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Low-temperature direct diffusion bonding of hydrogenated TC4 alloy and GH3128 superalloy
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Low-temperature direct diffusion bonding of hydrogenated TC4 alloy and GH3128 superalloy Lixia Zhang, Zhan Sun*, Junmiao Shi, Xiaofeng Ye, Zhiye Yang, Jicai Feng State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
article info
abstract
Article history:
Bonding at high temperatures can cause many problems, such as an induction of high
Received 1 September 2018
stress, grain coarsening and strict requirements for bonding equipments, etc.. In this
Received in revised form
paper, a hydrogenation approach was utilized for the TC4 alloy before the dissimilar ma-
23 November 2018
terials bonding process. Effects of hydrogen contents on the diffusion behavior of the TC4/
Accepted 27 November 2018
GH3128 joints were investigated. Particularly, the mechanism that the hydrogenation
Available online 8 January 2019
affected the low-temperature bonding process of the TC4/GH3128 joints was discussed. By regulating the bonding temperatures, holding durations and hydrogen contents, a
Keywords:
maximum of 92 MPa was achieved. The formation mechanism of the diffusion bonded
Hydrogenation
TC4/GH3128 joint was proposed. This novel metal hydrogenation idea can offer new in-
Diffusion bonding
sights on the development of the low-temperature joining particularly suitable for dis-
Superalloy
similar materials joining.
TC4 alloy
© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction TC4 alloy is an ideal structural material for aerospace engineering, automotive technology etc. due to its high strength, low density, good corrosion resistance and excellent toughness [1e3]. GH3128, as a kind of Co-free Ni-based superalloy, was developed as an ideal candidate for the intermediate heat exchanger, the burner chamber of the aeronautic engine, etc. [4]. A study on the reliable joining of the TC4 and GH3128 superalloy is particularly significant since it can save the cost and lower the weight of the structural component. Normally, the diffusion bonding temperature for TC4 alloy was higher than 850 C [5,6]. However, high temperatures will induce large thermal stresses on the joints and great thermal damages to base substrates [7,8]. So lowering bonding temperatures and maintaining a sound bonding of the TC4/GH3128 couples has always been a research target.
Fortunately, hydrogenation has been reported to be an effective approach to improve the plasticity and the atomic diffusion of metals [9e13], which can be utilized to achieve a low-temperature bonding process. Dong et al. [14] studied the effect of the hydrogen addition on the plasticity of the Zr-based metallic glasses. Results showed that the plastic strain increased from 1.2% to 10% after the hydrogenation treatment. The mechanism was attributed to the large amount of free volume induced by the hydrogen addition. Compared with the hydrogenation treatment on metallic glasses, more studies have been performed on the hydrogenation of the Ti-based alloy. Yuan et al. [15] conducted a hydrogenation treatment on the TC21 alloy at 750 C for 2 h. They found an increase of the b phase content happened. When the hydrogen content increased to the optimal value of 0.9 wt%, the hydrogenated TC21 alloy exhibited a much better plasticity and a lower flow stress of 150e200 MPa induced by the increased b phase content. The plasticity improvement of the hydrogenated
* Corresponding author. E-mail address: [email protected] (Z. Sun). https://doi.org/10.1016/j.ijhydene.2018.11.206 0360-3199/© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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titanium alloy will facilitate a sound contact during the diffusion bonding process. Li et al. [16] studied the gas bulging of the Ti-55 alloy using a hydrogenation approach. They found when a maxim ratio of height to radius reached 1.58 for the Ti-55 alloy, the one without the hydrogen addition formed at 925 C. While the one with the hydrogen addition formed only at 825 C, which was a 100 C reduction. Afterwards, they tested the elongation of the Ti-55 alloy with different hydrogen additions [17]. Results showed that a 706.3% elongation for the Ti55 alloy with 0.1 wt% hydrogen addition was achieved at 825 C. While the elongation of the Ti-55 alloy without the hydrogen addition was only 361.3% under the same test condition. As to the hydrogenation treatment in bonding researches, some progress has been made. Wang et al. [18] developed a novel hydrogenated Zr interlayer to conduct the self-bonding of the TC4 alloy. They found that the diffusion layer thickness increased by 5 and 1.2 times respectively when bonding at 650 C and 750 C. Feng et al. [19] studied the self-bonding of the TC4 alloy with a 0.3 wt% hydrogen addition. They found the fine needle a0 martensite and lamellar (a þ b) structure formed instead of the a0 martensite and bH phase. A shear strength of 450 MPa was achieved at 810 C for 8 min under a pressure of 6 MPa. However, the direct diffusion bonding of dissimilar materials, namely the TC4 alloy and the GH3128 superalloy at low temperatures (＜800 C) remains a challenge. In this paper, a hydrogenated TC4 alloy was prepared first, which laid a solid foundation for the subsequent lowtemperature bonding. The mechanism that the hydrogenation process affected the bonding process was discussed and proposed.
Experimental procedures Two kinds of TC4 alloys were prepared, i.e., a normal TC4 alloy and a hydrogenated TC4 alloy, which was hydrogenated at 750 C for 1 h. The TC4 alloy was heated in Ar atmosphere until it reached 750 C, when H2 started to fill into the furnace. Different hydrogenation contents were controlled by setting different H2 pressures, which were 0.05e0.15 MPa. The hydrogen contents were 0.1 wt%, 0.3 wt% and 0.5 wt%, which were weighed using an electronic balance with an accuracy of 0.01 mg. The GH3128 sheet was obtained by hot rolling. Chemical compositions of the two base metals were shown in Table 1. Samples for microstructural observation and shear tests were 10 10 1.3 mm (GH3128), 5 5 2 mm (TC4 alloy) and 20 10 1.3 mm（GH3128）, 5 5 2 mm (TC4). All the samples were polished using SiC abrasive papers from 240 to 1500# followed by the acetone cleaning for 10 min. The schematic of the assembly was shown in Fig. 1. The samples were then put into a M60 vacuum brazing furnace. The heating process did not start until the vacuum level reached 104 Torr. The heating rate was 25 C/min and the cooling rate was controlled
as 5 C/min. The program stopped until the temperature was lower than 200 C. Interfacial microstructure was observed using a Quanta 200FEG SEM. Phase structures of the typical phases were detected using a rotary anode X-ray diffractometer (XRD, D/max-rB type) manufactured by Rigaku Corporation. Hydrogenated TC4 alloys were cut into 3 2 2 mm and then analyzed using a STA 449C/6/G equipment. Ar was selected as the protective air and the heating rate was 25 C/min.
Results and discussion Typical microstructure of the TC4/GH3128 joint Fig. 2 shows the typical microstructure of the direct bonded TC4/GH3128 joint at 720 C/90 min under 15 MPa with 0.3 wt% hydrogen. The joint was well bonded without any unbonded region. The joint can be divided into 4 regions in total as shown in Fig. 2(a) and (b) shows the microstructure of the TC4 alloy. b phase became denser when it came close to the bonded interface circled by the dashed box. Ni atoms diffused into TC4 alloy, leading to a Ni rich area. Ni is a common b stabilizing element [20], which can lower the a/b transition temperature. Another point was the hydrogenation treatment, which could further facilitate the a/b transition of the TC4 alloy and stabilize the b phase [21,22]. A line scan was conducted to characterize the linear distribution of different elements. It was shown in Fig. 2(c) that the W, Mo, Cr, Mn and Fe still existed on the GH3128 side and a sharp decrease happened when it came to Layer Ⅱ. Al and V only existed in the diffusion layer of the TC4 alloy. Mutual diffusion of Ni and Ti happened in the bonded area. The accumulation of Cr, Mo and W atoms in Layer Ⅲ and Ⅳ inhibited the diffusion of Ni from the GH3128 to the TC4 side. At the same time, the formation of TieNi compounds could facilitate the diffusion of Ni into the TC4 alloy. These two factors worked together, resulting in a Ni-deficient area in Layer Ⅲ. To further confirm the types of the phases in the bonded area, typical spots were detected as shown in Fig. 2(a) and the data were shown in Table 2. For Layer Ⅰ and Ⅱ, the atomic ratios of Ti and Ni were 2: 1 and 1: 1. According to the TieNi binary phase diagram, these two layers were mainly composed of Ti2Ni and TiNi intermetallics. For Layer Ⅳ, it mainly consisted of Ni and Cr, indicating the existence of (Ni,Cr)ss. For Layer Ⅲ, Ti, Ni and Cr were the main elements. (Ni,Cr)ss should exist in this layer. Besides, TiNi and Ti2Ni should also exist in this layer according to the atomic ratio of Ti and Ni.
Effect of hydrogen contents on microstructure of the bonded joint Fig. 3(a)e(d) show the microstructural evolution of the TC4/ GH3128 joints with different hydrogen contents at 720 C/ 90 min under 15 MPa.
Table 1 e Chemical compositions of TC4 and GH3128 alloys (at.%).
TC4 GH3128
Ti
Al
V
Ni
Cr
Mo
W
Fe
Si
Bal. 0.4e0.8
5.5e6.8 0.4e0.8
3.5e4.5 e
e Bal.
e 19.0e22.0
e 7.5e9.0
e 7.5e9.0
0.3 2.0
0.15 e
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Fig. 1 e Schematic of the assembly of the samples for shear tests and microstructural observations. (a) Sample for shear tests; (b) Sample for microstructural observations.
Fig. 2 e Typical microstructure of the direct bonded TC4/GH3128 joint at 720 C/90/15 MPa with 0.3 wt% hydrogenation. (a) The whole joint; (b) On the TC4 alloy side; (c) Linear concentration distribution of different elements
Table 2 e Chemical compositions of each layer (at.%).
A B C D
Ti
Ni
Cr
Al
V
Mo
W
Possible Phases
61.01 46.84 22.33 9.90
28.38 47.49 23.63 27.43
0.45 1.22 36.01 45.54
8.30 4.14 4.51 1.30
1.35 0 0.96 0.25
0.23 0.25 7.76 9.84
0.28 0.26 4.80 5.74
Ti2Ni TiNi TiNi、Ti2Ni and (Ni,Cr)ss (Ni,Cr)ss
The joints were always composed of 4 diffusion layers regardless of the different hydrogen contents. However, the whole reaction layer thickened with the increase of the hydrogen content, which indicates that the hydrogenation can promote the diffusion of the atoms at the interface.
Fig. 3(e) shows the thickness variation of the whole reaction layer, which demonstrates that the thickening of the whole reaction layer gradually slowed down with the increase of the hydrogenation contents. The whole reaction layer reached 8.8 mm in average, which was 33% thicker than the case
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Fig. 3 e SEM images and reaction layer thickness of the diffusion bonded joints (720 C, 90 min, 15 MPa) with different hydrogen contents.(a) 0.0 wt%; (b) 0.1 wt%; (c) 0.3 wt%; (d) 0.5 wt%; (e) Reaction layer thickness variation with different hydrogen contents.
without the hydrogenation treatment. The thickness of the whole reaction layer has a great impact on the shear strength of the joints, which will be discussed later in this paper.
Microstructural evolution of the TC4/GH3128 joints bonded with various bonding parameters Fig. 4 shows the effects of the bonding temperatures on microstructural evolution of the TC4/GH3128 joints. Each single reaction layer as well as the whole reaction layer thickened with the elevation of the bonding temperatures. When bonding temperatures were low (620 and 650 C), unclosed voids could be seen at the Layer Ⅱ/Layer Ⅲ interface. When the bonding temperature further elevated, the reaction layer became denser without any unclosed voids. Another point was that the number of gray b phases significantly increased and the sizes became larger, connecting into a mesh with the elevation of the bonding temperatures. This was probably because of the sufficient diffusion of Ni from the GH3128 into the TC4 alloy, leading to a higher volume of the b phase. From the linear concentration distribution of Ni shown in Fig. 2(c), it seems that the Ni concentration in the TC4 alloy is very low. Whether Ni atoms have diffused into the TC4 alloy or not is still questionable. To confirm that, area distributions of different elements were shown in Fig. 5. From Fig. 5(c), we can see that there indeed exists Ni in the TC4 substrate. However, the distribution does not connect into a whole area. Instead, the Ni in the TC4 substrate connected into a “Mesh”. Put Fig. 5(a) and (c) into comparison, we can see that the Ni
distribution corresponds well with the b phase distribution, indicating that Ni probably locates inside the b phase. Since Ni is a b phase stabilizing element, it's reasonable to deduce that a higher bonding temperature leads to a more sufficient diffusion of Ni into the TC4 substrate, which, as a result, results in a higher volume of the b phase. Fig. 6(a) shows the reaction layer thickness variation versus different bonding temperatures. The thickness of the whole reaction layer, increased from 2.52 mm (0.63 mm for Ti2Ni layer, 1.15 mm for TiNi layer) to 10.6 mm (1.36 mm for the Ti2Ni layer, 4.69 mm for the TiNi layer) when the bonding temperature elevated from 620 to 780 C. The thickness increase for the whole reaction layer, the Ti2Ni layer and the TiNi layer were 320%, 116% and 310% respectively. Fig. 6(b) shows the shear strength of the TC4/GH3128 joints bonded at various temperatures. The shear strength increased first with the elevation of the bonding temperature and reached the maximum value of 92 MPa at 680 C for 60 min under 15 MPa with 0.3 wt% hydrogen. When the bonding temperature was low, the plasticity of the TC4 alloy was poor and a sound contact of the TC4/GH3129 interface could hardly be achieved, indicated by the unclosed voids shown in Fig. 4(a) and (b). Besides, the diffusion driving force was weak at low temperatures, leading to thinner reaction layers, which could hardly bear the high shear force. When the bonding temperature was high, the brittle Ti2Ni layer was too thick and became the preferential position to crack, which will be discussed later in this paper. Fig. 7 shows the microstructural evolution of the TC4/ GH3128 joints bonded for different holding durations.
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Fig. 4 e Microstructural evolution of the TC4/GH3128 joints at various bonding temperatures for 60 min under 15 MPa with 0.3 wt% hydrogen.(a) 620 C; (b) 650 C; (c) 680 C; (d) 720 C; (e) 750 C; (f) 780 C.
Fig. 5 e Microstructure (a) and elemental area distribution of the TC4/GH3128 joints (bef). The reaction layer thickened with the increase of the holding durations. When the holding duration was 20 min, unclosed voids could be seen at the TC4/GH3128 interface. When holding for 60 and 90 min, the reaction layers became dense. From Fig. 8(a), it shows that the growth of the reaction layers sped up first from 40 to 60 min and slowed down when
the holding duration extended from 60 to 90 min. The highest shear strength was achieved when bonding for 60 min as shown in Fig. 8(b). When the holding duration was too short, the reaction layer was quite thin. In particular, unclosed voids could be easily found when holding for 20 min, which should be responsible for the low shear strength.
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Fig. 6 e Reaction layer thickness (a) and shear strength (b) versus different bonding temperatures.
Fig. 7 e Microstructure evolution of the TC4/GH3128 joints bonded for different holding durations. (a) 20 min; (b) 40 min; (c) 60 min; (d) 90 min.
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Fig. 8 e Effect of holding durations on reaction layer thickness (a) and shear strength (b) of the TC4/GH3128 joints.
Fig. 9 shows the fracture morphology of the TC4/GH3128 joint bonded at 680 C for 60 min with 0.3 wt% hydrogen addition under 15 MPa. The fracture was quite flat with tiny steps, which indicated a typical brittle failure. From the magnified SEM image, there were 2 typical spots, i.e., Spot A and B. From the EDS data shown in Table 3, Spot A should be Ti2Ni and Spot B should be TiNi. So the fracture of the TC4/GH3128 joint was composed of Ti2Ni and TiNi and among them, Ti2Ni was the main phase. To further confirm the crack position, XRD was employed to analyze the phases of the fracture. It is shown in Fig. 10 that TiNi and Ti2Ni were the main phases of the fracture. Ti2Ni was a typical phase with high hardness and brittleness [25,26], which became the preferential location to crack under the thermal stress during the cooling process. When these two layers grow too thick, i.e., both the TiNi and the Ti2Ni layer further grow thicker when the holding duration extends from 60 to 90 min, it can hardly accommodate the thermal stress by a sufficient plastic deformation, which leads to a stress concentration inside these two layers. The joint thus failed with greater probability and the shear strength decreased as a result. Fig. 11 shows a through-wall crack formed in the Ti2Ni layer of the TC4/GH3128 joint bonded at 780 C for 90 min. Although there were 4 reaction layers formed in the TC4/ GH3128 joint. Cracks preferentially formed inside the Ti2Ni layer. The brittle Ti2Ni layer under the bonding parameter was quite thick which was around 2.3 mm. The crack mainly propagated inside the Ti2Ni layer and deflected into the TiNi layer at some local spots, which was quite consistent with the fracture analysis above. The TiNi phase showed a better
plasticity, evidenced by the uneven morphology of the Spot B in Fig. 9. While the morphology of the Ti2Ni phase (Spot A) was quite flat in contrast as shown in Fig. 9.
Formation mechanism of the direct bonded TC4/GH3128 joint From the analysis above, the Ti2Ni and TiNi layers formed during the direct diffusion bonding. The formation free energies of these two reaction products were described as follows: 2TiðsÞ þ NiðsÞ ¼ Ti2 NiðsÞ DGðTi2 NiÞ ¼ 49:12 þ 17:208
(1) 103 TðkJÞ mol
(2)
TiðsÞ þ NiðsÞ ¼ TiNiðsÞ
(3)
kJ DGðTiNiÞ ¼ 54:60 þ 18:133 103 T mol
(4)
Fig. 12 shows the relationship between the formation free energies of the reaction products and the bonding temperatures. The temperatures selected in this paper were between 600 and 800 C (873e1073 K). In this range, the formation free energies of Ti2Ni and TiNi were far below 0 and the value for TiNi was even lower than that of the Ti2Ni phase shown in Fig. 12(a), indicating the formation of the TiNi phase was more thermodynamically favorable. From the diffusion point of view, Bastin et al. [27] pointed out that the diffusion of Ni was much faster than Ti. So a large number of Ni atoms would diffuse into the TC4 alloy and the
Fig. 9 e Fracture morphology of the TC4/GH3128 joint.
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Table 3 e Chemical compositions of Spot A and B (at.%).
A B
Ti
Ni
Cr
Al
V
Possible phases
53.76 44.91
37.20 48.27
0.85 1.77
7.68 4.48
0.51 0.57
Ti2Ni TiNi
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became closer with the elevation of the heating temperature under the applied pressure as shown in Fig. 13(a). When the heating temperature was high enough to drive the diffusion, mutual diffusion of Ti and Ni atoms happened. The diffusion of Ni in the TC4 alloy was much faster than Ti. Ni concentration reached 33 at.% in the TC4 alloy first, resulting in the formation of the Ti2Ni layer as shown in Fig. 13(b). Then Ni further diffused into the TC4 alloy and reached 50 at.%, resulting in the formation of the TiNi layer as shown in Fig. 13(c). Ti2Ni layer was pushed towards the TC4 alloy and became continuous. Later, Formula (5) happened. The thickness of the Ti2Ni layer changed little and the TiNi layer further thickened. On the GH3128 side, Ni concentration decreased, resulting in the formation of a layer rich in Cr, Mo and W, i.e., a layer composed of (Ni,Cr)ss formed as shown in Fig. 13(d). Ti further diffused towards the (Ni,Cr)ss and a layer composed of TiNi þ Ti2Ni þ (Ni,Cr)ss formed as shown in Fig. 13(e).
Effect of hydrogenation treatment on low-temperature diffusion bonding of the TC4/GH3128 joints
Fig. 10 e XRD spectrum of the fracture.
Fig. 11 e Cracks formation of the TC4/GH3128 joint. Ti2Ni phase formed first with the increase of the Ni content, followed by the formation of the TiNi phase. Some of the Ti2Ni would react with Ni, leading to the formation of the TiNi phase, which was described as follows: Ti2 NiðsÞ þ NiðsÞ ¼ 2TiNiðsÞ
(5)
The formation free energy of Formula (5) was shown in Fig. 12(b). This reaction was also thermodynamically favorable, which explained the slow growth of the Ti2Ni layer. Based on the analysis above, the formation mechanism of the TC4/GH3128 joints can be summarized as follows: After the assembly of the TC4/GH3128 samples, the physical contact between the TC4 and the GH3128 alloy gradually
From the microstructure shown in Fig. 3, we found that the hydrogenation treatment had a great impact on the reaction layer thickness, which would further affect the joint property. To study the effect of the hydrogenation treatment on the low-temperature diffusion bonding, dehydration characteristics of the TC4 alloy were investigated. Fig. 14 shows the differential thermal analysis (DTA) of the hydrogenated TC4 alloy during a continuous heating process (from room temperature to 1300 C at a heating rate of 25 C/min). It shows that little change happened for the thermal gravity (TG) curves of the TC4 alloy with 0.0 wt% and 0.5 wt% hydrogen before the heating temperature reached 600 C. When further increasing the heating temperature, the TG curve of the TC4 alloy with 0.0 wt% hydrogen ramped up rapidly. Since the TC4 alloy possesses poor high temperature oxidation resistance [28,29], the rapid rise of the TG curve indicates a strong oxidation of the TC4 alloy. While for the TC4 alloy with 0.5 wt% hydrogen, the TG curve sharply dropped down at around 620 C and reached the minimum value at 980 C, indicating a strong hydride decomposition and the hydrogen release. For the DTA curves, there was a small endothermic peak at around 990 C for the TC4 alloy with 0.0 wt% hydrogen, which corresponds well with the b transus reported by Peng et al. [30]. While according to the variation of the TG curve shown in Fig. 14(b), the hydride decomposition temperature should range from 600 to 800 C. The endothermic peak of the hydride decomposition in this range was not obvious. The endothermic peak ranged from 800 to 900 C should be the a/b phase transition point, indicating that the a/b phase transition temperature was greatly reduced after the hydrogenation treatment. Fig. 15 shows the TG and DTA curves of the TC4 alloy with 0.5 wt% hydrogen at various holding temperatures. Sharp peaks appeared when the TG curve started to drop down, indicating the hydrogen release happened when holding for 24 min at various temperatures. However, the hydrogen release duration was 13 min shorter when holding temperature increased from 680 C to 750 C, indicating that a stronger hydride decomposition happened at higher temperatures. The effect of the
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Fig. 12 e Formation free energies of the (a) Ti2Ni and TiNi phases; (b) TiNi phase.
Fig. 13 e Schematic of the formation mechanism of the TC4/GH3128 joint. (a) Physical contact; (b) Formation of the Ti2Ni layer; (c) Formation of the TiNi layer; (d) Formation of the (Ni, Cr)ss and (e) Formation of the TiNi þ Ti2Ni þ (Ni, Cr)ss.
Fig. 14 e DTA of the hydrogenated TC4 alloy during a continuous heating process. (a) Without hydrogenation; (b) 0.5 wt% hydrogenation. hydrogenation on the direct diffusion bonding of the TC4/ GH3128 joint can be summarized as follows: The hydrogenation process can effectively facilitate the formation of the b phase with better plasticity at relatively low temperatures [31]. The TC4 alloy with a higher content of b phase can easily deform under the applied pressure, laying a
solid foundation for the subsequent bonding process. Besides, the hydrogenation treatment can clean the surface by weakening the oxidation of the TC4 alloy. From Fig. 14(a), it demonstrated that the TC4 alloy got easily oxidized at high temperatures with 0.0 wt% hydrogen. In this regard, an oxide film formed on the TC4 surface, which was harmful for the
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Fig. 15 e Isothermal DTA and TG analysis of the TC4 alloy with 0.5 wt% hydrogen at various temperatures. (a) 680 C; (b) 750 C.
Fig. 16 e Schematic of hydrogenation effect on promoting the diffusion of the dissolved atoms.
diffusion process and weakened the bonding property of the joint. While a hydrogen release happened for the TC4 alloy with 0.5 wt% hydrogen during the heating stage, which will protect the TC4 surface from severe oxidation. The diffusion was thus sufficient and the bonding was strong. The mechanism that the hydrogen promotes the diffusion of the dissolved atoms is proposed as shown in Fig. 16. After the hydrogenation of the TC4 alloy, a large number of the hydrogen atoms have dissolved into the TC4 substrate as interstitial atoms. As a result, a crystal lattice distortion happened for the TC4 substrate. Due to the weak-bonding effect of the interstitial hydrogen stated by Troiano [32]. The formation energy of the vacancy and migration energy were greatly reduced. As a result, more vacancies formed. Then the atoms from the substrate in close contact with the hydrogenated TC4 alloy dissolved. Due to the large number of vacancies created by the hydrogenation treatment, the diffusion of the dissolved atoms became easier and faster. When the heating temperature was high and the holding duration was long enough, the dissolved hydrogen sufficiently escaped from the TC4 alloy. The concentration of vacancies decreased and a crystal relaxation happened as a result. So the hydrogenation primarily worked on the diffusion process during the 4 stages as described in Fig. 16.
Conclusions (1) Typical microstructure of the TC4/GH3128 joint was analyzed, i.e., TC4(aþb)/Ti2Ni/TiNi/TiNi þ Ti2Ni þ (Ni,Cr) ss/(Ni,Cr)ss/GH3128. Hydrogenation can help promote
the diffusion of the interfacial atoms, which led to a 33% increase in the whole reaction layer thickness after the TC4 alloy was treated with 0.3% hydrogen. (2) Effect of bonding temperatures and durations on joint microstructure and properties were investigated. The maximum shear strength was 92 MPa, which was obtained at 680 C for 60 min under 15 MPa with 0.3 wt% hydrogen. Fracture mainly located at the Ti2Ni and TiNi layers. Ti2Ni was quite brittle and was the preferential layer to crack. (3) Formation mechanism of the direct bonded TC4/ GH3128 joint was proposed. Ti2Ni layer and TiNi formed in sequence when Ni reached 33 at.% and 50 at.% in the TC4 alloy. Then on the GH3128 side, (Ni,Cr)ss formed as a result of the reduction of Ni content. The TiNi þ Ti2Ni þ (Ni,Cr)ss layer finally formed as Ti further diffused towards the GH3128 alloy. (4) The effects of the hydrogenation on the lowtemperature diffusion bonding can be summarized as follows. The hydrides decomposed between 620 and 980 C. Higher percent of the b phase was achieved after the TC4 alloy was hydrogenated, which led to a better plasticity and a more close contact between base metals. The hydrogenation treatment can prevent the TC4 alloy from oxidation, achieving a clean and activated surface. Finally, the interstitial hydrogen can lead to a weak-bonging effect of the TC4 lattice. The formation energy of the vacancy thus lowered and the diffusion activation energy of the atoms was reduced, which enabled a reliable joining of the TC4/ GH3128.
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Acknowledgements This work was supported by the China Postdoctoral Science Foundation funded project [Grant No. 2018M631923], Fundamental Research Funds for the Central Universities (Grant No.HIT.NSRIF.2019006), National Natural Science Foundation of China [Grant No. 51805113, 51522404 and 51775142].
references
[1] Godfrey T, Wisbey A, Goodwin P, Bagnall K, Ward-Close C. Microstructure and tensile properties of mechanically alloyed Tie6A1e4V with boron additions. Mater Sci Eng A 2000;282:240e50. [2] Ma X, Li F, Cao J, Li J, Chen H, Zhao C. Vickers microhardness and microstructure relationship of Ti-6Al-4V alloy under cyclic forward-reverse torsion and monotonic torsion loading. Mater Des 2017;114:271e81. [3] Venkatesh B, Chen D, Bhole S. Effect of heat treatment on mechanical properties of Tie6Ale4V ELI alloy. Mater Sci Eng A 2009;506:117e24. [4] Shi Y, Yuan K, Zhao X, Wu Y. Study on comparison between Inconel 617 and Gh3128 as structural material candidates for intermediate heat exchanger. In: 2013 21st International Conference on Nuclear Engineering. American Society of Mechanical Engineers; 2013. V002T03A36-VT03A36. [5] Xue Z, Yang Q, Gu L, Hao X, Ren Y, Geng Y. Diffusion bonding of TiAl based alloy to Tie6Ale4V alloy using amorphous interlayer. Mater Sci Eng Tecnol 2015;46:40e6. [6] Zhang H, Li J, Ma P, Xiong J, Zhang F. Study on microstructure and impact toughness of TC4 titanium alloy diffusion bonding joint. Vacuum 2018;152:272e7. € [7] Cam G, Ozdemir U, Ventzke V, Kocak M. Microstructural and mechanical characterization of diffusion bonded hybrid joints. J Mater Sci 2008;43:3491e9. [8] Wang X, Yang Y, Luo X, Zhang W, Zhao G, Huang B. An investigation of Ti-43Al-9V/Ti-6Al-4V interface by diffusion bonding. Intermetallics 2013;36:127e32. [9] Hongo T, Edalati K, Iwaoka H, Arita M, Matsuda J, Akiba E, et al. High-pressure torsion of palladium: hydrogen-induced softening and plasticity in ultrafine grains and hydrogeninduced hardening and embrittlement in coarse grains. Mater Sci Eng A 2014;618:1e8. [10] Ma T, Chen R, Zheng D, Guo J, Ding H, Su Y, et al. Hydrogeninduced softening of Tie44Ale6Nbe1Cre2V alloy during hot deformation. Int J Hydrogen Energy 2017;42:8329e37. [11] Ma T, Chen R, Zheng D, Guo J, Ding H, Su Y, et al. Effect of bphase stabilizing elements and high temperature (1373e1693 K) on hydrogen absorption in TiAl alloys. Int J Hydrogen Energy 2017;42:86e95. [12] Tao P, Zhang W, Tu Q, Yang Y. Response of largedeformation behavior to strain rates in a bulk-metallic glass by hydrogen. Int J Hydrogen Energy 2016;41:9690e5.
[13] Zhang X, Zheng Y, Yuan B. Dehydrogenation technology of hydrogenated TC21 alloy. J Alloy Compd 2017;695:1444e7. [14] Dong F, Su Y, Luo L, Wang L, Wang S, Guo J, et al. Enhanced plasticity in Zr-based bulk metallic glasses by hydrogen. Int J Hydrogen Energy 2012;37:14697e701. [15] Yuan B, Zheng Y, Gong L, Chen Q, Lv M. Effect of hydrogen content on microstructures and room-temperature compressive properties of TC21 alloy. Mater Des 2016;94:330e7. [16] Li X, Chen N, Wu H, Chen J, Qu F. Low-temperature superplastic gas bulging of Ti-55 alloy by hydrogen addition. Int J Hydrogen Energy 2018;43:12455e9. [17] Li X, Jiang J, Wang S, Chen J, Wang Y. Effect of hydrogen on the microstructure and superplasticity of Ti-55 alloy. Int J Hydrogen Energy 2017;42:6338e49. [18] Wang D, Cao J, Li W, Dai X, Feng J. Zr hydrogenation by cathodic charging and its application in TC4 alloy diffusion bonding. Int J Hydrogen Energy 2017;42:6350e9. [19] Feng J, Liu H, He P, Cao J. Effects of hydrogen on diffusion bonding of hydrogenated Ti6Al4V alloy containing 0.3 wt% hydrogen at fast heating rate. Int J Hydrogen Energy 2007;32:3054e8. [20] Wang Y, Jiao M, Yang Z, Wang D, Liu Y. Vacuum brazing of Ti 2 AlNb and TC4 alloys using TieZreCueNi and TieZreCueNi þ Mo filler metals: microstructural evolution and mechanical properties. Arch Civ Mech Eng 2018;18:546e56. [21] Zhu T, Li M. Effect of hydrogen addition on the microstructure of TC21 alloy. Mater Sci Eng A 2010;527:7080e5. [22] Zhu T, Li M. Effect of hydrogen on the b transus temperature of TC21 alloy. Mater Char 2011;62:852e6. [25] Hashi K, Ishikawa K, Matsuda T, Aoki K. Hydrogen permeation characteristics of (V, Ta)eTieNi alloys. J Alloy Compd 2005;404:273e8. [26] Zoeram AS, Mousavi SAAA. Laser welding of Tie6Ale4V to nitinol. Mater Des 2014;61:185e90. [27] Bastin GF, Rieck GD. Diffusion in the titanium-nickel system: II. Calculations of chemical and intrinsic diffusion coefficients. Metall Trans 1974;5:1827e31. [28] Wei DB, Zhang PZ, Yao ZJ, Liang WP, Miao Q, Xu Z. Oxidation of double-glow plasma chromising coating on TC4 titanium alloys. Corros Sci 2013;66:43e50. [29] Zhang ZG, Peng YP, Mao YL, Pang CJ, Lu LY. Effect of hot-dip aluminizing on the oxidation resistance of Tie6Ale4V alloy at high temperatures. Corros Sci 2012;55:187e93. [30] Peng X, Guo H, Wang T, Yao Z. Effects of b treatments on microstructures and mechanical properties of TC4-DT titanium alloy. Mater Sci Eng A 2012;533:55e63. [31] Liu H, Cao J, He P, Feng JC. Effect of hydrogen on diffusion bonding of commercially pure titanium and hydrogenated Ti6Al4V alloys. Int J Hydrogen Energy 2009;34:1108e13. [32] Troiano AR. The role of hydrogen and other interstitials in the mechanical behavior of metals. Metallogr Microstruct Anal 2016;5:557e69.
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Low-temperature direct diffusion bonding of hydrogenated TC4 alloy and GH3128 superalloy Lixia Zhang, Zhan Sun*, Junmiao Shi, Xiaofeng Ye, Zhiye Yang, Jicai Feng State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
article info
abstract
Article history:
Bonding at high temperatures can cause many problems, such as an induction of high
Received 1 September 2018
stress, grain coarsening and strict requirements for bonding equipments, etc.. In this
Received in revised form
paper, a hydrogenation approach was utilized for the TC4 alloy before the dissimilar ma-
23 November 2018
terials bonding process. Effects of hydrogen contents on the diffusion behavior of the TC4/
Accepted 27 November 2018
GH3128 joints were investigated. Particularly, the mechanism that the hydrogenation
Available online 8 January 2019
affected the low-temperature bonding process of the TC4/GH3128 joints was discussed. By regulating the bonding temperatures, holding durations and hydrogen contents, a
Keywords:
maximum of 92 MPa was achieved. The formation mechanism of the diffusion bonded
Hydrogenation
TC4/GH3128 joint was proposed. This novel metal hydrogenation idea can offer new in-
Diffusion bonding
sights on the development of the low-temperature joining particularly suitable for dis-
Superalloy
similar materials joining.
TC4 alloy
© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Introduction TC4 alloy is an ideal structural material for aerospace engineering, automotive technology etc. due to its high strength, low density, good corrosion resistance and excellent toughness [1e3]. GH3128, as a kind of Co-free Ni-based superalloy, was developed as an ideal candidate for the intermediate heat exchanger, the burner chamber of the aeronautic engine, etc. [4]. A study on the reliable joining of the TC4 and GH3128 superalloy is particularly significant since it can save the cost and lower the weight of the structural component. Normally, the diffusion bonding temperature for TC4 alloy was higher than 850 C [5,6]. However, high temperatures will induce large thermal stresses on the joints and great thermal damages to base substrates [7,8]. So lowering bonding temperatures and maintaining a sound bonding of the TC4/GH3128 couples has always been a research target.
Fortunately, hydrogenation has been reported to be an effective approach to improve the plasticity and the atomic diffusion of metals [9e13], which can be utilized to achieve a low-temperature bonding process. Dong et al. [14] studied the effect of the hydrogen addition on the plasticity of the Zr-based metallic glasses. Results showed that the plastic strain increased from 1.2% to 10% after the hydrogenation treatment. The mechanism was attributed to the large amount of free volume induced by the hydrogen addition. Compared with the hydrogenation treatment on metallic glasses, more studies have been performed on the hydrogenation of the Ti-based alloy. Yuan et al. [15] conducted a hydrogenation treatment on the TC21 alloy at 750 C for 2 h. They found an increase of the b phase content happened. When the hydrogen content increased to the optimal value of 0.9 wt%, the hydrogenated TC21 alloy exhibited a much better plasticity and a lower flow stress of 150e200 MPa induced by the increased b phase content. The plasticity improvement of the hydrogenated
* Corresponding author. E-mail address: [email protected] (Z. Sun). https://doi.org/10.1016/j.ijhydene.2018.11.206 0360-3199/© 2018 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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titanium alloy will facilitate a sound contact during the diffusion bonding process. Li et al. [16] studied the gas bulging of the Ti-55 alloy using a hydrogenation approach. They found when a maxim ratio of height to radius reached 1.58 for the Ti-55 alloy, the one without the hydrogen addition formed at 925 C. While the one with the hydrogen addition formed only at 825 C, which was a 100 C reduction. Afterwards, they tested the elongation of the Ti-55 alloy with different hydrogen additions [17]. Results showed that a 706.3% elongation for the Ti55 alloy with 0.1 wt% hydrogen addition was achieved at 825 C. While the elongation of the Ti-55 alloy without the hydrogen addition was only 361.3% under the same test condition. As to the hydrogenation treatment in bonding researches, some progress has been made. Wang et al. [18] developed a novel hydrogenated Zr interlayer to conduct the self-bonding of the TC4 alloy. They found that the diffusion layer thickness increased by 5 and 1.2 times respectively when bonding at 650 C and 750 C. Feng et al. [19] studied the self-bonding of the TC4 alloy with a 0.3 wt% hydrogen addition. They found the fine needle a0 martensite and lamellar (a þ b) structure formed instead of the a0 martensite and bH phase. A shear strength of 450 MPa was achieved at 810 C for 8 min under a pressure of 6 MPa. However, the direct diffusion bonding of dissimilar materials, namely the TC4 alloy and the GH3128 superalloy at low temperatures (＜800 C) remains a challenge. In this paper, a hydrogenated TC4 alloy was prepared first, which laid a solid foundation for the subsequent lowtemperature bonding. The mechanism that the hydrogenation process affected the bonding process was discussed and proposed.
Experimental procedures Two kinds of TC4 alloys were prepared, i.e., a normal TC4 alloy and a hydrogenated TC4 alloy, which was hydrogenated at 750 C for 1 h. The TC4 alloy was heated in Ar atmosphere until it reached 750 C, when H2 started to fill into the furnace. Different hydrogenation contents were controlled by setting different H2 pressures, which were 0.05e0.15 MPa. The hydrogen contents were 0.1 wt%, 0.3 wt% and 0.5 wt%, which were weighed using an electronic balance with an accuracy of 0.01 mg. The GH3128 sheet was obtained by hot rolling. Chemical compositions of the two base metals were shown in Table 1. Samples for microstructural observation and shear tests were 10 10 1.3 mm (GH3128), 5 5 2 mm (TC4 alloy) and 20 10 1.3 mm（GH3128）, 5 5 2 mm (TC4). All the samples were polished using SiC abrasive papers from 240 to 1500# followed by the acetone cleaning for 10 min. The schematic of the assembly was shown in Fig. 1. The samples were then put into a M60 vacuum brazing furnace. The heating process did not start until the vacuum level reached 104 Torr. The heating rate was 25 C/min and the cooling rate was controlled
as 5 C/min. The program stopped until the temperature was lower than 200 C. Interfacial microstructure was observed using a Quanta 200FEG SEM. Phase structures of the typical phases were detected using a rotary anode X-ray diffractometer (XRD, D/max-rB type) manufactured by Rigaku Corporation. Hydrogenated TC4 alloys were cut into 3 2 2 mm and then analyzed using a STA 449C/6/G equipment. Ar was selected as the protective air and the heating rate was 25 C/min.
Results and discussion Typical microstructure of the TC4/GH3128 joint Fig. 2 shows the typical microstructure of the direct bonded TC4/GH3128 joint at 720 C/90 min under 15 MPa with 0.3 wt% hydrogen. The joint was well bonded without any unbonded region. The joint can be divided into 4 regions in total as shown in Fig. 2(a) and (b) shows the microstructure of the TC4 alloy. b phase became denser when it came close to the bonded interface circled by the dashed box. Ni atoms diffused into TC4 alloy, leading to a Ni rich area. Ni is a common b stabilizing element [20], which can lower the a/b transition temperature. Another point was the hydrogenation treatment, which could further facilitate the a/b transition of the TC4 alloy and stabilize the b phase [21,22]. A line scan was conducted to characterize the linear distribution of different elements. It was shown in Fig. 2(c) that the W, Mo, Cr, Mn and Fe still existed on the GH3128 side and a sharp decrease happened when it came to Layer Ⅱ. Al and V only existed in the diffusion layer of the TC4 alloy. Mutual diffusion of Ni and Ti happened in the bonded area. The accumulation of Cr, Mo and W atoms in Layer Ⅲ and Ⅳ inhibited the diffusion of Ni from the GH3128 to the TC4 side. At the same time, the formation of TieNi compounds could facilitate the diffusion of Ni into the TC4 alloy. These two factors worked together, resulting in a Ni-deficient area in Layer Ⅲ. To further confirm the types of the phases in the bonded area, typical spots were detected as shown in Fig. 2(a) and the data were shown in Table 2. For Layer Ⅰ and Ⅱ, the atomic ratios of Ti and Ni were 2: 1 and 1: 1. According to the TieNi binary phase diagram, these two layers were mainly composed of Ti2Ni and TiNi intermetallics. For Layer Ⅳ, it mainly consisted of Ni and Cr, indicating the existence of (Ni,Cr)ss. For Layer Ⅲ, Ti, Ni and Cr were the main elements. (Ni,Cr)ss should exist in this layer. Besides, TiNi and Ti2Ni should also exist in this layer according to the atomic ratio of Ti and Ni.
Effect of hydrogen contents on microstructure of the bonded joint Fig. 3(a)e(d) show the microstructural evolution of the TC4/ GH3128 joints with different hydrogen contents at 720 C/ 90 min under 15 MPa.
Table 1 e Chemical compositions of TC4 and GH3128 alloys (at.%).
TC4 GH3128
Ti
Al
V
Ni
Cr
Mo
W
Fe
Si
Bal. 0.4e0.8
5.5e6.8 0.4e0.8
3.5e4.5 e
e Bal.
e 19.0e22.0
e 7.5e9.0
e 7.5e9.0
0.3 2.0
0.15 e
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Fig. 1 e Schematic of the assembly of the samples for shear tests and microstructural observations. (a) Sample for shear tests; (b) Sample for microstructural observations.
Fig. 2 e Typical microstructure of the direct bonded TC4/GH3128 joint at 720 C/90/15 MPa with 0.3 wt% hydrogenation. (a) The whole joint; (b) On the TC4 alloy side; (c) Linear concentration distribution of different elements
Table 2 e Chemical compositions of each layer (at.%).
A B C D
Ti
Ni
Cr
Al
V
Mo
W
Possible Phases
61.01 46.84 22.33 9.90
28.38 47.49 23.63 27.43
0.45 1.22 36.01 45.54
8.30 4.14 4.51 1.30
1.35 0 0.96 0.25
0.23 0.25 7.76 9.84
0.28 0.26 4.80 5.74
Ti2Ni TiNi TiNi、Ti2Ni and (Ni,Cr)ss (Ni,Cr)ss
The joints were always composed of 4 diffusion layers regardless of the different hydrogen contents. However, the whole reaction layer thickened with the increase of the hydrogen content, which indicates that the hydrogenation can promote the diffusion of the atoms at the interface.
Fig. 3(e) shows the thickness variation of the whole reaction layer, which demonstrates that the thickening of the whole reaction layer gradually slowed down with the increase of the hydrogenation contents. The whole reaction layer reached 8.8 mm in average, which was 33% thicker than the case
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Fig. 3 e SEM images and reaction layer thickness of the diffusion bonded joints (720 C, 90 min, 15 MPa) with different hydrogen contents.(a) 0.0 wt%; (b) 0.1 wt%; (c) 0.3 wt%; (d) 0.5 wt%; (e) Reaction layer thickness variation with different hydrogen contents.
without the hydrogenation treatment. The thickness of the whole reaction layer has a great impact on the shear strength of the joints, which will be discussed later in this paper.
Microstructural evolution of the TC4/GH3128 joints bonded with various bonding parameters Fig. 4 shows the effects of the bonding temperatures on microstructural evolution of the TC4/GH3128 joints. Each single reaction layer as well as the whole reaction layer thickened with the elevation of the bonding temperatures. When bonding temperatures were low (620 and 650 C), unclosed voids could be seen at the Layer Ⅱ/Layer Ⅲ interface. When the bonding temperature further elevated, the reaction layer became denser without any unclosed voids. Another point was that the number of gray b phases significantly increased and the sizes became larger, connecting into a mesh with the elevation of the bonding temperatures. This was probably because of the sufficient diffusion of Ni from the GH3128 into the TC4 alloy, leading to a higher volume of the b phase. From the linear concentration distribution of Ni shown in Fig. 2(c), it seems that the Ni concentration in the TC4 alloy is very low. Whether Ni atoms have diffused into the TC4 alloy or not is still questionable. To confirm that, area distributions of different elements were shown in Fig. 5. From Fig. 5(c), we can see that there indeed exists Ni in the TC4 substrate. However, the distribution does not connect into a whole area. Instead, the Ni in the TC4 substrate connected into a “Mesh”. Put Fig. 5(a) and (c) into comparison, we can see that the Ni
distribution corresponds well with the b phase distribution, indicating that Ni probably locates inside the b phase. Since Ni is a b phase stabilizing element, it's reasonable to deduce that a higher bonding temperature leads to a more sufficient diffusion of Ni into the TC4 substrate, which, as a result, results in a higher volume of the b phase. Fig. 6(a) shows the reaction layer thickness variation versus different bonding temperatures. The thickness of the whole reaction layer, increased from 2.52 mm (0.63 mm for Ti2Ni layer, 1.15 mm for TiNi layer) to 10.6 mm (1.36 mm for the Ti2Ni layer, 4.69 mm for the TiNi layer) when the bonding temperature elevated from 620 to 780 C. The thickness increase for the whole reaction layer, the Ti2Ni layer and the TiNi layer were 320%, 116% and 310% respectively. Fig. 6(b) shows the shear strength of the TC4/GH3128 joints bonded at various temperatures. The shear strength increased first with the elevation of the bonding temperature and reached the maximum value of 92 MPa at 680 C for 60 min under 15 MPa with 0.3 wt% hydrogen. When the bonding temperature was low, the plasticity of the TC4 alloy was poor and a sound contact of the TC4/GH3129 interface could hardly be achieved, indicated by the unclosed voids shown in Fig. 4(a) and (b). Besides, the diffusion driving force was weak at low temperatures, leading to thinner reaction layers, which could hardly bear the high shear force. When the bonding temperature was high, the brittle Ti2Ni layer was too thick and became the preferential position to crack, which will be discussed later in this paper. Fig. 7 shows the microstructural evolution of the TC4/ GH3128 joints bonded for different holding durations.
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Fig. 4 e Microstructural evolution of the TC4/GH3128 joints at various bonding temperatures for 60 min under 15 MPa with 0.3 wt% hydrogen.(a) 620 C; (b) 650 C; (c) 680 C; (d) 720 C; (e) 750 C; (f) 780 C.
Fig. 5 e Microstructure (a) and elemental area distribution of the TC4/GH3128 joints (bef). The reaction layer thickened with the increase of the holding durations. When the holding duration was 20 min, unclosed voids could be seen at the TC4/GH3128 interface. When holding for 60 and 90 min, the reaction layers became dense. From Fig. 8(a), it shows that the growth of the reaction layers sped up first from 40 to 60 min and slowed down when
the holding duration extended from 60 to 90 min. The highest shear strength was achieved when bonding for 60 min as shown in Fig. 8(b). When the holding duration was too short, the reaction layer was quite thin. In particular, unclosed voids could be easily found when holding for 20 min, which should be responsible for the low shear strength.
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Fig. 6 e Reaction layer thickness (a) and shear strength (b) versus different bonding temperatures.
Fig. 7 e Microstructure evolution of the TC4/GH3128 joints bonded for different holding durations. (a) 20 min; (b) 40 min; (c) 60 min; (d) 90 min.
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Fig. 8 e Effect of holding durations on reaction layer thickness (a) and shear strength (b) of the TC4/GH3128 joints.
Fig. 9 shows the fracture morphology of the TC4/GH3128 joint bonded at 680 C for 60 min with 0.3 wt% hydrogen addition under 15 MPa. The fracture was quite flat with tiny steps, which indicated a typical brittle failure. From the magnified SEM image, there were 2 typical spots, i.e., Spot A and B. From the EDS data shown in Table 3, Spot A should be Ti2Ni and Spot B should be TiNi. So the fracture of the TC4/GH3128 joint was composed of Ti2Ni and TiNi and among them, Ti2Ni was the main phase. To further confirm the crack position, XRD was employed to analyze the phases of the fracture. It is shown in Fig. 10 that TiNi and Ti2Ni were the main phases of the fracture. Ti2Ni was a typical phase with high hardness and brittleness [25,26], which became the preferential location to crack under the thermal stress during the cooling process. When these two layers grow too thick, i.e., both the TiNi and the Ti2Ni layer further grow thicker when the holding duration extends from 60 to 90 min, it can hardly accommodate the thermal stress by a sufficient plastic deformation, which leads to a stress concentration inside these two layers. The joint thus failed with greater probability and the shear strength decreased as a result. Fig. 11 shows a through-wall crack formed in the Ti2Ni layer of the TC4/GH3128 joint bonded at 780 C for 90 min. Although there were 4 reaction layers formed in the TC4/ GH3128 joint. Cracks preferentially formed inside the Ti2Ni layer. The brittle Ti2Ni layer under the bonding parameter was quite thick which was around 2.3 mm. The crack mainly propagated inside the Ti2Ni layer and deflected into the TiNi layer at some local spots, which was quite consistent with the fracture analysis above. The TiNi phase showed a better
plasticity, evidenced by the uneven morphology of the Spot B in Fig. 9. While the morphology of the Ti2Ni phase (Spot A) was quite flat in contrast as shown in Fig. 9.
Formation mechanism of the direct bonded TC4/GH3128 joint From the analysis above, the Ti2Ni and TiNi layers formed during the direct diffusion bonding. The formation free energies of these two reaction products were described as follows: 2TiðsÞ þ NiðsÞ ¼ Ti2 NiðsÞ DGðTi2 NiÞ ¼ 49:12 þ 17:208
(1) 103 TðkJÞ mol
(2)
TiðsÞ þ NiðsÞ ¼ TiNiðsÞ
(3)
kJ DGðTiNiÞ ¼ 54:60 þ 18:133 103 T mol
(4)
Fig. 12 shows the relationship between the formation free energies of the reaction products and the bonding temperatures. The temperatures selected in this paper were between 600 and 800 C (873e1073 K). In this range, the formation free energies of Ti2Ni and TiNi were far below 0 and the value for TiNi was even lower than that of the Ti2Ni phase shown in Fig. 12(a), indicating the formation of the TiNi phase was more thermodynamically favorable. From the diffusion point of view, Bastin et al. [27] pointed out that the diffusion of Ni was much faster than Ti. So a large number of Ni atoms would diffuse into the TC4 alloy and the
Fig. 9 e Fracture morphology of the TC4/GH3128 joint.
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Table 3 e Chemical compositions of Spot A and B (at.%).
A B
Ti
Ni
Cr
Al
V
Possible phases
53.76 44.91
37.20 48.27
0.85 1.77
7.68 4.48
0.51 0.57
Ti2Ni TiNi
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became closer with the elevation of the heating temperature under the applied pressure as shown in Fig. 13(a). When the heating temperature was high enough to drive the diffusion, mutual diffusion of Ti and Ni atoms happened. The diffusion of Ni in the TC4 alloy was much faster than Ti. Ni concentration reached 33 at.% in the TC4 alloy first, resulting in the formation of the Ti2Ni layer as shown in Fig. 13(b). Then Ni further diffused into the TC4 alloy and reached 50 at.%, resulting in the formation of the TiNi layer as shown in Fig. 13(c). Ti2Ni layer was pushed towards the TC4 alloy and became continuous. Later, Formula (5) happened. The thickness of the Ti2Ni layer changed little and the TiNi layer further thickened. On the GH3128 side, Ni concentration decreased, resulting in the formation of a layer rich in Cr, Mo and W, i.e., a layer composed of (Ni,Cr)ss formed as shown in Fig. 13(d). Ti further diffused towards the (Ni,Cr)ss and a layer composed of TiNi þ Ti2Ni þ (Ni,Cr)ss formed as shown in Fig. 13(e).
Effect of hydrogenation treatment on low-temperature diffusion bonding of the TC4/GH3128 joints
Fig. 10 e XRD spectrum of the fracture.
Fig. 11 e Cracks formation of the TC4/GH3128 joint. Ti2Ni phase formed first with the increase of the Ni content, followed by the formation of the TiNi phase. Some of the Ti2Ni would react with Ni, leading to the formation of the TiNi phase, which was described as follows: Ti2 NiðsÞ þ NiðsÞ ¼ 2TiNiðsÞ
(5)
The formation free energy of Formula (5) was shown in Fig. 12(b). This reaction was also thermodynamically favorable, which explained the slow growth of the Ti2Ni layer. Based on the analysis above, the formation mechanism of the TC4/GH3128 joints can be summarized as follows: After the assembly of the TC4/GH3128 samples, the physical contact between the TC4 and the GH3128 alloy gradually
From the microstructure shown in Fig. 3, we found that the hydrogenation treatment had a great impact on the reaction layer thickness, which would further affect the joint property. To study the effect of the hydrogenation treatment on the low-temperature diffusion bonding, dehydration characteristics of the TC4 alloy were investigated. Fig. 14 shows the differential thermal analysis (DTA) of the hydrogenated TC4 alloy during a continuous heating process (from room temperature to 1300 C at a heating rate of 25 C/min). It shows that little change happened for the thermal gravity (TG) curves of the TC4 alloy with 0.0 wt% and 0.5 wt% hydrogen before the heating temperature reached 600 C. When further increasing the heating temperature, the TG curve of the TC4 alloy with 0.0 wt% hydrogen ramped up rapidly. Since the TC4 alloy possesses poor high temperature oxidation resistance [28,29], the rapid rise of the TG curve indicates a strong oxidation of the TC4 alloy. While for the TC4 alloy with 0.5 wt% hydrogen, the TG curve sharply dropped down at around 620 C and reached the minimum value at 980 C, indicating a strong hydride decomposition and the hydrogen release. For the DTA curves, there was a small endothermic peak at around 990 C for the TC4 alloy with 0.0 wt% hydrogen, which corresponds well with the b transus reported by Peng et al. [30]. While according to the variation of the TG curve shown in Fig. 14(b), the hydride decomposition temperature should range from 600 to 800 C. The endothermic peak of the hydride decomposition in this range was not obvious. The endothermic peak ranged from 800 to 900 C should be the a/b phase transition point, indicating that the a/b phase transition temperature was greatly reduced after the hydrogenation treatment. Fig. 15 shows the TG and DTA curves of the TC4 alloy with 0.5 wt% hydrogen at various holding temperatures. Sharp peaks appeared when the TG curve started to drop down, indicating the hydrogen release happened when holding for 24 min at various temperatures. However, the hydrogen release duration was 13 min shorter when holding temperature increased from 680 C to 750 C, indicating that a stronger hydride decomposition happened at higher temperatures. The effect of the
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Fig. 12 e Formation free energies of the (a) Ti2Ni and TiNi phases; (b) TiNi phase.
Fig. 13 e Schematic of the formation mechanism of the TC4/GH3128 joint. (a) Physical contact; (b) Formation of the Ti2Ni layer; (c) Formation of the TiNi layer; (d) Formation of the (Ni, Cr)ss and (e) Formation of the TiNi þ Ti2Ni þ (Ni, Cr)ss.
Fig. 14 e DTA of the hydrogenated TC4 alloy during a continuous heating process. (a) Without hydrogenation; (b) 0.5 wt% hydrogenation. hydrogenation on the direct diffusion bonding of the TC4/ GH3128 joint can be summarized as follows: The hydrogenation process can effectively facilitate the formation of the b phase with better plasticity at relatively low temperatures [31]. The TC4 alloy with a higher content of b phase can easily deform under the applied pressure, laying a
solid foundation for the subsequent bonding process. Besides, the hydrogenation treatment can clean the surface by weakening the oxidation of the TC4 alloy. From Fig. 14(a), it demonstrated that the TC4 alloy got easily oxidized at high temperatures with 0.0 wt% hydrogen. In this regard, an oxide film formed on the TC4 surface, which was harmful for the
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Fig. 15 e Isothermal DTA and TG analysis of the TC4 alloy with 0.5 wt% hydrogen at various temperatures. (a) 680 C; (b) 750 C.
Fig. 16 e Schematic of hydrogenation effect on promoting the diffusion of the dissolved atoms.
diffusion process and weakened the bonding property of the joint. While a hydrogen release happened for the TC4 alloy with 0.5 wt% hydrogen during the heating stage, which will protect the TC4 surface from severe oxidation. The diffusion was thus sufficient and the bonding was strong. The mechanism that the hydrogen promotes the diffusion of the dissolved atoms is proposed as shown in Fig. 16. After the hydrogenation of the TC4 alloy, a large number of the hydrogen atoms have dissolved into the TC4 substrate as interstitial atoms. As a result, a crystal lattice distortion happened for the TC4 substrate. Due to the weak-bonding effect of the interstitial hydrogen stated by Troiano [32]. The formation energy of the vacancy and migration energy were greatly reduced. As a result, more vacancies formed. Then the atoms from the substrate in close contact with the hydrogenated TC4 alloy dissolved. Due to the large number of vacancies created by the hydrogenation treatment, the diffusion of the dissolved atoms became easier and faster. When the heating temperature was high and the holding duration was long enough, the dissolved hydrogen sufficiently escaped from the TC4 alloy. The concentration of vacancies decreased and a crystal relaxation happened as a result. So the hydrogenation primarily worked on the diffusion process during the 4 stages as described in Fig. 16.
Conclusions (1) Typical microstructure of the TC4/GH3128 joint was analyzed, i.e., TC4(aþb)/Ti2Ni/TiNi/TiNi þ Ti2Ni þ (Ni,Cr) ss/(Ni,Cr)ss/GH3128. Hydrogenation can help promote
the diffusion of the interfacial atoms, which led to a 33% increase in the whole reaction layer thickness after the TC4 alloy was treated with 0.3% hydrogen. (2) Effect of bonding temperatures and durations on joint microstructure and properties were investigated. The maximum shear strength was 92 MPa, which was obtained at 680 C for 60 min under 15 MPa with 0.3 wt% hydrogen. Fracture mainly located at the Ti2Ni and TiNi layers. Ti2Ni was quite brittle and was the preferential layer to crack. (3) Formation mechanism of the direct bonded TC4/ GH3128 joint was proposed. Ti2Ni layer and TiNi formed in sequence when Ni reached 33 at.% and 50 at.% in the TC4 alloy. Then on the GH3128 side, (Ni,Cr)ss formed as a result of the reduction of Ni content. The TiNi þ Ti2Ni þ (Ni,Cr)ss layer finally formed as Ti further diffused towards the GH3128 alloy. (4) The effects of the hydrogenation on the lowtemperature diffusion bonding can be summarized as follows. The hydrides decomposed between 620 and 980 C. Higher percent of the b phase was achieved after the TC4 alloy was hydrogenated, which led to a better plasticity and a more close contact between base metals. The hydrogenation treatment can prevent the TC4 alloy from oxidation, achieving a clean and activated surface. Finally, the interstitial hydrogen can lead to a weak-bonging effect of the TC4 lattice. The formation energy of the vacancy thus lowered and the diffusion activation energy of the atoms was reduced, which enabled a reliable joining of the TC4/ GH3128.
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Acknowledgements This work was supported by the China Postdoctoral Science Foundation funded project [Grant No. 2018M631923], Fundamental Research Funds for the Central Universities (Grant No.HIT.NSRIF.2019006), National Natural Science Foundation of China [Grant No. 51805113, 51522404 and 51775142].
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