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Martensitic transformation in zirconia containing ceramics and its applications
Current Opinion in Solid State and Materials Science 9 (2005) 313–318
Martensitic transformation in zirconia containing ceramics and its applications Xue-Jun Jin School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, China
Abstract An introduction to tetragonal (t) ! monoclinic (m) martensitic transformation in zirconia containing ceramics, especially tetragonal zirconia polycrystalline (TZP) was presented. Thermodynamics, crystallographics and kinetics of t ! m martensitic transformation in TZP were emphasized. Transformation toughening and shape memory effect (SME) associated with t ! m martensitic transformation in the TZP were reviewed. Perspective of future challenges was briefly mentioned at the end. 2006 Published by Elsevier Ltd. Keywords: Martensitic transformation; TZP; Transformation toughening; Shape memory effect
1. Introduction Zirconia containing ceramics are materials of imparting toughness (known as transformation toughening [1]) while maintaining strength and chemical inertness, and of exhibiting new functions such as shape memory effect [2] by manipulating new microstructure. These properties are mainly dominated by the structure transformation from tetragonal (t) to monoclinic (m). Zirconia containing ceramics can be classified into three categories: tetragonal zirconia polycrystalline (TZP), partially stabilized zirconia (PSZ) and zirconia toughened/ dispersed ceramics (ZTC/ZDC). Tetragonal zirconia polycrystalline (TZP) is a material with nearly 100% t-ZrO2 phase, stabilized by yttria or ceria additions [3]. An alternative way to stabilize the tetragonal phase is to decrease the grain size of tetragonal phase to nanoscale [*4,**5]. Grain sizes of TZP ceramics are typically in the range of 0.2–1 lm [6]. These ceramics are often designated with the prefix with Ce- or Y- to denote ceria- or yttria-stabilized, for example, 8Ce–0.5Y–TZP represents a 8 mol% CeO2 and 0.5 mol% Y2O3 stabilized zirconia.
E-mail address: [email protected] 1359-0286/$ - see front matter 2006 Published by Elsevier Ltd. doi:10.1016/j.cossms.2006.02.012
Undoped zirconia exhibits the following phase transitions under ambient during thermal cycling [*7]: 1170 C
2370 C
2680 C
m-ZrO2 ) * t-ZrO2 ) * c-ZrO2 ) * liquid 950 C
It has been well documented that the t ! m transformation is a athermal martensitic transformation, associated with a large temperature hysteresis (several hundred K), a volume change or dilation component of transformation strain (4– 5%) and a large shear strain (14–15% or 9) [6–*8]. This leads to disintegration of sintered undoped zirconia parts. Dopants (yittria, ceria, etc.) are added to stabilize the high temperature tetragonal and/or cubic phase in the sintered microstructure [3]. In the view of the potential commercial applications [9] (typically room temperature) of high temperature polymorphs (tetragonal and cubic) of ZrO2, the issues associated with t ! m martensitic transformation, related mechanism of transformation toughening and stabilization of metastable tetragonal phase at lower temperatures have drawn much attention in both ceramic research and martensitic transformation worlds for three decades [1–**11]. In the present review, the author discusses characteristics of tetragonal (t) ! monoclinic (m) martensitic transformation in TZP, the shape memory effect and the
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transformation toughening closely related to the transformation, room temperature stabilization of metastable tetragonal phase, as well as perspective of challenges. 2. Crystal structures and crystallographics of tetragonal (t) ! monoclinic (m) martensitic transformation Zirconia has three allotropes: cubic (c), tetragonal (t) and monoclinic (m). The tetragonal (t) and monoclinic (m) polymorphs have distorted fluorite structures shown in Fig. 1 from [*7].
Fig. 1. Crystal structures of tetragonal (t) and monoclinic (m) phases [*7].
The strong covalent nature of the Zr–O bond favours a sevenfold co-ordination number, and as a result, monoclinic ZrO2 is thermodynamically stable at lower temperatures, whereas the co-ordination number of Zr4+ cations in tetragonal and cubic-ZrO2 is 8. ‘‘To accommodate the thermally generated oxygen ion vacancies’’ at higher temperature, the structure of ZrO2 changes to the structure having eightfold co-ordination (t or c) while it still maintains an effective co-ordination number close to 7 owing to the association of Zr4+ ions with the oxygen ion vacancies [**5,12,13]. The detailed crystallographic information on different ZrO2 polymorphs can be found in Ref. [*7]. The crystallography of t ! m martensitic transformations has been evaluated by a phenomenological theory in lots of zirconia-containing ceramics [*8,14–17], especially a recent comprehensive review by Kelly and Francis Rose [*8]. The phenomenological theory is believed to be capable of explaining all reported microstructural and crystallographic features of the t ! m martensitic transformation in zirconia containing ceramics. The agreement between the experimental results and theoretical prediction demonstrates that the theory can be applied to make reliable, quantitative predictions for the martensitic transformation in those ceramic systems, even better than its application in steels where it was developed.
Fig. 2. In situ TEM observation for stress-induced martensitic transformation in 8Ce–0.25Y–TZP. Reversible motion of boundary between thermal stress-induced monoclinic and tetragonal phases was observed (marked by an arrow in (a)) and grows (b,c) when focusing the electron beam, shrinks (d,e) and disappears (f) when defocusing the electron beam.
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The thermoelastic behavior and crystallography of the t ! m martensitic transformation in Ce–Y–TZP ceramics were investigated by means of in situ TEM observation and Wechsler–Lieberman–Read (W–L–R) phenomenological theory. In situ TEM observations [17] showed that in Ce–Y–TZP the t/m interface can move freely with the change of thermal stress generated by beam illumination shown in Fig. 2, whereas it was not found in thermal cycles. Based on the features of reversibility of interface motion, large thermal hysteresis and high critical driving force for Ce–Y–TZP, the t ! m transformation in Ce– Y–TZP was suggested as a semi-thermoelastic one. The habit plane and the lattice correspondence were determined as (1 3 0)t and ([0 0 1]tk[0 1 0]m), which is in agreement with the calculated results by the phenomenological theory. 3. Thermodynamics of tetragonal (t) ! monoclinic (m) martensitic transformation
Fig. 3. Thermodynamic evaluation of Gibbs free energies for tetragonal and monoclinic phases of 8Ce–0.5Y–TZP.
The change in total Gibbs free energy associated with the athermal martensitic transformation t ! m can be expressed as [18–21] DGt!m ¼ V ðDGch þ DGstr Þ þ SDGsur ;
ð1Þ
where subscripts ‘‘ch’’, ‘‘str’’, and ‘‘sur’’ refer to the chemical free energy, the strain energy including both shear and dilatational energy, and the surface energy including the surface free energy, twinning energy and micro-cracking energy, respectively. V refers to the volume and S, the area associated with the transformation. The equilibrium temperature between the t ! m transformation, T0, is the temperature at which DGch = 0, and the Ms is defined as the temperature at which DGt!m = 0. The chemical energy difference between two phases (DGch) of a multi-element system can be calculated from the related binary systems by means of thermodynamic models [22]. The other required parameters in the right hand side of Eq. (1) can be derived through estimation either from some available data or by experiments. Thus the Ms temperature can be calculated according to Eq. (1). The difference of Gibbs free energy between tetragonal and monoclinic phases in ZrO2–CeO2–Y2O3 as a function of composition and temperature was thermodynamically calculated from the three related binary systems [21]. In 8 mol% CeO2–0.5 mol% Y2O3–TZP, the equilibrium temperature between tetragonal and monoclinic phases, T0, was evaluated as 832.5 K and the Ms temperature of this alloy with a mean grain size of 0.90 lm was calculated as 249.9 K using the approach, which is in good agreement with the experimental one of 253 K by dilation measurement (Fig. 3). Related to the application of external stress in the case of transformation toughening(see Section 5), the total tree energy change per unit volume required for constrained transformation [18] can be expressed as
S DGsur DGext V ¼ DGch DGext þ DGbarrier ;
DGt!m ¼ DGch þ DGstr þ
ð2Þ
where Gext is the interaction energy density due to the external stress; Gbarrier is the sum of the changes in surface and strain free energy. Considering the phenomenology of stress induced phase transformation, a critical transformation stress may be defined as rc ¼ ðDGch þ DGbarrier Þ=et ;
ð3Þ
where et is the resultant dilational transformation strain, localized in the transformation zone around the crack tip. Application of only the resultant dilational transformation strain in explaining the transformation toughening is still in the controversy (see Section 5 [*8]). It is clear that critical stress is reduced as the temperature approaches Ms, since the difference of chemical Gibbs free energy between t and m phases increases and contribute more to the driving force. Also, from Eq. (3), the total free energy change can be increased and t-ZrO2 is retained by, decreasing the chemical free energy change by stabilizing with the addition of dopant oxide (e.g. yttria, ceria); increasing the strain free energy change by dispersing the tetragonal phase in a constraining elastic matrix (e.g. alumina, cubic zirconia); increasing the surface free energy (e.g. by reducing the tetragonal grain size) [**11]. 4. Kinetics of t ! m martensitic transformation in TZP Generally, athermal diffusionless t ! m martensitic transformation takes place quickly, with the motion of phase boundary as high as the sound speed [*23]. The overall transformation proceeds in two major stages [24]. First,
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transition of the lattice structure from tetragonal to monoclinic occurs by shearing displacement of zirconium ions. The second stage involves migration of oxygen ions to oxygen sites in the monoclinic lattice. The displacement of the oxygen ions from the ideal fluorite positions along the caxis has been investigated by X-ray diffraction (XRD) [25]. It was proposed that, ‘‘while the rapid shear displacement of the zirconium ions is the rate-controlling factor for nucleation and longitudinal growth of the monoclinic plates, the migration of the oxygen ions controls the lateral growth of the plates’’. In the reverse m ! t transformation, migration of the Zr4+ and O2 ions to their respective positions is diffusion controlled. Strongly time dependent behavior was also observed in Ce–TZP [26]. It was also documented that the length of compressed Ce–Y–TZP specimens increases continuously in the aging at room temperature [27], resulting from gradual reverse m ! t transformation. This unusual anelasticity may be suggested as pseudo-anelasticity phenomenon associated with transformation, differing from normal anelasticity [28]. 5. Transformation toughening Zirconia containing ceramic is one of only two classes of materials exhibiting transformation toughening. The other one is transformation induced plasticity/TRIP steels. The martensitic t ! m transformation can be induced by cooling or by external loading under isothermal conditions [1,29]. Both transformation routes are of importance [**11]. ‘‘While thermally induced transformation will control the amount of tetragonal phase that can be retained after thermal cycling, the stress induced martensitic transformation enhances the toughness of zirconia ceramics’’. Martensitic transformation exhibits high speed and a change of shape of the transformed volume, both of which are essential for transformation toughening. Transformation toughening occurs when metastable retained t-ZrO2 transforms to the stable m-ZrO2 phase in the tensile stress field around a propagating crack [29]. The volume expansion (4–5%) characteristic of the t ! m martensitic transformation introduces a net compressive stress in the process zone around the crack tip [30,31]. This reduces the local crack tip stress intensity and hence the driving force for crack propagation, so increasing the effective toughness of the ceramics (Fig. 4). Following [32,33], PM Kelly and LR Francis Rose suggested [*8] a model of ‘decoupling’ the nucleation strain from the final strain—the net transformation strain and allowing the final transformation strains to include a shear component. Nucleation strain determines whether or not the stress-induced martensitic transformation can occur at the tip of a potentially dangerous crack. ‘‘It is the net transformation strain left behind in the transformed region that provides toughening by hindering crack growth’’.
Fig. 4. Schematic presentation of stress induced phase transformation of metastable tetragonal zirconia particles in crack tip stress field; arrow indicate generation of compressive residual stress by transformation induced volume expansion and microstructural constrain [31].
The characteristics of an ideal transformation toughened ceramic such as TZP are summarized as [*8]: • Martensitic transformation is suppressed with transformation start temperature Ms just below operating room temperature and metastable parent phase will be stressinduced transformed at the crack tip resulting in a positive volume change (dilation). • The shape strain has a relatively large shear component, which is of great importance to ensure that the transformation is easy to stress-induce at a crack tip and to have the shear accommodated by means of the formation of correlated variants. The underlying physical mechanisms of transformation toughening can be conveniently considered to involve either a process zone or a bridging zone [6]. 6. Shape memory effect Shape memory behavior originated from martensitic and its reverse transformation, tetragonal (t) M monoclinic (m), was first found in zirconia ceramics partially stabilized with magnesia (Mg–PSZ) in 1986 [2] and observed in ceria–TZP [*34] as well as ceria–yittria–TZP stabilized tetragonal zirconia polycrystals several years later [35]. Though the relative low recoverable strain, i.e., <1%, and the brittleness limit their practical application, the high operating temperature(a few hundred degrees higher than Nitinol shape memory alloys), high strength and chemical inertness make ZrO2 containing shape memory ceramics attractive. By comparison of the shape memory effect (SME) and related properties among specimens with different contents of (8–12 mol%) CeO2 and (0.25–0.75 mil%) Y2O3 fabricated by different processes, it was found that the 8Ce–0.5Y–TZP sintered for 6 h at 1773 K demonstrated excellent SME [36], i.e., a complete shape recovery rate with a strain of 1.2% shown in Fig. 5. No microcracks were found after shape recovery in 8Ce–0.5Y–TZP.
X.-J. Jin / Current Opinion in Solid State and Materials Science 9 (2005) 313–318
317
Stress (MPa)
2000
1500
1000
500
Temperature (˚C)
0
0
1
2
100
3
4
5
Strain (%)
200 300 400 500
Fig. 5. Morphology and shape memory effect of 8Ce–0.5Y–TZP ceramics manufactured by superfine particle and sintered at ambient, which shows 1.2% recovery strain upon pre-strained at room temperature followed by heating above 550 C.
7. Stabilization of metastable tetragonal phase
Tetragonal
Monoclinic
1/12 nm-1
Interface
Monoclinic
Tetragonal
Incoherent Partially coherent Coherent
1.46 0.73 0.29
1.1 0.55 0.22
of tetragonal structure in an isolated, strain free, spherically shaped ZrO2 nanocrystallite is because of the generation of excess oxygen ion vacancies as a result of the ‘‘nanoparticle size effect’’ [43]. Therefore, ‘‘the mechanism of tetragonal phase stabilization in nanocrystalline ZrO2 appears to be the same as that in doped ZrO2 at room temperature and undoped ZrO2 at higher temperature’’. The excess oxygen ion vacancies may correlate with the excess volume, a parameter used in the evaluation the contribution of surface layer to the whole Gibbs free energy by dilation model [44,45]. The contribution becomes significant as the grain size of ZrO2 nanoparticle below the critical value. The concept that the presence of vacancies are essential for stabilization of the tetragonal phase was not accepted by all and more evidence is required [*7]. 8. Conclusions
300 K
Temperature
1443 K
A detail review about mechanisms of room temperature metastable tetragonal phase stabilization in zirconia was recently presented by Shukla and Seal [**5]. The tetragonal phase can be stabilized at room temperature in an isolated, single, strain free nanoparticle below critical size of 10 nm on account of the surface energy difference between tetragonal and monoclinic phases (Table 1) besides to a large specific surface area on the nanoscale [*4,**5,37]. An schematic phase diagram including the effect of grain size can be schematically shown as Fig. 6. The critical size increases to 33 nm owing to aggregation of ZrO2 nanocrystallines [*4]. Various factors such as hydrostatic strain energy, non-hydrostatic strain energy, structural similarity, foreign surface oxides, water vapor and anionic impurities, significantly affect the tetragonal phase stability at room temperature [38–42]. It is suggested that the stabilization
Table 1 Interfacial energies (J m2) for monoclinic and tetragonal ZrO2 under different conditions reported by Garvie [*4,**5]
1/D
Fig. 6. Schematic presentation of the effect of nanocrystallinite size on the t ! m martensitic transformation temperature under ambient.
Characteristics (thermodynamics, kinetics and crystallographics) of t ! m martensitic transformation, related mechanism of transformation toughening and stabilization of metastable tetragonal phase at lower temperatures have been briefly reviewed. Combination of toughness and new functions makes TZP very attractive. Attention will be continuously paid to the mechanism of t ! m transformation in bulk as well as in nanocrystalline such as the role of vacancies in the stabilization of metastable tetragonal phase and the structural similarities between the evolving
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phases and the amorphous phase. Detail physical mechanism in the vicinity of the crack tip is needed to advance our understanding of transformation toughening. The relationship between microstructure resulting from alloying or optimization of processes and the new functions such as shape memory effect are not well understood. Acknowledgements The author thanks the National Natural Science Foundation of China (NNSFC 59971029) and Emerson Electric Corporation for the partial financial support. References The papers of particular interest have been highlighted as: * of special interest; ** of very special interest. [1] Gavie RC, Hannink RH, Pascoe RT. Ceramic steel. Nature 1975; 258:703–4. [2] Swain MV. Shape memory behavior in partially stabilized zirconia ceramics. Nature 1986;322:234–6. [3] Li P, Chen IW, Penner-Hahn JE. Effect of dopants on zirconia stabilization—an X-ray absorption study: 1, trivalent dopants. J Am Ceram Soc 1994;77:118–28. [*4] Garvie RC. The occurrence of metastable tetragonal zirconia as a crystallite size effect. J Phys Chem 1965;69:1238–43. [**5] Shukla S, Seal S. Mechanisms of room temperature metastable tetragonal phase stabilization in zirconia. Inter Mater Rev 2005;50: 45–64. [6] Evans AG. Perspective on the development of high-toughness ceramics. J Am Ceram Soc 1990;73:187–206. [*7] Kisi EH, Howard CJ. Crystal structures of zirconia phases and their inter-relation. Key Eng Mater 1998;153–154:1–36. [*8] Kelly PM, Francis Rose LR. The martensitic transformation in ceramics—its role in transformation toughening. Prog Mater Sci 2002;47:463–557. [9] Hench LL. Bioceramics. J Am Ceram Soc 1998;81:1705–28. [10] Hannink RHJ, Swain MV. Progress in transformation toughening in ceramics. Annu Rev Mater Sci 1994;24:359–408. [**11] Basu B. Toughening of yttria-stabilised tetragonal zirconia ceramics. Inter Mater Rev 2005;50:239–56. [12] Stefanic G, Music S. Factors influencing the stability of low temperature tetragonal ZrO2. Croat Chem Acta 2002;75:727–67. [13] Gibson IR, Irvine JTS. Qualitative X-ray diffraction analysis of metastable tetragonal (t 0 ) zirconia. J Am Ceram Soc 2001;84:615–8. [14] Hayakawa M, Adachi K, Oka M. Tween contrast with (2 2 3) habit in arc-melted zirconia–yttria alloys. Acta Metall Mater 1990;38: 1761–7. [15] Kelly PM, Ball CJ. Crystallography of stress-induced martensitic transformations in partially stabilized zirconia. J Am Ceram Soc 1986;69:259. [16] Kelly PM, Wauchope CJ. The tetragonal to monoclinic martensitic transformation in zirconia. Key Eng Mater 1998;153–154:97–124. [17] Zhang YL, Jin XJ, Rong YH, et al. On the t ! m martensitic transformation in Ce–Y–TZP ceramics. Acta Mater 2006;54: 1289–95. [18] Becher PF. Toughening behavior in ceramics associated with the transformation of tetragonal ZrO2. Acta Metall 1986;34:1885–91. [19] Zhou XW, Hsu (Xu Zuyao) TY. Thermodynamics of the martensitic transformation in Cu–Zn–Al alloys. Acta Metall Mater 1991;39: 1045–51.
[20] Jiang BH, Li L, Hsu (Xu Zuyao) TY. Thermodynamic calculation of the Ms temperature in 8 mol% CeO2–ZrO2. Mater Trans JIM 1996;37:1284–6. [21] Zhang YL, Jin XJ, Hsu TY. Thermodynamic calculation of Ms in ZrO2–CeO2–Y2O3 system. J Eur Ceram Soc 2003;23:685–90. [22] Hoch M. Calculation of ternary, quaternary, and higher-order phase diagrams and binary thermodynamic data. J Phase Equilib 1993;14:710–7. [*23] Cohen M, Wayman CM. Fundamentals of martensitic reactions. In: Tien JK, Elliott JF, editors. Metallurgical treatises. Warrendale, PA: TMS-AIME; 1983. p. 445–68. [24] Ge QL, Lei TC, Mao JF, et al. In situ transmission electron microscopy observations of the tetragonal-to-monoclinic phase transformation of zirconia in Al2O3–ZrO2 (2 mol%Y2O3) composite. J Mater Sci Lett 1993;12:819–22. [25] Lamas DG, Reca NEWD. Oxygen ion displacement in compositionally homogeneous, nanocrystalline ZrO2–12 mol% Y2O3 powders. Mater Lett 1999;41:204–8. [26] Kisi EH, Kennedy SJ, Howard CJ. Neutron diffraction observations of ferroelastic domain switching and t ! m transformation in Ce–TZP. J Am Ceram Soc 1997;80:621–8. [27] Zhang YL, Jin XJ, Hsu TY, et al. Time-dependent transformation in zirconia-based ceramics. Script Mater 2001;45:621–4. [28] Nowick AS, Berry BS. Anelastic relaxation in crystalline solids. New York and London: Academic Press; 1972. [29] Evans AG, Cannon RM. Toughening of brittle solids by martensitic transformations. Acta Mater 1986;34:761–800. [30] Green DJ, Hannink RHJ, Swain MV. Transformation toughening of ceramics. Boca Bacon, FL: CRC Press; 1989. [31] Ceramics and glasses. Engineering materials handbook, vol. 4. USA: ASM International, The Materials Information Society; 1991. [32] Lambropoulos JC. Effect of nucleation on transformation toughening. J Am Ceram Soc 1986;69:218–22. [33] Sun QP, Hwang KC, Yu SW. A micromechanics constitutive model of. Transformation lasticity with shear and dilatation effect. J Mech Phys Solids 1991;39:507–24. [*34] Reyes-Morel PE, Cherng JS, Chen IW. Transformation plasticity of CeO2-stabilized tetragonal zirconia polycrystals: II, Pseudoelasticity and shape memory effect. J Am Ceram Soc 1988;71:648–57. [35] Jiang BH, Tu JB, Hsu TY, et al. The effect of some factors on Ms and SME in Ce–TZP ceramics. Mater Res Soc Symp Proc 1992; 246:213–6. [36] Zhang YL, Jin XJ, Hsu TY, et al. Shape memory effect in Y–Ce– TZP. Mater Sci Forum 2002;392–394:573–6. [37] Chraska T, King AH, Berndt CC. On the size-dependent phase transformation in nanoparticulate zirconia. Mater Sci Eng 2000; A286:169–78. [38] Noh HJ, Seo DS, Kim H, et al. Synthesis and crystallization of anisotropic shaped ZrO2 nanocrystalline powders by hydrothermal process. Mater Lett 2003;57:2425–31. [39] Ramamoorthy R, Sundararaman D, Ramasamy S. X-ray diffraction study of phase transformation in hydrolyzed zirconia nanoparticles. J Eur Cearm Soc 1999;19:1827–33. [40] Winterer M, Nitsche R, Redfern SAR, et al. Phase stability in nanostructured and coarse grained zirconia at high pressures. Nanostruct Mater 1995;5:679–88. [41] Igawa N, Ishi Y. Crystal structure of metastable tetragonal zirconia up to 1473 K. J Am Ceram Soc 2001;84:1169–71. [42] Zhou XD, Huebner W. Size-induced lattice relaxation in CeO2 nanoparticles. Appl Phys Lett 2001;79:3512–4. [43] Shukla S, Seal S, Vanfleet K. Sol–gel synthesis and phase evolution behavior of sterically stabilized nanocrystalline zirconia. J Sol–Gel Sci Technol 2003;27:119–36. [44] Meng QP, Rong YH, Hsu TY. Nucleation barrier of phase transformation in nano-sized crystals. Phys Rev B 2002;65: 174118–24. [45] Zhang YL, Jin XJ, Hsu TY, et al., unpublished data.
Martensitic transformation in zirconia containing ceramics and its applications Xue-Jun Jin School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, China
Abstract An introduction to tetragonal (t) ! monoclinic (m) martensitic transformation in zirconia containing ceramics, especially tetragonal zirconia polycrystalline (TZP) was presented. Thermodynamics, crystallographics and kinetics of t ! m martensitic transformation in TZP were emphasized. Transformation toughening and shape memory effect (SME) associated with t ! m martensitic transformation in the TZP were reviewed. Perspective of future challenges was briefly mentioned at the end. 2006 Published by Elsevier Ltd. Keywords: Martensitic transformation; TZP; Transformation toughening; Shape memory effect
1. Introduction Zirconia containing ceramics are materials of imparting toughness (known as transformation toughening [1]) while maintaining strength and chemical inertness, and of exhibiting new functions such as shape memory effect [2] by manipulating new microstructure. These properties are mainly dominated by the structure transformation from tetragonal (t) to monoclinic (m). Zirconia containing ceramics can be classified into three categories: tetragonal zirconia polycrystalline (TZP), partially stabilized zirconia (PSZ) and zirconia toughened/ dispersed ceramics (ZTC/ZDC). Tetragonal zirconia polycrystalline (TZP) is a material with nearly 100% t-ZrO2 phase, stabilized by yttria or ceria additions [3]. An alternative way to stabilize the tetragonal phase is to decrease the grain size of tetragonal phase to nanoscale [*4,**5]. Grain sizes of TZP ceramics are typically in the range of 0.2–1 lm [6]. These ceramics are often designated with the prefix with Ce- or Y- to denote ceria- or yttria-stabilized, for example, 8Ce–0.5Y–TZP represents a 8 mol% CeO2 and 0.5 mol% Y2O3 stabilized zirconia.
E-mail address: [email protected] 1359-0286/$ - see front matter 2006 Published by Elsevier Ltd. doi:10.1016/j.cossms.2006.02.012
Undoped zirconia exhibits the following phase transitions under ambient during thermal cycling [*7]: 1170 C
2370 C
2680 C
m-ZrO2 ) * t-ZrO2 ) * c-ZrO2 ) * liquid 950 C
It has been well documented that the t ! m transformation is a athermal martensitic transformation, associated with a large temperature hysteresis (several hundred K), a volume change or dilation component of transformation strain (4– 5%) and a large shear strain (14–15% or 9) [6–*8]. This leads to disintegration of sintered undoped zirconia parts. Dopants (yittria, ceria, etc.) are added to stabilize the high temperature tetragonal and/or cubic phase in the sintered microstructure [3]. In the view of the potential commercial applications [9] (typically room temperature) of high temperature polymorphs (tetragonal and cubic) of ZrO2, the issues associated with t ! m martensitic transformation, related mechanism of transformation toughening and stabilization of metastable tetragonal phase at lower temperatures have drawn much attention in both ceramic research and martensitic transformation worlds for three decades [1–**11]. In the present review, the author discusses characteristics of tetragonal (t) ! monoclinic (m) martensitic transformation in TZP, the shape memory effect and the
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transformation toughening closely related to the transformation, room temperature stabilization of metastable tetragonal phase, as well as perspective of challenges. 2. Crystal structures and crystallographics of tetragonal (t) ! monoclinic (m) martensitic transformation Zirconia has three allotropes: cubic (c), tetragonal (t) and monoclinic (m). The tetragonal (t) and monoclinic (m) polymorphs have distorted fluorite structures shown in Fig. 1 from [*7].
Fig. 1. Crystal structures of tetragonal (t) and monoclinic (m) phases [*7].
The strong covalent nature of the Zr–O bond favours a sevenfold co-ordination number, and as a result, monoclinic ZrO2 is thermodynamically stable at lower temperatures, whereas the co-ordination number of Zr4+ cations in tetragonal and cubic-ZrO2 is 8. ‘‘To accommodate the thermally generated oxygen ion vacancies’’ at higher temperature, the structure of ZrO2 changes to the structure having eightfold co-ordination (t or c) while it still maintains an effective co-ordination number close to 7 owing to the association of Zr4+ ions with the oxygen ion vacancies [**5,12,13]. The detailed crystallographic information on different ZrO2 polymorphs can be found in Ref. [*7]. The crystallography of t ! m martensitic transformations has been evaluated by a phenomenological theory in lots of zirconia-containing ceramics [*8,14–17], especially a recent comprehensive review by Kelly and Francis Rose [*8]. The phenomenological theory is believed to be capable of explaining all reported microstructural and crystallographic features of the t ! m martensitic transformation in zirconia containing ceramics. The agreement between the experimental results and theoretical prediction demonstrates that the theory can be applied to make reliable, quantitative predictions for the martensitic transformation in those ceramic systems, even better than its application in steels where it was developed.
Fig. 2. In situ TEM observation for stress-induced martensitic transformation in 8Ce–0.25Y–TZP. Reversible motion of boundary between thermal stress-induced monoclinic and tetragonal phases was observed (marked by an arrow in (a)) and grows (b,c) when focusing the electron beam, shrinks (d,e) and disappears (f) when defocusing the electron beam.
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The thermoelastic behavior and crystallography of the t ! m martensitic transformation in Ce–Y–TZP ceramics were investigated by means of in situ TEM observation and Wechsler–Lieberman–Read (W–L–R) phenomenological theory. In situ TEM observations [17] showed that in Ce–Y–TZP the t/m interface can move freely with the change of thermal stress generated by beam illumination shown in Fig. 2, whereas it was not found in thermal cycles. Based on the features of reversibility of interface motion, large thermal hysteresis and high critical driving force for Ce–Y–TZP, the t ! m transformation in Ce– Y–TZP was suggested as a semi-thermoelastic one. The habit plane and the lattice correspondence were determined as (1 3 0)t and ([0 0 1]tk[0 1 0]m), which is in agreement with the calculated results by the phenomenological theory. 3. Thermodynamics of tetragonal (t) ! monoclinic (m) martensitic transformation
Fig. 3. Thermodynamic evaluation of Gibbs free energies for tetragonal and monoclinic phases of 8Ce–0.5Y–TZP.
The change in total Gibbs free energy associated with the athermal martensitic transformation t ! m can be expressed as [18–21] DGt!m ¼ V ðDGch þ DGstr Þ þ SDGsur ;
ð1Þ
where subscripts ‘‘ch’’, ‘‘str’’, and ‘‘sur’’ refer to the chemical free energy, the strain energy including both shear and dilatational energy, and the surface energy including the surface free energy, twinning energy and micro-cracking energy, respectively. V refers to the volume and S, the area associated with the transformation. The equilibrium temperature between the t ! m transformation, T0, is the temperature at which DGch = 0, and the Ms is defined as the temperature at which DGt!m = 0. The chemical energy difference between two phases (DGch) of a multi-element system can be calculated from the related binary systems by means of thermodynamic models [22]. The other required parameters in the right hand side of Eq. (1) can be derived through estimation either from some available data or by experiments. Thus the Ms temperature can be calculated according to Eq. (1). The difference of Gibbs free energy between tetragonal and monoclinic phases in ZrO2–CeO2–Y2O3 as a function of composition and temperature was thermodynamically calculated from the three related binary systems [21]. In 8 mol% CeO2–0.5 mol% Y2O3–TZP, the equilibrium temperature between tetragonal and monoclinic phases, T0, was evaluated as 832.5 K and the Ms temperature of this alloy with a mean grain size of 0.90 lm was calculated as 249.9 K using the approach, which is in good agreement with the experimental one of 253 K by dilation measurement (Fig. 3). Related to the application of external stress in the case of transformation toughening(see Section 5), the total tree energy change per unit volume required for constrained transformation [18] can be expressed as
S DGsur DGext V ¼ DGch DGext þ DGbarrier ;
DGt!m ¼ DGch þ DGstr þ
ð2Þ
where Gext is the interaction energy density due to the external stress; Gbarrier is the sum of the changes in surface and strain free energy. Considering the phenomenology of stress induced phase transformation, a critical transformation stress may be defined as rc ¼ ðDGch þ DGbarrier Þ=et ;
ð3Þ
where et is the resultant dilational transformation strain, localized in the transformation zone around the crack tip. Application of only the resultant dilational transformation strain in explaining the transformation toughening is still in the controversy (see Section 5 [*8]). It is clear that critical stress is reduced as the temperature approaches Ms, since the difference of chemical Gibbs free energy between t and m phases increases and contribute more to the driving force. Also, from Eq. (3), the total free energy change can be increased and t-ZrO2 is retained by, decreasing the chemical free energy change by stabilizing with the addition of dopant oxide (e.g. yttria, ceria); increasing the strain free energy change by dispersing the tetragonal phase in a constraining elastic matrix (e.g. alumina, cubic zirconia); increasing the surface free energy (e.g. by reducing the tetragonal grain size) [**11]. 4. Kinetics of t ! m martensitic transformation in TZP Generally, athermal diffusionless t ! m martensitic transformation takes place quickly, with the motion of phase boundary as high as the sound speed [*23]. The overall transformation proceeds in two major stages [24]. First,
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transition of the lattice structure from tetragonal to monoclinic occurs by shearing displacement of zirconium ions. The second stage involves migration of oxygen ions to oxygen sites in the monoclinic lattice. The displacement of the oxygen ions from the ideal fluorite positions along the caxis has been investigated by X-ray diffraction (XRD) [25]. It was proposed that, ‘‘while the rapid shear displacement of the zirconium ions is the rate-controlling factor for nucleation and longitudinal growth of the monoclinic plates, the migration of the oxygen ions controls the lateral growth of the plates’’. In the reverse m ! t transformation, migration of the Zr4+ and O2 ions to their respective positions is diffusion controlled. Strongly time dependent behavior was also observed in Ce–TZP [26]. It was also documented that the length of compressed Ce–Y–TZP specimens increases continuously in the aging at room temperature [27], resulting from gradual reverse m ! t transformation. This unusual anelasticity may be suggested as pseudo-anelasticity phenomenon associated with transformation, differing from normal anelasticity [28]. 5. Transformation toughening Zirconia containing ceramic is one of only two classes of materials exhibiting transformation toughening. The other one is transformation induced plasticity/TRIP steels. The martensitic t ! m transformation can be induced by cooling or by external loading under isothermal conditions [1,29]. Both transformation routes are of importance [**11]. ‘‘While thermally induced transformation will control the amount of tetragonal phase that can be retained after thermal cycling, the stress induced martensitic transformation enhances the toughness of zirconia ceramics’’. Martensitic transformation exhibits high speed and a change of shape of the transformed volume, both of which are essential for transformation toughening. Transformation toughening occurs when metastable retained t-ZrO2 transforms to the stable m-ZrO2 phase in the tensile stress field around a propagating crack [29]. The volume expansion (4–5%) characteristic of the t ! m martensitic transformation introduces a net compressive stress in the process zone around the crack tip [30,31]. This reduces the local crack tip stress intensity and hence the driving force for crack propagation, so increasing the effective toughness of the ceramics (Fig. 4). Following [32,33], PM Kelly and LR Francis Rose suggested [*8] a model of ‘decoupling’ the nucleation strain from the final strain—the net transformation strain and allowing the final transformation strains to include a shear component. Nucleation strain determines whether or not the stress-induced martensitic transformation can occur at the tip of a potentially dangerous crack. ‘‘It is the net transformation strain left behind in the transformed region that provides toughening by hindering crack growth’’.
Fig. 4. Schematic presentation of stress induced phase transformation of metastable tetragonal zirconia particles in crack tip stress field; arrow indicate generation of compressive residual stress by transformation induced volume expansion and microstructural constrain [31].
The characteristics of an ideal transformation toughened ceramic such as TZP are summarized as [*8]: • Martensitic transformation is suppressed with transformation start temperature Ms just below operating room temperature and metastable parent phase will be stressinduced transformed at the crack tip resulting in a positive volume change (dilation). • The shape strain has a relatively large shear component, which is of great importance to ensure that the transformation is easy to stress-induce at a crack tip and to have the shear accommodated by means of the formation of correlated variants. The underlying physical mechanisms of transformation toughening can be conveniently considered to involve either a process zone or a bridging zone [6]. 6. Shape memory effect Shape memory behavior originated from martensitic and its reverse transformation, tetragonal (t) M monoclinic (m), was first found in zirconia ceramics partially stabilized with magnesia (Mg–PSZ) in 1986 [2] and observed in ceria–TZP [*34] as well as ceria–yittria–TZP stabilized tetragonal zirconia polycrystals several years later [35]. Though the relative low recoverable strain, i.e., <1%, and the brittleness limit their practical application, the high operating temperature(a few hundred degrees higher than Nitinol shape memory alloys), high strength and chemical inertness make ZrO2 containing shape memory ceramics attractive. By comparison of the shape memory effect (SME) and related properties among specimens with different contents of (8–12 mol%) CeO2 and (0.25–0.75 mil%) Y2O3 fabricated by different processes, it was found that the 8Ce–0.5Y–TZP sintered for 6 h at 1773 K demonstrated excellent SME [36], i.e., a complete shape recovery rate with a strain of 1.2% shown in Fig. 5. No microcracks were found after shape recovery in 8Ce–0.5Y–TZP.
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317
Stress (MPa)
2000
1500
1000
500
Temperature (˚C)
0
0
1
2
100
3
4
5
Strain (%)
200 300 400 500
Fig. 5. Morphology and shape memory effect of 8Ce–0.5Y–TZP ceramics manufactured by superfine particle and sintered at ambient, which shows 1.2% recovery strain upon pre-strained at room temperature followed by heating above 550 C.
7. Stabilization of metastable tetragonal phase
Tetragonal
Monoclinic
1/12 nm-1
Interface
Monoclinic
Tetragonal
Incoherent Partially coherent Coherent
1.46 0.73 0.29
1.1 0.55 0.22
of tetragonal structure in an isolated, strain free, spherically shaped ZrO2 nanocrystallite is because of the generation of excess oxygen ion vacancies as a result of the ‘‘nanoparticle size effect’’ [43]. Therefore, ‘‘the mechanism of tetragonal phase stabilization in nanocrystalline ZrO2 appears to be the same as that in doped ZrO2 at room temperature and undoped ZrO2 at higher temperature’’. The excess oxygen ion vacancies may correlate with the excess volume, a parameter used in the evaluation the contribution of surface layer to the whole Gibbs free energy by dilation model [44,45]. The contribution becomes significant as the grain size of ZrO2 nanoparticle below the critical value. The concept that the presence of vacancies are essential for stabilization of the tetragonal phase was not accepted by all and more evidence is required [*7]. 8. Conclusions
300 K
Temperature
1443 K
A detail review about mechanisms of room temperature metastable tetragonal phase stabilization in zirconia was recently presented by Shukla and Seal [**5]. The tetragonal phase can be stabilized at room temperature in an isolated, single, strain free nanoparticle below critical size of 10 nm on account of the surface energy difference between tetragonal and monoclinic phases (Table 1) besides to a large specific surface area on the nanoscale [*4,**5,37]. An schematic phase diagram including the effect of grain size can be schematically shown as Fig. 6. The critical size increases to 33 nm owing to aggregation of ZrO2 nanocrystallines [*4]. Various factors such as hydrostatic strain energy, non-hydrostatic strain energy, structural similarity, foreign surface oxides, water vapor and anionic impurities, significantly affect the tetragonal phase stability at room temperature [38–42]. It is suggested that the stabilization
Table 1 Interfacial energies (J m2) for monoclinic and tetragonal ZrO2 under different conditions reported by Garvie [*4,**5]
1/D
Fig. 6. Schematic presentation of the effect of nanocrystallinite size on the t ! m martensitic transformation temperature under ambient.
Characteristics (thermodynamics, kinetics and crystallographics) of t ! m martensitic transformation, related mechanism of transformation toughening and stabilization of metastable tetragonal phase at lower temperatures have been briefly reviewed. Combination of toughness and new functions makes TZP very attractive. Attention will be continuously paid to the mechanism of t ! m transformation in bulk as well as in nanocrystalline such as the role of vacancies in the stabilization of metastable tetragonal phase and the structural similarities between the evolving
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phases and the amorphous phase. Detail physical mechanism in the vicinity of the crack tip is needed to advance our understanding of transformation toughening. The relationship between microstructure resulting from alloying or optimization of processes and the new functions such as shape memory effect are not well understood. Acknowledgements The author thanks the National Natural Science Foundation of China (NNSFC 59971029) and Emerson Electric Corporation for the partial financial support. References The papers of particular interest have been highlighted as: * of special interest; ** of very special interest. [1] Gavie RC, Hannink RH, Pascoe RT. Ceramic steel. Nature 1975; 258:703–4. [2] Swain MV. Shape memory behavior in partially stabilized zirconia ceramics. Nature 1986;322:234–6. [3] Li P, Chen IW, Penner-Hahn JE. Effect of dopants on zirconia stabilization—an X-ray absorption study: 1, trivalent dopants. J Am Ceram Soc 1994;77:118–28. [*4] Garvie RC. The occurrence of metastable tetragonal zirconia as a crystallite size effect. J Phys Chem 1965;69:1238–43. [**5] Shukla S, Seal S. Mechanisms of room temperature metastable tetragonal phase stabilization in zirconia. Inter Mater Rev 2005;50: 45–64. [6] Evans AG. Perspective on the development of high-toughness ceramics. J Am Ceram Soc 1990;73:187–206. [*7] Kisi EH, Howard CJ. Crystal structures of zirconia phases and their inter-relation. Key Eng Mater 1998;153–154:1–36. [*8] Kelly PM, Francis Rose LR. The martensitic transformation in ceramics—its role in transformation toughening. Prog Mater Sci 2002;47:463–557. [9] Hench LL. Bioceramics. J Am Ceram Soc 1998;81:1705–28. [10] Hannink RHJ, Swain MV. Progress in transformation toughening in ceramics. Annu Rev Mater Sci 1994;24:359–408. [**11] Basu B. Toughening of yttria-stabilised tetragonal zirconia ceramics. Inter Mater Rev 2005;50:239–56. [12] Stefanic G, Music S. Factors influencing the stability of low temperature tetragonal ZrO2. Croat Chem Acta 2002;75:727–67. [13] Gibson IR, Irvine JTS. Qualitative X-ray diffraction analysis of metastable tetragonal (t 0 ) zirconia. J Am Ceram Soc 2001;84:615–8. [14] Hayakawa M, Adachi K, Oka M. Tween contrast with (2 2 3) habit in arc-melted zirconia–yttria alloys. Acta Metall Mater 1990;38: 1761–7. [15] Kelly PM, Ball CJ. Crystallography of stress-induced martensitic transformations in partially stabilized zirconia. J Am Ceram Soc 1986;69:259. [16] Kelly PM, Wauchope CJ. The tetragonal to monoclinic martensitic transformation in zirconia. Key Eng Mater 1998;153–154:97–124. [17] Zhang YL, Jin XJ, Rong YH, et al. On the t ! m martensitic transformation in Ce–Y–TZP ceramics. Acta Mater 2006;54: 1289–95. [18] Becher PF. Toughening behavior in ceramics associated with the transformation of tetragonal ZrO2. Acta Metall 1986;34:1885–91. [19] Zhou XW, Hsu (Xu Zuyao) TY. Thermodynamics of the martensitic transformation in Cu–Zn–Al alloys. Acta Metall Mater 1991;39: 1045–51.
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