Thermal Diffusivity of Advanced Ceramics under Harsh Environments

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Thermal Diffusivity of Advanced Ceramics under Harsh Environments
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Moore, Donald T
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University of Florida
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Master's ( M.S.)
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University of Florida
Degree Disciplines:
Materials Science and Engineering
Committee Chair:
Nino, Juan C
Committee Members:
Phillpot, Simon R
Baney, Ronald H

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ceramic -- irradiation -- thermal
Materials Science and Engineering -- Dissertations, Academic -- UF
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Abstract:
Materials capable of withstanding higher temperatures, pressures, stresses, and radiation environments for longer life cycles are an ever increasing desire.  These extremes cause defects that take a toll on materials structure, properties, and performance.  In order to improve materials and establish an operating range in harsh environments, the defect mechanism needs to be determined and how defects affect materials properties needs to be understood.  An example of one such harsh environment is within a nuclear reactor, especially the nuclear fuel which combines the extremes mentioned above.  An approach for reducing nuclear waste while utilizing their energetic value is by using a mixed oxide fuel (MOX) or an inert matrixfuel (IMF) for the transmutation of waste in light water reactors (LWRs) or fast reactors.  Transmutation is a process that converts the radioactive constituents of the waste into more stable elements.  Mixed oxide fuels contain uranium and plutonium or other actinides, but these fertile matrices result in neutron capture generating plutonium and actinides.  An inert matrix fuel consists of a non-fertile inert matrix (IM) that supports a fissile phase (Pu or minor actinides) for efficient transmutation.  MgO and MgAl2O4 were irradiated in the Advanced Test Reactor at Idaho National Lab.  The thermal diffusivity in harsh environments of these potential inert matrix materials, in addition to SiC, is the focus ofthis thesis.
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by Donald T Moore.
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Thesis (M.S.)--University of Florida, 2012.
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Adviser: Nino, Juan C.
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1 THERMAL DIFFUSIVITY OF ADVANCED CERAMICS UNDER HARSH ENVIRONMENTS By DONALD THOMAS MOORE A THESIS PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF MASTER OF SCIENCE UNIVERSITY OF FLORIDA 2012

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2 2012 Donald Thomas Moore

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3 To my wonderful girlfriend and family

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4 ACKNOWLEDGMENTS Most importantly, I would like to thank my pare nts for their love and support. To my brothers, al though you ar e taller than me, I will always be the oldest. M y only regret of going to college is not being at home to go to more basketball game s and playing basketball with the both of you. I express my love for my girlfriend for supporting me a nd to my girlfriend A very special recognition is given to my advisor Dr. Juan C. Nino who gave me the opportunity to start research with his group when I first joined the University. That opportunity defin e d my education as a student researcher or researcher student. I would also like to acknowledge past members of the NRG research group s tarting with Lu Cai and Peng Xu who provided guidance to me when I began research Also, Wei Qui and Beverly Hinojosa who shared group responsibilities and were great officemates I thank Satyajit Phadke, Samantha Yates, Shobit Omar, Luarel Wurcherer, Roberto Esquivel and many others for their assistance In addition I would like to thank my friend Adam Wilk for making classes, projects, and research m o re entertaining as well as Paul Johns for helping me with my research at the university I would also like to thank c urrent members Robert Kasse, Edgar Duarte, Hyuksu Has, Luping Li, Chris Turner, Sasmit Gokhale, Trey Da vis, Brittnee Mound Sara Bermudez and man y others for putting up with me; also, you can no longer blame me for Friday meetings. I also want to thank Todd Allen, James Cole, Randall Fielding, Bryan Forsmann, Mary Catherine Thelen, and others at Idaho National Lab for their contributions to the project. Finally, I want to acknowledge my committee members Dr. Simon Phillpot and Dr. Ronald Baney.

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5 TABLE OF CONTENTS page ACKNOW LEDGMENTS ................................ ................................ ................................ .. 4 LIST OF TABLES ................................ ................................ ................................ ............ 7 LIST OF FIGURES ................................ ................................ ................................ .......... 8 LIST OF ABBREVIAT IONS ................................ ................................ ........................... 11 CHAPTER 1 INTRODUCTION ................................ ................................ ................................ ........ 14 1.1 Statement of Problem and Motivation ................................ ............................... 14 1.2 Scientific Approach ................................ ................................ ........................... 16 1.3 Organization of the Thesis ................................ ................................ ................ 17 1.4 Contributions to the Field ................................ ................................ .................. 18 2 BACKGROUND ................................ ................................ ................................ .......... 19 2.1 Nuclear Fuel Cycle and Materials Testing ................................ ........................ 19 2.1.1 Nuclear Fuel C ycle ................................ ................................ .................. 19 2.1.2 Inert Matrix Materials ................................ ................................ ............... 20 2.1.3 Advanced Test Reactor ................................ ................................ ........... 21 2.2 Laser Flash Method ................................ ................................ .......................... 23 2.2.1 Thermal Diffusivity ................................ ................................ ................... 23 2.2.2 Specific Heat Capacity ................................ ................................ ............ 25 2.3 Thermal Conduction ................................ ................................ .......................... 25 2.3.1 Thermal Conductivity ................................ ................................ ............... 26 2.3.2 Phonon Mean Free Path ................................ ................................ ......... 27 3 THERMAL DIFFUSIVITY OF MAGNESIUM OXIDE ................................ .................. 30 3.1 Magnesium Oxide Structure and Properties ................................ ..................... 30 3.2 Experimental Procedure ................................ ................................ ................... 31 3.2.1 Pellet Fabrication ................................ ................................ ..................... 31 3.2.2 Irradiation Conditions ................................ ................................ ............... 32 3.3 Post irradiation Examination ................................ ................................ ............. 33 3.3.1 X Ray Diffraction and Scanning Electron Microscopy ............................. 33 3.3.2 Focused Ion Beam and Transmission Electron Microscopy .................... 34 3.3.3 Laser Flash Thermal Diffusivity ................................ ............................... 35 3.4 Results a nd Discussion ................................ ................................ ..................... 36 3.5 Conclusions ................................ ................................ ................................ ...... 49

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6 4 THERMAL DIFFUSIVITY OF MAGNESIUM ALUMINATE ................................ ......... 51 4.1 Spinel Crystal Structure ................................ ................................ .................... 51 4.1.1 Review of Spinel ................................ ................................ ...................... 51 4.1.2 Spinel Crystal Structure Parameters ................................ ....................... 52 4.1.3 Crystal Structure Models ................................ ................................ ......... 53 ................................ ................................ ........................ 55 4.2 Struc ture Calculations ................................ ................................ ....................... 57 4.2.1 Theoretical Density ................................ ................................ .................. 57 4.2.2 Global Instability Index ................................ ................................ ............ 58 4.3 Irradiation of MgAl 2 O 4 ................................ ................................ ....................... 61 4.3.1 Synthesis and Fabrication of MgAl 2 O 4 ................................ ..................... 61 4.3.2 Capsule Loading, Disassemb ly, and Irradiation Conditions ..................... 63 4.4 Results and Discussion ................................ ................................ ..................... 64 4.5 Chapter Summary ................................ ................................ ............................. 71 5 THERMAL PROPERTIES OF SILICON CARBIDE ................................ .................... 72 5.1 SiC Structure and Properties ................................ ................................ ............ 72 5.1.1 Review of SiC ................................ ................................ .......................... 72 5.1.2 Crystal Structure and Density ................................ ................................ .. 72 5.1.3 Sample Preparation ................................ ................................ ................. 73 5.1.4 Thermal Properties Characterization ................................ ....................... 74 5.2 Results and Discussion ................................ ................................ ..................... 75 5.3 Chapter Summary ................................ ................................ ............................. 80 6 SUMMARY AND FUTURE WORK ................................ ................................ ............. 81 6.1 Summary ................................ ................................ ................................ .......... 81 6.2 Future Work ................................ ................................ ................................ ...... 83 APPENDIX A MECHANICAL PROPERTIES OF SIC ................................ ................................ ...... 85 B EXPERIENCE AT IDAHO NATIONAL LABOR A TORY ................................ .............. 88 LIST OF REFERENCES ................................ ................................ ............................... 91 BIOGRAPHICAL SKETCH ................................ ................................ ............................ 98

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7 LIST OF TABLES Table page 3 1 MgO i rradiatio n conditions and identification. ................................ ..................... 33 3 2 Thermal diffusivity fitting parameters and diffusivity at irradiation temperature. 45 3 3 Inverse diffusivity slope and intercept of irradiated MgO. ................................ ... 46 4 1 Crystal structure information for MgAl 2 O 4 at 299 K ................................ ............. 52 4 2 Atomic data for MgAl 2 O 4 at 299 K ................................ ................................ ....... 53 4 3 Ionic crystal radius from Shannon ..................... 55 4 4 Mg O and Al O bond information. ................................ ................................ ....... 57 4 5 Theoretical density for MgAl 2 O 4 ................................ ................................ ......... 57 4 6 BSI and GII calculated for MgAl 2 O 4 with no antisit e disorder .............................. 58 4 7 BSI and GII calculated for MgAl 2 O 4 with inversion of x = 0.218. ......................... 59 5 1 Thermal diffusivity fitting paramet ers and diffusivity at 372 K. ............................ 76 5 2 Inv erse diffusivity slope and intercept of irradiated MgO. ................................ ... 76 A 1 Mechanical properties of the SiC pellets and other inert matrix material candidates ................................ ................................ ................................ .......... 87

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8 LIST OF FIGURES Figure page 1 1 2005 h uman development index verse energy consu mption .............................. 14 2 1 Constituents of used LWR fuel ................................ ................................ ........... 19 2 2 Possible fuel cycle options ................................ ................................ ................. 20 2 3 Schematic of the ATR and B1 position wh ere the samples were irradiated ....... 22 2 4 Schemati c of the laser flash experiment ................................ ............................. 23 2 5 Laser flash thermogram signal of a graphite reference ................................ ...... 24 2 6 Heat flow through a cylinder ................................ ................................ ............... 26 2 7 Th ree phonon scattering processes ................................ ................................ ... 28 2 8 Thermal conductivity regions in a solid ................................ ............................... 29 3 1 Crystal structure of MgO ................................ ................................ ..................... 30 3 2 Irradiation temperature and dpa calculations for MgO in capsules I1, I2, and I3 ................................ ................................ ................................ ........................ 33 3 3 Flow c hart for FIB sample preparation ................................ ................................ 35 3 4 Optical images of N and I1 irradiated MgO pellets ................................ ............. 36 3 5 XRD of MgO showing that the material did not amorphize of ch ang e phase during irradiation ................................ ................................ ................................ 37 3 6 The SEM of thermally etched and fracture surface SEM of MgO microstructure. ................................ ................................ ................................ .... 38 3 7 Thermal diffusivity of non irradiated MgO compared to MgO at d ifferent irradiation conditions ................................ ................................ ........................... 41 3 8 Under and over focus TEM images of I3 MgO shows voids at the grain boundary. ................................ ................................ ................................ ........... 42 3 9 TEM of non irradiated I1, I2, and I3 MgO ................................ ......................... 43 3 10 Inverse thermal diffusivity with linear fit to 573 and 873 K. ................................ 45 3 11 TEM image of I2 MgO after being annealed at 1673 K for 1h in flowing argon .. 47

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9 3 12 Thermal diffusivity verse dpa for MgO measured at 296 an d 569 K. .................. 48 3 13 Thermal conductivity of MgO pellets compare d to thin film MgO and bulk MgO ................................ ................................ ................................ .................... 49 4 1 MgO Al 2 O 3 phase diagram ................................ ................................ ................. 51 4 2 Ideal crystal structure model for spinel. ................................ .............................. 54 4 3 Layer sequence along the [100] for spinel. ................................ ......................... 54 4 4 Spinel cystal struc ture model with Mg tetrahedra and Al octahehra ................... 55 4 5 Mg tetrahedron, Al octahedron, and O tetrahedron. ................................ ........... 56 4 6 MgAl 2 O 4 BSI as a function of inversion parameter x ................................ ........... 60 4 7 MgAl 2 O 4 BSI as a function of oxygen parameter u ................................ ............. 61 4 8 Processing flow chart MgAl 2 O 4 ................................ ................................ ........... 62 4 9 Capsule assembly photographs ................................ ................................ ......... 63 4 10 Capsule disassembly in the hot cell using robotic manipu lators ......................... 63 4 11 Irradiation temperatures and dpa calculations for MgAl 2 O 4 in capsules I1, I2, and I3. ................................ ................................ ................................ ................ 64 4 12 Optical images of non irradiated and I1 irradiated MgAl 2 O 4 pellets. ................... 65 4 13 The SEM of thermally etched and fracture surface for MgAl 2 O 4 microstructure 65 4 14 Thermal diffusivity of MgAl 2 O 4 I1 on heating and cooling ................................ ... 67 4 15 TEM of non irradiated I1, I2, and I3 MgAl 2 O 4 ................................ .................... 69 4 16 TEM of MgAl 2 O 4 I3 at lower magnification and higher magnification shows stacking faults ................................ ................................ ................................ ..... 70 5 1 SiC ................................ ................................ ................... 73 5 2 Thermal di ffusivity of the SiC pellets ................................ ................................ ... 75 5 3 Inverse thermal diffusivity of SiC pellets with the same slope at high temp eratures and with increasing firing temperatures, a decrease in intercept indicating less defec t scattering ................................ ................................ .......... 77 5 4 Specific h eat capacity of the SiC pellets ................................ ............................. 78

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10 5 5 Thermal conductivity of the SiC pellets and UO 2 ................................ ................ 78 5 6 Raman spectra of pure AHPCS fired at 1203 and 1323 K temperatures, fired pellets with 10 wt% AH PCS, and SiC powder ................................ ................. 79 A 1 SEM images of fractured pellets fired at 1203 K with 1/16.9 m SiC powder in a 50/50 wt% ratio and 10 wt% AHPCS. ................................ .......................... 87 B 1 Timeline of the ART NSUF irradiation experiment. ................................ ............. 90

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11 LIST OF ABBREVIATION S ATR Advanced Test Reactor BF Bright Field BS Bond Strength BSI Bond Strain Index BVS Bond Valence Sum CN Coordination Number FIB Focused Ion Beam GHG Green House Gas GII G lobal Instability Index IM Inert Matrix IMF Inert Matrix Fuel INL Idaho National Laboratory LO Longitudinal Optical LWR Light Water Reactor MA Minor Actinide MOX Mixed Oxide Fuel NSUF National Scientific User Faci lity PIE Post Irradiation Examination PIP Polymer Infiltration and Pyrolysis SAED Selected Area Electron Diffraction SEM Scanning Electron Microscopy TEM Transmission Electron Microscopy TO Transverse Optical XRD X ray Diffraction

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12 Abstract of Thesis Prese nted to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Master of Science THERMAL DIFFUSIVITY OF ADVANCED CERAMICS UNDER HARSH ENVIRONMENTS By Donald Thomas Moore August 2012 Chair: Juan C. N ino Major: Materials Science and Engineering Materials capable of withstanding higher temperatures, pressures, stresses, and radiation environments for longer life cycles are an ever increasing desire. These extremes cause defects that take a toll on mat erials structure, properties, and performance. In order to improve materials and establish an operating range in harsh environments, the defect mechanism needs to be determined and how defects affect materials properties needs to be understood. An exampl e of one such harsh environment is within a nuclear reactor, especially the nuclear fuel which combines the extremes mentioned above. An approach for reducing nuclear waste while utilizing their energetic value is by using a mixed oxide fuel (MOX) or an in ert matrix fuel (IMF) for the transmutation of waste in light water reactors (LWRs) or fast reactors Transmutation is a process that converts the radioactive constituents of the waste into more stable elements. Mixed oxide fuels contain uranium and plut onium or other actinides, but these fertile matrices result in neutron capture generating plutonium and actinides. An inert matrix fuel consists of a non fertile inert matrix (IM) that supports a fissile phase (Pu or minor

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13 actinides) for efficient transmu tation. MgO and MgAl 2 O 4 were irradiated in the Advanced Test Reactor at Idaho National Lab. The thermal diffusivity in harsh environments of these potential inert matrix materials, in addition to SiC, is the focus of this thesis.

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14 CHAPTER 1 INTRODUCTIO N 1.1 Statement of Problem and Motivation As countries become more developed their energy consumption per capita rapidly increases; this trend is well captured by a human development index by the United Nations Development Programme. 1 Figure 1 1 shows the human development index verse energy consumption. 2 Figure 1 1. 20 05 h uman development index verse energy consumption. 2 Even in the United States, electricity demand is expected to grow 20 % higher by 203 5 3 Meanwhile, the Department of Energy has set the goal of decreasing greenhouse gas (GHG) emissions by 8 3 % by 2050. 2 In the U S two thirds of elec tricity produced comes from greenhouse gas contributor s such as coal and natural gas 4 Nuclear energy accounts for 20% of the total U .S. energy por tfolio; however, compared

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15 to non GHG sources, nuclear energy accounts for 70%. 4 Nuclear energy has many advantages such as carbon free electricity competitive prices and is a domestic energy source reducing dependence on foreign oil 5 N uclear energy will have a significant contribution to the reduction of greenhouse gas emissions Some nuclear energy challenges include waste management, high capital including research and development cost, resource availability, and now more than ever safety. From the waste management perspective there is currently no integrated and permanent solution to hig h 2 As such, t here is an increasing radiotoxic inventory of nuclear waste including plutonium from both spent nuclear fuels and dismantled nuclear weapons and m inor actinides (MA) such as neptunium, americium, and curium from spent nuclear fuel. 6 Geological disposal requires storage for thou sands of years due to long half and minor actinides. 7 An approach for reducing nuclear waste while utilizing their energetic value can be achieved using a mixed oxide fuel (MOX) or an inert matrix fuel (IMF) for the transmutation of waste in light water reactors (LWRs) or f ast reactors 8 Transmutation is a process that converts the radioactive constituents of the waste into more stable elements With transmutation, the length of time required for the nuclear waste to decay to a safe level is significantly reduced. Mixed oxide fuels contain uranium and plutonium or other actinides, bu t these fertile matrices result in neutron capture generating plutonium and actinides. 9 An iner t matrix fuel consists of a non fertile inert matrix (IM) that supports a fissile phase (Pu or minor actinid es) for efficient transmutation Within the inert matrix fuel the inert matrices will experience sever e operating conditions.

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16 M aterials capable of withstanding higher temperatures, pressures, stresses, and radiation environments for longer life cycles are an ever increasing desire. These extremes cause defects that take a toll on materials structure, properties, and performance. In order to improv e materials and establish an operating range in harsh environments, the defect mechanism needs to be determined and how defects affect materials properties needs to be understood. 10 An example of one such harsh environment is within a nuclear reactor, especially the nuclear fuel which combines the extremes mentioned above. 1.2 Scientific Approach The inert matrix material selection is guided by the neutronic properties according to transparency for neutrons. Physical properties of the neutron transparent material are then evaluated. Desired properties include high thermal conductivity, high melting point, compatibility with cladding and reactor coolant and good radiation toleranc e. The IM material needs to be extensively evaluated before it can be considered an option for nuclear applications. The IM is compared to UO 2 as a reference with most properties needing to meet or exceed those of UO 2 After screening material candidate s experimental work is performed to gain the required knowledge of the materials. Ion accelerators are used to study microstructural changes and those materials with low swelling and high amorphization dose are selected. Selected materials are then irra diated in a test reactor to determine the effect of irradiation on the materials properties. The final step is to fabricate the inert matrix fuel and proceed with evaluating its properties before and after irradiation. In this study, six potential inert matrix materials and ceramics were irradiated at Idaho National Laboratory (INL) to evaluate the performance of the materials with an

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17 emphasis on the thermophysical properties P ost irradiation examinations (PIE) were performed to understand the behavior of these materials under irradiation Objectives of this project were to i nvestigate the behavior of single phase Mg based compounds as an inert matrix in irradiation environments as well as c haracterize the effects of irradiation on the microstructure a nd thermophysical properties of the irradiated materials Post irradiation analyses were Complex. Post irradiation examination s included scanning electron microscopy, transmission electron microscopy, and thermal di ffusivity measurements. MgO and MgAl 2 O 4 were irradiated in the Advanced Test Reactor at Idaho National Lab. The thermal diffusivity in harsh environments of these potential inert matrix materials, in addition to SiC, are the focus of this thesis. SiC is another potential inert matrix material; however, s intering represents a major drawback In order to lower the sintering temperature SiC pellets were reaction bonded u sing a preceramic polymer as a binder The thermophysical properties of SiC pellets we re characterized The specific heat capacity and thermal diffusivity were measured and the thermal conductivity calculated. T he proces s is su itable for applications such catalysis or heat exchangers where open porosity is tolerable while preserving high thermal conductivity for the inert matrix application 1.3 Organization of the Thesis In chapter 2 background concepts useful for later chapters will be covered. Subjects covered in the background include inert matrix fuels, irradiation testing and an introduction to the laser flash technique for measuring thermal diffusivity and specific heat. Concepts of thermal conductivity, phonon scattering, and mean free path are also covered. Chapter 3 is comprised of post irradiation examination perfo r med on MgO,

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18 Chapter 4 is PIE work on MgAl 2 O 4 and Chapter 5 is on thermophysical properties of SiC. Conclusions and future work are presented in Chapter 6. Appendix A includes the mechanical properties of SiC discussed in Chapter 5 and Appendix B includes the a 1.4 Contributions to the Field This work investigates the microstructure and thermal diffusivity of in reactor ( in pile ) neutron irradiated ceramics and the thermal properties of SiC The contributions of t his investigation to the field of material science and engineering are summarized below. Neutron irradiation results in a significant reduction in the thermal diffusivity of MgO due to irradiation induced defects, but not due to a change in phonon phonon s cattering. T he number of defects increased with irradiation dose and decreased with the higher irradiation temperature The lower irradiation temperature had a more significant reduction in thermal diffusivity due to the higher density of irradiation dam age. The higher irradiation temperature lead to aggregation of defects and thus did not show as significant reduction in thermal diffusivity. In MgAl 2 O 4 Al is and is confirmed by bond valence sum. For the bond strain index, the oxygen position has a larger effect on the bond strain index than the inversion parameter. The l ow irradiation temperature leads to significant reduction in thermal diffusivity due to point defects and very small defe ct clusters ; however, the high irradiation temperature does not show reduction in thermal diffusivity because of strong vacancy interstitial recombination. The high temperature has areas with a high density of dislocations areas of few dislocations cause d by stacking faults and denuded areas at the grain boundaries. SiC pellets fabricated using a preceramic polymer as binder have a higher thermal diffusivity with increasing firing temperature because of reducing defect scattering, ordering of the amorp hous SiC, and increasing the bonding strength between the SiC particles.

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19 CHAPTER 2 BACKGROUND This chapter summarizes background and fundamentals required for the understanding the research work discussed in the following chapters. 2.1 Nuclear Fuel Cycle and Materi als Testing The following section covers the basics of the nuclear fuel cycle and materials testing of inert matrix materials. 2.1.1 Nuclear Fuel Cycle The c urrent fuel cycle in the U.S. is a once through cycle in which the ore is mined, fuels the reactor, and the spent nuclear fuel from light water reactors (LWR s ) is disposed of in a geological repository after a single use (1) as shown in Figure 2 2 1 1 Spent LWR fuel consists of 95% uranium, 4% of fission products, 1% of p lutonium and minor actinides (MA) as shown in Figure 2 1 2 However, this 1% dominates high level waste disposal issues because of their long term radiotoxicity for hundreds of thousands of years Figure 2 1. Constituents of used LWR fuel. 2 Options for creating a sustainable fuel cycle include improv ing uranium utilization, minimiz ing waste generation, and maximiz ing energy generation. 2 To transition to a clo sed fuel cycle options (2) and (3) as sho wn in Figure 2 2, c ould be included where

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20 spent nuclear fuel would be separated into useable and waste parts Nuclear waste would be deposited in a geological repository and usable components would be recycled. U O 2 would be recycled for enrichment or mixed oxide fuel ( MOX ) fabrication P lutonium and MA could be used in a MOX fuel or in an inert matrix fuel ( IMF ) in LWRs or in fast reactors as shown by (2) and (3) in Figure 2 1. By transmuting the Pu and MA into shorter half live isotopes, the radiotoxicity time frame can be significantly reduced. Figure 2 2 Possible f uel cycle options. (1) Current through fuel cycle, (2) recycling Pu and MA in LWR s (3) burning Pu and MA in fast reactor s 11 2.1.2 Inert Matrix Materials An inert matrix fuel consists of a non fertile inert matrix (IM) that support s a fissile phase (Pu or minor actinides) for efficient transmutation The inert matrix material experiences and needs to withstand the harsh irradiation environment inside the reactor including high temp eratures pressure s and radiation Damage in cera mics may occur

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21 ballistically by h igh energy neutrons that displace atoms and radiolytica l ly by electronic excitation 12 The responses to the damage include defect accumulation, recombination, aggregation, and amo rphization This leads to significant changes in mechanical thermal and optical properties as well as structure and dimensions Required properties of an inert matrix material include l ow neutron absorption cross section, good thermophysical properties such as high melting point and thermal conductivity, chemical stability with cladding and reactor coolant, good irradiation stability and behavior, g ood mechanical properties such as strength and hardness, and high density. 13 The IM need s to be capable of withstanding large displacements per atom (dpa) without becoming amorphous. Displacements per atom is the number of times ever y atom in the material in knocked from its lattice position for a given flue nce. The IM can be a metal or ceramic and the IMF can be single or multiphase. Burnable poisons such as Er and Gd have large neutron cross sections and are added to absorb neutrons based on the neutronic requirements. 13 2.1.3 Advanced Test Reactor The Advanced Test Reactor (ATR) at Idaho National Laboratory is primarily used for research and development as well as isotope production. Samples in this study were irradiated in position B1 as shown in Figu re 2 3 When operated at 110 MW th the fast flux (E>1 MeV) was 8.1x10 13 n/cm 2 s. The ne utron flux is the length travel ed by all neutrons per unit time and volume while the neutron fluence is the flux integrated over time. The n eutron energy represents th e kinetic energy in eV and can be distribute d into energy ranges such as fast neutrons (typically energies greater the 1 MeV), thermal neutrons, or others. The axial flux through the ATR core height is rat her flat compared to a typical reactor. 14 This is beneficial for materia ls testing, because a sample at the

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22 bottom of the core receiv es a similar flu x as a sample in the center of the core To ensure that all the sample s received the same fluence, capsules were inverted midway through the experiment. Figure 2 3 Schemat ic of the ATR and B1 position where the samples were irradiated. 14 The ATR has a higher flux than most commercial reactors such that accelerated materials testing can be performed. A light water reactor uses water as a primary coolant and neutron moderator to reduce the speed of the fast neutrons to thermal neutrons capable of s us taining the fission reaction of uranium 235. Fast neutron reactors do not use a neutron moderator and thus the fission reaction is sustained by fast neutrons.

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23 2.2 Laser Flash Method The next section reviews the laser flash method for measuring thermal diffusivity and specific heat capacity. 2.2.1 Thermal Diffusivity Thermal diffusivity, (m 2 /s), of a solid material is the related to the thermal conductivity by the following equation : (2 1 ) where (W/m K), is the thermal conductivity, c p (J/kg K) is the specific heat capacity at constant pressure and (kg/m 3 ) is the density. The standard method for measuring thermal diffusivity of ceramics is the laser flash method according to ASTM E1461 A small disc is subjected to a short but high ly intens e energy pulse. The energy pulse is absorbed on the front surface of the sample and the temperature rise of the opposite face is recorded as shown by the sc hematic in Figure 2 2 Figure 2 4 Schematic of the laser flash experiment.

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24 The signal measured of a graphite reference is shown in Figure 2 5 The thermal diffusivity is calculated by the Parker equation: 15 (2 2 ) where L is the disc thickness and t 1/2 is the time required for the back surface to reach half of the maximum temperature rise. The Parker equation is based on many assumptions such as no heat loss, infinitesimally short lase r pulse, and uniform absorption of the laser in a very thin layer. Since most of these assumptions are not valid during the experiment, there have been several corrections for different cases. nd takes into account radiative heat losses from the sample. 16 Figure 2 5 Laser flash thermogram signal of a graphite reference.

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25 2.2.2 Specific Heat Capacity Specific heat capacity c p (J/kg K), is the am ount of energy required to raise a unit mass of material by one unit of temperature at constant pressure described by the following equation: (2 3 ) where Q is energy, m is the mass, and is the change in temperature From the F igure 2 5 the specific heat is calculated by: (2 4) where T max is the maximum temperature rise (proportional to the maximum signal rise (V max ) of the infrared detector in Figure 2 5 ) and L is the disc thickness Using the laser flash technique, a reference of known c p ,ref is fl ashed with the same laser pulse (same Q ) as the sample so the c p sample becomes: 17 (2 5 ) Specific heat is measured with the laser flash method by comparing the maximum temperature rise of the sample to the maximum temperature rise of a referenc e sample under the sample measuring conditions. The reference sample should be similar in size, proportions, emissivity, opacity and diffusivity. The reference used in this thesis is thermographite from Anter Corp. 18 2.3 Thermal Conduction The following section is a brief review on the thermal conductivity in non metallic solids.

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26 2.3.1 Ther mal C onductivity eat flow per unit time ( Q / t ) is directly proportional to the temperature gradient ( T / x ) as shown in Figure 2 6 Thermal conductivity, (W/mK), of the material is the constant that relates the two variables in the following equation: (2 6 ) where A is the cross sectional area and the negative sign indicates that heat flows from the hot to cold Ceramics with high thermal conductivities generally have strong interatomic bonding, simple crystal structure, and light elements close together in the periodic table. The propert ies of polycrystalline materials are isotropic so the directio nal dependence is not covered. Figure 2 6. Heat flow through a cylinder. The thermal conductivity of a material can be calculated by: (2 7 ) Heat is conducted throu gh the lattice of non metallic solids by phonons and can be described from the theory of thermal conduction in a classical gas The lattice thermal conductivity of a material can be described by: 19

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27 (2 8 ) By setting equation 2 7 and 2 8 equal by assuming c p c v (generally valid to within 10% of the melting point of a material) 19 the thermal diffusivity can be described by: ( 2 9 ) where is the average phonon velocity, which is approximately constant with temperature and l tot is the total mean free path of the phonons. The average phonon velocity can be approximated by: (2 10 ) where E material and is the density of the material Therefore, the temperature dependence of thermal diffusivity is dominated by the temperature dependence of the phonon mean free path. 2.3. 2 Phonon Mean Free Path The mean free path of a phonon is the distan ce a lattice vibration travels before colliding with another phonon or defect The mean free path of a phonon is limited by collisions between other phonons and lattice im perfections such as grain boundaries, impurities, porosity, dislocations, stacking f aults, and other defects. The total mean free path of phonons, l tot is the sum of mean free path of defects causing scattering: 20 ( 2 1 1 ) w here l pp is the mean free path (m) due to thermal p honon phonon scattering, l pd i s the mean free path due to p hono n defect scattering l gb is phonon grain boundary scattering and l x is mean free path due to all other scattering mechanisms. When the

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28 temp erature is greater than the Debye temperature ( D ) l tot is mainly dependent on l pp and decreases close to the value of interplanar spacing. Phonon scattering that limits thermal conductivity must involve three or more phonons T hree phonon scattering can have the form of Figure 2 7a, k 1 + k 2 = k 3 G where two phonons collide to form one phonon or Figure 2 7b, k 1 G = k 2 + k 3 were one phonon annihilates to form two phonons In normal phonon process es, G = 0 and energy is conserved. I n Umklapp phonon process es, G 0 and energy is not conserved Th e number of phonons is given by the Bose Einstein statistical factor 21 and is proportional to temperature at high temperature s leading to l ~ 1/ T Umklapp processes are the dominate thermal resistance me chanism where three phonon process leads to a l ~ 1/ T relationship and four phonon process leads to l ~ 1/ T 2 relationship. T he heat capacity at high temperatures is approximated by 3k B per atom; therefore, varies with temperature by 1/ T in region C of F igure 2 8 In r egion A the thermal conductivity is limited by the specific heat which yields ~ T 3 In region B the maximum occurs when mean free path of phonon phonon scattering is approximately equal to the mean free path of phonon defect scattering In region D, the thermal conductivity saturates because the mean free path cannot be shorter than the distance between two atoms. Figure 2 7. Three phonon scattering processes

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29 Figure 2 8. Thermal conductivity regions in a solid. 21

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30 CHAPTER 3 THERMAL DIFFUSIVITY OF MAGNESIUM OXIDE This chapter covers p ost irradiation examination of neutron irradiated magnesium oxide and how defects affect the thermal diffusivity 3.1 Magnesium Oxide Structur e and Properties Magnesium oxide (MgO) has many of the p roperties desired for an inert matrix fuel ( IMF ) such as high thermal conductivity (30 W/mK at 773 K) 22 high melting point (3 073 K) 23 low neutron absorption cross section, and good radiation tolerance. Magn esium oxide, also known as m agnesia or periclase, has a rocksalt crystal structure shown in Figure 3 1 represented by interpenetrating face center cubic lattices and has a space group of Fm m (No. 225). Figure 3 1 Crystal struct ure of MgO Light anions and cations of s imilar atomic weight with a simple crystal structure typically leads to a high thermal conductivity and MgO has 30 W/mK at 773 K 22 More now than ever, a critical requirement in the search of new materials is that of accident tolerant fuels for light water reactors ( LWRs ) 24 For example, i n the event of a fuel pin failure the inert matrix ( IM ) needs to be compatible with th e reactor coolant. Mg O

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31 Unfortunately, MgO is soluble in hot water thus limiting it s use as a single phase IM material due to its poor hydration resistance for use in LWR s However, it has been demonstrated that MgO can be made into multiphase composites wit h sufficient hydration resistance for use in LWRs for example MgO ZrO 2 25 and MgO Nd 2 Zr 2 O 7 26 MgO is also being studied as a promising matrix for burning PuO 2 and actinides such as AmO 2 x in fast reactors 27 which do not use water as the primary coolant. Since MgO c an be used as a composite in IMFs for LWRs or as the inert matrix in IMFs for fast reactors i t is important to understand the effect s of irradiation damage on the thermophysical properties under different irradiation conditions There are several studies on the low temperature thermal conductivity of neutron irradiated MgO, 2 8 30 but t here is limited information on the effect of in pile irradiation damage on the thermal diffusivity or conductivity at high temperatures of MgO. Therefore, additional research is needed to understand how irradiation damage affects the thermophy sical properties of MgO. This experiment provides a constant comparison of MgO before and after in pile irradiation at 623 and 973 K This chapter will discuss the effects of irradiation damage on the thermal diffusivity of MgO. 3 .2 Experimental Proced ure 3.2.1 Pellet Fabrication C ommercial MgO (Cerac 99.9%) was ball milled with 70 ml of anhydrous ethanol (Fisher A405) The slurry was milled for 24 h then dried in a fume hood at ambient conditions overnight. In a porcelain mortar and pestle, 2 wt% of binder (Celvol 103 Polyvinyl Alcohol, PVA) was added to the MgO powder and ground to combine thoroughly then sieved through a 212 m mesh. The sieved powder was then dried at

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32 393 K for 5 min to evaporate the water. The powder was uniaxially pressed at 18 0 MPa then sintered in air at 1923 K for 12 h For thermal diffusivity samples, pellets were sintered into 6 mm diameter pellets ~1.8 mm thick. For transmission electron microscopy ( TEM ) samples, pellets 3 mm in diameter were cut using a n ultra sonic drill from an 11 mm diameter and ~1.0 mm thick sintered pellet. The weight and dimension s of the pellets were measured using a balance, caliper, and micrometer The geometric density was then calculated. 3.2.2 Irradiation Conditions To control the temperatu re at which the samples were irradiated, the capsules were back filled with either 100% h e lium for the low temperature (~ 623 K ) or 15% h e lium and 85% a r gon for the high temperature (~ 973 K ). The irradiation temperature was calculated by Peng Xu and Paul M urray at INL using ABAQUS 6.7 3 31 The samples were then irradiated in position B 1 of the Adva nced Test Reactor at Idaho National Laboratory to fast neutron fluencies of ~1x10 25 ( 145 effective full power days ) and ~2x10 25 n/m 2 ( 290 days ) 14 32 The displacements per atom w ere calculated by INL employees using MCNP and determined to be approximately 1.5 and 3 dpa for a displacement energy of 55 eV 33 used for MgO. The displacement energy is the minimum amount of energy in a collision in order for an atom to be displaced from its lattice site. Table 3 1 has the sample identifications for the different irradiation conditions (I1 3) and non irradiated (N) used throughout the thesis and Figure 3 2 has the results from the irradiation temperature and displacement calculations for MgO The slight differences in temperature ( 50 K ) and displacemen ts per atom ( 0.1 dpa) between the TEM and diffusivity samples are due to the different sample holders and

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33 height within the reactor core. For the purposes of the analysis, it was assumed that irradiation conditions of the TEM and diffusivity samples were equivalent. Table 3 1 MgO i rradiation conditions and identification. Sample Fluence (n/m 2 ) Displacements (dpa) T irr (K) N I1 1x10 25 1.5 623 I2 2x10 25 3.0 623 I3 2x10 25 3.0 973 Figure 3 2 Irradiation temperatures (green and red) and dpa (blue) calculations for MgO in capsules I1, I2, and I3. 3.3 Post irradiation Examination 3.3.1 X Ray Diffraction and Scanning Electron Microscopy The crystal structure of as sintered and out of pile pellets w as characterized by X ray diffraction ( XRD, In el Equinox 1000 Artenay FR ). The XRD was conducted using

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34 k V and 30 mA. Polished and thermal etched N MgO and f ractured cross sections of N and I1 pellet s were examined using s cannin g e lectron m icroscope (SEM) at 15 kV ( JEOL JCM 5000, Tokyo, JP; Zeiss 1455 VP Maple Grove MN) 3.3.2 F ocused I on B eam and Transmission Electron Microscopy The TEM specimens were prepared by dual beam focused ion beam (FIB, FEI Quanta 3D FEG FIB/SEM Hills boro, OR) to final lamella dimensions of 10 x 8 um by ~ 1 00 nm thickness. A 2 um thick platinum layer was deposited and trenches were then milled. The specimen was coarse milled at an accelerating voltage of 30 kV and a beam current of 7 nA to a thickness of 2 um then at 3 nA to an 1 um thickness. The specimen was cut an d lifted out using an om n iprobe then fixed to a TEM grid using Pt deposition. Care was taken not to image the sample with the ion beam and to use a low energy cleaning to minimize ion da mage. The sample was thinned with 30kV and 300 pA to 120 nm. The sample was then cleaned using 5 kV and 150 pA, 5 kV and 77 pA, and a final cleaning using 2 kV and 86 pA to the final thickness Figure 3 3 has the flow chart for the preparation of the TE M lamella. Bright field (BF) images two beam images, and selected area electron diffraction (SAED) patterns were recorded using TEM at room temperature (JEOL 2010F Tokyo, JP) operated at 300 kV. Pellets from I2 and I3 were anne aled in a graphite furnace (RDWEBB RD G, Natick, MA) at 1673 K in a f lowing argon atmosphere for 1 h, and then milled using the FIB to prepare a TEM lamella.

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35 Figure 3 3 Flow chart for FIB sample preparation 3.3. 3 Laser Flash Thermal Diffusivity The thermal diffusivity of th e samples was measured according to ASTM standard E1461 07 using the laser flash technique (Anter Flashline 5000 installed in a glove box, Pittsburgh, PA) All samples were thinned with 320 grit SiC paper to 1 mm. Since MgO is translucent the surfaces w ere sputter coated with gold so the sample was opaque and since gold is reflective, the sample was then sprayed with a thin coat of graphite. One sample of each irradiation condition w as measured under flowing a rgon during the same experiment with a multi ple sample carousel Three measurements of each sample were taken at 297 and 373 K to 873 K at 100 K intervals. The diffusivity was calculated following the Clark and Taylor correction for all samples. 16

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36 3.4 Results and Discussion Optical examination. Of the twelve 6 mm and eight 3mm diameter MgO pellets i rradiated, none cracked and the pellets changed color from white to black as shown in Figure 3 4 Samples were fractured and since the scattering length of neutrons in MgO is greater than the sample dimensions, the color change was uniform throughout the thickness of the sample as expected The observed color change is consistent with the occurrence of F+ centers ( which are oxygen vacancies occupied by one or more electron s ) in MgO, whose concentration increases with increasing neutron fluence until a saturation level of 1.3x10 23 m 2 34 After annealing in flowing argon at 1673 for 1 h (same atmosphere as thermal diffusivity measurements and similar to the 85%Ar/15% He I3 backfill gas), the I 2 and I3 sample color changed to a grey color indicating remaining F+ centers. The thermal anneal was under Argon and therefore reoxidation is not expected, as evidenced by the remaining grey color. The recovery was vacancy inte rstitial recombination, interstitial dislocation growth, and coalescence of vacancies into voids which were seen by TEM. Figure 3 4 Optical images of N (white) and I1 irradiated (black) MgO pellets. X ray diffraction. X ray diffraction normally a llows for calculating dislocation density from strain induced peak broadening. 35 However, a t low doses in MgO, damage is mainly Frenkel defects ( a vacancy interstitial pair ) while at higher doses or upon

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37 annealing, interstitial defects co alesce into dislocation loops. Large size or high number of interstitial loops causes the Bragg peak to disappear at 2x10 24 n/m 2 for MgO and a rather sharp peak forms due to diffuse scattering near the location of the Bragg peak. 36 As it will be discussed later, TEM shows that there is a high number of large dislocation loops and thus as previously described the peaks are due to diffuse scattering and not Bragg scattering. Therefore, the peak broadening and shift in the Figure 3 5 inset is due to diffuse scattering. Figur e 3 5 XRD of MgO showing that the material did not amorphize o r change phase during irradiation. The inset shows peak broadening and shift due to diffuse scattering caused by irradiation damage. Since the peak is due to diffuse scattering and not a Br agg scattering, strain, lattice parameter change, or swelling from XRD was not calculated. However, as a first approximation, the XRD patterns in Figure 3 5 would indicate that there was no phase

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38 change or extensive amorphization due to irradiation. This is corroborated by the slope of the inverse thermal diffusivity shown later, which remained unchanged with irradiation meaning there was no change in the intrinsic thermal diffusivity. Microstructure. The average grain diameter for the samples before irr adiation using the ASTM E112 10 intercept procedure was determined to be 11 1 from Figure 3 5 ( top ) There are some large grains due to irregular grain growth. Grain growth was not observed in the irradiated samples. In pile neutron irradiated MgO exhibits transgranular fracture suggesting a decrease in fracture toughness 12 The non irradiated MgO exhibited intergranular fracture through pores on the grain boundaries. The irradiated I1 MgO sample exhibited transgranular fracture and the Figure 3 6 SEM (top) of thermally etched, non irradiated MgO microstructure. The SEM (left) of non irradiated MgO shows intergranular fracture and the SEM (right) of I1 MgO shows transgranular fracture.

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39 change can be attributed to the grains becoming more brittle than the grain boundary due t o the high density of dislocations caused by irradiation. Swelling. A n eutron irradiation study of MgO reported s welling of 2.6 3 vol % for 30 dpa at 430 K due to voids, 37 which is an order of magnitude greater dpa than in this study Sample dimensions were measured geometrically for the thermal diffusivity measurements. All of the samples had an average density of 98% before irradiation and after irradiation. Geometric densities were not accurate enough to determine swelling since the error in geometric densities was calculated to be 2 % for the 6 mm samples Small voids were seen in TEM along grain boundaries in I3 and not in I2 indic a t ing that swelling is greater at the higher irradiation temperature Thermal diffusivity. In Chapter 2, thermal diffusivity was described by equation 2 9 which show s that the mean free p ath controls the thermal diffusivity of a material. The total mean free path of the phonons, l tot in equation 2 1 1 can be simplified to the following equation for irradiated materials: ( 3 1 ) where l pp is the mean free path due to phonon phonon scattering (intrinsic lattice diffusivity), l n is mean free path due to defects present before irradiation, l i is the mean free path due to irradiation induced defects. 38 The inverse of measured thermal diffusivity results in a linear relationship with temperature: ( 3 2 ) where A is related to temperature dependent phonon phonon scattering and B is related to the temperature independent phonon scattering by defects such as grain boundaries,

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40 impurities, dislocations, etc. 20 Neutron irradiation produces large defects, which can be seen with the electron microscope, and point defects, which account for large changes in thermal conductivity. 39 In Chapter 2, it was shown that the varies with temperature by 1/ T due to phonon phonon scattering. Therefore, the t hermal diffusivity of neutron ir radiated ceramics depends on T (K) by the following empirical fit : ( 3 3 ) where k is a constant related the absolute value, and n is a constant that represents the state of the induced defects 40 In Figure 3 7 the thermal diffusivity of non irradiated MgO and ir radiated MgO was fit by equation 3 3 Both MgO I1 and I 2 have the most significant reduction in thermal diffusivity due to the high density of small dislocations Nonetheless, there is little change with the increased irradiation dose. The saturation do se for thermal diffusivity or conductivity is not known for MgO. The saturation dose is important since the diffusivity will not decrease further with increasing dose. Since there is a small reduction in thermal diffusivity between I1 and I2, the saturat ion dose has not been reached for MgO in this study and would require irradiating to higher displacements per atom in order to determine MgO I3 irradiated at high temperature has an intermediate decrease in thermal diffusivity compared to N and I1 becau se of a lower density and larger size of dislocations. The d istance between defect clusters increases with increasing irradiation temperature due to recombination. 41 The MgO samples were measured on heating and upon reaching higher temperatures, the I1 and I 2 began to anneal and r ecover towards the thermal diffusivity of the I3 MgO.

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41 Figure 3 7 Thermal diffusivity of non irradiated MgO (blue) compared to MgO at different irradiation conditions. Both MgO I1 (green) and I2 (orange) have the most significant reduction in therma l diffusivity due to the high density of small dislocations. I3 (red) irradiated at high temperature has an intermediate decrease in thermal diffusivity because of the lower density and larger size of the dislocations. Transmission electron microscopy. Dislocation fo rmation in MgO is well studied and is characterized by dislocation loops and tangles 42 Inte rstitial loops that lie on the { 1 10} planes have <110> Burg ers vectors and often intersect form ing dislocation tangles. 42 Low temperature irradiations of MgO results in unfaulted, elongated interstitial loops at higher doses. 37 TEM images of MgO in Figure 3 9 show that at the low irradiation temperatures ( I1 and I2 ) there is a high density of small dislocations, which increa sed with the higher irradiation dose. The higher irradiation temperature ( I3 ) leads to a lower density of larger dislocations. The higher temperature increases interstitial diffusion leading to aggregation of dislocations which can be seen by the lack

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42 of damage around dislocation structures. The higher temperature also increases recombination resulting in a lower dislocation density. I3 MgO has tangled dislocations and a few dislocation loops as circled in Figure 3 9 At the low irradiation temperature s, point defects were not mobile enough to recombine or make larger defects and led to the high defect density, which caused the most significant decrease in thermal diffusivity. The density of the dislocations was calculated by the line intercept method 43 and determined to be 1.0, 2.4, and 4.5x10 14 m 2 for I3, I1, and I2 respectively. the presence of ~1 nm voids at the grain boundary of I3 MgO but voids in the bulk were not resolvable as shown in Figure 3 8 Using under and over focus BF images the voids appear white then black due to Fresnel contrast. There were no resolvable voids at the grain boundaries or in the bulk of I2 MgO. Two beam imaging was perfo rmed to give an overall representation of the total irradiation damage. Figure 3 8 Under and over focus TEM images of I3 MgO show s voids at the grain boundary.

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43 Figure 3 9 TEM of ( N ) non irradiated MgO has minor ion damage due to the FIB and ( I1 ) MgO has a high density of small dislocation s due to neutron irradiation. TEM of ( I2 ) MgO has a higher density of dislocations compared to (I1) MgO because of the larger irradiation dose. TEM of ( I3 ) MgO has a lower density of larger dislocations due to the higher irradiation temperature. Each TEM image is in the g = [200] two beam condition. Thermal diffusivity fitting The thermal diffusivity of the measured samples in Figure 3 7 was fit with equation 3 3 and constants are listed in Table 3 2 The fit was

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44 performed twice to compare fitting the measurements up to 573 K (bellow the irradiation temperature of 623 K) and up to 873 K (above the irradiation temperature of 623 K) for I1 and I2 samples. The annealing during the measurement had a slight dec rease in n for the I1 and I2 samples, which incorrectly overestimated the induced defects caused by irradiation. Therefore, to calculate the diffusivity at the irradiation temperature, irr the n 873 was used for the I3 and n 573 was used for I1 and I2, s uch that the fit was to measured values below the irradiation temperature s Therefore, to more accurately (and thus safely) predict the thermal diffusivity at operating conditions; the thermal diffusivity should be measured up to the irradiation temperatu re, and not above. This would incorrectly increase the calculated thermal diffusivity due to annealing during the measurement. The samples should also be irradiated as close to the operating conditions to determine correctly the thermal diffusivity beca use defects vary greatly with irradiation temperature and dose. Since the thermal diffusivity is measured out of pile, it is assumed that the post irradiated samples have the same amount of defects during irradiation, that additional defects were not intr oduced at lower temperature while cooling down the reactor, and that defects remain stable during the measurements up to the irradiation temperature. The thermal diffusivity at the irradiation temperature was calculated by the following equation and liste d in Table 3 2 : ( 3 4 )

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45 Table 3 2 Thermal diffusivity fitting parameters and diffusivity at irradiation temperature. Sample Dose (10 25 n/m 2 ) T irr n 573 n 873 297 (10 6 m 2 /s) irr (10 6 m 2 /s) from n 573 irr (10 6 m 2 /s) from n 873 N on irradiated 1.48(5) 17.78 I3 2 973 0.993(7) 0.999(4) 9.60 2.95 2.93 I1 1 623 0.71(3) 0.65(2) 5.52 3.26 3.41 I2 2 623 0.61(7) 0.56(3) 4.90 3.12 3.24 Figure 3 10 Inverse thermal diffusivity with linear fit to 573 and 873 K. Fitting measure d diffusivity values above the irradiation temperature results in an apparent decrease in the intrinsic lattice diffusivity (slope). Inverse thermal diffusivity. To better understand the change in n l inear extrapolation of the measured inverse diffusiv ity versus absolute temperature gives a slope that is determined by the lattice and the intercept at 0 K by defects. 20 The inverse thermal diffusivity linear fit ( equation 3 2 was performed twice with the same

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46 temperatures as before ) For I1 and I2 in Figure 3 10 when fitting using the data above the irradiation temperature, this led to an apparent decrease in the slope, which is the intrinsic lattice diffusivity, and an apparent increase in the intercept, which is due to defects. Table 3 3 has the results of the inverse fit up to but not above the irradiation temperature By fitting up to the irradiation temperature, the slope remains unchanged due to irradiatio n conditions, indicating no change in phonon phonon scattering. The intercept, which indicates the number of defects, increased with irradiation dose and decreased with the higher irradiation temperature because p oint defects are more effective as scatter ing than aggregates Higher irradiation temperatures result in greater aggregation and thus les s scattering. 12 Table 3 3 Inverse diffusivity slope (lattice) and intercept (defects) of irradiated MgO Sample Do se (10 25 n/m 2 ) T irr Slope (s/m 2 K) Intercept (s/m 2 ) Adj. R Square Non irradiated 363(6) 52(3) x 10 3 0.9985 I3 2 973 350(2) 0(1) x 10 3 0.9999 I1 1 623 386(23) 70(10) x 10 3 0.9893 I2 2 623 353(42) 106(19) x 10 3 0.9571 Mean 363 TEM of a nnealed MgO MgO I2 and I3 samples were thermally annealed at 1 673 K for 1 hour in argon prepared by FIB, and then examined by TEM. After annealing I3 (high irradiation temperature) MgO, by visual inspection there were many greater than 100 nm large faceted voids present in the bulk, which was suggested to be formed by vacancy condensation. 44 Voids less than 50 nm were present in I2 on the grain boundary. Annealing of I2 and I3 caused the vacancies to coalesce into voids approximately 4 nm in the bulk shown in Figure 3 11 By visual inspection, t he number of dislocations decreased and the length of dislocations i ncreased for both I2 and I3 It has been reported that loops disappear after annealing at 1623 K in air 45 and at 1773 K

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47 in argon 46 These TEM results of the annealed samples indicate tha t the thermal diffusivity would recover but not to the non irradiated MgO because of the presence of voids and dislocations. Figure 3 1 1 200 two beam TEM image of I2 MgO after being annealed at 1 673 K for 1h in flowing argon showing that the dislocati ons aggregated. Bright field TEM image of voids ~4nm voids in the bulk. Figure 3 12 shows the thermal diffusivity verse displacements per atom for Mg O measured at 296 and 569 K. At 296 K as the dpa increases, defects cause a greater decrease in the the rmal diffusivity than at 569 K. It is also clear that the saturation dose has not been reached for Mg O because there still a decrease in diffusivity from I1 to I2. At higher temperatures, the difference in thermal diffusivity between irradiation conditio ns decreases.

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48 Figure 3 12. Thermal diffusivity verse dpa for MgO measured at 296 and 569 K. Thermal Conductivity The t hermal diffusivit y and specific heat capacity of MgO were measured from 4 73 to 1173 K at 100 K increments in nitrogen using the la ser flash technique ( Anter Flashline 4010 Pittsburgh, PA). Pellets were ground flat and parallel with 320 grit SiC paper so the surfaces were dull. The surfaces were sputter coated with gold so the sample was opaque and since gold is reflective, the sam ple was then sprayed with a thin coat of graphite. One pellet w as measured three times per temperature and the values averaged For the thermal diffusivity measurements, a Clark and Taylor correction 16 was u sed and for heat capacity, a certified thermographite (Anter Corp., Pittsburgh, PA) sample was taken as a reference for the specific heat 18 The thermal conductivity was then calculated using equation 2 7. The results were compared to thin film samples measured using the laser reflectance technique by

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49 Cahil l 47 and compared to bulk MgO also measured using the laser flash technique and differential scanning calorimetry by Ronchi et al 22 The measured thermal conductivity of MgO matches well with the published data. Figure 3 13. Thermal conductivity of MgO pellets compared to thin film MgO and bulk MgO. 22 47 3.5 Conclusions MgO is being investigated as a possible inert matrix material because of high thermal diffusivity and radiation tolerance. MgO pellets were irradiated at 62 3 and 973 K to fast neutron fluenc ies of 1x10 25 (1.5 dpa) and 2x10 25 n/m 2 (3 dpa) in pile of the Advanced Test Reactor at Idaho National Laboratory The effects of in pile neutron irradiation temperature and dose on MgO properties are compared between non irradiated and irradiated MgO samples. Neutron irradiation results in significant

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50 reduction in the thermal diffusivity of MgO due to irradiation induced defects, but not due to a change in phonon phonon scattering. The lower irradiation temperature had a more significant reduction in thermal diffusivity due to the higher density of irradiation damage. The higher irradiation temperature increases interstitial diffusion, leading to aggregation of dislocations which can be seen by the lack of damage around dislocation structures. The higher irradiation temperature also increases recombination, resulting in a lower dislocation density. The intercept of the inverse diffusivity which indicates the number of defects, increased with irradiation dose and decre ased with the higher irradiation temperature because p oint defects are more effective as scattering than aggregates Annealing of MgO caused vacancy interstitial recombination, interstitial dislocation growth, and coalescence of vacancies into voids.

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51 CH APTER 4 THERMAL DIFFUSIVITY OF MAGNESIUM ALUMINATE This chapter covers how changes in the spinel structure affect the bond strain index of MgAl 2 O 4 This chapter also covers the post irradiation examination of neutron irradiated magnesium oxide and how de fects affect the thermal diffusivity. 4.1 Spinel Crystal Structure 4.1.1 Review of Spinel The spinel structure has the general formula AB 2 X 4 which contains two different cations and one anion that is typically oxygen. Many compounds with the spinel struc ture have applications such as electronic, magnetic, and refractory applications. The mineral spinel MgAl 2 O 4 for which the crystal structure is named has a face centered cubic unit cell. Bragg and Nishikawa determined the crystal structure independently in 1915. 48 49 A review of the spinel structure by Sickafus et al covers compounds with this general formula in great detail. 50 Figure 4 1 MgO Al 2 O 3 phase diagram. 51

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52 Figure 4 1 is of the MgO Al 2 O 3 phase diagram. The melting point of spinel is 2105 C and another important detail from the diagram is the ability of spinel to accommodate nonst oichiometry written as AO n B 2 O 3 where n describes the nonequimolarity as shown by MgO 1.5Al 2 O 3 4.1.2 Spinel Crystal Structure Parameters The unit cell of spinel contains eight formula units and a total of 32 anions and 24 cations. The A type cation occ upies a tetrahedral site with a coordination of four and the B type cation occupies an octahedral site with a coordination of six. The cations are arranged such that the lattice parameter of spinel is close to twice of MgO or eight unit cells of MgO could fit in the unit cell of spinel. Table 4 1 and 4 2 have the symmetry and atomic positions for MgAl 2 O 4 spinel from the International Tables for Crystallography and from Redfern et al 52 53 Table 4 1 Crystal structure information for MgAl 2 O 4 at 299 K. 52 53 Space Group Symmetry Point Group Atoms per Unit Cell Lattice Parameter a = b = c () Fd m (No. 227) m m 56 8.08360 The O 2 ions are generally di splaced from the f ace c entered cubic lattice sites which lead to the oxygen positional parameter ( u ) with an ideal position of 3/8. In normal spinel the A 2+ ions would be located on the tetrahedral sites and the B 3+ ions located on the octahedral sites. In inverse spinel, all the A 2+ ions and half of the B 3+ ions sit on the octahedral sites and the other half of the B 3+ ions occupy the tetrahedral sites. However, in most cases are not normal and there is some antisite disorder which leads to th e Al and Mg occupancy on the other cation site, which can be represented by [ IV ] (A 1 x B x ) [VI ] (B 2 x A x )O 4 where x is the cation inversion parameter. The

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53 inversion parameter specifies the fraction of A sites occupied by B site ions and is equal to 0 for norma l spinel, 2/3 for random arrangement of cations, and 1 for inverse spinel. As temperature increases, the inversion parameter increases asymptotically towards 2/3 and pressure has little effect on the disorder. 54 The atomic positions, cation distribution, and thermal displacement parameters from Rietveld refinement in stoichiometric synthetic spinel was determined using in situ time of flight neutron powder diffrac tion collected under vacuum at 26 C. The inversion parameter (equal to the occupancy of the Al in the tetrahedral site) is x = 0.218 and the oxygen positional parameter is u = 0.3867. The sample has a large inversion parameter for room temperature becaus e the sample was quenched from 1500 C. Redfern et al. reported atom positions in origin 2 where the unit cell origin is on a n octahedral vacancy but were converted to origin 1 with the A site cation as the unit cell origin for Table 4 2. Table 4 2 Atomi c data for MgAl 2 O 4 at 299 K. 52 53 Site Wyckoff position Site Symmetry Atomic Positions Occupancy U iso ( 2 ) x y z Mg 8a 3m 0 0 0 Mg = 0.7800 Al = 0.2180 0.219 Al 16c m 5/8 5/8 5/8 Al = 0.8910 Mg = 0.1090 0.288 O 32e m u = 0.3867 u = 0.3867 u = 0.3867 O 0.505 4.1.3 Crystal Structure Mode ls CrystalMaker was used to generate the models of spinel Figure 4 2 is the ideal crystal structure model for spinel highlighting the Mg tetrahedral and Al octahedral in opposite corners of the unit cell. Figure 4 3 is the layer sequence along the [100] for

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54 spinel showing the alternating pattern for the cation locations. Figure 4 4 is the spinel crystal structure model with Mg tetrahedra (green) and Al octahe d ra (purple). Figure 4 2 Ideal crystal structure model for s pinel. Z = 0 Z = 1/8 Z = 1/4 Z = 3/8 Z = 1/2 Z = 5/8 Z = 3/4 Z = 7/8 Figure 4 3 Layer sequence along the [100] for spinel. Mg Al O B type Cation A type Cation Anion

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55 Figure 4 4 Spinel cystal structure model with Mg tetrahedra (green) and Al octahehra (purple). 4. 1.4 stability of a crystal structure. Table 4 3 has information for each of the atoms in spinel which are used for the following calculations The crystal rad ius is for an ion with a certain charge and coordination number is from Shannon. 55 Table 4 3 55 56 Ion Coordination Number Crystal Radius ( ) r M /r O Electronegativity Mg 2+ IV 0.71 0.57 1.31 Al 3+ VI 0.675 0.54 1.61 O 2 IV 1.24 3.44 Ceramics do not have purely ionic or covalent bonding but a mixture of the two that can be calculated by the following empirical equation:

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56 Percent Ionic = 1 exp[ 0.25( X M X X ) 2 ] (4 1) where X m,x is the electronegativity for X and M. 57 A bond is more i onic with increasing difference s in electronegativity. The Mg O bond is 67.8% ionic and the Al O bond is 56.7% ionic. anions and the number of ions can be determined by the cation rad ius ratio. For the coordination number of six which is an octahedron, Rule #1 predicts a radius ratio range of 0.414 to 0.732 mathematically derived for that geometry From Rule #1 it is predicted that both Mg and Al should be in an octahedron; however in the ideal spinel crystal structure, Al is in an octahedron but Mg is in a tetrahedron. In MgO the Mg is in a six coordinated octahedron and in Al 2 O 3 the Al is in an octahedron. This indicates that Al is underbonded and Mg is overbonded in the spine l structure and is confirmed later by bond valence sum (BVS) in section 4.2.3 Figure 4 5 is the O and Mg tetrahedra and the Al octahedron for spinel. Figure 4 5 Mg tetrahedron, Al octahedron, and O tetrahedron. Tab le 4 4 has the Mg O and Al O bond information. Bond strength (BS) is the states that the coordinating polyhedral is arranged in three dimensions in a way that preserves charge Mg Al O

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57 neutr ality. Therefore, the sum of the strength will be equal to the charge of that anion. For oxygen in spinel: n Mg (BS Mg) + n Al (BS Al) = 1(1/2) + 3(1/2) = 2 (4 2) where n Mg,Al is the number Mg or Al bonded to oxygen and their respective bond strength (BS). Table 4 4 Mg O and Al O bond information. Percent Iconicity (%) Actual Bond Length ( ) Ideal Bond Length Ideal Bond Length (%) Bond Strength Mg O 67.8 1.914 1.950 98.15 +2/4 = Al O 56.7 1.931 1.915 100.84 +3/6 = Comparing the actual bond leng th to the ideal hard sphere bond length that is the sum of the Mg 2+ and Al 3+ radius, the Mg O bond is shorter than ideal and the Al O is slightly greater than ideal. The percent ideal bond length was calculated by the following equation: % Ideal Bond Len gth = (4 3) where d is the actual bond length and r A,B are the atomic radii. 4.2 Structure Calculations 4.2.1 Theoretical Density Table 4 5 Theoretical density for MgAl 2 O 4 Formula Units Formula Weight (g/mol) Unit Cell Volume ( 3 ) Density (g/cm 3 ) 8 142.27 528.22 3.58 The density of spinel listed in Table 4 5 is nearly the same as MgO which is 3.57 g/cm 3 The density of spinel incre ases with increasing disorder.

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58 4.2. 2 Global Instability Index Bond valence model is a tool for measuring the strain and estimating the stability of a crystal structure. Increasing the bond valence between two ions reduces the distance between the atoms and is used to calculate the bond strain index (BSI) : (4 4) (4 5) wh ere R o and B are constants, R ij is the bond length, and s is the bond strength. Typically, BSI values greater than 0.05 valence units (vu) indicate that the structure is strained. The sum of the valences of all the bonds formed by an ion is equal to the valence of the ion V i,calc and is used to calculate the g lobal instability index (GII) : (4 6) (4 7) where V i is atomic valence. GII values less than 0.05 vu suggest that little or no strain is present. GII values greater than 0.2 vu indicate that a structure is so strained that it is found to be unstable. Table 4 6 BSI and GII calculated for MgAl 2 O 4 with no antisite disorder. 58 Bond R o B S ij BSI V i,calc GII 1 (Mg O) 1.636 0.42 0.516 0.0262 2.063 0.135 2 (Al O) 1.644 0.38 0.470 2.819 The BSI and GII calculated in Table 4 6 were calculated assumin g that the Mg and Al atoms were sitting in their correct positions meaning the inversion parameter x (equal to the occupancy of the Al in the tetrahedral site) equals 0. The relatively large

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59 GII value for spinel is due to the Al O bond being less than it s expected valence of +3 for Al. This indicates that Mg is overbonded and the Al is underbonded as precisely Table 4 7 has the BSI and GII calculated with inversion of x = 0.218. A bond length of 1.914 was assumed for the tetrahedral position and 1.931 for the octahedral position regardless of the atom in the position. This result ed in a decreased BSI and GII meaning that antisite disorder increases the stability of the structure. Table 4 7 BSI and GII calcul ated for MgAl 2 O 4 with inversion of x = 0.218. Bond Bond Length Occupancy S ij BSI V i,calc GII 1 (Mg O) IV 1.914 0.78 0.516 0.0262 2.063 0.128 1 (Mg O) VI 1.931 0.11 0.495 2.972 2 (Al O) IV 1.931 0.89 0.470 2.819 2 (Al O) VI 1.914 0.22 0.491 1.966 The GII was calculated varying the inversion parameter from 0 (normal) to 2/3 (disordered) for room temperature (RT) spinel quenched from 1500C ( u = 0. 386 7 x = 0.218) and for RT spinel slowly cooled ( u = 0. 387 0 x = 0.143). Inversion increases with temperature and by quenching, the high temperature inversion was frozen in the structure. The importance of the samples being measured and BSI calculated at room temperature is that there is no change in lattice parameter due to thermal expansion and the BVS constants are for room temperature. Both structures had the same lattice parameter but had different x and u parameters.

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60 Figure 4 6 MgAl 2 O 4 BSI as a function of inversion parameter x (equal to the occupancy of the Al in the tetrahedral site) fo r RT quenched sample with oxygen positional parameter (O 32e atomic position u u u ) u = 0. 386 7 (red) and RT slow cooled sample with u = 0. 387 0 (blue). Circles are the BSI calculated from experimental values and squares are calculated by varying the inve rsion parameter. BSI decreases with increasing inversion for the given u values Figure 4 6 shows that the BSI decreases with increasing inversion and shows that the oxygen parameter u has a larger effect on the BSI than the x parameter. Therefore, BSI was calculated varying the oxygen positional parameter from u from 0.36 to 0.41 for inversion of x = 0, 0.218, and 2/3. Figure 4 7 shows that the BSI is minimized when the A O and B O bond lengths are equal.

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61 Figure 4 7 MgAl 2 O 4 BSI as a function of o xygen parameter u for inversion parameter of x = 0, 0.218, and 2/3. The BSI is minimized when the A O and B O bond lengths are equal. When the A site to oxygen (A O) bond length and B site to oxygen (B O) bond length are equal the BSI is minimized. 4.3 Irradiation of MgAl 2 O 4 4.3.1 Synthesis and Fabrication of MgAl 2 O 4 Conventional solid state synthesis was followed for MgAl 2 O 4 and the flow chart is presented in Figure 4 8 Stoichiometric amounts corrected for purity, of MgO ( C erac 99.9 % ) and Al 2 O 3 ( Alph a Aesar 99.997% ) were b all mill ed with 1 mL of ammonium polyacrylate dispersant (PAA, Darvan 821A) and 250 m L of deionized (DI) water. Ball milling was performed in a high density polyethylene bottle with 100 g of 3 m m alumina media on a unitary ball mill The slurry was milled at 85 rpm for 48 hours poured onto a Teflon sheet lining a square glass dish, covered with aluminum foil and subsequently

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62 dried at 120 C overnight The powder was then calcined at 1550 C for 12 hours and the n ball milled for an other 72 hours. The powder was dried ground in a porcelain mortar and pestle, and sieved. P owder and ~2 wt% of binder (Celvol 103 Polyvinyl Alcohol, PVA) were added to a mortar and pestle and ground to combine thoroughly. The powder was sieved trough a 212 m mesh and dried in oven at 120C for 5 min. Approximately 0.2 or 0.5 g of t he powder was added to either a 7 or 13 mm punch and die set cleaned with acetone, lubricated with WD 40, and pressed at 18 0 MPa on a Carver press. The pellet was removed from t he die and examined for cracks and surface finish. Pellets were sintered at 1650 C for 12 hours MgO 1.5Al 2 O 3 was sintered at the same conditions except that the samples were furnace quench ed phase pure samples From the phase diagram in Figure 4 1 ther e is a larger phase stability region for non stoichiometric spinel at high temperature but not at low temperature so that was the reason for the quenching. Figure 4 8 Processing flow chart MgAl 2 O 4

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63 4.3.2 Capsule Loading Disassembly and Irradiation Conditions Sample loadi ng of the 3 mm and 6 mm pellets into their respective sample holder was performed in January 2009. Figure 4 9 has images of the sample loading at the Willow Creek Building at Idaho National Laboratory ( INL ) Figure 4 9 C aps ule assembly. 3mm pellets for TEM (left). S tack of 6 mm discs (middle). Samples being inserted into capsule (right). Capsule I1 was disassembled and samples categorized in January 2010 and c apsule s I2 and I3 in December 2010 at the Hot Fuel Examinati on Facility by INL employs under the guidance of Gregg Wachs. Figure 4 10 has images from the disassembly in the hot cell using robotic manipulators. Figure 4 10 Capsule disassembly in the hot cell using robotic manipulators. Removing the sample holders from the capsule (left). S ample holder with the 3mm disc and stack of 6 mm pellets (middle) The 3mm pellets in sample holder (right).

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64 Irradiation conditions were the same as MgO in section 3.2.2 There is a slight difference in irradiation t emperature and displacements per atom ( dpa ) of MgAl 2 O 4 compared to MgO. This is attributed to the different composition and thermal properties of MgAl 2 O 4 and different d isplacement energy of 60 kV for MgAl 2 O 4 33 The calculated displacements per atom w ere approximately 1.25 and 2.5 dpa. The irradiation temperature for the thermal diffusivity discs was ~ 50 K lower than the TEM samples. Figure 4 11 Irradiation temperatures (green and red) and dpa (blue) calculations for MgAl 2 O 4 in capsules I1, I2, and I3. 4.4 Results and Discussion Optical examination. Of the twelve 6 mm and eight 3mm diameter MgAl 2 O 4 pelle ts irradiated, none cracked and the pellets changed color from white to reddish brown as shown in Figure 4 12 Samples were fractured and the color change was

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65 uniform throughout the thickness of the sample. Just as MgO, t he observed color change is due t o F and F+ centers in MgAl 2 O 4 59 The average geometric density for all of the MgAl 2 O 4 samples was 98% Figure 4 12 Optical images of n on irradiated (white) and I1 irradiated ( reddish brown) MgAl 2 O 4 pellets. Figure 4 13 The SEM (top) of thermally etched, non irradiated MgAl 2 O 4 microstructure. The non irradiated MgAl 2 O 4 SEM (left) shows intergranular fracture along the grain bounda ries. The I1 MgAl 2 O 4 SEM (right) shows transgranular fracture

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66 Microstructure. Non irradiated MgAl 2 O 4 exhibits int ergran ular fracture and transgranu lar fracture after irradiation. 60 It has also been show n that the tensile strength of MgAl 2 O 4 increases 24% with a fluence of 2x10 26 n/m 2 at 660K 60 The increase in fracture strength has been attributed to the increase in strain around dislocation s 61 Figure 4 13 is the microstructure of thermally etched MgAl 2 O 4 (top), and fracture surface of non irradiated and I1. T he change from intergran ular to transgranular fracture can be attributed to the grains becoming more brittle than the grain boundary due to irradiation. Thermal diffusivity. The thermal diffusivity of single crystal MgAl 2 O 4 is not degraded with irradiation because response to da mage is strong vacancy interstitial recombination. 61 Polycrystalline Mg Al 2 O 4 does show reduction in thermal diffusivity due to swe lling caused by voids at grain boundaries and possibly by intragranular point defects 62 Neutron irradiation studies of polycrystalline MgAl 2 O 4 have reported swelling of 0.4 vol% for 3 dpa at 1015 K caused by voids adjacent to grain boundaries. 37 The thermal diffusivity of one I1 MgAl 2 O 4 sample measured three times per temperature is shown in F igure 4 14 MgAl 2 O 4 shows significant reduction in the thermal diffusivity for low irradiation temperature. The thermal diffusivity was measured on heating and on cooling to show the effect of thermal annealing during the measur ement. Defects start annealing once the measurement exceeds the irradiation temperature. Once the measurement reaches 959 K, there is a noticeable increase in the thermal diffusivity. This is interesting to note because this is close to the I3 irradiati on temperature. Vacancy interstitial recombination causes the recovery of the thermal diffusivity. Even as the measurement increases to 1273 K there is not significant annealing of the defects causing further recovery of the thermal diffusivity. This i s shown by the

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67 measurement on cooling where the 1000 K value is relatively unchanged. Dislocations do not unfault and are stable up to above 1100 K 37 T his is why there is no additional recovery in thermal diffusivity. This indicates that the thermal diffusivity of the low temperature I1 and I2 irradiation conditions will reco ver, to the level of I3 but not further. In summary, low temperature irradiation leads to significant reduction in thermal diffusivity d ue to point defects; however for the high irradiation temperature, the thermal diffusivity does not show a reduction. This is because strong vacanc y interstitial recombination is greater at high temperatur e and p oint defects have a larger cross section per defect than aggregated defects. Figure 4 14 Thermal diffusivity of MgAl 2 O 4 I1 on heating and cooling show ing th e effect of thermal annealing during the measurement.

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68 Transmission electron microscopy. Dislocation formation in Mg Al 2 O 4 is characterized by interstitial F rank loops. 63 Low tempera ture ir radiations result ed in a high density of very small < 5 nm dislocation centers or loops. 37 High temperature irradiations result ed in large faulted dislocation loops on { 100 } and [111]. 37 Stacking faults in MgAl 2 O 4 result from interstitial condensation which results in a cation fault. 37 TEM images of MgAl 2 O 4 in Figure 4 15 show that N MgAl 2 O 4 has bending contours due to deformation of the TEM lamella I1 MgAl 2 O 4 has a high density of black spot damage due to neutron irradiation. I2 MgAl 2 O 4 has a higher density of black spot damage compar ed to I1 MgAl 2 O 4 because of the large r irradiation dose. TEM of I3 MgAl 2 O 4 has areas with high density of larger dislocations and areas with few dislocations caused by stacking fault s of due to the higher irradiation temperature. Each TEM image is in the g = [ 4 00] two beam condition. The low temperature irradiation show s very little aggregate damage and a dense distribution of small defect clusters. MgAl 2 O 4 has been shown to decom pose under irradiation at temperatures greater than 1700 K and under a high thermal gradient. 64 From the diffraction patterns the material remained crystalline and there were no addition al spots observed from different phases At high temperature irradiations, [110] loops form clusters of faulted loops, which are termed rosettes. 63

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69 Figure 4 15 TEM of (N) non irradiated MgAl 2 O 4 has bending contours due to deformation of the lamella and (I1) MgAl 2 O 4 has a high density of black spot damage to neutron irradiation. TEM of (I2) MgAl 2 O 4 has a higher density of dislocations compared to (I1) because of the larger irra diation dose. TEM of (I3) MgAl 2 O 4 has areas of high density of larger dislocations and areas of few dislocations caused by stacking faults of due to the higher irradiation temperature Each TEM image is in the g = [ 4 00] two beam condition.

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70 Figure 4 16 TEM of (top left) MgAl 2 O 4 I3 at lower magnification show s area s with lesser damage and areas with high density of large dislocations and (top right) higher magnification of same sample show s that the areas of lesser damage are cause d by stacking fault s. TEM of (bottom right) MgAl 2 O 4 I2 show s damage at the grain boundaries. TEM of (bottom left) MgAl 2 O 4 I3 show s denuded area s at the grain boundaries for the higher irradiation temperature. Figure 4 16 is the TEM of I3 MgAl 2 O 4 at a lower magnification and show s regions of lesser damage and regions of high density of large dislocations At a higher

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71 ma gnification of the same sample, the areas of lesser damage are cause d by stacking faults. In Figure 4 16, I2 MgAl 2 O 4 has damage at the grain boundaries wh ereas I3 MgAl 2 O 4 has denuded area s at the grain boundaries for the higher irradiation temperature. The denuded areas do not contain any dislocations because the grain boundaries act as sinks that reduce damage in this area. From visual inspection, the de nuded areas are approximately 200 nm in width. At higher irradiation temperatures, there is higher diffusion of dislocations and the grain boundaries, just like the stacking faults, act as sinks that reduce damage in the nearby regions. 4. 5 Chapter Summar y In MgAl 2 O 4 Al is underbonded and Mg is overbonded and is confirmed by bond valence sum. Bond strain index shows that the oxygen position has a larger effect on the BSI than the inversion parameter. Bond strain index is min imized when the A O and B O bond lengths are equal. Mg Al 2 O 4 pellets were irradiated at 623 and 973 K to fast neutron fluenc ies of 1x10 25 (1.25 dpa) and 2x10 25 n/m 2 (2.5 dpa) in pile of the Advanced Test Reactor at Idaho National Laboratory T he performan ce of the material with emphasis on the thermal diffusivity was evaluated In summary, the low irradiation temperature leads to a significant reduction in thermal diffusivity due to point defects; however, the high irradiation temperature does not show a reduction in thermal diffusivity The high irr adiation temperature MgAl 2 O 4 has areas with high density of larger dislocations areas with few dislocations caused by stacking faults and denuded area s at the grain boundaries. The low temperature irradiati on shows very little aggregate damage and a dense distribution of small defect clusters.

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72 CHAPTER 5 THERMAL PROPERTIES O F SILICON CARBIDE This chapter covers the thermal properties of reaction bonded silicon carbide using a preceramic polymer as a binder f or SiC powder. 5.1 SiC Structure and Properties 5.1 .1 Review of SiC Silicon carbide SiC also known as carborundum, has high hardness and strength, excellent creep and thermal shock resistance, oxidation and corrosion resistance, low density, and good th ermal conductivity. 65 67 SiC has been studied for applications such as an inert matrix fuel, high temperature catalysis supports, and high temperature heat exchan gers. S intering represents a major drawback due to the high temperatures (usually above 1 973 K ) pressures, and the utilization of sintering aids (TiO 2 Al 2 O 3 or Y 2 O 3 ). 67 An approach used to lower the sintering temperature of S iC is using a preceramic such as polycarbosilane as a binder for SiC powders 65 Upon heating, the preceramic polymer polycarbosilane crosslinks losing mass, converts from the organic polymer to inorganic amorphous SiC, then crystallizes into SiC. 68 72 A llylhydridopolycarbo silane ( AHPCS commercially available as StarPCS TM SMP 10) has been used to make SiC composites 73 74 and has b ee n proposed for inert matrix fuel purposes. 75 76 5.1.2 Crystal Structure and Density SiC (3C polytype) has a cu bic zinc blende crystal structure with a space group of F 3m (No. 216) SiC is covalently bonded SiC 4 and CSi 4 tetrahedra The theoretical density of SiC is 3.21 g/cm 3 compared to 2.4 g/cm 3 for amorphous SiC. The crystal structu re of SiC is presented in Figure 5 1

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73 Figure 5 SiC. 5. 1. 3 Sample Preparation The sample synthesis described in this section was performed by Roberto Esquivel as part of his Masters 77 work and therefore only a brief summary is given In order to prepare the samples, 99.8 SiC powder of nominal sizes 1 m (Alfa Aesar, Ward Hill, MA) and 16.9 m (Superior Graphite, Chicago, IL) in a 50/50 wt% ratio w ere mixed with 10 wt% of AHPCS (Starfire Systems Inc, Schenectady, NY) A 1:1 weight percent ratio of n Hexane 99+% (Acros Organics, New Jersey, US) to powder was added as a solvent The slurry was then manually shaken for 10 min in a 2 cm 3 stainless steel comminution vial with a 5 mm stainless steel ball (Chemplex Industries Inc, Palm City, FL) The slurry was dried overnight in a fume hood for the solvent evaporation The dried mixture was gr ound with a corundum mortar and pestle and sieved through a 150 m stainless steel wire mesh The powder was uniaxially pressed at 200 MPa and then isostatically pressed at 250 MPa The obtained pellets were placed on an alumina tray

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74 with sacrific ial powder and fired under a vacuum of 15 mTorr in a tube furnace (CM Inc, Bloomfield, NJ) with the following schedule: room temperature to 523 K (2 K /min), 523 to 923 K (1 K /min), 923 K to max temperature of 1203 and 1323 K (2 K /min), hold for 1 hour, max to room temperature (3 K /min). A pellet fired at 1203 K was heat treated to 1773 K in an argon atmosphere for 2 hours to fully recrystallize the amorphous SiC. U sing a preceramic polymer as a binder, the SiC pellets reaction bonded at 1203 and 1323 K and had a theoretical density of 77% utilizing the bimodal particle distributio n T he fracture toughness obtained is in Appendix A and is comparable to sintered MgAl 2 O 4 ceramic ; therefore, the process presented here is su itable for applications such as cata lysis or heat exchangers where open porosity is tolerable, or even desirable, while pres erving high mechanical strength. Nonetheless, at th ose processing temperature s only amorphous SiC resulted from the precursor. This might have a deleterious effect o n the thermophysi cal properties of the pellets. 5.1.4 Thermal Properties Characterization T hermal diffusivit y and specific heat capacity of pel lets fired at 1203, 1323, and 1773 K were measured from 373 to 1173 K at 100 K increments in nitrogen using the laser flash technique ( Anter Flashline 4010 Pittsburgh, PA). Pellets were ground flat and parallel with 320 grit SiC paper so the surfaces were dull. One of the highest density pellets of each firing temperature were measured three times per temperatur e and the values averaged For the thermal diffusivity measurements, a Clark and Taylor correction 16 was used and for heat capacity, a certified thermographite (Anter Corp., Pittsburgh, PA) sample was taken a s a reference 18

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75 5. 2 Results and Discussion Thermal d iffusivity It has been shown for p ellets with multiple polymer infiltration and pyrolysis cycles that the thermal conductivity increases with increasing firing temperature 78 and the same trend is obs erved using the preceramic as a binder. The observed behavior is consistent with the increasing atomic ordering from 1203 to 1323 K as indicated in Figure 5 6 and subsequent crystallization of SiC when fired at 1773K Fracture strength 77 suggests SiC and amorphous SiC with the higher firing temperature decreasing the boundary scattering. The 1773 SiC. In Figure 5 2 the measured thermal diffusivity of the SiC samp les was fit by equation 3 3 and the fitting parameters are listed in Table 5 1 Figure 5 2 Thermal diffusivity of the SiC pellets.

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76 Table 5 1 Thermal diffusivity fitting parameters and diffusivity at 372 K. Sample n Adj. R Square 372 (10 6 m 2 /s) SiC 1203 K 0.53(2) 0.999 5.15 SiC 1323 K 0.62(1) 0.996 7.06 SiC 1773K 0.988(9) 0.993 18.04 SiC CVD 19 34 Inverse thermal diffusivity. To better understand the change in n l inear extrapolation of the measured inverse diffusivity versus absolute temperature gives a slope that is determined by the lattice and the intercept at 0 K by defects. 20 The inverse diffusivity in Figure 5 3 appears to have two slopes for the 1203 and 1323 K so the data was fit by a segmented linear regression. The average slopes above 773 K for the three samples are 146 s/m 2 K indicating there was no change in the phonon phonon SiC particles. Since there are not two slopes for the 1773 sample and amorphous SiC was crystallized, the slope change below 773 K is believed to be caused by the amorphous SiC. The slope below 7 73 K was determined to be 230 s/m 2 K for 1323 K and 245 s/m 2 K for 1203 K. The decrease in intercept is due to a decrease in boundary resistance between particles with higher firing temperature. Completely recrystallizing the amorphous SiC above 1473 K is desirable to decrease boundary resistance resulting in the highest thermal diffusivity. Table 5 2 Inverse diffusivity slope (lattice) and intercept (defects) of irradiated MgO. Sample Low T Slope (s/m 2 K) X o High T Slope (s/m 2 K) Intercept (s/m 2 ) R Square SiC 1203 K 245(13) 783(60) 145(19) 108(8) x 10 3 0.996 SiC 1323 K 229(11) 695(40) 149(7) 59(6) x 10 3 0.999 SiC 1773K 144(1) 3(1) x 10 3 0.999 Mean 146

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77 Figure 5 3 Inverse thermal diffusivity of SiC pellets with the same slope at high tempe ratures and with increasing firing temperatures, a decrease in intercept indicating less defect scattering. Specific heat capacity and thermal conductivity. Figure 5 4 is the measured specific heat and Figure 5 5 is the calculated thermal conductivity u sing equation 2 6 The specific heat capacity presented matches well with literature 67 which is expected given that the pellets are 93 SiC regardle ss of the heat treatment employed It should b e noted that pellets with large cracks resulted in a dramatic decrease in the measured thermal diffusivity and increase in the specific heat with the laser flash technique. In Figure 5 5 the t hermal conductivity of t he SiC samples is compared to UO 2 79 Even though the SiC pellets are low density and its thermal conductivity is considerably lower than dense SiC, 67 it is still higher than that of UO 2 By using polymer infiltration of the SiC pellets the density and conductivity could be increased further. Shih et al. reported an increase to 86% of theoretical density with one PIP cycle. 80

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78 Fi gure 5 4 Specific heat capacity of the SiC pellets. Figure 5 5 Thermal conductivity of the SiC pellets and UO 2

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79 Raman spectroscopy. Raman spectroscopy give s information about the vibrational modes in the material and has been used to study the or dering of amorphous SiC. 81 The Raman spectra in Figure 5 6 are of pure AHPCS (fired at both temperatures), fired pellets with 10 wt% AHPCS, and SiC powder. Raman spectra were collected at room temperature using a Horiba LabRAM Aramis Raman spectrometer with a 532 nm diode lase r. A straight line background subtract was performed on the spectra. Figure 5 6 Raman spectra of pure AHPCS fire d at 1203 and 1323 K temperatures, fired pellets with 10 wt% AHPCS, and SiC powder. SiC has a peak reported at ~795 cm 1 and is attributed to the Si C transverse optical (TO) mode. There are also peaks at ~1315 cm 1 for disordered carbon (D) and at ~1 600 cm 1 for ordered graphite (G). 81 82 SiC powder has the Si C TO peak whereas the pure AHPCS only has the D and G peaks meaning the firing temperatures

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80 did not form crystalline SiC from the AHPCS. Raman spectra of the pure AHPCS confirms the amorphous SiC after firing at 1203 and 1323 K with increased intensity of the D and G bands at 1323 K The same effe ct is seen for the pellets with AHPCS except that the Si SiC powder. The ordering is consistent with the increase in thermal diffusivity with higher irradiation temperatures. 5. 3 Chapter Summary Sintered SiC pellets fired at 1203 K resulted in a density of 2.42 g/cm 3 (77% theoretical) and the measured thermophysical properties of the SiC pellets resulted in higher thermal conductivity 9 W/m K, than UO 2 4.5 W/m K at 673 K A higher firing temperature increases the thermal diffusi vity by reducing defect scattering, ordering of the amorphous SiC, and increas ing the bonding strength between the amorphous SiC SiC particles These results indicate that the thermophysical properties of the porous pellets composed of amorphous and polycrystalline SiC prepared in this work, are acceptable for inert matrix fu el application.

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81 CHAPTER 6 SUMMARY AND FUTURE WORK 6.1 Summary Materials capable of withstanding higher temperatures, pressures, stresses, and radiation environments for longer life cycles are an ever increasing desire. These extremes cause defects that t ake a toll on materials structure, properties, and performance. In order to improve materials and establish an operating range in harsh environments, the defect mechanism needs to be determined and how defects affect materials properties needs to be under stood. An example of one such harsh environment is within a nuclear reactor, especially the nuclear fuel which combines the extremes mentioned above. An approach for reducing nuclear waste while utilizing their energetic value is to us e a mixed oxide fu el (MOX) or an inert matrix fuel (IMF) for the transmutation of waste in light water reactors (LWRs) or fast reactors Transmutation is a process that converts the radioactive constituents of the waste into more stable elements. Mixed oxide fuels contain uranium and plutonium or other actinides, but these fertile matrices result in neutron capture generating plutonium and actinides. An inert matrix fuel consists of a non fertile inert matrix (IM) that supports a fissile phase (Pu or minor actinides) for efficient transmutation. MgO and MgAl 2 O 4 were irradiated in the Advanced Test Reactor at Idaho National Lab. P ost irradiation examination was performed to understand the behavior of these materials under irradiation Objectives of this project were to i nvestigate the behavior of single phase Mg based compounds as an inert matrix in irradiation environments as well as c haracterize the effects of irradiation on the microstructure and thermophysical

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82 properties of the irradiated materials The thermal diff usivity in harsh environments of these potential inert matrix materials, in addition to SiC, was investigated. MgO is being investigated as a possible inert matrix material because of high thermal conductivity and radiation tolerance. MgO pellets were i rradiated at 623 and 973 K to fast neutron fluenc ies of 1x10 25 (1.5 dpa) and 2x10 25 n/m 2 (3 dpa) in pile of the Advanced Test Reactor at Idaho National Laboratory The effects of in pile neutron irradiation temperature and dose on MgO properties are compa red between non irradiated and irradiated MgO samples. Neutron irradiation results in significant reduction in the thermal diffusivity of MgO due to irradiation induced defects, but not due to a change in phonon phonon scattering. The lower irradiation t emperature had a more significant reduction in thermal diffusivity due to the higher density of irradiation damage. The higher irradiation temperature increases interstitial diffusion, leading to aggregation of dislocations and also increases recombinatio n. T he number of defects, increased with irradiation dose and decreased with the higher irradiation temperature because p oint defects are more effective as scattering than aggregates Annealing of MgO caused vacancy interstitial recombination, interstiti al dislocation growth, and coalescence of vacancies into voids. In addition to MgO, MgAl 2 O 4 was evaluated as a potential inert matrix material. In the spinel structure Al is rules and is confirmed by bond valence sum. For the b ond strain index the oxygen position has a larger effect on the BSI than the inversion parameter. Bond strain index is minimized when the A O and B O bond lengths are equal. Mg Al 2 O 4 pellets were irradiated under the same c onditions as MgO, except displacements per atom of 1.25

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83 and 2.5 dpa was smaller to evaluate the performance of the material with an emphasis on the thermal diffusivity. In summary, low temperature irradiation leads to significant reduction in thermal dif fusivity due to point defects; however, high temperature irradiation does not show reduction in thermal diffusivity because of strong vacancy interstitial recombination. The high irradiation temperature MgAl 2 O 4 has areas with high density of larger disloc ations areas with few dislocations caused by stacking faults and denuded area s at the grain boundaries. The low temperature irradiation shows very little aggregate damage and dense distribution of small defect clusters. Although SiC was not irradiated, SiC pellets were reaction bonded u sing a preceramic polymer as a binder The specific heat capacity and thermal diffusivity were measured and the thermal conductivity calculated. T he process presented here is su itable for applications such as catalysis o r heat exchangers where open porosity is tolerable, or even desirable, while preserving high thermal conductivity A higher firing temperature increases the thermal diffusivity by reducing defect scattering, ordering of the amorphous SiC, and increasing t he bonding strength between the amorphous SiC SiC particles. T he measured thermophysical properties of the SiC pellets resulted in higher thermal conductivity (9 W/m K at 673 K) than UO 2 These results indicate that the thermophysical properties of the porous pellets composed of amorphous and polycrystalline SiC prepared in this work are acceptable for the inert matrix fuel application. 6.2 Future Work Other potential inert matrix materials that were irradiated in the advanced test reactor include M g 2 SnO 4 70/30 vol% MgO Nd 2 Zr 2 O 7 Nd 2 Zr 2 O 7 2 O 3 and p ost irradiation examination will continue with characterizing dislocation formation in

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84 Mg 2 SnO 4 to determine i f the atomic number and inverse spinel crystal structure are key parameters to i rradiation resistance as compared to MgAl 2 O 4 Samples from the University of Florida irradiation experiment will be added to Advanced Test Reactor National Scientific User Facility sample library and made available for future experiments. Completing pos t irradiation examination of MgO Nd 2 Zr 2 O 7 from capsules I2 and I3 would provide comparisons for irradiation temperature on the thermal diffusivity of the composite The irradiation behavior of MgO should be compared to the ceramic ceramic composite MgO Nd 2 Zr 2 O 7 Testing the mechanical properties of the samples irradiated is also important for MgO Nd 2 Zr 2 O 7 and Mg 2 SnO 4 because this has not been evaluated. Irradiating the samples to higher fluences would be valuable in assessing the potential inert matrix be cause in the actual fuel, the inert matrix would experience greater displacements per atom. Irradiating to higher displacements per atom would cause more swelling and larger decrease in the thermal diffusivity since the saturation dose of the thermal diff usivity was not reac h ed. The next progression in evaluating the inert matrix would be to fabricate the inert matrix fuel and test the properties before and after irradiation. Testing additional parameters such as the chemical compatibility with cladding, chemical stability with fission products, and fission gas solubility and retention is also important. Polymer infiltration and pyrolysis of SiC should be performed to see if the increase in thermal diffusivity at low temperatures is caused by amorphous Si C. Irradiation testing to determine the effect on thermal diffusivity of the SiC pellets with and without polymer infiltration and pyrolysis should also be performed.

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85 APPENDIX A MECHANICAL PROPERTIE S OF SIC The mechanical properties described in this sec tion was performed by Roberto Esquivel as part of his Masters 77 work and the results are discussed in Chapter 5. For SiC sintered with polycarbosilane binder, the bending strength was shown to increase with sintering temperature up to 1373 K then became constant and the bending strength increased with the number of polymer infiltration and pyrolysis ( PIP ) cycles. 65 Reaction bonded, dense SiC by liquid silicon infiltration of bimodal SiC powders have shown to have higher strength and toughness than unimodal SiC powders. 83 Utilizing bimodal SiC powders, a higher density was achieved and effect of the amorphous SiC the mechanical properties were measured for the different firing conditions. P ellets were polished to a 1200 grit finish with aid of a polishing wheel. T he Vickers hardness was then measured with a conventional Vickers indentation method using a load of 1 kg and applying equation 1: ( 4 ) where F is the load (kg) and d is the average length ( millimeters ) of the indentation diagonals Three indents per sample were measured (4 samples per firing temperature) For the fracture strength and fracture toughness tests a Vickers hardness indent of 10 kg was mad e on the center of the pellets to following the indentation/strength technique. 84 The samples were later loaded on a three ball fixture with the indent pointing down towards the three ball arrangement ( piston on 3 ball method) A test frame (Instron model 1350, Instron C orporation, Norwood, MA) was employed and a pin speed of 0.3 mm/min was maintained (1.6 mm in diameter) until the specimen failed.

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86 The load at the rupture was recorded and the fracture strength was calculated using equations 2 4 : 85 ( 5 ) ( 6 ) ( 7 ) where of the support circle, C is the radius (m) of the specimen, P is the load (N) and d is the thickness (m) of the specimen Four pellets were measured per firing temperature. The fracture toughness was obtained from equation 5 using the strength indentation method : 84 ( 8 ) where S is the fracture strength (Pa) and P is the indent load (N). In Table A 1, the SiC pellets fired at the two temperatures exhibit similar values of both fracture strength and toughness. The Vickers hardness of SiC fired at 1203 K was noticeably higher than the 1323 K indicating that the hardness increases inversely wit h firing temperature. This is explained by the densification of the amorphous SiC, resulting in porosity that decreases the boundary strength. SiC pellets were heat treated to 1773 K in an argon atmosphere for 2 hours. The results showed a decrease in h ardness and lower values of fracture strength and toughness, which could be

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87 explained by the increased porosity at interfaces due to the crystallization of amorphous SiC. Figure A 1 is of the fracture surface of the pellet showing intergranular fracture. In spite of these morphology changes, the properties remain superior to other materials. With respect to fully dense SiC, the hardness and fracture strength values of the pellets in this study are around four times smaller, due to the low density. In con trast, the fracture toughness was found to be very close, albeit, the testing methods vary in literature. For different sintering additives and methods, values ranging from 3.5 to 4.5 MPa m 1/2 are reported. 67 Table A 1 Mechanical properties of the SiC pellets and other inert matrix material candidates. Indentation mass: # 1000 g. Hv (kgf/mm 2 ) S (MPa) K IC (MPa m 1/2 ) SiC 1203 K # 542.71 46.7 89.52 4.90 2.55 0.10 SiC 1323 K # 446.65 66.4 93.31 10.60 2.63 0.22 SiC 1773K # 344.57 84.1 73.59 3.68 2.20 0.08 SiC 67 2150.87 2436.29 362 2.6 3 UO 2 86 632.00 20.39 76 0.87 0.14 Figure A 1 SEM images of fractured pellets fired at 1203 K with 1/16.9 m SiC powder in a 50/50 wt% ratio and 10 wt% AHPCS.

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88 APPENDIX B EXPERIENCE AT IDAHO NATIONAL LABOR A TORY earch group during my first year as an undergraduate through a research experience in materials program with the Materials Science and Engineering department. Beginning in January 2008, I began helping Peng Xu with solid state powder synthesis of the comp ounds irradiated at ATR which included synthesizing materials and fabricating pellets under a NERI project While Peng was at I daho National Lab (INL) from August to December 2008 designing the irradiation experiment, I was at UF fabricating the 120 pell ets for the irradiation experiment. In January 2009, I traveled to INL and assisted Pavel Medvedev and Gregg Wachs in loading those pellets into the capsules in 2009 for the ATR NSUF program The capsules were then sealed and inserted into advanced test reactor ( ATR ) where they were irradiated for 145 and 290 days. and began the post irradiation examination (PIE) of the samples that I had prepared as an undergraduate. Figure B 1 has the timeline of the project. I was t rained on radiation safety and h ow to handle radioactive sample s using as low as reasonably achievable (ALARA) principles by maximizing distance, minimizing time, using shielding, and avoid un wanted dose. I was trained to work in trans actinide glove box es and hoods. I learned h ow to decontaminate samples after being in the hot cell. My responsibilities were to lead and assist with the post irradiation examination at INL during the summers of 2010 and 20 11 With the help of Brandon Miller and Cynthia Papesch and many others at INL, we look ed at the effect of irradiation on inert matrix materials structure and thermal

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89 diffusivity During my second summer of research at INL, I had the opportunity to compar e the behavior of materials under different irradiation conditions. It is critical to develop safer ways of disposing of nuclear waste and to understand how irradiation damage affect s the material properties of these inert matrix candidates. Working at the INL has provided me with numerous opportunities to further my knowledge in nuclear science and materials science. The ATR NSUF Users Week during the summers has been a valuable learning experience for how irradiation affects materi als properties. Not only have I learned more about my project, but I have learned equally as much on other research projects at INL. Being able to experience firsthand the irradiation fuel looked at in the focused ion beam i s something I could not experience anywhere else. I took a class on atom probe tomography at the Center for Advanced Energy Studies. With the Advanced Test Reactor National Scientific User Facility I had the opportunity to go DC as a representative for the User Facility Organization to meet with Cong ress and the House of Representatives. I also went to Las Vegas with the ATR NSUF to share my experience at the American Nuclear Society student conference. With a materials science background, my great experience over the summer performing PIE has change d my perspective on how irradiation affects material microstructure and properties and has also taught me the importance of nuclear energy. Being part of the new generation of engineers and scientists in working nuclear materials with hands on experience at a national l aboratory is very rewarding and fulfilling. Having the opportunity to work at a national laboratory has given me an opportunity that many students rarely e xperience in a university setting.

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90 Figure B 1. Timeline of the ART NSUF irradi ation experiment.

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98 BIOGRAPHICAL SKETCH Donald Moore was born in San Diego, California in 1988. He i s oldest of three brothers, while also being the shortest at 6 feet 4 inches. In 1991, his family moved to Antioch, California, which is located in the bay area. He attended Jack London Elementary and Mno Grant Elementary then Black Diamond Middle School while in California. In 2002, his family moved across the country and settled in Cape Coral, Florida. He finished the remainder of his 7 th grade year at Gulf Middle School. While From 2004 to 2007, he attended Mariner High School, where he was a member of the varsity basketball team for three years. After graduating from high school, he began attending the University of Florida in August of 2007. During his freshman year, he jo His research included fabricating samples for the Advanced Test Reactor irradiation experiment and characterizing ceramics o f electronic applications. During the summers of 2010 and 2011 he spent three months as an intern at Idaho National Lab in Idaho Falls, Idaho. Here, he assisted in post irradiation examination of inert matrix ceramics under the guidance of Pavel Medvedev. In pursuing a combined BS/MS degree, he received his undergraduate degree in Materials Science and Engineering in May of 2012. He received his m aster s in Materials Science and Engineering from the University of Florida in the summer of 2012. After graduation, Donald relocat ed to Morristown, T N where he beg an his career as a Development Engineer w ith Alcoa.