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The flow behavior and the deformation mechanisms of Ti–6Al–2Zr–2Sn–2Mo–1.5Cr–2Nb alloy during isothermal compression
Journal of Alloys and Compounds 667 (2016) 44e52
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
The flow behavior and the deformation mechanisms of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression J. Luo*, J. Gao, L. Li, M.Q. Li School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, 710072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 September 2015 Received in revised form 24 December 2015 Accepted 21 January 2016 Available online 22 January 2016
The effects of the processing parameters on the shapes of flow curves, the microstructural evolution and the strain rate sensitivity are analyzed via the isothermal compression tests of Tie6Ale2Zre2Sne2Mo e1.5Cre2Nb alloy. The isothermal compression is performed on a Gleeble-1500 thermal simulator in the deformation temperature range of 1103e1243 K, strain rate range of 1.0 102 to 5.0 s1 and strain range of 0.2e1.2. The softening mechanisms are investigated thoroughly in the aþb phase region and b phase region through the experiments of optical microscopy, scanning electron microscopy and transmission electron microscopy. Then, the correlation between the flow behavior and the microstructural evolution is discussed. The results show that more noticeable flow softening at a high strain rate (5.0 s1) in the aþb phase region arises from the thermal softening, the dynamic recovery and the dynamic recrystallization of alpha phase. However, the thermal softening can no longer be considered to be the major softening mechanism at a low strain rate (1.0 102 s1). In the b phase region, the dynamic recovery and the dynamic recrystallization of b phase are main softening mechanisms. Moreover, the maximum m value of 0.3 occurs at a deformation temperature of 1163 K, a strain rate of 0.1 s1 and a strain of 0.7, in which the microstructure is equiaxed and uniform. © 2016 Elsevier B.V. All rights reserved.
Keywords: Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy Dynamic recovery Dynamic recrystallization Flow softening Strain rate sensitivity
1. Introduction Tie6Ale2Zre2Sne2Moe1.5Cre2Nb (TC21) alloy, as a type of two phase aþb titanium alloy, has recently been considered as a structure material for aircraft applications. This alloy was developed by adding some Nb elements to Ti-6-22-22S in Northwest Institute for Non-ferrous Metal Research, and exhibited the excellent combinations of ductility, fracture toughness and strength superior to those of Tie6Ale4V alloy [1,2]. In the past ten years, numerous investigations were performed to optimize the mechanical properties of Ti-6-22-22S alloy. Zhang et al. [2e4] investigated the microstructural development of Ti-622-22S alloy during heating treatment and its effect on the mechanical properties, and proposed that long term aging at 723e923 K had significant effect on fracture toughness. Long time aging at 923 K for 1000 h stabilized the a2 precipitations which grew to 5 ± 10 nm in diameter and 20 ± 50 nm in length. Wood
* Corresponding author. E-mail address: [email protected] (J. Luo). http://dx.doi.org/10.1016/j.jallcom.2016.01.164 0925-8388/© 2016 Elsevier B.V. All rights reserved.
et al. [5] reviewed various thermomechanical processing (TMP) routes and heat treatments on Ti-6-22-22S alloy so as to optimize the strength and the ductility. In addition, many researchers have focused on the deformation behavior, the microstructural evolution and the mechanical properties of TC21 alloy [6e9]. Fei et al. [10] investigated the effect of solution temperature, cooling rate and cooling mode on the phase and the microstructure of TC21 alloy, and found two precipitated phases (a2-Ti3Al and B2eTi2AlNb) through X-ray diffraction (XRD) analysis. Wang et al. [11] observed the microstructure at different stages of the transformation by scanning electron microscopy (SEM). Results showed that phase transformation in TC21 alloy during continuous heating included three stages, retained b/acicular a, acicular a/b and equiaxed a/b. Shao et al. [12] investigated room-temperature tensile deformation process of the lamellar microstructure in TC21 alloy by in-situ scanning electron microscopy (SEM). Zhu et al. [13] investigated the effect of deformation temperature, strain rate and height reduction on the microstructure evolution of as-cast TC21 alloy, and established the processing maps in which the optimal processing parameters were 1.0 102 s1/1423 K. Many uncertainties still remain regarding the correlation
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
between the flow behavior and the microstructural evolution during high temperature deformation of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy. And, the deformation mechanisms of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy in the aþb phase region and b phase region are not yet clarified. It is well known that revealing the relationship between the micromechanisms of deformation and observed plastic-flow response is very important to optimize the forging process and improve the workability of alloys, so further more investigations on Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy are needed. In this study, the shapes of flow curves are firstly analyzed in different processing parameter range. The microstructural evolution and the softening mechanisms are investigated thoroughly in the aþb phase region and b phase region through the experiments of optical microscopy, scanning electron microscopy and transmission electron microscopy. Then, the interaction between the flow stress and the microstructural evolution is discussed. The effect of the processing parameters on the strain rate sensitivity of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is analyzed based on the microstructural observations.
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Fig. 1. OM micrograph of as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at room temperature.
2. Material and experimental procedures
2.3. Microstructural observations
2.1. Material
The specimens for optical microscopy (OM) and scanning electron microscopy (SEM) were axially sectioned, electropolished and chemically etched in a solution of 10 ml HF, 15 ml HNO3 and 75 ml H2O. Photomicrographs were taken with a Leica DMI 3000M microscope. A SUPRA 55 SEM operating at 15 kV was used to examine the morphology of microstructure. The thin discs for transmission electron microscope (TEM) investigation were prepared by electrical-discharge machining followed by mechanical grinding to a thickness of less than 50 mm prior to ion beam thinning. TEM micrographs and corresponding selected area diffraction (SAD) pattern of each sample were obtained at 300 kV in a Technai F30G2.
The Tie6Ale2Zre2Sne2Moe1.5Cre2Nb forging bar with a 350 mm in diameter was used in present study. The chemical composition of as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is shown in Table 1. The microstructure of this alloy as revealed by optical metallography is shown in Fig. 1, and exhibits a typical duplex morphology. The grain size and the volume fraction of equiaxed alpha are measured to be about 7.8 mm and 48.0% using the quantitative metallography image analysis software. The b transus temperature for this alloy was about 1238 K and was in a good agreement with that reported by Shi et al. [9]. The mechanical properties of this alloy at room temperature were measured, and the yield strength, tensile strength, elongation and reduction were determined to be 1030 ± 13 MPa, 1103 ± 22 MPa, 14 ± 1%, and 51 ± 4%, respectively.
2.2. Isothermal compression tests Cylindrical compression specimens with 8.0 mm in diameter and 12.0 mm in height were machined from the as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb bar. The isothermal compression was performed on a Gleeble-1500 thermal simulator in the deformation temperature range of 1103e1243 K, strain rate range of 1.0 102 to 5.0 s1 and strain range of 0.2e1.2. Graphite powder was put between the specimens and the anvils so as to reduce the friction during isothermal compression. Prior to compression, the specimens were heated to a given deformation temperature at a linear rate of 20 K s1 and then held for 5 min. After isothermal compression, the isothermally compressed Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy was cooled in air to room temperature.
3. Results and discussion 3.1. Flow stress Selected flow stressestrain curves of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression are given in Fig. 2. For all deformation temperatures and strain rates, the flow softening is observed for Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy after the peak stress; but it at higher strain rates (0.1 s1) and lower deformation temperatures (1163 K) is more significant than that at lower strain rates and higher deformation temperatures. At a strain rate of 5.0 s1 and in the deformation temperature range of 1103e1223 K, the flow curves drop continuously and the steady state conditions are not observed, even up to a strain of 0.9, as illustrated in Fig. 2(a)e(d). The noticeable flow softening at higher strain rates and lower deformation temperatures is possibly attributed to the deformation heating and the microstructural changes. It is well known that most of plastic power converts to heat during high temperature deformation of metals or alloys. However, the thermal conductivity of
Table 1 Chemical composition of the as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy (mass fraction in %). Al
Mo
Nb
Sn
Zr
Cr
Si
Fe
C
N
O
Ti
6.4
2.8
2.0
2.0
2.2
1.6
6.8 102
2.1 102
5.0 103
4.0 103
0.1
Bal.
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Fig. 2. Selected flow stressestrain curves of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression.
titanium alloy is very low that makes that there is not enough time for the heat generated by plastic deformations to diffuse away. Finally, the adiabatic heating generated during deformation will raise the actual temperature of the samples and also the proportion of soft b phase [13,14]; so the loss of strength caused by the increase in temperature will, at a certain point, exceed the increase in strength due to working-hardening [15]. At this point, the stressestrain curve starts to go down, and thermal softening sets in. In addition, the microstructural changes discussed in Section 3.2 are also considered to be the principal source of flow softening. The microstructural characteristics of the processes occurring in these domains are discussed below. Moreover, the flow curves with
broad oscillations are observed at a strain rate of 5.0 s1 and in the deformation temperature range of 1103e1223 K. The broad oscillations are possibly attributed to unstable deformation. Similarly, Zhu et al. [16] also proposed that the oscillation phenomenon of ascast TC21 titanium alloy was indicative of instability such as flow localization or cracking. From Fig. 2, it is also seen that the continuous oscillations exist all the strain rates (5.0 s1, 1.0 s1, 0.1 s1 1.0 102 s1). The continuous oscillations have no relation to the deformation mechanisms (such as dynamic recrystallization). The continuous oscillations arise from the experimental procedure so as to obtain a constant strain rate during isothermal compression.
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
At a lower strain rate (0.01 s1) in the aþb phase region, the flow stress reaches a peak at a critical strain and then decreases to a steady state value at strains greater than 0.6. For deformation in the b phase region, the flow stress appears to reach a steady state value at a small strain, which implies that the effect of the strain on the flow stress is negligible, as shown in Fig. 2(e). It is also seen in Fig. 2 that the flow stress of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is sensitive to the deformation temperature and the strain rate. The flow stress sharply decreases with increasing deformation temperature at a fixed strain rate. The phenomenon is possibly attributed to the enhancement of dislocation motion and the grain boundary diffusion and a rapid increase of b phase at higher deformation temperatures due to phase transformation. The flow stress increases with increasing strain rate at a fixed deformation temperature. The main reason is that the rate of dislocation generation increases with increasing strain rate and the pre-exist nuclei cannot provide enough softening [17] because high strain rates promote less time for recovery processes. In other words, the rate of strain hardening increases with increasing strain rate. The tangled dislocation structures hinder the dislocation movement, leading to the increase in flow stress. In order to clarify the flow softening mechanism in different processing parameter range, the microstructural observations are implemented in the following section. And, the interaction between the flow stress and the microstructural evolution is discussed. 3.2. Microstructural evolution 3.2.1. Deformation in the aþb phase region Fig. 3 shows that the effect of the strain on the microstructural evolution at a deformation temperature of 1163 K and a strain rate of 1.0 102 s1. When the strain is 0.2, a lot of dislocations gather at the grain boundaries, as illustrated in Fig. 3(a). The SAD pattern (location A in Fig. 3(a)) with beam direction parallel to the ½1213 zone axis reveals the alpha grain (Fig. 3(b)). It is well known that the grain boundaries have a strong impediment for the dislocation motion because the grain orientation is different among adjacent grains. The dislocations will pile up at grain boundaries, and consequently, there is an appreciable work-hardening at the beginning of strain. Moreover, at the beginning of strain the rapid generation of mobile dislocations will promote the interaction between dislocations and dislocations, giving rise to dislocation dipoles and loops, and finally resulting in local dislocation tangles. Thus, the flow stress value at a lower strain (0.2) is high during isothermal compression of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy, as illustrated in Fig. 2(b). In addition, some acicular alpha is observed in transformed beta matrix. The distribution of acicular alpha is mixed and disorderly. When the strain is up to 0.5, the dislocation network and even the sub-boundaries are observed at the grains interiors of equiaxed alpha (Fig. 3(c)). Generally, multiple slip occurs for titanium alloys with increasing strain; primary dislocations interact with secondary dislocations, giving rise to dislocation dipoles and loops which result in local dislocation tangles and, eventually, the network of sub-boundaries. It is an accepted postulate that the dislocations agglomerate into cells during deformation because this configuration represents a minimum strain energy for a given dislocation content [18]. Then the dislocation density in the dislocation cells would decrease continuously, and the dislocations at boundaries would elongate and condense to decrease the thickness of cell boundary, which finally forms the dislocation network and the sub-boundaries. Therefore, there is no doubt that the evolution of dislocation network or the sub-boundaries is a type of the release of distortion energy and a kind of softening mechanisms. It is also seen from Fig. 3(d) that the lamellar alpha bending occurs at this deformation condition; and
47
the dislocations in cell walls appear in the bended lamellar alpha interior. The walls of cells formed under current circumstances are rather thick, with a small misorientation. Consequently, the walls will constitute relatively sharp sub-boundaries which are beneficial for the globularization process of lamellar alpha. Semiatin et al. [19] had proposed that the lamellar alpha bending was responsible for the flow softening, thus the lamellar alpha bending in present study is also a kind of softening mechanisms. The dislocation network, the walls of cells and sub-boundaries indicate a typical feature of dynamic recovery (DRV). Therefore, the DRV is main softening mechanism at a strain of 0.5. When the strain is 0.9, the more noticeable sub-boundaries are observed in Fig. 3(e)e(f). In general, the sub-boundaries exhibit straight which reveals the existence of DRV structure. Moreover, the penetration of b into lamellar a can be observed (Fig. 3(g)), the lamellar a would finally be separated to form the globules through interfaces migration to minimize the surface energy. It is also observed that the dynamic recrystallization (DRX) appears at this condition (Fig. 3(h)); the size of small recrystallized a grain is about 0.7 mm. This demonstrates that DRX is a type of softening mechanisms at a strain of 0.9. According to above-mentioned analysis, it is concluded that the heavily tangled dislocations at grain boundaries and in grain interiors lead to the working hardening during early deformation. However, the DRV structure including the lamellar alpha bending and the formation of sub-boundaries will be beneficial for the softening of materials. Finally, the globularization of lamellar a and DRX for a phase will participate in the softening of materials. So the flow stress firstly represents the tendency of a sharp increase at the beginning of strain, reaches a peak at a critical strain and then continuously decreases to a steady state. At a low strain rate (1.0 102 s1), there is sufficient time for the heat to escape from the specimens during deformation, the material is considered to be isothermal. Thus, thermal softening can no longer be considered to be the major softening mechanism. Fig. 4 shows that the microstructural observations at a deformation temperature of 1163 K, a strain rate of 5.0 s1 and a strain of 0.9. At a higher strain rate, there is insufficient time for the globularization for Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy. The morphology of a phase exhibits two characteristics: elongated and equiaxed grains, as illustrated in Fig. 4(a)-(b). The fine b grains marked in black arrow are also observed. The elongated a grains are perpendicular to the compression axis direction and have an obviously oriented characteristic. This is a typical dynamic recovery structure, implying that the dominant softening mechanism is dynamic recovery at this strain rate. Detailed TEM study also reveals the presence of elongated a grains at this strain rate, as illustrated in Fig. 4(c), (e) and (f). Fig. 4(d) is SAD pattern obtained from Fig. 4(c) (location A) with beam direction of ½1213a. Compared to TEM micrographs in Fig. 3 at a low strain rate, more dislocations in cell walls appear in the elongated alpha interior because of an increase of the dislocation generation rate at a high strain rate. The dislocations distributing in array will rearrange and finally form the dislocation network or the sub-boundaries during the consequent hot working or heat treatment processes as the author has outlined above. Some acicular a with a thickness of about 0.1 mm are observed in transformed b matrix. The acicular a precipitates from transformed b matrix during the air-cooling process after isothermal compression. Moreover, some fine DRX a grains appear at this condition (Fig. 4(g)e(h)). DRX would also decrease the distorted energy and work as another softening mechanism. At a high strain rate, because there is insufficient time for the heat to escape from the specimen during deformation, the material can no longer be considered be isothermal, and the loss of strength caused by the increase in temperature will, at a certain point, exceed the increase in strength due to work-hardening [15].
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Fig. 3. Microstructural observations of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1163 K/1.0 102 s1: (a) TEM micrograph after deformed at a strain of 0.2; (b) is SAD pattern obtained from (a) (location A) with beam direction on ½1213a ; (c and d) TEM micrographs after deformed at a strain of 0.5; (e, f, g and h) TEM micrographs after deformed at a strain of 0.9.
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
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Fig. 4. Microstructural observations of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1163 K/5.0 s1/0.9: (a and b) SEM micrographs in adjacent visual fields; (c, e, f, g and h) TEM micrographs; (d) is SAD pattern obtained from (c) (location A) with beam direction on ½1213a .The compression axis is vertical in SEM micrographs.
Thus, the flow curves show a sharp decrease beyond the peak
stress. In addition, the occurrence of DRV and DRX will enhance the
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softening effect during hot working of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy as above-mentioned analysis. So the thermal softening, the DRV and the DRX will result in that the flow curves at a higher strain rate drop continuously and the steady state conditions are not observed, even up to a strain of 0.9. 3.2.2. Deformation in the b phase region The effect of the strain rate on the microstructure of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at a deformation temperature of 1243 K and a strain of 0.9 is shown in Fig. 5. It is seen that the morphology of b phase exhibits two characteristics: fine recrystallized and coarse elongated b grains. The references [20,21] reported that a mixed b grain structure containing fine and coarse grains was the result of selective recrystallization in regions of high strain along beta grain boundaries during the b phase working and the reduced driving force for recrystallization in the grain interiors, where dynamic recovery dominated. This demonstrates that the DRV is still a type of softening mechanisms in the b phase region. In addition, the amount of recrystallized b grains increases with increasing strain rate, but apparently the size of recrystallized b grains tends to decrease. The main reason is that the occurrence of recrystallization nuclei is primarily dependent on the substructure cell size and the related cell-wall density [17]. More dislocations in cell walls appear at a higher strain rate compared to those at a lower strain rate as above-mentioned in Section 3.2.1, which promotes the occurrence of b recrystallization. However, there is insufficient time for the b grain growth at a higher strain rate, resulting in the small grain size at this strain rate. According to above-mentioned analysis, it is concluded that the DRV and the DRX of b phase are main softening mechanisms in the b phase region. Because the b phase has more slip systems than the a phase, the working hardening is not serious in the b phase region. This promotes that the dynamic softening is sufficient to counteract the work-hardening of the material at a lower strain. So the flow stress
in the b phase region appears to reach a steady state value at a small strain. 3.3. The strain rate sensitivity The strain rate sensitivity is usually used to determine the deformation mechanisms of a material. Values of the strain rate sensitivity are calculated using the following expression [22]:
m¼
d log s , d log ε
(1) ε;T
where s is the flow stress at a fixed strain and a fixed deformation : temperature (MPa), ε is the strain rate(s1), ε is the strain, and T is the absolute deformation temperature ( C). The strain rate sensitivity is estimated from the flow stressestrain curves during isothermal compression of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy and is shown in Fig. 6. From Fig. 6, it is observed that the values of the strain rate sensitivity reveal a noticeable dependence on the strain, the strain rate and the deformation temperature. At a deformation temperature of 1163 K and a strain rate of 0.1 s1 (Fig. 6(a)), the m values are in the range of 0.2e0.3 suggesting that the deformation mechanism is likely dislocation glide/climb [23]. And generally, m tends to increase with increasing strain, the value of m increases to 0.3 at a strain of 0.7. The microstructural observation at this strain is shown in Fig. 7(a). It is seen that the a grains are equiaxed and uniform, which is beneficial for the grain boundary sliding and accommodation. At a strain rate of 1.0 s1, the m values are approximately 0.2 or less, which are typical of deformation controlled by dislocation glide process characteristic. The a grains show two morphology characteristics: equiaxed and elongated, as illustrated in Fig. 7(b). The grain boundaries of elongated alpha are observed to become jagged and serrated at this strain rate. At a strain rate of 5.0 s1, the
Fig. 5. Optical micrographs of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1243 K/0.9: (a) 1.0 102 s1; (b) 0.1 s1; (c) 1.0 s1. The compression axis is vertical in all micrographs.
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
51
Fig. 6. Variation of the strain-rate sensitivity of the flow stress as a function of strain and strain rate at test temperatures of (a) 1163 K; (b) 1223 K.
Fig. 7. Optical micrographs of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at a strain of 0.7: (a) 1163 K/0.1 s1; (b) 1163 K/1.0 s1; (c) 1163 K/5.0 s1; (d) 1223 K/0.1 s1. The compression axis is vertical in all micrographs.
curve of m tends to oscillatory, and the m values are below 0.1, at which the deformation also appears to be dominated by dislocation glide process. The microstructural observation at this strain rate is shown in Fig. 7(c). It is observed that the microstructure is nonuniform and has slight oriented characteristics. The elongated a grains are perpendicular to the compression axis direction. So the non-uniformity of microstructure will finally lead to that the m values are small. At a deformation temperature of 1223 K, the variations of m with strain rate are similar to those at a deformation temperature of 1163 K. But, the m values decreases compared to that at a lower deformation temperature, which is attributed to a sharp decrease of equiaxed a phase. According to previous superplastic theory [24], it was concluded that approximately equal volume fractions of a phase and b phase in the alloy contributed to
grain boundary sliding and accommodation. Thus only 8.0% volume fraction of a phase at 1223 K will finally result in the smaller m values, as illustrated in Fig. 7(d).
4. Conclusions The flow curves and the strain rate sensitivity of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy are analyzed with the help of the microstructural observations. The microstructural evolution and the softening mechanisms are investigated thoroughly in the aþb phase region and b phase region, respectively. The following conclusions are obtained from the present investigation.
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(1) At a high strain rates (5.0 s1) in the aþb phase region, more noticeable flow softening arises from the thermal softening, the dynamic recovery and the dynamic recrystallization of alpha phase. (2) At a low strain rate (0.01 s1) in the aþb phase region, the dynamic recovery structure including the lamellar alpha bending and the formation of sub-boundaries, and the globularization of lamellar alpha and the DRX for alpha phase will participate in the softening of materials. (3) In the b phase region, the working hardening is not serious, which will promote that the dynamic softening effect (the DRV and the DRX of b phase) is sufficient to counteract the work-hardening at a lower strain. So the flow stress in the b phase region appears to reach a steady state value at a small strain. (4) The maximum m value of 0.3 occurs at 1163 K/0.1 s1/0.7 during the isothermal compression of this alloy. The microstructure at this processing condition is equiaxed and uniform, which is beneficial for the grain boundary sliding and accommodation. Acknowledgments
[6]
[7]
[8]
[9]
[10] [11]
[12]
[13]
[14] [15] [16]
The authors thank the financial supports from the National Natural Science Foundation of China with Grant No. 51205318 and No. 51275416, the Innovation Fund of Science and Technology in Northwestern Polytechnical University with Grant No. 2012KJ02002.
[17]
[18] [19]
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and heat treatment of Ti-6Al-2Sn-2Zr-2Cr-2Mo-Si for structural applications, Mater. Sci. Eng. A 243 (1998) 109e118. Y.Y. Zong, Y.C. Liang, Z.W. Yin, D.B. Shan, Effects of hydrogen addition on the high temperature deformation behavior of TC21 titanium alloy, Int. J. Hydrogen Energ 37 (2012) 13631e13637. Y.Q. Su, S.J. Wang, L.S. Luo, L. Wang, F.Y. Dong, J.J. Guo, H.Z. Fu, Gradient microstructure of TC21 alloy induced by hydrogen during hydrogenation, Int. J. Hydrogen Energ 37 (2012) 19210e19218. X.L. Wang, Y.Q. Zhao, H.L. Hou, Y.Q. Wang, Effect of hydrogen content on superplastic forming/diffusion bonding of TC21 alloys, J. Alloy. Compd. 503 (2010) 151e154. Z.F. Shi, H.Z. Guo, R. Liu, X.C. Wang, Z.K. Yao, Microstructure and mechanical properties of TC21 titanium alloy by near-isothermal forging, T. Nonferr. Metal. Soc. China 25 (2015) 72e79. Y.H. Fei, L. Zhou, H.L. Qu, Y.Q. Zhao, C.Z. Huang, The phase and microstructure of TC21 alloy, Mater. Sci. Eng. A 494 (2008) 166e172. Y.H. Wang, H.C. Kou, H. Chang, Z.S. Zhu, X.F. Su, J.S. Li, L. Zhou, Phase transformation in TC21 alloy during continuous heating, J. Alloy. Compd. 472 (2009) 252e256. H. Shao, Y.Q. Zhao, P. Ge, W.D. Zeng, In-situ SEM observations of tensile deformation of the lamellar microstructure in TC21 titanium alloy, Mater. Sci. Eng. A 559 (2013) 515e519. Y.C. Zhu, W.D. Zeng, J.L. Liu, Y.Q. Zhao, Y.G. Zhou, H.Q. Yu, Effect of processing parameters on the hot deformation behavior of as-cast TC21 titanium alloy, Mater. Des. 33 (2012) 264e272. R. Ding, Z.X. Guo, A. Wilson, Microstructural evolution of a Ti-6Al-4V alloy during thermomechanical processing, Mater. Sci. Eng. A 327 (2002) 233e245. M.A. Meyers, K.K. Chawla, Mechanical Behavior of Materials, Cambridge University Press, New York, 2009, pp. 392e395. Y.C. Zhu, W.D. Zeng, F. Feng, Y. Sun, Y.F. Han, Y.G. Zhou, Characterization of hot deformation behavior of as-cast TC21 titanium alloy using processing map, Mater. Sci. Eng. A 528 (2011) 1757e1763. H.J. McQueen, The role of dynamically recovered substructure in dynamic recrystallization, in: G. Krauss (Ed.), Deformation, Processing, and Structure, American Society for Metals, Metal Park, Ohio, 1982, p. 235. J.H. Gittus, Theoretical equation for steady-state dislocation creep: effects of jog drag and cell formation, Phil. Mag. 31 (1975) 465e470. S.L. Semiatin, V. Seetharaman, I. Weiss, Flow behavior and globularization kinetics during hot working of Ti-6Al-4V with a colony alpha microstructure, Mater. Sci. Eng. A 263 (1999) 257e271. I. Weiss, S.L. Semiatin, Thermomechanical processing of beta titanium alloysan overview, Mater. Sci. Eng. A 243 (1998) 46e65. P. Wanjara, M. Jahazi, H. Monajati, S. Yue, Influence of thermomechanical processing on microstructural evolution in near-a alloy IMI834, Mater. Sci. Eng. A 416 (2006), 300e300. J. Luo, M.Q. Li, W.X. Yu, H. Li, The variation of strain rate sensitivity exponent and strain hardening exponent in isothermal compression of Ti-6Al-4V alloy, Mater. Des. 31 (2010) 741e748. S.L. Semiatin, T.R. Bieler, The effect of alpha platelet thickness on plastic flow during hot working of Ti-6Al-4V with a transformed microstructure, Acta Mater 49 (2001) 3565e3573. J.W.D. Patterson, N. Ridley, Effect of phase proportions on deformation and cavitation of superplastic a/b brass, J. Mater. Sci. 16 (1981) 457e464.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
The flow behavior and the deformation mechanisms of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression J. Luo*, J. Gao, L. Li, M.Q. Li School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, 710072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 September 2015 Received in revised form 24 December 2015 Accepted 21 January 2016 Available online 22 January 2016
The effects of the processing parameters on the shapes of flow curves, the microstructural evolution and the strain rate sensitivity are analyzed via the isothermal compression tests of Tie6Ale2Zre2Sne2Mo e1.5Cre2Nb alloy. The isothermal compression is performed on a Gleeble-1500 thermal simulator in the deformation temperature range of 1103e1243 K, strain rate range of 1.0 102 to 5.0 s1 and strain range of 0.2e1.2. The softening mechanisms are investigated thoroughly in the aþb phase region and b phase region through the experiments of optical microscopy, scanning electron microscopy and transmission electron microscopy. Then, the correlation between the flow behavior and the microstructural evolution is discussed. The results show that more noticeable flow softening at a high strain rate (5.0 s1) in the aþb phase region arises from the thermal softening, the dynamic recovery and the dynamic recrystallization of alpha phase. However, the thermal softening can no longer be considered to be the major softening mechanism at a low strain rate (1.0 102 s1). In the b phase region, the dynamic recovery and the dynamic recrystallization of b phase are main softening mechanisms. Moreover, the maximum m value of 0.3 occurs at a deformation temperature of 1163 K, a strain rate of 0.1 s1 and a strain of 0.7, in which the microstructure is equiaxed and uniform. © 2016 Elsevier B.V. All rights reserved.
Keywords: Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy Dynamic recovery Dynamic recrystallization Flow softening Strain rate sensitivity
1. Introduction Tie6Ale2Zre2Sne2Moe1.5Cre2Nb (TC21) alloy, as a type of two phase aþb titanium alloy, has recently been considered as a structure material for aircraft applications. This alloy was developed by adding some Nb elements to Ti-6-22-22S in Northwest Institute for Non-ferrous Metal Research, and exhibited the excellent combinations of ductility, fracture toughness and strength superior to those of Tie6Ale4V alloy [1,2]. In the past ten years, numerous investigations were performed to optimize the mechanical properties of Ti-6-22-22S alloy. Zhang et al. [2e4] investigated the microstructural development of Ti-622-22S alloy during heating treatment and its effect on the mechanical properties, and proposed that long term aging at 723e923 K had significant effect on fracture toughness. Long time aging at 923 K for 1000 h stabilized the a2 precipitations which grew to 5 ± 10 nm in diameter and 20 ± 50 nm in length. Wood
* Corresponding author. E-mail address: [email protected] (J. Luo). http://dx.doi.org/10.1016/j.jallcom.2016.01.164 0925-8388/© 2016 Elsevier B.V. All rights reserved.
et al. [5] reviewed various thermomechanical processing (TMP) routes and heat treatments on Ti-6-22-22S alloy so as to optimize the strength and the ductility. In addition, many researchers have focused on the deformation behavior, the microstructural evolution and the mechanical properties of TC21 alloy [6e9]. Fei et al. [10] investigated the effect of solution temperature, cooling rate and cooling mode on the phase and the microstructure of TC21 alloy, and found two precipitated phases (a2-Ti3Al and B2eTi2AlNb) through X-ray diffraction (XRD) analysis. Wang et al. [11] observed the microstructure at different stages of the transformation by scanning electron microscopy (SEM). Results showed that phase transformation in TC21 alloy during continuous heating included three stages, retained b/acicular a, acicular a/b and equiaxed a/b. Shao et al. [12] investigated room-temperature tensile deformation process of the lamellar microstructure in TC21 alloy by in-situ scanning electron microscopy (SEM). Zhu et al. [13] investigated the effect of deformation temperature, strain rate and height reduction on the microstructure evolution of as-cast TC21 alloy, and established the processing maps in which the optimal processing parameters were 1.0 102 s1/1423 K. Many uncertainties still remain regarding the correlation
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
between the flow behavior and the microstructural evolution during high temperature deformation of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy. And, the deformation mechanisms of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy in the aþb phase region and b phase region are not yet clarified. It is well known that revealing the relationship between the micromechanisms of deformation and observed plastic-flow response is very important to optimize the forging process and improve the workability of alloys, so further more investigations on Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy are needed. In this study, the shapes of flow curves are firstly analyzed in different processing parameter range. The microstructural evolution and the softening mechanisms are investigated thoroughly in the aþb phase region and b phase region through the experiments of optical microscopy, scanning electron microscopy and transmission electron microscopy. Then, the interaction between the flow stress and the microstructural evolution is discussed. The effect of the processing parameters on the strain rate sensitivity of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is analyzed based on the microstructural observations.
45
Fig. 1. OM micrograph of as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at room temperature.
2. Material and experimental procedures
2.3. Microstructural observations
2.1. Material
The specimens for optical microscopy (OM) and scanning electron microscopy (SEM) were axially sectioned, electropolished and chemically etched in a solution of 10 ml HF, 15 ml HNO3 and 75 ml H2O. Photomicrographs were taken with a Leica DMI 3000M microscope. A SUPRA 55 SEM operating at 15 kV was used to examine the morphology of microstructure. The thin discs for transmission electron microscope (TEM) investigation were prepared by electrical-discharge machining followed by mechanical grinding to a thickness of less than 50 mm prior to ion beam thinning. TEM micrographs and corresponding selected area diffraction (SAD) pattern of each sample were obtained at 300 kV in a Technai F30G2.
The Tie6Ale2Zre2Sne2Moe1.5Cre2Nb forging bar with a 350 mm in diameter was used in present study. The chemical composition of as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is shown in Table 1. The microstructure of this alloy as revealed by optical metallography is shown in Fig. 1, and exhibits a typical duplex morphology. The grain size and the volume fraction of equiaxed alpha are measured to be about 7.8 mm and 48.0% using the quantitative metallography image analysis software. The b transus temperature for this alloy was about 1238 K and was in a good agreement with that reported by Shi et al. [9]. The mechanical properties of this alloy at room temperature were measured, and the yield strength, tensile strength, elongation and reduction were determined to be 1030 ± 13 MPa, 1103 ± 22 MPa, 14 ± 1%, and 51 ± 4%, respectively.
2.2. Isothermal compression tests Cylindrical compression specimens with 8.0 mm in diameter and 12.0 mm in height were machined from the as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb bar. The isothermal compression was performed on a Gleeble-1500 thermal simulator in the deformation temperature range of 1103e1243 K, strain rate range of 1.0 102 to 5.0 s1 and strain range of 0.2e1.2. Graphite powder was put between the specimens and the anvils so as to reduce the friction during isothermal compression. Prior to compression, the specimens were heated to a given deformation temperature at a linear rate of 20 K s1 and then held for 5 min. After isothermal compression, the isothermally compressed Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy was cooled in air to room temperature.
3. Results and discussion 3.1. Flow stress Selected flow stressestrain curves of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression are given in Fig. 2. For all deformation temperatures and strain rates, the flow softening is observed for Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy after the peak stress; but it at higher strain rates (0.1 s1) and lower deformation temperatures (1163 K) is more significant than that at lower strain rates and higher deformation temperatures. At a strain rate of 5.0 s1 and in the deformation temperature range of 1103e1223 K, the flow curves drop continuously and the steady state conditions are not observed, even up to a strain of 0.9, as illustrated in Fig. 2(a)e(d). The noticeable flow softening at higher strain rates and lower deformation temperatures is possibly attributed to the deformation heating and the microstructural changes. It is well known that most of plastic power converts to heat during high temperature deformation of metals or alloys. However, the thermal conductivity of
Table 1 Chemical composition of the as-received Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy (mass fraction in %). Al
Mo
Nb
Sn
Zr
Cr
Si
Fe
C
N
O
Ti
6.4
2.8
2.0
2.0
2.2
1.6
6.8 102
2.1 102
5.0 103
4.0 103
0.1
Bal.
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Fig. 2. Selected flow stressestrain curves of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy during isothermal compression.
titanium alloy is very low that makes that there is not enough time for the heat generated by plastic deformations to diffuse away. Finally, the adiabatic heating generated during deformation will raise the actual temperature of the samples and also the proportion of soft b phase [13,14]; so the loss of strength caused by the increase in temperature will, at a certain point, exceed the increase in strength due to working-hardening [15]. At this point, the stressestrain curve starts to go down, and thermal softening sets in. In addition, the microstructural changes discussed in Section 3.2 are also considered to be the principal source of flow softening. The microstructural characteristics of the processes occurring in these domains are discussed below. Moreover, the flow curves with
broad oscillations are observed at a strain rate of 5.0 s1 and in the deformation temperature range of 1103e1223 K. The broad oscillations are possibly attributed to unstable deformation. Similarly, Zhu et al. [16] also proposed that the oscillation phenomenon of ascast TC21 titanium alloy was indicative of instability such as flow localization or cracking. From Fig. 2, it is also seen that the continuous oscillations exist all the strain rates (5.0 s1, 1.0 s1, 0.1 s1 1.0 102 s1). The continuous oscillations have no relation to the deformation mechanisms (such as dynamic recrystallization). The continuous oscillations arise from the experimental procedure so as to obtain a constant strain rate during isothermal compression.
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At a lower strain rate (0.01 s1) in the aþb phase region, the flow stress reaches a peak at a critical strain and then decreases to a steady state value at strains greater than 0.6. For deformation in the b phase region, the flow stress appears to reach a steady state value at a small strain, which implies that the effect of the strain on the flow stress is negligible, as shown in Fig. 2(e). It is also seen in Fig. 2 that the flow stress of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy is sensitive to the deformation temperature and the strain rate. The flow stress sharply decreases with increasing deformation temperature at a fixed strain rate. The phenomenon is possibly attributed to the enhancement of dislocation motion and the grain boundary diffusion and a rapid increase of b phase at higher deformation temperatures due to phase transformation. The flow stress increases with increasing strain rate at a fixed deformation temperature. The main reason is that the rate of dislocation generation increases with increasing strain rate and the pre-exist nuclei cannot provide enough softening [17] because high strain rates promote less time for recovery processes. In other words, the rate of strain hardening increases with increasing strain rate. The tangled dislocation structures hinder the dislocation movement, leading to the increase in flow stress. In order to clarify the flow softening mechanism in different processing parameter range, the microstructural observations are implemented in the following section. And, the interaction between the flow stress and the microstructural evolution is discussed. 3.2. Microstructural evolution 3.2.1. Deformation in the aþb phase region Fig. 3 shows that the effect of the strain on the microstructural evolution at a deformation temperature of 1163 K and a strain rate of 1.0 102 s1. When the strain is 0.2, a lot of dislocations gather at the grain boundaries, as illustrated in Fig. 3(a). The SAD pattern (location A in Fig. 3(a)) with beam direction parallel to the ½1213 zone axis reveals the alpha grain (Fig. 3(b)). It is well known that the grain boundaries have a strong impediment for the dislocation motion because the grain orientation is different among adjacent grains. The dislocations will pile up at grain boundaries, and consequently, there is an appreciable work-hardening at the beginning of strain. Moreover, at the beginning of strain the rapid generation of mobile dislocations will promote the interaction between dislocations and dislocations, giving rise to dislocation dipoles and loops, and finally resulting in local dislocation tangles. Thus, the flow stress value at a lower strain (0.2) is high during isothermal compression of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy, as illustrated in Fig. 2(b). In addition, some acicular alpha is observed in transformed beta matrix. The distribution of acicular alpha is mixed and disorderly. When the strain is up to 0.5, the dislocation network and even the sub-boundaries are observed at the grains interiors of equiaxed alpha (Fig. 3(c)). Generally, multiple slip occurs for titanium alloys with increasing strain; primary dislocations interact with secondary dislocations, giving rise to dislocation dipoles and loops which result in local dislocation tangles and, eventually, the network of sub-boundaries. It is an accepted postulate that the dislocations agglomerate into cells during deformation because this configuration represents a minimum strain energy for a given dislocation content [18]. Then the dislocation density in the dislocation cells would decrease continuously, and the dislocations at boundaries would elongate and condense to decrease the thickness of cell boundary, which finally forms the dislocation network and the sub-boundaries. Therefore, there is no doubt that the evolution of dislocation network or the sub-boundaries is a type of the release of distortion energy and a kind of softening mechanisms. It is also seen from Fig. 3(d) that the lamellar alpha bending occurs at this deformation condition; and
47
the dislocations in cell walls appear in the bended lamellar alpha interior. The walls of cells formed under current circumstances are rather thick, with a small misorientation. Consequently, the walls will constitute relatively sharp sub-boundaries which are beneficial for the globularization process of lamellar alpha. Semiatin et al. [19] had proposed that the lamellar alpha bending was responsible for the flow softening, thus the lamellar alpha bending in present study is also a kind of softening mechanisms. The dislocation network, the walls of cells and sub-boundaries indicate a typical feature of dynamic recovery (DRV). Therefore, the DRV is main softening mechanism at a strain of 0.5. When the strain is 0.9, the more noticeable sub-boundaries are observed in Fig. 3(e)e(f). In general, the sub-boundaries exhibit straight which reveals the existence of DRV structure. Moreover, the penetration of b into lamellar a can be observed (Fig. 3(g)), the lamellar a would finally be separated to form the globules through interfaces migration to minimize the surface energy. It is also observed that the dynamic recrystallization (DRX) appears at this condition (Fig. 3(h)); the size of small recrystallized a grain is about 0.7 mm. This demonstrates that DRX is a type of softening mechanisms at a strain of 0.9. According to above-mentioned analysis, it is concluded that the heavily tangled dislocations at grain boundaries and in grain interiors lead to the working hardening during early deformation. However, the DRV structure including the lamellar alpha bending and the formation of sub-boundaries will be beneficial for the softening of materials. Finally, the globularization of lamellar a and DRX for a phase will participate in the softening of materials. So the flow stress firstly represents the tendency of a sharp increase at the beginning of strain, reaches a peak at a critical strain and then continuously decreases to a steady state. At a low strain rate (1.0 102 s1), there is sufficient time for the heat to escape from the specimens during deformation, the material is considered to be isothermal. Thus, thermal softening can no longer be considered to be the major softening mechanism. Fig. 4 shows that the microstructural observations at a deformation temperature of 1163 K, a strain rate of 5.0 s1 and a strain of 0.9. At a higher strain rate, there is insufficient time for the globularization for Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy. The morphology of a phase exhibits two characteristics: elongated and equiaxed grains, as illustrated in Fig. 4(a)-(b). The fine b grains marked in black arrow are also observed. The elongated a grains are perpendicular to the compression axis direction and have an obviously oriented characteristic. This is a typical dynamic recovery structure, implying that the dominant softening mechanism is dynamic recovery at this strain rate. Detailed TEM study also reveals the presence of elongated a grains at this strain rate, as illustrated in Fig. 4(c), (e) and (f). Fig. 4(d) is SAD pattern obtained from Fig. 4(c) (location A) with beam direction of ½1213a. Compared to TEM micrographs in Fig. 3 at a low strain rate, more dislocations in cell walls appear in the elongated alpha interior because of an increase of the dislocation generation rate at a high strain rate. The dislocations distributing in array will rearrange and finally form the dislocation network or the sub-boundaries during the consequent hot working or heat treatment processes as the author has outlined above. Some acicular a with a thickness of about 0.1 mm are observed in transformed b matrix. The acicular a precipitates from transformed b matrix during the air-cooling process after isothermal compression. Moreover, some fine DRX a grains appear at this condition (Fig. 4(g)e(h)). DRX would also decrease the distorted energy and work as another softening mechanism. At a high strain rate, because there is insufficient time for the heat to escape from the specimen during deformation, the material can no longer be considered be isothermal, and the loss of strength caused by the increase in temperature will, at a certain point, exceed the increase in strength due to work-hardening [15].
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Fig. 3. Microstructural observations of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1163 K/1.0 102 s1: (a) TEM micrograph after deformed at a strain of 0.2; (b) is SAD pattern obtained from (a) (location A) with beam direction on ½1213a ; (c and d) TEM micrographs after deformed at a strain of 0.5; (e, f, g and h) TEM micrographs after deformed at a strain of 0.9.
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
49
Fig. 4. Microstructural observations of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1163 K/5.0 s1/0.9: (a and b) SEM micrographs in adjacent visual fields; (c, e, f, g and h) TEM micrographs; (d) is SAD pattern obtained from (c) (location A) with beam direction on ½1213a .The compression axis is vertical in SEM micrographs.
Thus, the flow curves show a sharp decrease beyond the peak
stress. In addition, the occurrence of DRV and DRX will enhance the
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softening effect during hot working of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy as above-mentioned analysis. So the thermal softening, the DRV and the DRX will result in that the flow curves at a higher strain rate drop continuously and the steady state conditions are not observed, even up to a strain of 0.9. 3.2.2. Deformation in the b phase region The effect of the strain rate on the microstructure of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at a deformation temperature of 1243 K and a strain of 0.9 is shown in Fig. 5. It is seen that the morphology of b phase exhibits two characteristics: fine recrystallized and coarse elongated b grains. The references [20,21] reported that a mixed b grain structure containing fine and coarse grains was the result of selective recrystallization in regions of high strain along beta grain boundaries during the b phase working and the reduced driving force for recrystallization in the grain interiors, where dynamic recovery dominated. This demonstrates that the DRV is still a type of softening mechanisms in the b phase region. In addition, the amount of recrystallized b grains increases with increasing strain rate, but apparently the size of recrystallized b grains tends to decrease. The main reason is that the occurrence of recrystallization nuclei is primarily dependent on the substructure cell size and the related cell-wall density [17]. More dislocations in cell walls appear at a higher strain rate compared to those at a lower strain rate as above-mentioned in Section 3.2.1, which promotes the occurrence of b recrystallization. However, there is insufficient time for the b grain growth at a higher strain rate, resulting in the small grain size at this strain rate. According to above-mentioned analysis, it is concluded that the DRV and the DRX of b phase are main softening mechanisms in the b phase region. Because the b phase has more slip systems than the a phase, the working hardening is not serious in the b phase region. This promotes that the dynamic softening is sufficient to counteract the work-hardening of the material at a lower strain. So the flow stress
in the b phase region appears to reach a steady state value at a small strain. 3.3. The strain rate sensitivity The strain rate sensitivity is usually used to determine the deformation mechanisms of a material. Values of the strain rate sensitivity are calculated using the following expression [22]:
m¼
d log s , d log ε
(1) ε;T
where s is the flow stress at a fixed strain and a fixed deformation : temperature (MPa), ε is the strain rate(s1), ε is the strain, and T is the absolute deformation temperature ( C). The strain rate sensitivity is estimated from the flow stressestrain curves during isothermal compression of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy and is shown in Fig. 6. From Fig. 6, it is observed that the values of the strain rate sensitivity reveal a noticeable dependence on the strain, the strain rate and the deformation temperature. At a deformation temperature of 1163 K and a strain rate of 0.1 s1 (Fig. 6(a)), the m values are in the range of 0.2e0.3 suggesting that the deformation mechanism is likely dislocation glide/climb [23]. And generally, m tends to increase with increasing strain, the value of m increases to 0.3 at a strain of 0.7. The microstructural observation at this strain is shown in Fig. 7(a). It is seen that the a grains are equiaxed and uniform, which is beneficial for the grain boundary sliding and accommodation. At a strain rate of 1.0 s1, the m values are approximately 0.2 or less, which are typical of deformation controlled by dislocation glide process characteristic. The a grains show two morphology characteristics: equiaxed and elongated, as illustrated in Fig. 7(b). The grain boundaries of elongated alpha are observed to become jagged and serrated at this strain rate. At a strain rate of 5.0 s1, the
Fig. 5. Optical micrographs of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at 1243 K/0.9: (a) 1.0 102 s1; (b) 0.1 s1; (c) 1.0 s1. The compression axis is vertical in all micrographs.
J. Luo et al. / Journal of Alloys and Compounds 667 (2016) 44e52
51
Fig. 6. Variation of the strain-rate sensitivity of the flow stress as a function of strain and strain rate at test temperatures of (a) 1163 K; (b) 1223 K.
Fig. 7. Optical micrographs of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy at a strain of 0.7: (a) 1163 K/0.1 s1; (b) 1163 K/1.0 s1; (c) 1163 K/5.0 s1; (d) 1223 K/0.1 s1. The compression axis is vertical in all micrographs.
curve of m tends to oscillatory, and the m values are below 0.1, at which the deformation also appears to be dominated by dislocation glide process. The microstructural observation at this strain rate is shown in Fig. 7(c). It is observed that the microstructure is nonuniform and has slight oriented characteristics. The elongated a grains are perpendicular to the compression axis direction. So the non-uniformity of microstructure will finally lead to that the m values are small. At a deformation temperature of 1223 K, the variations of m with strain rate are similar to those at a deformation temperature of 1163 K. But, the m values decreases compared to that at a lower deformation temperature, which is attributed to a sharp decrease of equiaxed a phase. According to previous superplastic theory [24], it was concluded that approximately equal volume fractions of a phase and b phase in the alloy contributed to
grain boundary sliding and accommodation. Thus only 8.0% volume fraction of a phase at 1223 K will finally result in the smaller m values, as illustrated in Fig. 7(d).
4. Conclusions The flow curves and the strain rate sensitivity of Tie6Ale2Zre2Sne2Moe1.5Cre2Nb alloy are analyzed with the help of the microstructural observations. The microstructural evolution and the softening mechanisms are investigated thoroughly in the aþb phase region and b phase region, respectively. The following conclusions are obtained from the present investigation.
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(1) At a high strain rates (5.0 s1) in the aþb phase region, more noticeable flow softening arises from the thermal softening, the dynamic recovery and the dynamic recrystallization of alpha phase. (2) At a low strain rate (0.01 s1) in the aþb phase region, the dynamic recovery structure including the lamellar alpha bending and the formation of sub-boundaries, and the globularization of lamellar alpha and the DRX for alpha phase will participate in the softening of materials. (3) In the b phase region, the working hardening is not serious, which will promote that the dynamic softening effect (the DRV and the DRX of b phase) is sufficient to counteract the work-hardening at a lower strain. So the flow stress in the b phase region appears to reach a steady state value at a small strain. (4) The maximum m value of 0.3 occurs at 1163 K/0.1 s1/0.7 during the isothermal compression of this alloy. The microstructure at this processing condition is equiaxed and uniform, which is beneficial for the grain boundary sliding and accommodation. Acknowledgments
[6]
[7]
[8]
[9]
[10] [11]
[12]
[13]
[14] [15] [16]
The authors thank the financial supports from the National Natural Science Foundation of China with Grant No. 51205318 and No. 51275416, the Innovation Fund of Science and Technology in Northwestern Polytechnical University with Grant No. 2012KJ02002.
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