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Superplasticity in a large-grained TiAl alloy
Intermetallics 12 (2004) 875–883 www.elsevier.com/locate/intermet
Superplasticity in a large-grained TiAl alloy Dongliang Lin*, Feng Sun School of Materials Science and Engineering, and Open laboratory of Education Ministry of China for High Temperature Materials and Tests, Shanghai Jiao Tong University, Shanghai 200030, China Available online 9 April 2004
Abstract The superplastic behaviour was systematically investigated in a large-grained Ti– 47Al – 2Mn – 2Nb– B alloy having nearly equiaxed g-phase with grain size of 95 mm, in which a small amount of fine particles of a2 distribute uniformly. Superplastic deformation was examined at a temperature range of 1025–1100 8C and strain rates range of 4 £ 1025 – 1.28 £ 1023 s21. The large-grained TiAl alloy exhibits all deformation characteristics of conventionally fine-grained superplastic alloys without the prerequisites, fine grain size and grain boundary sliding. All the values of strain rate sensitivity, m are larger than 0.3. In most cases, an elongation over 200% was gained. A maximum elongation of 287.5% with an m value of 0.39 was obtained at 1100 8C and an initial strain rate of 4 £ 1025 s21. Microstructure evolution during superplastic deformation was characterized by optical microscopy, orientation imaging microscopy and transmission electron microscopy (TEM). Metallographic examination has shown that the average grain size of large-grained TiAl alloy decreased during superplastic deformation, after that a much finer grain size of 10 to 3 –5 mm could be obtained. Electron back-scattered diffraction analysis revealed that significant grain refinement was obtained at different levels with an increase in the density of low and high angle grain boundaries. A direct evidence for dynamic formation of grain boundaries with misorientation of 15 – 308 was found, which was evolved from subboundaries. The evidence of subboundary formation and dislocation glide in the interior of grains was revealed by TEM observation. A continuous recovery and recrystallization process similar to that in FeAl and Fe3Al alloys was proposed as the superplastic deformation mechanism in the large-grained TiAl alloy. q 2004 Elsevier Ltd. All rights reserved. Keywords: B. Superplastic behavior; C. Recrystallization and recovery (including grain growth)
1. Introduction g-TiAl-based alloys are powerful candidates for light-weight high temperature structural materials because of their excellent high temperature strength, low density and good oxidation resistance [1]. Unfortunately, TiAl alloys are difficult for machine and hot working due to ordered structure, which impede large-scale application. Great efforts have been exerted to improve workability. As a promising technology for near-net shaping, superplastic forming has gained extensive research. However, most investigations have been focused on fine-grained superplasticity in intermetallics [2]. Recently, superplasticity was found in Fe3Al and FeAl alloys with coarse grain sizes (. 60– 600 mm) in our laboratory [3,4]. It was found [5,6] that large-grained aluminides exhibit all deformation characteristics of conventionally fine-grained superplastic materials without the prerequisites of fine grain size and grain boundary sliding. A * Corresponding author. Tel.: þ 86-21-62932544; fax: þ 86-2162820892. E-mail address: [email protected] (D. Lin). 0966-9795/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2004.02.039
continuous recovery and recrystallization mechanism was proposed [5,6] for superplasticity in large-grained FeAl and Fe3Al alloys. The model is based on the dynamic formation and evolution of unstable subgrain boundaries, in that appropriate dislocation motion velocity and excess vacancy (or other point defects) concentration are two decisive factors for the formation of subgrain boundaries and their evolution to low and high angle grain boundaries. It can be inferred that large-grained superplasticity should be common in most intermetallics having intrinsically low dislocation mobility. This inference has been partly proved. More recently, superplastic phenomena in large-grained TiAl [7] and NiAl [8] alloys were also found in our laboratory. Other workers also reported similar phenomena in Fe3Si [9] and NiAl alloys [10]. In this paper, the superplastic deformation characteristics in a large-grained TiAl alloy was reported that the emphases were placed in the effects of testing temperature and strain rate on the superplastic deformation behaviour. The microstructure evolution during and after deformation was revealed and the mechanism of superplastic deformation was discussed.
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2. Experimental procedure
3. Results
Starting material was supplied in the form of 18 mm thick forged cake with a nominal composition of Ti –47Al – 2Mn – 2Nb – 1B (at%). Several sample pieces cut from the cake were heat-treated at 1380 8C for 1 h, and then rapidly quenched by liquid nitrogen to gain massive g structure. Finally, the pieces were annealed at 1050 8C for 12 h to obtain the final microstructure of nearly equiaxed g, in which a small amount of very fine particles of a2 distribute uniformly. Tensile specimens with a gauge section of 8 mm £ 4 mm £ 1.3 mm were prepared from the above pieces by electrical discharge machining. The flat surface of all plate specimens was parallel to the forging plane of the cake. The superplastic tensile tests were performed in air by SHIMAZU AG-100KNA test machine equipped with a three-zone resistor furnace. All specimens were coated by enamel to resist oxidation. An initial strain rate range of 2 £ 1025 –1.28 £ 1023 s21 and five temperatures of 1035, 1035, 1050, 1075 and 1100 8C were chosen for the tensile test. Specimens were held at test temperature for 15 min after about 75 min of heating to desired temperature. Optical microscopy (OM) was used to reveal microstructure evolution during and after deformation. Detailed substructure analyses were performed by transmission electron microscopy (TEM). Specimens prepared for TEM analysis were obtained by water quenching after deformed to a predefined strain level. Thin plates cut from gauge section of quenched specimens were ground to , 70 mm, then thinned electrolytically by standard twin-jet method to gain TEM foils. The electrolyte was a 5% perchloric acid solution in methanol, where the temperature and voltage were kept at 2 40 8C and 15 V, respectively. Orientation imaging microscopy (OIM) analyses were performed on specimens deformed to 30, 98 and 170% at 1075 8C and 80% at 1025 8C. Automatic generation and indexing of electron back-scattered diffraction (EBSD) patterns were carried out on an orientation imaging microscope produced by HKL Technology Inc., Denmark, which equipped a back-scattered electron detector and Channel 4 analysis software. Beam scan mode was adapted with a step spacing of 1 mm over an area of 115 £ 115 –161 £ 161 mm2. The image of microstructure was reconstructed by creating grain boundary maps from the EBSD pattern measurements. Designation of grain boundaries was based on a grain boundary criteria, v; given by the researcher. Misorientation angle u is calculated between grid points in the scan field and compared with v: In the paper, three criteria, 28 , v , 108; 108 # v , 158; and v $ 158 were considered. By employing these criteria during the generation of the grain boundary maps, different images of the microstructure were constructed.
3.1. Superplastic deformation characteristics It is shown that the large-grained TiAl alloy possesses a good ability of superplastic deformation. At 1050– 1100 8C, in most cases, an elongation of 200% and over were obtained, as shown in Fig. 1. It is clear that the elongation to fracture ðdf Þ varied with strain rate ð1Þ: _ There is no monotonous but fluctuant relationship between df and 1: _ For example, at 1075 8C, a maximum elongation was gained under a strain rate of 4 £ 1025 s21, while a second maximum was gained under a strain rate of 6.4 £ 1023 s21. The phenomenon is owing to the complex effect of the enamel coating and surface oxidation. The enamel coating was used to protect Ti alloys at hot forging, and may be not suitable for small-stress tensile tests. Examination of specimens after deformation showed that, at low temperature and high strain rate, the enamel coating still keeps thick, while at high temperature and low strain rate, the coating is watery. It is noted that the enamel coating confines to superplastic deformation of TiAl large-grained alloy. In other words, the coating depresses necking resistance of the matrix metal. At high temperature and low strain rate such as at 1075 8C and 4 £ 1025 s21, the enamel could have a proper viscosity to maintain flow synchronously with superplastic deformation, while at low temperature or high strain rate, microcracks may occur at coating layer due to incompatibility between the enamel and the TiAl alloy, then surface oxidation will be accelerated to cause premature failure of tensile specimens. The profiles of specimens fracture at different temperatures and under a fixed strain rate of 4 £ 1025 s21 were illustrated in Fig. 2. It will be worthwhile to emphasise here that the gauge section of the specimen deformed at 1100 8C still keeps uniform even at an elongation of 287.5%. The fact powerfully proved the potential of superplastic deformation for the large-grained TiAl alloy. The strain rate sensitivity index, m; was measured by a method [11] based on the data analyses from tensile tests at different strain rates. A true strain level of 20% was selected
Fig. 1. Effects of testing temperature and initial strain rate on elongation of the TiAl alloy.
D. Lin, F. Sun / Intermetallics 12 (2004) 875–883
Fig. 2. Specimens strained to failure at 1075 and 1100 8C under an initial strain rate of 4 £ 1025 s21.
Fig. 3. Stress–strain rate relationships for the large-grained TiAl alloy.
as checking point. In Fig. 3, the flow stress at 1 ¼ 0:2; were plotted against the strain rate at a temperature range of 1025 –1100 8C. The slopes at any point on curves can be calculated as the m value for that point. It is seen from Fig. 3 that the m values at temperatures of 1050 –1100 8C and under the strain rate of 1.6 £ 1024 s21 are higher than 0.3, which is a critical value for conventionally superplastic deformation. A maximum m value of 0.39 was gained at 1100 8C, which corresponds to an elongation of 287.5%. At a fixed temperature, the m value decreases with increasing strain rate. However, the m value increases with increasing temperature at low strain rate.
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The effects of testing temperature and strain rate on the flow behaviour are illustrated in Fig. 4a and b, respectively. In general, the flow stress increases with increasing strain until a peak stress is reached and then decreases gradually. As shown in Fig. 4, the peak flow stress increases with increasing strain rate, while vice versa for testing temperature. At a strain rate of 4 £ 1025 s21, softening was observed at 1075 8C and below, and strain hardening was observed at 1100 8C. At a fixed temperature of 1100 8C, the alloy showed hardening at low strain rates and softening at relatively high strain rates. Even under higher strain rates, up to 3.2 £ 1024 s21, a narrow plateau is exhibited, which means a balance between strain hardening and softening was reached. Softening is usually owing to recrystallization or geometrical softening caused by necking, and in that case the presumption of uniform deformation is not available. However, stress relaxation caused by the gradual decrease of the effective strain rate is possibly more important factor to induce softening. Since all the tests were conducted under the condition of constant crosshead speed, the curves of flow stress versus strain may be somewhat different from those under the condition of constant strain rate. Based on the stress versus strain relation in superplasticity, s ¼ k1; _ where k is a constant, a correlation to transform the stress at non-constant strain rate to that at constant strain rate was developed as
scs ¼ exp½mð1 2 1yp Þscv
ð1Þ
where scs is the equivalent stress under the condition of constant strain rate, 1yp is true strain corresponding to yielding point, scv is the stress under the condition of constant cross head speed. According to the above correlation, the curves of flow stress versus strain were replotted. As displayed in Fig. 5, it is obvious that a common trend of strain hardening exists in the investigated range when 1 , 0:4: Furthermore, the stage of strain hardening is broadened with increasing temperature or decreasing strain rate, accompanied by a weakening of strain hardening trend, and then quasi steady state will be gained. It is reasonable to think that there is a correspondent relationship between the length of hardening stage and the elongation to fracture,
Fig. 4. Effects of testing temperature and initial strain rate on the flow behaviour of the TiAl alloy: (a) initial strain rate; (b) temperature.
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Fig. 5. Calculated stress–strain curves corresponding to Fig. 4.
in which df will increase with increasing temperature or decreasing strain rate disregarding surface oxidation and in material homogeneity. Because of the similarity between superplastic deformation and creep, the usual power law creep relation was used to describe the steady-state flow deformation in superplsticity, 1_ ¼ Asn expð2Q=RTÞ
ð2Þ
where 1_ is the steady-state strain rate, n is the stress exponent ðn ¼ 1=mÞ; Q is the activation energy for the process, and R and T keep their usual meaning. The apparent activation energy for superplastic deformation can be derived from Eq. (2) as below Q ¼ n ›ln s=›ð1=TÞ
ð3Þ 25
Fig. 6 presents the data of 1025 –1100 8C and 4 £ 10 , 8 £ 1025 and 1.6 £ 1024 s21, in which the slopes of fitting lines correspond to the partial differential item in Eq. (3) wherein the Q value can be easily calculated. The average Q value was determined to be about 331.4 kJ/mol, which is quite close to the activation energy for self-diffusion of Ti in single phase Y-TiAl which is 295 kJ/mol [12], and also close to the activation energy for inter-diffusion of Ti and Al, namely 295 kJ/mol [13]. It is suggested that the rate process in superplastic deformation be controlled by lattice diffusion.
of grain size is still obvious (see Fig. 7c and d). No cavity was found in the microstructure after deformation, even at the failed end of specimen. Fig. 8 exhibits EBSD orientation mapping images of deformed specimens based on the grain boundary criteria, in which (a) – (c) correspond to the strain of 39, 98 and 170% at 1075 8C, respectively, and (d) to that of 80% at 1025 8C. In these maps, red lines represent grain boundary with a misorientation angle u higher than 158, which should be highangle grain boundary (HAGB). Lime green lines represent grain boundaries with misorientation angle u of 10 –158, whereas black lines for u between 2 and 108, which should be low-angle grain boundaries (LAGBs). Grain refinement and the increase in the amount of LAGB during deformation are clearly shown. At the strain of 30%, a few LAGBs were appeared in the microstructure, and they distributed mainly in the adjacent areas of HAGB. When strain increased up to 98%, more LAGBs appeared in microstructure. There are two different distribution morphologies for them. One is that LAGBs were distributed uniformly in some areas in grains, and the other is that they were distributed near HAGBs. Similar feature was observed in 170% deformed specimens (Fig. 8c) and 80% deformed specimens at 1025 8C (Fig. 8d). Moreover, the amount of HAGBs and small areas enclosed by HAGBs increased significantly with the increase of strain.
3.2. Microstructure evolution Fig. 7 shows the microstructure of the TiAl alloy before and after deformation. As shown in Fig. 7a, the microstructure before tensile test have nearly equiaxed g grains with a mean size of about 95 mm. Compared to the generally accepted fine grain size of no more than 10 mm for superplasticity, this grain size belongs to a large grain size category. After deformation, it is evident that the average grain size decreases. At low temperature of 1025 8C, the average grain size was decreased to about 10 mm, and finer grains about 3 – 5 mm can easily be observed, as shown in Fig. 7b. At 1050 8C or higher temperature, although the thermal-activated grain growth is enhanced, the refinement
Fig. 6. Linear relationships between ln s and lnð1=TÞ:
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Fig. 7. Optical micrographs of the TiAl-based alloy: (a) before deformation; (b) after deformation, at 1025 8C, 1_ ¼ 1:6 £ 1024 s21, df ¼ 188:4%; (c) and (d) after deformation, at 1075 8C, 1_ ¼ 1:6 £ 1024 s21, df ¼ 257:6%; at the area of uniform deformation and near the failed end, respectively.
It should be pointed out that not all HAGBs depicted are ordinary grain boundaries, in which some are twin boundaries, which can be led to explain why the depicted grain size is finer than that observed by OM. When comparing the microstructure of deformed specimens at different temperatures, it is clearly shown that the grain refinement is more significant at lower temperature, in which a high density of LAGBs and HAGBs can be observed. Considered the fact that the superplsticity at 1025 8C is more inferior than that at 1075 8C, which led to the suggestion that a slow reduction of grain size is needed to maintain superplastic flow to proceed. Fig. 9 shows the distribution of misorientation angles in EBSD orientation mapping images of deformed specimens in Fig. 8, which is expressed as relative frequency in total points. The correlated plot displays the misorientation data between neighbouring points in an orientation map, whereas an uncorrelated plot shows the misorientation between randomly chosen points in the data set. In the correlated data, a minimum angle of 18 was chosen for statistics. In addition, the data at u ¼ 85 – 908 have been filtered, which are abnormally high and have no large variations in different specimens. It is believed to be caused by anti-phase domain boundaries. All maps show a high fraction of misorientation angle at u ¼ 1 – 58; compared with a minimum angle of 28 (as seen in Fig. 9a), where it can be found that the high fraction is mainly contributed by u ¼ 1 – 28: This part is difficult to carry out quantitative analysis due to the resolution limit of 28. At the deformation of 39%, the correlated distribution shows lack of grain boundaries with misorientation of 15 –308, which can be seen in the uncorrelated distribution.
Moreover, the fraction of LAGBs with u ¼ 5 – 108 is very low, too, as seen in the orientation mappings of Fig. 8. It is indicated that the initial microstructure only has HAGBs with misorientation angles above 308 and strain of 30% is not enough to form uniform substructure. Probably, most immobile dislocations still exist in the form of dislocation walls and networks, and LAGBs with u , 58 at such low deformation level. It is consistent with the results mentioned in Section 3.1 of this paper that apparent strain softening occurs at true strain of , 0.4. With increasing strain, the fraction of LAGBs with u ¼ 5 – 108 increased, and more prominently, the grain boundaries with u ¼ 15 – 308 appeared, which also increased with the strain. The fraction of grain boundaries with u ¼ 10 – 158 increased when strain increased to 98%, but changed very little between the strain of 98 and 170%, or it could be considered to remain constant. This phenomenon was resulted only by the evolution of subboundaries. During deformation, LAGBs with u , 108 transformed continuously to those with u . 108; so the fraction of grain boundaries with u . 108 increased, while the grain boundaries with u ¼ 10 – 158 are extremely unstable, having a strong tendency to transform into HAGBs. At large strain, a dynamical balance will be reached between the evolution of LAGBs from u , 108 towards u . 108; and then the fraction of LAGBs with u ¼ 10 – 158 will be kept constant. Misorientation distribution at a high angle range of above 308 also shows a change during deformation, among which the fraction of HAGBs around u ¼ 608; probably being twin boundaries, was highest. As deformation increased, the fraction of 608 HAGB decreased, while other HAGBs
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Fig. 8. Orientation mapping images of deformed specimens under different strains at an initial strain rate of 1.6 £ 1024 s21: (a) 30, (b) 98, (c) 170 and (d) 80%; (a) –(c) deformed at 1075 8C and (d) at 1025 8C. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
especially those with u . 758; a maximum peak appeared at u ¼ 80 – 858: Such a change corresponded to the formation of new grains and the wipe-off of original twin boundaries. At lower temperature such as 1025 8C, a higher density of LAGBs was observed in the specimen deformed to 80%, while the distribution of HAGBs is similar to that deformed to 170% at 1075 8C. It indicates that at lower temperature dislocations have stronger tendency to form LAGBs and to transform to HAGBs. TEM images of substructure in the deformed samples are shown in Fig. 10. The results of TEM observation on the dislocation configurations in Y grains showed that there exist a great number of subgrain boundaries, which divided the initial large grains into fine subgrains during deformation. The common feature of these boundaries is that they are composed of dislocation walls and dislocation networks. Fig. 11 shows a forming subgrain boundary. Two dislocations labelled as A intersect and react with each other. It is interesting to view that the dislocation at label A are originated from the dislocation in front of the original
grain boundary labelled as B. It seems that dislocation pile-ups are favourable to form subboundary. Moreover, long dislocation segments and deformation twins across subboundaries were observed too in g grains in Fig. 10. It is noticed that subgrain in Fig. 10c were elongated significantly, which are typical characteristics of dislocation gliding and also a signal of incompatible deformation. In contrast with the high density in subgrain boundaries, the dislocation density is low in the interior of grains. It suggests that at the steady-state (or quasi steady-state) flow stage, the dislocation in grain will remain in relatively low density due to the formation of grain boundaries. In Fig. 12, a series of TEM micrographs were used to identify dislocation character in the subgrain boundary labelled as C in Fig. 11. As seen from Fig. 12a, at operating vector g ¼ 220; the subboundary shows a dislocation net, which is composed by two sets of intersecting dislocation and g ¼ 11 1; one set of dislocation arrays. When g ¼ 111 arrays is invisible, alternately. At g ¼ 020; two arrays become all visible. Using the g·b ¼ 0 criteria for
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Fig. 9. Distribution of misorientation angles in deformed microstructure presented in Fig. 1: (a) 30, (b) 98 and (c) 170% at 1050 8C, and (d) 80% at 1025 8C.
the invisibility of dislocations, where b is Burgers vector, it is easy to deduce that the subgrain boundary is composed by type superdislocations. There is no visible [101] and ½101 dissociation of superdislocations to be observed. Apart from type dislocations, other type dislothe [101] and ½101 cations are rare to find. The interaction between dislocations and subgrain boundaries and the evolution of boundaries are still inexplicit, which necessitates further research to be carried out.
4. Discussions The flow behaviour in superplastic deformation of the large-grained TiAl alloy is similar to those in conventionally superplastic alloys. However, there exist some significant differences in microstructural evolution
and mechanism between the large-grained TiAl alloy and conventionally fine-grained superplastic materials. It was shown in this paper that superplastic behaviour was found in large-grained TiAl alloy without the usual prerequisites of fine grain size and grain boundary sliding [14]. Moreover, metallographic examination has shown that the average grain size of the large-grained TiAl alloy decreased during superplastic deformation and a much finer grain size could be obtained after superplastic deformation, while in finegrained superplastic alloys the grain size keeps nearly constant or has a little growth. The phenomenon of grain refinement is confirmed further by OIM analyses. In the present study, the microstructure evolution during superplastic deformation underwent the same process in both Fe3Al and FeAl alloys [5,6]. The initial large grains in large-grained TiAl alloys were refined during deformation. Most of the refined grains are bordered by subgrain
Fig. 10. Typical substructures seen after deformation of the TiAl alloy at 1075 8C and 1.6 £ 1024 s21: (a) deformed to 60%; (b) deformed to 98%; (c) deformed to 170%.
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Fig. 11. A forming subboundary and dislocation pile-ups in front of an original grain boundary in the sample deformed to 60% at 1075 8C under 1.6 £ 1024 s21.
boundaries or low angle grain boundaries. The density of subgrain boundaries and LAGBs, as well as HAGBs increased with the strain. Rotation between subgrains occurred during deformation as seen from the standard orientation triangles. In our previous work [5,6], we have proposed a continuous recovery and recrystallization model to account for the superplastic behaviour of initially large-grained Fe 3Al and FeAl alloys. Microstructure evolution observed in this study obviously indicates that a continuous recovery and recrystallization process appeared during superplastic deformation in large-grained TiAl alloy.
At the early stage of deformation, dislocation sliding is dominant. With further deformation, climb process occurred and the dislocations arrange themselves as subgrain boundaries. The subgrain boundaries divide the initial large grains into fine subgrains. The subgrain boundaries absorb the dislocations coming from the subgrains and thus the misorientation between subgrains increases. The conclusion that the subgrain boundaries can absorb the dislocations was based on not only the analysis on the dislocations in the grain boundaries but also the fact that the density of dislocations in the grains or subgrains is very low. Based on our previous
1; B ¼ ½110; Fig. 12. Diffraction contrast analysis of the dislocations in subgrain boundary: (a) g ¼ 220; B ¼ ½110; (b) g ¼ 111; B ¼ ½110; (c) g ¼ 11 (d) g ¼ 020; B ¼ ½101:
D. Lin, F. Sun / Intermetallics 12 (2004) 875–883
TEM observations, there exists high density of subgrain boundaries with the misorientation lower than 28, which cannot be detected by present EBSD technique. If we take the misorientation between subgrains lower than 28 into account, the density of subgrain boundaries and low angle grain boundaries in the samples would be even higher than that detected by EBSD. Based on our EBSD and TEM study, we believe that the superplastic deformation of both Fe3Al and FeAl alloys as well as TiAl alloy is controlled by dislocation gliding and climbing. The gliding dislocations interacted and arranged themselves as subgrain boundaries through their climb process. Superplastic deformation needs not only dislocation sliding but also the movement between subgrains, such as subgrain boundary sliding and migration, which should be controlled by dislocation climbing in the subgrain boundaries. The movement between subgrains results in the migration of the initial high angle grain boundaries of the alloys. During deformation, gliding dislocations can be trapped and absorbed by subboundaries, and then subboundaries evolve continuously towards LAGBs, in further, HAGBs. A constant dislocation density in the interior of grains is maintained under a constant stress at given temperature, which corresponds to the steady-state flow stage of superplastic deformation. The experimental results support our previous model of continuous recovery and recrystallization for the mechanism of superplastic deformation in large-grained iron aluminides. A dislocation glide and climb process accommodated by subgrain boundary sliding, migration and rotation allows the superplastic flow to proceed. However, there is more complexity in large-grained TiAl alloy due to the existence of twin boundaries and anti-phase domain boundaries. The superplastic deformation can be accommodated by deformation twinning, which is induced by elastic incompatibility. Moreover, deformation twins can act as dislocation sources or as effective obstacles to help in the formation of subgrain boundaries. In this consideration, deformation twinning may insert a beneficial effect in superplasticity in large-grained TiAl alloy while it is harmful in fine-grained TiAl alloy [15].
5. Conclusions The large-grained TiAl alloy exhibits superplasticity in a temperature range of 1025 –1100 8C and a strain rate range of 4 £ 1025 – 1.28 £ 1025 s21. A maximum elongation of
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287.5% with an m value of 0.39 was obtained at 1100 8C and 1.28 £ 1025 s21. A significant refinement of average grain size with an increase in the density of LAGBs and HAGBs during superplastic deformation was found, which is beneficial to improve mechanical properties of TiAl alloy. Evolution of grain boundaries including LAGBs and HAGBs was depicted by distribution of misorientation angles, a direct evidence of dynamical formation of grain boundaries with misorientation of 15– 308 by OIM analyses, which is a result of subboundary evolution. TEM observations have shown the formation of subboundaries and also the evidence of dislocation gliding in the interior of grains. The experimental results support our previous model of continuous recovery and recrystallization for the mechanism of superplastic deformation in large-grained iron aluminides. A dislocation glide and climb process accommodated by subgrain boundary sliding, migration and rotation allows the superplastic flow to proceed.
Acknowledgements This work was supported by the National Natural Science Foundation of China. The authors would like to thank Professor Q. Liu and Mr J. Meng of Tsinghua University for kind help on EBSD experiments.
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Superplasticity in a large-grained TiAl alloy Dongliang Lin*, Feng Sun School of Materials Science and Engineering, and Open laboratory of Education Ministry of China for High Temperature Materials and Tests, Shanghai Jiao Tong University, Shanghai 200030, China Available online 9 April 2004
Abstract The superplastic behaviour was systematically investigated in a large-grained Ti– 47Al – 2Mn – 2Nb– B alloy having nearly equiaxed g-phase with grain size of 95 mm, in which a small amount of fine particles of a2 distribute uniformly. Superplastic deformation was examined at a temperature range of 1025–1100 8C and strain rates range of 4 £ 1025 – 1.28 £ 1023 s21. The large-grained TiAl alloy exhibits all deformation characteristics of conventionally fine-grained superplastic alloys without the prerequisites, fine grain size and grain boundary sliding. All the values of strain rate sensitivity, m are larger than 0.3. In most cases, an elongation over 200% was gained. A maximum elongation of 287.5% with an m value of 0.39 was obtained at 1100 8C and an initial strain rate of 4 £ 1025 s21. Microstructure evolution during superplastic deformation was characterized by optical microscopy, orientation imaging microscopy and transmission electron microscopy (TEM). Metallographic examination has shown that the average grain size of large-grained TiAl alloy decreased during superplastic deformation, after that a much finer grain size of 10 to 3 –5 mm could be obtained. Electron back-scattered diffraction analysis revealed that significant grain refinement was obtained at different levels with an increase in the density of low and high angle grain boundaries. A direct evidence for dynamic formation of grain boundaries with misorientation of 15 – 308 was found, which was evolved from subboundaries. The evidence of subboundary formation and dislocation glide in the interior of grains was revealed by TEM observation. A continuous recovery and recrystallization process similar to that in FeAl and Fe3Al alloys was proposed as the superplastic deformation mechanism in the large-grained TiAl alloy. q 2004 Elsevier Ltd. All rights reserved. Keywords: B. Superplastic behavior; C. Recrystallization and recovery (including grain growth)
1. Introduction g-TiAl-based alloys are powerful candidates for light-weight high temperature structural materials because of their excellent high temperature strength, low density and good oxidation resistance [1]. Unfortunately, TiAl alloys are difficult for machine and hot working due to ordered structure, which impede large-scale application. Great efforts have been exerted to improve workability. As a promising technology for near-net shaping, superplastic forming has gained extensive research. However, most investigations have been focused on fine-grained superplasticity in intermetallics [2]. Recently, superplasticity was found in Fe3Al and FeAl alloys with coarse grain sizes (. 60– 600 mm) in our laboratory [3,4]. It was found [5,6] that large-grained aluminides exhibit all deformation characteristics of conventionally fine-grained superplastic materials without the prerequisites of fine grain size and grain boundary sliding. A * Corresponding author. Tel.: þ 86-21-62932544; fax: þ 86-2162820892. E-mail address: [email protected] (D. Lin). 0966-9795/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2004.02.039
continuous recovery and recrystallization mechanism was proposed [5,6] for superplasticity in large-grained FeAl and Fe3Al alloys. The model is based on the dynamic formation and evolution of unstable subgrain boundaries, in that appropriate dislocation motion velocity and excess vacancy (or other point defects) concentration are two decisive factors for the formation of subgrain boundaries and their evolution to low and high angle grain boundaries. It can be inferred that large-grained superplasticity should be common in most intermetallics having intrinsically low dislocation mobility. This inference has been partly proved. More recently, superplastic phenomena in large-grained TiAl [7] and NiAl [8] alloys were also found in our laboratory. Other workers also reported similar phenomena in Fe3Si [9] and NiAl alloys [10]. In this paper, the superplastic deformation characteristics in a large-grained TiAl alloy was reported that the emphases were placed in the effects of testing temperature and strain rate on the superplastic deformation behaviour. The microstructure evolution during and after deformation was revealed and the mechanism of superplastic deformation was discussed.
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2. Experimental procedure
3. Results
Starting material was supplied in the form of 18 mm thick forged cake with a nominal composition of Ti –47Al – 2Mn – 2Nb – 1B (at%). Several sample pieces cut from the cake were heat-treated at 1380 8C for 1 h, and then rapidly quenched by liquid nitrogen to gain massive g structure. Finally, the pieces were annealed at 1050 8C for 12 h to obtain the final microstructure of nearly equiaxed g, in which a small amount of very fine particles of a2 distribute uniformly. Tensile specimens with a gauge section of 8 mm £ 4 mm £ 1.3 mm were prepared from the above pieces by electrical discharge machining. The flat surface of all plate specimens was parallel to the forging plane of the cake. The superplastic tensile tests were performed in air by SHIMAZU AG-100KNA test machine equipped with a three-zone resistor furnace. All specimens were coated by enamel to resist oxidation. An initial strain rate range of 2 £ 1025 –1.28 £ 1023 s21 and five temperatures of 1035, 1035, 1050, 1075 and 1100 8C were chosen for the tensile test. Specimens were held at test temperature for 15 min after about 75 min of heating to desired temperature. Optical microscopy (OM) was used to reveal microstructure evolution during and after deformation. Detailed substructure analyses were performed by transmission electron microscopy (TEM). Specimens prepared for TEM analysis were obtained by water quenching after deformed to a predefined strain level. Thin plates cut from gauge section of quenched specimens were ground to , 70 mm, then thinned electrolytically by standard twin-jet method to gain TEM foils. The electrolyte was a 5% perchloric acid solution in methanol, where the temperature and voltage were kept at 2 40 8C and 15 V, respectively. Orientation imaging microscopy (OIM) analyses were performed on specimens deformed to 30, 98 and 170% at 1075 8C and 80% at 1025 8C. Automatic generation and indexing of electron back-scattered diffraction (EBSD) patterns were carried out on an orientation imaging microscope produced by HKL Technology Inc., Denmark, which equipped a back-scattered electron detector and Channel 4 analysis software. Beam scan mode was adapted with a step spacing of 1 mm over an area of 115 £ 115 –161 £ 161 mm2. The image of microstructure was reconstructed by creating grain boundary maps from the EBSD pattern measurements. Designation of grain boundaries was based on a grain boundary criteria, v; given by the researcher. Misorientation angle u is calculated between grid points in the scan field and compared with v: In the paper, three criteria, 28 , v , 108; 108 # v , 158; and v $ 158 were considered. By employing these criteria during the generation of the grain boundary maps, different images of the microstructure were constructed.
3.1. Superplastic deformation characteristics It is shown that the large-grained TiAl alloy possesses a good ability of superplastic deformation. At 1050– 1100 8C, in most cases, an elongation of 200% and over were obtained, as shown in Fig. 1. It is clear that the elongation to fracture ðdf Þ varied with strain rate ð1Þ: _ There is no monotonous but fluctuant relationship between df and 1: _ For example, at 1075 8C, a maximum elongation was gained under a strain rate of 4 £ 1025 s21, while a second maximum was gained under a strain rate of 6.4 £ 1023 s21. The phenomenon is owing to the complex effect of the enamel coating and surface oxidation. The enamel coating was used to protect Ti alloys at hot forging, and may be not suitable for small-stress tensile tests. Examination of specimens after deformation showed that, at low temperature and high strain rate, the enamel coating still keeps thick, while at high temperature and low strain rate, the coating is watery. It is noted that the enamel coating confines to superplastic deformation of TiAl large-grained alloy. In other words, the coating depresses necking resistance of the matrix metal. At high temperature and low strain rate such as at 1075 8C and 4 £ 1025 s21, the enamel could have a proper viscosity to maintain flow synchronously with superplastic deformation, while at low temperature or high strain rate, microcracks may occur at coating layer due to incompatibility between the enamel and the TiAl alloy, then surface oxidation will be accelerated to cause premature failure of tensile specimens. The profiles of specimens fracture at different temperatures and under a fixed strain rate of 4 £ 1025 s21 were illustrated in Fig. 2. It will be worthwhile to emphasise here that the gauge section of the specimen deformed at 1100 8C still keeps uniform even at an elongation of 287.5%. The fact powerfully proved the potential of superplastic deformation for the large-grained TiAl alloy. The strain rate sensitivity index, m; was measured by a method [11] based on the data analyses from tensile tests at different strain rates. A true strain level of 20% was selected
Fig. 1. Effects of testing temperature and initial strain rate on elongation of the TiAl alloy.
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Fig. 2. Specimens strained to failure at 1075 and 1100 8C under an initial strain rate of 4 £ 1025 s21.
Fig. 3. Stress–strain rate relationships for the large-grained TiAl alloy.
as checking point. In Fig. 3, the flow stress at 1 ¼ 0:2; were plotted against the strain rate at a temperature range of 1025 –1100 8C. The slopes at any point on curves can be calculated as the m value for that point. It is seen from Fig. 3 that the m values at temperatures of 1050 –1100 8C and under the strain rate of 1.6 £ 1024 s21 are higher than 0.3, which is a critical value for conventionally superplastic deformation. A maximum m value of 0.39 was gained at 1100 8C, which corresponds to an elongation of 287.5%. At a fixed temperature, the m value decreases with increasing strain rate. However, the m value increases with increasing temperature at low strain rate.
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The effects of testing temperature and strain rate on the flow behaviour are illustrated in Fig. 4a and b, respectively. In general, the flow stress increases with increasing strain until a peak stress is reached and then decreases gradually. As shown in Fig. 4, the peak flow stress increases with increasing strain rate, while vice versa for testing temperature. At a strain rate of 4 £ 1025 s21, softening was observed at 1075 8C and below, and strain hardening was observed at 1100 8C. At a fixed temperature of 1100 8C, the alloy showed hardening at low strain rates and softening at relatively high strain rates. Even under higher strain rates, up to 3.2 £ 1024 s21, a narrow plateau is exhibited, which means a balance between strain hardening and softening was reached. Softening is usually owing to recrystallization or geometrical softening caused by necking, and in that case the presumption of uniform deformation is not available. However, stress relaxation caused by the gradual decrease of the effective strain rate is possibly more important factor to induce softening. Since all the tests were conducted under the condition of constant crosshead speed, the curves of flow stress versus strain may be somewhat different from those under the condition of constant strain rate. Based on the stress versus strain relation in superplasticity, s ¼ k1; _ where k is a constant, a correlation to transform the stress at non-constant strain rate to that at constant strain rate was developed as
scs ¼ exp½mð1 2 1yp Þscv
ð1Þ
where scs is the equivalent stress under the condition of constant strain rate, 1yp is true strain corresponding to yielding point, scv is the stress under the condition of constant cross head speed. According to the above correlation, the curves of flow stress versus strain were replotted. As displayed in Fig. 5, it is obvious that a common trend of strain hardening exists in the investigated range when 1 , 0:4: Furthermore, the stage of strain hardening is broadened with increasing temperature or decreasing strain rate, accompanied by a weakening of strain hardening trend, and then quasi steady state will be gained. It is reasonable to think that there is a correspondent relationship between the length of hardening stage and the elongation to fracture,
Fig. 4. Effects of testing temperature and initial strain rate on the flow behaviour of the TiAl alloy: (a) initial strain rate; (b) temperature.
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Fig. 5. Calculated stress–strain curves corresponding to Fig. 4.
in which df will increase with increasing temperature or decreasing strain rate disregarding surface oxidation and in material homogeneity. Because of the similarity between superplastic deformation and creep, the usual power law creep relation was used to describe the steady-state flow deformation in superplsticity, 1_ ¼ Asn expð2Q=RTÞ
ð2Þ
where 1_ is the steady-state strain rate, n is the stress exponent ðn ¼ 1=mÞ; Q is the activation energy for the process, and R and T keep their usual meaning. The apparent activation energy for superplastic deformation can be derived from Eq. (2) as below Q ¼ n ›ln s=›ð1=TÞ
ð3Þ 25
Fig. 6 presents the data of 1025 –1100 8C and 4 £ 10 , 8 £ 1025 and 1.6 £ 1024 s21, in which the slopes of fitting lines correspond to the partial differential item in Eq. (3) wherein the Q value can be easily calculated. The average Q value was determined to be about 331.4 kJ/mol, which is quite close to the activation energy for self-diffusion of Ti in single phase Y-TiAl which is 295 kJ/mol [12], and also close to the activation energy for inter-diffusion of Ti and Al, namely 295 kJ/mol [13]. It is suggested that the rate process in superplastic deformation be controlled by lattice diffusion.
of grain size is still obvious (see Fig. 7c and d). No cavity was found in the microstructure after deformation, even at the failed end of specimen. Fig. 8 exhibits EBSD orientation mapping images of deformed specimens based on the grain boundary criteria, in which (a) – (c) correspond to the strain of 39, 98 and 170% at 1075 8C, respectively, and (d) to that of 80% at 1025 8C. In these maps, red lines represent grain boundary with a misorientation angle u higher than 158, which should be highangle grain boundary (HAGB). Lime green lines represent grain boundaries with misorientation angle u of 10 –158, whereas black lines for u between 2 and 108, which should be low-angle grain boundaries (LAGBs). Grain refinement and the increase in the amount of LAGB during deformation are clearly shown. At the strain of 30%, a few LAGBs were appeared in the microstructure, and they distributed mainly in the adjacent areas of HAGB. When strain increased up to 98%, more LAGBs appeared in microstructure. There are two different distribution morphologies for them. One is that LAGBs were distributed uniformly in some areas in grains, and the other is that they were distributed near HAGBs. Similar feature was observed in 170% deformed specimens (Fig. 8c) and 80% deformed specimens at 1025 8C (Fig. 8d). Moreover, the amount of HAGBs and small areas enclosed by HAGBs increased significantly with the increase of strain.
3.2. Microstructure evolution Fig. 7 shows the microstructure of the TiAl alloy before and after deformation. As shown in Fig. 7a, the microstructure before tensile test have nearly equiaxed g grains with a mean size of about 95 mm. Compared to the generally accepted fine grain size of no more than 10 mm for superplasticity, this grain size belongs to a large grain size category. After deformation, it is evident that the average grain size decreases. At low temperature of 1025 8C, the average grain size was decreased to about 10 mm, and finer grains about 3 – 5 mm can easily be observed, as shown in Fig. 7b. At 1050 8C or higher temperature, although the thermal-activated grain growth is enhanced, the refinement
Fig. 6. Linear relationships between ln s and lnð1=TÞ:
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Fig. 7. Optical micrographs of the TiAl-based alloy: (a) before deformation; (b) after deformation, at 1025 8C, 1_ ¼ 1:6 £ 1024 s21, df ¼ 188:4%; (c) and (d) after deformation, at 1075 8C, 1_ ¼ 1:6 £ 1024 s21, df ¼ 257:6%; at the area of uniform deformation and near the failed end, respectively.
It should be pointed out that not all HAGBs depicted are ordinary grain boundaries, in which some are twin boundaries, which can be led to explain why the depicted grain size is finer than that observed by OM. When comparing the microstructure of deformed specimens at different temperatures, it is clearly shown that the grain refinement is more significant at lower temperature, in which a high density of LAGBs and HAGBs can be observed. Considered the fact that the superplsticity at 1025 8C is more inferior than that at 1075 8C, which led to the suggestion that a slow reduction of grain size is needed to maintain superplastic flow to proceed. Fig. 9 shows the distribution of misorientation angles in EBSD orientation mapping images of deformed specimens in Fig. 8, which is expressed as relative frequency in total points. The correlated plot displays the misorientation data between neighbouring points in an orientation map, whereas an uncorrelated plot shows the misorientation between randomly chosen points in the data set. In the correlated data, a minimum angle of 18 was chosen for statistics. In addition, the data at u ¼ 85 – 908 have been filtered, which are abnormally high and have no large variations in different specimens. It is believed to be caused by anti-phase domain boundaries. All maps show a high fraction of misorientation angle at u ¼ 1 – 58; compared with a minimum angle of 28 (as seen in Fig. 9a), where it can be found that the high fraction is mainly contributed by u ¼ 1 – 28: This part is difficult to carry out quantitative analysis due to the resolution limit of 28. At the deformation of 39%, the correlated distribution shows lack of grain boundaries with misorientation of 15 –308, which can be seen in the uncorrelated distribution.
Moreover, the fraction of LAGBs with u ¼ 5 – 108 is very low, too, as seen in the orientation mappings of Fig. 8. It is indicated that the initial microstructure only has HAGBs with misorientation angles above 308 and strain of 30% is not enough to form uniform substructure. Probably, most immobile dislocations still exist in the form of dislocation walls and networks, and LAGBs with u , 58 at such low deformation level. It is consistent with the results mentioned in Section 3.1 of this paper that apparent strain softening occurs at true strain of , 0.4. With increasing strain, the fraction of LAGBs with u ¼ 5 – 108 increased, and more prominently, the grain boundaries with u ¼ 15 – 308 appeared, which also increased with the strain. The fraction of grain boundaries with u ¼ 10 – 158 increased when strain increased to 98%, but changed very little between the strain of 98 and 170%, or it could be considered to remain constant. This phenomenon was resulted only by the evolution of subboundaries. During deformation, LAGBs with u , 108 transformed continuously to those with u . 108; so the fraction of grain boundaries with u . 108 increased, while the grain boundaries with u ¼ 10 – 158 are extremely unstable, having a strong tendency to transform into HAGBs. At large strain, a dynamical balance will be reached between the evolution of LAGBs from u , 108 towards u . 108; and then the fraction of LAGBs with u ¼ 10 – 158 will be kept constant. Misorientation distribution at a high angle range of above 308 also shows a change during deformation, among which the fraction of HAGBs around u ¼ 608; probably being twin boundaries, was highest. As deformation increased, the fraction of 608 HAGB decreased, while other HAGBs
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Fig. 8. Orientation mapping images of deformed specimens under different strains at an initial strain rate of 1.6 £ 1024 s21: (a) 30, (b) 98, (c) 170 and (d) 80%; (a) –(c) deformed at 1075 8C and (d) at 1025 8C. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
especially those with u . 758; a maximum peak appeared at u ¼ 80 – 858: Such a change corresponded to the formation of new grains and the wipe-off of original twin boundaries. At lower temperature such as 1025 8C, a higher density of LAGBs was observed in the specimen deformed to 80%, while the distribution of HAGBs is similar to that deformed to 170% at 1075 8C. It indicates that at lower temperature dislocations have stronger tendency to form LAGBs and to transform to HAGBs. TEM images of substructure in the deformed samples are shown in Fig. 10. The results of TEM observation on the dislocation configurations in Y grains showed that there exist a great number of subgrain boundaries, which divided the initial large grains into fine subgrains during deformation. The common feature of these boundaries is that they are composed of dislocation walls and dislocation networks. Fig. 11 shows a forming subgrain boundary. Two dislocations labelled as A intersect and react with each other. It is interesting to view that the dislocation at label A are originated from the dislocation in front of the original
grain boundary labelled as B. It seems that dislocation pile-ups are favourable to form subboundary. Moreover, long dislocation segments and deformation twins across subboundaries were observed too in g grains in Fig. 10. It is noticed that subgrain in Fig. 10c were elongated significantly, which are typical characteristics of dislocation gliding and also a signal of incompatible deformation. In contrast with the high density in subgrain boundaries, the dislocation density is low in the interior of grains. It suggests that at the steady-state (or quasi steady-state) flow stage, the dislocation in grain will remain in relatively low density due to the formation of grain boundaries. In Fig. 12, a series of TEM micrographs were used to identify dislocation character in the subgrain boundary labelled as C in Fig. 11. As seen from Fig. 12a, at operating vector g ¼ 220; the subboundary shows a dislocation net, which is composed by two sets of intersecting dislocation and g ¼ 11 1; one set of dislocation arrays. When g ¼ 111 arrays is invisible, alternately. At g ¼ 020; two arrays become all visible. Using the g·b ¼ 0 criteria for
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Fig. 9. Distribution of misorientation angles in deformed microstructure presented in Fig. 1: (a) 30, (b) 98 and (c) 170% at 1050 8C, and (d) 80% at 1025 8C.
the invisibility of dislocations, where b is Burgers vector, it is easy to deduce that the subgrain boundary is composed by type superdislocations. There is no visible [101] and ½101 dissociation of superdislocations to be observed. Apart from type dislocations, other type dislothe [101] and ½101 cations are rare to find. The interaction between dislocations and subgrain boundaries and the evolution of boundaries are still inexplicit, which necessitates further research to be carried out.
4. Discussions The flow behaviour in superplastic deformation of the large-grained TiAl alloy is similar to those in conventionally superplastic alloys. However, there exist some significant differences in microstructural evolution
and mechanism between the large-grained TiAl alloy and conventionally fine-grained superplastic materials. It was shown in this paper that superplastic behaviour was found in large-grained TiAl alloy without the usual prerequisites of fine grain size and grain boundary sliding [14]. Moreover, metallographic examination has shown that the average grain size of the large-grained TiAl alloy decreased during superplastic deformation and a much finer grain size could be obtained after superplastic deformation, while in finegrained superplastic alloys the grain size keeps nearly constant or has a little growth. The phenomenon of grain refinement is confirmed further by OIM analyses. In the present study, the microstructure evolution during superplastic deformation underwent the same process in both Fe3Al and FeAl alloys [5,6]. The initial large grains in large-grained TiAl alloys were refined during deformation. Most of the refined grains are bordered by subgrain
Fig. 10. Typical substructures seen after deformation of the TiAl alloy at 1075 8C and 1.6 £ 1024 s21: (a) deformed to 60%; (b) deformed to 98%; (c) deformed to 170%.
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Fig. 11. A forming subboundary and dislocation pile-ups in front of an original grain boundary in the sample deformed to 60% at 1075 8C under 1.6 £ 1024 s21.
boundaries or low angle grain boundaries. The density of subgrain boundaries and LAGBs, as well as HAGBs increased with the strain. Rotation between subgrains occurred during deformation as seen from the standard orientation triangles. In our previous work [5,6], we have proposed a continuous recovery and recrystallization model to account for the superplastic behaviour of initially large-grained Fe 3Al and FeAl alloys. Microstructure evolution observed in this study obviously indicates that a continuous recovery and recrystallization process appeared during superplastic deformation in large-grained TiAl alloy.
At the early stage of deformation, dislocation sliding is dominant. With further deformation, climb process occurred and the dislocations arrange themselves as subgrain boundaries. The subgrain boundaries divide the initial large grains into fine subgrains. The subgrain boundaries absorb the dislocations coming from the subgrains and thus the misorientation between subgrains increases. The conclusion that the subgrain boundaries can absorb the dislocations was based on not only the analysis on the dislocations in the grain boundaries but also the fact that the density of dislocations in the grains or subgrains is very low. Based on our previous
1; B ¼ ½110; Fig. 12. Diffraction contrast analysis of the dislocations in subgrain boundary: (a) g ¼ 220; B ¼ ½110; (b) g ¼ 111; B ¼ ½110; (c) g ¼ 11 (d) g ¼ 020; B ¼ ½101:
D. Lin, F. Sun / Intermetallics 12 (2004) 875–883
TEM observations, there exists high density of subgrain boundaries with the misorientation lower than 28, which cannot be detected by present EBSD technique. If we take the misorientation between subgrains lower than 28 into account, the density of subgrain boundaries and low angle grain boundaries in the samples would be even higher than that detected by EBSD. Based on our EBSD and TEM study, we believe that the superplastic deformation of both Fe3Al and FeAl alloys as well as TiAl alloy is controlled by dislocation gliding and climbing. The gliding dislocations interacted and arranged themselves as subgrain boundaries through their climb process. Superplastic deformation needs not only dislocation sliding but also the movement between subgrains, such as subgrain boundary sliding and migration, which should be controlled by dislocation climbing in the subgrain boundaries. The movement between subgrains results in the migration of the initial high angle grain boundaries of the alloys. During deformation, gliding dislocations can be trapped and absorbed by subboundaries, and then subboundaries evolve continuously towards LAGBs, in further, HAGBs. A constant dislocation density in the interior of grains is maintained under a constant stress at given temperature, which corresponds to the steady-state flow stage of superplastic deformation. The experimental results support our previous model of continuous recovery and recrystallization for the mechanism of superplastic deformation in large-grained iron aluminides. A dislocation glide and climb process accommodated by subgrain boundary sliding, migration and rotation allows the superplastic flow to proceed. However, there is more complexity in large-grained TiAl alloy due to the existence of twin boundaries and anti-phase domain boundaries. The superplastic deformation can be accommodated by deformation twinning, which is induced by elastic incompatibility. Moreover, deformation twins can act as dislocation sources or as effective obstacles to help in the formation of subgrain boundaries. In this consideration, deformation twinning may insert a beneficial effect in superplasticity in large-grained TiAl alloy while it is harmful in fine-grained TiAl alloy [15].
5. Conclusions The large-grained TiAl alloy exhibits superplasticity in a temperature range of 1025 –1100 8C and a strain rate range of 4 £ 1025 – 1.28 £ 1025 s21. A maximum elongation of
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287.5% with an m value of 0.39 was obtained at 1100 8C and 1.28 £ 1025 s21. A significant refinement of average grain size with an increase in the density of LAGBs and HAGBs during superplastic deformation was found, which is beneficial to improve mechanical properties of TiAl alloy. Evolution of grain boundaries including LAGBs and HAGBs was depicted by distribution of misorientation angles, a direct evidence of dynamical formation of grain boundaries with misorientation of 15– 308 by OIM analyses, which is a result of subboundary evolution. TEM observations have shown the formation of subboundaries and also the evidence of dislocation gliding in the interior of grains. The experimental results support our previous model of continuous recovery and recrystallization for the mechanism of superplastic deformation in large-grained iron aluminides. A dislocation glide and climb process accommodated by subgrain boundary sliding, migration and rotation allows the superplastic flow to proceed.
Acknowledgements This work was supported by the National Natural Science Foundation of China. The authors would like to thank Professor Q. Liu and Mr J. Meng of Tsinghua University for kind help on EBSD experiments.
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