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Effects of Sub-Eutectoid Heat Treatment on Microstructure of Ti-Al-Cr Intermetallic Alloys

H. Ghasemi, H. Jafarian, S. Heshmati-Manesh* and M. Nili-Ahmadabadi

Department of Metallurgy and Materials, Faculty of Engineering, University of Tehran, Tehran, Iran

*Author for Correspondence, E-mail: sheshmat@ut.ac.ir

ABSTRACT In this paper, heat treatment of intermetallic alloys based on γ phase with the aim of reducing interdenritic segregation has been investigated. The results indicated that heat treatment of Ti-Al-Cr alloys just below the eutectoid transformation line may induce nucleation of gamma grains and discontinuous coarsening of α2 and γ layers. The growth of these layers via grain boundary diffusion mechanism, results in a compositional homogeneity of the alloys. Also, X-ray diffraction results on the

homogenized alloys revealed the presence of small amount of B2 phase in the microstructures.

1. Introduction Intermetallic alloys based on γ-TiAl are potential candidates for substitution of nickel-based super alloys because of their low density, strength retention, and creep and oxidation resistance at high temperatures. During the last two decades a great deal of effort has been devoted for development of

these alloys [1]. Gamma alloys contain 46-52 %at. Aluminum, together with small additions of Nb, W, Mn, V, Cr and Ta elements [2]. Figure 1 shows the partial Ti-Al phase diagram at high temperature range, which includes γ, α2, α and β phases. γ-TiAl is an ordered phase with fct crystal structure. The α phase has hcp crystal structure and α2-Ti3Al is simply the ordered version of the α phase. The β phase is also a disordered phase with bcc crystal structure. According to figure 1 if an alloy with C0 composition is cooled slowly from T1 temperature a lamellar α2+γ structure would be obtained. If such an alloy is heat treated at T2, T3 or T4 temperatures, upon slow cooling a duplex structure composed of equiaxed γ and α2+γ grains would result. In the lamellar structure the γ and α2 phases have a coherent interface with (0001)║{111} orientation relationship[1]. This microstructure due to its low interface energy has a good thermal stability. It has been reported that the best mechanical properties such as ductility, strength and creep resistance is obtained in a fine lamellar structure with small interlamellar spacing[3]. Coarsening phenomena in these alloys means increment of interlamellar spacing between α2 and γ layers. Growth of the layers occurs by continuous and/or discontinuous coarsening mechanisms. In continuous coarsening mechanism the layers grow due to Gibbs-Thompson effect in a way to decrease the interface area between phases, which in turn results a reduction in surface energy. On the other hand, discontinuous coarsening is a process during which new nuclei are formed on grain boundaries or other crystal imperfections and grow further by grain boundary immigration of atoms. Reduction of surface energy and chemical free energy are the driving forces for such a process[4,5]. Discontinuous coarsening has already been reported in Ti-Al alloys with 40-47 %at. Al, and its effects on grain refinement and ductility, which is an important issue in these alloys, has been discussed[6]. Homogenizing heat treatment on cast alloys reduces the interdendritic segregations and results in a better ductility and hot workability of intermetallic alloys. During the homogenizing process, coarsening of the layers occurs which results a better homogeneity of the alloy. In this paper the mutual

effects of discontinuous coarsening and homogenization has been investigated.

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2. Materials and Methods Samples of Ti-47.7Al-2Cr alloys were prepared from high purity elements using a vacuum arc remelting furnace with a non consumable tungsten electrode. Melting was repeated (at least 5 times) in order to ensure chemical homogeneity of the alloys. Weight losses during melting were less than 0.2 %wt., indicating that the actual compositions were close to the nominal ones. Samples were cut from the ingots for microstructural analysis in the as-cast condition. Samples for heat treatment were wrapped in molybdenum foil and encapsulated in argon filled quartz tubes. Course of heat treatment for homogenization is shown in figure 2. The microstructure and phase relationships were characterized by metallography, X-ray diffraction and electron microscopy. Phase compositions were determined by count analysis using an EDS system on an SEM.

3. Results and discussion

Figure 3 shows the as cast structure of the alloy which consists of cellular grains with fully lamellar structure. Homogenization heat treatment on this alloy below the eutectoid transformation line, results in a duplex microstructure, Figure 4. Figure 5 indicates XRD pattern of the homogenized sample. In this figure α2 and γ phases as well as traces of B2 phase have been detected. The following transformations take place during solidification in equilibrium condition: L→L+β→α→α+γ→α2+γ (1)

Figure 1. Partial Ti-Al phase diagram. Figure 2. Course of homogenization heat treatment.

(hours) Time

FC

1150 0C1100 0C

Te

hours72 γ+2α

γ+α

Tα

Tem

pret

ure

α

Figure 3. The as cast microstructure.

Mµ200 mµ40

Figure 4. Microstructure of duplex

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According to the above relation, the final equilibrium microstructure should consist of the α2 and γ phases. Presence of B2 phase as detected by XRD in the homogenized structure is an indication of a relatively high solidification rate. In fact some Β2 phase has remained in the microstructure during solidification in L+β region, which transforms to the ordered B2 at lower temperatures. It should be mentioned that Cr has been reported to be a B2 phase stabilizer in Ti-Al alloys [7, 8]. The B2 phase manifests itself by the presence of the strongest {011} reflection at 2θ = 39.82 position. Using x-ray diffraction results, the lattice parameter of the B2 phase was found to be a=0.3186 nm, which correlates well with literature [8-10]. Annealing heat treatment in the α phase region may results in the dissolution of this phase. XRD results of the sample heat treated above the α-transus line, confirms the dissolution of the B2 phase after annealing heat treatment, Figure 6.

The lattice parameters of the γ phase calculated from x-ray diffraction results are presented in table 1. It should be noted that homogenization has affected both “a” and tetragonality. Table 1 shows that tetragonality of the γ phase in the as cast sample is less than the homogenized sample. This could be due to the difference in the chemical compositions of the samples. Comparing with other alloying elements, addition of Cr has been reported to decreases tetragonality of γ phase. As an example, the c/a ratio in the as cast sample is much smaller than that reported for a typical Mn containing alloy. [8]. The microstructure of the homogenized sample indicated that sub-eutectoid heat treatment causes grain refinement. Figure 7 shows the microstructures of this sample.

Figure 5. XRD pattern of the homogenized sample.

(020

) α2

(021

) α2

(011

) B2

(1

11) γ

(002

) γ

(020

) γ

2θ

α2 (

023)

γ (2

22)

α2 (

202)

γ (2

21)

γ (1

13)

B2 (

112)

0

5

10

15

20

25

30

35

40

45

50

55

60

65

70

75

80

66 67 68 69 70 71 72 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90

intensity

θ2

α2 {0

22}

γ {22

2}γ{

131}

γ{11

3}

γ {02

2}

γ{22

1}

γ{02

1}

γ{12

1}

α2 {0

23}

2θ

Inte

nsity

Figure 6. XRD of the sample heat treated at 1400°C for 15 min and then quenched in oil.

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(a) (b)

Figure 7. The microstructures of the homogenized sample.

As seen in figure 7a, because of a planar interface ”S”, a new coarse lamellar grain has probably nucleated in the boundary of γ and lamellar grains and has grown within the parent grains by grain boundary migration mechanism. As is seen, layers of the previous grain have not completely converted to the coarse layers of the secondary grain. As the boundary of two grains is clearly visible in all directions and direction of layers is not the same within these two grains, it may be concluded that the

new grain has nucleated independently. Figure 7b shows layers of a new secondary grain and in this case it has completely grown inside the parent grain through discontinuous coarsening mechanism. Figure 8 is a secondary electron image (SEI) of this sample, again indicating a growing coarse lamellar grain, which seems that its nucleation has occurred within a lamellar grain. EDS analysis from different parts of this sample revealed that homogenization has been indeed achieved in this alloy. In fact, grain boundary migration and rapid diffusion of elements via grain boundary diffusion, which occurs during discontinuous

ParameterSample Lattice Parameter “a” (nm) Tetragonality c/a

As cast 0.4012 1.012 Homogenized 0.3997 1.020

S

γ

mµ10

Transformation front

γ DC S F

mµ10

γ

α

FL

DC

Figure 8. SEI of the homogenized sample.

Table1. Lattice parameter “a” and tetragonality “c/a” for the γ phase in different samples.

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coarsening, might have resulted in a better homogeneity.

4. Conclusions 1. In samples prepared by vacuum arc remelting furnace, the B2 phase is retained in the final

microstructure because of the rapid cooling rate during solidification and presence of Cr as a B2 phase stabilizer.

2. The B2 phase dissolves in the matrix upon annealing above the Tα temperature. 3. Annealing in the (α2 +γ) two-phase region has little effect on the as-cast fully lamellar morphology.

However, it results in a considerable continuous and discontinuous coarsening. 4. Discontinuous coarsening during heat treatment of the alloy at 1373˚K results in a relatively

homogeneous composition throughout the alloy texture. 5. References

[1] Y-W Kim, JOM, July 1989, Pages 24-30. [2] Y-W Kim and D.M. Dimiduk, JOM, August 1991, pages 40-47. [3] Guoxin Cao, et al., Intermetallics, 2000, Vol. 8, pages 647-653.

[4] S. Mitao and L.A. Bendersky, Acta mater., 1997, Vol. 45, No. 11, pages 4475-4489. [5] A. Denquin and S. Naka, Acta mater., 1996, Vol. 44, No. 1, pages 353-3658.

[6] Jie Yang, J.N. Wang, Yong Wong, Qiangfei Xia, Intermetallics, 2003, Vol. 11, pages 971-974. [7] D.M Dimiduk, P.L. Martin, Y-W. Kim, Mat. Sci. & Eng. A, 1998, Vol. 243, Page 66.

[8]. B.J. Inkson, H. Clemens and J. Marien, Scripta Material, 1998, Vol. 38(9), Page 1377. [9]. T.J. Jewett, B. Ahrens and M. Dahms, Mat. Sci. & Eng. A, 1997, Vol. 225, Page 29.

[10]. P.R. Bhowal, H.F. Marrick and D.E. Larsen, Mat. Sci. & Eng. A, 1995, Vol. 192(3), Page 685.

6. Acknowledgement Authors would like to acknowledge University of Tehran for financial support.