Opportunities for TBCs in the ZrO2–YO1.5–TaO2.5 system

Surface & Coatings Technology 201 (2007) 6044 – 6050 www.elsevier.com/locate/surfcoat

Opportunities for TBCs in the ZrO2–YO1.5–TaO2.5 system Felicia M. Pitek ⁎, Carlos G. Levi Materials Department, University of California, Santa Barbara, Santa Barbara, CA, 93106-5050, USA Received 8 August 2006; accepted in revised form 13 November 2006 Available online 3 January 2007

Abstract An examination of the ZrO2–YO1.5–TaO2.5 system reveals several promising attributes for use in thermal barrier coating applications. The rather unique presence of a stable, non-transformable tetragonal region in this ternary oxide system allows for phase stability to high temperatures (1500 °C). Selected compositions with high levels of yttria and tantala have also shown superior resistance to vanadate corrosion than the commercially utilized 7YSZ. In addition, Y + Ta stabilized zirconia compositions within the non-transformable tetragonal phase field exhibit toughness values comparable or somewhat higher than those of 7YSZ, which bodes well for their durability as TBCs. These promising attributes are discussed in this paper in the context of recent experimental work. © 2006 Elsevier B.V. All rights reserved. Keywords: Thermal barrier coatings; Hot corrosion; Phase stability; Toughness; Zirconia

1. Introduction State-of-the-art thermal barrier coatings (TBCs) for gas turbines have a typical composition of ZrO2–7.6 ± 1%YO1.5 (heretofore 7YSZ).1 Attractive properties of this oxide include low intrinsic thermal conductivity (k ∼ 2.3 ± 0.1 W m− 1 K− 1 for 200 °C ≤ T ≤ 1000 °C) [1], a high coefficient of thermal expansion (CTE ∼ 10.7 ppm K− 1 for 20 °C ≤ T ≤ 1500 °C) [2], thermochemical stability in moisture-laden combustion environments [3], compatibility with the underlying thermally grown oxide (TGO) [4] and, most notably, strain tolerance that derives from a combination of adequate fracture toughness and a porous/microcracked architecture that promotes in-plane compliance [5]. Since its introduction 7YSZ has performed remarkably as a coating material and while its potential has yet to be fully exploited, it is also evident that it will eventually be limited in at least two important aspects [6]. One relates to the phase constitution of 7YSZ, which is based on a metastable nontransformable tetragonal form (t′) [7] whose stability is increasingly compromised by the relentless drive toward higher operating temperatures [6,8]. The second limitation relates to the resistance of 7YSZ to corrosion by molten deposits [9,10]. ⁎ Corresponding author. Tel.: +1 805 893 4723; fax: +1 805 893 8971. E-mail address: [email protected] (F.M. Pitek). 1 All compositions in mole percent unless specified otherwise. 0257-8972/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2006.11.011

Of interest to this work is the chemical de-stabilization of 7YSZ by molten sulfate–vanadate salts that result from Na ingestion with the intake air, especially in marine environments, combined with S/V impurities in the fuel [11–15]. The problem can be circumvented by tightly controlling the fuel quality, but that constrains the ability to expand the use of gas turbines for power generation, as well as the flexibility to re-fuel marine vessels in times of emergency [16]. At the root of both the desirable attributes as well as the limitations of 7YSZ as a TBC material is its rather unique phase constitution and associated chemical composition. The Y content must be high enough to render the TBC immune to the disruptive tetragonal ↔ monoclinic transformation upon thermal cycling [8], but low enough to retain the tetragonal structure and avoid its transformation to cubic at the projected use temperatures. The rationale for the latter derives from the absence of significant toughening mechanisms in the cubic form, whereas a non-transformable t′ phase can, in principle, exhibit ferroelastic toughening [17]. More importantly, the latter mechanism can operate at high temperatures [18], where transformation toughening is not effective [19]. Because the Y content required to yield a non-transformable t′ phase is beyond the equilibrium solubility of Y in ZrO2 at all temperatures of interest (∼1200 ± 200 °C) [4], the supersaturated t′ is susceptible to decomposition into Y-rich cubic and Ylean tetragonal phases at high temperature, rendering the latter

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transformable to monoclinic upon cooling with deleterious consequences to the durability of the coating [8,20,21]. The problem is severely aggravated when molten sulfate/vanadate salts deposit on the surface of the TBC, whereby de-stabilization of t′ may occur at much lower temperatures (650–950 °C) by selective depletion of Y according to the reactions [9,13,22]: Y2 O3 ðt′  ZrO2 Þ þ V2 O5 ðmeltÞ→2YVO4 ↓

ð1Þ

Y2 O3 ðt′  ZrO2 Þ þ SO3 ðmeltÞ→Y2 ðSO4 Þ3 ðmeltÞ

ð2Þ

The mechanisms of these reactions are not well understood, but the effects are well documented. In general, YSZ is claimed to be more sensitive to degradation by formation of YVO4 than of Y2(SO4)3 [9], arguably because the latter requires a relatively high SO3 partial pressure (102 Pa at ∼ 900 °C) [23,24]. The search for TBC materials that are resistant to S/V melts and/or phase stable well above 1200 °C has been of interest for some time. A comparative analysis of various stabilizers including MgO, Y2O3, Sc2O3, In2O3, CeO2, SnO2 and TiO2 appeared in [16]. Among these, Sc2O3 was highlighted as the most promising [16,25,26], but the concentrations proposed are likely to render the structure cubic at high temperature, with concomitantly lower toughness. A review of the literature on the ZrO2–CeO2 system suggests a narrow range of compositions (N15% and b19%) that are non-transformable and intrinsically phase-stable (i.e. within the equilibrium tetragonal field) at 1400–1500 °C but not at lower or higher temperatures because of the retrograde shape of the tetragonal solvus [27,28]. Compositions in this range are also reported to be resistant to hot corrosion [29] but can be subject to melt-mediated phase separation (t′ → c + m) at 700–900 °C [13,16], and may exhibit formation of Ce vanadate and/or sulfate in burner rig tests using fuels contaminated with high levels of S and V [30]. Interestingly, CeSZ is generally acknowledged to have lower erosion resistance than 7YSZ [31] but some studies claim a higher toughness for the former [32], which might be expected from the higher tetragonality of its unit cell [33]. Conversely, trivalent rare-earth oxides added to 7YSZ to further reduce its thermal conductivity generally decrease cyclic life and erosion resistance [6], consistent with a decrease in tetragonality [34], and have higher affinity for V2O5 than yttria [5]. An alternate approach to TBC design is offered by the ZrO2– YO1.5–TaO2.5 system (Fig. 1). The solubility of Ta5+ in ZrO2 is quite modest, but there is strong synergism when co-doping with Y3+ so that the solubility of YTaO4 in tetragonal ZrO2 is much higher than that of Y3+ or Ta5+ alone [35,36]. This strong interaction between Y and Ta in solid solution suggests that their activities are mutually reduced, which implies that the tendency of Y to react with V in the melt could be substantially lower than in 7YSZ. The system is also rather unique in that the YTaO4rich end of the tetragonal field in Fig. 1 is both stable up to 1500 °C and non-transformable to monoclinic upon cooling [35], at variance with most other ZrO2 systems for TBCs wherein non-transformable t′ phases are typically metastable. Finally, compositions in this range could also be amenable to ferroelastic toughening owing to their tetragonal structure,

Fig. 1. Isothermal section of the ZrO2–YO1.5–TaO2.5 phase diagram at 1500 °C showing the various regions of interest [35]. The “transformable” region is delineated by the intersection of the T0(t/m) surface with ambient, which moves to higher Y contents with increasing Ta substitution. The metastable region within the t + c field is conceptually equivalent to the t′/c′ range in the binary system and is susceptible to de-stabilization by partitioning into the equilibrium t (transformable) and c phases. The circles represent the compositions under investigation: (a) 7YSZ and (b) 16YTaSZ.

offering in principle a tough, phase stable composition that may resist S/V corrosion. Preliminary work elsewhere [37,38] is supportive of this hypothesis. Sintered pellets of ZrO2–20% YTaO4 exhibited k ∼ 1.8–2.3 W m− 1 K− 1 at 100–800 °C and CTE ∼ 10.5 ppm K− 1 at 1000 °C [38], as well as higher resistance to corrosion by NaVO3–V2O5 melts than 7YSZ [37]. Air plasma spray (APS) coatings of the same material were also reported to show promising durability in burner rig testing [38]. This investigation is part a broader effort to elucidate the mechanisms underlying the behavior of compositions in the ZrO2–YO1.5–TaO2.5 system relevant to TBC applications, and to provide guidelines for materials design. The present manuscript discusses preliminary findings regarding three key parameters, namely phase stability, corrosion and toughness, which are discussed in separate sections below. 2. Experimental procedure Stabilized zirconia powders were prepared by reverse coprecipitation from mixed solutions of precursor salts [39]. The starting materials were ZrOCl2·8H2O (99.9% pure), Y (NO3)3·6H2O (99.9%) and TaCl5 (99.8%) (all from Alfa Aesar, Ward Hill, MA). The solution precursors were prepared by dissolving the individual chemicals in de-ionized water, in the case of ZrOCl2·8H2O and Y(NO3)3·6H2O, or in ethanol for the TaCl5. The desired compositions were prepared by mixing the assayed solutions in the necessary proportions, which were then precipitated by adding them drop-wise to aqueous ammonium

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hydroxide. Single-phase solid solutions were obtained after drying the precipitates and then pyrolizing at 900 °C for 2 h. The primary compositions studied were 7YSZ as a baseline, and a co-doped sample with 16.6%YO1.5 + 16.6%TaO2.5 (16YTaSZ), which lies in the non-transformable tetragonal, single-phase regime. Two additional compositions with 14.5%YO1.5 + 14.5% TaO2.5 (14YTaSZ) and 17.6%YO1.5 + 17.6%TaO2.5 (17YTaSZ) were prepared for additional toughness measurements. For phase stability assessment the specimens were heat treated as loose particulates in order to minimize constraint effects that could hinder the tetragonal–monoclinic transformation. For consistency with a broader group of studies, the heat treatment schedule involved cumulative steps on the same specimen, with intermediate characterization by X-ray diffractometry (XRD) at ambient temperature, as discussed in detail elsewhere [40]. The cycles were performed in air and comprised a ramp up at 5 K/min, followed by a hold time typically of 24 h, and then cooling at 10 K/min. The treatment sequence was: 24 h/1200 °C, 72 h/1200 °C, and then 24 h/(Ti = Ti−1 + 50 °C) up to 1500 °C or until destabilization was detected, as manifested by the appearance of distinct (–111) and (111) monoclinic peaks at 2θ ∼ 28° and ∼ 31.5°, respectively, in the XRD pattern. For corrosion and toughness testing the pyrolized zirconia powders were compacted into pellets and sintered at 1200 °C/ 4 h for 7YSZ and 1500 °C/2 h for 16YTaSZ. X-ray diffraction (XRD) was performed on the pellets to ascertain that their phase constitution was still single phase (t′). Na2SO4–30%NaVO3 was selected as a “model deposit” for corrosion testing. Several pellets from each composition were run in parallel. The specimens were first heated to ∼ 220 °C and the corrodent in the form of aqueous slurry was added drop wise to the surface of the pellets to yield a dose of 25–35 mg/cm2. The salted pellets were then placed in a tube furnace and exposed at 900 °C under flowing air for 50 h. A crucible with additional corrodent was placed upstream from the sample to help reduce the rate of volatilization of salt from the specimen. After each cycle the pellet was removed from the furnace, washed with DI water, re-salted with an equivalent amount of corrodent, and exposed for another 50 h cycle. Every other cycle (100 h), a pellet of each composition was removed, washed, and analyzed using XRD and Raman spectroscopy. Testing was halted after significant degradation occurred, manifested by the presence of monoclinic zirconia or up to a total exposure time of 500 h. Additional analysis of the corroded surface was performed by scanning (SEM) and transmission (TEM) electron microscopy. Sintered pellets were polished to 3 μm using diamond lapping films and their toughness values were determined using a standard diamond shaped microhardness indenter to produce cracks from the corners of the indent [41,42] using loads in the range of 0.5–2 kg. The toughness (J/m2) required for crack propagation in the ceramic material was calculated from ∼ 40 indents on average by:
ð3Þ C ¼ n2 2d 2 P=c3 where P is the load, d the half diamond length of the Vicker's impression, c the crack length from the center of the indent,

and ξ = 0.016 a constant derived from a calibration of the test [42]. The lattice parameters and tetragonality of the toughness specimens were determined from the positions of the {400} XRD peaks prior to being indented. 3. Phase stability The phase evolution of 7YSZ and 16YTaSZ as a function of temperature is presented in Fig. 2. Both compositions are initially single-phase after pyrolysis at 900 °C/2 h. The tetragonality is not immediately evident in XRD because of nanocrystalline structure of the powder broadens the peaks and obscures the split of the {200)/(002) reflections at ∼ 34.5°. However, Raman spectroscopy confirms the tetragonal nature of the structure and the pattern persists after the first 1200 °C/24 h stage, wherein the {200)/(002) split is now clearly evident (Fig. 2). After further exposure to high temperatures the behaviors of the two compositions differ. 7YSZ is initially t′, but the cubic (200) reflection appears between the {200)/(002)t peaks after the 1300 °C step, signaling the onset of partitioning. (These observations are further supported by analysis of the {400)/ (004) peaks, but these are not shown in Fig. 2 in the interest of space). The relative intensity of the cubic (200) peak continues to increase with exposure to higher temperatures, consistent with the progress of partitioning, but the first hint of monoclinic formation upon cooling does not appear until ∼1450°, and the {111}m peaks at 28.5° and 31° are clearly detectable by XRD only after the 1500 °C treatment. Conversely, no evidence of monoclinic formation was found at any stage of the heat treatment for 16YTaSZ, and the co-doped material remained tetragonal up to 1500 °C, in agreement with the proposed thermodynamic stability of the tetragonal phase in this system at 1500 °C [35]. The extra reflections appearing between 1250 °C and 1450 °C in Fig. 2(b) correspond to a transient YTaO4 phase reported to form in materials synthesized from precursors [43], but its evolution is not well understood. The phase is indeed transient as it dissolves above 1450 °C and would not reappear in the present samples even after prolonged (1 week) additional exposure to 1300 °C, where it was originally observed. On first impression the difference in phase stability between 7YSZ and 16YTaSZ does not appear critical, as 1500 °C is well beyond current operating temperatures for TBCs (≤ 1200 °C). While the test is not a rigorously quantitative measure of phase stability, it is evident that the temperatures involved in destabilizing 7YSZ are substantially higher than the 1200 °C limit suggested in earlier works [8]. This result has been confirmed by tests in other similarly prepared materials [40,44] as well as in electron-beam physical vapor deposition (EB-PVD) coatings [45]. Nevertheless, it is clear that 7YSZ is thermodynamically metastable at temperatures relevant to TBC applications and indeed partitions readily, whereas 16YTaSZ is immune to decomposition into c + t even at 1500 °C. The absence of monoclinic upon cooling in 7YSZ is a result of sluggish kinetics [21,46], but 16YTaSZ is actually non-transformable as indicated by Fig. 1. It is worth noting that small additions of Ta5+ to 7YSZ should be detrimental rather than beneficial to phase stability.

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Fig. 2. XRD patterns of (a) 7YSZ and (b) 16YTa SZ revealing the phase evolution behavior of the materials during exposure to increasingly higher temperatures. 7YSZ undergoes partitioning into the cubic + tetragonal phases, where the tetragonal phase is Y-lean, leading to the eventual transformation into monoclinic zirconia upon cooling from 1500 °C. The 16YTaSZ sample remains tetragonal throughout the treatment, except for a minor transient YTaO4 phase that forms at intermediate temperatures but dissolves with further heating.

As Ta5+ substitutes for Zr4+ in 7YSZ the anion vacancies responsible for the stabilization of the tetragonal phase against the monoclinic transformation are gradually annihilated. When the addition of TaO2.5 is above ∼ 4% the structure moves from the c + t field into the stable t field becoming insensitive to partitioning, but its T0(t/m) rises above ambient making it transformable and therefore inadequate for TBCs. Intermediate compositions are metastable t′, as is 7YSZ, but the effect of Ta5+ on their kinetic stability remains to be assessed. With sufficiently high levels of YTaO4 the synergistic effect of the local distortions associated with the substitution of larger (Y3+) and smaller (Ta5+) cations for Zr4+ result in stabilization of the tetragonal phase even in the absence of anion vacancies [36]. Concomitantly, the T0(t/m) temperature falls below ambient, although the reasons for this are less understood.

This results in the unique “non-transformable” but stable t phase at the tip of the tetragonal field in Fig. 1. The implications of this region are important, as it gives rise to a composition domain wherein one could synthesize a metastable t′ phase that would remain non-transformable even if it were to partition during high temperature exposure. The phase stability of ZrO2–YO1.5– TaO2.5 compositions in the c + t field is further examined in a forthcoming publication. 4. Corrosion Substantial amounts of research have been done on the corrosion behavior of YSZ in S/V melts [11–15], but the usual description of the mechanism as one of “leaching” of stabilizer is rather inconsistent with the temperatures of the reaction

Fig. 3. Views of the surface of a 7YSZ sample after 100 h of exposure at 900 °C in air to a Na2SO4-30%NaVO3 melt. The “crust” of a large monoclinic ZrO2 and YVO4 crystals in (a) is shown in greater detail in (b) and (c).

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Fig. 4. Views of the surface of a 16YTaSZ sample after 500 h of exposure at 900 °C in air to a Na2SO4-30%NaVO3 melt. A few crystals of NaTaO3 developed on the surface of the pellet, as seen in (a) and (b), along with some degradation to the surface (c), as explained further in Fig. 5.

(b1000 °C), wherein cation diffusion within ZrO2 is extremely sluggish [47]. The more likely mechanism is one of dissolution of 7YSZ in the melt followed by re-precipitation of both YVO4 and Y-depleted zirconia, but details in the literature are rather scarce and the issue was further explored as part of this work. Clear evidence of the proposed dissolution–reprecipitation mechanism is shown in Fig. 3. Monoclinic zirconia was first detected on the surface of the pellet after the first corrosion cycle (50 h/900 °C). A 20- to 40-μm-thick crust of YVO4 and Ydepleted m-ZrO2 was observed after 100 h along with extensive spallation (Fig. 3a). The grain size of the residual 7YSZ on the spalled surface is an order of magnitude smaller (b 0.5 μm) than that of either of the reaction products. The YVO4 crystals appear to be ribbon-like and lie closer to the plane of the surface (Fig. 3b), while the m-ZrO2 crystals are acicular, with roughly equiaxed cross-sections, and protrude away from the surface (Fig. 3c). TEM analysis of the m-ZrO2 revealed large single

crystals with no evidence of solid-state transformation, suggesting that they precipitated directly as monoclinic from the melt, rather than transforming from a high temperature tetragonal phase upon cooling. In this sense the phenomenon is different from thermal de-stabilization since there would not be disruptive stresses associated with the t → m transformation upon cooling (although other deleterious strains clearly arise as evidenced by the extensive spallation). The orientation of the mZrO2 crystals is consistent with growth into the melt, although it is not clear whether this growth occurred at temperature or upon cooling. The predominant in-plane orientation of the YVO4 crystals is not well understood and remains to be elucidated. Minor amounts of monoclinic were observed in the 16YTaSZ co-doped composition in both XRD and Raman spectroscopy but only after 500 h of exposure to the corrodent. SEM revealed evidence of some surface degradation and a minimal amount of NaTaO3 formation, as shown in Fig. 4, but no significant spallation. TEM analysis of this specimen revealed a thin surface layer (b500 nm) with a tetragonal structure that appears to have grown epitaxially on the parent 16YTaSZ grain (Fig. 5); whether this layer develops on cooling or at temperature remains to be elucidated. The surface layer was significantly depleted in both Y (3.1%) and Ta (5%) but the diffraction patterns are consistent with tetragonal, rather than monoclinic zirconia. The detection of monoclinic by XRD and Raman spectroscopy suggests that both phases may be present, but the details remain to be determined. The 16YTaSZ sample appears to have dissolved slightly into the corrodent salt at temperature, where some of the Ta reacted with Na to form sodium tantalate crystals that deposited on the surface. Table 1 Toughness (J/m2), along with the corresponding standard deviation, and tetragonality values for 7YSZ and several compositions within the nontransformable tetragonal regime of the YO1.5–TaO2.5–ZrO2 system

Fig. 5. TEM micrograph of the 16YTaSZ composition after 500 h of exposure at 900 °C in air to a Na2SO4-30%NaVO3 melt. A thin layer of tetragonal zirconia (a), depleted in Y and Ta, was detected on the surface of the tetragonal parent grains (b).

Sample

YO1.5

TaO2.5

Γ (J/m2)

σ (J/m2)

c/a

7YSZ 14YTaSZ 16YTaSZ 17YTaSZ

7.6 14.5 16.6 17.6

– 14.5 16.6 17.6

41.8 52.5 41.4 39.3

8.4 13.3 12.4 10.4

1.013 1.0235 1.024 1.025

The concentrations of YO1.5 and TaO2.5 are noted for each composition studied.

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However, no YVO4 was observed. In spite of these signs of incipient corrosion, it is evident that the 16YTaSZ composition has significantly higher corrosion resistance than the standard 7YSZ. Additional research is underway to assess the effects of varying SO3 partial pressure on the relative corrosion resistance of these materials.

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tion measurements as 7YSZ, which bodes well for its eventual durability in TBC applications. Preliminary results suggest that higher toughness values are possible with small changes in composition, which should not degrade the phase stability or corrosion resistance of the material. Ongoing research will broaden the study to other compositions within this system, to be reported in subsequent publications.

5. Toughness Acknowledgments The toughness values measured on 7YSZ and the three Y + Ta co-doped compositions are listed in Table 1, together with values for the tetragonality calculated from the {400} peaks in the XRD pattern. Confidence in the results is given by the agreement between the measurement for 7YSZ and prior reports in the literature [48]. It is first noted that all the co-doped compositions tested have toughness values at least comparable or higher than 7YSZ, in spite of a significantly higher tetragonality that would suggest a larger improvement based on the hypothesis that ferroelastic switching is the underlying mechanism. Moreover, the toughness decreases with the small but systematic increase in tetragonality for the Y + Ta compositions, a trend also noted in an earlier study [35] but not yet understood. One possibly significant factor for the rather modest improvement in toughness between the 7YSZ and the codoped compositions is the substantially smaller grain size of the former. Because these compositions did not initially contain twin tetragonal domains, as characteristic of the other t′ structures in which ferroelastic toughening has been studied, one might anticipate that twin nucleation could be a critical step in enabling the mechanism, and the twinning process could be influenced by differences in the grain size as is the case in transformation toughening [49]. Indeed, twin nucleation has been observed around indentation cracks in 7YSZ samples originally twin-free and subjected to toughness measurements by the same technique used in this study [50]. Alternatively, it is possible that the coercive stress for domain boundary motion increases with tetragonality, and thus the requisite stresses may be achieved in a smaller fraction of the grains around the crack. Elucidation of these issues is in progress and will be reported in a future publication. 6. Conclusions The ZrO2–YO1.5–TaO2.5 system offers remarkable opportunities for developing novel TBC compositions with improved phase stability, resistance to corrosion by sulfate/vanadate melts and toughness at least comparable with that of the state-of-theart 7YSZ material. In particular, a 16.6%YO1.5 + 16.6%TaO2.5 stabilized zirconia composition is tetragonal, stable against phase partitioning up to at least 1500 °C, and insensitive to the tetragonal–monoclinic transformation upon thermal cycling. This material also showed only slight evidence of corrosion in S/V melts after 500 h, compared with extensive attack and spallation at much shorter times (50–100 h) for 7YSZ. Properly designed co-doping thus circumvents the problems of thermal and chemical de-stabilization in 7YSZ. In bulk form, the same composition shows essentially the same toughness by indenta-

The Office of Naval Research sponsored this investigation under grant N00014-99-1-0471, originally monitored by Dr. Steven Fishman and currently by Dr. David Shifler. Helpful discussions with Professors G.H. Meier and F.S. Pettit (Univ. Pittsburgh), as well as Prof. A.G. Evans and Dr. S. Krämer (UCSB), are gratefully acknowledged. The authors would also like to thank Ms. M. Gentleman, Messrs. D. Stave and B. Gilbert, and Drs. J.P.A. Löfvander and C. Mercer, for assistance with various aspects of the experimental techniques. References [1] R. Mévrel, J.C. Laizet, A. Azzopardi, B. Leclercq, M. Poulain, O. Lavigne, D. Demange, J. Eur. Ceram. Soc. 24 (2004) 3081. [2] J.W. Adams, H.H. Nakamura, J. Am. Ceram. Soc. 68 (9) (1985) C228. [3] K.N. Lee, N.S. Jacobson, R.A. Miller, MRS Bull. 14 (10) (1994). [4] O. Fabrichnaya, F. Aldinger, Z. Met.kd. 95 (1) (2004) 27. [5] D.R. Clarke, C.G. Levi, Ann. Rev. Mater. Res. 33 (2003) 383. [6] C.G. Levi, Curr. Opin. Solid State Mater. Sci. 8 (1) (2004) 77. [7] A.H. Heuer, R. Chaim, V. Lanteri, in: S. Somiya, N. Yamamoto, H. Yanagida (Eds.), Science and Technology of Zirconia, vol. III, Amer. Ceram. Soc, Westerville, OH, 1988, p. 3. [8] R.A. Miller, J.L. Smialek, R.G. Garlick, in: A.H. Heuer, L.W. Hobbs (Eds.), Science and Technology of Zirconia, Amer. Ceram. Soc., Inc., Columbus, OH, 1981, p. 241. [9] R.L. Jones, in: K.H. Stern (Ed.), Metallurgical and Ceramic Protective Coatings, Chapman and Hall, London, 1996, p. 194. [10] M.P. Borom, C.A. Johnson, L.A. Peluso, Surf. Coat. Technol. 86–87 (1996) 116. [11] A.S. Nagelberg, J. Electrochem. Soc. 132 (1) (1985) 2502. [12] R.L. Jones, C.E. Williams, S.R. Jones, J. Electrochem. Soc. 133 (1) (1986) 227. [13] R.L. Jones, C.E. Williams, Surf. Coat. Technol. (1987) 349. [14] R.L. Jones, High Temp. Technol. 6 (4) (1988) 187. [15] W. Hertl, J. Appl. Phys. 63 (11) (1988) 5514. [16] R.L. Jones, NRL/MR/6170-96-7841, Naval Research Lab., Washington, D.C., 1996 [17] A.V. Virkar, Key Eng. Mater. 153–154 (1998) 183. [18] D. Baither, M. Bartsch, B. Baufeld, A. Tikhonovsky, F.A.M. Rühle, U. Messerschmidt, J. Am. Ceram. Soc. 84 (8) (2001) 1755. [19] R.H.J. Hannink, P.M. Kelly, B.C. Muddle, J. Am. Ceram. Soc. 83 (3) (2000) 461. [20] N.R. Rebollo, Materials, University of California, Santa Barbara, CA, 2005. [21] V. Lughi, D.R. Clarke, Surf. Coat. Technol. 200 (5–6) (2005) 1287. [22] R.L. Jones, Mater. High Temp. 9 (4) (1991) 228. [23] R.L. Jones, J. Therm. Spray Technol. 6 (1) (1997) 77. [24] R.H. Barkalow, F.S. Pettit, Proceedings of the First Conf. on Advanced Materials for Alternative Fuel Capable Directly Fired Heat Engines (CONF-790749), U.S. Department of Energy, 1979. [25] R.L. Jones, D. Mess, Surf. Coat. Technol. 86–87 (1996) 94. [26] R.L. Jones, R.F. Reidy, D. Mess, Surf. Coat. Technol. 82 (1996) 70. [27] E. Tani, M. Yoshimura, S. Sömiya, J. Am. Ceram. Soc. 66 (7) (1983) 506. [28] J.R. Brandon, R. Taylor, Surf. Coat. Technol. 46 (1) (1991) 91.

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