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Microstructure evolution and brazing mechanism of Ti2AlC–Ti2AlC joint by using pure-silver filler metal
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CERAMICS INTERNATIONAL
Ceramics International 41 (2015) 8203–8210 www.elsevier.com/locate/ceramint
Microstructure evolution and brazing mechanism of Ti2AlC–Ti2AlC joint by using pure-silver filler metal Guochao Wang, Guohua Fan, Jie Zhangn, Tianpeng Wang, Xiaowen Liu School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China Received 11 January 2015; received in revised form 7 March 2015; accepted 8 March 2015 Available online 14 March 2015
Abstract This paper investigated the microstructure evolution of a Ti2AlC–Ti2AlC joint brazed using pure-silver filler metal at 1030 1C for 5 min, characterized the newly formed products in the joint and elucidated their formation mechanism. Furthermore, it reveals the effect of Ag on the crystal structure of Ti2AlC. Ti2AlC is perfectly bonded by using Ag filler, with no crack or pore being observed. During brazing, the molten Ag filler metal infiltrates into the Ti2AlC along the grain boundaries. Al and Ti migrate out of Ti2AlC and dissolve into the Ag, forming the Ag[Al, Ti] solid solution. Ag3[Ti, Al] is produced on the basis of Ag[Al, Ti] through ordering transition during the cooling process. Simultaneously, Ag element diffuses into the crystal structure of Ti2AlC. Nevertheless, the Ti2AlC retains its crystal structure under the partial loss of Al and Ti. The shear strength of the joint was 13278 MPa, nearly 73% of that of the Ti2AlC substrate. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Ti2AlC; Ag; Microstructure; Brazing; Shear strength
1. Introduction In recent decades, the nano-laminated ternary compounds with a formula of Mn þ 1AXn (where M is an early transition metal, A is an A-group (mostly groups of 13 or 14) element, and X is C and/or N) are intensively studied. The so-called MAX phases have layered hexagonal structure with space group P63/mmc. Nearclose-packed Mn þ 1Xn layers are interleaved with hexagonal nets of A element. The unique structure endows the compounds with an unusual combination of metallic and ceramic behaviors [1–3]. As a fascinating member of this series, Ti2AlC is of particular interest since this compound has the lowest density (4.11 g cm 3) and the best oxidation resistance in the family of MAX phases [4–6]. Besides, it has high Young's modulus [7], good thermal and electrical conductivity [8,9], self-lubricating function [10–13], high thermal stability and excellent mechanical properties in the radiation environment [6], which make it a suitable candidate for many structural and functional applications, such as sliding
n
Corresponding author. Tel./fax: þ 86 451 86414234. E-mail address: [email protected] (J. Zhang).
http://dx.doi.org/10.1016/j.ceramint.2015.03.040 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
electrical contact, conducting ceramic in harsh environment and reinforcement for soft metals [8,10–13]. Similar to most other ceramics, large-dimensioned Ti2AlC ceramic is difficult to synthesize due to its narrow phase range in the Ti–Al–C ternary phase diagram [14]. This problem is usually solved by joining ceramics, a method which allows the large, complex, multifunctional ceramic components to be manufactured through the controlled integration of smaller, simpler, more easily manufactured parts [15]. Therefore, studies devoted to joining Ti2AlC are significant for broadening their applications. In the past decades, researchers have been able to join Ti–Al–C and Ti–Si–C ceramics due to their structural characteristics and the ease with which Al3O2 or Ti3Al(Si)C2 solid solutions are formed [16–20]. These joints are obtained by solid phase diffusion bonding, by which the crystal structure of the ceramic matrix can be well sustained. In our previous research [21,22], brazing was adopted to join MAX phases and metallic parts because of its simplicity, good repetitiveness and excellent adaptability of joint size and shape. Ti2AlC/Cu brazed joint can be successfully fabricated using Ag–Cu eutectic filler alloy. Due to the weak Ti–Al bond in Ti2AlC, Al was apt to be divorced from the crystal structure and react with Cu. As one of the substrates to be joined, the parent copper partially
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dissolved into the molten Ag–Cu filler during brazing. The reaction was intensified by increasing the content of Cu in the filler alloy, resulting in the further out-migration of Al and Ti from Ti2AlC. After a certain extent of reaction was reached, the crystal structure of Ti2AlC would decompose, leading to the change in intrinsic properties of the phases. It was also found that Ag is much less reactive to Ti2AlC comparing with Cu. No stable compound can be found in the Ag–Al binary phase diagram, which is an indication of a much weaker tendency in Ag–Al reaction [23]. Therefore, puresilver filler was applied to join the Ti2AlC ceramics in this study. Wang et al. proposed that the decomposition of Ti2AlC is triggered by the outward diffusion of Al atoms [24]. It was also calculated that Al vacancy shows lower formation energy of mono-vacancy than that of Ti and C vacancy. Furthermore, defective Ti2AlC may maintain phase stability down to a substoichiometry of Ti2Al0.5C. The structural stability of Ti2AlC is determined by its depletion of Al at high temperatures in an oxygen-containing atmosphere or in species having a high affinity to Al [25]. Therefore, the structural stability of Ti2AlC should be concerned during the brazing of Ti2AlC. The purpose of this work is threefold: firstly, to obtain the well-bonded joint of Ti2AlC and clarify the microstructure evolution of the joint during brazing; secondly, to characterize the microstructure and elucidate the brazing mechanism of the Ti2AlC/Ag/Ti2AlC joint; and finally, to explain the effect of Ag on the crystal structure of Ti2AlC.
was 180 MPa. Ti2AlC ceramic was tailored to rectangular specimens which are 3 4 3 mm3 in size. The brazed surface (3 4 mm2) was coarsely grounded with SiC sandpapers and then polished down to 0.5 μm diamond paste. Brazing experiments were carried out at 1030 1C for 5 min in a vacuum furnace with the vacuum level up to 1.0 10 3–3.0 10 3 Pa. The microstructure of the joint was analyzed using a scanning electron microscope (Quanta 200FEG, FEI Instruments Co., Ltd., USA) equipped with an energy-dispersive spectroscope (EDS). The phases in the joint were identified using glancing incidence X-ray diffraction with Cu-Kα radiation at a scanning speed of 0.041 step 1. Thin foil specimen for TEM observation was sliced from the brazed joint, being prepared by mechanical grinding to 40 μm dimpled down to 10 μm and then ion beam milled at 5 kV. A field emission transmission electron microscope (Tecnai G2 F30, FEI Instruments Co., Ltd., USA) was used for selected-area electron diffraction (SAED) analysis and high-resolution TEM observations. Fast Fourier Transformation (FFT) and Inverse Fast Fourier transformation (IFFT) were carried out using a digital micrograph software package (Gatan, USA). The shear strength of joint was measured by the Instron5569 electronic universal testing machine using a specially designed jig, as illustrated in Fig. 2. To minimize the effect of friction of Ti2AlC on the shear testing results, two steel rollers were set in the mold. Additionally, by adjusting the bolt, the gap width could be set accordingly for placing the Ti2AlC sample. The crosshead speed was 0.5 mm/min during testing.
2. Materials and experimental methods 3. Results and discussion The high-purity Ti2AlC ceramic was provided by the State Key Laboratory of Advanced Technology for Synthesis and Processing, Wuhan University of Technology. The microstructure of Ti2AlC ceramic is presented in Fig. 1a, with no defects like pores or cracks being observed. X-ray diffraction patterns of the Ti2AlC ceramic are shown in Fig. 1b, from which no other phases than Ti2AlC can be detected. The silver foil filler was a commercial material (99.99 wt% Ag) with a thickness of 50 μm (General Research Institute of Nonferrous Metals, Beijing, China). The average shear strength of the Ti2AlC substrate
3.1. Microstructure of the Ti2AlC/Ag/Ti2AlC brazed joint Fig. 3 shows the microstructure of a Ti2AlC–Ti2AlC joint vacuum brazed at 1030 1C for 5 min with pure-silver filler metal. Fig. 3a presents the sound bonding joint without any interlayer or defect. The brazing layer was homogeneous and mainly composed of phase with bright contrast, in which very few phases D can be found. There was a 50 μm wide area in the Ti2AlC substrate close to the brazing layer. In this area,
Fig. 1. Microstructure characterization of the pure-phase Ti2AlC ceramic. (a) Back-scattered electron (BSE) image and (b) XRD patterns for Ti2AlC.
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phase B with bright contrast (Fig. 3b) can be found existing at the grain boundaries of Ti2AlC. It is supposed that the microstructure evolution can be attributed to the interaction between Ag and Ti2AlC. Accordingly, this area is labeled as interaction area. The interface marked by the black square in Fig. 3a is enlarged and shown in Fig. 3b. The intimate contact at the interface reveals good wetting of the Ti2AlC by the Ag
Fig. 2. Schematic description of shear test configuration.
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filler metal. Table 1 gives the chemical composition of each area to confirm the possible phases marked in the joint. The result shows that the dark lamella marked by A is Ti2AlC, in which a small amount of Ag can be found. The phase marked by B is Ag [Al, Ti] solid solution. Both the tiny particles marked by C and the phases with dark contrast marked by D (in Fig. 3a) are considered as Ti2AlC. The area marked by E in the brazing layer is determined to be Ag[Al] solid solution. To further confirm the phases in the joint, the joint was cut close to the brazing layer and then grounded until the interaction area was exposed. Fig. 3c shows the XRD patterns for the interaction area. For comparison, the patterns of the pure silver and the Ti2AlC ceramic substrate are also shown. Ti2AlC and Ag were found in the XRD pattern for interaction area without any other phase being detected. By comparing these patterns, it can be seen that the peak positions of Ti2AlC reflections in the interaction area remained unchanged. In contrast to the pattern originally obtained for the Ag filler alloy, the peak positions of Ag reflections in the patterns of interaction area shifted slightly toward low angles. By combining these results with those obtained from compositional analysis, Ag is determined to exist in the form of Ag[Al, Ti] solid solution in the interaction area. Both the atomic radius of Ti (rTi ¼ 2.0 Å) and Al (rAl ¼ 1.82 Å) are a little larger than that of Ag (rAg ¼ 1.75 Å), which may be the reason for the lattice expansion. Fig. 3d shows the results of composition analysis across the scanning lines that are labeled in Fig. 3b. Because EDS has
Fig. 3. (a) Backscattered electron image of the Ti2AlC–Ti2AlC joint; (b) magnified image of the Ti2AlC/brazing layer interface; (c) XRD patterns for Ti2AlC, Ag and interaction area and (d) elemental composition fluctuation across the black line in (b).
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low sensitivity to elements lighter than Na [26], the concentration of carbon cannot be measured with sufficient degree of accuracy. In another word, the concentration fluctuation of C cannot help to understand the composition fluctuation, thus it was not presented in Fig. 3d. From the scan result, it can be seen that the Ag element concentration changed rapidly at the interface, where the Ti and Al element concentration also changed sharply. Thus, it is difficult for Ag to diffuse very deep into the Ti2AlC grains. Fig. 4a shows a TEM bright-field micrograph of the interaction area. According to SAED patterns along different zone axes, the relatively lager blocks were identified as Ti2AlC, being presented in Fig. 4b. One noteworthy phenomenon is the appearance of (0001) reflection at the extinction position of Ti2AlC. Interestingly, this phenomenon could only be found at the area in Ti2AlC substrate that was close to the Ag particles. It is supposed that the appearance of (0001) reflection might be caused by the interaction between Ti2AlC and Ag. For comparison, the SAED of Ti2AlC far away from the Ag particles is shown in Fig. 4c, in which the (0001) reflection was forbidden. Moreover, the value of the (0002) interplanar spacing (d(0002)) were measured for both the Ti2AlC close to Ag particles and the Ti2AlC far away from Ag particles. Statistics suggests that the average d(0002) of the former (0.674 nm) is larger than that of the latter (0.643 nm), which indicates that the lattice expansion along the c-axis of Ti2AlC Table 1 EDS results for various points marked in Fig. 3. Positions
A B C D E
Element (at%)
Possible phase
Ti
Al
Ag
C
49.98 04.03 54.73 52.67 01.44
24.07 09.64 20.25 17.85 11.91
01.38 86.33 4.98 07.25 86.66
24.57 – 21.39 22.23 –
Ti2AlC Ag[Al,Ti] Ti2AlC Ti2AlC Ag[Al]
may be induced by the incorporation of Ag into the Ti2AlC. The dark spots scattered in Ti2AlC along the grain boundary were Ag, and the dark triangle phase at the grain boundary was also confirmed as Ag. The strips with bright contrast lying between the Ti2AlC grains and the triangular Ag was confirmed as Ag3[Ti, Al]. The SAED of Ag3[Ti, Al] (Fig. 4e) was similar to that of Ag (Fig. 4d), only the diffraction spots such as (110) and (001) appeared at the extinction position of Ag. This might be caused by the periodically substitution of Ti and Al for Ag at the vertex position of the cubic structure of Ag crystal. In other words, Ag3[Ti,Al] is a kind of ordered phase that is formed on the basis of Ag[Ti,Al] solid solution. Ag3[Ti,Al] has the same Ll2-ordered structure as that of the (Al,Ag)3Ti, which has been systematically investigated by Oh and Han [27]. From the above results, we now know that Ag can permeate into Ti2AlC along the grain boundaries. However, it is still unclear that in what ways Ag exist in Ti2AlC. Thus, it is necessary to disclose the existence of Ag in the Ti2AlC. Fig. 5 presents the high-angle annular dark field (HADDF) scanning transmission electron microscopy (STEM) of the interaction area. From Fig. 5a we can see that this area was constituted by several Ti2AlC grains. It can be clearly found that some of the Ag with bright contrast distributed along the grain boundaries, while some other Ag scattered inside the Ti2AlC grains (labeled as Grain I & II). The magnified micrograph of Grain I is shown in Fig. 5b, in which most of the Ag distributes along a certain direction and parallel to each other. Fig. 5c shows the magnified view of Grain II, in which Ag can be found at the grain boundaries, and some tiny Ag particles of irregular shapes dispersing randomly within Ti2AlC grain. Fig. 5d displays the EDS line scanning across the areas marked by the blue solid line in Fig. 5c. It can be seen that Ti and Al concentrations increased gradually from grain boundaries to the interior of Ti2AlC grains. Correspondingly, the Ag concentration was nearly zero within Ti2AlC grains while jumped dramatically at the grain boundary and the sites of particles with bright contrast. By comparing the morphology in Fig. 5b
Fig. 4. (a) A TEM bright-field micrograph of the interaction area in the brazed Ti2AlC–Ag–Ti2AlC joint; (b)–(e) SAED patterns for the different phases in (a). The Ti2AlC of (b) is close to the Ag particles, while Ti2AlC of (c) is far away from the Ag particles.
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Fig. 5. (a) HADDF-STEM image of Ti2AlC grains in interaction area; (b) and (c) magnified image of Grain I and Grain II, respectively; (d) an EDS line scan across the grain boundary marked by blue solid line in (c); (e) and (f) SAED patterns for both Ti2AlC and Ag in (b) and (c), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
and c, we can see that Ag behaved differently in different Ti2AlC grains. To clarify this difference, SAED patterns for both Ti2AlC and Ag were taken, and the corresponding results are presented in Fig. 5e and f, respectively. As shown in Fig. 5e, the SAED pattern was indexed as Ti2AlC, while no diffraction pattern of Ag was found. This phenomenon was confirmed by SAED patterns for various similar positions in Fig. 5b. A possible explanation for this is that Ag diffuses into the crystal of Ti2AlC and substitutes for some lattice atoms, forming the Ti2AlC(Ag) solid solution. This will be further discussed in Section 3.2. In Fig. 5f, the diffraction patterns of both Ti2AlC and Ag were obtained, yet no specific crystallographic relationship was found between these two phases. 3.2. Microstructure characterization of the Ti2AlC(Ag) solid solution From the above analysis, we now know that Ag can be found both at the grain boundaries and inside the Ti2AlC grains. The grain boundaries in Ti2AlC can serve as channels for Ag to infiltrate and as locations for Ag to exist. In addition, Ag is likely to diffuse into Ti2AlC crystal, and form the Ti2AlC(Ag) solid solution. This conclusion was mentioned in the former section and demonstrated in Fig. 5. Corresponding to Fig. 5b, Fig. 6a shows a similar image of Ti2AlC, in which Ag can be seen in the form of scattered particles. Fig. 6b shows the magnified image of the region in the blue square. According to the Fast Fourier
Transform (FFT) result we know that the interface of Ti2AlC (Ag) area was parallel to the (0001) plane of the Ti2AlC grains. The yellow square in Fig. 6b is the central part of the Ti2AlC (Ag) area, of which the FFT result shows that the main diffraction pattern came from Ti2AlC. Therefore, we could know that why Ag behaved differently in different Ti2AlC grains. Firstly, at the area where Ag concentration was very low, Ag distributed in Ti2AlC in the form of the atoms rather than the individual crystal structure. The morphology of Ag might be constrained by the nano-laminated structure of Ti2AlC. Hence, the interfaces of Ag particles were parallel to the (0001) plane, which can be seen in Fig. 5b or Fig. 6. Secondly, at the area where Ag concentration was higher, its morphology could not be constrained by the structure of Ti2AlC anymore. Thus the Ag particles with cubic crystal structure (Fig. 5f) might be formed by random nucleation, namely the Ag particles of irregular shapes dispersing randomly within Ti2AlC grain (Fig. 5c). Importantly, both the existence ways of Ag in Ti2AlC had little effect on the intrinsic structure of Ti2AlC, which could be confirmed by the inserted FFT results in Fig. 5b. 3.3. Brazing mechanism and shear strength of the Ti2AlC–Ag–Ti2AlC joint Fig. 7 schematically shows the whole process of transformation in the joint, which can be described as a 2-step procedure: (1) intergranular penetration and (2) intracrystalline penetration.
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Fig. 6. (a) A TEM bright-field image of Ti2AlC grain and (b) HRTEM image of the blue square area marked in (a). Inserts are FFT results for the white square and the yellow square in (b), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 7. Schematic description for the forming process of the Ti2AlC/Ag/Ti2AlC joint. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
Firstly, as demonstrated in Fig. 7a, the parent Ti2AlC ceramic and Ag filler were assembled into sandwich structure for brazing. During the heating process, Ag atoms preferentially diffused along the Ti2AlC grain boundaries, being demonstrated in Fig. 7b. Secondly, as shown in Fig. 7c and d, with increasing the heating temperature, Ti and Al elements migrated out of Ti2AlC and dissolved into Ag, forming the Ag[Al, Ti] solid solutions both in the brazing layer and at the grain boundaries. Simultaneously, Ag diffused into the crystal structure of Ti2AlC along the passways left by the Al and Ti vacancies, giving rise to the Ti2AlC(Ag) solid solution. When the heating temperature reached the melting point of the Ag filler (about 961 1C), both the intergranular penetration and intracrystalline penetration procedures were significantly intensified. Some of the tiny grains with relatively small size were exfoliated and wrapped in the molten Ag filler, as shown in Fig. 7e. Besides that, a phase transformation was found during cooling-the Ag[Al, Ti] solid solution transformed to an ordered phase Ag3[Ti, Al], labeled as blue blocks in Fig. 7e. The mechanical properties of the joint are beneficial for understanding the brazing mechanism. Fig. 8 shows the shear strength and fracture surface of the joint brazed at 1030 1C for 5 min with pure-silver filler metal. The average shear strength of the joint was 13278 MPa, nearly 73% of that of the Ti2AlC substrate. It can be seen from Fig. 8a that the fractured surface was mainly composed of the Ti2AlC ceramic. At the lower left corner of the fracture, partial filler metal was observed. A typical morphology
of the fracture was marked by the yellow rectangular in Fig. 8a, which was magnified and demonstrated in Fig. 8b. Ti2AlC lamellae with different crystalline orientations could be found in this area, some Ti2AlC particles gathered in the gap between the Ti2AlC lamellae, and Ag existed as a kind of network structure adhering to the Ti2AlC grains. From the morphology observation we could know that the fracture occurred in the Ti2AlC ceramic adjacent to the brazing layer, namely the interaction area. This type of fracture indicates that the interfacial strength between interaction area and brazing layer was high, while the interaction area was the relatively weaker position in the joint. The decreased strength of the Ti2AlC ceramic in the interaction area might be caused by the intergranular penetration of Ag into the grain boundaries of the Ti2AlC. 4. Conclusion In this study, Ti2AlC was successfully brazed at 1030 1C for 5 min using pure-silver filler metal. The microstructure evolution, brazing mechanism and the effect of Ag on the crystal structure of Ti2AlC were investigated. Besides, the shear strength of the joint was characterized. And the following conclusions can be drawn: 1. The typical structure of the joint was Ti2AlC ceramic/ interaction area/brazing layer/interaction area/Ti2AlC ceramic. The brazing layer was mainly composed of Ag[Al]
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Fig. 8. (a) Shear strength and fractography of the Ti2AlC/Ag/Ti2AlC joint and (b) magnified view for the yellow rectangular area in (a). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
2.
3.
4.
5.
solid solution. The interaction area was mainly composed of Ti2AlC grains, the exfoliated Ti2AlC particles and the Ag [Al, Ti] solid solution at the grain boundaries of Ti2AlC. During the heating process, Ag atoms preferentially diffused into the Ti2AlC substrates along the grain boundaries. Moreover, Ti and Al elements migrated out of Ti2AlC and dissolved into Ag, forming the Ag[Al, Ti] solid solutions. Simultaneously, Ag diffused into the crystal structure of Ti2AlC along the passways left by the Al and Ti vacancies, giving rise to the Ti2AlC(Ag) solid solution. Some of the Ti2AlC grains were exfoliated and wrapped by the molten Ag. Besides, partial Ag[Al, Ti] solid solution transformed into the ordered phase Ag3[Ti,Al] during the cooling process. Ti2AlC retained its crystal structure although the Ti2AlC (Ag) solid solution were formed. In a word, neither crystal structure transformation nor decomposition of Ti2AlC has been caused by its interaction with Ag. The average shear strength of the Ti2AlC–Ag–Ti2AlC brazed joint was nearly 73% of that of the Ti2AlC substrate. The fracture in the joint occurred in the interaction area, where the strength of Ti2AlC substrate was weakened by the intergranular penetration of Ag into the grain boundaries. Both the intergranular penetration and intracrystalline penetration of Ag into Ti2AlC were expected to improve the electrical and thermal conductivity of the Ti2AlC/Ag/Ti2AlC joint. This will be further investigated experimentally and theoretically in our future study.
Acknowledgments The authors thank Ms. Y. Wang for revising the language. This work was supported by the National Natural Science Foundation of China under the number of 51372049 and 51321061. References [1] M.W. Barsoum, The MN þ 1AXN phases: a new class of solids: thermodynamically stable nanolaminates, Prog. Solid State Chem. 28 (2000) 201–281.
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CERAMICS INTERNATIONAL
Ceramics International 41 (2015) 8203–8210 www.elsevier.com/locate/ceramint
Microstructure evolution and brazing mechanism of Ti2AlC–Ti2AlC joint by using pure-silver filler metal Guochao Wang, Guohua Fan, Jie Zhangn, Tianpeng Wang, Xiaowen Liu School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, China Received 11 January 2015; received in revised form 7 March 2015; accepted 8 March 2015 Available online 14 March 2015
Abstract This paper investigated the microstructure evolution of a Ti2AlC–Ti2AlC joint brazed using pure-silver filler metal at 1030 1C for 5 min, characterized the newly formed products in the joint and elucidated their formation mechanism. Furthermore, it reveals the effect of Ag on the crystal structure of Ti2AlC. Ti2AlC is perfectly bonded by using Ag filler, with no crack or pore being observed. During brazing, the molten Ag filler metal infiltrates into the Ti2AlC along the grain boundaries. Al and Ti migrate out of Ti2AlC and dissolve into the Ag, forming the Ag[Al, Ti] solid solution. Ag3[Ti, Al] is produced on the basis of Ag[Al, Ti] through ordering transition during the cooling process. Simultaneously, Ag element diffuses into the crystal structure of Ti2AlC. Nevertheless, the Ti2AlC retains its crystal structure under the partial loss of Al and Ti. The shear strength of the joint was 13278 MPa, nearly 73% of that of the Ti2AlC substrate. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: Ti2AlC; Ag; Microstructure; Brazing; Shear strength
1. Introduction In recent decades, the nano-laminated ternary compounds with a formula of Mn þ 1AXn (where M is an early transition metal, A is an A-group (mostly groups of 13 or 14) element, and X is C and/or N) are intensively studied. The so-called MAX phases have layered hexagonal structure with space group P63/mmc. Nearclose-packed Mn þ 1Xn layers are interleaved with hexagonal nets of A element. The unique structure endows the compounds with an unusual combination of metallic and ceramic behaviors [1–3]. As a fascinating member of this series, Ti2AlC is of particular interest since this compound has the lowest density (4.11 g cm 3) and the best oxidation resistance in the family of MAX phases [4–6]. Besides, it has high Young's modulus [7], good thermal and electrical conductivity [8,9], self-lubricating function [10–13], high thermal stability and excellent mechanical properties in the radiation environment [6], which make it a suitable candidate for many structural and functional applications, such as sliding
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Corresponding author. Tel./fax: þ 86 451 86414234. E-mail address: [email protected] (J. Zhang).
http://dx.doi.org/10.1016/j.ceramint.2015.03.040 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
electrical contact, conducting ceramic in harsh environment and reinforcement for soft metals [8,10–13]. Similar to most other ceramics, large-dimensioned Ti2AlC ceramic is difficult to synthesize due to its narrow phase range in the Ti–Al–C ternary phase diagram [14]. This problem is usually solved by joining ceramics, a method which allows the large, complex, multifunctional ceramic components to be manufactured through the controlled integration of smaller, simpler, more easily manufactured parts [15]. Therefore, studies devoted to joining Ti2AlC are significant for broadening their applications. In the past decades, researchers have been able to join Ti–Al–C and Ti–Si–C ceramics due to their structural characteristics and the ease with which Al3O2 or Ti3Al(Si)C2 solid solutions are formed [16–20]. These joints are obtained by solid phase diffusion bonding, by which the crystal structure of the ceramic matrix can be well sustained. In our previous research [21,22], brazing was adopted to join MAX phases and metallic parts because of its simplicity, good repetitiveness and excellent adaptability of joint size and shape. Ti2AlC/Cu brazed joint can be successfully fabricated using Ag–Cu eutectic filler alloy. Due to the weak Ti–Al bond in Ti2AlC, Al was apt to be divorced from the crystal structure and react with Cu. As one of the substrates to be joined, the parent copper partially
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dissolved into the molten Ag–Cu filler during brazing. The reaction was intensified by increasing the content of Cu in the filler alloy, resulting in the further out-migration of Al and Ti from Ti2AlC. After a certain extent of reaction was reached, the crystal structure of Ti2AlC would decompose, leading to the change in intrinsic properties of the phases. It was also found that Ag is much less reactive to Ti2AlC comparing with Cu. No stable compound can be found in the Ag–Al binary phase diagram, which is an indication of a much weaker tendency in Ag–Al reaction [23]. Therefore, puresilver filler was applied to join the Ti2AlC ceramics in this study. Wang et al. proposed that the decomposition of Ti2AlC is triggered by the outward diffusion of Al atoms [24]. It was also calculated that Al vacancy shows lower formation energy of mono-vacancy than that of Ti and C vacancy. Furthermore, defective Ti2AlC may maintain phase stability down to a substoichiometry of Ti2Al0.5C. The structural stability of Ti2AlC is determined by its depletion of Al at high temperatures in an oxygen-containing atmosphere or in species having a high affinity to Al [25]. Therefore, the structural stability of Ti2AlC should be concerned during the brazing of Ti2AlC. The purpose of this work is threefold: firstly, to obtain the well-bonded joint of Ti2AlC and clarify the microstructure evolution of the joint during brazing; secondly, to characterize the microstructure and elucidate the brazing mechanism of the Ti2AlC/Ag/Ti2AlC joint; and finally, to explain the effect of Ag on the crystal structure of Ti2AlC.
was 180 MPa. Ti2AlC ceramic was tailored to rectangular specimens which are 3 4 3 mm3 in size. The brazed surface (3 4 mm2) was coarsely grounded with SiC sandpapers and then polished down to 0.5 μm diamond paste. Brazing experiments were carried out at 1030 1C for 5 min in a vacuum furnace with the vacuum level up to 1.0 10 3–3.0 10 3 Pa. The microstructure of the joint was analyzed using a scanning electron microscope (Quanta 200FEG, FEI Instruments Co., Ltd., USA) equipped with an energy-dispersive spectroscope (EDS). The phases in the joint were identified using glancing incidence X-ray diffraction with Cu-Kα radiation at a scanning speed of 0.041 step 1. Thin foil specimen for TEM observation was sliced from the brazed joint, being prepared by mechanical grinding to 40 μm dimpled down to 10 μm and then ion beam milled at 5 kV. A field emission transmission electron microscope (Tecnai G2 F30, FEI Instruments Co., Ltd., USA) was used for selected-area electron diffraction (SAED) analysis and high-resolution TEM observations. Fast Fourier Transformation (FFT) and Inverse Fast Fourier transformation (IFFT) were carried out using a digital micrograph software package (Gatan, USA). The shear strength of joint was measured by the Instron5569 electronic universal testing machine using a specially designed jig, as illustrated in Fig. 2. To minimize the effect of friction of Ti2AlC on the shear testing results, two steel rollers were set in the mold. Additionally, by adjusting the bolt, the gap width could be set accordingly for placing the Ti2AlC sample. The crosshead speed was 0.5 mm/min during testing.
2. Materials and experimental methods 3. Results and discussion The high-purity Ti2AlC ceramic was provided by the State Key Laboratory of Advanced Technology for Synthesis and Processing, Wuhan University of Technology. The microstructure of Ti2AlC ceramic is presented in Fig. 1a, with no defects like pores or cracks being observed. X-ray diffraction patterns of the Ti2AlC ceramic are shown in Fig. 1b, from which no other phases than Ti2AlC can be detected. The silver foil filler was a commercial material (99.99 wt% Ag) with a thickness of 50 μm (General Research Institute of Nonferrous Metals, Beijing, China). The average shear strength of the Ti2AlC substrate
3.1. Microstructure of the Ti2AlC/Ag/Ti2AlC brazed joint Fig. 3 shows the microstructure of a Ti2AlC–Ti2AlC joint vacuum brazed at 1030 1C for 5 min with pure-silver filler metal. Fig. 3a presents the sound bonding joint without any interlayer or defect. The brazing layer was homogeneous and mainly composed of phase with bright contrast, in which very few phases D can be found. There was a 50 μm wide area in the Ti2AlC substrate close to the brazing layer. In this area,
Fig. 1. Microstructure characterization of the pure-phase Ti2AlC ceramic. (a) Back-scattered electron (BSE) image and (b) XRD patterns for Ti2AlC.
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phase B with bright contrast (Fig. 3b) can be found existing at the grain boundaries of Ti2AlC. It is supposed that the microstructure evolution can be attributed to the interaction between Ag and Ti2AlC. Accordingly, this area is labeled as interaction area. The interface marked by the black square in Fig. 3a is enlarged and shown in Fig. 3b. The intimate contact at the interface reveals good wetting of the Ti2AlC by the Ag
Fig. 2. Schematic description of shear test configuration.
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filler metal. Table 1 gives the chemical composition of each area to confirm the possible phases marked in the joint. The result shows that the dark lamella marked by A is Ti2AlC, in which a small amount of Ag can be found. The phase marked by B is Ag [Al, Ti] solid solution. Both the tiny particles marked by C and the phases with dark contrast marked by D (in Fig. 3a) are considered as Ti2AlC. The area marked by E in the brazing layer is determined to be Ag[Al] solid solution. To further confirm the phases in the joint, the joint was cut close to the brazing layer and then grounded until the interaction area was exposed. Fig. 3c shows the XRD patterns for the interaction area. For comparison, the patterns of the pure silver and the Ti2AlC ceramic substrate are also shown. Ti2AlC and Ag were found in the XRD pattern for interaction area without any other phase being detected. By comparing these patterns, it can be seen that the peak positions of Ti2AlC reflections in the interaction area remained unchanged. In contrast to the pattern originally obtained for the Ag filler alloy, the peak positions of Ag reflections in the patterns of interaction area shifted slightly toward low angles. By combining these results with those obtained from compositional analysis, Ag is determined to exist in the form of Ag[Al, Ti] solid solution in the interaction area. Both the atomic radius of Ti (rTi ¼ 2.0 Å) and Al (rAl ¼ 1.82 Å) are a little larger than that of Ag (rAg ¼ 1.75 Å), which may be the reason for the lattice expansion. Fig. 3d shows the results of composition analysis across the scanning lines that are labeled in Fig. 3b. Because EDS has
Fig. 3. (a) Backscattered electron image of the Ti2AlC–Ti2AlC joint; (b) magnified image of the Ti2AlC/brazing layer interface; (c) XRD patterns for Ti2AlC, Ag and interaction area and (d) elemental composition fluctuation across the black line in (b).
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low sensitivity to elements lighter than Na [26], the concentration of carbon cannot be measured with sufficient degree of accuracy. In another word, the concentration fluctuation of C cannot help to understand the composition fluctuation, thus it was not presented in Fig. 3d. From the scan result, it can be seen that the Ag element concentration changed rapidly at the interface, where the Ti and Al element concentration also changed sharply. Thus, it is difficult for Ag to diffuse very deep into the Ti2AlC grains. Fig. 4a shows a TEM bright-field micrograph of the interaction area. According to SAED patterns along different zone axes, the relatively lager blocks were identified as Ti2AlC, being presented in Fig. 4b. One noteworthy phenomenon is the appearance of (0001) reflection at the extinction position of Ti2AlC. Interestingly, this phenomenon could only be found at the area in Ti2AlC substrate that was close to the Ag particles. It is supposed that the appearance of (0001) reflection might be caused by the interaction between Ti2AlC and Ag. For comparison, the SAED of Ti2AlC far away from the Ag particles is shown in Fig. 4c, in which the (0001) reflection was forbidden. Moreover, the value of the (0002) interplanar spacing (d(0002)) were measured for both the Ti2AlC close to Ag particles and the Ti2AlC far away from Ag particles. Statistics suggests that the average d(0002) of the former (0.674 nm) is larger than that of the latter (0.643 nm), which indicates that the lattice expansion along the c-axis of Ti2AlC Table 1 EDS results for various points marked in Fig. 3. Positions
A B C D E
Element (at%)
Possible phase
Ti
Al
Ag
C
49.98 04.03 54.73 52.67 01.44
24.07 09.64 20.25 17.85 11.91
01.38 86.33 4.98 07.25 86.66
24.57 – 21.39 22.23 –
Ti2AlC Ag[Al,Ti] Ti2AlC Ti2AlC Ag[Al]
may be induced by the incorporation of Ag into the Ti2AlC. The dark spots scattered in Ti2AlC along the grain boundary were Ag, and the dark triangle phase at the grain boundary was also confirmed as Ag. The strips with bright contrast lying between the Ti2AlC grains and the triangular Ag was confirmed as Ag3[Ti, Al]. The SAED of Ag3[Ti, Al] (Fig. 4e) was similar to that of Ag (Fig. 4d), only the diffraction spots such as (110) and (001) appeared at the extinction position of Ag. This might be caused by the periodically substitution of Ti and Al for Ag at the vertex position of the cubic structure of Ag crystal. In other words, Ag3[Ti,Al] is a kind of ordered phase that is formed on the basis of Ag[Ti,Al] solid solution. Ag3[Ti,Al] has the same Ll2-ordered structure as that of the (Al,Ag)3Ti, which has been systematically investigated by Oh and Han [27]. From the above results, we now know that Ag can permeate into Ti2AlC along the grain boundaries. However, it is still unclear that in what ways Ag exist in Ti2AlC. Thus, it is necessary to disclose the existence of Ag in the Ti2AlC. Fig. 5 presents the high-angle annular dark field (HADDF) scanning transmission electron microscopy (STEM) of the interaction area. From Fig. 5a we can see that this area was constituted by several Ti2AlC grains. It can be clearly found that some of the Ag with bright contrast distributed along the grain boundaries, while some other Ag scattered inside the Ti2AlC grains (labeled as Grain I & II). The magnified micrograph of Grain I is shown in Fig. 5b, in which most of the Ag distributes along a certain direction and parallel to each other. Fig. 5c shows the magnified view of Grain II, in which Ag can be found at the grain boundaries, and some tiny Ag particles of irregular shapes dispersing randomly within Ti2AlC grain. Fig. 5d displays the EDS line scanning across the areas marked by the blue solid line in Fig. 5c. It can be seen that Ti and Al concentrations increased gradually from grain boundaries to the interior of Ti2AlC grains. Correspondingly, the Ag concentration was nearly zero within Ti2AlC grains while jumped dramatically at the grain boundary and the sites of particles with bright contrast. By comparing the morphology in Fig. 5b
Fig. 4. (a) A TEM bright-field micrograph of the interaction area in the brazed Ti2AlC–Ag–Ti2AlC joint; (b)–(e) SAED patterns for the different phases in (a). The Ti2AlC of (b) is close to the Ag particles, while Ti2AlC of (c) is far away from the Ag particles.
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Fig. 5. (a) HADDF-STEM image of Ti2AlC grains in interaction area; (b) and (c) magnified image of Grain I and Grain II, respectively; (d) an EDS line scan across the grain boundary marked by blue solid line in (c); (e) and (f) SAED patterns for both Ti2AlC and Ag in (b) and (c), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
and c, we can see that Ag behaved differently in different Ti2AlC grains. To clarify this difference, SAED patterns for both Ti2AlC and Ag were taken, and the corresponding results are presented in Fig. 5e and f, respectively. As shown in Fig. 5e, the SAED pattern was indexed as Ti2AlC, while no diffraction pattern of Ag was found. This phenomenon was confirmed by SAED patterns for various similar positions in Fig. 5b. A possible explanation for this is that Ag diffuses into the crystal of Ti2AlC and substitutes for some lattice atoms, forming the Ti2AlC(Ag) solid solution. This will be further discussed in Section 3.2. In Fig. 5f, the diffraction patterns of both Ti2AlC and Ag were obtained, yet no specific crystallographic relationship was found between these two phases. 3.2. Microstructure characterization of the Ti2AlC(Ag) solid solution From the above analysis, we now know that Ag can be found both at the grain boundaries and inside the Ti2AlC grains. The grain boundaries in Ti2AlC can serve as channels for Ag to infiltrate and as locations for Ag to exist. In addition, Ag is likely to diffuse into Ti2AlC crystal, and form the Ti2AlC(Ag) solid solution. This conclusion was mentioned in the former section and demonstrated in Fig. 5. Corresponding to Fig. 5b, Fig. 6a shows a similar image of Ti2AlC, in which Ag can be seen in the form of scattered particles. Fig. 6b shows the magnified image of the region in the blue square. According to the Fast Fourier
Transform (FFT) result we know that the interface of Ti2AlC (Ag) area was parallel to the (0001) plane of the Ti2AlC grains. The yellow square in Fig. 6b is the central part of the Ti2AlC (Ag) area, of which the FFT result shows that the main diffraction pattern came from Ti2AlC. Therefore, we could know that why Ag behaved differently in different Ti2AlC grains. Firstly, at the area where Ag concentration was very low, Ag distributed in Ti2AlC in the form of the atoms rather than the individual crystal structure. The morphology of Ag might be constrained by the nano-laminated structure of Ti2AlC. Hence, the interfaces of Ag particles were parallel to the (0001) plane, which can be seen in Fig. 5b or Fig. 6. Secondly, at the area where Ag concentration was higher, its morphology could not be constrained by the structure of Ti2AlC anymore. Thus the Ag particles with cubic crystal structure (Fig. 5f) might be formed by random nucleation, namely the Ag particles of irregular shapes dispersing randomly within Ti2AlC grain (Fig. 5c). Importantly, both the existence ways of Ag in Ti2AlC had little effect on the intrinsic structure of Ti2AlC, which could be confirmed by the inserted FFT results in Fig. 5b. 3.3. Brazing mechanism and shear strength of the Ti2AlC–Ag–Ti2AlC joint Fig. 7 schematically shows the whole process of transformation in the joint, which can be described as a 2-step procedure: (1) intergranular penetration and (2) intracrystalline penetration.
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Fig. 6. (a) A TEM bright-field image of Ti2AlC grain and (b) HRTEM image of the blue square area marked in (a). Inserts are FFT results for the white square and the yellow square in (b), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 7. Schematic description for the forming process of the Ti2AlC/Ag/Ti2AlC joint. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
Firstly, as demonstrated in Fig. 7a, the parent Ti2AlC ceramic and Ag filler were assembled into sandwich structure for brazing. During the heating process, Ag atoms preferentially diffused along the Ti2AlC grain boundaries, being demonstrated in Fig. 7b. Secondly, as shown in Fig. 7c and d, with increasing the heating temperature, Ti and Al elements migrated out of Ti2AlC and dissolved into Ag, forming the Ag[Al, Ti] solid solutions both in the brazing layer and at the grain boundaries. Simultaneously, Ag diffused into the crystal structure of Ti2AlC along the passways left by the Al and Ti vacancies, giving rise to the Ti2AlC(Ag) solid solution. When the heating temperature reached the melting point of the Ag filler (about 961 1C), both the intergranular penetration and intracrystalline penetration procedures were significantly intensified. Some of the tiny grains with relatively small size were exfoliated and wrapped in the molten Ag filler, as shown in Fig. 7e. Besides that, a phase transformation was found during cooling-the Ag[Al, Ti] solid solution transformed to an ordered phase Ag3[Ti, Al], labeled as blue blocks in Fig. 7e. The mechanical properties of the joint are beneficial for understanding the brazing mechanism. Fig. 8 shows the shear strength and fracture surface of the joint brazed at 1030 1C for 5 min with pure-silver filler metal. The average shear strength of the joint was 13278 MPa, nearly 73% of that of the Ti2AlC substrate. It can be seen from Fig. 8a that the fractured surface was mainly composed of the Ti2AlC ceramic. At the lower left corner of the fracture, partial filler metal was observed. A typical morphology
of the fracture was marked by the yellow rectangular in Fig. 8a, which was magnified and demonstrated in Fig. 8b. Ti2AlC lamellae with different crystalline orientations could be found in this area, some Ti2AlC particles gathered in the gap between the Ti2AlC lamellae, and Ag existed as a kind of network structure adhering to the Ti2AlC grains. From the morphology observation we could know that the fracture occurred in the Ti2AlC ceramic adjacent to the brazing layer, namely the interaction area. This type of fracture indicates that the interfacial strength between interaction area and brazing layer was high, while the interaction area was the relatively weaker position in the joint. The decreased strength of the Ti2AlC ceramic in the interaction area might be caused by the intergranular penetration of Ag into the grain boundaries of the Ti2AlC. 4. Conclusion In this study, Ti2AlC was successfully brazed at 1030 1C for 5 min using pure-silver filler metal. The microstructure evolution, brazing mechanism and the effect of Ag on the crystal structure of Ti2AlC were investigated. Besides, the shear strength of the joint was characterized. And the following conclusions can be drawn: 1. The typical structure of the joint was Ti2AlC ceramic/ interaction area/brazing layer/interaction area/Ti2AlC ceramic. The brazing layer was mainly composed of Ag[Al]
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Fig. 8. (a) Shear strength and fractography of the Ti2AlC/Ag/Ti2AlC joint and (b) magnified view for the yellow rectangular area in (a). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
2.
3.
4.
5.
solid solution. The interaction area was mainly composed of Ti2AlC grains, the exfoliated Ti2AlC particles and the Ag [Al, Ti] solid solution at the grain boundaries of Ti2AlC. During the heating process, Ag atoms preferentially diffused into the Ti2AlC substrates along the grain boundaries. Moreover, Ti and Al elements migrated out of Ti2AlC and dissolved into Ag, forming the Ag[Al, Ti] solid solutions. Simultaneously, Ag diffused into the crystal structure of Ti2AlC along the passways left by the Al and Ti vacancies, giving rise to the Ti2AlC(Ag) solid solution. Some of the Ti2AlC grains were exfoliated and wrapped by the molten Ag. Besides, partial Ag[Al, Ti] solid solution transformed into the ordered phase Ag3[Ti,Al] during the cooling process. Ti2AlC retained its crystal structure although the Ti2AlC (Ag) solid solution were formed. In a word, neither crystal structure transformation nor decomposition of Ti2AlC has been caused by its interaction with Ag. The average shear strength of the Ti2AlC–Ag–Ti2AlC brazed joint was nearly 73% of that of the Ti2AlC substrate. The fracture in the joint occurred in the interaction area, where the strength of Ti2AlC substrate was weakened by the intergranular penetration of Ag into the grain boundaries. Both the intergranular penetration and intracrystalline penetration of Ag into Ti2AlC were expected to improve the electrical and thermal conductivity of the Ti2AlC/Ag/Ti2AlC joint. This will be further investigated experimentally and theoretically in our future study.
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