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Static coarsening behavior of a pre-deformed Ti2AlNb-based alloy during heat treatment
Vacuum 169 (2019) 108934
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Static coarsening behavior of a pre-deformed Ti2AlNb-based alloy during heat treatment
T
Hongyu Zhang, Chong Li∗∗, Zongqing Ma, Yuan Huang, Liming Yu, Yongchang Liu∗ State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Ti2AlNb-Based alloy Aging O precipitates Coarsening
Microstructure evolution, static coarsening mechanism and coarsening kinetics of a pre-deformed Ti2AlNb-based alloy during aging at 945 °C and 970 °C were investigated in this study. The corresponding microstructures were characterized by scanning electron microscopy and transmission electron microscopy. The results show that the pre-deformed alloy exhibited different morphologies aged at 945 °C and 970 °C. When the alloy was aged at 970 °C, the coarsening mechanism of the O precipitates was mainly controlled by the Ostwald ripening mechanism at first aging stage; after that, it was controlled by the boundary splitting mechanism and termination migration mechanism. Due to the difference in interface energy between the terminals and the long-axis direction of the O lath, diffusion process of solute atoms leads to the coarsening and fragmenting of O lath. The dimension of the O precipitates in the studied alloy continued to grow with the prolonged time during aging at 945 °C or 970 °C while the coarsening rate decreased. The static coarsening kinetic of the studied alloy was established in term of the modified Lifshitz-Slyozov-Wagner (LSW) theory. Static coarsening behavior of the predeformed alloy was mainly controlled by bulk diffusion at 945 °C and interfacial diffusion at 970 °C, respectively.
1. Introduction In recent decades, Ti–Al–Nb intermetallics were widely developed for aerospace and automotive application due to the wonderful elevated temperature properties, such as good creep resistance, high specific strength and improved oxidation resistance with respect to the conventional titanium alloys [1–6]. O + β/B2 Ti2AlNb-based alloy is one of the most popular titanium alloys for potential candidate used in the temperature range of 650–750 °C because of its desirable combinations of strength, toughness and ductility. It is of great significance to reduce the self-weight of the aircraft, improve fuel efficiency and high-temperature service performance [7–14]. The acceptable properties of Ti2AlNb-based alloy are mainly attributed to the O phase which possesses the ordered orthorhombic structure (Cmcm symmetry based on Ti2AlNb) [15–19]. For Ti2AlNb-based alloys, the mechanical properties depend strongly on their microstructure, especially morphology and size of the O phase. For instance, the lamellar morphology which is formed during furnace cooling from β/B2-phase region exhibits a moderate high specific strength and great creep resistance but bad ductility. By contrast, globular or equiaxed morphology owns good ductility and toughness but low strength [20–23]. Hence, how to regulate the morphology and control microstructure evolution through ∗
thermal processing and heat treatment processes is critical. Microstructure coarsening and globularization for conventional titanium alloys (i.e. Ti-17 and Ti–6Al–4V) has received considerable attentions due to their great influence on mechanical performance during service process [24–30]. Semiatin et al. [24] researched the coarsening behavior of Ti–6Al–4V alloy with an untra-fine-grain (UFG) microstructure, and discovered that the static coarsening behavior of alpha phase follows the r 3 − r03 vst rule. Zherebtsov et al. [25] studied the globularization of the lamellar microstructure in Ti–6Al–4V alloy during deformation and following annealing at 600 °C and 800 °C, and summarized that the spheroidization of the lamellar morphology is controlled by boundary splitting mechanism first and then by means of termination migration. Stefansson et al. [26,27] investigated the kinetics and mechanisms of spheroidization of Ti–6Al–4V alloy after deformation and subsequent static heat treatment. Their results showed that the static spheroidization is dependent on the formation and evolution of dislocation substructure, and the kinetics of globularization of alpha phase depends on the degree of pre-deformation and the temperature of following heat treatment. Xu et al. [28–30] systematic investigated the static coarsening and globularization behaviors of the Ti17 alloy with lamellar and equiaxed alpha phase during sstatic heat treatment at 820 °C and 860 °C. They found that the static coarsening
Corresponding author. Co-corresponding authors. E-mail addresses: [email protected] (C. Li), [email protected] (Y. Liu).
∗∗
https://doi.org/10.1016/j.vacuum.2019.108934 Received 20 August 2019; Received in revised form 1 September 2019; Accepted 7 September 2019 Available online 09 September 2019 0042-207X/ © 2019 Elsevier Ltd. All rights reserved.
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examined using a Hitachi S-4800 SEM operating at 15 kV. Transmission electron microscope (TEM, JEM-2100 at 200 kV) observations were employed to characterize more details of the precipitates and coarsening behavior. The slices for TEM observations were firstly machined to thin discs of about 300 μm and mechanically ground to 50–70 μm, then core-drilled to 3 mm diameter wafers. A twin-jet electro-polishing device was used to thin the specimens at −40 °C and 70 V in a solution of 34 vol% n-butanol, 6 vol% perchloric acid, and 60 vol% carbinol. The dimension of phases was measured by image analysis on SEM images using Image Pro Plus 6.0 software. For each specimen, 10 images were chosen to measure the dimension of O precipitates by manual measurement. In addition, O precipitates were considered to be a globular one if it owned an aspect of 2 or less. On the contrary, it was considered to be a lamellar or lath one [32].
behavior of the lamellar and equiaxed alpha phase in Ti-17 alloy was interpreted by the modified Lifshitz-Slyozov-Wagner (LSW) theory and the kinetics of static spheroidization was depend mainly on the strain, heating temperature and heating time. The coarsening mechanism and coarsening coefficient were also determined. From the above all, a lot of studies have been carried out for static coarsening behavior of titanium alloy while limited work like this has been done for Ti2AlNb-based alloy. Since the coarsening behavior is important for engineering of such materials, the objective of the current study was to investigate and make clear the static coarsening behavior and microstructure evolution of a Ti2AlNb-based alloy after pre-deformation and subsequent heat treatment. This information is useful and guided for the microstructure control and optimization of heat treatment process of these alloys from both scientific viewpoint and industrial perspective.
3. Results and discussion
2. Materials and experimental method
3.1. Microstructure observations
As-received material in this investigation was a Ti2AlNb-based alloy fabricated at Central Iron and Steel Research Institute (CISRI). Before the investigation, the studied alloy was prepared by quadruple consumable vacuum arc remelting. It had a measured chemical composition (in atom percent) of 22.3Al, 25.5Nb, and balance Ti, which was consistent with the nominal composition of Ti–22Al–25Nb, and the gas impurity was lower (i.e. nitrogen ≈82 ppm, and oxygen ≈600 ppm). Thermal compression test specimens with a gauge dimension of Ø8 × 12 mm were cut off from the center of the as-forged alloy by wire electro discharge machining (WEDM). Isothermal compression test was implemented at 1040 °C under the strain rate of 0.01 s−1 and height reduction of 20% on a thermo mechanical simulator (Gleeble-3500) according to the authors’ previous study about hot deformation behavior of Ti–22Al–25Nb [31]. Before the test, all specimens were heated to the determined temperature and kept for 180 s to reduce the thermal gradient. To determine the phase region and heat treatment procedures, pre-deformed specimen was heated up from 25 °C to 1200 °C in the differential scanning calorimetry (Mettler ToledoTGA/DSC 1) with heating rate of 20 °C·min−1. The heat flow curve was demonstrated in Fig. 1. In terms of the phase transition temperature from the DSC curve, the pre-deformed Ti2AlNb-based alloys were put into the vacuum furnace for 1 h, 6 h, 18 h, and 42 h at 945 °C (i.e., in the O + β/B2-phase region) and 970 °C (i.e., in the O + β/B2 + α2-phase region) followed by water quenching (WQ). After that, all specimens were machined along the compression axis section, standard metallographic polished and then etched with a corrosive of HF: HNO3: H2O2 = 1:3:5 at normal temperature. The microstructure and morphologies of O phase were
3.1.1. True stress-strain curve and pre-deformed microstructure True stress-strain curve at the deformation temperature of 1040 °C is shown in Fig. 2. At the initial stage of strain, the flow stress is increased rapidly with increasing straining near to peak stress and then it slowly rises to the peak stress and gradually drops to a relatively stable state until the deformation ends. The reason for this trend is generally due to the competition between work hardening and dynamic softening mechanisms (e.g. dynamic recovery and dynamic recrystallization). At the early stage, the role of work hardening is remarkable, which is due to the rapid accumulation of dislocation. With the further increase of strain, the slightly decline of flow stress after peak stress can be attributed to the evolution of dislocation substructure [33,34]. The peak stress and peak strain are determined as 99.27 MPa and 0.085, respectively. Fig. 3(a) illustrates the pre-deformed microstructure after isothermal compression at 1040 °C (i.e. α2+B2/β phase region), it can be seen that most of the composition was matrix B2/β phase, besides, some O phase were detected from grain boundary. The energy dispersive spectrometer (EDS) analysis reveals the chemical compositions of grain boundary O phases. The EDS result of O phase is exhibited in Fig. 3(b). It is worth mentioning that the existence of the α2 phase was not found in the matrix, which can be confirmed by XRD diffraction peak shown in Fig. 3(c). Previous study [35] has also confirmed that deformation can inhibit the precipitation of α2 phase.
Fig. 1. DSC curve of the pre-deformed Ti2AlNb-based alloy with the heating rate of 20 °C· min−1.
Fig. 2. True stress-strain curve of the explored Ti2AlNb-based alloy under the strain rate of 0.01s−1. 2
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Fig. 3. (a) SEM image of the pre-deformed Ti2AlNb-based alloy, (b) energy dispersive spectrum (EDS) of grain boundary O and (c) XRD pattern.
Furthermore, as can be seen from Fig. 5(d), the interface between O and B2/β phase becomes unclear due to the continuous transition zone.
3.1.2. Microstructure evolution after aging treatment Microstructural evolution of the O precipitates during isothermal heat treatment in pre-deformed Ti2AlNb-based alloys is investigated by SEM characterizations. Fig. 4 demonstrates the microstructure evolution of pre-deformed Ti2AlNb-based alloy during aging at 945 °C for different time. As mentioned above, there are no lamellar O precipitates prior to heat treatment (seen in Fig. 3(a)). After aging for 1 h, terminals of some lamellar O phase were bent to some extent, as shown in Fig. 4(a), the average thickness of lamellar O was measured as 0.51 ± 0.07 μm. With the prolongation of heating time, dimension of the lamellar O phase increased continually and the interface between O phase and B2/β phase became clearer. Moreover, the amount of O phase precipitates did not decrease with the increase of heat treatment time, which was reflected in the fact that there was still a relatively dense lamellar O precipitates distributed in the matrix after heating for 42 h (seen in Fig. 4(d)). SEM images of the pre-deformed Ti2AlNb-based alloy during aging at 970 °C for different time are shown in Fig. 5. One could see that the morphology of 970 °C-aging-treatment is different from that of 945 °C. Some spherical O phase existed in the matrix when the heating time is less than 6 h (seen in Fig. 5(a) and (b)) while lath O appeared when the heating time is more than 18 h (seen in Fig. 5(c) and (d)). It also can be found that the coarsening of O phase was obvious with the increase of aging time and the coarsening rate was faster intuitively on account of the diffusion of solid atoms increases distinctly with the increase of temperature. In addition, unlike the 945 °C-aging process, more spherical O precipitates dissolved in the matrix at 970 °C. The reason for this phenomenon may be that 970 °C is close to the phase transition temperature from (O + α2 + B2/β) region to (α2 + B2/β) region.
3.2. Coarsening mechanism at 970 °C Microstructural evolution of the pre-deformed Ti2AlNb-based alloy at 970 °C is more obvious than that at 945 °C. Therefore, the morphology evolution and coarsening mechanism of Ti2AlNb-based alloy are discussed in detail by taking aging at 970 °C as an example. Fig. 6 shows the SEM images of the pre-deformed Ti2AlNb-based alloy aged at 970 °C for 1 h and 6 h. Compared with the specimen with longer aging time (> 6 h), a considerable amount of spherical O phase exists in the matrix when the aging time is less than 6 h. When the aging time is 6 h, the average particle size of spherical O phase is slightly smaller than that of 1 h, and the count of O phase particles decreases with aging time. The size of the overall O precipitates, especially the width of the O lath, increases with the prolonged aging time, while the size of the spherical O decreases, which could be explained by the Ostwald ripening mechanism [27,36]. The mechanism has been extensively used to interpret the diffusion process of elements in solid solutions [37–40]. It is substantially a coarsening mechanism that represents a process that reduces free energy of the system [30,41]. The small-sized O phase grains possess higher Gibbs free energy, while the large-sized grains have a lower Gibbs free energy. The energy gradient between the different O phase grains provides a driving force for the behavior of large-scale particles merged with small-scale particles. The coalescence process reduced the system free energy and makes the microstructure tend to be steady. The trace of microstructure evolution can be detected from the residual diffusion zones which are labeled by 3
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Fig. 4. SEM micrographs of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h, (b) 6 h, (c) 18 h, and (d) 42 h at 945 °C.
(large-sized lath). This produces the observed phenomenon that the number of spherical O phase grains decreases but the dimension of O increases. As discussed above, Ostwald ripening mechanism played an important role in the coarsening of the O phase precipitates during the initial aging stage of the alloy at 970 °C. However, with the prolongation of aging time, the effects of other mechanisms began to emerge. The coarsening mechanism and morphological evolution of O precipitates during aged at 970 °C for 18 h is investigated in detail by TEM
black dashed circles in Fig. 6. For example, some small-sized O phase grains (labeled by black dashed circles in Fig. 6(a)) are dissolving and meanwhile, some spherical O phase grains have been almost dissolved (labeled by black dashed circles in Fig. 6(b)). Schematic illustration of Ostwald ripening mechanism in the studied Ti2AlNb alloy is shown in Fig. 7. The small-sized spherical O phase grains are dissolved in the matrix B2/β by a phase transformation process, and then large-sized lath O phase grow through the phase equilibrium. The process can be summarized as follows: O (small-sized spherical) → B2/β (matrix) → O
Fig. 5. SEM micrographs of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h, (b) 6 h, (c) 18 h, and (d) 42 h at 970 °C. 4
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Fig. 6. SEM images of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h and (b) 6 h at 970 °C.
which is manifested that some lath O phase has been completed and segmented (seen in Fig. 8(b)). This process can be interpreted by the boundary splitting mechanism. At the same time, separation of the terminal material was observed, which was characterized by the fragmenting and local dissolution of the terminals of the lath O (seen in Fig. 8(c)). Isothermal deformation process prior to heat treatment causes the local disturbances and formation of defects, resulting in larger curvature and higher chemical potential gradients [27,29,42]. Due to the difference in interface energy between the terminals of the lath O phase and the long-axis direction, the diffusion process of the solute atoms is more remarkable at the terminals and defect regions. Eventually, the B2/β matrix is constantly embedded, leading to the fragmenting of the lath O. This phenomenon is consistent with the termination migration mechanism, which is an important coarsening mechanism based on diffusion theory. It was first presented in the research of the spheroidization and coarsening behavior of pearlite in steel [43,44], and has been applied to several titanium-aluminum alloys with success [24,25,30,45]. The termination migration mechanism is a diffusion process of solute atoms which is influenced by the difference in the curvature of terminals and the flat interfaces. The diffusion process of solute atoms can be verified by energy dispersive spectrometer (EDS) analysis, as shown in Fig. 9 and Table 1. From position 1 to 3, the content of β-stabilizer element “Nb” gradually increases from the
Fig. 7. Schematic illustration of Ostwald ripening mechanism in Ti2AlNb alloy.
characterizations, as presented in Fig. 8. Fig. 8(a)–(c) reflected different coarsening mechanisms, and the precipitates can be determined by the SAED pattern in Fig. 8(d). The O phase precipitated in the matrix undergoes local shear deformation inside the lath (shown by the red arrow in Fig. 8(a)), which is caused by the instability of new interface inside the O precipitates which were formed during the pre-deformation process. In addition, the defects generated by the deformation stored higher energy, and the energy storage in the subsequent aging process will stimulate the distortion energy and promote the diffusion of the elements [30]. Due to the inhomogeneous deformation and shorter aging time, the local shear deformation process is not synchronous,
Fig. 8. Different coarsening behaviors (i.e. (a) (b) boundary splitting mechanism, (c) termination migration mechanism) of the pre-deformed Ti2AlNb-based alloy and (d) selected area electron diffraction (SAED) pattern of the O precipitates. 5
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prolonging from 1 h to 42 h at 945 °C. In contrast, the coarsening kinetic was more remarkable at a higher temperature. At 970 °C, the average dimension of O precipitates increased from 1.24 ± 0.19 μm to 1.95 ± 0.32 μm after 1 h–42 h of aging. In addition, it is worth noting that the slope of the curve decreases with time, indicating that the coarsening rate is decreasing. After the nucleation and precipitation process of the O precipitates in the supersaturated solid solution is completed, the aging process will lead to the coarsening of particles. The driving force for coarsening is derived from the decreasing of free energy of the interface between the precipitates and matrix or the diffusion of the solute atoms. The coarsening kinetic depends on the rate-controlled mode of different processes, and different rate-control modes lead to coarsening and growth rate of different precipitates. Previous studies used LSW theory to explain the coarsening behavior of precipitates controlled by diffusion. However, the LSW theory was based on the assumption that the volume fraction of precipitates is infinitely small and the spherical particles are completely coherent with the matrix. Despite some deviations in the rate constant and particle size distribution [46], the general form of coarsening kinetics predicted by LSW theory has been well documented. For example, some commercial nickel-based superalloys and two-phase titanium alloys followed the linear law of LSW theory [47–49], which is called modified LSW theory and is in the following form [24,50]:
Fig. 9. SEM micrograph of pre-deformed Ti2AlNb-based alloy aged for 42 h at 970 °C.
Table 1 The EDS analysis results performed in different positions shown in Fig. 9 on predeformed Ti2AlNb-based alloy heat treated for 42 h at 970 °C. Position
Ti, at.%
Al, at.%
Nb, at.%
1 2 3 4
57.65 51.43 48.69 51.27
27.70 27.19 26.69 22.20
14.65 21.38 24.62 26.52
(1)
d = Kt n
in which d denotes the average dimension of O precipitates, K is the proportional constant, t is aging time, and n is the coarsening coefficient. Assuming an initial dimension d0 in this study, Eq. (1) would be the following form [25]:
d inside to outside of the O lath, which confirms the diffusion behavior between the O precipitates and the matrix. Besides, the element content of the residual diffusion zone in the matrix (see position 4 in Fig. 9) is substantially consistent with the nominal composition of the alloy, indicating that it is almost completely dissolved in the B2/β matrix.
1
1
n
− d 0 n = K ′t
(2)
in which K ′ is still a proportional constant. The values of K ′ and n relied on the coarsening mechanism, morphology and volume fraction of the precipitates. Previous researches on titanium alloys have shown that the coarsening process is controlled by bulk diffusion when n = 0.33 and by interfacial diffusion when n = 0.5 [27,50,51]. The coarsening coefficient n can be determined by drawing the results according to
3.3. Static coarsening kinetics
d
1
1
n
− d 0 n vs t. For this, assuming several different values of coarsening 1
1
coefficient n, variation curves of d n − d 0 n versus aging time at two different heating temperatures were exhibited in Fig. 11. It can be observed that the curvature of these two curves gradually changes from concave downward to concave upward with the coarsening coefficient
As an isothermal diffusion process, the static coarsening behavior is directly related to the heat treatment time and temperature. Fig. 10 shows the variation of the dimension of O precipitates with the aging time at two different temperatures. It can be seen that the dimension of O precipitates at both temperatures is significantly incremental with the increase of aging time, and the coarsening behavior is more obvious at the beginning of aging. The average dimension of O precipitates increased from 0.51 ± 0.07 μm to 0.98 ± 0.24 μm when heating time
1
1
n increasing. The curve of d n − d 0 n vs t was nearly linear at 945 °C when the coarsening coefficient n was 0.25 or 0.33 and when the 1
1
coarsening coefficient n was 0.4, 0.45 or 0.5, the curve of d n − d 0 n vs t was close to the linear at 970 °C. Therefore, the pre-deformed Ti2AlNbbased alloy was determined to have a coarsening coefficient of 0.25–0.33 for 945 °C and 0.4–0.5 for 970 °C, respectively. The above discussion revealed that the static coarsening behavior of pre-deformed Ti2AlNb alloy is not controlled independently by bulk diffusion or interfacial diffusion. When the heating temperature is 945 °C, bulk diffusion dominated; while interfacial diffusion is dominant at 970 °C. These results are consistent with the studies of conventional titanium alloys [24,26,29,52]. 4. Conclusions Static coarsening behavior and microstructure evolution of the studied alloy after pre-deformation and subsequent heat treatments at 945 °C and 970 °C have been investigated. The coarsening mechanism and kinetic of a pre-deformed Ti2AlNb-based alloy were discussed through sufficient metallographic analysis. The follow conclusions have been reached:
Fig. 10. Variation curves of the O precipitates dimension with aging time at 945 °C and 970 °C. 6
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Fig. 11. Curves of d (f) 1.
1
1
n
− d 0 n vs t for the explored material heat treated at 945 °C and 970 °C, and values of n equal to (a) 0.25, (b) 0.33, (c) 0.4, (d) 0.45, (e) 0.5, and
theory. The coarsening coefficient was determined to be 0.25–0.33 at 945 °C and the coarsening process was mainly controlled by bulk diffusion; the coarsening coefficient was determined to be 0.4–0.5 at 970 °C and the coarsening process was principally controlled by interfacial diffusion.
(1) The pre-deformed Ti2AlNb-based alloy exhibited completely different morphologies aging at 945 °C and 970 °C. The microstructure consisted of lamellar O at 945 °C while changed from spherical O to lath O with increasing aging time at 970 °C. (2) When the pre-deformed Ti2AlNb-based alloy was aged at 970 °C, the size of spherical O phase grains decreased or even disappeared with the increase of aging time, and the dimension of the lath O was gradually increased. When the aging time was less than 6 h, the coarsening mechanism of the precipitates was mainly controlled by Ostwald ripening mechanism; after 6 h, it was controlled by the boundary splitting mechanism and termination migration mechanism. (3) The dimension of O precipitates in the pre-deformed Ti2AlNb-based alloy continuously increase with aging time during static heat treatment at 945 °C or 970 °C and the coarsening rate of O precipitates decreases with the aging time. Static coarsening kinetic was more pronounced at 970 °C. (4) Static coarsening kinetic of O precipitates of the pre-deformed Ti2AlNb-based alloy was established according to the modified LSW
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Static coarsening behavior of a pre-deformed Ti2AlNb-based alloy during heat treatment
T
Hongyu Zhang, Chong Li∗∗, Zongqing Ma, Yuan Huang, Liming Yu, Yongchang Liu∗ State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Ti2AlNb-Based alloy Aging O precipitates Coarsening
Microstructure evolution, static coarsening mechanism and coarsening kinetics of a pre-deformed Ti2AlNb-based alloy during aging at 945 °C and 970 °C were investigated in this study. The corresponding microstructures were characterized by scanning electron microscopy and transmission electron microscopy. The results show that the pre-deformed alloy exhibited different morphologies aged at 945 °C and 970 °C. When the alloy was aged at 970 °C, the coarsening mechanism of the O precipitates was mainly controlled by the Ostwald ripening mechanism at first aging stage; after that, it was controlled by the boundary splitting mechanism and termination migration mechanism. Due to the difference in interface energy between the terminals and the long-axis direction of the O lath, diffusion process of solute atoms leads to the coarsening and fragmenting of O lath. The dimension of the O precipitates in the studied alloy continued to grow with the prolonged time during aging at 945 °C or 970 °C while the coarsening rate decreased. The static coarsening kinetic of the studied alloy was established in term of the modified Lifshitz-Slyozov-Wagner (LSW) theory. Static coarsening behavior of the predeformed alloy was mainly controlled by bulk diffusion at 945 °C and interfacial diffusion at 970 °C, respectively.
1. Introduction In recent decades, Ti–Al–Nb intermetallics were widely developed for aerospace and automotive application due to the wonderful elevated temperature properties, such as good creep resistance, high specific strength and improved oxidation resistance with respect to the conventional titanium alloys [1–6]. O + β/B2 Ti2AlNb-based alloy is one of the most popular titanium alloys for potential candidate used in the temperature range of 650–750 °C because of its desirable combinations of strength, toughness and ductility. It is of great significance to reduce the self-weight of the aircraft, improve fuel efficiency and high-temperature service performance [7–14]. The acceptable properties of Ti2AlNb-based alloy are mainly attributed to the O phase which possesses the ordered orthorhombic structure (Cmcm symmetry based on Ti2AlNb) [15–19]. For Ti2AlNb-based alloys, the mechanical properties depend strongly on their microstructure, especially morphology and size of the O phase. For instance, the lamellar morphology which is formed during furnace cooling from β/B2-phase region exhibits a moderate high specific strength and great creep resistance but bad ductility. By contrast, globular or equiaxed morphology owns good ductility and toughness but low strength [20–23]. Hence, how to regulate the morphology and control microstructure evolution through ∗
thermal processing and heat treatment processes is critical. Microstructure coarsening and globularization for conventional titanium alloys (i.e. Ti-17 and Ti–6Al–4V) has received considerable attentions due to their great influence on mechanical performance during service process [24–30]. Semiatin et al. [24] researched the coarsening behavior of Ti–6Al–4V alloy with an untra-fine-grain (UFG) microstructure, and discovered that the static coarsening behavior of alpha phase follows the r 3 − r03 vst rule. Zherebtsov et al. [25] studied the globularization of the lamellar microstructure in Ti–6Al–4V alloy during deformation and following annealing at 600 °C and 800 °C, and summarized that the spheroidization of the lamellar morphology is controlled by boundary splitting mechanism first and then by means of termination migration. Stefansson et al. [26,27] investigated the kinetics and mechanisms of spheroidization of Ti–6Al–4V alloy after deformation and subsequent static heat treatment. Their results showed that the static spheroidization is dependent on the formation and evolution of dislocation substructure, and the kinetics of globularization of alpha phase depends on the degree of pre-deformation and the temperature of following heat treatment. Xu et al. [28–30] systematic investigated the static coarsening and globularization behaviors of the Ti17 alloy with lamellar and equiaxed alpha phase during sstatic heat treatment at 820 °C and 860 °C. They found that the static coarsening
Corresponding author. Co-corresponding authors. E-mail addresses: [email protected] (C. Li), [email protected] (Y. Liu).
∗∗
https://doi.org/10.1016/j.vacuum.2019.108934 Received 20 August 2019; Received in revised form 1 September 2019; Accepted 7 September 2019 Available online 09 September 2019 0042-207X/ © 2019 Elsevier Ltd. All rights reserved.
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examined using a Hitachi S-4800 SEM operating at 15 kV. Transmission electron microscope (TEM, JEM-2100 at 200 kV) observations were employed to characterize more details of the precipitates and coarsening behavior. The slices for TEM observations were firstly machined to thin discs of about 300 μm and mechanically ground to 50–70 μm, then core-drilled to 3 mm diameter wafers. A twin-jet electro-polishing device was used to thin the specimens at −40 °C and 70 V in a solution of 34 vol% n-butanol, 6 vol% perchloric acid, and 60 vol% carbinol. The dimension of phases was measured by image analysis on SEM images using Image Pro Plus 6.0 software. For each specimen, 10 images were chosen to measure the dimension of O precipitates by manual measurement. In addition, O precipitates were considered to be a globular one if it owned an aspect of 2 or less. On the contrary, it was considered to be a lamellar or lath one [32].
behavior of the lamellar and equiaxed alpha phase in Ti-17 alloy was interpreted by the modified Lifshitz-Slyozov-Wagner (LSW) theory and the kinetics of static spheroidization was depend mainly on the strain, heating temperature and heating time. The coarsening mechanism and coarsening coefficient were also determined. From the above all, a lot of studies have been carried out for static coarsening behavior of titanium alloy while limited work like this has been done for Ti2AlNb-based alloy. Since the coarsening behavior is important for engineering of such materials, the objective of the current study was to investigate and make clear the static coarsening behavior and microstructure evolution of a Ti2AlNb-based alloy after pre-deformation and subsequent heat treatment. This information is useful and guided for the microstructure control and optimization of heat treatment process of these alloys from both scientific viewpoint and industrial perspective.
3. Results and discussion
2. Materials and experimental method
3.1. Microstructure observations
As-received material in this investigation was a Ti2AlNb-based alloy fabricated at Central Iron and Steel Research Institute (CISRI). Before the investigation, the studied alloy was prepared by quadruple consumable vacuum arc remelting. It had a measured chemical composition (in atom percent) of 22.3Al, 25.5Nb, and balance Ti, which was consistent with the nominal composition of Ti–22Al–25Nb, and the gas impurity was lower (i.e. nitrogen ≈82 ppm, and oxygen ≈600 ppm). Thermal compression test specimens with a gauge dimension of Ø8 × 12 mm were cut off from the center of the as-forged alloy by wire electro discharge machining (WEDM). Isothermal compression test was implemented at 1040 °C under the strain rate of 0.01 s−1 and height reduction of 20% on a thermo mechanical simulator (Gleeble-3500) according to the authors’ previous study about hot deformation behavior of Ti–22Al–25Nb [31]. Before the test, all specimens were heated to the determined temperature and kept for 180 s to reduce the thermal gradient. To determine the phase region and heat treatment procedures, pre-deformed specimen was heated up from 25 °C to 1200 °C in the differential scanning calorimetry (Mettler ToledoTGA/DSC 1) with heating rate of 20 °C·min−1. The heat flow curve was demonstrated in Fig. 1. In terms of the phase transition temperature from the DSC curve, the pre-deformed Ti2AlNb-based alloys were put into the vacuum furnace for 1 h, 6 h, 18 h, and 42 h at 945 °C (i.e., in the O + β/B2-phase region) and 970 °C (i.e., in the O + β/B2 + α2-phase region) followed by water quenching (WQ). After that, all specimens were machined along the compression axis section, standard metallographic polished and then etched with a corrosive of HF: HNO3: H2O2 = 1:3:5 at normal temperature. The microstructure and morphologies of O phase were
3.1.1. True stress-strain curve and pre-deformed microstructure True stress-strain curve at the deformation temperature of 1040 °C is shown in Fig. 2. At the initial stage of strain, the flow stress is increased rapidly with increasing straining near to peak stress and then it slowly rises to the peak stress and gradually drops to a relatively stable state until the deformation ends. The reason for this trend is generally due to the competition between work hardening and dynamic softening mechanisms (e.g. dynamic recovery and dynamic recrystallization). At the early stage, the role of work hardening is remarkable, which is due to the rapid accumulation of dislocation. With the further increase of strain, the slightly decline of flow stress after peak stress can be attributed to the evolution of dislocation substructure [33,34]. The peak stress and peak strain are determined as 99.27 MPa and 0.085, respectively. Fig. 3(a) illustrates the pre-deformed microstructure after isothermal compression at 1040 °C (i.e. α2+B2/β phase region), it can be seen that most of the composition was matrix B2/β phase, besides, some O phase were detected from grain boundary. The energy dispersive spectrometer (EDS) analysis reveals the chemical compositions of grain boundary O phases. The EDS result of O phase is exhibited in Fig. 3(b). It is worth mentioning that the existence of the α2 phase was not found in the matrix, which can be confirmed by XRD diffraction peak shown in Fig. 3(c). Previous study [35] has also confirmed that deformation can inhibit the precipitation of α2 phase.
Fig. 1. DSC curve of the pre-deformed Ti2AlNb-based alloy with the heating rate of 20 °C· min−1.
Fig. 2. True stress-strain curve of the explored Ti2AlNb-based alloy under the strain rate of 0.01s−1. 2
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Fig. 3. (a) SEM image of the pre-deformed Ti2AlNb-based alloy, (b) energy dispersive spectrum (EDS) of grain boundary O and (c) XRD pattern.
Furthermore, as can be seen from Fig. 5(d), the interface between O and B2/β phase becomes unclear due to the continuous transition zone.
3.1.2. Microstructure evolution after aging treatment Microstructural evolution of the O precipitates during isothermal heat treatment in pre-deformed Ti2AlNb-based alloys is investigated by SEM characterizations. Fig. 4 demonstrates the microstructure evolution of pre-deformed Ti2AlNb-based alloy during aging at 945 °C for different time. As mentioned above, there are no lamellar O precipitates prior to heat treatment (seen in Fig. 3(a)). After aging for 1 h, terminals of some lamellar O phase were bent to some extent, as shown in Fig. 4(a), the average thickness of lamellar O was measured as 0.51 ± 0.07 μm. With the prolongation of heating time, dimension of the lamellar O phase increased continually and the interface between O phase and B2/β phase became clearer. Moreover, the amount of O phase precipitates did not decrease with the increase of heat treatment time, which was reflected in the fact that there was still a relatively dense lamellar O precipitates distributed in the matrix after heating for 42 h (seen in Fig. 4(d)). SEM images of the pre-deformed Ti2AlNb-based alloy during aging at 970 °C for different time are shown in Fig. 5. One could see that the morphology of 970 °C-aging-treatment is different from that of 945 °C. Some spherical O phase existed in the matrix when the heating time is less than 6 h (seen in Fig. 5(a) and (b)) while lath O appeared when the heating time is more than 18 h (seen in Fig. 5(c) and (d)). It also can be found that the coarsening of O phase was obvious with the increase of aging time and the coarsening rate was faster intuitively on account of the diffusion of solid atoms increases distinctly with the increase of temperature. In addition, unlike the 945 °C-aging process, more spherical O precipitates dissolved in the matrix at 970 °C. The reason for this phenomenon may be that 970 °C is close to the phase transition temperature from (O + α2 + B2/β) region to (α2 + B2/β) region.
3.2. Coarsening mechanism at 970 °C Microstructural evolution of the pre-deformed Ti2AlNb-based alloy at 970 °C is more obvious than that at 945 °C. Therefore, the morphology evolution and coarsening mechanism of Ti2AlNb-based alloy are discussed in detail by taking aging at 970 °C as an example. Fig. 6 shows the SEM images of the pre-deformed Ti2AlNb-based alloy aged at 970 °C for 1 h and 6 h. Compared with the specimen with longer aging time (> 6 h), a considerable amount of spherical O phase exists in the matrix when the aging time is less than 6 h. When the aging time is 6 h, the average particle size of spherical O phase is slightly smaller than that of 1 h, and the count of O phase particles decreases with aging time. The size of the overall O precipitates, especially the width of the O lath, increases with the prolonged aging time, while the size of the spherical O decreases, which could be explained by the Ostwald ripening mechanism [27,36]. The mechanism has been extensively used to interpret the diffusion process of elements in solid solutions [37–40]. It is substantially a coarsening mechanism that represents a process that reduces free energy of the system [30,41]. The small-sized O phase grains possess higher Gibbs free energy, while the large-sized grains have a lower Gibbs free energy. The energy gradient between the different O phase grains provides a driving force for the behavior of large-scale particles merged with small-scale particles. The coalescence process reduced the system free energy and makes the microstructure tend to be steady. The trace of microstructure evolution can be detected from the residual diffusion zones which are labeled by 3
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Fig. 4. SEM micrographs of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h, (b) 6 h, (c) 18 h, and (d) 42 h at 945 °C.
(large-sized lath). This produces the observed phenomenon that the number of spherical O phase grains decreases but the dimension of O increases. As discussed above, Ostwald ripening mechanism played an important role in the coarsening of the O phase precipitates during the initial aging stage of the alloy at 970 °C. However, with the prolongation of aging time, the effects of other mechanisms began to emerge. The coarsening mechanism and morphological evolution of O precipitates during aged at 970 °C for 18 h is investigated in detail by TEM
black dashed circles in Fig. 6. For example, some small-sized O phase grains (labeled by black dashed circles in Fig. 6(a)) are dissolving and meanwhile, some spherical O phase grains have been almost dissolved (labeled by black dashed circles in Fig. 6(b)). Schematic illustration of Ostwald ripening mechanism in the studied Ti2AlNb alloy is shown in Fig. 7. The small-sized spherical O phase grains are dissolved in the matrix B2/β by a phase transformation process, and then large-sized lath O phase grow through the phase equilibrium. The process can be summarized as follows: O (small-sized spherical) → B2/β (matrix) → O
Fig. 5. SEM micrographs of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h, (b) 6 h, (c) 18 h, and (d) 42 h at 970 °C. 4
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Fig. 6. SEM images of the pre-deformed Ti2AlNb-based alloy aged for (a) 1 h and (b) 6 h at 970 °C.
which is manifested that some lath O phase has been completed and segmented (seen in Fig. 8(b)). This process can be interpreted by the boundary splitting mechanism. At the same time, separation of the terminal material was observed, which was characterized by the fragmenting and local dissolution of the terminals of the lath O (seen in Fig. 8(c)). Isothermal deformation process prior to heat treatment causes the local disturbances and formation of defects, resulting in larger curvature and higher chemical potential gradients [27,29,42]. Due to the difference in interface energy between the terminals of the lath O phase and the long-axis direction, the diffusion process of the solute atoms is more remarkable at the terminals and defect regions. Eventually, the B2/β matrix is constantly embedded, leading to the fragmenting of the lath O. This phenomenon is consistent with the termination migration mechanism, which is an important coarsening mechanism based on diffusion theory. It was first presented in the research of the spheroidization and coarsening behavior of pearlite in steel [43,44], and has been applied to several titanium-aluminum alloys with success [24,25,30,45]. The termination migration mechanism is a diffusion process of solute atoms which is influenced by the difference in the curvature of terminals and the flat interfaces. The diffusion process of solute atoms can be verified by energy dispersive spectrometer (EDS) analysis, as shown in Fig. 9 and Table 1. From position 1 to 3, the content of β-stabilizer element “Nb” gradually increases from the
Fig. 7. Schematic illustration of Ostwald ripening mechanism in Ti2AlNb alloy.
characterizations, as presented in Fig. 8. Fig. 8(a)–(c) reflected different coarsening mechanisms, and the precipitates can be determined by the SAED pattern in Fig. 8(d). The O phase precipitated in the matrix undergoes local shear deformation inside the lath (shown by the red arrow in Fig. 8(a)), which is caused by the instability of new interface inside the O precipitates which were formed during the pre-deformation process. In addition, the defects generated by the deformation stored higher energy, and the energy storage in the subsequent aging process will stimulate the distortion energy and promote the diffusion of the elements [30]. Due to the inhomogeneous deformation and shorter aging time, the local shear deformation process is not synchronous,
Fig. 8. Different coarsening behaviors (i.e. (a) (b) boundary splitting mechanism, (c) termination migration mechanism) of the pre-deformed Ti2AlNb-based alloy and (d) selected area electron diffraction (SAED) pattern of the O precipitates. 5
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prolonging from 1 h to 42 h at 945 °C. In contrast, the coarsening kinetic was more remarkable at a higher temperature. At 970 °C, the average dimension of O precipitates increased from 1.24 ± 0.19 μm to 1.95 ± 0.32 μm after 1 h–42 h of aging. In addition, it is worth noting that the slope of the curve decreases with time, indicating that the coarsening rate is decreasing. After the nucleation and precipitation process of the O precipitates in the supersaturated solid solution is completed, the aging process will lead to the coarsening of particles. The driving force for coarsening is derived from the decreasing of free energy of the interface between the precipitates and matrix or the diffusion of the solute atoms. The coarsening kinetic depends on the rate-controlled mode of different processes, and different rate-control modes lead to coarsening and growth rate of different precipitates. Previous studies used LSW theory to explain the coarsening behavior of precipitates controlled by diffusion. However, the LSW theory was based on the assumption that the volume fraction of precipitates is infinitely small and the spherical particles are completely coherent with the matrix. Despite some deviations in the rate constant and particle size distribution [46], the general form of coarsening kinetics predicted by LSW theory has been well documented. For example, some commercial nickel-based superalloys and two-phase titanium alloys followed the linear law of LSW theory [47–49], which is called modified LSW theory and is in the following form [24,50]:
Fig. 9. SEM micrograph of pre-deformed Ti2AlNb-based alloy aged for 42 h at 970 °C.
Table 1 The EDS analysis results performed in different positions shown in Fig. 9 on predeformed Ti2AlNb-based alloy heat treated for 42 h at 970 °C. Position
Ti, at.%
Al, at.%
Nb, at.%
1 2 3 4
57.65 51.43 48.69 51.27
27.70 27.19 26.69 22.20
14.65 21.38 24.62 26.52
(1)
d = Kt n
in which d denotes the average dimension of O precipitates, K is the proportional constant, t is aging time, and n is the coarsening coefficient. Assuming an initial dimension d0 in this study, Eq. (1) would be the following form [25]:
d inside to outside of the O lath, which confirms the diffusion behavior between the O precipitates and the matrix. Besides, the element content of the residual diffusion zone in the matrix (see position 4 in Fig. 9) is substantially consistent with the nominal composition of the alloy, indicating that it is almost completely dissolved in the B2/β matrix.
1
1
n
− d 0 n = K ′t
(2)
in which K ′ is still a proportional constant. The values of K ′ and n relied on the coarsening mechanism, morphology and volume fraction of the precipitates. Previous researches on titanium alloys have shown that the coarsening process is controlled by bulk diffusion when n = 0.33 and by interfacial diffusion when n = 0.5 [27,50,51]. The coarsening coefficient n can be determined by drawing the results according to
3.3. Static coarsening kinetics
d
1
1
n
− d 0 n vs t. For this, assuming several different values of coarsening 1
1
coefficient n, variation curves of d n − d 0 n versus aging time at two different heating temperatures were exhibited in Fig. 11. It can be observed that the curvature of these two curves gradually changes from concave downward to concave upward with the coarsening coefficient
As an isothermal diffusion process, the static coarsening behavior is directly related to the heat treatment time and temperature. Fig. 10 shows the variation of the dimension of O precipitates with the aging time at two different temperatures. It can be seen that the dimension of O precipitates at both temperatures is significantly incremental with the increase of aging time, and the coarsening behavior is more obvious at the beginning of aging. The average dimension of O precipitates increased from 0.51 ± 0.07 μm to 0.98 ± 0.24 μm when heating time
1
1
n increasing. The curve of d n − d 0 n vs t was nearly linear at 945 °C when the coarsening coefficient n was 0.25 or 0.33 and when the 1
1
coarsening coefficient n was 0.4, 0.45 or 0.5, the curve of d n − d 0 n vs t was close to the linear at 970 °C. Therefore, the pre-deformed Ti2AlNbbased alloy was determined to have a coarsening coefficient of 0.25–0.33 for 945 °C and 0.4–0.5 for 970 °C, respectively. The above discussion revealed that the static coarsening behavior of pre-deformed Ti2AlNb alloy is not controlled independently by bulk diffusion or interfacial diffusion. When the heating temperature is 945 °C, bulk diffusion dominated; while interfacial diffusion is dominant at 970 °C. These results are consistent with the studies of conventional titanium alloys [24,26,29,52]. 4. Conclusions Static coarsening behavior and microstructure evolution of the studied alloy after pre-deformation and subsequent heat treatments at 945 °C and 970 °C have been investigated. The coarsening mechanism and kinetic of a pre-deformed Ti2AlNb-based alloy were discussed through sufficient metallographic analysis. The follow conclusions have been reached:
Fig. 10. Variation curves of the O precipitates dimension with aging time at 945 °C and 970 °C. 6
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Fig. 11. Curves of d (f) 1.
1
1
n
− d 0 n vs t for the explored material heat treated at 945 °C and 970 °C, and values of n equal to (a) 0.25, (b) 0.33, (c) 0.4, (d) 0.45, (e) 0.5, and
theory. The coarsening coefficient was determined to be 0.25–0.33 at 945 °C and the coarsening process was mainly controlled by bulk diffusion; the coarsening coefficient was determined to be 0.4–0.5 at 970 °C and the coarsening process was principally controlled by interfacial diffusion.
(1) The pre-deformed Ti2AlNb-based alloy exhibited completely different morphologies aging at 945 °C and 970 °C. The microstructure consisted of lamellar O at 945 °C while changed from spherical O to lath O with increasing aging time at 970 °C. (2) When the pre-deformed Ti2AlNb-based alloy was aged at 970 °C, the size of spherical O phase grains decreased or even disappeared with the increase of aging time, and the dimension of the lath O was gradually increased. When the aging time was less than 6 h, the coarsening mechanism of the precipitates was mainly controlled by Ostwald ripening mechanism; after 6 h, it was controlled by the boundary splitting mechanism and termination migration mechanism. (3) The dimension of O precipitates in the pre-deformed Ti2AlNb-based alloy continuously increase with aging time during static heat treatment at 945 °C or 970 °C and the coarsening rate of O precipitates decreases with the aging time. Static coarsening kinetic was more pronounced at 970 °C. (4) Static coarsening kinetic of O precipitates of the pre-deformed Ti2AlNb-based alloy was established according to the modified LSW
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