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Thermal stability of TiAl-based intermetallic alloys subjected to high pressure torsion
Materials Science & Engineering A 651 (2016) 306–310
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of TiAl-based intermetallic alloys subjected to high pressure torsion M.A. Nikitina n, R.K. Islamgaliev, V.D. Sitdikov Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, K. Marx str. 12, Ufa 450000, Russia
art ic l e i nf o
a b s t r a c t
Article history: Received 17 August 2015 Received in revised form 30 October 2015 Accepted 31 October 2015 Available online 2 November 2015
This paper deals with a study of the structure and thermal stability of intermetallic TiAl-based alloys subjected to high-pressure torsion (HPT). The structural features of the HPT-processed samples have been studied by transmission electron microscopy and X-ray diffraction. In order to study the thermal stability of ultrafine-grained (UFG) structure, the dependence of microhardness on the annealing temperature in a temperature range of 150–1000 °С has been investigated. The effect of heat treatment on the grain structure stability and microhardness of the HPT-processed samples is discussed. & 2015 Elsevier B.V. All rights reserved.
Keywords: Electron microscopy Hardness measurement Intermetallic Bulk deformation Hardening
1. Introduction It is known that intermetallics based on the TiAl γ-phase are of special interest, because they are promising for application as lightweight and high-strength structural materials operating at temperatures of 600–800 °С. By the present time, it has been established that intermetallics consisting of a mixture of γ-TiAl and α2-Ti3Al possess the best combination of properties [1]. On their basis, intermetalic alloys additionally alloyed with niobium and molybdenum have been developed, called the TNM alloys [2,3]. Hot isothermal forging at high temperature of 1200–1300 °С has been developed for deformation of cast billets from these alloys, enabling the formation of a grain structure with a grain size of several tens of microns [4]. The TNM alloys can be subjected to heat treatment at temperatures of 900–1100 °С and deformation at temperatures of 800–1000 °С to produce a globular fine-grained structure [5,6]. Nevertheless, these approaches do not eliminate the drawbacks inherent in TiAl-based intermetallic alloys, a namely, brittleness at room temperature and high temperature of the brittle–ductile transition. On the other hand, it is known that the formation of ultrafinegrained (UFG) structure contributes to a decrease of the temperature of brittle–ductile transition in hard-to-deform materials [7,8]. For example, UFG samples of Ni3Al can exhibit the effect of n
Corresponding author. E-mail address: [email protected] (M.A. Nikitina).
http://dx.doi.org/10.1016/j.msea.2015.10.121 0921-5093/& 2015 Elsevier B.V. All rights reserved.
low-temperature superplasticity, retaining high values of ultimate tensile strength [9]. At the same time, new approaches have been developed recently, leading to the enhancement of ductility at room temperature in various UFG metals as result of the formation of a bimodal structure and an increase in the parameter of strain-rate sensitivity [10]. All these examples allow expecting the drawbacks inherent in the TNM intermetallics to be eliminated through the formation of a UFG structure. In recent years, to produce the UFG structure in metallic materials, various techniques of severe plastic deformation (SPD) have been developed, based on the application of large strains at high pressures and low homologous temperatures [11–13]. These techniques have become a widely-known approach to produce UFG metals and alloys, since they have a number of advantages over other ways of fabrication of bulk nanomaterials [14]. In particular, they enable the fabrication of non-porous, chemically pure ultrafine-grained samples from various materials for investigation of the structure and mechanical properties. Therefore, these approaches have been selected in the present work for producing the UFG structure in intermetallic TiAl-based alloys. It is known that the structure of TiAl-based alloys is characterized by a complex phase composition, including γ-TiAl with the crystalline lattice L1o, α2-Ti3Al with the lattice DO19 and βо with the lattice В2 [4]. The UFG structure formation can lead to a change in the phase composition because of enhanced grain boundary diffusion. It is also expected that the effects of ordering and precipitation of particles will inhibit the grain growth during
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heating. Therefore, a study of the structure and thermal stability of UFG samples from TiAl-based alloys after additional annealing at various temperatures arouses a special interest, and is the aim of the present work.
2. Experimental TiAl-based alloys alloyed with niobium, molybdenum or chromium were selected as initial material for investigations (Table 1). Prior to HPT processing, upsetting to 30% was conducted at a temperature of 850 °С. HPT processing of samples with an initial diameter of 10 mm and a thickness of 2.5 mm was conducted under a pressure of 0.5 GPa at a temperature of 850 °С using anvils from superalloys. Annealing of the HPT-processed samples was performed in a temperature range of 150–1000 °С, with a step of 50 °С for 30 min. Structural studies were performed using JEM-2100 transmission electron microscope. The average grain size D was calculated from the measurements of the diameters of at least 100 grains. Diffraction patterns were taken from an area of 2 mm2. Microhardness measurements were carried out at a load of 100 g for 10 s, using a Micromet-5101 tester equipped with software for image analysis. X-ray diffraction analysis (XRD) was conducted on a Rigaku Ultima IV diffractometer, with the goniometer focusing according to the Bragg–Brentano method, using copper filtered radiation. The images were taken at a voltage of 40 kV and a current of 40 mA. The wave length λКα1 ¼ 1.54060 Å was used for the calculation. The general view of the X-ray patterns was taken with a scanning step of 0.02° and an exposure time of 2 s in each point in the 2Θ angles range of 20–155°. The quantitative phase analysis, the estimation of the lattice parameter, size of coherent scattering domains and root-mean-square microdistortions were made using the PDXL software package (www.rigaku.com). All the investigations of the structure and mechanical properties were performed immediately in the regions located at a half radius distance of the HPT-processed samples.
3. Results and discussion Colonies with an average size of 30 mm, consisting of alternating lamellae of the γ- and α2-phases have been observed in both initial alloys (Fig. 1). At the boundaries of the colonies, aggregates of particles were found, representing particles of the βo-phase (bright contrast) and γ-phase (dark contrast) [4]. As a result of HPT processing, a bimodal structure consisting of equiaxed α2-phase grains with an average size of 250 nm was observed in alloy №1 (Fig. 2). Colonies of the γ- and α2-phase lamellae with an average interlamellar distance of 150–200 nm, whose volume fraction did not exceed 70%, were observed as well. UFG structure with an average α2-phase grain size of 300 nm was typical for microstructure of alloy №2 (Fig. 3a). However, 30% of the sample's surface consisted of colonies of the γ- and α2-phase lamellae with an interlamellar distance of 100–300 nm (Fig. 3b). It should be noted that the UFG structure in the TNM Table 1 Chemical composition of intermetallic TiAl-based alloys, at%, maximum.
1 2
Ti
Al
Nb
Mo
Cr
B
Base Base
42.3 44.97
3.98 1.28
0.96 –
– 1.7
0.1 0.2
Fig. 1. Structure of intermetallic alloys №1 in the initial state.
intermetallic alloys was already observed after upsetting to 60% at a temperature of 800–900 °С [5]. At that, the volume fraction of recrystallized grains after such treatment was within the range of 40–50%, while the remaining part was occupied by colonies with a lamellar structure [6]. In other works, hot deformation was conducted in a temperature range of 1250–1300 °С, as a result of which a fine-grained structure was produced, containing grains of different phases with an average size of 5–10 mm [4,15]. In our case, the application of HPT processing at a temperature of 850 °С contributed to increased volume fraction of the UFG structure up to 70%. The retention of the lamellar structure for the remaining parts of the HPT-processed samples may be associated with an unfavorable orientation of the slip planes in these lamellae due to the presence of a crystallographic texture. After annealing of alloy №1 at 800 °C the mean grain size in the UFG structure slightly increased from 250 nm to 350 nm (Fig. 4a). One can see also beginning of recrystallization on example of a grain with a size of 1 mm in the annealed sample consisted of subgrains with a size of 300 nm equal to a size of grains observed in alloy №2 directly after HPT processing (Fig. 4b). The study of the structure by X-ray diffraction demonstrated that in all the samples the phases γ-TiAl, α2-Ti3Al and β-Ti, typical for these alloys, were observed (Fig. 5). The quantitative phase analysis showed that in the initial state of alloy №1 the volume fraction of β-Ti is 5.3 times larger, α2-Ti3Al is 1.2 times larger, and γ-TiAl is 1.1 times smaller, than in alloy №2 (Table 2). As a result of HPT processing, both in alloy №1 and alloy №2 phase transformations take place, leading to increased volume fraction of the γTiAl and β-Ti phases. At the same time, HPT processing results in a reduction of the volume fraction of the α2-Ti3Al phase. In the tetragonal lattice of γ-TiAl, titanium atoms are surrounded by four atoms of aluminum, while in the hcp lattice of α2Ti3Al, the number of surrounding aluminum atoms equals 5. As a result of a reduction of the α2-Ti3Al phase after the increase in the volume fraction of the γ-TiAl phase, the released titanium atoms increase the volume fraction of the β-phase. Besides, as a result of the activation of diffusion processes during HPT processing, niobium atoms may interact with titanium atoms, which, as known, leads to the stabilization of the β-phase. Probably, this process leads to an increase of the β-Ti phase content in the alloys as a result of HPT processing. On the other hand, as can be seen from the X-ray diffraction patterns, HPT processing leads to the broadening of X-ray peaks, which is visible distinctly at the 2θ angle equal to 45° (two peaks become unresolved as a result of broadening). Analysis of the X-ray peak broadening within the Halder–Wagner approach [16] showed that HPT processing led to significant decrease of the coherent scattering domains (CSD) sizes in the γ-TiAl phase. At the same time, the root-mean-square microdistortions of the
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Fig. 2. Microstructure of intermetallic alloy №1 after HPT processing: (a, b) regions with equiaxed α2-phase grains; (c,d) bimodal structure; (a, c) bright field image; (b, d) dark field image.
crystalline lattice increased 1.5-fold to 2-fold. The indicated changes in the microstructure can be related to a strong grain refinement and an increased density of introduced defects. In the initial alloys №1 and №2, the microhardness had a value of 3700 and 3500 MPa, respectively. After annealing at temperatures of 300–650 °С, the increase of microhardness up to 4000 MPa was observed (Fig. 6). Obviously the main reason of the enhanced microhardness in the UFG samples is the grain refinement. In particular as shown in [17] yield strength of γ-based TiAl intermetallics increase with a decrease in a grain size in consistent with the Hall–Petch relationship. Additional hardenining can be introduced by reducing of average interlamellar distance to 100–300 nm in colonies of the γ- and α2-phase lamellae as result of HPT processing. One can assume that a decrease of interlammelar distance in γ-based TiAl intermetallics leads to additional restriction in dislocation movement by γ/α2 interfaces.
In the UFG samples of both alloys №1 and №2, higher microhardness values were revealed, 6000 MPa and 4900 MPa, respectively (Fig. 6), which became unchanged after annealing at temperatures of 800 °С. At the same time, after annealing at temperatures of 300 °С and 750 °С, local rises of microhardness were observed. The first peak at 300 °С is probably related to the growth of internal elastic stresses in the conditions of a thermally stable grain structure at this temperature. The second peak at the annealing temperatures of 750–800 °С is in the vicinity of the temperature of transition from a globular structure to a lamellar one [4], at which the content of the γ-phase increases. As shown in Fig. 4, a mean grain size of γ-phase in the samples annealed at 800 °C changed not so much. Therefore a change in the grain size cannot be responsible for the decreasing of microhardness after annealing at 800 °C. Since the microhardness of the γ-phase, equal to 5000 MPa, is much smaller as compared with the microhardness of the α2-
Fig. 3. Microstructure of intermetallic alloy №2 after HPT processing: (а) regions with equiaxed α2-phase grains; (b) colony of γ- and α2-phase lamellae.
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Fig. 4. Microstructure after HPT processing and additional annealing at 800 °C for 30 min: (a) alloy №1; (b) alloy №2. The arrows indicate to subgrains in the annealed sample with a size equal to a size of grains in the HPT sample before annealing.
Fig. 5. X-ray patterns of intermetallic alloys №1 and №2 after HPT.
4. Conclusions
Table 2 Microstructure parameters obtained by X-ray analysis. State
Volume fraction, %
AlTi Sample №1 initial
Ti3Al
β-Ti
Lattice parameter a and c, Ǻ
ε , % CSD size, nm
AlTi
AlTi
AlTi
0.13
112
0.26
21
0.14
112
68.6(11) 25.9(5) 5.52(18) 2.8383(8) 4.061(4)
Sample №1 after HPT
82.8(10)
10.0(3)
7.17(15) 2.8407(5) 4.059(2)
Sample №2 initial
78.0(16)
21.0(4)
1.03(8)
2.8344(4)
85.9(17)
11.5(4) 2.53(10)
2.8377(8)
The possibility of applying HPT processing at elevated temperatures for grain refinement in intermetallic TiAl-based alloys has been demonstrated. It has been found that as a result of HPT processing, a UFG structure is formed with a volume fraction of 70% and an average α2-phase grain size of 250–300 nm, while in the remaining parts of the samples colonies of the γ- and α2-phase lamellae are preserved, with an interlamellar distance in a range of 100–300 nm. It has been shown that the UFG structure formation contributes to a significant increase in the microhardness, by 40– 50% in comparison with the initial samples. The increased microhardness values are retained in a temperature range of 20–850 °С, which testifies to high thermal stability of the UFG structure.
Acknowledgments
4.0540(17) Sample №2 after HPT
Fig. 6. The dependence of the microhardness of intermetallic alloys on the annealing temperature for 30 min.
0.22
26
4.048(4)
phase ( 7500 MPa) [4], a change in the microhardness of the UFG alloy after annealing at a temperature above 800 °С may be associated with an increase in the γ-phase content. It should be noted that a high thermal stability of the UFG structure in intermetallics alloys was already observed in [15,18], where the important role of ordering of the crystalline lattice in inhibiting grain growth at high temperatures was discussed.
M.A. would like to acknowledge the support from the Russian Foundation for Basic Research under Project no. 15-08-06163. R.K. would like to acknowledge the support from the Russian Ministry of Education and science within the basic of the program for universities under Progect no. 2014/240.
References [1] D.G. Morris, M.A. Muñoz-Morris, Rev. Met. Madrid Vol. Extr. (2005) 498–501. [2] H. Clemens, S. Mayer, Adv. Eng. Mater. 15 (4) (2013) 191–215. [3] W. Wallgram, T. Schmölzer, L. Cha, G. Das, V. Güther, H. Clemens, Int. J. Mater. Res. 100 (8) (2009) 1021–1030. [4] M. Schloffer, F. Iqbal, H. Gabrisch, E. Schwaighofer, F.-P. Schimansky, S. Mayer, A. Stark, T. Lippmann, M. Goeken, F. Pyczak, H. Clemens, Intermetallics 22
310
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(2008) 231–240. [5] V. Imayev, R. Imayev, T. Khismatullin, V. Güther, W. Beck, H.-J. Fecht, Scr. Mater. 57 (3) (2007) 193–196. [6] A.V. Kuznetsov, G.S. Dyakonov, D.G. Shaysultanov, V.S. Sokolovsky, G. A. Salishchev, Belgorod State University Scientific Bulletin. Series: Mathematics & Physics, (2013) 132–141 , in Russian. [7] N. Tsuji, S. Okuno, Y. Koizumi, Y. Minamino, Mater. Trans. 45 (7) (2004) 2272–2281. [8] Y. Zhang, A.V. Ganeev, J.T. Wang, J.Q. Liu, I.V. Alexandrov, Mater. Sci. Eng. A 503 (1–2) (2009) 37–40. [9] R.K. Islamgaliev, R.Z. Valiev, R.S. Mishra, A.K. Mukherjee, Mat. Sci. Eng. A 304– 306 (2001) 206–210. [10] Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699–1707.
[11] [12] [13] [14] [15] [16] [17] [18]
R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Progr. Mater. Sci. 45 (2000) 103. R.Z. Valiev, T.G. Langdon, Prog. Mater. Sci. 51 (2006) 881. A.P. Zhilyaev, T.G. Langdon, Prog. Mater. Sci. 53 (2008) 893. L.S. Toth, Proceedings of international conference NanoSPD6. JOP Conf. Ser.: Mater. Sci. Eng., 63, 2014. V.M. Imaev, R.M. Imaev, T.I. Oleneva, T.G. Khismatullin, Phys. Met. Metallogr. 106 (6) (2008) 641–648. N.C. Halder, C.N.J. Wagner, Acta Cryst. 20 (1966) 312–313. M. Usta, H. Wolfe, D.J. Duquette, N.S. Stoloff, R.N. Wright, Mat. Sci. Eng. A 359 (2003) 168–177. C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, J. Mater. Sci. 43 (2008) 7264–7272.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of TiAl-based intermetallic alloys subjected to high pressure torsion M.A. Nikitina n, R.K. Islamgaliev, V.D. Sitdikov Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, K. Marx str. 12, Ufa 450000, Russia
art ic l e i nf o
a b s t r a c t
Article history: Received 17 August 2015 Received in revised form 30 October 2015 Accepted 31 October 2015 Available online 2 November 2015
This paper deals with a study of the structure and thermal stability of intermetallic TiAl-based alloys subjected to high-pressure torsion (HPT). The structural features of the HPT-processed samples have been studied by transmission electron microscopy and X-ray diffraction. In order to study the thermal stability of ultrafine-grained (UFG) structure, the dependence of microhardness on the annealing temperature in a temperature range of 150–1000 °С has been investigated. The effect of heat treatment on the grain structure stability and microhardness of the HPT-processed samples is discussed. & 2015 Elsevier B.V. All rights reserved.
Keywords: Electron microscopy Hardness measurement Intermetallic Bulk deformation Hardening
1. Introduction It is known that intermetallics based on the TiAl γ-phase are of special interest, because they are promising for application as lightweight and high-strength structural materials operating at temperatures of 600–800 °С. By the present time, it has been established that intermetallics consisting of a mixture of γ-TiAl and α2-Ti3Al possess the best combination of properties [1]. On their basis, intermetalic alloys additionally alloyed with niobium and molybdenum have been developed, called the TNM alloys [2,3]. Hot isothermal forging at high temperature of 1200–1300 °С has been developed for deformation of cast billets from these alloys, enabling the formation of a grain structure with a grain size of several tens of microns [4]. The TNM alloys can be subjected to heat treatment at temperatures of 900–1100 °С and deformation at temperatures of 800–1000 °С to produce a globular fine-grained structure [5,6]. Nevertheless, these approaches do not eliminate the drawbacks inherent in TiAl-based intermetallic alloys, a namely, brittleness at room temperature and high temperature of the brittle–ductile transition. On the other hand, it is known that the formation of ultrafinegrained (UFG) structure contributes to a decrease of the temperature of brittle–ductile transition in hard-to-deform materials [7,8]. For example, UFG samples of Ni3Al can exhibit the effect of n
Corresponding author. E-mail address: [email protected] (M.A. Nikitina).
http://dx.doi.org/10.1016/j.msea.2015.10.121 0921-5093/& 2015 Elsevier B.V. All rights reserved.
low-temperature superplasticity, retaining high values of ultimate tensile strength [9]. At the same time, new approaches have been developed recently, leading to the enhancement of ductility at room temperature in various UFG metals as result of the formation of a bimodal structure and an increase in the parameter of strain-rate sensitivity [10]. All these examples allow expecting the drawbacks inherent in the TNM intermetallics to be eliminated through the formation of a UFG structure. In recent years, to produce the UFG structure in metallic materials, various techniques of severe plastic deformation (SPD) have been developed, based on the application of large strains at high pressures and low homologous temperatures [11–13]. These techniques have become a widely-known approach to produce UFG metals and alloys, since they have a number of advantages over other ways of fabrication of bulk nanomaterials [14]. In particular, they enable the fabrication of non-porous, chemically pure ultrafine-grained samples from various materials for investigation of the structure and mechanical properties. Therefore, these approaches have been selected in the present work for producing the UFG structure in intermetallic TiAl-based alloys. It is known that the structure of TiAl-based alloys is characterized by a complex phase composition, including γ-TiAl with the crystalline lattice L1o, α2-Ti3Al with the lattice DO19 and βо with the lattice В2 [4]. The UFG structure formation can lead to a change in the phase composition because of enhanced grain boundary diffusion. It is also expected that the effects of ordering and precipitation of particles will inhibit the grain growth during
M.A. Nikitina et al. / Materials Science & Engineering A 651 (2016) 306–310
307
heating. Therefore, a study of the structure and thermal stability of UFG samples from TiAl-based alloys after additional annealing at various temperatures arouses a special interest, and is the aim of the present work.
2. Experimental TiAl-based alloys alloyed with niobium, molybdenum or chromium were selected as initial material for investigations (Table 1). Prior to HPT processing, upsetting to 30% was conducted at a temperature of 850 °С. HPT processing of samples with an initial diameter of 10 mm and a thickness of 2.5 mm was conducted under a pressure of 0.5 GPa at a temperature of 850 °С using anvils from superalloys. Annealing of the HPT-processed samples was performed in a temperature range of 150–1000 °С, with a step of 50 °С for 30 min. Structural studies were performed using JEM-2100 transmission electron microscope. The average grain size D was calculated from the measurements of the diameters of at least 100 grains. Diffraction patterns were taken from an area of 2 mm2. Microhardness measurements were carried out at a load of 100 g for 10 s, using a Micromet-5101 tester equipped with software for image analysis. X-ray diffraction analysis (XRD) was conducted on a Rigaku Ultima IV diffractometer, with the goniometer focusing according to the Bragg–Brentano method, using copper filtered radiation. The images were taken at a voltage of 40 kV and a current of 40 mA. The wave length λКα1 ¼ 1.54060 Å was used for the calculation. The general view of the X-ray patterns was taken with a scanning step of 0.02° and an exposure time of 2 s in each point in the 2Θ angles range of 20–155°. The quantitative phase analysis, the estimation of the lattice parameter, size of coherent scattering domains and root-mean-square microdistortions were made using the PDXL software package (www.rigaku.com). All the investigations of the structure and mechanical properties were performed immediately in the regions located at a half radius distance of the HPT-processed samples.
3. Results and discussion Colonies with an average size of 30 mm, consisting of alternating lamellae of the γ- and α2-phases have been observed in both initial alloys (Fig. 1). At the boundaries of the colonies, aggregates of particles were found, representing particles of the βo-phase (bright contrast) and γ-phase (dark contrast) [4]. As a result of HPT processing, a bimodal structure consisting of equiaxed α2-phase grains with an average size of 250 nm was observed in alloy №1 (Fig. 2). Colonies of the γ- and α2-phase lamellae with an average interlamellar distance of 150–200 nm, whose volume fraction did not exceed 70%, were observed as well. UFG structure with an average α2-phase grain size of 300 nm was typical for microstructure of alloy №2 (Fig. 3a). However, 30% of the sample's surface consisted of colonies of the γ- and α2-phase lamellae with an interlamellar distance of 100–300 nm (Fig. 3b). It should be noted that the UFG structure in the TNM Table 1 Chemical composition of intermetallic TiAl-based alloys, at%, maximum.
1 2
Ti
Al
Nb
Mo
Cr
B
Base Base
42.3 44.97
3.98 1.28
0.96 –
– 1.7
0.1 0.2
Fig. 1. Structure of intermetallic alloys №1 in the initial state.
intermetallic alloys was already observed after upsetting to 60% at a temperature of 800–900 °С [5]. At that, the volume fraction of recrystallized grains after such treatment was within the range of 40–50%, while the remaining part was occupied by colonies with a lamellar structure [6]. In other works, hot deformation was conducted in a temperature range of 1250–1300 °С, as a result of which a fine-grained structure was produced, containing grains of different phases with an average size of 5–10 mm [4,15]. In our case, the application of HPT processing at a temperature of 850 °С contributed to increased volume fraction of the UFG structure up to 70%. The retention of the lamellar structure for the remaining parts of the HPT-processed samples may be associated with an unfavorable orientation of the slip planes in these lamellae due to the presence of a crystallographic texture. After annealing of alloy №1 at 800 °C the mean grain size in the UFG structure slightly increased from 250 nm to 350 nm (Fig. 4a). One can see also beginning of recrystallization on example of a grain with a size of 1 mm in the annealed sample consisted of subgrains with a size of 300 nm equal to a size of grains observed in alloy №2 directly after HPT processing (Fig. 4b). The study of the structure by X-ray diffraction demonstrated that in all the samples the phases γ-TiAl, α2-Ti3Al and β-Ti, typical for these alloys, were observed (Fig. 5). The quantitative phase analysis showed that in the initial state of alloy №1 the volume fraction of β-Ti is 5.3 times larger, α2-Ti3Al is 1.2 times larger, and γ-TiAl is 1.1 times smaller, than in alloy №2 (Table 2). As a result of HPT processing, both in alloy №1 and alloy №2 phase transformations take place, leading to increased volume fraction of the γTiAl and β-Ti phases. At the same time, HPT processing results in a reduction of the volume fraction of the α2-Ti3Al phase. In the tetragonal lattice of γ-TiAl, titanium atoms are surrounded by four atoms of aluminum, while in the hcp lattice of α2Ti3Al, the number of surrounding aluminum atoms equals 5. As a result of a reduction of the α2-Ti3Al phase after the increase in the volume fraction of the γ-TiAl phase, the released titanium atoms increase the volume fraction of the β-phase. Besides, as a result of the activation of diffusion processes during HPT processing, niobium atoms may interact with titanium atoms, which, as known, leads to the stabilization of the β-phase. Probably, this process leads to an increase of the β-Ti phase content in the alloys as a result of HPT processing. On the other hand, as can be seen from the X-ray diffraction patterns, HPT processing leads to the broadening of X-ray peaks, which is visible distinctly at the 2θ angle equal to 45° (two peaks become unresolved as a result of broadening). Analysis of the X-ray peak broadening within the Halder–Wagner approach [16] showed that HPT processing led to significant decrease of the coherent scattering domains (CSD) sizes in the γ-TiAl phase. At the same time, the root-mean-square microdistortions of the
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Fig. 2. Microstructure of intermetallic alloy №1 after HPT processing: (a, b) regions with equiaxed α2-phase grains; (c,d) bimodal structure; (a, c) bright field image; (b, d) dark field image.
crystalline lattice increased 1.5-fold to 2-fold. The indicated changes in the microstructure can be related to a strong grain refinement and an increased density of introduced defects. In the initial alloys №1 and №2, the microhardness had a value of 3700 and 3500 MPa, respectively. After annealing at temperatures of 300–650 °С, the increase of microhardness up to 4000 MPa was observed (Fig. 6). Obviously the main reason of the enhanced microhardness in the UFG samples is the grain refinement. In particular as shown in [17] yield strength of γ-based TiAl intermetallics increase with a decrease in a grain size in consistent with the Hall–Petch relationship. Additional hardenining can be introduced by reducing of average interlamellar distance to 100–300 nm in colonies of the γ- and α2-phase lamellae as result of HPT processing. One can assume that a decrease of interlammelar distance in γ-based TiAl intermetallics leads to additional restriction in dislocation movement by γ/α2 interfaces.
In the UFG samples of both alloys №1 and №2, higher microhardness values were revealed, 6000 MPa and 4900 MPa, respectively (Fig. 6), which became unchanged after annealing at temperatures of 800 °С. At the same time, after annealing at temperatures of 300 °С and 750 °С, local rises of microhardness were observed. The first peak at 300 °С is probably related to the growth of internal elastic stresses in the conditions of a thermally stable grain structure at this temperature. The second peak at the annealing temperatures of 750–800 °С is in the vicinity of the temperature of transition from a globular structure to a lamellar one [4], at which the content of the γ-phase increases. As shown in Fig. 4, a mean grain size of γ-phase in the samples annealed at 800 °C changed not so much. Therefore a change in the grain size cannot be responsible for the decreasing of microhardness after annealing at 800 °C. Since the microhardness of the γ-phase, equal to 5000 MPa, is much smaller as compared with the microhardness of the α2-
Fig. 3. Microstructure of intermetallic alloy №2 after HPT processing: (а) regions with equiaxed α2-phase grains; (b) colony of γ- and α2-phase lamellae.
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Fig. 4. Microstructure after HPT processing and additional annealing at 800 °C for 30 min: (a) alloy №1; (b) alloy №2. The arrows indicate to subgrains in the annealed sample with a size equal to a size of grains in the HPT sample before annealing.
Fig. 5. X-ray patterns of intermetallic alloys №1 and №2 after HPT.
4. Conclusions
Table 2 Microstructure parameters obtained by X-ray analysis. State
Volume fraction, %
AlTi Sample №1 initial
Ti3Al
β-Ti
Lattice parameter a and c, Ǻ
ε , % CSD size, nm
AlTi
AlTi
AlTi
0.13
112
0.26
21
0.14
112
68.6(11) 25.9(5) 5.52(18) 2.8383(8) 4.061(4)
Sample №1 after HPT
82.8(10)
10.0(3)
7.17(15) 2.8407(5) 4.059(2)
Sample №2 initial
78.0(16)
21.0(4)
1.03(8)
2.8344(4)
85.9(17)
11.5(4) 2.53(10)
2.8377(8)
The possibility of applying HPT processing at elevated temperatures for grain refinement in intermetallic TiAl-based alloys has been demonstrated. It has been found that as a result of HPT processing, a UFG structure is formed with a volume fraction of 70% and an average α2-phase grain size of 250–300 nm, while in the remaining parts of the samples colonies of the γ- and α2-phase lamellae are preserved, with an interlamellar distance in a range of 100–300 nm. It has been shown that the UFG structure formation contributes to a significant increase in the microhardness, by 40– 50% in comparison with the initial samples. The increased microhardness values are retained in a temperature range of 20–850 °С, which testifies to high thermal stability of the UFG structure.
Acknowledgments
4.0540(17) Sample №2 after HPT
Fig. 6. The dependence of the microhardness of intermetallic alloys on the annealing temperature for 30 min.
0.22
26
4.048(4)
phase ( 7500 MPa) [4], a change in the microhardness of the UFG alloy after annealing at a temperature above 800 °С may be associated with an increase in the γ-phase content. It should be noted that a high thermal stability of the UFG structure in intermetallics alloys was already observed in [15,18], where the important role of ordering of the crystalline lattice in inhibiting grain growth at high temperatures was discussed.
M.A. would like to acknowledge the support from the Russian Foundation for Basic Research under Project no. 15-08-06163. R.K. would like to acknowledge the support from the Russian Ministry of Education and science within the basic of the program for universities under Progect no. 2014/240.
References [1] D.G. Morris, M.A. Muñoz-Morris, Rev. Met. Madrid Vol. Extr. (2005) 498–501. [2] H. Clemens, S. Mayer, Adv. Eng. Mater. 15 (4) (2013) 191–215. [3] W. Wallgram, T. Schmölzer, L. Cha, G. Das, V. Güther, H. Clemens, Int. J. Mater. Res. 100 (8) (2009) 1021–1030. [4] M. Schloffer, F. Iqbal, H. Gabrisch, E. Schwaighofer, F.-P. Schimansky, S. Mayer, A. Stark, T. Lippmann, M. Goeken, F. Pyczak, H. Clemens, Intermetallics 22
310
M.A. Nikitina et al. / Materials Science & Engineering A 651 (2016) 306–310
(2008) 231–240. [5] V. Imayev, R. Imayev, T. Khismatullin, V. Güther, W. Beck, H.-J. Fecht, Scr. Mater. 57 (3) (2007) 193–196. [6] A.V. Kuznetsov, G.S. Dyakonov, D.G. Shaysultanov, V.S. Sokolovsky, G. A. Salishchev, Belgorod State University Scientific Bulletin. Series: Mathematics & Physics, (2013) 132–141 , in Russian. [7] N. Tsuji, S. Okuno, Y. Koizumi, Y. Minamino, Mater. Trans. 45 (7) (2004) 2272–2281. [8] Y. Zhang, A.V. Ganeev, J.T. Wang, J.Q. Liu, I.V. Alexandrov, Mater. Sci. Eng. A 503 (1–2) (2009) 37–40. [9] R.K. Islamgaliev, R.Z. Valiev, R.S. Mishra, A.K. Mukherjee, Mat. Sci. Eng. A 304– 306 (2001) 206–210. [10] Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699–1707.
[11] [12] [13] [14] [15] [16] [17] [18]
R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Progr. Mater. Sci. 45 (2000) 103. R.Z. Valiev, T.G. Langdon, Prog. Mater. Sci. 51 (2006) 881. A.P. Zhilyaev, T.G. Langdon, Prog. Mater. Sci. 53 (2008) 893. L.S. Toth, Proceedings of international conference NanoSPD6. JOP Conf. Ser.: Mater. Sci. Eng., 63, 2014. V.M. Imaev, R.M. Imaev, T.I. Oleneva, T.G. Khismatullin, Phys. Met. Metallogr. 106 (6) (2008) 641–648. N.C. Halder, C.N.J. Wagner, Acta Cryst. 20 (1966) 312–313. M. Usta, H. Wolfe, D.J. Duquette, N.S. Stoloff, R.N. Wright, Mat. Sci. Eng. A 359 (2003) 168–177. C.C. Koch, R.O. Scattergood, K.A. Darling, J.E. Semones, J. Mater. Sci. 43 (2008) 7264–7272.