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Phase transformation in reaction-sintered Gd3Fe5O12–ZrO2 (II)
Materials Science and Engineering A297 (2001) 124 – 131 www.elsevier.com/locate/msea
Phase transformation in reaction-sintered Gd3Fe5O12 –ZrO2 (II) M.L. Wu, P. Shen *, D. Gan Institute of Materials Science and Engineering, National Sun Yat-sen Uni6ersity, 80424 Kaohsiung, Taiwan Received 16 March 2000; received in revised form 9 June 2000
Abstract Gadolinium iron garnet/zirconia (GIG/ZrO2) samples of 1:9 wt. ratio are reaction sintered at 1300°C for 8 h and studied by analytical electron microscopy. The GIG grain was found to transform into a Zr4 + -doped hematite core surrounded by Gd-stabilized cubic-ZrO2 polycrystals shell. The hematite core is a single crystal nearly free of Gd3 + . Iron zirconate was found to precipitate within the hematite core, and it is the same as that first found in the aged a-Fe2O3 – ZrO2 composite. The size of the iron zirconate precipitates decrease from core center toward the edge. The probable process leading to the microstructure and phase changes is discussed. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Gadolinium iron garnet; Zirconia; Hematite; Iron zirconate precipitate; AEM
1. Introduction Recently in a Zr4 + -oversaturated a-Fe2O3 system aged at 850°C in air, G.P. zones were found to nucleate homogeneously and then disk-like precipitates of iron zirconate were formed [1]. Since no compound is present in the equilibrium phase diagram, this precipitate must be a metastable phase. The composition, the habit plane and the possible precipitation mechanism were studied [1]. On the other hand, the introduction of chemical component through diffusion at a fixed temperature may also cause solute oversaturation and hence precipitation of phases. For example, this typically happened for the precipitation of various phases from b-NiAl in aluminized coating on nickel based superalloys [2,3] and oxyexsolution of oxide minerals in TiO2 –FeO– Fe2O3 system [4,5]. In this work, the isothermal reaction and interdiffusion of the gadolinium iron garnet (GIG) and ZrO2 system were studied via a reactive-sintering route. The phase transformations of GIG garnet and ZrO2 were investigated in detail. The same iron zirconate precipitate were found in the Zr-doped hematite formed by the * Corresponding author. Tel.: + 886-7-5252000, ext 4051; fax: +886-7-5254099. E-mail address: [email protected] (P. Shen).
phase transformation of GIG by interdiffusion. Such results shed light not only on the phase boundaries of the pseudoternary GIG–ZrO2 –Fe2O3 system at high temperature, but also on the interdiffusion process and the possible mechanisms of phase transformations of GIG and ZrO2. For examples: diffusion induced polygonization (DIP) [6], and diffusion induced dislocation migration (DIDM) [7]. Finally phase change to form core-shell structure in the GIG/ZrO2 diffusion couples with a curved rather than a flat interface is discussed
2. Experimental The GIG (Gd3Fe5O12) powder was prepared by mixing the stoichiometric powders of Gd2O3 (Cerac, 99.9%) and Fe2O3 (Cerac, 99.9%) and firing at 1300°C for 2 h in an open-air furnace. The ZrO2 (Cerac, 99.9%) powder was then mixed with GIG powders in 1:9 wt. ratio and die-pressed at 80 MPa to form pellets about 10 mm in diameter. Reactive sintering was conducted at 1300°C for 8 h in air and cooled in furnace. X-ray diffractometer (Cu Ka, 40 kV, 20 mA, Diano 8536) was used to detect major phases in sintered specimens. Scanning electron microscopy (SEM, using JEOL, JSM 6400 at 25 kV in secondary electron image (SEI) and back-scattered electron image (BEI) mode)
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was used to reveal the distribution of various phases in the composites. Thin-sections of the samples were Arion milled and then studied by analytical electron microscopy (AEM, using JEOL 3010 at 300 kV for imaging and 200 kV for energy dispersive X-ray (EDX) analysis. The K shell counts of Fe, and L shell counts of Zr and Gd without absorption correction were used for the determination of composition and the concentration profiles across the diffusion couples. The present EDX study was done on a semiquantitative basis because of electron beam broadening effect and thin-film effect, which interfered somehow the characterization of tiny precipitates in Fe2O3 − x grain as addressed in the interrelated paper [1]. (Methods employed for the analysis are clarified in [1].) Bright field image (BFI), dark field image (DFI), lattice image and selected area diffraction (SAD) pattern were used to identify phases.
3. Results
3.1. XRD and SEM XRD results indicated the presence of hematite, garnet and zirconia phases in the sample. SEI and BEI (Fig. 1a and b, respectively) results indicate the GIG grain has transformed into a dark core surrounded by a bright shell. The composition profiles based on pointcount EDX line scan across the AA% trace in Fig. 1b show that the dark core is Fe-rich and the bright shell is Zr-rich and contains more Gd than the core. Thermal etching at 1200°C for 10 min revealed more dark Fe-rich cores below the polished surface. Fig. 1. SEM (a) SEI, (b) BEI, and EDX profiles along the trace AA% across the hematite core (dark) and zirconia shell (bright).
Fig. 2. (a) TEM (BFI) of the hematite core with plate-like precipitates and c-ZrO2 polycrystals around it. Surrounding c-ZrO2 is the smallgrained m-ZrO2. (b) SAD pattern of hematite core tilted to [0001] zone axis. Three variants (labeled 1–3) of iron zirconate are present.
3.2. AEM The phase behavior observed and compiled in the following figures are typical to specific heating treatments adopted. Transmission electron microscopic (TEM) study showed the dark core is hematite with some plate-like precipitates, which are identified to be iron zirconate (Fig. 2a). Electron diffraction of the core showed the hematite and three variants of the precipitate (Fig. 2b). These iron zirconate precipitates have the same crystal structure and about the same composition as that precipitated in Zr-oversaturated a-Fe2O3 in Fe2O3 –ZrO2 system [1]. The iron zirconate precipitates are generally larger in size near the center of hematite core (Fig. 2a). They also have the same habit plane (steps following {1014} of hematite host) and ledgegrowth behavior. It is noteworthy that the precipitate has rather coherent steps, but the lattice fringes of hematite are slightly distorted near the incoherent ledge front (Fig. 3). The Zr-rich shell surrounding the hematite core is identified to be cubic ZrO2 (c-ZrO2) by electron diffrac-
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Point count EDX analysis results of various phases are compiled in Table 1. The m-ZrO2 particles were found to contain ca. 2 at.% Fe and 1 at.% Gd, while the c-ZrO2 shell has a much higher solid solubility of Fe and Gd, nearly 7 and 13 at.%, respectively. The hematite matrix free of iron zirconate precipitate contains nearly 1 at.% Zr and negligible Gd. The iron zirconate is nearly 1:1 in the atomic ratio of Fe:Zr, similar to previous results in aged Fe2O3 –ZrO2 composite [1].
4. Discussion
4.1. Hematite core and c-ZrO2 shell
Fig. 3. Lattice image of iron zirconate showing ledges (arrow). The hematite lattice fringes near incoherent ledge front are slightly disturbed.
tion (Fig. 4). The c-ZrO2 shell is free of twins and any precipitates. Further, SAD patterns in [111], [233] and [011] zone axes (Fig. 4a – c, respectively) showed diffuse diffraction intensities, indicating the c-ZrO2 has defect clusters and/or ordering as discussed in Section 4.2. Occasionally, diffraction spots of tetragonal-ZrO2 (tZrO2) were found to be rather well developed (Fig. 4d). Outside the c-ZrO2 shell are the moniclinic-ZrO2 (mZrO2) grains showing the characteristic twin feature (Fig. 2a). The c-ZrO2 shell actually consists of faceted grains with dislocations and grows at the expense of the outer much smaller m-ZrO2 grains, as indicated by the outward concave interface (Fig. 2a). In the process, some ZrO2 particles are occasionally incorporated into the c-ZrO2 grain (Fig. 5). The c-ZrO2 shell also showed dislocation arrays and cracks along the c/m-ZrO2 interface (Fig. 6). It is noteworthy that the m-ZrO2 phase in this sample shows a specific {111} twin boundary (Fig. 7), in contrast to the mosaic {100} and {110} twin variants common to other partially stabilized zirconia (PSZ) systems.
The GIG/ZrO2 specimen reaction-sintered at 1300°C for 8 h showed that the GIG grain transformed into a single crystal Zr4 + -doped hematite core surrounded by Gd-stabilized c-ZrO2 polycrystals. The iron zirconate precipitated within the hematite core has the same crystal structure, composition, shape and growth habit plane as that in a supercooled Zr-doped a-Fe2O3 after aging [1]. The precipitates are however larger near the core center and decrease in size toward the edge with a precipitate free zone (PFZ) at the edge. Since this phase is not present in the equilibrium phase diagram [8], it must be an intermediate metastable phase with lower activation energy barrier than its end members, i.e. ZrO2 and hematite. The iron zirconate that precipitates in the hematite is formed during isothermal reaction-sintering process rather than during the cooling of the specimen, otherwise it should be of approximate the same size and density inside hematite. The larger precipitates near the core center is formed first so that they have time to grow, while the ones near the edge are formed later and are smaller in size. In the PFZ at the edge the precipitates are not formed yet, as previous results indicate that an incubation time is necessary for its nucleation [1]. Therefore, from the size and distribution of iron zirconate it can be deduced that the hematite core starts growing from the center toward the edge. The solubility of ZrO2 in hematite is low, about 3 wt.% at 1300°C [8]. Since the precipitation occurs during reaction sintering, the extra amount of Zr must already be there when hematite first forms and any extra amount precipitates as iron zirconate. It therefore occurs by the diffusion of Zr into GIG after the start of reaction sintering but before the formation of hematite, as when it is formed the Zr content is already supersaturated. The diffusion rate of Zr in GIG at 1300°C, however, is not available to support this point. After reaction sintering starts, in addition to Zr diffusion, Gd and Fe start diffusion into ZrO2, which at 1300°C is originally tetragonal in structure [9]. However, since the solubility of Fe2O3 in t-ZrO2 is also low,
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Fig. 4. SAD patterns of the c-ZrO2 shell in (a) [111], (b) [233], (c) [011] zone axes, with diffuse diffraction. Occasionally t-ZrO2 diffraction spots are shown, as in (d) of [011] zone axis.
Fig. 5. (a) TEM (DFI) and (b) SAD pattern (Z= [011]) of a twinned m-ZrO2 particle incorporated in c-ZrO2 grain.
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M.L. Wu et al. / Materials Science and Engineering A297 (2001) 124–131 Table 1 Composition (at.%) of phases in reaction-sintered Gd3Fe5O12–ZrO2 Sample
Hematite Iron zirconate c-ZrO2 m-ZrO2
Fig. 6. TEM (BFI) showing dislocation arrays in c-ZrO2 shell (Z = [011]) near the c-/m-ZrO2 boundary.
about 3 wt.% at 1300°C [8], not much Fe can diffuse into ZrO2. However, although the solubility of Gd2O3 in t-ZrO2 is low, less than 3 at.% [9 – 11], Gd can stabilize t-ZrO2 into c-ZrO2 and the solubility of Gd2O3 in c-ZrO2 increases quickly, in the range from approximately 8 at.% to more than 30 at.% till Gd2Zr2O7 forms
Phase Fe
Zr
Gd
98.6 50 6.7 1.7
1.4 50 80.0 97.4
– – 13.3 0.9
[9–11]. Therefore significant amount of Gd then diffuse into ZrO2, as evidenced by the large c-ZrO2 shell around hematite core (Figs. 1 and 2). In Figs. 1 and 2, a sharp boundary between c-ZrO2 and m-ZrO2 is noticed. (In Fig. 1 the c-ZrO2 shell appears bright due to the large nucleus of Gd atoms.) It is the original boundary between c-ZrO2 and t-ZrO2 at 1300°C, i.e. the diffusion front of Gd, as the solubility drops significantly across the interface. This is also confirmed by the EDX analysis data in Table 1. Upon cooling, the Gd-stabilized c-ZrO2 maintains its structure to room temperature but the t-ZrO2 is transformed into m-ZrO2, as shown in Fig. 2a. The c-ZrO2 grains are further polygonized into subgrains separated by misfit dislocation arrays to accommodate the composition difference. In this regard, chemical zoning of a phase in the diffusion couples may be due to diffusion induced polygonization [6,7,12,13] and diffusion induced dislocation migration [14]. Both phenomena have been verified to be valid for retrograde zoning of sili-
Fig. 7. TEM images of m-ZrO2 grain with a specific {111} twin boundary: (a) BFI, (b) SAD pattern (Z=[011]) with diffraction streaks in [111] direction.
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cate garnet formed at relatively low temperature for effective short-circuit diffusion [15]. The formation of hematite core is an interesting question. The original GIG grains are ca. 2 mm in diameter according to SEM observations of the slightly sintered GIG polycrystals (not shown) and are converted into core-shell structure of comparable size (Figs. 1 and 2). The hematite core is believed to first start forming near the core center by the following reasons. First, the size and distribution of iron zirconate precipitate strongly suggests it, as discussed previously. Second, there is the symmetry nature of the core-shell structure around the core center. And, third, the hematite is a single crystal. The hematite core may not be formed from several nuclei, near edge or at center, and then grow into a single crystal, as the size and distribution of iron zirconate precipitate must be different. Once hematite is formed near core center, the diffusion picture becomes different. The Zr content is supersaturated in hematite and results in the precipitation of iron zirconate. However, since hematite and Gd2O3 are nearly completely mutual insoluble [16], Gd must be nearly completely expelled from the hematite core, as confirmed in EDX analysis in Table 1. With the large solubility of Gd in c-ZrO2, the process continues till nearly all Gd diffuses into surrounding c-ZrO2, with a smaller hematite core remains in the center. Why the hematite core starts forming near the core center is puzzling. The diffusion data are scarce in this system. The diffusion of Gd in GIG and c-ZrO2 can be relatively easy while a large Gd concentration drop occurs across the c-ZrO2/t-ZrO2 interface. Therefore Gd concentration gradient may be low in the GIG and c-ZrO2 grains, the latter case is confirmed in the EDX analysis of c-ZrO2 grains. However, it is not clear what triggers the decomposition of GIG near core center. Future study of GIG/ZrO2 diffusion couples with a flat rather than a curved interface would prove worthwhile to clarify this point.
4.2. Stabilization of c-ZrO2 The (Gd,Fe)-codissolved zirconia (typically 13 at.% Gd, 7 at.% Fe and 80 at.% Zr) was found to stabilize as c-fluorite structure. The effective ion radius is 0.078 nm for Fe3 + in coordination number (C.N.) 8 and 0.084 nm for Zr4 + in C.N. 8 [17]. In Fe3 + -dissolved ZrO2 the undersized Fe3 + cannot stabilize the ZrO2 upon cooling to room temperature, unless under the influence of the volume constraint of a host [18,19]. Since Gd3 + and Fe3 + are both trivalent, the charge-compensating oxygen vacancy effect must be the same. However, the much larger Gd3 + ion (0.1053 nm in C.N. 8 [17]) in ZrO2 can effectively increase the cation/anion radii ratio and stabilize the higher C.N. structure of c-ZrO2
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to room temperature [20], as shown in the corresponding phase diagram [9–11]. Local atomic structure of Zr and dopant cations in zirconia solid solutions with Fe2O3 and Gd2O3 have been determined by X-ray absorption spectroscopy using synchrotron radiation light source [21]. The dopant cations were found to substitute for Zr ions and their sizes determine the preferred locations of oxygen vacancies. It was also found that the oversized dopant Gd3 + ions are located as nearest neighbors to Zr atoms, leaving 8-fold oxygen coordination to dopant cations; while the undersized dopant Fe3 + ion competes with Zr ions for the oxygen vacancies in zirconia, resulting in 6-fold oxygen coordination and a large disturbance to the surrounding next nearest neighbors. Therefore the oversized Gd3 + dopant is more effective than undersized Fe3 + dopant in stabilizing the c-ZrO2, as indeed observed in the present study. Defect clusters and/or ordering for (Gd,Fe)-codissolved zirconia are indicated by its significant diffuse diffraction intensities (Fig. 4). This diffraction phenomenon is common to anion-deficient fully stabilized zirconia (FSZ) of c-ZrO2 and could be due to short range ordering of defects [22]. Detailed diffuse X-ray scattering measurements recently made of yttria-stabilized c-ZrO2 (YSZ) [23,24] and a modulation wave approach [25] suggested that vacancy pairs oriented along a given 111 direction repel each other, but do not interact strongly with unlike-oriented vacancy pairs. This appears to be able to account satisfactorily for practically all of the diffuse scattering features observed in YSZ [23,24]. It is probable that this model can be extended to Gd-stabilized c-ZrO2 as the Gd3 + (0.1053 nm) is a little larger than Y3 + (0.1019 nm) in C.N. 8 [17].
4.3. Twin boundaries in m-ZrO2 The t-ZrO2 transforms to m-ZrO2 at ca. 1170°C by a martensitic transformation. The m-ZrO2 thus formed has twin variants, which were characterized [20,26–29] to be {100} and/or {110} type mosaic twins for pure ZrO2 and a number of partially stabilized zirconia (PSZ). The m-ZrO2 in the present GIG/ZrO2 samples, surprisingly, was found to have specific {111} twin boundary. Crystallographically, there are three possible choices of lattice correspondences, i.e. a-, b-, or c-axis of m-ZrO2 to be parallel to the c-axis of t-ZrO2. The resultant three habit planes of twinning on (100)m, (001)m, and (110)m were observed [27]. The orientation relationship established between the m- and t-ZrO2 phases are of the general form (100)m//{100}t and [100]m//100t [20,26,29,30], where t-ZrO2 is indexed as face-centered tetragonal cell. The habit plane has been suggested to be (671)m and (761)m for lenticular-type
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product and (100)m for plate-shaped product [28,29]. However, none with {111} twin plane was reported in contrast to the present case. It may be possible that the point defects caused by dissolution of Fe3 + and Gd3 + ion in t-ZrO2 form clusters to weaken {111} plane of the t-ZrO2 phase, a distorted fluorite lattice with cations more or less in f.c.c.-type array, making it feasible for crystallographic shear to occur along {111}. However, this remains to be studied.
4.4. Fe2O3 –ZrO2 – GIG pseudoternary system To study the Fe2O3 – ZrO2 – GIG pseudoternary system at 1300°C, a specimen with GIG/ZrO2 ratio of 9:1 is prepared under the same condition, i.e. reaction sintered at 1300°C for 8 h in air and furnace cooled. This sample showed the same phase change and microstructure characteristics as the previous sample except that more garnet survived the reaction. Thus at 1300°C, the a-Fe2O3 – ZrO2 – GIG pseudoternary system is divided into H+ t, H+c + t; H+c; and H+c + GIG regions, where H, c, and t denote hematite, c- and t-ZrO2, respectively (Fig. 8) and iron zirconate exsolved from Zr-oversaturated hematite lattice. There is no liquid phase for the a-Fe2O3 – ZrO2 – GIG pseudoternary system at 1300°C. The eutectic temperatures are 1525°C, \ 2000°C (probably) and 1500°C for the binary iron oxide – ZrO2 [8], Gd2O3 – ZrO2 [9 –11] and Gd2O3 –Fe2O3 [16], respectively. The phase assemblages of the two diffusion couples can then be rationalized by specific diffusion path and bulk compositions as shown in Fig. 8. In the above discussion, local equilibrium was assumed at incoherent interfaces between c- and t-ZrO2, between hematite and c-ZrO2, and between hematite and iron zirconate plate at its edge.
Fig. 8. Fe2O3 – ZrO2 –GIG pseudoternary isothermal section (mol.%) at 1300°C. Point I is the specimen with GIG/ZrO2 weight ratio of 1:9 and point II of 9:1. The arrows indicate the diffusion path, and H denotes hematite.
5. Conclusions In a GIG/ZrO2 sample reaction-sintered at 1300°C, the GIG grain is transformed into a single crystal hematite core with iron zirconate precipitates and a Gd-stabilized c-ZrO2 polycrystals shell around it. The iron zirconate precipitated in the hematite core has the same crystal structure, composition, shape and habit plane as that in Zr-supersaturated a-Fe2O3 after aging. The precipitates decrease in size from core center to edge with a PFZ near the edge. The dissolution of Gd3 + can effectively stabilize the c-ZrO2. The (Gd,Fe)-codoped m-ZrO2 has a specific (111) twin boundary in contrast to mosaic {100} and {110} in other PSZ’s. At 1300°C, the a-Fe2O3 –ZrO2 –GIG pseudoternary system is divided into H+ t, H+ c+ t; H+c; and H+c +GIG regions, where H, c, and t denote hematite, c-ZrO2 and t-ZrO2, respectively.
References [1] M.L. Wu, D. Gan, P. Shen, Mater. Sci. Eng. A (2001) in press. [2] A. Chien, Ph.D. Thesis, National Sun Yat-Sen University, Taiwan, 1995. [3] A. Chien, D. Gan, P. Shen, Mater. Sci. Eng. A206 (1996) 215. [4] A.E. Goresy, S.E. Haggerty, J.S. Huebner, D.H. Lindsley, D. Rumble, III, in: D. Rumble, III (Ed.), Oxide Minerals, Mineralogical Society of America Short Course Notes, Southern, Blacksburg, VA, 1976. [5] A. Putnis, J.D.C. McConnell, Principles of Mineral Behavior, Elsevier, New York, 1980. [6] V.Y. Doo, R.W. Balluffi, Acta Metall. 6 (1958) 428. [7] F.J.A. Den Broeder, S. Nakahara, Scripta Metall. 17 (1983) 399. [8] T.S. Jones, S. Kimura, A. Muan, J. Am. Ceram. Soc. 50 (1967) 137. [9] J. Lefe`vre, Ann. Chim. 8 (1963) 128. [10] M. Perez, Y. Jorba, Ann. Chim. 7 (1962) 509. [11] A. Rouanet, Rev. Int. Hautes Temp. Refract. 8 (1971) 161. [12] C.R.M. Grovenor, D.A. Smith, M.J. Goringe, Thin Solid Films 74 (1980) 257. [13] C.R.M. Grovenor, Acta Metall. 33 (1985) 579. [14] C.P. Chang, M.H. Loretto, Acta. Metall. 36 (1988) 805. [15] S.L. Hwang, P. Shen, T.F. Yui, H.T. Chu, Proceedings of the Annual Meeting, Geological Society of China, Taipei, Taiwan, 1999. [16] I. Warshaw, R. Roy, J. Am. Ceram. Soc. 42 (1959) 434. [17] R.D. Shannon, Acta Crystallogr. A32 (1976) 751. [18] M. Ru¨hle, A.H. Heuer, Adv. Ceram. 12 (1984) 14. [19] P. Shen, S. Chen, W.S. Chang, Mater. Sci. Eng. A184 (1994) L5. [20] D.J. Green, R.H.J. Hannink, M.V. Swain, Transformation Toughening of Ceramics, CRC Press, Boca Raton, FL, 1989 references cited therein. [21] P. Li, I.W. Chen, J.E. Penner-Hanh, J. Am. Ceram. Soc. 77 (1994) 118. [22] R. Chain, D.G. Brandon, Adv. Ceram. 12 (1983) 86–95. [23] T.R. Welberry, R.L. Withers, J.G. Thompson, B.D. Butler, J. Solid State Chem. 100 (1992) 71. [24] T.R. Welberry, B.D. Butler, J.G. Thompson, R.L. Withers, J. Solid State Chem. 106 (1993) 461.
M.L. Wu et al. / Materials Science and Engineering A297 (2001) 124–131 [25] T.R. Welberry, R.L. Withers, S.C. Mayo, J. Solid State Chem. 115 (1995) 43. [26] G.M. Wolten, J. Am. Ceram. Soc. 46 (9) (1963) 418. [27] J.E. Bailey, Proc. R. Soc. London Ser. A 279 (1378) (1964) 395.
.
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[28] G.K. Bansal, A.H. Heuer, Acta Metall. 20 (11) (1972) 1281. [29] G.K. Bansal, A.H. Heuer, Acta Metall. 22 (4) (1974) 409. [30] D.K. Smith, H.W. Newkirk, Acta Crystallogr. 18 (6) (1965) 983.
Phase transformation in reaction-sintered Gd3Fe5O12 –ZrO2 (II) M.L. Wu, P. Shen *, D. Gan Institute of Materials Science and Engineering, National Sun Yat-sen Uni6ersity, 80424 Kaohsiung, Taiwan Received 16 March 2000; received in revised form 9 June 2000
Abstract Gadolinium iron garnet/zirconia (GIG/ZrO2) samples of 1:9 wt. ratio are reaction sintered at 1300°C for 8 h and studied by analytical electron microscopy. The GIG grain was found to transform into a Zr4 + -doped hematite core surrounded by Gd-stabilized cubic-ZrO2 polycrystals shell. The hematite core is a single crystal nearly free of Gd3 + . Iron zirconate was found to precipitate within the hematite core, and it is the same as that first found in the aged a-Fe2O3 – ZrO2 composite. The size of the iron zirconate precipitates decrease from core center toward the edge. The probable process leading to the microstructure and phase changes is discussed. © 2001 Elsevier Science B.V. All rights reserved. Keywords: Gadolinium iron garnet; Zirconia; Hematite; Iron zirconate precipitate; AEM
1. Introduction Recently in a Zr4 + -oversaturated a-Fe2O3 system aged at 850°C in air, G.P. zones were found to nucleate homogeneously and then disk-like precipitates of iron zirconate were formed [1]. Since no compound is present in the equilibrium phase diagram, this precipitate must be a metastable phase. The composition, the habit plane and the possible precipitation mechanism were studied [1]. On the other hand, the introduction of chemical component through diffusion at a fixed temperature may also cause solute oversaturation and hence precipitation of phases. For example, this typically happened for the precipitation of various phases from b-NiAl in aluminized coating on nickel based superalloys [2,3] and oxyexsolution of oxide minerals in TiO2 –FeO– Fe2O3 system [4,5]. In this work, the isothermal reaction and interdiffusion of the gadolinium iron garnet (GIG) and ZrO2 system were studied via a reactive-sintering route. The phase transformations of GIG garnet and ZrO2 were investigated in detail. The same iron zirconate precipitate were found in the Zr-doped hematite formed by the * Corresponding author. Tel.: + 886-7-5252000, ext 4051; fax: +886-7-5254099. E-mail address: [email protected] (P. Shen).
phase transformation of GIG by interdiffusion. Such results shed light not only on the phase boundaries of the pseudoternary GIG–ZrO2 –Fe2O3 system at high temperature, but also on the interdiffusion process and the possible mechanisms of phase transformations of GIG and ZrO2. For examples: diffusion induced polygonization (DIP) [6], and diffusion induced dislocation migration (DIDM) [7]. Finally phase change to form core-shell structure in the GIG/ZrO2 diffusion couples with a curved rather than a flat interface is discussed
2. Experimental The GIG (Gd3Fe5O12) powder was prepared by mixing the stoichiometric powders of Gd2O3 (Cerac, 99.9%) and Fe2O3 (Cerac, 99.9%) and firing at 1300°C for 2 h in an open-air furnace. The ZrO2 (Cerac, 99.9%) powder was then mixed with GIG powders in 1:9 wt. ratio and die-pressed at 80 MPa to form pellets about 10 mm in diameter. Reactive sintering was conducted at 1300°C for 8 h in air and cooled in furnace. X-ray diffractometer (Cu Ka, 40 kV, 20 mA, Diano 8536) was used to detect major phases in sintered specimens. Scanning electron microscopy (SEM, using JEOL, JSM 6400 at 25 kV in secondary electron image (SEI) and back-scattered electron image (BEI) mode)
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was used to reveal the distribution of various phases in the composites. Thin-sections of the samples were Arion milled and then studied by analytical electron microscopy (AEM, using JEOL 3010 at 300 kV for imaging and 200 kV for energy dispersive X-ray (EDX) analysis. The K shell counts of Fe, and L shell counts of Zr and Gd without absorption correction were used for the determination of composition and the concentration profiles across the diffusion couples. The present EDX study was done on a semiquantitative basis because of electron beam broadening effect and thin-film effect, which interfered somehow the characterization of tiny precipitates in Fe2O3 − x grain as addressed in the interrelated paper [1]. (Methods employed for the analysis are clarified in [1].) Bright field image (BFI), dark field image (DFI), lattice image and selected area diffraction (SAD) pattern were used to identify phases.
3. Results
3.1. XRD and SEM XRD results indicated the presence of hematite, garnet and zirconia phases in the sample. SEI and BEI (Fig. 1a and b, respectively) results indicate the GIG grain has transformed into a dark core surrounded by a bright shell. The composition profiles based on pointcount EDX line scan across the AA% trace in Fig. 1b show that the dark core is Fe-rich and the bright shell is Zr-rich and contains more Gd than the core. Thermal etching at 1200°C for 10 min revealed more dark Fe-rich cores below the polished surface. Fig. 1. SEM (a) SEI, (b) BEI, and EDX profiles along the trace AA% across the hematite core (dark) and zirconia shell (bright).
Fig. 2. (a) TEM (BFI) of the hematite core with plate-like precipitates and c-ZrO2 polycrystals around it. Surrounding c-ZrO2 is the smallgrained m-ZrO2. (b) SAD pattern of hematite core tilted to [0001] zone axis. Three variants (labeled 1–3) of iron zirconate are present.
3.2. AEM The phase behavior observed and compiled in the following figures are typical to specific heating treatments adopted. Transmission electron microscopic (TEM) study showed the dark core is hematite with some plate-like precipitates, which are identified to be iron zirconate (Fig. 2a). Electron diffraction of the core showed the hematite and three variants of the precipitate (Fig. 2b). These iron zirconate precipitates have the same crystal structure and about the same composition as that precipitated in Zr-oversaturated a-Fe2O3 in Fe2O3 –ZrO2 system [1]. The iron zirconate precipitates are generally larger in size near the center of hematite core (Fig. 2a). They also have the same habit plane (steps following {1014} of hematite host) and ledgegrowth behavior. It is noteworthy that the precipitate has rather coherent steps, but the lattice fringes of hematite are slightly distorted near the incoherent ledge front (Fig. 3). The Zr-rich shell surrounding the hematite core is identified to be cubic ZrO2 (c-ZrO2) by electron diffrac-
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Point count EDX analysis results of various phases are compiled in Table 1. The m-ZrO2 particles were found to contain ca. 2 at.% Fe and 1 at.% Gd, while the c-ZrO2 shell has a much higher solid solubility of Fe and Gd, nearly 7 and 13 at.%, respectively. The hematite matrix free of iron zirconate precipitate contains nearly 1 at.% Zr and negligible Gd. The iron zirconate is nearly 1:1 in the atomic ratio of Fe:Zr, similar to previous results in aged Fe2O3 –ZrO2 composite [1].
4. Discussion
4.1. Hematite core and c-ZrO2 shell
Fig. 3. Lattice image of iron zirconate showing ledges (arrow). The hematite lattice fringes near incoherent ledge front are slightly disturbed.
tion (Fig. 4). The c-ZrO2 shell is free of twins and any precipitates. Further, SAD patterns in [111], [233] and [011] zone axes (Fig. 4a – c, respectively) showed diffuse diffraction intensities, indicating the c-ZrO2 has defect clusters and/or ordering as discussed in Section 4.2. Occasionally, diffraction spots of tetragonal-ZrO2 (tZrO2) were found to be rather well developed (Fig. 4d). Outside the c-ZrO2 shell are the moniclinic-ZrO2 (mZrO2) grains showing the characteristic twin feature (Fig. 2a). The c-ZrO2 shell actually consists of faceted grains with dislocations and grows at the expense of the outer much smaller m-ZrO2 grains, as indicated by the outward concave interface (Fig. 2a). In the process, some ZrO2 particles are occasionally incorporated into the c-ZrO2 grain (Fig. 5). The c-ZrO2 shell also showed dislocation arrays and cracks along the c/m-ZrO2 interface (Fig. 6). It is noteworthy that the m-ZrO2 phase in this sample shows a specific {111} twin boundary (Fig. 7), in contrast to the mosaic {100} and {110} twin variants common to other partially stabilized zirconia (PSZ) systems.
The GIG/ZrO2 specimen reaction-sintered at 1300°C for 8 h showed that the GIG grain transformed into a single crystal Zr4 + -doped hematite core surrounded by Gd-stabilized c-ZrO2 polycrystals. The iron zirconate precipitated within the hematite core has the same crystal structure, composition, shape and growth habit plane as that in a supercooled Zr-doped a-Fe2O3 after aging [1]. The precipitates are however larger near the core center and decrease in size toward the edge with a precipitate free zone (PFZ) at the edge. Since this phase is not present in the equilibrium phase diagram [8], it must be an intermediate metastable phase with lower activation energy barrier than its end members, i.e. ZrO2 and hematite. The iron zirconate that precipitates in the hematite is formed during isothermal reaction-sintering process rather than during the cooling of the specimen, otherwise it should be of approximate the same size and density inside hematite. The larger precipitates near the core center is formed first so that they have time to grow, while the ones near the edge are formed later and are smaller in size. In the PFZ at the edge the precipitates are not formed yet, as previous results indicate that an incubation time is necessary for its nucleation [1]. Therefore, from the size and distribution of iron zirconate it can be deduced that the hematite core starts growing from the center toward the edge. The solubility of ZrO2 in hematite is low, about 3 wt.% at 1300°C [8]. Since the precipitation occurs during reaction sintering, the extra amount of Zr must already be there when hematite first forms and any extra amount precipitates as iron zirconate. It therefore occurs by the diffusion of Zr into GIG after the start of reaction sintering but before the formation of hematite, as when it is formed the Zr content is already supersaturated. The diffusion rate of Zr in GIG at 1300°C, however, is not available to support this point. After reaction sintering starts, in addition to Zr diffusion, Gd and Fe start diffusion into ZrO2, which at 1300°C is originally tetragonal in structure [9]. However, since the solubility of Fe2O3 in t-ZrO2 is also low,
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Fig. 4. SAD patterns of the c-ZrO2 shell in (a) [111], (b) [233], (c) [011] zone axes, with diffuse diffraction. Occasionally t-ZrO2 diffraction spots are shown, as in (d) of [011] zone axis.
Fig. 5. (a) TEM (DFI) and (b) SAD pattern (Z= [011]) of a twinned m-ZrO2 particle incorporated in c-ZrO2 grain.
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M.L. Wu et al. / Materials Science and Engineering A297 (2001) 124–131 Table 1 Composition (at.%) of phases in reaction-sintered Gd3Fe5O12–ZrO2 Sample
Hematite Iron zirconate c-ZrO2 m-ZrO2
Fig. 6. TEM (BFI) showing dislocation arrays in c-ZrO2 shell (Z = [011]) near the c-/m-ZrO2 boundary.
about 3 wt.% at 1300°C [8], not much Fe can diffuse into ZrO2. However, although the solubility of Gd2O3 in t-ZrO2 is low, less than 3 at.% [9 – 11], Gd can stabilize t-ZrO2 into c-ZrO2 and the solubility of Gd2O3 in c-ZrO2 increases quickly, in the range from approximately 8 at.% to more than 30 at.% till Gd2Zr2O7 forms
Phase Fe
Zr
Gd
98.6 50 6.7 1.7
1.4 50 80.0 97.4
– – 13.3 0.9
[9–11]. Therefore significant amount of Gd then diffuse into ZrO2, as evidenced by the large c-ZrO2 shell around hematite core (Figs. 1 and 2). In Figs. 1 and 2, a sharp boundary between c-ZrO2 and m-ZrO2 is noticed. (In Fig. 1 the c-ZrO2 shell appears bright due to the large nucleus of Gd atoms.) It is the original boundary between c-ZrO2 and t-ZrO2 at 1300°C, i.e. the diffusion front of Gd, as the solubility drops significantly across the interface. This is also confirmed by the EDX analysis data in Table 1. Upon cooling, the Gd-stabilized c-ZrO2 maintains its structure to room temperature but the t-ZrO2 is transformed into m-ZrO2, as shown in Fig. 2a. The c-ZrO2 grains are further polygonized into subgrains separated by misfit dislocation arrays to accommodate the composition difference. In this regard, chemical zoning of a phase in the diffusion couples may be due to diffusion induced polygonization [6,7,12,13] and diffusion induced dislocation migration [14]. Both phenomena have been verified to be valid for retrograde zoning of sili-
Fig. 7. TEM images of m-ZrO2 grain with a specific {111} twin boundary: (a) BFI, (b) SAD pattern (Z=[011]) with diffraction streaks in [111] direction.
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cate garnet formed at relatively low temperature for effective short-circuit diffusion [15]. The formation of hematite core is an interesting question. The original GIG grains are ca. 2 mm in diameter according to SEM observations of the slightly sintered GIG polycrystals (not shown) and are converted into core-shell structure of comparable size (Figs. 1 and 2). The hematite core is believed to first start forming near the core center by the following reasons. First, the size and distribution of iron zirconate precipitate strongly suggests it, as discussed previously. Second, there is the symmetry nature of the core-shell structure around the core center. And, third, the hematite is a single crystal. The hematite core may not be formed from several nuclei, near edge or at center, and then grow into a single crystal, as the size and distribution of iron zirconate precipitate must be different. Once hematite is formed near core center, the diffusion picture becomes different. The Zr content is supersaturated in hematite and results in the precipitation of iron zirconate. However, since hematite and Gd2O3 are nearly completely mutual insoluble [16], Gd must be nearly completely expelled from the hematite core, as confirmed in EDX analysis in Table 1. With the large solubility of Gd in c-ZrO2, the process continues till nearly all Gd diffuses into surrounding c-ZrO2, with a smaller hematite core remains in the center. Why the hematite core starts forming near the core center is puzzling. The diffusion data are scarce in this system. The diffusion of Gd in GIG and c-ZrO2 can be relatively easy while a large Gd concentration drop occurs across the c-ZrO2/t-ZrO2 interface. Therefore Gd concentration gradient may be low in the GIG and c-ZrO2 grains, the latter case is confirmed in the EDX analysis of c-ZrO2 grains. However, it is not clear what triggers the decomposition of GIG near core center. Future study of GIG/ZrO2 diffusion couples with a flat rather than a curved interface would prove worthwhile to clarify this point.
4.2. Stabilization of c-ZrO2 The (Gd,Fe)-codissolved zirconia (typically 13 at.% Gd, 7 at.% Fe and 80 at.% Zr) was found to stabilize as c-fluorite structure. The effective ion radius is 0.078 nm for Fe3 + in coordination number (C.N.) 8 and 0.084 nm for Zr4 + in C.N. 8 [17]. In Fe3 + -dissolved ZrO2 the undersized Fe3 + cannot stabilize the ZrO2 upon cooling to room temperature, unless under the influence of the volume constraint of a host [18,19]. Since Gd3 + and Fe3 + are both trivalent, the charge-compensating oxygen vacancy effect must be the same. However, the much larger Gd3 + ion (0.1053 nm in C.N. 8 [17]) in ZrO2 can effectively increase the cation/anion radii ratio and stabilize the higher C.N. structure of c-ZrO2
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to room temperature [20], as shown in the corresponding phase diagram [9–11]. Local atomic structure of Zr and dopant cations in zirconia solid solutions with Fe2O3 and Gd2O3 have been determined by X-ray absorption spectroscopy using synchrotron radiation light source [21]. The dopant cations were found to substitute for Zr ions and their sizes determine the preferred locations of oxygen vacancies. It was also found that the oversized dopant Gd3 + ions are located as nearest neighbors to Zr atoms, leaving 8-fold oxygen coordination to dopant cations; while the undersized dopant Fe3 + ion competes with Zr ions for the oxygen vacancies in zirconia, resulting in 6-fold oxygen coordination and a large disturbance to the surrounding next nearest neighbors. Therefore the oversized Gd3 + dopant is more effective than undersized Fe3 + dopant in stabilizing the c-ZrO2, as indeed observed in the present study. Defect clusters and/or ordering for (Gd,Fe)-codissolved zirconia are indicated by its significant diffuse diffraction intensities (Fig. 4). This diffraction phenomenon is common to anion-deficient fully stabilized zirconia (FSZ) of c-ZrO2 and could be due to short range ordering of defects [22]. Detailed diffuse X-ray scattering measurements recently made of yttria-stabilized c-ZrO2 (YSZ) [23,24] and a modulation wave approach [25] suggested that vacancy pairs oriented along a given 111 direction repel each other, but do not interact strongly with unlike-oriented vacancy pairs. This appears to be able to account satisfactorily for practically all of the diffuse scattering features observed in YSZ [23,24]. It is probable that this model can be extended to Gd-stabilized c-ZrO2 as the Gd3 + (0.1053 nm) is a little larger than Y3 + (0.1019 nm) in C.N. 8 [17].
4.3. Twin boundaries in m-ZrO2 The t-ZrO2 transforms to m-ZrO2 at ca. 1170°C by a martensitic transformation. The m-ZrO2 thus formed has twin variants, which were characterized [20,26–29] to be {100} and/or {110} type mosaic twins for pure ZrO2 and a number of partially stabilized zirconia (PSZ). The m-ZrO2 in the present GIG/ZrO2 samples, surprisingly, was found to have specific {111} twin boundary. Crystallographically, there are three possible choices of lattice correspondences, i.e. a-, b-, or c-axis of m-ZrO2 to be parallel to the c-axis of t-ZrO2. The resultant three habit planes of twinning on (100)m, (001)m, and (110)m were observed [27]. The orientation relationship established between the m- and t-ZrO2 phases are of the general form (100)m//{100}t and [100]m//100t [20,26,29,30], where t-ZrO2 is indexed as face-centered tetragonal cell. The habit plane has been suggested to be (671)m and (761)m for lenticular-type
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product and (100)m for plate-shaped product [28,29]. However, none with {111} twin plane was reported in contrast to the present case. It may be possible that the point defects caused by dissolution of Fe3 + and Gd3 + ion in t-ZrO2 form clusters to weaken {111} plane of the t-ZrO2 phase, a distorted fluorite lattice with cations more or less in f.c.c.-type array, making it feasible for crystallographic shear to occur along {111}. However, this remains to be studied.
4.4. Fe2O3 –ZrO2 – GIG pseudoternary system To study the Fe2O3 – ZrO2 – GIG pseudoternary system at 1300°C, a specimen with GIG/ZrO2 ratio of 9:1 is prepared under the same condition, i.e. reaction sintered at 1300°C for 8 h in air and furnace cooled. This sample showed the same phase change and microstructure characteristics as the previous sample except that more garnet survived the reaction. Thus at 1300°C, the a-Fe2O3 – ZrO2 – GIG pseudoternary system is divided into H+ t, H+c + t; H+c; and H+c + GIG regions, where H, c, and t denote hematite, c- and t-ZrO2, respectively (Fig. 8) and iron zirconate exsolved from Zr-oversaturated hematite lattice. There is no liquid phase for the a-Fe2O3 – ZrO2 – GIG pseudoternary system at 1300°C. The eutectic temperatures are 1525°C, \ 2000°C (probably) and 1500°C for the binary iron oxide – ZrO2 [8], Gd2O3 – ZrO2 [9 –11] and Gd2O3 –Fe2O3 [16], respectively. The phase assemblages of the two diffusion couples can then be rationalized by specific diffusion path and bulk compositions as shown in Fig. 8. In the above discussion, local equilibrium was assumed at incoherent interfaces between c- and t-ZrO2, between hematite and c-ZrO2, and between hematite and iron zirconate plate at its edge.
Fig. 8. Fe2O3 – ZrO2 –GIG pseudoternary isothermal section (mol.%) at 1300°C. Point I is the specimen with GIG/ZrO2 weight ratio of 1:9 and point II of 9:1. The arrows indicate the diffusion path, and H denotes hematite.
5. Conclusions In a GIG/ZrO2 sample reaction-sintered at 1300°C, the GIG grain is transformed into a single crystal hematite core with iron zirconate precipitates and a Gd-stabilized c-ZrO2 polycrystals shell around it. The iron zirconate precipitated in the hematite core has the same crystal structure, composition, shape and habit plane as that in Zr-supersaturated a-Fe2O3 after aging. The precipitates decrease in size from core center to edge with a PFZ near the edge. The dissolution of Gd3 + can effectively stabilize the c-ZrO2. The (Gd,Fe)-codoped m-ZrO2 has a specific (111) twin boundary in contrast to mosaic {100} and {110} in other PSZ’s. At 1300°C, the a-Fe2O3 –ZrO2 –GIG pseudoternary system is divided into H+ t, H+ c+ t; H+c; and H+c +GIG regions, where H, c, and t denote hematite, c-ZrO2 and t-ZrO2, respectively.
References [1] M.L. Wu, D. Gan, P. Shen, Mater. Sci. Eng. A (2001) in press. [2] A. Chien, Ph.D. Thesis, National Sun Yat-Sen University, Taiwan, 1995. [3] A. Chien, D. Gan, P. Shen, Mater. Sci. Eng. A206 (1996) 215. [4] A.E. Goresy, S.E. Haggerty, J.S. Huebner, D.H. Lindsley, D. Rumble, III, in: D. Rumble, III (Ed.), Oxide Minerals, Mineralogical Society of America Short Course Notes, Southern, Blacksburg, VA, 1976. [5] A. Putnis, J.D.C. McConnell, Principles of Mineral Behavior, Elsevier, New York, 1980. [6] V.Y. Doo, R.W. Balluffi, Acta Metall. 6 (1958) 428. [7] F.J.A. Den Broeder, S. Nakahara, Scripta Metall. 17 (1983) 399. [8] T.S. Jones, S. Kimura, A. Muan, J. Am. Ceram. Soc. 50 (1967) 137. [9] J. Lefe`vre, Ann. Chim. 8 (1963) 128. [10] M. Perez, Y. Jorba, Ann. Chim. 7 (1962) 509. [11] A. Rouanet, Rev. Int. Hautes Temp. Refract. 8 (1971) 161. [12] C.R.M. Grovenor, D.A. Smith, M.J. Goringe, Thin Solid Films 74 (1980) 257. [13] C.R.M. Grovenor, Acta Metall. 33 (1985) 579. [14] C.P. Chang, M.H. Loretto, Acta. Metall. 36 (1988) 805. [15] S.L. Hwang, P. Shen, T.F. Yui, H.T. Chu, Proceedings of the Annual Meeting, Geological Society of China, Taipei, Taiwan, 1999. [16] I. Warshaw, R. Roy, J. Am. Ceram. Soc. 42 (1959) 434. [17] R.D. Shannon, Acta Crystallogr. A32 (1976) 751. [18] M. Ru¨hle, A.H. Heuer, Adv. Ceram. 12 (1984) 14. [19] P. Shen, S. Chen, W.S. Chang, Mater. Sci. Eng. A184 (1994) L5. [20] D.J. Green, R.H.J. Hannink, M.V. Swain, Transformation Toughening of Ceramics, CRC Press, Boca Raton, FL, 1989 references cited therein. [21] P. Li, I.W. Chen, J.E. Penner-Hanh, J. Am. Ceram. Soc. 77 (1994) 118. [22] R. Chain, D.G. Brandon, Adv. Ceram. 12 (1983) 86–95. [23] T.R. Welberry, R.L. Withers, J.G. Thompson, B.D. Butler, J. Solid State Chem. 100 (1992) 71. [24] T.R. Welberry, B.D. Butler, J.G. Thompson, R.L. Withers, J. Solid State Chem. 106 (1993) 461.
M.L. Wu et al. / Materials Science and Engineering A297 (2001) 124–131 [25] T.R. Welberry, R.L. Withers, S.C. Mayo, J. Solid State Chem. 115 (1995) 43. [26] G.M. Wolten, J. Am. Ceram. Soc. 46 (9) (1963) 418. [27] J.E. Bailey, Proc. R. Soc. London Ser. A 279 (1378) (1964) 395.
.
131
[28] G.K. Bansal, A.H. Heuer, Acta Metall. 20 (11) (1972) 1281. [29] G.K. Bansal, A.H. Heuer, Acta Metall. 22 (4) (1974) 409. [30] D.K. Smith, H.W. Newkirk, Acta Crystallogr. 18 (6) (1965) 983.