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Effects of creep on the microstructural features of an oxide-dispersion-strengthened superalloy
Materials Science and Engineering, A 131 ( 1991 ) 1-7
I
Effects of Creep on the Microstructural Features of an Oxide-dispersionstrengthened Superailoy T. KEHAGIAS and L. COHEUR
Materials Development Department, SCK-CEN, B-2400 Mol (Belgium) E DELAVIGNETTE
Service of Surface Physics, ULB, B-1050 Brussels (Belgium) (Received February 26, 1990; in revised form May 30, 1990)
Abstract
The effects of high temperature creep on the microstructure and the grain boundaries of the oxide-dispersion-strengthened superalloy M A 6000 are investigated by means of transmission electron microscopy. Rafting of the y' phase is observed in the elongated grains, whereas in the small nonrecrystallized grains y' rafting does not occur. Dislocations form networks around ~' precipitates only in particular areas of the coarse grains where the oxide particles are absent. In the rest of the areas, where oxide dispersion is homogeneous, dislocations are pinned to the hard oxide particles instead of accumulating at Y-Y' interfaces. Dislocation networks around the y' phase are not observed in the small grains. There is no detectable influence of high temperature creep on the grain boundaries of the material. The majority of them are of the small angle type and exhibit a very disordered morphology. Large angle grain boundaries are randomly oriented and they are mainly decorated by chromium carbides.
I. Introduction
During the last decade modern oxide-dispersion-strengthened (ODS) superalloys have received a lot of interest for use in high temperature applications. This is due to the exceptional high temperature mechanical properties of these alloys, compared with other conventional superalloys. Precipitation and dispersion hardening are two of the most important strengthening mechanisms operating in ODS alloys, which provide them an excellent strength over a wide range of temperatures. 0921-5093/91/$3.50
The nickel-base Inconel (Inconel is a trademark of the INCO family of companies) superalloy MA6000 has undergone a zone annealing treatment in order to obtain a coarse elongated grain structure. On the basis of numerous determinations, a preferential grain orientation was established [1]: the coarse grains share a (110)_+ 5 ° axis parallel to the extrusion direction. The determinations were carried out using a method developed for the characterization of grain boundaries (GBs) by transmission electron microscopy (TEM) [2]. Along this direction the grains are several millimetres long and the typical grain aspect ratio (GAR) of the alloy investigated is of the order of 10 [3]. This high G A R accounts for an additional contribution to the resistance of the material at elevated temperatures [4]. The MA6000 contains a high volume fraction of 7' phase (L12 structure), about 50% in the asreceived material, which is coherent with the disordered 7 matrix. The dispersion of the incoherent, hard oxide particles is, usually, homogeneous throughout the metallic matrix. The MA6000 exhibits an anisotropic mechanical behaviour. During high temperature creep, the alloy sustains high stresses and can be plastically deformed along the longitudinal (extrusion) direction. On the contrary, along the long-transverse direction (normal to the elongation of the grains) it exhibits a rather brittle behaviour and fails at much lower stresses. An extensive study was undertaken in order to correlate this anisotropic behaviour with the microstructure and the GBs of the material. The study of GBs especially is of high interest, since they constitute potential damaging sources of the alloy during high temperature deformation. In this paper we present © Elsevier Sequoia/Printed in The Netherlands
observations on the effects of creep on the microstructural features of MA6000, by means of TEM.
2. Experimental details The Inconel MA6000 is manufactured, by IncoMAP, in flat billets of rectangular crosssection, by extrusion and hot rolling (Fig. 1). The nominal composition of the alloy is 15Cr4.5A1- 4 W - 2.5Ti- 2 M o - 2Ta- 0.15Zr- 0.05C0.01B-l.1Y203, in weight per cent, with the balance nickel. Three cylindrical bars were machined from a billet, having their axes parallel to the longitudinal (l) and the long-transverse (lt) directions, as shown in Fig. 1. These specimens were creep tested to failure in an air environment. All the tests were performed at 1223 K. The conditions and the results of creep testing are summarized in Table 1. In order to obtain specimens for TEM, thin foils were cut from the failed bars parallel to the T (transverse) and L (longitudinal) planes of the billet (Fig. 1). The foils were electropolished in a twin-jet Tenupol apparatus at 72 V, in medium flow rate, at room temperature. The electrolyte used for the polish-
ing was 5% perchloric acid in acetic acid. The perforated discs were examined in a JEM 200 CX JEOL microscope, equipped with a goniometer stage. An E D A X SW9100 instrument with a beryllium window was used for the energy-dispersive X-ray (EDX) analyses; the minimum probe size is of the order of 0.5/tm. The calculations of the characteristics of GBs and dislocations are supported by software developed for the analysis of GB coincidences [5].
3. Microstructural changes after creep
3.1. y' rafting A number of considerable changes occur after creep testing in the elongated grains. The shape of the ~' phase is gradually altered. The originally cuboidal y' particles (Fig. 2(a)) coalesce and form lamellae along a direction approximately perpendicular to the direction of the applied stress (Figs. 2(b) and 2(c)). The phenomenon has already been reported in the literature as y' rafting [4, 6-8], but the details of the formation of these rafts are yet to be revealed. It is rather evident from Fig. 2 that the lamellae of ~,' in the specimens where the elongation at failure was 6% (Fig. 2(c)) are longer and more irregular than those of the specimens which were elongated only 2.9% (Fig. 2(b)). In the case of the LT specimens (0.1% elongation at rupture), probably because of the rapid failure and the low elongation at failure, ~' rafting was not observed. However, the observations indicated an alignment of cuboidal y' particles, as they approach each other in order to form a lamellar structure.
3.2. Dislocation networks at Y-7' interfaces [tt] Fig. 1. Schematic diagram of a billet of the as-received MA6000: 1, longitudinal direction; lt, long-transverse direction; L, longitudinal plane; T, transverse plane; the shape of a specimen prepared for creep testing is shown on the T plane. TABLE 1 The creep testing conditions and results of the investigated alloy Specimen type a
Applied stress (MPa)
Time to failure (h)
Elongation at failure (%)
L1 (1) L2 (1) LT (It)
230 230 170
183.4 108.6 5.1
2.9 6.0 0.1
aDirections of the external applied stress in parentheses.
Several small oxide-depleted zones are frequently observed both in the as-received [1] and in the crept material. These zones are, generally, situated in the vicinity of GBs. Inside the oxidedepleted zones dislocation networks around Y' precipitates are observed (Fig. 3). Characterization of these dislocations shows that they are of a pure edge or predominantly edge type with Burgers' vectors of the ½(110) type. A method described by Amelinckx [9] was used as a basis for all the determinations of the dislocation characteristics. Similar observations by other researchers [10-14] have been reported to occur in many other nickel-base superalloys. Nevertheless, the kinetics and the role of these interracial dislocation networks are not yet clear.
Fig. 3. Dislocation network formation around 7' precipitates in an oxide-depleted zone; limited shearing of the 7' phase by dislocations is observed as well.
Fig. 4. Weak beam image (g-3g) of y' cutting by pairs of ~(110) screw dislocations; the originating array of dislocations can also be observed in the 7 matrix.
Fig. 2. Shape of the 7' phase in the coarse grains: (a) asreceived material; (b) specimen L1; (c) specimen L2; for (b) and (c) the direction of the external applied stress is perpendicular to the y' rafts.
Four types of dislocation interaction with second-phase particles can be considered [15]: Orowan bowing between the precipitates, dislocation bypass by particle shearing, cross-slip between particles and climb of dislocations over particles. The first two mechanisms are predominant during low temperature deformation,
whereas dislocation motion in high temperature creep is dominated by the last two mechanisms [1 1, 12]. In the present case no Orowan loops are detected among the failed specimens, in the oxide-depleted zones. Shearing of y' particles, although the tests were performed at high temperature, is occasionally observed (Fig. 3) and in some cases, in particular, pairs of ½(110) screw dislocations are cutting the 7' phase (Fig. 4). As shown in Fig. 4, a dislocation array produced by a dislocation source in the 7 matrix is paired just after it crosses the 7-7' interface. The dislocations of each pair have Burgers' vectors of the ½(110) type, the same as that of the originating array. The arrangement of dislocations around Y' in the micrographs indicates that cross-slip and dislocation climb are the two most probable
mechanisms of dislocation interaction with second-phase particles, during high temperature creep testing of MA6000. It is still unclear whether the dislocation networks around Y' precipitates are generated only in order to accommodate the misfit stress between the two phases [10-13]. It is claimed that the external applied stress contributes to their formation, as well [14]. In the present investigation the characterization of these dislocations did not lead to any firm conclusion concerning the influence of creep on the formation of the interfacial networks. The loss of coherency between 7 and 7' phases is a direct effect of these dislocation networks and it may affect the strength of the alloy [12-14]. The microhardness, for instance, is decreasing as a result of the increasing incoherency between Y and 7' [13]. Thus, the creep strength of the material in the oxide-depleted zones might be considerably lower. In particular, when these zones are close to GBs they become potential areas for an easier crack propagation. In the areas where oxide particles are present dislocations are interacting with these particles rather than forming networks at 7-Y' interfaces (Fig. 5). The refractory oxide particles act as barriers to the dislocation motion during high temperature deformation. It seems that, in addition to the direct interaction with the dislocations, the oxide particles can also contribute to the strength of the material at high temperatures, by retaining the 7' phase coherent with the 7 matrix. In Fig. 5 an oxide-depleted zone is crossing an area of an elongated grain in an L1 specimen. Within the limits of the zone, the 7' precipitates are surrounded by dislocations, whereas in the adjacent areas dislocations are pinned to the oxide particles.
3.3. The structure of small grains Rows of small grains are detected in the crept as well as in the as-received material [1]. They are residues of an incomplete secondary recrystallization of the alloy [16] and their sizes do not exceed a few micrometres in diameter. The small non-recrystallized grains can lead to crack initiation during high temperature creep [16]. In these grains, for all types of specimens, 7' rafting is not observed. The 7' particles retain their original cuboidal shape, but their apparent linear dimensions are increased by at least a factor of 2 (Fig. 6). Large parts of the small grains are denuded from oxide particles. In the crept state especially
Fig. 5. A small oxide-depleted zone in specimen L1; within the zone, dislocations accumulate at Y-7' interfaces, whereas in the adjacent areas the dislocations are interacting with the refractory oxide particles.
(Fig. 6(b)) a lack of fine oxide particles with mean diameters less than 15 nm is observed. In addition, no 7' bypass by dislocations occurs in the small grains. Generally, the dislocation density in these grains is far lower than that in the coarse grains. EDX analysis was used in order to determine any differences in the chemical composition between small and elongated grains. A rather high chromium content is detected in the small grains, compared with the quantity of chromium in the coarse grains. This result is most probably a consequence of the large number of chromium carbides observed in the GBs between the small grains, as will be described in Sections 4.1 and 4.2. 4. Grain boundary structure after creep
The influence of creep on the GBs of the material in both the elongated and the small grains is studied. Highly disordered morphologies of GBs are observed after creep testing. Although Timmins and Arzt [17] reported the formation of
Fig. 7. Morphology of a small angle GB in specimen L2; the direction of the external applied stress is parallel to the mean average direction of the boundary.
Fig. 6. Shape of the 7' phase in the small grains: (a) asreceived material; (b) specimen L2; the increase in the linear dimensions of the 7' particles after creep is evident, although their cuboidal shape is retained.
7'-free zones during high temperature creep, in the vicinity of GBs, none of these zones has been detected in this study.
Fig. 8. Morphology of a large angle GB in specimen LT; chromium carbides and small oxide-depleted grains are decorating the boundary.
4.1. Elongated grains
dislocations appear to have no preferential orientation direction, resulting in very disordered morphologies. GB plane orientation is not significant owing to the twisted form of the GBs. Nevertheless, an average GB plane can be determined and in this case was found parallel to the {111} crystallographic planes. All the small angle GBs which have been examined had a mixed character. The determinations of the low-angle GB characteristics were carried out in accordance with the method of Polychroniadis and Delavignette [18]. Large angle GBs between elongated grains are easily recognizable from their morphology: a large amount of carbides of various sizes and small oxide-depleted grains are present along them (Fig. 8). Selected area diffraction experi-
The great majority (about 80%) of the GBs between elongated grains are of the small angle type with rotation angles between 1° and 6 °. The same was also observed in the as-received material [1]. Few large angle GBs are detected, only among LT specimens. The GB structure of small angle GBs is highly perturbed (Fig. 7). They appear to be twisted all along their visible, in TEM, length. Almost none of the expected dislocation configurations are observed. However, in LT and L1 specimens limited dislocation networks, belonging to more than one family of dislocations, are present. On the basis of a limited amount of characterization these dislocations are found to be of the mixed type with Burgers' vectors b=½(ll0). In the majority of the GBs,
ments show that the carbides are mostly of the E D X analysis was used to determine M. It is mainly chromium, whereas rather high quantities of molybdenum and tungsten suggested that some of the chromium atoms are substituted by atoms of these two elements. M23C 6 type.
4.2. Small grains All the GBs between small grains or between a small grain and an elongated grain are of the large angle type. No coincidence site lattice boundaries are determined among them, except in some cases where twins are observed. The morphology of these large angle GBs is the same as that of the large angle GBs between elongated grains. Again a large amount of M23C 6 carbides (M mainly being chromium) decorates the GBs. The role of the chromium carbides appears rather important owing to their contradictory effects on the mechanical properties of the alloy. On the one hand, they improve the rupture strength by impeding grain boundary sliding during high temperature deformation. On the other hand, however, it has been suggested that frequently failure was initiated either by decohesion of the M23C6-~/ interface or by fracture of the M23C 6 particle [19]. Thus, M23C 6 carbides might behave as another possible source of crack initiation in addition to the sensitive small grains and the oxide-depleted zones. In this study no porosities or any other features which may give rise to crack initiation were observed.
5. Conclusions The occurrence of Y' rafting is observed only in the coarse, elongated grains. No formation of y' rafts is detected among the small grains, although the chemical composition appears to be the same in both cases. A n increase in the apparent linear dimensions of the cuboidal Y' particles in the small grains is observed. Dislocation networks are formed, during high temperature creep, around 7' particles at 7-7' interfaces only within oxide-depleted zones. These dislocations are predominantly of the edge type with Burgers' vectors b-- ½(110). Occasionally, sheafing of y' by pairs of ½(110) screw dislocations occurs as well. In the areas with oxide dispersion (more than 99% of the total volume) dislocations are interacting with the oxide particles and the loss of coherency between y and 7'
is delayed. In the small grains no formation of dislocation networks around the cuboidal y' particles is observed. There is no evident influence of the high temperature deformation on the GB structure of the material. The main characteristic of the small angle GBs is the highly disordered morphology, resulting from dislocations with no preferential orientation direction. Mixed type ½(110) dislocations are determined in the few detectable dislocation arrays. Small angle GBs have mixed character and an average {111 } GB plane. High concentrations of chromium carbides are detected along the large angle GBs between elongated or small grains. The large angle character of the GBs is easily identified by the accumulated carbides. No chromium carbides are observed on the small angle GBs of the material.
Acknowledgment This study was performed in the framework of the CEC Stimulation Contract ST 2J-0289-C.
References 1 Th. Kehagias, L. Coheur and P. Delavignette, J. Phys. Paris, Colloq. C1, 51 (1990)C1-855 2 P. Delavignette, Th. Karakostas, G. Nouet and E W. Schapink, Phys. Status Solidi A, 107 (1988) 551. 3 J. K. Tien, T. E. Howson and D. E. Matejczyk, in J. S. Benjamin (ed.), Frontiers of High Temperature Materials I, IncoMAP, New York, 1981, p. 13. 4 R. E Singer, R. C. Benn and S. K. Kang, in J. S. Benjamin and R. C. Benn (eds.), Frontiers of High Temperature MaterialslI, IncoMAE New York, 1983, p. 336. 5 E Delavignette, J. Phys. Paris, Colloq. C6, 43 (1982) C6-1. 6 M.V. Nathal and L. J. Ebert, MetalL Trans. A,, 16 (1985) 427. 7 R. A. MacKay and L. J. Ebert, Metall. Trans. A, 16 (1985) 1969. 8 M.V. Nathal, Metall. Trans. A, 18 (1987) 1961. 9 S. Amelinckx, J. Cryst. Growth, 24-25 (1974) 6. 10 A. Lasalmonie and J. L. Strudel, Philos. Mag., 32 (1975) 937. 11 W.W. Milligan and S. D. Antolovich, Metall. Trans. A, 18 (1987) 85. 12 A. K. Singh, N. Louat and K. Sadananda, Metall. Trans. A, 19 (1988) 2965. 13 M. Feller-Kniepmeier and T. Link, Metall. Trans. A, 20 (1989) 1233. 14 T. P. Gabb, S. L. Draper, R. A. MacKay and M. V. Nathal, Mater. Sci. Eng., A l l 8 (1989) 59.
15 B. A. Wilcox and A. H. Clawer, in C. T. Sims and W. C. Hagel (eds.), The Superalloys, Wiley, New York, 1972, pp. 199-206. 16 H. Zeizinger and E. Arzt, Z. Metallkd., 79 (1988) 774. 17 R. Timmins and E. Arzt, Scr. Metall., 22 (1988) 1353.
18 E. K. Polychroniadis and P. Delavignette, Phys. Status SolidiA, 77(1983)291. 19 R. E Decker and C. T. Sims, in C. T. Sims and W. C. Hagel (eds.), The Superalloys, Wiley, New York, 1972, p. 55.
I
Effects of Creep on the Microstructural Features of an Oxide-dispersionstrengthened Superailoy T. KEHAGIAS and L. COHEUR
Materials Development Department, SCK-CEN, B-2400 Mol (Belgium) E DELAVIGNETTE
Service of Surface Physics, ULB, B-1050 Brussels (Belgium) (Received February 26, 1990; in revised form May 30, 1990)
Abstract
The effects of high temperature creep on the microstructure and the grain boundaries of the oxide-dispersion-strengthened superalloy M A 6000 are investigated by means of transmission electron microscopy. Rafting of the y' phase is observed in the elongated grains, whereas in the small nonrecrystallized grains y' rafting does not occur. Dislocations form networks around ~' precipitates only in particular areas of the coarse grains where the oxide particles are absent. In the rest of the areas, where oxide dispersion is homogeneous, dislocations are pinned to the hard oxide particles instead of accumulating at Y-Y' interfaces. Dislocation networks around the y' phase are not observed in the small grains. There is no detectable influence of high temperature creep on the grain boundaries of the material. The majority of them are of the small angle type and exhibit a very disordered morphology. Large angle grain boundaries are randomly oriented and they are mainly decorated by chromium carbides.
I. Introduction
During the last decade modern oxide-dispersion-strengthened (ODS) superalloys have received a lot of interest for use in high temperature applications. This is due to the exceptional high temperature mechanical properties of these alloys, compared with other conventional superalloys. Precipitation and dispersion hardening are two of the most important strengthening mechanisms operating in ODS alloys, which provide them an excellent strength over a wide range of temperatures. 0921-5093/91/$3.50
The nickel-base Inconel (Inconel is a trademark of the INCO family of companies) superalloy MA6000 has undergone a zone annealing treatment in order to obtain a coarse elongated grain structure. On the basis of numerous determinations, a preferential grain orientation was established [1]: the coarse grains share a (110)_+ 5 ° axis parallel to the extrusion direction. The determinations were carried out using a method developed for the characterization of grain boundaries (GBs) by transmission electron microscopy (TEM) [2]. Along this direction the grains are several millimetres long and the typical grain aspect ratio (GAR) of the alloy investigated is of the order of 10 [3]. This high G A R accounts for an additional contribution to the resistance of the material at elevated temperatures [4]. The MA6000 contains a high volume fraction of 7' phase (L12 structure), about 50% in the asreceived material, which is coherent with the disordered 7 matrix. The dispersion of the incoherent, hard oxide particles is, usually, homogeneous throughout the metallic matrix. The MA6000 exhibits an anisotropic mechanical behaviour. During high temperature creep, the alloy sustains high stresses and can be plastically deformed along the longitudinal (extrusion) direction. On the contrary, along the long-transverse direction (normal to the elongation of the grains) it exhibits a rather brittle behaviour and fails at much lower stresses. An extensive study was undertaken in order to correlate this anisotropic behaviour with the microstructure and the GBs of the material. The study of GBs especially is of high interest, since they constitute potential damaging sources of the alloy during high temperature deformation. In this paper we present © Elsevier Sequoia/Printed in The Netherlands
observations on the effects of creep on the microstructural features of MA6000, by means of TEM.
2. Experimental details The Inconel MA6000 is manufactured, by IncoMAP, in flat billets of rectangular crosssection, by extrusion and hot rolling (Fig. 1). The nominal composition of the alloy is 15Cr4.5A1- 4 W - 2.5Ti- 2 M o - 2Ta- 0.15Zr- 0.05C0.01B-l.1Y203, in weight per cent, with the balance nickel. Three cylindrical bars were machined from a billet, having their axes parallel to the longitudinal (l) and the long-transverse (lt) directions, as shown in Fig. 1. These specimens were creep tested to failure in an air environment. All the tests were performed at 1223 K. The conditions and the results of creep testing are summarized in Table 1. In order to obtain specimens for TEM, thin foils were cut from the failed bars parallel to the T (transverse) and L (longitudinal) planes of the billet (Fig. 1). The foils were electropolished in a twin-jet Tenupol apparatus at 72 V, in medium flow rate, at room temperature. The electrolyte used for the polish-
ing was 5% perchloric acid in acetic acid. The perforated discs were examined in a JEM 200 CX JEOL microscope, equipped with a goniometer stage. An E D A X SW9100 instrument with a beryllium window was used for the energy-dispersive X-ray (EDX) analyses; the minimum probe size is of the order of 0.5/tm. The calculations of the characteristics of GBs and dislocations are supported by software developed for the analysis of GB coincidences [5].
3. Microstructural changes after creep
3.1. y' rafting A number of considerable changes occur after creep testing in the elongated grains. The shape of the ~' phase is gradually altered. The originally cuboidal y' particles (Fig. 2(a)) coalesce and form lamellae along a direction approximately perpendicular to the direction of the applied stress (Figs. 2(b) and 2(c)). The phenomenon has already been reported in the literature as y' rafting [4, 6-8], but the details of the formation of these rafts are yet to be revealed. It is rather evident from Fig. 2 that the lamellae of ~,' in the specimens where the elongation at failure was 6% (Fig. 2(c)) are longer and more irregular than those of the specimens which were elongated only 2.9% (Fig. 2(b)). In the case of the LT specimens (0.1% elongation at rupture), probably because of the rapid failure and the low elongation at failure, ~' rafting was not observed. However, the observations indicated an alignment of cuboidal y' particles, as they approach each other in order to form a lamellar structure.
3.2. Dislocation networks at Y-7' interfaces [tt] Fig. 1. Schematic diagram of a billet of the as-received MA6000: 1, longitudinal direction; lt, long-transverse direction; L, longitudinal plane; T, transverse plane; the shape of a specimen prepared for creep testing is shown on the T plane. TABLE 1 The creep testing conditions and results of the investigated alloy Specimen type a
Applied stress (MPa)
Time to failure (h)
Elongation at failure (%)
L1 (1) L2 (1) LT (It)
230 230 170
183.4 108.6 5.1
2.9 6.0 0.1
aDirections of the external applied stress in parentheses.
Several small oxide-depleted zones are frequently observed both in the as-received [1] and in the crept material. These zones are, generally, situated in the vicinity of GBs. Inside the oxidedepleted zones dislocation networks around Y' precipitates are observed (Fig. 3). Characterization of these dislocations shows that they are of a pure edge or predominantly edge type with Burgers' vectors of the ½(110) type. A method described by Amelinckx [9] was used as a basis for all the determinations of the dislocation characteristics. Similar observations by other researchers [10-14] have been reported to occur in many other nickel-base superalloys. Nevertheless, the kinetics and the role of these interracial dislocation networks are not yet clear.
Fig. 3. Dislocation network formation around 7' precipitates in an oxide-depleted zone; limited shearing of the 7' phase by dislocations is observed as well.
Fig. 4. Weak beam image (g-3g) of y' cutting by pairs of ~(110) screw dislocations; the originating array of dislocations can also be observed in the 7 matrix.
Fig. 2. Shape of the 7' phase in the coarse grains: (a) asreceived material; (b) specimen L1; (c) specimen L2; for (b) and (c) the direction of the external applied stress is perpendicular to the y' rafts.
Four types of dislocation interaction with second-phase particles can be considered [15]: Orowan bowing between the precipitates, dislocation bypass by particle shearing, cross-slip between particles and climb of dislocations over particles. The first two mechanisms are predominant during low temperature deformation,
whereas dislocation motion in high temperature creep is dominated by the last two mechanisms [1 1, 12]. In the present case no Orowan loops are detected among the failed specimens, in the oxide-depleted zones. Shearing of y' particles, although the tests were performed at high temperature, is occasionally observed (Fig. 3) and in some cases, in particular, pairs of ½(110) screw dislocations are cutting the 7' phase (Fig. 4). As shown in Fig. 4, a dislocation array produced by a dislocation source in the 7 matrix is paired just after it crosses the 7-7' interface. The dislocations of each pair have Burgers' vectors of the ½(110) type, the same as that of the originating array. The arrangement of dislocations around Y' in the micrographs indicates that cross-slip and dislocation climb are the two most probable
mechanisms of dislocation interaction with second-phase particles, during high temperature creep testing of MA6000. It is still unclear whether the dislocation networks around Y' precipitates are generated only in order to accommodate the misfit stress between the two phases [10-13]. It is claimed that the external applied stress contributes to their formation, as well [14]. In the present investigation the characterization of these dislocations did not lead to any firm conclusion concerning the influence of creep on the formation of the interfacial networks. The loss of coherency between 7 and 7' phases is a direct effect of these dislocation networks and it may affect the strength of the alloy [12-14]. The microhardness, for instance, is decreasing as a result of the increasing incoherency between Y and 7' [13]. Thus, the creep strength of the material in the oxide-depleted zones might be considerably lower. In particular, when these zones are close to GBs they become potential areas for an easier crack propagation. In the areas where oxide particles are present dislocations are interacting with these particles rather than forming networks at 7-Y' interfaces (Fig. 5). The refractory oxide particles act as barriers to the dislocation motion during high temperature deformation. It seems that, in addition to the direct interaction with the dislocations, the oxide particles can also contribute to the strength of the material at high temperatures, by retaining the 7' phase coherent with the 7 matrix. In Fig. 5 an oxide-depleted zone is crossing an area of an elongated grain in an L1 specimen. Within the limits of the zone, the 7' precipitates are surrounded by dislocations, whereas in the adjacent areas dislocations are pinned to the oxide particles.
3.3. The structure of small grains Rows of small grains are detected in the crept as well as in the as-received material [1]. They are residues of an incomplete secondary recrystallization of the alloy [16] and their sizes do not exceed a few micrometres in diameter. The small non-recrystallized grains can lead to crack initiation during high temperature creep [16]. In these grains, for all types of specimens, 7' rafting is not observed. The 7' particles retain their original cuboidal shape, but their apparent linear dimensions are increased by at least a factor of 2 (Fig. 6). Large parts of the small grains are denuded from oxide particles. In the crept state especially
Fig. 5. A small oxide-depleted zone in specimen L1; within the zone, dislocations accumulate at Y-7' interfaces, whereas in the adjacent areas the dislocations are interacting with the refractory oxide particles.
(Fig. 6(b)) a lack of fine oxide particles with mean diameters less than 15 nm is observed. In addition, no 7' bypass by dislocations occurs in the small grains. Generally, the dislocation density in these grains is far lower than that in the coarse grains. EDX analysis was used in order to determine any differences in the chemical composition between small and elongated grains. A rather high chromium content is detected in the small grains, compared with the quantity of chromium in the coarse grains. This result is most probably a consequence of the large number of chromium carbides observed in the GBs between the small grains, as will be described in Sections 4.1 and 4.2. 4. Grain boundary structure after creep
The influence of creep on the GBs of the material in both the elongated and the small grains is studied. Highly disordered morphologies of GBs are observed after creep testing. Although Timmins and Arzt [17] reported the formation of
Fig. 7. Morphology of a small angle GB in specimen L2; the direction of the external applied stress is parallel to the mean average direction of the boundary.
Fig. 6. Shape of the 7' phase in the small grains: (a) asreceived material; (b) specimen L2; the increase in the linear dimensions of the 7' particles after creep is evident, although their cuboidal shape is retained.
7'-free zones during high temperature creep, in the vicinity of GBs, none of these zones has been detected in this study.
Fig. 8. Morphology of a large angle GB in specimen LT; chromium carbides and small oxide-depleted grains are decorating the boundary.
4.1. Elongated grains
dislocations appear to have no preferential orientation direction, resulting in very disordered morphologies. GB plane orientation is not significant owing to the twisted form of the GBs. Nevertheless, an average GB plane can be determined and in this case was found parallel to the {111} crystallographic planes. All the small angle GBs which have been examined had a mixed character. The determinations of the low-angle GB characteristics were carried out in accordance with the method of Polychroniadis and Delavignette [18]. Large angle GBs between elongated grains are easily recognizable from their morphology: a large amount of carbides of various sizes and small oxide-depleted grains are present along them (Fig. 8). Selected area diffraction experi-
The great majority (about 80%) of the GBs between elongated grains are of the small angle type with rotation angles between 1° and 6 °. The same was also observed in the as-received material [1]. Few large angle GBs are detected, only among LT specimens. The GB structure of small angle GBs is highly perturbed (Fig. 7). They appear to be twisted all along their visible, in TEM, length. Almost none of the expected dislocation configurations are observed. However, in LT and L1 specimens limited dislocation networks, belonging to more than one family of dislocations, are present. On the basis of a limited amount of characterization these dislocations are found to be of the mixed type with Burgers' vectors b=½(ll0). In the majority of the GBs,
ments show that the carbides are mostly of the E D X analysis was used to determine M. It is mainly chromium, whereas rather high quantities of molybdenum and tungsten suggested that some of the chromium atoms are substituted by atoms of these two elements. M23C 6 type.
4.2. Small grains All the GBs between small grains or between a small grain and an elongated grain are of the large angle type. No coincidence site lattice boundaries are determined among them, except in some cases where twins are observed. The morphology of these large angle GBs is the same as that of the large angle GBs between elongated grains. Again a large amount of M23C 6 carbides (M mainly being chromium) decorates the GBs. The role of the chromium carbides appears rather important owing to their contradictory effects on the mechanical properties of the alloy. On the one hand, they improve the rupture strength by impeding grain boundary sliding during high temperature deformation. On the other hand, however, it has been suggested that frequently failure was initiated either by decohesion of the M23C6-~/ interface or by fracture of the M23C 6 particle [19]. Thus, M23C 6 carbides might behave as another possible source of crack initiation in addition to the sensitive small grains and the oxide-depleted zones. In this study no porosities or any other features which may give rise to crack initiation were observed.
5. Conclusions The occurrence of Y' rafting is observed only in the coarse, elongated grains. No formation of y' rafts is detected among the small grains, although the chemical composition appears to be the same in both cases. A n increase in the apparent linear dimensions of the cuboidal Y' particles in the small grains is observed. Dislocation networks are formed, during high temperature creep, around 7' particles at 7-7' interfaces only within oxide-depleted zones. These dislocations are predominantly of the edge type with Burgers' vectors b-- ½(110). Occasionally, sheafing of y' by pairs of ½(110) screw dislocations occurs as well. In the areas with oxide dispersion (more than 99% of the total volume) dislocations are interacting with the oxide particles and the loss of coherency between y and 7'
is delayed. In the small grains no formation of dislocation networks around the cuboidal y' particles is observed. There is no evident influence of the high temperature deformation on the GB structure of the material. The main characteristic of the small angle GBs is the highly disordered morphology, resulting from dislocations with no preferential orientation direction. Mixed type ½(110) dislocations are determined in the few detectable dislocation arrays. Small angle GBs have mixed character and an average {111 } GB plane. High concentrations of chromium carbides are detected along the large angle GBs between elongated or small grains. The large angle character of the GBs is easily identified by the accumulated carbides. No chromium carbides are observed on the small angle GBs of the material.
Acknowledgment This study was performed in the framework of the CEC Stimulation Contract ST 2J-0289-C.
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