Mechanical property of induction brazing TiAl-based intermetallics to steel 35CrMo using AgCuTi filler metal

Materials Science and Engineering A 418 (2006) 45–52

Mechanical property of induction brazing TiAl-based intermetallics to steel 35CrMo using AgCuTi filler metal P. He ∗ , J.C. Feng, W. Xu National Key Laboratory of Advanced Welding Production Technology, Harbin Institute of Technology, Harbin 150001, People’s Republic of China Accepted 5 November 2005

Abstract TiAl-based intermetallics joined by vacuum induction brazing using Ag–Cu35.2–Ti1.8 filler alloy at 850–970 ◦ C for 1–10 min was investigated. The optimum brazing parameters are as follows: brazing temperature is 870–880 ◦ C, brazing time is 4–6 min. The maximum tensile strength of the joint is 320 MPa for the specimen induction brazed at for 870 ◦ C for 5 min, and at this time, the thickness of Al–Cu–Ti intermetallic compound layer is about one third of that of the whole interface. Interfacial microstructure of TiAl/Ag–Cu–Ti/35CrMo joint brazed at 870 ◦ C for 5 min is instable at 400 ◦ C. The instantaneous high temperature strength is 248 MPa at 400 ◦ C, the fracture location of joint is mainly at Al–Cu–Ti intermetallic compound layer, the fracture pattern of brazed joint deviate from transcrystalline crack at room temperature into intergranular crack. © 2005 Elsevier B.V. All rights reserved. Keywords: Brazed joints; Mechanical property; TiAl-based intermetallics; AgCuTi filler metal

1. Introduction The TiAl-based intermetallics is one of the most advanced intermetallics, and has successfully demonstrated its potential in aerospace and automotive applications for both military and civil purposes such as turbine blades, exhaust valves and turbo charger rotors [1,2]. The primary advantages of TiAl-based intermetallics are low specific gravity, high specific strength as well as good oxidation resistance, high stiffness and strength at elevated temperature compared to conventional titanium alloys. Therefore, they are considered as potential replacements for superalloys in aircraft turbine engines, airframes and automotive engines [3]. TiAl-based intermetallics can be fabricated by traditional casting and ingot metallurgy technology, but the processing cost tends to be high due to a high degree of segregation that occurs during the solidification. In addition to the conventional casting, there are many other manufacturing processes available for production of TiAl-based intermetallics, e.g. powder metallurgy, hot isostatic pressing (HIP), near net shape technology, direct laser fabrication, etc. [4–6].

∗

Corresponding author. Tel.: +86 451 8641 8146; fax: +86 451 8641 8146. E-mail address: [email protected] (P. He).

0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.11.005

Besides these manufacturing techniques for TiAl-based intermetallics, the effective utilization of TiAl-based intermetallics requires reliable joining techniques, especially to a variety of other materials [3,4]. The concept of utilizing TiAl-based intermetallics and steel to attain missile and tank engine turbo components by a bonding process is a recent approach [5,6]. Since the mechanical properties of TiAl-based intermetallics are sensitive to their microstructure, and their thermal expansion coefficient is fairly low under certain circumstances, and large internal stresses can be easily developed in their joints with other materials, the conventional fusion welding techniques cause complicated composition, a severe thermal cracking tendency and easy formation of brittle intermetallic compounds in the joints produced, and thus make the performance of joints very poor, and make it very difficult to obtain satisfactory joints [7–11]. Researches recommend diffusion bonding [12,13]. However, where TiAl intermetallic was diffusion bonded to steel directly, brittle intermetallic compounds FeAl, FeAl2 and TiC were formed at the interface of TiAl/steel joints, and the maximum tensile strength of the joint was only 170–185 MPa [14]. Vacuum brazing is a good choice for bonding materials difficult to join by traditional welding process. However, both microstructural evolution and strength evaluation of the brazed TiAl joint need further study. Compared with conventional vacuum furnace brazing, induction vacuum brazing is characterized by very rapid thermal cycles [15]. There are many successful joints that are reported

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in the literature using this novel technique [15–17]. The heating rate of a traditional vacuum brazing furnace is usually below 50 ◦ C/min, which is much slower than that of an induction furnace. The molten braze can react simultaneously with the substrate at temperatures above the liquidus of braze alloy and below the brazing temperature. Consequently, the initial stage of reaction at the interface between the braze alloy and substrate cannot be analyzed precisely using furnace brazing due to the insufficient heating rate during brazing. On the contrary, the maximum heating rate of an induction brazing furnace is as high as 3000 ◦ C/min [15]. With the aid of induction brazing, very early stage interfacial reaction kinetics of the brazed joint can be accurately examined by the experiment. Therefore, induction brazing is a powerful tool for the analysis of early stages of interfacial reaction in joints. Complicated components and mass production also require more convenient and easy to apply joining technology, such as induction brazing. This paper aims to demonstrate the feasibility of induction brazing a TiAl-based alloy with an Ag–Cu–Ti filler, and focusing on the fracture characteristic across the brazed joint and mechanical property of the brazed joints. 2. Experimental procedure The typical mechanical performance and nominal composition of the TiAl-based intermetallics and the steel used in this study are given in Table 1. The filler alloy was a 50 ␮m Ag–Cu35.2–Ti1.8 (wt%) foil, its melting range is 800–850 ◦ C, σ b is 343 MPa. The TiAl-based intermetallics was cut into specimens cylinders Φ 10 mm (diameter) × 30 mm (length). Brazed surfaces were ground flat using SiC grit papers (grit 800) followed by cleaning in ethanol and acetone prior to brazing. As mentioned above, compared with conventional furnace brazing, induction brazing has much faster thermal cycles. In addition, higher induction brazing temperature will do less damage to the base alloy. Since a high brazing temperature can significantly speed up the microstructural evolution of the brazed joint, higher brazing temperatures (such as 970 ◦ C) were chosen for this study. Induction brazing experiments were conducted in a vacuum of 3 × 10−4 Pa, and the heating rate was set at 2000 ◦ C/min throughout the experiment. All specimens were preheated to 600 ◦ C for 2 min before they were heated up to the brazing temperature. Because there was a time delay between the actual specimen temperature and programmer temperature, so time compensation was necessary in the experiment. The brazing time specified in the test was the actual specimen holding time. Hereby, a brazing temperature was from 850 ◦ C to 970 ◦ C and the brazing time was from 1 min to 10 min. The filler alloy was sandwiched between two the base alloy, and a K-type ther-

mal couple was inserted into the steel cylinder, in contact with the brazed interface. The brazed TiAl joints were cross-sectioned, perpendicular to the brazed interfaces, using a low-speed diamond saw, and the cross-sections of these joints were metallographically polished to a 0.5 ␮m diamond paste finish and cleaned in acetone. The fracture surfaces of the joints were observed by scanning electron microscopy (SEM, Cambridge Instruments S-570). The compositional analyses of the reaction products were performed by electron probe microanalysis (EPMA, JEOL JXA-8600 and GEOL superrobe 733). The microstructures of the brazed TiAl joints were examined by metallographic microscopy (NEUPHIOLYMPUS) and scanning electron microscopy (S-570). The crystal structures of the reaction products were identified from the reaction layer and the fracture surfaces of the joints by X-ray diffraction (XRD, JDX-3530M). The room temperature tensile strengths of the TiAl/steel joints were evaluated by means of a testing machine (INSTRON MODEL 1186), and the brazed specimen was drawed with a constant speed of 1 mm/min. Microhardness measurements were made using a microhardness tester (Nano-Indenter XP) with a load of 10 g and a duration time of 10 s. The experimental data were averaged from five measurements of each brazing condition. 3. Results and discussion 3.1. Tensile performance at room temperature The brazing temperature and the brazing time influence the quantity of atomic diffusion and the dissolution reaction between the melting filler metal and the master alloy, therefore they influenced the homogeneity of the composition and microstructure of the joint. Fig. 1 shows the effect of the brazing temperature on the tensile strength of the joint and the thickness of Al–Cu–Ti intermetallic compounds formed at the TiAl/Ag–Cu35.2–Ti1.8 interface with the brazing time 5 min (the thickness of intermetallic compounds denotes the extent of atomic interdiffusion). It can be seen that when the brazing temperature T is 850 ◦ C, the tensile strength of the joint is very low (only 154 MPa), this is because the amount of atomic diffusion is low and the dissolution reaction is insufficient between the melting filler metal and the master alloy when the brazing temperature is lower. With increasing T from 850 ◦ C to 870 ◦ C, the atomic diffusivity increases which results in an easier and more speedy diffusion reaction and chemical joining. Therefore, the tensile strength of the brazed joint increases with increasing T. When T is 870 ◦ C, the tensile strength of the joint is a maximum (320 MPa). When the brazing temperature increases more, more Al–Cu–Ti intermetallic compounds are formed at the interface, and the tensile

Table 1 Mechanical properties of the TiAl base alloy and steel 35CrMo Alloy

σ s (MPa)

σ b (MPa)

δ (%)

KIC (MPa m1/2 )

Ti–47.5Al–2.5V–1Cr (at%) Ti–48Al–2Cr–2Nb (at%) Fe–0.35C–0.27Si–0.55Mn–0.1Cr–0.2Mo (wt%)

400–450 414–480 830–840

560–620 559–668 970–990

3.8–4.2 2.6–3.2 11–13

18.0 16.8 –

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Fig. 1. Effect of brazing temperature on tensile strength of TiAl/AgCuTi/35CrMo joint and the thickness of Al–Ti–Cu intermetallic compounds at the interface of TiAl/AgCuTi (t = 5 min).

Fig. 2. Effect of brazing time on shear strength of tensile strength of TiAl/AgCuTi/35CrMo joint and the thickness of Al–Ti–Cu intermetallic compounds at the interface of TiAl/AgCuTi (T = 870 ◦ C).

strength of the joint will decrease. The selection of brazing time relies on the value of T. Under optimum brazing parameter (T is 870 ◦ C), which is concluded from the above experiment, the effect of brazing time on the tensile strength of the joint and the thickness of Al–Cu–Ti intermetallic compounds is shown in Fig. 2. It can be seen that when t is 1–5 min, the tensile strength increases rapidly, when t is more than 5 min, the thickness of Al–Cu–Ti intermetallic compounds increases markedly and the tensile strength of the brazed joint decreases. We show that the quick increase of Al–Cu–Ti intermetallic compounds leads to a decrease in the strength and an increase in the brittleness of the joint. In conclusion, the higher strength of joint brazed at 870 ◦ C for 5 min is obtained to be 320 MPa by using Ag–Cu–Ti braze alloy. SEM images in Fig. 3 show the fracture faces of a specimen that failed at the TiAl/Ag–Cu–Ti/35CrMo interface. The fracture is found to be flat and compact, and a majority of fracture area is occupied by part I, whose fracture morphology

is displayed in Fig. 3(b). River pattern and ligule’s pattern of cleavage fracture are observed as well as the torn arris of gliding fracture in Fig. 3(b), so it is classified as brittle fracture with some toughness. Results of energy spectrum analysis for part I show the composition is Ti32.7–Cu45.6–Al20.1–Cr1.1–Ag0.5 in at pct, based on this result, the fracture location is on Al–Cu–Ti intermetallic compound layer. X-ray analysis of the fracture surface also confirms it, as shown in Fig. 4. Fig. 5 shows microhardness measurements of interfacial microstructure, based on which relative hardness of interfacial microstructure can be determined. As shown in Fig. 5, AlCu2 Ti* phase close to TiAl substrate is harder, which is sensitive to defects. While Ag–Cu solid solution in the joint is softer, this can help releasing some stress. So the thickness of AlCu2 Ti* reaction layer and its proportion in the brazing interface must be controlled in order to obtain the high-strength joint. Because of the high strength got at 870 ◦ C for 5 min, the thickness ratio of interfacial reaction layer under this technological standard is considered to be suitable. The high-strength joint can be obtained when the thickness of AlCu2 Ti* phase is about one third of the width of the whole brazing interface. 3.2. Effect of different TiAl substrate on the microstructure and property of the joint

Fig. 3. X-ray analysis of the fracture of the specimen brazed at 870 ◦ C for 5 min.

Fig. 6 shows three different TiAl base alloys, and (b), (d), (f) are the enlargement of the part marked in (a), (c), (e) correspondingly. As displayed in the photos, the microstructure of all the three base alloys is binary-configuration microstructure, which is composed of ␥-TiAl single crystal and ␥-TiAl/␣2 -Ti3 Al lamellar colony, and the light phase is ␥-TiAl. While, because

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Fig. 4. Fracture pattern at 870 ◦ C for 5 min. (a) Macroscopic view; (b) microscopic view.

Fig. 5. Microhardness measurements of interfacial microstructure.

␣2 -Ti3 Al phase is too thin to recognize in the picture, the dark line is considered to be the interface between Ti3 Al phase and ␥-TiAl/␣2 -Ti3 Al phase. As displayed in Fig. 6, the volume fraction of␥-TiAl/␣2 -Ti3 Al lamellar colony is decreased along 1#, 3#, 2#, and that of TiAl single crystal is increased along 1#, 3#, 2#, while amount of TiAl/Ti3 Al interface is increased along 1#, 2#, 3#. Fig. 7 shows the interfacial microstructure of different TiAl base alloys brazed with Ag–Cu–Ti filler alloy. Specimens brazed with 1#, 2#, 3#, TiAl base alloy are also marked by 1#, 2#, 3#, correspondingly. In Fig. 7, Al–Cu–Ti intermetallic compound layers of a certain thickness are observed close to all the three TiAl base alloys, but the morphology of the intermetallic compound layers are different. First, AlCuTi* layer is observed obviously at the interface of specimen 1#, which extends to TiAl base alloy being acicular. The phenomena also appears in specimen 3#, but not as obvious as in specimen 1#. While this phenom-

Fig. 6. Microstructure of different TiAl base alloys. (a) 1# substrate; (b) 1# substrate; (c) 2# substrate; (d) 2# substrate; (e) 3# substrate; (f) 3# substrate.

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Fig. 7. Interfacial microstructure of different TiAl substrates brazed to steel 35CrMo. (a) Specimen 1#; (b) Specimen 2#; (c) Specimen 3#.

ena is hardly found in specimen 2#. Second, configurations of AlCu2 Ti* intermetallic compound reaction layers are different. AlCu2 Ti* at the interface of specimen 1# is nearly agglomerate, and AlCu2 Ti* at the interface of specimen 2# is nearly columnar, while that of specimen 3# is columnar. The two kinds of phenomenon are explained below. According to the aforementioned analysis results of reaction paths, AlCuTi* phase grows along Ti3 Al phase, while Ti3 Al phase exists in TiAl/Ti3 Al lamellar colony, so the growth configuration and size of the intermetallic compound are related to the morphology and amount of TiAl/Ti3 Al lamellar colony. The amount of TiAl/Ti3 Al lamellar colony is larger in 1#TiAl substrate, and the lamellar TiAl phase is thin, which means the volume fraction of Ti3 Al phase in the lamellar colony is large, consequently, AlCuTi* phase in substrate 1# grows obviously in the way of being acicular. Compared with substrate 2#, amount of TiAl/Ti3 Al lamellar colony in substrate 3# is larger, so the growth of AlCuTi* phase in substrate 3# is more obvious than in substrate 2#. The growth configuration of AlCu2 Ti* phase is related to the interface of TiAl/Ti3 Al in TiAl substrate. The reaction degree of the brazing seam can be approximately judged by the amount of eutectic pattern in the braze, the reason for which is the formation of intermetallic compound layer results in the consumption of Cu in the filler alloy, while the amount of eutectic pattern can reflect the amount of Cu indirectly. The more the eutectic pattern is, the more Cu exists in the residual filler alloy, the smaller the reaction degree of braze and TiAl substrate is. Consequently, the reaction degree of 1#, 2#, 3#, braze and TiAl substrate is decreased one by one,

resulting in different configurations of AlCu2 Ti* phase varying from agglomerate to columnar. The reason for the largest reaction degree of specimen 1# is that 1# TiAl substrate contains the most phase interfaces and more Ti, Al with high energy, which makes the interdiffusion of braze and base alloy easier. Fig. 8 shows fracture pattern of specimen 3#. As displayed in Fig. 8(a), the fracture is flat, and a majority of the fracture area is occupied by part I, whose fracture morphology is displayed in Fig. 8(b). According to the fracture morphology, the fracture remains a quasi-cleavage crack. Results of energy spectrum analysis for part I show the composition is Al19.6–Ti37–Cu41.3–Cr1.4–Ag0.6 in at pct, based on this result, the fracture location is still on Al–Cu–Ti intermetallic compound layer. Compared with the fracture of specimen 1# (Fig. 2), the macrofracture of specimen 3# is not as compact as that of specimen 1#. According to the microfracture, neither the torn arris nor the ligule’s pattern of cleavage fracture is observed in 1# and 3#, so the strength of specimen 1# is confirmed to be higher than that of specimen 3#. Results of tensile test also confirm it, which show the tensile strength of 1# and 3# are 320 MPa and 270 MPa, respectively. Lots of minute white flakes are observed on the fracture of specimen 1# when observing the microfracture attentively, and the white flake is considered to be the pattern of acicular AlCuTi* intermetallic compound fractured based on the size of the white flake. Although small white flakes also appear on the fracture of specimen 3#, the size of white flake is larger than that of specimen 1#, and the amount is not as large as that of specimen 1#. Hereby,

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Fig. 8. Fracture pattern of specimen 3#. (a) Macroscopic view; (b) microscopic view.

according to the result of joint strength, pattern of microfracture and the configuration of interfacial microstructure, formation of acicular AlCuTi* phase is considered to strengthen the bonding of brazing interface close to TiAl substrate. Based on the aforementioned information, degree of interfacial reaction is related to the configuration of AlCu2 Ti* reaction layer and amount of TiAl/Ti3 Al phase interface in the TiAl substrate under the same technological conditions. The more the TiAl/Ti3 Al phase interface is, the more severely brazing filler alloy reacts to TiAl substrate. The configuration and the size of AlCuTi* phase have relation to the configuration and amount of

TiAl/Ti3 Al lamellar colony. Visible acicular AlCuTi* phase can be generated by the approximately lamellar structure with low Al content, which is beneficial to the joining of TiAl substrate. 3.3. Elevated temperature property of the brazed joint To assess the elevated temperature property of the brazing joint, thermal retardation test at 400 ◦ C is performed at the brazing interface, and microstructure alteration of the brazing interface is analyzed. Fig. 9 shows the microstructure of brazing interface with different TiAl substrates and before and after ther-

Fig. 9. Microstructure of brazing interface before and after thermal retardation at 400 ◦ C. (a) Microstructure of specimen 2# before thermal retardation; (b) microstructure of specimen 2# after thermal retardation for 12 h; (c) microstructure of specimen 3# before thermal retardation; (d) microstructure of specimen 3# after thermal retardation for 20 h.

P. He et al. / Materials Science and Engineering A 418 (2006) 45–52

mal retardation at 400 ◦ C for different time. As shown in Fig. 9, interfacial microstructure changes after thermal retardation at 400 ◦ C. First, the straight boundary of intermetallic compound reaction layer and residual braze microstructure has disappeared (marked by 1). Second, recrystallization occurs on the residual braze microstructure layer, and there is visible grain boundary in the microstructure (marked by 2). Third, formation configuration of microstructure on the intermetallic compound layer changes (marked by 3) in a sort of way. In addition, intermetallic compound layer at TiAl substrate interface blows-up close to TiAl substrate (marked by 4). According to the aforementioned phenomenon, it is confirmed residual brazing filler alloy layer, phase interface between residual brazing filler alloy layer and intermetallic compound layer, and phase interface between intermetallic compound layer and TiAl substrate are the instable areas of the brazing interface. Therefore, interfacial microstructure of TiAl/Ag–Cu–Ti/35CrMo joint brazed at 870 ◦ C for 5 min is instable at 400 ◦ C. The tensile strength of brazed joint at 400 ◦ C is 248 MPa. Fig. 10 shows the high temperature fracture at 400 ◦ C and 300 ◦ C. The fracture can be divided into

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Table 2 Energy spectrum analysis of high temperature fracture at 400 ◦ C (at%) Location

Al

Ag

Ti

Cu

C

I II III

27.16 6.72 5.07

0.9 75.09 44.87

34.43 – 6.97

36.05 18.29 15.4

– – 27.69

three areas based on the macrostructure of the fracture, whose energy spectrum analysis results are displayed in Table 2. Part III mainly consists of Ag, Cu, C elements, and the unwelded defect is observed at the edge of specimen in Fig. 10(b). Part II is Ag base solid solution layer, whose cracking pattern is identified to be intergranular crack according to Fig. 10(c), and cavity crack is found at the grain boundary, which can be easily generated with low stress when the temperature is higher comparatively. Therefore, Ag–Cu solid solution structure in the brazing seam can do harm to elevated temperature property. The fracture location of joint is mainly in part I, and it is identified to be Al–Cu–Ti intermetallic compound based on the energy spectrum analy-

Fig. 10. High temperature fracture. (a) Macroscopic fracture at 400 ◦ C; (b) microscopic of part I fracture at 400 ◦ C; (c) microscopic of part II fracture at 400 ◦ C; (d) microscopic of part III fracture at 400 ◦ C; (e) macroscopic fracture at 300 ◦ C; (f) microscopic of part I fracture at 300 ◦ C.

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sis. The high temperature fracture at 300 ◦ C also occurs on the Al–Cu–Ti intermetallic compound layer, which is displayed in Fig. 10(e), (f). Comparing the fracture of part I at room temperature, 300 ◦ C and 400 ◦ C, the fracture pattern of brazed joint is found to deviate from transcrystalline fracture into intergranular crack gradually. 4. Summary and conclusions The following conclusions can be obtained by investigating vacuum induction brazing TiAl-based intermetallics to steel 35CrMo with AgCuTi filler alloy: (1) The optimum brazing parameters are as follows: brazing temperature is 870–880 ◦ C, brazing time is 4–6 min. The maximum tensile strength of the joint is 320 MPa for the specimen induction brazed at for 870 ◦ C for 5 min, and at this time, the thickness of Al–Cu–Ti intermetallic compound layer is about one third of that of the whole interface. (2) Interfacial microstructure of TiAl/Ag–Cu–Ti/35CrMo joint brazed at 870 ◦ C for 5 min is instable at 400 ◦ C. The instantaneous high temperature intensity is 248 MPa at 400 ◦ C, the fracture location of joint is mainly at Al–Cu–Ti intermetallic compound layer, the fracture pattern of brazed joint is found to deviate from transcrystalline crack at room temperature into intergranular crack. Acknowledgement The research is sponsored by the National Natural Science Foundation of China (No.50505008).

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