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Ti-6Al-4V interface by diffusion bonding
Intermetallics 36 (2013) 127e132
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
An investigation of Ti-43Al-9V/Ti-6Al-4V interface by diffusion bonding X.R. Wang, Y.Q. Yang*, X. Luo, W. Zhang, G.M. Zhao, B. Huang State Key Lab of Solidification Processing, Northwestern Polytechnical University, Friendship Western Road No. 127, Xi’an, Shanxi 710072, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 6 August 2012 Received in revised form 24 December 2012 Accepted 28 December 2012 Available online 16 February 2013
Ti-43Al-9V/Ti-6Al-4V joints were fabricated by vacuum hot pressing under 920 C/45 MPa/2 h. After fabrication the joints were thermally exposed in vacuum at 920 C for 5 h and 17 h, and at 1100 C for 1 h. The interfaces of both the as-prepared and the thermally exposed joints were analyzed by means of optical microscope, scanning electron microscope, electron backscatter diffraction system and energy dispersive spectrometer. The results indicate that sound joints can be achieved under 920 C/45 MPa/2 h, and the interfacial phase sequence of the as-prepared joints was identified as Ti-43Al-9V/g(TiAl)/B2 (the ordered form of b-Ti phase)/a2(Ti3Al)/a(Ti)/Ti-6Al-4V. The thickness of the interfacial zone increased with the duration of thermal exposure at 920 C. Furthermore, the interfacial zone thickened rapidly when the joints were thermally exposed at temperature above the phase transformation point (a-Ti / b-Ti) of Ti-6Al-4V. In this paper, the interface growth way was also discussed based on the locations of the refractory phase YAl2 and microvoids. The analysis of the thermally exposed joints confirmed the correctness of the interface growth way. Moreover, elemental diffusion mechanism was discussed in this paper. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl C. Joining D. Microstructure B. Thermal stability E. Phase stability, prediction
1. Introduction
g-TiAl intermetallics are winning more attention because of their low density, high elastic modulus, better oxidation resistance and good creep properties at elevated temperature [1e3]. Simultaneously, the welding performance of g-TiAl alloy must be studied before their extensive application in aerospace field. For the welding between Ti-6Al-4V (called as Ti64 hereafter) and TiAl alloy, the welding technologies used in previous studies mainly belong to fusion welding [4,5]. However, fusion welding can affect the structure of base materials, and it can lead to voids and hot cracking. Diffusion bonding can effectively avoid these defects and needs not to melt base materials [6e8]. Many studies have been carried out concerning g-TiAl/Ti64 joint by diffusion bonding. Holmquist et al. studied the interfacial microstructure, the distribution of elements and tensile property of Ti64/TiAl joint which was achieved by hot isostatic pressing [1]. According to their results, sound joints can be achieved under 900 C/200 MPa/1 h. The interfacial thickness increased with the increase of temperature, and the interfacial phase sequence is gTiAl/g þ a2/a2 þ a/Ti64 (in which g, a2 and a are TiAl, Ti3Al and a-Ti
* Corresponding author. Tel./fax: þ86 29 88460499. E-mail address: [email protected] (Y.Q. Yang). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.12.018
phase, respectively). However, tensile and fracture properties have no obvious distinction between the joints and g-TiAl IHI alloy. Wang et al. [2] successfully joined Ti-46.5Al-2.5V-2Cr-1.5Nb and Ti64 via hot isostatic pressing under different parameters. They found that sound joints can be obtained at 800e900 C/100 MPa/ 2 h. Their tensile results indicated that the joints fabricated under 880 C/100 MPa/2 h have the highest fracture strength. G. Çam et al. studied the microstructural and mechanical characterization of diffusion bonded hybrid TiAl/Ti64 joints [7]. According to the measured shear strength values of the joints, the optimum pressure and bonding temperature are 850 C and 5 MPa, respectively. They thought that the formed interfacial product Ti3Al is detrimental to the bond strength of the joints, which can be concluded by the results of the shear strength test of the produced joints. Kong et al. also studied the interfacial microstructure and shear strength of Ti64/Ti-43Al-9V laminate composite sheet fabricated by hot packed rolling [8]. Combining the results of energy dispersive spectrometer (EDS), X-ray diffraction (XRD) and transmission electron microscope (TEM) analysis, they thought that the structure of the interfacial region was TiAl/g þ B2/a2/Ti64 (in which B2 is the ordered form of b-Ti). The average shear strength of the produced joints was 335.7 MPa, and one of the reasons for this value is that the strength of the formed a2 is better than that of the g-TiAl intermetallic compound. Several other investigations have also been done by researchers [9e12]. Although the compositions and statuses of the
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Fig. 1. OM photographs of the as-prepared joint with a lower magnification in (a) and a higher magnification in (b).
intermetallic alloys based on TieAl system are different in these studies, similar promising results have been achieved. And the previous studies mainly focused on the processing and mechanical property of the joints. However, the interface growth way has not been addressed, and the element diffusion behavior has not been deeply studied, either. Therefore, the interfacial microstructure of the Ti-43Al-9V/Ti64 joint, the interface growth way and element diffusion behavior will be studied in this work. In this study, Ti-43Al-9V/Ti64 joints were fabricated via vacuum hot pressing under 920 C/45 MPa/2 h. Part of the as-prepared joints was then thermally exposed at 920 C for 5 h (written as 920 C/5 h hereafter), 920 C/17 h and 1100 C/1 h, respectively. Then the interfacial microstructure and element distribution characteristic were analyzed by means of optical microscope (OM), scanning electron microscope (SEM), EDS and electron backscatter diffraction (EBSD). 2. Experimental Ti64 is a typical a þ b duplex phase alloy, and the thickness of Ti64 foil utilized in this study is 0.35 mm. The nominal composition of g-TiAl alloy is Ti-43Al-9V-0.3Y (at%) which is composed of massive g-TiAl phase, a small amount of lamellar a2/g, stripe-like B2, YAl2 and Y2O3 phase. Before diffusion bonding, the two base materials were successively ground on silicon carbide abrasive papers from no. 600 to 1500. In order to remove the superficial oil and oxidation contamination, they were then cleaned with acetone and etchant (10% HF; 5% HNO3; 85% H2O in volume fraction) sequentially. After mounted in a mould, the materials were heated up to 920 C from
room temperature through 3 h. The pressure, 45 MPa, was kept 2 h at 920 C. The cooling process took 5 h to room temperature. During the hot pressing process, the vacuum pressure in the furnace was kept below 3.5 102 Pa. By vacuum seal welding machine, part of the as-prepared joints was sealed in the stainless steel tubes which supply a vacuum atmosphere (<5 103 Pa). Then these sealed joints were thermally exposed under 920 C/5 h, 920 C/17 h and 1100 C/1 h in a box type resistance furnace, respectively. Both the as-prepared and the thermally exposed joints were cut perpendicular to the bonding interface with a low speed diamond saw. Then metallographic specimens were prepared by conventional preparation methods of metallographic sample. After that the interfacial microstructure and element distribution characteristic were analyzed by an optical microscope and a SUPRA 55 field emission scanning electron microscope equipped with an Oxford INCA energy-dispersive spectrometer and HKL electron backscatter diffraction system. 3. Results and discussion 3.1. Microstructural analysis of the interface for the as-prepared joint Fig. 1 shows the OM micrographs of the as-prepared joint. It is seen that the interfacial zone is uniform with no obvious defects, such as voids or gaps. Typical equiaxial structure appears in Ti64 alloy, which can be attributed to the distortion happened in the duplex phase field. Obvious stratification phenomenon can also be found at the interfacial zone.
Fig. 2. (a) SEI of the interfacial zone of the as-prepared joint, (b) The corresponding line distribution of elements.
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Fig. 3. (a) EBSD analysis field of the as-prepared joint interface, (b) the front backscatter electron micrograph of the EBSD analysis field in (a), (c) the schema of phase distribution in the EBSD analysis field.
Fig. 4. SEI of the as-prepared joint with (a) showing the location of EDS point analysis at the interfacial zone; and (b) showing the location of the YAl2 phase.
The stratification phenomenon and element distribution characteristic in the interfacial zone were further studied by SEM and EDS analysis. Fig. 2 shows the secondary electron imaging (SEI) of the interface and the corresponding line analysis curves of elements. It can be seen that the interfacial zone can be divided into two regions as marked in Fig. 2a. The region close to Ti-43Al-9V is marked as region 1, while the region close to Ti64 is marked as region 2. The thickness of the interfacial zone, region 1 and 2 are about 11.2 mm, 6.7 mm, 4.5 mm, respectively. Obviously, it can also be seen that the grains in the interfacial zone are refined. The condition of 920 C and 45 MPa is adequate to arouse dynamic recovery and recrystallization for these two alloys, which can then refine the interfacial grains [13].
Table 1 The EDS analysis result of the interfacial grey white phase in Fig. 4a. Element
O
Ti
Al
V
Y
Content (at%)
9.32
8.13
50.85
0.52
31.18
Fig. 5. SEI showing the discontinuous microvoids in the interfacial zone of the asprepared joint.
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X.R. Wang et al. / Intermetallics 36 (2013) 127e132
3.2. Interface growth way
Fig. 6. The established schema of the interface growth way (DB is short for diffusion bonding).
There is a gradient variation of Al, Ti and V at the interfacial zone as shown in Fig. 2b, which indicates that evident atomic diffusion happened. Fig. 3a shows the SEI of the as-prepared joint interface and the EBSD analysis field, and Fig. 3b is the front backscatter electron micrograph of the analysis field. From Fig. 3b, the morphology of grains can be seen clearly. Fig. 3c is the distribution of phases in the analyzed field which was partly drawn by hand according to the synthesized analysis of EDS and EBSD analysis results. From Fig. 3c, it can be seen that the interfacial phase sequence is Ti-43Al-9V/g/B2/a2/a/Ti64. The position of B2 phase can be confirmed by the peak of V content curve in Fig. 2b. The ratios of Ti and Al can also illustrate the appearance of a2 and a. This result is consistent with the previous TEM analysis in reference [8]. Therefore, there is the mixture phase of B2 and a2 in region 1 while only a phase in region 2.
Rare-earth element Y was added into the Ti-43Al-9V alloy to refine grains and structure. Element Y will pile up along grain boundary to form the refractory phase YAl2 during the casting process of the alloy. Therefore, those YAl2 which are exactly located at the surface can be regarded as a tag to observe the interface growth way. Fig. 4 shows the EDS point analysis position of the grey white phase, and the relevant results are listed in Table 1. It indicates that the analyzed phase is YAl2 which was also detected by other researchers using XRD analysis in Ti-43Al-9V alloy [14]. The location of the YAl2 can imply the interface growth way, i.e. the original interface intrudes into both Ti-43Al-9V and Ti64 simultaneously, and the interface growth rate toward Ti-43Al-9V is higher. Discontinuous microvoids can also be observed in very few regions of the interfacial zone, as shown in Fig. 5. Since a few regions of the bonding surface are not very flat, some microvoids appeared after diffusion bonding. However, what is no doubt is that the location of these microvoids is the original interface of the two base materials. The location of these microvoids can also be regarded as a tag to confirm the interface growth way described above. Therefore, a simplistic schema of the interface growth way can be established, as shown in Fig. 6. From the schema, it can be seen that the original interface grows into both Ti-43Al-9V and Ti64 simultaneously, and the interface growth rate toward Ti-43Al-9V is higher. Therefore, the thickness of region 1 is bigger than that of region 2. 3.3. Vacuum thermal exposure In order to investigate the element diffusion characteristic at elevated temperature, part of the as-prepared joints was thermally exposed in vacuum. Fig. 7 shows the SEI of the interfacial zones which were thermally exposed under different conditions in vacuum. The average interfacial thickness of the joints thermally exposed under 920 C/5 h and 920 C/17 h are about 17 mm and 24 mm, respectively, which indicates that the interfacial zone
Fig. 7. SEI showing the interfacial zone of the joints thermally exposed to (a) 920 C/5 h; (b) 920 C/17 h; (c) 1100 C/1 h.
X.R. Wang et al. / Intermetallics 36 (2013) 127e132
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Table 2 EDS point analysis result of the white phase in Fig. 9. Element
O
Y
Content (at%)
58.73
41.27
Fig. 9 shows the interfacial zone of the joint thermally exposed at 920 C for 17 h. The white phase in the interfacial zone was analyzed by EDS, and the results are shown in Table 2. It can be seen that the white phase should be Y2O3 (melting point is 2690 C) which is similar to YAl2. The position of Y2O3 can also be regarded as a tag to confirm the correctness of the aforementioned interface growth way. Furthermore, the microvoids in Fig. 7b can also be potent evidence. 3.4. Element diffusion mechanism analysis Fig. 8. Thickening kinetics of the interface at 920 C.
increased with the duration of time. Moreover, the interfacial thickness can reach 27 mm after being thermally exposed at 1100 C for 1 h. It is known that the element diffusion coefficient becomes higher at elevated temperature. So the elements will diffuse with a higher speed both in Ti-43Al-9V and Ti64 at 1100 C than at low temperature. Moreover, the phase transformation point of Ti64 alloy is between 980 C and 990 C, so only b phase existed in Ti64 at 1100 C [15]. Element diffusion coefficient is higher in b-Ti (bcc) than a-Ti (hcp) due to the lower atomic stacking density [16], which will accelerate the diffusion of Ti and Al in Ti64 at 1100 C. The Ti43Al-9V alloy contains b and g phases when the temperature is 1100 C [17], and the b phase can accelerate the diffusion of the elements in the alloy. From the above analysis it can be easily understood why the interface thickened so quickly at 1100 C for only 1 h. Fig. 8 shows the growth kinetics curve of the interface at 920 C. It is seen that the interfacial thickening obeys parabolic law, i.e. H ¼ k$t1/2 þ b (where k is the rate constant, b and H are the original and final interfacial thickness, respectively). Therefore, the thickening of the interface is controlled by diffusion [18]. The formula fitted based on the experimental results is H ¼ 3.4 t1/2 þ 10.
According to the foregoing analysis, the interface thickening is controlled by diffusion. The diffusion phenomena during the whole fabrication process can be divided into three main stages: (a) the diffusion before hot pressing; (b) the diffusion during vacuum hot pressing; and (c) the diffusion after vacuum hot pressing. The element diffusion mechanism may be different for these stages. The atomic radius of Ti, Al and V are 0.145 nm, 0.143 nm and 0.135 nm, respectively, which implies that vacancy mechanism is the main diffusion way during the whole diffusion bonding process [19]. During the course of heating up to 920 C, the concentration gradient of each type of element is relative high at the original interface, which leads to a strong diffusion driving force because the diffusion in this experiment belongs to downhill diffusion. Therefore, elements should diffuse rapidly in this heating process. However, since the pressure had not been added, diffusion happened mainly among the contact points [20]. According to the radii of Ti, Al and V, they diffused only via vacancies during this heating stage. Vacancy mechanism still existed during the hot pressing. Moreover, dislocations formed and their density increased with the increase of pressure. As channels of diffusion, dislocations can facilitate the diffusion of elements. The distribution of elements tends to homogenize as the diffusion going on, which makes it possible to produce new phases. And then the high grain boundary energy of these metastable grains can further promote the element diffusion [2,20]. Therefore, many effective short circuit diffusion ways exist in this stage. Since the element diffusion coefficient is temperature dependent by Arrhenius function, it will drop quickly with the decrease of temperature [21]. Therefore, both the diffusion rate and quantity became less during the cooling to room temperature. The diffusion may proceed along the grain boundaries during the cooling stage. 4. Conclusion
Fig. 9. SEI showing the interfacial zone of the joint thermally exposed at 920 C for 17 h.
The interface of the Ti-43Al-9V/Ti64 joint was studied by diffusion bonding in this paper. Sound Ti-43Al-9V/Ti64 joints have been achieved under 920 C/45 MPa/2 h vacuum hot pressing. Based on the EDS and EBSD analysis, the interfacial phase sequence is Ti-43Al-9V/g/B2/a2/a/Ti64. The locations of the refractory YAl2 and microvoids in the as-prepared joints indicate that the original interface grows into both base materials simultaneously with a higher thickening speed toward Ti-43Al-9V. The microvoids and refractory Y2O3 phase located at the interface of the thermally exposed joints further confirm the interface growth way. Through vacuum thermal exposure, it is known that the interface growth is controlled by diffusion, which is well illustrated by the fitted
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formula H ¼ 3.4 t1/2 þ 10. Moreover, the elemental diffusion mechanism was complex during the fabrication process. The diffusion of Ti, Al and V mainly relies on vacancies, dislocations and grain boundaries during hot pressing. Acknowledgment Thanks are given to the financial supports of the Natural Science Foundation of China (no. 51071122), the Aviation Science Foundation of China (no. 2009ZF53062), the NPU Foundation for Fundamental Research (NPU-FFR-JC20100210 and JC201110), and the fund of the State Key Laboratory of Solidification Processing in NWPU (SKLSP201107). References [1] Holmquist M, Recina V, Ockborn J, Pettersson B, Zumalde E. Hot isostatic diffusion bonding of titanium alloy Ti-6Al-4V to gamma titanium aluminide IHI alloy 01A. Scripta Mater 1998;39:1101e6. [2] Wang XF, Ma M, Liu XB, Wu XQ, Tan CG, Shi RK, et al. Diffusion bonding of gTiAl alloy to Ti-6Al-4V alloy under hot pressure. Trans Nonferrous Met Soc China 2006;16:1059e63. [3] Blue CA, Lin. RY. Microstructural evolution in joining of TiAl with a liquid Ti alloy. Scr Metall Mater 1995;32:127e32. [4] Acoff VL, Wilkerson S, Arenas M. The effect of rolling direction on the weld structure and hardness of gamma-TiAl sheet material. Mater Sci Eng A 2002; 329:763e7. [5] Shiue RK, Wu SK, Chen YT, Shiue CY. Infrared brazing of Ti50Al50 and Ti-6Al-4V using two Ti-based filler metals. Intermetallics 2008;9:1083e9. [6] Yu WX, Li MQ, Hu YQ. Superplasticity and application of superplastic forming/ diffusion bonding technology. Mater Rev 2009;23:9e10 [in Chinese].
[7] Çam G, Özdemir U, Ventzke V, Koçak M. Microstructural and mechanical characterization of diffusion bonded hybrid joints. J Mater Sci 2008;43: 3491e9. [8] Kong FT, Chen YY, Zhang DL. Interfacial microstructure and shear strength of Ti-6Al-4V laminate composite sheet fabricated by hot packed rolling. Mater Des 2011;32:3167e72. [9] He P, Zhang JH, Feng JC, Qian YY. Diffusion bonding of TiAl to Ti and TC4 alloy. Acta Metall Sin 2000;13:162e7. [10] Glatz Wolfgang, Clemens Helmut. Diffusion bonding of intermetallic Ti-47Al2Cr-0.2Si sheet material and mechanical properties of joints at room temperature and elevated temperatures. Intermetallics 1997;5:415e23. [11] Wang XF, Luo ZC, Yu GS, Lin JG. Application of a diffusion bonding model to superplastic gamma-TiAl alloys. Mater Sci Technol 2006;22:186e92. [12] Yan P, Wallach ER. Diffusion-bonding of TiAl. Intermetallics 1993;1:83e97. [13] Çam G, Clemens H, Gerling R, Kocak M. Diffusion bonding of g-TiAl sheets. Intermetallics 1999;7:1025e31. [14] Chen GQ, He JS, Zhang BG, Feng JC. Investigations of microstructures and properties in electron beam welded joints of TiAl to TC4. China Welding 2007; 3:10e3 [in Chinese]. [15] Wang M, Guo HZ. Study on superplasticity and fine-grained of TC4 alloy. J Plasticity Eng 2008;15:155e8 [in Chinese]. [16] Leyens Christoph, Peters Manfred. Titanium and titanium alloys. KgaA, Weinheim: Wiley-VCH Verlag GmbH & Co; 2003. [17] Takeyama Masao, Kobayashi Satoru. Physical metallurgy for wrought gamma titanium aluminides microstructure control through phase transformations. Intermetallics 2005;13:993e9. [18] Yang YQ, Werner A, Dudek HJ, Kumpfert J. TEM investigations of interfacial processes in SCS-6 SiC/TiB2/Super a2 composites. Composites Part A 1999;30: 1209e14. [19] Mogilevsky P, Werner A, Dudek HJ. Application of diffusion barriers in composite materials. Mater Sci Eng A 1998;242:235e47. [20] Orhan N, Aksoy M, Eroglu M. A new model for diffusion bonding and its application to duplex alloys. Mater Sci Eng A 1999;271:458e68. [21] . Herzig Chr, Przeorski T, Mishin Y. Self-diffusion in g-TiAl: an experimental study and atomistic calculations. Intermetallics 1999;17:389e404.
Contents lists available at SciVerse ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
An investigation of Ti-43Al-9V/Ti-6Al-4V interface by diffusion bonding X.R. Wang, Y.Q. Yang*, X. Luo, W. Zhang, G.M. Zhao, B. Huang State Key Lab of Solidification Processing, Northwestern Polytechnical University, Friendship Western Road No. 127, Xi’an, Shanxi 710072, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 6 August 2012 Received in revised form 24 December 2012 Accepted 28 December 2012 Available online 16 February 2013
Ti-43Al-9V/Ti-6Al-4V joints were fabricated by vacuum hot pressing under 920 C/45 MPa/2 h. After fabrication the joints were thermally exposed in vacuum at 920 C for 5 h and 17 h, and at 1100 C for 1 h. The interfaces of both the as-prepared and the thermally exposed joints were analyzed by means of optical microscope, scanning electron microscope, electron backscatter diffraction system and energy dispersive spectrometer. The results indicate that sound joints can be achieved under 920 C/45 MPa/2 h, and the interfacial phase sequence of the as-prepared joints was identified as Ti-43Al-9V/g(TiAl)/B2 (the ordered form of b-Ti phase)/a2(Ti3Al)/a(Ti)/Ti-6Al-4V. The thickness of the interfacial zone increased with the duration of thermal exposure at 920 C. Furthermore, the interfacial zone thickened rapidly when the joints were thermally exposed at temperature above the phase transformation point (a-Ti / b-Ti) of Ti-6Al-4V. In this paper, the interface growth way was also discussed based on the locations of the refractory phase YAl2 and microvoids. The analysis of the thermally exposed joints confirmed the correctness of the interface growth way. Moreover, elemental diffusion mechanism was discussed in this paper. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: A. Titanium aluminides, based on TiAl C. Joining D. Microstructure B. Thermal stability E. Phase stability, prediction
1. Introduction
g-TiAl intermetallics are winning more attention because of their low density, high elastic modulus, better oxidation resistance and good creep properties at elevated temperature [1e3]. Simultaneously, the welding performance of g-TiAl alloy must be studied before their extensive application in aerospace field. For the welding between Ti-6Al-4V (called as Ti64 hereafter) and TiAl alloy, the welding technologies used in previous studies mainly belong to fusion welding [4,5]. However, fusion welding can affect the structure of base materials, and it can lead to voids and hot cracking. Diffusion bonding can effectively avoid these defects and needs not to melt base materials [6e8]. Many studies have been carried out concerning g-TiAl/Ti64 joint by diffusion bonding. Holmquist et al. studied the interfacial microstructure, the distribution of elements and tensile property of Ti64/TiAl joint which was achieved by hot isostatic pressing [1]. According to their results, sound joints can be achieved under 900 C/200 MPa/1 h. The interfacial thickness increased with the increase of temperature, and the interfacial phase sequence is gTiAl/g þ a2/a2 þ a/Ti64 (in which g, a2 and a are TiAl, Ti3Al and a-Ti
* Corresponding author. Tel./fax: þ86 29 88460499. E-mail address: [email protected] (Y.Q. Yang). 0966-9795/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.intermet.2012.12.018
phase, respectively). However, tensile and fracture properties have no obvious distinction between the joints and g-TiAl IHI alloy. Wang et al. [2] successfully joined Ti-46.5Al-2.5V-2Cr-1.5Nb and Ti64 via hot isostatic pressing under different parameters. They found that sound joints can be obtained at 800e900 C/100 MPa/ 2 h. Their tensile results indicated that the joints fabricated under 880 C/100 MPa/2 h have the highest fracture strength. G. Çam et al. studied the microstructural and mechanical characterization of diffusion bonded hybrid TiAl/Ti64 joints [7]. According to the measured shear strength values of the joints, the optimum pressure and bonding temperature are 850 C and 5 MPa, respectively. They thought that the formed interfacial product Ti3Al is detrimental to the bond strength of the joints, which can be concluded by the results of the shear strength test of the produced joints. Kong et al. also studied the interfacial microstructure and shear strength of Ti64/Ti-43Al-9V laminate composite sheet fabricated by hot packed rolling [8]. Combining the results of energy dispersive spectrometer (EDS), X-ray diffraction (XRD) and transmission electron microscope (TEM) analysis, they thought that the structure of the interfacial region was TiAl/g þ B2/a2/Ti64 (in which B2 is the ordered form of b-Ti). The average shear strength of the produced joints was 335.7 MPa, and one of the reasons for this value is that the strength of the formed a2 is better than that of the g-TiAl intermetallic compound. Several other investigations have also been done by researchers [9e12]. Although the compositions and statuses of the
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X.R. Wang et al. / Intermetallics 36 (2013) 127e132
Fig. 1. OM photographs of the as-prepared joint with a lower magnification in (a) and a higher magnification in (b).
intermetallic alloys based on TieAl system are different in these studies, similar promising results have been achieved. And the previous studies mainly focused on the processing and mechanical property of the joints. However, the interface growth way has not been addressed, and the element diffusion behavior has not been deeply studied, either. Therefore, the interfacial microstructure of the Ti-43Al-9V/Ti64 joint, the interface growth way and element diffusion behavior will be studied in this work. In this study, Ti-43Al-9V/Ti64 joints were fabricated via vacuum hot pressing under 920 C/45 MPa/2 h. Part of the as-prepared joints was then thermally exposed at 920 C for 5 h (written as 920 C/5 h hereafter), 920 C/17 h and 1100 C/1 h, respectively. Then the interfacial microstructure and element distribution characteristic were analyzed by means of optical microscope (OM), scanning electron microscope (SEM), EDS and electron backscatter diffraction (EBSD). 2. Experimental Ti64 is a typical a þ b duplex phase alloy, and the thickness of Ti64 foil utilized in this study is 0.35 mm. The nominal composition of g-TiAl alloy is Ti-43Al-9V-0.3Y (at%) which is composed of massive g-TiAl phase, a small amount of lamellar a2/g, stripe-like B2, YAl2 and Y2O3 phase. Before diffusion bonding, the two base materials were successively ground on silicon carbide abrasive papers from no. 600 to 1500. In order to remove the superficial oil and oxidation contamination, they were then cleaned with acetone and etchant (10% HF; 5% HNO3; 85% H2O in volume fraction) sequentially. After mounted in a mould, the materials were heated up to 920 C from
room temperature through 3 h. The pressure, 45 MPa, was kept 2 h at 920 C. The cooling process took 5 h to room temperature. During the hot pressing process, the vacuum pressure in the furnace was kept below 3.5 102 Pa. By vacuum seal welding machine, part of the as-prepared joints was sealed in the stainless steel tubes which supply a vacuum atmosphere (<5 103 Pa). Then these sealed joints were thermally exposed under 920 C/5 h, 920 C/17 h and 1100 C/1 h in a box type resistance furnace, respectively. Both the as-prepared and the thermally exposed joints were cut perpendicular to the bonding interface with a low speed diamond saw. Then metallographic specimens were prepared by conventional preparation methods of metallographic sample. After that the interfacial microstructure and element distribution characteristic were analyzed by an optical microscope and a SUPRA 55 field emission scanning electron microscope equipped with an Oxford INCA energy-dispersive spectrometer and HKL electron backscatter diffraction system. 3. Results and discussion 3.1. Microstructural analysis of the interface for the as-prepared joint Fig. 1 shows the OM micrographs of the as-prepared joint. It is seen that the interfacial zone is uniform with no obvious defects, such as voids or gaps. Typical equiaxial structure appears in Ti64 alloy, which can be attributed to the distortion happened in the duplex phase field. Obvious stratification phenomenon can also be found at the interfacial zone.
Fig. 2. (a) SEI of the interfacial zone of the as-prepared joint, (b) The corresponding line distribution of elements.
X.R. Wang et al. / Intermetallics 36 (2013) 127e132
129
Fig. 3. (a) EBSD analysis field of the as-prepared joint interface, (b) the front backscatter electron micrograph of the EBSD analysis field in (a), (c) the schema of phase distribution in the EBSD analysis field.
Fig. 4. SEI of the as-prepared joint with (a) showing the location of EDS point analysis at the interfacial zone; and (b) showing the location of the YAl2 phase.
The stratification phenomenon and element distribution characteristic in the interfacial zone were further studied by SEM and EDS analysis. Fig. 2 shows the secondary electron imaging (SEI) of the interface and the corresponding line analysis curves of elements. It can be seen that the interfacial zone can be divided into two regions as marked in Fig. 2a. The region close to Ti-43Al-9V is marked as region 1, while the region close to Ti64 is marked as region 2. The thickness of the interfacial zone, region 1 and 2 are about 11.2 mm, 6.7 mm, 4.5 mm, respectively. Obviously, it can also be seen that the grains in the interfacial zone are refined. The condition of 920 C and 45 MPa is adequate to arouse dynamic recovery and recrystallization for these two alloys, which can then refine the interfacial grains [13].
Table 1 The EDS analysis result of the interfacial grey white phase in Fig. 4a. Element
O
Ti
Al
V
Y
Content (at%)
9.32
8.13
50.85
0.52
31.18
Fig. 5. SEI showing the discontinuous microvoids in the interfacial zone of the asprepared joint.
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3.2. Interface growth way
Fig. 6. The established schema of the interface growth way (DB is short for diffusion bonding).
There is a gradient variation of Al, Ti and V at the interfacial zone as shown in Fig. 2b, which indicates that evident atomic diffusion happened. Fig. 3a shows the SEI of the as-prepared joint interface and the EBSD analysis field, and Fig. 3b is the front backscatter electron micrograph of the analysis field. From Fig. 3b, the morphology of grains can be seen clearly. Fig. 3c is the distribution of phases in the analyzed field which was partly drawn by hand according to the synthesized analysis of EDS and EBSD analysis results. From Fig. 3c, it can be seen that the interfacial phase sequence is Ti-43Al-9V/g/B2/a2/a/Ti64. The position of B2 phase can be confirmed by the peak of V content curve in Fig. 2b. The ratios of Ti and Al can also illustrate the appearance of a2 and a. This result is consistent with the previous TEM analysis in reference [8]. Therefore, there is the mixture phase of B2 and a2 in region 1 while only a phase in region 2.
Rare-earth element Y was added into the Ti-43Al-9V alloy to refine grains and structure. Element Y will pile up along grain boundary to form the refractory phase YAl2 during the casting process of the alloy. Therefore, those YAl2 which are exactly located at the surface can be regarded as a tag to observe the interface growth way. Fig. 4 shows the EDS point analysis position of the grey white phase, and the relevant results are listed in Table 1. It indicates that the analyzed phase is YAl2 which was also detected by other researchers using XRD analysis in Ti-43Al-9V alloy [14]. The location of the YAl2 can imply the interface growth way, i.e. the original interface intrudes into both Ti-43Al-9V and Ti64 simultaneously, and the interface growth rate toward Ti-43Al-9V is higher. Discontinuous microvoids can also be observed in very few regions of the interfacial zone, as shown in Fig. 5. Since a few regions of the bonding surface are not very flat, some microvoids appeared after diffusion bonding. However, what is no doubt is that the location of these microvoids is the original interface of the two base materials. The location of these microvoids can also be regarded as a tag to confirm the interface growth way described above. Therefore, a simplistic schema of the interface growth way can be established, as shown in Fig. 6. From the schema, it can be seen that the original interface grows into both Ti-43Al-9V and Ti64 simultaneously, and the interface growth rate toward Ti-43Al-9V is higher. Therefore, the thickness of region 1 is bigger than that of region 2. 3.3. Vacuum thermal exposure In order to investigate the element diffusion characteristic at elevated temperature, part of the as-prepared joints was thermally exposed in vacuum. Fig. 7 shows the SEI of the interfacial zones which were thermally exposed under different conditions in vacuum. The average interfacial thickness of the joints thermally exposed under 920 C/5 h and 920 C/17 h are about 17 mm and 24 mm, respectively, which indicates that the interfacial zone
Fig. 7. SEI showing the interfacial zone of the joints thermally exposed to (a) 920 C/5 h; (b) 920 C/17 h; (c) 1100 C/1 h.
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Table 2 EDS point analysis result of the white phase in Fig. 9. Element
O
Y
Content (at%)
58.73
41.27
Fig. 9 shows the interfacial zone of the joint thermally exposed at 920 C for 17 h. The white phase in the interfacial zone was analyzed by EDS, and the results are shown in Table 2. It can be seen that the white phase should be Y2O3 (melting point is 2690 C) which is similar to YAl2. The position of Y2O3 can also be regarded as a tag to confirm the correctness of the aforementioned interface growth way. Furthermore, the microvoids in Fig. 7b can also be potent evidence. 3.4. Element diffusion mechanism analysis Fig. 8. Thickening kinetics of the interface at 920 C.
increased with the duration of time. Moreover, the interfacial thickness can reach 27 mm after being thermally exposed at 1100 C for 1 h. It is known that the element diffusion coefficient becomes higher at elevated temperature. So the elements will diffuse with a higher speed both in Ti-43Al-9V and Ti64 at 1100 C than at low temperature. Moreover, the phase transformation point of Ti64 alloy is between 980 C and 990 C, so only b phase existed in Ti64 at 1100 C [15]. Element diffusion coefficient is higher in b-Ti (bcc) than a-Ti (hcp) due to the lower atomic stacking density [16], which will accelerate the diffusion of Ti and Al in Ti64 at 1100 C. The Ti43Al-9V alloy contains b and g phases when the temperature is 1100 C [17], and the b phase can accelerate the diffusion of the elements in the alloy. From the above analysis it can be easily understood why the interface thickened so quickly at 1100 C for only 1 h. Fig. 8 shows the growth kinetics curve of the interface at 920 C. It is seen that the interfacial thickening obeys parabolic law, i.e. H ¼ k$t1/2 þ b (where k is the rate constant, b and H are the original and final interfacial thickness, respectively). Therefore, the thickening of the interface is controlled by diffusion [18]. The formula fitted based on the experimental results is H ¼ 3.4 t1/2 þ 10.
According to the foregoing analysis, the interface thickening is controlled by diffusion. The diffusion phenomena during the whole fabrication process can be divided into three main stages: (a) the diffusion before hot pressing; (b) the diffusion during vacuum hot pressing; and (c) the diffusion after vacuum hot pressing. The element diffusion mechanism may be different for these stages. The atomic radius of Ti, Al and V are 0.145 nm, 0.143 nm and 0.135 nm, respectively, which implies that vacancy mechanism is the main diffusion way during the whole diffusion bonding process [19]. During the course of heating up to 920 C, the concentration gradient of each type of element is relative high at the original interface, which leads to a strong diffusion driving force because the diffusion in this experiment belongs to downhill diffusion. Therefore, elements should diffuse rapidly in this heating process. However, since the pressure had not been added, diffusion happened mainly among the contact points [20]. According to the radii of Ti, Al and V, they diffused only via vacancies during this heating stage. Vacancy mechanism still existed during the hot pressing. Moreover, dislocations formed and their density increased with the increase of pressure. As channels of diffusion, dislocations can facilitate the diffusion of elements. The distribution of elements tends to homogenize as the diffusion going on, which makes it possible to produce new phases. And then the high grain boundary energy of these metastable grains can further promote the element diffusion [2,20]. Therefore, many effective short circuit diffusion ways exist in this stage. Since the element diffusion coefficient is temperature dependent by Arrhenius function, it will drop quickly with the decrease of temperature [21]. Therefore, both the diffusion rate and quantity became less during the cooling to room temperature. The diffusion may proceed along the grain boundaries during the cooling stage. 4. Conclusion
Fig. 9. SEI showing the interfacial zone of the joint thermally exposed at 920 C for 17 h.
The interface of the Ti-43Al-9V/Ti64 joint was studied by diffusion bonding in this paper. Sound Ti-43Al-9V/Ti64 joints have been achieved under 920 C/45 MPa/2 h vacuum hot pressing. Based on the EDS and EBSD analysis, the interfacial phase sequence is Ti-43Al-9V/g/B2/a2/a/Ti64. The locations of the refractory YAl2 and microvoids in the as-prepared joints indicate that the original interface grows into both base materials simultaneously with a higher thickening speed toward Ti-43Al-9V. The microvoids and refractory Y2O3 phase located at the interface of the thermally exposed joints further confirm the interface growth way. Through vacuum thermal exposure, it is known that the interface growth is controlled by diffusion, which is well illustrated by the fitted
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formula H ¼ 3.4 t1/2 þ 10. Moreover, the elemental diffusion mechanism was complex during the fabrication process. The diffusion of Ti, Al and V mainly relies on vacancies, dislocations and grain boundaries during hot pressing. Acknowledgment Thanks are given to the financial supports of the Natural Science Foundation of China (no. 51071122), the Aviation Science Foundation of China (no. 2009ZF53062), the NPU Foundation for Fundamental Research (NPU-FFR-JC20100210 and JC201110), and the fund of the State Key Laboratory of Solidification Processing in NWPU (SKLSP201107). References [1] Holmquist M, Recina V, Ockborn J, Pettersson B, Zumalde E. Hot isostatic diffusion bonding of titanium alloy Ti-6Al-4V to gamma titanium aluminide IHI alloy 01A. Scripta Mater 1998;39:1101e6. [2] Wang XF, Ma M, Liu XB, Wu XQ, Tan CG, Shi RK, et al. Diffusion bonding of gTiAl alloy to Ti-6Al-4V alloy under hot pressure. Trans Nonferrous Met Soc China 2006;16:1059e63. [3] Blue CA, Lin. RY. Microstructural evolution in joining of TiAl with a liquid Ti alloy. Scr Metall Mater 1995;32:127e32. [4] Acoff VL, Wilkerson S, Arenas M. The effect of rolling direction on the weld structure and hardness of gamma-TiAl sheet material. Mater Sci Eng A 2002; 329:763e7. [5] Shiue RK, Wu SK, Chen YT, Shiue CY. Infrared brazing of Ti50Al50 and Ti-6Al-4V using two Ti-based filler metals. Intermetallics 2008;9:1083e9. [6] Yu WX, Li MQ, Hu YQ. Superplasticity and application of superplastic forming/ diffusion bonding technology. Mater Rev 2009;23:9e10 [in Chinese].
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