Microstructure evolution of Ti44Al alloy during directional induction heat treatment and its effect on mechanical properties

Journal Pre-proof Microstructure evolution of Ti44Al alloy during directional induction heat treatment and its effect on mechanical properties Yangli Liu, Xiang Xue, Hongze Fang, Ruirun Chen, Yingmei Tan, Yanqing Su, Jingjie Guo PII:

S0921-5093(19)31487-X

DOI:

https://doi.org/10.1016/j.msea.2019.138701

Reference:

MSA 138701

To appear in:

Materials Science & Engineering A

Received Date: 6 October 2019 Revised Date:

15 November 2019

Accepted Date: 17 November 2019

Please cite this article as: Y. Liu, X. Xue, H. Fang, R. Chen, Y. Tan, Y. Su, J. Guo, Microstructure evolution of Ti44Al alloy during directional induction heat treatment and its effect on mechanical properties, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/j.msea.2019.138701. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Author Contributions Ruirun Chen designed and supervised the project. Yangli Liu conducted the experiments, performed the microstructure observation and mechanical properties tests. Yingmei Tan and Hongze Fang prepared the samples. Chen Ruirun and Yangli Liu analyzed experimental data and wrote the study. All authors have discussed the results and commented on the manuscript.

Microstructure evolution of Ti44Al alloy during directional induction heat treatment and its effect on mechanical properties Yangli Liua, Xiang Xuea, Hongze Fanga*, Ruirun Chena,b*, Yingmei Tana, Yanqing Sua, Jingjie Guoa a

National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technique, Harbin 150001, China b School of Materials Science and Technique, Shandong University of Science and Technique, Qingdao, 266590, China ∗Corresponding author, Ruirun Chen, Hongze Fang Tel.: +86-0451-8641-2394; E-mail: [email protected], [email protected]

Abstract For realizing equiaxed to columnar transition, Ti44Al alloy was processed by directional induction heat treatment (DHT) and mechanical properties were tested. The results show that columnar grains can be obtained and elongation is improved at room temperature. During this heat treatment, Ti44Al alloy was heated in single β phase domain. The microstructure is consisted of fully lamellar colonies after DHT. The average lamellar spacing is decreased, and the orientation of lamellar colonies with angle of less than 60°is more than 76%. Moreover, grain boundary presents zigzag. Tensile testing shows that strength and elongation of heat-treated alloy are 217 MPa and 1.78% respectively, they are 253 MPa and 1.63% for as-cast Ti44Al alloy. The elongation is increased about 9.2% although tensile strength is decreased. Cleavage fracture is main way in as-cast Ti44Al alloy but a mixed fracture consisting of some cleavage planes and many tearing edges is in Ti44Al alloys after DHT. The fracture morphology shows that deterioration of strength is mainly resulted from the micropores that reserved from as-cast alloy, it means DHT could not eliminate the micropores. The improvement of ductility is caused by decreasing of lamellar spacing and changing orientation lamellar colonies. Key words: TiAl alloy; Microstructure; β transformation; Heat treatment 1. Introduction TiAl based alloy has been considered as the better candidate for replacing the Ni-based alloy in range of 873 K - 1273 K, due to low density, high specific strength and high creep resistance [1-5]. However, there are inherent defects in TiAl based alloy. 1

Generally, elongation of as-cast TiAl alloys at room temperature is low, which is difficult to meet the plasticity requirement of structural parts [6-10]. It is a reason why materials researchers have been seeking to improve the plasticity of TiAl based alloys at room temperature in recent years [11-14]. Cold crucible direction solidification (CCDS) technique is an effective method to melt refractory and reactive alloys with low contamination. Metals or alloys are melted by electromagnetic induction heating, and the melt will have soft contact with water-cooled copper crucible under action of electromagnetic force [15]. A directional microstructure will be obtained for single direction heat transfer during CCDS. So mechanical properties of TiAl based alloys have been improved greatly. At present, the elongation of Ti-46Al-0.5W-0.5Si alloy prepared by CCDS technique, can be 3.6% at room temperature [16]. However, in process of CCDS, the liquid-solid interface may be bent due to the contact between metal melt and water-cooled copper crucible. Thus, grains nearby edge of ingot deflect easily and mechanical properties of alloy decrease [17, 18]. DHT technique is proposed by us in order to solve the problem on deflection of columnar grains in the surface area of TiAl alloy ingot during CCDS. The core thought of DHT is to guarantee grains in effective heat treatment area undergoing dynamic continuous solid phase transformation. In the process of DHT, grains will grow up along axial direction by controlling loaded power, withdrawing rate and section size of test bar. At present, finite element analysis and numerical simulation have been carried out on temperature distribution of test bars with different section size under the same loaded power as well as with the same section size under different loaded power. The simulation results have been verified by experiment, as detailed in the reference [19]. In addition, Z. W. Zhang et al. studied microstructure evolution and mechanism of columnar grain growth in Fe-6.5wt%Si alloy and commercial pure iron during directional annealing and concluded that deformation is not a key factor in forming columnar grain structure. Surface energy and grain boundary energy play important roles in formation of the columnar grain structure [20-22]. I. Baker et al. conducted directional annealing on cold-rolled copper single crystals and hot-rolled ODS nickel-based superalloy MA 754 alloy and found that temperature gradient ahead of the hot zone is a key parameter for directional annealing 2

[23, 24]. In this study, effect of DHT on the microstructure and mechanical properties of Ti44Al alloy is investigated based on the previous experimental results. The main effects of DHT on microstructure include lamellar space, lamellar thickness and lamellar phase orientation. Furthermore, the relationship between DHT and mechanical properties is explored and analyzed, as well as fracture morphology and toughening mechanism of Ti44Al alloy in tensile process at room temperature. It is essential to optimize parameters of DHT. 2. Experiment procedures 2.1 Preparation of Ti44Al alloy Ti44Al alloy (at%, similar hereinafter) ingot with an actual composition of Ti-44.5Al was fabricated by vacuum induction furnace under argon atmosphere. The raw materials are sponge titanium (purity, 99.98%) and aluminium (purity, 99.98%). Samples were cut from 1/4 and 3/4 locations of TiAl cylindrical ingot for XRF (X Ray Fluorescence) composition analysis and the rest samples were used for DHT. The chemical composition results are shown in Table 1. Table 1. Chemical composition of TiAl ingot (at%).

location

Ti

Al

1/4

55.32

44.68

3/4

55.58

44.42

2.2 DHT for Ti44Al alloy Ti44Al alloy sample (diameter with 20 mm) was placed in a multi-function induction heating furnace for DHT surrounding with a vacuum of 300 Pa argon atmosphere. Loaded power is 21.6 kW and temperature for DHT is 1750 K. Withdrawing rate is 4.17 µm/s and range for treatment is 40 mm. Sample has been treated for 4 times to ensure grains grow up sufficiently along axial direction. 2.3 Microstructure of Ti44Al alloy after DHT Sample after DHT was split along axis direction, polished and corroded. The morphology was observed by Nikon D850 Digital Single Lens Reflex and ultra-depth optical microscope (OM). Lamellar space, lamellar thickness, fracture morphology and 3

phase component were determined by scanning electron microscopy in BSE mode (SEM-BSE) and energy dispersive X-ray spectroscopy (EDS). Lamellar spacing was measured by Nano-Measure software. 2.4 Mechanical properties of Ti44Al alloy At room temperature, tensile test was carried out on the 5569-Instron testing machine. Geometric dimension of tensile sample is shown in Fig. 1 and 3 samples were tested for each group. Strain rate is 2×10-4s-1.

Fig. 1. Tensile test specimen.

3. Results and analysis 3.1 Macrostructure of Ti44Al alloys Macrostructures of as-cast Ti44Al alloy and Ti44Al alloy treated by DHT (denoted as DHT-Ti44Al, hereinafter) are shown in Fig. 2. It can be seen that grain morphology changes obviously in DHT-Ti44Al alloy. Grains grow up significantly and there is a transition region about 15mm height for equiaxed grains transforming to columnar grains in early stage of DHT. The reason is that height of effective heat treatment area is about 15 mm for 5-turn square induction coil. In the initial stage of DHT, thermal insulation treatment should be carried out for 15 min before sample being pulled downwards when loaded power reaches the set value by stepping formula. At that time, grains in effective heat treatment area almost grow up along all direction to form equiaxed grains. Above the transition region, it is directional growth region where grains enter into effective heat treatment area continuously and transfer to single β phase (called as β phase transformation). Due to the same crystal structure, new transferring β phase will attach to existing β phase to grow up under the continuous movement of withdrawing device. The macrostructure of DHT-Ti44Al alloy is shown in Fig. 2 b. Length of columnar grain is measured by Nano-Measure software and results 4

indicate that the max length of columnar grain in DHT-Ti44Al alloy is about 25.88 mm. It is almost the same length as single crystal obtained by Prof. Chen using optical floating zone furnace [25]. The macrostructure also shows that columnar grains nearby edge of ingot have never deflected after DHT and β phase transformation almost occurs in all grains. This results verify the correctness of transformation and growth model proposed in reference [19] and illustrates the feasibility of DHT.

a

b

Directional growth region

Transition region

Growth direction Fig. 2. Macrostructure of Ti44Al alloy. a, as-cast Ti44Al; b, DHT-Ti44Al.

3.2 Microstructure of Ti44Al alloys Microstructure of as-cast Ti44Al alloy and DHT-Ti44Al alloy are shown in Fig. 3. Macrostructure of DHT-Ti44Al alloy in white dotted frame of Fig. 2 b is detected by SEM as Fig. 3 b. It can be seen that microstructure of as-cast Ti44Al alloy is fully lamellar morphology and this microstructure is kept in DHT-Ti44Al alloy. Fig. 3 b is magnified as Fig. 3 c. EDS results indicate that the black contrast lamellar is γ phase and the gray contrast lamellar is α2 phase. In Fig. 3 c, some granular black morphology can be found on the α2 phase. Combined with EDS results, it can be found that this morphology is formed by corrosion of α2 phase. Because the same sample is observed by OM and SEM in order to investigate morphology on the same region further. So the sample was corroded slightly before OM analysis. As well, it can be seen that corrosion degree is more obvious at the junction of different lamellar phase orientation (grain boundary), as arrows point in Fig. 3 b. It is possible that α2 phases at these positions are 5

subjected to great stress, so the corrosion pits are bigger and deeper.

a

b

c

Fig. 3. SEM-BSE of as-cast and heat-treated Ti44Al alloys. a, as-cast Ti44Al alloy; b, DHT-Ti44Al alloy; c, the magnified microstructure of b.

It shows obviously that grain boundary of as-cast Ti44Al alloy is relatively wide and the average lamellar spacing in two adjacent grains is large difference, seen in Fig. 3 a. The main reason is that in process of solidification, heat dissipation is non-uniform and adjacent grains grow up insufficiently, which leads to the area between dendrite grains becoming larger and eventually forms a wide grain boundary. The spacing difference of lamellar phases in adjacent grains is due to different heat dissipation, resulting in different diffusion rate of Al atom during solid phase transformation. Furthermore, nucleation and growth of γ phase is affected and finally spacing of lamellar phase is different. Solid phase transformation would occur and form single β phase when Ti44Al alloy is directional induction heat treated at 1750 K. Grains completed β phase transformation would move directly with pulling bar and be cooled slowly. So atoms can diffuse relatively sufficiently during DHT and an uniform fully lamellar microstructure has been obtained. The reduction of intergranular region improves morphology of grain boundary. Finally a zigzag boundary is formed as shown in Fig. 3 b, which is particular beneficial to creep properties of TiAl-based alloy at high temperature [26, 27]. 6

Spacing for α2 phase/ µm

1.2

1.0

0.8

0.6

0.4

0.2

b

1.4

Spacing for γ phase/ µ m

a

1.2 1.0 0.8 0.6 0.4

0

2

0.2

4

0

2

times

4

times

Fig. 4. The spacing of lamellar phase in Ti44Al alloys. a, α2 phase; b, γ phase.

Fig. 4 shows spacing distribution of lamellar phases in as-cast Ti44Al alloy and DHT-Ti44Al alloy. It can be seen that average spacing of α2 phase and γ phase is about 0.79 µm and 0.85 µm in as-cast Ti44Al alloy, respectively. After DHT, the average spacing of α2 phase and γ phase decreases. They are about 0.57 µm and 0.60 µm, respectively. In addition, Fig. 4 also shows that deviation between actual measured spacing value and average spacing value is large in as-cast Ti44Al alloy (standard deviation is 0.21 and 0.27, respectively). However, the deviation is small in DHT-Ti44Al alloy (both standard deviation are 0.15). Therefore, it can be concluded that spacing of lamellar phase would be smaller and lamellar phase would be more uniform in Ti44Al alloy after DHT. When sample is subjected to external force, a composite strengthening effect would occur due to lamellar phase being thin and uniform. It is effective to prevent stress concentration and crack propagation, which is very beneficial to improve mechanical properties of TiAl alloy [28].

a

b

50 As-cast

DHT

Percentage / %

40 30 20 10 0

Growth direction

0°~30°

30°~60° Angel / °

60°~90°

Fig. 5. Orientation distribution of lamellar phase after DHT. a, orientation of lamellar phase; b, statistical results.

Orientation distribution of lamellar phase in DHT-Ti44Al alloy is shown in Fig. 5. 7

The different color lines stand for direction difference between lamellar phase orientation and growth direction of columnar grains in OM. Green line means the direction difference within 0~30°, red line means within 30°~60° and purple line means within 60°~90°. It can be concluded that most columnar grains in Ti44Al alloy grow up along axial direction after DHT for 4 times at 1750 K. Then as the temperature decreases, a strict lattice corresponding relationship is maintain between α phase and β phase in the process of solid phase transformation. Therefore, 40% lamellar phase keeps in 0~30° orientation difference with growth direction of columnar grains. 37% of lamellar phase keeps in 30°~60° (the most is 45°) and only 23% lamellar phase keeps in 60°~90° (the most is 90°). The main reason is that Ti44Al alloy is in single β phase domain at 1750 K during DHT. As the sample moves downwards, grains would go through (β + α), α and (α + γ) phase domains in turns. Due to the lattice corresponding relationship of β // α and α // γ being (110)β // (0001)α and (0001)α // (111)γ, respectively, the angle are only 0° and 45° between lamellar phase orientation and growth direction of β grain [12, 29]. Therefore, the direction difference between lamellar phase orientation and growth direction of columnar grains is mainly 0° or 45° after DHT. However, it is probable that non-equilibrium solidification occurs during the preparation of Ti44Al ingot. Due to segregation of Al atoms, a little amount of α solidification exists in as-cast Ti44Al alloy. When as-cast Ti44Al alloy was treated at 1750 K for DHT, grains did not transfer β phase completely and γ phase directly precipitates from α phase in the cooling process. Because of the lattice corresponding relationship of α // γ, the angle difference is 90° between lamellar phase orientation and growth direction of columnar grains [4]. 3.3 Mechanical properties and fracture morphology of Ti44Al alloys Tensile test results indicate that tensile strength and elongation of C-Ti44Al alloy are 253 MPa and 1.63%. Under the same condition, they are 217 MPa and 1.78% for DHT-Ti44Al alloy. Tensile strength decreases at room temperature but the elongation increases by nearly 9.2% for Ti44Al alloy after DHT. By analyzing microstructure of Ti44Al alloys in Fig. 2 and Fig. 3, it can predict that mechanical properties of DHT-Ti44Al alloy would been improved significantly. But tensile test results are 8

unexpected by us. So fracture morphology of Ti44Al alloys were tested by SEM, like as Fig. 6.

a

b

c

d

Fig. 6. Fracture morphology of Ti44Al alloys. a and b for as-cast Ti44Al alloy; c and d for DHT-Ti44Al alloy.

Fig. 6 shows fracture morphology of as-cast Ti44Al alloy and DHT-Ti44Al alloy at room temperature. It can be seen that fracture mode of as-cast Ti44Al alloy is cleavage fracture but a mixed fracture consisting of some cleavage planes and many tearing edges is in DHT-Ti44Al alloy, as marked by arrows in Fig. 6 a and 6 c. The lamellar phase orientation is perpendicular or at a certain angle to the direction of tensile stress, the fracture presents cleavage plane. While the lamellar phase orientation is parallel to the direction of tensile stress, the fracture presents tearing edges. By comparing Fig. 6 b with 6 d, it shows dense tearing edges in DHT-Ti44Al alloy. As reasons are given above, atoms can diffuse enough during DHT and γ phase could precipitate from α phase enough as well. So the spacing of lamellar phase is small and uniform. Fig. 6 c shows that there are some micropores with a diameter about 20 µm in the fracture morphology of DHT-Ti44Al alloy. And there are cracks around these micropores, as marked by arrows in Fig. 6 d. Therefore, these micropores become the source of crack propagation during tensile test. They are extended rapidly to form cracks with large size in the 9

tensile process. The existence of cracks must worsen mechanical properties of DHT-Ti44Al alloy. So the tensile strength of DHT-Ti44Al alloy do not increase significantly. 4. Discussion 4.1 Microstructure of Ti44Al alloy during DHT As temperature decreases for Ti44Al alloy after solidification, a successive solid phase transformation β → (β + α) → α → (α + γ) → (α2 + γ) would occur. And in the process of solid phase transformation, Burgers and Blackburn orientation relationship exist among β, α and γ phases [30-32]. Fig. 2 b is magnified into Fig. 7. It is clear that columnar grain is composed of lamellar clusters and network microstructure, as shown in Fig. 7 a. The yellow parallel double arrows in Fig. 7 a indicate growth direction of lamellar cluster. The blue arrow indicates orientation of lamellar phase. The angle difference is near about 45° between growth direction of lamellar cluster and growth direction of columnar grain. While lamellar phase orientation is almost parallel to the growth direction of columnar grain. The magnified Fig. 7 b indicates that these lamellar clusters are actually connected each other, as marked by arrows, and the white network-like microstructure is transitional zone. According to the EDS results, the black contrast phase is γ containing rich Al element while the gray contrast phase is α2 containing rich Ti element. Therefore, it can be inferred that after β phase transition, large number of α phases precipitate simultaneously on the isothermal cross section of β grain when Ti44Al sample is pulled downwards gradually far away from effective heating range of coil. Due to solid solubility of Ti atom in α phase is less than that in β phase, a gray network-like zone with rich Ti atoms is formed between α phases, namely a transition zone, shown in Fig. 7 a. Moreover, due to the lattice corresponding relationship of β and α phase, the angle difference is 45° between growth direction of lamellar cluster and growth direction of columnar grain. Since grain re-precipitation does not occur but only the short-range diffusion of atoms takes place in the process of β → β + α phase transition, complete coherent relationship is in adjacent α phases. As temperature continues to decrease, α phase grows up firstly and then γ phase precipitates from α phase. Finally, (α2 + γ) lamellar microstructure is formed. Due to 10

complete coherent relationship between α phases as well as directional heat conduction during DHT, the orientation of (α2 + γ) lamellar cluster is the same, shown in Fig. 7 a.

a Growth direction

b

Fig. 7. OM of DHT-Ti44Al alloy. a, OM of columnar grain; b, magnified OM of a.

In addition，Chen et al. found that the lamellar phase orientation is related to the anisotropic interfacial energy when α phase precipitates from β phase [25]. Compared with orientation difference of 45°, the β / α interfacial energy is less when lamellar phase orientation is parallel to growth direction of columnar grain. According to solid phase transformation theory and literature [33], β / α interfacial energy is a resistance for precipitation but a driving force for growth. It is easier to precipitate α phase with 0° than that with 45° due to interfacial energy difference. Hence, if the precipitation driving force is enough for α phase with 0° but it may be insufficient for α phase with 45°. If the precipitation driving force is enough for α phase with 0° and 45° as well, α phase with 45° would grow up more easily and faster. In the process of solid phase transformation, the precipitation driving force is closely related to the supercooling effect and supercooling degree increases with the increase of cooling rate. The relationship among cooling rate (C), temperature gradient (G) and withdrawing rate (V) is as follows [33], C=G*V Therefore, temperature gradient should be controlled reasonably to limit the orientation difference between lamellar phase orientation and growth direction of columnar grain under a constant withdrawing rate. During DHT, temperature gradient of effective heat treatment area can be controlled by changing height distance between loaded coil and liquid Ga-In pool. When height distance is 58 mm, the average temperature gradient of effective heat treatment area is about 18 K/mm. Meanwhile, DHT technique can rapidly heat Ti44Al alloy and guarantee the temperature of effective heat treatment area in single β phase 11

domain. On the other hand, DHT technique can ensure β grains in effective heat treatment area to grow up continuously and along axial direction under the temperature gradient about 18 K/mm on average. Therefore, the morphology of Ti44Al alloy changes significantly after DHT. Grains are arranged along axial direction of sample. Due to relatively severe heat conduction during DHT, a large number of lamellar phases precipitate simultaneously and thickness of lamellar phase decreases greatly. Distribution of lamellar phase is more uniform. The angle difference between lamellar phase orientation and growth direction of columnar grains is mainly 0° or 45°. Furthermore, since Ti44Al sample contacts with nothing except Ar gas to avoid lateral heat dissipation effectively during DHT, there is no deflection for all grains in the surface area of sample. It also proves that at room temperature grain morphology of TiAl alloys via β solidification is derived from β grain morphology. By controlling heat transfer during DHT, not only columnar grains can be obtained but also it is easy to form lamellar phase parallel to growth direction of columnar grains. So mechanical properties of TiAl alloys would be improved. 4.2 Mechanical properties of Ti44Al alloy after DHT Mechanical properties can be greatly improved under a condition of the same direction between lamellar phase orientation and tensile stress. After DHT, not only columnar grains are arranged along axial direction of sample, but also the number of casting defects with little size is reduced. The spacing of lamellar phase is more uniform and the thickness of grain boundary decreases in DHT-Ti44Al alloy. Grain boundary presents zigzag, as shown in Fig. 3. The main reason is that after DHT a series of solid phase transformation, β → β + α → α → (α + γ) → (α2 + γ), would occur in the process of cooling. In particular, the α → (α + γ) → (α2 + γ) phase transition occurs resulting in fully lamellar microstructure at room temperature. Since supercooling degree is relative high and distribution of temperature gradient is uniform during DHT, a large number of γ phases precipitate simultaneously on the same isothermal cross section, seen in Fig. 8. Grains at the upper end of effective heat treatment area are equiaxial grains just completing β phase transition. As Ti44Al sample is pulled down, grains completing β 12

phase transition will attach to existing β grains. Under action of interface tension, grain boundary will migrate directionally against the direction of heat conduction during DHT. In this way, β grains can grow up directionally. Due to cooling action of liquid Ga-In alloy at the bottom of Ti44Al sample, an isothermal cross section for γ phase transition exists below effective heat treatment area, shown in Fig. 8. Under action of directional heat conduction, γ lamellar phase will grow up along axial direction of sample. Moreover, isothermal cross section of γ phase is relatively straight, γ phase precipitate at many locations of the isothermal cross section simultaneously. As a consequence, the spacing of lamellar phase decreases greatly and lamellar phase distributes uniformly.

Fig. 8. Schematic diagram of lamellar phase formation during DHT.

Moreover, when Ti44Al alloy is heated to 1750 K, internal shrinkage defects will be improved under gravity due to movement downwards with low speed. Therefore, when Ti44Al alloy is conducted tensile experiment at room temperature, stress concentration and crack initiation points reduce. It is beneficial to improve mechanical properties of DHT-Ti44Al alloy, especially for elongation. Fracture morphology shows that fracture model of DHT-Ti44Al alloy is a mixture including cleavage and tearing. Fracture morphology consists of many cleavage steps and tearing edges. Cleavage steps are formed due to orientation of lamellar phase perpendicular to tensile stress. The existence of tearing edges indicates that plasticity of Ti44Al alloy is improved after DHT. Because tearing edge is formed in the process of material transition from 13

brittleness to plasticity. In addition, fracture morphology also indicates that height of cleavage step decreases for the spacing of lamellar phase decreases after DHT. In this way, a large amount of energy is consumed during propagation of cracks through lamellar phases. So comprehensive mechanical properties of DHT-Ti44Al alloy will be improved. However, mechanical properties of DHT-Ti44Al alloy have not been improved significantly. The fracture morphology shows reasons to us. The fracture morphology of Ti44Al alloy shows that some defects with large size are almost not eliminated in the process of DHT, for example porosity and shrinkage cavity. We have measured the size of pores and shrinkage cavity in fracture of as-cast Ti44Al alloy and DHT-Ti44Al alloy. The results indicate that the average size of shrinkage cavity in as-cast Ti44Al alloy is about 13.94 µm and the maximum size of micropore is about 34.75 µm. After DHT, the average size of shrinkage cavity is about 34.01 µm, which increases by about 3 times. And the maximum size of pore is about 119.54 mm. Therefore, we speculate that some shrinkage cavity could merge and grow up during DHT, which leads to crack initiation worsening mechanical properties of DHT-Ti44Al alloy. Meanwhile, OM also shows that growth of β grain is impeded seriously due to the existence of porosity. Some directional grains are discontinuity. It can be concluded that the preparation of as-cast TiAl based alloys is crucial to completely present mechanical properties of DHT-TiAl based alloys. It offers a new idea to prepare as-cast TiAl based alloy by more excellent solidification technique before DHT. 5. Conclusion In this study, comparing microstructure and mechanical properties of as-cast Ti44Al alloy with that of DHT-Ti44Al alloy, conclusions are as follows, (1) The columnar grains are obtained in Ti44Al alloy after DHT for 4 times with parameters of 21.6 kW loaded power and 1750 K. The microstructure is consisted of fully lamellar colonies. Spacing of lamellar phase is smaller about 0.6 µm and more uniform due to sufficient diffusion of atoms during DHT. And the width of grain boundary reduces and the morphology is zigzag. (2) The plasticity of Ti44Al alloy is improved after DHT. The tensile strength and 14

elongation are 217 MPa and 1.78%, respectively, at room temperature. Compared with that of as-cast Ti44Al alloy, elongation increases about 9.2%. The deterioration of strength is mainly resulted from the micropores that reserved from as-cast alloy. The improvement of ductility is caused by decreasing lamellar spacing and changing orientation of lamellar colonies. (3) The cleavage fracture is main way in as-cast Ti44Al alloy but a mixed fracture consisted of some cleavage planes and many tearing edges is in Ti44Al alloys after DHT. And height of cleavage step decreases greatly for the spacing of lamellar phase decreases after DHT. Acknowledgement This work was supported by National Key Research and Development Program of China (2017YFA0403800) and National Natural Science Foundation of China (No. 51825401). References: [1] Y. Liu, R. Hu, J. Yang, J. Li, Tensile properties and fracture behavior of in-situ synthesized Ti2AlN/Ti48Al2Cr2Nb composites at room and elevated temperatures, Materials Science and Engineering: A, 679 (2017) 7-13. [2] S.W. Kim, J.K. Hong, Y.S. Na, J.T. Yeom, S.E. Kim, Development of TiAl alloys with excellent mechanical properties and oxidation resistance, Materials & Design, 54 (2014) 814-819. [3] H. Clemens, S. Mayer, Design, Processing, Microstructure, Properties, and Applications of Advanced Intermetallic TiAl Alloys, Advanced Engineering Materials, 15 (2013) 191-215. [4] K. Kothari, R. Radhakrishnan, N.M. Wereley, Advances in gamma titanium aluminides and their manufacturing techniques, Progress in Aerospace Sciences, 55 (2012) 1-16. [5] H. Clemens, W. Smarsly, Light-Weight Intermetallic Titanium Aluminides-Status of Research and Development, Advanced Materials Research, 278 (2011) 551-556. [6] X. Song, H. Cui, Y. Han, N. Hou, N. Wei, L. Ding, Q. Song, Effect of carbon reactant on microstructures and mechanical properties of TiAl/Ti2AlC composites, Materials Science and Engineering: A, 684 (2017) 406-412. [7] C.Y. Nam, D.M. Wee, P. Wang, K.S. Kumar, Microstructure and toughness of nitrogen-doped TiAl alloys, Intermetallics, 10 (2002) 113-127. [8] X. Wu, Review of alloy and process development of TiAl alloys, Intermetallics, 14 (2006) 1114-1122. [9] W. Luo, J. Shen, Z. Min, H. Fu, Lamellar phase orientation control of TiAl alloys under high temperature gradient with a Ti-43Al-3Si seed, Journal of Crystal Growth, 310 (2008) 5441-5446. [10] J. Fan, X. Li, Y. Su, R. Chen, J. Guo, H. Fu, Microstructure evolution of directionally solidified Ti-46Al-0.5W-0.5Si alloy, Journal of Crystal Growth, 337 (2011) 52-59. [11] X.F. Ding, J.P. Lin, L.Q. Zhang, Y.Q. Su, G.L. Chen, Microstructural control of TiAl-Nb alloys by directional solidification, Acta Materialia, 60 (2012) 498-506. [12] P. Erdely, P. Staron, E. Maawad, N. Schell, J. Klose, S. Mayer, H. Clemens, Effect of hot rolling and 15

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Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: