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Diffusion induced recrystallization of NiO
AC/U mm//. Vol. 32, No. I. pp. 29-33, 1984 Printed in Great Britain. All rights rescind
owl-6160/84s3.00+0.00 Copyright Q 1984 Pergamon Prow Ltd
DIFFUSION INDUCED RECRYSTALLIZATION OF NiO T. A. PARTHASARATHY and P. C. SHEWMON Department of Metallurgical Engineering, The Ohio State University, Columbus. OH 43210. U.S.A. (Receioed 3 Augu.rl 1983) Abstract-It is shown that changing the composition of a sample of NiO from that in equilibrium with air at I2WC to that in equilibrium with oxygen saturated Ni at 800-9OO’Cwill e the surface to a much l&r gnin size. Annealing back at 1200°C in air will again rcuystallizc the surface layer. This type of diffusion induced rw@aUi&oon has o&n been observed in metals, but has llcvcT before been reported in ceramics. It’s oawrewe in NiO is interpreted as a demonstration that Won induced grain boundary motion is driven directly by the free energy of mixing defects into the matrix instead of ind&tly as suggested by others. R&mn&Nous montrons quc k changuncnt de la composition d’un &hantillon de NiO de la composition en Wquilibrc dans l’air a I2WC &cclle qui cst en 6quilibrc dans Ni saturCen oxyg&neh 8W+OO“Cproduit une rwristallisation de la surface & beaucoup plus petite &hellc. Si I’on razuit ensuite i 12oooC d l’air, la couche supcrficielle ncristpllise de nouveau. Cc type de rccristallisation induite par diffbsion a Ctc souvent obse& dans lu m&w, mais il n’avait jamais &t&rapport& dans lcs &amlques. Nous intcrp&ons son &tencc dans NiO oomme la dhonstration que le mouvement da joints de graham induit par di&ion cst produit dimctcmcnt par l%ncrgiclibrc dcs dCfauts de mClangc darts la matrlcc et non indircctemcnt comme le pensent d’autres auteurs. -aalang-wd die zusammMsetxlm g eincr NiO-Robe ge&dcrt, indcm sic aus dun Glcicbge wicht mit Luft bell= in cin Gleichgewic.ht mit sauerstoffgc&ittigtemNi bci 800-9OOV g&racbtwird, dann t&istallisicrt die Ober&bc mit vie1 fcincrcr Komgr66c. Aushcilcn bci 1200°C an Luft bringt die ObcrWhe in den alten Zustand xmitck. l%sc Art dcr diffusionsinduzicrtcn R&istaUlsation wurde s&r OR in Metallen gcfundcn, wurdc jedoch no& nie Eir Kcramiken berichtet. Das Au&ten in NiO witd als ein Hinwcis darauf angesebu~, dal3 die diEusionsinduzierte Komgmnxwan denmg dir& durcbe die freie Energie bcwirkt wird, die dur& Miscbcn da Dcfekte in die Matrix gcwonncn wird, statt-wit sonst vorgescblagen-lndimkt.
carbon diffusion into Ni [12], and oxygen into Cu or Ag [13]. Heretofore the phenomena has not been reported to occur in ceramic systems, but herein we describe the results of experiments that show it can be a very effective means for recrystallizing NiO. Based on our analysis we would expect DIGM/DIR to be observable in a wide range of ceramic systems, given the proper heat treatment. While it is clear that DIGM could not occur in the absence of a concentration gradient, one of the 6rst suggestions was that DIGM could only occur when there was negligible diffusion of solute into the lattice ahead of the moving boundary 1141,that is, when the ratio of the lattice dilfuaion co&cient, D, to the boundary velocity, u, is less than the lattice parameter, “u”. However, DIGM has now been reported in substitutional alloys [3, Is], and interstitial alloys [12,13] at temperatures where D/v is orders of magnitude greater than “a”. All agree that the energy to expand and drive the grain boundaries undergoing DIGM is the free energy of mixing. The controversy arises in just how the free energy couples to the grain boundary motion. Three different models have been put forward, none of which is open to quantitative verification due to lack of data. The first mechanistic model (16,171
INTRODUCTION
lays, that is
that a concentration gradient in a metal can give rise not only to diffusion through the lattice and along stationary grain boundaries, but can also lead to the motion of grain botmdarim so that these high It has become cleat over the past scwraI years
diffusivity paths sweep through the lattice to aid the transport down the gradient (termed diffusion in-
duced grain boundary motion, DIGM). With larger gradients new grains are nucleated and grow under the same driving force (dilfusion induced recrystallization, DIR), to m accelcmte the contribution of boundary m [l-3]. The DIGM phenomenon was originally identified by den Broeder [4] in studying the diffusion of chromium into tungsten at temperature? far below the melting point of tungsten. Hillert and Purdy have shown that it could play a central role in the boundary motion essential to disconthmous precipitation [5,6]. More recently it has been observed in studies as diverse as diffusion in thin 6hns a, 81, the EpUttering of heated targets p, lo], and the annealing of radiation damage 11I]. These studies were all made in alloy systems with substitutional solutes, but the phenomenon has now been found in interstitial al29
30
PARTHASARATHYand SHEWMON:DIFFUSION INDUCED RECRYSTALLIZATION
suggested that as a result of the difference in the boundary diffusion coefficients of the two substitutional elements, diffusion down the gradient would induce a net pr~uction, or destruction, of sites in the boundary which in turn could advance the boundary through the climb of edge type dislocations intrinsic to the grain boundary. Hillert [5] has argued that the coherency strains set up by the change in lattice parameter resulting from the diffusion of solute into the lattice ahead of the moving boundary. He is silent on how this might cuzcurwhen D/v < u. In this model the driving force is proportional to the square of the misfit parameter, ?J(?J=dln a/dN, where N is the atom fraction). Finally, we have suggested [12,13] a “ratchet mechanism” whereby fluctuations in the position of the boundary in the “forward” direction lead to more solute mixing and a larger decrease in the system free energy than fluctuations in the backward direction. Thus fluctuations in position would tend to rat&et the boundary in one direction. The mechanism by which new grains are nucleated in DIR has been discussed very little. Most of the published pietures of boundaries undergoing DIGM are of samples exposed at relatively low supersaturation (low concentration gradients), since the microstructures are less confusing in that range. However, at higher ~~-~~tio~ a dense array of new grains eau be nucleated over the entire surface of foils (Cu-Zn [l], Fe-Zn [2J),or copiously near grain boundaries (Cu-Ni [3], Ni-Fe 1181).Shewmon [19] has discussed the volume and shape change that will accompany the composition change when DIGM occurs in alloys. This volume change, which is proportional to the misfit parameter r~,will resist the nucleation process. Indeed it seems to be able to completely inhibit tbe nudeation of new grains in thick iron samples exposed to Zn vapor eoneentrations that generate a thick layer of nuclei on the surface of &in foils [2]. Shewmon et al. [13] have suggested that a corollary of the “ratchet meehanism” for DIGM is that Diffusion Induced Dislocation Gli& (DIDG) should also oecnr, and argued that the copious and oddly shaped twin boundaries seen id Cu and Ag which undergo DIGM is due to the DIDG driven nucleation, and grow% of the glissile dislooations forming the incoherent boundary of disc shaped twin segments. A search for DIGM in NiO when the ~rn~~tion is driven from one side of its stoichiometry range to the other, was undertaken as a critical test of the various theories of DIGM nucleation and growth. The reasons are as follows: l there is a negligible change in lattice parameter across the stoichiomctry range of NiO, i.e. q = 0. Thus the coherency strain in front of the advancing interface would be negligible, and would play no role in driving DIGM, l there is a large free energy change for this small change in the Nil0 ratio (z24il kJfmot0 [28]), and
stoichiometry shift, though small, is great enough to give a free energy change per unit volume sufficient to ratchet the boundary forward against the curvature of a submicron grain diamter, 0 nickel diffuses much faster than oxygen in NiO, and the grain boundary diffusion coefficient of Ni is roughly five orders of magnitude greater than the lattice diffusion coefficient at 1000°C [21]. Thus reduction of NiO within the stoichiometric range would take place by diffusion of Ni alone along the grain boundaries. However, boundary motion through the elimb of edge-type boundary dislocations would require the diffusion of both nickel and oxygen. Thus dislocation climb could not be the driving mechanism for DIGM, l the negligible volume change accompanying the change in composition will minimize the barrier to nucleation of new grains. the
PROCEDUREAND RESULTS To make the NiO, 80 pm thick foils of 99.95% pure nickel were decarburized in wet hydrogeu at 95O’C for 10 h, then oxidized in air at 1200°Cfor 24 h. This. resulted in well faceted polycrystalline NiO (Fig. 1). To reduce the NiO to the low side of its stoichiometry rang% it was encapsulated in an evacuated vyeor tnbe
Fig. 1. Nickel (99.95%) foil oxidized in air at 1200°C for 24h. SEM image shows well faceted grain structure.
Fig. 2. Nickel oxide (NiO) reduced at 800°C for 24 h. showing new grains on surface.
PARTHASARATHY and SHEWMON: DIFFUSION INDUCED RECRYSTALLIZATION
(a)
(b) Fig. 3. NiO reduced at 850°C for 24 h. High density of new
grains covers entire surface. (a) 1000x . (b) 2000x .
31
(SSO’C) for 20 h produced a much higher density of new grains, covering the entire surface of the NiO sample [Fig. 3(a) and (b)]. The new grains were again not faceted. Various tests were made to check if this apparent recrystaliiition could possibly be the result of diffusion of some element other than oxygen into the sample. Examination under EDAX did not indicate the presence of any other element. The sample shown in Fig. 3 was encapsulated in quartz, whereas those in Figs 2 and 4 were in vycor. No diffbrence was observed. Electron channeling patterns obtained using a STEM showed clear patterns on samples like those shown in Fig. 1, but no coherent pattern from a sample like Fig. 3. This is interpreted to mean that there is not a strong epitaxial relationship between the new and old grains. Subsequently annealing the 800°C reduced NiO in air produced various structures. Annealing at 850°C in air for times up to 72 h did not produce any change in the morphology of the new grains, but resulted in an increase in the average grain size of the new grains, as can be seen by comparing Fig. 2 with Fig. 4. However, annealing the 800°C reduced NiO in air at 1200°C for 4 h produced faceted sub-micron sixed new grains on the surface pig. 5(a) and (b)]. Annealing the 800°C reduced NiO sample at 1200°C in air for much longer, 24 h, resulted in a mixture of coarse (old) grains, and fine (new) grams pg. a(a)].
along with a piece of decarburized nickel foil. The encapsulation was done by evacuating the tube with a mechanical forepump, and then sealing the tube. The encapsulated specimens were given differing heat treatments, and the resulting microstructures observed with the SEM. Annealing at 800°C for 24 h produced a grainy structure on the faceted surfaces of the NiO grains (Fig. 2). These sub-micron sixed new grains did not show any faceted Wucture visible at 5000x. Annealing the NiO/Ni couples at a higher temperature
00
Fig. 4. NiO reduced at 800°C for 24 h and wmcaicd in air at 850°C for 72 h. Notice grain growth compared to Fig. 2.
Fig. 5. NiO rcduozd at 800°C for 24 h and oxidized in air at 12OO’Cfor 2 h, showing faceted new grains. (a) 1000x . (b) 2ooox.
32
PARTHASARATHY
and SHEWMON: DIFFUSION
INDUCED
RECRYSTALLIZATION
All of the grains were well faceted [Fig. 6(b)]. The difference in the ~s~bution of the new grainsin Fig. 5(a) compared with that in Fig. 2 suggests that new grains formed during both the reduction and the
oxidation parts of the cycle. The depth of the reuystallimd layer was determined by examining a surf= normal to the free surface. Figure 7(a) shows the fracture section of a ~ampie similar to that in Fig 1, i.e. as-oxidixed at 12OO’C.Figure 7(b) shows a fracture section of the Sample in Fig. 6. Figure 7(b), clearly shows a diffusion rcuysm layer about one eighth of the sample thi&ws or 1Opm. It is clear that layem of new grains have nucleated repeatedly on the free surface, as has been observed in the Ni-Fe System PI. Essentially no di&sion induced motion of prior grain boundaries was observed. Conditions would have been better for 0-g it if we had had plane polished surfacts instead of deeply f-ted oneS. And, if we had worked at lower supersaturation, to decrease the nucleation rate, and at high temperatures where the boundary mobiity is higher. DEKUssION Fig. 6. NiO mduced at 800°C for 24 h and oxidii in air at 12OVC for 244 showing a mixture of coarse (old) and fine (new) grains. (a) 1000x. @) 2000x.
Fig. 7. Fracture sections of NiO samples. (a) Asoxidized at
12OWC(maneas Fig. I). (b) Following oxidation/~~tion cycle, as it. Fig. 6, showing layers of new grains resulting from cycle.
These experiments clearly show that recrystallization of NiO can be cmscd by a cycle of reduction and o~dation.~~n the range of stoichiometry of NiO. Indeed, the microstructural observations suggest that recrystallixation can be induced by both reduction and by oxidation. We believe the observation of DIR in NiO provides clear support for the ratchet mechanism for coupling the f&e energy of mixing to boundary motion, that is, neither coherency strain nor boundary edge dislocation climb can explain the coupling. However, these experiments also lead to the broader questions of what other ceramics, or compounds might show such behavior. And,-why the entire surface is covered with such fine recrystallized grains. Consider first the queStion of driving force. The exact free energy change per mole of oxygen, AC, for the oxidation or reduction of NiO from one Side of its range of stoichiometry to the other depends on the taco, ~d~or the NifO ratio establishedat the temperature of a prior axmeal.Howaver from the work of Rapp et al. [20] it is roughly 245 kJ/mol of 0. The difference in Ni/O ratio, a*, between NiO in equilibrium with air at 1200°C and that in equilibrium with oxygen saturated nickel in the range of &IO-1200°C is about 10”. For this sto~ome~ change in NiO, the free energy change is NNZx 10’J/m’ of NiO. Taking the boundary fret energy, y, to be 0.5 J/m? this volume f= energy change could just balance the surface tension fOrCeS of a grain of radius, t IE0.05 pm. Thus the 2nd Law of ~e~~~arni~ sets as a limit r > (2y/AG Ax).
PARTHASARATHY
and SHEWMON:
DIFFUSION
This would suggest that in compounds with a vanishingly small stoi~hiomet~ range, the Ax that could be introduced by a change in the annealing atmosphere would not be sufficient to support DIR. Compounds like aluminum oxide and magnesium oxide come to mind. However, the stoichiometry of a great majority of oxides, sulfides, etc. can be varied
more by changes in the annealing atmosphere than can NiO. Thus from the viewpoint of thermodynamic sufficiency a great many compounds would be amenable to such treatment. The copious nucleation of new grains on NiO is not
really different from the structures that can develop on Fe 121,or Cu [1] when the surface is exposed to levels of Zn approaching the maximum solubility for the solid solution. However, in these cases the free energy differences available to drive the nucleation is between one and two orders of magnitude larger than is available in NiO [l ,221.A significant barrier to the nucleation in these metal systems is the volume change that must accompany the nucleation, and the plastic deformation that must accompany the growth [2,13]. In the case of NiO, and probably most ceramics being worked within their range of stoichiometry, there is a negligible volume change associated the nucleation and growth of new grains, so a large fraction of the free energy available could go to produce the new grain boundaries. The observation of diffusion induced boundary motion in NiO may also shed new light on the cause of the duplex oxide layer that forms on nickel when it is oxidized at high temperatures (> 8OO”C),and in thicker foils (>9rm) [2l, 231. This dupiex layer consists of an outer layer of high density columnar grains, and a porous inner layer of e&axed grains. The mechanism for the generation of the scale is complex and controversial (e.g. see Refs [24-261).We would like to suggest that the columnar grains could develop by the DIGM movement of grain boundaries from the oxide/air interface into the layer. These boundaries would follow a roughly constant composition, and would sinter shut any porosity in their path. DlGM could also drive the grain growth observed in the columnar layer by Rhines and Connell[24]. Below 800°C the grain boundary mobilit)’ is too low for such motion, and the scale might be expected to form only quiaxe4i grains.
INDUCED
RECRYSTALLlZATlON
33
Ac~nowf~~~e~e~~ts-~is work was funded through a grant from the Division of Materials Research, OES, U.S. Departmentof Energy. The authors wouldalso like to thank Professors Robert Rapp and Dennis Readey for helpful discussions.
REFERENCES 1. Chongmo Li and M. Hillert, Acru m&u& 29, 1949 (1981). 2. Zong-sen Yu and P. G. Shewmon, Metal. Trans. 14A, 1579 (1983). 3. F. J. A. Den Brocder and S. Nakahara, Scripta metali. 17, 399 (1983). 4. F. J. A. den Broader, Acta m&I. 20, 319 (1972). 5. hi. Hillert. Scrinta metals. 17. 273 (1983). k. h. Hillert and 6. R. Furdy, ,&a n&Ii. &, 333 (1978). 7. J. D. Pan and R. W. Ballti, ACIU metall. 30,861(1982). 8. F. J. A. Den Breeder, M. Klerk, J. M. Vandenbergand
R. A. Hamm, Acta metaR. 31, 285(1983). 9. S. Nakahara and F. 1. A. Den Brocdcr, Scr&ta metall. 17, 607 (1983). IO. H. Shim&, N. Koyama and Y. Ishida, Jupun J. appl. Phys. 19, L67l (1980). 1 I. W. V. Vaidya, J. nucl. Mater. 113, 219-227, and 149-162 (1983). 12. T. Parthasarathy and P. G. Sheumon, Scripa meiaff. 17, 943 (1983). 13. P. G. Shewmon, G. Meyrick S. Mishra. T. A. Parthasarathy,Scripta nwcail. To be published. 14. J. Cahn and R. W. Ballulli, Scripla me&all.13, 499 (1979). 15. K. Tashiro and G. Purdy, Scripta metall. 17,455 (1983). 16. R. W. Ball& and J, Cahn, Acta metail. 29,493 (1981). 17. D. A. Smith and A. H. King, Phil. Mug. A44, 333 (1981). 18. T. A. Parthasarathy and P. G. Shewmon, Metail. Trans. To be published. 19. P. G. Shewmon.Acta metall. 29, 1567(1982). 20. This number comes from the equilibrium constant, and the cation vaeaney eone&ration data of Y. D. Trctyakov, R. A. Rapp, A.f.M.E.Bras 245,1235(1%9). 21. A. Atkinson and R. I. Taylor, PhB M&g. A 43, 979 (1981). 22. R. Hultgren, P. Dcsai, D. Hawkins, M. Gleiser and K. Kelley. Selected Values of tk Tkm&wmb Properties of&nary Alloys. Am. Sot. Metals, Metal Park. OH (197j). 23. A. Atkinson, R. I. Taylor and A. E. Hughes, Phil. Mug. A 45, 823 (1982). 24. F. N: Rhink and R. G. Connell Jr, 3. Wectrocktn. Sot. 124, 1122 (1977). 25. F. N. Rhines, R. 0. Connell Jr and M. S. Choi, 1. Efectrochem. Sot. 126, 1061 (1979). 26. A. Atkinson, R. I. Taylor and P. D. Goode, Oxidation Meta& 13, 519 (1979).
owl-6160/84s3.00+0.00 Copyright Q 1984 Pergamon Prow Ltd
DIFFUSION INDUCED RECRYSTALLIZATION OF NiO T. A. PARTHASARATHY and P. C. SHEWMON Department of Metallurgical Engineering, The Ohio State University, Columbus. OH 43210. U.S.A. (Receioed 3 Augu.rl 1983) Abstract-It is shown that changing the composition of a sample of NiO from that in equilibrium with air at I2WC to that in equilibrium with oxygen saturated Ni at 800-9OO’Cwill e the surface to a much l&r gnin size. Annealing back at 1200°C in air will again rcuystallizc the surface layer. This type of diffusion induced rw@aUi&oon has o&n been observed in metals, but has llcvcT before been reported in ceramics. It’s oawrewe in NiO is interpreted as a demonstration that Won induced grain boundary motion is driven directly by the free energy of mixing defects into the matrix instead of ind&tly as suggested by others. R&mn&Nous montrons quc k changuncnt de la composition d’un &hantillon de NiO de la composition en Wquilibrc dans l’air a I2WC &cclle qui cst en 6quilibrc dans Ni saturCen oxyg&neh 8W+OO“Cproduit une rwristallisation de la surface & beaucoup plus petite &hellc. Si I’on razuit ensuite i 12oooC d l’air, la couche supcrficielle ncristpllise de nouveau. Cc type de rccristallisation induite par diffbsion a Ctc souvent obse& dans lu m&w, mais il n’avait jamais &t&rapport& dans lcs &amlques. Nous intcrp&ons son &tencc dans NiO oomme la dhonstration que le mouvement da joints de graham induit par di&ion cst produit dimctcmcnt par l%ncrgiclibrc dcs dCfauts de mClangc darts la matrlcc et non indircctemcnt comme le pensent d’autres auteurs. -aalang-wd die zusammMsetxlm g eincr NiO-Robe ge&dcrt, indcm sic aus dun Glcicbge wicht mit Luft bell= in cin Gleichgewic.ht mit sauerstoffgc&ittigtemNi bci 800-9OOV g&racbtwird, dann t&istallisicrt die Ober&bc mit vie1 fcincrcr Komgr66c. Aushcilcn bci 1200°C an Luft bringt die ObcrWhe in den alten Zustand xmitck. l%sc Art dcr diffusionsinduzicrtcn R&istaUlsation wurde s&r OR in Metallen gcfundcn, wurdc jedoch no& nie Eir Kcramiken berichtet. Das Au&ten in NiO witd als ein Hinwcis darauf angesebu~, dal3 die diEusionsinduzierte Komgmnxwan denmg dir& durcbe die freie Energie bcwirkt wird, die dur& Miscbcn da Dcfekte in die Matrix gcwonncn wird, statt-wit sonst vorgescblagen-lndimkt.
carbon diffusion into Ni [12], and oxygen into Cu or Ag [13]. Heretofore the phenomena has not been reported to occur in ceramic systems, but herein we describe the results of experiments that show it can be a very effective means for recrystallizing NiO. Based on our analysis we would expect DIGM/DIR to be observable in a wide range of ceramic systems, given the proper heat treatment. While it is clear that DIGM could not occur in the absence of a concentration gradient, one of the 6rst suggestions was that DIGM could only occur when there was negligible diffusion of solute into the lattice ahead of the moving boundary 1141,that is, when the ratio of the lattice dilfuaion co&cient, D, to the boundary velocity, u, is less than the lattice parameter, “u”. However, DIGM has now been reported in substitutional alloys [3, Is], and interstitial alloys [12,13] at temperatures where D/v is orders of magnitude greater than “a”. All agree that the energy to expand and drive the grain boundaries undergoing DIGM is the free energy of mixing. The controversy arises in just how the free energy couples to the grain boundary motion. Three different models have been put forward, none of which is open to quantitative verification due to lack of data. The first mechanistic model (16,171
INTRODUCTION
lays, that is
that a concentration gradient in a metal can give rise not only to diffusion through the lattice and along stationary grain boundaries, but can also lead to the motion of grain botmdarim so that these high It has become cleat over the past scwraI years
diffusivity paths sweep through the lattice to aid the transport down the gradient (termed diffusion in-
duced grain boundary motion, DIGM). With larger gradients new grains are nucleated and grow under the same driving force (dilfusion induced recrystallization, DIR), to m accelcmte the contribution of boundary m [l-3]. The DIGM phenomenon was originally identified by den Broeder [4] in studying the diffusion of chromium into tungsten at temperature? far below the melting point of tungsten. Hillert and Purdy have shown that it could play a central role in the boundary motion essential to disconthmous precipitation [5,6]. More recently it has been observed in studies as diverse as diffusion in thin 6hns a, 81, the EpUttering of heated targets p, lo], and the annealing of radiation damage 11I]. These studies were all made in alloy systems with substitutional solutes, but the phenomenon has now been found in interstitial al29
30
PARTHASARATHYand SHEWMON:DIFFUSION INDUCED RECRYSTALLIZATION
suggested that as a result of the difference in the boundary diffusion coefficients of the two substitutional elements, diffusion down the gradient would induce a net pr~uction, or destruction, of sites in the boundary which in turn could advance the boundary through the climb of edge type dislocations intrinsic to the grain boundary. Hillert [5] has argued that the coherency strains set up by the change in lattice parameter resulting from the diffusion of solute into the lattice ahead of the moving boundary. He is silent on how this might cuzcurwhen D/v < u. In this model the driving force is proportional to the square of the misfit parameter, ?J(?J=dln a/dN, where N is the atom fraction). Finally, we have suggested [12,13] a “ratchet mechanism” whereby fluctuations in the position of the boundary in the “forward” direction lead to more solute mixing and a larger decrease in the system free energy than fluctuations in the backward direction. Thus fluctuations in position would tend to rat&et the boundary in one direction. The mechanism by which new grains are nucleated in DIR has been discussed very little. Most of the published pietures of boundaries undergoing DIGM are of samples exposed at relatively low supersaturation (low concentration gradients), since the microstructures are less confusing in that range. However, at higher ~~-~~tio~ a dense array of new grains eau be nucleated over the entire surface of foils (Cu-Zn [l], Fe-Zn [2J),or copiously near grain boundaries (Cu-Ni [3], Ni-Fe 1181).Shewmon [19] has discussed the volume and shape change that will accompany the composition change when DIGM occurs in alloys. This volume change, which is proportional to the misfit parameter r~,will resist the nucleation process. Indeed it seems to be able to completely inhibit tbe nudeation of new grains in thick iron samples exposed to Zn vapor eoneentrations that generate a thick layer of nuclei on the surface of &in foils [2]. Shewmon et al. [13] have suggested that a corollary of the “ratchet meehanism” for DIGM is that Diffusion Induced Dislocation Gli& (DIDG) should also oecnr, and argued that the copious and oddly shaped twin boundaries seen id Cu and Ag which undergo DIGM is due to the DIDG driven nucleation, and grow% of the glissile dislooations forming the incoherent boundary of disc shaped twin segments. A search for DIGM in NiO when the ~rn~~tion is driven from one side of its stoichiometry range to the other, was undertaken as a critical test of the various theories of DIGM nucleation and growth. The reasons are as follows: l there is a negligible change in lattice parameter across the stoichiomctry range of NiO, i.e. q = 0. Thus the coherency strain in front of the advancing interface would be negligible, and would play no role in driving DIGM, l there is a large free energy change for this small change in the Nil0 ratio (z24il kJfmot0 [28]), and
stoichiometry shift, though small, is great enough to give a free energy change per unit volume sufficient to ratchet the boundary forward against the curvature of a submicron grain diamter, 0 nickel diffuses much faster than oxygen in NiO, and the grain boundary diffusion coefficient of Ni is roughly five orders of magnitude greater than the lattice diffusion coefficient at 1000°C [21]. Thus reduction of NiO within the stoichiometric range would take place by diffusion of Ni alone along the grain boundaries. However, boundary motion through the elimb of edge-type boundary dislocations would require the diffusion of both nickel and oxygen. Thus dislocation climb could not be the driving mechanism for DIGM, l the negligible volume change accompanying the change in composition will minimize the barrier to nucleation of new grains. the
PROCEDUREAND RESULTS To make the NiO, 80 pm thick foils of 99.95% pure nickel were decarburized in wet hydrogeu at 95O’C for 10 h, then oxidized in air at 1200°Cfor 24 h. This. resulted in well faceted polycrystalline NiO (Fig. 1). To reduce the NiO to the low side of its stoichiometry rang% it was encapsulated in an evacuated vyeor tnbe
Fig. 1. Nickel (99.95%) foil oxidized in air at 1200°C for 24h. SEM image shows well faceted grain structure.
Fig. 2. Nickel oxide (NiO) reduced at 800°C for 24 h. showing new grains on surface.
PARTHASARATHY and SHEWMON: DIFFUSION INDUCED RECRYSTALLIZATION
(a)
(b) Fig. 3. NiO reduced at 850°C for 24 h. High density of new
grains covers entire surface. (a) 1000x . (b) 2000x .
31
(SSO’C) for 20 h produced a much higher density of new grains, covering the entire surface of the NiO sample [Fig. 3(a) and (b)]. The new grains were again not faceted. Various tests were made to check if this apparent recrystaliiition could possibly be the result of diffusion of some element other than oxygen into the sample. Examination under EDAX did not indicate the presence of any other element. The sample shown in Fig. 3 was encapsulated in quartz, whereas those in Figs 2 and 4 were in vycor. No diffbrence was observed. Electron channeling patterns obtained using a STEM showed clear patterns on samples like those shown in Fig. 1, but no coherent pattern from a sample like Fig. 3. This is interpreted to mean that there is not a strong epitaxial relationship between the new and old grains. Subsequently annealing the 800°C reduced NiO in air produced various structures. Annealing at 850°C in air for times up to 72 h did not produce any change in the morphology of the new grains, but resulted in an increase in the average grain size of the new grains, as can be seen by comparing Fig. 2 with Fig. 4. However, annealing the 800°C reduced NiO in air at 1200°C for 4 h produced faceted sub-micron sixed new grains on the surface pig. 5(a) and (b)]. Annealing the 800°C reduced NiO sample at 1200°C in air for much longer, 24 h, resulted in a mixture of coarse (old) grains, and fine (new) grams pg. a(a)].
along with a piece of decarburized nickel foil. The encapsulation was done by evacuating the tube with a mechanical forepump, and then sealing the tube. The encapsulated specimens were given differing heat treatments, and the resulting microstructures observed with the SEM. Annealing at 800°C for 24 h produced a grainy structure on the faceted surfaces of the NiO grains (Fig. 2). These sub-micron sixed new grains did not show any faceted Wucture visible at 5000x. Annealing the NiO/Ni couples at a higher temperature
00
Fig. 4. NiO reduced at 800°C for 24 h and wmcaicd in air at 850°C for 72 h. Notice grain growth compared to Fig. 2.
Fig. 5. NiO rcduozd at 800°C for 24 h and oxidized in air at 12OO’Cfor 2 h, showing faceted new grains. (a) 1000x . (b) 2ooox.
32
PARTHASARATHY
and SHEWMON: DIFFUSION
INDUCED
RECRYSTALLIZATION
All of the grains were well faceted [Fig. 6(b)]. The difference in the ~s~bution of the new grainsin Fig. 5(a) compared with that in Fig. 2 suggests that new grains formed during both the reduction and the
oxidation parts of the cycle. The depth of the reuystallimd layer was determined by examining a surf= normal to the free surface. Figure 7(a) shows the fracture section of a ~ampie similar to that in Fig 1, i.e. as-oxidixed at 12OO’C.Figure 7(b) shows a fracture section of the Sample in Fig. 6. Figure 7(b), clearly shows a diffusion rcuysm layer about one eighth of the sample thi&ws or 1Opm. It is clear that layem of new grains have nucleated repeatedly on the free surface, as has been observed in the Ni-Fe System PI. Essentially no di&sion induced motion of prior grain boundaries was observed. Conditions would have been better for 0-g it if we had had plane polished surfacts instead of deeply f-ted oneS. And, if we had worked at lower supersaturation, to decrease the nucleation rate, and at high temperatures where the boundary mobiity is higher. DEKUssION Fig. 6. NiO mduced at 800°C for 24 h and oxidii in air at 12OVC for 244 showing a mixture of coarse (old) and fine (new) grains. (a) 1000x. @) 2000x.
Fig. 7. Fracture sections of NiO samples. (a) Asoxidized at
12OWC(maneas Fig. I). (b) Following oxidation/~~tion cycle, as it. Fig. 6, showing layers of new grains resulting from cycle.
These experiments clearly show that recrystallization of NiO can be cmscd by a cycle of reduction and o~dation.~~n the range of stoichiometry of NiO. Indeed, the microstructural observations suggest that recrystallixation can be induced by both reduction and by oxidation. We believe the observation of DIR in NiO provides clear support for the ratchet mechanism for coupling the f&e energy of mixing to boundary motion, that is, neither coherency strain nor boundary edge dislocation climb can explain the coupling. However, these experiments also lead to the broader questions of what other ceramics, or compounds might show such behavior. And,-why the entire surface is covered with such fine recrystallized grains. Consider first the queStion of driving force. The exact free energy change per mole of oxygen, AC, for the oxidation or reduction of NiO from one Side of its range of stoichiometry to the other depends on the taco, ~d~or the NifO ratio establishedat the temperature of a prior axmeal.Howaver from the work of Rapp et al. [20] it is roughly 245 kJ/mol of 0. The difference in Ni/O ratio, a*, between NiO in equilibrium with air at 1200°C and that in equilibrium with oxygen saturated nickel in the range of &IO-1200°C is about 10”. For this sto~ome~ change in NiO, the free energy change is NNZx 10’J/m’ of NiO. Taking the boundary fret energy, y, to be 0.5 J/m? this volume f= energy change could just balance the surface tension fOrCeS of a grain of radius, t IE0.05 pm. Thus the 2nd Law of ~e~~~arni~ sets as a limit r > (2y/AG Ax).
PARTHASARATHY
and SHEWMON:
DIFFUSION
This would suggest that in compounds with a vanishingly small stoi~hiomet~ range, the Ax that could be introduced by a change in the annealing atmosphere would not be sufficient to support DIR. Compounds like aluminum oxide and magnesium oxide come to mind. However, the stoichiometry of a great majority of oxides, sulfides, etc. can be varied
more by changes in the annealing atmosphere than can NiO. Thus from the viewpoint of thermodynamic sufficiency a great many compounds would be amenable to such treatment. The copious nucleation of new grains on NiO is not
really different from the structures that can develop on Fe 121,or Cu [1] when the surface is exposed to levels of Zn approaching the maximum solubility for the solid solution. However, in these cases the free energy differences available to drive the nucleation is between one and two orders of magnitude larger than is available in NiO [l ,221.A significant barrier to the nucleation in these metal systems is the volume change that must accompany the nucleation, and the plastic deformation that must accompany the growth [2,13]. In the case of NiO, and probably most ceramics being worked within their range of stoichiometry, there is a negligible volume change associated the nucleation and growth of new grains, so a large fraction of the free energy available could go to produce the new grain boundaries. The observation of diffusion induced boundary motion in NiO may also shed new light on the cause of the duplex oxide layer that forms on nickel when it is oxidized at high temperatures (> 8OO”C),and in thicker foils (>9rm) [2l, 231. This dupiex layer consists of an outer layer of high density columnar grains, and a porous inner layer of e&axed grains. The mechanism for the generation of the scale is complex and controversial (e.g. see Refs [24-261).We would like to suggest that the columnar grains could develop by the DIGM movement of grain boundaries from the oxide/air interface into the layer. These boundaries would follow a roughly constant composition, and would sinter shut any porosity in their path. DlGM could also drive the grain growth observed in the columnar layer by Rhines and Connell[24]. Below 800°C the grain boundary mobilit)’ is too low for such motion, and the scale might be expected to form only quiaxe4i grains.
INDUCED
RECRYSTALLlZATlON
33
Ac~nowf~~~e~e~~ts-~is work was funded through a grant from the Division of Materials Research, OES, U.S. Departmentof Energy. The authors wouldalso like to thank Professors Robert Rapp and Dennis Readey for helpful discussions.
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