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II. Deformed microstructures during creep of TiAl alloys: role of mechanical twinning
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5 (1997)
339-354
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II. Deformed microstructures during creep of TiAl alloys: role of mechanical twinning M. A. Morris & M. Leboeuf Institute of Structural Metallurgy.
University of Neuchdtel, Av. Bellevaux, 51, 2000 Neuchritel. Switzerland
(Received 25 November 1996; accepted 6 January 1997)
The deformation mechanisms responsible for the creep behaviour of two TiAl alloys with duplex and lamellar structures have been analysed after carrying out microstructural observations as a function of creep strain and applied stress at 700°C. The duplex alloy exhibits extensive mechanical twinning and the new twin interfaces subdivide the y grains throughout the primary stage of creep. At the onset of the minimum creep rate, the twin interfaces in the duplex alloy behave in the same way as the y/y or the oa/y interfaces in the lamellar alloy. Thus, dislocation accumulation and emission from the interfaces occur under the action of the high local stresses. The high stress exponent, n = 19, measured in both alloys at the minimum strain rate has been interpreted in terms of a threshold stress representing the high local stresses required to liberate dislocations from the interfaces and initiate their propagation within the y matrix. 0 1997 Elsevier Science Limited Key words: A. titanium aluminides based on TiAI, B. creep, twinning, D. defects:
planar faults, F. electron microscopy,
Some studies of creep deformation of duplex microstructures1’~i2 have shown that a reduction in lamellar interface spacing occurs due to mechanical twinning. It has been suggestedI that if the twin interfaces act as do lamellar interfaces, they would provide barriers for dislocation motion and lead to similar hardening. However, mechanical twinning has been shown to decrease at low stressesi or high temperatures. i5 For two-phase alloys, an increased twinning activity has been observed at temperatures below 760°C and stresses over 200 MPa. The aim of the present study was to analyse and interpret the deformation structures produced by creep at a temperature of 700°C and high stresses (above 250MPa) in alloys with initial duplex and lamellar microstructures. The alloys chosen have previously been studied and the presence of twins confirmed after tensile (or compression) deformation at 700”C.‘6~17 The evolution of the microstructure during the primary stage of creep and, in particular, the increase of twinning activity with increasing strain have been studied. Also, the effect of pre-deformation on twinning activity has been evaluated in terms of its contribution to improvement of the creep resistance observed in the duplex
INTRODUCTION The study of deformed microstructures produced during creep of y-based TiAl alloys appears relevant to the understanding of the fundamental aspects that contribute to the primary stage of creep and lead to the onset of a minimum strain rate. In particular, the higher creep resistance of alloys with lamellar (CQ+ y) microstructures has been well establishedi but not much work has been carried out to understand the microstructural evolution during the primary stage of creep of alloys with initially duplex and/or lamellar microstructures. Also, it has been suggested that the creep resistance of TiAl alloys can be improved by addition of alloying elemNents5or introduction of a lamellar structure.6 The primary creep might also be reduced by pre-straining at high stress prior to creep-testing at a lower stress7 However, systematic studies of the pre-deformed microstructure before and after changing the stress have not been made. The better creep resistance of the fully lamellar structure is generally attributed to the CQlaths and r/r interfaces acting as barriers to slip8,9 and to the development of constraints at the o/a/y interfaces.” 339
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alloy. Similarly, the effect of stress on twinning activity and its influence on the creep parameters obtained from both alloys have been analysed.
EXPERIMENTAL
obtained by tilting the specimens to different known orientations (zone axes) from which different diffraction vectors were chosen to obtain invisibility under some contrast conditions. Weakbeam images were taken using the g: 3g condition.
DETAILS
The two alloys compared in the present study had compositions Ti-48Al-2Mn-2Nb and Ti-48Al (atoh), respectively. Details about alloy preparation and initial microstructures have been described elsewhere.16T17The creep tests were carried out in tension on cylindrical specimens machined to a diameter of 3mm and with 28mm gauge length. Their surface was subsequently mechanically and electro-polished using a solution of 5% perchloric acid in methanol. The creep machine used was equipped with an Andrade-type cam with a profile that allowed the stress to be maintained constant throughout the test. Tests were performed at 700°C for stresses ranging between 250 and 360MPa for the duplex alloy and between 360 and 430 MPa for the lamellar alloy. The temperature of the specimens was maintained constant with a precision of kO.2”C. The elongation of the specimens was monitored with a precision of about f 2p.m. Although some tests were performed to rupture (see previous paper), those for microstructural observations were arrested at 1 or 2% deformation, the latter corresponding to the strain just before the minimum creep rate was measured. At this stage, the specimens were rapidly cooled under load to avoid any dislocation rearrangement prior to analysis. Microstructural analysis of the deformed specimens was performed by transmission electron microscopy (TEM) using a Philips CM200 microscope equipped with EDS facilities, at 200KV. Also, to obtain information about any large scale substructure formation during the primary stage of creep, scanning electron microscopy (SEM) observations were made using crystallographic contrast from backscattered electrons. For these TEM and SEM observations, discs were cut perpendicular to the tensile axis from the deformed samples and electro-polished. The electro-polished samples were prepared by standard jet polishing techniques using a solution of 5% perchloric acid, 30% butan-l-01 and 65% methanol at -20°C and 50mA, but only the TEM samples were polished to produce a hole. At least two to three different foils were examined from each crept specimen and several y grains or lamellae were studied from each case. Dislocation analysis was carried out from projected images
RESULTS
AND COMPARISONS
After deformation to different strains during the primary stage of creep, the deformed microstructures were studied both by SEM observations made by crystallographic contrast (using backscattered electrons) and by TEM images obtained from the same deformed specimens. The primary stage of creep ends at the strain for which the minimum strain rate was reached and in our alloys this occurred generally at about E =0~02-0.025. Since the strains at which the microstructures could be compared were rather limited, it was important to take into account the microstructures of the undeformed samples as described and illustrated in the first part of this study (see previous paper). Although several specimens were examined after testing at different stresses, only the microstructures of the samples corresponding to the lowest and highest stresses are shown. The intermediate conditions exhibited the same features and have been omitted. Microstructures in the duplex alloy The main feature observed in the deformed microstructures of the duplex alloy was the extensive twinning activity which was already present after only 1% strain in the specimens crept at the lower stress of 280 MPa. Figure l(a) shows an example from this specimen, taken at low magnification in the TEM, of one of the large y grains that was completely traversed by several twins. Detailed diffraction analysis of these twins and others observed under different conditions of strain or applied stress confirmed that they all exhibited the true-twin relationship between the matrix and the twinned layer. In this sense, the twinned interfaces were exactly the same as the r/y interfaces of the true-twin type defined by Yamaguchi and Umakoshii* as producing a rotation of 180” around a < 1 1 1 > pole and analysed in some of our previous work.i6,i9 This twinned microstructure was confirmed to be produced by emission and glide of l/6 < 1 12 > dislocations on {1 1 l} planes, confirming our previous studies’6~‘7~19and those by other authors.20,2’ The extensive twinning activity
II. Deformed microstructures during creep of TiAl alloys was also responsible for the activation of large numbers of single l/2 < 1 10 > dislocations at intersections with some of the v/v interfaces present. Details from some of these intersections (labelled A and B) are s(hown in Figs l(b) and (c), respectively. In these figures we also note that some of the l/2 < 1 10 > dislocations remain on the {1 1 l} planes of the twin interfaces, as also evidented during in situ studies of twin propagation in
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yTi-Al by Couret et al. 22 Besides, some of the twins do not propagate across the entire grain (as seen in two of them within the area labelled B). Thus, it appears that the local stress concentration is limited and cannot provide the twinning driving force to cross the entire grain as suggested by Jin and Bieler.21 Since the number of large grains that could be examined in the TEM was rather limited, we
Fig. 1. (a) Example of extensive twinning traversing an equiaxed y grain observed by TEM from the duplex alloy after 1% creep strain under 280 MPa. (b), (c). Details from areas A and B in Fig. l(a) confirming emission of l/2 < 1 1 0 > dislocations.
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examined complete deformed sections by SEM from discs cut perpendicularly to the tensile axis from the gauge length of the samples. In order to examine these discs in the SEM, they were electropolished; the crystallographic contrast allowing the observation of twins is due to the fact that the thin twinned regions are rotated by 180” with respect to the matrix. Therefore, a thin slab of twinned region appears as a line due to its different crystallographic contrast. Figure 2 shows examples comparing grains taken from sections of the undeformed specimens (head of tensile sample) and
after creep to 1% deformation at low stress (280MPa). We note that the major difference between the grains of undeformed (Fig. 2(a)) and deformed (Fig. 2(b)) material is that while in the former there is a total absence of mechanical twins, in the latter the appearance of mechanical twins is similar to that observed by TEM under the same creep conditions (compare Figs l(a) and (b)). These SEM observations confirmed that although some of the grains did not exhibit twinning after 1% strain, the majority of them did and in all cases only one twin system was activated.
Fig. 2. Comparison between y grains observed by crystallographic contrast in the SEM from the duplex alloy (a) undeformed state (b) after 1% creep strain at 280 MPa. The latter confirms the extensive twinning observed by TEM in Fig. l(a).
II. Deformed microstructures during creep of TiAl alloys
After creep to 2% strain (under the same stress of 280MPa), the SEM observations revealed a more pronounced twinning activity, as illustrated in Fig. 3. Here we see that the twin transfer takes place across adjacent grains, leaving large steps at the grain boundaries (Fig. 3(a)) and also that two twin systems were active in some cases, as in Fig. 3(b) (note that the bright elongated globules are some of the CX~ phase). These features were also encountered in TEM observations, as illustrated in Fig. 4(a). Here we see a large area taken at low magnification traversed by two twin systems that intersect. This area has been analysed in detail and Figs 4(b) and (c) show detailed examples taken from this area (labelled ,4 and B) at higher magnification and under different diffraction conditions.
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Apart from the twins, single dislocations were active and they are invisible in Fig. 4(a) taken with g= 002. Their visibility in Figs 4(b) and (c) taken with g = 200 and 1i 1, respectively, confirms that the Burgers vector is 1/2[1iO]. Also, in Fig. 4(c), the twin system labelled 1 is invisible with g = 1i 1 and its analysis has confirmed it to be 1/6[i 121 (1 1 1). The cross-twin system has been analysed as 1/6[ 1 121 (i i 1). From the micrographs in which the twin interfaces were imaged parallel to the electron beam, the thickness of the twinned layers measured was about 30-40 nm. The 1/2[ 1iO] dislocations seen in Fig. 4(c) have a rather angular morphology and their two distinct directions have been analysed as [ 1121 and [loll, respectively. Although these directions lie on the same (7 11)
Fig. 3. SEM observations made from the duplex alloy after 2% creep strain at 280MPa. (a) Twinning transfer across adjacent y grains produces steps at grain boundary. (b) Example showing that several y grains deform by mechanical twinning, including two systems in some cases. Note that the bright globules are from the a2 phase.
M. A. Morris, M. Leboeuf
Fig. 4. Microstructures observed by TEM from the same specimen of Fig. 3, i.e. after 2% creep strain at 280 MPa. (a) T‘ypical y grain deformed by two intersecting twin systems (b), (c). Details from areas A and B in Fig. 4(a). (Note that all the active dislocations have Burgers vector 1/2[1iO] and are invisible in Fig. 4(a) taken with g = 002). (d) Details from the oz globules (seen in Fig. 4(a) as elongated dark particles). The weak-beam image shows that accumulation of matrix dislocations occurs at the ~Z/Y interface while the bright-field micrograph shows evidence of shear across the az phase.
II. Deformed microstructures during creep of TiAl alloys
(4
Fig. 4. Contd
plane, their Burgers vector does not belong on it, indicating that these dislocations have climbed (since one of them has pure edge character). Other details present in these images are the globular, elongated a2 particles whose interfaces seem to accumulate dislocations but which can also be sheared by the stress concentrations produced as the twins from the matrix intersect them. An example of these sheared particles is shown in Fig. 4(d) (the weak-beam image shows details of the (11~/yinterface). The effect produced on the deformed microstructure by the increase in applied stress is only that of accentuating the details already described for the sample crept at low stress. Figure 5 shows the typical microstructure observed by SEM from the sample crept to the same 2% strain previously shown in Fig. 3 but tested under a higher stress of 360MPa (the higher stress used for this alloy). We see that the crystallographic contrast obtained from these grains confirms the more pronounced twinning activity at the higher stress level, the two active twin systems be:ing emitted from grain boundaries in most cases. Under these high stresses
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it was rare to find grains that were not deformed by mechanical twinning. Figure 6 shows the comparison of a TEM microstructure seen from the same specimen crept 2% at 360 MPa (Fig. 6(a)), together with an SEM image (Fig. 6(b)) taken from the detailed area labelled A in Fig. 5. We note that for the same amount of strain, the twin density is much higher in this specimen crept at 360MPa than that observed at 280MPa. As a result, the subdivision that occurs in the y grains is much finer in the specimens crept at the higher stress, i.e. the spacing between twinned layers was about 5-6hm after 2% strain and 280 MPa (see Figs 3 and 4) and I-2pm in the case of the test performed at 360MPa (see Figs 5 and 6). The thickness of the twinned layers measured from this specimen tested at high stress was about 60-80nm, confirming that a larger number of l/6 < 1 12 > dislocations have contributed to their formation compared to those formed under the low stress regime. From Fig. 6(a) we also note that, apart from mechanical twins, the y matrix is also deformed by slip of dislocations, which have been analysed. Figure 7(a) shows details from these dislocation segments whose Burgers vector has been confirmed to be 1/2[1 TO] (as they are invisible with g = 002 but visible with g = 1 i 1) and appear elongated along their [l TO] direction, i.e. they have screw character. Therefore, at higher stresses the dislocations propagate by glide as observed in our previous study,t7 i.e. by movement of the edge segment leaving the screw segment behind. Finally, Fig. 7(b) shows details from these twin interfaces observed by weak-beam contrast in which the single dislocations appear bowed, being emitted in exactly the same way as they have been observed from y/y interfaces in some of our previous studies and those of other authors.23y24 Thus, we confirm that in the duplex alloy, after 2% strain, i.e. just before the minimum strain rate is reached, the twinning activity has produced a subdivision of the equiaxed grains into regions about 20 to 50 times smaller (depending on the applied stress). At the same time, the twin interfaces act as sites for dislocation emission and accumulation and their propagation within the y matrix produces deformation in a similar manner as observed in lamellar structures. I3 Microstructures in the lamellar alloy
The general results obtained from the observations of the deformed lamellar alloy do not differ greatly from those already seen in previous studies.25 SEM
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observations carried out from discs cut from the gauge length of the specimen indicated that twinning was prominent only after strains of more than 3%. Figure 8 shows a comparison of these observations made from the undeformed specimen and those crept to strains of 24% at a stress of 400MPa Although a high twin density was observed after 4% strain, this was not the case at 2% strain (compare Figs 8(b) and (c)), even in the case of the test carried out at the higher stress of 430 MPa. TEM observations made from this alloy in the undeformed state16 had confirmed that it had a higher density of dislocations than the duplex alloy17 and this was consistent with the observations carried out from the undeformed head of the tensile creep specimens (see previous paper). The latter were made using several thin foils from
which a large number of y grains or lamellae were studied. After small amounts of strain during the primary stage of creep (l-2%) little variations of the deformed microstructure were seen. This is probably because the dislocation density within the y lamellae was not very high and there was no evidence to distinguish between those dislocations present before and after the creep deformation. Similar comments have also been made by other authors who have studied the crept lamellar structures.27 However, as illustrated in Fig. 9, weakbeam images taken from the c~/y interfaces showed evidence that extrinsic dislocations from the y matrix were accumulating at the existing networks (Figs 9(a) and (b)) and that dislocation segments bowed out to be emitted into the y matrix (Fig. 9(c)), and even that at certain sites destruction of dislocation nodes led to imperfections in
activity observed by SEM in the duplex alloy after 2% creep strain under a sitress of 360 MPa. Note the much finer spacing between twins in all the grains.
Fig. 5. Example of increased twinning
II. Deformed microstructures during creep of TiAl alloys
those interface networks, as seen in the area marked by the arrows in Fig. 9(d). After 2% strain (just reaching the minimum creep rate), more evidence of deformation within the y matrix was observed. In some lamellae twinning was already present even though at this strain the twin contrast had not been seen by SEM. However, the twins were extremely thin and probably this was not sufficient to produce enough crystallographic contrast in the SEM, as was the case in the duplex alloy even after 1% strain. Figure 10 shows examples of lamellae deformed by twinning (Fig. 10(a)) and dislocation slip (Fig. IO(b)) observed after 2% strain at 360 MPa. Also here we show a detail (Fig. 10(c)) from one of the twin interfaces from which bowed dislocations are emitted. These, as well as those seen in Fig. 10(a),
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were confirmed to be l/2 < 1 10 > segments as they were invisible with g = 002. Figures 11 and 12 show different examples of lamellae deformed 2% at stresses of 370 and 400MPa, respectively. The common feature in these images is the observation of the destruction of the a2 laths that, after breaking down, appear as a2 particles aligned in the centre of the v lamellae. These particles are marked by arrows in Fig. 1l(a), taken with g= 002. Under this diffraction condition, only superdislocations are visible and they remain attached to the cz2particles as shown in the detail of Fig. 1l(c). Also, in Fig. 1l(a) we show the same deformed lamella taken with g = 020 and we note the high dislocation density of Burgers vector 1/2<11O>.Th e d es t ruction of the cz2laths occurs by shear, as illustrated in the detail of Fig. 1l(b).
Fig. 6. Comparison of the mic:rostructures observed by (a) TEM and (b) SEM from a specimen of the duplex alloy crept to 2% straiin under 360 MPa. The SEM image was taken from the area labelled A in Fig. 5. Note the extensive twinning activity and 1high dislocation density relative to the equivalent microstructures obtained at low stress and seen in Figs 3 and 4.
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Here we see that the extrinsic dislocations at the interface delineate the profile of the sheared a2 particle. A complete analysis of the dislocations present within this lamella has been made. The superdislocations seen attached to the (~2particles in Figs 1l(a) and (c) have Burgers vector [TO l] and two line directions have been analysed, namely [Oil] and [l lo], indicating that they lie in two glide planes, i.e. (1 1 1) and (1 i’ 1). From the total analysis of the dislocations seen within this )I lamella, only 10% correspond to the mentioned superdislocations; the rest are single l/2 < 1 10 > segments. Of the latter, 62% have Burgers vector 1/2[i 101 and two distinct line directions were analysed, namely [0 0 11, i.e. the segments had pure edge character, and [0 lo], i.e. at 45” from the direction of the Burgers vector. Those directions define the respective (1 10) and (0 0 1) glide planes containing the Burgers vector and the line direction. However, dislocations with line direction [0 0 l] might have climbed due to their pure edge character. Similarly, the remaining 28% dislocations have Burgers vector 1/2[1 lo] and major line directions [0 0 l] and [0 lo], defining the glide planes (110) and (0 0 l), respectively. From these analyses we conclude that climbing and gliding mechanisms of single dislocations are
operative within the y lamellae at the strain at which the minimum creep rate is reached. Also, enough accumulation of dislocations has occurred at the a2/y interfaces that high stress concentrations are produced and shear of the o2 laths occurs locally. From the example shown in Fig. 12(a), we note the general dislocation arrangements in contact with the 01~interface, and the small bowed dislocation segments seen by the side of the broken
Fig. 7. Details from dislocation activity seen in Fig. 6(a). (a) Screw 1/2[1 1 O] dislocatjon segments analysed within the y matrix. (b) Bowed 1/2[1 lo] segments being emitted from twin interfaces.
Fig. 8. Microstructures observed by SEM from the lamellar alloy. (a) Undeformed state; (b), (c) after 2 and 4% creep strain, respectively, under 400 MPa. Note that twinning across the y grains was pronounced only after more than 2% strain.
II. Deformed microstructures during creep of TiAl alloys
a2 lath, in the weak beam detail of Fig. 12(b), confirms the presence of high local stress concentrations. Therefore, we deduce that saturation of the number of dislocations that accumulate at the q/v interfaces leads to increased stress concentrations that lead to shear of the a2 laths as the minimum creep rate is reached. As a result, the continuous increase in strain rate observed thereafter is due to the slow destruction of the ~3 laths
Fig. 9. Examples of TEM obs,ervations made by weak-beam contrast from the a& interfaces of the lamellar alloy. (a), (b) Examples of dislocation networks at which extrinsic dislocations from the matrix accumulate. (c) Small bowed dislocation segments are emitted from the networks. (d) Some nodes of the dislocation networks (indicated by the arrows) appear destroyed.
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and the gradual increase of mean free path of moving dislocations.
DISCUSSION The information obtained from the deformed microstructures will be used to compare the relative creep resistance of the two alloys studied and, in particular, the hardening effect observed during the primary creep (see previous paper). The instantaneous strain observed in the duplex alloy is almost three times larger than that of the lamellar material but the hardening is more pronounced in the latter. On the other hand, the microstructures observed at 2% strain, i.e. just at the onset of the minimum creep rate, will help us deduce the deformation mechanisms responsible for the high stress exponents measured at this stage. The microstructural observations have provided evidence that the primary stage of creep of the duplex alloy is entirely influenced by the extensive activity of mechanical twinning. We have shown that even at low stresses (280MPa) the twinning
Fig. 10. Examples of deformed microstructures
observed from the lamellar alloy after 2% creep strain. (a) Some y lamellae deform by mechanical twinning. (b) Some y lamellae deform by dislocation activity. (c) Detail from (a) showing bowed dislocations being emitted from a twin interface.
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activity increases substantially between 1 and 2% strain (compare Figs 2 and 3) subdividing the initially large y grains during the primary stage of creep. Thus, the initial mean free path of mobile dislocations decreases from W-100 vrn, which was
the initial grain size, to about 5-6pm which is the distance between twin interfaces measured after 2% strain. During this period, the strain rate decreases continuously to reach a minimum value. At the same time as the density of twin interfaces increases, so does the density of l/2 < 1 10 > mobile dislocations which are also activated at twin intersections with one another or with grain boundaries. Emission of bowed l/2 < 1 I 0 > dislocations was already observed from these intersections at low stresses and 1% strain (see Fig. 1) as well as from twin interfaces at high stress and 2% strain (see Fig. 7(b)), just as the minimum creep rate was being reached. Thus, the mechanisms of deformation responsible for the hardening effect leading to the minimum creep rate (see Fig. 4, previous paper) might be the same in both low and high stress regime even though, in the former, a stress exponent n = 6 was obtained while in the
Fig. 11. conditions) from the lamellar alloy after 2% creep strain at 370MPa showing an example of an ~2 lath broken down (indicated by the arrows) as the y matrix exhibits a high dislocation activity. (b) Weak-beam detail from (a) showing detachment of a small particle from the long u2 lath. (c) Details of superdislocations attached to the small 012broken down particles.
Fig. 12. (a) Microstructure
observed from the lamellar alloy deformed to 2% creep strain under 400MPa confirming the breaking down of a a2 lath. (b) Detail from (a) seen by weak beam contrast. Note the small bowed dislocation segment being emitted from the ~zz/y interface that reflects the existence of stress concentrations.
II. Deformed microstructures during creep of TiAl alloys
latter a much larger value, n = 19, was measured (see Fig. 6(b), previous paper). The hardening effect provided by the twin interfaces has been confirmed by the increase in creep resistance observed after predeforming to 1% strain at low stress (see Figs 5(a) and (b), previous paper). We have seen that the deformed microstructure after 1% creep strain at 280 MPa consists of a sufficient density of mechanical twins (see Figs 1 and 2) to reduce the mean free path of mobile dislocations. This can account for the lower strain rate observed for the prestrained sample on reloading at 342 MPa relative to that measured without prior pre-strain. In this sense, the twin interfaces block the dislocation motion, acting in the same way as the y/y or c~/y interfaces in a lamellar alloy. For this reason, on reloading, the pre-strained duplex alloy approaches a creep resistance equivalent to that of the lamellar structure (see Fig. 5(b), previous paper) as had previously been predicted.i3 From the analysis made of the line directions of single dislocations active within the y matrix, we have evidence that at low stresses both climb and glide mechanisms were possible. At high stresses, however, there was no evidence for possible climb as the majority of the dislocation segments have screw character (see Fig. 7(a)). This indicates that the mechanism controlhng the mobility of single dislocations within the y matrix is determined by the applied stress. At low stresses, climb of edge segments is possible because the chemical driving force on the mobile dislocation density produces a climb velocity sufficiently fast to allow dislocation annihilation at the same rate as the dislocation production. At high stresses, the climbing rate for the corresponding mobile dislocation density is not sufficiently fast and the elimination of mobile dislocations cannot be achieved at the required rate relative to their production for the given stress. Then, the deformation takes place by glide only. As a consequence of the glide process, elongation of the screw segments occurs as the edge segments propagate. This argument is applicable only to the mobile l/2 < 1 10 > dislocations since the twinning dislocations propagate only by glide, providing a significant fraction of thle total strain of the material. The contribution that twinning dislocations make to the total strain of the material can be estimated for the two extreme values of stress used, from which measurements of the twin spacings and of the twinned layer thickness have been made. From the microstructure obtained after 2% strain
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at 360MPa, the spacing between twins in the matrix was measured as about dtwin= l-2 pm, and the twinned layer thickness, ftwin= 60-80 nm. Assuming that all the twins traverse the grains of diameter D = 100 pm, and section area, S= 8 x 1O-9m2, the number of twins per grain is given by Nt = D/dt,i, = 50 and the twin density (number of twins/area) is N,/S= 6x 109/m2. From our present and previous analysis,‘9 and in agreement with the mechanism proposed by Jin and Bieler,21 the twin propagation occurs by glide of l/6 < 1 12 > dislocations on every adjacent {1 1 1) plane and the number of twin dislocations can be obtained by dividing the twinned layer thickness by the interplanar spacing of (1 1 l} planes (dl11 = 0.2316 nm). Thus, for an average layer thickness of 70nm, we obtain a value of about 300 dislocations per twin. Multiplying this number by the twin density, we calculate a twinning dislocation density, Ptwin= 1.8 x 10i2/m2. Since these dislocations traverse the entire grains, the total shear strain produced by this twinning process is ytwin= Ptwinb D, where b is the Burgers vector of the twinning dislocation z 1.6x lo-iOm and D the grain size. This yields a contribution to the total creep strain ytwin, produced by the twin activity, ytwin= 0.0288. Th is value corresponds to a macroscopic strain &twin= ytwin/3 = 0.0096 ( w 1%) which is half the total strain of the specimens studied. It should be mentioned that in our duplex alloy about 20% of the grains have a lamellar structure in which the twinning activity was less pronounced. However, we have treated them in the same way as the equiaxed grains for our calculations, Even though our calculations might not be very precise and we might be overestimating the contribution from twinning to the total strain, it is evident that this contribution is certainly as important as that due to the mobility of single dislocations. Similar calculations made from the specimen crept at the lower stress of 280MPa, where the matrix spacing between twins was measured as about dtwin = 5-6pm and the twinned layer thickness, tt,in = 3@-40nm, yields a twin density of 2.5~ lo9 m2 and a twinning dislocation density, Ptwin= 4’5 X 10” m2. The total contribution of the twin activity to the shear strain is then ytwin= 0*0072 and this corresponds to a macroscopic creep, &twin= 0.0024. We note that this contribution is about a quarter of that obtained at the high stress, indicating that at lower stresses there is a much larger contribution from the mobile l/2 < 1 10 > dislocations to the total strain. To a first approximation we can conclude that the
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contribution from mechanical twinning to the total creep strain is about 10 and 50% in the low and high stress regimes, respectively. During the primary stage of creep, both alloys exhibit very pronounced hardening even though the deformed microstructures observed appear rather different. As already discussed by other authors,i3,27 the better initial creep resistance of the lamellar alloy has to be attributed to the presence of many interfaces that reduce the mean free path of mobile dislocations. We have confirmed that, in this alloy, the major deformation mechanism is due to the mobility of l/2 < 1 10 > dislocations and that twinning only appears more important at strains near the minimum creep rate and larger. Thus, the instantaneous strain in this alloy (see previous paper) was due to the initial activation of a higher dislocation density from the high density of interfaces acting as dislocation sources. On the other hand, the hardening represents a blocking process whereby the mobile dislocation density decreases. Since the latter did not appear to change significantly within the y lamellae between 1 and 2% strain, the hardening should be attributed to the immobilisation of dislocations at the interfaces. This has been confirmed by the presence of accumulated extrinsic dislocations at interface networks (see Fig. 9(b)). This dislocation accumulation is responsible for the increase in stress concentrations which will in turn lead to emission of bowed segments into the matrix. Figures 9(c) and (d) confirm this emission as well as the local destruction of some of the nodes of the network. Thus we propose that the hardening effect in this alloy is produced by the accumulation of dislocations at interfaces that leads to higher stress concentrations required to initiate mobility of dislocations with increasing strain. At the end of the hardening stage, in spite of the different deformation mechanisms involved, a minimum creep rate is reached in both alloys. At this stage, in both cases (and for the same stress regime) a high stress exponent, n = 19, has been measured (see previous paper). This exponent was calculated using the minimum creep rates and therefore represents the mechanism of deformation at that strain. The microstructures observed at that strain are described by the presence of a high density of twin interfaces subdividing the y grains in the duplex alloy and a high density of c& and v/v interfaces in the lamellar alloy. From all types of interfaces, dislocation emission has been observed under the action of high local stresses. The local stress concentration must be large enough to over-
come any backward force resulting from the line tension as the small segments bow out and expand into the matrix. Thus, in both alloys the high stress exponent represents the necessity to reach high stress concentrations to be able to liberate dislocations from the interfaces. We have shown in the previous paper that the high stress exponents are generally rationalised using the concept of a threshold stress,28y29 i.e. the existence of back stresses necessary to initiate dislocation mobility. Using this concept we have calculated the values of the threshold stresses as 264 and 3 16 MPa for the duplex and lamellar alloys, respectively. The higher threshold stress measured in the lamellar alloy reflects one of two possibilities; either a higher stress concentration is necessary to emit dislocations from the interfaces of this alloy or the higher density of interfaces producing this emission leads to a larger integrated value of the average stress required. We should remember that the behaviour deduced here, the thermally activated glide of dislocations in the matrix is controlled by the frequency of their emission from the interfaces, is also exhibited by other metals and alloys.30 In that case, the creep substructure formed during the primary stage, for example sub-boundary networks, represent the interfaces from which dislocation emission occurs during the secondary stage of creep. The existence of high local stress concentrations has also been confirmed and their values have been measured.31 The formalism used to describe the creep rate in materials where the controlling mechanism is the thermally activated glide of dislocations is generally:32
In this case, the term Eo represents a temperature independent parameter that will depend on the microstructure. The term inside the exponential will correspond to the equivalent apparent activation energy that we can express: Q = Q. - Abar,,. The value, Qo, corresponds to the true activation energy necessary to overcome the obstacle by a mechanism that controls the deformation process. The value, oioC, is the local stress necessary to help the thermal energy such that the thermally activated event can be made to occur. The term, Ab, is known as the activation volume, although it is better expressed as an activation area, A, over which the activated event occurs, multiplied by the Burgers vector in order to give the dimensions of a
II. Deformed microstructures during creep of TiAl alloys
volume. In this equation the apparent activation energy Q can be considered as being the sum of two terms: the thermal energy itself, Qo, and a mechanically produced energy corresponding to Abel,,. The use of such an equation requires that the values of oloC be known in order to separate the thermal part of the controlling deformation mechanism. This is only possible from direct measurements from the dislocation microstructures. On the other hand, the advantage of using such an equation is that we should also be able to apply it for other stress regimes in which the stress exponents are not so high. In the present case we have seen that, at low stresses, the duplex alloy exhibits a value of the stress exlponent n = 6. Also, in that case we have seen that the same deformed microstructures are observed, only the density of interfaces is lower and the extent of stress concentration less important. We conclude that the same formalism of thermally activated glide should be used for the low and high stress regimes of the duplex alloy and future measurements of comparative values of local stresses should provide complete distinction between the behaviour of our two alloys.
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mechanism, i.e. dislocation propagation across the y matrix with subsequent accumulation and emission from twin interfaces in the duplex alloy or from v/v or al/y interfaces in the lamellar alloy. In the latter, the existence of stress concentrations produces destruction of the a2 laths that are sheared and lead to the formation of small ct2particles along the y matrix. (5) The high stress exponent, II = 19, measured at the onset of the minimum creep rate in both alloys, has been interpreted as due to the existence of back stresses that need to be overcome to initiate dislocation mobility. In this sense, the use of a threshold stress concept represents the stress concentration necessary to emit dislocations from the interfaces. The lower stress exponent, n = 6, measured in the duplex alloy at low stresses, even though the deformed microstructures reflect the same mechanism of deformation, has been interpreted as due to the existence of a lower density of twin interfaces and lower stress concentration required for dislocation emission.
CONCLUSIONS ACKNOWLEDGEMENTS
(1) The creep behaviour of two TiAl alloys with espective duplex and lamellar structures has been examined after carrying out systematic observations of the deformed microstructures. (2) In the duplex alloy, extensive twinning activity has been observed within the y grains that increases with increasing strain and becomes more pronounced at higher stresses. The twin interfaces subdivide the y grains during the primary stage of creep leading to a decrease in the mean free path of mobile dislocations and a corresponding decrease in strain rate. During the primary stage of creep and up to a strain of 2%, the contribution from mechanical twinning to the total creep strain is about 50% at high stresses and only 10% at low stresses. The hardening effect produced by twin (3) interfaces has been used to increase the creep resistance of the duplex alloy by carrying out predeformation tests at low stress prior to the creep tests at higher stresses. At the strain where the minimum creep rate (4) occurs, the two alloys deform by the same
The authors wish to thank Professor D. G. Morris for reading the two manuscripts and for many useful discussions. The Swiss National Science Foundation is also acknowledged for its help in financing this study.
REFERENCES I. Kim, Y.-W, in Ordered Intermetallics, Part III: Gamma Titanium Aluminides, Journal of Metals, 1994, 46, 30. 2. Kim, S., Cho, W. and Hong, C.-P., Mat. Science Tech., 1995, 11, 1147. 3. Schwenker, S. W. and Kim, Y.-W., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi. TMS, Las Vegas, USA, 1995, 985. 4. Worth, B. D., Jones, J. W. and Allison, J. E., Met. Mat. Trans., 1995,26A, 2947. 5. Martin, P. L., Mendiratta, M. G. and Lipsitt, H. A., MetaN. Trans., 1970, 14A, 2170. 6. Huang, S.-C, MetaN. Trans., 1992, 23A, 375. 7. Wang, J. N, Schwartz, A. J., Nieh, T. G. and Clemens, D., Mat. Science Eng., 1996, A206, 63. 8. Bartholomeusz, M. F., Yang, Q. and Wert, J. A., Scripta Metall., 1993, 29, 389. 9. Huang, S.-C. and Shih, D. S., in Properties and Microstructures of High Temperature Materials, ed. Y.-W. Kim and R. Boyer, TMS, Warrendale, PA, USA, 1991, p. 105. 10. Bartholomeusz, M. F. and Wert, J. A., Met. Mat. Trans., 1994, 25A, 2161.
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11. Seo, D. Y., An, S. U., Bieler, T. R., Larsen, D. E., Bhowal, P. and Merrick, H., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi, TMS, Las Vegas, USA, 1995, p. 745. 12. Feng, C. R., Smith, H., Michel, D. J. and Crowe C. R., in Intermetallic Matrix Composites, ed. D. L. Anton et al., MRS, Pittsburgh, USA, 1990, pp. 194-219. 13. Beddoes, J., Wallace, W. and Zhao, L., Int. Mat. Rev., 199540, 197. 14. Wang, J. N., Schwartz, A. J., Nieh, T. G., Liu, C. T., Sikka, V. K. and Clemens, D., in Gamma Titanium Aluminides, ed. Y.-W. Kim, R. Wagner and M. Yamaguchi, TMS, Las Vegas, USA, 1995, p. 949. 15. Jin, Z. and Bieler, T. R., Scripta Metall. Mat., 1992, 27,
1301. 16. Morris, M. A., Phil. Mag. A, 1993, 68, 259. 17. Morris, M. A., Intermetallics, 1996, 4, 417. 18. Yamaguchi, M. and Umakoshi, Y., Progr. Mat. Science, 1990,34, 1. 19. Luster, J. and Morris, M. A., Metall. Mat. Trans., 1995, 26A, 1745.
20. Farenc, S., Coujou, A. and Couret, A., Phil. Mag., 1993, A67, 127. 21. Jin, Z. and Bieler, T. R., Phil. Mag., 1995,71A, 925. 22. Couret, A. et al., Private Communication at Workshop of
Intermetallics, Birmingham, UK, 17-18 April, 1996. 23. Morris, M. A. and Lipe, T., Scripta Metall. Mat., 1994, 31, 689. 24. Appel, F., Christoph, U. and Wagner, R., Phil. Mag., 1995, 72, 341. 25. Morris, M. A,, Phil. Mag A, 1994, 69, 129. 26. Morris, M. A., Phil. Mag A, 1993, 68, 237. 27. Es-Souni, M., Bartels, A. and Wagner, R., Mat. Sci. Eng.,
1993, A171, 127. 28. Lund, R. W. and Nix, W., Acta Metall., 1976, 24, 469. 29. Burt, H., Dennison, J. P. and Wilshire, B., Metals Sci., 1979, 13, 295. 30. Caillard, D. and Martin, J. L., Acta Metall., 1982, 30,437. 31. Morris, M. A. and Martin, J. L, Acta Metall., 1984, 32, 549. 32. Poirier, J. P., Plasticitt a Haute Temperature des Solides Cristallins. Editions Eyroles, Paris, 1976.
ELSEVIER
5 (1997)
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0 1997 Elsevier Science Limited in Great Britain. All rights reserved
SO966-9795(97)00002-2
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+ 0.00
II. Deformed microstructures during creep of TiAl alloys: role of mechanical twinning M. A. Morris & M. Leboeuf Institute of Structural Metallurgy.
University of Neuchdtel, Av. Bellevaux, 51, 2000 Neuchritel. Switzerland
(Received 25 November 1996; accepted 6 January 1997)
The deformation mechanisms responsible for the creep behaviour of two TiAl alloys with duplex and lamellar structures have been analysed after carrying out microstructural observations as a function of creep strain and applied stress at 700°C. The duplex alloy exhibits extensive mechanical twinning and the new twin interfaces subdivide the y grains throughout the primary stage of creep. At the onset of the minimum creep rate, the twin interfaces in the duplex alloy behave in the same way as the y/y or the oa/y interfaces in the lamellar alloy. Thus, dislocation accumulation and emission from the interfaces occur under the action of the high local stresses. The high stress exponent, n = 19, measured in both alloys at the minimum strain rate has been interpreted in terms of a threshold stress representing the high local stresses required to liberate dislocations from the interfaces and initiate their propagation within the y matrix. 0 1997 Elsevier Science Limited Key words: A. titanium aluminides based on TiAI, B. creep, twinning, D. defects:
planar faults, F. electron microscopy,
Some studies of creep deformation of duplex microstructures1’~i2 have shown that a reduction in lamellar interface spacing occurs due to mechanical twinning. It has been suggestedI that if the twin interfaces act as do lamellar interfaces, they would provide barriers for dislocation motion and lead to similar hardening. However, mechanical twinning has been shown to decrease at low stressesi or high temperatures. i5 For two-phase alloys, an increased twinning activity has been observed at temperatures below 760°C and stresses over 200 MPa. The aim of the present study was to analyse and interpret the deformation structures produced by creep at a temperature of 700°C and high stresses (above 250MPa) in alloys with initial duplex and lamellar microstructures. The alloys chosen have previously been studied and the presence of twins confirmed after tensile (or compression) deformation at 700”C.‘6~17 The evolution of the microstructure during the primary stage of creep and, in particular, the increase of twinning activity with increasing strain have been studied. Also, the effect of pre-deformation on twinning activity has been evaluated in terms of its contribution to improvement of the creep resistance observed in the duplex
INTRODUCTION The study of deformed microstructures produced during creep of y-based TiAl alloys appears relevant to the understanding of the fundamental aspects that contribute to the primary stage of creep and lead to the onset of a minimum strain rate. In particular, the higher creep resistance of alloys with lamellar (CQ+ y) microstructures has been well establishedi but not much work has been carried out to understand the microstructural evolution during the primary stage of creep of alloys with initially duplex and/or lamellar microstructures. Also, it has been suggested that the creep resistance of TiAl alloys can be improved by addition of alloying elemNents5or introduction of a lamellar structure.6 The primary creep might also be reduced by pre-straining at high stress prior to creep-testing at a lower stress7 However, systematic studies of the pre-deformed microstructure before and after changing the stress have not been made. The better creep resistance of the fully lamellar structure is generally attributed to the CQlaths and r/r interfaces acting as barriers to slip8,9 and to the development of constraints at the o/a/y interfaces.” 339
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alloy. Similarly, the effect of stress on twinning activity and its influence on the creep parameters obtained from both alloys have been analysed.
EXPERIMENTAL
obtained by tilting the specimens to different known orientations (zone axes) from which different diffraction vectors were chosen to obtain invisibility under some contrast conditions. Weakbeam images were taken using the g: 3g condition.
DETAILS
The two alloys compared in the present study had compositions Ti-48Al-2Mn-2Nb and Ti-48Al (atoh), respectively. Details about alloy preparation and initial microstructures have been described elsewhere.16T17The creep tests were carried out in tension on cylindrical specimens machined to a diameter of 3mm and with 28mm gauge length. Their surface was subsequently mechanically and electro-polished using a solution of 5% perchloric acid in methanol. The creep machine used was equipped with an Andrade-type cam with a profile that allowed the stress to be maintained constant throughout the test. Tests were performed at 700°C for stresses ranging between 250 and 360MPa for the duplex alloy and between 360 and 430 MPa for the lamellar alloy. The temperature of the specimens was maintained constant with a precision of kO.2”C. The elongation of the specimens was monitored with a precision of about f 2p.m. Although some tests were performed to rupture (see previous paper), those for microstructural observations were arrested at 1 or 2% deformation, the latter corresponding to the strain just before the minimum creep rate was measured. At this stage, the specimens were rapidly cooled under load to avoid any dislocation rearrangement prior to analysis. Microstructural analysis of the deformed specimens was performed by transmission electron microscopy (TEM) using a Philips CM200 microscope equipped with EDS facilities, at 200KV. Also, to obtain information about any large scale substructure formation during the primary stage of creep, scanning electron microscopy (SEM) observations were made using crystallographic contrast from backscattered electrons. For these TEM and SEM observations, discs were cut perpendicular to the tensile axis from the deformed samples and electro-polished. The electro-polished samples were prepared by standard jet polishing techniques using a solution of 5% perchloric acid, 30% butan-l-01 and 65% methanol at -20°C and 50mA, but only the TEM samples were polished to produce a hole. At least two to three different foils were examined from each crept specimen and several y grains or lamellae were studied from each case. Dislocation analysis was carried out from projected images
RESULTS
AND COMPARISONS
After deformation to different strains during the primary stage of creep, the deformed microstructures were studied both by SEM observations made by crystallographic contrast (using backscattered electrons) and by TEM images obtained from the same deformed specimens. The primary stage of creep ends at the strain for which the minimum strain rate was reached and in our alloys this occurred generally at about E =0~02-0.025. Since the strains at which the microstructures could be compared were rather limited, it was important to take into account the microstructures of the undeformed samples as described and illustrated in the first part of this study (see previous paper). Although several specimens were examined after testing at different stresses, only the microstructures of the samples corresponding to the lowest and highest stresses are shown. The intermediate conditions exhibited the same features and have been omitted. Microstructures in the duplex alloy The main feature observed in the deformed microstructures of the duplex alloy was the extensive twinning activity which was already present after only 1% strain in the specimens crept at the lower stress of 280 MPa. Figure l(a) shows an example from this specimen, taken at low magnification in the TEM, of one of the large y grains that was completely traversed by several twins. Detailed diffraction analysis of these twins and others observed under different conditions of strain or applied stress confirmed that they all exhibited the true-twin relationship between the matrix and the twinned layer. In this sense, the twinned interfaces were exactly the same as the r/y interfaces of the true-twin type defined by Yamaguchi and Umakoshii* as producing a rotation of 180” around a < 1 1 1 > pole and analysed in some of our previous work.i6,i9 This twinned microstructure was confirmed to be produced by emission and glide of l/6 < 1 12 > dislocations on {1 1 l} planes, confirming our previous studies’6~‘7~19and those by other authors.20,2’ The extensive twinning activity
II. Deformed microstructures during creep of TiAl alloys was also responsible for the activation of large numbers of single l/2 < 1 10 > dislocations at intersections with some of the v/v interfaces present. Details from some of these intersections (labelled A and B) are s(hown in Figs l(b) and (c), respectively. In these figures we also note that some of the l/2 < 1 10 > dislocations remain on the {1 1 l} planes of the twin interfaces, as also evidented during in situ studies of twin propagation in
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yTi-Al by Couret et al. 22 Besides, some of the twins do not propagate across the entire grain (as seen in two of them within the area labelled B). Thus, it appears that the local stress concentration is limited and cannot provide the twinning driving force to cross the entire grain as suggested by Jin and Bieler.21 Since the number of large grains that could be examined in the TEM was rather limited, we
Fig. 1. (a) Example of extensive twinning traversing an equiaxed y grain observed by TEM from the duplex alloy after 1% creep strain under 280 MPa. (b), (c). Details from areas A and B in Fig. l(a) confirming emission of l/2 < 1 1 0 > dislocations.
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examined complete deformed sections by SEM from discs cut perpendicularly to the tensile axis from the gauge length of the samples. In order to examine these discs in the SEM, they were electropolished; the crystallographic contrast allowing the observation of twins is due to the fact that the thin twinned regions are rotated by 180” with respect to the matrix. Therefore, a thin slab of twinned region appears as a line due to its different crystallographic contrast. Figure 2 shows examples comparing grains taken from sections of the undeformed specimens (head of tensile sample) and
after creep to 1% deformation at low stress (280MPa). We note that the major difference between the grains of undeformed (Fig. 2(a)) and deformed (Fig. 2(b)) material is that while in the former there is a total absence of mechanical twins, in the latter the appearance of mechanical twins is similar to that observed by TEM under the same creep conditions (compare Figs l(a) and (b)). These SEM observations confirmed that although some of the grains did not exhibit twinning after 1% strain, the majority of them did and in all cases only one twin system was activated.
Fig. 2. Comparison between y grains observed by crystallographic contrast in the SEM from the duplex alloy (a) undeformed state (b) after 1% creep strain at 280 MPa. The latter confirms the extensive twinning observed by TEM in Fig. l(a).
II. Deformed microstructures during creep of TiAl alloys
After creep to 2% strain (under the same stress of 280MPa), the SEM observations revealed a more pronounced twinning activity, as illustrated in Fig. 3. Here we see that the twin transfer takes place across adjacent grains, leaving large steps at the grain boundaries (Fig. 3(a)) and also that two twin systems were active in some cases, as in Fig. 3(b) (note that the bright elongated globules are some of the CX~ phase). These features were also encountered in TEM observations, as illustrated in Fig. 4(a). Here we see a large area taken at low magnification traversed by two twin systems that intersect. This area has been analysed in detail and Figs 4(b) and (c) show detailed examples taken from this area (labelled ,4 and B) at higher magnification and under different diffraction conditions.
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Apart from the twins, single dislocations were active and they are invisible in Fig. 4(a) taken with g= 002. Their visibility in Figs 4(b) and (c) taken with g = 200 and 1i 1, respectively, confirms that the Burgers vector is 1/2[1iO]. Also, in Fig. 4(c), the twin system labelled 1 is invisible with g = 1i 1 and its analysis has confirmed it to be 1/6[i 121 (1 1 1). The cross-twin system has been analysed as 1/6[ 1 121 (i i 1). From the micrographs in which the twin interfaces were imaged parallel to the electron beam, the thickness of the twinned layers measured was about 30-40 nm. The 1/2[ 1iO] dislocations seen in Fig. 4(c) have a rather angular morphology and their two distinct directions have been analysed as [ 1121 and [loll, respectively. Although these directions lie on the same (7 11)
Fig. 3. SEM observations made from the duplex alloy after 2% creep strain at 280MPa. (a) Twinning transfer across adjacent y grains produces steps at grain boundary. (b) Example showing that several y grains deform by mechanical twinning, including two systems in some cases. Note that the bright globules are from the a2 phase.
M. A. Morris, M. Leboeuf
Fig. 4. Microstructures observed by TEM from the same specimen of Fig. 3, i.e. after 2% creep strain at 280 MPa. (a) T‘ypical y grain deformed by two intersecting twin systems (b), (c). Details from areas A and B in Fig. 4(a). (Note that all the active dislocations have Burgers vector 1/2[1iO] and are invisible in Fig. 4(a) taken with g = 002). (d) Details from the oz globules (seen in Fig. 4(a) as elongated dark particles). The weak-beam image shows that accumulation of matrix dislocations occurs at the ~Z/Y interface while the bright-field micrograph shows evidence of shear across the az phase.
II. Deformed microstructures during creep of TiAl alloys
(4
Fig. 4. Contd
plane, their Burgers vector does not belong on it, indicating that these dislocations have climbed (since one of them has pure edge character). Other details present in these images are the globular, elongated a2 particles whose interfaces seem to accumulate dislocations but which can also be sheared by the stress concentrations produced as the twins from the matrix intersect them. An example of these sheared particles is shown in Fig. 4(d) (the weak-beam image shows details of the (11~/yinterface). The effect produced on the deformed microstructure by the increase in applied stress is only that of accentuating the details already described for the sample crept at low stress. Figure 5 shows the typical microstructure observed by SEM from the sample crept to the same 2% strain previously shown in Fig. 3 but tested under a higher stress of 360MPa (the higher stress used for this alloy). We see that the crystallographic contrast obtained from these grains confirms the more pronounced twinning activity at the higher stress level, the two active twin systems be:ing emitted from grain boundaries in most cases. Under these high stresses
345
it was rare to find grains that were not deformed by mechanical twinning. Figure 6 shows the comparison of a TEM microstructure seen from the same specimen crept 2% at 360 MPa (Fig. 6(a)), together with an SEM image (Fig. 6(b)) taken from the detailed area labelled A in Fig. 5. We note that for the same amount of strain, the twin density is much higher in this specimen crept at 360MPa than that observed at 280MPa. As a result, the subdivision that occurs in the y grains is much finer in the specimens crept at the higher stress, i.e. the spacing between twinned layers was about 5-6hm after 2% strain and 280 MPa (see Figs 3 and 4) and I-2pm in the case of the test performed at 360MPa (see Figs 5 and 6). The thickness of the twinned layers measured from this specimen tested at high stress was about 60-80nm, confirming that a larger number of l/6 < 1 12 > dislocations have contributed to their formation compared to those formed under the low stress regime. From Fig. 6(a) we also note that, apart from mechanical twins, the y matrix is also deformed by slip of dislocations, which have been analysed. Figure 7(a) shows details from these dislocation segments whose Burgers vector has been confirmed to be 1/2[1 TO] (as they are invisible with g = 002 but visible with g = 1 i 1) and appear elongated along their [l TO] direction, i.e. they have screw character. Therefore, at higher stresses the dislocations propagate by glide as observed in our previous study,t7 i.e. by movement of the edge segment leaving the screw segment behind. Finally, Fig. 7(b) shows details from these twin interfaces observed by weak-beam contrast in which the single dislocations appear bowed, being emitted in exactly the same way as they have been observed from y/y interfaces in some of our previous studies and those of other authors.23y24 Thus, we confirm that in the duplex alloy, after 2% strain, i.e. just before the minimum strain rate is reached, the twinning activity has produced a subdivision of the equiaxed grains into regions about 20 to 50 times smaller (depending on the applied stress). At the same time, the twin interfaces act as sites for dislocation emission and accumulation and their propagation within the y matrix produces deformation in a similar manner as observed in lamellar structures. I3 Microstructures in the lamellar alloy
The general results obtained from the observations of the deformed lamellar alloy do not differ greatly from those already seen in previous studies.25 SEM
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346
observations carried out from discs cut from the gauge length of the specimen indicated that twinning was prominent only after strains of more than 3%. Figure 8 shows a comparison of these observations made from the undeformed specimen and those crept to strains of 24% at a stress of 400MPa Although a high twin density was observed after 4% strain, this was not the case at 2% strain (compare Figs 8(b) and (c)), even in the case of the test carried out at the higher stress of 430 MPa. TEM observations made from this alloy in the undeformed state16 had confirmed that it had a higher density of dislocations than the duplex alloy17 and this was consistent with the observations carried out from the undeformed head of the tensile creep specimens (see previous paper). The latter were made using several thin foils from
which a large number of y grains or lamellae were studied. After small amounts of strain during the primary stage of creep (l-2%) little variations of the deformed microstructure were seen. This is probably because the dislocation density within the y lamellae was not very high and there was no evidence to distinguish between those dislocations present before and after the creep deformation. Similar comments have also been made by other authors who have studied the crept lamellar structures.27 However, as illustrated in Fig. 9, weakbeam images taken from the c~/y interfaces showed evidence that extrinsic dislocations from the y matrix were accumulating at the existing networks (Figs 9(a) and (b)) and that dislocation segments bowed out to be emitted into the y matrix (Fig. 9(c)), and even that at certain sites destruction of dislocation nodes led to imperfections in
activity observed by SEM in the duplex alloy after 2% creep strain under a sitress of 360 MPa. Note the much finer spacing between twins in all the grains.
Fig. 5. Example of increased twinning
II. Deformed microstructures during creep of TiAl alloys
those interface networks, as seen in the area marked by the arrows in Fig. 9(d). After 2% strain (just reaching the minimum creep rate), more evidence of deformation within the y matrix was observed. In some lamellae twinning was already present even though at this strain the twin contrast had not been seen by SEM. However, the twins were extremely thin and probably this was not sufficient to produce enough crystallographic contrast in the SEM, as was the case in the duplex alloy even after 1% strain. Figure 10 shows examples of lamellae deformed by twinning (Fig. 10(a)) and dislocation slip (Fig. IO(b)) observed after 2% strain at 360 MPa. Also here we show a detail (Fig. 10(c)) from one of the twin interfaces from which bowed dislocations are emitted. These, as well as those seen in Fig. 10(a),
347
were confirmed to be l/2 < 1 10 > segments as they were invisible with g = 002. Figures 11 and 12 show different examples of lamellae deformed 2% at stresses of 370 and 400MPa, respectively. The common feature in these images is the observation of the destruction of the a2 laths that, after breaking down, appear as a2 particles aligned in the centre of the v lamellae. These particles are marked by arrows in Fig. 1l(a), taken with g= 002. Under this diffraction condition, only superdislocations are visible and they remain attached to the cz2particles as shown in the detail of Fig. 1l(c). Also, in Fig. 1l(a) we show the same deformed lamella taken with g = 020 and we note the high dislocation density of Burgers vector 1/2<11O>.Th e d es t ruction of the cz2laths occurs by shear, as illustrated in the detail of Fig. 1l(b).
Fig. 6. Comparison of the mic:rostructures observed by (a) TEM and (b) SEM from a specimen of the duplex alloy crept to 2% straiin under 360 MPa. The SEM image was taken from the area labelled A in Fig. 5. Note the extensive twinning activity and 1high dislocation density relative to the equivalent microstructures obtained at low stress and seen in Figs 3 and 4.
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Here we see that the extrinsic dislocations at the interface delineate the profile of the sheared a2 particle. A complete analysis of the dislocations present within this lamella has been made. The superdislocations seen attached to the (~2particles in Figs 1l(a) and (c) have Burgers vector [TO l] and two line directions have been analysed, namely [Oil] and [l lo], indicating that they lie in two glide planes, i.e. (1 1 1) and (1 i’ 1). From the total analysis of the dislocations seen within this )I lamella, only 10% correspond to the mentioned superdislocations; the rest are single l/2 < 1 10 > segments. Of the latter, 62% have Burgers vector 1/2[i 101 and two distinct line directions were analysed, namely [0 0 11, i.e. the segments had pure edge character, and [0 lo], i.e. at 45” from the direction of the Burgers vector. Those directions define the respective (1 10) and (0 0 1) glide planes containing the Burgers vector and the line direction. However, dislocations with line direction [0 0 l] might have climbed due to their pure edge character. Similarly, the remaining 28% dislocations have Burgers vector 1/2[1 lo] and major line directions [0 0 l] and [0 lo], defining the glide planes (110) and (0 0 l), respectively. From these analyses we conclude that climbing and gliding mechanisms of single dislocations are
operative within the y lamellae at the strain at which the minimum creep rate is reached. Also, enough accumulation of dislocations has occurred at the a2/y interfaces that high stress concentrations are produced and shear of the o2 laths occurs locally. From the example shown in Fig. 12(a), we note the general dislocation arrangements in contact with the 01~interface, and the small bowed dislocation segments seen by the side of the broken
Fig. 7. Details from dislocation activity seen in Fig. 6(a). (a) Screw 1/2[1 1 O] dislocatjon segments analysed within the y matrix. (b) Bowed 1/2[1 lo] segments being emitted from twin interfaces.
Fig. 8. Microstructures observed by SEM from the lamellar alloy. (a) Undeformed state; (b), (c) after 2 and 4% creep strain, respectively, under 400 MPa. Note that twinning across the y grains was pronounced only after more than 2% strain.
II. Deformed microstructures during creep of TiAl alloys
a2 lath, in the weak beam detail of Fig. 12(b), confirms the presence of high local stress concentrations. Therefore, we deduce that saturation of the number of dislocations that accumulate at the q/v interfaces leads to increased stress concentrations that lead to shear of the a2 laths as the minimum creep rate is reached. As a result, the continuous increase in strain rate observed thereafter is due to the slow destruction of the ~3 laths
Fig. 9. Examples of TEM obs,ervations made by weak-beam contrast from the a& interfaces of the lamellar alloy. (a), (b) Examples of dislocation networks at which extrinsic dislocations from the matrix accumulate. (c) Small bowed dislocation segments are emitted from the networks. (d) Some nodes of the dislocation networks (indicated by the arrows) appear destroyed.
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and the gradual increase of mean free path of moving dislocations.
DISCUSSION The information obtained from the deformed microstructures will be used to compare the relative creep resistance of the two alloys studied and, in particular, the hardening effect observed during the primary creep (see previous paper). The instantaneous strain observed in the duplex alloy is almost three times larger than that of the lamellar material but the hardening is more pronounced in the latter. On the other hand, the microstructures observed at 2% strain, i.e. just at the onset of the minimum creep rate, will help us deduce the deformation mechanisms responsible for the high stress exponents measured at this stage. The microstructural observations have provided evidence that the primary stage of creep of the duplex alloy is entirely influenced by the extensive activity of mechanical twinning. We have shown that even at low stresses (280MPa) the twinning
Fig. 10. Examples of deformed microstructures
observed from the lamellar alloy after 2% creep strain. (a) Some y lamellae deform by mechanical twinning. (b) Some y lamellae deform by dislocation activity. (c) Detail from (a) showing bowed dislocations being emitted from a twin interface.
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activity increases substantially between 1 and 2% strain (compare Figs 2 and 3) subdividing the initially large y grains during the primary stage of creep. Thus, the initial mean free path of mobile dislocations decreases from W-100 vrn, which was
the initial grain size, to about 5-6pm which is the distance between twin interfaces measured after 2% strain. During this period, the strain rate decreases continuously to reach a minimum value. At the same time as the density of twin interfaces increases, so does the density of l/2 < 1 10 > mobile dislocations which are also activated at twin intersections with one another or with grain boundaries. Emission of bowed l/2 < 1 I 0 > dislocations was already observed from these intersections at low stresses and 1% strain (see Fig. 1) as well as from twin interfaces at high stress and 2% strain (see Fig. 7(b)), just as the minimum creep rate was being reached. Thus, the mechanisms of deformation responsible for the hardening effect leading to the minimum creep rate (see Fig. 4, previous paper) might be the same in both low and high stress regime even though, in the former, a stress exponent n = 6 was obtained while in the
Fig. 11. conditions) from the lamellar alloy after 2% creep strain at 370MPa showing an example of an ~2 lath broken down (indicated by the arrows) as the y matrix exhibits a high dislocation activity. (b) Weak-beam detail from (a) showing detachment of a small particle from the long u2 lath. (c) Details of superdislocations attached to the small 012broken down particles.
Fig. 12. (a) Microstructure
observed from the lamellar alloy deformed to 2% creep strain under 400MPa confirming the breaking down of a a2 lath. (b) Detail from (a) seen by weak beam contrast. Note the small bowed dislocation segment being emitted from the ~zz/y interface that reflects the existence of stress concentrations.
II. Deformed microstructures during creep of TiAl alloys
latter a much larger value, n = 19, was measured (see Fig. 6(b), previous paper). The hardening effect provided by the twin interfaces has been confirmed by the increase in creep resistance observed after predeforming to 1% strain at low stress (see Figs 5(a) and (b), previous paper). We have seen that the deformed microstructure after 1% creep strain at 280 MPa consists of a sufficient density of mechanical twins (see Figs 1 and 2) to reduce the mean free path of mobile dislocations. This can account for the lower strain rate observed for the prestrained sample on reloading at 342 MPa relative to that measured without prior pre-strain. In this sense, the twin interfaces block the dislocation motion, acting in the same way as the y/y or c~/y interfaces in a lamellar alloy. For this reason, on reloading, the pre-strained duplex alloy approaches a creep resistance equivalent to that of the lamellar structure (see Fig. 5(b), previous paper) as had previously been predicted.i3 From the analysis made of the line directions of single dislocations active within the y matrix, we have evidence that at low stresses both climb and glide mechanisms were possible. At high stresses, however, there was no evidence for possible climb as the majority of the dislocation segments have screw character (see Fig. 7(a)). This indicates that the mechanism controlhng the mobility of single dislocations within the y matrix is determined by the applied stress. At low stresses, climb of edge segments is possible because the chemical driving force on the mobile dislocation density produces a climb velocity sufficiently fast to allow dislocation annihilation at the same rate as the dislocation production. At high stresses, the climbing rate for the corresponding mobile dislocation density is not sufficiently fast and the elimination of mobile dislocations cannot be achieved at the required rate relative to their production for the given stress. Then, the deformation takes place by glide only. As a consequence of the glide process, elongation of the screw segments occurs as the edge segments propagate. This argument is applicable only to the mobile l/2 < 1 10 > dislocations since the twinning dislocations propagate only by glide, providing a significant fraction of thle total strain of the material. The contribution that twinning dislocations make to the total strain of the material can be estimated for the two extreme values of stress used, from which measurements of the twin spacings and of the twinned layer thickness have been made. From the microstructure obtained after 2% strain
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at 360MPa, the spacing between twins in the matrix was measured as about dtwin= l-2 pm, and the twinned layer thickness, ftwin= 60-80 nm. Assuming that all the twins traverse the grains of diameter D = 100 pm, and section area, S= 8 x 1O-9m2, the number of twins per grain is given by Nt = D/dt,i, = 50 and the twin density (number of twins/area) is N,/S= 6x 109/m2. From our present and previous analysis,‘9 and in agreement with the mechanism proposed by Jin and Bieler,21 the twin propagation occurs by glide of l/6 < 1 12 > dislocations on every adjacent {1 1 1) plane and the number of twin dislocations can be obtained by dividing the twinned layer thickness by the interplanar spacing of (1 1 l} planes (dl11 = 0.2316 nm). Thus, for an average layer thickness of 70nm, we obtain a value of about 300 dislocations per twin. Multiplying this number by the twin density, we calculate a twinning dislocation density, Ptwin= 1.8 x 10i2/m2. Since these dislocations traverse the entire grains, the total shear strain produced by this twinning process is ytwin= Ptwinb D, where b is the Burgers vector of the twinning dislocation z 1.6x lo-iOm and D the grain size. This yields a contribution to the total creep strain ytwin, produced by the twin activity, ytwin= 0.0288. Th is value corresponds to a macroscopic strain &twin= ytwin/3 = 0.0096 ( w 1%) which is half the total strain of the specimens studied. It should be mentioned that in our duplex alloy about 20% of the grains have a lamellar structure in which the twinning activity was less pronounced. However, we have treated them in the same way as the equiaxed grains for our calculations, Even though our calculations might not be very precise and we might be overestimating the contribution from twinning to the total strain, it is evident that this contribution is certainly as important as that due to the mobility of single dislocations. Similar calculations made from the specimen crept at the lower stress of 280MPa, where the matrix spacing between twins was measured as about dtwin = 5-6pm and the twinned layer thickness, tt,in = 3@-40nm, yields a twin density of 2.5~ lo9 m2 and a twinning dislocation density, Ptwin= 4’5 X 10” m2. The total contribution of the twin activity to the shear strain is then ytwin= 0*0072 and this corresponds to a macroscopic creep, &twin= 0.0024. We note that this contribution is about a quarter of that obtained at the high stress, indicating that at lower stresses there is a much larger contribution from the mobile l/2 < 1 10 > dislocations to the total strain. To a first approximation we can conclude that the
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contribution from mechanical twinning to the total creep strain is about 10 and 50% in the low and high stress regimes, respectively. During the primary stage of creep, both alloys exhibit very pronounced hardening even though the deformed microstructures observed appear rather different. As already discussed by other authors,i3,27 the better initial creep resistance of the lamellar alloy has to be attributed to the presence of many interfaces that reduce the mean free path of mobile dislocations. We have confirmed that, in this alloy, the major deformation mechanism is due to the mobility of l/2 < 1 10 > dislocations and that twinning only appears more important at strains near the minimum creep rate and larger. Thus, the instantaneous strain in this alloy (see previous paper) was due to the initial activation of a higher dislocation density from the high density of interfaces acting as dislocation sources. On the other hand, the hardening represents a blocking process whereby the mobile dislocation density decreases. Since the latter did not appear to change significantly within the y lamellae between 1 and 2% strain, the hardening should be attributed to the immobilisation of dislocations at the interfaces. This has been confirmed by the presence of accumulated extrinsic dislocations at interface networks (see Fig. 9(b)). This dislocation accumulation is responsible for the increase in stress concentrations which will in turn lead to emission of bowed segments into the matrix. Figures 9(c) and (d) confirm this emission as well as the local destruction of some of the nodes of the network. Thus we propose that the hardening effect in this alloy is produced by the accumulation of dislocations at interfaces that leads to higher stress concentrations required to initiate mobility of dislocations with increasing strain. At the end of the hardening stage, in spite of the different deformation mechanisms involved, a minimum creep rate is reached in both alloys. At this stage, in both cases (and for the same stress regime) a high stress exponent, n = 19, has been measured (see previous paper). This exponent was calculated using the minimum creep rates and therefore represents the mechanism of deformation at that strain. The microstructures observed at that strain are described by the presence of a high density of twin interfaces subdividing the y grains in the duplex alloy and a high density of c& and v/v interfaces in the lamellar alloy. From all types of interfaces, dislocation emission has been observed under the action of high local stresses. The local stress concentration must be large enough to over-
come any backward force resulting from the line tension as the small segments bow out and expand into the matrix. Thus, in both alloys the high stress exponent represents the necessity to reach high stress concentrations to be able to liberate dislocations from the interfaces. We have shown in the previous paper that the high stress exponents are generally rationalised using the concept of a threshold stress,28y29 i.e. the existence of back stresses necessary to initiate dislocation mobility. Using this concept we have calculated the values of the threshold stresses as 264 and 3 16 MPa for the duplex and lamellar alloys, respectively. The higher threshold stress measured in the lamellar alloy reflects one of two possibilities; either a higher stress concentration is necessary to emit dislocations from the interfaces of this alloy or the higher density of interfaces producing this emission leads to a larger integrated value of the average stress required. We should remember that the behaviour deduced here, the thermally activated glide of dislocations in the matrix is controlled by the frequency of their emission from the interfaces, is also exhibited by other metals and alloys.30 In that case, the creep substructure formed during the primary stage, for example sub-boundary networks, represent the interfaces from which dislocation emission occurs during the secondary stage of creep. The existence of high local stress concentrations has also been confirmed and their values have been measured.31 The formalism used to describe the creep rate in materials where the controlling mechanism is the thermally activated glide of dislocations is generally:32
In this case, the term Eo represents a temperature independent parameter that will depend on the microstructure. The term inside the exponential will correspond to the equivalent apparent activation energy that we can express: Q = Q. - Abar,,. The value, Qo, corresponds to the true activation energy necessary to overcome the obstacle by a mechanism that controls the deformation process. The value, oioC, is the local stress necessary to help the thermal energy such that the thermally activated event can be made to occur. The term, Ab, is known as the activation volume, although it is better expressed as an activation area, A, over which the activated event occurs, multiplied by the Burgers vector in order to give the dimensions of a
II. Deformed microstructures during creep of TiAl alloys
volume. In this equation the apparent activation energy Q can be considered as being the sum of two terms: the thermal energy itself, Qo, and a mechanically produced energy corresponding to Abel,,. The use of such an equation requires that the values of oloC be known in order to separate the thermal part of the controlling deformation mechanism. This is only possible from direct measurements from the dislocation microstructures. On the other hand, the advantage of using such an equation is that we should also be able to apply it for other stress regimes in which the stress exponents are not so high. In the present case we have seen that, at low stresses, the duplex alloy exhibits a value of the stress exlponent n = 6. Also, in that case we have seen that the same deformed microstructures are observed, only the density of interfaces is lower and the extent of stress concentration less important. We conclude that the same formalism of thermally activated glide should be used for the low and high stress regimes of the duplex alloy and future measurements of comparative values of local stresses should provide complete distinction between the behaviour of our two alloys.
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mechanism, i.e. dislocation propagation across the y matrix with subsequent accumulation and emission from twin interfaces in the duplex alloy or from v/v or al/y interfaces in the lamellar alloy. In the latter, the existence of stress concentrations produces destruction of the a2 laths that are sheared and lead to the formation of small ct2particles along the y matrix. (5) The high stress exponent, II = 19, measured at the onset of the minimum creep rate in both alloys, has been interpreted as due to the existence of back stresses that need to be overcome to initiate dislocation mobility. In this sense, the use of a threshold stress concept represents the stress concentration necessary to emit dislocations from the interfaces. The lower stress exponent, n = 6, measured in the duplex alloy at low stresses, even though the deformed microstructures reflect the same mechanism of deformation, has been interpreted as due to the existence of a lower density of twin interfaces and lower stress concentration required for dislocation emission.
CONCLUSIONS ACKNOWLEDGEMENTS
(1) The creep behaviour of two TiAl alloys with espective duplex and lamellar structures has been examined after carrying out systematic observations of the deformed microstructures. (2) In the duplex alloy, extensive twinning activity has been observed within the y grains that increases with increasing strain and becomes more pronounced at higher stresses. The twin interfaces subdivide the y grains during the primary stage of creep leading to a decrease in the mean free path of mobile dislocations and a corresponding decrease in strain rate. During the primary stage of creep and up to a strain of 2%, the contribution from mechanical twinning to the total creep strain is about 50% at high stresses and only 10% at low stresses. The hardening effect produced by twin (3) interfaces has been used to increase the creep resistance of the duplex alloy by carrying out predeformation tests at low stress prior to the creep tests at higher stresses. At the strain where the minimum creep rate (4) occurs, the two alloys deform by the same
The authors wish to thank Professor D. G. Morris for reading the two manuscripts and for many useful discussions. The Swiss National Science Foundation is also acknowledged for its help in financing this study.
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1993, A171, 127. 28. Lund, R. W. and Nix, W., Acta Metall., 1976, 24, 469. 29. Burt, H., Dennison, J. P. and Wilshire, B., Metals Sci., 1979, 13, 295. 30. Caillard, D. and Martin, J. L., Acta Metall., 1982, 30,437. 31. Morris, M. A. and Martin, J. L, Acta Metall., 1984, 32, 549. 32. Poirier, J. P., Plasticitt a Haute Temperature des Solides Cristallins. Editions Eyroles, Paris, 1976.