Microstructure control and mechanical properties of Ti44Al6Nb1.0Cr2.0V alloy by cold crucible directional solidification

Materials Science & Engineering A 614 (2014) 67–74

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Microstructure control and mechanical properties of Ti44Al6Nb1.0Cr2.0V alloy by cold crucible directional solidification Shulin Dong, Ruirun Chen n, Jingjie Guo n, Hongsheng Ding, Yanqing Su, Hengzhi Fu School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China

art ic l e i nf o

a b s t r a c t

Article history: Received 31 March 2014 Received in revised form 2 July 2014 Accepted 3 July 2014 Available online 10 July 2014

The as-cast Ti44Al6Nb1.0Cr2.0V alloy master ingot was prepared by vacuum consumable melting technology. Some bars were cut from this ingot and they were directionally solidified by cold crucible under different pulling velocities. The samples could be well directionally solidified when the power (P) was 45 kW and the pulling velocity (V) was 8.33 μm/s or 11.67 μm/s. The results show that the interlamellar space of the directionally solidified samples decreases from the average 1650 nm of ascast to less than 565 nm and is more homogeneous. The microcrack in the master ingot can be eliminated completely and the room temperature (RT) tensile property is also improved after cold crucible directional solidification (CCDS). The ultimate tensile strength (UTS) is 602.5 MPa and the elongation is 1.20% as P ¼45 kW and V ¼11.67 μm/s, compared with as-cast 499 MPa of UTS and 0.53% of elongation. Trans-granular and trans-lamella fractures are predominant modes. The relationship between CCDS interlamellar space (d) and the pulling velocity can be described as d ¼ 1783:2V  0:554 and r 21 ¼ 0:972, where r 21 is the corresponding regression coefficient. The CCDS interlamellar space and  0:145 and nanoindentation hardness (HN) in the lamella region can be described as H N ¼ 17:95d r 22 ¼ 0:986, and they are changed as H 0N ¼ 14:03d0  0:104 and r 23 ¼ 0:975 when the cast condition is considered. The nanoindentation hardness of the B2 phase and the block γ phase are about 8.89 GPa and 8.15 GPa, respectively; both of them keep almost the same in different conditions. & 2014 Elsevier B.V. All rights reserved.

Keywords: Titanium aluminum Cold crucible Directional solidification Tensile property Nanoindentation hardness Microstructure

1. Introduction TiAl-based alloys have attracted more and more attention in the last few decades owing to their excellent properties, such as low density, high specific strength, excellent high-temperature properties, high Young's modulus and so on [1–3]. It is recognized that TiAl-based alloy is one of the most promising materials to replace Ni-based high-temperature alloys in the aviation and aerospace field. However, the preparation and the processing of this alloy are always difficult due to their brittleness at room temperature (RT), high melting point and high activity, especially for the high-Nb-contained TiAl-based alloys [4–7]. The cold crucible directional solidification (CCDS) method is one way of solving these problems. It brings in no or less contamination and is efficient to prepare the DS samples with the industry size. The Ti43Al1.5Si1.5Mo (at%, similarly hereinafter except the special illustration) alloy has been prepared by CCDS and it shows an

n

Corresponding authors. Tel./fax: þ86 451 86412394. E-mail addresses: [email protected] (S. Dong), [email protected] (R. Chen), [email protected] (J. Guo), [email protected] (H. Ding), [email protected] (Y. Su), [email protected] (H. Fu). http://dx.doi.org/10.1016/j.msea.2014.07.003 0921-5093/& 2014 Elsevier B.V. All rights reserved.

improved strength by nearly 30% and the enhanced elongation by almost 200% [8]. The Ti46Al0.5W0.5Si alloy has been studied by Nie [9] and the strength and the elongation after CCDS are about 500 MPa and 2%. Yang [10] studied Ti46Al6Nb-(B) alloy and the strength at RT and high temperature (800 1C) was found to be 473 MPa and 611 MPa, respectively, after CCDS. These results show that TiAl-based alloys after CCDS can exhibit more excellent mechanical properties and can be used as an engineering component, especially suited for producing engine blade [11,12]. Not only TiAl-based alloys, but also many other kinds of materials including titanium alloys, NbSi alloys and solar-grade polycrystalline silicon have been prepared by the CCDS method successfully [13–15]. A number of researches on the cold crucible design, temperature field, flow field and CCDS technology have also been carried out in our research group [16–18]. However, there are still few detailed reports on the microstructure and properties of TiAl-based alloys by CCDS. In this paper, the CCDS method was used to prepare the Ti44Al6Nb1.0Cr2.0V alloy with the aim of enhancing its RT tensile properties. The microstructure and properties of as-cast master ingot and CCDS samples were compared and analyzed. Meanwhile, the function relationship among pulling velocity, interlamellar space (the thickness of α2 þγ) and nanoindentation hardness was

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established. These results may provide guidance for macro/microstructure control by CCDS and for applicable feasibility of TiAlbased alloys.

2. Experimental procedures The as-cast ingot with a nominal composition of Ti44Al6Nb1.0Cr2.0V was prepared by vacuum consumable melting technology. It was melted two times for homogenization with the size of ∅ 225 mm  320 mm. Then the ingot was cut into many smaller bars (∅ 19 mm) and these bars were directional solidified by cold crucible under different pulling velocities (V) and the power (P) of 45 kW. The experimental detail of CCDS was described in Ref. [19]. The CCDS samples were cut into two halves longitudinally for macro/microstructure observation. The macrostructure of the CCDS samples was photographed using a Nikon D7000 Digital Single Lens Reflex. The optical microscope (OM) was used to observe the microstructure in lower magnification and made the statistics of lamella orientation. The scanning electron microscopy in back-scattered electron (SEM-BSE) mode was used to investigate the microstructure. Transmission electron microscopy (TEM) was used to identify the different phases, measure the interlamellar space and investigate the microstructure in higher magnification. The interlamellar space (one α2 plus one γ) was acquired by the line interception method: drawing a line segment perpendicular to the lamellas firstly, counting the number of the lamellas passing through this line segment secondly, and calculating the average interlamellar space finally. The interlamellar space of each ingot was measured from 8 TEM images and calculated 20 times. The RT tensile testing was carried out on the 5569 Instron testing machine. The gage section of the specimen was 18 mm  6 mm  2 mm and the tensile speed was 0.5 mm/min. For each condition, at least 5 qualified samples were tested. The tensile fracture morphology was observed by SEM-secondary electron (SE). The size of the samples for nanoindentation hardness testing was 8 mm  10 mm  2 mm and the maximum indenter depth was 240 nm. The nanoindentation testing was completed obeying a certain array rule at some region, then the SEM-BSE mode was used to find out what phase every indentation was on. The specimens were polished and etched in the modified Kroll's reagent composed of 20 vol% HF, 2 vol% HNO3 and 78 vol% H2O. TEM foils were wire-electrode cut firstly by about 300 μm, followed by manual grinding to about 80 μm, and finally they were prepared by twin jet electro-polishing in a solution of 58% methanol, 37% butyl alcohol and 5% perchloric acid at 15 V and 30 1C. The microstructure and properties of as-cast master ingot were also studied for the comparison analysis.

3. Results 3.1. Macrostructure of the samples Fig. 1 shows the macrostructure of the CCDS Ti44Al6Nb1. 0Cr2.0V samples under different pulling velocities: (a) 8.33 μm/s; (b) 11.67 μm/s; (c) 16.67 μm/s; and (d) 21.67 μm/s. As shown in Fig. 1(a) and (b), they exhibit the well-DS macrostructure. While, as shown in Fig. 1(c), it exhibits a mixed macrostructure that includes most columnar crystals and some equiaxed crystals under V¼ 16.67 μm/s. As shown in Fig. 1(d), when the pulling velocity increases to 21.67 μm/s, it presents equiaxed crystals completely. So, with the increase in pulling velocity from 8.33 μm/s to 21.67 μm/s, the macrostructure transforms from columnar to equiaxed. For the well-DS macrostructure, taking Fig. 1(a) as an

example, it can be divided into three parts from the top down: they are liquid region (or final solidification region), stable DS region and initial solidification region. The initial solidification region is located in the bottom. Owing to the competitive growth of initial crystals at the beginning of the experiment, it exhibits a disordered macrostructure in this region. In the stable DS region, the columnar crystals are more homogeneous, which indicates that the heat and mass transfer are relatively equilibrium. The liquid region is located in the top and it is formed because the power is turned off instantaneously in the experiment. The solid/ liquid (S/L) interface is marked as “AA” in Fig. 1. 3.2. Microstructure of the samples Because of the instantaneous power cut in the experiment, the liquid in front of the S/L interface was solidified fast and the initial crystal growth morphology in the stable DS region was reserved. Fig. 2(a)–(d) exhibits the microstructure of the liquid regions bordering on the S/L interface, respectively, corresponding to Fig. 1(a)–(d). Fig. 2(a) and (b) exhibits the typical β-solidification characteristic, since the second dendritic arms (SDAS) are perpendicular to the primary dendritic arms (PDAS). While, as shown in Fig. 2(c) and (d), there are some crystals exhibiting α-solidification characteristic, since the SDAS presents a 601 angle with the PDAS, which are marked by the dotted line ellipse. For the Fig. 1(a)–(c) samples, the parts where the height is 45– 46 mm from the bottom (they are in the stable DS regions) are chosen to evaluate the lamella orientations by two-dimensional observation with OM. The statistics results are shown in Fig. 3, in which the angle means the lamella orientation deviates from the pulling direction. It can be found that the angle becomes bigger gradually with the increase in the pulling velocity. When the pulling velocity increases from 8.33 μm/s to 11.67 μm/s, this variation is not obvious. But when the velocity increases to 16.67 μm/s, the angle gets bigger obviously; meanwhile, it is noted that the percentage of 0–301 angle decreases and that of the 60–901 angle increases apparently. When the pulling velocity increases to 21.67 μm/s, the sample is mainly composed of equiaxed crystals and it is meaningless to evaluate the lamella orientation. Fig. 4 shows the microstructure of the as-cast ingot and the Fig. 1(b) sample that has the typical CCDS microstructure, Fig. 4 (a) is the as-cast condition, Fig. 4(b) is from the liquid region and Fig. 4(c) and (d) are from the stable DS region. As shown in this figure, for as-cast and CCDS samples, this kind of microstructure is defined as modified near lamellar (M-NL) structure [20], which is composed of mainly (α2 þγ) lamellae and some (B2 þγ) blocky morphology. It is different from the traditional NL structure that is composed of mainly (α2 þ γ) lamellae and some equiaxed γ grains. There are more B2 phases in this M-NL structure, which result from the addition of more β-stabilizers and relatively lower Al content of this alloy. It can be observed that there are some microcracks in the as-cast ingot, which mainly distributed in the blocky γ regions, as shown in Fig. 4(a). After CCDS, these microcracks can be eliminated completely. As shown in Fig. 4(b), the B2 phase content in the liquid region is less than that in the stable DS region. This phenomenon is caused by the higher Al content in the liquid region. The higher the Al content is, the lesser the B2 phase forms. In Fig. 4(c), some white segregation stripes distribute along the colony boundaries and present a large angle to the pulling direction, which is the so-called C-type β-segregation [20]. SEMEDS was employed to investigate the composition in different parts (marked as A–D in Fig. 4) and the results are listed in Table 1, from which it can be confirmed that there are more β-stabilizers in the white contrast B2 phase region and more Al in the liquid region and blocky γ phase region. Furthermore, it also contains more βstabilizers in the blocky γ phase due to the fact that it comes from

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Pulling direction

Growth direction

liquid region

A

A

stable DS region

initial solidification region

Pulling direction

Growth direction

Fig. 1. Macrostructure of CCDS samples under the power of 45 kW and different pulling velocities: (a) 8.33 μm/s; (b) 11.67 μm/s; (c) 16.67 μm/s; and (d) 21.67 μm/s.

Fig. 2. Microstructure observed by OM in the liquid regions of the samples corresponding to Fig. 1(a) and (b), respectively.

the B2 phase nearby. Fig. 5 shows the microstructure of Fig. 1 (b) samples in the stable DS region by TEM and the corresponding selected area diffraction patterns (SADPs). Fig. 5(a) shows B2

phase, γ phase and (α2 þγ) lamella region. The (α2 þγ) lamella structure near the blocky B2 phase is not as regular as that in other parts, which is because the B2 phase can heighten the α transus

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temperature in its neighboring area and then increase the diffusivity, consequently enhancing the mobility of the interfacial dislocations and ledge, so the coarser and irregular lamellar structure is formed [21,22]. These irregular (α2 þγ) lamella regions will not be counted in the related statistics including interlamellar space statistics and the nanoindentation hardness statistics. Fig. 5(b) and (c) are the images of the lamella region and the corresponding SADPs, in which the γ/γT structure can be found. Fig. 5(d) exhibits the lamella region by HREM, in which the α2 phase and the γ phase are distinguished by integral power region fast Fourier transform (FFT). 3.3. The interlamellar space and the mechanical properties The interlamellar space, nanoindentation hardness and tensile properties of all samples are listed in Table 2. Because the as-cast ingot is large, the region of 1/2 radius and 1/2 height is chosen for the study. For the CCDS samples, the regions for measuring the intelamellar space and testing the nanoindentation hardness are still about 45 mm height. After CCDS, the interlamellar space gets much smaller and more homogeneous. With the increase in

pulling velocity, the average interlamellar space decreases gradually. The nanoindentation hardness in the lamella region is mainly affected by the interlamellar space, and the nanoindentation hardness increases gradually with the decreasing of interlamellar space. The nanoindentation hardness of the B2 phase and block γ phase almost keep the same under different conditions. The CCDS method can improve the tensile property effectively. It exhibits better tensile properties under the condition that P ¼45 kW and V¼8.33 μm/s or V¼11.67 μm/s. Fig. 6 shows the SEM-BSE images of the nanoindentations in the lamella region of the Fig. 1(b) sample. This nanoindentation hardness tester can exactly test a certain region without the influence of other phases.

3.4. Fracture morphology Fig. 7 shows the fracture morphology and the profile of the tensile samples at RT. Fig. 7(a) is the as-cast condition and Fig. 7 (b)–(d) are the CCDS condition from Fig. 1(b) sample. In Fig. 7(a), the as-cast contraction cavities can be seen on the fracture surface, which will worsen the tensile property. In Fig. 7(b) and (c), it can be found that, after CCDS, the tensile fracture mainly exhibits the trans-granular trans-lamella fracture mode with the fracture passing through the colonies and tearing the lamellas. In Fig. 7 (b), there is a ladder shaped fracture morphology marked by the dotted line frame, which may be induced by the local transgranular interlamellar fracture. In this figure, the regions marked by the arrows may derive from the B2 phase being pulled out of the matrix during tensile testing. In Fig. 7(c), the regions marked Table 1 SEM-EDS results in different parts of the CCDS sample.

Fig. 3. Statistics results of lamella orientation in the chosen regions.

Sign

Regions

Ti

Al

Nb

Cr

V

A B C D

Liquid region-lamellae Stable DS region-lamellae Stable DS region-blocky γ Stable DS region-B2

42.02 46.22 42.34 49.21

50.48 46.57 48.44 38.51

5.21 5.05 6.18 6.81

0.77 0.82 1.20 2.33

1.51 1.34 1.84 3.14

A

microcrack

10μm

50μm

Growth direction

C C-type β-segregation

B

D B2+γ

50μm

10μm

Fig. 4. Microstructure of as-cast ingot and Fig. 1(b) sample: (a) as-cast condition; (b) liquid region; (c) stable region; and (d) the same region with (c).

S. Dong et al. / Materials Science & Engineering A 614 (2014) 67–74

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α2 +γ

B2

γ

0.232nm

0.465nm

Fig. 5. Microstructure of the Fig. 1(b) sample observed by TEM and the corresponding SADPs: (a) (B2 þγ) blocky morphology region; (b) lamellar structure region; (c) SADPs with B ¼ ½101  γ==½112  0α2 corresponding to the picture (b); and (d) HREM image and the corresponding FFT results. Table 2 Interlamellar space and nanoindentation hardness and tensile property of all samples. Pulling velocity (μm/s) Interlamellar space (nm) Nanoindentation hardness (GPa) (α2 þ γ) B2 Blocky γ Tensile property Ultimate tensile strength (UTS/MPa) Elongation (%)

As-cast

8.33

11.67

16.67

21.67

1650 7 1200

565 7 257

4377 288

381 7212

326 7 217

6.517 0.71 8.85 7 0.60 8.117 0.13

7.157 0.31 8.917 0.51 8.137 0.17

7.38 7 0.26 8.93 7 0.55 8.107 0.12

7.58 70.23 8.86 70.48 8.20 70.20

7.737 0.21 8.90 7 0.52 8.137 0.12

498.87 98 0.53 7 0.22

589.4 7 28 1.167 0.12

592.1 747 0.91 70.17

570.27 39 0.88 7 0.15

602.5 7 35 1.20 7 0.1

Blocky γ

B2 5μm

5μm Fig. 6. SEM-BSE images with nanoindentations on the lamellar region.

by the arrows exhibit the fishbone shape morphology that usually results from the cleavage fracture, in which the mid-ridge part is formed by the cleavage of {100}[100]γ and the two sides are

formed by the cleavage of {100}[110]γ and {112}[110]γ. In Fig. 7(d), there are some microcracks and small holes on the profile after the tensile test. They mainly distribute in the blocky γ regions and also

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Contraction cavities 200 μm

200 μm

Blocky γ

20 μm

50 μm

Fig. 7. Fracture morphology and the profile of the as-cast and the CCDS samples: (a) as-cast condition; and (b)–(d) CCDS condition in which panel (d) is the profile.

exist in the lamella region. These holes originate from the initial microcrack and they present narrow strip shapes paralleling each other.

4. Discussions 4.1. Effect of pulling velocity on the macrostructure, primary phase and lamella orientation With the increase in pulling velocity, the macrostructure of the CCDS samples changes from the columnar crystal to equiaxed crystal. It can be explained in general as follows: ΔT n ¼ T L ðC 0 Þ  T n

ð1Þ

ΔT t ¼ T L ðC 0 Þ T t

ð2Þ

where T L ðC 0 Þ is the liquidus line temperature of the C 0 content alloy, T n is nucleation temperature of equiaxed crystals and T t is the tip temperature of the columnar crystal. So, the equiaxed and columnar crystal nucleation undercooling degree ΔT n and ΔT t can be expressed as formulas (1) and (2). Formula (3) shows the condition criterion when the crystal transforms from columnar to equiaxed:
 ! ΔT n 3 1=3 G o 0:617N ΔT t 1  ð3Þ ΔT t where N is the nucleation rate and G is the temperature gradient. However, there is a relationship between ΔT t and solidification rate V (equal to pulling velocity or crystal growth rate) as shown in the following formula: ΔT t p V 1=2

ð4Þ

With the increasing of V , both ΔT t and N get bigger, With the increasing of V, ΔTt and N increase, then the right side of Eq. (3) is increased too, When it meets Eq. (3), there will be a

transformation from the columnar to the equiaxed. However, this criterion can only explain this phenomenon in general, because it is established based on many ideal assumed conditions. For instance, the liquid solute concentration should keep C 0 from the beginning to the end of the experiment and the temperature gradient G should present linear in the solidification front for the columnar crystals. Strictly speaking, these assumed conditions are obviously impossible in practice. As shown in Fig. 2, when the pulling velocity is lower, it exhibits the β-solidification characteristic. When the pulling velocity increases to 16.67 μm/s, α-solidification characteristic can be found. It can be explained from the interfacial temperature response function theory. Fig. 8 [23] is the schematic diagram of the interfacial temperature response function with growth rate (equal to pulling velocity). When the growth rate V is lower, namely lg V 1 o lg V o lg V 2 , the β phase will be the primary phase. When the growth rate V increases and meets lg V 2 o lg V, α phase will be the primary phase. This theory is in agreement with the experiment results. In the ideal DS condition, if the β phase is the primary phase, the final lamella orientation will be 01 or 451 with the pulling direction; if the α phase is the primary phase, the final lamella orientation should be 901 with the pulling direction [24,25]. However, during the CCDS experiments, the intense electromagnetic stirring and the lateral heat transfer towards the crucible wall will also influence the final lamella orientation to some extent. Anyhow, β-solidification will increase the percentage of small degree (0–301) lamellas, while the αsolidification will increase the percentage of large degree (60– 901) lamellas. As shown in Fig. 2, when the pulling velocity is lower, it exhibits the β-solidification characteristic. With the increase in the pulling velocity, it gradually exhibits an αsolidification characteristic. Therefore, as shown in Fig. 3, with the increasing in pulling velocity, the percentage of the large angle lamellas increases and the percentage of the small angle lamellas decreases gradually.

S. Dong et al. / Materials Science & Engineering A 614 (2014) 67–74

4.3. Effect of the interlamellar space on the nanoindentation hardness

β

β

TL

The relationship of the interlamellar space d and the nanoindentation hardness H N in the CCDS lamella region can be expressed as follows:

α

T(K)

TL

α

α

α

73

β

H N ¼ k2 d

β

b

ð8Þ

where k2 is a proportional constant and b is the interlamellar space exponent, respectively. Regression analysis of experimental data yields the equations as follows: lgV1

H N ¼ 17:95d

lgV2

lgV/(μm/s)

r 22 ¼ 0:986

Fig. 8. Schematic diagram of the interfacial temperature response function with growth rate.

Fig. 9. Linear fitted results between the pulling velocity and the interlamellar space after logarithm conversion.

4.2. Effect of pulling velocity on the interlamellar space When the power is 45 kW, the temperature gradient in front of the S/L interface keeps almost the same in this pulling velocity level [19,23]. With the increase in the pulling velocity V, the interlamellar space d decreases according to the following relationship: d ¼ k1 V a

 0:145

ð9Þ ð10Þ

r 22

is the regression coefficient. Fig. 10 shows the linear where relation between d and H N in the CCDS lamella region after logarithm conversion according to Eq. (9). This result is a little different from that of others using Vickers hardness testing [19,23,26–28], which may result from the different method to test the hardness. It should be noted that although the alloy is mainly composed of a lamellar structure, there are numerous B2 phases with higher hardness and some blocky γ phases. The Vickers hardness indentation is larger [27,28]; when the Vickers hardness tester presses on them (B2 or blocky γ), the test results will certainly suffer error, and hence it cannot test the hardness of the lamella region accurately. Compared with the Vickers hardness test, the nanoindentation hardness method is more appropriate to test the hardness of the lamella because of its smaller tester (the indentation is shown in Fig. 6), especially for obtaining the relationship between the hardness and the interlamellar space. Furthermore, for testing the results of nanoindentation hardness in one region, they will be given automatically and almost keep constant after the tester presses in a certain depth, which is better able to enhance the testing accuracy. The nanoindentation hardness in the lamella region is mainly decided by the interlamellar space. They should still meet Eq. (8) relationship if the as-cast condition is taken into account. Regression analysis of experimental data yields the equations as following when including the as-cast condition: H 0N ¼ 14:03d0  0:104

ð11Þ

r 23 ¼ 0:975

ð12Þ

ð5Þ

Here, k1 is a proportional constant and a is the velocity exponent. Regression analysis of experimental data yields an equation in this study as follows: d ¼ 1783:2V  0:554

ð6Þ

r 21 ¼ 0:972

ð7Þ

where r 21 is the regression coefficient. Fig. 9 shows the linear relation between d and V after logarithm conversion according to Eq. (6). The results indicate that the pulling velocity is the main factor affecting the interlamellar space in this experiment. The interlamellar space decreases with the increase in pulling velocity. These results are similar to the ones by Nie and Fan [19,26].

Fig. 10. Linear fitted results between the interlamellar space and the nanoindentation hardness in the CCDS lamella region after logarithm conversion.

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UTS 499 MPa and elongation 0.53%. The main fracture mode is trans-granular trans-lamella fracture for the CCDS samples under lower pulling velocity. (3) The CCDS interlamellar space is mainly affected by the pulling velocity, they have the relationship of d ¼ 1783:2V  0:554 and r 21 ¼ 0:972, where r 21 is the corresponding regression coefficient. (4) The CCDS nanoindentation hardness in the lamella region and its interlamellar space have the relationship of  0:145 H N ¼ 17:95d and r 22 ¼ 0:986. If the as-cast condition is considered, the relationship is H 0N ¼ 14:03d0  0:104 and r 23 ¼ 0:975, which implies that the nanoindentation hardness in the lamella region is mainly affected by interlamellar space. The nanoindentation hardness of the B2 phase and block γ phase almost keep the same under different conditions.

Fig. 11. Linear fitted results between the interlamellar space and the nanoindentation hardness after logarithm conversion including the as-cast condition. 0

where r 23 is the regression coefficient. H 0N and d are the nanoindentation hardness and the interlamellar space, respectively. The regression coefficient of this fitting is better, which verifies the conclusion that the nanoindentation hardness in the lamella region is mainly decided by the interlamellar space. Fig. 11 shows 0 the linear relation between d and H 0N in the lamella region after logarithm conversion according to Eq. (11). 4.4. Tensile property at room temperature As a kind of brittle material, TiAl-based alloy is more sensitive to microcracks [29]. After CCDS, the as-cast microcrack that mainly distributed in the blocky γ can be eliminated completely. The interlamellar space gets smaller and more homogeneous, which will reduce the distance of dislocation pileup and is helpful for the compatibility of deformation under stress conditions. The columnar crystal boundaries distribute longitudinally too. Furthermore, when the pulling velocity is lower, the percentage of small-angle lamellae is higher. All of them are helpful to improve the tensile property. The tensile test samples are wire-electrode cut down parallel to the columnar crystal direction, and the fracture surface is always perpendicular to the columnar crystal direction and the tensile direction. As a result, the RT tensile fracture mainly exhibits trans-granular trans-lamella fracture mode for the CCDS samples under lower pulling velocity. From the discussion above, it can be seen that the CCDS technique is an efficient method to prepare nocontamination TiAl-based alloys ingot with industry size. The manufacture processing and the final microstructure can be well controlled by proper technological parameters (power and pulling velocity). The well-DS sample possesses better mechanical properties [30–32], especially the high-temperature property; thus, it is suitable for the engine blade in the area of aviation. 5. Conclusions (1) Ti44Al6Nb1.0Cr2.0V alloys are well directional solidified by a cold crucible when the power is 45 kW and the pulling velocity is 8.33 or 11.67 μm/s and they show β-solidification characteristics. When the pulling velocity increases to 16.67 μm/s, they begin to show αsolidification characteristics. After CCDS, the interlamellar space gets smaller and more homogeneous and the microcrack in the cast ingot can be eliminated completely. (2) The RT tensile property is improved after CCDS; it has a UTS 602.5 MPa and elongation 1.20% compared with the as-cast

Acknowledgments This research was supported by National Basic Research Program of China (2011CB605504) and the Program of New Century Excellent Talents in University (NCET-12-0153) and National Natural Science Foundation of China (51274076).

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