Phase precipitation behavior of a quenched β-solidifying TiAl alloy with a fully-B2 microstructure during annealing at 800°C

Phase precipitation behavior of a quenched β-solidifying TiAl alloy with a fully-B2 microstructure during annealing at 800°C

Journal of Alloys and Compounds 812 (2020) 152118 Contents lists available at ScienceDirect Journal of Alloys and Compounds journal homepage: http:/...

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Journal of Alloys and Compounds 812 (2020) 152118

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Phase precipitation behavior of a quenched b-solidifying TiAl alloy with a fully-B2 microstructure during annealing at 800 C Guang Yang a, *, Xiaoxiao Yang a, Yifei Wang a, Liang Cheng b, Hongchao Kou c, Yanhui Liu a, Yuzhi Li a, Pengyi Wang a, Wei Ren a a b c

College of Mechanical and Electrical Engineering, Shaanxi University of Science and Technology, Xi'an, Shaanxi, 710021, PR China School of Materials Engineering, Jiangsu University of Technology, Changzhou, Jiangsu, 213001, PR China State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, Shaanxi, 710072, PR China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 28 June 2019 Received in revised form 26 August 2019 Accepted 31 August 2019 Available online 3 September 2019

In this work, a fully-B2 microstructure was obtained in a Tie40Ale8Nb (at. %) alloy by water-quenching from the single b phase region, and the decomposition behavior of the fully-B2 microstructure during annealing at 800  C was investigated. Results showed that the fully-B2 microstructure was thermodynamically unstable and numerous lath-liked a2 phases precipitated from the B2 matrix following the Burgers orientation relationships after annealing for 3 min. With the increase of isothermal holding time, these lath-liked a2 phases transformed into ultra-thin (a2þg) lamellar structures with a lamellar spacing of 4 nm. Due to the perpendicular decomposition of the a2 lamellar and coarsening of the g lamellar, u0 and g phases formed in the lamellar structures. Besides, the decomposition of the B2 phase directly into fine g and u0 particles was also observed and the orientation relationships among these three phases were (1e10)B2//(1-1-1)g//(11e20)u0 and [111]B2//[01-1]g//[0001]u0. After tempering for 3000 min, the growth of the g and u0 phases occurred, leading to a u0þgþB2 microstructure with phase size of 300 e500 nm. © 2019 Elsevier B.V. All rights reserved.

Keywords: Intermetallics Phase transitions Microstructure transmission electron microscopy TEM

1. Introduction

b-solidifying TiAl alloys have been considered as a promising candidate to replace Ni-based superalloys in gas turbine engines because of their low density and excellent performance at elevated temperatures [1e3]. Compared with the traditional peritectically solidifying TiAl alloys, there are much more b-stabilizing elements, such as Nb, Mo, W and Cr, in the b-solidifying TiAl alloys [4,5]. As a result, there is a single b phase region at high temperature in these alloys [6e8]. The disordered b phase with a body-centred cubic (BCC) structure has a large number of independent slip systems; therefore, the hot-workability of the b-solidifying TiAl alloys is significantly improved [5,9]. However, the b phase is soft due to its open BCC structure, which results in a reduced creep resistance [10,11]. During the cooling process, martensite, massive and Wid€tten a phases can precipitate from the b matrix following mannsta

* Corresponding author. E-mail address: [email protected] (G. Yang). https://doi.org/10.1016/j.jallcom.2019.152118 0925-8388/© 2019 Elsevier B.V. All rights reserved.

the Burgers orientation relationships, i.e., (0001)ak{110}b and <11e20>ak<1e11>b, and these a phases are usually used to refine the microstructure of the b-solidifying TiAl alloys by subsequent heat treatments [12,13]. Besides the transformed a phases, about 10% of the b phase can transform into ordered B2 phase. The B2 phase is intrinsically brittle; therefore, the room-temperature ductility and workability of these alloys are deteriorated [14,15]. In addition, the B2 phase is much harder than a2 and g phases, which results in crack generation in the b-solidifying TiAl alloys [6,16]. Previous studies confirmed that the B2 phase was thermodynamically metastable and prone to decompose into u-related phases, such as B82-u (u0), D88-u and trigonal-u՛՛. Song et al. found that the volume fraction and size of u0 phase were strongly affected by the alloying elements and heat treatment process, which was closely related to the diffusion of the b-stabilizing elements [17,18]. Huang et al. revealed that the u0 phase could transform into D88-u phase during long time annealing [19]. Wang et al. reported that the formation of the u՛՛ phase was inevitable even by water quenching [20,21]. Moreover, the B2/g transition is also observed, including the discontinuous coarsening of the adjacent g lamellar and the direct

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nucleation of g phase in B2 matrix [22e24]. Therefore, a microstructure composed of u-related and g phases is usually obtained in the B2 region. Although the decomposition behavior of the B2 phase and related mechanisms were extensively studied, the existing researches mainly focus on nearly fully lamellar microstructures with a small amount of B2 phase (approximately 10% in phase fraction). In this work, a fully-B2 microstructure was obtained in Tie40Ale8Nb alloys (all compositions are given in at. %) by water quenching from 1400  C, and the decomposition process of the fully-B2 microstructure during annealing at 800  C was investigated in detail by using transmission electron microscopy analysis.

2. Experimental details The starting material with a nominal composition of Tie40Ale8Nb (at. %) was fabricated by an induction skull melting device with a water-cooled crucible. In order to remove the casting porosity and micro-segregation, a hot isostatic pressing (HIP) at 1180  C for 4 h was performed under an argon pressure of 150 MPa. The actual chemical composition of the present alloy is Ti-39.51Al7.89Nb containing 460 ppm of oxygen and 80 ppm of nitrogen. According to the references [25e27], the b transus temperature of this alloy is about 1250  C. Samples with size of 10  10  10 mm3 were cut from the HIPed ingot. These cube samples were sealed in vacuum quartz tubes back-filled with argon and then solution-treated at 1400  C. After the isothermal holding for 5 min, the samples were quenched in a boiling water. The selection of the quenching medium, i.e., boiling water, can prevent the formation of severe cracks in these specimens. Hereafter, we refer to the sample quenched from 1400  C as the “b-quenched sample”. After that, the b-quenched samples were sealed again and tempered at 800  C for 3, 10, 30, 60, 300 and 3000 min, respectively, and then water-quenched to room temperature. It should be noted here that the heating treatment furnace had been heated to 800  C and held for 15 min before the annealing treatment. The symbol “S-x” was used to denote the

samples heat-treated with different annealing time, and “x” means the isothermal holding time. For microstructural characterization, all of the samples were wire-cut from the center. The cutting surfaces were mechanically grounded and then electrolytically polished with a solution composed of 240 ml methanol, 140 ml butanol and 20 ml perchloric acid at 20 V and 30  C. Phase constituents were analyzed by X-Ray Diffraction (XRD, DX-2700) with Cu Ka radiation at 40 kV and 30 mA. The microstructural features were examined by using a P scanning electron microscopy (SEM, Zeiss- SIGMA500) with back scattered electron (SEM-BSE) mode and electron backscatter diffraction (EBSD) with a step size of 50 nm. An energy dispersive spectrometer (EDS) was used to detect the chemical compositions of the various constituent phases. To gain detailed microstructural information, transmission electron microscopy (TEM) observations were conducted on the Tecnai G2 F20 S-TWIN (FEI company) fieldmission TEM with an operating voltage of 300 kV. TEM slices were cut from the samples and then mechanically polished to 60 mm. The TEM foils were prepared by double-jet polishing with the same electrolyte as that used in the electrolytic polishing.

3. Results and discussion 3.1. b-quenched sample Fig. 1 shows the microstructure of the b-quenched sample. It is evident from Fig. 1(a) that the b-quenched sample is composed of the single B2 phase with an equiaxed morphology, and the average size of the B2 grains is approximately 4 mm. B2 phase is the ordered structure of the b phase [26e28]; therefore, the excessive coarsening of the B2 grains is attributed to the high solutiontemperature, i.e., 1400  C, which is about 150  C higher than the b transus of this alloy. The XRD pattern shown in Fig. 1(b) further confirms the phase constitutes of this sample, i.e., single B2 phase. Fig. 1(c) shows the TEM bright field image and it can be seen that no other phases can be found, remaining only B2 phase. However, from the selected area diffraction pattern (SADP) taken from the

Fig. 1. Microstructure characteristics of the b-quenched sample: (a) IPF map; (b) XRD pattern; (c) TEM bright field image; (d) the corresponding SADP of (c).

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Fig. 2. Microstructure characteristics of the S-3: (a) SEM-BSE image; (b) XRD pattern; (c) TEM bright field image; (d) the SADP of the region marked by the red circle in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

[111]B2 zone axis as shown in Fig. 1(d), u-related weak diffraction spots (labeled by red circles) are obviously observed indicating that some u-related structures exist in the B2 matrix. 3.2. Decomposition of the B2 phase at initial stage Fig. 2 shows the microstructure of the sample after annealing at 800  C for 3 min. Combining the SEM-BSE image (Fig. 2(a)) and XRD pattern (Fig. 2(b)), one can see that numerous lath-liked a2 phases precipitate from the B2 matrix, and these a2 phases show a characteristic of triangular clusters (indicated by red rectangles). The chemical compositions of the a2 phase and B2 matrix are shown in Table 1. It is obvious that the B2 matrix has a lower Al concentration than that of a2 phase, 38.29 at. % vs. 39.62 at. %, but enriched in Nb, 8.21 at. % vs. 7.54 at. %; therefore, the B2 matrix shows a relatively bright contrast. Fig. 2(c) shows the TEM bright field image of S-3 and the SADP of the circled region is given in Fig. 2(d). Three a2 variants (V1, V2, V3) can be observed in Fig. 2(c), and these a2 phases constitute a triangular configuration, which is consistent with the SEM-BSE image shown in Fig. 2(a). It is obvious from Fig. 2(d) that these a2 phases respect the Burgers orientation relationships with the B2 matrix, i.e., (0001)a2//{110}B2 and <11e20>a2//<1e11>B2, and these a2 variants are interrelated by 60 rotation around a common <11-20> axis. Previous researches confirmed that the formation of triangular clusters could minimize the elastic energy concentration in local regions; therefore, such a microstructure characteristic was usually found in pure Ti, Ti alloys

Table 1 Chemical compositions of the B2 matrix and a2 phase in S-3.

B2 matrix a2 phase

Al (at. %)

Nb (at. %)

Ti (at. %)

38.29 39.62

8.21 7.54

53.50 52.84

and TiAl alloys after the b/a transformation [13,29e31]. Fig. 3 shows the TEM results of the remaining B2 regions in S-3. The microstructure characteristics of these regions can be classified into two types. Type (I): the decomposition of the B2 phase does not occur (Fig. 3(a) and Fig. 3(b)), showing the similar microstructure to that of the b-quenched sample. Type (II): the B2 matrix transforms into a2, g and u0 phases with size of only a few nanometers; therefore, it is difficult to distinguish the morphologies of these phase constitutes in Fig. 3(c). However, the SADP shown in Fig. 3(d) indicates the precipitation of a2, g and u0 phases in the B2 regions, and the orientation relationships among these phases are (1e10)B2//(1-1-1)g//(11e20)u0//(0002)a2 and [111]B2//[01-1]g// [0001]u0//[11e20] a2. In order to clarify the decomposition behavior of the fully-B2 microstructure, the phase fraction diagram with temperature of the Tie40Ale8Nb alloy was calculated using the Calphad Software with the database from Ref. [25], and the result is shown in Fig. 4. It can be seen that equilibrium phase constitutes at 800  C of the present alloy are a2, g and u0 phases, indicating that the B2 phase at 800  C is thermodynamically unstable and readily decomposes into more stable phases, such as a2, g and u0 phases. The direct nucleation of u0 and g phases in the B2 matrix as shown in Fig. 3(c) and Fig. 3(d) confirms the B2/gþu0 transformation, which is frequently reported in previous studies [17,18,22,24,32e34]. However, an interesting phenomenon was found that the B2/a2 phase transformation occurs during the annealing process as shown in Fig. 2, which is observed for the first time in TiAl alloys. Such a transition in the present alloy is analogous to the precipitation of a phase during annealing in near b Ti alloys [35,36]. Moreover, numerous a2 phases with the size of a few microns precipitate within a very short annealing time, i.e., 3 min, demonstrating that explosive nucleation of the a2 phases happened in S-3. According to the calculated phase fraction diagram (Fig. 4), the lowest temperature at which a stable fully-B2 microstructure exists is 1244  C;

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Fig. 3. TEM results of the remaining B2 regions in S-3: (a, b) TEM bright image of the type (I) B2 region and the corresponding SADP of the circled area in (a); (c, d) TEM bright image of the type (II) B2 region and the corresponding SADP of the circled region in (c).

Fig. 4. Phase fraction diagram with temperature of the present alloy.

therefore, the B2 phase at 800  C is highly undercooled (DT ¼ 444  C), providing a huge driving force for the B2/a2 transformation. Fig. 5 shows the microstructure of S-10. It is evident from Fig. 5(a) that g lamellar precipitates from the a2 lath, resulting in ultra-thin (a2þg) lamellar structures. A high-resolution transmission electron microscopy (HRTEM) observation was performed so as to examine the detailed information. As shown in Fig. 5(b), ultra-thin (a2þg) lamellar structures form in S-10 with the average

lamellar spacing of 4 nm. On the basis of the fast Fourier transformation (FFT) graph given in Fig. 5(c), the orientation relationships between g and a2 phases are (11-1)g//(0002)a2 and [1e10]g// <11e20>a2. The diffraction spots of a2 phase can be clearly seen in Fig. 5(c), while the difficulty is to distinguish the diffraction spots belonging to g lamellar because that the corresponding diffraction spots are diffusive as a line. The local HRTEM pictures of the squared and circled regions in Fig. 5(b) are shown in Fig. 5(d) and Fig. 5(e). It can be perceived from Fig. 5(d) that an isolated g lamellar precipitated in a2 matrix, and the atomic arrangements further confirm the Blackburn orientation relationships between the a2 and g phases. Besides, two g variants with twin-relation, g and gT, are observed in Fig. 5(e), indicating the interfacial nucleation of twin-related g lamellar during annealing process [37]. Fig. 6 shows the TEM results of a region adjacent to lamellar structures in S-10. There are two g variants with twin-relation and u0 phase in the B2 region. The orientation relationships among these three phases are (1e10)B2//(11e20)u0//(1-1-1)g and [111]B2// [0001]u0//[01-1]g, which is the same as that depicted in Fig. 3(d). Compared with the microstructure shown in Fig. 3(c), the size of the g and u0 phases increases, which is resulted from the growth of the precipitated phases with the extension of annealing time. 3.3. Microstructure evolution during the subsequent annealing Fig. 7 shows the XRD patterns of the samples with different tempering time. It can be seen that the S-30, S-60, S-300 samples have the same phase constitutes, i.e., a2, u0 and g phases. The primary diffraction peak of B2 phase appears at 39.9499 , and angles of characteristic peaks belonging to the u0 phase are 39.3237 and 39.7585 ; therefore, it is hard to distinguish these two phases

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Fig. 5. Microstrcuture of S-10: (a, b) TEM bright and HRTEM images of the lamellar structures; (c) FFT graph of (b); (d, e) local HRTEM pictures of the square and circular regions in (b).

Fig. 6. (a) TEM bright image of a region adjacent to lamellar structures in S-10; (b) the SADP of the circled region in (a).

preciously. In the light of the XRD patterns shown in the red squares, two diffraction peaks were detected at the angle range of 39e40 . Thus, these two peaks were labeled as u0 phase in XRD analysis. However, the B2 phase was also observed in TEM observation (Fig. 8(b)), indicating that the actual phase constitutes of these samples are a2, u0, g and B2 phases. Besides, the XRD results show that the fraction of a2 phase decreases and the contents of u0 and g phases increase with the extension of isothermal holding time, leading to a (u0þgþB2) microstructure in S-3000 as shown in Fig. 7(d). The precipitation process of the u0 phase can be classified into two types. Type (I): u0 phase precipitates from the B2 phase following the orientation relationships of (1e10)B2//(11e20)u0 and [111]B2//[0001]u0, as shown in Fig. 8(a, b, c). Type (II): u0 phase forms in the ultra-thin (a2þg) lamellar structures respecting the orientation relationships of (1-1-1)g//(11e20)u0//(0002)a2 and [01-

1]g//[0001]u0//[11e20]a2, as shown in Fig. 8(d, e, f). The direct nucleation of u0 phase from the B2 matrix (type I) is caused by the thermodynamic instability of the B2 phase [20,21]. The formation of u0 phase in lamellar structures (type II) is probably resulted from the perpendicular decomposition of a2 lamellar [38,39]. As shown in Fig. 5(d), the interface between a2 and g lamellar is coherent and no interface dislocation exists at a2/g boundary; therefore, a large elastic strain might exist because of the misfit between lattice parameters of a2 and g phases. In addition, the lamellar spacing is only about 4 nm, indicating the interfacial energy of the ultra-thin lamellar structures is considerably high. In such a case, the lamellar structures are extremely unstable. Huang et al. [38] and Song et al. [39] reported that the u0 phase would precipitate from the single a2 lamellar, showing a sandwich feature between a2 sections. The nucleation of u0 phase usually occurs at a2/g interfaces because that these interfaces can provide preferred energy

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Fig. 7. XRD patterns of the samples after annealing for different time: (a) 30 min; (b) 60min; (c) 300 min and (d) 3000 min.

Fig. 8. (a, b, c) Direct nucleation of the u0 phase from B2 phase in S-30: (a) TEM bright image; (b) the magnified TEM bright image of the squared region in (a); (c) the corresponding SADP; (d, e, f) precipitation of the u0 phase from lamellar structures in S-60: (d) TEM bright image; (e) the SADP of squared region in (d) taken from the [1000] axis of the u0 phase; (f) the SADPs of the circled region in (d).

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and chemical composition conditions [38]. Thus, it is reasonably deduced that the nucleation of u0 phase occurs at the a2/g interfaces because of the large stored energy (elastic energy þ interface energy) in nano-lamellar structures and microsegregation of chemical composition at these interfaces. It is worth noting that the size of the u0 phase in Fig. 8(d) is much lager than the lamellar spacing, indicating that the growth of the u0 phase is not confined to a single a2 lamellar, which is probably caused by the large stored energy of the ultra-thin lamellar structures. Several types of g formation were observed during the annealing process and they can be simplified into two types. Type (I): the direct nucleation and growth of g phase within the B2 phase. As shown in Fig. 3 and 6, the direct nucleation of the g phase always accompanies the formation of the u0 phase in order to balance the distribution of Nb element. Previous investigations reported the u0 phase was a Nb-enriched phase, while the g phase was associated with lower Nb concentration; therefore, the formation of the u0 phase can provide an appropriate chemical component for the nucleation of the g phase [34,40]. Moreover, two g variants with twin-relation are always observed in the B2 regions. This phenomenon is probably resulted from that the formation of g twins is beneficial to spread the elastic strains in certain orientation, resulting in a minimum system elastic energy [33,34,40,41]. Type (II): the coarsen of the pre-existing g phases in (a2þg) lamellar structures through interface migration. Fig. 9 shows the microstructure of the S-300. It is evident from Fig. 9(a) that the growth of the g lamellar in the lamellar structures occurs by the way of consuming the adjacent B2 phase. No obvious interface was observed between the grown g grain and the g lamellar, indicating the coarsened g phase originates from

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the g lamellar without a new nucleation process. The SADPs given in Fig. 9(b) and Fig. 9(c) further confirm this kind of B2/g transition. As shown in Fig. 9(d), it is widely observed that the growth of the pre-existing g twins in lamellar structures occurs at some locations indicated by blue arrows, which results in the coarsened g phase with a twin feature. Fig. 10 shows the EBSD results of the S-3000. From the phase composition map shown in Fig. 10(a), it can be seen that the phase constitutes of S-3000 are u0, g and B2 phases with the phase fractions of 48.61%, 46.7% and 4.69%, respectively. The pole figures shown in Fig. 10 (b) indicate that there are 24 g variants and 4 u0 variants in this region. In the pole figures of g phase, the main reflexes (indicated by black square and triangle symbols) are surrounded by six satellites, and the formation of these satellites is attributed to the crystallographic symmetry of the B2 and g phases [42]. From the pole figures shown in Fig. 10(b), it can be seen that the orientation relationships among B2, u0 and g phases are {110}B2//{111}g//{11e20}u0 and <1e11>B2//<1e10>g//[0001]u0. Precisely speaking, there should be 72 g variants derived from one single B2 grain due to the difference of lattice parameters in g phase (lattice constant c is slightly larger than that of a and b). However, g phase was treated as a normal face-centred-cubic structure in this work; therefore, only 24 g variants were detected. As mentioned above, there are two different ways for the formation of u0 phase, i.e., directly nucleation from the B2 matrix and precipitation in the ultra-thin lamellar structures. Nevertheless, the amounts of the u0 variants are the same regardless of the precipitation process, which is resulted from that the u0 phases are originated from the same B2 grain. The g phase is in the same situation. In order to examine the morphologies of the g and u0 phases in S-3000, the IPF figures of the g phase and one u0 variant (V1 in Fig. 10(b)) were given. As

Fig. 9. (a) TEM image showing the growth of the g lamellar by consuming the adjacent B2 phase; (b) and (c) the SADPs of the squared and circled region in (a); (d) coarsen of the pre-existing g twins in lamellar structures.

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Fig. 10. EBSD results of the S-3000: (a) phase composition map; (b) pole figures of B2, u0 and g phases; (c) IPF map of the g phase; (d) misorientation distribution map of the g phase; (e) IPF map of one u0 variants, V1 indicated in (b).

shown in Fig. 10(c), a variety of g phases present as the twins with the size about 460 nm so that there is a sharp peak at 60 in the misorientation distribution map as shown in Fig. 10(d). According to the phase transition process, it can be deduced that one part of g twins are resulted from direct nucleation from B2 matrix (Fig. 6) and others are inherited from the lamellar structure as shown in Fig. 9(d). Moreover, it is obvious from Fig. 10(e) that the u0 phases present irregular shapes with the average size of 340 nm. In summary, a submicron microstructure composed of u0, g and B2 phases was obtained after annealing for 3000 min. The microstructure refinement is mainly caused by the considerably large driving force for the decomposition of the fully-B2 microstructure and the multiple phase transitions during annealing. However, the phase fractions in S-3000 are far from identical to that predicted by the calculated phase diagram shown in Fig. 4. Many researches confirmed that the microstructure of b-solidifying TiAl alloys was composed of B2 and g phases [13,43e46]. For example, a fully-B2 microstructure was obtained in Tie44Ale7Mo alloys by waterquenching from the single b phase region, and a submicron microstructure composed of B2 and g phases was produced after subsequent heating process [47]. The size of the phase constitutes in Ref. [47] is similar to that in the present study, which is caused by the high instability of the B2 phase. However, the phase fractions are absolutely different, which is closely related to the chemical composition of the two alloys, i.e., Mo element can suppress the formation of u0 phase [48] and the B2 phase is prone to transform into u0 phase in high Nb containing TiAl alloys [17e24]. Based on an assumption that most of the B2 phase in the (B2þg) microstructure decomposes into u0 phase, the phase constitutes of S-3000 seem to be possible. Besides, it is quite curious that the a2 phase precipitates from the B2 matrix at the initial stage, and then disappears during subsequent annealing. This phenomenon may be closely related to the energy barrier of the B2/g and B2/u0 transitions, i.e., the presence of a2 phase probably reduces the energy barrier of these

transformations, and the detailed mechanism need to be further explained. 4. Conclusions In this work, the phase precipitation behavior of a quenched bsolidifying TiAl alloy with a fully-B2 microstructure during annealing at 800  C was investigated. The main conclusions can be drawn as follows: (1) A fully-B2 microstructure was obtained in a Tie40Ale8Nb alloy by water-quenching from 1400  C. The B2 phase is extremely unstable and readily decomposes into a2, u0 and g phases during annealing at 800  C. (2) At the initial stage of the annealing, numerous lath-liked a2 phases precipitated from the B2 matrix following the Burgers orientation relationships, and these a2 phases transformed into ultra-thin lamellar structures with the lamellar spacing of 4 nm in the subsequent annealing process. (3) After annealing at 800  C for 3000 min, a submicron microstructure with the phase constitutes of u0, g and B2 phases was produced, and the phase fractions of these three phases were 48.61%, 46.7% and 4.69%, respectively. (4) Two formation processes of the u0 and g phases were observed, including the direct nucleation from the B2 matrix and the precipitation from the ultra-thin lamellar structures. Acknowledgement The authors thank the National Natural Science Foundation of China (No. 51701107, No. 51571162, No. 51805308, No. 51805309), the Natural Science Basic Research Plan in Shaanxi Province of China (No. 2018JQ5161, No. 2019JQ-303) for their financial support.

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References [1] K. Kothari, R. Radhakrishnan, N.M. Wereley, Advances in gamma titanium aluminides and their manufacturing techniques, Prog. Aerosp. Sci. 55 (2012) 1e16. [2] G. Chen, Y.B. Peng, M.Z. Wang, H.C. Yu, C.L. Dong, C.T. Liu, Polysynthetic twinned TiAl single crystals for high-temperature applications, Nat. Mater. 15 (2016) 876e881. [3] F. Apple, H. Clemens, F.D. Fischer, Modeling concepts for intermetallic titanium aluminides, Prog. Mater. Sci. 81 (2016), 22-124. [4] Y. Jin, J.N. Wang, J. Yang, Y. Wang, Microstructure refinement of cast TiAl alloys by b solidification, Scr. Mater. 51 (2004) 113e117. [5] R.M. Imayev, V.M. Imayev, M. Oehring, F. Appel, Alloy design concepts for refined gamma titanium aluminide based alloys, Intermetallics 15 (2007) 451e460. [6] G. Yang, H.C. Kou, J.R. Yang, J.S. Li, H.Z. Fu, Microstructure control of Ti45Al8. 5Nb (W, B, Y) alloy during the solidification process, Acta Mater. 112 (2016) 121e131. [7] G. Yang, H.C. Kou, J.R. Yang, J.S. Li, H.Z. Fu, In-situ investigation on the b to a phase transformation in Ti-45Al-8.5 Nb-(W, B, Y) alloy, J. Alloy. Comp. 663 (2015) 594e600. [8] X. Wu, Review of alloy and process development of TiAl alloys, Intermetallics 14 (2006) 1114e1122. [9] B. Tang, L. Cheng, H.C. Kou, J.S. Li, Hot forging design and microstructure evolution of a high Nb containing TiAl alloy, Intermetallics 58 (2015) 7e14. [10] J.G. Wang, T.G. Nieh, Creep of a beta phase-containing TiAl alloy, Intermetallics 8 (2000) 737e748. [11] T. Ye, L. Song, Y.F. Liang, M.H. Quan, J.P. He, J.P. Lin, Precipitation behavior of uo phase and texture evolution of a forged Ti-45Al-8.5Nb-(W, B, Y) alloy during creep, Mater. Char. 136 (2018) 41e51. [12] L. Chen, J.P. Lin, X.J. Xu, C. Li, Y. Xu, Y.F. Liang, Microstructure refinement via martensitic transformation in TiAl alloys, J. Alloy. Comp. 714 (2018) 1175e1182. [13] Y. Chen, L. Cheng, L.Y. Sun, Y.L. Lu, G. Yang, H.C. Kou, E. Bouzy, Characterization of a new microstructure in a b-solidifying TiAl alloy after air cooling from a b phase field and subsequent tempering, Metals 8 (2018) 156e166. [14] L. Cheng, H. Chang, B. Tang, H.C. Kou, J.S. Li, Deformation and dynamic recrystallization behavior of a high Nb containing TiAl alloy, J. Alloy. Comp. 552 (2013) 363e369. [15] H.Z. Niu, Y.Y. Chen, S.L. Xiao, L.J. Xu, Microstructure evolution and mechanical properties of a novel beta g-TiAl alloy, Intermetallics 31 (2012) 225e231. [16] M. Schloffer, F. Iqbal, H. Gabrisch, E. Schwaighofer, F.P. Schimansky, S. Mayer, €ken, F. Pyczakc, H. Clemens, Microstructure A. Stark, T. Lippmann, M. Go development and hardness of a powder metallurgical multi phase g-TiAl based alloy, Intermetallics 22 (2012) 231e240. [17] L. Song, X.J. Xu, J. Sun, J.P. Lin, Cooling rate effects on the microstructure evolution in the bo zones of cast Ti-45Al-8.5Nb-(W, B, Y) alloy, Mater. Char. 93 (2014) 62e67. [18] L. Song, J.P. Lin, J.S. Li, Effects of trace alloying elements on the phase transformation behaviors of ordered u phases in high Nb-TiAl alloys, Mater. Des. 113 (2017) 47e53. [19] Z.W. Huang, Ordered u phases in a 4Zr-4Nb-containing TiAl-based alloy, Acta Mater. 56 (2008) 1689e1700. [20] X.Y. Wang, J.R. Yang, K.R. Zhang, R. Hu, L. Song, H.Z. Fu, Atomic-scale observations of B2/u-related phases transition in high-Nb containing TiAl alloy, Mater. Char. 130 (2017) 135e138. [21] X.Y. Wang, J.R. Yang, L. Song, H.C. Kou, J.S. Li, H.Z. Fu, Evolution of B2(u) region in high-Nb containing TiAl alloy in intermediate temperature range, Intermetallics 82 (2017) 32e39. [22] T.T. Cheng, M.H. Loretto, The decomposition of the beta phase in Ti-44Al-8Nb and Ti-44Al-4Nb-4Zr-0.2Si alloys, Acta Mater. 46 (1998) 4801e4819. [23] Z.W. Huang, W.E. Voice, P. Bowen, Thermal stability of Ti-46Al-5Nb-1W alloy, Mater. Sci. Eng. A 329 (2002) 435e445. [24] L. Song, X.J. Xu, L. You, Y.F. Liang, J.P. Lin, Ordered u phase transformations in Ti-45Al-8.5Nb-0.2B alloy, Intermetallics 65 (2015) 22e28. [25] V.T. Witusiewicz, A.A. Bondar, U. Hecht, T.Y. Velikanova, The Al-B-Nb-Ti system: IV. Experimental study and thermodynamic re-evaluation of the binary Al-Nb and ternary Al-Nb-Ti systems, J. Alloy. Comp. 472 (2009) 133e161. [26] G. Yang, W. Ren, Y.H. Liu, W.J. Song, F.B. Han, Y. Chen, L. Cheng, Effect of predeformation in the b phase field on the microstructure and texture of the a

[27]

[28]

[29] [30]

[31]

[32]

[33]

[34]

[35]

[36]

[37] [38]

[39]

[40]

[41]

[42]

[43]

[44]

[45]

[46]

[47]

[48]

9

phase in a boron-added b-solidifying TiAl alloy, J. Alloy. Comp. 742 (2018) 304e311. P.J. Tian, G. Yang, Z.H. Ge, Y.F. Wang, L. Cheng, Y.H. Liu, H.C. Kou, Responses of microstructure and texture of a phase to boron addition in Ti-40Al-8Nb-xB alloys modified by hot deformation above the b transus, Mater. Char. 153 (2019) 148e156. T. Klein, S. Niknafs, R. Dippenaar, H. Clemens, S. Mayer, Grain growth and b to a transformation behavior of a b-solidifying TiAl alloy, Adv. Eng. Mater. 17 (2015) 786e790. S.C. Wang, M. Aindow, M.J. Starink, Effect of self-accommodation on a/a boundary populations in pure titanium, Acta Mater. 51 (2013) 2485e2503. K. Hua, Y.D. Zhang, W.M. Gan, H.C. Kou, J.S. Li, C. Esling, Correlation between imposed deformation and transformation lattice strain on a variant selection in metastable b-Ti alloy under isothermal compression, Acta Mater. 161 (2018) 150e160. Y. Chen, H.C. Kou, L. Cheng, K. Hua, L.Y. Sun, Y.L. Lu, E. Bouzy, Crystallography of phase transformation during quenching from b phase field of a V-rich TiAl alloy, J. Mater. Sci. 54 (2019) 1844e1856. B. Tang, W.C. Ou, H.C. Kou, J.S. Li, Experimental study of the dissolution and reprecipitation of u0 phase in high Nb containing TiAl alloy, Mater. Char. 109 (2015) 122e127. L. Song, J.P. Lin, J.S. Li, Phase transformation mechanisms in a quenched Ti45Al-8.5Nb-0.2W-0.2B-0.02Y alloy after subsequent annealing at 800 oC, J. Alloy. Comp. 691 (2017) 60e66. L. Song, X.J. Xu, L. You, Y.F. Liang, J.P. Lin, Phase transformation and decomposition mechanisms of the b0(u) phase in cast high Nb containing TiAl alloy, J. Alloy. Comp. 616 (2014) 483e491. J.K. Fan, J.S. Li, H.C. Kou, K. Hua, B. Tang, Y.D. Zhang, Microstructure and mechanical property correlation and property optimization of a near b titanium alloy Ti-7333, J. Alloy. Comp. 628 (2016) 517e524. R.F. Dong, J.S. Li, J.K. Fan, H.C. Kou, B. Tang, Precipitation behavior of a phase during aging treatment in a b-quenched Ti-7333, Mater. Char. 140 (2018) 275e280. S. Zghal, M. Thomas, S. Naka, A. Finel, A. Couret, Phase transformations in TiAl based alloys, Acta Mater. 53 (2005) 2653e2664. Z.W. Huang, W. Voice, P. Bowen, Thermal exposure induced a2þg/B2(u) and a2/B2(u) phase transformation in a high Nb containing fully lamellar TiAl alloy, Scr. Mater. 48 (2003) 79e84. L. Song, X.J. Xu, L. You, Y.F. Liang, Y.L. Wang, J.P. Lin, Ordered a2 to u0 phase transformations in high Nb-containing TiAl alloys, Acta Mater. 91 (2015) 330e339. T. Ye, L. Song, S.B. Gao, Y.F. Liang, Y.L. Wang, J.P. Lin, Precipitation behavior of the u0 phase in an annealed high Nb-TiAl alloy, J. Alloy. Comp. 701 (2017) 882e891. A. Denquin, S. Naka, Phase transformation mechanisms involved in two-phase TiAl-based alloys-I. Lambellar structure formation, Acta Mater. 44 (1996) 343e352. M. Michael, S. Mayer, C. Pauly, H. Clemens, F. Mücklich, 3 D characterization of an intermetallic b/g-Titanium aluminide alloy, Adv. Eng. Mater. 15 (2013) 1125e1128. X.B. Li, H. Xu, W.W. Xing, B. Chen, Y.C. Ma, K. Liu, Phase transformation behavior of a b-solidifying g-TiAl-based alloy from different phase regions with various cooling methods, Metals 8 (2018) 731e740. Y.J. Su, F.T. Kong, Y.Y. Chen, N. Gao, D.L. Zhang, Microstructure and mechanical properties of large size Ti-43Al-9V-0.2Y alloy pancake produced by packforging, Intermetallics 34 (2019) 29e34. F.T. Kong, N. Cui, Y.Y. Chen, X.P. Wang, A novel composition design method for beta-gamma TiAl alloys with excellent hot workability, Metall. Mater. Trans. A 49 (2018) 5574e5584. L. Cheng, J.S. Li, X.Y. Xue, B. Tang, H.C. Kou, O. Perroud, E. Bouzy, Effect of b/B2 phase on cavitation behavior during superplastic deformation of TiAl alloys, J. Alloy. Comp. 693 (2017) 749e759. P. Erdely, P. Staron, A. Stark, T. Klein, H. Clemens, S. Mayer, In situ and atomicscale investigations of the early stages of g precipitate growth in a supersaturated intermetallic Ti-44Al-7Mo (at.%) solid solution, Acta Mater. 164 (2019) 110e121. € berl, E. Schwaighofer, Z.L. Zhang, H. Clemens, M. Schloffer, B. Rashkova, T. Scho S. Mayer, Evolution of the u0 phase in a b-stabilized multi-phase TiAl alloy and its effect on hardness, Acta Mater. 64 (2014) 241e252.