J.Am. Ceran.So,9同1901-1920(2009) DOI:10.11111551-2916.200903278.x urna The Tetragonal-Monoclinic Transformation in Zirconia Lessons learned and future trends Jerome Chevalier and Laurent gremillard' University of Lyon. INSA-Lyon, MATEIS, Villeurbanne FR-69621, France Anil v.Ⅴirka Department of Material Science Engineering. University of Utah, Salt Lake City, Utah 84112 David R. Clarke School of Engineering and Applied Sciences, Harvard University, Cambridge, Massachusetts 02138 Zirconia ceramics have found broad applications in a variety L. Introduction energy and biomedical applications because of their unusua combination of strength, fracture toughness, ionic conductivity ZiRals fo hwsll eterne oe the m ost ime disct era moir trans. and low thermal conductivity. These attractive characteristics formation toughening in 1975 heralded new visions for new are largely associated with the stabilization of the tetragonal and cubic phases through alloying with aliovalent ions. The large high-performance applications of zirconia, ranging from bearing and wear applications to thermal barrier coatings (TBCs)to the aliovalent alloying is responsible for both the exceptionally most recently, biomedical applications. The subsequent discov high ionic conductivity and the unusually low, and temperature ery that zirconia could also be toughened by ferroelastic switch- independent, thermal conductivity. The high fracture toughnes g gave further confidence in the application of zirconia exhibited by many of zirconia ceramics is attributed to the con- ceramics in critical applications. Nevertheless, despite the suc- straint of the tetragonal-to-monoclinic phase transformation cess of zirconia in many new applications, it has become appar- and its release during crack propagation. In other zirconia ce- ent that certain zirconia compositions can also have an achilles the tett phase, the high fracture tough heel, namely their propensity to low-temperature degradation ness is associated with ferroelastic domain switching. However (LTD)in the presence of moisture. This is a kinetic phenomenon many of these attractive features of zirconia, especially fracture in which polycrystalline tetragona toughness and strength, are compromised after prolonged expo monoclinic zirconia over a rather narrow but important tem In p iess referm t intermediate temperatures(- c) perature range, typically room temperature to around 400oC, sure to water v to as low-temperature degradation (LTD), depending on the stabilizer, its concentration, and the grain size and initially identified over two decades ago. This is particularly of the ceramic. The transformation occurs by a nucleation and growth process and typically begins at the surfaces of polycrys for zirconia in biomedical applications, such as hip implants talline ceramics. It has all the characteristics of being an iso- and dental restorations. Less well substantiated is the possibility that the same process can also occur in zirconia used in other thermal martensite. Also. although there continues to remain applications, for instance, zirconia thermal barrier coatings af- some uncertainty as to the precise mechanism by which moisture ter long exposure at high temperature. Based on experience with causes destabilization of the tetragonal phase, the observation the failure of zirconia femoral heads as well as studies of LtD. diffusion suggests that the transformation occurs by the hcy that the kinetics of Ltd are similar to those of oxygen vaca it is shown that many of the problems of ltd can be mitigate by the appropriate choice of alloying and or process control diffusion of a moisture species with an activation energy similar to that of oxygen vacancy diffusion. In practical terms, LTD is in effect, an alternative to crack propagation, stress-induced transformation for the transformation from metastable tetra onal to monoclinic(I-m)zirconia(see Fig. 1). In this feature article we describe the role of phase trans- rmations responsible for the impressive combination of me- chanical properties of zirconia, their relationship to equilibriu and metastable phase diagrams, and the phenomenon of Ltd ript No 26208 Received April 27, 2009, ap uthor to whom correspondence she We include the effects of transformations at free surfaces be- cause these affect the surface finish that is important for many Feature
The Tetragonal-Monoclinic Transformation in Zirconia: Lessons Learned and Future Trends Je´roˆme Chevalier and Laurent Gremillardw University of Lyon, INSA-Lyon, MATEIS, Villeurbanne FR-69621, France Anil V. Virkar Department of Material Science & Engineering, University of Utah, Salt Lake City, Utah 84112 David R. Clarke School of Engineering and Applied Sciences, Harvard University, Cambridge, Massachusetts 02138 Zirconia ceramics have found broad applications in a variety of energy and biomedical applications because of their unusual combination of strength, fracture toughness, ionic conductivity, and low thermal conductivity. These attractive characteristics are largely associated with the stabilization of the tetragonal and cubic phases through alloying with aliovalent ions. The large concentration of vacancies introduced to charge compensate of the aliovalent alloying is responsible for both the exceptionally high ionic conductivity and the unusually low, and temperature independent, thermal conductivity. The high fracture toughness exhibited by many of zirconia ceramics is attributed to the constraint of the tetragonal-to-monoclinic phase transformation and its release during crack propagation. In other zirconia ceramics containing the tetragonal phase, the high fracture toughness is associated with ferroelastic domain switching. However, many of these attractive features of zirconia, especially fracture toughness and strength, are compromised after prolonged exposure to water vapor at intermediate temperatures (B301–3001C) in a process referred to as low-temperature degradation (LTD), and initially identified over two decades ago. This is particularly so for zirconia in biomedical applications, such as hip implants and dental restorations. Less well substantiated is the possibility that the same process can also occur in zirconia used in other applications, for instance, zirconia thermal barrier coatings after long exposure at high temperature. Based on experience with the failure of zirconia femoral heads, as well as studies of LTD, it is shown that many of the problems of LTD can be mitigated by the appropriate choice of alloying and/or process control. I. Introduction ZIRCONIA has been one of the most important ceramic materials for well over a century but the discovery of transformation toughening in 19751 heralded new visions for new high-performance applications of zirconia, ranging from bearing and wear applications to thermal barrier coatings (TBCs) to, most recently, biomedical applications. The subsequent discovery that zirconia could also be toughened by ferroelastic switching2 gave further confidence in the application of zirconia ceramics in critical applications. Nevertheless, despite the success of zirconia in many new applications, it has become apparent that certain zirconia compositions can also have an Achilles heel, namely their propensity to low-temperature degradation (LTD) in the presence of moisture. This is a kinetic phenomenon in which polycrystalline tetragonal material slowly transforms to monoclinic zirconia over a rather narrow but important temperature range, typically room temperature to around 4001C, depending on the stabilizer, its concentration, and the grain size of the ceramic. The transformation occurs by a nucleation and growth process and typically begins at the surfaces of polycrystalline ceramics. It has all the characteristics of being an isothermal martensite. Also, although there continues to remain some uncertainty as to the precise mechanism by which moisture causes destabilization of the tetragonal phase, the observation that the kinetics of LTD are similar to those of oxygen vacancy diffusion suggests that the transformation occurs by the indiffusion of a moisture species with an activation energy similar to that of oxygen vacancy diffusion. In practical terms, LTD is, in effect, an alternative to crack propagation, stress-induced transformation for the transformation from metastable tetragonal to monoclinic (t–m) zirconia (see Fig. 1). In this feature article we describe the role of phase transformations responsible for the impressive combination of mechanical properties of zirconia, their relationship to equilibrium and metastable phase diagrams, and the phenomenon of LTD. We include the effects of transformations at free surfaces because these affect the surface finish that is important for many Feature D. J. Green—contributing editor w Author to whom correspondence should be addressed. e-mail: laurent.gremillard@ insa-lyon.fr Manuscript No. 26208. Received April 27, 2009; approved 1 July 2009. Journal J. Am. Ceram. Soc., 92 [9] 1901–1920 (2009) DOI: 10.1111/j.1551-2916.2009.03278.x r 2009 The American Ceramic Society
1902 Journal of the American Ceramic Society--Chevalier et al Vol. 92. No 9 zirconia polycrystals. In FSZ zirconia is in its cubic form, the Tetragonal zirconia ceramics form most commonly used in oxygen sensors and fuel cell elec- olytes. It is generally obtained with large concentration of sta bilizers (ie, more than 8 mol% Y2O3). The PSZ consists of anosized tetragonal or monoclinic particles that have precipi tated out in a cubic matrix. Such zirconia ceramics are generally Moist Atmosphere obtained with the addition of lime or magnesia. TZPs are often onsidered as monoliths of tetragonal phase, although they may contain a secondary cubic phase( see Panel A). The majority of TZPs that have been investigated are those stabilized with yttria or ceria ow Temperature Toughening Il. Stabilization and Transformation of the Tetragonal phase Fig 1. Crack propagation-induced transformation and intermediate temperature exposure to moisture are two alternative means by which As mentioned above, stabilization of powders and sintered metastable tetragonal phase can transform to monoclinic phase. ceramics can be achieved by alloying pure zirconia with other oxides. Investigations of the stability of different phases wit applic atages as wvo a thr bine s mals desease app ressons phase diagrams such as those compiled by the American Ce- learnt from investigation of LTD in femoral implants and the ramic Society-NIST project Is echanisms that control it. Much of this article is concerned Alloy stabilization not only enables fabrication of crack-free with the properties of the zirconia-yttria material system be- zrconia but as demonstrated by Gupta et al., sintered bodies cause the majority of the research and development that has of polycrystalline tetragonal zirconia can be prepared and the been performed on zirconia in the last three decades has been tetragonal phase retained down to room temperature even n yttria-stabilized zirconia (YSZ). YSZ ceramics have the best ough the equilibrium phase is monoclinic. These metastable combination of toughness and strength of any of the stabi- trigonal ceramics exhibited tional fracture toughness when the transformation to the monoclinic phase was triggered lized zirconias. Also, and this cannot be overemphasized, it is by a propagating crack. The toughening that can be achieved for undoubtedly due to the early and continued availability of high purity, uniform submicrometer particle size powders from com- different concentrations of yttria is illustrated in Fig. 2, which panies such as Tosoh in Japan summarizes data for transformation-toughened zirconia and ferroelastic toughening. Also shown, superimposed, is the sen- sitivity to aging as a function of yttria concentration. The frac- II. Experimental Procedure ture toughness of monoclinic and cubic zirconia, which is similar o that of window glass, provides a refe The principal properties of zirconia are well known and various toughening through the I-m phase transformation can be co aspects have been reviewed in detail many times since the dis- pared. It is emphasized that the data are obtained from standard covery of transformation toughening by garvie et al. in calcia acture toughness tests such as indentation and double canti stabilized zirconia. For this reason in this section we summarize ever beam tests. in fast fracture conditions under which ltd is and ea principal features of the stabilization of zirconia, the avoided. As indicated in Fig. 2, the attainable fracture toughness raphy of stabilized zirconia, and the relationship be- decreases as the yttria concentration increases. In the context of rmations the metastable phase diagram, the toughening is proportional to The panels describe the essentials of phase equilibria the magnitude of the undercooling below the To(t/m) temper ature(see Panel A for a more detailed explanation). Further Pure zirconium oxide exhibits three allotropes: monoclin more. transformation toughening is restricted to moderate m), which is the stable phase up to 1170.C, where it transforms nperatures, becoming ineffective when the stable phase is to tetragonal (0), and then cubic (c) at temperatures above tragonal and not monoclinic. for these reasons the most 2370C. A comprehensive review of the different structures for tractive compositions for transformation toughening are those zirconia can be found in green et al. The t-m transformation, with low yttria concentrations(but high enough to prevent which is martensitic, has been the subject of the most careful spontaneous f-m transformation during cooling), typically 2- attention, because it usually occurs during the sintering and on 3 mol%Y2O3 both heating and cooling. The t-m transformation is accompa In the ce of mo he transformation of metastable nied by a large shear strain and a large volume increase(see -m can alternatively occur without the passage of a crack In Panel B). Together these create large internal stresses on cooling. this sense, LTD is a competing process to transformation tough large, in fact, that pure zirconia sintered above 1170C inev- ning with the two providing limiting behaviors by which the can either sinter at low temperature for it to remain monoclinic a propagating crack, then one can get enhanced toughening(see during sintering-which leads to a low-strength and toughness "Section Ill(2)"). On the other hand, the transformation ma ceramic--or stabilize the tetragonal or the cubic phases at room be triggered"chemically "by the infusion of water-derived spe temperature by alloying, thereby avoiding the f-m transforma- cies from the surface. The process on a surface is complex tion during cooling. The fundamental approach to the engineer (Fig 3)and results not only in the undesirable transformation ng use of zirconia and avoiding the transformation-induced but also surface roughening, microcracking, and grain pull-out cracking described by Ruff and Ebert" almost a century ago re- Is well as loss of strength-all processes detrimental to struc- mains valid today alloying pure zirconia with another oxide to tural applications. The dilemma facing the alloy designer is that ully or partially stabilize the tetragonal and or the cubic pha the Ysz compositions that are the most attractive for their Calcium, magnesium, yttrium, and cerium oxides have been the fracture toughness are also those that are most susceptible to most widely used stabilizers and lead to a number of different LTD. This is illustrated by the comparison of the fracture microstructures In general, zirconia ceramics may conveniently toughness data with the ltd data in Fig. 2. be classified into three major types according to their micro- The stabilization of the tetragonal phase in polycrystalline structure: FSZ, PSZ, and TZP, standing, respectively, for fully ceramics is, btedly, largely dependent on the mutual elas- stabilized zirconia, partially stabilized zirconia, and tetragonal vided by the surrounding, untransformed
applications as well as the kinetics. We also describe approaches being taken to avoid LTD, or minimize it, based on lessons learnt from investigation of LTD in femoral implants and the mechanisms that control it. Much of this article is concerned with the properties of the zirconia–yttria material system because the majority of the research and development that has been performed on zirconia in the last three decades has been on yttria-stabilized zirconia (YSZ). YSZ ceramics have the best combination of toughness and strength of any of the stabilized zirconias. Also, and this cannot be overemphasized, it is undoubtedly due to the early and continued availability of highpurity, uniform submicrometer particle size powders from companies such as Tosoh in Japan. II. Experimental Procedure The principal properties of zirconia are well known and various aspects have been reviewed in detail many times since the discovery of transformation toughening by Garvie et al. 1 in calciastabilized zirconia. For this reason, in this section we summarize only the principal features of the stabilization of zirconia, the crystallography of stabilized zirconia, and the relationship between mechanical properties and the phase transformations in zirconia. The panels describe the essentials of phase equilibria and the transformation crystallography. Pure zirconium oxide exhibits three allotropes: monoclinic (m), which is the stable phase up to 11701C, where it transforms to tetragonal (t), and then cubic (c) at temperatures above 23701C. A comprehensive review of the different structures for zirconia can be found in Green et al. 3 The tm transformation, which is martensitic, has been the subject of the most careful attention, because it usually occurs during the sintering and on both heating and cooling. The tm transformation is accompanied by a large shear strain and a large volume increase (see Panel B). Together these create large internal stresses on cooling. So large, in fact, that pure zirconia sintered above 11701C inevitably disintegrates by cracking upon cooling. To maintain the integrity of sintered zirconia bodies at room temperature, one can either sinter at low temperature for it to remain monoclinic during sintering—which leads to a low-strength and toughness ceramic—or stabilize the tetragonal or the cubic phases at room temperature by alloying, thereby avoiding the tm transformation during cooling. The fundamental approach to the engineering use of zirconia and avoiding the transformation-induced cracking described by Ruff and Ebert4 almost a century ago remains valid today: alloying pure zirconia with another oxide to fully or partially stabilize the tetragonal and/or the cubic phase. Calcium, magnesium, yttrium, and cerium oxides have been the most widely used stabilizers and lead to a number of different microstructures. In general, zirconia ceramics may conveniently be classified into three major types according to their microstructure: FSZ, PSZ, and TZP, standing, respectively, for fully stabilized zirconia, partially stabilized zirconia, and tetragonal zirconia polycrystals. In FSZ zirconia is in its cubic form, the form most commonly used in oxygen sensors and fuel cell electrolytes. It is generally obtained with large concentration of stabilizers (i.e., more than 8 mol% Y2O3). The PSZ consists of nanosized tetragonal or monoclinic particles that have precipitated out in a cubic matrix. Such zirconia ceramics are generally obtained with the addition of lime or magnesia. TZPs are often considered as monoliths of tetragonal phase, although they may contain a secondary cubic phase (see Panel A). The majority of TZPs that have been investigated are those stabilized with yttria or ceria. III. Stabilization and Transformation of the Tetragonal Phase As mentioned above, stabilization of powders and sintered ceramics can be achieved by alloying pure zirconia with other oxides. Investigations of the stability of different phases with different stabilizers led to the development of the equilibrium phase diagrams such as those compiled by the American Ceramic Society—NIST project.15 Alloy stabilization not only enables fabrication of crack-free zirconia but as demonstrated by Gupta et al.,16 sintered bodies of polycrystalline tetragonal zirconia can be prepared and the tetragonal phase retained down to room temperature even though the equilibrium phase is monoclinic. These metastable tetragonal ceramics exhibited exceptional fracture toughness when the transformation to the monoclinic phase was triggered by a propagating crack. The toughening that can be achieved for different concentrations of yttria is illustrated in Fig. 2, which summarizes data for transformation-toughened zirconia and ferroelastic toughening. Also shown, superimposed, is the sensitivity to aging as a function of yttria concentration. The fracture toughness of monoclinic and cubic zirconia, which is similar to that of window glass, provides a reference against which the toughening through the tm phase transformation can be compared. It is emphasized that the data are obtained from standard fracture toughness tests, such as indentation and double cantilever beam tests, in fast fracture conditions under which LTD is avoided. As indicated in Fig. 2, the attainable fracture toughness decreases as the yttria concentration increases. In the context of the metastable phase diagram, the toughening is proportional to the magnitude of the undercooling below the T0 (t/m) temperature (see Panel A for a more detailed explanation). Furthermore, transformation toughening is restricted to moderate temperatures, becoming ineffective when the stable phase is tetragonal and not monoclinic. For these reasons, the most attractive compositions for transformation toughening are those with low yttria concentrations (but high enough to prevent spontaneous tm transformation during cooling), typically 2– 3 mol% Y2O3. In the presence of moisture, the transformation of metastable t–m can alternatively occur without the passage of a crack. In this sense, LTD is a competing process to transformation toughening with the two providing limiting behaviors by which the metastable tetragonal phase transforms to the more stable monoclinic phase (Fig. 1). If the transformation is triggered by a propagating crack, then one can get enhanced toughening (see ‘‘Section III(2)’’). On the other hand, the transformation may be triggered ‘‘chemically’’ by the infusion of water-derived species from the surface. The process on a surface is complex (Fig. 3) and results not only in the undesirable transformation but also surface roughening, microcracking, and grain pull-out as well as loss of strength—all processes detrimental to structural applications. The dilemma facing the alloy designer is that the YSZ compositions that are the most attractive for their fracture toughness are also those that are most susceptible to LTD. This is illustrated by the comparison of the fracture toughness data with the LTD data in Fig. 2. The stabilization of the tetragonal phase in polycrystalline ceramics is, undoubtedly, largely dependent on the mutual elastic constraint provided by the surrounding, untransformed Fig. 1. Crack propagation-induced transformation and intermediate temperature exposure to moisture are two alternative means by which metastable tetragonal phase can transform to monoclinic phase. 1902 Journal of the American Ceramic Society—Chevalier et al. Vol. 92, No. 9
September 2009 The Tetragonal-Monoclinic Transformation in Zirconia 1903 Panel a. Zirconia-Yttria phase diagram The zirconia-yttria phase diagram has been significantly refined many times since it was first introduced in 1951 by duwez et al (The phase diagram book devoted to zirconia includes 30 different variations in the phase diagram. ) At first, the succession revisions might be a surprise but the essential difficulty is that cation diffusion in zirconia is so slow that it has proved Furthermore, the slow diffusion kinetics means that metastable extensions of the phases can readily occur. Interestingly, it was the prospect of diffusion-limited processes in zirconia that led Pol Duwez to investigate the system in his pioneering studies of rapid solidification and glass formation. In addition, the characteristic features of the martensitic transformation, such as the start and finish temperatures, have further complicated the interpretation of the diagram and the interpretation of microstructures. This uncertainty is shown in Scott's phase diagram by the hatched region. As a result, there has been considerable confusion in the literature about the details of the phase diagram. Yashima et al. present a graphic illustration of particularly pronounced for the region pertinent to transformation toughening and the low-temperature degradation are he confusion by superimposing many of the published diagrams on one another in their Fig. Al. The disagreements are 2000 2000 T 1500 02 YO Mol fraction Y, O Mole Fraction Y,O, Mole Fraction 0026 0053 111 0026 0053 0081 0.111 2400 2400 2100 C 51500 T+C g T+C CICm = 900 600M C+M T(T/M) 25 25 0 YO,s Mole Fraction YO O. Mole Fraction Oxygen Site Fraction Vacancies Oxygen Site Fraction Vacancies Fig. Al. Evolution in our knowledge of the zirconia-yttria phase diagram: (a) original diagram by Duwez in 1951,(b)diagram presented by Scott in 1975.(c)and (d) most recent diagrams.(d)is the metastable phase diagram. The evolution in our knowledge of the phase diagram can be illustrated by the diagrams in Fig. Al, showing the original diagram by Duwez, the diagram presented by Scott in 1975 and the most recent diagrams that combines experimental and putational studies. At the time of writing, there are indications that even this version may not be quite correct and that the tetragonal boundary may exhibit retrograde curvature, as occurs in the ZrOz-CeO system. Considera ble clarification has been obtained from computer determination of the phase diagram, particularly the position of the metastable To lines. However, it has to be emphasized that the metastable lines themselves, as well as the phase boundary lines are obtained from optimization
Panel A. Zirconia-Yttria Phase Diagram5,6 The zirconia–yttria phase diagram has been significantly refined many times since it was first introduced in 1951 by Duwez et al. 7 (The phase diagram book devoted to zirconia includes 30 different variations in the phase diagram.) At first, the succession of revisions might be a surprise but the essential difficulty is that cation diffusion in zirconia is so slow8 that it has proved particularly difficult to establish equilibrium and hence the phase boundaries at temperatures below about 14001C. Furthermore, the slow diffusion kinetics means that metastable extensions of the phases can readily occur. Interestingly, it was the prospect of diffusion-limited processes in zirconia that led Pol Duwez to investigate the system in his pioneering studies of rapid solidification and glass formation. In addition, the characteristic features of the martensitic transformation, such as the start and finish temperatures, have further complicated the interpretation of the diagram and the interpretation of microstructures. This uncertainty is shown in Scott’s phase diagram by the hatched region.9 As a result, there has been considerable confusion in the literature about the details of the phase diagram. Yashima et al. 5 present a graphic illustration of the confusion by superimposing many of the published diagrams on one another in their Fig. A1. The disagreements are particularly pronounced for the region pertinent to transformation toughening and the low-temperature degradation. The evolution in our knowledge of the phase diagram can be illustrated by the diagrams in Fig. A1, showing the original diagram by Duwez, the diagram presented by Scott in 19758 and the most recent diagrams that combines experimental and computational studies.10,11 At the time of writing, there are indications that even this version may not be quite correct and that the tetragonal boundary may exhibit retrograde curvature, as occurs in the ZrO2–CeO2 system.12 Considerable clarification has been obtained from computer determination of the phase diagram, particularly the position of the metastable T0 lines. However, it has to be emphasized that the metastable lines themselves, as well as the phase boundary lines are obtained from optimization Fig. A1. Evolution in our knowledge of the zirconia–yttria phase diagram: (a) original diagram by Duwez in 1951,6 (b) diagram presented by Scott in 1975,8 (c) and (d) most recent diagrams9,10 (d) is the metastable phase diagram. September 2009 The Tetragonal-Monoclinic Transformation in Zirconia 1903
1904 Journal of the American Ceramic Society--Chevalier et al Vol. 92. No 9 Panel A. Continued procedures that use the experimentally determined equilibrium phase boundaries as input parameters. For instance, the temperature of the monoclinic to tetragonal and the tetragonal to cubic, as well as the t/c boundary line are used to determine the metastable To(/m) boundary so if the t/e boundary line is inaccurate, then the computed To(/m) may not be fully correct The necessity of considering the metastable phase diagram is that cation diffusion in zirconia is exceptionally slow at all but the highest temperature. This is illustrated by the graph in Fig. A2, which shows the estimated time for diffusion to occur to homogenize the y content of a 0. 5-or 3-Hm-diameter grain at different temperatures(calculated from Kilo et al. and ).Even at a temperature of 1500C, a sintering temperature commonly used for zirconia ceramics in the past, it is estimated to take several weeks to achieve compositional homogeneity for the 3-um-grain size material. For a 0.5-um-grain size, this would represent days/hours instead of weeks( time roughly divided by 36 compared with the 3 um case). This explains why phase and yttria partitioning was observed in 3Y-TZP sintered for 5 h at 1550C in previous work 2O, Mole Fraction 0025 005 0075 Cation (Zr, Y 2400 2000 01500+-9 900 600 300 MOnoclinic. Time(s) 25 Estimated time for diffusion to homogenize the y 005 0.15 of 0.5.and 3-Hm-diameter grains at different temperatures in YO, Mole Fraction 3Y-TZP. TZP, tetragonal zirconia polycrystals; Y, yttria. Fig A3. Metastable zirconia-yttria phase diagram. To illustrate the consequences of the very slow cation diffusion, and the crucial importance of the metastable phase diagram in understanding LTD, we take as an example, a 3 mlo Y,,(6.0 m/o YO,s)material sintered at 1500C, composition C in Fig. A3. At equilibrium at this temperature, the sintered sample should consist of two phases, a tetragonal phase of composition 2.4 mo Y,O3(4.5 YO, 5)(point A)and a cubic phase of composition of 7.5 m/o Y,O,(point B). At room temperature. the equilibrium phases, from Fig. Al, would be a monoclinic phase with a yttria concentration of almost zero and a cubic phase with a yttria concentration of 18 m/o Y203. However, for this equilibrium condition to occur the yttrium ions must diffuse to partition into the yttria-poor monoclinic and yttria-rich cubic phases as shown by the horizontal arrows in the figure. As the ndicated by the diffusion distance figure, this would take many years. Instead, under typical cooling conditions, little or no yttrium ion partitioning occurs and the compositions obtained on cooling to room temperature, will be given by the intersection of the vertical dashed lines with the composition axis. At temperatures below the To (t/m). the tetragonal phase is metastable with respect to transformation to the monoclinic but the transformation is kinetically limited For instance, if there has been no diffusion, the To(/m) temperature is given by the intersection point E whereas if diffusional partitioning is complete at the sintering temperature, then the To (t/m)temperature is given by the intersection point D. Consequently, before low-temperature aging, the phases observed will then be a metastable tetragonal phase and a cubic phase, both with the same yttria concentrations as they have at the sintering temperature. Then, as the transformation occurs below the To(t/m) temperature, the nonoclinic phase will have the same yttria concentration as the tetragonal phase, namely 2.4 m/o Y,O3 (4.5 YO,5)in this example. Interestingly, as the yttria stabilizer concentration is increased and no partitioning occurs, the To(t/m) temperature decreases until at about 3.6 mo Y2O3(7.0 m/o YO1.5), it falls to below room temperature. So, unless the material is first transformed to the equilibrium cubic and tetragonal phases during sintering and subsequent heat treatment, it will not be susceptible to transformation until lower temperature is attained A further consequence of the equilibrium t/m phase boundary is that its steep slope implies that the composition of the tetragonal phase formed by diffusional partitioning at high temperatures is not very sensitive to the average composition of the powders, and hence there is little variation in the To (/m) temperature with the yttria content. What does change are the relative volume fractions of the tetragonal and cubic phases at the sintering temperature, and hence the maximum volume fraction of tetragonal that can transform to monoclinic by either transformation toughening or moisture-mediated LTD
Panel A. Continued procedures that use the experimentally determined equilibrium phase boundaries as input parameters. For instance, the temperature of the monoclinic to tetragonal and the tetragonal to cubic, as well as the t/c boundary line are used to determine the metastable T0 (t/m) boundary so if the t/c boundary line is inaccurate, then the computed T0 (t/m) may not be fully correct. The necessity of considering the metastable phase diagram is that cation diffusion in zirconia is exceptionally slow at all but the highest temperature. This is illustrated by the graph in Fig. A2, which shows the estimated time for diffusion to occur to homogenize the Y content of a 0.5- or 3-mm-diameter grain at different temperatures (calculated from Kilo et al. 8 and13). Even at a temperature of 15001C, a sintering temperature commonly used for zirconia ceramics in the past, it is estimated to take several weeks to achieve compositional homogeneity for the 3-mm-grain size material. For a 0.5-mm-grain size, this would represent days/hours instead of weeks (time roughly divided by 36 compared with the 3 mm case). This explains why phase and yttria partitioning was observed in 3Y-TZP sintered for 5 h at 15501C in previous work.14 To illustrate the consequences of the very slow cation diffusion, and the crucial importance of the metastable phase diagram in understanding LTD, we take as an example, a 3 m/o Y2O3 (6.0 m/o YO1.5) material sintered at 15001C, composition C in Fig. A3. At equilibrium at this temperature, the sintered sample should consist of two phases, a tetragonal phase of composition 2.4 m/o Y2O3 (4.5 YO1.5) (point A) and a cubic phase of composition of 7.5 m/o Y2O3 (point B). At room temperature, the equilibrium phases, from Fig. A1, would be a monoclinic phase with a yttria concentration of almost zero and a cubic phase with a yttria concentration of B18 m/o Y2O3. However, for this equilibrium condition to occur the yttrium ions must diffuse to partition into the yttria-poor monoclinic and yttria-rich cubic phases as shown by the horizontal arrows in the figure. As the indicated by the diffusion distance figure, this would take many years. Instead, under typical cooling conditions, little or no yttrium ion partitioning occurs and the compositions obtained on cooling to room temperature, will be given by the intersection of the vertical dashed lines with the composition axis. At temperatures below the T0 (t/m), the tetragonal phase is metastable with respect to transformation to the monoclinic but the transformation is kinetically limited. For instance, if there has been no diffusion, the T0 (t/m) temperature is given by the intersection point E whereas if diffusional partitioning is complete at the sintering temperature, then the T0 (t/m) temperature is given by the intersection point D. Consequently, before low-temperature aging, the phases observed will then be a metastable tetragonal phase and a cubic phase, both with the same yttria concentrations as they have at the sintering temperature. Then, as the transformation occurs below the T0 (t/m) temperature, the monoclinic phase will have the same yttria concentration as the tetragonal phase, namely 2.4 m/o Y2O3 (4.5 YO1.5) in this example. Interestingly, as the yttria stabilizer concentration is increased and no partitioning occurs, the T0 (t/m) temperature decreases until at about 3.6 m/o Y2O3 (7.0 m/o YO1.5), it falls to below room temperature. So, unless the material is first transformed to the equilibrium cubic and tetragonal phases during sintering and subsequent heat treatment, it will not be susceptible to transformation until lower temperature is attained. A further consequence of the equilibrium t/m phase boundary is that its steep slope implies that the composition of the tetragonal phase formed by diffusional partitioning at high temperatures is not very sensitive to the average composition of the powders, and hence there is little variation in the T0 (t/m) temperature with the yttria content. What does change are the relative volume fractions of the tetragonal and cubic phases at the sintering temperature, and hence the maximum volume fraction of tetragonal that can transform to monoclinic by either transformation toughening or moisture-mediated LTD. Fig. A2. Estimated time for diffusion to occur to homogenize the Y content of 0.5- and 3-mm-diameter grains at different temperatures in 3Y-TZP.7,11 TZP, tetragonal zirconia polycrystals; Y, yttria. Fig. A3. Metastable zirconia–yttria phase diagram. 1904 Journal of the American Ceramic Society—Chevalier et al. Vol. 92, No. 9
eptember 2009 The Tetragonal-Monoclinic Transformation in Zirconia 1905 While the stabilization of pure zirconia can be understood in erms of the balance between chemical and surface energy, the IP ceram reason that different aliovalent ions are effective as stabilizers nd also, perhaps, why moisture causes destabilization and iso- thermal transformation from tetragonal to monoclinic remains 6 oo to be understood. Many researchers have argued that stabilize- EEyY 9 Kc TZP ceramics tion is a direct consequence of the presence of the oxygen va cancies introduced by the aliovalent alloying element rather than the aliovalent dopant itself. 0.2 However, this hypothesis could 4 not be tested until the advent of large-scale computations. Now it has been shown computationally that the tetragonal phase 20 can be produced with lower energy than the monoclinic phase by introducing oxygen vacancies and without any aliovalent ions into the unit cell. This form of stabilization alone is unlikely 2 Kr Cubic to be the complete explanation because the solubility of the te- tragonal and cubic phases at different temperatures depends or the alloying dopant. Otherwise, all the phase diagrams for differ Aging sensitivity ent stabilizers would collapse onto one when plotted as a func- ion of oxygen vacancies. Nevertheless, it attractive explanation that has been used to rationalize the pre- 8 vailing explanation of LTD(see""Section IV): that moisture Yttria content(mol %Y,O3 pecies enter into the tetragonal lattice, annihilating oxygen ion Fig. 2. Fracture toughness and aging sensitivity of yttria-stabilized vacancies and thereby destabilizing the tetragonal(and cubic zirconia as a function of yttria stabilizer concentration. Toughness phases) data for cubic and monoclinic zirconia is indicated together with the The most systematic study of the role of different stabilizer rroelastic toughness in a densified thermal barrier coating. aging (dopant) ions on the stability of tetragonal and cubic zirconia sensitivity is here described by the degree of tetragonal phase transfo has been performed using X-ray absorption spectroscopy and mation results published in a series of papers by li et al. 20.23 They ex tetragonal fraction) after 3 h at 134.C in water vapor. mined the effect of trivalent and tetravalent dopant ions, and the effect of undersized and oversized dopants on the local en grains, whether these are also tetragonal grains, as in TzP, or ronment of zirconium ions Local atomic structures around the cubic matrix material in which the tetragonal precipitates are the Zr and around dopant cations in zirconia solid solution mbedded for the psz ceramics. as the I-mm transformation is were determined. These included undersized(Fe, GaT) and ccompanied by a volume increase, the transformation is con 3+ Gd trivalent ions as well as undersized strained by the surrounding grains. In these cases, the therme (Ge)and oversized(Ce)tetravalent ions dynamic framework that includes mechanical work arguments In the case of trivalent dopants, oxyge provides a rationale for the stabilization. The most general de- ated for charge com vas concluded tha scription of the energetics involved is reproduced below cies are associated with the zr cations in the case of oversized Recently with detailed study of zirconia nanoparticles, Gar dopants, and with the two dopant cations in the case of under- ie's claimthat pure, zirconia powders could be retained in the zed dopants. Both configurations favor sevenfold coordinated tetragonal state provided that their size was below a critical size xygen ions around the Zr cations and stabilize the tetragonal or has been extensively validated. Garvie's concept was that stabi lization could occur by surface energy alone and a series of en- gen vacancies to Zr is responsible for the more effective stabi ergy cross-overs between monoclinic, tetragonal, and cubic ation effects of oversized trivalent dopants. In essence, the been demonstrated as being consistent with stabilization of tetragonal zirconia with oversized trivalent cat- the dominant role of surface energy at nanometer particle sizes For example, a 4-24 nm stability range of tetragonal zirconia at From the results and the analysis performed by Li and col- room temperature can be extrapolated from Pitcher et al.s leagues, it is evident that doping by trivalent oversized cations, work,while monoclinic phase is stable above this size. This such as Y3+, is most efficient in relieving the overcrowd ole of surface energy is confirmed by Suresh et al.,who found ing(via both oxygen vacancies generation and dilatation of the a decrease of the f-m transformation temperature upon cooling tion network). The conceptual idea being that oxygen over- crowding around the small zirconium Zr cation is responsib for th the tetragonal phase may be stabilized by oxygen vacancies ad jacent to the Zr t ion and introduced by aliovalent doping. this lso provides a conceptual explanation for the stabilization by dilatation of the cation structure, and explains why 1.5 mol% of Y,O3 is sufficient to stabilize tetragonal zirconia, whereas Stresses Roughening Micro-cracking Path for wate 10 mol% of CeO2 is needed to achieve the same stability. Recent k- using Y NMR has provided more direct experimen support for the preference of oxygen vacancies to reside in lattice sites adjacent to the zr ion in YSz alloy Increased wear Slow Crack Growth (1) The Energetics of Transformation The foregoing discussion describes the stabilization of tetrago nal and cubic phases of zirconia under stress-free conditions Wear debris Loss in strength For polycrystalline ceramics, where the tetragonal phase is retained, transformation is mechanically constrained under Fig 3. Potential effect of aging on the integrity of zirconia devices. metastable conditions. The condition for transformation can ontributions to the mechanisms
grains, whether these are also tetragonal grains, as in TZP, or the cubic matrix material in which the tetragonal precipitates are embedded for the PSZ ceramics. As the tm transformation is accompanied by a volume increase, the transformation is constrained by the surrounding grains. In these cases, the thermodynamic framework that includes mechanical work arguments provides a rationale for the stabilization. The most general description of the energetics involved is reproduced below. Recently with detailed study of zirconia nanoparticles, Garvie’s claim17 that pure, zirconia powders could be retained in the tetragonal state provided that their size was below a critical size has been extensively validated. Garvie’s concept was that stabilization could occur by surface energy alone and a series of energy cross-overs between monoclinic, tetragonal, and cubic phases has since been demonstrated as being consistent with the dominant role of surface energy at nanometer particle sizes. For example, a 4–24 nm stability range of tetragonal zirconia at room temperature can be extrapolated from Pitcher et al.’s work,18 while monoclinic phase is stable above this size. This role of surface energy is confirmed by Suresh et al.,19 who found a decrease of the tm transformation temperature upon cooling with grain size. While the stabilization of pure zirconia can be understood in terms of the balance between chemical and surface energy, the reason that different aliovalent ions are effective as stabilizers and also, perhaps, why moisture causes destabilization and isothermal transformation from tetragonal to monoclinic remains to be understood. Many researchers have argued that stabilization is a direct consequence of the presence of the oxygen vacancies introduced by the aliovalent alloying element rather than the aliovalent dopant itself.20,21 However, this hypothesis could not be tested until the advent of large-scale computations. Now, it has been shown computationally22 that the tetragonal phase can be produced with lower energy than the monoclinic phase by introducing oxygen vacancies and without any aliovalent ions into the unit cell. This form of stabilization alone is unlikely to be the complete explanation because the solubility of the tetragonal and cubic phases at different temperatures depends on the alloying dopant. Otherwise, all the phase diagrams for different stabilizers would collapse onto one when plotted as a function of concentration of oxygen vacancies. Nevertheless, it is an attractive explanation that has been used to rationalize the prevailing explanation of LTD (see ‘‘Section IV’’): that moisture species enter into the tetragonal lattice, annihilating oxygen ion vacancies and thereby destabilizing the tetragonal (and cubic phases). The most systematic study of the role of different stabilizer (dopant) ions on the stability of tetragonal and cubic zirconia has been performed using X-ray absorption spectroscopy and results published in a series of papers by Li et al. 20,23 They examined the effect of trivalent and tetravalent dopant ions, and the effect of undersized and oversized dopants on the local environment of zirconium ions. Local atomic structures around the Zr41 and around dopant cations in zirconia solid solutions were determined. These included undersized (Fe31, Ga31) and oversized (Y31, Gd31) trivalent ions as well as undersized (Ge41) and oversized (Ce41) tetravalent ions. In the case of trivalent dopants, oxygen vacancies are generated for charge compensation. It was concluded that the vacancies are associated with the Zr cations in the case of oversized dopants, and with the two dopant cations in the case of undersized dopants. Both configurations favor sevenfold coordinated oxygen ions around the Zr cations and stabilize the tetragonal or even the cubic phases. However, the different availability of oxygen vacancies to Zr is responsible for the more effective stabilization effects of oversized trivalent dopants. In essence, the stabilization of tetragonal zirconia with oversized trivalent cations is twice as efficient as with undersized trivalent cations. From the results and the analysis performed by Li and colleagues, it is evident that doping by trivalent oversized cations, such as Y31, is most efficient in relieving the oxygen overcrowding (via both oxygen vacancies generation and dilatation of the cation network). The conceptual idea being that oxygen overcrowding around the small zirconium Zr41 cation is responsible for the poor stability of undoped tetragonal zirconia, and that the tetragonal phase may be stabilized by oxygen vacancies adjacent to the Zr41 ion and introduced by aliovalent doping. This also provides a conceptual explanation for the stabilization by dilatation of the cation structure, and explains why 1.5 mol% of Y2O3 is sufficient to stabilize tetragonal zirconia, whereas 10 mol% of CeO2 is needed to achieve the same stability. Recent work24 using 89Y NMR has provided more direct experimental support for the preference of oxygen vacancies to reside in lattice sites adjacent to the Zr41 ion in YSZ alloys. (1) The Energetics of Transformation The foregoing discussion describes the stabilization of tetragonal and cubic phases of zirconia under stress-free conditions. For polycrystalline ceramics, where the tetragonal phase is retained, transformation is mechanically constrained under metastable conditions. The condition for transformation can be expressed in terms of the different energy contributions to the overall energy, as discussed by Lange,25 who considered the Fig. 2. Fracture toughness and aging sensitivity of yttria-stabilized zirconia as a function of yttria stabilizer concentration. Toughness data for cubic and monoclinic zirconia is indicated together with the ferroelastic toughness in a densified thermal barrier coating. Aging sensitivity is here described by the degree of tetragonal phase transformation toward monoclinic (i.e., the ratio monoclinic fraction/initial tetragonal fraction) after 3 h at 1341C in water vapor. Fig. 3. Potential effect of aging on the integrity of zirconia devices. Arrows indicate coupling of aging with crack propagation and wear mechanisms. September 2009 The Tetragonal-Monoclinic Transformation in Zirconia 1905