On the shearing mechanism of g 0 precipitates by a single (a/6) h112i Shockley partial in Ni-based superalloys

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1 Philosophical Magazine, 1January 004 Vol. 84, No. 1, On the shearing mechanism of g 0 precipitates by a single (a/6) h11i Shockley partial in Ni-based superalloys B. De campsy Laboratoire de Chimie Me tallurgique des Terres Rares, Unite Propre de Recherche associe e au CNRS 09, 8 rue Henri Dunant, 940 Thiais Cedex, France S. Raujol, A. Coujou, F. Pettinari-Sturmel, N. Cle ment Centre d Elaboration de Mate riaux et d Etudes Structurales, Laboratoire d Optique Electronique, CNRS, 9 rue Jeanne Marvig, 1055 Toulouse Cedex, France D. Locq and P. Caron Office National d Etudes et de Recherches Ae rospatiales, BP 7, 9 Chaˆ tillon, France [Received 9 May 00 and accepted in revised form 0 August 00] Abstract The weak-beam technique of transmission electron microscopy has been used to analyse a new shearing configuration of g 0 precipitates after creep at 700 C of a Ni-based superalloy for gas turbine discs. The shearing configurations are made up of superlattice extrinsic stacking faults, matrix stacking faults and individual (a/6) h11i Shockley dislocations. This mechanism is initiated by the decorrelated movement of the two Shockley partials of a single (a/) h110i matrix dislocation. The propagation of the leading partial creates this shearing process. This phenomenon that occurs in small g channels owing to the flexibility of dislocations can be used to evaluate microstructural evolutions during ageing in the alloy. } 1. Introduction The gas turbine discs for future supersonic aircraft require the use of Ni-based superalloys having long-term and elevated creep properties at temperatures up to 700 C. Because of the large dimensions of the discs, only the powder metallurgy route allows one to obtain components with the required microstructure and mechanical properties. The NR powder metallurgical Ni-based superalloy was thus designed with the objective of attaining high tensile and creep properties at 700 C, while avoiding the precipitation of undesirable brittle topologically close-packed phases during long-term exposures at high temperatures (Locq et al. 000). y brigitte.decamps@glvt-cnrs.fr. Philosophical Magazine ISSN print/issn online # 004 Taylor & Francis Ltd DOI: /

2 9 B. De camps et al. This alloy is strengthened by a bimodal dispersion of g 0 -Ni (Al, Ti) particles (L1 structure) which precipitate within a fcc Ni-based solid solution. Such a dispersion of g 0 precipitates is very sensitive to ageing during service at temperatures above 650 C. Coarsening of the largest precipitates and partial solutioning of the finest precipitates can thus influence the nature of the creep-controlling deformation mechanisms. So the analysis of the micromechanisms controlling the deformation can bring some useful information about the microstructural evolution of the alloy during long-term creep. A number of studies devoted to the Ni-based superalloys suited for single-crystal blade applications have shown that, at temperatures around 700 C, the principal mode of shearing of g 0 precipitates results in the generation of superlattice stacking faults within these strengthening particles. Kear and co-workers (Kear et al. 1968, 1969a, b, c, Leverant and Kear 1970, Kear and Oblak 1974) have proposed a mechanism for the formation of extended configurations of intrinsic extrinsic fault pairs with a net ah11i Burgers vector during primary creep at 760 C of Mar-M00 single crystals. Such shearing configurations were also observed during creep at 760 C of CMSX- single crystals containing a monomodal distribution of g 0 precipitates with a mean size of 0 nm (Caron et al. 1988b), and more recently in Re-containing single-crystal superalloys creep tested at 750 C (Rae et al. 00). On the other hand, on the basis of observations made using the weak-beam technique of electron microscopy (Huis in t Veld et al. 1985, De camps and Condat 1986, Condat and Décamps 1987, Caron et al. 1988, De camps et al. 1991a, b, 1994), (a/) h110i matrix dislocations, moving in g matrix channels between g 0 precipitates, are shown to be involved in the shearing configurations of g 0 particles. On the basis of a complete analysis of shearing configurations, Condat and Décamps (1987) have proposed an original mechanism of shearing of precipitates by a single (a/) h110i matrix dislocation involving nucleation of a Shockley dislocation within the antiphase boundary (APB) to form a superlattice intrinsic stacking fault (SISF). Later, to explain the dependence of the shearing configurations (SISFs and/or superlattice extrinsic stacking faults (SESFs)) on the mechanical testing, De camps et al. (1991b) proposed an evolution of the shearing mechanism starting from a dissociated matrix dislocation crossing the interface with or without inversion of the partials to explain the formation of a SISF or SESF by nucleation of a Shockley dislocation within the APB. It was only with high-resolution experiments that Décamps et al. (1994) showed that shearing configurations involving SESFs requires the nucleation of the Shockley dislocation in an adjacent plane above or under the complex fault (CF) plane and not within the APB. The present paper gives some experimental evidence of another mechanism derived from this previous model initiated by (a/6) h11imatrix Shockley dislocations involving SESFs. The origin of this mechanism with respect to the precipitate distribution and morphology, and its use in the analysis of the microstructural evolution of the NR superalloy, will be discussed. }. Material and experimental procedure The samples and specimens used in this study were machined from pancakes of NR superalloy produced by Tecphy and Snecma using the following industrial route: (i) vacuum-induction-melted ingot; (ii) argon atomization;

3 Shearing mechanism of g 0 precipitates 9 Table 1. Weight and atomic percentage of the NR alloy. Ni Co Cr Ti Al Mo Hf Zr C B Amount (wt.%) Amount (at.%) (iii) powder sieving ( 00 mesh); (iv) hot extrusion; (v) isothermal forging. The composition of the NR alloy is given in table 1. A g 0 supersolvus heat treatment was applied for h at 110 C, with air cooling at 100 C min 1, to produce a coarse-grained microstructure that is known to be beneficial to the high temperature creep strength (Raisson and Davidson 1990). The final ageing treatments were for 4 h at 700 C, and then for 4 h at 800 C, both with air cooling. The resulting standard microstructure is characterized by a mean grain size of about 50 mm (figure 1 (a)), and a bimodal distribution of g 0 precipitates (figure 1 (b)). The coarse secondary precipitates have a mean size of 40 nm, while the fine tertiary precipitates have an average size of 5 nm. Primary g 0 particles were dissolved during the supersolvus solution heat treatment. Tensile creep tests were performed on cylindrical specimens in air at 700 C and 650 MPa. For the transmission electron microscopy (TEM) study, the tests were interrupted after a creep strain of about 0.%. The specimens were then forced air cooled under stress. Some creep tests and TEM analyses were also performed on NR specimens on which were applied different ageing heat treatment procedures in order to modify the g g 0 microstructure and to evaluate the resulting change in creep-controlling mechanism. In one case, the two-stage standard ageing heat treatment was replaced by a single ageing heat treatment for 4 h at 700 C. In comparison with the standard microstructure, the main difference is a reduction of the size of the tertiary g 0 precipitates down to 7 nm. In another case, the standard ageing heat treatments were followed by a long-term exposure for 500 h at 800 C aimed at solutionizing all the tertiary g 0 precipitates (Raujol et al. 00). Samples were cut normally to the tensile axis of the crept specimens and thinned by electropolishing to be observed in a transmission electron microscope (JEOL 010 and 000 EX) under bright-field (Hirsch et al. 1977) and weak-beam (Cockayne et al. 1969) conditions. The Burgers vectors of dislocations both in the matrix and in the g 0 precipitates are described throughout the paper with the finish start/right-handed convention (Bilby et al. 1955) using the Thompson (195) notation. The Thompson tetrahedron used in the paper is represented in figure. In the Thompson notation for the fcc structure, a dissociated dislocation BD in the a plane (dissociated into the partials Ba þ ad separated by an intrinsic fault) has a right-hand component Ba if the dislocation is viewed from outside the tetrahedron and along the direction of the dislocation line (figure (a)). In the L1 structure, the Thompson notation applies but the rules are inverted (Kear and Oblak 1974) for a superlattice dislocation BD dissociated into two partial dislocations Ba þ ad

4 94 B. De camps et al. Figure 1. Microstructures after the standard heat treatment: (a) Optical micrograph showing the microstructure. Large grains (average size of 50 mm) are visible. (b) Transmission electron micrograph showing the bimodal distribution with secondary precipitates (about 40 nm) and tertiary precipitates (about 5 nm). separated by a SISF, the right-hand component is ad (figure (b)). For a SESF, the right-hand component is Ba (figure (c)). On this basis, the nature of the superlattice stacking fault is determined when the Burgers vector of the dislocation bounding the fault is known. Burgers vectors are analysed using the invisibility criterion: jg bj ¼ 0 (Howie and Whelan 196).

5 Shearing mechanism of g 0 precipitates 95 Figure. Schematic representation of the Thompson tetrahedron used in the paper. D B A below C αd I Bα SISF Bα αd αd SESF Bα (a) (b) (c) Figure. Dissociation of dislocations in the fcc structure and the L1 structure according to the Thompson notation: (a) dissociation of a perfect (a/)½101š matrix dislocation with an intrinsic fault I; (b) dissociation of a perfect a½101š dislocation of the L1 structure with a SISF; (c) dissociation of a perfect a½101š dislocation of the L1 structure with a SESF.

6 96 B. De camps et al. The sense of the Burgers vectors with respect to the dislocation lines is determined in bright-field for jg bj ¼ by using the apparent sense of displacement of the image for linear dislocations (Hirsch et al. 1977). For dislocation loops, the identification of the sense of the Burgers vector is made from inside outside contrast using g images which have non-integral values of jg bj (Loretto and Smallman 1975). }. Experimental results Figure 4 (a) is typical of the microstructure encountered after a tensile creep test at 700 C under 650 MPa up to 0.% of plastic deformation. This figure clearly demonstrates that the preponderant shearing process involves superlattice stacking faults in the four gliding planes. The observed intersections of such shearing bands are known to produce strong strain hardening. Looking at the detail, the microstructure reveals different shearing configurations as shown in figures 4 (b) and (c). For secondary precipitates, figure 4 (b) is typical of the Condat De camps (1987) mechanism where precipitates are sheared by superlattice stacking faults while the matrix remains unsheared. In figure 4 (c), on the contrary, both the precipitates and the matrix appear to be sheared, which cannot be explained by the crossing of the interface by a single (a/) h110i matrix dislocation. Considering now the tertiary precipitates, the same shearing configuration appears clearly in the first case (figure 5 (a)) while it is not so obvious when precipitates are embedded in the sheared matrix (figure 5 (b)). In order to determine precisely the succession of steps leading to this new shearing configuration, the crystallography of the two configurations A and B in figure 6 (a) is analysed in detail..1. Step 1: decorrelation of the movement of two Shockley partials belonging to one matrix dislocation (configuration A) According to figures 6 (a) (f), configuration A is consistent with a matrix dislocation BD lying in the a plane. One part of the dislocation is widely dissociated in the a plane, forming an intrinsic fault. A schematic representation of the configuration is given in figure 6 (g). Table gives the jg bj values for all perfect dislocations and Shockley partials encountered in the (111) plane using the reflections presented in the paper. For example the BD dislocation is visible for all reflections of figure 6, which is consistent with table. The fault is visible for g ¼ 11 1 (figure 6 (a)) and g ¼00 (figures 6 (e) and (f)) and is out of contrast for g ¼ 0 (figure 6 (b)), g ¼ 0 (figure 6 (c)) and g ¼ 0 (figure 6 (d)), which means that it is located in the a plane. The sense of the Burgers vector BD according to the scheme in figure 6 (g) is obtained using the displacement of image contrast for g ¼ 0 (figure 6 (c)). The Burgers vectors analysis of the two Shockley partials that border it allows the determination of the nature of the fault. Ba and ad are out of contrast for g ¼ 0 (figure 6 (d )) and g ¼ 0 (figure 6 (b)) respectively according to table and to the scheme in figure 6 (g). The fault is then shown to be intrinsic. Finally, in configuration A, the wide extent of the fault ribbon is produced by the decorrelation of the movement of the two Shockley partials of a matrix dislocation.

7 Shearing mechanism of g 0 precipitates 97 Figure 4. (a) Dislocation structure in a grain of a foil creep-strained at 700 C under 650 MPa. (b) Shearing configuration of g 0 precipitates involving stacking faults only in the precipitates according to the De camps et al. (1994) model. (c) New shearing configuration of g 0 precipitates involving stacking faults both in precipitates and in the matrix... Step : Shearing configuration of g 0 precipitates (configuration B) Configuration B (figure 6 (a)) is an example of the new shearing process involving stacking faults both in the g 0 precipitates and in the matrix. In this configuration, both secondary and tertiary precipitates are part of the process. The configuration is analysed using reflections of figures 6 (a) (f) and others not presented here. A schematic

8 98 B. De camps et al. Figure 5. Shearing configuration of tertiary g 0 precipitates: (a) according to the De camps et al. (1994) model; (b) involving the new shearing mechanism with faults in the matrix and in the precipitates. representation of the configuration is given in figure 6 (h). All stacking faults, which lie in the a plane, are visible for g ¼ 11 1 (figure 6 (a)) and g ¼00 (figures 6 (e) and ( f )) and out of contrast for g ¼ 0 (figure 6 (b)), g ¼ 0 (figure 6 (c)) and g ¼ 0 (figure 6 (d )). All dislocations that form the configuration have ad or ad Burgers vectors as they are invisible for g ¼ 0 (figure 6 (b)) according to table. The weak-beam micrograph in figure 6 (d) shows clearly that the dislocation located at the top of the configuration is dissociated in two ad dislocations. The sense of the Burgers vector of this ad dislocation with respect to its line is determined in bright-field using the displacement of image contrast (figure 6 (c), g ¼ 0 ). The shearing configuration is then shown, according to figure, to involve only

9 Shearing mechanism of g 0 precipitates 99 Figure 6. (a) (c), (e) Bright-field and (d), (f) weak-beam micrographs of the alloy NR after a tension creep test at 700 C under 650 MPa: (a) g ¼ 111 (the studied configurations are denoted A and B); (b) g ¼ 0; (c) g ¼ 0; (d) g ¼ 0; (e) g ¼ 00; (f) g ¼ 00. (g) Schematic representation of the configuration A. (h) Schematic representation of the shearing configuration B.

10 100 B. De camps et al. Figure 6. Continued.

11 Shearing mechanism of g 0 precipitates 101 D BD I Bα (g) αd αd αd αd αd SESF αd I Dα SESF X Y Y αd Surface (h) Figure 6. Continued. Table. Values of g b for dislocations of the a plane. g b value for the following g b CB ¼ DC ¼ BD ¼ ab ¼ ab ¼ Ca ¼ Ca ¼ ad ¼ ad ¼ SESFs. By applying the node law, the schematic representation of the configuration B displayed in figure 6 (h) is obtained. Regarding dislocations denoted Y in figure 6 (h), their contrast is consistent with the ad Burgers vector according to all reflections of figure 6. The sense of the Burgers vector of dislocation Y is determined by looking at the inner contrast and outside contrast respectively of the loop for g ¼ 00 (figure 6 (e)) and g ¼ 00

12 10 B. De camps et al. Figure 7. Model of shear of a g 0 precipitate by one decorrelated Shockley partial of a single dissociated matrix dislocation forming a SESF. (a) Decorrelation of the two Shockley partials. (b) Crossing of the interface by the single Shockley partial forming a CF and nucleation of an ad Shockley dislocation in an adjacent plane above or under the CF according to the Décamps et al. (1994) model forming the SESF. (c) Resulting configuration during the shearing process. (figures 6 ( f )). Such inner outside contrast is clearly visible in the inserts of figure 6 (e) (inner contrast) and figure 6 (f) (outside contrast). Thus, the Burgers vector of dislocation Y is Da according to figure 6 (h). The nature of the X dislocation of figure 6 (h) (Da or ad according to the contrast of all reflections of figure 6) will be discussed later. } 4. Interpretation and discussion The observations show that all shearing configurations presented above involve superstacking faults homogeneously distributed within the grain. Many of these present faults in the matrix. If the configurations involving faults located only in the precipitates may be simply explained by the mechanism proposed by De camps et al. (1994), it is not the case for those displaying faults in the matrix. On the basis of the existence of the decorrelated movement of the Shockley partials of matrix dislocation as shown above (}.1), a new mechanism derived from the previous mechanism may be proposed (figure 7). Because of the decorrelated movement of Shockley partials (figure 7 (a)), only the Da partial crosses the interface, forming a CF (figure 7 (b)). Then nucleation of a Shockley partial occurs above or under the CF, forming a SESF (figure 7 (b)) according to the following reaction: Da þ Da ¼ Da: The resultant shearing configuration is shown in figure 7 (c). Finally, the coherency existing between these two different shearing models must be noted; in both cases the mechanism is initiated by the leading partial that enters the precipitate. Thus, after nucleation of a Shockley partial above or under the CF, the precipitate is sheared by

13 Shearing mechanism of g 0 precipitates 10 a SESF. It is important to note that, in the second case, owing to the decorrelation process, the formation of SISFs that requires the inversion of the two incoming partials according to De camps et al. (1991b, 1994) is not possible. Such a model easily explains the shearing configuration described in figure 6, taking into account the fact that both original perfect DC or DB dislocations would produce the same configuration as only the Da partial creates the process. Nevertheless, owing to the presence of a perfect BD dislocation (see configuration A) moving in the opposite sense to that required to produce the shearing configuration, it can be concluded that the original matrix dislocation was DC whose movement is compatible with the applied stress (according to the Schmid factor). Close to dislocation X, the removal of the matrix fault in the channel of figure 6 (h) may be explained by the interaction of the trailing Ca partial with a perfect BD dislocation, forming the X dislocation according to the following reaction: Da þ Ca þ ad þ Ba ¼ ad ¼ X: Thus only ad or Da dislocations constitute the configuration as observed. So, this new shearing process is strictly dependent upon the decorrelated movement of the two Shockley partials of a dissociated matrix dislocation. Such a decorrelation process is the result of the interplay between three effects: (i) Orowan stress; (ii) matrix stacking-fault energy (SFE); (iii) channel width. A decrease in the matrix SFE and/or the channel width and/or an increase in the Orowan stress would favour this mechanism. By evaluating the effective stress, as well as the respective flexibility on each one of these partial dislocations, the possible occurrence of such movements for some particular dislocation characters b and channel widths d is accounted for. The main idea is that a dislocation is able to enter a particular matrix channel, if the effective resulting stress on the considered dislocation segment is in excess compared with the threshold stress or characterizing the local geometry of the obstacle. This so-called Orowan stress Or ¼ =bd depends on the line tension which is inversely proportional to the flexibility. These effective forces can be different considering a perfect dislocation or each one of its partials, inducing then the decorrelation of their movements (Coujou et al. 00). It is also noteworthy to compare the effects of these two shearing mechanisms (with and without faults in the matrix) on the mechanical behaviour of the superalloy. It is obvious that the mechanism requiring a decorrelation process would demand a higher energy to operate than the other. So, any heat treatment procedure leading to a microstructure requiring a decorrelation process to occur before shearing would improve the creep strength of the superalloy. Furthermore, the more energy the decorrelation process requires, the better the creep resistance will be. A high matrix SFE would increase the energy required for the decorrelation process to occur but it must not be too high with respect to the channel width. The decorrelation process can also be favoured by the Suzuki effect. Indeed, the stacking fault situated between the two partials constitutes a hc block inserted in the fcc matrix. As the chemical potential of the solutes is different in the two phases, a concentration variation is generated to compensate this difference. This phenomenon proposed by Suzuki (195) has been observed in several alloys (Coujou 1979), gen-

14 104 B. De camps et al. Figure 8. Bright-field micrograph displaying both shearing processes (with faults only present in precipitates and involving faults both in the matrix and the precipitates) in the same shearing configuration. erally accompanied by a SFE decrease, which favours decorrelation at high temperatures. The contributions of these mechanisms have been compared quantitatively for the grain presented in figure 4 (a), which is slightly deformed. In this case, an evaluation of the deformation in the different active gliding systems p for the two mechanisms m (m ¼ 1 represents the Condat De camps process and m ¼ the new shearing process) has been made according to the Orowan relation with " m ¼ X4 p¼1 p, m b m L p, m p, m ¼ l p, m V where l is the dislocation length, L is their displacement, b is the Burgers vector corresponding to the higher Schmid factor and projected in the tensile direction and

15 Shearing mechanism of g 0 precipitates 105 V is the volume considered. The calculation reveals that " 1 =" ¼ 1:1. Thus, although the number of faults is more important for the new configuration, its contribution to the deformation is close to that generated by the Condat De camps mechanism because it is proportional to the Burgers vector modulus, weaker in the new configuration. As far as the ageing of such a bimodal g g 0 microstructure is concerned, identification of these two types of shearing configuration widely distributed in the crept specimens can be of great help in determining the evolution of the g-channel dimensions. The observations of shearing configurations involving a faulted matrix may be considered as the signature of the presence of narrow channels impeding the easy propagation of matrix dislocations and causing their decorrelation. As shown in figure 8, the shearing process can start, on the same glide plane, with a classical shearing configuration as proposed by De camps et al. (1994) followed suddenly by the second, which expands on a neighbouring area. This can be analysed as follows: owing to the deformation, sources are preferentially operative in the soft phase and dissociated matrix dislocations are created and begin to propagate in g channels. When the width of the channel is sufficient, the first mechanism (shearing configuration involving a total (a/) h110i dissociated dislocation as proposed by De camps et al. (1994) takes place preferentially and the shearing of g 0 precipitates is observed, the matrix remaining unsheared. If, on the same plane, channels become suddenly narrower, a decorrelation of the two partials is then necessary to allow the propagation to continue. In such configurations, as it is clear that the observed configuration is located on the same glide plane, a direct measurement of the channel width between tertiary g 0 particles can be made. For example, statistically in figure 5 (a), tertiary g 0 precipitates are separated by channels of width close to 68 nm while they are close to 45 nm in figure 5 (b). The physical reason for such a decorrelation is linked to the dislocation flexibility which depends on its character and is lower for a partial dislocation than for a perfect dislocation (Benyoucef et al. 1995). As a consequence, in narrow channels, the first Shockley partial alone is able to propagate, creating a stacking fault in the matrix and giving rise to the second shearing process when encountering the g g 0 interfaces (Coujou et al. 00). Shearing configuration analyses were also performed in NR crept specimens containing very fine tertiary g 0 precipitates (ageing for 4 h at 700 C) and others without tertiary g 0 precipitates (overageing for 500 h at 800 C). Raujol et al. (00) have shown that the evolution of the microstructure during creep at intermediate temperatures is consistent with a decrease in the volume fraction of tertiary g 0 precipitates in correlation with an enlargement of the g channels. While TEM post-mortem observations performed on samples aged in the standard conditions (110 C for h þ 700 C for 4 h þ 800 C for 4 h) reveal a majority of shearing configurations with faulted precipitates (mean channel width, around 5 nm), shearing configurations with a faulted matrix and precipitates are prevalent in the material containing the finest tertiary g 0 precipitates (mean channel width, around nm). In the overaged specimen, the secondary precipitates are overcome by Orowan processes and matrix dislocations are free to move without any observed decorrelation of the partials. From the above observations, a quantitative measurement of the amount of shearing configurations populations present in each specimen can therefore be performed. The degree of ageing of the g g 0 microstructure could consequently be deduced from this analysis. Such a method presents the following advantages:

16 106 B. De camps et al. (i) The measurement of channel width, which is difficult owing to the threedimensional repartition of tertiary precipitates, is not necessary. (ii) An eventual local morphology evolution of tertiary precipitates could be easily followed through the respective percentage of each kind of faults (see for example figure 8, where on the same glide plane the shearing process begins by the first type of shearing followed by the second when the channels become narrower). } 5. Conclusions The combination of TEM bright-field and weak-beam techniques has been used to analyse a new shearing configuration of g 0 precipitates in crept specimens of a polycrystalline Ni-based superalloy, involving stacking faults both in the precipitates and in the matrix. Such a shearing configuration is shown to be initiated by the decorrelated movement of the two Shockley partials of an individual dissociated matrix dislocation. On the basis of these results, we propose an original model for the shearing of g 0 precipitates by a single (a/6) h11i decorrelated Shockley partial; the model is derived from the previous mechanism proposed by De camps et al. (1994) involving a single dissociated matrix dislocation. The decorrelation process, prerequisite for the new shearing model to operate, is shown to be the interplay between three effects: Orowan stress, matrix SFE and channel width. The quantitative analysis of the shearing configurations in the crept material could be a good tool to quantify the degree of ageing of its g g 0 microstructure. References Benyoucef, M., Cle ment, N., and Coujou, A., 1995, Phil. Mag., 7, 104. Bilby, B. H., Bullough, R., and Smith, E., 1955, Proc. R. Soc. A, 1, 6. Caron, P., Khan, T., and Veyssiere, P., 1988a, Phil. Mag. A, 57, 859. Caron, P., Ohta, Y., Nakagawa, Y. G., and Khan, T., 1988b, Superalloys 1988, edited by D. N. Duhl, G. Maurer, S. Antolovich, S. Lund and S. Reichman (Warrendale, Pennsylvania: Metallurgical Society of AIME), pp Cockayne, D. J. H., Ray, I. L. F., and Whelan, M. J., 1969, Phil. Mag., 0, 165. Condat, M., and De camps, B., 1987, Scripta metall., 1, 607. Coujou, A., Raujol, S., Pettinari-Sturmel, F., Cle ment, N., Locq, D. and Caron, P., 00 (to be published). Coujou, A., 1979, Phil. Mag., 40, 17. De camps, B., and Condat, M., 1986, J. Spectrosc. Electron., 11, 141. De camps, B., Condat, M., and Morton, A. J., 1991a, Microsc. Microanal. Microstruct.,, 60. De camps, B., Morton, A. J., and Condat, M., 1991b, Phil. Mag. A, 64, 641. De camps, B., Penisson, J. M., Condat, M., Guetaz, L., and Morton, A. J., 1994, Scripta metall., 0, 145. Hirsch, P. B., Howie, A., Nicholson, R. B., Pashley, D. W., and Whelan, M. J., 1977, Electron Microscopy of Thin Crystals (New York: Krieger). Howie, A., and Whelan, M. J., 196, Proc. R. Soc. A, 67, 06. Huis in t Veld, A. J., Boom, G., Bronsveld, P. M., and de Hosson,J.Th. M., 1985, Scripta metall., 19, 11. Kear, B. H., Giamei, A. F., Leverant, G. R., and Oblak, J. M., 1969a, Scripta metall.,, 1; 1969b, ibid.,, 455. Kear, B. H., Giamei, A. F., Silcock, J. M., and Ham, R. K., 1968, Scripta metall.,, 87. Kear, B. H., Leverant, G. R., and Oblak, J. M., 1969c, Trans. Am. Soc. Metals, 6, 69.

17 Shearing mechanism of g 0 precipitates 107 Kear, B. H., and Oblak, J. M., 1974, J. Phys., Paris, 5, C7 5. Leverant, G. R., and Kear, B. H., 1970, Metall. Trans., 1, 477. Locq, D., Marty, M., Walder, A., and Caron, P., 000, Intermetallics and Superalloys, Euromat 99, Vol. 10, edited by D. G. Morris, S. Naka and P. Caron (Weinhem: Wiley- VCH Verlag Gmbh), pp Loretto, M. H., and Smallman, R. E., 1975, Defect Analysis in Electron Microscopy (London: Chapman and Hall) (New York: Wiley). Rae, C. M. F., Kakehi, K., and Reed, R. C., 00, Materials for Advanced Power Engineering 00, Part I, edited by J. Lecomte-Beckers et al. (Forschungszentrum Ju lich), pp Raisson, G. and Davidson, J. H., 1990, High Temperature Materials for Power Engineering 1990, Part II, edited by E. Bachelet et al. (Dordrecht: Kluwer), pp Raujol, S., Pettinari-Sturmel, F., Locq, D., Caron, P., Coujou, A., and Cle ment, N., 00 (to be published). Suzuki, H., 195, Sci. Rep. Res. Inst. Tohoku Univ. A, 4, 45. Thompson, N., 195, Proc. Phys. Soc. B, 66, 481.