- Email: [email protected]
Microstructure evolution leading to high strains during high temperature deformation of a Ti–Al intermetallic
DTD=4.1.0
Intermetallics 7 (1999) 1069±1079
Microstructure evolution leading to high strains during high temperature deformation of a Ti±Al intermetallic Maria A. Morris-MunÄoz*, David G. Morris Department of Physical Metallurgy, Centro Nacional de Investigaciones Metalurgicas, CSIC Avda Gregorio del Amo, 28040 Madrid, Spain Received 24 February 1999; accepted 8 March 1999
Abstract The high strains achieved during high temperature deformation of a Ti±Al rolled sheet have been evaluated by following the microstructural evolution of tensile samples tested along the transverse and rolling directions. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. Microstructural re®nement occurs by subdivision of the grains either by the twin interfaces or by subgrain formation followed by recovery/recrystallization processes. The lower strains achieved in the samples deformed along the transverse direction are a result of a more inhomogeneous microstructure due to the dierent deformation mechanisms involved, that include activation of superdislocations at high strains. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; B. Twinning; D. Defects: dislocation geometry and arrangement; D. Microstructure
1. Introduction In the ®eld of new structural materials, the intermetallics based on the Ti±Al system attract much attention because of their speci®c applications as engineering materials. Their low density, high elastic modulus and good mechanical and oxidation properties up to temperatures of 750±800 C make them very attractive for high-temperature applications. Mechanical properties of these intermetallics such as the creep strength, the fatigue limit and the wear tolerance are very sensitive functions of the microstructure [1±3]. For this reason a certain amount of eort has been invested in developing processing techniques such as forging or high temperature rolling [4±6] to produce material with the precise control of microstructure needed to optimize the speci®c properties for each application. Understanding the correlations between such processing, the composition and microstructure and the mechanical properties is of primary importance to obtain the material with best performance for speci®c applications [7,8].
* Corresponding author. Tel.: +34-91-553-7425; fax: +34-91-5347425. E-mail address: [email protected] (M.A. Morris-MunÄoz)
Recent work [9] carried out on rolled sheet of a Ti± 48Al±2Cr alloy con®rmed that the material exhibited mechanical anisotropy of room-temperature tensile deformation: i.e. dierent behaviour when the tensile axis is parallel to the rolling or the transverse direction. At the same time, texture studies carried out in these samples [9] con®rmed that after annealing the rolled sheet exhibited a modi®ed cube texture with the transverse direction aligned along the c-axis of the tetragonal L10 structure, i.e. along the [001] axis, while the rolling and normal directions were aligned parallel to the a-axes of the tetragonal structure, i.e. the [100] and [010] directions. Examination of deformed tensile samples of the sheet material by electron microscopy [10] con®rmed that twinning and ordinary dislocations activated at dierent stages of hardening were responsible for the observed mechanical anisotropy. A similar alloy of composition Ti±46.5Al±4(Cr±Nb± Ta±B) processed by hot rolling and annealing has been examined in the present study to explain the dierent tensile strengths and ductilities measured when tested along the rolling and transverse directions at temperatures between 700 and 900 C. The evolution of the deformed microstructure as a function of strain has been analysed and the active dislocation and twinning mechanisms observed have been considered to explain the measured mechanical properties.
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(99)00019-9
1070
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
2. Experimental The alloy used in the present study had a composition (at%) Ti±46.5Al±4(Cr±Nb±Ta±B) and was supplied by Plansee AG, Reutte, Austria, as ¯at tensile samples (of dimensions 1.5 mm thickness, 4 mm width and 20 mm gauge length) deformed to fracture at temperatures of 700, 800 and 900 C. The ¯at tensile samples had been cut from sheet prepared by hot rolling from HIP powder compacts, followed by an anneal at 1000 C for 2 h. Details of the rolling process have been given elsewhere [11,12]. Two types of deformed samples were examined, namely, after tests performed with the tensile axis along the rolling and the transverse directions, respectively. To understand the deformation mechanisms activated during the early stages of straining, other samples deformed to 1% strain were also examined. Microstructural analysis of the deformed tensile specimens was performed by transmission electron microscopy (TEM) using a Philips CM200 microscope operated at 200 KV, and by scanning electron microscopy (SEM) using a Cambridge 360 Stereoscan microscope. The SEM studies were carried out using backscattered electrons to obtain crystallographic contrast that provides information about grain orientation distribution and any large-scale twinning activity in the samples. Quantitative measurements of the grain size and of the number of grains deformed by twinning were carried out using an image analyser directly attached to the SEM which allowed accumulation of a large number of images. At least 300 grains were measured for each condition. For all TEM and SEM observations, discs were cut out by spark erosion perpendicular to the tensile axis from the ¯at surface of the deformed samples (i.e. the normal directions of the discs were perpendicular to the tensile axis) and electropolished. The electropolished samples were prepared by standard jet-polishing techniques using a solution of 5% perchloric acid, 30% butan1-ol and 65% methanol at ÿ20 C and 50 mA, but only the TEM samples were polished to produce a hole. Dislocation analysis was carried out from projected images obtained by tilting the specimens to dierent known orientations (zone axes) from which dierent diraction vectors were chosen to obtain invisibility under some contrast conditions. Weak-beam images were taken using the g: 3 g condition. 3. Results The initial microstructure of the sheet samples was equiaxed and consisted of small g grains of size about 15±20 mm with the a2 phase as small particles distributed along grain boundaries. Texture studies by X-ray diraction con®rmed that, after annealing, the
heavily rolled sheet exhibited a modi®ed cube texture with the tetragonal c-axis aligned along the transverse direction and the a-axes aligned parallel to the rolling and normal directions [10]. After high-temperature deformation this microstructure evolved with increasing strain owing to the activation of dierent deformation mechanisms at dierent temperatures. At deformation temperatures between 700 and 900 C, the material behaves dierently when tested with the tensile axis along the rolling (RD) or the transverse (TD) directions. Table 1 shows that the the yield strength is higher in the TD sample at the three temperatures tested and that the total strains achieved increase with increasing temperature in the RD samples while they decrease in the TD samples. SEM and TEM observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. Fig. 1 shows examples of the microstructures observed by SEM in the RD samples from discs cut out adjacent to the the fracture surface and to the heads of the tensile samples. We note that the distance between mechanical twins, seen by crystallographic contrast, decreases with decreasing deformation temperature and increasing strain (i.e. the twin spacing is much smaller at 700 C and near the fracture surface than near the heads of the tensile samples). In Table 2 we give the twin spacing, dt , measured in samples deformed at 700 C from discs cut out at dierent positions of the gauge length where the corresponding total strains, "T , have been estimated by measuring the local section of the specimens, ("T Ao ÿ Af =Ao ). From the micrographs in Fig. 1(b) and (d) we note that the twin spacing measured near the fracture surface becomes also the ®nal grain size at these large strains. At 700 C the initial grain size of 10 mm becomes 100 nm after 67% strain in the RD sample and 150 nm after 54% strain in the TD sample. At these large strains the total fraction of grains that deform by mechanical twinning was 80% in the RD sample at both temperatures while in the TD sample it was only 48% at 700 C and 35% at 800 C. On the other hand, at 800 C the twin spacing was three times larger than that measured at 700 C, con®rming that the twin density and, therefore, the total strain produced by twinning dislocations decreased considerably with increasing temperature. Table 1 Flow stress and ductility of both types of samples RD sample
TD sample
Td ( C)
0:2 (MPa)
" (%)
0:2 (MPa)
" (%)
700 800 900
478 374 181
67 70 84
570 433 243
54 51 45
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1071
Fig. 1. Examples of deformed microstructures from the RD samples observed by crystallographic contrast using backscattered electrons in the SEM. (a) Area adjacent to the head of the sample where the twinned structure is much coarser than in the area (b) adjacent to the fracture surface after deformation at 700 C. After deformation at 800 C the twin spacing is larger than at 700 C, both in the area adjacent to the head of the sample (c) and near the fracture surface (d).
Fig. 2 shows a comparison between the microstructures observed in the two samples deformed to fracture at 700 and 900 C. We con®rm the lack of twinning activity during deformation at 900 C while the grain size changes much less as a function of the distance to the fracture surface. In spite of the dierent total strain achieved in the two samples at this temperature (900 C) there is no evidence of dynamic recrystallization and in both cases the reduction of grain size from 10 mm to 2±3 mm just behind the fracture surface appears to take place by subdivision of the Table 2 Twin spacing, dt, measured after deformation at 700 C in discs cut out at dierent positions along gauge length RD sample
TD sample
Foil no.
"T (%)
dt (nm)
"T (%)
dt (nm)
6 (next to head) 4 (middle gauge) 1 (next to fracture)
3.4 36 65
650 300 100
5.3 30 53
500 350 150
grains into smaller subgrains that become the new grains [examples of such subgrain contrast are clearly seen in Fig. 2(c)]. To determine whether mechanical twinning was active during the early stages of the deformation process, microstructural observations were made from specimens deformed at 700 and 800 C to 1% total strain. Fig. 3 shows examples con®rming that twinning was already activated after this small strain, however the twin density was much lower. Quantitative measurements indicated that the total fraction of grains that have deformed by mechanical twinning at this strain was 40± 30% in the RD sample and 38±25% in the TD sample at 700±800 C, respectively. 3.1. TEM microstructures of samples deformed at 700 C The microstructural observations made by TEM generally con®rmed the SEM studies and helped characterize the speci®c slip/twin systems present. In the RD sample these were mostly ordinary 1/2h110i dislocations and mechanical twins as shown in the examples of Fig. 4
1072
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 3. Microstructures observed by SEM after deformation to 1% strain in both types of sample. (a) TD sample deformed at 700 C, (b) RD sample deformed at 800 C. Note that at this small strain mechanical twins are already present in some of the grains.
Fig. 2. Microstructures observed by SEM from the centre of the gauge in both types of samples deformed to fracture. (a) TD sample and (b) RD sample deformed at 700 C; (c) RD sample deformed at 900 C. Note that at this higher temperature mechanical twins are not seen and subgrains have formed.
after deformation at 700 C to 1% strain. The presence of ®ne debris and pinning points along the screw segments, [seen in Fig. 4(a)] characteristic of the deformation process by ordinary dislocations in these alloys is also evident. Fig. 5 shows the increased twin density in
the same sample deformed to fracture but within a foil taken from the middle of the gauge length where the total strain measured was 36%. Similar microstructures were observed in the TD sample but superdislocations were also seen in some of the grains deformed by ordinary dislocations after 1% strain. There was also more evidence of ordinary dislocations being emitted from twin/twin and twin/boundary intersections, as illustrated in Fig. 6. Since the majority of the grains examined had the (010) orientation in the ¯at, untilted foil, it was possible to image a large number of grains using the g=002 diraction vector for which ordinary dislocations are invisible and only superdislocations appear visible. After 30% strain, in a foil taken from the middle of the gauge length of the sample deformed to fracture, the TD sample exhibited two distinct microstructures: grains deformed by both ordinary dislocations and a high density of twins, and grains deformed by
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1073
Fig. 5. (a) Microstructure observed by TEM from the centre of the gauge in the RD sample deformed to fracture with high densities of ordinary dislocations and twins. (b) Detail from (a).
superdislocations with only a few ordinary dislocations, as seen in Fig. 7. 3.2. TEM microstructures of samples deformed at 800 C
Fig. 4. Examples of microstructures observed by TEM after deformation to 1% strain at 700 C in the RD sample. Note that some grains deform by twinning activity and others by ordinary 1/2h110i dislocations (a), diraction vector g=200, zone (010).
From observations after deformation at 800 C, in those grains deformed by ordinary dislocations, there was evidence of recovery/climb processes taking place and large dislocation loops were present. Also the density of ®ne debris was lower and the pinning points along ordinary dislocation segments were less pronounced. This is shown in Fig. 8 for the RD sample deformed to 1% strain. At this stage the twin density was low in both samples with only one twin system activated in those grains that deformed with some contribution of twinning. Fig. 9 shows an example of twinned microstructures observed after 1% strain in the TD
1074
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 6. Microstructure observed at the centre of the gauge in the TD sample deformed to fracture where mechanical twins traverse the grain (a) and produce emission of ordinary dislocations at the intersection with the boundary (b). As a consequence the lower grain deforms by zone (011). ordinary dislocations only. Diraction vector g=1 11,
sample. In the TD sample some slip bands of superdislocations were also seen within grains that deformed by ordinary dislocations, as shown in Fig. 10. In both types of samples, the major dierence observed after the large strains to fracture was a substantial increase in dislocation and twin densities. 3.3. TEM microstructures of samples deformed at 900 C The deformation process at 900 C, as already con®rmed from SEM observations, was characterized by the absence of mechanical twinning in both types of samples. Fig. 11 shows an example of the microstructure seen in the RD sample deformed to fracture (84% strain), in a foil taken from the middle of the gauge length. We note that most of the original grains contain a high density of dislocations and are subdivided into smaller subgrains, while a few (like the one labelled A) appear as new recrystallized grains and with
Fig. 7. Examples of the two types of microstructure observed from the centre of the gauge in the TD sample deformed to fracture. (a) Grain deformed by mechanical twins and ordinary dislocations; Diraction zone (112). (b) Grain deformed by a majority of vector g=1 11, superdislocations; diraction vector g=002, zone (010).
a low density of defects. In Fig. 12 we show one of the grains from the same thin foil sample, imaged under two diraction conditions, con®rming that all the
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1075
Fig. 8. Microstructure observed by TEM from the RD sample deformed to 1% strain at 800 C. Note the large dislocation loops and the lower density of pinning points along the ordinary dislocations seen. Diraction vector g=200, zone (010).
Fig. 10. Example of microstructure observed under two diraction conditions from the TD sample deformed to 1% strain at 800 C. (a) Diraction vector g=111, zone (110); both ordinary and superdislocations are visible. (b) Diraction vector g=002, zone (110); only the bands of superdislocations are visible.
Fig. 9. Example of microstructure from the TD sample deformed to 1% strain at 800 C where twins and ordinary dislocations are present within the two grains observed.
dislocations are of the type 1/2h110i. In Fig. 13(a) we show a similar distribution of grains in the TD sample deformed to fracture (45% strain) while in Fig. 13(b) and (c) we show examples of grains from Fig. 13(a) deformed by a mixture of ordinary and superdislocations. Thus the major dierence observed in the TD sample deformed at 900 C at these large strains was the larger
contribution of superdislocations in a high fraction of grains. Also the material contained a large fraction of recrystallized grains (as opposed to the recovered grains seen in the RD sample) like the examples shown in Fig. 14. 4. Discussion The evolution of the microstructure has been interpreted in terms of the active deformation mechanisms observed in both samples. The RD material has a
1076
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 11. Microstructure observed by TEM from the centre of the gauge in the RD sample deformed to fracture at 900 C. Most of the grains contain a high dislocation density and are subdivided into smaller subgrains and only a few, as the one labelled A, appear as new recrystallized grains.
fraction of grains deformed by mechanical twins that increases from about 40 to 80% between 1 and 3.4% strain (measured adjacent to the head of the tensile sample). This fraction does not change signi®cantly with increasing strain but the intensity of twinning within grains increases substantially between 3.4% strain (near the head of the sample) and 67% strain (adjacent to the fracture surface). This leads to a major re®nement of the microstructure which is more pronounced at 700 C, as evidenced in Fig. 1. This eect is less pronounced in the TD sample because the total strain achieved was only 54% and the fraction of grains deformed by twinning changed from 38 to 48% between 1 and 5.3% (measured adjacent to the head of the tensile sample) at 700 C. The same behaviour is observed in both types of samples at 800 C, except that the fraction of grains that deform by twinning is slightly lower (30% in the RD sample and 25% in the TD sample at 1% strain and saturating at 80 and 35% at about 4% strain in the RD and TD samples, respectively). It seems, therefore, that the fraction of grains that deform by twinning is determined in the early stages of deformation while the total strain achieved must be determined by the evolution of the microstructure during the deformation process. The dierent behaviour of the materials both in terms of strength as well as ductility has been attributed to the inhomogeneity of deformation due to the extent of texture in the samples.
Fig. 12. Example of a grain, imaged under two diraction conditions, from the same area of Fig. 11 in which all dislocations are invisible with g=002, con®rming that they all are of the type 1/2h110i.
At 700 and 800 C in the RD sample the majority of grains (80%) deform by twinning after a strain as low as 3.4% and ordinary dislocations are also observed within the grains. The texture of the sample would allow two twinning systems and two slip systems for ordinary dislocations due to their favourable Schmid factor [10]. In the TD sample the number of grains that deform by twinning is lower, but still signi®cant (38% after 1% strain and 48% after 5.3% strain). This can be attributed to the less favourable texture, i.e. grains which are very close to the [001] direction are not able to twin
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1077
Fig. 14. Typical recrystallized grains observed from the TD sample deformed to fracture at 900 C.
Fig. 13. (a) General microstructure observed by TEM from the centre of the gauge in the TD sample deformed to fracture at 900 C. (b) Example of a grain from (a) deformed by superdislocations; diraction condition g=002, zone (010). (c) Example of a grain from (a) deformed by ordinary and superdislocations; diraction condition g=111, zone (110).
when stressed in tension. Ordinary dislocations are also observed while superdislocations are rare at low strains and presumably their activation is dicult due to their low mobility. Initial deformation therefore occurs inhomogeneously in those grains favourably oriented for twinning and ordinary dislocations. The higher yield strength of this sample should be attributed to the higher resolved shear stress needed to activate twinning in those grains [10]. Once mechanical twinning is active, the local stresses produced at twin/twin and twin/ boundary intersections lead to emission of ordinary
dislocations (see Figs. 6 and 9). The latter can thus be activated by the local stresses even in grains oriented along [001] directions for which these slip systems would be unfavourable according to the macroscopic Schmid factors if they were ideal single crystals factors [10]. At 700 and 800 C in the RD samples deformation is homogeneously distributed as a mixture of ordinary dislocation movement and twinning (80% of grains) while the rest deforms by ordinary dislocations only. The grains that deform by twinning are subdivided by the twin interfaces (see Fig. 1) and the scale of the subdivision is that of the twin spacing which decreases with increasing strain. Across the twin interfaces another subdivision is observed by crystallographic contrast in the SEM as small subgrains that become very ®ne grains at large strains before fracture. These are produced by the accumulation of ordinary dislocations within the deformed substructure as they propagate across the twins. The most important aspect of this structure and evolution is its uniformity which is imposed by the twin
1078
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
interfaces where dislocations can accumulate as deformation continues throughout. In the TD samples there are about 48 to 35% of the grains (depending on the temperature) that initially deform in the same uniform manner as that of the RD sample, i.e. with twins subdividing the grains and ordinary dislocations being produced at their intersections. As local stresses are built up, other grains start to deform by ordinary and superdislocations. This leads to two scales of the microstructure in which the distribution of ®nal grains is much more inhomogeneous in the TD samples. In the RD samples about 80% of the grains ®nally become about 100 nm in size after subdivision by twins and smaller subgrains, accompanied by recovery/recrystallization processes. Only 20% of the grains in the RD samples become 2±3 mm in size due to the subgrains produced by ordinary dislocations only. In the TD samples, however, only about 48±35% of the grains ®nally become 150 nm in size due to subdivision by twins and subgrains and recovery/recrystallization processes. There are 52±65% of the grains that remain larger than 2±3 mm because in the absence of mechanical twinning subdivision by twin interfaces does not occur. The lack of uniformity in these structures leads to inhomogeneity of the deformation process and results in the lower ductility observed in the TD samples. At the highest temperature (900 C), no twinning is observed in either sample, presumably because activation of ordinary dislocations occurs at a much lower stress. Such dislocations both glide and climb, and therefore the stresses necessary to activate twinning [10] are never reached (see Table 1). In the RD sample (at 900 C) the subdivision of the initial grains continues by a mechanism of subgrain formation leading to new grains by such recovery processes. This leads to a uniform and homogeneous microstructure throughout the material (see Fig. 11). In the TD sample (at 900 C) the subdivision of grains into subgrains is not so uniform due to the dierent deformation mechanisms from grain to grain (by ordinary dislocations or by superdislocations). Possibly the lack of mobility of the superdislocations makes recovery processes more dicult and these grains accumulate dislocations until enough driving force exists for recrystallization [see Figs. 13(a) and 14]. The result is a more inhomogeneous structure with some grains recovered and others recrystallized, i.e. with some grains the prior texture remains while others take totally new orientations. This inhomogeneous structure with dierent types of grain boundaries results in incompatibility of deformation across grains. It is such inhomogeneity that is believed to be responsible for the lower ductility than that observed in the RD sample in which most grains (80%) have recovered and accommodated local stresses in the process.
5. Conclusions . The in¯uence of texture on the strength and ductility achieved during tensile deformation of a rolled sheet of a Ti±Al alloy has been studied by following the microstructural evolution during the deformation process. . SEM and TEM observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. At the lower temperatures mechanical twinning was already activated at 1% strain. . In those grains that deform by twinning, the twin spacing decreases with decreasing deformation temperature and increasing strain in both samples. Twinning produces a subdivision of the grains during the deformation process and leads to a ®ner grain size in the RD sample before fracture. . The fraction of grains that deform by twinning is much lower in the TD sample which contains 50± 60% of the grains deforming by either superdislocations or by a mixture of ordinary and superdislocations. The lack of uniformity in these structures leads to inhomogeneity of the deformation process and results in the lower ductility observed in the TD samples. . At 900 C twinning does not occur in either sample and the re®nement of the microstructure is lesspronounced. In the RD sample the majority of grains are subdivided by a mechanism of subgrain formation leading to new grains by recovery processes. This subdivision is not so uniform in theTD sample where some grains deform by less mobile superdislocations that hinder recovery processes and accumulate until recrystallization becomes possible.
Acknowledgements The authors would like to thank Professor H. Clemens for providing the tensile samples and the mechanical data.
References [1] Kim Y-W, Dimiduk DM. Journal of Metals 1991;43:40. [2] Huang SC, Hall EL. Metall Trans A 1991;22:2619. [3] Nakano T, Kawanaka T, Yasuda HY, Umakoshi Y. Mater Science Engineering A 1995;194:43. [4] Liu CT, Maziasz PJ, Clemens DR, Schneibel JH, Sikka VK, Nieh TG, Wright J, Walker LR. In Kim Y-W, Wagner R, Yamaguchi M, editors. g Titanium aluminides. Warrendale (PA): TMS, 1995. p. 679. [5] Wang G, Dahms M. Journal of Metals 1993;45:53.
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079 [6] Watkins D, Piehler HR, Seetharaman V, Lombard CM, Semiatin SL. Metall Trans A 1992;23:2669. [7] Dimiduk DM. In: Kim Y-W, Wagner R, Yamaguchi M, editors. g Titanium aluminides. Warrendale (PA): TMS, 1995. p 3. [8] Huang SC, Shih DS. In: Kim Y-W, Boyer RR, editors. Microstructure/property relationships in titanium aluminides and alloys. Warrendale (PA): TMS, 1990. p 105. [9] Bartels A, Clemens H, Hartig C, Mecking H. In: Koch CC, Liu CT, Stolo NS, Wanner A. High temperature intermetallic alloys VII, Vol. 460. Pittsburgh (PA): MRS, 1997. p 141.
1079
[10] Morris MA, Clemens H, SchloÈg SM. Intermetallics 1998;6:511. [11] Clemens H, Glatz W, Schretter P, Koeppe C, Bartels, A, Behr R, Wanner A. In: Kim Y-W, Wagner R, Yamaguchi M, editors. g titanium aluminides. Warrendale (PA): TMS, 1995. p 717. [12] Clemens H, Glatz W, Eberhard N, Martinz H-P, Knabl W. In: Koch, CC, Liu CT, Stolo NS, Wanner A, editors. High temperature ordered intermetallics VII, vol. 460. Pittsburgh (PA): MRS, 1997. p 29.
Intermetallics 7 (1999) 1069±1079
Microstructure evolution leading to high strains during high temperature deformation of a Ti±Al intermetallic Maria A. Morris-MunÄoz*, David G. Morris Department of Physical Metallurgy, Centro Nacional de Investigaciones Metalurgicas, CSIC Avda Gregorio del Amo, 28040 Madrid, Spain Received 24 February 1999; accepted 8 March 1999
Abstract The high strains achieved during high temperature deformation of a Ti±Al rolled sheet have been evaluated by following the microstructural evolution of tensile samples tested along the transverse and rolling directions. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. Microstructural re®nement occurs by subdivision of the grains either by the twin interfaces or by subgrain formation followed by recovery/recrystallization processes. The lower strains achieved in the samples deformed along the transverse direction are a result of a more inhomogeneous microstructure due to the dierent deformation mechanisms involved, that include activation of superdislocations at high strains. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Titanium aluminides, based on TiAl; B. Twinning; D. Defects: dislocation geometry and arrangement; D. Microstructure
1. Introduction In the ®eld of new structural materials, the intermetallics based on the Ti±Al system attract much attention because of their speci®c applications as engineering materials. Their low density, high elastic modulus and good mechanical and oxidation properties up to temperatures of 750±800 C make them very attractive for high-temperature applications. Mechanical properties of these intermetallics such as the creep strength, the fatigue limit and the wear tolerance are very sensitive functions of the microstructure [1±3]. For this reason a certain amount of eort has been invested in developing processing techniques such as forging or high temperature rolling [4±6] to produce material with the precise control of microstructure needed to optimize the speci®c properties for each application. Understanding the correlations between such processing, the composition and microstructure and the mechanical properties is of primary importance to obtain the material with best performance for speci®c applications [7,8].
* Corresponding author. Tel.: +34-91-553-7425; fax: +34-91-5347425. E-mail address: [email protected] (M.A. Morris-MunÄoz)
Recent work [9] carried out on rolled sheet of a Ti± 48Al±2Cr alloy con®rmed that the material exhibited mechanical anisotropy of room-temperature tensile deformation: i.e. dierent behaviour when the tensile axis is parallel to the rolling or the transverse direction. At the same time, texture studies carried out in these samples [9] con®rmed that after annealing the rolled sheet exhibited a modi®ed cube texture with the transverse direction aligned along the c-axis of the tetragonal L10 structure, i.e. along the [001] axis, while the rolling and normal directions were aligned parallel to the a-axes of the tetragonal structure, i.e. the [100] and [010] directions. Examination of deformed tensile samples of the sheet material by electron microscopy [10] con®rmed that twinning and ordinary dislocations activated at dierent stages of hardening were responsible for the observed mechanical anisotropy. A similar alloy of composition Ti±46.5Al±4(Cr±Nb± Ta±B) processed by hot rolling and annealing has been examined in the present study to explain the dierent tensile strengths and ductilities measured when tested along the rolling and transverse directions at temperatures between 700 and 900 C. The evolution of the deformed microstructure as a function of strain has been analysed and the active dislocation and twinning mechanisms observed have been considered to explain the measured mechanical properties.
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(99)00019-9
1070
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
2. Experimental The alloy used in the present study had a composition (at%) Ti±46.5Al±4(Cr±Nb±Ta±B) and was supplied by Plansee AG, Reutte, Austria, as ¯at tensile samples (of dimensions 1.5 mm thickness, 4 mm width and 20 mm gauge length) deformed to fracture at temperatures of 700, 800 and 900 C. The ¯at tensile samples had been cut from sheet prepared by hot rolling from HIP powder compacts, followed by an anneal at 1000 C for 2 h. Details of the rolling process have been given elsewhere [11,12]. Two types of deformed samples were examined, namely, after tests performed with the tensile axis along the rolling and the transverse directions, respectively. To understand the deformation mechanisms activated during the early stages of straining, other samples deformed to 1% strain were also examined. Microstructural analysis of the deformed tensile specimens was performed by transmission electron microscopy (TEM) using a Philips CM200 microscope operated at 200 KV, and by scanning electron microscopy (SEM) using a Cambridge 360 Stereoscan microscope. The SEM studies were carried out using backscattered electrons to obtain crystallographic contrast that provides information about grain orientation distribution and any large-scale twinning activity in the samples. Quantitative measurements of the grain size and of the number of grains deformed by twinning were carried out using an image analyser directly attached to the SEM which allowed accumulation of a large number of images. At least 300 grains were measured for each condition. For all TEM and SEM observations, discs were cut out by spark erosion perpendicular to the tensile axis from the ¯at surface of the deformed samples (i.e. the normal directions of the discs were perpendicular to the tensile axis) and electropolished. The electropolished samples were prepared by standard jet-polishing techniques using a solution of 5% perchloric acid, 30% butan1-ol and 65% methanol at ÿ20 C and 50 mA, but only the TEM samples were polished to produce a hole. Dislocation analysis was carried out from projected images obtained by tilting the specimens to dierent known orientations (zone axes) from which dierent diraction vectors were chosen to obtain invisibility under some contrast conditions. Weak-beam images were taken using the g: 3 g condition. 3. Results The initial microstructure of the sheet samples was equiaxed and consisted of small g grains of size about 15±20 mm with the a2 phase as small particles distributed along grain boundaries. Texture studies by X-ray diraction con®rmed that, after annealing, the
heavily rolled sheet exhibited a modi®ed cube texture with the tetragonal c-axis aligned along the transverse direction and the a-axes aligned parallel to the rolling and normal directions [10]. After high-temperature deformation this microstructure evolved with increasing strain owing to the activation of dierent deformation mechanisms at dierent temperatures. At deformation temperatures between 700 and 900 C, the material behaves dierently when tested with the tensile axis along the rolling (RD) or the transverse (TD) directions. Table 1 shows that the the yield strength is higher in the TD sample at the three temperatures tested and that the total strains achieved increase with increasing temperature in the RD samples while they decrease in the TD samples. SEM and TEM observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. Fig. 1 shows examples of the microstructures observed by SEM in the RD samples from discs cut out adjacent to the the fracture surface and to the heads of the tensile samples. We note that the distance between mechanical twins, seen by crystallographic contrast, decreases with decreasing deformation temperature and increasing strain (i.e. the twin spacing is much smaller at 700 C and near the fracture surface than near the heads of the tensile samples). In Table 2 we give the twin spacing, dt , measured in samples deformed at 700 C from discs cut out at dierent positions of the gauge length where the corresponding total strains, "T , have been estimated by measuring the local section of the specimens, ("T Ao ÿ Af =Ao ). From the micrographs in Fig. 1(b) and (d) we note that the twin spacing measured near the fracture surface becomes also the ®nal grain size at these large strains. At 700 C the initial grain size of 10 mm becomes 100 nm after 67% strain in the RD sample and 150 nm after 54% strain in the TD sample. At these large strains the total fraction of grains that deform by mechanical twinning was 80% in the RD sample at both temperatures while in the TD sample it was only 48% at 700 C and 35% at 800 C. On the other hand, at 800 C the twin spacing was three times larger than that measured at 700 C, con®rming that the twin density and, therefore, the total strain produced by twinning dislocations decreased considerably with increasing temperature. Table 1 Flow stress and ductility of both types of samples RD sample
TD sample
Td ( C)
0:2 (MPa)
" (%)
0:2 (MPa)
" (%)
700 800 900
478 374 181
67 70 84
570 433 243
54 51 45
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1071
Fig. 1. Examples of deformed microstructures from the RD samples observed by crystallographic contrast using backscattered electrons in the SEM. (a) Area adjacent to the head of the sample where the twinned structure is much coarser than in the area (b) adjacent to the fracture surface after deformation at 700 C. After deformation at 800 C the twin spacing is larger than at 700 C, both in the area adjacent to the head of the sample (c) and near the fracture surface (d).
Fig. 2 shows a comparison between the microstructures observed in the two samples deformed to fracture at 700 and 900 C. We con®rm the lack of twinning activity during deformation at 900 C while the grain size changes much less as a function of the distance to the fracture surface. In spite of the dierent total strain achieved in the two samples at this temperature (900 C) there is no evidence of dynamic recrystallization and in both cases the reduction of grain size from 10 mm to 2±3 mm just behind the fracture surface appears to take place by subdivision of the Table 2 Twin spacing, dt, measured after deformation at 700 C in discs cut out at dierent positions along gauge length RD sample
TD sample
Foil no.
"T (%)
dt (nm)
"T (%)
dt (nm)
6 (next to head) 4 (middle gauge) 1 (next to fracture)
3.4 36 65
650 300 100
5.3 30 53
500 350 150
grains into smaller subgrains that become the new grains [examples of such subgrain contrast are clearly seen in Fig. 2(c)]. To determine whether mechanical twinning was active during the early stages of the deformation process, microstructural observations were made from specimens deformed at 700 and 800 C to 1% total strain. Fig. 3 shows examples con®rming that twinning was already activated after this small strain, however the twin density was much lower. Quantitative measurements indicated that the total fraction of grains that have deformed by mechanical twinning at this strain was 40± 30% in the RD sample and 38±25% in the TD sample at 700±800 C, respectively. 3.1. TEM microstructures of samples deformed at 700 C The microstructural observations made by TEM generally con®rmed the SEM studies and helped characterize the speci®c slip/twin systems present. In the RD sample these were mostly ordinary 1/2h110i dislocations and mechanical twins as shown in the examples of Fig. 4
1072
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 3. Microstructures observed by SEM after deformation to 1% strain in both types of sample. (a) TD sample deformed at 700 C, (b) RD sample deformed at 800 C. Note that at this small strain mechanical twins are already present in some of the grains.
Fig. 2. Microstructures observed by SEM from the centre of the gauge in both types of samples deformed to fracture. (a) TD sample and (b) RD sample deformed at 700 C; (c) RD sample deformed at 900 C. Note that at this higher temperature mechanical twins are not seen and subgrains have formed.
after deformation at 700 C to 1% strain. The presence of ®ne debris and pinning points along the screw segments, [seen in Fig. 4(a)] characteristic of the deformation process by ordinary dislocations in these alloys is also evident. Fig. 5 shows the increased twin density in
the same sample deformed to fracture but within a foil taken from the middle of the gauge length where the total strain measured was 36%. Similar microstructures were observed in the TD sample but superdislocations were also seen in some of the grains deformed by ordinary dislocations after 1% strain. There was also more evidence of ordinary dislocations being emitted from twin/twin and twin/boundary intersections, as illustrated in Fig. 6. Since the majority of the grains examined had the (010) orientation in the ¯at, untilted foil, it was possible to image a large number of grains using the g=002 diraction vector for which ordinary dislocations are invisible and only superdislocations appear visible. After 30% strain, in a foil taken from the middle of the gauge length of the sample deformed to fracture, the TD sample exhibited two distinct microstructures: grains deformed by both ordinary dislocations and a high density of twins, and grains deformed by
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1073
Fig. 5. (a) Microstructure observed by TEM from the centre of the gauge in the RD sample deformed to fracture with high densities of ordinary dislocations and twins. (b) Detail from (a).
superdislocations with only a few ordinary dislocations, as seen in Fig. 7. 3.2. TEM microstructures of samples deformed at 800 C
Fig. 4. Examples of microstructures observed by TEM after deformation to 1% strain at 700 C in the RD sample. Note that some grains deform by twinning activity and others by ordinary 1/2h110i dislocations (a), diraction vector g=200, zone (010).
From observations after deformation at 800 C, in those grains deformed by ordinary dislocations, there was evidence of recovery/climb processes taking place and large dislocation loops were present. Also the density of ®ne debris was lower and the pinning points along ordinary dislocation segments were less pronounced. This is shown in Fig. 8 for the RD sample deformed to 1% strain. At this stage the twin density was low in both samples with only one twin system activated in those grains that deformed with some contribution of twinning. Fig. 9 shows an example of twinned microstructures observed after 1% strain in the TD
1074
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 6. Microstructure observed at the centre of the gauge in the TD sample deformed to fracture where mechanical twins traverse the grain (a) and produce emission of ordinary dislocations at the intersection with the boundary (b). As a consequence the lower grain deforms by zone (011). ordinary dislocations only. Diraction vector g=1 11,
sample. In the TD sample some slip bands of superdislocations were also seen within grains that deformed by ordinary dislocations, as shown in Fig. 10. In both types of samples, the major dierence observed after the large strains to fracture was a substantial increase in dislocation and twin densities. 3.3. TEM microstructures of samples deformed at 900 C The deformation process at 900 C, as already con®rmed from SEM observations, was characterized by the absence of mechanical twinning in both types of samples. Fig. 11 shows an example of the microstructure seen in the RD sample deformed to fracture (84% strain), in a foil taken from the middle of the gauge length. We note that most of the original grains contain a high density of dislocations and are subdivided into smaller subgrains, while a few (like the one labelled A) appear as new recrystallized grains and with
Fig. 7. Examples of the two types of microstructure observed from the centre of the gauge in the TD sample deformed to fracture. (a) Grain deformed by mechanical twins and ordinary dislocations; Diraction zone (112). (b) Grain deformed by a majority of vector g=1 11, superdislocations; diraction vector g=002, zone (010).
a low density of defects. In Fig. 12 we show one of the grains from the same thin foil sample, imaged under two diraction conditions, con®rming that all the
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1075
Fig. 8. Microstructure observed by TEM from the RD sample deformed to 1% strain at 800 C. Note the large dislocation loops and the lower density of pinning points along the ordinary dislocations seen. Diraction vector g=200, zone (010).
Fig. 10. Example of microstructure observed under two diraction conditions from the TD sample deformed to 1% strain at 800 C. (a) Diraction vector g=111, zone (110); both ordinary and superdislocations are visible. (b) Diraction vector g=002, zone (110); only the bands of superdislocations are visible.
Fig. 9. Example of microstructure from the TD sample deformed to 1% strain at 800 C where twins and ordinary dislocations are present within the two grains observed.
dislocations are of the type 1/2h110i. In Fig. 13(a) we show a similar distribution of grains in the TD sample deformed to fracture (45% strain) while in Fig. 13(b) and (c) we show examples of grains from Fig. 13(a) deformed by a mixture of ordinary and superdislocations. Thus the major dierence observed in the TD sample deformed at 900 C at these large strains was the larger
contribution of superdislocations in a high fraction of grains. Also the material contained a large fraction of recrystallized grains (as opposed to the recovered grains seen in the RD sample) like the examples shown in Fig. 14. 4. Discussion The evolution of the microstructure has been interpreted in terms of the active deformation mechanisms observed in both samples. The RD material has a
1076
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
Fig. 11. Microstructure observed by TEM from the centre of the gauge in the RD sample deformed to fracture at 900 C. Most of the grains contain a high dislocation density and are subdivided into smaller subgrains and only a few, as the one labelled A, appear as new recrystallized grains.
fraction of grains deformed by mechanical twins that increases from about 40 to 80% between 1 and 3.4% strain (measured adjacent to the head of the tensile sample). This fraction does not change signi®cantly with increasing strain but the intensity of twinning within grains increases substantially between 3.4% strain (near the head of the sample) and 67% strain (adjacent to the fracture surface). This leads to a major re®nement of the microstructure which is more pronounced at 700 C, as evidenced in Fig. 1. This eect is less pronounced in the TD sample because the total strain achieved was only 54% and the fraction of grains deformed by twinning changed from 38 to 48% between 1 and 5.3% (measured adjacent to the head of the tensile sample) at 700 C. The same behaviour is observed in both types of samples at 800 C, except that the fraction of grains that deform by twinning is slightly lower (30% in the RD sample and 25% in the TD sample at 1% strain and saturating at 80 and 35% at about 4% strain in the RD and TD samples, respectively). It seems, therefore, that the fraction of grains that deform by twinning is determined in the early stages of deformation while the total strain achieved must be determined by the evolution of the microstructure during the deformation process. The dierent behaviour of the materials both in terms of strength as well as ductility has been attributed to the inhomogeneity of deformation due to the extent of texture in the samples.
Fig. 12. Example of a grain, imaged under two diraction conditions, from the same area of Fig. 11 in which all dislocations are invisible with g=002, con®rming that they all are of the type 1/2h110i.
At 700 and 800 C in the RD sample the majority of grains (80%) deform by twinning after a strain as low as 3.4% and ordinary dislocations are also observed within the grains. The texture of the sample would allow two twinning systems and two slip systems for ordinary dislocations due to their favourable Schmid factor [10]. In the TD sample the number of grains that deform by twinning is lower, but still signi®cant (38% after 1% strain and 48% after 5.3% strain). This can be attributed to the less favourable texture, i.e. grains which are very close to the [001] direction are not able to twin
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
1077
Fig. 14. Typical recrystallized grains observed from the TD sample deformed to fracture at 900 C.
Fig. 13. (a) General microstructure observed by TEM from the centre of the gauge in the TD sample deformed to fracture at 900 C. (b) Example of a grain from (a) deformed by superdislocations; diraction condition g=002, zone (010). (c) Example of a grain from (a) deformed by ordinary and superdislocations; diraction condition g=111, zone (110).
when stressed in tension. Ordinary dislocations are also observed while superdislocations are rare at low strains and presumably their activation is dicult due to their low mobility. Initial deformation therefore occurs inhomogeneously in those grains favourably oriented for twinning and ordinary dislocations. The higher yield strength of this sample should be attributed to the higher resolved shear stress needed to activate twinning in those grains [10]. Once mechanical twinning is active, the local stresses produced at twin/twin and twin/ boundary intersections lead to emission of ordinary
dislocations (see Figs. 6 and 9). The latter can thus be activated by the local stresses even in grains oriented along [001] directions for which these slip systems would be unfavourable according to the macroscopic Schmid factors if they were ideal single crystals factors [10]. At 700 and 800 C in the RD samples deformation is homogeneously distributed as a mixture of ordinary dislocation movement and twinning (80% of grains) while the rest deforms by ordinary dislocations only. The grains that deform by twinning are subdivided by the twin interfaces (see Fig. 1) and the scale of the subdivision is that of the twin spacing which decreases with increasing strain. Across the twin interfaces another subdivision is observed by crystallographic contrast in the SEM as small subgrains that become very ®ne grains at large strains before fracture. These are produced by the accumulation of ordinary dislocations within the deformed substructure as they propagate across the twins. The most important aspect of this structure and evolution is its uniformity which is imposed by the twin
1078
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079
interfaces where dislocations can accumulate as deformation continues throughout. In the TD samples there are about 48 to 35% of the grains (depending on the temperature) that initially deform in the same uniform manner as that of the RD sample, i.e. with twins subdividing the grains and ordinary dislocations being produced at their intersections. As local stresses are built up, other grains start to deform by ordinary and superdislocations. This leads to two scales of the microstructure in which the distribution of ®nal grains is much more inhomogeneous in the TD samples. In the RD samples about 80% of the grains ®nally become about 100 nm in size after subdivision by twins and smaller subgrains, accompanied by recovery/recrystallization processes. Only 20% of the grains in the RD samples become 2±3 mm in size due to the subgrains produced by ordinary dislocations only. In the TD samples, however, only about 48±35% of the grains ®nally become 150 nm in size due to subdivision by twins and subgrains and recovery/recrystallization processes. There are 52±65% of the grains that remain larger than 2±3 mm because in the absence of mechanical twinning subdivision by twin interfaces does not occur. The lack of uniformity in these structures leads to inhomogeneity of the deformation process and results in the lower ductility observed in the TD samples. At the highest temperature (900 C), no twinning is observed in either sample, presumably because activation of ordinary dislocations occurs at a much lower stress. Such dislocations both glide and climb, and therefore the stresses necessary to activate twinning [10] are never reached (see Table 1). In the RD sample (at 900 C) the subdivision of the initial grains continues by a mechanism of subgrain formation leading to new grains by such recovery processes. This leads to a uniform and homogeneous microstructure throughout the material (see Fig. 11). In the TD sample (at 900 C) the subdivision of grains into subgrains is not so uniform due to the dierent deformation mechanisms from grain to grain (by ordinary dislocations or by superdislocations). Possibly the lack of mobility of the superdislocations makes recovery processes more dicult and these grains accumulate dislocations until enough driving force exists for recrystallization [see Figs. 13(a) and 14]. The result is a more inhomogeneous structure with some grains recovered and others recrystallized, i.e. with some grains the prior texture remains while others take totally new orientations. This inhomogeneous structure with dierent types of grain boundaries results in incompatibility of deformation across grains. It is such inhomogeneity that is believed to be responsible for the lower ductility than that observed in the RD sample in which most grains (80%) have recovered and accommodated local stresses in the process.
5. Conclusions . The in¯uence of texture on the strength and ductility achieved during tensile deformation of a rolled sheet of a Ti±Al alloy has been studied by following the microstructural evolution during the deformation process. . SEM and TEM observations have con®rmed that in both types of samples extensive twinning activity occurs during deformation at temperatures of 700 and 800 C but not at 900 C. At the lower temperatures mechanical twinning was already activated at 1% strain. . In those grains that deform by twinning, the twin spacing decreases with decreasing deformation temperature and increasing strain in both samples. Twinning produces a subdivision of the grains during the deformation process and leads to a ®ner grain size in the RD sample before fracture. . The fraction of grains that deform by twinning is much lower in the TD sample which contains 50± 60% of the grains deforming by either superdislocations or by a mixture of ordinary and superdislocations. The lack of uniformity in these structures leads to inhomogeneity of the deformation process and results in the lower ductility observed in the TD samples. . At 900 C twinning does not occur in either sample and the re®nement of the microstructure is lesspronounced. In the RD sample the majority of grains are subdivided by a mechanism of subgrain formation leading to new grains by recovery processes. This subdivision is not so uniform in theTD sample where some grains deform by less mobile superdislocations that hinder recovery processes and accumulate until recrystallization becomes possible.
Acknowledgements The authors would like to thank Professor H. Clemens for providing the tensile samples and the mechanical data.
References [1] Kim Y-W, Dimiduk DM. Journal of Metals 1991;43:40. [2] Huang SC, Hall EL. Metall Trans A 1991;22:2619. [3] Nakano T, Kawanaka T, Yasuda HY, Umakoshi Y. Mater Science Engineering A 1995;194:43. [4] Liu CT, Maziasz PJ, Clemens DR, Schneibel JH, Sikka VK, Nieh TG, Wright J, Walker LR. In Kim Y-W, Wagner R, Yamaguchi M, editors. g Titanium aluminides. Warrendale (PA): TMS, 1995. p. 679. [5] Wang G, Dahms M. Journal of Metals 1993;45:53.
M.A. Morris-MunÄoz, D.G. Morris / Intermetallics 7 (1999) 1069±1079 [6] Watkins D, Piehler HR, Seetharaman V, Lombard CM, Semiatin SL. Metall Trans A 1992;23:2669. [7] Dimiduk DM. In: Kim Y-W, Wagner R, Yamaguchi M, editors. g Titanium aluminides. Warrendale (PA): TMS, 1995. p 3. [8] Huang SC, Shih DS. In: Kim Y-W, Boyer RR, editors. Microstructure/property relationships in titanium aluminides and alloys. Warrendale (PA): TMS, 1990. p 105. [9] Bartels A, Clemens H, Hartig C, Mecking H. In: Koch CC, Liu CT, Stolo NS, Wanner A. High temperature intermetallic alloys VII, Vol. 460. Pittsburgh (PA): MRS, 1997. p 141.
1079
[10] Morris MA, Clemens H, SchloÈg SM. Intermetallics 1998;6:511. [11] Clemens H, Glatz W, Schretter P, Koeppe C, Bartels, A, Behr R, Wanner A. In: Kim Y-W, Wagner R, Yamaguchi M, editors. g titanium aluminides. Warrendale (PA): TMS, 1995. p 717. [12] Clemens H, Glatz W, Eberhard N, Martinz H-P, Knabl W. In: Koch, CC, Liu CT, Stolo NS, Wanner A, editors. High temperature ordered intermetallics VII, vol. 460. Pittsburgh (PA): MRS, 1997. p 29.