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Materials Science & Engineering A 767 (2019) 138393
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The characteristics and mechanisms of creep brittle-ductile transition in TiAl alloys
T
Qi Wanga, Ruirun Chena,b,∗, Dezhi Chena, Yanqing Sua, Hongsheng Dinga, Jingjie Guoa, Hengzhi Fua a b
National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, 150001, PR China State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: TiAl alloy Brittle-ductile transition Creep Stress exponent Activation energy
This article presents the performance and mechanisms of brittle-ductile transition during creep in Ti-46Al-8Nb alloys. Experimental results show that the brittle-ductile transition temperature (BDTT) was determined to be 760–780 °C in Ti-46Al-8Nb alloy, and the creep lifetime and creep strain obviously changed at BDTT. Major dislocation slip systems of α2 lamellae are activated above BDTT, which promote the plastic deformation of α2/γ lamellae and contribute to overall creep strain. The thermal activation is very active that dislocations can overcome obstacle under a small effective stress during creep above BDTT, and the deformation depends on dislocation slip assisted by the thermal activation during creep below BDTT. The apparent activation energy is 402 kJ/mol during creep above BDTT, as the creep is controlled by dislocation climb. The liner relationship between 1/T and ln(ε) breaks down during creep below BDTT, as the interface sliding mechanism being dominant in this regime. There is a significant change in apparent activation energy value at BDTT. Moreover, the BDTT of Ti-46Al-8Nb alloy is 60 °C higher than binary Ti-44Al alloy.
1. Introduction TiAl alloys have gotten a lot of attention because of their attractive elevated temperature properties combined with the low density; especially the high-Nb TiAl alloys are candidates to be used at temperature higher than 800 °C in the advanced aerospace and automotive engines [1–4]. Low-pressure turbine blades prepared from TiAl alloys have been manufactured for GEnx 2B (Boeing 747-8) and Enx 1B (Boeing 787) turbine engine applications [1,5]. However, the service temperature of TiAl alloy is commonly below brittle-ductile transition temperature (BDTT), and the TiAl alloys usually exhibit the plastic deformation characteristics when the deformation temperature exceeds the BDTT [6–11]. In the nearly/fully TiAl alloys, the deformation relies on the elastoplastic code formation of γ and α2 lamellae. The α2 phase with the D019 superlattice structure is brittle phase at elevated temperature, and the plastic deformation of lamellar colony at elevated temperature is mainly influenced by the α2 phase [12,13]. Imayev [14] has proposed that the brittle-ductile transition is a thermal activation process where the stress in matrix is relaxed through the plastic deformation. Therefore, the creep test is suitable for researching the brittle-ductile transition. TiAl alloys usually have a high
∗
creep strain rate during creep at elevated temperature, however, the overall creep strain allowable in the aero-engine turbine blades is very small [7,8]. Therefore, it is essential to investigate the performance and mechanisms of brittle-ductile transition during creep of TiAl alloy. Moreover, the deformation mechanisms change at the BDTT, and the apparent activation energy depends on the deformation mechanisms [15,16]. It is necessary to investigate the change mechanisms of activation energy at the BDTT during creep. These aspects are addressed in the present study; where the creep tests were conducted at different temperature and stress in Ti-46Al-8Nb and Ti-44Al alloys to determine the BDTT. Creep tests were carried out based on the tensile testing results at elevated temperature. The performance and mechanisms of creep brittle-ductile transition are investigated and discussed. An attention is given to the change mechanism of apparent activation energy at the BDTT. 2. Materials and methods The as-cast ingots with nominal composition of Ti-44Al and Ti-46Al8Nb were prepared using the vacuum consumable melting technology, and the ingot was melted two times for composition homogenization.
Corresponding author. National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, 150001, PR China. E-mail address: [email protected] (R. Chen).
https://doi.org/10.1016/j.msea.2019.138393 Received 4 April 2019; Received in revised form 4 September 2019; Accepted 5 September 2019 Available online 06 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 767 (2019) 138393
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phases.
Table 1 Chemical composition of the investigated alloys (at.%). Alloy
Al
Nb
Ti-44Al Ti-46Al-8Nb
44.32 45.52
— 8.25
3.2. Brittle-ductile transition Fig. 2(a) shows the statistical results of elongation and yield stress versus temperature of Ti-46Al-8Nb alloy. The elevated temperature yield stress is 438 MPa and the elongation is 2.37% at testing temperature of 780 °C. The elevated temperature yield stress is 344 MPa and the elongation is 10.77% at the testing temperature of 800 °C. The elevated temperature yield strength decreases by 21.5% and the elongation increases 4.15 times as the testing tempreature increasing from 780 to 800 °C, which indicates that the BDTT is at 780–800 °C. Based on the elevated temperature tensile testing results, the samples were creep tested under 276 MPa at 740–800 °C, as shown in Fig. 2(b). Creep tests were terminated either at a strain well into the secondary stage (740 °C) or at fracture (760–800 °C). This alloy exhibits a low minimum creep rate during creep at 740–760 °C, and the creep strain is 2.62% and creep lifetime is 529 h after creep fracture test at 760 °C. This alloy exhibits a large creep strain and a short creep lifetime during creep at 780–800 °C, and the creep strain is 17.71% and the creep lifetime is 125 h after creep fracture test at 780 °C. The creep strain increases 6.8 times with the testing tempreature increasing from 760 to 780 °C, which suggests the creep BDTT is at 760–780 °C. The minimum creep rate significantly increases with the testing tempreature increasing from 760 to 780 °C. Moreover, the creep BDTT is 20 °C lower than the BDTT at elevated temperature tensile, which is related to the different strain rates between elevated temperature tensile tests and creep tests. As seen in the following discussion.
The creep testing simples were cut from the as-cast ingots. The chemical compositions of Ti-44Al and Ti-46Al-8Nb alloys were measured using the PW4400 X-ray fluorescence spectrum analysis, and they were summarized in Table 1. The transmission electron microscopy (TEM) foils of the investigated alloys were cut to about 350 μm by the wire-electrode firstly, and then they were grinded to approximately 70 μm, followed by twinjet electro-polishing in a solution of 35% butyl alcohol, acid 60% methanol and 5% perchloric at −35 °C and 15 V. The microstructure before/after creep tests were observed by the Quanta 200FEG scanning electron microscopy in the back-scattered electron mode (SEM-BSE). For the further analysis of α2/γ lamellar microstructure, the microstructure after creep test was observed by the Tecnai G2 F30 transmission electron microscope (TEM). The high-temperature tensile specimens with a gauge length of 18 mm and cross-sections of 5 mm × 2 mm were tested at a strain rate of 1×10−4/s on the 5569 Instron testing machine, to determine the BDTT. The creep tests were carried out based on the high-temperature tensile testing results. The creep specimens with a gauge length of 20 mm and cross-sections of 4.5 mm × 2.5 mm were tested in the air under the constant load using the GWT504 creep tester. The thermoelectric couple was used to measure the creep temperature, and the two linear variable differential transducers were used to measure the displacement.
3.3. Microstructure deformation Fig. 3(a) shows the SEM-BSE microstructure of Ti-46Al-8Nb alloy after creep test at 760 °C (below the BDTT), and the black arrows mark the direction of stress axis. The crack at colony boundary is perpendicular to the direction of applied loading, and no obvious microstructure deformation can be found in α2/γ lamellae. Fig. 3(b) shows the SEMBSE microstructure of Ti-46Al-8Nb alloy after creep test at 780 °C (above the BDTT). The crack is observed at colony boundary and the red line marks the bowed lamellae, which suggests that the lamellae are obviously deformed during creep above the BDTT. Fig. 4 shows the TEM images of α2/γ lamellae after creep testing at 780 °C under 276 MPa. Many dislocations can be observed in α2 lamellae, which indicate that the major dislocation slip systems of α2 lamellae are thermally activated above the BDTT. In the nearly/fully lamellar TiAl alloys, the deformation of α2/γ lamellae is determined by α2 lamellae at elevated temperature, and the brittle-ductile transition is a thermally activated process [6,14]. In this study, when creep above the BDTT, the main dislocation slip systems of α2 lamellae are thermally activated, which promote the plastic deformation of α2/γ lamellae and contribute to the overall creep strain. The crack initiation and propagation along colony boundaries is the main creep failure mechanism at the creep temperature both below and above BDTT. When creep above BDTT, the significant deformation of microstructure reduces microstructural stability. The lamellar colonies on both sides of grain boundary usually have different lamellar directions in TiAl alloys, and the lamellar colonies with different lamellar directions usually have different deformation capacities during creep above the BDTT [20]. The plastic incompatibility of lamellar colonies improves the local stress concentration at colony boundaries, which promotes the formation of cavities and cracks at colony boundaries, and accelerates creep failure. Moreover, the rapid necking occurred in the specimen during creep above the BDTT; therefore, this alloy exhibits a short creep lifetime during creep above the BDTT. When the Ti-46Al8Nb alloy creep below the BDTT, the main dislocation slip systems of α2 lamellae cannot be thermally activated. The deformation of α2/γ lamellae is very small, and a little plastic incompatibility occurred in
3. Results and discussion 3.1. Initial microstructure Fig. 1 shows the initial SEM-BSE microstructure of Ti-46Al-8Nb alloy. The γ phases with dark contrast and B2 phases with bright contrast are observed in the matrix, as they are induced by the high Nb content in Ti-46Al-8Nb alloy. Chen et al. [17] have reported that the B2 phases formed by the phase transformation of β → α at the colony boundaries during cooling, or formed via the phase transition of α → γ + α2 + β in lamellar structures in the high Nb-containing TiAl alloys. However, the B2 phase content in this alloy is much lower than that in the low Al-containing high-Nb TiAl alloys [18,19], due to the Ti-46Al8Nb alloy with the high Al-containing can restrict the formation of B2
Fig. 1. The initial BSE-SEM microstructure of Ti-46Al-8Nb alloy. 2
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Fig. 2. Elevated temperature properties: (a) the elevated temperature tensile tests at 760–820 °C and (b) the tensile creep tests at 740–800 °C.
lamellar colonies. Moreover, the slow necking occurred in the specimen during creep above the BDTT; therefore, this alloy exhibits a high microstructure stability and a long creep lifetime during creep below the BDTT. 3.4. Stress exponent at the BDTT To compare the overall creep behaviors, the dependence of minimum stress (σ) and creep rate (εmin) can be formulated by the power-law creep equation:
Q εmin = A (σ )nexp ⎛− c ⎞ ⎝ RT ⎠
(1)
where, n corresponds to the stress exponent, A corresponds to the constant, Qc corresponds to the activation energy, T corresponds to the absolute temperature and R corresponds to the gas constant. The logarithm of the stress versus minimum strain rate at creep temperature of 760 and 780 °C, are plotted in Fig. 5. The stress exponent n is about 1.4 during creep at 760 °C under 246–306 MPa, which suggests that the creep is controlled by interface sliding [21]. The stress exponent is about 4.3 during creep under 246–306 MPa at 780 °C. The stress exponent at range of 3–5 suggests that the creep is controlled by dislocation-climb [22]. This result further validates that the BDTT is controlled by the main dislocation slip system of α2 lamellae.
Fig. 4. TEM bright-field images of α2/γ lamellae after creep testing at 780 °C under 276 MPa.
lamellae. The elevated temperature deformation process is mainly influenced by the distribution, characteristics, and intensity of obstacles. According to a report by Zhang et al. [23], the obstacles can hinder the dislocation glide in a couple of ways: (1) the long-range internal stress τi, the dislocations cannot overcome long-rang obstacles merely via the thermal activation. (2) The local short-range obstacles, the obstacles can be overcome mainly by thermal activation. When the applied stress is greater than internal stress τi, the τ can be divided into two parts:
3.5. The brittle-ductile transition mechanisms The main dislocation slip of α2 lamellae can be activated during creep above the BDTT under the action of applying stress and thermal activation. It is obvious that the creep is a thermally activated process, and the creep mechanisms change at the BDTT, and the deformation is mainly controlled by the dislocation slip system of α2 lamellae in α2/γ
τ = τi + τe
(2)
Here, the τe corresponds to the effective stress. Fig. 6 shows the resistance curve corresponding to a local obstacle. The y-axis corresponds to the force f(x) acting on the dislocation line.
Fig. 3. SEM-BSE microstructure of colony boundaries after creep tests at (a) 760 °C and (b) 780 °C. 3
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Fig. 5. The logarithm of the minimum strain rate versus stress, (a) at 760 °C and (b) 780 °C.
Here, the T0 corresponds to the critical temperature. When T > T0 (above the BDTT), τ = τi, the flow stress does not depend on strain rate ε or temperature T. The thermal activation is very active that the dislocations can overcome local obstacle with the help of a small effective stress. Therefore, the applied stress just requires to overcome the long-range internal stress. The main dislocation slip system of α2 lamellae can be thermally activated during creep above the BDTT under a low stress. When T ≤ T0 (below BDTT), the deformation is dependent on the dislocation slip assisted by the thermal activation, and the flow stress is dependent on both strain rate ε and temperature Τ. Therefore, the main dislocation slip of α2 lamellae can be activated under very a high-applied stress and it cannot be activated during creep at 760 °C under 276 MPa. The BDTT measured by creep tests is 20 °C lower than that measured by elevated temperature tensile tests, as shown in Fig. 2. According to Eq. (7), the activation frequency at T0 matches the dislocation velocity required for deforming TiAl alloys at a strain rate of ε. The thermal activation frequency must be increased by the increasing temperature during TiAl alloys deformation at a higher strain rate. In this case, the BDTT depends on strain rate and the higher strain rate leads to a higher BDTT. The strain rate during creep tests is much lower than the strain rate during elevated temperature tensile tests, therefore, the BDTT measured by creep testing is lower than that measured by elevated temperature tensile testing.
Fig. 6. Resistance curve of an obstacle.
The total area below the resistance curve corresponds to the total energy required for a dislocation to overcome a local obstacle; and the shaded region corresponds to the energy supplied by thermal activation. The region below the shaded region corresponds to the energy supplied by effective stress. Therefore, the variation of Gibbs free energy during thermally activated process can be expressed as: x2
Δ G=
∫ [f (x ) − τbl] dx = ΔF − τe bΔa x1
(3)
x2
Here, the
∫ [f (x ) − τi bl] dx = ΔF corresponds to the total energy,
3.6. Apparent activation energy at the BDTT
x1
the Δa = lΔx corresponds to the region swept by the dislocation slip under the action of the thermal activation. The frequency of dislocations overcoming the local obstacle can be expressed as:
ΔG Δ F− τe bΔa ⎞ ⎞ = vo exp ⎛− v= vo exp ⎛− kT ⎝ kT ⎠ ⎠ ⎝
According to equation (1), the apparent activation energy can be described as:
d (lnε ) ⎤ Q c = −R ⎡ ⎢ d ⎣ (1/ T ) ⎥ ⎦σ / E , S
(4)
Here, the vo corresponds to the vibration frequency of dislocations. The relationship between dislocation slip and macroscopic deformation follows the Orowan equation:
ε0 = ρm bv
Here, σ/E corresponds to the elastic modulus normalized stress and S corresponds to the dislocation structure. The apparent activation energy Qc for Ti-46Al-8Nb alloy is calculated by the plotting ln(ε) and 1/T, as shown in Fig. 7, where the apparent activation energy was calculated through the slope of each line. From this figure, a decided change in value of apparent activation energy can be found (ranging from 205 to 653 kJ/mol). It can be seen from Fig. 7 that the liner relationship between 1/T and ln(ε) breaks down when the temperature is below BDTT. Previous studies have shown that the power law breakdown usually took place when the creep temperature was below 0.5 or 0.6 Tm (Tm corresponds to the melting temperature) [23]. The activation energy for self-diffusion is equal to the activation energy for creep when the creep temperature is above 0.5 or 0.6 Tm, and the activation energy for creep is below the activation energy for self-diffusion when the creep temperature is below 0.5 or 0.6 Tm [23–26]. The BDTT of TiAl alloys is usually at the temperature range of 0.5–0.6 Tm [6,9,14]. Lapin et al.
(5)
Here, the ρm corresponds to the density of mobile dislocation; the ν corresponds to the mean velocity of dislocation slip. According to Eqs. (3) and (4), the strain rate supplied by dislocations slip under the action of thermally activated as follows:
ΔG ⎞ ΔF − τe bΔa ⎞ ε = ε0 exp ⎛− = ε0 exp ⎛− kT ⎝ kT ⎠ ⎠ ⎝
(6)
Here ε0 = bρλv0, the λ corresponds to the distance between two adjacent obstacles. The flow stress as follows:
τ=
ΔF − kTln (ε0 / ε ) bΔa
⎧ τi + ⎨ τi ⎩
T ≤ T0
T > T0
(8)
(7) 4
Materials Science & Engineering A 767 (2019) 138393
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relationship between ln(ε) and 1/T, at constant stresses in Ti-44Al alloy, and the slope of the lines represents the apparent activation energy value Qc. There is a significant change in the apparent activation energy value at creep temperature of 700–720 °C. The 1/T and ln(ε) have a linear relation during creep above 720 °C. The apparent activation energy value at 660–700 °C is below the apparent activation energy value at 720–760 °C. The results indicate that the BDTT of Ti-44Al alloy is at 700–720 °C. The BDTT of Ti-46Al-8Nb alloy (760–780 °C) is 60 °C higher than Ti44Al alloy (700-700 °C), which well agrees with the previous results that the creep strength retention of high-Nb TiAl alloy is 60–100 °C higher than that of the binary TiAl alloys [31]. The activation parameters of elevated temperature deformation have been determined in both binary and the equivalent Al-containing high-Nb ternary TiAl alloys in previous studies [32]. These results agree with the researches in Nb diffusion in TiAl alloys, and imply that it is more difficult for the diffusion-assisted deformation in high-Nb TiAl alloys, which benefits the creep resistance and elevated temperature strength [33].
Fig. 7. The relationship between ln(ε) and 1/T, at constant stresses in Ti-46Al8Nb alloy. The slope of the lines represents the apparent activation energy value of QC.
4. Conclusions 1. The BDTT is determined to be 780–800 °C by high-temperature tensile tests and 760–780 °C by creep tests. The creep strain and creep life significantly change during creep at the BDTT. 2. The main dislocation slip systems of α2 lamellae are thermally activated above the BDTT, which promotes the plastic deformation of α2/γ lamellae and improves the overall creep strain. The plastic incompatibility of lamellar colonies promotes the formation of cavities and cracks at colony boundaries and accelerates creep failure. 3. When creep above the BDTT, the thermal activation is very active that dislocations can overcome local obstacle under acting a small effective stress, and the applied stress only needs to overcome longrange internal stress. When creep below the BDTT, the deformation is dependent on dislocation slip assisted by thermal activation, and the flow stress is dependent on both strain rate and temperature. Therefore, the main dislocation slip system α2 lamellae can be activated under very a high-applied stress. 4. The 1/T and ln(ε) have a linear relation during creep above the BDTT, due to the creep is controlled by the dislocation climb, and the apparent activation energy is 402 kJ/mol. The liner relationship between 1/T and ln(ε) breaks down during creep below the BDTT, due to the interface sliding mechanism being dominant in this regime. Moreover, there is a significant change in the apparent activation energy value at the BDTT, since the main dislocation slip system of α2 lamellae is thermally activated at the BDTT.
Fig. 8. The relationship between (lnε) and 1/T, at constant stresses in Ti-44Al alloy. The slope of the lines represents the apparent activation energy value of Qc.
[25,26] have reported that the liquidus temperature of Ti-46Al-8Nb alloy is 1570 °C. The BDTT measured by creep tests at the temperature range of 780–800 °C is at 0.5–0.6 Tm in this study [7,17]. According to Section 3.4, the creep is controlled by dislocation climb above the BDTT, and the creep is controlled by interface sliding below the BDTT. The 1/T and ln(ε) have a linear relation when the creep temperature is above the BDTT, combining the results of Section 3.4, which suggests that the of 1/T and ln(ε) have a linear relation when creep is controlled by dislocation climb, and the apparent activation energy is 402 kJ/mol. This result agrees to those of previous studies [27–30]. The liner relationship between 1/T and ln(ε) breaks down when the temperature is below the BDTT, combining the results of Section 3.4, which suggests that the liner relationship between 1/T and ln(ε) breaks down when the interface sliding mechanism being dominant in this regime. Moreover, there is a significant change in apparent activation energy value at the BDTT, since the main dislocation slip system of α2 lamellae is thermally activated at the BDTT, which leads to a significant increase in the creep rate and a large apparent activation energy value at the BDTT. It is obvious that the value of apparent activation energy depends on the creep mechanisms. The 1/T and ln(ε) have a linear relation during creep above the BDTT, there is a significant change in the apparent activation energy at the BDTT, and the apparent activation energy value is small during creep below the BDTT. Therefore, the apparent activation energy value can be used to determine the BDTT, and the apparent activation energy value is calculated by the minimum strain rate. TiAl alloys usually exhibit a short primary creep stage and rapidly decrease to a minimum value, therefore, the BDTT measured by the minimum strain rate can reduce creep testing time. Fig. 8 shows the
Acknowledgements The authors gratefully acknowledge support from the National Key Research and Development Programme of China (2017YFA0403804), China Postdoctoral Science Foundation funded project (2019TQ0076) and the National Natural Science Foundation of China Nos. 51741404. References [1] G. Chen, Y. Peng, G. Zheng, Z. Qi, M. Wang, H. Yu, C. Dong, G.T. Liu, Polysynthetic twinned TiAl single crystals for high-temperature applications, Nat. Mater. 15 (2016) 876. [2] Q. Wang, R.R. Chen, X. Gong, J.J. Guo, Y.Q. Su, H.S. Ding, H.Z. Fu, Microstructure, mechanical properties and crack propagation behaviour in high-Nb TiAl alloys by directional solidification, Metall. Mater. Trans. A 49A (2018) 4555–4564. [3] B. Cao, J.J. Yang, X.Y. Wang, Y.L. Wu, R. Hu, Microstructure evolution and mechanical properties of a Ti-45Al-8.5Nb-(W, B, Y) alloy obtained by controlled cooling from a single β region, J. Alloy. Comp. 740 (2017) 1140–1148. [4] Z. Wu, R. Hu, T. Zhang, H. Zhou, J. Li, Moisture induced environmental embrittlement of a high Nb containing TiAl alloy, Mater. Sci. Eng. A 701 (2017) 214–220. [5] H. Clemens, S. Mayer, Intermetallic titanium aluminides in aerospace applications-
5
Materials Science & Engineering A 767 (2019) 138393
Q. Wang, et al.
processing, microstructure and properties, Mater, High Temp. 33 (2016) 1–11. [6] R.R. Chen, Q. Wang, Y.H. Yang, J.J. Guo, Y.Q. Su, H.S. Ding, H.Z. Fu, Brittle–ductile transition during creep in nearly and fully lamellar high-Nb TiAl alloys, Intermetallics 93 (2018) 47–54. [7] F. Appel, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminide Alloys: Science and Technology, Wiley-VCH, 2011. [8] J. Lapin, T. Pelachová, M. Dománková, Long-term creep behaviour of cast TiAl-Ta alloy, Intermetallics 95 (2018) 24–32. [9] Y.W. Kim, S.L. Kim, Effects of microstructure and C and Si additions on elevated temperature creep and fatigue of gamma TiAl alloys, Intermetallics 53 (2014) 92–101. [10] M. Kastenhuber, B. Rashkova, H. Clemens, S. Mayer, Effect of microstructural instability on the creep resistance of an advanced intermetallic γ-TiAl based alloy, Intermetallics 80 (2017) 1–9. [11] M. Kastenhuber, T. Klein, H. Clemens, S. Mayer, Tailoring microstructure and chemical composition of advanced γ-TiAl based alloys for improved creep resistance, Intermetallics 97 (2018) 27–33. [12] S.Z. Zhang, C.J. Zhang, Z.X. Du, Z.P. Hou, P. Lin, F.T. Kongd, Y.Y. Chen, Deformation behavior of high Nb containing TiAl based alloy in α + γ two phase field region, Mater. Des. 90 (2016) 225–229. [13] L. Cheng, J. Li, X. Xue, B. Tang, H. Kou, E. Bouzy, Superplastic deformation mechanisms of high Nb containing TiAl alloy with (α2 +γ) microstructure, Intermetallics 75 (2016) 62–71. [14] V.M. Imayev, R.M. Imayev, G.A. Salishchev, On two stages of brittle-to-ductile transition in TiAl intermetallic, Intermetallics 8 (2000) 1–6. [15] D. Dorner, K. Röller, B. Skrotzki, B. Stöckhert, G. Eggeler, Creep of a TiAl alloy: a comparison of indentation and tensile testing, Mater. Sci. Eng. A 357 (2003) 346–354. [16] W. Wang, W. Zeng, C. Xue, X. Liang, J. Zhang, Microstructural evolution, creep, and tensile behavior of a Ti–22Al–25Nb (at%) orthorhombic alloy, Mater. Sci. Eng. A 603 (2014) 176–184. [17] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Microsegregation in high Nb containing TiAl alloy ingots beyond laboratory scale, Intermetallics 15 (2007) 625–631. [18] B. Tang, L. Cheng, H. Kou, J. Li, Hot forging design and microstructure evolution of a high Nb containing TiAl alloy, Intermetallics 58 (2015) 7–14. [19] J.R. Yang, X.Y. Wang, B. Cao, Y.L. Wu, K.R. Zhang, R. Hu, Tailoring the
[20] [21] [22]
[23] [24]
[25]
[26]
[27] [28] [29]
[30]
[31]
[32] [33]
6
microstructure of a β-Solidifying TiAl alloy by controlled post-solidification isothermal holding and cooling, Metall. Mater. Trans. A 48 (2017) 1–11. X.W. Du, J. Zhu, Y.W. Kim, Microstructural characterization of creep cavitation in a fully-lamellar TiAl alloy, Intermetallics 9 (2001) 137–146. A.M. Hodge, L.M. Hsiung, T.G. Nieh, Creep of nearly lamellar TiAl alloy containing W, Scr. Mater. 51 (2004) 411–415. S. Bystrzanowski, A. Bartels, H. Clemens, R. Gerling, F.P. Schimansky, G. Dehmd, H. Kestler, Creep behaviour and related high temperature microstructural stability of Ti-46Al-9Nb sheet material, Intermetallics 13 (2005) 515–524. J.S. Zhang, Elevated Temperature Deformation and Fracture of Materials, (2010). M. Sujata, D.H. Sastry, C. Ramachandra, Microstructural characterization and creep behaviour of as-cast titanium aluminide Ti-48Al-2V, Intermetallics 12 (2004) 691–697. Z. Gabalcová, J. Lapin, Estimation of high temperature phase equilibria in directionally solidified intermetallic Ti-45.9Al-8Nb alloy, Kovove Mater 45 (2007) 231–240. Z. Gabalcová, J. Lapin, T. Pelachová, Effect of Y2O3 crucible on contamination of directionally solidified intermetallic Ti-46Al-8Nb alloy, Intermetallics 19 (2011) 396–403. L.M. Hsiung, Interfacial control of creep deformation in ultrafine lamellar TiAl, Mrs Proc. (2002) 740. Y. Ishikawa, K. Maruyama, H. Oikawa, Three regions in the power-law creep regime of TiAl, Mater. Trans. 33 (2007) 1182–1184. M.E. Kassner, The rate dependence and microstructure of high-purity silver deformed to large strains between 0.16 and 0.30Tm, Metall. Trans. 20A (1989) 2001–2010. H. Luthy, A.K. Miller, O.D. Sherby, The stress and temperature dependence of steady state flow at intermediate temperature for polycrystalline aluminum, Acta Metall. 28 (1980) 169–178. Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Effects of Nb and Al on the microstructures and mechanical properties of high Nb containing TiAl base alloys, Intermetallics 10 (2002) 653–659. Y. Mishin Y, C. Herzig, Diffusion in the Ti–Al system, Acta Mater. 48 (2000) 589–623. F. Appel, U. Lorenz, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminides 1999, TMS, (1999), p. 381.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The characteristics and mechanisms of creep brittle-ductile transition in TiAl alloys
T
Qi Wanga, Ruirun Chena,b,∗, Dezhi Chena, Yanqing Sua, Hongsheng Dinga, Jingjie Guoa, Hengzhi Fua a b
National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, 150001, PR China State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: TiAl alloy Brittle-ductile transition Creep Stress exponent Activation energy
This article presents the performance and mechanisms of brittle-ductile transition during creep in Ti-46Al-8Nb alloys. Experimental results show that the brittle-ductile transition temperature (BDTT) was determined to be 760–780 °C in Ti-46Al-8Nb alloy, and the creep lifetime and creep strain obviously changed at BDTT. Major dislocation slip systems of α2 lamellae are activated above BDTT, which promote the plastic deformation of α2/γ lamellae and contribute to overall creep strain. The thermal activation is very active that dislocations can overcome obstacle under a small effective stress during creep above BDTT, and the deformation depends on dislocation slip assisted by the thermal activation during creep below BDTT. The apparent activation energy is 402 kJ/mol during creep above BDTT, as the creep is controlled by dislocation climb. The liner relationship between 1/T and ln(ε) breaks down during creep below BDTT, as the interface sliding mechanism being dominant in this regime. There is a significant change in apparent activation energy value at BDTT. Moreover, the BDTT of Ti-46Al-8Nb alloy is 60 °C higher than binary Ti-44Al alloy.
1. Introduction TiAl alloys have gotten a lot of attention because of their attractive elevated temperature properties combined with the low density; especially the high-Nb TiAl alloys are candidates to be used at temperature higher than 800 °C in the advanced aerospace and automotive engines [1–4]. Low-pressure turbine blades prepared from TiAl alloys have been manufactured for GEnx 2B (Boeing 747-8) and Enx 1B (Boeing 787) turbine engine applications [1,5]. However, the service temperature of TiAl alloy is commonly below brittle-ductile transition temperature (BDTT), and the TiAl alloys usually exhibit the plastic deformation characteristics when the deformation temperature exceeds the BDTT [6–11]. In the nearly/fully TiAl alloys, the deformation relies on the elastoplastic code formation of γ and α2 lamellae. The α2 phase with the D019 superlattice structure is brittle phase at elevated temperature, and the plastic deformation of lamellar colony at elevated temperature is mainly influenced by the α2 phase [12,13]. Imayev [14] has proposed that the brittle-ductile transition is a thermal activation process where the stress in matrix is relaxed through the plastic deformation. Therefore, the creep test is suitable for researching the brittle-ductile transition. TiAl alloys usually have a high
∗
creep strain rate during creep at elevated temperature, however, the overall creep strain allowable in the aero-engine turbine blades is very small [7,8]. Therefore, it is essential to investigate the performance and mechanisms of brittle-ductile transition during creep of TiAl alloy. Moreover, the deformation mechanisms change at the BDTT, and the apparent activation energy depends on the deformation mechanisms [15,16]. It is necessary to investigate the change mechanisms of activation energy at the BDTT during creep. These aspects are addressed in the present study; where the creep tests were conducted at different temperature and stress in Ti-46Al-8Nb and Ti-44Al alloys to determine the BDTT. Creep tests were carried out based on the tensile testing results at elevated temperature. The performance and mechanisms of creep brittle-ductile transition are investigated and discussed. An attention is given to the change mechanism of apparent activation energy at the BDTT. 2. Materials and methods The as-cast ingots with nominal composition of Ti-44Al and Ti-46Al8Nb were prepared using the vacuum consumable melting technology, and the ingot was melted two times for composition homogenization.
Corresponding author. National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin, 150001, PR China. E-mail address: [email protected] (R. Chen).
https://doi.org/10.1016/j.msea.2019.138393 Received 4 April 2019; Received in revised form 4 September 2019; Accepted 5 September 2019 Available online 06 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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phases.
Table 1 Chemical composition of the investigated alloys (at.%). Alloy
Al
Nb
Ti-44Al Ti-46Al-8Nb
44.32 45.52
— 8.25
3.2. Brittle-ductile transition Fig. 2(a) shows the statistical results of elongation and yield stress versus temperature of Ti-46Al-8Nb alloy. The elevated temperature yield stress is 438 MPa and the elongation is 2.37% at testing temperature of 780 °C. The elevated temperature yield stress is 344 MPa and the elongation is 10.77% at the testing temperature of 800 °C. The elevated temperature yield strength decreases by 21.5% and the elongation increases 4.15 times as the testing tempreature increasing from 780 to 800 °C, which indicates that the BDTT is at 780–800 °C. Based on the elevated temperature tensile testing results, the samples were creep tested under 276 MPa at 740–800 °C, as shown in Fig. 2(b). Creep tests were terminated either at a strain well into the secondary stage (740 °C) or at fracture (760–800 °C). This alloy exhibits a low minimum creep rate during creep at 740–760 °C, and the creep strain is 2.62% and creep lifetime is 529 h after creep fracture test at 760 °C. This alloy exhibits a large creep strain and a short creep lifetime during creep at 780–800 °C, and the creep strain is 17.71% and the creep lifetime is 125 h after creep fracture test at 780 °C. The creep strain increases 6.8 times with the testing tempreature increasing from 760 to 780 °C, which suggests the creep BDTT is at 760–780 °C. The minimum creep rate significantly increases with the testing tempreature increasing from 760 to 780 °C. Moreover, the creep BDTT is 20 °C lower than the BDTT at elevated temperature tensile, which is related to the different strain rates between elevated temperature tensile tests and creep tests. As seen in the following discussion.
The creep testing simples were cut from the as-cast ingots. The chemical compositions of Ti-44Al and Ti-46Al-8Nb alloys were measured using the PW4400 X-ray fluorescence spectrum analysis, and they were summarized in Table 1. The transmission electron microscopy (TEM) foils of the investigated alloys were cut to about 350 μm by the wire-electrode firstly, and then they were grinded to approximately 70 μm, followed by twinjet electro-polishing in a solution of 35% butyl alcohol, acid 60% methanol and 5% perchloric at −35 °C and 15 V. The microstructure before/after creep tests were observed by the Quanta 200FEG scanning electron microscopy in the back-scattered electron mode (SEM-BSE). For the further analysis of α2/γ lamellar microstructure, the microstructure after creep test was observed by the Tecnai G2 F30 transmission electron microscope (TEM). The high-temperature tensile specimens with a gauge length of 18 mm and cross-sections of 5 mm × 2 mm were tested at a strain rate of 1×10−4/s on the 5569 Instron testing machine, to determine the BDTT. The creep tests were carried out based on the high-temperature tensile testing results. The creep specimens with a gauge length of 20 mm and cross-sections of 4.5 mm × 2.5 mm were tested in the air under the constant load using the GWT504 creep tester. The thermoelectric couple was used to measure the creep temperature, and the two linear variable differential transducers were used to measure the displacement.
3.3. Microstructure deformation Fig. 3(a) shows the SEM-BSE microstructure of Ti-46Al-8Nb alloy after creep test at 760 °C (below the BDTT), and the black arrows mark the direction of stress axis. The crack at colony boundary is perpendicular to the direction of applied loading, and no obvious microstructure deformation can be found in α2/γ lamellae. Fig. 3(b) shows the SEMBSE microstructure of Ti-46Al-8Nb alloy after creep test at 780 °C (above the BDTT). The crack is observed at colony boundary and the red line marks the bowed lamellae, which suggests that the lamellae are obviously deformed during creep above the BDTT. Fig. 4 shows the TEM images of α2/γ lamellae after creep testing at 780 °C under 276 MPa. Many dislocations can be observed in α2 lamellae, which indicate that the major dislocation slip systems of α2 lamellae are thermally activated above the BDTT. In the nearly/fully lamellar TiAl alloys, the deformation of α2/γ lamellae is determined by α2 lamellae at elevated temperature, and the brittle-ductile transition is a thermally activated process [6,14]. In this study, when creep above the BDTT, the main dislocation slip systems of α2 lamellae are thermally activated, which promote the plastic deformation of α2/γ lamellae and contribute to the overall creep strain. The crack initiation and propagation along colony boundaries is the main creep failure mechanism at the creep temperature both below and above BDTT. When creep above BDTT, the significant deformation of microstructure reduces microstructural stability. The lamellar colonies on both sides of grain boundary usually have different lamellar directions in TiAl alloys, and the lamellar colonies with different lamellar directions usually have different deformation capacities during creep above the BDTT [20]. The plastic incompatibility of lamellar colonies improves the local stress concentration at colony boundaries, which promotes the formation of cavities and cracks at colony boundaries, and accelerates creep failure. Moreover, the rapid necking occurred in the specimen during creep above the BDTT; therefore, this alloy exhibits a short creep lifetime during creep above the BDTT. When the Ti-46Al8Nb alloy creep below the BDTT, the main dislocation slip systems of α2 lamellae cannot be thermally activated. The deformation of α2/γ lamellae is very small, and a little plastic incompatibility occurred in
3. Results and discussion 3.1. Initial microstructure Fig. 1 shows the initial SEM-BSE microstructure of Ti-46Al-8Nb alloy. The γ phases with dark contrast and B2 phases with bright contrast are observed in the matrix, as they are induced by the high Nb content in Ti-46Al-8Nb alloy. Chen et al. [17] have reported that the B2 phases formed by the phase transformation of β → α at the colony boundaries during cooling, or formed via the phase transition of α → γ + α2 + β in lamellar structures in the high Nb-containing TiAl alloys. However, the B2 phase content in this alloy is much lower than that in the low Al-containing high-Nb TiAl alloys [18,19], due to the Ti-46Al8Nb alloy with the high Al-containing can restrict the formation of B2
Fig. 1. The initial BSE-SEM microstructure of Ti-46Al-8Nb alloy. 2
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Fig. 2. Elevated temperature properties: (a) the elevated temperature tensile tests at 760–820 °C and (b) the tensile creep tests at 740–800 °C.
lamellar colonies. Moreover, the slow necking occurred in the specimen during creep above the BDTT; therefore, this alloy exhibits a high microstructure stability and a long creep lifetime during creep below the BDTT. 3.4. Stress exponent at the BDTT To compare the overall creep behaviors, the dependence of minimum stress (σ) and creep rate (εmin) can be formulated by the power-law creep equation:
Q εmin = A (σ )nexp ⎛− c ⎞ ⎝ RT ⎠
(1)
where, n corresponds to the stress exponent, A corresponds to the constant, Qc corresponds to the activation energy, T corresponds to the absolute temperature and R corresponds to the gas constant. The logarithm of the stress versus minimum strain rate at creep temperature of 760 and 780 °C, are plotted in Fig. 5. The stress exponent n is about 1.4 during creep at 760 °C under 246–306 MPa, which suggests that the creep is controlled by interface sliding [21]. The stress exponent is about 4.3 during creep under 246–306 MPa at 780 °C. The stress exponent at range of 3–5 suggests that the creep is controlled by dislocation-climb [22]. This result further validates that the BDTT is controlled by the main dislocation slip system of α2 lamellae.
Fig. 4. TEM bright-field images of α2/γ lamellae after creep testing at 780 °C under 276 MPa.
lamellae. The elevated temperature deformation process is mainly influenced by the distribution, characteristics, and intensity of obstacles. According to a report by Zhang et al. [23], the obstacles can hinder the dislocation glide in a couple of ways: (1) the long-range internal stress τi, the dislocations cannot overcome long-rang obstacles merely via the thermal activation. (2) The local short-range obstacles, the obstacles can be overcome mainly by thermal activation. When the applied stress is greater than internal stress τi, the τ can be divided into two parts:
3.5. The brittle-ductile transition mechanisms The main dislocation slip of α2 lamellae can be activated during creep above the BDTT under the action of applying stress and thermal activation. It is obvious that the creep is a thermally activated process, and the creep mechanisms change at the BDTT, and the deformation is mainly controlled by the dislocation slip system of α2 lamellae in α2/γ
τ = τi + τe
(2)
Here, the τe corresponds to the effective stress. Fig. 6 shows the resistance curve corresponding to a local obstacle. The y-axis corresponds to the force f(x) acting on the dislocation line.
Fig. 3. SEM-BSE microstructure of colony boundaries after creep tests at (a) 760 °C and (b) 780 °C. 3
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Fig. 5. The logarithm of the minimum strain rate versus stress, (a) at 760 °C and (b) 780 °C.
Here, the T0 corresponds to the critical temperature. When T > T0 (above the BDTT), τ = τi, the flow stress does not depend on strain rate ε or temperature T. The thermal activation is very active that the dislocations can overcome local obstacle with the help of a small effective stress. Therefore, the applied stress just requires to overcome the long-range internal stress. The main dislocation slip system of α2 lamellae can be thermally activated during creep above the BDTT under a low stress. When T ≤ T0 (below BDTT), the deformation is dependent on the dislocation slip assisted by the thermal activation, and the flow stress is dependent on both strain rate ε and temperature Τ. Therefore, the main dislocation slip of α2 lamellae can be activated under very a high-applied stress and it cannot be activated during creep at 760 °C under 276 MPa. The BDTT measured by creep tests is 20 °C lower than that measured by elevated temperature tensile tests, as shown in Fig. 2. According to Eq. (7), the activation frequency at T0 matches the dislocation velocity required for deforming TiAl alloys at a strain rate of ε. The thermal activation frequency must be increased by the increasing temperature during TiAl alloys deformation at a higher strain rate. In this case, the BDTT depends on strain rate and the higher strain rate leads to a higher BDTT. The strain rate during creep tests is much lower than the strain rate during elevated temperature tensile tests, therefore, the BDTT measured by creep testing is lower than that measured by elevated temperature tensile testing.
Fig. 6. Resistance curve of an obstacle.
The total area below the resistance curve corresponds to the total energy required for a dislocation to overcome a local obstacle; and the shaded region corresponds to the energy supplied by thermal activation. The region below the shaded region corresponds to the energy supplied by effective stress. Therefore, the variation of Gibbs free energy during thermally activated process can be expressed as: x2
Δ G=
∫ [f (x ) − τbl] dx = ΔF − τe bΔa x1
(3)
x2
Here, the
∫ [f (x ) − τi bl] dx = ΔF corresponds to the total energy,
3.6. Apparent activation energy at the BDTT
x1
the Δa = lΔx corresponds to the region swept by the dislocation slip under the action of the thermal activation. The frequency of dislocations overcoming the local obstacle can be expressed as:
ΔG Δ F− τe bΔa ⎞ ⎞ = vo exp ⎛− v= vo exp ⎛− kT ⎝ kT ⎠ ⎠ ⎝
According to equation (1), the apparent activation energy can be described as:
d (lnε ) ⎤ Q c = −R ⎡ ⎢ d ⎣ (1/ T ) ⎥ ⎦σ / E , S
(4)
Here, the vo corresponds to the vibration frequency of dislocations. The relationship between dislocation slip and macroscopic deformation follows the Orowan equation:
ε0 = ρm bv
Here, σ/E corresponds to the elastic modulus normalized stress and S corresponds to the dislocation structure. The apparent activation energy Qc for Ti-46Al-8Nb alloy is calculated by the plotting ln(ε) and 1/T, as shown in Fig. 7, where the apparent activation energy was calculated through the slope of each line. From this figure, a decided change in value of apparent activation energy can be found (ranging from 205 to 653 kJ/mol). It can be seen from Fig. 7 that the liner relationship between 1/T and ln(ε) breaks down when the temperature is below BDTT. Previous studies have shown that the power law breakdown usually took place when the creep temperature was below 0.5 or 0.6 Tm (Tm corresponds to the melting temperature) [23]. The activation energy for self-diffusion is equal to the activation energy for creep when the creep temperature is above 0.5 or 0.6 Tm, and the activation energy for creep is below the activation energy for self-diffusion when the creep temperature is below 0.5 or 0.6 Tm [23–26]. The BDTT of TiAl alloys is usually at the temperature range of 0.5–0.6 Tm [6,9,14]. Lapin et al.
(5)
Here, the ρm corresponds to the density of mobile dislocation; the ν corresponds to the mean velocity of dislocation slip. According to Eqs. (3) and (4), the strain rate supplied by dislocations slip under the action of thermally activated as follows:
ΔG ⎞ ΔF − τe bΔa ⎞ ε = ε0 exp ⎛− = ε0 exp ⎛− kT ⎝ kT ⎠ ⎠ ⎝
(6)
Here ε0 = bρλv0, the λ corresponds to the distance between two adjacent obstacles. The flow stress as follows:
τ=
ΔF − kTln (ε0 / ε ) bΔa
⎧ τi + ⎨ τi ⎩
T ≤ T0
T > T0
(8)
(7) 4
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relationship between ln(ε) and 1/T, at constant stresses in Ti-44Al alloy, and the slope of the lines represents the apparent activation energy value Qc. There is a significant change in the apparent activation energy value at creep temperature of 700–720 °C. The 1/T and ln(ε) have a linear relation during creep above 720 °C. The apparent activation energy value at 660–700 °C is below the apparent activation energy value at 720–760 °C. The results indicate that the BDTT of Ti-44Al alloy is at 700–720 °C. The BDTT of Ti-46Al-8Nb alloy (760–780 °C) is 60 °C higher than Ti44Al alloy (700-700 °C), which well agrees with the previous results that the creep strength retention of high-Nb TiAl alloy is 60–100 °C higher than that of the binary TiAl alloys [31]. The activation parameters of elevated temperature deformation have been determined in both binary and the equivalent Al-containing high-Nb ternary TiAl alloys in previous studies [32]. These results agree with the researches in Nb diffusion in TiAl alloys, and imply that it is more difficult for the diffusion-assisted deformation in high-Nb TiAl alloys, which benefits the creep resistance and elevated temperature strength [33].
Fig. 7. The relationship between ln(ε) and 1/T, at constant stresses in Ti-46Al8Nb alloy. The slope of the lines represents the apparent activation energy value of QC.
4. Conclusions 1. The BDTT is determined to be 780–800 °C by high-temperature tensile tests and 760–780 °C by creep tests. The creep strain and creep life significantly change during creep at the BDTT. 2. The main dislocation slip systems of α2 lamellae are thermally activated above the BDTT, which promotes the plastic deformation of α2/γ lamellae and improves the overall creep strain. The plastic incompatibility of lamellar colonies promotes the formation of cavities and cracks at colony boundaries and accelerates creep failure. 3. When creep above the BDTT, the thermal activation is very active that dislocations can overcome local obstacle under acting a small effective stress, and the applied stress only needs to overcome longrange internal stress. When creep below the BDTT, the deformation is dependent on dislocation slip assisted by thermal activation, and the flow stress is dependent on both strain rate and temperature. Therefore, the main dislocation slip system α2 lamellae can be activated under very a high-applied stress. 4. The 1/T and ln(ε) have a linear relation during creep above the BDTT, due to the creep is controlled by the dislocation climb, and the apparent activation energy is 402 kJ/mol. The liner relationship between 1/T and ln(ε) breaks down during creep below the BDTT, due to the interface sliding mechanism being dominant in this regime. Moreover, there is a significant change in the apparent activation energy value at the BDTT, since the main dislocation slip system of α2 lamellae is thermally activated at the BDTT.
Fig. 8. The relationship between (lnε) and 1/T, at constant stresses in Ti-44Al alloy. The slope of the lines represents the apparent activation energy value of Qc.
[25,26] have reported that the liquidus temperature of Ti-46Al-8Nb alloy is 1570 °C. The BDTT measured by creep tests at the temperature range of 780–800 °C is at 0.5–0.6 Tm in this study [7,17]. According to Section 3.4, the creep is controlled by dislocation climb above the BDTT, and the creep is controlled by interface sliding below the BDTT. The 1/T and ln(ε) have a linear relation when the creep temperature is above the BDTT, combining the results of Section 3.4, which suggests that the of 1/T and ln(ε) have a linear relation when creep is controlled by dislocation climb, and the apparent activation energy is 402 kJ/mol. This result agrees to those of previous studies [27–30]. The liner relationship between 1/T and ln(ε) breaks down when the temperature is below the BDTT, combining the results of Section 3.4, which suggests that the liner relationship between 1/T and ln(ε) breaks down when the interface sliding mechanism being dominant in this regime. Moreover, there is a significant change in apparent activation energy value at the BDTT, since the main dislocation slip system of α2 lamellae is thermally activated at the BDTT, which leads to a significant increase in the creep rate and a large apparent activation energy value at the BDTT. It is obvious that the value of apparent activation energy depends on the creep mechanisms. The 1/T and ln(ε) have a linear relation during creep above the BDTT, there is a significant change in the apparent activation energy at the BDTT, and the apparent activation energy value is small during creep below the BDTT. Therefore, the apparent activation energy value can be used to determine the BDTT, and the apparent activation energy value is calculated by the minimum strain rate. TiAl alloys usually exhibit a short primary creep stage and rapidly decrease to a minimum value, therefore, the BDTT measured by the minimum strain rate can reduce creep testing time. Fig. 8 shows the
Acknowledgements The authors gratefully acknowledge support from the National Key Research and Development Programme of China (2017YFA0403804), China Postdoctoral Science Foundation funded project (2019TQ0076) and the National Natural Science Foundation of China Nos. 51741404. References [1] G. Chen, Y. Peng, G. Zheng, Z. Qi, M. Wang, H. Yu, C. Dong, G.T. Liu, Polysynthetic twinned TiAl single crystals for high-temperature applications, Nat. Mater. 15 (2016) 876. [2] Q. Wang, R.R. Chen, X. Gong, J.J. Guo, Y.Q. Su, H.S. Ding, H.Z. Fu, Microstructure, mechanical properties and crack propagation behaviour in high-Nb TiAl alloys by directional solidification, Metall. Mater. Trans. A 49A (2018) 4555–4564. [3] B. Cao, J.J. Yang, X.Y. Wang, Y.L. Wu, R. Hu, Microstructure evolution and mechanical properties of a Ti-45Al-8.5Nb-(W, B, Y) alloy obtained by controlled cooling from a single β region, J. Alloy. Comp. 740 (2017) 1140–1148. [4] Z. Wu, R. Hu, T. Zhang, H. Zhou, J. Li, Moisture induced environmental embrittlement of a high Nb containing TiAl alloy, Mater. Sci. Eng. A 701 (2017) 214–220. [5] H. Clemens, S. Mayer, Intermetallic titanium aluminides in aerospace applications-
5
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Q. Wang, et al.
processing, microstructure and properties, Mater, High Temp. 33 (2016) 1–11. [6] R.R. Chen, Q. Wang, Y.H. Yang, J.J. Guo, Y.Q. Su, H.S. Ding, H.Z. Fu, Brittle–ductile transition during creep in nearly and fully lamellar high-Nb TiAl alloys, Intermetallics 93 (2018) 47–54. [7] F. Appel, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminide Alloys: Science and Technology, Wiley-VCH, 2011. [8] J. Lapin, T. Pelachová, M. Dománková, Long-term creep behaviour of cast TiAl-Ta alloy, Intermetallics 95 (2018) 24–32. [9] Y.W. Kim, S.L. Kim, Effects of microstructure and C and Si additions on elevated temperature creep and fatigue of gamma TiAl alloys, Intermetallics 53 (2014) 92–101. [10] M. Kastenhuber, B. Rashkova, H. Clemens, S. Mayer, Effect of microstructural instability on the creep resistance of an advanced intermetallic γ-TiAl based alloy, Intermetallics 80 (2017) 1–9. [11] M. Kastenhuber, T. Klein, H. Clemens, S. Mayer, Tailoring microstructure and chemical composition of advanced γ-TiAl based alloys for improved creep resistance, Intermetallics 97 (2018) 27–33. [12] S.Z. Zhang, C.J. Zhang, Z.X. Du, Z.P. Hou, P. Lin, F.T. Kongd, Y.Y. Chen, Deformation behavior of high Nb containing TiAl based alloy in α + γ two phase field region, Mater. Des. 90 (2016) 225–229. [13] L. Cheng, J. Li, X. Xue, B. Tang, H. Kou, E. Bouzy, Superplastic deformation mechanisms of high Nb containing TiAl alloy with (α2 +γ) microstructure, Intermetallics 75 (2016) 62–71. [14] V.M. Imayev, R.M. Imayev, G.A. Salishchev, On two stages of brittle-to-ductile transition in TiAl intermetallic, Intermetallics 8 (2000) 1–6. [15] D. Dorner, K. Röller, B. Skrotzki, B. Stöckhert, G. Eggeler, Creep of a TiAl alloy: a comparison of indentation and tensile testing, Mater. Sci. Eng. A 357 (2003) 346–354. [16] W. Wang, W. Zeng, C. Xue, X. Liang, J. Zhang, Microstructural evolution, creep, and tensile behavior of a Ti–22Al–25Nb (at%) orthorhombic alloy, Mater. Sci. Eng. A 603 (2014) 176–184. [17] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Microsegregation in high Nb containing TiAl alloy ingots beyond laboratory scale, Intermetallics 15 (2007) 625–631. [18] B. Tang, L. Cheng, H. Kou, J. Li, Hot forging design and microstructure evolution of a high Nb containing TiAl alloy, Intermetallics 58 (2015) 7–14. [19] J.R. Yang, X.Y. Wang, B. Cao, Y.L. Wu, K.R. Zhang, R. Hu, Tailoring the
[20] [21] [22]
[23] [24]
[25]
[26]
[27] [28] [29]
[30]
[31]
[32] [33]
6
microstructure of a β-Solidifying TiAl alloy by controlled post-solidification isothermal holding and cooling, Metall. Mater. Trans. A 48 (2017) 1–11. X.W. Du, J. Zhu, Y.W. Kim, Microstructural characterization of creep cavitation in a fully-lamellar TiAl alloy, Intermetallics 9 (2001) 137–146. A.M. Hodge, L.M. Hsiung, T.G. Nieh, Creep of nearly lamellar TiAl alloy containing W, Scr. Mater. 51 (2004) 411–415. S. Bystrzanowski, A. Bartels, H. Clemens, R. Gerling, F.P. Schimansky, G. Dehmd, H. Kestler, Creep behaviour and related high temperature microstructural stability of Ti-46Al-9Nb sheet material, Intermetallics 13 (2005) 515–524. J.S. Zhang, Elevated Temperature Deformation and Fracture of Materials, (2010). M. Sujata, D.H. Sastry, C. Ramachandra, Microstructural characterization and creep behaviour of as-cast titanium aluminide Ti-48Al-2V, Intermetallics 12 (2004) 691–697. Z. Gabalcová, J. Lapin, Estimation of high temperature phase equilibria in directionally solidified intermetallic Ti-45.9Al-8Nb alloy, Kovove Mater 45 (2007) 231–240. Z. Gabalcová, J. Lapin, T. Pelachová, Effect of Y2O3 crucible on contamination of directionally solidified intermetallic Ti-46Al-8Nb alloy, Intermetallics 19 (2011) 396–403. L.M. Hsiung, Interfacial control of creep deformation in ultrafine lamellar TiAl, Mrs Proc. (2002) 740. Y. Ishikawa, K. Maruyama, H. Oikawa, Three regions in the power-law creep regime of TiAl, Mater. Trans. 33 (2007) 1182–1184. M.E. Kassner, The rate dependence and microstructure of high-purity silver deformed to large strains between 0.16 and 0.30Tm, Metall. Trans. 20A (1989) 2001–2010. H. Luthy, A.K. Miller, O.D. Sherby, The stress and temperature dependence of steady state flow at intermediate temperature for polycrystalline aluminum, Acta Metall. 28 (1980) 169–178. Z.C. Liu, J.P. Lin, S.J. Li, G.L. Chen, Effects of Nb and Al on the microstructures and mechanical properties of high Nb containing TiAl base alloys, Intermetallics 10 (2002) 653–659. Y. Mishin Y, C. Herzig, Diffusion in the Ti–Al system, Acta Mater. 48 (2000) 589–623. F. Appel, U. Lorenz, J.D.H. Paul, M. Oehring, Gamma Titanium Aluminides 1999, TMS, (1999), p. 381.