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The microstructure, mechanical properties, and oxidation behavior of beta-gamma TiAl alloy with excellent hot workability
Materials Science & Engineering A 700 (2017) 366–373
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The microstructure, mechanical properties, and oxidation behavior of betagamma TiAl alloy with excellent hot workability ⁎
MARK
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S.Z. Zhanga,b, , Y.B. Zhaoa,b, C.J. Zhanga,b, J.C. Hanc, , M.J. Sund, M. Xua,b a
School of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China Shanxi Key Laboratory of Advanced Magnesium-based Materials, Taiyuan University of Technology, Taiyuan 030024, PR China c School of Mechanical Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China d Shanghai Space flight Precision Machinery Institute, Shanghai 200000, PR China b
A R T I C L E I N F O
A B S T R A C T
Keywords: TiAl alloys Hot deformation Microstructure Oxidation Mechanical properties
New type beta-gamma TiAl alloy of nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y (at%) was fabricated by induction skull melting, which was designed based on the study of the microstructural effects of Mo, Cr, V, and Nb on TiAl alloy. The microstructure, oxidation resistance, hot workability, and tensile properties were studied by SEM in BSE mode, XRD, Gleeble 3800, as well as Instron 5500 R, respectively. The oxidation resistance of this new designed Mo containing beta-gamma TiAl alloy was better than that of only Nb and V containing betagamma TiAl alloy. With a better hot workability at temperature under 1200 °C, Ti-44Al-4Nb-4V-0.3Mo-Y alloy was successfully hot forged at 1150 °C with a traditional hydraulic machine without thermal insulation measure. Though containing abound order β0 phase, the new fabricated beta-gamma TiAl alloy still has a high elongation at 850 °C, because of less α2+γ lamellar colonies containing. Benefits of partly replacing Mo with V was avoiding omega phase precipitation, and not damaging the hot workability.
1. Introduction TiAl-based alloy possessed outstanding properties such as low density, high specific strength and modulus, superior oxidation resistance, and creep resistance, which was considered to be one of the most promising high-temperature structural materials used as turbine blades of aero-engines, turbocharger turbine wheels of automotive combustion engines, etc. [1,2]. However, TiAl alloy was difficult to be processed into component parts due to their unsatisfactory room temperature ductility and poor hot deformability, which seriously restricted its practical application [3]. Lots of work have been done to improve the ductility and workability of TiAl alloy, especially the alloy design [4–8]. Kim et al. [6] came up with a concept of novel beta-gamma TiAl alloy and its composition range was: Ti-(40−45)Al-(2−7)Nb-(1−9) M-(0–0.5) N, where M represents β phase stabilizing elements, such as Nb, Cr, Mn, V, and Mo, while N represents grain refiner elements, such as C, B, and Y elements. As BCC crystal structure had sufficient number of independent slip systems, disordered β phase was softer than γ and α phase at elevated temperature and improved the deformability of TiAl alloy (the ordered phase of β is β0) [9,10]. Moreover, beta-gamma TiAl alloy could effectively eliminate composition segregation and ⁎
microstructure inhomogeneity, reduce cast texture, and refine microstructure via β phase solidification, which avoided peritectic reaction [11]. An alloy design concept had been put forward for beta-gamma TiAl alloy by Clemens [8,12,13]. Based on this concept, TNM alloy with composition of Ti-43Al-4Nb-1Mo-0.1B was fabricated [12]. TNM alloy exhibited an appropriate volume fraction of β phase, solidified completely via β phase, possessed balanced mechanical properties, and had been machined into a turbine blade component successfully after a three-step forging process [11–13]. Besides, Ti-43Al-9V and Ti-42Al5Mn were also typical beta-gamma TiAl alloys and were confirmed with excellent hot working properties [14,15]. Alloying elements, such as Nb, Mo, and V elements, influenced different properties of TiAl alloy [16]. With development of TiAl alloy towards higher temperature and higher performance, Nb became the indispensable alloying element of TiAl alloy [17]. However, Nb could raise the diffusion activation energy thereby damaging the hot deformability of the TiAl alloy [18]. The Mo element addition improved the room temperature ductility, microstructure stability, and oxidation resistance of the TiAl alloy [19]. However, omega phase with lower crystal symmetry, which was extremely brittle and was detrimental to mechanical performance of TiAl, could be produced during the decomposition of β/β0 phase in high Nb and Mo containing TiAl alloy
Corresponding author at: School of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China. Corresponding author. E-mail addresses: [email protected] (S.Z. Zhang), [email protected] (J.C. Han).
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http://dx.doi.org/10.1016/j.msea.2017.06.025 Received 1 February 2017; Received in revised form 6 June 2017; Accepted 7 June 2017 Available online 09 June 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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Fig. 1. Scanning electron micrographs taken in backscattering electron mode: (a) Ti-44Al alloy; (b) Ti44Al-2Mo alloy; (c) Ti-44Al-4Mo alloy; (d) Ti-44Al4Cr alloy; (e) Ti-44Al-4V alloy; (f) Ti-44Al-4Nb alloy.
were studied. And a novel beta-gamma TiAl alloy with a nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y was designed and fabricated. The microstructure characteristic of this alloy co-added with Nb, V, and Mo was discussed. As Mo was partly replaced of V, the hot workability and oxidation resistance of Ti-44Al-4Nb-4V-0.3Mo-Y were investigated to understand the effect of co-addition of V and Mo. Moreover, high temperature tensile properties of this alloy were investigated as well.
such as TNM alloy and high Nb containing TiAl alloy (TNB alloy) [20–22]. Then the total content of Nb and Mo in beta-gamma TiAl alloy should be decreased further. As a β phase stabilizer, V improved the hot workability of TiAl dramatically but damaged its oxidation resistance [23–26]. Omega phase cannot be founded in high V containing TiAl alloy (Ti-43Al-9V and Ti-42Al-10V) when the content of Al was no less than 40% [26,27]. Moreover, the volume fraction of the stabilized β phase should be in an appropriate value for the sake of being removed by heat-treatment after hot-working, since the appearance of ordered β0 phase deteriorates the creep resistance, room-temperature strengths, and ductility [12]. In order to avoid the decrease of mechanical properties induced by omega phase and to maintain a better hot workability, the total containing of Nb and Mo should be partly replaced by other β stabilizing elements [21]. Hence, the β phase stabilizing ability of element needs be considered, and the effect of co-addition of Nb, V, and Mo on microstructure and mechanical properties has been still unknown. In this paper, the microstructural effects and β phase stabilizing ability of Mo, Cr, V, Nb on Ti-44Al (in at%, similarly hereinafter) alloy
2. Experimental procedure Six sets of alloys with nominal compositions of Ti-44Al-(0, 2Mo, 4Mo, 4Cr, 4 V, 4 Nb) were prepared using nonconsumable vacuum arc melting. Pancake ingots with 60 g were melted four times to ensure chemical homogeneity. Samples for microstructural analysis were cut from the center of as-cast pancakes. Ingot of Ti-44Al-4Nb-4V-0.3Mo-Y alloy were fabricated by induction skull melting (ISM) technique. Cubic specimens with a dimension of 10 × 10 × 10 mm3 were cut for oxidation tests by using electrical-discharge method from the ingots of Ti-44Al-4Nb-4V-0.3Mo-Y alloy. The specimen surfaces were ground 367
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to 1000 grit using SiC papers, followed by ultrasonic cleaning for 15 min in acetone and methanol and dried for oxidation tests (ultrasonic cleaning was performed by JP-020S ultrasonic cleaning machines from SKYMEN made in China). The short-time oxidation tests were performed at 800 °C, in static air of 5–80 h using a high temperature furnace (KSL-1200X-J from HF-Kejing made in China). Each specimen weight was measured by Ruipin FA1004A electronic balance (China, accuracy, ± 0.1 mg). Samples for high temperature compression tests were cut with dimensions of φ8 mm×12 mm and the surfaces were ground. High temperature compression experiments were carried out on Gleeble 3800 machine at temperatures 800 °C, 1190 °C, and 1250 °C up to a high reduction of 70% with strain rates of 1–0.01 s−1. Samples for hot forged were cut from the cast ingot and heated at 1150 °C for about 40 min. Then samples were hot forged in a conventional hydraulic press with a strain rate of 0.02 s−1 (the strain rate was calculated with the . . equation of ε = ε / t , where ε is strain rate, ε is strain, and t is deformation time). Tensile test samples with a gauge section of 5 × 2 mm2 cross-section and a gauge length of 20 mm were wire-cut parallel to the length of ingot. High temperature tensile tests were performed at temperatures of 700 °C, 800 °C, and 850 °C at strain rates of 5 × 10−4 s−1 using an Instron 5500 R testing machine. The microstructural characterization was performed by Quanta 200FEG field-emission scanning electron microscopy (FESEM) using backscattered electrons mode (BSE). Microstructural specimens were mechanical ground and polished. Phase analyses were performed by Xray diffraction (XRD) on D/max−2500 pc (RIGAKU, Japan) in the range 20–100°(2θ) using a step size of 0.02°. 3. Results and discussion
located both inside and at the boundaries of the lamellar colony was shown in Fig. 1(d). The mixed structure was produced by the cellular reaction β→β+γ during solidification [32]. As shown in Fig. 1(e), there exists a little amount of β/β0 phase in the microstructure of Ti44Al-4V alloy, and small amount of β/β0 phase could not refine the microstructure of as-cast TiAl alloy and then could not improve its hot workability. It was revealed that the β phase stabilizing ability was weaker than that of Mo element. As shown in Fig. 1(f), the microstructure of Ti-44Al-4Nb alloy was full lamellar structure without β/β0 phase containing, which illustrated that Nb was the weakest β phase stabilizer among Mo, V, Cr and Nb. According to the given in analysis, it was reasonable to surmise that the β phase stabilizing ability was Mo > Cr > V > Nb, which was in accordance with the previous research [33]. It was also obvious that the microstructure and β/β0 phase containing had no significant change comparing with Ti-44Al alloy, when the single addition of Nb and V was no more than 4% in TiAl alloy. And the result also revealed that the high containing of β/β0 phase in the Widmannstatten structure of Ti44Al-4Mo alloy did not meet the design concepts of beta-gamma phase TiAl alloy. Since, Nb could dramatically improve the high temperature strength and oxidation resistance of TiAl alloy, it was a necessity for TiAl alloy [34]. As was presented in literature [35], the ratio of β phase stabilizing ability between Mo, V, Nb was 4.25:1.6:1. To obtain the same β/β0 phase content in TiAl ally with an equal Al containing, the atomic ratio of Mo, V, and Nb was 4.25:1.6:1 when they were single added. In order to enhance the high temperature properties of TiAl alloy, avoid omega phase producing, and improve hot workability, a new beta-gamma TiAl alloy was designed with a nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y which has an equivalent weight of about 2Mo with TNM alloy (the equivalence formula is 1Mo=4.25 Nb, and 1Mo=2.66 V according to literature [33]).
3.1. Microstructure of Ti-44Al-xM alloy
3.2. As-cast microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy
In order to evaluate the β phase stabilizing ability of β stabilizing elements and their effects on the microstructure of TiAl alloy, the ascast microstructure of Ti-44Al-xM (where, x=0,4; M represented Mo, Cr, V, Nb) alloys were studied (Fig. 1). In BSE mode, β/β0 phase has the brightest contract, α/α2 phase is grey while the contrast of γ phase is black [28]. It was not difficult to distinguish these phases by contrast. As shown in Fig. 1(a), Ti-44Al alloy with a larger colony size had a full lamellar structure and was mainly composed of (α2+γ) colonies which had an average size of about 100 µm. The small colony size was the result of β phase solidification which avoided the peritectic transformation in traditional TiAl alloy which even had a colony size of 1000 µm [29,30]. Fig. 1(b) and (c) showed the microstructures of Ti44Al-2Mo and Ti-44Al-4Mo alloy. It was shown that a large amount of β/β0 phase distributed at (α2+γ) lamellar colony boundaries and inside the (α2+γ) colonies, which divided the large lamellar colonies into many fine lamellar grains, when the Mo content was 2%. It was noticed that there existed three orientations of the lamellar structure which were at 90°, 45°, and parallel to the primary growth direction of the β dendrite (the red line, blue line and black line in Fig. 1(b) represented the angle of 0°, 45°, and 90 °, respectively). As the precipitation of α phase at β phase followed an as-called Burgers relationship which was and [001]α //[110]β , then the (α2+γ) lamellar structure formed by the reaction of α→α2+γ following Blackburn relationship has the above three angles with the remnant β/β0 phase [31]. However, with the Mo content increased to 4%, Widmannstatten β/β0 phase distributed around the needle-like proeutectoid γ phase and became the dominating phase [23]. This microstructure damaged the mechanical property of alloy, and the content of β/β0 phase exceeded the betagamma TiAl alloy design concept and could not be eliminated by subsequent process [12]. Therefore, the content of Mo element could be no more than 4% when design a beta-gamma TiAl alloy. In the microstructure of the Ti-44Al-4Cr alloy, the mixture structure of γ and β/β0
Fig. 2(a) was the as-cast microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy. In this condition, the alloy exhibited a nearly lamellar microstructure which consisted of relatively small α2+γ lamellar colonies (grey contrast) with the size of 30–100 µm, and mixture structure of γ (black contrast) and β/β0 (light grey contrast) grains located along the lamellar colony boundaries. As shown in Fig. 2(b), the α2+γ lamellar colony was marked with green letters. XRD result was shown in Fig. 2(c) also confirmed that the alloy was mainly composed of γ, β/β0, and α2 phases. On the one hand, fine lamellar colonies were ascribed to the completely β phase solidification which could form as many as 12 possible orientation variants of α phase from β according to Burgers relationship [8]. On the other hand, the remnant β phase stabilized by Mo, V, and Nb elements was rejected to the interface of the different orientation α phase and then was separated along the boundaries, which impeded the growth of the Burgers α grains [8]. A substantial amount of β phase which was stabilized to room temperature by the addition of Nb, V, and Mo with a volume fraction of 37% was equal to that of TNM alloy and Ti-43Al-9V alloy [36,37]. It was illustrated by the Ti-Al-(4−10)% Nb-(V, Mo) polythermal sectional diagram that the solidifying pathway of the present alloy was: L→L+β→β→β+α→ α+β+γ→β+γ→α2+β/β0+γ, with no single-phase α-region existing [34]. As shown in Fig. 2(b), except for the fine spacing α2+γ lamellar colonies, there was also large amount of mixed structures of coarse β/β0 and γ, as well as crooked β/β0+γ lamellar morphology colonies. It was noted that two kinds of β/β0+γ structure morphology were shown in Fig. 2(b) marked with red and yellow letters. The β/ β0+γ colony structure marked with red letters was formed by the transformation of α→β+γ in a pearlitic mode procedure [23]. Whereas, the mixed structure of β/β0 and γ marked with yellow letters was the result of the cellular reaction of β→β+γ whose driving force was the chemical disequilibrium when β phase stabilizing elements segregated to β/α-interface in the course of the β→α transformation 368
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Fig. 2. As-cast microstructure and XRD result of Ti44Al-4Nb-4V-0.3Mo-Y alloy: (a) and (b) as-cast microstructure in BSE mode; (c) XRD result.
literatures [39–42]. Five alloys showed similar oxidation kinetics including linear growth stage and parabolic growth stage. Generally, Nb containing alloys exhibited a slower mass gain than the Nb free alloy at the same oxidation time, and V had the opposite effect on mass gain. In the initial oxidation stage, the mass gain rate of Ti-44Al-4Nb-4V-0.3MoY alloy was slow, but this stage can continue for 15 h, indicating that the oxidation scale yielded in this stage was protective for the alloy. The linear growth stage of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was almost on the same level as Ti-45Al-8Nb and Ti-43.5Al-4Nb-1Mo-B alloy. After 80 h, the mass gain of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was about 1.92 mg/ cm2, which was a little higher than that of Ti-45Al-8Nb and Ti-43.5Al4Nb-1Mo-B alloy, whereas oxidation resistance was significantly improved comparing with Ti-45Al-4.5Nb-4.5V-Y and Ti-45Al-3Nb-6V-Y alloy. Though Ti-44Al-4Nb-4V-0.3Mo-Y alloy was a V containing alloy, oxidation resistance ability was improved dramatically by the amount of 0.3 at% Mo element addition. It was pointed out that the continuous Al3O2 layer could hinder the diffusion of oxygen and prevent the formation of TiO2 which deteriorated oxidation resistance [43]. Networked Al2O3 layer was formed on the external surface of Mo addition TiAl alloy, and the second Al2O3 layer from the outer surface was more continuous and stable than that of Mo free TiAl alloy [44]. It was obvious that the amount of 0.3 at% Mo element addition in Ti-43Al-4Nb4V-0.3Mo-Y alloy improved the formation of Al2O3 layer which enhanced the oxidation resistance dramatically than free Mo containing TiAl alloy.
boundaries [13]. Bright particles appeared in Fig. 2(b) may be determined as the YAl2 phase according to the previous studies [38]. The YAl2 phase cannot be found by XRD because the phase is too rare to be detected. 3.3. High-temperature oxidation behavior In order to assess the high temperature oxidation resistance of the Ti-44Al-4Nb-4V-0.3Mo-Y alloy, its cyclic oxidation kinetics curve was plotted together with other alloys in Fig. 3 for comparison. The oxidation data of Ti-45Al-8Nb, Ti-45Al-4.5Nb-4.5V-Y, Ti-45Al-5.4V3.6Nb-0.3Y, Ti-43.5Al-4Nb-1Mo-B alloy were published in the existed
3.4. Hot-working behavior In order to evaluate the hot workability of Ti-44Al-4Nb-4V-0.3Mo-Y alloy, high temperature compression tests were conducted, and the true stress-true strain curves were shown in Fig. 4. As shown in Fig. 4(a) and (b), the highest stress of 260 MPa was obtained during hot compressed at 1190 °C with a strain rate of 1 s−1. The peak stresses of this alloy and TNM alloy for the given temperature were summarized in Table 1
Fig. 3. Cyclic oxidation kinetics curves for five different TiAl alloys. The added data of Ti45Al-8Nb, Ti-45Al-4.5Nb-4.5V-Y, Ti-45Al-5.4V-3.6Nb-0.3Y, Ti-43.5Al-4Nb-1Mo-B alloy was from Refs. [39–42] for comparison.
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Fig. 4. True stress-true strain curves of Ti-44Al-4Nb-4V-0.3Mo-Y alloy at different deformation temperature: (a) 1190 °C with strain rates of 0.01 s−1, 0.05 s−1, 0.1 s−1, 0.5 s−1, and 1 s−1; (b) 1250 °C with different strain rate of 0.01 s−1, 0.05 s−1, 0.1 s−1, 0.5 s−1, and 1 s−1; (c) 800 °C with 0.05 s−1.
(1190 °C) and 1250 °C were listed in Table 1. However, peak stress could reach as high as 740 MPa when deformed at 800 °C with 0.05 s−1 shown in Fig. 4(c). The stress curves shown in Fig. 4 indicated that Ti44Al-4Nb-4V-0.3Mo-Y alloy possessed good high temperature deformability and high strength at lower temperature, especially at the temperature around and even lower than 1200 °C. And the alloy could be hot deformed with lower deformation resistance and good deformability at temperatures around 1200 °C. The microstructure of samples hot compressed at 800 °C, 1190 °C and 1250 °C with strain rate of 0.05 s−1 was shown in Fig. 5. As shown in Fig. 5(a) and (c), microstructures of samples hot compressed at 1190 °C and 800 °C were typical deformed structure with flow lines which were perpendicular to external force axis (external force axis was indicated by red arrow in Fig. 5). β/β0 phase was stretched more seriously and was broken down into discontinuous structure somewhere in Fig. 5(c). Acicular γ phase precipitated in β/β0 phase when deformation temperature elevated to 1190 °C. When temperature increase to 1250 °C, block γ phase and acicular γ phase were observed in Fig. 5(b), and flow lines were not as clear as Fig. 5(a). Obviously, the flow stress was highly sensitive to strain, temperature, and strain rate. The higher the strain rates and the lower the
Table 1 The peak stress of TNM alloy and Ti-44Al-4Nb-4V-0.3Mo-Y alloy. Strain rate
Temperature 1200 °C*
0.5 s−1 0.05 s−1 0.005 s−1*
275 MPa (TNM) 180 MPa (TNM) 110 MPa (TNM)
1250 °C 183 MPa 114 MPa 90 MPa
160 MPa (TNM) 60 MPa (TNM) 30 MPa (TNM)
111 MPa 67 MPa 36 MPa
State: 1200 °C* represented that TNM alloy was deformed at 1200 °C, while Ti-44Al-4Nb4V-0.3Mo-Y alloy was deformed at 1190 °C; 0.005 s−1* represented that TNM alloy was deformed at 0.005 s−1, while Ti-44Al-4Nb-4V-0.3Mo-Y alloy was deformed at 0.01 s−1.
[45–47]. It was obvious that the peak stress at 1190 °C of this alloy was lower than that of TNM at 1200 °C [47]. For instance, the peak stress of TNM alloy deformed at 1200 °C with strain rate of 0.5 s−1 was 275 MPa [47], while the peak stress of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was 183 MPa which was two thirds of TNM alloy at 1200 °C with strain rate of 0.5 s−1. Compared to TNM alloy, the peak stress of Ti-44Al-4Nb4V-0.3Mo-Y alloy was for a maximum of 33% decrease. The peak stresses of Ti-44Al-4Nb-4V-0.3Mo-Y alloy and TNM alloy at 1200 °C
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Fig. 5. Microstructure of samples hot compressed with strain rate of 0.05 s−1 at different temperatures: (a) 1190 °C; (b) 1250 °C; (c) 800 °C. The arrow indicated the orientation of external force which was perpendicular to flow lines in (a) and (c).
0.3Mo-Y alloy were shown a smooth and crack-free appearance. As the equipment was a conventional hydraulic press without any protective atmosphere and the forged sample was not protected by other technique, such as canned with stainless steel to acquire near-isothermally and prevent oxidation, the surface of the pancakes were coated with thin oxidation layer during heated at1150°C for 40 min. Successfully hot forging at 1150 °C without any protection on a conventional hydraulic press revealed that the as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was easy to be hot forged. As shown in Fig. 6, it was mainly composed of β phase and γ phase when the alloy was heated to the temperature of 1150 °C. It was reported that β phase improved hot workability of TiAl alloy dramatically for its disordered BCC structure and the dynamic softening mechanism of γ phase was DRX [9,53]. As the α2+γ lamellar colony had higher yield stress, the absence of α2 phase which always formed α2+γ lamellar colony further reduced the deformation resistance [54]. The excellent hot workability was the result of large amount of disordered β phase and the absence of α2+γ lamellar colony at 1150 °C. Fig. 7 represents the tensile stress-strain curves at different temperatures with a constant strain rate of 1 × 10−4 s−1. Tensile strength of Ti-44Al-4Nb-4V-0.3Mo-Y alloy at 700 °C was about 550 MPa, while the temperature increase to 850 °C, tensile strength decreased to 400 MPa and elongation increased to 13%. The ductile-brittle transition temperature (DBTT) of as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy with a strain rate of 1 × 10−4 s−1 was between 800 °C and 850 °C. Under the DBTT, the deformation mechanism was dislocation glide which was difficult to be operated in ordered intermetallics alloy. While with the temperature increase to DBTT, the activated atomic diffusivity which improved dislocation climb and grain boundary sliding. These were reasons why the elongation increased significantly when the tensile temperature was higher than DBTT. In addition, the DBTT of TiAl was significantly influenced by the alloy component and microstructure, e.g. the DBTT of duplex structure was lower than that of full lamellar structure and near lamellar structure [55]. In other words, the higher the volume fraction of α2+γ lamellar colony, the higher the DBTT was. As indicated by Fig. 2, the microstructure of as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy was mainly composed of β/β0 phase, γ phase, and small amount of α2+γ lamellar colonies. Compared with TNB and TNM alloy,
temperature, the higher value of the peak stress is. The true stress-true strain curves exhibited peak stress at the initial deformation stage and then decreased gradually to a steady-state with increasing strain, which was the outcome of competition of work hardening and dynamic softening. In the preliminary stage, the density of dislocations increased dramatically with the increase of strain, and work hardening was the predominant factor, showing in the curve of flow stress of that the stress rapidly increased with small strain increase. The dynamic softening, such as dynamic recovery (DRV) and dynamic recrystallization (DRX), occurred with the increase of dislocation density under the condition of the thermal activation at high temperature [47]. The curves exhibited peak stress when the effectiveness of work hardening was equivalent to dynamic softening. With the increase of strain, dynamic softening become dominant and the curve decreased till reaching a steady-state flow. The lower the temperature, the harder the diffusion was, and then the harder the dynamic softening occurred. So, the highest peak stress was obtained at 800 °C. As observed in Fig. 5(c), β/β0 phase was broken down during stretching progress deformed at 800 °C. As we known, grain fragmentation was in favor of DRX [48]. When β/β0 phase stretched to a certain degree, it would be broken down and dynamic crystallized, and released stress. And then with the increase of strain, stress increases again. The fluctuation of the curve reflected the repeatedly broken down of β/β0 phase. The dynamic crystallized β/β0 phase were still in its fine grain size for the temperature was only 800 °C. When temperature was higher than 1500 °C, β phase became dominant phase, shown in Fig. 6 (the order-disorder transformation temperature of β phase is about 1210 °C [49]). The disordered β phase is helpful to improve hot workability and decrease deformation resistance [50]. The transformation of β→γ in β phase also reduced stress concentration [51]. These were why the curves at 1190 °C and 1250 °C were steady and have lower peak stresses than that of 800 °C. The influence of strain rate was on the opposite of temperature. The dislocation density increased rapidly, and the effect of work hardening was noticeable at high strain rate [52]. Then the peak stress increased with the increase of strain rate. The as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was successfully hot forged at 1150 °C with a calculated strain rate of 0.02 s−1 by a conventional hydraulic press. Pancakes of as-forged Ti-44Al-4Nb-4V371
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Fig. 6. Microstructure of as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy heat treated at 1150 °C holding 2 h and water cooled: (a) low-magnification microstructure in BSE mode; (b) high-magnification microstructure in BSE mode.
Foundation of China (Nos. 51604191, 51504163, 51604159), the State Key Laboratory for Advanced Metal and Materials foundation (Nos. 2014-ZD06) and the Special/Youth Foundation of Taiyuan University of Technology (Nos. 2015QN014, 2013T004 and 2013T003). We also thank the financially supporting of the Qualified Personnel Foundation of Taiyuan University of Technology (Nos. tyutrc201342a and tyutrc201343a). References [1] [2] [3] [4] [5] [6]
[7] [8]
Fig. 7. Typical tensile stress-strain curves of as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy at different temperature with a constant strain rate of 1 × 10−4 s−1.
[9]
the microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy consisted of less α2+γ lamellar colony which strongly increased the DBTT, and more pearlite-like crooked β/β0+γ structure [56]. Then the DBTT of this ascast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was lower, and the elongation was relatively larger than the other as-cast alloy at 850 °C. As illustrated by Fig. 2, the pearlite-like β/β0+γ structure was composed of lots of fine β/β0 laths and γ laths, which was usual as one component in as-forged beta-gamma TiAl alloy [27]. The pearlite-like β/β0+γ structure may be the reason why the strength was higher than the other beta-gamma alloy in as-cast condition such as Ti-43Al-9V alloy [57].
[10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusion
[20]
1. Based on the study of the β phase stabilizing ability of Mo, Cr, V, and Nb element and the equivalence formulas of 1Mo=4.25 Nb, and 1Mo=2.66 V, a new beta-gamma TiAl alloy of Ti-44Al-4Nb-4V0.3Mo-Y alloy was fabricated. 2. The mixture structure of γ and β/β0 was the main component of the as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy. The addition of Mo, into a V and Nb element containing beta-gamma TiAl alloy improved its oxidation resistance at 800 °C even on the level of TNM alloy. 3. Possessing perfect hot workability, the as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy could be hot forged at 1150 °C by a conventional hydraulic press without any protection easily. The elongation of ascast microstructure was 13% at 850 °C and the tensile strength was about 550 MPa at 700 °C.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]
[33] [34]
Acknowledgments
[35]
This work was financially supported by the National Natural Science 372
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The microstructure, mechanical properties, and oxidation behavior of betagamma TiAl alloy with excellent hot workability ⁎
MARK
⁎⁎
S.Z. Zhanga,b, , Y.B. Zhaoa,b, C.J. Zhanga,b, J.C. Hanc, , M.J. Sund, M. Xua,b a
School of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China Shanxi Key Laboratory of Advanced Magnesium-based Materials, Taiyuan University of Technology, Taiyuan 030024, PR China c School of Mechanical Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China d Shanghai Space flight Precision Machinery Institute, Shanghai 200000, PR China b
A R T I C L E I N F O
A B S T R A C T
Keywords: TiAl alloys Hot deformation Microstructure Oxidation Mechanical properties
New type beta-gamma TiAl alloy of nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y (at%) was fabricated by induction skull melting, which was designed based on the study of the microstructural effects of Mo, Cr, V, and Nb on TiAl alloy. The microstructure, oxidation resistance, hot workability, and tensile properties were studied by SEM in BSE mode, XRD, Gleeble 3800, as well as Instron 5500 R, respectively. The oxidation resistance of this new designed Mo containing beta-gamma TiAl alloy was better than that of only Nb and V containing betagamma TiAl alloy. With a better hot workability at temperature under 1200 °C, Ti-44Al-4Nb-4V-0.3Mo-Y alloy was successfully hot forged at 1150 °C with a traditional hydraulic machine without thermal insulation measure. Though containing abound order β0 phase, the new fabricated beta-gamma TiAl alloy still has a high elongation at 850 °C, because of less α2+γ lamellar colonies containing. Benefits of partly replacing Mo with V was avoiding omega phase precipitation, and not damaging the hot workability.
1. Introduction TiAl-based alloy possessed outstanding properties such as low density, high specific strength and modulus, superior oxidation resistance, and creep resistance, which was considered to be one of the most promising high-temperature structural materials used as turbine blades of aero-engines, turbocharger turbine wheels of automotive combustion engines, etc. [1,2]. However, TiAl alloy was difficult to be processed into component parts due to their unsatisfactory room temperature ductility and poor hot deformability, which seriously restricted its practical application [3]. Lots of work have been done to improve the ductility and workability of TiAl alloy, especially the alloy design [4–8]. Kim et al. [6] came up with a concept of novel beta-gamma TiAl alloy and its composition range was: Ti-(40−45)Al-(2−7)Nb-(1−9) M-(0–0.5) N, where M represents β phase stabilizing elements, such as Nb, Cr, Mn, V, and Mo, while N represents grain refiner elements, such as C, B, and Y elements. As BCC crystal structure had sufficient number of independent slip systems, disordered β phase was softer than γ and α phase at elevated temperature and improved the deformability of TiAl alloy (the ordered phase of β is β0) [9,10]. Moreover, beta-gamma TiAl alloy could effectively eliminate composition segregation and ⁎
microstructure inhomogeneity, reduce cast texture, and refine microstructure via β phase solidification, which avoided peritectic reaction [11]. An alloy design concept had been put forward for beta-gamma TiAl alloy by Clemens [8,12,13]. Based on this concept, TNM alloy with composition of Ti-43Al-4Nb-1Mo-0.1B was fabricated [12]. TNM alloy exhibited an appropriate volume fraction of β phase, solidified completely via β phase, possessed balanced mechanical properties, and had been machined into a turbine blade component successfully after a three-step forging process [11–13]. Besides, Ti-43Al-9V and Ti-42Al5Mn were also typical beta-gamma TiAl alloys and were confirmed with excellent hot working properties [14,15]. Alloying elements, such as Nb, Mo, and V elements, influenced different properties of TiAl alloy [16]. With development of TiAl alloy towards higher temperature and higher performance, Nb became the indispensable alloying element of TiAl alloy [17]. However, Nb could raise the diffusion activation energy thereby damaging the hot deformability of the TiAl alloy [18]. The Mo element addition improved the room temperature ductility, microstructure stability, and oxidation resistance of the TiAl alloy [19]. However, omega phase with lower crystal symmetry, which was extremely brittle and was detrimental to mechanical performance of TiAl, could be produced during the decomposition of β/β0 phase in high Nb and Mo containing TiAl alloy
Corresponding author at: School of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, PR China. Corresponding author. E-mail addresses: [email protected] (S.Z. Zhang), [email protected] (J.C. Han).
⁎⁎
http://dx.doi.org/10.1016/j.msea.2017.06.025 Received 1 February 2017; Received in revised form 6 June 2017; Accepted 7 June 2017 Available online 09 June 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
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Fig. 1. Scanning electron micrographs taken in backscattering electron mode: (a) Ti-44Al alloy; (b) Ti44Al-2Mo alloy; (c) Ti-44Al-4Mo alloy; (d) Ti-44Al4Cr alloy; (e) Ti-44Al-4V alloy; (f) Ti-44Al-4Nb alloy.
were studied. And a novel beta-gamma TiAl alloy with a nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y was designed and fabricated. The microstructure characteristic of this alloy co-added with Nb, V, and Mo was discussed. As Mo was partly replaced of V, the hot workability and oxidation resistance of Ti-44Al-4Nb-4V-0.3Mo-Y were investigated to understand the effect of co-addition of V and Mo. Moreover, high temperature tensile properties of this alloy were investigated as well.
such as TNM alloy and high Nb containing TiAl alloy (TNB alloy) [20–22]. Then the total content of Nb and Mo in beta-gamma TiAl alloy should be decreased further. As a β phase stabilizer, V improved the hot workability of TiAl dramatically but damaged its oxidation resistance [23–26]. Omega phase cannot be founded in high V containing TiAl alloy (Ti-43Al-9V and Ti-42Al-10V) when the content of Al was no less than 40% [26,27]. Moreover, the volume fraction of the stabilized β phase should be in an appropriate value for the sake of being removed by heat-treatment after hot-working, since the appearance of ordered β0 phase deteriorates the creep resistance, room-temperature strengths, and ductility [12]. In order to avoid the decrease of mechanical properties induced by omega phase and to maintain a better hot workability, the total containing of Nb and Mo should be partly replaced by other β stabilizing elements [21]. Hence, the β phase stabilizing ability of element needs be considered, and the effect of co-addition of Nb, V, and Mo on microstructure and mechanical properties has been still unknown. In this paper, the microstructural effects and β phase stabilizing ability of Mo, Cr, V, Nb on Ti-44Al (in at%, similarly hereinafter) alloy
2. Experimental procedure Six sets of alloys with nominal compositions of Ti-44Al-(0, 2Mo, 4Mo, 4Cr, 4 V, 4 Nb) were prepared using nonconsumable vacuum arc melting. Pancake ingots with 60 g were melted four times to ensure chemical homogeneity. Samples for microstructural analysis were cut from the center of as-cast pancakes. Ingot of Ti-44Al-4Nb-4V-0.3Mo-Y alloy were fabricated by induction skull melting (ISM) technique. Cubic specimens with a dimension of 10 × 10 × 10 mm3 were cut for oxidation tests by using electrical-discharge method from the ingots of Ti-44Al-4Nb-4V-0.3Mo-Y alloy. The specimen surfaces were ground 367
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to 1000 grit using SiC papers, followed by ultrasonic cleaning for 15 min in acetone and methanol and dried for oxidation tests (ultrasonic cleaning was performed by JP-020S ultrasonic cleaning machines from SKYMEN made in China). The short-time oxidation tests were performed at 800 °C, in static air of 5–80 h using a high temperature furnace (KSL-1200X-J from HF-Kejing made in China). Each specimen weight was measured by Ruipin FA1004A electronic balance (China, accuracy, ± 0.1 mg). Samples for high temperature compression tests were cut with dimensions of φ8 mm×12 mm and the surfaces were ground. High temperature compression experiments were carried out on Gleeble 3800 machine at temperatures 800 °C, 1190 °C, and 1250 °C up to a high reduction of 70% with strain rates of 1–0.01 s−1. Samples for hot forged were cut from the cast ingot and heated at 1150 °C for about 40 min. Then samples were hot forged in a conventional hydraulic press with a strain rate of 0.02 s−1 (the strain rate was calculated with the . . equation of ε = ε / t , where ε is strain rate, ε is strain, and t is deformation time). Tensile test samples with a gauge section of 5 × 2 mm2 cross-section and a gauge length of 20 mm were wire-cut parallel to the length of ingot. High temperature tensile tests were performed at temperatures of 700 °C, 800 °C, and 850 °C at strain rates of 5 × 10−4 s−1 using an Instron 5500 R testing machine. The microstructural characterization was performed by Quanta 200FEG field-emission scanning electron microscopy (FESEM) using backscattered electrons mode (BSE). Microstructural specimens were mechanical ground and polished. Phase analyses were performed by Xray diffraction (XRD) on D/max−2500 pc (RIGAKU, Japan) in the range 20–100°(2θ) using a step size of 0.02°. 3. Results and discussion
located both inside and at the boundaries of the lamellar colony was shown in Fig. 1(d). The mixed structure was produced by the cellular reaction β→β+γ during solidification [32]. As shown in Fig. 1(e), there exists a little amount of β/β0 phase in the microstructure of Ti44Al-4V alloy, and small amount of β/β0 phase could not refine the microstructure of as-cast TiAl alloy and then could not improve its hot workability. It was revealed that the β phase stabilizing ability was weaker than that of Mo element. As shown in Fig. 1(f), the microstructure of Ti-44Al-4Nb alloy was full lamellar structure without β/β0 phase containing, which illustrated that Nb was the weakest β phase stabilizer among Mo, V, Cr and Nb. According to the given in analysis, it was reasonable to surmise that the β phase stabilizing ability was Mo > Cr > V > Nb, which was in accordance with the previous research [33]. It was also obvious that the microstructure and β/β0 phase containing had no significant change comparing with Ti-44Al alloy, when the single addition of Nb and V was no more than 4% in TiAl alloy. And the result also revealed that the high containing of β/β0 phase in the Widmannstatten structure of Ti44Al-4Mo alloy did not meet the design concepts of beta-gamma phase TiAl alloy. Since, Nb could dramatically improve the high temperature strength and oxidation resistance of TiAl alloy, it was a necessity for TiAl alloy [34]. As was presented in literature [35], the ratio of β phase stabilizing ability between Mo, V, Nb was 4.25:1.6:1. To obtain the same β/β0 phase content in TiAl ally with an equal Al containing, the atomic ratio of Mo, V, and Nb was 4.25:1.6:1 when they were single added. In order to enhance the high temperature properties of TiAl alloy, avoid omega phase producing, and improve hot workability, a new beta-gamma TiAl alloy was designed with a nominal composition of Ti-44Al-4Nb-4V-0.3Mo-Y which has an equivalent weight of about 2Mo with TNM alloy (the equivalence formula is 1Mo=4.25 Nb, and 1Mo=2.66 V according to literature [33]).
3.1. Microstructure of Ti-44Al-xM alloy
3.2. As-cast microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy
In order to evaluate the β phase stabilizing ability of β stabilizing elements and their effects on the microstructure of TiAl alloy, the ascast microstructure of Ti-44Al-xM (where, x=0,4; M represented Mo, Cr, V, Nb) alloys were studied (Fig. 1). In BSE mode, β/β0 phase has the brightest contract, α/α2 phase is grey while the contrast of γ phase is black [28]. It was not difficult to distinguish these phases by contrast. As shown in Fig. 1(a), Ti-44Al alloy with a larger colony size had a full lamellar structure and was mainly composed of (α2+γ) colonies which had an average size of about 100 µm. The small colony size was the result of β phase solidification which avoided the peritectic transformation in traditional TiAl alloy which even had a colony size of 1000 µm [29,30]. Fig. 1(b) and (c) showed the microstructures of Ti44Al-2Mo and Ti-44Al-4Mo alloy. It was shown that a large amount of β/β0 phase distributed at (α2+γ) lamellar colony boundaries and inside the (α2+γ) colonies, which divided the large lamellar colonies into many fine lamellar grains, when the Mo content was 2%. It was noticed that there existed three orientations of the lamellar structure which were at 90°, 45°, and parallel to the primary growth direction of the β dendrite (the red line, blue line and black line in Fig. 1(b) represented the angle of 0°, 45°, and 90 °, respectively). As the precipitation of α phase at β phase followed an as-called Burgers relationship which was and [001]α //[110]β , then the (α2+γ) lamellar structure formed by the reaction of α→α2+γ following Blackburn relationship has the above three angles with the remnant β/β0 phase [31]. However, with the Mo content increased to 4%, Widmannstatten β/β0 phase distributed around the needle-like proeutectoid γ phase and became the dominating phase [23]. This microstructure damaged the mechanical property of alloy, and the content of β/β0 phase exceeded the betagamma TiAl alloy design concept and could not be eliminated by subsequent process [12]. Therefore, the content of Mo element could be no more than 4% when design a beta-gamma TiAl alloy. In the microstructure of the Ti-44Al-4Cr alloy, the mixture structure of γ and β/β0
Fig. 2(a) was the as-cast microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy. In this condition, the alloy exhibited a nearly lamellar microstructure which consisted of relatively small α2+γ lamellar colonies (grey contrast) with the size of 30–100 µm, and mixture structure of γ (black contrast) and β/β0 (light grey contrast) grains located along the lamellar colony boundaries. As shown in Fig. 2(b), the α2+γ lamellar colony was marked with green letters. XRD result was shown in Fig. 2(c) also confirmed that the alloy was mainly composed of γ, β/β0, and α2 phases. On the one hand, fine lamellar colonies were ascribed to the completely β phase solidification which could form as many as 12 possible orientation variants of α phase from β according to Burgers relationship [8]. On the other hand, the remnant β phase stabilized by Mo, V, and Nb elements was rejected to the interface of the different orientation α phase and then was separated along the boundaries, which impeded the growth of the Burgers α grains [8]. A substantial amount of β phase which was stabilized to room temperature by the addition of Nb, V, and Mo with a volume fraction of 37% was equal to that of TNM alloy and Ti-43Al-9V alloy [36,37]. It was illustrated by the Ti-Al-(4−10)% Nb-(V, Mo) polythermal sectional diagram that the solidifying pathway of the present alloy was: L→L+β→β→β+α→ α+β+γ→β+γ→α2+β/β0+γ, with no single-phase α-region existing [34]. As shown in Fig. 2(b), except for the fine spacing α2+γ lamellar colonies, there was also large amount of mixed structures of coarse β/β0 and γ, as well as crooked β/β0+γ lamellar morphology colonies. It was noted that two kinds of β/β0+γ structure morphology were shown in Fig. 2(b) marked with red and yellow letters. The β/ β0+γ colony structure marked with red letters was formed by the transformation of α→β+γ in a pearlitic mode procedure [23]. Whereas, the mixed structure of β/β0 and γ marked with yellow letters was the result of the cellular reaction of β→β+γ whose driving force was the chemical disequilibrium when β phase stabilizing elements segregated to β/α-interface in the course of the β→α transformation 368
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Fig. 2. As-cast microstructure and XRD result of Ti44Al-4Nb-4V-0.3Mo-Y alloy: (a) and (b) as-cast microstructure in BSE mode; (c) XRD result.
literatures [39–42]. Five alloys showed similar oxidation kinetics including linear growth stage and parabolic growth stage. Generally, Nb containing alloys exhibited a slower mass gain than the Nb free alloy at the same oxidation time, and V had the opposite effect on mass gain. In the initial oxidation stage, the mass gain rate of Ti-44Al-4Nb-4V-0.3MoY alloy was slow, but this stage can continue for 15 h, indicating that the oxidation scale yielded in this stage was protective for the alloy. The linear growth stage of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was almost on the same level as Ti-45Al-8Nb and Ti-43.5Al-4Nb-1Mo-B alloy. After 80 h, the mass gain of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was about 1.92 mg/ cm2, which was a little higher than that of Ti-45Al-8Nb and Ti-43.5Al4Nb-1Mo-B alloy, whereas oxidation resistance was significantly improved comparing with Ti-45Al-4.5Nb-4.5V-Y and Ti-45Al-3Nb-6V-Y alloy. Though Ti-44Al-4Nb-4V-0.3Mo-Y alloy was a V containing alloy, oxidation resistance ability was improved dramatically by the amount of 0.3 at% Mo element addition. It was pointed out that the continuous Al3O2 layer could hinder the diffusion of oxygen and prevent the formation of TiO2 which deteriorated oxidation resistance [43]. Networked Al2O3 layer was formed on the external surface of Mo addition TiAl alloy, and the second Al2O3 layer from the outer surface was more continuous and stable than that of Mo free TiAl alloy [44]. It was obvious that the amount of 0.3 at% Mo element addition in Ti-43Al-4Nb4V-0.3Mo-Y alloy improved the formation of Al2O3 layer which enhanced the oxidation resistance dramatically than free Mo containing TiAl alloy.
boundaries [13]. Bright particles appeared in Fig. 2(b) may be determined as the YAl2 phase according to the previous studies [38]. The YAl2 phase cannot be found by XRD because the phase is too rare to be detected. 3.3. High-temperature oxidation behavior In order to assess the high temperature oxidation resistance of the Ti-44Al-4Nb-4V-0.3Mo-Y alloy, its cyclic oxidation kinetics curve was plotted together with other alloys in Fig. 3 for comparison. The oxidation data of Ti-45Al-8Nb, Ti-45Al-4.5Nb-4.5V-Y, Ti-45Al-5.4V3.6Nb-0.3Y, Ti-43.5Al-4Nb-1Mo-B alloy were published in the existed
3.4. Hot-working behavior In order to evaluate the hot workability of Ti-44Al-4Nb-4V-0.3Mo-Y alloy, high temperature compression tests were conducted, and the true stress-true strain curves were shown in Fig. 4. As shown in Fig. 4(a) and (b), the highest stress of 260 MPa was obtained during hot compressed at 1190 °C with a strain rate of 1 s−1. The peak stresses of this alloy and TNM alloy for the given temperature were summarized in Table 1
Fig. 3. Cyclic oxidation kinetics curves for five different TiAl alloys. The added data of Ti45Al-8Nb, Ti-45Al-4.5Nb-4.5V-Y, Ti-45Al-5.4V-3.6Nb-0.3Y, Ti-43.5Al-4Nb-1Mo-B alloy was from Refs. [39–42] for comparison.
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Fig. 4. True stress-true strain curves of Ti-44Al-4Nb-4V-0.3Mo-Y alloy at different deformation temperature: (a) 1190 °C with strain rates of 0.01 s−1, 0.05 s−1, 0.1 s−1, 0.5 s−1, and 1 s−1; (b) 1250 °C with different strain rate of 0.01 s−1, 0.05 s−1, 0.1 s−1, 0.5 s−1, and 1 s−1; (c) 800 °C with 0.05 s−1.
(1190 °C) and 1250 °C were listed in Table 1. However, peak stress could reach as high as 740 MPa when deformed at 800 °C with 0.05 s−1 shown in Fig. 4(c). The stress curves shown in Fig. 4 indicated that Ti44Al-4Nb-4V-0.3Mo-Y alloy possessed good high temperature deformability and high strength at lower temperature, especially at the temperature around and even lower than 1200 °C. And the alloy could be hot deformed with lower deformation resistance and good deformability at temperatures around 1200 °C. The microstructure of samples hot compressed at 800 °C, 1190 °C and 1250 °C with strain rate of 0.05 s−1 was shown in Fig. 5. As shown in Fig. 5(a) and (c), microstructures of samples hot compressed at 1190 °C and 800 °C were typical deformed structure with flow lines which were perpendicular to external force axis (external force axis was indicated by red arrow in Fig. 5). β/β0 phase was stretched more seriously and was broken down into discontinuous structure somewhere in Fig. 5(c). Acicular γ phase precipitated in β/β0 phase when deformation temperature elevated to 1190 °C. When temperature increase to 1250 °C, block γ phase and acicular γ phase were observed in Fig. 5(b), and flow lines were not as clear as Fig. 5(a). Obviously, the flow stress was highly sensitive to strain, temperature, and strain rate. The higher the strain rates and the lower the
Table 1 The peak stress of TNM alloy and Ti-44Al-4Nb-4V-0.3Mo-Y alloy. Strain rate
Temperature 1200 °C*
0.5 s−1 0.05 s−1 0.005 s−1*
275 MPa (TNM) 180 MPa (TNM) 110 MPa (TNM)
1250 °C 183 MPa 114 MPa 90 MPa
160 MPa (TNM) 60 MPa (TNM) 30 MPa (TNM)
111 MPa 67 MPa 36 MPa
State: 1200 °C* represented that TNM alloy was deformed at 1200 °C, while Ti-44Al-4Nb4V-0.3Mo-Y alloy was deformed at 1190 °C; 0.005 s−1* represented that TNM alloy was deformed at 0.005 s−1, while Ti-44Al-4Nb-4V-0.3Mo-Y alloy was deformed at 0.01 s−1.
[45–47]. It was obvious that the peak stress at 1190 °C of this alloy was lower than that of TNM at 1200 °C [47]. For instance, the peak stress of TNM alloy deformed at 1200 °C with strain rate of 0.5 s−1 was 275 MPa [47], while the peak stress of Ti-44Al-4Nb-4V-0.3Mo-Y alloy was 183 MPa which was two thirds of TNM alloy at 1200 °C with strain rate of 0.5 s−1. Compared to TNM alloy, the peak stress of Ti-44Al-4Nb4V-0.3Mo-Y alloy was for a maximum of 33% decrease. The peak stresses of Ti-44Al-4Nb-4V-0.3Mo-Y alloy and TNM alloy at 1200 °C
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Fig. 5. Microstructure of samples hot compressed with strain rate of 0.05 s−1 at different temperatures: (a) 1190 °C; (b) 1250 °C; (c) 800 °C. The arrow indicated the orientation of external force which was perpendicular to flow lines in (a) and (c).
0.3Mo-Y alloy were shown a smooth and crack-free appearance. As the equipment was a conventional hydraulic press without any protective atmosphere and the forged sample was not protected by other technique, such as canned with stainless steel to acquire near-isothermally and prevent oxidation, the surface of the pancakes were coated with thin oxidation layer during heated at1150°C for 40 min. Successfully hot forging at 1150 °C without any protection on a conventional hydraulic press revealed that the as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was easy to be hot forged. As shown in Fig. 6, it was mainly composed of β phase and γ phase when the alloy was heated to the temperature of 1150 °C. It was reported that β phase improved hot workability of TiAl alloy dramatically for its disordered BCC structure and the dynamic softening mechanism of γ phase was DRX [9,53]. As the α2+γ lamellar colony had higher yield stress, the absence of α2 phase which always formed α2+γ lamellar colony further reduced the deformation resistance [54]. The excellent hot workability was the result of large amount of disordered β phase and the absence of α2+γ lamellar colony at 1150 °C. Fig. 7 represents the tensile stress-strain curves at different temperatures with a constant strain rate of 1 × 10−4 s−1. Tensile strength of Ti-44Al-4Nb-4V-0.3Mo-Y alloy at 700 °C was about 550 MPa, while the temperature increase to 850 °C, tensile strength decreased to 400 MPa and elongation increased to 13%. The ductile-brittle transition temperature (DBTT) of as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy with a strain rate of 1 × 10−4 s−1 was between 800 °C and 850 °C. Under the DBTT, the deformation mechanism was dislocation glide which was difficult to be operated in ordered intermetallics alloy. While with the temperature increase to DBTT, the activated atomic diffusivity which improved dislocation climb and grain boundary sliding. These were reasons why the elongation increased significantly when the tensile temperature was higher than DBTT. In addition, the DBTT of TiAl was significantly influenced by the alloy component and microstructure, e.g. the DBTT of duplex structure was lower than that of full lamellar structure and near lamellar structure [55]. In other words, the higher the volume fraction of α2+γ lamellar colony, the higher the DBTT was. As indicated by Fig. 2, the microstructure of as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy was mainly composed of β/β0 phase, γ phase, and small amount of α2+γ lamellar colonies. Compared with TNB and TNM alloy,
temperature, the higher value of the peak stress is. The true stress-true strain curves exhibited peak stress at the initial deformation stage and then decreased gradually to a steady-state with increasing strain, which was the outcome of competition of work hardening and dynamic softening. In the preliminary stage, the density of dislocations increased dramatically with the increase of strain, and work hardening was the predominant factor, showing in the curve of flow stress of that the stress rapidly increased with small strain increase. The dynamic softening, such as dynamic recovery (DRV) and dynamic recrystallization (DRX), occurred with the increase of dislocation density under the condition of the thermal activation at high temperature [47]. The curves exhibited peak stress when the effectiveness of work hardening was equivalent to dynamic softening. With the increase of strain, dynamic softening become dominant and the curve decreased till reaching a steady-state flow. The lower the temperature, the harder the diffusion was, and then the harder the dynamic softening occurred. So, the highest peak stress was obtained at 800 °C. As observed in Fig. 5(c), β/β0 phase was broken down during stretching progress deformed at 800 °C. As we known, grain fragmentation was in favor of DRX [48]. When β/β0 phase stretched to a certain degree, it would be broken down and dynamic crystallized, and released stress. And then with the increase of strain, stress increases again. The fluctuation of the curve reflected the repeatedly broken down of β/β0 phase. The dynamic crystallized β/β0 phase were still in its fine grain size for the temperature was only 800 °C. When temperature was higher than 1500 °C, β phase became dominant phase, shown in Fig. 6 (the order-disorder transformation temperature of β phase is about 1210 °C [49]). The disordered β phase is helpful to improve hot workability and decrease deformation resistance [50]. The transformation of β→γ in β phase also reduced stress concentration [51]. These were why the curves at 1190 °C and 1250 °C were steady and have lower peak stresses than that of 800 °C. The influence of strain rate was on the opposite of temperature. The dislocation density increased rapidly, and the effect of work hardening was noticeable at high strain rate [52]. Then the peak stress increased with the increase of strain rate. The as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was successfully hot forged at 1150 °C with a calculated strain rate of 0.02 s−1 by a conventional hydraulic press. Pancakes of as-forged Ti-44Al-4Nb-4V371
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Fig. 6. Microstructure of as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy heat treated at 1150 °C holding 2 h and water cooled: (a) low-magnification microstructure in BSE mode; (b) high-magnification microstructure in BSE mode.
Foundation of China (Nos. 51604191, 51504163, 51604159), the State Key Laboratory for Advanced Metal and Materials foundation (Nos. 2014-ZD06) and the Special/Youth Foundation of Taiyuan University of Technology (Nos. 2015QN014, 2013T004 and 2013T003). We also thank the financially supporting of the Qualified Personnel Foundation of Taiyuan University of Technology (Nos. tyutrc201342a and tyutrc201343a). References [1] [2] [3] [4] [5] [6]
[7] [8]
Fig. 7. Typical tensile stress-strain curves of as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy at different temperature with a constant strain rate of 1 × 10−4 s−1.
[9]
the microstructure of Ti-44Al-4Nb-4V-0.3Mo-Y alloy consisted of less α2+γ lamellar colony which strongly increased the DBTT, and more pearlite-like crooked β/β0+γ structure [56]. Then the DBTT of this ascast Ti-44Al-4Nb-4V-0.3Mo-Y alloy was lower, and the elongation was relatively larger than the other as-cast alloy at 850 °C. As illustrated by Fig. 2, the pearlite-like β/β0+γ structure was composed of lots of fine β/β0 laths and γ laths, which was usual as one component in as-forged beta-gamma TiAl alloy [27]. The pearlite-like β/β0+γ structure may be the reason why the strength was higher than the other beta-gamma alloy in as-cast condition such as Ti-43Al-9V alloy [57].
[10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
4. Conclusion
[20]
1. Based on the study of the β phase stabilizing ability of Mo, Cr, V, and Nb element and the equivalence formulas of 1Mo=4.25 Nb, and 1Mo=2.66 V, a new beta-gamma TiAl alloy of Ti-44Al-4Nb-4V0.3Mo-Y alloy was fabricated. 2. The mixture structure of γ and β/β0 was the main component of the as-cast Ti-44Al-4Nb-4V-0.3Mo-Y alloy. The addition of Mo, into a V and Nb element containing beta-gamma TiAl alloy improved its oxidation resistance at 800 °C even on the level of TNM alloy. 3. Possessing perfect hot workability, the as-cast Ti-44Al-4Nb-4V0.3Mo-Y alloy could be hot forged at 1150 °C by a conventional hydraulic press without any protection easily. The elongation of ascast microstructure was 13% at 850 °C and the tensile strength was about 550 MPa at 700 °C.
[21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32]
[33] [34]
Acknowledgments
[35]
This work was financially supported by the National Natural Science 372
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