equal volume fraction of the two phases. This result is consisted with the effective medium percolation theory and the optimization of the TPB. Figure 7-8 shows the impedance plot for the 40/60 Y1.5Pr0.5Ru2O7-ESB composite on GDC electrolyte while figure 7-9 compar es the impedance plot for the Y1.5Pr0.5Ru2O7 single phase electrode and the 40/60 composite measured in air at 700oC. The introduction of the ESB into the

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98 electrode resulted in a significant impedance reduction for both high frequency and low frequency arcs. Figure 7-10 compares the electro de ASR temperature variation when applying Y1.5Pr0.5Ru2O7 single phase electrode or the 40/60 Y1.5Pr0.5Ru2O7-ESB composite on GDC electrolyte. The activation energy of the composite cathode drasti cally decreased (Ea = 1.1 eV), and this is indicative of a change of the rate limiting step in the electrochemical process. Therefore, the enhanced ionic conductivity of ESB decreased th e interfacial resistance and accelerated the electrochemical reaction at the TPB. Finally, the electrochemical performance of Y1.5Pr0.5Ru2O7 single phase electrode was compared with both 80/20 Y1.5Pr0.5Ru2O7-GDC and 40/60 Y1.5Pr0.5Ru2O7-ESB composites. Figure 7-11 shows the ASR temper ature variation of each symmet ric cell. Both the composite cathodes allowed a reduced electr ode resistance and an increa se of the electrochemical performance. At higher temperatures (~700oC) the electrochemical performance of 80/20 Y1.5Pr0.5Ru2O7-GDC and 40/60 Y1.5Pr0.5Ru2O7-ESB composites electrodes were comparable. However, the 40/60 Y1.5Pr0.5Ru2O7-ESB composites showed lower activation energy, which is indicative of a faster electrochemical process. Th is can be explained considering the fact that ESB has an ionic conductivity higher then GDC, thus a composite made with ESB can allow an increase of the width of the electrochemical reactive zone where th e oxygen reduction reaction occurs. 7.5 Conclusions Reduced electrode resistance was obtain ed using a composite cathode made of Y1.5Pr0.5Ru2O7 and GDC or Y1.5Pr0.5Ru2O7 and ESB materials. The best electrochemical performance was obtained using Y1.5Pr0.5Ru2O7-ESB composite cathode, with an ASR value of 0.713 cm2 at 700 oC for 60 vol% of ESB. The enhanced performance derived from the

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99 extension of the reaction ac tive zone through the electrode thickness which improved the cathodic kinetics. Moreover, the higher ESB ionic conductivity allowed a faster kinetics at the reaction zone. Therefore, combining Y1.5Pr0.5Ru2O7 and ESB materials in a composite electrode is an effective strategy to enhance the cathode electrochemical performa nce on GDC electrolyte.

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100 Figure 7-1. Variation of the ASR of Y1.5Pr0.5Ru2O7/GDC composite electrodes as a function of vol % GDC. Figure 7-2. The FE-SEM image of the 80/20 Y1.5Pr0.5Ru2O7-GDC composite electrode.

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101 Figure 7-3. Impedance plots of 80/20 Y1.5Pr0.5Ru2O7-GDC composite electrodes on GDC electrolyte in air. A) At 500oC. B) At 700oC. Figure 7-4. Impedance plots for Y1.5Pr0.5Ru2O7 single phase and composite 80/20 Y1.5Pr0.5Ru2O7GDC on GDC electrolyte, measured in air at 700oC. A) Imaginary vs real impedance. B) Imaginary impedance vs frequency semi-log scale.

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102 Figure 7-5. Arrhenius plots of the AS R for single phase and composite 80/20 Y1.5Pr0.5Ru2O7GDC electrodes. Figure 7-6. Variation of the ASR of Y1.5Pr0.5Ru2O7/ESB composite electrodes as a function of vol % ESB.

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103 Figure 7-7. The FE-SEM image of 40/60 Y1.5Pr0.5Ru2O7-ESB composite electrodes. Figure 7-8. Impedance plot for the 40/60 Y1.5Pr0.5Ru2O7-ESB composite on GDC electrolyte measured in air at 700oC.

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104 Figure 7-9. Impedance plots for Y1.5Pr0.5Ru2O7 single phase and composite with 60 vol% of ESB on GDC electrolyte, measured in air at 700oC. A) Imaginary vs real impedance. B) Imaginary impedance vs frequency semi-log scale. Figure 7-10. Arrhenius plots of th e ASR for single phase and 40/60 Y1.5Pr0.5Ru2O7-ESB composite electrodes.

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105 Figure 7-11. Arrhenius plots of the ASR for single phase and composites electrodes on GDC electrolyte.

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106 CHAPTER 8 CONCLUSIONS Solid Oxide Fuel Cells are a promising altern ative source of energy because of their high conversion efficiency and low pollutant emission s. In order for them to be commercialized, materials and processing must be cost effective, thus their operation temperature need to be reduced to around 600oC. Performance of IT-SOFCs strongl y depends on the choice of cathode material, because the oxygen re duction process is thermally ac tivated and cathodic polarization increases rapidly with decrea sing temperatures. Ruthenates are promising candidates for cathodes in IT-SOFCs due to their high elec tro-catalytic ac tivity towards oxygen reduction reaction and high electrical conductivity. This disse rtation research wanted to investigate how to optimize a novel ruthenate materi al for cathodic application. Powders with Y1-xPrxRu2O7 (0 x 2) compositions were synthesized by a soft precipitation route, which allowed an efficient control of the composition and morphology. Each composition was a single pyrochlore pha se that was stable up to 1200oC. The lattice size of the pyrochlores increased linearly wi th the Pr content, following the Vegard’s law. Nanocrystalline particles were obtained for Pr c oncentrations up to 25 mol %, while larger concentrations led to micro-sized particles. The temperature depe ndence of the electri cal conductivity of Y1-xPrxRu2O7 (0 x 2) showed a typical semiconducting behavi or for all the compositions. The electrical conductivity increased linearly with increasing the Pr amount. At 700 C, increased from 27 S/cm to 38 S/cm to 60 S/cm for samples with 0, 10, 25 mol% of Pr respectively, and reached the value of 124 S/cm for Pr2Ru2O7. The observed change in electri cal conductivity can be explained by structural changes within the pyrochlore. The introduction of Pr into th e A-site increased the A-site ionic radius without destabilizing the py rochlore structure. As a consequence of the structural change the ove rlap between the Ru4d t2g and O2p orbitals increased favoring the

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107 material electrical conductivity 54. Therefore, introducing Pr was an effective strategy to improve the electrical performance. Investigations were further conti nued by applying the single phase Y1-xPrxRu2O7 (0 x 2) materials as cathodes on two electrolytes, GDC or ESB, to compare their electro-catalytic response. Both systems, Y2-xPrxRu2O7/GDC and Y2-xPrxRu2O7/ESB, showed a similar variation of electrode ASR depending on th e amount of praseodymium in th e cathode material. The best performance was obtained for the elec trode containing 25 mol% of Pr (Y1.5Pr0.5Ru2O7), which showed an ASR of 4.23 cm2 on GDC and 0.19 cm2 on ESB at 700oC. Increasing Pr content the ASR values increased again. This trend wa s due to a large cathode microstructure that reduced the numbers of the electrochemical r eaction sites for the redu ction of the oxygen gas molecules. Moreover, it is interesting to note that the electrochemical performance of Y1.5Pr0.5Ru2O7 on ESB was similar with those of Pb2Ru2O6.5 and Bi2Ru2O7 cathode materials despite the fact that Y1.5Pr0.5Ru2O7 showed a semiconducting behaviour and Bi2Ru2O7 and Pb2Ru2O6.5 are metallic conductors. This can be due to the fact that the ca talytic activity of ruthenate materials is governed by surface properties, which can be improved by a fine microstructure. Further research activities a bout these electro-catalyt ic materials can be dedicated to their surface struct ure and properties, si nce they have been shown a predominant influence in the cathodic elec trochemical performance. The best results were obtained using Y1.5Pr0.5Ru2O7 on ESB electrolyte. Since the oxygen mass diffusion is the limiting step of the electro de kinetics, the low va lue of resistivity of Y1.5Pr0.5Ru2O7/ESB/Y1.5Pr0.5Ru2O7 cell resulted from a reduced ch arge transfer resistance if compared to the Y1.5Pr0.5Ru2O7/GDC/Y1.5Pr0.5Ru2O7 system. The enhanced performance on ESB electrolyte can be also favored by partial soli d diffusion at the electr ode/electrolyte interface

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108 which contributed to lower the interfacial po larization in this system. Such evidence was observed by microstructural analys is by FE-SEM. In conclusion, the electrochemical properties of Y1.5Pr0.5Ru2O7 electrode suggested that the nanosized Y1.5Pr0.5Ru2O7 material is promising as cathodes on ESB-based electrolytes for intermedia te temperature solid oxide fuel cells (ITSOFCs).

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109 APPENDIX A DATA REPRODUCIBILITY Each of the experimental techniques used in this work involved ma ny complexities which can be sources of experimental error. To mini mize data uncertainty and obtain reproducible results, each measurement was repeated at l east three times on different replicate samples. Repeating the measurements allowed a first evaluati on of the experimental errors to be obtained, which confirm the reliability of the experimental results. In this study, the experimental work was carri ed out in two different laboratories, one at the University of Florida and one at the Univer sity of Rome, Tor Vergata. Working at the two locations caused a slight variati on in the experimental set up, t hus repeating measurements were necessary. However, data were collected using the same instrument as frequency response analyzer (FRA Solartron SI 1260). Figure A-1 shows the temperatur e dependence of the ASR for Y2Ru2O7 electrode on ESB electrolyte, while figure A2 shows the temperature dependence of the ASR for Y1.5Pr0.5Ru2O7 electrode on ESB electrolyte. The measurements were taken on two different samples, one prepared in the laboratory at the University of Fl orida (UF) and one prepar ed in the laboratory at the University of Rome. The data collected in th e laboratory at UF and in the laboratory at Rome differed in ASR by a maximum of about 26%. In addi tion, the data collected in Rome exhibited a higher ASR temperature dependence. However, other factors can influence an elec trochemical measurement, and, for instance, the electrode microstructure plays an important ro le. In fact, reduced elec trode particles size can enhance the electrochemical performance b ecause a large surface area per volume can be available for the electro-catalytic process. In addition, electrode porosity influences the oxygen diffusion through the electrode up to the electro chemical active zone. Usually, the electrode

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110 morphology can be controlled during the synthesis of the material and the electrode porosity during the electrode slurry preparation. However, each of these steps can be sources of experimental variability that influence the re sult of an electrochemical measurement. In conclusion, considering the multitude of steps duri ng the sample preparation, and the variability of the experimental set up between the two laborat ories, the data collected in the laboratory at UF and at Rome can be considered reproducible.

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111 Figure A-1. Arrhenius plots of the ASR for Y2Ru2O7 electrode on ESB electrolyte, tested in the laboratories of the University of Rome and the University of Florida. Figure A-2. Arrhenius plots of the ASR for Y1.8Pr0.5Ru2O7 electrode on ESB electrolyte, tested in the laboratories of the University of Rome and the University of Florida

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117 BIOGRAPHICAL SKETCH Chiara Abate was born in Castelfranco Veneto a town in north east Italy (40 km from Venice), in 1976. She grew up in the Venetian countryside, where she had the occasion to appreciate art and history from th e Veneto region. She decided to use her love for art and history when she took part in a cultura l association and volunt eer as a tourist guide in Castelfranco Veneto. In 1995 she was enrolled at the Univer sity of Padua, and she obtained the Laurea degree in chemical engineering in 2002. In the begi nning of 2003 she obtained a 12-month grant from the Department of Chemical Science and Technology of the University of Rome, Tor Vergata, in collaboration with ENEA (Italian National La bs for Energy and Environment). During this period she worked in the group of Dr. Mauro Falc onieri and Dr. Elisabetta Borsella on “Studies on Chemical and Physical Techniques for th e Disagglomeration of Semiconducting Nanophasic Powders Based on Silicon in Suspensions” and sh e began to work in collaboration with the group of Prof. Enrico Traversa. In 2004 she decide d to begin a Ph.D program in Prof. Traversa’s research group. In 2005 she arrived at the Univers ity of Florida as visiti ng student, and she had the occasion to meet Prof. Eric Wachsman and to work in his Solid Oxide Fuel Cell research group. She was enthusiastic obout this experience, and in 2006 she enrolled at the University of Florida in a joint Ph.D program between UF and the University of Rome, Tor Vergata. In these years she formed a vast and solid background of knowledge and experiences. She received a Ph.D in materials science and engineering in December 2008.