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Influence of heat treatment on microstructure and tensile property of a new high strength beta alloy Ti–2Al–9.2Mo–2Fe
Materials Science & Engineering A 580 (2013) 250–256
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Influence of heat treatment on microstructure and tensile property of a new high strength beta alloy Ti–2Al–9.2Mo–2Fe Cheng-Lin Li a,b, Xu-Jun Mi a, Wen-Jun Ye a, Song-Xiao Hui a, Dong-Geun Lee b,n, Yong-Tai Lee b a b
State Key Laboratory for Nonferrous Metals & Process, General Research Institute for Nonferrous Metals, Beijing 100088, PR China Titanium Department, Korea Institute of Materials Science, Changwon 642831, South Korea
art ic l e i nf o
a b s t r a c t
Article history: Received 23 December 2012 Received in revised form 27 April 2013 Accepted 29 April 2013 Available online 17 May 2013
An investigation is given on the influence of heat treatment by the microstructural characteristics and tensile properties of a new high strength alloy Ti–2Al–9.2Mo–2Fe. Both of the α/β and β solution treatment (α/β-ST and β-ST), then aged at temperatures ranging from 400 1C to 600 1C, were prepared. The primary α phases having 2–5 μm are formed during the α/β-ST, and restrain the size of β grains below 10 μm. As a result of the fine β grains, the α/β-ST contributes a higher strength than the β-ST. The coexistence of α″ and athermal ω phase is found in the β-ST and water quenched samples. However, these have little influence on the alloy hardening. After aging, the alloys in the α/β-ST and β-ST condition reveal the phase transformation of β to isothermal ω, and β to α depending on the aging temperature. Although the primary α phase formed during the α/β-ST increases the stability in the β matrix, and the isothermal ω phase also appears to occur during aging at 400 1C and 450 1C for 2 h. These phenomena are less common in beta titanium alloys, when treated in the α/β-ST and aged at lower temperatures. The isothermal ω phase formed in both conditions results in high strength levels (1600 MPa of ultimate strength) with much ductility loss (2.5–4.5% of elongation) as a result of the superior hardening effect and brittleness of ω phase. However, the secondary α phase with the size of 1–3 μm leads to attractive combinations of strength and ductility (1200–1400 MPa of ultimate strength with 7.5–12.5% of elongation). The reason for that is too fine α phase below 1 μm tends to result in ultra-high strength with much ductility loss. As a whole, the alloy can be heat treated to obtain excellent balances of strength and ductility, and provided abundant stress levels with optional ductility as a usable material. & 2013 Elsevier B.V. All rights reserved.
Keywords: Omega and alpha phase Age strengthening Tensile property Beta titanium alloy
1. Introduction Beta titanium alloys have the highest strength-to-weight ratio among all titanium alloys (alpha, alpha/beta and beta alloys). The beta alloys are attractive as a structural material for applications, such as automotive parts and surgical orthopedic implants besides the conventional use in the aircraft industry [1–3]. The beta titanium alloys are typically formed or fabricated in the solution treated (ST) condition because they are ductile with low yield strength ranging from 700 MPa to 1000 MPa. By the heat treatment of aging they will show high strength (1200–1700 MPa). In general, most of the increased strength due to aging results from the precipitation of fine α phase in the β matrix [4]. The fine α phases are hard and un-deformable particles to hinder the movement of dislocations and thus strengthen the alloys. This process is known as precipitation hardening in the STA condition [5,6]. The
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Corresponding author. Tel.: +82 55 280 3367; fax: +82 55 280 3255. E-mail address: [email protected] (D.-G. Lee).
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size and distribution of the α phases affect the strength level and depend on the aging time at temperature. While the prior β grain size determines the ductility [7–9]. The optimum combination of mechanical properties of beta titanium alloys can be altered via heat treatment and controlling their microstructures for desired mechanical properties [10,11]. Although the beta titanium alloys have an advantage in a unique combination of high specific strength and excellent corrosion resistance, the application is significantly restricted by the high cost of the titanium itself during melting or processing [12,13]. For these reasons, recently, the development of the new Ti alloys for non-aerospace applications have led to design of lower cost alloys, such as Timetal LCB (Ti–6.8Mo–4.5Fe–1.5Al) alloy for automotive and other industrial applications [14]. The alloy has a ultimate strength level of 1000 MPa and 80–90 GPa of elastic modulus with 18% elongation in a typical solution treated condition [15,16]. While it can be heat-treated (or aged) to over 1400 MPa of ultimate strength, 110–117 GPa of elastic modulus with 13% of elongation [9,17]. It is the low cost Mo–Fe master alloy that highlights the LCB alloy. However, it readily produces so
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called “beta flecks” due to the high content of Fe addition. The beta flecks are solute rich regions that exhibit a different microstructure from the surrounding material in the final product. Therefore, they may result in inhomogeneous aging products and unstable mechanical properties [18–20]. Given all of that, a new beta titanium alloy with low cost, high strength and low elastic modulus should be designed and expected to overcome the Timetal LCB alloy and/or other newly developed LCB alloys. A Ti–2Al–9.2Mo–2Fe alloy is a new beta titanium alloy developed to be used as structural materials such as automotive spring application. It was designed on the basis of Bo–Md molecular orbital method (Bo, the bond order; Md, the metal d-orbital energy level) [21]. However, as a newly designed alloy, the microstructural features and mechanical properties of the Ti–2Al–9.2Mo–2Fe alloy are not well understood yet. Furthermore, it is of extreme importance to understand its microstructure– mechanical properties relationship by several heat treatment methods. Therefore, the present work is undertaken to understand the relationship of microstructural characteristics and tensile properties of the Ti–2Al–9.2Mo–2Fe alloy during the α/β or β solution treatment and subsequent aging, and concurrently, try to reveal how the microstructures of the new alloy change by heat treatments and how the changes affect the tensile properties.
2. Material and experimental methods The Ti–2Al–9.2Mo–2Fe alloy was melted by multiple vacuum arc re-melting (VAR) method. The ingot was forged at around 1100 1C to a square billet with 400 400 L mm in dimensions. The square billet was then hot rolled at 800 1C to round bars with 12 mm in diameter. The chemical composition of the alloy is listed in Table 1. The molybdenum equivalent ([Mo]eq) of the alloy is about 15.2 according to its chemical composition and an empirical equation proposed by Collings [22]. The β transus temperature of the alloy is measured to be approximately 815 1C with a metallographic method. The microstructure of the as rolled alloy is shown in Fig. 1. It mainly consists of small deformed β grains and fine nano-scaled acicular α precipitates due to the rolling temperature below the β transus temperature. Samples from the as-rolled alloy were solution treated in a preheated air furnace at 790 1C and 850 1C (below and above the β transus temperature) and then held for 1 h followed by water quench. All the samples were solution treated in an air box furnace. In view of the chemical activity of titanium alloys at high temperatures, all the samples are coated with Acheson's Deltaglaze 151 protective coating to minimize the formation of α case before heat treatment. Following the solution treatment, the samples were subjected to aging at temperatures ranging from 400 1C to 600 1C for 2 h in a pre-heated air furnace and followed by water quench. Specimens for microstructure observation were mechanically polished with the standard metallographic procedure and etched in the Kroll's reagent (5% HF+15% HNO3+80% H2O) to reveal the grain boundaries. Microstructure characterization was performed on a JEOL JSM-5800 scanning electron microscopy (SEM) and a JEOL JEM-2100F transmission electron microscopy (TEM) operated at 20 kV and 200 kV, respectively. For the TEM thin foil preparation it was firstly grinded on waterproof abrasive papers to 60 μm Table 1 Chemical composition of the Ti–2Al–9.2Mo–2Fe alloy (wt%). Al
Mo
Fe
Cu
Si
C
H
O
N
Ti
2.33
9.33
1.98
0.015
0.037
0.0091
0.0013
0.054
0.0087
Bal.
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in thickness and then prepared with the MTP-1A twin jet electrochemical polishing method at −30 1C in an electrolyte consisting of perchloric acid, n-butyl alcohol and methanol (volume ratio 5%:35%:60%). Standard cylindrical tensile specimens with gauge length of 30 mm and gauge diameter of 6 mm were machined from both the solution treated and aged materials. The tensile tests were performed at room temperature with a nominal strain rate of 5.5 10–4 s−1, using a screw driven INSTRON machine of 50 kN capacity. At least three samples of each condition were tested for an average value of ultimate tensile strength (Rm), yield strength (Rp0.2) (that is, 0.2% proof stress), elongation (A) and elastic modulus (E).
3. Results 3.1. Microstructural characteristics The microstructural features of the Ti–2Al–9.2Mo–2Fe alloy, which was solution treated at below and above the β transus temperature for 1 h, are shown in Fig. 2. As seen in Fig. 2(a), there are stubby and globular primary α phases precipitated in the β matrix, as well as discontinuous primary α phases formed on and along the prior β grain boundaries. The stubby and globular primary α phases have 2–5 μm length or diameter. However, the remaining β grains exhibit a size within 10 μm. In case of the solution treatment at 850 1C condition, it results in equiaxed and coarse β grains with a mean size of about 50 μm due to recrystallization, as seen in Fig. 2(b). A feature in Fig. 2(b) shows that plate-like precipitations emerge in some β grains, which refer to martensites formed during water quench. Consequently, the martenistes are difficult to be seen with the TEM. However, a selected area diffraction (SAD) pattern from the β matrix of Fig. 2 (d) is shown in Fig. 2(c), which is consistently indexed as [113] zone axis of β phase. Additional weaker spots are visible at the 1/3 (121) and 2/3(121) positions of the β reflection where they have straight or curved lines of diffuse intensity (as known as diffuse streaking) instead of sharp spots. That refers to as the “diffuse ω phase” relating the early stages of ω phase formation [22,23]. Microstructures of the Ti–2Al–9.2Mo–2Fe alloy after the α/β-ST and aged at different temperatures are shown in Fig. 3. The first heat treatment carried out on the as-quenched sample is aged at 400 1C for 2 h. The ω precipitates are imaged from the 1/3 (or 2/3) (112) β reflections in the [113] β zone axis. As shown in Fig. 3(a), the size scale of the ω precipitates is in the range of 10–20 nm. Further aging treatments are conducted at 450 1C for 2 h. The ω precipitates clearly exhibit an ellipsoidal shape with a size scale of about 30 nm seen in Fig. 3(b). It is observed that the ω precipitates have coarsened during aging at 450 1C. However, by aging this alloy at a higher temperature like 500 1C, it is observed in Fig. 3 (c) that the fine scale α precipitates with a size of 100–200 nm shapes a more lenticular morphology. There is no evidence of ω precipitate presented from the dark field image. The microstructures of the α/β solution treated alloy aged at 550 1C and 600 1C for 2 h are shown in Fig. 3(d) and (e), respectively. The α platelet clearly appears to be present in the β matrix, and exhibits a size of about 0.5 μm for aging at 550 1C and about 1–2 μm for 600 1C. The microstructures of the Ti–2Al–9.2Mo–2Fe alloy after the β-ST plus aged at different temperatures for 2 h are shown in Fig. 4. From these figures it can be seen that the ω phases with an ellipsoidal shape form in the β matrix when the alloy was aged at low temperatures of 400 1C and 450 1C. Moreover, the ω phases obtained at 450 1C exhibits a coarser size (about 50 nm) than that at 400 1C (about 25 nm) during the same aging time. No α phase is
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Fig. 1. Microstructures of the as rolled Ti–2Al–9.2Mo–2Fe alloy: (a) OM, and (b) SEM.
Fig. 2. Microstructures of the alloy after solution treated at: (a) SEM 790 1C, (b) OM 850 1C, (c) 850 1C, an [113]β selected area diffraction pattern having extra reflections at 1/3 and 2/3 (112)β positions corresponding to athermal ω phases, (d) 850 1C dark-field image from the extra reflections to show nano-scaled ω phases.
Fig. 3. Microstructures of the alloy solution treated at 790 1C and aged for 2 h at: (a) 400 1C, (b) 450 1C, (c) 500 1C, (d) 550 1C and (e) 600 1C.
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Fig. 4. Microstructures of the alloy solution treated at 850 1C and aged for 2 h at: (a) 400 1C, (b) 450 1C, (c) 500 1C (d) 550 1C, and (e) 600 1C.
Table 2 Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy in solution treated condition. ST
Rm (MPa)
Rp0.2 (MPa)
A (%)
E (GPa)
790 1C/1 h/WQ 850 1C/1 h/WQ
957 833
853 673
20 42
61 67
detected by TEM micrographs when the alloy is aged at the two temperatures. However, only the α phase can be found in Fig. 4 (c) which corresponds to the alloy aged at 500 1C for 2 h. The lamellar α phase has the size of about 500 nm. As seen in Fig. 4 (d) and (e), very fine α platelet distributes in the β matrix, the size itself exhibits 1 μm for 550 1C in Fig. 4(d) and 3 μm for 600 1C in Fig. 4(e). In other words, the α phase has a slight increase in size as the aging temperature increased from 500 1C to 600 1C within aging for 2 h.
aged at 450 1C and 500 1C for 2 h shows an ultimate strength as high as 1600 MPa with poor ductility of below 5% of the elongation, but an increase in the aging temperature leads to a rapid decrease in strength with an increase in ductility. Aging at 600 1C tends to produce an interesting property balance (above 1200 MPa in ultimate strength with 12.5% of elongation). Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy after the β-ST plus aging at different temperatures are shown in Fig. 8. When the alloy was aged at 450 1C, it does result in an ultimate strength about 1400 MPa with a low ductility (1.5% of the elongation). Putting aside of the ω-aged alloy in terms of 450 1C, the strength basically decreases sharply with increasing the aging temperature. The highest strength level is obtained at 500 1C aging with a remarkable ductility loss. On the other hand, a good balance of strength and ductility is obtained at 550 1C and 600 1C.
4. Discussion 3.2. Tensile properties Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy, which was solution treated at 790 1C and 850 1C for 1 h, are listed in Table 2. The relative stress–strain curves of the Ti–2Al–9.2Mo–2Fe alloy in the α/β and β solution condition are shown in Fig. 5. Both of the two solution treatments specimens behave an elastic modulus of 61–67 GPa. The alloy solution treated at 790 1C for 1 h shows high yield strength of 853 MPa, and a relative small elongation of 20%. In addition, it shows a weaker work hardening phenomenon as seen in Fig. 5(a). On the contrary, the alloy solution treated at 850 1C for 1 h shows a low yield strength level of 673 MPa with a larger elongation of 42%. But in this condition, it shows a significant work hardening as seen in Fig. 5(b). The microstructures from the region near the fracture of the tensile specimens, which were pre-solution treated at 790 1C and 850 1C for 1 h, are shown in Fig. 6. The micrograph in Fig. 6(b), which refers to the alloy solution treated at 850 1C, clearly shows the presence of α″ and some other plate-like features. The former refers to the stress or strain induced martensite (SIM) and the latter corresponds to slip bands. However, the absence of SIM or twinning in Fig. 6(a), corresponding to the alloy solution treated at 790 1C, is supposed that only the slip mode dominates the deformation mechanism. Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy after the α/βST plus aging are illustrated in Fig. 7. It can be seen that the alloy
The microstructural characteristics of the alloy were studied mainly by a SEM, in some cases the SEM observations were augmented by TEM studies. The various microstructures showed pronounced differences as well as tensile behavior and were thus difficult to discuss concurrently. Therefore, this section will be divided into three parts: (1) solution treatment (ST); (2) α/β solution treatment plus aging (α/β-STA); and (3) β solution treatment plus aging (β-STA). 4.1. Solution treatment The Ti–2Al–9.2Mo–2Fe alloy is a metastable beta titanium alloy with the β transus temperature of 815 1C. In case of the beta titanium alloys, solution treatment above the β transus temperature results in coarse β grains. While, solution treatment slightly below the β transus temperature leads to precipitations of primary α phase [24]. The primary α phases obtained in the condition α/β-ST (Fig. 2 (a)) exhibit stubby and globular shapes. It is supposed that the primary α phases tend to form on the basis of the growth or globalization of the nano-scaled α phases in the as-rolled microstructure (Fig. 1), accompanying with some of them dissolved. It is well known that the important characteristic of the primary α phase concerns its ability to pin the prior β grain boundaries and reduce their mobility, which therefore limits the recrystallization
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Fig. 5. Nominal stress–strain curves of the alloys after solution treatment: (a) 790 1C and (b) 850 1C.
Fig. 6. Microstructures from the region near the fracture of the tensile deformed specimens post-solution treated at: (a) 790 1C, and (b) 850 1C. The tensile axis is perpendicular to the paper.
Fig. 7. Tensile properties of the alloy solution treated at 790 1C and aged at different temperatures for 2 h.
Fig. 8. Tensile properties of the alloy solution treated at 850 1C and aged at different temperatures for 2 h.
and grain growth of the β grain [25]. In case of β-ST, besides the equiaxed and coarse β grains obtained due to recrystallization, discontinuous distributed martensites formed in some β grains during water quench. The reason for the discontinuous distribution would be that local chemical variations or composition segregation. The segregation allows some grains to have an uneven of beta stabilizers, which leads to the difference of stability. It has been proven by an EPMA method [26–28]. It should be noted that the small volume fraction of ω phase found in the β matrix has little direct influence on the strength of the alloy, but indeed, which acts as nucleation sites for isothermal ω and/or α precipitation and leads to homogenously and finely distributed α particles throughout the β matrix during subsequent aging [29,30]. The strength of the alloy in the α/β solution condition is somewhat superior to that of in the β solution condition, a 180 MPa difference in yield strength being mentioned above. From the aspect of microstructure, the relative higher strength perhaps is supposed to originate from two reasons. The first is that defects retained during hot working are not completely dissolved during the α/β solution treatment. The second is because of the stubby shape α phase in the β matrix, which can lead to an increase in strength to some extent. One of the important characteristic of the α phase concerns its ability to limit recrystallization and grain growth. It can also be said that the globular shape α phase has very little influence upon the strength of beta Ti alloys but has a direct influence on the ductility [31]. However, the β solution treatment leads to a reduction in strength with an improved ductility. It can be attributed to the recrystallized strain-free β grains [32]. The tensile properties of the alloy in the β solution condition are superior to that of the common alloy BETA C (Ti–3Al–8V–6Cr– 4Mo–4Zr) alloy in condition of 900 1C/30 min/WQ, which shows 879 MPa of the yield strength with only 23% of the elongation [33].
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The alloy in both of the two solution treatments condition behaves an elastic modulus of 61–67 GPa, which is much lower than the other commercial beta titanium alloys. Comparing to the Ti–15Mo alloy, which has a similar [Mo]eq value with the programmed alloy in this research, has an elastic modulus of 78 GPa in the solution annealed condition according to ASTM F 2066-2001 [34]. The reason for such a low elastic modulus is probably that the synergistic effect of 9.2% Mo and 2% Fe is more effective to lower the elastic modulus than that of 15% Mo. In general, the tensile behavior of the solution treated beta alloy significantly depends on the deformation mode which is also strongly dependent on the stability of the alloy itself. The formation of SIM found in Fig. 6(b) leads to lower apparent yield stress with a better work-hardening response as seen in Fig. 5(b). The combination of the slip with the formation of SIM has also been reported that often leads to a good compromise in terms of enhanced ductility [35,36]. However, the absence of SIM or twinning in Fig. 6(a), corresponding to the alloy solution treated at 790 1C, is supposed that only the slip mode dominates the deformation mechanism. The reason is that the deformation mode changes from the martensite formation to the slip because the stability of the β matrix is increased due to the precipitation of primary α phase [37–39]. Therefore, very little work hardening and poor ductility is observed in Fig. 5(a) when the deformation becomes activated only by the slip mode. 4.2. α/β-Solution treatment plus aging In general, the aging response depends on whether the solution treatment temperature is above or below the β transus temperature, which allows the stability of the remained β phase or the driving force for α precipitation to be controllable. Undergoing the α/β solution treatment, the remained β phase has an improved stability due to the precipitation of primary α phase and it shows relatively low strength levels. The α/β solution treated alloy thus needs to be aged to obtain a higher strength level. Generally, the primary α phase precipitated during the α/β-ST results in a more stable remained β matrix. As a result, there is little possibility of the ω formation in the α/β-ST and water quenched beta alloys during isothermally aged at lower temperature. Thus, there are few researches reported that the isothermal ω phase is able to be formed in this condition. However, this kind of isothermal ω phase is found indeed in this research when the alloy solution treated in the α/β phase region and aged at 400–450 1C for 2 h. It is well known that the ω phase formation tends to accommodate a critical composition range or the β phase stability. Here we should argue that the remained β phase stability that experienced the α/β-ST is still in the range though it should be superior to that of the β-ST. In addition, it seems to that the ω precipitates have mostly dissolved at 5001C during 2 h. Therefore, it can be concluded that the ω dissolving temperature of this alloy is likely to be in the range of 450–500 1C. The fine α precipitates in Fig. 3(d) and (e) is possibly attributed to the effect of the α/β solution treatment, which is to enrich the β matrix by β stabilizers and somewhat deplete in α stabilizers. Therefore the α/β-ST would weaken its propensity to decompose and thus reduce the driving force for the α precipitates during subsequent aging. Furthermore, the width of the prior highly elongated β grains (∼10 μm) would also limit the maximum size of the secondary α phases. Due to the microstructural evolution during aging at different temperatures as mentioned above, it appears that increasing aging temperature results in a sharp decline of strength. The microstructure produced from small ω to coarse ω to nano sized α to small α to coarsen α when the alloy experienced aging temperature from 400 1C to 600 1C for 2 h. It should be pointed out that in aging at 400 1C, it is failed and ruptured during its elastic tensile
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stage. Perhaps it derives from the brittleness of ω phase described above. The presence of brittleness or super hardening is also involved in the α/β-ST alloy when aged at 450 1C for 2 h, which leads to quite a high strength above 1600 MPa with about 2.5% of the elongation. In case of aging at 500 1C, the alloy also obtains quite a high strength above 1600 MPa with about 4.5% of the elongation, which is more acceptable than that of the alloy produced the ω phase during aging. However, due to the α phase with 0.5–2 μm when the alloy aged at 550 1C and 600 1C, it shows attractive balance of strength and ductility (1200–1400 MPa of the ultimate strength with 7.5–12.5% of the elongation). The properties are similar to Timetal LCB [18] and Timetal 555 [40] alloy but quite superior to Beta C [9] and VT 22 alloy [8] in a similar heat treatment condition. 4.3. β-Solution treatment plus aging For beta titanium alloys, however, solution treated above the β transus temperature results in coarse β grains, and the least stable β matrix and the greatest driving force for its decomposition during subsequent aging [20]. Thus, the α phase nucleation kinetics are kinetically faster in this condition. It means that a shorter aging period may be required to reach a peak strength. As a result, the α sizes about 1 μm in Fig. 4(d) and about 3 μm in Fig. 4(e) appear to be coarser than that in Fig. 3(d) and (e) which corresponds to the α/β-STA. In addition, the ω phases in Fig. 4 (a) and (b) also shows a litter coarser than that in Fig. 3(a) and (b) for the same reason. Here should be pointed that, the tensile test of the sample aged at 400 1C also failed and ruptured within the elastic stage due to the brittle ω. When the alloy was aged at 450 1C, it results in a moderate strength about 1400 MPa but with a low ductility (1.5% of the elongation). Both of the strength and ductility is not so attractive as it occurred in the α/β-ST condition when aged at the same temperature. It is suggested that the reason for that is the finer ω phase obtained in the α/β-STA at 450 1C than that in the β-STA (30 nm in Fig. 3(b) and about 50 nm in Fig. 4(b)). However, the α phases with about 1–3 μm obtained at 550 1C and 600 1C result in a good balance of strength and ductility. However, the α phase about 500 nm obtained at 500 1C aging leads to an excessive strengthening with too low ductility. It means that too fine α phase smaller than 1 μm is responsible for low ductility and not beneficial to property combination. This problem can be solved by an extreme over aging time and thus coarsen the α phase to lower the strength with an improved ductility. Overall, however, the tensile properties of the alloy seem to be independent on its previous solution treatment (α/β-ST or β-ST) when sharing the same aging treatment. The reason is not clear yet. It is suggested that the variety in the stability of β matrix results in different aging kinetics and needs different aging time to reach a peak strength itself. 4.4. Comparisons to other beta alloys As mentioned above, the Ti–2Al–9.2Mo–2Fe alloy is a beta alloy with the molybdenum equivalent of 15.2 and the β transus temperature of 815 1C. The molybdenum equivalent and β transus temperature are similar to the Ti–15Mo alloy, Timetal LCB alloy and the Beta C alloy. However, this new alloy is melted with cheap Mo–Fe master alloy which is commonly found in steel industry. Therefore, this alloy is much cheaper than the other beta titanium alloys having high and expensive alloying elements, such as Mo and V. Additionally, the alloy in solution condition has lower modulus (61–67 GPa) and higher percent of elongation (∼40%), which is superior to the Ti–15Mo, Timetal LCB and Beta C alloy. Generally, these alloys have higher elastic modulus of about 78–
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90 GPa in solution condition and elongation of about 20%. In addition, this new alloy after aging can be strengthened to 1400– 1600 MPa with an accepted ductility (7.5–12.5% in the elongation), these properties are also better than the Beta C and Timetal LCB alloy. It means that this new alloy is quite suitable for a spring usage because of the requirement of high strength and low modulus for spring applications. 5. Conclusions The microstructural characteristic and tensile properties of the Ti–2Al–9.2Mo–2Fe alloy were investigated systematically. The alloy was solution treated below and above the β transus temperature, and aged at temperatures ranging from 400 1C to 600 1C afterwards. The main results are summarized as follows: (1) The Ti–2Al–9.2Mo–2Fe alloy, solution treated below and above the β transus temperature, has a low elastic modulus (61– 67 GPa) due to the synergistic effect of Mo and Fe. The low modulus should be attractive in some spring applications. (2) Although the isothermal ω phase is less common in beta alloys when treated in the α/β-ST and aged at lower temperatures, it is detected in the alloy with the α/β-ST and aged at low temperatures (400–450 1C). (3) The ω phase contributes to very high strength levels (1600 MPa of ultimate strength), with low ductility (2.5–4.5% of elongation) or even brittle fracture during the elastic stage of tensile test. It shows a poor balance of strength and ductility in the ω-aged alloy. Therefore, the improvement of the ductility needs more detailed investigations. (4) The α phase with 1–3 μm obtained at high aging temperatures leads to attractive combinations of strength and ductility (1200–1400 MPa of ultimate strength with 7.5–12.5% of elongation). It can be assumed that the alloy may be a usable structural material. (5) The tensile properties of the alloy seem to be independent of its previous solution treatment when sharing the same aging treatment. The reason is not clear yet and is needed to be understood further. Acknowledgments The research was sponsored by the Korea Institute of Materials Science (KIMS, Korea) and the General Institute Research for Nonferrous Metals (GRINM, China). The authors would like to acknowledge this financial support from the Ministry of Education, Science and Technology (MEST), the Ministry of Knowledge Economy (MKE) of Korea, and the Ministry of Science and Technology of China under Grant no. 2013DFG52920. References [1] S. Ankem, C.A. Greene, Mater. Sci. Eng. A 265 (1999) 127–131. [2] Y.T. Lee, Titanium, Korea Metal Journal, first ed., Korea Metal Journal News 2009. [3] K. Wang, Mater. Sci. Eng. A 213 (1996) 134–137.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Influence of heat treatment on microstructure and tensile property of a new high strength beta alloy Ti–2Al–9.2Mo–2Fe Cheng-Lin Li a,b, Xu-Jun Mi a, Wen-Jun Ye a, Song-Xiao Hui a, Dong-Geun Lee b,n, Yong-Tai Lee b a b
State Key Laboratory for Nonferrous Metals & Process, General Research Institute for Nonferrous Metals, Beijing 100088, PR China Titanium Department, Korea Institute of Materials Science, Changwon 642831, South Korea
art ic l e i nf o
a b s t r a c t
Article history: Received 23 December 2012 Received in revised form 27 April 2013 Accepted 29 April 2013 Available online 17 May 2013
An investigation is given on the influence of heat treatment by the microstructural characteristics and tensile properties of a new high strength alloy Ti–2Al–9.2Mo–2Fe. Both of the α/β and β solution treatment (α/β-ST and β-ST), then aged at temperatures ranging from 400 1C to 600 1C, were prepared. The primary α phases having 2–5 μm are formed during the α/β-ST, and restrain the size of β grains below 10 μm. As a result of the fine β grains, the α/β-ST contributes a higher strength than the β-ST. The coexistence of α″ and athermal ω phase is found in the β-ST and water quenched samples. However, these have little influence on the alloy hardening. After aging, the alloys in the α/β-ST and β-ST condition reveal the phase transformation of β to isothermal ω, and β to α depending on the aging temperature. Although the primary α phase formed during the α/β-ST increases the stability in the β matrix, and the isothermal ω phase also appears to occur during aging at 400 1C and 450 1C for 2 h. These phenomena are less common in beta titanium alloys, when treated in the α/β-ST and aged at lower temperatures. The isothermal ω phase formed in both conditions results in high strength levels (1600 MPa of ultimate strength) with much ductility loss (2.5–4.5% of elongation) as a result of the superior hardening effect and brittleness of ω phase. However, the secondary α phase with the size of 1–3 μm leads to attractive combinations of strength and ductility (1200–1400 MPa of ultimate strength with 7.5–12.5% of elongation). The reason for that is too fine α phase below 1 μm tends to result in ultra-high strength with much ductility loss. As a whole, the alloy can be heat treated to obtain excellent balances of strength and ductility, and provided abundant stress levels with optional ductility as a usable material. & 2013 Elsevier B.V. All rights reserved.
Keywords: Omega and alpha phase Age strengthening Tensile property Beta titanium alloy
1. Introduction Beta titanium alloys have the highest strength-to-weight ratio among all titanium alloys (alpha, alpha/beta and beta alloys). The beta alloys are attractive as a structural material for applications, such as automotive parts and surgical orthopedic implants besides the conventional use in the aircraft industry [1–3]. The beta titanium alloys are typically formed or fabricated in the solution treated (ST) condition because they are ductile with low yield strength ranging from 700 MPa to 1000 MPa. By the heat treatment of aging they will show high strength (1200–1700 MPa). In general, most of the increased strength due to aging results from the precipitation of fine α phase in the β matrix [4]. The fine α phases are hard and un-deformable particles to hinder the movement of dislocations and thus strengthen the alloys. This process is known as precipitation hardening in the STA condition [5,6]. The
n
Corresponding author. Tel.: +82 55 280 3367; fax: +82 55 280 3255. E-mail address: [email protected] (D.-G. Lee).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.04.118
size and distribution of the α phases affect the strength level and depend on the aging time at temperature. While the prior β grain size determines the ductility [7–9]. The optimum combination of mechanical properties of beta titanium alloys can be altered via heat treatment and controlling their microstructures for desired mechanical properties [10,11]. Although the beta titanium alloys have an advantage in a unique combination of high specific strength and excellent corrosion resistance, the application is significantly restricted by the high cost of the titanium itself during melting or processing [12,13]. For these reasons, recently, the development of the new Ti alloys for non-aerospace applications have led to design of lower cost alloys, such as Timetal LCB (Ti–6.8Mo–4.5Fe–1.5Al) alloy for automotive and other industrial applications [14]. The alloy has a ultimate strength level of 1000 MPa and 80–90 GPa of elastic modulus with 18% elongation in a typical solution treated condition [15,16]. While it can be heat-treated (or aged) to over 1400 MPa of ultimate strength, 110–117 GPa of elastic modulus with 13% of elongation [9,17]. It is the low cost Mo–Fe master alloy that highlights the LCB alloy. However, it readily produces so
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called “beta flecks” due to the high content of Fe addition. The beta flecks are solute rich regions that exhibit a different microstructure from the surrounding material in the final product. Therefore, they may result in inhomogeneous aging products and unstable mechanical properties [18–20]. Given all of that, a new beta titanium alloy with low cost, high strength and low elastic modulus should be designed and expected to overcome the Timetal LCB alloy and/or other newly developed LCB alloys. A Ti–2Al–9.2Mo–2Fe alloy is a new beta titanium alloy developed to be used as structural materials such as automotive spring application. It was designed on the basis of Bo–Md molecular orbital method (Bo, the bond order; Md, the metal d-orbital energy level) [21]. However, as a newly designed alloy, the microstructural features and mechanical properties of the Ti–2Al–9.2Mo–2Fe alloy are not well understood yet. Furthermore, it is of extreme importance to understand its microstructure– mechanical properties relationship by several heat treatment methods. Therefore, the present work is undertaken to understand the relationship of microstructural characteristics and tensile properties of the Ti–2Al–9.2Mo–2Fe alloy during the α/β or β solution treatment and subsequent aging, and concurrently, try to reveal how the microstructures of the new alloy change by heat treatments and how the changes affect the tensile properties.
2. Material and experimental methods The Ti–2Al–9.2Mo–2Fe alloy was melted by multiple vacuum arc re-melting (VAR) method. The ingot was forged at around 1100 1C to a square billet with 400 400 L mm in dimensions. The square billet was then hot rolled at 800 1C to round bars with 12 mm in diameter. The chemical composition of the alloy is listed in Table 1. The molybdenum equivalent ([Mo]eq) of the alloy is about 15.2 according to its chemical composition and an empirical equation proposed by Collings [22]. The β transus temperature of the alloy is measured to be approximately 815 1C with a metallographic method. The microstructure of the as rolled alloy is shown in Fig. 1. It mainly consists of small deformed β grains and fine nano-scaled acicular α precipitates due to the rolling temperature below the β transus temperature. Samples from the as-rolled alloy were solution treated in a preheated air furnace at 790 1C and 850 1C (below and above the β transus temperature) and then held for 1 h followed by water quench. All the samples were solution treated in an air box furnace. In view of the chemical activity of titanium alloys at high temperatures, all the samples are coated with Acheson's Deltaglaze 151 protective coating to minimize the formation of α case before heat treatment. Following the solution treatment, the samples were subjected to aging at temperatures ranging from 400 1C to 600 1C for 2 h in a pre-heated air furnace and followed by water quench. Specimens for microstructure observation were mechanically polished with the standard metallographic procedure and etched in the Kroll's reagent (5% HF+15% HNO3+80% H2O) to reveal the grain boundaries. Microstructure characterization was performed on a JEOL JSM-5800 scanning electron microscopy (SEM) and a JEOL JEM-2100F transmission electron microscopy (TEM) operated at 20 kV and 200 kV, respectively. For the TEM thin foil preparation it was firstly grinded on waterproof abrasive papers to 60 μm Table 1 Chemical composition of the Ti–2Al–9.2Mo–2Fe alloy (wt%). Al
Mo
Fe
Cu
Si
C
H
O
N
Ti
2.33
9.33
1.98
0.015
0.037
0.0091
0.0013
0.054
0.0087
Bal.
251
in thickness and then prepared with the MTP-1A twin jet electrochemical polishing method at −30 1C in an electrolyte consisting of perchloric acid, n-butyl alcohol and methanol (volume ratio 5%:35%:60%). Standard cylindrical tensile specimens with gauge length of 30 mm and gauge diameter of 6 mm were machined from both the solution treated and aged materials. The tensile tests were performed at room temperature with a nominal strain rate of 5.5 10–4 s−1, using a screw driven INSTRON machine of 50 kN capacity. At least three samples of each condition were tested for an average value of ultimate tensile strength (Rm), yield strength (Rp0.2) (that is, 0.2% proof stress), elongation (A) and elastic modulus (E).
3. Results 3.1. Microstructural characteristics The microstructural features of the Ti–2Al–9.2Mo–2Fe alloy, which was solution treated at below and above the β transus temperature for 1 h, are shown in Fig. 2. As seen in Fig. 2(a), there are stubby and globular primary α phases precipitated in the β matrix, as well as discontinuous primary α phases formed on and along the prior β grain boundaries. The stubby and globular primary α phases have 2–5 μm length or diameter. However, the remaining β grains exhibit a size within 10 μm. In case of the solution treatment at 850 1C condition, it results in equiaxed and coarse β grains with a mean size of about 50 μm due to recrystallization, as seen in Fig. 2(b). A feature in Fig. 2(b) shows that plate-like precipitations emerge in some β grains, which refer to martensites formed during water quench. Consequently, the martenistes are difficult to be seen with the TEM. However, a selected area diffraction (SAD) pattern from the β matrix of Fig. 2 (d) is shown in Fig. 2(c), which is consistently indexed as [113] zone axis of β phase. Additional weaker spots are visible at the 1/3 (121) and 2/3(121) positions of the β reflection where they have straight or curved lines of diffuse intensity (as known as diffuse streaking) instead of sharp spots. That refers to as the “diffuse ω phase” relating the early stages of ω phase formation [22,23]. Microstructures of the Ti–2Al–9.2Mo–2Fe alloy after the α/β-ST and aged at different temperatures are shown in Fig. 3. The first heat treatment carried out on the as-quenched sample is aged at 400 1C for 2 h. The ω precipitates are imaged from the 1/3 (or 2/3) (112) β reflections in the [113] β zone axis. As shown in Fig. 3(a), the size scale of the ω precipitates is in the range of 10–20 nm. Further aging treatments are conducted at 450 1C for 2 h. The ω precipitates clearly exhibit an ellipsoidal shape with a size scale of about 30 nm seen in Fig. 3(b). It is observed that the ω precipitates have coarsened during aging at 450 1C. However, by aging this alloy at a higher temperature like 500 1C, it is observed in Fig. 3 (c) that the fine scale α precipitates with a size of 100–200 nm shapes a more lenticular morphology. There is no evidence of ω precipitate presented from the dark field image. The microstructures of the α/β solution treated alloy aged at 550 1C and 600 1C for 2 h are shown in Fig. 3(d) and (e), respectively. The α platelet clearly appears to be present in the β matrix, and exhibits a size of about 0.5 μm for aging at 550 1C and about 1–2 μm for 600 1C. The microstructures of the Ti–2Al–9.2Mo–2Fe alloy after the β-ST plus aged at different temperatures for 2 h are shown in Fig. 4. From these figures it can be seen that the ω phases with an ellipsoidal shape form in the β matrix when the alloy was aged at low temperatures of 400 1C and 450 1C. Moreover, the ω phases obtained at 450 1C exhibits a coarser size (about 50 nm) than that at 400 1C (about 25 nm) during the same aging time. No α phase is
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Fig. 1. Microstructures of the as rolled Ti–2Al–9.2Mo–2Fe alloy: (a) OM, and (b) SEM.
Fig. 2. Microstructures of the alloy after solution treated at: (a) SEM 790 1C, (b) OM 850 1C, (c) 850 1C, an [113]β selected area diffraction pattern having extra reflections at 1/3 and 2/3 (112)β positions corresponding to athermal ω phases, (d) 850 1C dark-field image from the extra reflections to show nano-scaled ω phases.
Fig. 3. Microstructures of the alloy solution treated at 790 1C and aged for 2 h at: (a) 400 1C, (b) 450 1C, (c) 500 1C, (d) 550 1C and (e) 600 1C.
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253
Fig. 4. Microstructures of the alloy solution treated at 850 1C and aged for 2 h at: (a) 400 1C, (b) 450 1C, (c) 500 1C (d) 550 1C, and (e) 600 1C.
Table 2 Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy in solution treated condition. ST
Rm (MPa)
Rp0.2 (MPa)
A (%)
E (GPa)
790 1C/1 h/WQ 850 1C/1 h/WQ
957 833
853 673
20 42
61 67
detected by TEM micrographs when the alloy is aged at the two temperatures. However, only the α phase can be found in Fig. 4 (c) which corresponds to the alloy aged at 500 1C for 2 h. The lamellar α phase has the size of about 500 nm. As seen in Fig. 4 (d) and (e), very fine α platelet distributes in the β matrix, the size itself exhibits 1 μm for 550 1C in Fig. 4(d) and 3 μm for 600 1C in Fig. 4(e). In other words, the α phase has a slight increase in size as the aging temperature increased from 500 1C to 600 1C within aging for 2 h.
aged at 450 1C and 500 1C for 2 h shows an ultimate strength as high as 1600 MPa with poor ductility of below 5% of the elongation, but an increase in the aging temperature leads to a rapid decrease in strength with an increase in ductility. Aging at 600 1C tends to produce an interesting property balance (above 1200 MPa in ultimate strength with 12.5% of elongation). Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy after the β-ST plus aging at different temperatures are shown in Fig. 8. When the alloy was aged at 450 1C, it does result in an ultimate strength about 1400 MPa with a low ductility (1.5% of the elongation). Putting aside of the ω-aged alloy in terms of 450 1C, the strength basically decreases sharply with increasing the aging temperature. The highest strength level is obtained at 500 1C aging with a remarkable ductility loss. On the other hand, a good balance of strength and ductility is obtained at 550 1C and 600 1C.
4. Discussion 3.2. Tensile properties Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy, which was solution treated at 790 1C and 850 1C for 1 h, are listed in Table 2. The relative stress–strain curves of the Ti–2Al–9.2Mo–2Fe alloy in the α/β and β solution condition are shown in Fig. 5. Both of the two solution treatments specimens behave an elastic modulus of 61–67 GPa. The alloy solution treated at 790 1C for 1 h shows high yield strength of 853 MPa, and a relative small elongation of 20%. In addition, it shows a weaker work hardening phenomenon as seen in Fig. 5(a). On the contrary, the alloy solution treated at 850 1C for 1 h shows a low yield strength level of 673 MPa with a larger elongation of 42%. But in this condition, it shows a significant work hardening as seen in Fig. 5(b). The microstructures from the region near the fracture of the tensile specimens, which were pre-solution treated at 790 1C and 850 1C for 1 h, are shown in Fig. 6. The micrograph in Fig. 6(b), which refers to the alloy solution treated at 850 1C, clearly shows the presence of α″ and some other plate-like features. The former refers to the stress or strain induced martensite (SIM) and the latter corresponds to slip bands. However, the absence of SIM or twinning in Fig. 6(a), corresponding to the alloy solution treated at 790 1C, is supposed that only the slip mode dominates the deformation mechanism. Tensile properties of the Ti–2Al–9.2Mo–2Fe alloy after the α/βST plus aging are illustrated in Fig. 7. It can be seen that the alloy
The microstructural characteristics of the alloy were studied mainly by a SEM, in some cases the SEM observations were augmented by TEM studies. The various microstructures showed pronounced differences as well as tensile behavior and were thus difficult to discuss concurrently. Therefore, this section will be divided into three parts: (1) solution treatment (ST); (2) α/β solution treatment plus aging (α/β-STA); and (3) β solution treatment plus aging (β-STA). 4.1. Solution treatment The Ti–2Al–9.2Mo–2Fe alloy is a metastable beta titanium alloy with the β transus temperature of 815 1C. In case of the beta titanium alloys, solution treatment above the β transus temperature results in coarse β grains. While, solution treatment slightly below the β transus temperature leads to precipitations of primary α phase [24]. The primary α phases obtained in the condition α/β-ST (Fig. 2 (a)) exhibit stubby and globular shapes. It is supposed that the primary α phases tend to form on the basis of the growth or globalization of the nano-scaled α phases in the as-rolled microstructure (Fig. 1), accompanying with some of them dissolved. It is well known that the important characteristic of the primary α phase concerns its ability to pin the prior β grain boundaries and reduce their mobility, which therefore limits the recrystallization
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Fig. 5. Nominal stress–strain curves of the alloys after solution treatment: (a) 790 1C and (b) 850 1C.
Fig. 6. Microstructures from the region near the fracture of the tensile deformed specimens post-solution treated at: (a) 790 1C, and (b) 850 1C. The tensile axis is perpendicular to the paper.
Fig. 7. Tensile properties of the alloy solution treated at 790 1C and aged at different temperatures for 2 h.
Fig. 8. Tensile properties of the alloy solution treated at 850 1C and aged at different temperatures for 2 h.
and grain growth of the β grain [25]. In case of β-ST, besides the equiaxed and coarse β grains obtained due to recrystallization, discontinuous distributed martensites formed in some β grains during water quench. The reason for the discontinuous distribution would be that local chemical variations or composition segregation. The segregation allows some grains to have an uneven of beta stabilizers, which leads to the difference of stability. It has been proven by an EPMA method [26–28]. It should be noted that the small volume fraction of ω phase found in the β matrix has little direct influence on the strength of the alloy, but indeed, which acts as nucleation sites for isothermal ω and/or α precipitation and leads to homogenously and finely distributed α particles throughout the β matrix during subsequent aging [29,30]. The strength of the alloy in the α/β solution condition is somewhat superior to that of in the β solution condition, a 180 MPa difference in yield strength being mentioned above. From the aspect of microstructure, the relative higher strength perhaps is supposed to originate from two reasons. The first is that defects retained during hot working are not completely dissolved during the α/β solution treatment. The second is because of the stubby shape α phase in the β matrix, which can lead to an increase in strength to some extent. One of the important characteristic of the α phase concerns its ability to limit recrystallization and grain growth. It can also be said that the globular shape α phase has very little influence upon the strength of beta Ti alloys but has a direct influence on the ductility [31]. However, the β solution treatment leads to a reduction in strength with an improved ductility. It can be attributed to the recrystallized strain-free β grains [32]. The tensile properties of the alloy in the β solution condition are superior to that of the common alloy BETA C (Ti–3Al–8V–6Cr– 4Mo–4Zr) alloy in condition of 900 1C/30 min/WQ, which shows 879 MPa of the yield strength with only 23% of the elongation [33].
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The alloy in both of the two solution treatments condition behaves an elastic modulus of 61–67 GPa, which is much lower than the other commercial beta titanium alloys. Comparing to the Ti–15Mo alloy, which has a similar [Mo]eq value with the programmed alloy in this research, has an elastic modulus of 78 GPa in the solution annealed condition according to ASTM F 2066-2001 [34]. The reason for such a low elastic modulus is probably that the synergistic effect of 9.2% Mo and 2% Fe is more effective to lower the elastic modulus than that of 15% Mo. In general, the tensile behavior of the solution treated beta alloy significantly depends on the deformation mode which is also strongly dependent on the stability of the alloy itself. The formation of SIM found in Fig. 6(b) leads to lower apparent yield stress with a better work-hardening response as seen in Fig. 5(b). The combination of the slip with the formation of SIM has also been reported that often leads to a good compromise in terms of enhanced ductility [35,36]. However, the absence of SIM or twinning in Fig. 6(a), corresponding to the alloy solution treated at 790 1C, is supposed that only the slip mode dominates the deformation mechanism. The reason is that the deformation mode changes from the martensite formation to the slip because the stability of the β matrix is increased due to the precipitation of primary α phase [37–39]. Therefore, very little work hardening and poor ductility is observed in Fig. 5(a) when the deformation becomes activated only by the slip mode. 4.2. α/β-Solution treatment plus aging In general, the aging response depends on whether the solution treatment temperature is above or below the β transus temperature, which allows the stability of the remained β phase or the driving force for α precipitation to be controllable. Undergoing the α/β solution treatment, the remained β phase has an improved stability due to the precipitation of primary α phase and it shows relatively low strength levels. The α/β solution treated alloy thus needs to be aged to obtain a higher strength level. Generally, the primary α phase precipitated during the α/β-ST results in a more stable remained β matrix. As a result, there is little possibility of the ω formation in the α/β-ST and water quenched beta alloys during isothermally aged at lower temperature. Thus, there are few researches reported that the isothermal ω phase is able to be formed in this condition. However, this kind of isothermal ω phase is found indeed in this research when the alloy solution treated in the α/β phase region and aged at 400–450 1C for 2 h. It is well known that the ω phase formation tends to accommodate a critical composition range or the β phase stability. Here we should argue that the remained β phase stability that experienced the α/β-ST is still in the range though it should be superior to that of the β-ST. In addition, it seems to that the ω precipitates have mostly dissolved at 5001C during 2 h. Therefore, it can be concluded that the ω dissolving temperature of this alloy is likely to be in the range of 450–500 1C. The fine α precipitates in Fig. 3(d) and (e) is possibly attributed to the effect of the α/β solution treatment, which is to enrich the β matrix by β stabilizers and somewhat deplete in α stabilizers. Therefore the α/β-ST would weaken its propensity to decompose and thus reduce the driving force for the α precipitates during subsequent aging. Furthermore, the width of the prior highly elongated β grains (∼10 μm) would also limit the maximum size of the secondary α phases. Due to the microstructural evolution during aging at different temperatures as mentioned above, it appears that increasing aging temperature results in a sharp decline of strength. The microstructure produced from small ω to coarse ω to nano sized α to small α to coarsen α when the alloy experienced aging temperature from 400 1C to 600 1C for 2 h. It should be pointed out that in aging at 400 1C, it is failed and ruptured during its elastic tensile
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stage. Perhaps it derives from the brittleness of ω phase described above. The presence of brittleness or super hardening is also involved in the α/β-ST alloy when aged at 450 1C for 2 h, which leads to quite a high strength above 1600 MPa with about 2.5% of the elongation. In case of aging at 500 1C, the alloy also obtains quite a high strength above 1600 MPa with about 4.5% of the elongation, which is more acceptable than that of the alloy produced the ω phase during aging. However, due to the α phase with 0.5–2 μm when the alloy aged at 550 1C and 600 1C, it shows attractive balance of strength and ductility (1200–1400 MPa of the ultimate strength with 7.5–12.5% of the elongation). The properties are similar to Timetal LCB [18] and Timetal 555 [40] alloy but quite superior to Beta C [9] and VT 22 alloy [8] in a similar heat treatment condition. 4.3. β-Solution treatment plus aging For beta titanium alloys, however, solution treated above the β transus temperature results in coarse β grains, and the least stable β matrix and the greatest driving force for its decomposition during subsequent aging [20]. Thus, the α phase nucleation kinetics are kinetically faster in this condition. It means that a shorter aging period may be required to reach a peak strength. As a result, the α sizes about 1 μm in Fig. 4(d) and about 3 μm in Fig. 4(e) appear to be coarser than that in Fig. 3(d) and (e) which corresponds to the α/β-STA. In addition, the ω phases in Fig. 4 (a) and (b) also shows a litter coarser than that in Fig. 3(a) and (b) for the same reason. Here should be pointed that, the tensile test of the sample aged at 400 1C also failed and ruptured within the elastic stage due to the brittle ω. When the alloy was aged at 450 1C, it results in a moderate strength about 1400 MPa but with a low ductility (1.5% of the elongation). Both of the strength and ductility is not so attractive as it occurred in the α/β-ST condition when aged at the same temperature. It is suggested that the reason for that is the finer ω phase obtained in the α/β-STA at 450 1C than that in the β-STA (30 nm in Fig. 3(b) and about 50 nm in Fig. 4(b)). However, the α phases with about 1–3 μm obtained at 550 1C and 600 1C result in a good balance of strength and ductility. However, the α phase about 500 nm obtained at 500 1C aging leads to an excessive strengthening with too low ductility. It means that too fine α phase smaller than 1 μm is responsible for low ductility and not beneficial to property combination. This problem can be solved by an extreme over aging time and thus coarsen the α phase to lower the strength with an improved ductility. Overall, however, the tensile properties of the alloy seem to be independent on its previous solution treatment (α/β-ST or β-ST) when sharing the same aging treatment. The reason is not clear yet. It is suggested that the variety in the stability of β matrix results in different aging kinetics and needs different aging time to reach a peak strength itself. 4.4. Comparisons to other beta alloys As mentioned above, the Ti–2Al–9.2Mo–2Fe alloy is a beta alloy with the molybdenum equivalent of 15.2 and the β transus temperature of 815 1C. The molybdenum equivalent and β transus temperature are similar to the Ti–15Mo alloy, Timetal LCB alloy and the Beta C alloy. However, this new alloy is melted with cheap Mo–Fe master alloy which is commonly found in steel industry. Therefore, this alloy is much cheaper than the other beta titanium alloys having high and expensive alloying elements, such as Mo and V. Additionally, the alloy in solution condition has lower modulus (61–67 GPa) and higher percent of elongation (∼40%), which is superior to the Ti–15Mo, Timetal LCB and Beta C alloy. Generally, these alloys have higher elastic modulus of about 78–
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90 GPa in solution condition and elongation of about 20%. In addition, this new alloy after aging can be strengthened to 1400– 1600 MPa with an accepted ductility (7.5–12.5% in the elongation), these properties are also better than the Beta C and Timetal LCB alloy. It means that this new alloy is quite suitable for a spring usage because of the requirement of high strength and low modulus for spring applications. 5. Conclusions The microstructural characteristic and tensile properties of the Ti–2Al–9.2Mo–2Fe alloy were investigated systematically. The alloy was solution treated below and above the β transus temperature, and aged at temperatures ranging from 400 1C to 600 1C afterwards. The main results are summarized as follows: (1) The Ti–2Al–9.2Mo–2Fe alloy, solution treated below and above the β transus temperature, has a low elastic modulus (61– 67 GPa) due to the synergistic effect of Mo and Fe. The low modulus should be attractive in some spring applications. (2) Although the isothermal ω phase is less common in beta alloys when treated in the α/β-ST and aged at lower temperatures, it is detected in the alloy with the α/β-ST and aged at low temperatures (400–450 1C). (3) The ω phase contributes to very high strength levels (1600 MPa of ultimate strength), with low ductility (2.5–4.5% of elongation) or even brittle fracture during the elastic stage of tensile test. It shows a poor balance of strength and ductility in the ω-aged alloy. Therefore, the improvement of the ductility needs more detailed investigations. (4) The α phase with 1–3 μm obtained at high aging temperatures leads to attractive combinations of strength and ductility (1200–1400 MPa of ultimate strength with 7.5–12.5% of elongation). It can be assumed that the alloy may be a usable structural material. (5) The tensile properties of the alloy seem to be independent of its previous solution treatment when sharing the same aging treatment. The reason is not clear yet and is needed to be understood further. Acknowledgments The research was sponsored by the Korea Institute of Materials Science (KIMS, Korea) and the General Institute Research for Nonferrous Metals (GRINM, China). The authors would like to acknowledge this financial support from the Ministry of Education, Science and Technology (MEST), the Ministry of Knowledge Economy (MKE) of Korea, and the Ministry of Science and Technology of China under Grant no. 2013DFG52920. References [1] S. Ankem, C.A. Greene, Mater. Sci. Eng. A 265 (1999) 127–131. [2] Y.T. Lee, Titanium, Korea Metal Journal, first ed., Korea Metal Journal News 2009. [3] K. Wang, Mater. Sci. Eng. A 213 (1996) 134–137.
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