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Chemically oriented γ′ plate development in a nickel base superalloy
Scripta METALLURGICA
Vol. 23, pp. 1877-1882, 1989 Printed in the U.S.A.
Pergamon Press plc All rights reserved
CHEMICALLY ORIENTED ~ PLATE DEVELOPMENT IN A NICKEL BASE SUPERALLOY. Alain HAZOTTE and Jacques LACAZE Laboratoire de Science et G~nie des MatOriaux M~talliques Ecole des Mines -Parc de Saurupt - 54042 NANCY - FRANCE (Received August l, 1989) (Revised August 29, 1989)
Introduction The 1'-7' microstructure of nickel base superalloys can encounter drastic changes during high temperature heat treatments. The morphology of the 1" precipitates evolves from different mechanisms: i) competitive coarsening in order to reduce the specific area of the 1'-7' interface (Oswald ripening) (1-5) and ii) shape changes in order to minimize the sum of interfacial and elastic interaction energies. Different theoretical approaches forecast successive equilibrium shapes of the 7' precipitates depending not only on their size but also on the characteristics of the 1'-7' interface (6-14). Most of these works led to the conclusion that the precipitates should evolve towards a plate-like morphology when reaching a size of several microns. Although these predictions have been numerically established in the case of isolated groups of precipitates, there is some experimental evidence of this "rafting" in the case of alloys with a high amount of strengthening phase (6,15-17). The largest faces of the plates develop along the {100} planes of the crystallographic lattice common to both the 1'and y' phases. In the absence of any external applied stress the orientation of the 11' plates would have been thought to be evenly distributed along the three possible {100} planes. This report deals with experimental observations of preferentially oriented plate development of the 7' precipitates that take place without any applied stress. The experiments reported here concern a single grain commercial superalloy which has been submitted to long time aging treatments at 1100°C. When the material was not perfectly chemically homogeneous initially, it was noticed that locally oriented 1~ rafts developed on the dendritic scale, and that their orientation was related to the dendritic structure. It is proposed that this phenomenon is induced by the local chemical heterogeneities due to the usual solidification process. Some mechanisms which could link together chemical gradient and plate development are discussed. Experimental Procedure Experiments were performed on the commercial alloy AM1 ( Ni base-7.5wt% Cr, 5.2 AI, 5.5 W, 8.6 Ta, 6.5 Co, 1.2 Ti, 2 Mo )(18). Single grain bars with a diameter of 14 mm were produced by directional solidification in the SNECMA foundry house. The growth direction was almost parallel to the [100] crystallographic direction. The withdrawal rate used was the usual one for manufacturing turbine blades (4 mm/min.). It involves a dendritic solidification front and leads to final mlcrosegregatlons In the as-cast product, that is chemical heterogeneities on the scale of the dendrites trunks and their dendritic arms. One usually performs high temperature heat treatments to reduce these chemical heterogeneities. In the case of nickel base superalloys of the second generation as is AM1, there exists a "window", that is a temperature interval between the solvus temperature of the 7' precipitates and the incipient melting
1877 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc
1878
~' PLATES
IN A SUPERALLOY
Vol.
23, No.
ii
temperature. The homogenizing treatments are made in this window, which is 1295-1310°C for alloy AM1 (lg). In the present work the effect of the initial chemical microsegregations was investigated by using materials homogenized for different times : i) as cast, without any homogenizing treatment, ii) after a heat treatment at 1300°C for 3 hours then air cooled, iii) after 100 hours at 1300°C then air cooled. It has been verified by microprobe analysis that microsegregations have completely disappeared after the third treatment while they are still detected after the second one which corresponds to the standard industrial treatment. Fig. 1 presents optical micrographs of samples when as-cast (l-a) and after the homogenizing treatment of 3 hours (l-b). The chemical etching allows the solidification structure to be seen thanks to the residual microsegregations : in Fig. 1-a it is possible to note the presence of tertiary dendrite arms that have a width of about 60 p.m. In the case of the samples homogenized 100 hours, the dendritic structure does not appear any more. About 60% of ~,' phase precipitates during the last cooling stage of the heat treatments. The mean size of ~ precipitates was always found to be around 0.4 t ~ . However, their actual size distribution obviously depended on the chemical homogeneity of the material since the local composition is known to strongly determine the precipitation kinetics (20,21). As a matter of fact the size of the precipitate increased from the cores of the dendrites to the interdendritic regions in the case of materials not fully homogenized. Their morphology also evolved from cuboidal in the cores to more erratic shapes in the interdendritic regions. The aging treatments were performed on samples cut from different bars and sealed under secondary vacuum in silica tubes. They were held at 1100°C _+ 5°C for different times ranging from 2 up to 1500 hours and then air cooled. This temperature was chosen since it was the highest possible without promoting a drastic dissolution of the ~' phase. Only the observations made on the samples which were subjected to the longest aging treatment will be presented in this paper, i.e. 1500 hours at 1100°C. The samples aged for shorter times showed a similar but less marked trend. They have been used for a study concerned with the kinetics of coarsening of the precipitates that is presented elsewhere (22).
a~ b
FIG. 1. Optical observations on sections perpendicular to the drawing axis of single crystal AM1 alloy, as cast ( a ) or after a standard homogenization treatment for 3 h. at 1300°C ( b )
Vo].
25,
No.
ii
y' PLATES IN A SUPERALLOY
/
interdendrltlc
1879
zone
) dendrite
core ~.~_....~"~secondary dendritic a r m ~ d
\
\
\
I
/
/
./
FIG. 2. Microstructures resulting from an aging treatment at 1100°0 for 1500 hours in the case of the initially as cast alloy. The different fields were observed perpendicular to the growth axis of the bar. Their location is indicated with respect to the solidification structure. l=xDerimental results The aging treatment at 1100°C for 1500 hours led to both coarsening and rafting of the 7' phase. However the -/' plate orientation was noted to be not always randomly distributed. Fig. 2 gives some examples of the microstructures observed in the case of the as cast alloy aged for 1500 hours at 1100°C. Observations were done at different places of a section perpendicular to the growth axis of the bar. As the chemical etching used also permitted observation of the initial primary trunks of the solidification structure, it was possible to detect that the local orientation of the "y' platelets was strong;y dependent on the position of the observed field with respect to the dendritic arms: in the middle of the secondary dendrite arms the 7' lamellae develop parallel to the growth direclion of the arm, the two families of plates coexist in the core of the primary axis while the precipitate shape appears to be much more erratic elsewhere. This phenomenon was observed for a lot of different
1880
Y' PLATES IN A SUPERALLOY
Vol.
23, No. ii
dendrite trunks on several sections of the sample. As it was previously mentioned, the same trends were already detected in the case of the samples aged for shorter times. Observations on sections parallel to the growth axis confirmed that the/' plates develop parallel to the secondary arms although the fields observed were less easily located with regard to the dendritic array than in the case of transverse sections.
a
b
FiG. 3. Microstructures resulting from an aging treatment at 11000C for 1500 hours in the case of the alloy initlal',y homogenized 3 hours at 1300°C; the fields a and b have the same location as fields b and d of Fig. 2 respectively. Fig. 3 presents the microstructure of the sample which was homogenized for 3 hours then aged. Local orientation of the plates was less marked but still noticeable. However, the preferential directions of the precipitates along the secondary dendritic arms were found to be just perpendicular to those observed in the previous case. This unexpected observation will be discussed later in this paper. In the case of the alloy homogenized for 100 hours, it was no longer possible to detect the solidification structure after the aging treatment. However in any field of observation, a microstructure similar to the one shown in Fig. 4 was observed: the precipitates also developed a plate-like morphology but without any preferential orientation. It has to be pointed out that the same type of aging treatment was also applied to two other superalloys of the AM1 family. A similar trend to local orientation was observed with identical relation to the solidification structure. This suggests that this phenomenon must be general in these alloys although it is not easy to detect because of the very different scales of the related structures.
FIG. 4. Microstructure resulting from an aging treatment at 1100°C for 1500 hours in the case of the alloy initially homogenized for 100 hours at 1300°C.
Vol.
23, No.
ii
y' PLATES
IN A S U P E R A L L O Y
1881
Discussion The reported observations clearly prove that an oriented development of the ~' plates can be obtained even in the absence of any mechanical driving force. Indeed with regard to the "soft" thermal history of the material used it seems unlikely that the phenomenon observed could be due to any internal residual stresses. Since the aged samples only differed by their initial degree of homogeneity, the development of locally oriented lamellae must be related to the internal chemical segregations remaining at the beginning of the aging treatment. As far as we know it is the first time that such a ~' rafting under the action of a chemical driving force is observed, although some authors have already noted such an influence of a diffusion flow on the ~' coalescence in the case of thermomechanical treatments (23). However the relation between segregations and rafting appears to be complex in regard with the apparently opposite microstructures of Fig. 2 and 3. These observations can however be understood if one considers that the effect of the short-time homogenization treatment is not only to smooth out the internal chemical segregations but also to change the shape of the isoconcentration curves. As it can be seen on Fig. 1, the segregations related to the tertiary dendritic arms completely disappeared during the 3 hours homogenizing treatment. This is confirmed by estimating the time t needed for microsegregations to be smoothed out through the use of the following relation : Ds • t/L2 = 1 [1] where Ds is the diffusion coefficient of the segregated element and L is the length over which the heterogeneities take place. Using the diffusion coefficient for a medium rate diffusing species as chromium which is Ds=6.10-5.exp(-30,900/T) m2.s -1 (24) and half the tertiary dendrite arm spacing for L, the time t for homogenization is found to be about 70 mn. As the secondary and tertiary dendrite arms grow perpendicular to each other, it is suggested that the 3 hour homogenizing treatment could have changed the direction of the main local chemical gradients and consequently the orientation of the y' platelets. However, confirmation of this assumption will require very precise characterization of the evolution of the microsegregations during the homogenizing treatment, using fine microprobe analysis and/or modeling of the diffusion path. Moreover it is surprising that the oriented rafting is preferentialy observed along the secondary dendrite arms while the internal chemical gradients are expected to be much higher elsewhere, for instance in the last solidified regions. This could be due to the fact that in these regions the main direction of the gradient was certainly not parallel to a <100> crystallographic direction or to the fact the y precipitate shapes were already erratic at the beginning of the aging treatment. Finally the question is to know in what way the internal chemical gradients can orientate the development of the ~' plates. On the one hand it may be that the diffusion flow, which should be operative during aging of an inhomogeneous alloy, could directly modify the equilibrium shape of the precipitates or lead to an anisotropy in their growth kinetics during Oswald ripening. Unfortunately, due to the lack of sufficient knowledge concerning the diffusion mecanisms in two-phases materials, it is difficult to give further developments to these approaches. On the other hand it can also be suggested that the chemical heterogeneities could indirectly induce the oriented rafting via their influence on the initial microstructure. As already noted, the local composition strongly determines the precipitation kinetics and morphology. In our case this results in an initial size, and perhaps also spatial, distribution of the precipitates related to the solidification structure. As the diffusion kinetics during Oswald ripening are directly related to the size difference between precipitates, it can be thought that the growth process would be enhanced in the direction of the main size gradient, i.e. of the main local chemical gradient. This suggests the possibility of progressive elongation in this direction. This assumption is supported by a recent numerical simulation of morphological development during 2-D Oswald ripening (25), where it is shown that initially circular precipitates can elongate while
1882
y'
PLATES
IN A SUPERALLOY
Vol.
23, No, ii
coarsening in the presence of a non isotropic size repartition. As a matter of fact, whatever the actual mechanism involved, is is possible that it would only need to be operative during the first stage of the aging treatment since the arrangement of plate-like precipitates parallel to each other has been shown to be very stable (9,10). Acknowledoements The authors are indebted to Ms E. GAUTIER, to Mr G. LESOULT and A. SIMON for many stimulating discussions during the preparation of this paper. Thanks are also due to SNECMA for providing the material. This study has been financially supported by the Centre National de la Recherche Sclentifique. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.
A.J. Ardell, Acta Metall. 20,61 (1972) C.K.L. Davies, P. Nash and R.N. Stevens, Acta Metail. 28,179 (1980) P.K. Footner and B.P. Richards, J. of Mat. Sci. 17, 2141 (1982) D. Mac Lean, Metal Scl. 18, 249 (1984) Y. Enemoto, M. Tokuyama and K. Kawazaki, Acta Metall. 34, 2119 (1986) S. Sadiq and D.R.F. West, Scripta Metall. 19, 833 (1985) T. Miyazaki, H. Imamura and T. Kozakai, Mat. Sci. Eng. 54, 9 (1982) M. Doi, T. Miyazaki and T. Wakatsuki, Mat. Sci. Eng. 67, 247 (1984) M. Doi and T. Miyazaki, Mat. Sci. Eng. 78, 87 (1986) T. Miyazaki, H. Imamura, H. Mori and T. Kozakai, Jour. Mat. Sci. 16, 1197 (1981) R.A. Ricks, A.J. Porter and R.C. Ecob, Acta Metall. 31, 43 (1983) W.C. Johnson and P.W. Voorhees, J. of Appl. Phys. 61 (4), 1610 (1987) M. Cornet and G. Martin, Scripta Metall. 21, 1091 (1987) A.G. Khachaturyan, S.V. Semenovskaya and J.W. Morris Jr, Acta Metall. 36, 1563 (1988) S.M. Foster, T.A. Nielsen and P. Nagy, "Superalloys 88" Proc. of the 6th Syrup. on Superalloys, Seven Springs, 245 (1988) 16. M.V. Nathal,'Metall. Trans. 18A, 1961 (1987) 17. M.V. Nathal, R.A. Mac Kay and R.G. Garlick, Scripta Metallo, 22, 1421 (1988) 18. French Patent n° 2.557.598, SNECMA-Imphy-Armines-ONERA (1983) 19. A. Fredholm, Doc. Eng. Thesis, Ecole des Mines de Paris, Jan. 1987 20. T. Grosdidier, D.E.A. Institut Polytechnique de Lorraine, June 1988 21. A. Hazotte and A. Simon, Proc. of the 1st ASM Europe Tech. Conf. on "Advanced Techniques for Structural Applications", Paris, 83 (1987) 22. A. Hazotte, T. Grosdidier and A. Simon, Proc. of the "Europ. Conf. on Advanced Materials and Process" (EUROMAT 89), Aachen, Nov. 22-24 1989 23. T. Massart, A. Morrison, JoL. Koutny and Y. Bienvenu, Proc. of the Conf. "4e Journ~es Nationales du Soudage', 164 (1989) 24. Smithells Metals Reference Book, sixth edition, E.A. BRANDES Editor, Butterworths 25. P.W. Voorhees, G.B. Mac Fadden, R.F. Boisvert and D.I. Meiron, Acta Metail.,36, 207(1988)
Vol. 23, pp. 1877-1882, 1989 Printed in the U.S.A.
Pergamon Press plc All rights reserved
CHEMICALLY ORIENTED ~ PLATE DEVELOPMENT IN A NICKEL BASE SUPERALLOY. Alain HAZOTTE and Jacques LACAZE Laboratoire de Science et G~nie des MatOriaux M~talliques Ecole des Mines -Parc de Saurupt - 54042 NANCY - FRANCE (Received August l, 1989) (Revised August 29, 1989)
Introduction The 1'-7' microstructure of nickel base superalloys can encounter drastic changes during high temperature heat treatments. The morphology of the 1" precipitates evolves from different mechanisms: i) competitive coarsening in order to reduce the specific area of the 1'-7' interface (Oswald ripening) (1-5) and ii) shape changes in order to minimize the sum of interfacial and elastic interaction energies. Different theoretical approaches forecast successive equilibrium shapes of the 7' precipitates depending not only on their size but also on the characteristics of the 1'-7' interface (6-14). Most of these works led to the conclusion that the precipitates should evolve towards a plate-like morphology when reaching a size of several microns. Although these predictions have been numerically established in the case of isolated groups of precipitates, there is some experimental evidence of this "rafting" in the case of alloys with a high amount of strengthening phase (6,15-17). The largest faces of the plates develop along the {100} planes of the crystallographic lattice common to both the 1'and y' phases. In the absence of any external applied stress the orientation of the 11' plates would have been thought to be evenly distributed along the three possible {100} planes. This report deals with experimental observations of preferentially oriented plate development of the 7' precipitates that take place without any applied stress. The experiments reported here concern a single grain commercial superalloy which has been submitted to long time aging treatments at 1100°C. When the material was not perfectly chemically homogeneous initially, it was noticed that locally oriented 1~ rafts developed on the dendritic scale, and that their orientation was related to the dendritic structure. It is proposed that this phenomenon is induced by the local chemical heterogeneities due to the usual solidification process. Some mechanisms which could link together chemical gradient and plate development are discussed. Experimental Procedure Experiments were performed on the commercial alloy AM1 ( Ni base-7.5wt% Cr, 5.2 AI, 5.5 W, 8.6 Ta, 6.5 Co, 1.2 Ti, 2 Mo )(18). Single grain bars with a diameter of 14 mm were produced by directional solidification in the SNECMA foundry house. The growth direction was almost parallel to the [100] crystallographic direction. The withdrawal rate used was the usual one for manufacturing turbine blades (4 mm/min.). It involves a dendritic solidification front and leads to final mlcrosegregatlons In the as-cast product, that is chemical heterogeneities on the scale of the dendrites trunks and their dendritic arms. One usually performs high temperature heat treatments to reduce these chemical heterogeneities. In the case of nickel base superalloys of the second generation as is AM1, there exists a "window", that is a temperature interval between the solvus temperature of the 7' precipitates and the incipient melting
1877 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc
1878
~' PLATES
IN A SUPERALLOY
Vol.
23, No.
ii
temperature. The homogenizing treatments are made in this window, which is 1295-1310°C for alloy AM1 (lg). In the present work the effect of the initial chemical microsegregations was investigated by using materials homogenized for different times : i) as cast, without any homogenizing treatment, ii) after a heat treatment at 1300°C for 3 hours then air cooled, iii) after 100 hours at 1300°C then air cooled. It has been verified by microprobe analysis that microsegregations have completely disappeared after the third treatment while they are still detected after the second one which corresponds to the standard industrial treatment. Fig. 1 presents optical micrographs of samples when as-cast (l-a) and after the homogenizing treatment of 3 hours (l-b). The chemical etching allows the solidification structure to be seen thanks to the residual microsegregations : in Fig. 1-a it is possible to note the presence of tertiary dendrite arms that have a width of about 60 p.m. In the case of the samples homogenized 100 hours, the dendritic structure does not appear any more. About 60% of ~,' phase precipitates during the last cooling stage of the heat treatments. The mean size of ~ precipitates was always found to be around 0.4 t ~ . However, their actual size distribution obviously depended on the chemical homogeneity of the material since the local composition is known to strongly determine the precipitation kinetics (20,21). As a matter of fact the size of the precipitate increased from the cores of the dendrites to the interdendritic regions in the case of materials not fully homogenized. Their morphology also evolved from cuboidal in the cores to more erratic shapes in the interdendritic regions. The aging treatments were performed on samples cut from different bars and sealed under secondary vacuum in silica tubes. They were held at 1100°C _+ 5°C for different times ranging from 2 up to 1500 hours and then air cooled. This temperature was chosen since it was the highest possible without promoting a drastic dissolution of the ~' phase. Only the observations made on the samples which were subjected to the longest aging treatment will be presented in this paper, i.e. 1500 hours at 1100°C. The samples aged for shorter times showed a similar but less marked trend. They have been used for a study concerned with the kinetics of coarsening of the precipitates that is presented elsewhere (22).
a~ b
FIG. 1. Optical observations on sections perpendicular to the drawing axis of single crystal AM1 alloy, as cast ( a ) or after a standard homogenization treatment for 3 h. at 1300°C ( b )
Vo].
25,
No.
ii
y' PLATES IN A SUPERALLOY
/
interdendrltlc
1879
zone
) dendrite
core ~.~_....~"~secondary dendritic a r m ~ d
\
\
\
I
/
/
./
FIG. 2. Microstructures resulting from an aging treatment at 1100°0 for 1500 hours in the case of the initially as cast alloy. The different fields were observed perpendicular to the growth axis of the bar. Their location is indicated with respect to the solidification structure. l=xDerimental results The aging treatment at 1100°C for 1500 hours led to both coarsening and rafting of the 7' phase. However the -/' plate orientation was noted to be not always randomly distributed. Fig. 2 gives some examples of the microstructures observed in the case of the as cast alloy aged for 1500 hours at 1100°C. Observations were done at different places of a section perpendicular to the growth axis of the bar. As the chemical etching used also permitted observation of the initial primary trunks of the solidification structure, it was possible to detect that the local orientation of the "y' platelets was strong;y dependent on the position of the observed field with respect to the dendritic arms: in the middle of the secondary dendrite arms the 7' lamellae develop parallel to the growth direclion of the arm, the two families of plates coexist in the core of the primary axis while the precipitate shape appears to be much more erratic elsewhere. This phenomenon was observed for a lot of different
1880
Y' PLATES IN A SUPERALLOY
Vol.
23, No. ii
dendrite trunks on several sections of the sample. As it was previously mentioned, the same trends were already detected in the case of the samples aged for shorter times. Observations on sections parallel to the growth axis confirmed that the/' plates develop parallel to the secondary arms although the fields observed were less easily located with regard to the dendritic array than in the case of transverse sections.
a
b
FiG. 3. Microstructures resulting from an aging treatment at 11000C for 1500 hours in the case of the alloy initlal',y homogenized 3 hours at 1300°C; the fields a and b have the same location as fields b and d of Fig. 2 respectively. Fig. 3 presents the microstructure of the sample which was homogenized for 3 hours then aged. Local orientation of the plates was less marked but still noticeable. However, the preferential directions of the precipitates along the secondary dendritic arms were found to be just perpendicular to those observed in the previous case. This unexpected observation will be discussed later in this paper. In the case of the alloy homogenized for 100 hours, it was no longer possible to detect the solidification structure after the aging treatment. However in any field of observation, a microstructure similar to the one shown in Fig. 4 was observed: the precipitates also developed a plate-like morphology but without any preferential orientation. It has to be pointed out that the same type of aging treatment was also applied to two other superalloys of the AM1 family. A similar trend to local orientation was observed with identical relation to the solidification structure. This suggests that this phenomenon must be general in these alloys although it is not easy to detect because of the very different scales of the related structures.
FIG. 4. Microstructure resulting from an aging treatment at 1100°C for 1500 hours in the case of the alloy initially homogenized for 100 hours at 1300°C.
Vol.
23, No.
ii
y' PLATES
IN A S U P E R A L L O Y
1881
Discussion The reported observations clearly prove that an oriented development of the ~' plates can be obtained even in the absence of any mechanical driving force. Indeed with regard to the "soft" thermal history of the material used it seems unlikely that the phenomenon observed could be due to any internal residual stresses. Since the aged samples only differed by their initial degree of homogeneity, the development of locally oriented lamellae must be related to the internal chemical segregations remaining at the beginning of the aging treatment. As far as we know it is the first time that such a ~' rafting under the action of a chemical driving force is observed, although some authors have already noted such an influence of a diffusion flow on the ~' coalescence in the case of thermomechanical treatments (23). However the relation between segregations and rafting appears to be complex in regard with the apparently opposite microstructures of Fig. 2 and 3. These observations can however be understood if one considers that the effect of the short-time homogenization treatment is not only to smooth out the internal chemical segregations but also to change the shape of the isoconcentration curves. As it can be seen on Fig. 1, the segregations related to the tertiary dendritic arms completely disappeared during the 3 hours homogenizing treatment. This is confirmed by estimating the time t needed for microsegregations to be smoothed out through the use of the following relation : Ds • t/L2 = 1 [1] where Ds is the diffusion coefficient of the segregated element and L is the length over which the heterogeneities take place. Using the diffusion coefficient for a medium rate diffusing species as chromium which is Ds=6.10-5.exp(-30,900/T) m2.s -1 (24) and half the tertiary dendrite arm spacing for L, the time t for homogenization is found to be about 70 mn. As the secondary and tertiary dendrite arms grow perpendicular to each other, it is suggested that the 3 hour homogenizing treatment could have changed the direction of the main local chemical gradients and consequently the orientation of the y' platelets. However, confirmation of this assumption will require very precise characterization of the evolution of the microsegregations during the homogenizing treatment, using fine microprobe analysis and/or modeling of the diffusion path. Moreover it is surprising that the oriented rafting is preferentialy observed along the secondary dendrite arms while the internal chemical gradients are expected to be much higher elsewhere, for instance in the last solidified regions. This could be due to the fact that in these regions the main direction of the gradient was certainly not parallel to a <100> crystallographic direction or to the fact the y precipitate shapes were already erratic at the beginning of the aging treatment. Finally the question is to know in what way the internal chemical gradients can orientate the development of the ~' plates. On the one hand it may be that the diffusion flow, which should be operative during aging of an inhomogeneous alloy, could directly modify the equilibrium shape of the precipitates or lead to an anisotropy in their growth kinetics during Oswald ripening. Unfortunately, due to the lack of sufficient knowledge concerning the diffusion mecanisms in two-phases materials, it is difficult to give further developments to these approaches. On the other hand it can also be suggested that the chemical heterogeneities could indirectly induce the oriented rafting via their influence on the initial microstructure. As already noted, the local composition strongly determines the precipitation kinetics and morphology. In our case this results in an initial size, and perhaps also spatial, distribution of the precipitates related to the solidification structure. As the diffusion kinetics during Oswald ripening are directly related to the size difference between precipitates, it can be thought that the growth process would be enhanced in the direction of the main size gradient, i.e. of the main local chemical gradient. This suggests the possibility of progressive elongation in this direction. This assumption is supported by a recent numerical simulation of morphological development during 2-D Oswald ripening (25), where it is shown that initially circular precipitates can elongate while
1882
y'
PLATES
IN A SUPERALLOY
Vol.
23, No, ii
coarsening in the presence of a non isotropic size repartition. As a matter of fact, whatever the actual mechanism involved, is is possible that it would only need to be operative during the first stage of the aging treatment since the arrangement of plate-like precipitates parallel to each other has been shown to be very stable (9,10). Acknowledoements The authors are indebted to Ms E. GAUTIER, to Mr G. LESOULT and A. SIMON for many stimulating discussions during the preparation of this paper. Thanks are also due to SNECMA for providing the material. This study has been financially supported by the Centre National de la Recherche Sclentifique. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.
A.J. Ardell, Acta Metall. 20,61 (1972) C.K.L. Davies, P. Nash and R.N. Stevens, Acta Metail. 28,179 (1980) P.K. Footner and B.P. Richards, J. of Mat. Sci. 17, 2141 (1982) D. Mac Lean, Metal Scl. 18, 249 (1984) Y. Enemoto, M. Tokuyama and K. Kawazaki, Acta Metall. 34, 2119 (1986) S. Sadiq and D.R.F. West, Scripta Metall. 19, 833 (1985) T. Miyazaki, H. Imamura and T. Kozakai, Mat. Sci. Eng. 54, 9 (1982) M. Doi, T. Miyazaki and T. Wakatsuki, Mat. Sci. Eng. 67, 247 (1984) M. Doi and T. Miyazaki, Mat. Sci. Eng. 78, 87 (1986) T. Miyazaki, H. Imamura, H. Mori and T. Kozakai, Jour. Mat. Sci. 16, 1197 (1981) R.A. Ricks, A.J. Porter and R.C. Ecob, Acta Metall. 31, 43 (1983) W.C. Johnson and P.W. Voorhees, J. of Appl. Phys. 61 (4), 1610 (1987) M. Cornet and G. Martin, Scripta Metall. 21, 1091 (1987) A.G. Khachaturyan, S.V. Semenovskaya and J.W. Morris Jr, Acta Metall. 36, 1563 (1988) S.M. Foster, T.A. Nielsen and P. Nagy, "Superalloys 88" Proc. of the 6th Syrup. on Superalloys, Seven Springs, 245 (1988) 16. M.V. Nathal,'Metall. Trans. 18A, 1961 (1987) 17. M.V. Nathal, R.A. Mac Kay and R.G. Garlick, Scripta Metallo, 22, 1421 (1988) 18. French Patent n° 2.557.598, SNECMA-Imphy-Armines-ONERA (1983) 19. A. Fredholm, Doc. Eng. Thesis, Ecole des Mines de Paris, Jan. 1987 20. T. Grosdidier, D.E.A. Institut Polytechnique de Lorraine, June 1988 21. A. Hazotte and A. Simon, Proc. of the 1st ASM Europe Tech. Conf. on "Advanced Techniques for Structural Applications", Paris, 83 (1987) 22. A. Hazotte, T. Grosdidier and A. Simon, Proc. of the "Europ. Conf. on Advanced Materials and Process" (EUROMAT 89), Aachen, Nov. 22-24 1989 23. T. Massart, A. Morrison, JoL. Koutny and Y. Bienvenu, Proc. of the Conf. "4e Journ~es Nationales du Soudage', 164 (1989) 24. Smithells Metals Reference Book, sixth edition, E.A. BRANDES Editor, Butterworths 25. P.W. Voorhees, G.B. Mac Fadden, R.F. Boisvert and D.I. Meiron, Acta Metail.,36, 207(1988)