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Development, characterization and degradation behavior of cubic zirconia doped with rare earth oxides… Siebert-Timmer, Audrey Joy Corrine 2014

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DEVELOPMENT, CHARACTERIZATION AND DEGRADATION BEHAVIOUR OFCUBIC ZIRCONIA DOPED WITH RARE EARTH OXIDES PROCESSED THROUGHSPARK PLASMA SINTERINGbyAudrey Joy Corrine Siebert-TimmerB.A.Sc., The University of British Columbia, 2012A THESIS SUBMITTED IN PARTIAL FULFILLMENT OFTHE REQUIREMENTS FOR THE DEGREE OFMASTER OF APPLIED SCIENCEinTHE COLLEGE OF GRADUATE STUDIES(Mechanical Engineering)THE UNIVERSITY OF BRITISH COLUMBIA(Okanagan)July 2014© Audrey Joy Corrine Siebert -Timmer, 2014iiAbstractStabilized zirconia ceramics are candidate materials for high temperature applications due totheir chemical inertness, high thermal stability and low thermal conductivity. These ceramicsalso exhibit superior stability under radiation exposure and thus there has been a rising interest inutilizing these materials in nuclear applications, particularly Supercritical Water Reactors(SCWRs). Compared to present day nuclear technologies, SCWRs boast improved thermalefficiencies, lower operating costs and enhanced safety. The fundamental idea behind thispromising technology is the use of supercritical water as a coolant. Existing at temperatures andpressures above 374 °C and 22.1 MPa, respectively, supercritical water is an exceptionallyaggressive and corrosive environment which many current nuclear materials cannot withstand.However, preliminary studies demonstrate that zirconia ceramics have improved stability under asupercritical water environment and, as such, there has been interest in using these ceramics as athermal and protective insulating material in the SCWR system. However, the applications ofzirconia ceramics are severely limited by their degradation behaviour. To date, there has beenlittle progress in effectively controlling this behaviour and, as such, further development isessential for the application of zirconia ceramics in supercritical water.In this research, 8 mol% Yttria Stabilized Zirconia (YSZ) as well as YSZ ceramicscontaining various amounts of CeO2 and Nd2O3 additions were fabricated via a novel techniquecalled Spark Plasma Sintering (SPS). The sintered ceramics were then subjected to a supercriticalwater environment at a temperature and pressure of 400 °C and 31 MPa, respectively. Weightloss and microstructural changes associated with supercritical water exposure were characterizedwith a Scanning Electron Microscope X-ray Energy Dispersive Spectroscopy (SEM-XEDS), X-ray Diffraction (XRD) and Optical Microscopy (OM) systems. It was found that both CeO2 andNd2O3 additives improved the degradation resistance of the YSZ ceramic. Microstructureobservations suggested that this improvement was related to the prevention of a phase change ofthe zirconia polymorphs in the CeO2 and Nd2O3 containing YSZ ceramics.iiiPrefaceThe results and discussion presented in Chapter 4 and Chapter 5 have been published in thefollowing papers:A. Siebert-Timmer, L.Bichler (2013). Degradation of SPS Fabricated YSZ and 5 mol% CeO2 +YSZ Composite in Supercritical Water. Materials Science & Technology 2013 CollectedProceedings: Unedited Section.A. Siebert-Timmer, L. Bichler (2014). Degradation of Spark Plasma Sintered (0, 5, 10 and 15mol%) CeO2 +YSZ Composites in Supercritical Water. Materials Science Forum, vol. 783-786,pp. 1093-1098.A. Siebert-Timmer, L. Bichler. Degradation Behaviour of SPS Processed Neodymia DopedCubic YSZ in Supercritical Water. Materials Science & Technology 2014 CollectedProceedings. (Submitted April 2014).A. Siebert-Timmer, K. Mondal, L. Bichler (2014). Degradation of SPS Fabricated YSZ andCeO2-YSZ Ceramics in Supercritical Water. [90% Complete]A. Siebert-Timmer, L. Bichler (2014). Degradation of SPS Fabricated YSZ and Nd2O3-YSZCeramics in Supercritical Water. [70% Complete]The relative contributions of all collaborators and listed authors are as follows. Ceramicsamples were fabricated either at the Indian Institute of Technology (Chennai, India) under thesupervision of Prof. B.S. Murty or at Quad City Manufacturing (Rock Island, USA). Furthersample preparation was conducted at the University of British Columbia– Okanagan (Kelowna,BC) by K. Akkiraju, S. Ipe Varhese, K. Robles, M. van Hanegem and A. Siebert-Timmer.Microstructure analysis as well as hardness testing of as-sintered and degraded specimens wasperformed by A. Siebert-Timmer at University of British Columbia – Okanagan (Kelowna, BC).Degradation testing was performed by A. Siebert-Timmer at the University of British Columbia(Vancouver, BC) under the supervision of Prof. E. Asselin. Inductive Coupled Plasma-MassSpectroscopy (ICP-MS) was performed by D. Arkinstall at the University of British Columbia –Okanagan (Kelowna, BC) while XRD analysis was performed by A. Lam at the University ofBritish Columbia (Vancouver, BC). Manuscripts were compiled and written by A. Siebert-ivTimmer while editing and technical input was provided by Prof. L. Bichler from the Universityof British Columbia – Okanagan (Kelowna, BC), Prof. E. Asselin from the University of BritishColumbia (Vancouver, BC) and Prof. K. Mondal from the Indian Institute of Technology(Kanpur, India).vTable of ContentsAbstract ........................................................................................................................................... iiPreface............................................................................................................................................ iiiTable of Contents ............................................................................................................................ vList of Tables .................................................................................................................................. xList of Figures ............................................................................................................................... xiiList of Abbreviations .................................................................................................................. xviiList of Symbols .......................................................................................................................... xviiiAcknowledgements....................................................................................................................... xxDedication .................................................................................................................................... xxiChapter 1: Introduction ................................................................................................................... 1Chapter 2: Literature Review.......................................................................................................... 62.1 Introduction to Zirconia ........................................................................................................ 62.2 Monolithic Zirconia .............................................................................................................. 62.3 Stabilization of Zirconia ....................................................................................................... 72.4 Stabilized Zirconia Ceramics ................................................................................................ 82.4.1 Partially Stabilized Zirconia .......................................................................................... 82.4.2 Tetragonal Zirconia Polycrystals ................................................................................. 102.4.3 Fully Stabilized Zirconia.............................................................................................. 122.5 Degradation Behaviour of Stabilized Zirconia Ceramics ................................................... 132.6 Degradation Mechanisms of Stabilized Zirconia Ceramics................................................ 142.6.1 Destabilization Mechanism.......................................................................................... 152.6.2 Stress Corrosion Mechanism ....................................................................................... 172.6.3 Oxygen Vacancy Annihilation Mechanism ................................................................. 182.7 Controlling Degradation ..................................................................................................... 19vi2.7.1 Effect of Grain Size on Degradation............................................................................ 192.7.2 Effect of Density on Degradation ................................................................................ 212.7.3 Effect of Secondary Stabilizers on Degradation.......................................................... 222.7.3.1 Effect of CeO2 Secondary Stabilizers on Degradation ......................................... 232.7.3.2 Effect of Nd2O3 Secondary Stabilizers on Degradation ....................................... 252.7.4 Effect of Secondary Phases on Degradation ................................................................ 262.7.5 Effect of Environment on Degradation........................................................................ 282.7.5.1 Effect of Temperature on Degradation ................................................................. 292.7.5.2 Effect of Pressure on Degradation ........................................................................ 312.7.5.3 Effect of pH on Degradation................................................................................. 322.8 Zirconia Ceramics for Supercritical Water Applications.................................................... 352.8.1 Properties of Supercritical water.................................................................................. 352.8.2 Degradation Behaviour of Zirconia Ceramics in Supercritical Water......................... 372.9 Fabrication of Materials through Spark Plasma Sintering .................................................. 392.9.1 Spark Plasma Sintering Methodology ......................................................................... 402.9.2 Zirconia Ceramics Fabricated through Spark Plasma Sintering .................................. 412.10 Summary........................................................................................................................... 43Chapter 3: Experimental Procedure .............................................................................................. 453.1 Precursor Powders .............................................................................................................. 453.2 Precursor Powder Blending ................................................................................................ 463.3 Spark Plasma Sintering Processing..................................................................................... 473.3.1 Sintering of YSZ and CeO2-YSZ Powders .................................................................. 473.3.2 Sintering of Nd2O3-YSZ Powders ............................................................................... 483.4 Polishing, Thermal Etching and Density Measurements .................................................... 493.5 Microhardness..................................................................................................................... 50vii3.6 X-ray Diffraction ................................................................................................................ 503.7 Microscopy, Chemical and Image Analysis ....................................................................... 513.8 Degradation Testing............................................................................................................ 513.8.1 Sample Preparation ...................................................................................................... 513.8.2 Vessel Design............................................................................................................... 523.8.3 Degradation Testing Procedure.................................................................................... 543.8.4 Weight Loss and Degradation Rate Calculations ........................................................ 553.9 Test Solution Analysis ........................................................................................................ 55Chapter 4: Results ......................................................................................................................... 564.1 Densification ....................................................................................................................... 564.1.1 Densification of YSZ Ceramics ................................................................................... 564.1.2 Densification of CeO2-YSZ Ceramics ......................................................................... 584.1.3 Densification of Nd2O3-YSZ Ceramics ....................................................................... 614.2 Degradation Rates............................................................................................................... 654.2.1 Degradation of YSZ Ceramics..................................................................................... 654.2.2 Degradation of CeO2-YSZ Ceramics........................................................................... 664.2.3 Degradation of Nd2O3-YSZ Ceramics ......................................................................... 684.3 Microhardness..................................................................................................................... 704.4 Phase Evolution .................................................................................................................. 714.4.1 Phase Evolution of YSZ Ceramics .............................................................................. 714.4.2 Phase Evolution of CeO2-YSZ Ceramics .................................................................... 724.4.3 Phase Evolution of Nd2O3-YSZ Ceramics................................................................... 754.5 Microstructure Evolution .................................................................................................... 774.5.1 Microstructure Evolution of YSZ Ceramics ................................................................ 774.5.2 Microstructure Evolution of CeO2-YSZ Ceramics ...................................................... 79viii4.5.2.1 Matrix Evolution of CeO2 -YSZ Ceramics ........................................................... 794.5.2.2 Band Evolution of CeO2 -YSZ Ceramics.............................................................. 854.5.3 Microstructure Evolution of Nd2O3-YSZ Ceramics .................................................... 924.5.3.1 Matrix Evolution of Nd2O3 -YSZ Ceramics ......................................................... 924.5.3.2 Band Evolution of Nd2O3 -YSZ Ceramics............................................................ 964.6 Solution Analysis .............................................................................................................. 102Chapter 5: Discussion ................................................................................................................. 1045.1 Spark Plasma Sintering Fabrication.................................................................................. 1045.1.1 Spark Plasma Sintering of YSZ Ceramics ................................................................. 1045.1.2 Spark Plasma Sintering of CeO2-YSZ Ceramics ....................................................... 1055.1.2.1 Effect of CeO2 on Sintered Grain Size and Densification .................................. 1055.1.2.2 Effect of CeO2 on Phase Stabilization................................................................. 1065.1.3 Spark Plasma Sintering of Nd2O3-YSZ Ceramics ..................................................... 1075.1.3.1 Effect of Nd2O3 on Sintered Grain Size and Densification ................................ 1085.1.3.2 Effect of Nd2O3 on Phase Stabilization............................................................... 1095.2 Degradation Behaviour ..................................................................................................... 1105.2.1 Degradation Behaviour of YSZ Ceramics ................................................................. 1105.2.2 Degradation Behaviour of CeO2-YSZ Ceramics ....................................................... 1125.2.2.1 Effect of CeO2 on Degradation Rates ................................................................. 1125.2.2.2 Effect of CeO2 on Phase Stability....................................................................... 1135.2.3 Degradation Behaviour of Nd2O3-YSZ Ceramics ..................................................... 1145.2.3.1 Effect of Nd2O3 on Degradation Rates ............................................................... 1145.2.3.2 Effect of Nd2O3 on Phase Stability ..................................................................... 116Chapter 6: Conclusion................................................................................................................. 1186.1 Fabrication of Novel Zirconia Ceramics........................................................................... 118ix6.2 Degradation Behaviour of Novel Zirconia Ceramics ....................................................... 118Chapter 7: Future Work .............................................................................................................. 1207.1 Spark Plasma Sintering of Novel Zirconia Ceramics ....................................................... 1207.2 Degradation of Zirconia Ceramics in Supercritical Water ............................................... 1207.3 Future Studies for SCWR Applications............................................................................ 121References................................................................................................................................... 122Appendices.................................................................................................................................. 130Appendix A: Lee-Kesler Diagram .......................................................................................... 130xList of TablesTable 2.1 Volume percent of monoclinic phase present in 3 mol% yttria stabilized zirconiasamples after degradation testing in (1) a vacuum, (2) a vacuum after an initialtreatment in water vapour and (3) water vapour [19]. ................................................ 15Table 2.2 Formation of monoclinic zirconia during degradation testing of 3 mol% yttriastabilized zirconia in various solvents at 95 °C for 120 hrs [46]. ............................... 33Table 2.3 Formation of monoclinic zirconia on the surface of various zirconia ceramics afterdegradation testing in acidic solutions at a temperature of 390 °C and a pressure of27 MPa for 50 hrs [68]................................................................................................ 34Table 3.1 Chemical composition and particle size of precursor powders. ................................... 45Table 3.2 Equivalent wt% of CeO2 and Nd2O3 additions in the CeO2-YSZ and Nd2O3-YSZpowder blends. ............................................................................................................ 47Table 3.3 SPS processing parameters for YSZ and CeO2-YSZ samples...................................... 48Table 3.4 SPS processing parameters for Nd2O3-YSZ samples. .................................................. 49Table 3.5 Non-converging pressure values calculated for supercritical water degradationtesting at 400 °C in the designed test vessel containing 20 mL of water. ................... 54Table 4.1 Measured grain size and density of as-sintered YSZ ceramics. ................................... 57Table 4.2 Measured grain size and density of as-sintered CeO2-YSZ ceramics. ......................... 59Table 4.3 Measured grain size and density of as-sintered Nd2O3-YSZ ceramics......................... 62Table 4.4 Calculated weight loss and degradation rates for YSZ ceramics.................................. 66Table 4.5 Calculated weight loss and degradation rates for CeO2-YSZ ceramics........................ 67Table 4.6 Calculated weight loss and degradation rates for Nd2O3-YSZ ceramics...................... 69Table 4.7 Lattice parameter measurement of the YSZ ceramics before and after degradationtesting. ......................................................................................................................... 72Table 4.8 Lattice parameter measurement of the CeO2-YSZ ceramics before and afterdegradation testing. ..................................................................................................... 74Table 4.9 Lattice parameter measurement of the Nd2O3-YSZ ceramics before and afterdegradation testing. ..................................................................................................... 76Table 4.10 Chemical analysis of the YSZ ceramic before and after degradation testing. ............ 79Table 4.11 Chemical analysis of the as-sintered matrix in the CeO2-YSZ ceramics. .................. 80Table 4.12 Chemical analysis of CeO2-YSZ ceramics after degradation testing. ........................ 82xiTable 4.13 Chemical analysis of the Ce-rich oxide layer formed on the CeO2-YSZ ceramicsafter degradation testing.............................................................................................. 84Table 4.14 Chemical analysis of the spalled surface in CeO2-YSZ ceramics after degradationtesting. ......................................................................................................................... 85Table 4.15 Chemical analysis of Ce-rich bands in the as-sintered CeO2-YSZ ceramics. ............ 87Table 4.16 Chemical analysis of remaining Ce-rich bands in 15 mol% CeO2-YSZ ceramicafter degradation testing.............................................................................................. 91Table 4.17 Chemical analysis of the Ce-rich areas found in the CeO2-YSZ ceramics afterdegradation testing. ..................................................................................................... 92Table 4.18 Chemical analysis of the as-sintered microstructure of the Nd2O3-YSZ ceramics..... 93Table 4.19 Chemical analysis of the Nd2O3-YSZ ceramics after degradation testing.................. 96Table 4.20 Chemical analysis of crystals formed in the 10 mol% Nd2O3-YSZ ceramic afterdegradation testing. ..................................................................................................... 96Table 4.21 Chemical analysis of Nd bands found in the as-sintered 10 mol% Nd2O3-YSZceramic. ....................................................................................................................... 98Table 4.22 Chemical analysis of white bands found in 10 mol% Nd2O3-YSZ ceramic afterdegradation testing. ................................................................................................... 101Table 4.23 ICP-MS results from solutions retrieved after degradation testing. ......................... 103xiiList of FiguresFigure 1.1 Proposed HEC design for SCWRs [4]........................................................................... 3Figure 1.2 Research overview......................................................................................................... 5Figure 2.1 Crystal polymorphs of zirconia [20].............................................................................. 7Figure 2.2 Phase relations in 0-30 mol% MgO-ZrO2 stabilized system [18]. ................................ 9Figure 2.3 Transformation toughening mechanism in PSZ ceramics [15]. .................................. 10Figure 2.4 Phase relations in 0-10 mol% Y2O3-ZrO2 stabilized system [15]. .............................. 12Figure 2.5 Self-accommodating variants and autocatalytic propagation of the monoclinicphase in tetragonal zirconia [19]................................................................................. 14Figure 2.6 Yttrium-rich crystallites (indicated by arrows) formed in a 6 mol% yttria stabilizedzirconia ceramic after degradation testing at 250 °C for 18 hrs in water vapour[35].............................................................................................................................. 16Figure 2.7 Stress corrosion mechanism through the formation of Zr-O-Zr bonds [46]. .............. 17Figure 2.8 Effect of grain size on the monoclinic phase formation in 3 mol% Y2O3-ZrO2ceramics after degradation testing at a temperature and pressure of 200 °C and16 bar, respectively [11]. ............................................................................................ 21Figure 2.9 Effect of density and grain size on the degradation resistance of 3 mol% Y2O3-ZrO2after degradation testing in water at 260 °C for 8 hrs [12].......................................... 22Figure 2.10 Effect of CeO2 on the formation of monoclinic zirconia after degradation testingin air at 200 °C [8]....................................................................................................... 24Figure 2.11 Effect of Nd2O3 on the formation of monoclinic zirconia on the surface of yttriastabilized zirconia ceramics after degradation testing at a temperature and pressureof 200 °C and 1.55 MPa, respectively [9]. .................................................................. 26Figure 2.12 Effect of secondary alumina additions on the formation of monoclinic zirconia inzirconia toughened alumina ceramics [61]. ................................................................ 28Figure 2.13 Effect of temperature on the degradation of yttria stabilized zirconia ceramics[62].............................................................................................................................. 30Figure 2.14 Effect of pressure on the tetragonal to monoclinic transformation in 3 mol%Y2O3-ZrO2 ceramics [66]. ........................................................................................... 32Figure 2.15 Effect of LiOH concentrations on the degradation of 3 mol% yttria stabilizedzirconia ceramics in water at a temperature of 200 °C for 2 hrs [69]. ........................ 35xiiiFigure 2.16 Effect of temperature and pressure on the density and ionic product of water [3]. .. 37Figure 2.17 Formation of needle-like particles (circled) on the surface of yttria stabilizedzirconia ceramics with 15 wt% of TiO2 additions after supercritical waterdegradation testing [71]. ............................................................................................. 39Figure 2.18 Schematic of the SPS process [73]. ........................................................................... 41Figure 2.19 Resulting density of 8 mol% yttria stabilized zirconia ceramics sintered withdifferent sintering techniques and at different temperatures [78]. .............................. 42Figure 2.20 Average grain size of 8 mol% yttria stabilized zirconia ceramics sintered withdifferent sintering techniques and at different temperatures [78]. .............................. 42Figure 3.1 SEM micrographs of precursor powders: a) YSZ, b) CeO2 and c) Nd2O3.................. 46Figure 3.2 As-sintered YSZ and CeO2-YSZ ceramics: a) pure YSZ, b) 5 mol% CeO2-YSZ,c) 10 mol% CeO2-YSZ and d) 15 mol% CeO2-YSZ. ................................................. 48Figure 3.3 As-sintered Nd2O3-YSZ ceramics: a) 5 mol% Nd2O3-YSZ and b) 10 mol%Nd2O3-YSZ. ................................................................................................................ 49Figure 3.4 Supercritical water degradation test vessel.................................................................. 52Figure 4.1 Fractured as-sintered YSZ ceramic. ............................................................................ 57Figure 4.2 SEM micrograph of grain structure in the as-sintered YSZ ceramic. ......................... 57Figure 4.3 XRD plots for the YSZ powder and as-sintered YSZ ceramic. .................................. 58Figure 4.4 Fractured disk showing dark and orange discolorations in as-sintered 10 mol%CeO2-YSZ ceramic. .................................................................................................... 59Figure 4.5 SEM micrographs of grain structure in as-sintered YSZ ceramics with: a) 5 mol%CeO2, b) 10 mol% CeO2 and c) 15 mol% CeO2 additions.......................................... 60Figure 4.6 XRD plots for YSZ powder, CeO2 powder and as-sintered CeO2-YSZ ceramics. ..... 61Figure 4.7 As-sintered 10 mol% Nd2O3-YSZ ceramic. ................................................................ 62Figure 4.8 SEM micrographs of grain structure in as-sintered 5 mol% Nd2O3-YSZ ceramic:a) matrix grains and b) textured grains. ...................................................................... 63Figure 4.9 SEM micrographs of the grain structure in as-sintered 10 mol% Nd2O3-YSZceramic: a) matrix grains, b) nano grains (indicated by arrows) and c) texturedgrains........................................................................................................................... 64Figure 4.10 XRD plots for YSZ powder, Nd2O3 powder and as-sintered Nd2O3-YSZ ceramics..................................................................................................................................... 65xivFigure 4.11 Images of the sectioned YSZ sample before and after degradation testing. ............. 66Figure 4.12 Images of the sectioned CeO2-YSZ samples before and after degradation testing.Areas of spallation are indicated by an arrow............................................................. 68Figure 4.13 Images of the sectioned Nd2O3-YSZ samples before and after degradation testing..................................................................................................................................... 69Figure 4.14 Measured microhardness for the various sintered ceramics before and afterdegradation testing. ..................................................................................................... 70Figure 4.15 XRD plots for YSZ ceramics before and after degradation testing........................... 71Figure 4.16 XRD plots of CeO2-YSZ ceramics before and after degradation testing.................. 74Figure 4.17 XRD plots for Nd2O3-YSZ ceramics before and after degradation testing............... 76Figure 4.18 General microstructure of the YSZ ceramic: a) before and b) after degradationtesting. ......................................................................................................................... 78Figure 4.19 SEM micrographs of the YSZ ceramic: a) before and b) after degradation testing. . 78Figure 4.20 SEM micrographs of the as-sintered CeO2-YSZ ceramics: a) 5 mol% CeO2-YSZ,b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. ................................................. 80Figure 4.21 Linescan across a Ce-rich agglomerate in the 5 mol% CeO2-YSZ ceramic ............. 81Figure 4.22 SEM micrographs of CeO2-YSZ ceramics after degradation testing: a) 5 mol%CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ............................... 82Figure 4.23 SEM micrographs of the Ce-rich oxide layer formed on the CeO2-YSZ ceramicsafter degradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15mol% CeO2-YSZ.. ...................................................................................................... 83Figure 4.24 SEM micrographs of spalled surfaces in the CeO2-YSZ ceramics afterdegradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ andc) 15 mol% CeO2-YSZ ............................................................................................. 85Figure 4.25 General microstructure of the as-sintered CeO2-YSZ ceramics: a) 5 mol%CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ............................... 86Figure 4.26 Interface region surrounding Ce-rich bands in as-sintered CeO2-YSZ ceramics:a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ............. 87Figure 4.27 Linescan across a CeO2 band in the as-sintered 10 mol% CeO2-YSZ ceramic. ....... 88Figure 4.28 Void formation (arrows) at the tips of an CeO2 band in the as-sintered 15 mol%CeO2-YSZ ceramic. .................................................................................................... 89xvFigure 4.29 Ce-rich bands (rectangles) and Ce-rich areas (arrows) found in CeO2-YSZceramics after degradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZand c) 15 mol% CeO2-YSZ......................................................................................... 90Figure 4.30 Crack formation from Ce-rich bands in 15 mol% CeO2-YSZ ceramic afterdegradation testing. ..................................................................................................... 90Figure 4.31 Linescan across a Ce-rich band in the 15 mol% CeO2-YSZ ceramic afterdegradation testing. ..................................................................................................... 91Figure 4.32 Ce-rich areas in CeO2-YSZ ceramics after degradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ......................................... 92Figure 4.33 SEM micrographs of as-sintered Nd2O3-YSZ ceramics: a) 5 mol% Nd2O3-YSZand b) 10 mol% Nd2O3-YSZ....................................................................................... 93Figure 4.34 Linescan across a Nd-rich agglomerate in the matrix of the as-sintered 5 mol%Nd2O3-YSZ ceramic.................................................................................................... 94Figure 4.35 SEM micrographs of Nd2O3-YSZ ceramics after degradation testing: a) 5 mol%Nd2O3-YSZ and b) 10 mol% Nd2O3-YSZ. ................................................................. 95Figure 4.36 Crystal formation in 10 mol% Nd2O3-YSZ ceramic after degradation testing:a) along pores and b) along crack interphases. ........................................................... 96Figure 4.37 General microstructure of the as-sintered Nd2O3-YSZ ceramics: a) 5 mol%Nd2O3-YSZ and b) 10 mol% Nd2O3-YSZ.. ................................................................ 97Figure 4.38 Linescan across a white band in the as-sintered 10 mol% Nd2O3-YSZ ceramic. ..... 98Figure 4.39 Linescan across a grey band found in the as-sintered 10 mol% Nd2O3-YSZceramic. ....................................................................................................................... 99Figure 4.40 SEM micrographs of 10 mol% Nd2O3-YSZ ceramics after degradation testing..... 100Figure 4.41 Linescan across a white band in a 10 mol% Nd2O3-YSZ ceramic afterdegradation testing. ................................................................................................... 101Figure 4.42 Retrieved solutions after degradation testing: a) YSZ, b) CeO2-YSZ andc) Nd2O3-YSZ ceramics. ........................................................................................... 103Figure 5.1 Lattice shift present in the as-sintered CeO2-YSZ ceramics. .................................... 107Figure 5.2 Lattice shift present in the as-sintered Nd2O3-YSZ ceramics. .................................. 110Figure 5.3 Crack propagation in the 10 mol% Nd2O3-YSZ ceramic after degradation testing:a) general overview and b) formation of microcracks from Nd-rich white bands.... 115xviFigure A.0.1 Fluid compressibility factor as a function of reduced temperature and reducedpressure [98].............................................................................................................. 130xviiList of AbbreviationsAECL Atomic Energy Canada Ltd.AES Auger Electron SpectroscopyEPR Electron Paramagnetic ResonanceF/M Ferritic-martensiticFSZ Fully Stabilized ZirconiaHEC High Efficiency ChannelHIP Hot Isostatically PressedICP-MS Inductively Coupled Plasma – Mass SpectroscopyN-G Nucleation and GrowthOM Optical MicroscopyPSZ Partially Stabilized ZirconiaSAD Selected-Area DiffractionSCWR Supercritical Water ReactorSEM-XEDS Scanning Electron Microscope X-Ray Energy Dispersive SpectroscopySIMS Secondary Ion Mass SpectrometrySPS Spark Plasma SinteringTEM Transmission Electron MicroscopyTZP Tetragonal Zirconia PolycrystalsXPS X-ray Photoelectron SpectroscopyXRD X-ray DiffractionYSZ 8 mol% Yttria Stabilized ZirconiaZAF Atomic Number (Z) Absorption (A) and Fluorescence (F) CorrectionsxviiiList of SymbolsA Initial Surface AreaCe CeriumCe3+ Cerium Ion (+3 charge)Ce4+ Cerium Ion (+4 charge)CeO2 Cerium Dioxide or CeriaCl ChlorineHf HalfniniumK Unit Conversion Constantkw Ionic Productn MolesNd NeodymiumNd2O3 Neodymium Oxide or NeodymiaNd3+ Neodymium IonO OxygenO2- Oxygen IonP Pressurer Degradation RateR Gas ConstantR2 Linearity CoefficientT Absolute Temperaturet Timet’ Metastable Dopant-rich Tetragonal PhaseV VolumeW Final Weight LossxixY YttriumY2O3 Yttrium Oxide or YttriaY3+ Yttrium Ionz Compressibility FactorZr ZirconiumZr4+ Zirconium IonZrO2 Zirconium Dioxide or Zirconiaρ DensityxxAcknowledgementsFirst and foremost, I wish to acknowledge my supervisors Dr. Lukas Bichler and Dr. Ed Asselinfor their support and guidance over the past couple of years. To Dr. Lukas Bichler: thank-you foryour relentless support in writing recommendation letters, funding applications, and editingnumerous journal papers and thesis drafts. Without your expertise and attention to detail thiswork would not have been possible. To Dr. Ed Asselin: thank-you for your valuable inputthroughout the course of this study. You have helped enormously in developing the testingprocedures as well as refining several of the analysis techniques used.In addition to my supervisors, I would also like to acknowledge Dr. Kallol Mondal andthe research group at the Institute of India Madras particularly Karthik Akkiraju and Stephen IpeVarheese for their assistance in fabricating the experimental materials.Furthermore, I would like to recognize all the incredible lab technicians that assisted meover the course of my work particularly David Arkinstall for his assistance in SEM and ICP-MSanalysis as well as Anita Lam for performing XRD on my many, many, many samples.I would like to thank all of my colleagues who I have had the privilege to work with fortheir help, support and camaraderie. I would especially like to acknowledge the unwaveringsupport and encouragement from Karen Robles, Scott Forsyth and JD Fraser.On a more personal note, I am grateful to my family for their sacrifice, encouragementand love they provided me with during this long journey. Without their support, completion ofthis work would not be possible.And finally, I would like to acknowledge the financial support received from the NaturalScience and Engineering Research Council of Canada.xxiDedicationTo my Father,whose unfailing love and support hasmade me the woman I am today.I love you dad.1Chapter 1: IntroductionThis chapter provides an overview of Generation IV nuclear reactors with a focus on materialrequirements. A review of zirconia as a candidate material is presented followed by a briefsummary of the research carried out in this thesis.The continuous pursuit of new and innovative technologies relies on the availability ofadvanced materials. Progress in material development, synthesis and processing has created vasttechnological opportunities that were unthinkable a few short decades ago. Of particularimportance today is the development of efficient, sustainable and economical energy sources.The need for alternative energy is fuelled by the increasing number of environmental issues oursociety is facing. Our dependence on fossil fuels has led to issues including greenhouse gasemissions, global warming and resource shortages. Now, more than ever, the wealth of Canadarelies on finding a safe, sustainable and economical energy source. The new generation ofnuclear power plants being developed by Atomic Energy of Canada Ltd. (AECL) has thepotential to answer this demand; however, significant material research and development isnecessary.Supercritical Water Reactors (SCWRs) being developed by AECL represent the nextgeneration of nuclear power. Compared to present day reactors, SCWRs will operate at higherefficiencies and, due to simplified components, will have a 35% decrease in operating costs [1].Most importantly, SCWR technology will have the ability to reprocess existing nuclear wasteand will produce a byproduct with safe radioactive levels after only decades rather than millennia[2]. The fundamental method behind this promising new technology is the utilization ofsupercritical water as a coolant. Supercritical water exists at temperatures and pressures above374 °C and 22.1 MPa, respectively. Above this critical point, the liquid and gas phase of waterbecome indistinguishable from each other and a dense gas is formed [3]. The benefits of usingsupercritical water in nuclear reactors stem from a significant increase in coolant enthalpy aswell as the elimination of coolant phase changes. These advantages allow for higher systemefficiencies as well as simplified plant designs [1]. However, the combination of hightemperatures and pressures, along with strong doses of radiation, results in an exceptionallyaggressive and corrosive environment particularly for in-core components.2The reactor core of a SCWR plant consists of a series of fuel channels that runhorizontally through a moderator fluid (a medium which controls the speed of the neutrons toallow for a continuous nuclear reaction). Traditionally, these fuel channels are composed of twoconcentric tubes separated by an air gap: (1) the inner pressure tube and (2) the outer calandriatube. The inner pressure tube contains uranium fuel rods as well as the high pressure coolant.The air gap and the surrounding calandria tube act as a thermal barrier separating the highertemperature pressure tube (~100 °C) from the moderator fluid (which operates at an averagetemperature of 80 °C) [4]. However, due to the elevated temperatures and pressures ofsupercritical water, this current fuel channel design will be inadequate for future SCWRs. As aresult, a High Efficiency Channel (HEC) shown in Figure 1.1 has been proposed. This new fuelchannel design eliminates the calandria tube and, instead, an insulating material is placed alongthe inner wall of each pressure tube. This insulating material must provide sufficient temperatureinsulation from the supercritical coolant to the moderator as each pressure tube is in directcontact with the moderator fluid. Therefore, the selected insulating material must exhibit a lowthermal conductivity as well as be corrosion resistant to supercritical water at varioustemperatures. In addition, the material should be dimensionally stable during thermal cycling andirradiation exposure. The selected material does not need to be high strength as the coolantpressure is transferred to the pressure tube wall. However, the insulator must be able to withstandthe weight of the fuel bundles. The HEC is designed to have a life span of 30 years [4].Therefore, the integrity of the insulating material must be evaluated according to these long-termoperation requirements. Also, the fabrication and manufacturing procedures for the thermalinsulator must be considered.3Figure 1.1 Proposed HEC design for SCWRs [4].One candidate material for the HEC insulator is a ceramic called zirconia (ZrO2).Compared to the traditional nuclear materials, zirconia ceramics have a low thermal conductivityas well as a superior thermal and chemical stability in aggressive environments [5]. Furthermore,zirconia ceramics exhibit excellent stability under radiation exposure and have a low neutroncross-section [4]. However, the application of zirconia is severely limited by a detrimental phasetransformation, which results in the rapid deterioration of mechanical properties. Studies inliterature have found that the rate of this degradation behaviour is affected by a variety of factorsincluding the microstructure and chemical composition of the zirconia ceramic as well as the testenvironment [6]. Despite these discoveries, there has been little progress in effectivelycontrolling this behaviour and, as such, further development is essential for the application ofzirconia ceramics in supercritical water environments. Literature suggests that rare earth oxidesadditions can significantly improve the degradation resistance of zirconia ceramics [7]-[10].  Inaddition, it has been found that controlling grain size and porosity through specialized processingtechniques can also delay the degradation process [11], [12].In the present research, moderately dense 8 mol% Yttria Stabilized Zirconia (YSZ) aswell as YSZ ceramics containing various amounts of either cerium dioxide (CeO2) orneodymium oxide (Nd2O3) were fabricated via Spark Plasma Sintering (SPS). Sintered ceramicspecimens were then subjected to static degradation testing in supercritical water at 400 °C and31 MPa. Weight loss and microstructural changes associated with supercritical water exposure4were characterized with a Scanning Electron Microscope X-Ray Energy DispersiveSpectroscopy (SEM-XEDS), X-Ray Diffraction (XRD) and Optical Microscopy (OM) systems.This thesis document is structured as follows:Chapter 2 presents a literature review on the status quo of zirconia ceramics includingproposed degradation mechanisms as well as a brief overview of the SPS process.Chapter 3 describes the experimental procedures used to fabricate the various ceramicspecimens. Also, methods used to characterize materials and perform static supercritical watertests are described.Chapter 4 presents the results of SPS processing and supercritical degradation testingincluding the phase and microstructure evolution of the materials before and after degradationtesting.Chapter 5 discusses and correlates the results to other findings in literature.Chapter 6 provides the conclusions from this research.Chapter 7 presents topics and ideas suitable for future research.Chapter 8 provides the references used throughout this thesis.A schematic of the scope of this research is provided in Figure 1.2.5Spark PlasmaSintering (SPS)Yttria StabilizedZirconia (YSZ)CeramicsCeO2-YSZ CeramicsCharacterizationDensificationBehaviour During SPSDegradation RatesDuring SupercriticalWater ExposurePhase Evolution DuringSupercritical WaterExposureMicrostructureEvolution DuringSupercritical WaterExposureTest SolutionAnalysisNd2O3-YSZ CeramicsFigure 1.2 Research overview.6Chapter 2: Literature ReviewThis chapter introduces the properties of zirconia ceramics including a comprehensive literaturereview on the stabilization and degradation of these ceramics. In addition, a section on thefabrication of zirconia ceramics through Spark Plasma Sintering (SPS) is included.2.1 Introduction to ZirconiaCeramics are known for their superior stability in aggressive environments and, as such, havebecome of particular interest for use in supercritical water applications [13]. Boukis et al. [14]evaluated the performance of various ceramic materials exposed to supercritical water. HotIsostatically Pressed (HIP) BN, B4C, TiB2 ceramics were found to disintegrate completely whileSiC- and Si3Ni4-based ceramics showed significant weight loss of up to 90% [14]. Incomparison, alumina- and zirconia-based ceramics were found to have improved stability undera supercritical water environment and did not degrade severely. Compared to alumina, zirconiahas been found to exhibit enhanced fracture toughness properties with values up to 10 MPa m0.5which is double the toughness of most ceramic materials [15]. This superior fracture toughnessof zirconia has led to numerous developments in characterizing this strength. Furthermore, due toadditional properties of low neutron cross-section, low thermal conductivity and high stability insupercritical water, there has been increasing interest in using zirconia in SCWRs as aninsulating material [4].2.2 Monolithic ZirconiaZirconia (ZrO2) is the oxide form of the transitional metal zirconium. At room temperaturezirconia favours a 7-fold oxygen coordination number forming a monoclinic crystal structure asshown in Figure 2.1a [16]. Often referred to as Baddeleyite, the formation of this monocliniccrystal structure is attributed to large ionic size difference between the Zr4+ and O2- ions whichresults in a combination of ionic and covalent bonding between the Zr4+ cations and O2- anions[5], [16]. With increasing temperature above ~1100 °C, monoclinic zirconia undergoes atransformation to a tetragonal structure, followed by a transformation to a cubic structure above2300 °C, and finally ending with melting at 2690 °C [5], [13]. Classified as martensitic, thesetransformations involve shifts in lattice positions rather than the diffusion of atoms. The highertemperature tetragonal and cubic polymorphs have been found to favor a more symmetric 8-fold7oxygen coordination number which, due to oxygen overcrowding, strains the cation network andchanges the volume of the unit cell [17]. Of particular importance is the 4.7% volume increaseduring tetragonal to monoclinic transformations, which occurs upon cooling [18]. This volumechange results in intergranular microcracking and ultimately the disintegration of the ceramicthrough grain fall-out and fracture, limiting the practical use of pure zirconia [19]. A deleteriousvolume change is not observed in cubic to tetragonal transformations, since the volume of theunit cells are relatively similar in these two phases. The crystal structures for the three phases ofzirconia are shown in Figure 2.1 where the corresponding densities for the monoclinic, tetragonaland cubic phases are 5.65, 6.10 and 6.27 g/cm3, respectively [13].a) Monoclinic b) Tetragonal c) CubicFigure 2.1 Crystal polymorphs of zirconia [20].2.3 Stabilization of ZirconiaBy doping zirconia with oxides of lower-valency cations, the tetragonal and cubic polymorphs ofzirconia can be stabilized at room temperature [5]. During stabilization, the added cationssubstitute for the Zr4+ ions in the lattice and, to maintain electrostatic neutrality, generate oxygenvacancies in the sublattice [16]. It has been found that it is the generation of these oxygenvacancies that control the room temperature phase stabilization of tetragonal and cubic zirconiarather than the substitutional dopant cation [21], [22]. Li et al. [17] explained the mechanism ofstabilization in terms of the relief of the internal strain of zirconia’s higher temperaturepolymorphs. In Section 2.2, it was discussed that tetragonal and cubic zirconia have an oxygencoordination number of 8, which strains the cation lattice. Through introducing oxygenvacancies, the effects of oxygen overcrowding are reduced and the internal strain in the Zr4+cation lattice is relieved, allowing for room temperature stabilization [17]. Recent studies by8Fabris et al. [22] determined that doping the zirconia matrix with oxygen vacancies only cantheoretically stabilize the tetragonal and cubic phase, where stabilized cubic zirconia had fourtimes the amount of vacancies compared to stabilized tetragonal zirconia. These results indicatethat it is the defect crystal structure surrounding the oxygen vacancies that controls the roomtemperature phase stabilization of tetragonal and cubic zirconia, rather than the defect structuresurrounding the substitutional dopant cations.Several types of stabilized zirconia systems with low valency dopants have been studiedin the literature [16]-[18], [21]. It has been found that in addition to the required low valencestate and high solid solubility in zirconia, suitable dopant cations must also have a larger ionicradius larger than Zr4+ ions (0.84 Å) [23]. Li et al. [21] reported that oversized trivalent dopantsgenerated oxygen vacancies in the Zr4+ cation lattice, while undersized trivalent dopantsgenerated oxygen vacancies within the dopant lattice. These constraints limit the number ofsuitable dopants to trivalent rare earth oxides. However, less expensive divalent additives such asCaO and MgO have also been found to be able to stabilize the higher temperature polymorphs ofzirconia [16], [18].2.4 Stabilized Zirconia CeramicsCompared to the monoclinic phase, stabilized tetragonal and cubic phases of zirconia haveenhanced mechanical properties, since the volume changes associated with heating and coolingare eliminated. As a result, stabilized zirconia ceramics have a wide range of applicationsincluding biomedical implants, solid oxide fuel cells, thermal barrier coatings as well as oxygenpumps and sensors [5]. In general, stabilized zirconia ceramics can be classified into threecategories: (1) Partially Stabilized Zirconia (2), Tetragonal Zirconia Polycrystals and (3) FullyStabilized Zirconia.2.4.1 Partially Stabilized ZirconiaPartially Stabilized Zirconia (PSZ) ceramics are often associated with ZrO2 systems stabilizedwith either CaO or MgO [24]-[26]. These stabilized zirconia ceramics exhibit a multi-phasemicrostructure consisting of tetragonal, cubic and sometimes monoclinic phases. The number,amount, size and distribution of phases present is determined by the amount of dopant added aswell as the processing parameters. For illustrative purposes the MgO stabilized zirconia system isreviewed.9The phase diagram of the zirconia-rich, MgO-ZrO2 system is presented in Figure 2.2. Asshown, cubic zirconia is stable at temperatures above 1400 °C. Tetragonal zirconia can bereached at lower temperatures of 1240 °C; however, there is little solubility of MgO in tetragonalzirconia. Studies performed by Sim and Stubican [27] showed that maximum MgO solubility intetragonal zirconia was 1 mol% at a temperature of 1350 °C. Further cooling to temperaturesbelow 1240 °C resulted in the formation of monoclinic zirconia, causing MgO to precipitate(supporting that there was limited solubility of MgO in monoclinic zirconia).Figure 2.2 Phase relations in 0-30 mol% MgO-ZrO2 stabilized system [18].An important feature of the MgO-ZrO2 phase diagram is the phase field containing bothcubic and tetragonal zirconia. Many commercially available MgO stabilized zirconia ceramicsuse this phase field to produce PSZ ceramics. MgO-PSZ ceramics are generally stabilized with 6to 8.5 mol% of MgO and are initially sintered at temperatures above 1700 °C to produce ahomogenous cubic solid solution [18]. The MgO stabilized ceramic is then rapidly cooled intothe cubic and tetragonal phase field, resulting in the nucleation and growth of tetragonalprecipitates. After a specific dwell time, the ceramic is quenched resulting in a cubic matrix withtetragonal precipitates. Similar multi-phase microstructures can be obtained with zirconia10ceramics stabilized with CaO dopants [28]. It has been found that these tetragonal precipitatesare metastable and transform to monoclinic zirconia in the presence of propagating crack tips.Due to the volume expansion, the tetragonal to monoclinic transformations apply a netcompressive stress around the crack tip as shown in Figure 2.3. This reduces the local crack tipstress field, decreases the driving force for further crack propagation, and ultimately results in anincrease of fracture toughness [15]. It is this transformation toughening mechanism that has beencredited for the high fracture toughness of PSZ ceramics.Figure 2.3 Transformation toughening mechanism in PSZ ceramics [15].2.4.2 Tetragonal Zirconia PolycrystalsTetragonal Zirconia Polycrystals (TZP) are prepared by sintering at temperatures correspondingto the tetragonal phase field. TZP are generally formed with low dopant concentrations (around 2to 2.5 mol% for zirconia ceramics stabilized with rare earth oxides) at temperatures rangingbetween 1200 to 1500 °C [29]. Rapid quenching to room temperature results in the formation ofa metastable tetragonal microstructure. Similar to the tetragonal precipitates in PSZ, thesetetragonal grains in sintered TZP ceramics are metastable and, under stress, can transform to themonoclinic phase inducing a transformation toughening effect, analogous to what was describedin Section 2.4.3.1.  In general, TZPs contain yttrium oxide (Y2O3), or yttria, as a stabilizer [18],[25], [26]. Compared to MgO and CaO, Y2O3 dopants have a higher solubility in zirconia and no11precipitates form in the cubic, tetragonal and monoclinic phase fields, as shown in Figure 2.4.Furthermore, in comparison to other rare earth stabilized systems, yttria stabilized zirconiaceramics (YSZ) systems are well established and, as a result, are commercially produced for arange of applications [30].The zirconia rich Y2O3-ZrO2 phase diagram is presented in Figure 2.4. In comparison tothe MgO-ZrO2 system, tetragonal transformations in the Y2O3-ZrO2 system occur at a muchlower temperature (~600 °C). With increasing temperature a small ellipsoidal monoclinic andtetragonal region is encountered after which a tetragonal phase field exists. It is within thisregion that TZP are sintered, generally consisting of dopant level between 1.5 to 2.5 mol% Y2O3and at temperatures between 1200 °C to 1500 °C (Figure 2.4). Increasing the Y2O3 concentrationto values larger than 2.5 mol% results in a dual phase region consisting of tetragonal and cubiczirconia. Similar to MgO and CaO stabilized zirconia systems, this tetragonal and cubic phaseregion can be used to form PSZ ceramics. Y2O3 stabilized PSZ ceramics are generally formedwith dopant concentrations ranging from 2.5 to 6 mol% and at sintering temperatures between1750 °C to 2000 °C (Figure 2.4). Further increasing Y2O3 content results in the stabilization ofcubic zirconia with room temperature stabilization achieved at concentrations greater than10 mol%.12Figure 2.4 Phase relations in 0-10 mol% Y2O3-ZrO2 stabilized system [15].2.4.3 Fully Stabilized ZirconiaIncreasing the amount of dopant results in the stabilization of cubic zirconia, often referred to asFully Stabilized Zirconia (FSZ). In the Y2O3-ZrO2 system, FSZ forms when dopantconcentrations are equal to or greater than 8 mol%. Sintering of these ceramics is performed inthe cubic phase field, generally at temperatures around 1400 °C. The sintered ceramic is thenrapidly cooled. Depending on the composition, the resulting cubic matrix may contain metastabletetragonal grains. Using the lever rule for an 8 mol% yttria stabilized zirconia ceramic, and thephase diagram presented in Figure 2.4, the resulting ceramic is expected to consist of a dopant-rich cubic matrix containing 21% of dopant-depleted tetragonal grains. This microstructure,however, is based on equilibrium conditions and in reality is rarely observed. Instead, rapidcooling inhibits the diffusion of dopant ions and as a result a dopant-rich tetragonal phase isgenerated. Due to the increase in dopant concentration this tetragonal phase is oftendistinguished as t’.  The formation of t’ phase has been found in several other trivalent stabilized13zirconia systems [31]-[33]. The t’ phase has the same crystallography as tetragonal zirconia;however, in contrast, no stress induced transformations have been observed for this dopant richphase [15], [33]. As a result this phase is often referred to as untransformable tetragonal zirconia.Due to the inability for this t’ phase transform, FSZ ceramics exhibit low fracture toughnessvalues around 2.8 MPa m0.5 [5]. However, in comparison to TZP, FSZ ceramics exhibitproperties of high ionic conductivity [5]. This is attributed to the large amount of oxygenvacancies generated in the zirconia matrix due to the high dopant concentrations. As a result,FSZ ceramics are often used in solid-oxide fuel cells, oxygen pumps and electrochemical reactorapplications [5], [24].2.5 Degradation Behaviour of Stabilized Zirconia CeramicsDegradation of stabilized zirconia involves the unwanted transformation of the stabilized phaseback to the monoclinic polymorph. To distinguish from the transformation tougheningmechanism, where transformations are induced by stress fields generated from propagatingcracks, destabilization refers to transformations initiated by environmental conditions. As firstobserved by Kobayashi et al. [34], this degradation behaviour can occur at ambient temperaturesand is enhanced by the presence of water or humidity [19], [30], [35]. Due to the rapid loss ofmechanical properties and, ultimately, to the catastrophic failure of the ceramic, thesespontaneous reverse transformations severely limit the applicability of zirconia ceramics. As aresult, numerous papers in literature have been dedicated to explaining and controlling thisbehaviour. Most of the research on degradation has focused on studying stabilized tetragonalzirconia systems as the large volume changes associated with tetragonal to monoclinictransformation significantly contributes to the disintegration of the ceramic.An in-depth analysis by Muñoz-Tabares et al. [19] proposed that degradation oftetragonal zirconia is initiated by the nucleation of a single monoclinic variant on the surface ofthe sintered ceramic.  This monoclinic variant progressively spreads through the grain until it isstopped by the grain boundary which, due to the volume difference, imposes a localized highstress region resulting in the formation of an intergranular microcrack. Further propagation of themonoclinic phase may then occur through two different methods: (1) self-accommodatingvariants and (2) autocatalytic propagation (Figure 2.5). Propagation through self-accommodatingvariants is characterized through the formation of martensitic plates arranged in bands where the14shear deformation of each variant is opposed by the equal and opposite deformation of theadjacent variant. In contrast, propagation through autocatalytic propagation involves theformation of a monoclinic variant in the adjacent grains which nucleates due to residual stressesinduced on the secondary grain by the monoclinic variants formed in the initial grain. Throughthis sequential nucleation and growth (N-G) behaviour, transformation continues to proceed fromthe surface of the sintered of the ceramic to the interior, with degradation depth increasing withtime [19].Figure 2.5 Self-accommodating variants and autocatalytic propagation of the monoclinic phase in tetragonalzirconia [19].Many studies in literature support the degradation process proposed by Muñoz-Tabares etal. [19] as common characteristics of degradation include: intergranular microcracking, self-accommodating variants and the sequential progression of degradation from the surface of theceramic to the interior [6], [30], [36], [37]. However, despite this agreement, the exactmechanism behind the nucleation of the initial monoclinic variant as well as the role of water inthe degradation process rates remains debated.2.6 Degradation Mechanisms of Stabilized Zirconia CeramicsDegradation of zirconia has been shown to be strongly enhanced by the presence of water orwater vapour. Muñoz-Tabares et al. [19] demonstrated that the presence of water facilitates thedestabilization of the tetragonal phase. It was found that specimens initially treated with water15vapour and then placed in a vacuum had slower transformation rates than specimens testedpurely in water vapour, as shown in Table 2.1. In contrast, only 5% of monoclinic phase wasfound for specimens tested under a vacuum without an initial water vapour treatment.Table 2.1 Volume percent of monoclinic phase present in 3 mol% yttria stabilized zirconia samples afterdegradation testing in (1) a vacuum, (2) a vacuum after an initial treatment in water vapour and (3) watervapour [19].Sample Environment*ExposureTime(hours)InitialMonoclinicVolume (%)FinalMonoclinicVolume (%)TransformationRate(% per hour)1 Vacuum 96 0 5 0.052 Water Vapour 10 0 25 2.50Vacuum 96 25 50 0.263 Water Vapour 60 0 100 1.70*all tests were performed at a temperature of 131°C. Testing performed under a vacuum was at a pressureof ~10-3 atm while testing performed in water vapour was at a pressure of 0.2 MPa.In comparison to tetragonal zirconia, the degradation of cubic zirconia seems to beunaffected by the presence of water [38]. Secondary Ion Mass Spectrometry (SIMS) studiesperformed by Duong et al. [39] demonstrated that during exposure to water moisture-relatedspecies (O2-, H+ , or OH- ions) diffuse into yttria stabilized zirconia ceramics where the diffusionrates of these species depend on the stabilized phase. The presented results suggested that themoisture diffusivity into cubic zirconia was lower than diffusivity into tetragonal zirconia which,in turn, was lower than the diffusivity into monoclinic zirconia [39]. Duong et al. [39] linked thisbehaviour to the improved stability of cubic stabilized zirconia ceramics indicating that waterplays a key role in the degradation of zirconia ceramics. Several proposed degradationmechanisms are reviewed in the following sections.2.6.1 Destabilization MechanismTransmission Electron Microscopy (TEM) and Selected-Area Diffraction (SAD) observations byLange et al. [35] showed the formation of yttrium-rich crystallites in a 6 mol% Y2O3-ZrO2ceramic after degradation testing at 250 °C in water vapour (Figure 2.6). Lange et al. [35]attributed the formation of these Y(OH)3 or Y(O)OH crystallites to the reaction of water with theyttrium stabilizer found at the surface of the sintered ceramic. The formation of these crystallitesreduced the content of yttrium in the tetragonal matrix initiating the nucleation of the monoclinicphase. Further depletion of yttrium continued the growth of the monoclinic nuclei until they16reached a critical size after which spontaneous tetragonal to monoclinic transformations wouldoccur without further yttrium depletion. The volume expansion associated with thesespontaneous tetragonal to monoclinic transformations resulted in the formation of microcracksalong grain boundaries which opened up new surfaces, allowing the further reaction of theceramic with water. Several separate studies involving Auger Electron Spectroscopy (AES),Electron Paramagnetic Resonance (EPR) and X-ray Photoelectron Spectroscopy (XPS) supportthis proposed destabilization mechanism [40]-[43]. However, some researchers claim theformation of yttrium hydroxides is a complementary reaction rather than the degradationinitiation. For example, Houges et al. [44] only found the presence of Y(OH)3 crystallites aftersevere hydrothermal treatment  in an autoclave for 170 hrs at 170°C.Figure 2.6 Yttrium-rich crystallites (indicated by arrows) formed in a 6 mol% yttria stabilized zirconiaceramic after degradation testing at 250 °C for 18 hrs in water vapour [35].172.6.2 Stress Corrosion MechanismStress corrosion of glasses and ceramics was initially suggested by Michalske and Freiman [45]while studying the interaction of water with crack-tip bonds in amorphous silica. It was proposedthat, during exposure to ambient environments, the gradual strength loss and resulting delayedfailure of most glasses and oxide ceramics could be attributed to the slow growth of pre-existingflaws on the surface due to the dissolution of water at the at the flaw crack tip.  This mechanismwas adapted by Sato and Shimada [46] who suggested that water vapour reacted with Zr-O-Zrbonds at the surface of the sintered ceramic to form Zr-OH bonds (Figure 2.7). The formation ofZr-OH bonds generates localized stress sites, which could be released either though crackpropagation or a monoclinic transformation.Figure 2.7 Stress corrosion mechanism through the formation of Zr-O-Zr bonds [46].A similar stress corrosion mechanism was proposed by Yoshiumura et al. [47]. XRDstudies performed on tetragonal zirconia ceramics stabilized with 2.66 mol% of yttria revealedthat during degradation the crystal lattice of the sintered ceramic increased. Restoration of theinitial lattice parameters could be achieved after annealing the sample in air at a temperature of1300 °C for a period of 1 h. Infrared spectroscopy, as well as Raman spectroscopy, identified this18reversible characteristic to be related to the inclusion and exclusion, respectively, of the OH- ionsin the lattice [47]. Degradation due to yttrium depletion was dismissed that as the authors arguedthat the restoration of the lattice would not be observed at low temperatures due to the slowdiffusion of Y3+ ions compared to the migration of OH- ions. Based on these results Yoshimuraet al. [47] suggested that degradation of the stabilized zirconia ceramic occurred due to thedissociation of water at the surface of the ceramic and the formation Zr-OH and/or Y-OH bonds.Similar to Sato and Shimada [46], Yoshimura et al. [47] proposed that the formation of thesebonds resulted in the generation of local stressed sites. These local stress sites contributed(simultaneously with the motion of OH- ions in the lattice by occupying oxygen vacancies) to thenucleation of the monoclinic phase.  However, Guo [48] argued that, if water vapour reactedwith either yttrium or zirconium, basic and acid additions to the test solution would significantlyvary the resulting degradation rates. Nevertheless, both Y2O3 and ZrO2 have been proven to bevery stable under basic and acidic conditions [48].2.6.3 Oxygen Vacancy Annihilation MechanismTo maintain the electrostatic neutrality of oxide doped zirconia, oxygen vacancies are generated[16]. For example, considering Y2O3 stabilized zirconia, one oxygen vacancy is generated forevery two yttrium ions. These oxygen vacancies (Vo'') may be annihilated by water molecules(H2O) resulting in the formation of proton defects ((OH)o) as shown in Equation 2.1 [49].+ + → 2( ) Equation 2.1Based on Equation 2.1 a new sequential degradation mechanism was proposed by Guo [48]: (1)chemical absorption of H2O on the surface, (2) reaction of H2O with surface O2- ions forminghydroxyl groups of OH-, (3) penetration of OH- ions into the matrix through diffusion resultingin (4) the annihilation of oxygen vacancies. The change in oxygen vacancy concentration resultsin (5) monoclinic transformations and the formation of both micro- and macrocracks whichexposes new surfaces allowing for further degradation. Many experimental observations supportthe proposed oxygen vacancy annihilation mechanism. For example, XPS studies performed byGuo [48] detected hydroxyl OH- groups on the surface of tetragonal zirconia after exposing thesintered ceramics to boiling water for 40 hrs. In addition, XPS detected higher OH-concentrations along grain boundaries suggesting that OH- ions migrate into the ceramic throughgrain boundary diffusion [50]. Separate impedance spectroscopy studies reported similar results19as grain boundary resistivity measurements were found to increase in 3 mol% yttria stabilizedzirconia ceramics when exposed to water vapour at 252 °C. The measured increase in resistivitywas correlated to the annihilation of oxygen vacancies at the grain boundaries by the diffusion ofOH- ions [50]. Alternatively, SIMS experiments performed by Duong et al. [39] indicated that,rather than the migration of OH- ions, O2- and H+ ions are the diffusing species. It was suggestedthat the diffusion of O2- and H+ ions are independent of each other, but remain linked by theproposed oxygen vacancy mechanism [39].2.7 Controlling DegradationControlling degradation is a complex matter as numerous factors contribute to the stability of thezirconia ceramic. As a result, degradation rates reported in the literature have a large degreescatter. The following section reviews the effect of five key factors on the degradation behaviourof zirconia ceramics including: grain size, porosity, secondary stabilizers, secondary phases andenvironmental conditions.2.7.1 Effect of Grain Size on DegradationGarvie [51] proposed that alternative to dopants tetragonal zirconia could be stabilized at roomtemperature though maintaining a critical crystallite size. Pure zirconia crystallites of varioussizes were prepared though annealing oxide nanopowders for 24 hrs at temperatures rangingfrom 0 to 1000 °C [51]. XRD analysis was then performed to determine both the size and phaseof the resulting crystallites. Initial oxide nanopowders were found to be stabilized in thetetragonal phase and had a mean crystallite size of 5.4 nm and 15.5 nm for powders prepared viaprecipitation and calcination, respectively. However, during annealing, these tetragonalcrystallites grew and when the crystallite size reached ~30 nm a tetragonal to monoclinic phasetransformation occurred [51]. The stabilization of the zirconia nanopowders was related to thefree energy of the crystallites, where a decrease in free energy (through a smaller crystallite size)promoted the stabilization of the tetragonal phase. Similar results were observed by Clearfield[52] where a critical crystallite size of 12 nm was determined for the stabilization of thetetragonal phase. Garvie [51] attributed the difference in critical crystallite size to theexperimental procedure as Clearfield [52] annealed his experimental materials in an aqueousenvironment.20Bulk zirconia ceramics exhibit a similar critical size behaviour, however, rather than areduction of crystallite size, a reduction of grain size promotes the stabilization the tetragonalphase. Studies performed by Maglia et al. [53] used a high-pressure field-assisted rapid sinteringmethod to produce bulk zirconia samples with grains size between 45 to 50 nm. XRD analysisfurther revealed the grains were stabilized in the tetragonal phase supporting that tetragonalzirconia ceramics can be sintered without the use of dopants through grain refinement. Themeasured critical grain size was 15-20 nm larger than the critical crystallite size, as determinedby Garvie [51]. This difference was attributed to the effect of the surrounding sintered matrixwhich applied rigid constraints along the sintered grain boundaries and opposed the volumeexpansion of transforming grains, increasing the tetragonal to monoclinic transformation energybarrier [54].Numerous reports in the literature regarding the grain size stabilization effect in zirconiaceramics were summarized by Li et al. [55]. Reported values show a large degree of scatter as,for tetragonal zirconia, reported critical crystallite sizes range from 2-35 nm, while the criticalgrain sizes range from 30-600 nm [55]. The large variation in critical sizes was attributed to theeffects of different processing and analysis techniques [55]. However, in general, a refinement ofgrain size promotes the stabilization of the tetragonal phase. Furthermore, preliminary studiesperformed by Roy and Ghose [56] found that pure zirconia powders with a mean crystallite sizeof ~20 nm were able to retain the cubic phase without additives. These results suggest that grainrefinement contributed to the stabilization of the cubic phase.The ability to stabilize the tetragonal and cubic phases of zirconia through grainrefinement can also be related to the improved degradation resistance of fine-grained zirconiamaterials. Studies performed by Muñoz-Saldaña et al. [12] indicated that fully dense 3 mol%Y2O3-ZrO2 ceramics with grain sizes <360 nm exhibited superior degradation resistance afteraging in water with temperatures ranging from 0 – 250 °C for 8 hrs. Eichler et al. [11] observedthe propagation of the monoclinic phase into the bulk ceramic was reduced for 3 mol% yttriastabilized zirconia with grain sizes of 110 and 210 nm compared to sintered specimens withgrain sizes of 380 and 480 nm grain sizes, as shown in Figure 2.8.  These results demonstratethat sintering zirconia ceramics with nanostructured grain size may further improve degradationresistance.21Figure 2.8 Effect of grain size on the monoclinic phase formation in 3 mol% Y2O3-ZrO2 ceramics afterdegradation testing at a temperature and pressure of 200 °C and 16 bar, respectively [11].2.7.2 Effect of Density on DegradationIn addition to studying the effect of grain size, Muñoz-Saldaña et al. [12] also considered theeffect of density on the degradation resistance of sintered 3 mol% Y2O3-ZrO2 ceramics. Resultsrevealed that the degradation behaviour of 3 mol% Y2O3-ZrO2 ceramics was a function of bothdensity and grain size. In general, sintered ceramics with a relative density close to 100% and agrain size smaller than 360 nm were able to resist degradation in water at 260 °C for 8 hrs.However, a decrease in the relative density of the ceramics did not necessarily result in thedegradation of the ceramic. Muñoz-Saldaña et al. [12] found that degradation was resisted for 3mol% Y2O3-ZrO2 ceramics sintered with a relative density of 94.2%. The degradation resistanceof this ceramic was attributed to the small grain size (~320 nm). Ceramics with grain sizes ≤370nm were found to completely disintegrate during testing regardless of the sintered relativedensity [12]. Furthermore, microstructure analysis of tested specimens revealed that pores actedas preferential nucleation sites for the monoclinic phase. This was attributed to the high stressconcentrations found in these areas as well as the localized increase in surface area [12]. Similar22results were observed by Masaki [57]. Zirconia ceramics stabilized with 2.5, 3 and 4 mol% Y2O3were degraded in air at temperatures ranging between 200-250 °C for 2000 hrs. Degradation wasmeasured as a decrease in strength of the sintered ceramic after testing. For the various Y2O3-ZrO2 ceramics tested, Masaki [57] plotted curves of critical densities and grain sizes, wherezirconia ceramics fabricated with values above this curve would be able to resist degradation inair for 2000 hrs at temperatures ranging between 200-250 °C. The plotted curve for 3 mol%Y2O3-ZrO2 system is included in Figure 2.9. In general, for all the Y2O3-ZrO2 systems, adecrease in grain size resulted in a decrease of critical density.  A similar plot was presented byMuñoz-Saldaña et al. [12] for the degradation resistance of 3 mol% Y2O3-ZrO2 ceramics inwater at a temperature of 260 °C for a time of 8 hrs, which is also presented in Figure 2.9. Theseresults show that degradation resistance is improved with increasing density and decreasing grainsize.Figure 2.9 Effect of density and grain size on the degradation resistance of 3 mol% Y2O3-ZrO2 afterdegradation testing in water at 260 °C for 8 hrs [12].2.7.3 Effect of Secondary Stabilizers on DegradationWhen doping ZrO2 with divalent or trivalent oxides, oxygen vacancy defects are generated tomaintain the electrostatic neutrality of the ceramic. As mentioned previously in Section 2.4.2,23yttria is considered a superior stabilizer due to both the high solubility of yttria in zirconia, aswell as the segregation of yttrium at the grain boundaries, which inhibits grain growth resultingin a fine grained, high strength ceramic [5]. However, it has been reported that degradation rateof yttria stabilized zirconia can be further reduced by doping the stabilized matrix with asecondary additive. Available literature describes the ability to enhance the mechanicalproperties of yttria stabilized zirconia by co-stabilizing the matrix with secondary stabilizers suchas CeO2 [7], [8], [10], [41], [58] or Nd2O3 [9], [59].2.7.3.1 Effect of CeO2 Secondary Stabilizers on DegradationSato et al. [10] examined the effects of several oxide additions to various yttria stabilizedzirconia systems. They found that CeO2 additions were effective in reducing the degradationbehaviour of tetragonal zirconia without compromising the mechanical properties of the ceramic.Furthermore, the amount of monoclinic phase formed for the various CeO2-Y2O3-ZrO2 systemsduring degradation testing in water at 100 °C for 7 days was found to decrease with increasingamounts of both Y2O3 and CeO2 additions. Lin et al. [8] performed a similar study on variousCeO2-Y2O3-ZrO2 ceramics with grain sizes ranging between 0.7 to 1.1 µm. It was found thatyttria stabilized zirconia systems with 4 mol% or 2 mol% of Y2O3 were able to resist degradationfor 1000 hrs in air at a temperature of 200 °C if they contained CeO2 additions ≥5.2 or ≥7.5mol%, respectively. The monoclinic content found on the surface of the various CeO2-Y2O3-ZrO2 systems is presented in Figure 2.10.  Similar results were found by Duh and Wan [7], whoobserved that both 9.5 mol% CeO2 -1.5 mol% Y2O3- ZrO2 and 9 mol% CeO2 -2 mol% Y2O3-ZrO2 ceramic systems were able to resist degradation during ageing in water at 200 °C for 21days. These results indicate that the degradation of stabilized zirconia ceramics can be controlledthrough the addition of CeO2 additives.24Figure 2.10 Effect of CeO2 on the formation of monoclinic zirconia after degradation testing in air at 200 °C.The first number in the sample name correlates to the mol% of Y2O3 while the second number correlates tothe mol% of CeO2 [8].Leach and Khan [58] proposed that the improved degradation resistance of yttriastabilized zirconia ceramics with ceria additions is due to the strong segregation of ceria alonggrain boundaries which forms a protective Ce-rich layer. This layer inhibits the reaction of waterwith the stabilizing yttrium dopant, preventing the destabilization of the zirconia matrix.Hernandez et al. [41] studied the effect of CeO2 additions on the degradation resistance oftetragonal zirconia ceramics stabilized with 2 and 3 mol% of yttria during exposure to water at170 °C for 1000 hrs. They found that the monoclinic phase formation was prevented if the 2 and3 mol% yttria stabilized zirconia ceramics contained CeO2 additions ≥6 and ≥2 mol%,respectively. XPS studies of the tested samples revealed the atomic ratio of Y/Zr in the yttriastabilized zirconia ceramics containing ceria was three times larger than the yttria stabilizedzirconia ceramics without ceria additions. The higher Y/Zr ratio indicated that CeO2 additionsprevented the dissolution of yttria from the zirconia matrix.25The increased stability of CeO2 containing yttria stabilized zirconia ceramics mayalternatively be attributed to the substitution of stabilizing Y3+ ions with Ce4+ ions in the lattice.It has been found that similar to trivalent dopants, tetravalent dopants may also stabilize thetetragonal the cubic phases of zirconia. Studies performed by Li et al. [17] observed thatoversized dopants, such as CeO2, can stabilize both the tetragonal and cubic phase through thedilation of the cation network.  In comparison to trivalent dopants, these tetravalent dopants donot generate oxygen vacancies, in the ZrO2 lattice. Thus, the improvement of degradationresistance of yttria stabilized zirconia ceramics containing CeO2 may be due to the decrease inavailable oxygen vacancies which impairs diffusion of moisture related species into the latticeand, ultimately, reduces the degradation rate.2.7.3.2 Effect of Nd2O3 Secondary Stabilizers on DegradationSalehi et al. [9] investigated the influence of Nd2O3 additions in yttria stabilized zirconiaceramics with sintered grain sizes ranging between 0.5 to 1 µm. The degradation behaviour wasstudied for two types of stabilized zirconia ceramics: (1) 1 mol% yttria zirconia ceramicscontaining either 1 or 2 mol% of Nd2O3 additions, and (2) pure yttria stabilized zirconia ceramicscontaining either 2 or 3 mol% of Y2O3. Degradation testing was performed in an autoclave at200 °C at a saturated H2O pressure of 1.55 MPa. It was found that during degradation testing theyttria stabilized zirconia ceramics containing Nd2O3 additions had a higher vol% of monocliniczirconia on the surface, however, despite the rapid formation of this phase, the Nd2O3 containingceramics demonstrated a higher stability as complete disintegration did not occur until 50 and1250 min for the yttria stabilized ceramic containing 1 and 2 mol% of Nd2O3, respectively(Figure 2.11). In comparison, the pure yttria stabilized zirconia ceramics containing 2 and 3mol% Y2O3 disintegrated after 10 and 75 min, respectively.  The improved stability of the Nd2O3containing zirconia ceramics was attributed to a reduction of in-depth degradation of the ceramicmaterial [9]. Furthermore, XRD analysis indicated that lattice parameter of the stabilized ceramicincreased with Nd2O3 additions. This increase in cell dimensions was attributed to thesubstitution of Y3+ ions with larger Nd3+ ions (the ionic radius of Y3+ and Nd3+ ions is 1.019 Åand 1.109 Å, respectively) [23]. Salehi et al. [9] proposed that the improved stability of the yttriastabilized zirconia ceramics containing Nd2O3 additions was due to the larger lattice dimensionswhich could accommodate the internal strains applied during the diffusion of moisture speciesinto the lattice. Similar results were reported by Xu et al. [59] where 1-2 mol% of Nd2O326additions in zirconia ceramics containing 1 mol% Y2O3 ceramics was found to improve thestability of the sintered tetragonal phase.Figure 2.11 Effect of Nd2O3 on the formation of monoclinic zirconia on the surface of yttria stabilized zirconiaceramics after degradation testing at a temperature and pressure of 200 °C and 1.55 MPa, respectively [9].2.7.4 Effect of Secondary Phases on DegradationThe degradation behaviour of stabilized zirconia ceramics may be further controlled byintroducing a secondary phase which constrains stabilized zirconia grains. Studies performed byBasu et al. [60] reviewed the effect of the alumina matrix on the degradation behaviour oftetragonal zirconia additions. Alumina ceramics containing 5, 10, 12, 15 and 20 vol% of 3 mol%yttria stabilized zirconia additions were conventionally sintered for 1 h at a temperature andpressure of 1550 °C and 100 MPa, respectively. Analysis of the sintered specimens revealed thatthe average grain size of the alumina matrix ranged from 2.7-5.5 µm while embedded tetragonalzirconia grains had sizes varying from 0.3-0.5 µm. The refinement of both the alumina andtetragonal zirconia grains was associated with increasing vol% of the 3 mol% yttria stabilizedzirconia additions [60]. Degradation testing of the zirconia toughened alumina ceramic wasperformed in an air-steam mixture at a constant humidity level at temperatures of 150, 200, and250 °C at a flow rate of 10 mL/min. All specimens revealed rapid surface degradation within the27first 6 hours of testing. However, compared to the pure 3 mol% yttria stabilized zirconia, it wasfound the transformation rates of the zirconia grains embedded in an alumina matrix were muchslower. This was attributed the rigid alumina matrix, which opposed the volume changeassociated with tetragonal to monoclinic transformations and applied a compressive forces ontransforming zirconia grains [60]. With increasing tetragonal zirconia additions this constraintwas found to decrease which corresponded to the increased degradation rates for zirconiaceramics containing high vol% of zirconia additions.Similar results were found by Deville et al. [61] where alumina samples containing 6.7,20 and 40 vol% of 3 mol% yttria stabilized zirconia were exposed to a steam atmosphere in anautoclave at a temperature and pressure of 140 °C and 2 bars, respectively. The measured vol%of monoclinic zirconia formed on the surface of the tested samples was compared to a pure 3mol% yttria stabilized zirconia ceramic, as shown in Figure 2.12. Degradation rates increased foralumina samples containing higher vol% of the 3 mol% yttria stabilized zirconia additions. Theauthors attributed this behaviour to two factors: (1) the decreasing constraint of the aluminamatrix with increasing vol% of stabilized zirconia additions and (2) the percolation threshold ofthe stabilized zirconia additions [61]. It was suggested that in alumina samples with stabilizedzirconia additions greater than 25 vol% the zirconia grains formed a connected networkthroughout the matrix. This network assisted the degradation of the stabilized zirconia phase asthe monoclinic transformation of a single grain may easily propagate to neighbouring grains.However, alumina samples with stabilized zirconia addition less than 25 vol% the zirconia grainsdo not form this detrimental network. Instead, the stabilized zirconia grains are separated fromeach other and therefore, the monoclinic transformation of a single grain is contained and doesnot result in the further degradation of the stabilized zirconia phase.28Figure 2.12 Effect of secondary alumina additions on the formation of monoclinic zirconia in zirconiatoughened alumina ceramics [61].These results suggest that degradation may be reduced by introducing a second phase,which constrains the mobility of zirconia grains preventing monoclinic transformations.However, the addition of these secondary phases is challenging as the size and distribution of thetwo phases must be carefully controlled. Furthermore, sintering a two phase ceramic structure isdifficult due to large residual stresses which may form during sintering. Basu et al. [60] revealedthat alumina ceramics with increasing amounts of 3 mol% yttria stabilized zirconia additionsexhibited an increase of as-sintered monoclinic content. This may be attributed to the differentthermal expansion coefficients of the two materials which, during cooling, results in a buildup ofresidual stresses and ultimately initiates tetragonal to monoclinic phase transformations.2.7.5 Effect of Environment on DegradationIn addition to the discussed material factors (grain size, porosity, secondary stabilizers andadditives), the degradation behaviour of zirconia ceramics is also dependent on severalenvironmental factors. This section reviews key environmental factors including the effect of29temperature, pressure and the composition of the aqueous solvent used on the degradationbehaviour of zirconia ceramics.2.7.5.1 Effect of Temperature on DegradationNumerous reports in literature define a strong temperature dependence on degradation rates ofzirconia. Sato et al. [62] reviewed the degradation behaviour of various yttria stabilized zirconiasamples after testing in air for 50 hours at temperatures ranging from 150-800 °C. Themonoclinic content found on the surface of the various samples after testing was plotted as afunction of annealing temperature, and is shown in Figure 2.13. The authors found the amount ofmonoclinic phase of all the sintered specimens increased with increasing annealing temperatureup to a maximum at ~200 °C. Above this temperature, the monoclinic content was found todecrease where no transformations was observed at temperatures above 400 °C for zirconiaceramics stabilized with 3 and 4 mol% of Y2O3 and at temperatures above 600 °C for zirconiaceramics stabilized with 2 mol% Y2O3.  The authors attributed the decrease of monoclinic phaseformed at temperatures above 200 °C to the improved stability of the tetragonal phase (since foryttria stabilized zirconia ceramics the phase boundary between the monoclinic to tetragonalphase lies around ~500 °C) as reviewed in Section 2.4.2, Figure 2.4. In addition, it was foundthat the width and the height of the monoclinic peak decreased with increasing dopant levels aswell as decreasing grain size. These results confirm that the stability of the tetragonal phaseincreases with increasing dopant levels and decreasing grain size.30Figure 2.13 Effect of temperature on the degradation of yttria stabilized zirconia ceramics [62].Matsui et al. [63] correlated the temperature dependence of degradation rates to theinternal friction of the sintered ceramic. The internal friction, or the measurement of forcebetween elements in a solid while it undergoes deformation, was measured by Shimada et al.[64] for three structural ceramics: alumina, silicon nitride and 3 mol% of yttria stabilized31zirconia. Shimada et al. [64] found that the internal friction of the stabilized zirconia ceramicincreased rapidly with increasing temperature reaching a maximum point around 200 °C.Comparatively, the internal friction of both the alumina and silicon nitride ceramics remainedconstant. Bases on these results, Matsui et al. [63] concluded that the rise in internal friction wasdue to the deformation associated with tetragonal to monoclinic transformations. At atemperature of 200 °C the rate of these transformations reaches a maximum indicating thatmonoclinic phase transformation occurs most rapidly at 200 °C for the 3 mol% of yttriastabilized zirconia ceramic. Similar internal friction peaks were observed by Kondoh and Shiota[65] for various tetragonal and cubic zirconia polycrystals stabilized with yttria over a range oftemperatures (25-300 °C). Kondoh and Shiota [65] also measured the activation energy for theinternal friction and compared it to the activation energy for ionic conductivity for the variousstabilized zirconia systems. The authors found these activation energies were very similar andconcluded that the internal friction peaks were derived from rearrangement of point defectsassociated with oxygen vacancies. These results support that (1) tetragonal to monoclinictransformations are temperature dependent and (2) oxygen vacancies are involved in degradationprocess.2.7.5.2 Effect of Pressure on DegradationIn addition to the presence of water, vapour pressure also influences degradation rates. Sato et al.[66] studied the effect of water pressure on the degradation behaviour of zirconia ceramicsstabilized with 2, 3 and 4 mol% of yttria. Degradation testing was performed in a stream ofhumid air (flow rate of 10 mL/min) at temperatures ranging between 100-600 °C for 1 to 50 hrs.The partial pressure of water was adjusted by bubbling air through various saturated aqueoussolutions (where tested partial pressures were 462, 746 1540, 2330, 3350 and 7350 Pa). Attemperatures below 200 °C, it was found that monoclinic transformation rates increased withincreasing pressure until a saturation point was reached, as shown in Figure 2.14. This indicatedthat water vapour pressure did not change the reaction equilibrium, but simply accelerated therate of tetragonal to monoclinic transformations. No effect of pressure was observed for testtemperatures above 200 °C which was attributed to the improved stability of tetragonal zirconiadue to the higher temperatures [66]. Similar results were reported by Guo and Schober [67]where an increase in water vapour pressure enhanced tetragonal to monoclinic transformations as32the measured monoclinic volume increased from 0.6 to 73% for water vapour pressures of 0.026and 1 bar, respectively (temperature remained constant at 250 °C).Figure 2.14 Effect of pressure on the tetragonal to monoclinic transformation in 3 mol% Y2O3-ZrO2 ceramics[66].2.7.5.3 Effect of pH on DegradationIt is well known that the degradation of stabilized zirconia ceramics is enhanced by the presenceof water; however, some studies in the literature also consider the effect of pH on degradationrates. Sato and Shimada [46] reviewed the effect of various acidic and basic aqueous solvents onthe degradation behaviour of tetragonal stabilized zirconia ceramic containing 3 mol% of yttria.The sintered ceramic was found to be fully stabilized in the tetragonal phase and had a grain sizeof 1.1 µm as well as a relative density of 99.9%. Degradation testing was performed by placingthe ceramic material and the selected solvent in a sealed tube and heated a temperature of 95 °Cfor 120 hrs. It was found that the mol% of monoclinic zirconia formed on the surface of theceramic remained unaffected despite the solution medium, as shown in Table 2.2.33Table 2.2 Formation of monoclinic zirconia during degradation testing of 3 mol% yttria stabilized zirconia invarious solvents at 95 °C for 120 hrs [46].Environment Monoclinic ZirconiaPresent (mol%)H2O 48.85M HCl 53.35M H3PO4 52.65M HClO4 51.15M HNO3 45.75M H2SO4 50.95M NaOH 53.830wt% NH3 49.7A similar study was performed by Schacht et al. [68] who studied various stabilizedzirconia systems in diluted acidic aqueous solutions (HCl, H2SO4 or H3PO4) at a temperature andpressure of 390 °C and 27 MPa, respectively. Three stabilized zirconia systems were tested: (1)tetragonal zirconia ceramics stabilized with 12 mol% CeO2, (2) 3mol% yttria stabilized zirconiaceramics containing 10.2 mol% of MgO and (3) commercially available MgO PartiallyStabilized Zirconia (PSZ). Out of the three acidic aqueous mediums, highest degradation rateswere observed in tests performed in the aqueous HCl and H2SO4 solutions, while the lowestdegradation rates occurred in the H2PO4 solution. Comparing the different materials tested, itwas found that the CeO2 stabilized zirconia ceramic constantly had lower degradation ratesregardless of the test medium, as seen in Table 2.3. In the HCl solution, the degradation in theMgO containing zirconia ceramics was attributed to the localized chemical attack of partiallyglassy grain boundaries due to the formation of the Mg2SiO4 -rich grain boundary phase formedin the sintered ceramic [68]. In the H2SO4 solution, monoclinic transformations found in thetested ceramics were attributed to the dissolution of the stabilizing additive (either MgO orCeO2). The authors noted that the lower monoclinic content found on the CeO2 containingceramic after testing in the H2SO4 solution was due to the significant material loss (~10 timesgreater than the MgO containing ceramics). This indicated that in-depth degradation occurredrapidly in the CeO2 containing ceramic while this behaviour was limited in the zirconia ceramicsstabilized with MgO [68].   Finally, degradation in H3PO4 solutions was much lower for all thetested ceramics. This behaviour was attributed to the formation of zirconium phosphates, whichinhibited further degradation [68]. Compared to Sato et al. [46], these results suggest that the34combined factors of both the composition of the sintered ceramic as well as the composition ofthe test environment contribute to the degradation performance of stabilized zirconia ceramic.Table 2.3 Formation of monoclinic zirconia on the surface of various zirconia ceramics after degradationtesting in acidic solutions at a temperature of 390 °C and a pressure of 27 MPa for 50 hrs [68].EnvironmentMonoclinic Zirconia (vol%)12 mol% CeO2-ZrO210.2 mol% MgO -3 mol% Y2O3-ZrO2CommercialMgO-PSZ0.1 mol kg-1 HCl <2 31 930.1 mol kg-1 H2SO4 33 92 980.1 mol kg-1 H3PO4 <2 18 11Kim et al. [69] studied the degradation of tetragonal polycrystals stabilized with 3 mol%of yttria in pure water as well as water containing LiOH concentrations (ranging from 3.5 to 350ppm). Degradation testing was performed in a stainless steel autoclave at a temperature of 200°C for a period of 2 hrs. It was found degradation depth increased in solutions containing higherconcentrations of LiOH, as shown in Figure 2.15. The authors correlated this behaviour to therapid dissolution of LiOH on the surface of the ceramic forming OH- ions, which diffused intothe matrix and caused the buildup of stresses and, ultimately, facilitated tetragonal to monoclinictransformations. These results support that degradation involves the diffusion of OH- ions intothe lattice however, further research is necessary to fully understand the role of pH as well as thecomposition of the test solution on the resulting degradation behaviour of zirconia ceramics.35Figure 2.15 Effect of LiOH concentrations on the degradation of 3 mol% yttria stabilized zirconia ceramics inwater at a temperature of 200 °C for 2 hrs [69].2.8 Zirconia Ceramics for Supercritical Water ApplicationsZirconia ceramics have desirable properties of low thermal conductivity, high thermal stabilityand good chemical stability making it a potential material for insulating and protecting structuralmaterials used in supercritical water applications [5], [4]. This section reviews: (1) the propertiesof supercritical water and (2) the available literature on the performance of zirconia ceramics insupercritical water.2.8.1 Properties of Supercritical waterSupercritical water is defined as the state phase of water that exists above temperatures andpressures of 374 °C and 22.1 MPa, respectively. Above this point, the phase boundary betweenliquid and vapour become indistinguishable resulting in the formation of a dense, gas-like,substance [3]. At these high temperatures and pressures, water exhibits both liquid-like andvapour-like characteristics, resulting in excellent heat-transporting properties (ie., highdiffusivity and enthalpy). Below this critical point, water acts as a polar solvent and, at room36temperature, has a density and ionic product of 1000 kg m-3 and 10-14 mol2 L-2, respectively [3].The ionic product of water is simply a measure of the self-dissociation of water into H3O+ andOH- ions, as shown by the reaction outlined in Equation 2.2 [70].2 ⇆ + Equation 2.2Where the ionic product (kw) is simply calculated as the multiplication of the concentration of thetwo ionic species:= [ ][ ] Equation 2.3Based on Equation 2.2, for every H3O+ ion one OH- ion is formed and thus the concentration ofH3O+ and OH- ionic species will be equal. Therefore, the pH can simply be calculated by takingthe negative log of the ionic product and dividing it in half. Thus, for water at room temperatureand at atmospheric pressure we find the pH of water is 7, confirming that water is a neutralsolvent. Heating water under a constant pressure shifts the equilibrium position of the reactionpresented in Equation 2.2 to the right side resulting in an increase in ionic product and a decreasein pH. This behaviour continues until the supercritical point of water is reached (at temperatures> 374 °C and pressures > 22.1 MPa). Above this critical point, water exhibits a drastic change inphysical properties, namely, a significant drop in density and ionic product. These significantchanges in physical properties results in supercritical water to act like a non-polar solvent andhas complete solvency for organic compounds as well as complete miscibility of gases such asoxygen. Thus, depending on what species are present, supercritical water can become a veryaggressive oxidizing environment. Numerous numerical models have been developed in order todetermine the properties of supercritical water as a function of temperature and pressure, asillustrated in Figure 2.16 [70]. As shown, the temperature at which this drop of physicalproperties occurs is dependent on pressure as with rising pressures the drop of ionic conductivityand density shifts to higher temperatures.37Figure 2.16 Effect of temperature and pressure on the density and ionic product of water [3].2.8.2 Degradation Behaviour of Zirconia Ceramics in Supercritical WaterThere are limited studies in literature reviewing the performance of zirconia ceramics in asupercritical water environment. Boukis et al. [14] reviewed the degradation performance ofvarious ceramics in supercritical water containing 0.44 mol kg-1 oxygen and 0.05 mol kg-1hydrochloric acid at a temperature of 465 °C and a pressure of 25 MPa. Three classes of zirconiaceramics were tested: (1) zirconia ceramics stabilized with various amounts of yttria (3, 3.5 and10.5 mol% Y2O3), (2) zirconia ceramics stabilized with 3.5 mol% of MgO and (4) zirconiaceramics containing various amounts of yttria, MgO and an Mg-Al compound called spinel(MgAl2O4). Exposure to a supercritical water environment revealed the zirconia ceramicsstabilized with just yttria disintegrated rapidly which was attributed to tetragonal to monoclinictransformations [14]. The zirconia ceramics containing various amounts of yttria, MgO andspinel exhibited improved degradation resistance; however, the zirconia ceramic stabilized with3.5 mol% of MgO exhibited the lowest weight loss. Further analysis revealed the as-sinteredsample of the 3.5 mol% MgO zirconia system contained partially stabilized cubic matrix with7% monoclinic zirconia and 30% of tetragonal zirconia. After testing it was revealed that theamount of monoclinic zirconia was 37% indicating that all the tetragonal zirconia was38transformed while no transformation of the cubic zirconia occurred [14]. These results supportthe improved stability of the cubic phase when compared to the stabilized tetragonal phase.Supercritical testing by Schacht et al. [68] found contradicting results. In their work,zirconia ceramics stabilized with MgO had the lowest stability in supercritical aqueous solutionscontaining 0.1mol kg-1 hydrochloric acid [68]. The performance of zirconia ceramics stabilizedwith either CeO2, a combination of Y2O3 and MgO or only MgO was exposed to a supercriticalenvironment at a temperature and pressure of 390 °C and 27 MPa, respectively. Out of the threeceramic systems tested the zirconia ceramic stabilized with CeO2 exhibited the strongest stabilityand contained the lowest monoclinic content after 50 hrs of exposure. The improved resistanceof zirconia ceramics with CeO2 additions was attributed to the absence of silicate-rich grainboundary phase found in the zirconia ceramics containing MgO additions [68].Barrett et al. [71] performed supercritical water degradation testing on various ceramicmaterials including alumina, titanium oxide (TiO2) and 7.5 mol% yttria stabilized zirconiacontaining 0, 5, 10 and 15 wt% of TiO2 additions. XRD analysis revealed at the yttria stabilizedzirconia ceramic containing TiO2 additions were stabilized in the tetragonal phase and containedless than 2.5 mol% of monoclinic phase. Degradation testing was performed in an autoclave attemperatures varying from 500-540 °C and at pressures varying from 22.4-27.8 MPa. Tests wereperformed for times ranging between 940 to 1176 hrs. Furthermore, the dissolved oxygen wasmaintained between 0.5-1 ppm while the measured pH both before and after testing was found tobe 5. Weight measurements of the specimens before and after testing revealed the pure and TiO2containing yttria stabilized zirconia ceramics had a minimal weight change when compared tothe tested alumina ceramic. The larger weight gain of the alumina sample was attributed to theformation of Al2O3(H2O)1.5 due to the hydration of alumina. Microstructure analysis of thedegraded specimens revealed the presence needle-shaped particles on the surface of the TiO2containing yttria stabilized zirconia ceramics, as shown in Figure 2.17. The formation of theseparticles was attributed TiO2 rich areas which formed in the sintered specimens due to either (1)the incomplete dissolution of TiO2 during sintering, or, (2) exceeding the solid solution limit ofTiO2 into 7.5 mol% yttria stabilized zirconia during cooling of the sintered specimen. Cross-section analysis revealed that these particles only formed on the surface of the specimens.39Furthermore, no scale formation or pitting was observed on the surface of any of the testedceramic materials indicating that, in general, these materials were stable in supercritical water.Figure 2.17 Formation of needle-like particles (circled) on the surface of yttria stabilized zirconia ceramicswith 15 wt% of TiO2 additions after supercritical water degradation testing [71]. Surface contaminants areindicated by a square.2.9 Fabrication of Materials through Spark Plasma SinteringSpark Plasma Sintering (SPS) is a novel powder metallurgy technique for the fabrication ofceramic materials. It utilizes uniaxial pressure and simultaneous pulsed DC current toconsolidate and sinter powders. Compared to conventional sintering, SPS has two mainadvantages: (1) lower sintering temperatures and (2) shorter processing times [72]. Sinteringtemperatures of SPS range between 200 to 500 ˚C lower than conventional sintering techniques[72]. These lower sintering temperatures allow for the control of material crystallization which isof particular interest in the formation of dense amorphous materials, as undesirable phasetransformations can be avoided [73]. Furthermore, processing times of SPS can be up to 20 timesfaster than in conventional sintering [72]. This rapid sintering enables tight control of graingrowth allowing for the fabrication of highly dense nanocrystalline materials [73]. Shortprocessing times also ensure that heat is simultaneously and evenly distributed throughout thepowders. This uniform heating greatly reduces the residual stresses of the final sintered material40[72]. Numerous metallurgical and ceramic materials, including composites, have beensuccessfully sintered to nearly 100% density by SPS, often without the use of binders andadditives [74]. Alternatively, SPS can produce materials with specific and uniform porosity byadjusting sintering pressure and temperature [72]. The unique processing capabilities of SPShave resulted in the formation of a variety of materials including nanocrystalline metals/alloys,nanoceramics, bulk metallic glasses, nanocomposites as well as amorphous and functionallygraded materials [74]. SPS processing of these materials not only simplified material preparationand processing (reducing manufacturing cost), but also improved the material mechanicalproperties due to the absence of weaker binder materials as well as eliminated binder removalprocesses. These attributes have led to the application of SPS materials in a variety of industries. Forexample, high density materials including wear resistant cobalt-less tungsten carbide tooling, high-purity ceramic fuel cells and ceramic optics have all been produced through SPS [72]. This sectionreviews the basic methodology behind SPS, as well as, literature results on the fabrication ofzirconia ceramics through this novel technique.2.9.1 Spark Plasma Sintering MethodologyThe basic components involved in SPS are shown in Figure 2.18. First, powders are loaded in agraphite die enclosed by two concentric graphite punches. A pulsed alternating DC current isthen applied to the punches along with a uniaxial pressure and the powders are compressed. Thepulsing DC current initiates sparks, or momentary high temperature points, on and between thesurfaces of the powder particles. These high-temperature states cause the localized vapourizationand incipient melting of the powder particle surface [72]. The applied uniaxial pressurecontributes to the even distribution of the spark points throughout the powder whileconsolidating the particles [72]. The entire sintering process is performed in either a low pressurevacuum or inert gas environment. After a specified amount of sintering time the pulsed DCcurrent and uniaxial pressure are removed and the final sintered material is then released fromthe die. The challenge of SPS is the optimization of the process parameters in order to obtain thedesired sintered properties for different materials and die configurations [74]. Generally, thisoptimization is performed empirically; however, significant efforts are being directed tounderstanding the exact mechanism(s) involved in SPS and developing a mathematic model[75]-[77].41Figure 2.18 Schematic of the SPS process [73].2.9.2 Zirconia Ceramics Fabricated through Spark Plasma SinteringDahl et al. [78] fabricated 8 mol% yttria stabilized zirconia ceramics through three processingtechniques: (1) SPS, (2) hot pressing and (3) conventional sintering. For each process, thesintering parameters (temperature, pressure and dwell time) were varied and the resulting densityand grain sizes of the sintered ceramics were compared. It was found that compared to hotpressing and conventional sintering, zirconia ceramics fabricated through SPS were able tomaintain a high relative density despite low sintering temperatures, as shown in Figure 2.19. Inaddition, the dwell time of the ceramics sintered through SPS was one twelfth of the dwell timeused for the conventional and hot pressed ceramics (5min compared to 1 hour, respectively).These results indicated that SPS is a rapid and highly efficient sintering technique. Furthermore,SPS samples generally exhibited smaller grain sizes, however, compared to hot pressing it wasnoticed that rapid grain growth would occur at sintering temperature greater than 1200 °C, asshown in Figure 2.20.42Figure 2.19 Resulting density of 8 mol% yttria stabilized zirconia ceramics sintered with different sinteringtechniques and at different temperatures [78].Figure 2.20 Average grain size of 8 mol% yttria stabilized zirconia ceramics sintered with different sinteringtechniques and at different temperatures [78].43Banghao et al. [79] compared the microstructure of 8 mol% yttria stabilized zirconiaceramics fabricated through SPS and conventional sintering. SPS samples were sintered under auniaxle pressure of 28 MPa at a temperature of 1350 °C (heating rate of 160 °C/min). Thesintering temperature was held for 10 min. In contrast, the conventional samples were sintered at1450 °C for 4 hrs in an electrical resistance furnace. After sintering both types of samples wereallowed to cool naturally. The SPS samples had a higher relative density and a smaller grain sizein comparison to the conventionally sintered samples. The measure density and grain size for theSPS sintered ceramics were 99% and 0.8 µm, respectively, while for the conventionally sinteredceramics the measured values were 97% and >1 µm, respectively.Dense 8 mol% yttria stabilized zirconia ceramics were fabricated by Anselmi-Tamburiniet al. [80] through a high pressure SPS technique. Sintering was performed for a dwell time of 5min at pressures ranging from 35 to 1000 MPa and temperatures ranging between 930 and1350 °C (with a heating rate of 200 °C/min). The resulting ceramics were found to exhibitrelative densities greater than 98% as well as grain sizes ranging around 10 nm [80]. Theseresults demonstrate that nanograined stabilized zirconia ceramics can be rapidly fabricatedthrough SPS; however, careful selection of sintering parameters must be performed to controlgrain growth as well as to avoid phase transformations.2.10 SummaryAt room temperature pure zirconia has a monoclinic structure. Upon heating of zirconia, thisstructure undergoes several martensitic transformations: first, to a tetragonal structure at1170 °C, followed by a cubic transformation at 2370 °C [5], [13]. The structure of the hightemperature polymorphs can be controlled by doping the zirconia matrix with divalent ortrivalent oxides [5], [21]. In this way, the properties of zirconia ceramics can be tailored for avariety of applications including biomedical implants, solid oxide fuel cells, thermal barriercoatings as well as oxygen pumps and sensors [5]. However, the application of zirconia islimited by a degradation behaviour [30]. Tetragonal to monoclinic transformation results in largestresses within the ceramic, causing the formation of microcracks along grain boundaries and,ultimately, results in the loss of the material properties [19]. This phase transformation has been44extensively studied and many degradation mechanisms as well as methods to control thistransformation have been proposed [35], [46], [48]. In contrast, very little research has beenperformed on the degradation behaviour of cubic zirconia.  In general, cubic zirconia is regardedas an untransformable, stable material, since no degradation been has been observed for cubiczirconia in similar environments that caused tetragonal degradation. However, several studieshave shown that cubic to tetragonal transformations do occur, but at a much slower rate [38].Preliminary degradation tests in supercritical water suggested that cubic zirconia can withstandthis aggressive environment while tetragonal zirconia ceramics exhibit rapid disintegration [14].Despite these findings, there are many factors which contribute to the degradation behaviour ofstabilized zirconia ceramics which can be correlated to the high discrepancy of degradation ratesavailable in the literature. Therefore, characterizing and preventing the degradation of bothstabilized tetragonal and cubic zirconia is necessary for the increased application of this ceramicsystem.45Chapter 3: Experimental ProcedureThis chapter provides an overview of the experimental procedures involved in materialpreparation and processing, degradation testing, and subsequent evaluation andcharacterization.3.1 Precursor PowdersHigh purity (99.9%+) precursor powders of 8 mol% Yttria Stabilized Zirconia (YSZ), ceriumdioxide (CeO2) and neodymium oxide (Nd2O3) were obtained from Inframat AdvancedMaterials, US (product #4039OR-8601, #58-0803 and #60N-0801, respectively). Material datasheets provided by Inframat Advanced Materials indicated that the YSZ powders had a D50 value~0.5 µm while the rare earth oxide powders had a D50 value < 1.0 µm [81]-[83]. These reportedparticle distribution values were consistent with the observed particles sizes presented in Figure3.1. Furthermore, chemical analysis showed that the powders retained a homogenous chemicalcomposition, as summarized in Table 3.1. Traces of Hf and Cl were attributed to the wetchemical co-precipitation method used to fabricate the precursor powders.Table 3.1 Chemical composition and particle size of precursor powders.PrecursorPowderConcentration (wt%) ParticleSize(nm)Zr O Y Ce Nd Hf ClYSZ 63.40 24.58 9.93 -- -- 1.63 0.46 500CeO2 -- 14.55 -- 84.90 -- -- 0.52 50-80Nd2O3 -- 13.95 -- -- 83.53 -- 2.57 ~8046Figure 3.1 SEM micrographs of precursor powders: a) YSZ, b) CeO2 and c) Nd2O3.3.2 Precursor Powder BlendingBlending of the pure YSZ and CeO2-YSZ powders was carried out at the Indian Institute ofTechnology Madras (Chennai, India) under the supervision of Prof. B.S. Murty. Powders wereweighed with an electronic scale and blended for 15 min at 300 rpm using a Fritsch P-5 highenergy planetary ball mill without milling balls or binders. Final powder blends of YSZ withadditions of 0, 5, 10 and 15 mol% CeO2 were prepared. Blended Nd2O3-YSZ powders wereprepared at the University of British Columbia (Kelowna, Canada) by mixing the measuredpowders in a beaker with alcohol for 4 min using a rotating magnet on a hot plate. The powderswere then allowed to dry for 24 hrs at room temperature after which they were manually groundusing a porcelain pestle for 20 min. Final powder blends of YSZ with additions of 5 and 10mol% Nd2O3 were prepared. YSZ plus 15 mol% Nd2O3 powder blends were not fabricated dueto budget limitations. Equivalent wt% of the CeO2 and Nd2O3 additives are provided in Table3.2.47Table 3.2 Equivalent wt% of CeO2 and Nd2O3 additions in the CeO2-YSZ and Nd2O3-YSZ powder blends.Additive mol% wt%CeO25.0 6.410.0 12.715.0 18.8Nd2O35.0 12.010.0 22.23.3 Spark Plasma Sintering ProcessingSPS processing of the blended powders was performed at two locations: the Indian Institute ofTechnology Madras (Chennai, India) and Quad City Manufacturing Lab (Rock Island, USA).Sintering was performed by qualified technicians at each institution. As the densificationbehaviour of the different powder blends depended on dopant type and level, the sinteringparameters (i.e. temperature, pressure, time) had to be varied. An overview of the sinteringprocedures performed for the respective ceramics are provided in the following sections.3.3.1 Sintering of YSZ and CeO2-YSZ PowdersPure YSZ and blended CeO2-YSZ powders were sintered using the Dr. Sinter SPS-625 machineat the Indian Institute of Technology Madras (Chennai, India). For each specimen, 10 g of thepowder blend was weighed and loaded into a 20 mm graphite die. Grafoil sheets were used toline the die to enable an easy removal of the sintered material after SPS processing. The loadeddie was then placed into the SPS machine and an initial uniaxial pressure of 5 MPa (ie. preload)was applied along with a vacuum of 5x10-2 Torr. During sintering, the uniaxial compressivepressure was increased to 30 MPa. A heating rate of 100 °C/min was applied until the diereached the required sintering temperature. After a defined dwell time at the desired sinteringtemperature, the sintering pressure was released and the material was allowed to cool. To avoidoxidation of the graphite die and the sintered ceramic, the vacuum was only released when thesample cooled below 600 °C. Sintering temperature and dwell time was varied for the variouspowder blends in order to obtain solid samples which did not fracture upon ejection from thegraphite die. The final SPS process parameters for the powder blends are listed in Table 3.3. Thesintered ceramic materials were cylindrical in shape and had a diameter of 20 mm with a heightof ~4 mm as seen in Figure 3.2. Two sintered disks were fabricated for the pure YSZ sampleswhile four sintered disks were fabricated for each of the CeO2-YSZ samples. Duplicate disks48were fabricated as, despite the optimized sintering parameters, ~50% of the sintered specimensfractured upon ejection from the die.Table 3.3 SPS processing parameters for YSZ and CeO2-YSZ samples.Sample SinteringTemperature (°C)SinteringTime (min)SinteringPressure (MPa)Pure YSZ 1400 7305 mol% CeO2 - YSZ 1300 510 mol% CeO2 - YSZ 1400 715 mol% CeO2 - YSZ 1400 7Figure 3.2 As-sintered YSZ and CeO2-YSZ ceramics: a) pure YSZ, b) 5 mol% CeO2-YSZ, c) 10 mol% CeO2-YSZ and d) 15 mol% CeO2-YSZ.3.3.2 Sintering of Nd2O3-YSZ PowdersBlended Nd2O3-YSZ powders were sintered using a Thermal Technology Spark PlasmaSintering machine (10-3 model) at Quad City Manufacturing Lab (Rock Island, USA). For eachsintered ceramic, 10 g of the blended powder was placed in a 19 mm diameter graphite die with a49Grafoil lining. An initial load of 5 MPa was applied, and the sintering chamber was pumpeddown to < 2x10-2 Torr vacuum. The uniaxial pressure was then increased and a heating rate of200 °C/min was applied until the temperature reached 600 °C. This temperature was sustainedfor a 10 s dwell time to enable accurate temperature measurements using an optical pyrometerduring subsequent sintering temperatures.  Heating at 200 °C/min was then maintained until thedesired sintering temperature was reached. After sintering, the furnace was turned off and thesamples were allowed to cool naturally in the graphite die. The SPS process conditions (pressure,temperature and dwell time) were varied extensively to determine the optimal process parameterswhich resulted in solid samples that did not fracture when ejected from the graphite die. Theselected SPS processing parameters for the Nd2O3-YSZ samples are provided in Table 3.4. Asshown in Figure 3.3, resulting ceramics measured 19 mm in diameter with a height of ~4 mm.Only one sintered disk of each specimen was fabricated.Table 3.4 SPS processing parameters for Nd2O3-YSZ samples.Sample SinteringTemperature (°C)SinteringTime (min)SinteringPressure (MPa)5 mol% Nd2O3 - YSZ 1400 5 4010 mol% Nd2O3 - YSZ 1300Figure 3.3 As-sintered Nd2O3-YSZ ceramics: a) 5 mol% Nd2O3-YSZ and b) 10 mol% Nd2O3-YSZ.3.4 Polishing, Thermal Etching and Density MeasurementsTo reveal the microstructure of the ceramic materials, representative samples were sectionedusing a diamond precision saw (LECO VS-50) and mounted in Fast Cure Acrylic epoxy(MetLab, Cat# M142). Manual polishing of the mounted samples was performed in six stages:(1) 240 grit Al2O3 paper for 2 min, (2) 600 grit SiC paper for 10 min, (3) 1200 grit SiC paper for5020 min, (4) 9 µm diamond paste for 30 min, (5) 6 µm diamond paste for 30 min and (6) 1 µmdiamond paste for 30 min. Distilled water was used as a coolant in stages 1-3, and a syntheticdiamond extender fluid (MetaDi) was used in stages 4-6.To reveal the grain structure, each of the polished ceramics was then thermally etched inair for 1 h at a temperature 50 °C below its corresponding sintering temperature. The furnaceused was a MHI furnace equipped with 7-MoSi2 heating elements. The heating rate wasmaintained at 10 °C/min, while the cooling rate was set to 20 °C/min.The specific gravity of the sintered materials was measured using the Archimedes’method. An electronic balance with a beaker of distilled water at room temperature was used.The relative density was calculated by dividing the measured density by the theoretical density.Theoretical densities were obtained from online material data sheets provided by InframatAdvanced Materials [81]-[83].3.5 MicrohardnessMicrohardness of the sintered ceramics was performed by applying a load of 9.807 N for a dwelltime of 10 s using a MicroVicker hardness tester (model SXHV-1000TA). These defined loadingconditions resulted in indentations with an average diagonal length of 44 ± 1µm for all theceramic specimens tested. Measurements were repeated until the 95% confidence interval waswithin 5% of the average value. Hardness measurements were performed in the center of thecross-sectioned face of the ceramics (both before and after degradation testing).3.6 X-ray DiffractionTo characterize the phases present in the studied materials, XRD analysis of precursor powders,as-sintered and tested ceramics were performed. Precursor powders were characterized using andX’pert Pro (PANalytical) X-Ray Diffractometer with Bragg-Brentano set-up, equipped withX’Celerator detector. Diffraction peak measurements were performed in a range of 20-90° at a0.02° step size with a 20 s per step along with Cu-Kα radiation (45 kV, 30 mA). As-sintered andtested ceramics were analyzed using a Bruker D8 Advanced-ray Diffactometer in Bragg-Brentano configuration with a Lynxeye silicon strip detector (3.7 degree window opening, 1 mmdivergence slit, 8 mm anti-scatter). Diffraction peak measurements were performed from 5-90°51using a 0.039° step size with 163 s per step. Cu-Kα radiation (40 mA, 40 kV) was used for allsamples.3.7 Microscopy, Chemical and Image AnalysisHigh resolution imaging and chemical analysis was performed using a Tescan MIRA3 XMUScanning Electron Microscope (SEM) with Oxford X-max X-ray Energy DispersiveSpectroscopy (XEDS) detector (without ZAF correction). All samples were coated with a 5 nmplatinum/palladium alloy using Cressington 208HR Sputter Coater. SEM imaging and analysisof the as-sintered and tested samples were performed using the Back Scatter Electron (BSE)detector with the microscope parameters set at a working distance of 14 mm along with a beamvoltage of 20 kV. The beam spot size was controlled by a unitless parameter called the beamintensity which was set to 15. This parameter was optimized for the specified working distanceand beam voltage in order to retain a reasonable image resolution while producing sufficientnumber of x-rays for analysis. Analysis of the as-sintered ceramics was performed on specimenspolished and etched by the processes reviewed in Section 3.4. Comparatively, as separatesamples were prepared for degradation testing, the analysis of the ceramics after testing wasperformed on samples polished only with 600 grit SiC paper (test sample preparation is providedin Section 3.8.1). For both the as-sintered and tested ceramics the cross-section surface of thesintered disc was analyzed. Due to the short working distances required, SEM images of sinteredgrain sizes were performed using the InBeam detector along with a beam voltage of 20 kV and abeam intensity of 10. Grain size was measured via the linear intercept method using BuehlerOmnimet image analysis software.3.8 Degradation TestingDegradation testing of fabricated ceramic samples in supercritical water was carried out at theDepartment of Materials Engineering at the University of British Columbia, (Vancouver,Canada) by the author. The following text provides an overview of the instruments used,measurements performed, and subsequent calculations implemented to determine degradationrates.3.8.1 Sample PreparationSintered ceramic disks were sectioned into 4x4x4 mm3 cubes. If possible, three sectionedspecimens were obtained from each sintered ceramic disk. To remove traces of the Grafoil sheet,52all sides of the sectioned cubes were manually polished with 600 grit SiC paper with distilledwater as lubricant. Samples were then ultrasonically cleaned for 5 min in distilled water, andwere allowed to dry before the weight and dimensions of each individual test specimen wasrecorded.3.8.2 Vessel DesignSupercritical water degradation testing was performed using a simple tube vessel which wascapped at each end as shown in Figure 3.4. The 316L tubing, fabricated according to ASTM A269[84], was purchased from Unified Alloys, Canada while the end caps were purchased fromSwagelok, Canada (product # SS-1610-C).Figure 3.4 Supercritical water degradation test vessel.Test vessels were designed according to ASTM G31 Standard Practice for LaboratoryImmersion Corrosion Testing of Metals which states the “minimum solution volume to specimenarea ratio is 0.20 mL/mm2 (125 mL/in.2) of specimen surface” [85]. For square specimens with aside length of 4 mm, the corresponding vessel design was a 25.4 mm (1 in.) outer diameter tube,254 mm (10 in.) in length with a wall thickness of 3.05 mm (0.120 in.). Assuming the solution ofsupercritical water expands to the volume of the vessel minus volume of the specimen thesolution to volume ratio was calculated to be 0.77 mL/mm2, well above the required ASTMstandard.53The wall thickness of the vessel was selected based on the calculated pressure expectedfor water sealed in the designed vessel heated to temperature above the supercritical point (>374 °C). Using the ideal gas law, the pressure (P) was calculated using Equation 3.1:= ( ) Equation 3.1Where V was the volume of the solution which is the free volume of the vessel (as supercriticalwater is assumed to expand to the free volume of the vessel), T was the absolute temperature inKelvin, R was the gas constant (9.81 J K-1 mol-1), n was the moles of water present and z was thecompressibility factor.The selected parameters for test temperature as well as volume of water sealed in thevessel was optimized based on two main factors: (1) that the resulting temperatures and pressureswould result in the formation of supercritical water and (2) the maximum pressure within thevessel was within 10 MPa of the allowable working pressure of the designed vessel. Through aniterative design process a test temperature of 400 °C and water volume of 20 mL was determinedto meet the required specifications. Knowing the temperature that the vessel was heated to andthe volume of water sealed in the vessel, the resulting pressure could be calculated. It isimportant to note that at temperatures above the supercritical water point (temperatures andpressures greater than 374 °C and 22.1 MPa, respectively) the compressibility factor is no longerat unity. Instead, the compressibility factor is a function of both the temperature and pressure andcan be determined through the Lee Kesler diagram (included in Appendix A). Since the pressureis unknown, the compressibility factor was determined through trial and error as shown in Table3.5. For a defined test temperature of 400 °C and a water volume of 20 mL, the minimumdifference between the assumed and actual compressibility factor was found when the z wasequal to 0.4. Decreasing the assumed compressibility factor to 0.2 resulted in a negativedifference, indicating that the actual value was between 0.4 and 0.2. By interpolating betweenthese two points, the final compressibility factor was determined to be 0.37. Point of operation isindicated in Figure A.0.1 in Appendix A. Using this value for z in Equation 3.1 the final pressurewas calculated to be 31 MPa (4.5 kpsi). At room temperature, both the tubing and caps had apressure rating of 32.4 MPa (4.7 kpsi) as defined by ASTM A269 [84]. However, to determine theallowable working pressure at an elevated temperature of 400 °C, an safety factor of 0.79 was54applied reducing the pressure rating to 25.5 MPa (3.7 kpsi). This value was lower than thecalculated pressure however, the stated pressure rating assumed continuous use while in realitythe test vessels would only be exposed to the supercritical water environment for a short periodof time.  Therefore the use of 316L tubing with a wall thickness 3.05 mm (0.120 inches) wasconsidered acceptable for the designed experiment.Table 3.5 Non-converging pressure values calculated for supercritical water degradation testing at 400 °C inthe designed test vessel containing 20 mL of water.Iteration Assumed z Calculated Pressure (MPa) Actual z Differenceϕ1 1 83.66 0.55 0.452 0.8 66.92 0.45 0.353 0.6 50.19 0.39 0.214 0.4 33.46 0.3 0.15 0.2 16.73 0.75 -0.55ϕcalculated as the difference between assumed z minus the actual z which was determined from thecompressibility chartAs discussed in Section 2.8.1, heating water past the supercritical point changes thephysical properties of water. During degradation testing, distilled water with an assumed pH of 7was heated past the supercritical point to a temperature of 400 °C and a pressure of 31 MPa.These environmental conditions caused a drop in ionic product to10-17 mol2 L-2 (based on thenumerical model presented in Section 2.8.1) resulting in an increase of pH to 8.5 [3].3.8.3 Degradation Testing ProcedureDegradation testing of the various ceramics was performed by placing one sectioned cube of theas-sintered specimen along with 20 mL of distilled water into a sealed vessel. Vessels wereplaced horizontally in a furnace (Q5C HeatTreat) and heated at a rate of 2.6 °C/min until atemperature of 404 °C was reached. This temperature was held for a dwell time of 2.5 hrs afterwhich the furnace was switched off and the samples were allowed to cool. The holdingtemperature of 404 °C was selected to ensure the temperature feedback loop did not allow thefurnace temperature to drop below 400 °C. The vessels were removed from the furnace once thetemperature dropped below 60 °C (after ~17 hrs). Vessels were then allowed to naturally cooluntil they reached room temperature (~40 min). This ensured that no water would evaporateupon opening the vessels, and accurate water loss could be measured. Degradation testing was55repeated three times for each type of sintered ceramic with the exception of the 10 mol% CeO2-YSZ ceramic as only one test sample could be sectioned due to the brittle nature of the material.For each degradation test, the water loss was no more than 2 mL.3.8.4 Weight Loss and Degradation Rate CalculationsWeight loss was measured as the difference in weight between the sectioned cube before andafter testing. Ultrasonic cleaning of the tested samples was performed for 5 min to remove anyloose scale from the surface. The difference in weight before and after ultrasonic cleaning wasfound to be less than 0.01 g for all tested samples. Degradation rates were then calculated usingthe following equation [85]:= ( ∙ )( ∙ ∙ ) Equation 3.2Where W was the final weight loss in g, A was the initial surface area of the specimen in cm2, ρwas the measured density in g/cm3 and t was the total time of exposure in hrs. Parameter K wasthe unit conversion constant which, for the final degradation rates to be expressed in mm/year,was set to equal 8.76x104 [85].3.9 Test Solution AnalysisInductively Coupled Plasma-Mass Spectroscopy (ICP-MS) analysis was performed to determinethe presence of dissolved species in the 20 mL of water after supercritical water degradationtesting. Tests were performed at the University of British Columbia (Kelowna, Canada) usingThermoFisher Element XR ICP-MS. ICP-MS samples were collected after allowing thesupercritical water solutions to cool and settle. To avoid the inclusion of spalled ceramic pieces,the retrieved ICP samples were then centrifuged at a speed of 5000 rpm for 10 min in a ThermoBenchmate to isolate the precipitates from the bulk solution. A volume of 0.1 mL was thenwithdrawn and diluted 100 times into nano-pure water containing 1% of trace-grade nitric acidand 1.0 ppb of Indium as the internal standard. The multi-element calibration standards, alsomade in 1% nitric acid and 1.0 ppb Indium, yielded linearity coefficient (R2) values of ≥0.99991for all detected elements. Sample results were averaged from interpolation data obtained fromthree passes of each sample, then corrected for the dilution factor.56Chapter 4: ResultsThis section presents the results obtained during this research. For each ceramic, thedensification behaviour, degradation rates, microhardness measurements, phase andmicrostructure evolution as well as the composition of the retrieved test solutions are discussed.4.1 DensificationThe observed densification behaviour, including density and grain size for the various sinteredceramics are summarized in the following sections.4.1.1 Densification of YSZ CeramicsUpon ejection from the graphite sintering die, the YSZ ceramics were found to fracture easily, asseen in Figure 4.1. Surface cracks were common in the sintered discs suggesting that significantthermal strains were likely introduced during sintering at the selected processing conditions.Furthermore, sintered specimens exhibited a dark grey discoloration which was attributed to thediffusion of carbon from the graphite tooling. Similar observations of carbon contaminationduring SPS processing have been reported in the literature [74], [78], [79]. Thermal annealing inair for 1 hour at 1000 °C changed the color of the YSZ disk to a uniform white color, indicatingthe removal of carbon contaminants. The presence of these carbon contaminates were consideredto have an insignificant influence on the degradation behaviour of the as-sintered ceramics as thetemperature and pressures during degradation testing would enable carbon diffusion out of theas-sintered specimen. Therefore, sectioned test specimens for the as-sintered disks were notannealed before degradation testing. The results of the relative density and grain sizemeasurements for the sintered YSZ ceramics are presented in Table 4.1. An image of the sinteredgrain structure is presented in Figure 4.2.57Figure 4.1 Fractured as-sintered YSZ ceramic.Table 4.1 Measured grain size and density of as-sintered YSZ ceramics.SampleSinteringTemperature(°C)SinteringTime(min)SinteringPressure(MPa)GrainSizea (µm)Density(g/cm3)RelativeDensitya(%)YSZ 1400 7 30 9.84±0.53 5.83 95.6±1.1aerror shown taken from a 95% confidence intervalFigure 4.2 SEM micrograph of grain structure in the as-sintered YSZ ceramic.58XRD analysis of the pure YSZ powders revealed that trace amounts of monocliniczirconia was present. However, after sintering it was observed that this monoclinic phase fullytransformed into cubic zirconia, as shown by the diffraction pattern in Figure 4.3. In addition,zirconium carbides (ZrC) formed within the as-sintered ceramic. The formation of these carbidepeaks was attributed to the carbon diffusion from the graphite die.Figure 4.3 XRD plots for the YSZ powder and as-sintered YSZ ceramic.4.1.2 Densification of CeO2-YSZ CeramicsSimilar to the pure YSZ ceramics, the as-sintered CeO2-YSZ ceramics also had a strongtendency to fracture upon ejection from the graphite die. Fractured surfaces exhibited grey andorange colorations, as shown in Figure 4.4. These color variations were attributed to the carbondiffusion from the graphite tooling which, along with the applied vacuum during sintering,enabled the reduction of the Ce4+ ions to Ce3+. Due to the change of valence state, the additionsof CeO2 changed from a pale yellow to an orange color. Similar color changes were observed byHuang et al. [86] during SPS of zirconia ceramics containing Y2O3 and CeO2 additions.Furthermore, due to the ionic radius difference (Ce4+ = 0.97 Å, Ce3+ = 1.11 Å, Zr4+ = 0.84 Å)59[23], the reduction of Ce4+ likely resulted in the development of significant internal strainscontributing to the cracking of the sintered sample. Thermal annealing of the as-sintered sampleat 1000 °C for 1 h resulted in a uniform pale white-yellow coloration, indicating Ce3+ ions werereverted back to the Ce4+ state.Figure 4.4 Fractured disk showing dark and orange discolorations in as-sintered 10 mol% CeO2-YSZceramic.The measured grain size and density for the as-sintered CeO2-YSZ samples are providedin Table 4.2. As seen from the table, the selected SPS parameters resulted in moderately denseceramics with a relative density around ~87.5%. Representative images of grains found in the as-sintered ceramics are shown in Figure 4.5.Table 4.2 Measured grain size and density of as-sintered CeO2-YSZ ceramics.SampleSinteringTemperature(°C)SinteringTime(min)SinteringPressure(MPa)GrainSizea (µm)Density(g/cm3)RelativeDensitya(%)5 mol%CeO2 - YSZ1300 5 30 1.58±0.26 5.37 87.0±0.810 mol%CeO2 - YSZ1400 7 30 3.78±0.67 5.53 88.3±0.915 mol%CeO2 - YSZ1400 7 30 2.07±0.17 5.53 87.3±1.3aerror shown taken from a 95% confidence interval60Figure 4.5 SEM micrographs of grain structure in as-sintered YSZ ceramics with: a) 5 mol% CeO2, b) 10mol% CeO2 and c) 15 mol% CeO2 additions.XRD analysis showed that fully stabilized cubic zirconia was present after sintering,despite the trace amounts of monoclinic zirconia found in the precursor YSZ powders. XRDplots for the various CeO2-YSZ ceramics are shown in Figure 4.6.61Figure 4.6 XRD plots for YSZ powder, CeO2 powder and as-sintered CeO2-YSZ ceramics.4.1.3 Densification of Nd2O3-YSZ CeramicsSintered YSZ ceramics with Nd2O3 additions were found to have higher densities in comparisonto the YSZ ceramics with CeO2 additions. Furthermore, as-sintered Nd2O3-YSZ ceramic disksdid not fracture after ejection from the graphite die, as seen in Figure 4.7. Density measurementsrevealed that the sintered ceramics had a high relative density of ~92%.62Figure 4.7 As-sintered 10 mol% Nd2O3-YSZ ceramic.Microstructure analysis revealed that the as-sintered Nd2O3-YSZ ceramics displayedeither a bi- or tri-modal grain structure. Depending on the doping level, the types of grainsobserved in the as-sintered ceramics were: (1) general matrix grains, (2) larger textured grainsand (3) irregular nanosized grains. Average grain sizes for these grains are presented in Table 4.3for the different ceramics. Representative micrographs of the different grain structures areprovided in Figure 4.8 and Figure 4.9.Table 4.3 Measured grain size and density of as-sintered Nd2O3-YSZ ceramics.SampleSinteringTemperature(°C)SinteringTime(min)SinteringPressure(MPa)GrainSizea (µm)Density(g/cm3)RelativeDensitya(%)5 mol%Nd2O3 - YSZ1300 5 402.37€±0.215.71 91.8±0.12.50ϕ±0.1510 mol%Nd2O3 - YSZ1400 5 400.83€±0.055.83 92.3±0.10.35δ±0.051.52ϕ±0.10aerror shown taken from a 95% confidence interval€matrix grains, ϕ textured grains, δ nano grainsIn the 5 mol% Nd2O3-YSZ ceramic, only two types of grains were observed (Figure 4.8):(1) general matrix grains with an average grain size of 2.37 µm and (2) larger textured grainswith a grain size of 2.50 µm. The textured grains were found to exist around Nd2O3 agglomeratesand often were present in regions of higher porosity. Chemical analysis of the sintered matrix ispresented in Section 4.5.3.63Figure 4.8 SEM micrographs of grain structure in as-sintered 5 mol% Nd2O3-YSZ ceramic: a) matrix grainsand b) textured grains.SEM micrographs of the 10 mol% Nd2O3-YSZ ceramic in Figure 4.9 revealed that inaddition to the matrix grains and textured grains the sintered ceramic contained a third type ofgrain with an irregular morphology and an average grain size in the nm range (~350 nm). Thesenanograins were seen to form within grey band regions randomly distributed throughout thematrix. The larger textured grains (Figure 4.9c) were observed to form predominately in Nd-richareas and bands of the ceramic. Chemical analysis of the sintered matrix is presented in Section4.5.3.64Figure 4.9 SEM micrographs of the grain structure in as-sintered 10 mol% Nd2O3-YSZ ceramic: a) matrixgrains, b) nano grains (indicated by arrows) and c) textured grains.XRD analysis confirmed that fully stabilized cubic zirconia was present after sintering,despite the trace amounts of monoclinic zirconia found in the precursor YSZ powders. In the 10mol% Nd2O3-YSZ ceramic, additional peaks were found. These peaks were attributed to anunknown trace phase. XRD plots for the various Nd2O3-YSZ ceramics are presented in Figure4.10.65Figure 4.10 XRD plots for YSZ powder, Nd2O3 powder and as-sintered Nd2O3-YSZ ceramics with (*) peaksindicating an unknown trace phase.4.2 Degradation RatesSectioned as-sintered test specimens were weighed and photographed both before and afterdegradation testing (for clarity images are labelled as pre- and post-degradation, respectively).The corresponding degradation rates for each sample were then calculated and are summarizedin the following sections.4.2.1 Degradation of YSZ CeramicsThe calculated weight loss and degradation rates of the test samples for the pure YSZ ceramicare outlined in Table 4.4. Representative images of the sectioned specimens before and afterdegradation testing are shown in Figure 4.11. Out of the three test specimens, only one specimenremained after testing. The degradation rate of the two disintegrated specimens was calculatedassuming that complete disintegration occurred only at the end of the test (i.e the total time ofexposure (t) equaled the entire experimental timeframe, ~ 23 hrs).66Table 4.4 Calculated weight loss and degradation rates for YSZ ceramics.Composition SampleNumber Weight Loss (g)DegradationRate (mm/y) CommentsYSZ1 0.192 120 PreferentialDegradation2 0.399 260 Disintegrated3 0.170 166 DisintegratedPre-Degradation Post-DegradationFigure 4.11 Images of the sectioned YSZ sample before and after degradation testing.4.2.2 Degradation of CeO2-YSZ CeramicsThe calculated weight loss and degradation rates of the test samples for the CeO2-YSZ ceramicsare tabulated in Table 4.5. Representative images of the sectioned specimens before and afterdegradation testing are shown in Figure 4.12. It was found that during degradation the CeO2-YSZ ceramics exhibited a spallation behaviour where, rather than uniform material loss, largefragments of material were found to be ejected from the test specimen, as indicated with arrowsin Figure 4.12. Due to the brittle nature of the 10 mol% CeO2-YSZ ceramic, only one 4x4x4mm3 test specimen could be sectioned. Degradation testing of this specimen resulted in adegradation rate of 51.1 mm/year, however, this calculated rate may vary significantly asdemonstrated by the large variation in degradation rates observed for the YSZ ceramicscontaining 5 and 10 mol% of CeO2.67Table 4.5 Calculated weight loss and degradation rates for CeO2-YSZ ceramics.Composition SampleNumber Weight Loss (g)DegradationRate (mm/y) Comments5 mol%CeO2 - YSZ1 0.017 11.6 Spallation2 0.002 1.77 Intact3 0.005 3.90 Spallation10 mol%CeO2 - YSZ1 0.068 51.1 Spallation15 mol%CeO2 - YSZ1 0.002 1.53 Spallation2 0.124 84.5 Spallation3 0.005 4.13 Spallation68mol% CeO2Additions Pre- Degradation Post-Degradation51015Figure 4.12 Images of the sectioned CeO2-YSZ samples before and after degradation testing. Areas ofspallation are indicated by an arrow.4.2.3 Degradation of Nd2O3-YSZ CeramicsThe calculated weight loss and degradation rates of the test samples for the Nd2O3-YSZ ceramicsare outlined in Table 4.6. Representative images of the sectioned specimens before and afterdegradation testing are shown in Figure 4.13. It was found that during degradation, the 5 mol%Nd2O3-YSZ ceramics remained intact and had an average degradation rate of ~1.5 mm/y. YSZceramics with 10 mol% Nd2O3 exhibited a spallation behaviour similar to that found in the69CeO2-YSZ ceramics. In addition, the degraded samples had numerous large cracks across thesurface of the sample, as seen in Figure 4.13d.Table 4.6 Calculated weight loss and degradation rates for Nd2O3-YSZ ceramics.Composition SampleNumber Weight Loss (g)DegradationRate (mm/y) Comments5 mol%Nd2O3 - YSZ1 0.002 1.67 Intact2 0.002 1.49 Intact3 0.002 1.35 Intact10 mol%Nd2O3 - YSZ1 0.012 10.1 Cracking andSpallation2 0.004 2.76 Cracking andSpallation3 0.025 17.8 Cracking andSpallationmol% Nd2O3Additions Pre-Degradation Post-Degradation510Figure 4.13 Images of the sectioned Nd2O3-YSZ samples before and after degradation testing.704.3 MicrohardnessMicrohardness of the various ceramics was measured both before and after degradation testing.The results are presented in Figure 4.14. Measured indentations were observed to have anaverage diagonal length of 44.0 ± 0.9 µm for all the ceramic specimens tested. Comparison ofthis indentation size to the measured grain sizes reported in Section 4.1 indicated that themeasured hardness values represented the average hardness of several grains in the as-sinteredand degraded ceramics. Before degradation testing, the pure YSZ ceramic was found to have anas-sintered hardness of 10.27 GPa. Due to severe degradation during supercritical waterexposure no hardness measurements were available for the tested YSZ samples. Additions ofCeO2 decreased the hardness of the sintered ceramic; however, after degradation testing,measured hardness was found to increase, as seen in Figure 4.14. Similar results were found forthe YSZ ceramics containing 10 mol% of Nd2O3 additions. In contrast, the 5 mol% Nd2O3-YSZceramic exhibited an increase in hardness compared to the pure YSZ ceramic. The hardnessdecreased after degradation testing.Figure 4.14 Measured microhardness for the various sintered ceramics before and after degradation testing.02468101214YSZ 5mol% CeO2-YSZ10mol% CeO2-YSZ15mol% CeO2-YSZ5mol% Nd2O3-YSZ10mol% Nd2O3-YSZHV (Gpa)Pre-Degradation Post-Degradation714.4 Phase EvolutionXRD analysis of the sintered ceramics before and after degradation testing (indicated as pre- andpost-degradation, respectively) was performed to analyze the phase evolution of the testedspecimens during supercritical water exposure.4.4.1 Phase Evolution of YSZ CeramicsXRD plots of the YSZ ceramic before and after degradation testing are shown in Figure 4.15.Due to severe degradation of the sample, XRD could only be performed on one post-degradedYSZ sample. After testing it was found that the zirconium carbides observed in the sinteredceramic disappeared. No new peaks were observed; however, a significant peak shift wasnoticed, indicating a change in lattice parameters which may be due to a possible transformationof a cubic system to a tetragonal system, as seen in Table 4.7. Both the measured cubic andtetragonal lattice parameters correspond closely to other findings in literature [87]-[88].Figure 4.15 XRD plots for YSZ ceramics before and after degradation testing.72Table 4.7 Lattice parameter measurement of the YSZ ceramics before and after degradation testing.Specimen Sample#Lattice Parameter (Ȧ)Pre-Degradationθ Post-DegradationϕYSZ1 5.1274at = 5.1390ct = 3.62632 5.1245 disintegrated3 5.1329 disintegratedθcubic lattice (ac), space group= Fm3mϕ tetragonal lattice (at,ct), space group = P42/nmc4.4.2 Phase Evolution of CeO2-YSZ CeramicsXRD plots of the various CeO2-YSZ ceramics before and after degradation testing are shown inFigure 4.16. It was found that during supercritical water exposure the cubic phase of the matrixwas maintained; however, analysis of the tested specimens also revealed the evolution of a newphase corresponding to the formation of CeO2. Lattice parameter measurements showed anincrease in the unit cell length with increasing CeO2 content, as seen in Table 4.8.  This indicatedthat the CeO2-ZrO2 solid solutions formed. After degradation testing, it was found that the unitcell length decreased, suggesting that CeO2 diffused out of the lattice.7374Figure 4.16 XRD plots of CeO2-YSZ ceramics before and after degradation testing with (*) peaks indicatingthe CeO2 phase.Table 4.8 Lattice parameter measurement of the CeO2-YSZ ceramics before and after degradation testing.Specimen Sample#Lattice Parameter (Ȧ) % Lattice ParameterDifferencePre-Degradationθ Post-Degradationθ5 mol%CeO2 -YSZ1 5.1557 5.1502 0.10712 5.1544 5.1517 0.05243 5.1558 5.1493 0.126110 mol%CeO2 - YSZ1 5.1748 5.1619 0.239315 mol%CeO2 - YSZ1 5.1902 5.1732 0.32752 5.1905 5.1748 0.30253 5.1882 5.1729 0.2949θcubic lattice (ac), space group= Fm3m754.4.3 Phase Evolution of Nd2O3-YSZ CeramicsXRD plots of the various Nd2O3-YSZ ceramics before and after degradation testing are shown inFigure 4.17. It was found that the cubic phase of the matrix was retained both before and afterdegradation testing. Lattice parameter measurements in the sintered specimens revealed anincrease in the unit cell length when Nd2O3 was added to YSZ, as shown in Table 4.9. Thisindicated that the Nd2O3 additions were incorporated into the ZrO2 lattice. After degradationtesting, it was found that the measured unit cell lengths of all Nd2O3-YSZ ceramics remainedwithin ±0.1% of the as-sintered lattice dimensions. This suggested that Nd2O3 diffusion out ofthe lattice did not occur at the conditions (30 MPa, 400 °C) investigated. In addition, XRDanalysis revealed that the trace phase found in the sintered 10 mol% Nd2O3-YSZ ceramicsdisappeared after testing.76Figure 4.17 XRD plots for Nd2O3-YSZ ceramics before and after degradation testing with (*) peaks indicatingthe trace phase.Table 4.9 Lattice parameter measurement of the Nd2O3-YSZ ceramics before and after degradation testing.Specimen Sample #Lattice Parameter (Ȧ)% DifferencePre-Degradationθ Post-Degradationθ5 mol%Nd2O3 -YSZ1 5.1627 5.1615 0.02322 5.1622 5.1614 0.01553 5.1614 5.1615 -0.001910 mol%Nd2O3 - YSZ1 5.1834 5.1792 0.08102 5.1793 5.1809 -0.03093 5.1809 5.1824 -0.0290θcubic lattice (ac), space group= Fm3m774.5 Microstructure EvolutionThe following sections review the microstructure analysis performed on the sintered ceramicsbefore and after degradation testing.4.5.1 Microstructure Evolution of YSZ CeramicsThe general microstructure of the YSZ ceramic before and after degradation testing exhibited ahomogenous matrix, as seen in Figure 4.18. Porosity was observed to increase in the degradedspecimen which was attributed to significant grain fall-out. As seen in Figure 4.19, grain fall-outwas present in both the as-sintered and degraded microstructures; however, grain fall-out wasmore visible on the surface of the degraded specimen. SEM-XEDS scans revealed that the levelsof the primary solutes in the as-sintered matrix were in agreement with the theoretical chemicalcomposition of the YSZ ceramic, as seen in Spectrum I in Table 4.10. After degradation testing,the matrix composition remained near the as-sintered values as shown in Table 4.10, SpectrumII. However, small particles were found on the surface of the degraded YSZ samples (indicatedas Spectrum III in Figure 4.19). Chemical analysis revealed that these particles had acomposition similar to pure ZrO2 (74 wt% Zr and 26 wt% O). This suggested that duringsupercritical water degradation testing, reverse transformation of the stabilized cubic phase to themonoclinic phase likely occurred. This process is clearly seen in the lower left corner of Figure4.19b, where an YSZ grain has transformed completely into small fragments of ZrO2 particles.These observations support XRD analysis results, where the tetragonal phase was detected in theYSZ ceramics after degradation testing (Section 4.4.1).78Figure 4.18 General microstructure of the YSZ ceramic: a) before and b) after degradation testing.Figure 4.19 SEM micrographs of the YSZ ceramic: a) before and b) after degradation testing. SEM-XEDSscans are provided in Table 4.10..IIIGrain Fall-outIII79Table 4.10 Chemical analysis of the YSZ ceramic before and after degradation testing.Sample SpectrumElement wt%Zr O Y Hf ClYSZTheoretical 63.7 25.3 11.0 - -I 63.8 25.2 9.3 1.7 0II 64.3 25.2 8.9 1.6 0III 73.3 25.9 0 0.8 04.5.2 Microstructure Evolution of CeO2-YSZ CeramicsThe microstructure of as-sintered CeO2-YSZ ceramics had two distinct phases: (1) generalmatrix and (2) large Ce-rich bands. The evolution of these phases during exposure tosupercritical water is discussed in the following sections.4.5.2.1 Matrix Evolution of CeO2 -YSZ CeramicsRepresentative images of the CeO2-YSZ ceramics before degradation testing are presented inFigure 4.20. As shown by the SEM-XEDS scans summarized in Table 4.11 (Spectrum I), the Cecontent in the matrix of the sintered CeO2-YSZ ceramics were within ~0.3, 1.1 and 2.5 wt% ofthe predicted theoretical compositions for the YSZ ceramics containing 5, 10 and 15 mol% ofCeO2, respectively. Examination of the as-sintered matrix revealed the presence of small Ce-richagglomerates dispersed throughout the matrix (as indicated by Spectrum II in Figure 4.20. Theformation of these agglomerates was attributed to the presence of powder agglomerations in theprecursor powder blends. These agglomerates were found to have significant concentrations ofCl (~15 wt%). Since no Cl was found in the surrounding matrix, this Cl content was associatedwith the Ce-rich agglomerates. X-ray linescan at the agglomerate interface in Figure 4.21showed a gradual decrease of Ce concentration near the interface that Ce was diffusing into thematrix. Therefore, during sintering, sufficient interdiffusion within the precursor powders likelyoccurred.80Figure 4.20 SEM micrographs of the as-sintered CeO2-YSZ ceramics: a) 5 mol% CeO2-YSZ, b) 10 mol%CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM-XEDS scans are provided in Table 4.11.Table 4.11 Chemical analysis of the as-sintered matrix in the CeO2-YSZ ceramics.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZTheoretical 59.6 24.9 10.3 5.2 - -I 60.2 24.5 9.0 4.9 1.0 0.4II 11.5 11.7 3.1 57.5 0 16.210 mol%CeO2-YSZTheoretical 55.6 24.5 9.6 10.3 - -I 59.9 20.5 9.0 9.2 1.4 0.0II 12.4 13.4 2.6 57.0 0 14.615 mol%CeO2-YSZTheoretical 51.7 24.0 9.0 15.3 - -I 58.2 18.8 8.8 12.9 1.3 0.0II 13.7 8.1 2.6 60.7 0.1 14.8I I III IIII81Figure 4.21 Linescan across a Ce-rich agglomerate in the 5 mol% CeO2-YSZ ceramicRepresentative images of the CeO2-YSZ ceramic microstructures after degradationtesting are provided in Figure 4.22. Compared to the as-sintered matrix, the Ce content in thedegraded matrix was found to increase by 15, 3 and 1 wt% for the 5, 10 and 15 mol% CeO2-YSZ0.4 µm82ceramics, respectively (Table 4.12, Spectrum III). In addition, the Ce-rich agglomerationsdisappeared after degradation testing and extensive porosity was observed in the degradedspecimens, especially at grain triple points. These observations suggest that the preferentialdissolution of the Ce-rich agglomerates occurred in the vicinity of grain boundaries duringdegradation testing.Figure 4.22 SEM micrographs of CeO2-YSZ ceramics after degradation testing: a) 5 mol% CeO2-YSZ, b) 10mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM-XEDS scans are provided in Table 4.12.Table 4.12 Chemical analysis of CeO2-YSZ ceramics after degradation testing.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZTheoretical 59.6 24.9 10.3 5.2 - -III 43.3 21.8 6.8 26.3 1.7 0.110 mol%CeO2-YSZTheoretical 55.6 24.5 9.6 10.3 - -III 55.5 23.8 7.8 11.9 1.0 015 mol%CeO2-YSZTheoretical 51.7 24.0 9.0 15.3 - -III 53.0 23.4 8.1 14.2 1.3 0Closer analysis revealed that the large increase in Ce content in the 5 mol% CeO2-YSZceramic matrix was due to the formation of a Ce-rich layer on the surface of the tested sample, asindicated by Spectrum V in Figure 4.23a. Remnants of this Ce-rich layer could also be found onthe surface of the 10 and 15 mol% CeO2-YSZ ceramics, as shown in Figure 4.23b and Figure4.23c, respectively. This suggested that during degradation testing a Ce-rich layer formed on theIIIIIIIII83surface of the CeO2-YSZ specimens; however, in the 10 and 15 mol% CeO2-YSZ ceramics, thislayer would exhibit a spallation behaviour. SEM-XEDS scans of the matrix beneath this layerrevealed a composition close to the theoretical values, as seen in Table 4.13, Spectrum IV. Thissuggested that this oxide layer likely formed from Ce-rich agglomerates and bands found withinthe as-sintered matrix.Figure 4.23 SEM micrographs of the Ce-rich oxide layer formed on the CeO2-YSZ ceramics afterdegradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM-XEDSscans are provided in Table 4.13.VIVVIVIVV84Table 4.13 Chemical analysis of the Ce-rich oxide layer formed on the CeO2-YSZ ceramics after degradationtesting.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZTheoretical 59.6 24.9 10.3 5.2 - -IV 59.7 24.3 9.1 5.9 1.0 0V 35.8 30.2 6.2 26.9 0.9 010 mol%CeO2-YSZTheoretical 55.6 24.5 9.6 10.3 - -IV 62.7 16.7 9.1 10.0 1.5 0V 44.8 27.5 6.7 20.2 0.7 0.115 mol%CeO2-YSZTheoretical 51.7 24.0 9.0 15.3 - -IV 52.7 23.3 7.7 15.1 1.0 0V 36.8 30.5 5.6 26.5 0 0As discussed in Section 4.2.2, a macroscopic spallation behaviour was observed in theCeO2-YSZ ceramics during supercritical water testing. Analysis of the spalled surfaces revealedthat small agglomerates of Ce-rich particles were present between grains as indicated bySpectrum VII in Figure 4.24. This suggested that the Ce-rich agglomerates observed on thesurface of the as-sintered ceramic likely extended below the surface of the specimen.Furthermore, the spalled surfaces showed significant grain fall-out similar to what was found inthe pure YSZ sample after degradation testing. Chemical analysis of the exposed grains revealedthat the chemical composition was similar to the as-sintered matrix values (as indicated bySpectrum VI in Table 4.14). In contrast, chemical analysis of the Ce-rich particles revealedvariable amount of Ce (as shown by Spectrum VII in Table 4.14). This was attributed to theuneven surface of the spalled regions which resulted in unwanted x-rays from protruding grainswhich surrounded the Ce-rich particles.85Figure 4.24 SEM micrographs of spalled surfaces in the CeO2-YSZ ceramics after degradation testing: a) 5mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM-XEDS scans are provided inTable 4.14.Table 4.14 Chemical analysis of the spalled surface in CeO2-YSZ ceramics after degradation testing.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZTheoretical 59.6 24.9 10.3 5.2 - -VI 59.7 24.3 9.1 5.9 1.0 0VII 26.6 25.1 4.4 43.5 0.4 010 mol%CeO2-YSZTheoretical 55.6 24.5 9.6 10.3 - -VI 55.4 23.8 8.0 11.2 1.5 0.1VII 49.6 22.4 4.7 23.9 0.9 0.315 mol%CeO2-YSZTheoretical 51.7 24.0 9.0 15.3 - -VI 53.8 23.5 7.5 13.7 1.5 0VII 19.1 17.8 3.0 59.5 0.5 04.5.2.2 Band Evolution of CeO2 -YSZ CeramicsThe general microstructure of the CeO2-YSZ ceramics before degradation testing is shown inFigure 4.25. The as-sintered CeO2-YSZ ceramics consisted of a matrix containing large bandsorientated perpendicular to the applied sintering pressure. Chemical analysis revealed that thesebands contained high concentrations of Ce (~ 76 wt%), Spectrum IX in Table 4.15. Thus, theformation of these Ce-rich bands was attributed to the improper mixing of the precursorpowders, resulting in large CeO2 powder agglomerations in the powder blends. Linescan analysisrevealed a sharp increase in Ce content along the band interface indicating that Ce was notVIVIIVIVIVIIVII86gradually diffusing into the matrix, as shown in Figure 4.27. However, closer analysis revealedthat these Ce-rich bands were surrounded by a Ce-rich interface region where localized elevatedconcentrations of Ce were detected, as indicated by Spectrum VIII in Figure 4.26. The presenceof these bands contributed to the discrepancy observed in the matrix composition presented inSection 4.5.2.1. Ce-rich bands were found in all three compositions (5, 10 and 15 mol% CeO2);however the number as well as the size of the bands were found to increase with increasing CeO2concentration, as shown in Figure 4.25. This supported why a larger composition discrepancywas observed for the higher dopant levels of CeO2 (15 mol%) compared to the lower dopantlevels (5 mol%). In addition, large voids were often observed at the tips of these bands indicatingthat these bands hindered densification (Figure 4.28).Figure 4.25 General microstructure of the as-sintered CeO2-YSZ ceramics: a) 5 mol% CeO2-YSZ, b) 10mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ.87Figure 4.26 Interface region surrounding Ce-rich bands in as-sintered CeO2-YSZ ceramics: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM –XEDS scans are provided in Table 4.15.Table 4.15 Chemical analysis of Ce-rich bands in the as-sintered CeO2-YSZ ceramics.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZTheoretical 59.6 24.9 10.3 5.2 - -VIII 13.6 16.2 2.2 52.1 0.3 15.6IX 0 3.8 0.4 79.8 0 16.110 mol%CeO2-YSZTheoretical 55.6 24.5 9.6 10.3 - -VIII 11.9 10.3 2.4 59.9 0.1 15.4IX 0 7.3 0.1 74.0 0 18.615 mol%CeO2-YSZTheoretical 51.7 24.0 9.0 15.3 - -VIII 31.6 21.2 4.7 41.9 0.6 0IX 0 7.5 0 73.8 0 18.7VIIIVIIIVIIIIXIXIX88Figure 4.27 Linescan across a CeO2 band in the as-sintered 10 mol% CeO2-YSZ ceramic.10 µm89Figure 4.28 Void formation (arrows) at the tips of an CeO2 band in the as-sintered 15 mol% CeO2-YSZceramic.After degradation testing it was found that the number of Ce-rich bands remaining in thematrix decreased and elongated voids were observed, as shown in Figure 4.29. Compared to thechemical composition of the observed as-sintered bands, SEM-XEDS scans of the bands in thetested specimens revealed a higher Ce concentration which was due to the absence of Cl, asshown in Table 4.16, Spectrum X. The removal of Cl from the Ce-rich bands was attributed tothe dissolution of this element into the supercritical water medium. In the 15 mol% CeO2-YSZceramic, it was found that cracks would generally initiate from the remaining Ce-rich bands.These cracks often propagated towards the sample surface, where spallation was subsequentlyobserved (Figure 4.30). Linescan analysis across remaining Ce-rich bands in the degradedspecimens revealed that Ce was not diffusing into the matrix (Figure 4.31). This indicated thatthe disappearance of bands was due to the dissolution of Ce into the supercritical water medium.Furthermore, large Ce-rich areas formed on the surface of tested CeO2-YSZ ceramics, as seen inFigure 4.32. SEM-XEDS scans (Table 4.17, Spectrum XI) revealed high counts of Ce (rangingbetween 60-80 wt%) in these areas suggesting the formation of a Ce-rich layer. Theseobservations correspond to the CeO2 phase found in the XRD results of the degraded CeO2-YSZceramics (Section 4.4.2).90Figure 4.29 Ce-rich bands (rectangles) and Ce-rich areas (arrows) found in CeO2-YSZ ceramics afterdegradation testing: a) 5 mol% CeO2-YSZ, b) 10 mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ.Figure 4.30 Crack formation from Ce-rich bands in 15 mol% CeO2-YSZ ceramic after degradation testing.Labelled SEM-XEDS scan is provided in Table 4.16.SpalledSurfaceCrackPropagationX91Table 4.16 Chemical analysis of remaining Ce-rich bands in 15 mol% CeO2-YSZ ceramic after degradationtesting.Sample SpectrumElement wt%Zr O Y Ce Hf Cl15 mol%CeO2 -YSZX 0 16.8 0.3 82.8 0 0.1Figure 4.31 Linescan across a Ce-rich band in the 15 mol% CeO2-YSZ ceramic after degradation testing.20 µm92Figure 4.32 Ce-rich areas in CeO2-YSZ ceramics after degradation testing: a) 5 mol% CeO2-YSZ, b) 10mol% CeO2-YSZ and c) 15 mol% CeO2-YSZ. SEM-XEDS scans are included in Table 4.17.Table 4.17 Chemical analysis of the Ce-rich areas found in the CeO2-YSZ ceramics after degradation testing.Sample SpectrumElement wt%Zr O Y Ce Hf Cl5 mol%CeO2 -YSZXI 0 14.8 2.7 81.7 0.7 0.110 mol%CeO2-YSZXI 16.3 21.2 4.6 57.5 0.2 0.215 mol%CeO2-YSZXI 1.1 14.8 0.3 83.5 0 0.34.5.3 Microstructure Evolution of Nd2O3-YSZ CeramicsThe following sections discuss the microstructure evolution of the Nd2O3-YSZ ceramics duringsupercritical water degradation testing.4.5.3.1 Matrix Evolution of Nd2O3 -YSZ CeramicsRepresentative SEM micrographs of the as-sintered matrix microstructure for both the 5 and 10mol% Nd2O3-YSZ ceramics are provided in Figure 4.33. Before degradation testing, the matrixof the as-sintered Nd2O3-YSZ ceramics was found to have slightly lower Nd content (3 and 7wt% for the 5 and 10 mol% Nd2O3-YSZ ceramics, respectively) than the predicted theoreticalvalues, as seen in Table 4.18, Spectrum I. Furthermore, as indicated by Spectrum II in Figure4.33, Nd-rich agglomerates were present in the matrix. The formation of these matrixagglomerates was attributed to the presence of small Nd2O3 agglomerations in the precursorpowders. However, SEM-XEDS scans revealed that these Nd-rich agglomerates also had a highXIXIXI93concentration of Cl (~12.75 wt%). As seen in Spectrum I in Table 4.18, no Cl was present in theceramic matrix thus; the Cl concentration was associated with the Nd-rich agglomerates.Linescan analysis across these agglomerates revealed that Nd was diffusing into the matrix, asshown by the gradual increase in Nd concentration across the agglomerate interface in Figure4.34.Figure 4.33 SEM micrographs of as-sintered Nd2O3-YSZ ceramics: a) 5 mol% Nd2O3-YSZ and b) 10 mol%Nd2O3-YSZ. SEM-XEDS scans are presented in Table 4.18.Table 4.18 Chemical analysis of the as-sintered microstructure of the Nd2O3-YSZ ceramics.Sample SpectrumElement wt%Zr O Y Nd Hf Cl5 mol%Nd2O3 -YSZTheoretical 56.0 24.0 9.7 10.3 - -I 62.2 19.5 9.2 7.5 1.6 0II 20.4 10.7 5.2 50.4 0.8 12.510 mol%Nd2O3 -YSZTheoretical 49.5 22.9 8.6 19.0 - -I 57.1 20.2 8.8 12.8 1.0 0II 19.8 13.0 4.1 49.8 0.3 13.0IIIIII94Figure 4.34 Linescan across a Nd-rich agglomerate in the matrix of the as-sintered 5 mol% Nd2O3-YSZceramic.After degradation testing, it was found that the Nd-rich agglomerations disappeared and aporosity network was observed, as shown in Figure 4.35. This indicated that during supercriticalwater degradation testing preferential dissolution of the Nd-rich agglomerates occurred. SEM-1 µm95XEDS analysis of the degraded matrix revealed that the Nd content was lower than thetheoretical composition, but it remained within ~1% of the as-sintered value, as shown in Table4.19, Spectrum III, suggesting that during degradation testing Nd diffusion out of the lattice wasnot significant. However, degraded 10 mol% Nd2O3-YSZ ceramics showed the formation of Nd-rich precipitates (crystals). These crystals were located in regions of high porosity and alongcrack interfaces, as shown in Figure 4.36. SEM-XEDS analysis revealed that these crystals hadhigh concentrations of Nd and O, but in contrast to the Nd-rich agglomerates observed in the as-sintered specimen, these crystals exhibited low counts of Cl (Table 4.20, Spectrum IV). Theformation of these crystals was attributed to the high pressure and temperature environmentwhich may have resulted in transformation of metastable Nd-rich phases in the as-sinteredspecimen. Barrett et al. [71] observed a similar crystal formation behaviour during supercriticalwater exposure of TiO2-doped zirconia ceramics stabilized with 7 mol% of yttria.Figure 4.35 SEM micrographs of Nd2O3-YSZ ceramics after degradation testing: a) 5 mol%Nd2O3-YSZ and b) 10 mol% Nd2O3-YSZ. SEM-XEDS scans are provided in Table 4.19..III III96Table 4.19 Chemical analysis of the Nd2O3-YSZ ceramics after degradation testing.Sample SpectrumElement wt%Zr O Y Nd Hf Cl5 mol%Nd2O3 -YSZTheoretical 56.0 24.0 9.7 10.3 - -III 54.6 29.9 7.8 6.7 1.0 010 mol%Nd2O3 -YSZTheoretical 49.5 22.9 8.6 19.0 - -III 54.2 25.1 7.8 11.6 1.3 0Figure 4.36 Crystal formation in 10 mol% Nd2O3-YSZ ceramic after degradation testing: a) along pores andb) along crack interphases. SEM-XEDS scans of the crystals are summarized in Table 4.20.Table 4.20 Chemical analysis of crystals formed in the 10 mol% Nd2O3-YSZ ceramic after degradationtesting.Sample SpectrumElement wt%Zr O Y Nd Hf Cl10 mol%Nd2O3 -YSZIV 7.9 36.1 6.8 48.6 0.6 0V 9.3 38.6 4.9 46.3 0.6 0.34.5.3.2 Band Evolution of Nd2O3 -YSZ CeramicsRepresentative images of the Nd2O3-YSZ ceramics before testing are shown in Figure 4.37. Theas-sintered microstructure of the 5 mol% Nd2O3-YSZ ceramic had a relatively homogenousmatrix, with areas of elevated porosity as indicated in Figure 4.37a. However, the microstructureVIV97of the 10 mol% Nd2O3-YSZ ceramic showed the formation of two distinct band types: (1)smaller white bands (minority) and (2) larger grey bands (majority) as indicated in Figure 4.37b.The formation of these bands in supported the lower Nd content observed in the matrix of 10mol% Nd2O3-YSZ ceramic as reviewed in Section 4.5.3.1. Both bands were alignedperpendicular to the applied pressure during sintering which indicated that these bands formeddue to the presence of Nd2O3 powder agglomerations in the precursor powder blends. Chemicalanalysis revealed that the smaller white bands were composed primarily of Nd (~72 wt%), asseen in Table 4.21 (Spectrum VI). Linescan analysis across the interface of the white bandsshowed Nd was not diffusing into the matrix as indicated by the sudden jump in Nd content atthe band boundary in Figure 4.41. In contrast, larger grey bands were found to consist of an Nd-rich (~38 wt%) matrix, as shown in Spectrum VII in Table 4.21. Linescan analysis across thegrey bands revealed that Nd was diffusing into the matrix from the band as shown by the gradualincrease in Nd concentration along the band boundary in Figure 4.39.Figure 4.37 General microstructure of the as-sintered Nd2O3-YSZ ceramics: a) 5 mol% Nd2O3-YSZ and b) 10mol% Nd2O3-YSZ. Labelled SEM-XEDS are provided in Table 4.21.PorosityRegionVIVII98Table 4.21 Chemical analysis of Nd bands found in the as-sintered 10 mol% Nd2O3-YSZ ceramic.Sample SpectrumElement wt%Zr O Y Nd Hf Cl10 mol%Nd2O3 -YSZVI 1.6 8.9 0.5 71.6 0 17.4VII 34.9 15.6 4.4 38.4 0.4 6.3Figure 4.38 Linescan across a white band in the as-sintered 10 mol% Nd2O3-YSZ ceramic.10 µm99Figure 4.39 Linescan across a grey band found in the as-sintered 10 mol% Nd2O3-YSZ ceramic.After degradation testing, the large grey bands were found to disappear and instead, asshown in Figure 4.40a, porosity was observed. The formation of porosity bands was attributed tothe preferential dissolution of the Nd2O3 band agglomerates. In contrast, the white bands werefound to remain in the degraded matrix and were often fractured, as seen in Figure 4.40b.Chemical analysis of the fractured bands revealed that no Cl remained in the white bands after80 µm100degradation testing. Furthermore, compared to the chemical composition of the as-sintered bandsthe Nd content in the remaining white bands was found to decrease by ~26 wt% afterdegradation testing, as seen in Table 4.22, Spectrum VIII. Linescan analysis across the whitebands after degradation testing showed that Nd was not diffusing into the matrix (Figure 4.41).This indicated that the decrease in Nd content within the white band after degradation testingwas likely due to the dissolution of Nd into the supercritical water medium. Furthermore,extensive matrix cracking was also observed in the tested 10 mol% Nd2O3-YSZ ceramics. Thesecracks seemed to originate at the white bands and then propagated through the weaker porosityregions found in the tested specimens (Figure 4.40). This suggested that banding contributed tocrack initiation and propagation.Figure 4.40 SEM micrographs of 10 mol% Nd2O3-YSZ ceramics after degradation testing. Labelled SEM-XEDS scans are provided in Table 4.22.VIIIPorosityBandWhiteBandPropagatingCracks101Table 4.22 Chemical analysis of white bands found in 10 mol% Nd2O3-YSZ ceramic after degradation testing.Sample SpectrumElement wt%Zr O Y Nd Hf Cl10 mol%Nd2O3 -YSZVIII 41.9 25.3 5.1 26.5 1.2 0Figure 4.41 Linescan across a white band in a 10 mol% Nd2O3-YSZ ceramic after degradation testing.20 µm1024.6 Solution AnalysisRepresentative photos of the solutions retrieved after supercritical water testing of YSZ, CeO2-YSZ and Nd2O3-YSZ ceramics are shown in Figure 4.42. As seen in the figure, distinct redcolorations were present in the doped YSZ ceramics. Preliminary ICP testing revealed that inaddition to Y3+ ions, corresponding dopant ions (either Ce4+ or Nd3+) were present in the coloredsolutions. A brief summary of the ICP results are provided in Table 4.23. In general, theconcentration of rare earth ions found in the solution increased with increasing dopantconcentration. Furthermore, the concentration of Y3+ in the solutions retrieved from either theCeO2-YSZ or Nd2O3-YSZ ceramics was found to be greater than the Y3+ concentration present inthe solution retrieved from the pure YSZ ceramic. A similar trend was observed for Zr4+ ions.However, overall measured values were found to be significantly lower than the theoreticalconcentration values which were calculated assuming the weight loss of the tested sample wasdue entirely to the dissolution of the ceramic. This difference was likely attributed to two factors:first, retrieved solutions often contained solid particulates which were associated with eitherspalled ceramic pieces that did not dissolve during supercritical water exposure or possiblyprecipitates that formed during cooling. Second, due to improper solution acidification aftertesting the measured ICP-MS results may not be accurate since, between solution retrieval andtesting, the solution may have changed by forming precipitates, dissolving ceramic remnants orthrough evaporation. In addition, the retrieved solutions possibly contained contaminants fromsamples tested previously in the 316L steel vessels. This was likely attributed to experimentalchallenges during vessel cleaning indicating that ultrasonic cleaning of the vessels between testswas inadequate and rather an acid wash in nitric acid should be used. Thus, further testing isrequired to confirm the ICP-MS results.103Figure 4.42 Retrieved solutions after degradation testing: a) YSZ, b) CeO2-YSZ and c) Nd2O3-YSZ ceramics.Table 4.23 ICP-MS results from solutions retrieved after degradation testing. Due to cross contamination theresults of the mass balance are not presented.SampleElement Concentration (ppb)Zr4+ Y3+ Ce4+ Nd3+YSZTheoretical 10,718,165 1,855,876 -- --Measured 2E-1 9 -- --5 mol%CeO2-YSZTheoretical 131,108 22,694 11,462 --Measured 16 38 121 --10 mol%CeO2-YSZTheoretical 1,826,351 315,575 339,631 --Measured 45 126 9,621 --15 mol%CeO2-YSZTheoretical 3,121,054 542,188 924,408 --Measured 15 37 2,022 --5 mol%Nd2O3-YSZTheoretical 56,041 9,685 -- 10,288Measured 57 132 -- 2,23710 mol%Nd2O3-YSZTheoretical 406,116 70,380 -- 156,071Measured 60 485 -- 14,189104Chapter 5: DiscussionThis section provides a critical discussion on the results outlined in Chapter 4. The discussion isdivided into two categories: (1) fabrication and (2) degradation behaviour of the novel ceramicmaterials. For each category, the effect of either CeO2 or Nd2O3 addition is compared to thebehaviour of the pure YSZ sample.5.1 Spark Plasma Sintering FabricationThe following sections discuss the fabrication of the various sintered ceramic materials with afocus on the grain size, densification behaviour and phase stabilization.5.1.1 Spark Plasma Sintering of YSZ CeramicsSPS of pure YSZ powders was performed at a temperature of 1400 °C, pressure of 30 MPa for adwell time of 7 min. The resulting ceramics were found to have a high relative density (~96%)with an average grain size of 9.84 µm. A similar densification behaviour was observed byAnselmi-Tamburini et al. [80] where it was predicted that 8 mol% yttria stabilized zirconiapowders processed by SPS for 5 min under a pressure of 30 MPa and at a temperature of 1375 °Cwould result in a ceramic with a relative density of 95%. In addition, Anselmi-Tamburini et al.[80] predicted that the resulting ceramic would have a grain size of 210 nm which correlated to a97% grain growth when comparing the as-sintered grain size to the initial precursor powders(210 and 6.6 nm, respectively). A similar increase in grain size was observed in the present work,where the measured as-sintered grain size correlated to a 95% grain growth (9.84 µm from 0.5µm, respectively). These observations suggest that the SPS performed in this research wascomparable to those published in the literature. Also, the microhardness of the sintered pure YSZceramic was 10.27 GPa, which is in agreement with the work of Abrham and Grizner [89] wherecubic zirconia stabilized with 8 mol% of yttria (theoretical density of 96%) was found to have ahardness of 9.2 GPa.XRD analysis of the as-sintered ceramic revealed that despite traces of monocliniczirconia found in the precursor powders, fully stabilized cubic zirconia was present aftersintering. Measured lattice parameter was ~ 5.1282 Å, which closely corresponded to valuesreported in the literature [88]. SEM-XEDS scans of the sintered matrix revealed a homogenous105composition close to the predicted theoretical values. Traces of Hf in the matrix were likely theresult of the presence of Hf in the precursor powders. No traces of Cl were found in the as-sintered YSZ ceramics.5.1.2 Spark Plasma Sintering of CeO2-YSZ CeramicsDepending on the composition, SPS of CeO2-YSZ ceramics was performed at a temperature ofeither 1300 or 1400 °C, and at a dwell time of either 5 or 7 min. Variable sintering temperatureand dwell times were necessary in order to fabricate bulk ceramic disks that did not fracture uponejection from the sintering die. All CeO2-YSZ ceramics were sintered under a constant pressureof 30 MPa.5.1.2.1 Effect of CeO2 on Sintered Grain Size and DensificationWhen 5, 10 and 15 mol% of CeO2 was added to YSZ, the grain size decreased by 84, 62 and79%, respectively. Similar refinement of YSZ microstructure with CeO2 additions was alsoobserved by Wang et al. [90] where cubic stabilized zirconia containing 8 mol% of yttria wasfabricated with various amount of CeO2 additions (0-25 mol%). It was found that increasingCeO2 additions from 0 to 25 mol% resulted in a decrease in the crystallite size from 9 to 4 nm.The authors attributed this behaviour to two mechanisms: (1) solute drag and (2) defectclustering as well as lattice distortion due to the incorporation of Ce4+ ions in the lattice.However, in the presented work, the strongest grain refinement was observed for the 5 mol%CeO2-YSZ ceramic, which differs from the findings presented by Wang et al. [90]. This may beattributed to the difference in sintering parameters used in this work to fabricate the YSZ andCeO2-YSZ ceramics. For the pure YSZ ceramic as well as the 10 and 15 wt% CeO2-YSZceramics, the selected sintering parameters remained constant (at a temperature, dwell time andpressure of 1400 °C, 7 min and 30 MPa, respectively). However, in order to fabricate a wholespecimen that did not fracture upon ejection from the die, the 5 mol% CeO2-YSZ ceramic wassintered at a lower temperature (1300 °C) and a shorter dwell time (5 min). SPS densificationstudies of 8 mol% yttria stabilized zirconia powders by Dahl et al. [78] revealed that slightchanges in SPS temperature and dwell time affect the resulting grain size. For example, Dahl etal. [78] observed that grain size decreased by 87% when sintering temperature was reduced from1250 to 1200 °C (while the sintering parameters of pressure and dwell time remained constant at70 MPa and 5 min, respectively). A similar relationship was observed with a reduction of dwell106time, where a decrease from 10 to 5 min reduced the grain size by 11% (while the sinteringpressure and temperature remained constant at 70 MPa and 1200 °C, respectively) [78].Therefore the variation of the SPS parameters in the 5 mol% CeO2-YSZ ceramic likelycontributed to the observed grain size reduction.Along with grain size reduction, a decrease in the relative density with CeO2 additionswas also observed. Thus, CeO2 additions hindered densification, resulting in a decrease in thedensity of the sintered CeO2-YSZ ceramics. For example, the CeO2-YSZ ceramics had anaverage relative density of ~87.5%, corresponding to a decrease of 8% when compared to the as-sintered YSZ ceramic. Similar results were reported by Bekale et al. [91] where, despite theselected processing parameters as well as technique used (cold isostatic pressing or hot isostaticpressing), CeO2 additions to 8 mol% yttria stabilized zirconia consistently reduced the as-sintered density of the resulting ceramic. It was found that the fabricated ceramics all had adensity greater than 96%; however, ceramics containing 10 mol% of CeO2 had a relative densitythat ranged between 1-3% below the as-sintered density of the pure 8 mol% yttria stabilizedzirconia [91].When 5, 10 and 15 mol% of CeO2 was added to YSZ, the hardness decreased by 11, 19and 26%, respectively. It is important to note that the refinement of grain size with CeO2additions did not necessarily increase the as-sintered hardness of the ceramic. This behaviour issupported by SPS densification studies performed by Dahl et al. [78] where for 8 mol% yttriastabilized cubic zirconia the grain size (ranging between 0.2 to 12 µm) did not have an impact onthe as-sintered hardness. Thus, the observed decrease in hardness in the CeO2-YSZ ceramics inthe present work was likely associated with the combination of two factors: (1) the increase(~8%) in porosity in the sintered CeO2-YSZ specimens and (2) the development of Ce-rich bandsin the sintered matrix. After degradation testing a slight hardness increase (~1 GPa) wasobserved for all three of the CeO2-YSZ ceramic compositions. This increase in hardness wasassociated to the disappearance of Ce-rich bands in the degraded specimens.5.1.2.2 Effect of CeO2 on Phase StabilizationXRD analysis revealed that while the cubic YSZ was present in all as-sintered CeO2-YSZceramics, a shift of the cubic peaks could be observed at the higher angles when CeO2concentration increased in the ceramic, as shown in Figure 5.1. This peak shift was found to be107more prominent at elevated CeO2 content. With respect to the as-sintered YSZ ceramic, CeO2additions of 5, 10 and 15 mol% increased the measured lattice parameter by 0.5, 0.9 and 1.2%,respectively (as previously shown in Table 4.8). This indicated that Ce4+ ions were effectivelyincorporated into the ZrO2 lattice. Similar peak shift with increasing CeO2 additions to YSZceramics was reported by Wang et al. [90]. Measured lattice parameters reported by Wang et al.[90] closely corresponded to the unit cell dimensions measured for the CeO2-YSZ ceramics withthe largest discrepancy observed for the 15 mol% CeO2-YSZ ceramic (~0.3% difference).Figure 5.1 Lattice shift present in the as-sintered CeO2-YSZ ceramics.5.1.3 Spark Plasma Sintering of Nd2O3-YSZ CeramicsTo ensure whole specimens upon ejection form the graphite die, two different sinteringtemperatures were selected for the 5 and 10 mol% Nd2O3-YSZ ceramics (1400 and 1300 °C,respectively). Sintering pressure and dwell time were selected to remain constant at 40 MPa and5 min, respectively, for both the Nd2O3 specimens.1085.1.3.1 Effect of Nd2O3 on Sintered Grain Size and DensificationWith respect to the as-sintered pure YSZ ceramic, 5 and 10 mol% of Nd2O3 additions reducedthe matrix grain size by 76 and 92%, respectively. To the best of the author’s knowledge, there isno published literature on the effect of Nd2O3 additions on the microstructure and densificationof cubic stabilized zirconia ceramics. However, grain refinement of partially stabilized zirconiaby Nd2O3 additions has been reported by Salehi et al. [9]. Partially stabilized zirconia powderscontaining 1 mol% of Y2O3 were blended with either 1 or 2 mol% of Nd2O3. Powders wereconsolidated through SPS at a pressure of 62 MPa for 3 min at a temperature of 1450 °C. It wasfound that with increasing Nd2O3 content the grain size of the sintered ceramics decreased from0.4 to 0.2 µm. Furthermore, the grain size of the Nd2O3 containing ceramics was found to beconsistently smaller when compared to zirconia ceramics only containing either 2 or 3 mol% ofY2O3 (where the observed grain size was 0.7 and 0.5 µm for the zirconia ceramics stabilized with2 and 3 mol% Y2O3, respectively). These results are in agreement with the observed grainrefinement in the as-sintered Nd2O3-YSZ ceramics. Thus, similar to CeO2 additions, Nd2O3additions have a tendency to hinder grain growth during sintering of YSZ. However, the effect ofthe selected SPS processing parameters must also be considered as a contributing factor, sincethe Nd2O3-YSZ ceramics were sintered with shorter dwell times and, in addition, the 10 mol%Nd2O3 –YSZ ceramic was sintered at a lower temperature.When 5 and 10 mol% of Nd2O3 was added to YSZ, the average relative density of the as-sintered specimens decreased by ~4%. This suggests that Nd2O3 addition likely hindereddensification, as was also observed for the CeO2-YSZ ceramics. However, compared to thesintered YSZ ceramics, Nd2O3-YSZ ceramics were fabricated for shorter dwell times and, for the10 mol% Nd2O3-YSZ ceramic, sintering was performed at a temperature 100 °C lower than theYSZ ceramic. Both these parameters are known to affect grain growth and density of the SPSprocessed ceramic [78]. Furthermore, sintering of the Nd2O3-YSZ ceramics was performed undera higher pressure of 40 MPa. Studies by Anselmi-Tamburini et al. [92] revealed that the finaldensity of the sintered ceramic is strongly influenced by the applied pressure during SPSprocessing where an increase in pressure increases the density of the sintered ceramic. Zirconiananopowders stabilized with 8 mol% of yttria were SPS at a temperature of 1200 °C for 5 min atvariable pressures. Anselmi-Tamburini et al. [92] found that increasing the sintering pressure109from 30 to 40 MPa resulted in a 9% increase in relative density. Higher pressures furtherincreased the relative density and sintered specimens with a 99% relative density were achievedat a pressure of 140 MPa [92]. However, in the present work, Nd2O3-YSZ ceramics were foundto have a lower density than the pure YSZ ceramics despite the increase in sintering pressure.This contradiction further supports that the Nd2O3 additions hindered densification duringsintering.When 5 mol% of Nd2O3 was added to YSZ, the hardness increased by 22%, while itdecreased by 38% for YSZ ceramics containing 10 mol% of Nd2O3 additions. Both the Nd2O3-YSZ ceramics exhibited a similar relative density; however, microstructure analysis revealed theformation of large Nd-rich bands in the matrix of the 10 mol% Nd2O3-YSZ ceramic. Thus, thedecrease in the as-sintered hardness of the 10 mol% Nd2O3-YSZ ceramic was possibly related tothe presence of these bands. Compared to the as-sintered values, the hardness of the degradedNd2O3-YSZ ceramics were found to decrease by 18% for the 5 mol% Nd2O3-YSZ ceramic, butincrease by 64% for the 10 mol% Nd2O3-YSZ ceramic. Similar to the CeO2-YSZ ceramics, theobserved hardness increase in the 10 mol% Nd2O3-YSZ ceramic was attributed to thedisappearance of large Nd-rich bands. However, the hardness decrease observed in the degraded5 mol% Nd2O3-YSZ ceramic was attributed to the localized dissolution of Nd-rich agglomerateswhich increased the surface porosity in the degraded specimen. This behaviour was alsoobserved in the 10 mol% Nd2O3-YSZ ceramic; however, the dissolution of large Nd-rich bandscontributed to the increase of the hardness more so than the increase in porosity.5.1.3.2 Effect of Nd2O3 on Phase StabilizationXRD analysis revealed that after sintering, fully stabilized cubic zirconia matrix was retained inthe Nd2O3-YSZ ceramics despite the formation of an unknown trace phase in the 10 mol%Nd2O3-YSZ ceramic. Compared to the CeO2-YSZ ceramics, the Nd2O3-YSZ ceramics exhibiteda similar peak shift behaviour, suggesting that Nd3+ ions were effectively incorporated into thelattice, as shown in Figure 5.2. Lattice parameter measurements revealed a 0.6 and 1% increasefor the 5 and 10 mol% Nd2O3-YSZ ceramics, respectively, when compared to the as-sinteredYSZ ceramic. This lattice parameter increase was slightly greater (0.1% for both 5 and 10 mol%dopant levels) than the observed lattice increase in the YSZ doped with CeO2. This may be110attributed to the slightly larger ionic size of the Nd3+ ion when compared to the Ce4+ ion (1.109 Åand 0.97 Å, respectively) [23].Figure 5.2 Lattice shift present in the as-sintered Nd2O3-YSZ ceramics.5.2 Degradation BehaviourThis section reviews the degradation behaviour of the various sintered ceramics particularlyfocusing on the observed degradation rates and phase stability during exposure to supercriticalwater.5.2.1 Degradation Behaviour of YSZ CeramicsPure YSZ ceramic specimens were found to degrade rapidly during supercritical water exposure.Out of the three tested specimens, only one sample remained as one piece after testing. Thisresult was attributed to the preferential initiation of degradation in localized regions of highporosity which formed in the specimen during sintering. Non-homogenous density has beenobserved in other SPS materials and was attributed to the presence of a temperature gradient inthe sample during sintering [93], [94]. SPS studies by Mazzia et al. [93] found that temperature111gradients were related to the current intensity applied during sintering (ie. the applied heatingrate). It was found that higher heating rates resulted in an uneven current distribution leading tomicrostructure and porosity variation along the radial direction of the sintered samples [93].These results suggest that the variable degradation behaviour of the YSZ specimens may beattributed to a microstructure and density gradient in the sectioned specimens as both thesefactors have a direct influence on degradation rates [12].XRD analysis of the degraded sample revealed that a tetragonal phase transformationoccurred during testing. This suggested that the rapid degradation of the pure YSZ ceramic wasalso facilitated by the volumetric lattice change associated with a monoclinic phasetransformation. SEM analysis confirmed this observation as small ZrO2 flakes found dispersedalong the surface of the degraded matrix, supporting that tetragonal to monoclinictransformations were likely present in the degraded specimen. The monoclinic phase was notfound during XRD analysis, because the associated volume expansion enabled grain fall-out ofthe matrix. This likely resulted in the volume of the monoclinic phase remaining in the degradedmatrix to fall below the required detection limit of the XRD instrument.Despite the presence of tetragonal and monoclinic transformations, chemical analysisrevealed that the yttrium content remained relatively constant after exposure of YSZ tosupercritical water. This suggested that cubic to tetragonal transformation occurred due to anoxygen vacancy annihilation mechanism, rather than due to a destabilization mechanism.However, the presence of pure ZrO2 particles indicated that yttrium depletion of the sinteredmatrix also occurred. This observation was further confirmed by preliminary ICP-MS analysiswhich revealed the presence of yttrium ions in the retrieved solution, indicating that thedissolution of stabilizing yttrium ions into the test solution took place. Furthermore, severe grainfall-out observed in the degraded specimen also may have contributed to the similar yttriumconcentration found in the as-sintered and degraded matrix as grain fall out would removeyttrium depleted grains and expose new surfaces that would react with water. However, despitethese observations, further analysis is required to understand the exact mechanism(s) activebehind this degradation behaviour.1125.2.2 Degradation Behaviour of CeO2-YSZ CeramicsThe following sections review the degradation behaviour of YSZ ceramics containing CeO2additions.5.2.2.1 Effect of CeO2 on Degradation RatesDuring degradation testing, it was found that the addition of CeO2 to YSZ reduced thedegradation rates. The CeO2-YSZ ceramics had degradation rates ranging from 1.58 to84.48 mm/y (while the pure YSZ ceramic were generally found to completely disintegrate). Thisimproved degradation behaviour may be attributed to the small grain sizes observed in the as-sintered CeO2-YSZ ceramics as a refinement of microstructure has been reported to improve thedegradation resistance of stabilized zirconia materials [11],[12]. However, it has been proposedthat degradation, or the reversion of the stabilized matrix to the monoclinic phase, is due to theannihilation of oxygen vacancies though the diffusion OH- ions into the matrix [48]. Literaturestudies suggest that the migration of the OH- ions possibly occurs through a grain boundarydiffusion mechanism [50]. Therefore, as a reduction in grain size increases the grain boundaryarea, there may be opposing factors contributing to the degradation behaviour of the ceramic.The large variation in the measured degradation rates was attributed to a spallationbehaviour initiated by the formation of large cracks. These cracks formed from the Ce-rich bandspresent in the as-sintered matrix. The formation of these cracks was correlated to the differencein thermal expansion between the zirconia matrix and CeO2. Studies in the literature [95], [96]indicate that at room temperature the thermal expansion coefficient for pure CeO2 and 8 mol%yttria stabilized zirconia is 9.33x10-6 K-1 and 7.6x10-6 K-1, respectively, corresponding to a 19%difference. With increasing temperature, the difference between the constituents further increasesto 21% at 400 °C. Therefore, large Ce-rich bands likely promoted the nucleation of cracks duringheating and cooling resulting in spallation and variable degradation rates. Comparison of themaximum measured degradation rates for the various CeO2-YSZ ceramics suggested that as thenumber and size of the CeO2 bands increased (with increasing amount of CeO2 doping) theobserved degradation rates also increased. The maximum degradation rates for the 5, 10 and 15mol% CeO2-YSZ ceramic were 11.62, 51.06 and 84.48 mm/y, respectively. Only one samplecontaining 5 mol% CeO2 did not exhibit any surface spallation where the measured degradationrate was 1.77 mm/y.113Despite the improved stability of YSZ ceramics with the addition of CeO2, the observeddegradation behaviour of these novel ceramics remains problematic for the application of thesematerials in SCWRs. As discussed in Chapter 1, due to the high temperatures and pressures ofthe supercritical water coolant, an insulating material must be present along the inner walls of thepressure tube in the fuel channel component. Considering the requirement for a 30 year lifespan,the insulating material must provide sufficient thermal insulation, while remaining structurallysound [4]. Zirconia ceramics have been selected as a candidate material due to their desirablethermal properties along with their high stability in aggressive and irradiated environments.However, as demonstrated by the observed degradation behaviour of the pure YSZ ceramics,these ceramics undergo a rapid loss of material when exposed to supercritical water. Theaddition of CeO2 to YSZ significantly reduces this detrimental behaviour, however; even thelowest measured degradation rates would require 53 mm of sacrificial material in addition to theinsulator thickness in order to meet the 30 year lifespan requirement. This indicates that, thoughthis work provides some insight on methods to control degradation, further research on materialoptimization is necessary.5.2.2.2 Effect of CeO2 on Phase StabilityXRD analysis of the CeO2-YSZ specimens before and after degradation testing revealed that theas-sintered cubic phase of the matrix was retained during supercritical water exposure. However,despite retaining this stabilized phase, the measured lattice parameters were found to decreaseafter testing. This indicated that during supercritical water exposure, Ce4+ or Y3+ diffused out ofthe lattice. The average percent decrease in lattice parameter after degradation testing wascalculated to be 0.1, 0.2 and 0.3% for the 5, 10 and 15 mol% CeO2-YSZ ceramics, respectively.ICP-MS analysis revealed that the presence of both Ce4+ or Y3+ ions in the solution retrievedafter testing, further supporting the dissolution of the stabilizing additives into the surroundingenvironment. However, the presence of Ce4+ ions may be also due to the preferential dissolutionof Ce-rich agglomerates and bands.Analysis of the CeO2-YSZ ceramics also revealed the formation of a secondary phase inthe degraded specimens. SEM-XEDS scans of the matrix before and after testing revealed thatthe Ce content on the surface of the 5 mol% CeO2-YSZ increased significantly (~15 wt% Ce).This was attributed to the formation of a Ce-rich layer on the surface of the tested sample. This114layer also possibly formed on the 10 and 15 mol% CeO2-YSZ samples; however, due to thehigher Ce content in these ceramics the mass of this layer was enhanced which caused it to spall.Thus, only localized remnants of this layer were observed on the 10 and 15 mol% CeO2-YSZspecimens. In addition to Ce-rich layers, SEM-XEDS analysis also revealed the presence of largeCe-rich areas on the surface of all the degraded CeO2-YSZ specimens. These Ce-rich areasexhibited a nodular texture and generally had a composition close to the theoretical compositionof CeO2 (81% Ce and 19% O). The development of these Ce-rich phases was found to onlyoccur on the surface of the ceramic. Hernandez et al. [41] suggested that the improveddegradation resistance of CeO2 containing yttria stabilized ceramics was due to the formation ofa thin CeO2 film on the surface of the sintered grains which prevented water attack at yttriumactive points. These findings are in agreement with the observed formation of Ce-rich layers andCe-rich areas on the surface of the degraded CeO2-YSZ specimens. However, the improvedstability of the CeO2 containing ceramics may also be related to the removal of oxygen vacanciesby substituting stabilizing Y3+ ions with Ce4+ ions, as was discussed in Section 2.7.3.1.5.2.3 Degradation Behaviour of Nd2O3-YSZ CeramicsThe following sections review the degradation behaviour of YSZ ceramics containing Nd2O3additions.5.2.3.1 Effect of Nd2O3 on Degradation RatesSimilar to the CeO2-YSZ ceramics, Nd2O3 additions reduced the degradation rate of the pureYSZ ceramic. The measured degradation rates for the 5 mol% Nd2O3-YSZ ceramic remainedrelatively constant with an average degradation rate of 1.5 mm/year. In contrast, the degradationrates of the 10 mol% Nd2O3-YSZ ceramic ranged from 2.76 to 17.80 mm/year. It was found thatthe 10 mol% Nd2O3-YSZ ceramic exhibited a cracking and spallation behaviour similar to thatobserved in the CeO2-YSZ ceramics.Microstructure analysis revealed that band formation was present in the sintered 10 mol%Nd2O3-YSZ ceramics where two types of bands were identified: (1) smaller white bands whichwere generally surrounded by (2) larger grey bands. Both bands were found to have a highcontent of Nd when compared to the matrix. After degradation testing, grey bands disappearedand left large porosity bands. However, the smaller white Nd bands were present in the degradedspecimens and often fractured by propagating cracks. Closer examination revealed the formation115of small microcracks propagating away from the white bands into the matrix. These cracks werefound to propagate through areas of high porosity and often merge to span the entire specimen,as shown in Figure 5.3. SEM-XEDS scans revealed the composition of the matrix before andafter degradation testing remained similar and therefore the propagation of these cracks was notdiffusion based. Rather, crack propagation likely occurred due to the difference in thermalexpansion between the sintered cubic zirconia matrix and the Nd-rich white bands. At roomtemperature, Nd2O3 has a hexagonal structure with a thermal expansion coefficient of 6.9x10-6 K-1 and 14.7x10-6 K-1 along the a- and c-axis, respectively [97]. Thus, at room temperature, thethermal expansion along the c-axis of Nd2O3 is nearly double the thermal expansion of 8 mol%yttria stabilized zirconia. In addition, the thermal expansion coefficients of Nd2O3 increases withtemperature where, at a temperature to ~424 °C, the coefficient for the c-axis increases by 7%,while the coefficient for the a-axis increases by 33% [97]. Therefore during heating and coolingthe difference in thermal expansion behaviour between Nd-rich bands and the cubic zirconiamatrix likely promoted the nucleation of cracks. This observation was further supported by theabsence of cracking and spallation in the 5 mol% Nd2O3-YSZ ceramic which was attributed tothe homogenous microstructure and the absence of Nd-rich bands.Figure 5.3 Crack propagation in the 10 mol% Nd2O3-YSZ ceramic after degradation testing: a) generaloverview and b) formation of microcracks from Nd-rich white bands.1165.2.3.2 Effect of Nd2O3 on Phase StabilityXRD analysis of the Nd2O3-YSZ samples before and after degradation testing revealed that thecubic phase of the matrix was retained during testing. However, contrary to the CeO2-YSZsamples, the unit cell lengths of the Nd2O3-YSZ ceramics were found to remain constant beforeand after degradation testing. The lattice remained within ±1% of the original as-sintered latticeparameter. This suggested that Nd3+ or Y3+ diffusion out of the lattice was not active duringsupercritical water exposure. However, ICP-MS analysis revealed the presence of both Nd3+ andY3+ ions in the retrieved test solutions. The presence of the Nd3+ ions may due to the preferentialdissolution of the Nd-rich phases in the as-sintered ceramic, since the tested specimens oftenexhibited porosity networks in areas where grey Nd-rich bands once existed. The presence of Y3+may be attributed to cross contamination as the designed test vessels were reused numeroustimes.Studies performed by Salehi et al. [9] also reported improved degradation resistance ofyttria stabilized zirconia ceramics with Nd2O3 additives. Salehi et al. [9] fabricated two differentclasses of stabilized zirconia ceramics: (1) zirconia ceramics stabilized with either 2 or 3 mol%of yttria, or (2) 1 mol% yttria stabilized zirconia ceramics with either 1 or 2 mol% Nd2O3.  Rawprecursor powders were SPS at a temperature of 1450 °C for a dwell time of 3 min and at apressure of 62 MPa. Applied heating rate was 400 °C/min until a temperature of 1050 °C wasreached, with a heating rate of 200 °C/min thereafter until the desired sintering temperature of1450 °C was reached. The resulting sintered ceramics were then tested in a stainless steelautoclave at 200 °C under a saturated H2O pressure of 1.55 MPa.  The evolution of themonoclinic phase on the surface of the specimen was measured as a function of time. It wasfound that the zirconia ceramics stabilized with only yttria completely degraded after 20 min oftesting, while zirconia ceramic containing both yttria and Nd2O3 additions remained intact.However, despite this improved stability, the measured surface monoclinic content was found torapidly increase in the yttria stabilized zirconia ceramics with either 1 or 2 mol% Nd2O3. Thissuggested that though Nd2O3 additions reduced the in-depth progression of degradation, yttriastabilized zirconia had a stronger initial degradation resistance. Salehi et al.[9] proposed that theimproved degradation resistance in yttria stabilized zirconia ceramics containing Nd2O3 additionswas due the increase in unit cell dimensions of the stabilized tetragonal phase, attributed toincorporation of larger Nd3+ ions in the lattice (where the ionic radius for the Nd3+ and Y3+ ions117is 1.109 Å and 1.019 Å, respectively).  The increased unit cell dimension lowered the tetragonalto monoclinic transformation temperature, because the internal strain applied to the lattice due tooxygen overcrowding was reduced.  Salehi et al. [9] further suggested that during oxygenannihilation, the larger lattice parameter accommodated the residual tensile stresses which areapplied to the lattice during the diffusion of OH- ions.118Chapter 6: ConclusionThis chapter reviews the key findings obtained from this thesis.  Conclusions are categorized intotwo sections: (1) fabrication and (2) degradation behaviour of novel zirconia ceramics.6.1 Fabrication of Novel Zirconia CeramicsIn this research novel YSZ ceramics were fabricated with various rare earth additives throughSpark Plasma Sintering. From the results and discussion outlined in Chapter 4 and Chapter 5 thefollowing conclusions can be summarized: Precursor powders must be properly blended to avoid agglomerations and the formationof large bands in the sintered matrix Proper selection of SPS parameters (temperature, dwell time and pressure) must be takento ensure sufficient interdiffusion of rare earth additives into the matrix while controllinggrain growth and densification During SPS carbon diffused from the graphite tooling into the as-sintered ceramics CeO2 and Nd2O3 additions were effectively incorporated into the YSZ lattice throughSPS processing CeO2 and Nd2O3 additions hindered densification and grain growth6.2 Degradation Behaviour of Novel Zirconia CeramicsDegradation testing of the fabricated zirconia ceramics revealed the following conclusions: Degradation rates of the pure YSZ ceramic were significantly reduced with CeO2 andNd2O3 additions, however, despite the improved stability of these rare earth doped YSZceramics the observed degradation rates remained unacceptable for the application ofthese materials in SCWRs CeO2 and Nd2O3 agglomerates and bands provided preferential dissolution points in thematrix resulting in the formation of a porosity network and elongated voids duringexposure to supercritical water CeO2 and Nd2O3 bands promoted cracking and spallation in the ceramics due to thedifference in thermal expansion between the band and the sintered YSZ matrix CeO2-YSZ ceramics exhibited a decrease in lattice parameter suggesting the diffusion ofeither Ce4+ or Y3+ ions out of the lattice occurred at the conditions investigated (31 MPaand 400 °C)119 No change in lattice parameter was measured for Nd2O3-YSZ ceramics after supercriticalwater exposure, indicating Nd3+ or Y3+ diffusion out of the lattice was not active120Chapter 7: Future WorkThis chapter provides suggestions for future work focusing on procedures and methods toimprove the degradation resistance of cubic zirconia ceramics. The recommended future work iscategorized into three general areas: (1) fabrication and (2) degradation behaviour of novelzirconia ceramics as well as (3) future studies required for SCWR applications of thesematerials.7.1 Spark Plasma Sintering of Novel Zirconia CeramicsRecommended future work in processing and fabricating novel yttria stabilized zirconia ceramicscontaining CeO2 and Nd2O3 additions are as follows: Apply techniques to effectively mix precursor powders to eliminate the formation oflarge agglomerations Study the effect of Spark Plasma Sintering processing parameters (pressure, temperature,dwell time) on the densification and resulting grain size of the sintered ceramics Optimize SPS parameters to allow for complete interdiffusion of rare earth additives intothe matrix Perform further in-depth studies on the role of rare earths additives on the grain growthand densification behaviour of zirconia ceramics during SPS7.2 Degradation of Zirconia Ceramics in Supercritical WaterThe following recommendations suggest work to advance the control of the degradationbehaviour of cubic zirconia ceramics: Perform degradation testing on sintered zirconia ceramics that exhibit a homogenousmatrix without large bands or agglomerates Study the effect of porosity and grain size on degradation rates Perform in-depth studies on the phase and microstructure evolution during degradationtesting and the role of rare earth additives on preventing phase transformations Record real-time degradation rates during sample heating and cooling to identify theamount of degradation that occurs during supercritical water exposure1217.3 Future Studies for SCWR ApplicationsSuggested future research areas aimed at further understanding the performance of zirconiaceramics in a SCWR environment: Study the effect of temperature, pressure and water composition on the properties ofsupercritical water and the resulting effect on degradation rates of zirconia ceramics Perform degradation testing in a dynamic environment Review the effect of radiation exposure on microstructure and degradation resistance ofcubic zirconia ceramics.122References[1] Nuclear Energy Research Advisory Committee and the Generation IV InternationalForum, “A technology roadmap for generation IV nuclear energy systems,”Department of Energy - Office of Scientific and Technical Information, 2002.[2] International Energy Agency and Nuclear Energy Agency (2010, June 16).Technology roadmap nuclear energy [online]. 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