chemically induced grain boundary migration and recrystallization in al2o3

Acta metall, mater. Vol. 38, No. 7, pp. 1307-1312, 1990 0956-7151/90 S3.00 + 0.00 Printed in Great Britain. All rights reserved Copyright © 1990 Pergamon Press plc

CHEMICALLY INDUCED GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN A1203

HO YONG LEE and SUK-JOONG L. KANG Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology,

Cheongryang, Seoul 130-650, Korea

(Received 27 October 1989)

Abstract--The chemically induced grain boundary migration in A12 03 has been observed during depletion of Fe203 from AI203-Fe203 solid solution as well as additon of Fe203 in A1203. Many migrating boundaries are faceted. In some receding grains, the faceted planes of different grain boundaries advancing into the same grain are parallel to each other. The faceting is attributed to the anisotropy of driving force in thin diffusion layer of the receding grain; this result provides an experimental support for coherency strain energy as the major driving force for the migration. When the supply of the solute is excessive during alloying, recrystallization occurs in A1203 grains.

Rrsum~-La migration des joints de grains induite chimiquement dans AI203 est observre au cours de la deplrtion du Fe 203/t partir de la solution solide A12 O3-Fe 2 O 3 ainsi qu'au cours de l'addition de Fe 203 dans A1203. De nombreux joints en migration prdsentent des facettes. Dans certains grains qui reculent, les plans ~i facettes des diffrrents joints de grains avancant dans le mrme grain sont parallrles les uns aux autres. Le facettage est attribu6 d l'anisotropie de la force motrice dans la couche mince de diffusion du grain qui recule: ce rrsultat prouve exprrimentalement que l'rnergie de drformation de cohrrence est la principale force motrice de la migration. Lorsque la quantit6 de solut6 est en excrs pendant la formation de l'alliage, il se produit une recristallisation dans les grains de Al203 .

Zusammenfassung--Die chemisch induzierte Korngrenzwanderung in A1203 wurde w/ihrend der Verar- mung oder Anreicherung von Fe203 in dem Mischkristall Al203-Fe203 beobachtet. Viele der wandernden Korngrenzen sing facettiert. In einigen verschwindenden K6rnern sind die facettierten Teile verschiedener in dasselbe Korn vordringender Korngrenzen zueinander parallel. Die Facettierung wird der Anisotropie der treibenden Kraft in der diinnen Diffusionschicht des verschwindenden Kornes zugeschrieben. Dieses Ergebnis ist eine experimentelle Stiitze daffir, dab die wesentliche treibende Kraft der Korngrenzwan- derung die KohS.renzverzerrungsenergie ist. Ist das Angebot an gel6sten Stoffen w/ihrend des Legierens aul3ergew6hnlich hoch, dann tritt in Al203-K6rnern Rekristallisation auf.

1. INTRODUCTION

Under the difference in the chemical potential of solute atoms in grains and surrounding phase or atmosphere, the interface migration, concomitant with the formation of a new solid solution, has been observed in various metallic alloy systems [1--6]. In some alloys under large chemical potential differences, recrystallization of grains has also been re- ported [7, 8]. The chemical potential difference can be provided by either solid, liquid or vapor phase. In case of the migration of liquid films between grains [9], its driving force has been found to be coherency strain energy in thin diffusion layer of solute atoms at the interface of receding grains [10, ll]. For the grain boundary migration (CIGM), the coherency strain energy has also been revealed to be a dominant driving force for some alloys [12, 13].

Compared to the extensive study and the knowl- edge of the interface migration for metal systems, investigations for ceramic systems are very limited [14-19]. In some ceramic systems, grain boundaries

have also been observed to migrate during depletion or enrichment of solute component, as in metal systems. Recrystallization of grains under chemical potential difference in solute component has only been observed in PLZT [18]. These investigations [14-18], however, were focused on the demonstration of the phenomena under chemical potential variation in the system, and did not provide detailed analyses, such as the determination of compositional change in migrated or recrystallized region.

In the present investigation, chemically induced grain boundary migration (CIGM) and recrystallization (CIR) have been studied in A1203, one of the most common oxides of practical importance. As a solute compound, Fe20a is selected to take advantage of its high solid solubility in AI:O3 [20] and high vapor pressure [21] at usual annealing temperatures of A12Oa. In addition, possible valency problems, which can arise from the difference in valencies between solute and matrix atoms, can be avoided. Fe203 can be depleted or solutioned by vapor phase transport.

1307

1308 LEE and KANG: GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN A1203

2. EXPERIMENTAL

Two series of specimens were prepared from AI203 (99.99 wt% in purity) and Fe203 (99.9 wt%) powders: one for removing of Fe203 from A1203-Fe203 solid solution and the other for adding of Fe203 in pure AI203 . For a dealloying experiment, 93A1203-7Fe203 (in wt%) powder mixture was wet- milled for 12h in alcohol. The dried slurry was isostatically pressed under 100MPa into disks of about 13 mm in diameter and 7 mm in height. The compacts were packed in powders of same composition and sintered at 1600°C for 20 h in air in order to obtain coarse grains. For an alloying experiment, sintered A1203 compacts were prepared following a similar procedure to that for sintered AI203-Fe203 compacts. In this case, in order to obtain coarse A1203 grains in a reasonable sintering time span, the compacts were sintered at 1850°C for 2.5 h in vacuum. (In case of AI203-Fe 2 03, the compacts were sintered at lower temperature of 1600°C, at which sufficient grain growth occurred, in order to reduce the volatilization of Fe203.)

All the sintered specimens were cut, polished and then annealed at 1600°C in air. In the dealloying experiment, the sintered specimens were placed at the hot zone of the furnace in air. In the alloying experiment, Fe203 was supplied to sintered A1203 by placing Fe203 powder near the compact on a Pt foil. The compact and the Fe203 powder were covered by a cap made of refractory brick. Since the vapor pressure of Fe203 was high enough (over 10 -5 atm at 1600°C [21 ]), Fe2 03 was transported and added to the specimen through the vapor phase. For some alloying specimens Fe203 was directly supplied by placing them on an Fe203 powder bed. The microstructural change was observed on the prepolished sections without further thermal treatments. The cation composition was determined by WDX analysis.

3. RESULTS AND DISCUSSION

3.1. Dealloying of Fe203from 93A120 3- 7Fe 20 3

After sintering of 93A1203-7Fe203 at 1600°C for 20 h, specimens consist of large grains of about 180#m in average size, and large pores at grain boundaries and grain corners, as shown in Fig. 1. The grains contain numerous fine pores, indicating that the fast grain growth occurred before final densification of the compact.

When the sintered specimen has been annealed in air after cutting and polishing, the grain boundary has migrated, as shown in Fig. 2. The thin smooth boundaries (indicated by arrows with "b" ) in Fig. 2 are their original position in the sintered specimen before annealing and the thick curved ones are those (indicated by arrows with "a" ) positioned after the annealing. Thermal grooving of the original position of boundaries appears to be fast enough to reveal the position.

Fig. 1. Microstructure of 93AI203-7Fe203 (wt%) specimen sintered at 1600°C for 20 h in air.

The grain boundary migrates in a way to increase its area, independent of the boundary curvature, as in the chemically induced grain boundary migration in various metal systems [1-6]. Some boundaries seem to migrate unevenly, resulting in a wavy appearance of the migrated region, such as the boundary indicated by an arrow with "u" in Fig. 2. The uneven motion of the boundaries is believed to result from the dragging effect of pores on the migration. In fact, the shape of the boundaries revealed in the migrated region appears to be pinned by relatively large pores. Corrugation of the boundaries is also observed, for example, the migrated boundary positioned at the upper left on Fig. 2. Such a corrugation has been observed for almost all the grain boundaries at the early stage of annealing (about 30 min), suggesting that the corrugation is more related to the local nucleation of new solid solution layer of the two facing grains.

Figure 3 shows the variation of Fe203 content between two adjacent original grains and migrated region. The migrated region contains less Fe203 than the original grain by about 1.5 wt%, indicating that the migration has been accompanied by a loss of

Fig. 2. Grain boundary migration at the surface of sintered 93AI203-7Fe203 (wt%) specimen during annealing at 1600°C for 2 h in air. Arrows with "b" and "a" indicate grain boundaries before (b) and after (a) annealing, and an

arrow with "u" indicates uneven boundary motion.

LEE and KANG: GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN AI203 1309

8

"7 0

¢ t

~6

d 5 E 0

0

4;

• I I i • i i i

1 Io i

m

" i o ' / o ' ~ ' ~o

Distance (ku'n)

Fig. 3. Measured concentrations of Fe203 in migrated region and original A],O 3 Fe203 grain. (The dotted and solid lines, respectively, show the initial and final positions

of the grain boundary.)

Fe203. The concentration of Fe203 in migrated region, however, is not uniform. The nonuniformity seems to result from uneven volatilization of Fe203 due to uneven grain-boundary migration. Since Fe203 is lost by volatilization through the grain boundary during migration, similar to the dezincifica- tion in an Fe-Zn specimen [2], the phenomenon is indeed a, so-called, diffusion induced grain boundary migration (DIGM) in A1203.

In addition to the dealloying of Fe203, the decrease of the area of the grain boundaries in contact with the specimen surface, i.e. the tendency for grain boundaries to become vertical to the specimen surface, can also contribute to the migration. This contribution, however, has been found to be marginal in an annealing experiment of pure A1203 (see next section). Moreover, the boundary corrugation cannot occur by the reduction of grain boundary area.

A closer examination of the morphology of migrated boundaries in Fig. 2 reveals some faceted ones. Figure 4 is another example of faceted grain boundaries resulting from their migration. The two faceted boundaries of grain G show a parallelism in their shape [A to a. B to b (b') and C to c (c')], implying that the crystallographic orientations of corresponding parallel parts of the receding grain G are identical.t The faceting is, therefore, believed to be essentially related to the crystallographic orientation of the receding grain and reflects, in turn, an anisotrop) in the grain boundary mobility and the driving force for the migration.

The grain boundary mobility, however, may not be a dominant factor in our case. Since the boundary mobility is proportional to the diffusivity of atoms perpendicular to the boundary, the mobility would be determined by crystallographic orientations of both gro~ing and receding grains. Therefore, an anisotropy in the driving force, which arises from the

tSince such a parallelism has been observed in some receding grains of the same specimen, the possibility of identical crystallographic orientations of the growing grains into one receding grain in three dimension can be discarded.

Fig. 4. Microstructure displaying the faceting of migrated boundaries in 93A1203-7Fe203 specimen annealed at 1600°C for 2h in air. The shape of the two faceted boundaries (upper and lower) of a grain G showing a

parallelism: A to a, B to b (b'), C to c (c').

difference in chemical potential of atoms adjacent to grain boundaries of both receding and growing grains, is thought to determine the formation of parallel faceting in same receding grains.

The anisotropy in the driving force, related to the crystallographic orientation of receding grains, cannot be expected in other models for the driving force in CIGM except in the coherency strain energy model [22], which is well accepted in various metal systems [10--13].

According to the coherency strain energy model, the strain energy in coherent diffusion layer is anisotropic in any crystal system, and its anisotropy is related to the crystallographic orientation of the receding grain (see Appendix). The resultant faceted orientations of moving boundaries should be determined by the orientations which provide low coherency strain energies during solute diffusion. Therefore the observed parallelism in faceted boundaries in Fig. 4 appears to provide an experimental support for the coherency strain energy as being the dominant driving force for the chemically induced grain boundary migration in AI203 base ceramics.

3.2. Alloying of Fe:03 in Al203

The sintered AI 2 03 specimen showed the typical polycrystalline microstructure with large grains of about 200 #m in average size. When the specimen was annealed at 1600°C in Fe203 vapor containing air, the grain boundary migrated to increase its area, as shown in Fig. 5. Boundary corrugation and faceting are also observed. To check the contribution of decrease in grain boundary area to the migration during annealing, the sintered specimen was annealed at the same temperature of 1600°C for 5h in air without Fe203 vapor. Unidirectional boundary migration has frequently been observed; however, its maximum distance is limited to less than a few micron. Indeed the migration observed in Fig. 5 results mainly from a chemical effect.

1310 LEE and KANG: GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN A1203

Fig. 5. Grain boundary migration at the surface of sintered AI203 specimen during annealing at 1600°C for 2 h in Fe203

vapor containing air.

Figure 6 plots the measured increase of Fe203 content in the migrated region. The migration is accompanied by alloying of Fe20 a of about 4 wt% in A1203, contrary to the former case of dealloying of Fe203 from A1203-Fe203 solid solution. (The solutioning of Fe203 of about 1 wt% in original Al203 grain seems to be due to the lattice diffusion of Fe203 during alloying.) Fe203 has been supplied as a vapor phase to the grain boundary during its migration, similar to the alloying of Zn in the Fe-Zn system [2, 3].

When a sintered A1203 specimen was in direct contact with an Fe203 powder bed during its annealing, refinement of Al 203 grains as well as grain boundary migration occurred, as shown in Fig. 7(a). The specimen surface on the figure was vertically placed on the Fe203 powder bed. From the end of the surface in contact with the bed, a melted Fe203 phase on the polished alumina surface (region A in the figure), fine grain structure of Al203 [region B and its SEM micrograph, Fig. 7(b)] and grain boundary migration (region C) are progressively observed.

The refinement of A1203 grains is a result of chemically induced recrystallization, as observed in various metals [7, 8] and PLZT ceramics [18]. Such a grain refinement is known to be possible under the

5

~4

~3 &z.

~2

0

06 " lb "a'o " A "~ "sb "eo Distance ( /an)

Fig. 6. Measured concentrations of Fe203 in migrated region and original A1203 grain. (The dotted and solid lines respectively show the initial and final positions of

the grain boundary).

• • w

- i

Fig. 7. (a) Surface microstructure of sintered A1203 specimen after an annealing on an Fe203 powder bed at 1600°C for 2 h and (b) SEM micrograph of recrystallized region [B].

condition of excessive supply of solute atoms to the crystal [3, 8]. The Fe203 content in such fine grains has in fact been measured to be high, about 6.3 wt%, near the limit of solid-solubility (about 8 wt%) at the annealing temperature. By contrast, in the region where only the boundary migration is observed, Fe203 content is relatively low, about 4wt%. The variation of the microstructure in Fig. 7 depending on the distance from the Fe203 powder bed, is therefore due to the difference in solute content: high solute content resulting in chemically induced recrystallization and low content in chemically induced grain boundary migration. The difference in the solute content should stem from differences in the diffusion distance of solute atoms.

4. CONCLUSIONS

The grain boundary migration and the grain recrystallization in A12 03 (Fe2 O3) have occurred by the variation of Fe203 solute content during annealing. Since the vapor pressure of Fe2 03 is sufficiently high, Fe2 03 is transported as vapor phase; the experimental condition can easily be monitored, as for some metal systems [3, 4]. Faceting of the migrated boundary is found to be a general feature in A1203. Some faceted boundaries for the same receding grain show a

parallel ism in their shape. An i so t ropy in dr iving force is t hough t to be the main cause for the faceting and to provide an exper imental suppor t for the coherency s t ra in energy model in the bounda ry migra t ion of AI2 03 . Recrystai l izat ion of AI 2 03 grains results f rom a supply of excessive a m o u n t of F e 2 0 3. The present exper imenta l results demons t ra te tha t the chemically induced grain bounda ry migra t ion and recrystalliz- a t ion can be c o m m o n p h e n o m e n a not only in metals but in oxide ceramics under appropr ia te experimental condit ions. Coherency s t ra in energy seems to be the ma jo r driving force for the bounda ry migra t ion in bo th metallic and ceramic materials.

As far as the shape of migrated grain boundar ies is concerned, in addi t ion to faceted boundaries , cor- rugated boundar ies are also observed. The bounda ry cor rugat ion , however, is favored at the early stage of anneal ing. Since the increase of grain bounda ry area can improve the s t rength of the mater ia l [23], surface s t rengthening of AI20 3 may also be expected by inducing the grain boundary corrugat ion. Appropr i - ate solute elements and exper imental condi t ions need to be explored.

Acknowledgements--The authors are grateful to Professor D. N. Yoon for invaluable comments and discussions. This work was supported jointly by the Ministries of Science and Technology of the Republic of Korea (MOST, contract number 1U00500) and of the Federal Republic of Germany (BMFT, contract number 03M1026).

REFERENCES

o"

1. F. J. A. den Broeder, Acta metall. 20, 319 (1972). 2. M. Hillert and G. R. Purdy, Acta metall. 26, 333 (1978). 3. Li Chongmo and M. Hillert, Acta metall. 29, 1949

(1981). 4. Li Chongmo and M. Hillert, Acta metall. 30, 1133

(1982). 5. Y.J. Baik and D. N. Yoon, Acta metall. 34, 2039 (1986). 6. D. N. Yoon, Ann. Rev. Mater. Sci. 19, 43 (1989). 7. V. Y. Doo and R. W. Balluffi, Acta metall. 6, 428 (1958). 8. F. J. A. den Broeder, Thin Solid Films 124, 135 (1985). 9. D. N. Yoon and W. J. Huppmann, Acta metall. 27, 973

(1979). 10. W. H. Rhee. Y. D. Song and D. N. Yoon, Acta metall.

35, 57 (1987). 11. J. K. Kim, M.S. thesis, Korea Advanced Inst. Sci.

Tech.. Seoul (1987). 12. W. H. Rhee and D. N. Yoon, Acta metall. 37, 221

(1989). 13. K. R. Lee, Y. J. Baik and D. N. Yoon, Acta metall. 35,

2145 (1987). 14. R. Chaim, A. H. Heuer and D. G. Brandon, J. Am.

Ceram. Soc. 69, 243 (1986). 15. J. E. Blendell, C. A. Handwerker, C. A. Shen and

Ngoc-Duyen Dang, in Ceramic Microstructures '86 [Mater. Sci. Res.. Vol. 21] (edited by J. A. Pask and A. G. Evans), p. 541. Plenum Press, New York (1988).

16. R. S. Hay and B. Evans, Acta metall. 35, 2049 (1987). 17. K. J. Yoon, D. N. Yoon and S.-J. L. Kang, Ceramics

Int. In press. 18. J. J. Kim, B. M. Song, D. Y. Kim and D. N. Yoon, Am.

Ceram. Soc. Bull. 65, 1390 (1986). 19. J. W. Jeong, D. N. Yoon and D. Y. Kim, J. Am. Ceram.

Soc. In press. 20. J. C. Willshee and J. White, Trans. Br. Ceram. Soc. 67,

271 (1968).

21. Toshiyuki Sata, Refractories 30, 383 (1972). 22. M. Hillert, Metall. Trans. 3A, 2729 (1972). 23. J. J. Kim, Ph.D. thesis, Seoul National Univ., Seoul

(1988). 24. J. W. Cahn, Acta metall. 9, 795 (1961). 25. J. E. Hilliard, in Phase Transformations (edited by H. I.

Aaronson), p. 509. Am. Soc. Metals (1968). 26. J. F. Nye, in Physical Properties of Crystals. Clarendon

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(1966).

A P P E N D I X

Calculation of the Coherency Strain Energy in the [Jm] Trigonal System

A complete description of the coherency strain energy for crystal systems is not available yet except for the cubic system [24]. In order to explain the faceting of moving boundaries in AI203 during the chemically induced grain boundary migration, we will derive the coherency strain energy for the []m] trigonal system by taking advantage of Hilliard's derivation [25] for the cubic system.

Consider, first, a bar and a slice cut from the bar as in Fig. AI. When Solute atoms whose size is different from that of matrix atoms diffuse into the slice within the solubility limit, hydrostatic stress builds up, if the dimension of the slice is not changed and the atoms behave as rigid balls. Provided that no plastic deformation results, the elastic strain energy per unit volume, E(1), in crystals of []m] trigonal symmetry is

E(1) = (~)Za:j

= ( ½ ) z s : ~

= (~)a~(2Sll + 2S12 + 4S13 + $33) (AI)

where a h is hydrostatic stress in the slice and S U elastic compliances, a h can be expressed as the product of the linear compressibility, fl [26], and the normal strain, E

~h=fl X E

= [(S. + $12 + Si3)

- (Sil + St2 - Si3 - $33)l~] [(~t0 - ct)/~t]. (A2)

Here ~t 0 and ~t are the equilibrium lattice parameter of chemical compositions of the bar and the slice, respectively. Next, if we attach reversibly the slice to the bar and make

LEE and KANG: GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN Al203 1311

Fig. A1. Schematic showing the procedure for the calculation of coherency strain energy in a thin diffusion layer depending on crystallographic orientation of receding grain. If a slice having a composition different to that of a slab is coherent with it, the slice is subjected to the stress, a, in y

and z directions.

1312 LEE and KANG: GRAIN BOUNDARY MIGRATION AND RECRYSTALLIZATION IN AI203

the stress in x direction of the slice released, the released normal strain, e,, is

E., --- cr.,l Y,

1/Yx = (1 - 12)2S,, + 1](1 - 12)(2S,3 + S44)

+ t~s33 + 2/d3(3/~ - /Ds , , (A3)

where a., is the released stress, 1/y, the reciprocal of Young's modulus in the stress direction [26] and l~ the cosine value of the direction in the Jim] trigonal system. If the amount of relaxed shear strain due to the stress releasing is small, the released strain energy per unit volume, E(2), can be approximated as

E(2) = (½)o.,~ x = (½)(I/y,)a~. (A4)

The elastic strain energy per unit volume, Ec, remaining in the slice is then

E c = E(1) - E(2) = ~ba2h (A5)

where

I// = (~)(2Sll + 2SI2 + 4SI3 + $33 -- l /y , ) .

The resultant energy change of the imaginary processes so far assimilates with the strain energy of a coherent diffusion layer on a matrix crystal, i.e. coherency strain energy in the layer. The coherency strain energy is, therefore, determined by the crystallographic direction of stress release.

If we know the elastic compliance values for the material concerned, the variation of relative coherency strain energy with the crystallographic orientation can easily be calculated from equation (A5) by using a personal computer. For alumina, from an available S O data obtained at 900 K [27] the energy minima have been found to exist for directions perpendicular to (102), (0II), (I22) plane, etc. These planes correspond to (1012), (0111) and (1"21"2) planes in the hexagonal axis, respectively. The calculated energies for the planes relative to the maximum value is almost the same of about 0.43. If the coherency strain energy in thin diffusion layer of receding grain ahead of moving boundary is the major driving force for the chemically induced grain boundary migration in the alumina system, these planes should be major planes of faceted boundaries of the receding crystal.

chemically induced grain boundary migration and recrystallization in al2o3

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