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Dislocation evolution with creep strain and dislocation emission related with α2-phase dissolution
Materials Science and Engineering A239 – 240 (1997) 457 – 463
Dislocation evolution with creep strain and dislocation emission related with a2-phase dissolution Soo Woo Nam 1,a,*, Han Seo Cho a, Sun-Keun Hwang1, b, Nack J. Kim c a
Department of Materials Science and Engineering, Korea Ad6anced Institute of Science and Technology, 373 -1 Kusong-dong Yusong-gu, Taejon, 305 -701, South Korea b Department of Metall. Engineering, Inha Uni6ersity, Yonghyun-dong Nam-gu, Inchon, 402 -751, South Korea c Center for Ad6anced Aerospace Materials, Pohang Uni6. of Science and Technology, San 31 Hyoja-dong Nam-gu, Pohang, 790 -784, South Korea
Abstract Microstructural evolution of the lamellar structured Ti – 46.6Al – 1.4Mn – 2Mo (at.%) alloy (made by elemental powder metallurgy, EPM) during primary creep deformation under the condition of 800°C/200 MPa is investigated using transmission electron microscopy (TEM). As the primary creep deformation is progressing, the density of the initial dislocation is found to decrease very significantly. Two kinds of dislocation generation modes related to the lamellar interfaces were observed; cross slip and bowing of the interfacial dislocations. And it is suggested that the applied stress play a significant role in the dissolution of a2-phase, and the interfacial dislocation emission and dissolution of a2-phase are closely related with each other in controlling creep deformation rate. © 1997 Elsevier Science S.A. Keywords: Tensile creep; Titanium aluminide; Lamellar interface; Primary creep; Elemental powder metallurgy; Dislocation; a2-dissolution
1. Introduction Alloys of g-TiA1 have been developed for the use of structural components such as turbine blade of the advanced gas turbine jet engine, and engine valves in the automobile, which are operated under sustained stress and high temperature [1]. Considering the fact that the size tolerance in design of these kinds of structural components is usually limited within the primary creep strain under ordinary condition, it is very important to understand the primary creep deformation mechanism of these alloys. The reports concerning to the creep deformation mechanism of the g-TiAl alloys have been concentrated on the study of deformation mechanism through the measurement of the activation energy or the stress exponent using the steady state creep rates (or minimum creep rates) [2 – 6]. However, so far, there have been limited reports on the investigation of the primary creep behaviors of g-TiAl alloys [7,8]. Although the * Corresponding author. 1 Jointly Appointed at the Center for the Advanced Aerospace Materials 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 6 1 7 - 5
substructure in crept lamellar TiAl alloys has been reported [9], detailed mechanism of dislocation structure evolution with increasing creep strain in primary regime is not understood yet. The g-TiAl alloy used in this present study was processed by elemental powder metallurgy (EPM), hereafter, the test alloy is referred to as ‘EPM alloy’. The creep deformation mechanism of this material during secondary creep, and its fracture mode in air over the temperature range of 775–900°C under constant stress ranging from 150 to 250 MPa have been discussed elsewhere [10,11]. Authors of the present investigation had measured the apparent activation energy for the primary creep deformation of this alloy at the strain of 0.2, 0.5 and 0.8% and at the steady state [12,13]. The apparent activation energies measured for the different creep strain within the primary region were found to be increased slightly with creep strain from 340 kT mol − 1 at 0.2% creep strain and saturated to the value of steady state (1.7% creep strain), 380 kJ mol − 1, and it was suggested that the reduction of the effective stress with creep strain within the primary region is responsible for the increase of the apparent activation energy [12,13]. However, the substructural variation of disloca-
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2. Experimental
Fig. 1. The variation of the creep deformation rate with creep strain under the condition of 800°C/200 MPa.
tion structure with increased creep strain was not fully understood and so the microscopic creep deformation mechanism of lamellar TiAl alloy has not been investigated. The main purpose of the present study is to examine the evolution of dislocation structure with creep strain in primary creep regime and to reveal the dislocation generation mechanism related with a2-dissolution. Microscopic creep deformation mechanism in primary regime related with dislocation behavior examined in this study will be published later.
Ti–46.6Al–1.4Mn–2Mo (at.%) extruded bar of 20 mm in diameter, was stabilized at 1000°C for 20 h followed by lamellar forming heat treatment of 1380°C/ 30 min/air cooling and 950°C/5 h/air cooling. Detailed fabrication procedures for the extruded bar are described elsewhere [14]. The microstructure of the EPM alloy before creep test is lamellar structure with grain size of 150 mm [11]. The creep specimens of 3.2 mm in diameter and 15 mm in gauge length were machined from the extruded bar. Constant stress tensile creep tests were performed using the creep machine equipped with the constant stress Andrade-Chalmers cam. Linear variable differential transformer (LVDT) and LVDT signal conditioner were used to measure the creep strain and the displacement measurement accuracy was 5× 10 − 7 m. In order to observe the microstructural change with creep strain in primary region, five creep specimens were deformed up to the creep strain of 0.1, 0.2, 0.5, 0.8 and 1.7% (steady state), respectively, under the constant stress of 200 MPa at 800°C. The final TEM foils were prepared by the twinjet electropolishing with a solution of 90vol.% methanol+ 10vol.% perchloric acid at 90 mA and −45 − 55°C. TEM observation was conducted in Philips CM-20 and in JEOL 2000-EX microscope operating at 200 kV.
Fig. 2. Bright field image of the lamellar structure before creep test. The a2-lath is not perfectly straight but has a curved step.
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Fig. 3. Evolution of the dislocation structures with creep strain at the condition of 800°C/200 MPa. (a) Before creep test; (b) 0.1%; (c) 0.2%; (d) 0.5%; (e) 0.8% and (f) 1.7% creep strain (on set of the secondary stage).
3. Results and discussion The EPM alloy shows the normal primary creep behavior in which the creep deformation rate decreases with the creep strain. Fig. 1 represents the variation of the creep deformation rate with creep strain under the condition of 800°C/200 MPa. The open circle symbols in Fig. 1 indicate the creep strains at which TEM observations were conducted.
3.1. Shape of a2 -lath before creep test Because there is a orientation relationship between the a2 and g lath of (111)g //(0001)a2, [1( 10]g //112( 0a2, in lamellar grain [15], straight interface of a2/g lath is more stable than curved interface. However, it is observed that the most of the a2 laths are not perfectly straight and have curved step as shown in Fig. 2,
therefore, there must be abundant micro ledges at the a2/g interface. The role of these ledges are discussed in Section 3.3 in this report.
3.2. Dislocation e6olution with creep strain in primary creep regime The dislocation structure in the specimens deformed up to different amount of creep strains are represented in Fig. 3(a)–(f) to show the dislocation structure before creep test, after 0.1, 0.2, 0.5, 0.8 and 1.7% creep strain (on set of the secondary stage), respectively, under the condition of 800°C/200 MPa. Abundant initial matrix dislocations, like a dislocation debris in the g-lamellar lath, are observed before the creep test as in Fig. 3(a), and these dislocations are considered to be formed while air cooling of lamellar forming heat treatment. For the 0.1% crept specimen,
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though the dislocation structures are different in each lamellae, the density of the initial matrix dislocation is observed to be decreased drastically, and twinning, considered to be occurred during creep deformation, is observed at some lamellar grains as in Fig. 3(b). As the creep strain increases, almost all the initial matrix dislocations disappear, and the dislocation structure is changed to the shape of slightly curved array of dislocation as in Fig. 3(c) – (f). Moreover, the tangled dislocations or dislocations interacting with each other are not observed frequently in these crept specimens.
3.3. Dislocation generation and dissolution of a2 -lamellae during primary creep deformation It has been reported that the lamellar interface takes a significant role for the dislocation generation [16,17], Wunderlich et al. [16] suggested that the lamellar interface dislocations can be emitted by planar or cross slip mode. In addition, Appel et al. [17] suggested that the lamellar interface dislocations can be bowed out easily under the action of coherency stress at high temperature and, as another mechanism, the dipole arms trailed at the jogs of screw dislocations can act as single ended dislocation sources. And it is reported that there are four kinds of interfaces, except the lamellar grain boundary, in TiAl lamellar structure; a2/g-phase boundary, g-twins, pseudo twins between the g-laths and low angle grain boundary in g [16]. It is important to note that the a2/g-phase boundary takes the 60 70% of these four kinds of interfaces, and this percent-
Fig. 4. Dislocations which are considered to be generated by the cross slip mode at the specimen of 0.2% creep strain. Slip plane of the observed dislocations is determined to be (1( 11) and upper interface plane at which the g- meet the a2-lath is (11( 1)g.
age is slightly dependent on the alloy composition and the heat treatment [16]. Therefore, a2/g-phase boundary has the most significant effect on the dislocation generation mode among the dislocation generation mechanism related with lamellar interfaces. In this study, though the dipole arms trailed at the jog of screw dislocations suggested by Appel et al. [17] are not observed, dislocations which are considered to be generated by the cross slip mode are observed in 0.2% crept specimen as shown in Fig. 4. The dislocations of Fig. 4 (marked as 1, 2 and 3) are almost parallel each other, and are determined to have the same Burgers vector of 1/2[110] with mixed character and the Schmid factor of 0.30. In order to estimate the Schmid factor, the tensile direction was determined to be [10 12 9] by calculation with reasonable assumption that the TEM foil normal is parallel to the tensile axis of the creep test. By the cross product of the line direction of dislocation and the Burgers vector, the slip plane of these three dislocations are determined to be (1( 11). The upper interface between the a2- and g-lath is not straight but has rounded steps (arrow in Fig. 4). Meanwhile, the interface plane between the g- and a2-laths is determined to be (11( 1)g -plane, therefore, the Burgers vector of 1/2[110] is contained to the interface boundary, (11( 1)g. So considering that the Schmid factor of dislocations are somehow high in g-lath, it seems that the interfacial dislocations are emitted by cross slip mode as suggested by Wunderlich et al. [16] with the help of thermal energy and applied stress during creep test from the upper interface between a2 and g. It has been reported that the lamellar structure is not thermally stable, thus, coarsening of the lamellar structure occurs by discontinuous growth or/and continuous growth when lamellar TiA1 alloys are exposed to elevated temperature without any applied stress [18–21]. Specially, Ramanujan et al. [20,21] observed that some lamellar were dissolved after aging even at 800°C for 168 h, and dislocation networks are left after dissolution of the lamellae after aging at 1000°C for 168 h. In addition, the fact that the stress induced phase transformation from a2- to g-phase at room temperature could be occurred was reported by Zhang et al. [22]. Considering the mentioned above reports [18–22] dealing with the dissolution of lamellar structure at high temperature and stress induced phase transformation at room temperature, when the lamellar TiAl alloys are exposed to high temperature with applied stress as in creep test, it is supposed that dissolution (or phase transformation) of a2-phase to g-phase is easier than when the thermal energy or stress acts separately. And, if the generated dislocation following a2-dissolution process mentioned by Ramanuian et al. [20,21] has high enough Schmid factor for glide, gliding of the generated dislocation is considered to contribute the creep strain. However, in the other way, it is possible to consider that the emis-
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Fig. 5. TEM photographs at the 0.2% crept specimen. (a) Dislocations which are considered to be generated by the bowing process at the specimen of 0.2% creep strain. (b) Diffraction pattern of single g-phase taken at the center of the g-lath (marked ‘A’ in (a)).
sion of the interfacial dislocation can accelerate the dissolution of a2-phase because that the emission of interfacial dislocation accommodating the misfit strain in g/a2-interface decreases the stability of lamellar structure. Therefore, the ledge part (arrow mark in Fig. 4) can be considered as being dissolving activated by dislocation emission. Among the four kinds of continuous coarsening and dissolution mode suggested by Ramanujan et al. [20] such as: (1) boundary migration; (2)
Fig. 6. Schematic diagram for the emission of the interfacial dislocations to adjacent g-lath (a) before local phase transformation, (b) phase transformation assisted by dislocation generation at the ledge.
edge migration; (3) ledge migration; and (4) interface dissolution, the mode of the edge migration is considered to be under progressing in Fig. 4 (arrow mark). Even though it is not clear at present which process causes the other one between the dissolution of a2phase and the emission of the interfacial dislocation, it is assumed that the dissolution of a2-phase and the emission of the interfacial dislocation during creep test are closely related with each other. It is important to note that while it takes only 30 min to obtain 0.2% creep strain (during which a2-phase has been dissolved) in a creep test at 800°C/200 MPa, the dissolution of a2 only by aging at 800°C took 168 h [20,21]. Thus, it can be assumed that the applied stress plays a significant role in the dissolution of a2-phase and emission of the interfacial dislocation. Another dislocation generation mode observed in this study is a bowing of the interfacial dislocations shown
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Fig. 7. Schematic diagram of the sequential process of the dislocation generation by the bowing mechanism at the a2/g-interface and disappearing of the thin a2-lath.
combination of climbed dislocations of opposite signs at the beginning stage of primary creep deformation. At the same time, new dislocations are considered to be generated by a2-phase dissolution process occurred at the a2/g-interfaces and creep strain seems to proceed by gliding of these newly generated dislocations. Consequently, it is considered that the creep strain at the beginning of the primary stage is proceeded mainly by gliding of the initial dislocations and the newly generated dislocations take major role in later stage of creep straining. However, considering that most of the observed dislocations at the steady state condition (Fig. 2(f)) are pinned at the lamellar interfaces, the gliding of newly generated dislocations may be difficult.
4. Conclusion in Fig. 5(a). Fig. 5(b) taken at the center of the g-lath (marked ‘A’ in Fig. 5(a)) shows the diffraction pattern of 011]g, and indicates that the both sides of the bowed dislocations are single g-phase which does not have any other phases like a a2-phase. The two dislocations marked as ‘1’ and ‘2’ are dislocation segments bowed out from the center of the g-lamellar lath. Accepting the proposed assumption that the dissolution of a2-phase and the emission of the interfacial dislocation are closely related with each other, dislocation generation modes shown in Fig. 4 (cross slip) and Fig. 5 (bowing) can be explained as follows. The ledge of a2-lamellar migrates while emitting the interfacial dislocations by cross slip mode as in Fig. 6. Fig. 6 shows schematically the process of a2-dissolution by ledge migration while emitting dislocations. And, as mentioned in Section 3.1, the ledge is considered to exist before creep test. Fig. 7 represents the schematic diagram of the sequential process of the dislocation generation by the bowing mechanism at the a2/g-interface and disappearing of the thin a2-lath by continuous coarsening. Although, it is not clear which kind of continuous coarsening mode (or dissolution mode) was operated in Fig. 5, it seems that a thin a2-lath was at the middle of g-lath before creep test and this thin a2-lath was disappeared by continuous coarsening with dislocation emission. On the basis of the above mentioned argument, it can be stressed again that the dislocation generation and dissolution of a2-phase are closely related with each other.
3.4. Creep deformation procedure within primary regime From the observations, mentioned in Section 3.2, that the density of the initial dislocation is found to decrease very significantly as creep strain increases, it is assumed that the initial matrix dislocations are annihilated by gliding to the interface boundaries and/or
1. The initial dislocation density is decreased with creep strain within the primary region. 2. Two kinds of dislocation generation modes related with lamellar interfaces were observed—cross slip and bowing of the interfacial dislocations. 3. It is suggested that applied stress plays a significant role in the dissolution of a2-phase, and the emission of the interfacial dislocation and dissolution of a2-phase are closely related to each other. 4. It is considered that the creep strain at the beginning of primary stage is proceeded mainly by gliding of initial dislocation and gliding of newly generated dislocations. Finally, creep deformation is controlled by gliding of the newly generated dislocations.
Acknowledgements This work was supported by the Korea Science and Engineering Foundation through the Center for Advanced Aerospace Materials, POSTECH.
References [1] Y.-W. Kim, J. Met. 46 (1994) 30. [2] M. Es-Souni, A. Bartels, R. Wagner, Mat. Sci. Eng. A 171 (1993) 127. [3] M. Es-Souni, A. Bartels, R. Wagner, Acta Metall. Mater. 43 (1995) 153. [4] G.B. Viswanthan, V.K. Vasudevan, in: Y.-W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, Warrendale, PA, 1996, pp. 967 [5] R.W. Hayes, P.A. MaQuay, Scr. Metall. Mater. 30 (1994) 259. [6] J. Trinatafillou, J. Beddoes, L. Zhao, W. Wallace, Scr. Metall. Mater. 31 (1994) 1387. [7] R.W. Hayes, Scr. Metall. Mater. 29 (1993) 1229. [8] J.N. Wang, A.J. Schwartz, T.G. Nieh, D. Clemens, Mater. Sci. Eng. A 206 (1996) 63. [9] L.H. Hsiung, T.G. Nieh, Scr. Mater. 36 (1997) 323.
S. W. Nam et al. / Materials Science and Engineering A239–240 (1997) 457–463 [10] H.S. Cho, S.W. Nam, S.-K. Hwang, N.J. Kim, J. Korean Inst. Met. Mater. 34 (1996) 1299. [11] H.S. Cho, S.W. Nam, S.-K. Hwang, N.J. Kim, Scr. Mater. 36 (1997) 1295. [12] S.W. Nam, H.S. Cho, S.-K. Hwang, N.J. Kim, in: Y.C. Yoo (Ed.), Proc. of the 10th Conf. Mechanical Behaviors of Materials, Ansan, KIMM, 1996, pp. 385. [13] H.S. Cho, S.W. Nam, S.-K. Hwang, N.J. Kim, Mat. Sci. Res. Int. 3 (1997) 166. [14] I.-S. Lee, S.-K. Hwang, W.K. Park, J.H. Lee, D.H. Park, H.M. Kim, Y.T. Lee, Scr. Metall. Mater. 31 (1994) 57. [15] M. Yamaguchi, Y. Umakoshi, Prog. Mater. Sci. 34 (1990) 1. [16] W. Wunderlich, Th. Kremser, G. Frommeyer, Acta Metall.
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Mater. 41 (1993) 1791. [17] F. Appel, R Wagner, in: Y.-W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, Warrendale, PA, 1996, pp. 231. [18] H.H. Tian, Z. Huang, C.Q. Chen, Scr. Metall. Mater. 30 (1994) 165. [19] R.V. Ramanujan, Acta Metall. Mater. 42 (1996) 2775. [20] R.V. Ramanujan, P.J. Maziasz, C.T. Liu, Acta Metall. Mater. 44 (1996) 2611. [21] P.J. Maziasz, R.V. Ramanujan, C.T. Liu, J.L. Wright, Intermetallics 5 (1997) 83. [22] Y.G. Zhang, F.D. Tichelaar, F.W. Schapink, Q. Xu, C.Q. Chen, Scr. Metall. Mater. 32 (1995) 981.
Dislocation evolution with creep strain and dislocation emission related with a2-phase dissolution Soo Woo Nam 1,a,*, Han Seo Cho a, Sun-Keun Hwang1, b, Nack J. Kim c a
Department of Materials Science and Engineering, Korea Ad6anced Institute of Science and Technology, 373 -1 Kusong-dong Yusong-gu, Taejon, 305 -701, South Korea b Department of Metall. Engineering, Inha Uni6ersity, Yonghyun-dong Nam-gu, Inchon, 402 -751, South Korea c Center for Ad6anced Aerospace Materials, Pohang Uni6. of Science and Technology, San 31 Hyoja-dong Nam-gu, Pohang, 790 -784, South Korea
Abstract Microstructural evolution of the lamellar structured Ti – 46.6Al – 1.4Mn – 2Mo (at.%) alloy (made by elemental powder metallurgy, EPM) during primary creep deformation under the condition of 800°C/200 MPa is investigated using transmission electron microscopy (TEM). As the primary creep deformation is progressing, the density of the initial dislocation is found to decrease very significantly. Two kinds of dislocation generation modes related to the lamellar interfaces were observed; cross slip and bowing of the interfacial dislocations. And it is suggested that the applied stress play a significant role in the dissolution of a2-phase, and the interfacial dislocation emission and dissolution of a2-phase are closely related with each other in controlling creep deformation rate. © 1997 Elsevier Science S.A. Keywords: Tensile creep; Titanium aluminide; Lamellar interface; Primary creep; Elemental powder metallurgy; Dislocation; a2-dissolution
1. Introduction Alloys of g-TiA1 have been developed for the use of structural components such as turbine blade of the advanced gas turbine jet engine, and engine valves in the automobile, which are operated under sustained stress and high temperature [1]. Considering the fact that the size tolerance in design of these kinds of structural components is usually limited within the primary creep strain under ordinary condition, it is very important to understand the primary creep deformation mechanism of these alloys. The reports concerning to the creep deformation mechanism of the g-TiAl alloys have been concentrated on the study of deformation mechanism through the measurement of the activation energy or the stress exponent using the steady state creep rates (or minimum creep rates) [2 – 6]. However, so far, there have been limited reports on the investigation of the primary creep behaviors of g-TiAl alloys [7,8]. Although the * Corresponding author. 1 Jointly Appointed at the Center for the Advanced Aerospace Materials 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 6 1 7 - 5
substructure in crept lamellar TiAl alloys has been reported [9], detailed mechanism of dislocation structure evolution with increasing creep strain in primary regime is not understood yet. The g-TiAl alloy used in this present study was processed by elemental powder metallurgy (EPM), hereafter, the test alloy is referred to as ‘EPM alloy’. The creep deformation mechanism of this material during secondary creep, and its fracture mode in air over the temperature range of 775–900°C under constant stress ranging from 150 to 250 MPa have been discussed elsewhere [10,11]. Authors of the present investigation had measured the apparent activation energy for the primary creep deformation of this alloy at the strain of 0.2, 0.5 and 0.8% and at the steady state [12,13]. The apparent activation energies measured for the different creep strain within the primary region were found to be increased slightly with creep strain from 340 kT mol − 1 at 0.2% creep strain and saturated to the value of steady state (1.7% creep strain), 380 kJ mol − 1, and it was suggested that the reduction of the effective stress with creep strain within the primary region is responsible for the increase of the apparent activation energy [12,13]. However, the substructural variation of disloca-
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2. Experimental
Fig. 1. The variation of the creep deformation rate with creep strain under the condition of 800°C/200 MPa.
tion structure with increased creep strain was not fully understood and so the microscopic creep deformation mechanism of lamellar TiAl alloy has not been investigated. The main purpose of the present study is to examine the evolution of dislocation structure with creep strain in primary creep regime and to reveal the dislocation generation mechanism related with a2-dissolution. Microscopic creep deformation mechanism in primary regime related with dislocation behavior examined in this study will be published later.
Ti–46.6Al–1.4Mn–2Mo (at.%) extruded bar of 20 mm in diameter, was stabilized at 1000°C for 20 h followed by lamellar forming heat treatment of 1380°C/ 30 min/air cooling and 950°C/5 h/air cooling. Detailed fabrication procedures for the extruded bar are described elsewhere [14]. The microstructure of the EPM alloy before creep test is lamellar structure with grain size of 150 mm [11]. The creep specimens of 3.2 mm in diameter and 15 mm in gauge length were machined from the extruded bar. Constant stress tensile creep tests were performed using the creep machine equipped with the constant stress Andrade-Chalmers cam. Linear variable differential transformer (LVDT) and LVDT signal conditioner were used to measure the creep strain and the displacement measurement accuracy was 5× 10 − 7 m. In order to observe the microstructural change with creep strain in primary region, five creep specimens were deformed up to the creep strain of 0.1, 0.2, 0.5, 0.8 and 1.7% (steady state), respectively, under the constant stress of 200 MPa at 800°C. The final TEM foils were prepared by the twinjet electropolishing with a solution of 90vol.% methanol+ 10vol.% perchloric acid at 90 mA and −45 − 55°C. TEM observation was conducted in Philips CM-20 and in JEOL 2000-EX microscope operating at 200 kV.
Fig. 2. Bright field image of the lamellar structure before creep test. The a2-lath is not perfectly straight but has a curved step.
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Fig. 3. Evolution of the dislocation structures with creep strain at the condition of 800°C/200 MPa. (a) Before creep test; (b) 0.1%; (c) 0.2%; (d) 0.5%; (e) 0.8% and (f) 1.7% creep strain (on set of the secondary stage).
3. Results and discussion The EPM alloy shows the normal primary creep behavior in which the creep deformation rate decreases with the creep strain. Fig. 1 represents the variation of the creep deformation rate with creep strain under the condition of 800°C/200 MPa. The open circle symbols in Fig. 1 indicate the creep strains at which TEM observations were conducted.
3.1. Shape of a2 -lath before creep test Because there is a orientation relationship between the a2 and g lath of (111)g //(0001)a2, [1( 10]g //112( 0a2, in lamellar grain [15], straight interface of a2/g lath is more stable than curved interface. However, it is observed that the most of the a2 laths are not perfectly straight and have curved step as shown in Fig. 2,
therefore, there must be abundant micro ledges at the a2/g interface. The role of these ledges are discussed in Section 3.3 in this report.
3.2. Dislocation e6olution with creep strain in primary creep regime The dislocation structure in the specimens deformed up to different amount of creep strains are represented in Fig. 3(a)–(f) to show the dislocation structure before creep test, after 0.1, 0.2, 0.5, 0.8 and 1.7% creep strain (on set of the secondary stage), respectively, under the condition of 800°C/200 MPa. Abundant initial matrix dislocations, like a dislocation debris in the g-lamellar lath, are observed before the creep test as in Fig. 3(a), and these dislocations are considered to be formed while air cooling of lamellar forming heat treatment. For the 0.1% crept specimen,
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though the dislocation structures are different in each lamellae, the density of the initial matrix dislocation is observed to be decreased drastically, and twinning, considered to be occurred during creep deformation, is observed at some lamellar grains as in Fig. 3(b). As the creep strain increases, almost all the initial matrix dislocations disappear, and the dislocation structure is changed to the shape of slightly curved array of dislocation as in Fig. 3(c) – (f). Moreover, the tangled dislocations or dislocations interacting with each other are not observed frequently in these crept specimens.
3.3. Dislocation generation and dissolution of a2 -lamellae during primary creep deformation It has been reported that the lamellar interface takes a significant role for the dislocation generation [16,17], Wunderlich et al. [16] suggested that the lamellar interface dislocations can be emitted by planar or cross slip mode. In addition, Appel et al. [17] suggested that the lamellar interface dislocations can be bowed out easily under the action of coherency stress at high temperature and, as another mechanism, the dipole arms trailed at the jogs of screw dislocations can act as single ended dislocation sources. And it is reported that there are four kinds of interfaces, except the lamellar grain boundary, in TiAl lamellar structure; a2/g-phase boundary, g-twins, pseudo twins between the g-laths and low angle grain boundary in g [16]. It is important to note that the a2/g-phase boundary takes the 60 70% of these four kinds of interfaces, and this percent-
Fig. 4. Dislocations which are considered to be generated by the cross slip mode at the specimen of 0.2% creep strain. Slip plane of the observed dislocations is determined to be (1( 11) and upper interface plane at which the g- meet the a2-lath is (11( 1)g.
age is slightly dependent on the alloy composition and the heat treatment [16]. Therefore, a2/g-phase boundary has the most significant effect on the dislocation generation mode among the dislocation generation mechanism related with lamellar interfaces. In this study, though the dipole arms trailed at the jog of screw dislocations suggested by Appel et al. [17] are not observed, dislocations which are considered to be generated by the cross slip mode are observed in 0.2% crept specimen as shown in Fig. 4. The dislocations of Fig. 4 (marked as 1, 2 and 3) are almost parallel each other, and are determined to have the same Burgers vector of 1/2[110] with mixed character and the Schmid factor of 0.30. In order to estimate the Schmid factor, the tensile direction was determined to be [10 12 9] by calculation with reasonable assumption that the TEM foil normal is parallel to the tensile axis of the creep test. By the cross product of the line direction of dislocation and the Burgers vector, the slip plane of these three dislocations are determined to be (1( 11). The upper interface between the a2- and g-lath is not straight but has rounded steps (arrow in Fig. 4). Meanwhile, the interface plane between the g- and a2-laths is determined to be (11( 1)g -plane, therefore, the Burgers vector of 1/2[110] is contained to the interface boundary, (11( 1)g. So considering that the Schmid factor of dislocations are somehow high in g-lath, it seems that the interfacial dislocations are emitted by cross slip mode as suggested by Wunderlich et al. [16] with the help of thermal energy and applied stress during creep test from the upper interface between a2 and g. It has been reported that the lamellar structure is not thermally stable, thus, coarsening of the lamellar structure occurs by discontinuous growth or/and continuous growth when lamellar TiA1 alloys are exposed to elevated temperature without any applied stress [18–21]. Specially, Ramanujan et al. [20,21] observed that some lamellar were dissolved after aging even at 800°C for 168 h, and dislocation networks are left after dissolution of the lamellae after aging at 1000°C for 168 h. In addition, the fact that the stress induced phase transformation from a2- to g-phase at room temperature could be occurred was reported by Zhang et al. [22]. Considering the mentioned above reports [18–22] dealing with the dissolution of lamellar structure at high temperature and stress induced phase transformation at room temperature, when the lamellar TiAl alloys are exposed to high temperature with applied stress as in creep test, it is supposed that dissolution (or phase transformation) of a2-phase to g-phase is easier than when the thermal energy or stress acts separately. And, if the generated dislocation following a2-dissolution process mentioned by Ramanuian et al. [20,21] has high enough Schmid factor for glide, gliding of the generated dislocation is considered to contribute the creep strain. However, in the other way, it is possible to consider that the emis-
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Fig. 5. TEM photographs at the 0.2% crept specimen. (a) Dislocations which are considered to be generated by the bowing process at the specimen of 0.2% creep strain. (b) Diffraction pattern of single g-phase taken at the center of the g-lath (marked ‘A’ in (a)).
sion of the interfacial dislocation can accelerate the dissolution of a2-phase because that the emission of interfacial dislocation accommodating the misfit strain in g/a2-interface decreases the stability of lamellar structure. Therefore, the ledge part (arrow mark in Fig. 4) can be considered as being dissolving activated by dislocation emission. Among the four kinds of continuous coarsening and dissolution mode suggested by Ramanujan et al. [20] such as: (1) boundary migration; (2)
Fig. 6. Schematic diagram for the emission of the interfacial dislocations to adjacent g-lath (a) before local phase transformation, (b) phase transformation assisted by dislocation generation at the ledge.
edge migration; (3) ledge migration; and (4) interface dissolution, the mode of the edge migration is considered to be under progressing in Fig. 4 (arrow mark). Even though it is not clear at present which process causes the other one between the dissolution of a2phase and the emission of the interfacial dislocation, it is assumed that the dissolution of a2-phase and the emission of the interfacial dislocation during creep test are closely related with each other. It is important to note that while it takes only 30 min to obtain 0.2% creep strain (during which a2-phase has been dissolved) in a creep test at 800°C/200 MPa, the dissolution of a2 only by aging at 800°C took 168 h [20,21]. Thus, it can be assumed that the applied stress plays a significant role in the dissolution of a2-phase and emission of the interfacial dislocation. Another dislocation generation mode observed in this study is a bowing of the interfacial dislocations shown
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Fig. 7. Schematic diagram of the sequential process of the dislocation generation by the bowing mechanism at the a2/g-interface and disappearing of the thin a2-lath.
combination of climbed dislocations of opposite signs at the beginning stage of primary creep deformation. At the same time, new dislocations are considered to be generated by a2-phase dissolution process occurred at the a2/g-interfaces and creep strain seems to proceed by gliding of these newly generated dislocations. Consequently, it is considered that the creep strain at the beginning of the primary stage is proceeded mainly by gliding of the initial dislocations and the newly generated dislocations take major role in later stage of creep straining. However, considering that most of the observed dislocations at the steady state condition (Fig. 2(f)) are pinned at the lamellar interfaces, the gliding of newly generated dislocations may be difficult.
4. Conclusion in Fig. 5(a). Fig. 5(b) taken at the center of the g-lath (marked ‘A’ in Fig. 5(a)) shows the diffraction pattern of 011]g, and indicates that the both sides of the bowed dislocations are single g-phase which does not have any other phases like a a2-phase. The two dislocations marked as ‘1’ and ‘2’ are dislocation segments bowed out from the center of the g-lamellar lath. Accepting the proposed assumption that the dissolution of a2-phase and the emission of the interfacial dislocation are closely related with each other, dislocation generation modes shown in Fig. 4 (cross slip) and Fig. 5 (bowing) can be explained as follows. The ledge of a2-lamellar migrates while emitting the interfacial dislocations by cross slip mode as in Fig. 6. Fig. 6 shows schematically the process of a2-dissolution by ledge migration while emitting dislocations. And, as mentioned in Section 3.1, the ledge is considered to exist before creep test. Fig. 7 represents the schematic diagram of the sequential process of the dislocation generation by the bowing mechanism at the a2/g-interface and disappearing of the thin a2-lath by continuous coarsening. Although, it is not clear which kind of continuous coarsening mode (or dissolution mode) was operated in Fig. 5, it seems that a thin a2-lath was at the middle of g-lath before creep test and this thin a2-lath was disappeared by continuous coarsening with dislocation emission. On the basis of the above mentioned argument, it can be stressed again that the dislocation generation and dissolution of a2-phase are closely related with each other.
3.4. Creep deformation procedure within primary regime From the observations, mentioned in Section 3.2, that the density of the initial dislocation is found to decrease very significantly as creep strain increases, it is assumed that the initial matrix dislocations are annihilated by gliding to the interface boundaries and/or
1. The initial dislocation density is decreased with creep strain within the primary region. 2. Two kinds of dislocation generation modes related with lamellar interfaces were observed—cross slip and bowing of the interfacial dislocations. 3. It is suggested that applied stress plays a significant role in the dissolution of a2-phase, and the emission of the interfacial dislocation and dissolution of a2-phase are closely related to each other. 4. It is considered that the creep strain at the beginning of primary stage is proceeded mainly by gliding of initial dislocation and gliding of newly generated dislocations. Finally, creep deformation is controlled by gliding of the newly generated dislocations.
Acknowledgements This work was supported by the Korea Science and Engineering Foundation through the Center for Advanced Aerospace Materials, POSTECH.
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