Cambridge, MA 02138, USA. Berkeley, CA 94720, USA

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1 This article was downloaded by:[harvard University] On: 30 November 2007 Access Details: [subscription number ] Publisher: Taylor & Francis Informa Ltd Registered in England and Wales Registered Number: Registered office: Mortimer House, Mortimer Street, London W1T 3JH, UK Philosophical Magazine First published in 1798 Publication details, including instructions for authors and subscription information: Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films M. Tsuchiya a ; A. M. Minor b ; S. Ramanathan a a Harvard School of Engineering and Applied Sciences, Harvard University, Cambridge, MA 02138, USA b National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA Online Publication Date: 01 December 2007 To cite this Article: Tsuchiya, M., Minor, A. M. and Ramanathan, S. (2007) 'Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films', Philosophical Magazine, 87:36, To link to this article: DOI: / URL: PLEASE SCROLL DOWN FOR ARTICLE Full terms and conditions of use: This article maybe used for research, teaching and private study purposes. Any substantial or systematic reproduction, re-distribution, re-selling, loan or sub-licensing, systematic supply or distribution in any form to anyone is expressly forbidden. The publisher does not give any warranty express or implied or make any representation that the contents will be complete or accurate or up to date. The accuracy of any instructions, formulae and drug doses should be independently verified with primary sources. The publisher shall not be liable for any loss, actions, claims, proceedings, demand or costs or damages whatsoever or howsoever caused arising directly or indirectly in connection with or arising out of the use of this material.

2 Philosophical Magazine, Vol. 87, No. 36, 21 December 2007, Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films M. TSUCHIYAy, A. M. MINORz and S. RAMANATHAN*y yharvard School of Engineering and Applied Sciences, Harvard University, Cambridge, MA 02138, USA znational Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA (Received 18 April 2007; in final form 24 September 2007) Phase stability in nanoscale pure zirconia and 9.5 mol.% yttria-doped zirconia (YDZ) thin films was studied by in-situ transmission electron microscopy. Oxygen vacancies are found to play a significant role in determining the microstructure and phase evolution. Pure zirconia thin films of 52 nm thickness were stabilized without any dopants at room temperature, whereas they transformed into a tetragonal phase upon heating to 400 C. On the other hand, 9.5% yttria doping enables stabilization of the cubic structure regardless of grain growth. Annealing of amorphous YDZ films in air (oxygen-rich) leads to tetragonal phase formation, whereas ultrahigh vacuum (oxygen-deficient) annealed samples display a cubic phase at high temperature. Detailed discussions on the effects of initial microstructure, oxygen deficiency, aliovalent doping and thickness are presented. 1. Introduction Among the class of fluorite oxides, zirconia is a particularly important material owing to its tremendous potential for technological applications [1]. Examples of potential applications for thin film zirconia include advanced nanoelectronic devices [2] and electrolyte membranes for hydrogen energy conversion devices [3]. The microstructure of zirconia controls many of its important material properties, including oxygen ionic conductivity, mechanical properties and electronic band gap [1, 4, 5]. Bulk pure zirconia has three phases; monoclinic from room temperature to 1443 K, tetragonal from 1443 K to 2643 K and a cubic structure from 2643K to its melting temperature of 2953 K (at atmospheric pressure) [1]. However, the phase stability in nanocrystalline zirconia is markedly different from that observed in the bulk. For example, it is known that the high-temperature tetragonal phase can be found at near room temperature in particles below about 10 nm in diameter [6]. Dopants such as yttria can also enable stabilization of the high-temperature cubic phase at lower temperatures [7]. Whereas there have been significant advances in *Corresponding author. shriram@seas.harvard.edu Philosophical Magazine ISSN print/issn online ß 2007 Taylor & Francis DOI: /

3 5674 M. Tsuchiya et al. understanding phase evolution in particles of pure and doped zirconia, little information is available for thin film systems [8]. Both pure and Y-doped thin film zirconia are particularly important owing to their significant applications in energy science and technologies, including intermediate-temperature fuel cells and separation membranes. Recently, there has been a surge in interest towards understanding the ion transport properties in nanoscale zirconia thin films owing to the interesting size dependence of electrochemical properties [9, 10]; however, the structural aspects are less well understood. Microstructural evolution studies are critical for interpreting the ion transport properties in such oxide systems [1, 11]. In this paper, we report systematic studies of the size dependence of phase transformations in pure and Y-doped thin film zirconia synthesized by electron beam evaporation. From detailed in-situ transmission electron microscopy (TEM) studies, we show that the phase transformation sequences in pure and Y-doped zirconia of comparable thicknesses are dramatically different from each other. Unlike bulk crystals, pure zirconia thin films undergo a cubic tetragonal phase transformation when heated from room temperature to about 400 C. Y-doped zirconia thin films can be grown in the cubic phase and remain stable during thermal treatments, although grain growth begins to occur around 500 C. It is possible to obtain amorphous Y-doped zirconia even in films less than 50 nm thickness. We show further that the phase transformations can also be affected by the annealing ambient through controlled ex-situ studies. These results are important for understanding phase stability in oxide ion conductors. 2. Experimental Pure zirconia and 9.5 mol.% yttria-doped zirconia (YDZ) pellets were synthesized by electron beam evaporation in a high-vacuum chamber (base pressure Torr) at room temperature to grow thin YDZ film on Ge (100) substrate. No substrate bias was applied during the film deposition. The estimated deposition rate was 0.5 Å s 1. Prior to deposition, substrates were cleaned with acetone and ethanol for 10 min, respectively, and then treated with buffered HF (2%) for 2 min followed by deionized water cleaning for 30 s before loading into the chamber. Ge substrate was chosen primarily to avoid significant interfacial reactions between the film and substrate. Extensive characterization was carried out by Rutherford backscattering and nuclear reaction analysis (using 16 O(d,p) 17 O reaction) on the various samples [12]. It was confirmed that interfacial layer formation between as-grown zirconia and the Ge substrate was negligible and the yttria content in the YDZ films was similar to the composition of the source material. The absence of interfacial layers was further confirmed by cross-section TEM. The cross-section micrograph of a representative 82 nm YDZ film used in this study is shown in figure 1. The film thicknesses were measured after growth by the X-ray low-angle reflectivity technique and also confirmed by ellipsometry. The film thicknesses investigated in this study ranged from about 10 nm to 100 nm. Plan view samples for TEM were first dimpled and then thinned to electron transparency by Ar þ ion beam milling. For the in-situ heating experiments, samples were mounted in a heating

4 Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films 5675 Ge YDZ Figure 1. Cross-sectional high-resolution TEM micrograph of 82 nm YDZ film on Ge (100) substrate along the h110i zone axis. The inset shows a selected area electron diffraction pattern. holder and heated inside the transmission electron microscope from room temperature to 1073 K. To investigate the effect of the ambient environment on phase transformations, identical samples were heated ex-situ in controlled atmospheres in a custom furnace. Heating and cooling rates were computer controlled to simulate nearly identical heat treatment conditions to the in-situ TEM studies for direct comparison. JEOL 3010 and JEOL 2100 transmission electron microscopes were used in this study. 3. Results and discussion 3.1. Pure zirconia thin films Figure 2 shows bright-field and dark-field images of microstructural evolution in a 52 nm thick pure ZrO 2 film. The room-temperature (as-grown) polycrystalline sample shown in figure 2a was heated in the transmission electron microscope at a rate of 5 C min 1. Grain growth occurs with increasing temperature. The bright- and dark-field images in figure 2b were taken at 500 C and 400 C, respectively. The grain size increased to 80 nm upon annealing to 800 C, as shown in figure 2c. The grain size defined in this work is the average of sampling in an area typically about 500 nm 500 nm per image. The measurements were performed in regions distributed over several microns. The grain size

5 5676 M. Tsuchiya et al. Figure 2. Bright- and dark-field images of the microstructural evolution in a pure ZrO 2 52 nm thin film grown on Ge substrate. (a) The room-temperature (as-grown) sample has an average grain size of 12 nm; (b) Grain growth occurs with increasing temperature. The bright-field image was taken at 500 C and dark-field image was taken at 400 C; (c) Grain size increased to 80 nm diameter upon annealing to 800 C. distribution (i.e. the number of grains with respect to grain size) was found to be Gaussian for nearly all samples. The dark-field images are a random sampling of grains from different orientations. As clearly seen in figure 2a, the as-grown film was crystalline with nanometre scale grains. The grain size had increased with annealing from 10 nm at room temperature (as-grown) to 80 nm at 800 C. Grain size increased gradually with temperature and no abnormal grain growth was observed. Figure 3 shows the phase evolution from electron diffraction. The as-grown cubic crystalline phase shown in figure 3a was transformed into a tetragonal phase at around 400 C. The diffraction patterns at the onset of phase transformation (400 C) and after the phase formation (450 C) are shown in figures 3b and 3c, respectively. No further grain growth was observed even at 800 C, as shown in figure 3d. The transformation temperature was found to be 400 C. The polycrystalline rings from the d 102 (d ¼ A ) plane is a clear indication of tetragonal phase formation. This result is consistent with grain size dependent phase transformation observed by Moulzolf and Lad [13]. They observed that the cubic phase was stable up to a grain size of 10 nm in zirconia films grown on amorphous silica substrate. The formation of cubic phase in pure zirconia is a particularly interesting result owing to the fact that it is almost never stable at room temperature without impurity doping [7].

6 Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films 5677 (a) (d) c (420) c (331) c (400) c (222) c (311) c (220) c (200) c (111) t (101) t (110) t (102) t (200) t (211) t (202) t (004) t (222) Figure 3. Phase evolution in a ZrO 2 52 nm thin film grown on a Ge substrate as shown by electron diffraction analysis: (a) room temperature cubic phase; (b) onset of phase change at 400 C; (c) tetragonal phase formation at 450 C; (d) DP showing the tetragonal phase was stable even at 800 C. The existence of metastable phases in zirconia nanoparticles has been suggested to be due to the competing volume and surface free energies of the various phases. Clearly, as the size decreases, the surface contributions become important and it is possible to calculate the critical radius when alternate phases with lower surface energy can become stable at room temperature [14]. Such analyses have predicted that tetragonal zirconia could be stable up to 30 nm grain size in particles. However, these results may not be extended directly to thin films owing to film substrate energetics and strain effects. Also it is important to note that the calculated surface energy of the cubic phase is higher than that of both tetragonal and monoclinic phases [5]. Further, the role of oxygen vacancies needs to be considered in detail to understand the phase evolution. For example, Mommer et al. [15] demonstrated that annealing tetragonal or monoclinic zirconia under low oxygen partial pressures results in cubic phase formation even in bulk systems. Similarly, Collins and Browman [16] reported that the tetragonal phase can be stabilized when nanocrystalline ZrO 2 was synthesized under low oxygen partial pressures. Theoretical studies on phase stability also indicate that oxygen vacancies may lead to cubic phase formation in the zirconia system [17]. In addition, the phase diagram of Zr ZrO 2 [18] and ZrO 2 Y 2 O 3 [7], where Y 2 O 3 substitution into ZrO 2 creates oxygen vacancies, clearly shows that the cubic phase stability has a strong correlation with oxygen vacancy concentration. It is generally known that surface and grain boundaries function as a sink or source for point defects owing to small defect formation energy. Recent studies show that vacancies in zirconia can be generated not only by oxygen reduction at high temperatures in ultrahigh vacuum or doping with trivalent impurities, but also by (b) (c)

7 5678 M. Tsuchiya et al. (a) 100 nm c (111) c (200) c (220) c (311) c (222) (b) 100 nm t (101) t (110) t (102) t (200) t (211) t (202) t (004) t (222) 100 nm increasing the surface or grain boundary areas. It is known that grain boundaries contain much higher oxygen vacancy concentration than bulk. Direct observation of higher oxygen vacancy concentration correlated with surface area has been reported by Liu et al. [19], in which the authors have studied the relationship between electron spin resonance (ESR) signal from zirconia nanopowders of nm size and the specific surface area. Their results indicate that the intensity of F-centre signal, arising from trapped electrons at oxygen vacancies, increases with an increase in specific surface area. Several reports that strongly support the model of increased oxygen vacancy concentration in nanocrystallites have been summarized by Shukla and Seal [6]. Interestingly, the oxygen vacancy concentration reduction due to grain growth is more significant than the oxygen incorporation from surrounding ambient. The as-grown 75 nm thick zirconia thin film shown in figure 4a was heated up ex-situ in ultrahigh vacuum (UHV) 10 8 Torr and air at 550 C for 1 h. The micrographs of the specimen after annealing in UHV and air are shown in figures 4b and 4c, respectively. In both cases, the grain size had increased from 10 nm to 100 nm. Similar to the in-situ heating result of 52 nm zirconia sample, tetragonal-to-cubic phase transformation due to grain growth is clearly seen in the diffraction patterns for both samples. It is expected that oxygen reduction takes place when samples are annealed in UHV, whereas oxygen incorporation takes place for annealing in air. Given the fact that cubic-to-tetragonal transformation was observed even in UHV annealed specimens, the oxygen vacancy concentration reduction due to grain growth appears to be more significant than the oxygen vacancy creation by the reduction reaction. The minimal effect of oxygen incorporation from surrounding ambient also supports this view. (c) t (101) t (110) t (102) t (200) t (211) t (202) t (004) t (222) Figure 4. Dark-field images and corresponding selected area diffraction patterns taken of a pure 75 nm ZrO 2 thin film grown on Ge substrate. Samples were annealed ex-situ at 550 C for 1 h in ultrahigh vacuum and air. (a) As-grown film; (b) annealed in ultrahigh vacuum (10 8 Torr) and (c) annealed in air.

8 Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films 5679 (a) (c) Dark Field Figure 5. Microstructural evolution of a 32 nm thick YDZ film on Ge: (a) at room temperature; (b) at 520 C; (c) bright-field image at 600 C; (d) dark-field image at 600 C. The film was amorphous to begin with, then instantaneously crystallized into the cubic fluorite structure with average grain size of 12 nm at 520 C. (b) (d) 3.2. Yttria-doped zirconia thin films Figure 5 shows the microstructural evolution in 32 nm thick 9.5 mol.% yttria-doped zirconia thin films synthesized under identical conditions. The as-grown film was amorphous and instantaneously crystallized into the cubic phase at 520 C and maintained the cubic structure even up to 800 C. The micrographs of as-grown amorphous structure at room temperature, onset of crystallization at 520 C, and fully crystallized structure at 600 C are shown in figures 5a 5c. The dark-field image taken at the same conditions as for figure 5c is shown in figure 5d. Diffraction patterns in figure 6 show the transformation sequence. The diffraction pattern of the as-grown room temperature amorphous phase, the onset of crystallization at 520 C, completion of crystallization at 520 C and the fully crystallized cubic phase are shown in figures 6a 6d. Detailed studies carried out on 42 nm thick films also showed similar results. In contrast, an identical as-grown amorphous YDZ film annealed in air showed significantly different behaviour: we observed formation of the tetragonal phase at nearly 600 C in these films, as shown in figure 7. The results correlate well with prior reports suggesting oxygen vacancies to be largely responsible for cubic phase formation [15]. A representative high-resolution image is shown in figure 8 indicating no amorphous phases formed at the grain boundaries. Size effects can also affect the crystallization process directly. In this study, we observed that 32 nm and 42 nm thick YDZ films had crystallized at 520 C. We also noticed that a 16 nm thick YDZ film grown on Si (100) substrate under identical

9 5680 M. Tsuchiya et al. Figure 6. Diffraction pattern for 32 nm thick YDZ film: (a) at room temperature; (b) onset of crystallization at 520 C; (c) completion of crystallization at 520 C; (d) at 600 C. (a) c (111) c (200) In-Situ 32 nm c (220) c (311) c (222) (b) Ex-Situ 32 nm t (111) t (200) t (112) t (220) t (311) t (222) Figure 7. Comparison of diffraction patterns for ex-situ (air) and in-situ (TEM) annealing. The ex-situ sample is tetragonal, whereas the sample annealed under ultrahigh vacuum is cubic. conditions retains an amorphous structure even at 800 C. The results show that the thinner the film, the higher the crystallization temperature. Gusev et al. [20] reported that the crystallization temperature of HfO 2 thin films increased by more than 100 C when film thickness is decreased from 40 nm to 5 nm. It has been reported that nanoscale oxides exists in amorphous phase below a critical thickness, which is considered an effect of the large surface/bulk ratio [21]. Presumably, the amorphous

10 Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films 5681 Figure 8. High-resolution image of inter-grain regions in the ex-situ annealed 32 nm thick YDZ film. No amorphous phase was found at the grain boundaries. phase has a smaller surface energy than the crystalline phases. This has been verified in studies on hafnia and zirconia nanoparticles: Ushakov et al. [22] found that crystallization temperature of hafnia particles increases by more than 300 C with an increase in surface/interface ratio of the amorphous phase. They have also reported stability crossovers with increasing surface area; thus, the amorphous structure becomes energetically favourable over other crystalline phases at high surface areas [23]. Interestingly, as-grown thicker YDZ films are crystalline. Figures 9 and 10 show plan view bright-field images and diffraction patterns, respectively, of an 82 nm thick YDZ film. Figures 9a 9d show the micrographs of as-grown nanocrystalline film, just prior to instantaneous grain growth, the onset of instantaneous grain growth and resultant 80 nm grains at 800 C. Corresponding diffraction patterns are shown in figures 10a 10d. The film is cubic and undergoes nearly instantaneous grain growth around 550 C. However, it is interesting to note that several intermediate phases formed during the process of grain growth and quickly disappeared. This is illustrated in figure 10, where arrows show the presence of monoclinic phases that appeared at the onset and disappeared following grain growth. This could possibly arise from segregation of yttrium atoms during the thermal treatment [24, 25]. It is interesting to note that instantaneous grain growth was observed in YDZ films, whereas normal grain growth was observed in pure zirconia films. In YDZ, Y atoms that segregate at grain boundaries strongly depress grain boundary mobility [26]. This suggests that the instantaneous grain growth seen at an elevated temperature (550 C) may be related to the mobility of Y atoms at the grain boundaries. 4. Conclusions We have shown that phase stability in nanoscale thin films of pure and Y-doped zirconia are dramatically different than in both bulk and nanoparticle forms.

11 5682 M. Tsuchiya et al. (a) 100nm (c) 100nm (b) 20nm 100nm Figure 9. Grain growth in 82 nm thick YDZ thin film grown on Ge: (a) at room temperature; (b) at 550 C just prior to instantaneous grain growth; (c) at the onset of instantaneous grain growth around 550 C; (d) at 800 C, where the average grain size is 80 nm. (d) c (400) c (420) c (311) c (200) c (331) c (220) c (222) c (111) Figure 10. Diffraction pattern for 82 nm thick YDZ film: (a) at room temperature; (b) onset of grain growth at 550 C; (c) during crystallization at 550 C; (d) at 600 C. Metastable monoclinic phases were found during grain growth that subsequently disappeared with further annealing, as indicated by arrows in (b) and (c).

12 Size-dependent phase transformations in nanoscale pure and Y-doped zirconia thin films 5683 Observed differences include the stability of the cubic phase in thin film pure zirconia at room temperature followed by cubic to tetragonal phase transformation at 400 C. Further, Y-doped zirconia that contains significant oxygen vacancies determined by extrinsic doping exhibits a room temperature phase that is either amorphous or cubic depending on the film thickness. This cubic phase remains stable at high temperatures even up to 800 C. However, annealing YDZ in an oxygen-rich environment leads to the formation of a tetragonal phase at 600 C. Oxygen vacancies in zirconia films, created either by the growth conditions, aliovalent doping or surface/grain boundaries effects, appear to play a significant role in determining the phase transformation and stability. It is well known that the different phases of zirconia possess different dielectric and ion transport properties. Thus, understanding phase stability in thin film zirconia systems is critical to tailoring microstructural features that may be useful for synthesis of fast-ion conducting solid electrolytes. Acknowledgements SR and MT are grateful to the School of Engineering and Applied Sciences, Harvard University, for financial support. AMM was supported by the Director, Office of Science, Office of Basic Energy Sciences, of the US Department of Energy under Contract No. DE-AC02-05CH Additionally, SR acknowledges the NCEM Visiting Scientist Fellowship that supported the microscopy done at NCEM, LBNL. References [1] A.H. Heuer and L.W. Hobbs (Editors), Advances in Ceramics, Vol. 3 (American Ceramic Society, Columbus, OH, 1981). [2] G.D. Wilk, R.M. Wallance and J.M. Anthony, J. Appl. Phys (2001). [3] H. Huang, M. Nakamura, P. Su, et al., J. Electrochem. Soc. 154 B20 (2007). [4] D. Ciuparu, A. Ensuque, G. Shafeev, et al., J. Mater. Sci. Lett (2000). [5] I.W. Chen and L.A. Xue, J. Am. Ceram. Soc (1990). [6] S. Shukla and S. Seal, Int. Mater. Rev (2005). [7] H.G. Scott, J. Mater. Sci (1975). [8] T. Kiguchi, N. Wakiya, K. Shiozaki, et al., J. Mater. Res (2005). [9] I. Kosacki, C.M. Rouleau, P.F. Becher, et al., Solid St. Ionics (2005). [10] A. Karthikeyan, C.L. Chang and S. Ramanathan, Appl. Phys. Lett (2006). [11] X.J. Chen, K.A. Khor, S.H. Chan, et al., Mater. Sci. Engng A (2002). [12] C.L. Chang, V. Shutthanandan, S.C. Singhal, et al., Proc. Mater. Res. Soc. Symp. (under review) (2007). [13] S.C. Moulzolf and R.J. Lad, J. Mater. Res (2000). [14] R.C. Garvie, J. Phys. Chem (1965). [15] N. Mommer, T. Lee and J.A. Gardner, J. Mater. Res (2000). [16] D.E. Collins and K.J. Bowman, J. Mater. Res (1998).

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