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C composites with high temperature TiSi eutectic alloy
Accepted Manuscript Pre-infiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy Zongjing He, Chun Li, Junlei Qi, Yongxian Huang, Jicai Feng, Jian Cao PII:
S0008-6223(18)30751-6
DOI:
10.1016/j.carbon.2018.08.021
Reference:
CARBON 13376
To appear in:
Carbon
Received Date: 27 May 2018 Revised Date:
22 July 2018
Accepted Date: 9 August 2018
Please cite this article as: Zongjing He, Chun Li, Junlei Qi, Yongxian Huang, Jicai Feng, Jian Cao, Preinfiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy, Carbon (2018), doi: 10.1016/j.carbon.2018.08.021 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Pre-infiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy Zongjing He, Chun Li, Junlei Qi*, Yongxian Huang, Jicai Feng, Jian Cao*
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State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
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Corresponding authors. Tel: +86-451-86418882. E-mail: [email protected] (J. Cao); Tel: +86-451-86418882. E-mail: [email protected] (J. Qi)
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Pre-infiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy Zongjing He, Chun Li, Junlei Qi*, Yongxian Huang, Jicai Feng, Jian Cao*
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State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
Abstract
A novel method is developed to join carbon fiber reinforced carbon composites (Cf/C
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composites) via a two-step approach using Ti14Si86 metal foil as the joining material. The
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infiltration behavior, joint microstructure, interface reaction, fracture characteristics and joining mechanism are investigated. The preparation of the almost compact deposit layer and the infiltration carbide can immediately heal the pores of Cf/C composites and form the gradient composite structure, which can effectively minimize the impact of the coefficient of
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thermal expansion (CTE) mismatch between the outer reaction layer and the carbon matrix and avoid the excessive filler infiltration in the brazing step. The brazing process reveals that
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the growing SiC interface and Ti-Si eutectic structure in the seam also behave as a graded CTE composite structure, which is beneficial for strengthening the joint. Compared with
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direct brazing, this method can obtain more reliable Cf/C composites joints due to strong interfacial bonding and the gradient structure. As a result, a favorable shear strength of 26 MPa is achieved within 10 min at a brazing temperature of 1400 °C. 1. Introduction Carbon fiber reinforced carbon composites are attractive materials in aerospace
*
Corresponding authors. Tel: +86-451-86418882. E-mail address: [email protected] (J. Cao). Tel: +86-451-86418882. E-mail address: [email protected] (J. L. Qi). 1
ACCEPTED MANUSCRIPT application such as space vehicle heat shields and aircraft brakes due to low thermal expansion and outstanding high-temperature mechanical properties at considerably reduced weight [1-5]. Along with extremely low neutron activation and low atomic number, great
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interest is drawn to the application of Cf/C composites in nuclear fusion reactors [6]. Cf/C composites are usually required to join to other substrates or to themselves to be fabricated into complex assemblies [7,8]. Among the current joining technologies, brazing has been
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extensively used due to its simplicity, high joint strength and excellent adaptability of sample
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configuration [9-12]. But the major problem in brazing Cf/C composites lies in its limited application at relatively low temperature [7,13,14]. It is essential to develop a new braze filler and improve high-temperature joint properties.
The judicious use of braze chemistry and application of innovative brazing strategies
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should combine reliable interfacial bonding and structure with graded CTE. Both Ti and Si elements show perfect wetting and adherence to the carbon material and equilibrium contact angles could reach near 0° in short brazing time [15,16]. Moreover, the reaction products
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between Ti, Si and carbon show excellent high temperature resistance [17]. In addition, Ti
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and Si are low-activation elements, which have shown great potential in thermonuclear fusion applications [6,18,19]. Dadras et al. [20] brazed Cf/C composites using TiSi2 and the achieved joint showed a maximum shear strength of 34.4 MPa when testing at 1164 °C, which indicates that Ti-Si alloy has potential applications in high temperature. With relatively low melting point and perfect fluidity, Ti-Si eutectic metal is much more beneficial to filler metal spreading and has exhibited good adhesion with silicon carbide [21]. However, the brazing of Cf/C composites with Ti-Si eutectic alloy has not been reported. In addition, the authors have 2
ACCEPTED MANUSCRIPT carried out some investigation of joining Cf/C composites. Because of the abundant porous structure of the low-density substrate, all attempts to directly braze the Cf/C joints have failed. Thus a necessary pretreatment of the Cf/C composites should be considered before brazing
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process. Chemical vapor infiltration (CVI), precursor infiltration and pyrolysis (PIP), slurry infiltration (SI) and reactive melt infiltration (RMI) have been applied to densify or modify
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fiber-reinforced ceramic-matrix composites (CMCs) or Cf/C composites [22-26]. SiC
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transition layer was prepared through CVI, chemical vapor deposition (CVD) and pack cementation processes to form a zigzag interface structure of Cf/C substrate to improve the poor wettability of LAS glass ceramic [27,28]. However, no practical methods can be found to moderate the residual stress in high-temperature brazed Cf/C joints. Compared with other
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technologies, RMI is an attractive modification method to introduce ceramics because of its time-saving preparation and relatively low cost [23,29,30]. Considering the relatively low density, small coefficient of thermal expansion as well as high working temperature in air of
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SiC [25,31], a novel method combined brazing with RMI was used to design a complex
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structure with graded CTE to mitigate residual stress among the joints. Furthermore, multiple effects of the simple method such as pinning the substrate and strengthening the Cf/C composites can also be expected. Taking advantage of the perfect wettability of the high-temperature Ti14Si86 eutectic metal and the porous structure of the substrate, a novel brazing method was developed to join Cf/C composites in argon gas atmosphere without extra pressure. The infiltration behavior and braze ability of Ti-Si for Cf/C composites were studied in this work. The joint 3
ACCEPTED MANUSCRIPT microstructure, shear fracture characteristics and interface bonding mechanism of the resulting joints were investigated in detail. 2. Experimental
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2.1. Materials Fig. 1(a) and (b) show the microstructure of the as-received Cf/C composites and Ti-Si eutectic alloy. The Cf/C composite used in this study had an open porosity of about 9%, and
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both the pores in the pyrocarbon and bundles could be observed, as shown in Fig. 1(a). The
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density and three point bending strength of the substrate were about 1.64 g/cm3 and 52.5±3 MPa, respectively. The Cf/C composites were cut into dimension of 5 mm × 5 mm × 5 mm for metallographic observation and 5 mm × 5 mm × 10 mm for shear tests. The surfaces perpendicular to the fiber planes were used as the joining planes. The eutectic filler alloy was
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prepared using the raw materials of sponge Ti (99.99 wt.% purity) and Si (99.99 wt.% purity) by vacuum arc-melting technology. As shown in Fig. 1(b), the alloy was homogeneous and mainly consisted of two phases, which were confirmed to be Si matrix phase (the gray phase)
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and TiSi2 intermetallic (the white phase) by X-ray diffraction (XRD) (Fig. 1(c)). The filler
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alloy was cut into 5 mm × 5 mm × 300 µm and 5 mm × 5 mm × 900 µm pieces, then ground by 800-grit silicon carbide paper. All the Cf/C composites and metal foils were ultrasonic cleaned in absolute ethanol for 5 min prior to brazing. 2.2. Pre-infiltration and brazing processes In this work, a novel two-step brazing method was designed to join the Cf/C composites. Fig. 1(d) and (e) show the schematics of the melt infiltration process and the brazing process. For the melt infiltration process, firstly, Ti-Si eutectic alloy foils of 300µm were placed on the 4
ACCEPTED MANUSCRIPT surface of the raw Cf/C composites (the area of 5 mm × 5 mm) used for brazing. Afterwards, the pre-treated assembly was heated to the holding temperature (1380 °C, 1400 °C, 1420 °C, 1440 °C) at a rate of 5 °C/min and held for 10 min, then slowly cooled down to room
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temperature. The brazing step was performed under the same heating cycle as the first step, and brazing assembly was like a sandwich structure with the filler alloy (300 µm) between two pre-infiltrated Cf/C substrates (PI-Cf/C). Particularly, the direct brazing of the Cf/C
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composites (using filler alloy of 900 µm) without per-infiltration process was served as a
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contrast experiment. The infiltration and brazing processes were carried out in argon atmospheric and no additional pressure was applied on each assembly except the dead weight.
2.3. Microstructural analysis and property measurement
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The microstructure of the PI-Cf/C samples and brazed joints was characterized by field emission scanning electron microscope (FE-SEM, Quanta 200FEG) coupled with energy dispersive spectroscope (EDS, INCA Energy 300) and transmission electron microscope
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(TEM, Talos F200X). The phases in the samples were identified by X-ray diffraction analysis
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(XRD, Philips X’Pert). The shear strength of the brazed joints at room temperature was measured through universal testing machine (Instron 5569) at a constant speed of 0.5mm/min, the schematic diagram is shown in Fig. 1(f). With the same method, the high temperature shear test was conducted at 1200 °C by radiation heating in the argon flow. Nano-indentation test (NanoIndenter G200) was also used to evaluate the mechanical behaviors of the brazed joints. The fracture characteristics was examined to determine the fracture propagation by SEM and XRD. Each value reported is the average of at least 10 tests. 5
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Fig. 1. The micrographs of (a) the Cf/C composites and (b) the Ti-Si alloy. (c) The XRD pattern of the Ti-Si alloy. The schematic diagrams of (d) the infiltration process, (e) the brazing process and (f) the shear test.
3. Results and discussion
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3.1. Analysis of infiltration behaviors
Fig. 2((a)-(c)) show the microstructure of the PI-Cf/C composites. As can be seen from
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Fig. 2(a), both the pores in bundles and among the pyrocarbon of the Cf/C substrate have been completely filled up with Ti-Si alloy [32], which might be attributed to the infiltration of
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melt. The transport of liquid melt through the pores is governed by capillarity and the resistance of the viscous forces, which can be represented by the modified Washburn model [33-35]:
2 3 B rt- At 2 3 A= 2D M
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C Si 一 Ti
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Si 一 Ti C
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(3-1) (3-2)
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Cσcosθ 2µ
(3-3)
where h is the infiltration height, t is the annealing time, µ is the viscosity, σ is the surface tension, θ is the contact angle-about 0° according to sessile drop technique, r is the
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mean pore radius, C is the tortuosity factor, D is the effective diffusivity of carbon, Mi and ρi are the atomic weight and density of component i. Eq. (3-1) and (3-2) indicate that the infiltration behavior depends on two factors: chemical reaction and properties of the liquid
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melt. In this research, A and B can be assumed as constants in Eq. (3-1). Hence, the
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infiltration height (h) depends on the annealing time (t) and the pore dimension (r), and the trend graph can be drawn as Fig. 2(e). As shown in Fig. 2(e), the X-direction represents the annealing time at the infiltration temperature; and the Y-direction represents the infiltration height. It can be concluded that carbides formed at the small opening pores would seal the
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pores immediately at the beginning of infiltration process [35]. When the melt metal infiltrates into the bigger opening pores, not only carbide layer will be formed during the interface reaction, but also a flow of liquid can be remained through the center of pores. The
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combined effects of both the chemical reaction and the melt infiltration give rise to the results
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that the initial infiltration of the melt metal is rapid and then slow down with the annealing time increasing. The height of the melt in the large opening pores could be higher than that in the small opening pores. R. Asthana [36-38] reported a new reactive infiltration model characterized by pore shrinkage and decaying contact angle due to interphase formation, which could qualitatively explain the sudden drop in the penetration velocity with increasing distance in Si/porous carbon system. The dynamic model in the articles could better predict the actual penetration lengths for Si/C samples than the other analyses. And our function 7
ACCEPTED MANUSCRIPT curves of infiltration height as annealing time agree well with the theoretical profiles for Si infiltration of a carbon capillary [36]. The infiltration cell was composed of pores along carbon fiber bundles and pores within
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pyrocarbon [39]. The dimensions of residual spaces in the interspaces were about 100 µm. During the infiltration process, the molten metal of titanium-silicon was drawn into the residual spaces and reacted with the previously deposited pyrocarbon to form the SiC layer.
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The melt penetrated into the Cf/C composites and the penetration depth depends on the types
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and dimensions of the pores [3]. It can be found in Fig. 2(b) that the molten metal of titanium-silicon infiltrated along the carbon fiber bundles and reacted with carbon fiber to form silicon carbide. In contrast, the interspaces among the pyrocarbon were too large to be completely filled by the Ti-Si melt and residual unreacted Ti-Si alloy could be found in the
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center of the pores (Fig. 2(c)). Therefore, as shown in the schematic diagram of Fig. 2(d), the molten metal of Ti-Si penetrated into the pores of the substrate and formed a tight combination at the interface. From the insert in Fig. 2(a), it can be seen that the hole in the
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substrate was not filled up by the filler alloy when a dense deposit layer has already formed,
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which could prevent further infiltration of the braze filler in the second step. The interlaced infiltration across the whole substrate could effectively heal the pores of Cf/C composites [40], which could avoid excessive braze filler infiltration and strengthen the Cf/C composites. The coefficient of thermal expansion of SiC is 4.4×10-6 K-1 [19], higher than that of Cf/C composites (2×10-6 K-1) [41] and Si (3.5×10-6 K-1) [42], but lower than TiSi2 (10.4×10-6 K-1) [43]. This gradient structure could effectively minimize the impact of the mismatch in thermal expansion coefficient between the reaction layer and the carbon matrix, which could 8
ACCEPTED MANUSCRIPT help to improve the joint strength during the second step of brazing process. Zeng et al. found the release of residual thermal stress in Zr-Ti infiltrated porous Cf/C composites using Raman
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spectroscopy analysis [44].
Fig. 2. (a), (b), (c) The micrographs and (d) the schematic diagram of the PI-Cf/C composites by Ti-Si alloy.
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(e) The function curves of infiltration height as annealing time.
Fig. 3 shows the typical interfacial microstructure of the PI-Cf/C composites under different temperatures. Fig. 3(a) indicates that there were three different phases distributed in
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solidified melt. EDS results reveal that the three phases were Si (Si: 99.85 at.%, Ti: 0.15
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at.%), TiSi2 (Si: 66.27 at.%, Ti: 33.73 at.%) and SiC (Si: 48.96 at.%, C: 51.04 at.%), respectively. The existence of Si element reflected an inadequate reaction between Si from the Ti-Si eutectic foil and Cf/C composites at such a low temperature. When the melt infiltrating temperature was increased to 1400 °C or above, the content of Si and TiSi2 rapidly decreased (Fig. 3(b)) and the flaky phase was exposed. It should be Ti3SiC2 according to EDS results (Ti: 48.59 at.%, Si: 17.26 at.%, C: 34.15 at.%). As the temperature was increased to 1440 °C, the new spherical-like phase was identified which should be further confirmed. The 9
ACCEPTED MANUSCRIPT morphology changes during pre-infiltration process can also be visually observed by the inserts in Fig. 3. The well-combined metallization layer got thinning as the temperature was increased, reflecting the fine wettability and permeability of Ti14Si86 metal foil on Cf/C
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composites. The infiltration depth could be higher than low temperatures due to the changing of series of factors such as carbon diffusion, viscosity, wetting angle, surface tension [34] and more melt would infiltrate into porous structure of Cf/C composites. Uneven melt was
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detected at 1380 °C (insert in Fig. 3(a)), which is mainly due to the relatively low reaction
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temperature. Uniform and nearly compact deposit layer could be formed when the temperature was increased. However, it became loose and porous when the reaction temperature was further increased to 1440 °C because of the perfect wettability of the
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Ti14Si86 braze filler on Cf/C composites.
Fig. 3. Interface microstructure (magnification of the yellow box) of PI-Cf/C composites at (a) 1380 °C, (b) 1400 °C, (c) 1420 °C, (d) 1440 °C. (The insert images show the interface microstructure of the plane and depth direction.) 10
ACCEPTED MANUSCRIPT The interface phase evolution can also be reflected by XRD patterns. As shown in Fig. 4(a), for PI-Cf/C composites system, SiC and TiSi2 formed at the interface between filler metal and Cf/C composites at 1380 °C, indicating that Si has reacted with carbon. It shows
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that TiSi2 disappeared when the reaction temperature was raised up to 1420 °C. Si was almost exhausted at 1400 °C, which indicated a further interface reaction between Si and carbon as well as greater liquid infiltration. At the same time, Ti3SiC2 phase was detected
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caused by the reaction among Ti-Si-C system. Racault et al. [45] reported that the compound
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Ti3SiC2 in contact with graphite can decompose at high temperature as follows: Ti3SiC2=Si(g)+3TiC0.67(s)
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Si(g)+C(s)=SiC(s)
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TiC0.67(s)+0.33C(s)=TiC(s)
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As shown in Fig. 4(b), Ti3SiC2 began to decrease as the reaction temperature increased and the phases of TiC and TiC0.98 obviously appeared at about 1420 °C then continually increased. Therefore, it could be deduced that the spherical phases existed among Ti3SiC2
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were TiC and TiC0.98 (Fig. 3(d)). It was reported these carbides had more metal-like character,
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which was considered to be beneficial to the adhesion between the carbides and the matrix [46]. When the temperature was raised up to 1440 °C, the main phases were SiC and small amount of TiCx, indicating the reaction during melt infiltration was almost complete. In conclusion, an almost compact deposit layer was achieved through the pre-infiltration process, which provided a convenient way of densification and modification of the substrate. To enrich the phase analysis, Fig. 5 shows the detailed microstructure of the deposit layer of the PI-Cf/C composites observed by TEM. Bright field images and Selected Area Electron 11
ACCEPTED MANUSCRIPT Diffraction (SAED) patterns of TiSi2, Si and SiC are exhibited in Fig. 5(a), (b) and (d), respectively. The nanocrystal in Fig. 5(a) was confirmed to be Ti3SiC2 by High Resolution Transmission Electron Microscopy (HRTEM) image and its fast Fourier transform (FFT) (Fig.
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5(c)). Therefore, the TEM analysis agrees well with the above SEM and XRD results.
Fig. 4. (a) Full and (b) portion of XRD patterns of the deposit layer of the PI-Cf/C composites at various
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infiltrating temperatures.
Fig. 5. TEM investigation of the deposit layer of the PI-Cf/C composites: (a), (b) and (d): bright field images of TiSi2, Si, SiC and their SAED patterns, respectively; (c) HRTEM image and its FFT of the area labeled in (a). 12
ACCEPTED MANUSCRIPT 3.2. Joint microstructure The perfect wetting of Ti-Si system on Cf/C composites indicates that the Ti14Si86 foil is suitable to braze Cf/C composites. In order to evaluate the two-step brazing procedure of
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Ti14Si86 foil on Cf/C composites, the backscattered electron (BSE) image and the electron probe micro-analyzer (EPMA) maps of the brazed joints at 1380 °C are shown in Fig. 6. A fine contact brazing seam formed without any defects at 1380 °C. According to the spectrum
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analysis results in table 1, it could be deduced that eutectic microstructure of Si and TiSi2
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made up the brazing seam (Fig. 6(a)). Strip-like TiSi2 uniformly located in and strengthened the Si matrix, besides, there was also SiC phase dispersively distributing in the joining zone and Cf/C composites substrate. The dispersed SiC may be attributed to the SiC reaction layer formed during the pre-infiltration procedure. When the PI-Cf/C composites were brazed at
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1380 °C, lower than the melting point of SiC, partial SiC layer would be soured into debris by the melt of Ti-Si system owing to uncompact combination with Cf/C composites at low temperature. Finally, the dispersed SiC phase distributed across the joining zone, conductive
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to reducing CTE mismatch between brazing seam and Cf/C composites. Adjacent to Cf/C
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composites, a thin and discontinuous reaction layer of SiC already formed at 1380 °C. The maps of elements show that most of Si distributed all over the brazing seam (Fig. 6(b)) and Ti mainly dispersed in TiSi2 phase location and the interface between the brazing seam and the Cf/C composites (Fig. 6(c)). Besides, the rest of Si and Ti elements had infiltrated into the pores of Cf/C composites. The element of carbon only concentrated within the substrate and the dispersed SiC phase (Fig. 6(d)).
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Fig. 6. (a) The SEM image of the two-step brazed joints at 1380 °C; the EPMA maps of (b) Si, (c) Ti, and (d) C.
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Table 1 EDS chemical analysis (at.%) of different positions in Fig. 6(a) Si 64.91 98.97 49.33
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Phase Label A B C D
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Ti 34.05 0.32 0.40
C 0.12 0.07 49.55
Possible Phases TiSi2 Si SiC
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Fig. 7 shows the microstructure of the brazed joints of the Cf/C composites at various temperatures. Similar to the micrograph of the brazed joint at 1380 °C (Fig. 6(a)), the brazing seam was mainly made up of eutectic Si matrix and distributed TiSi2 phase. Additionally, TiSi2 phases got refined and evident “nail” effect appeared as the brazing temperature increased (Fig. 7(a)-(d)). The thickness of brazing seam gradually reduced from about 270 µm (1380 °C) to 80 µm (1440 °C) and the reaction layer of SiC kept growing as temperature increased. During the brazing procedure, Ti-Si melt continually infiltrated 14
ACCEPTED MANUSCRIPT through the SiC layer designed in the first step to form the “nails” which could both strengthen the Cf/C joints and moderate the CTE mismatch [15,47]. Carbon atoms diffused through the interface to react with Ti-Si melt, which could further densify the SiC reaction
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layer. The insert in Fig. 7(b) shows that a dense and zigzag SiC layer which is more than 10µm has formed. The growing SiC layer compared with Fig. 3(b) illustrated a densification process of the reaction layer during the brazing step. Therefore, through the novel method
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which combined brazing process with pre-infiltration, the achieved brazed joint evolved into
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a graded composite structure which comprised modified and densified Cf/C substrate having closer CTE with metal filler, dense and continual SiC reaction layer as well as strip-like TiSi2 strengthened Si substrate. The gradual CTE transition effectively reduced the residual stress derived from the thermal mismatch among the brazed joints. The insert in Fig. 7(c) shows
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parallel interface cracks existed in the brazed joints at 1420 °C, which is detrimental to mechanical properties of the joints. A giant hole formed when the joint was brazed at 1440°C (insert in Fig. 7(d)), which can probably attribute to the excessive interfacial reaction and
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Fig. 7. The microstructure of two-step brazed joints at (a) 1380 °C; (b) 1400 °C; (c) 1420 °C; (d) 1440 °C. (The insert image in (b) shows the interface and the insert images in (c) and (d) show the defects.)
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3.3. Mechanical properties and fracture characters of the joint Shear strength test and nano-indentation test were both used to evaluate the mechanical behaviors of brazed joints. Fig. 8 shows the shear strength of the two-step Cf/C joints tested at
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room temperature. When the brazing temperature increased from 1380 °C to 1400 °C, the
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average shear strength raised dramatically from 11 MPa to 26 MPa. The poor mechanical performance of the joint achieved at lower temperature was probably ascribed to uncompact reaction layer (Fig. 6(a)) and incomplete melting of the foils. Cf/C composites joined at 1400 ° C and 1420 ° C showed similar shear strength, indicating sufficient interfacial chemical reaction and strong bonding between the substrate and the brazing seam. In the research of X. Zhou [5], a contact pressure of 30 MPa was applied on Cf/C joints to obtain a maximum shear strength of 26.3 MPa. In this research, the developed pressureless method 16
ACCEPTED MANUSCRIPT also obtained similar joint strength. The mechanical properties of the phases in the joint were investigated by nano-indentation. As shown in Fig. 9, all the reaction products of SiC, TiSi2 and Si displayed larger modulus and hardness than the Cf/C composites [48]. When the
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brazing temperature was elevated, the growing SiC reaction layer contributed to reinforcing the whole Cf/C joints. Nevertheless, the flaws among the brazing seam such as longitudinal cracks (1420 °C) and bulky holes (1440 °C) supposedly caused negative impact on the
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mechanical properties of the Cf/C joints. Thus, both these two factors should be responsible
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for the shear strength of the joints. In the early stage, the mechanical performance was determined by the dense SiC reaction layer and strong combination among the graded CTE composite structure. But when temperature exceeded 1420 °C, the defects existed in the brazing seam may play a negative role. Besides, during the high temperature shear test, the
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joints brazed at 1400 °C for 10 min showed a maximum shear strength of 36 MPa, which indicated a promising application in high temperature. 30
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1400
1420
1440
Brazing Temperature(°C) Fig. 8. The shear strength of the two-step Cf/C joints brazed at various temperatures.
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Fig. 9. (a) Load versus penetration depth in various phases. (b) Modulus and hardness of the phases.
SEM images of the fracture microstructure of the two-step Cf/C joints after shear
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strength test were presented in Fig. 10. The detailed feature information was expressed in high magnification and the overall perspectives were shown by the insert photographs. When the brazing temperature was 1380 °C, the SiC reaction layer seemed to be discontinuous and uncompact (Fig. 6(a)), making it easier for cracks initiating and propagating in the brazing
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seam near the interface. Hence, the phases of Si, TiSi2 and SiC were all exposed. The flat fracture appearance indicated low shear strength of the brazing joints. For the joints brazed at
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1400 °C and 1420 °C, strong adhesion formed between Cf/C substrate and brazing seam due to sufficient reaction and further densification of the SiC layer. Although with similar shear
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strength, the brazed joints at different temperatures showed different fracture behaviors. When the brazing temperature was 1400 °C, in addition to Si and SiC, pull-out TiSi2 and carbon fiber dispersed over the fracture surface, conductive to improving mechanical properties of the joints [14,49]. The mixed fracture paths ran across brazing seam, interface and Cf/C composites, which made the Cf/C brazed joints more reliable. However, as the brazing temperature was increased to 1420 °C, the cracks parallel to interface turned into failure initiation. Such failure mainly occurred in the interface and brazing seam without the 18
ACCEPTED MANUSCRIPT reinforcement effect of the carbon fibers (insert in Fig. 10(c)). With further increase of the temperature to 1440 °C, the phases exposed in the fracture surface did not change and the refinement of TiSi2 was apparently observed among the Si substrate. But the severe flaws
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within the brazed joints led to the bulk strength degradation despite the pull-out of TiSi2 and
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carbon fibers (insert in Fig. 10(d)).
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Fig. 10. Fracture morphology of the Cf/C joints brazed at different temperatures: (a) 1380 °C; (b) 1400 °C;
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(c) 1420 °C; (d) 1440 °C.
XRD was applied to analyze the phase distribution in the fracture surface. As shown in Fig. 11, the phases of Si, TiSi2 and SiC were all detected in each period from 1380 °C to 1440 °C, which was in accordance with the analysis results in Fig. 10.
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3.4. Reaction and joining mechanism
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Fig. 11. XRD pattern of the fracture surfaces of the joints brazed at various temperatures.
Fig. 12 reveals the reaction and joining mechanism of the two-step brazed Cf/C joints through a series of schematic diagrams. During the pre-infiltration step, the Ti-Si eutectic
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alloy exhibits excellent compatibility and wettability on the Cf/C composites. An almost tight and consecutive boundary zone mainly of SiC can be fabricated with the interaction among the elements of C, Si and Ti. Such modification treatment provides an economic and simple
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hole-sealing way to properly avoid the excessive liquid impregnation into the porous
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substrate, which highly benefits the reliable bonding of the Cf/C joints in the next brazing process. The designation of gradient structure of the near-surface Cf/C composites could reinforce the starting carbon matrix firstly, then effectively minimize the impact of the CTE mismatch between the outer reaction layer and the carbon matrix. The schematic of the PI-Cf/C sample and a Ti-Si foil are shown in Fig. 12(a). Next, in the non-pressure brazing procedure, the original unreacted phases of Si and TiSi2 attached to PI-Cf/C will dissolve into the liquid Ti-Si melt (Fig. 12(b)). The adherent 20
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unreacted Si and Ti elements form the eutectic structure in the brazed joints. The gradual CTE transition structure has formed, which can effectively reduce the residual stress derived from the thermal mismatch (Fig. 12(c)). As the brazing temperature is elevated, the growing
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SiC reaction zone, refining TiSi2 strips, stronger nail effect as well as the graded CTE
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structure contribute to the reliable bonding of the brazed joints. However, serious reaction among the C-Ti-Si system under higher temperature leads to interface degradation and enormous filler exhaustion, causing defects such as cracks and voids in the joints. An optimal
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fracture mode of the joint should contain the interface, the substrate and the brazing seam.
Fig. 12. The joining mechanism of two-step brazed Cf/C joints: (a) PI-Cf/C and the filler alloy; (b) the filler alloy melt during the second step; (c) the brazed joint.
In contrast, the direct brazing experiment was also carried out at various parameters. However, the shear strengths of the joints obtained by this process were all less than 10 MPa. All the Cf/C shear strength samples were confronted the same problem: most of the filler lost caused by the uncontrolled infiltration of the whole porous Cf/C substrate, leaving 21
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high temperature phases such as SiC, TiCx were produced in the first step (Fig. 3), which could serve as a barrier to prevent excess infiltration into the pores of Cf/C composites in the next step (insert in Fig. 2(a)). Therefore, strong bonding was achieved by the two-step
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As mentioned above, the key points to strong Cf/C brazed joints depend on both preparations of the nearly tight pre-infiltration layer and designation of the gradient-CTE composite joint structure with reliable bonding. Therefore, through the novel method which combined brazing process with pre-infiltration, the achieved brazed joint can effectively
4. Conclusion
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release thermal stress and strengthen the brazed joint and the intrinsic substrate.
In this work, reliable Cf/C joints have been successfully obtained with Ti-Si eutectic
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alloy through a novel two-step brazing method for the first time. The effects of the infiltration
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and brazing parameters on the phase compositions and mechanical behaviors of the joints were investigated. The joining mechanism was elaborated in detail. The infiltration structure effectively densified and strengthened the intrinsic Cf/C substrate firstly, and the PI-Cf/C further reacted with Ti-Si alloy to form a compact joint. Si and TiSi2 eutectic structure was retained in the brazing seam, and SiC phase was formed at the interface between Ti-Si structure and Cf/C composites. Therefore, both of the two steps are essential to the CTE-gradient bonding structure and excellent brazing joint of Cf/C composites. This research 22
ACCEPTED MANUSCRIPT provides a new high temperature braze filler (Ti-Si eutectic alloy) and technique to join the
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Cf/C composites.
Acknowledgment
The authors gratefully acknowledge the financial support from the National Natural
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Science Foundation of China under Grant Nos. 51622503, U1737205 and U1537206. References
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[1] S. Dietrich, J.M. Gebert, G. Stasiuk, A. Wanner, K.A. Weidenmann, O. Deutschmann, et al., Microstructure characterization of CVI-densified carbon/carbon composites with various fiber distributions, Compos. Sci. Technol. 72 (15) (2012) 1892-1900. [2] J. Yu, C. Zhou, H. Zhang, A micro-image based reconstructed finite element model of needle-punched C/C composite, Compos. Sci. Technol. 153 (2017) 48-61. [3] V. Casalegno, T. Koppitz, G. Pintsuk, M. Salvo, S. Rizzo, S. Perero, et al., Proposal for a modified non-active brazing alloy for joining CFC composites to copper, Compos. Part B 56 (2014) 882-888. [4] Y. Cai, X. Yin, S. Fan, L. Zhang, L. Cheng, Tribological behavior of three-dimensional needled ceramic modified carbon/carbon composites in seawater conditions, Compos. Sci. Technol. 87 (2013) 50-57. [5] X. Zhou, H. Yang, F. Chen, Y.-H. Han, J. Lee, S. du, et al., Joining of carbon fiber reinforced carbon composites with Ti3SiC2 tape film by electric field assisted sintering technique, Carbon 102 (2016) 106-115. [6] M. Salvo, P. Lemoine, M. Ferraris, M. Montorsi, Joining of carbon-carbon composites for thermonuclear fusion applications, J. Am. Ceram. Soc. 80 (1) (1997) 206-212. [7] X. Song, H. Li, V. Casalegno, M. Salvo, M. Ferraris, X. Zeng, Microstructure and mechanical properties of C/C composite/Ti6Al4V joints with a Cu/TiCuZrNi composite brazing alloy, Ceram. Int. 42 (5) (2016) 6347-6354. [8] M. Singh, R. Asthana, T.P. Shpargel, Brazing of carbon–carbon composites to Cu-clad molybdenum for thermal management applications, Mater. Sci. Eng. A 452-453 (2007) 699-704. [9] M. Singh, T.P. Shpargel, G.N. Morscher, R. Asthana, Active metal brazing and characterization of brazed joints in titanium to carbon–carbon composites, Mater. Sci. Eng. A 412 (1-2) (2005) 123-128. [10] W. Fu, S.P. Hu, X.G. Song, J.X. Li, J. Cao, J.C. Feng, et al., Wettability and bonding of graphite by Sn0.3Ag0.7Cu-Ti alloys, Carbon 121 (2017) 536-543. [11] S. Rizzo, S. Grasso, M. Salvo, V. Casalegno, M.J. Reece, M. Ferraris, Joining of C/SiC composites by spark plasma sintering technique, J. Eur. Ceram. Soc. 34 (4) (2014) 903-913. [12] Z. Wang, G. Wang, M. Li, J. Lin, Q. Ma, A. Zhang, et al., Three-dimensional graphene-reinforced Cu foam interlayer for brazing C/C composites and Nb, Carbon 118 (2017) 723-730. 23
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AC C
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[13] N.V. Moutis, C. Jimenez, X. Azpiroz, T. Speliotis, C. Wilhelmi, S. Messoloras, et al., Brazing of carbon–carbon composites to Nimonic alloys, J. Mater Sci. 45 (1) (2009) 74-81. [14] J. Cao, C. Li, J. Qi, Y. Shi, J. Feng, Combustion joining of carbon–carbon composites to TiAl intermetallics using a Ti–Al–C powder composite interlayer, Compos. Sci. Technol. 115 (2015) 72-79. [15] M. Singh, R. Asthana, Characterization of brazed joints of CC composite to Cu-clad-Molybdenum, Compos. Sci. Technol. 68 (14) (2008) 3010-3019. [16] V. Casalegno, M. Salvo, M. Ferraris, Surface modification of carbon/carbon composites to improve their wettability by copper, Carbon 50 (6) (2012) 2296-2306. [17] X.R. Song, H.J. Li, X. Zeng, Brazing of C/C composites to Ti6Al4V using multiwall carbon nanotubes reinforced TiCuZrNi brazing alloy, J. Alloy Compd. 664 (2016) 175-180. [18] S. Pellegrino, L. Thomé, A. Debelle, S. Miro, P. Trocellier, Radiation effects in carbides: TiC and ZrC versus SiC, Nucl. Instrum. Methods Phys. Res. Sect. B 327 (2014) 103-107. [19] M. Li, X. Zhou, H. Yang, S. Du, Q. Huang, The critical issues of SiC materials for future nuclear systems, Scr. Mater. 143 (2018) 149-153. [20] P. Dadras, T.T. Ngai, G.M. Mehrotra, Joining of carbon-carbon composites using boron and titanium disilicide interlayers, J. Am. Ceram. Soc. 80 (1) (1997) 125-132. [21] J. Li, L. Liu, Y. Wu, Z. Li, W. Zhang, W. Hu, Microstructure of high temperature Ti-based brazing alloys and wettability on SiC ceramic, Mater. Des. 30 (2) (2009) 275-279. [22] N.M. Larson, F.W. Zok, In-situ 3D visualization of composite microstructure during polymer-to-ceramic conversion, Acta Mater. 144 (2018) 579-589. [23] R. Naslain, Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview, Compos. Sci. Technol. 64 (2) (2004) 155-170. [24] H.A. Bale, A. Haboub, A.A. MacDowell, J.R. Nasiatka, D.Y. Parkinson, B.N. Cox, et al., Real-time quantitative imaging of failure events in materials under load at temperatures above 1,600°C, Nat. Mater. 12 (2012) 40. [25] D.L. Poerschke, A. Braithwaite, D. Park, F. Lauten, Crystallization behavior of polymer-derived Si-O-C for ceramic matrix composite processing, Acta Mater. 147 (2018) 329-341. [26] Y. Zeng, D. Wang, X. Xiong, X. Zhang, P. J. Withers, W. Sun, et al., Ablation-resistant carbide Zr0.8Ti0.2C0.74B0.26 for oxidizing environments up to 3,000°C, Nature commun. 8 (2017) 15836. [27] Q. Fu, F. Zhao, H. Li, H. Peng, X. Nan, A Multi-interlayer LMAS Joint of C/C–SiC Composites and LAS Glass Ceramics, J. Mater. Sci. Technol. 31 (5) (2015) 467-472. [28] Q.G. Fu, H. Peng, H.J. Li, J. Wang, Q. Zhu, A gradient LMAS interlayer joint of SiC coated C/C composites to LAS glass ceramics, Ceram. Int. 40 (1) (2014) 2461-2466. [29] S. Chen, C. Zhang, Y. Zhang, H. Hu, Preparation and properties of carbon fiber reinforced ZrC–ZrB2 based composites via reactive melt infiltration, Compos. Part B 60 (2014) 222-226. [30] F.H. Gern, R. Kochendörfer, Liquid silicon infiltration: description of infiltration dynamics and silicon carbide formation, Compos. Part A 28 (4) (1997) 355-364. [31] H. Zhou, R.N. Singh, Kinetics model for the growth of silicon carbide by the reaction of liquid silicon with carbon, J. Am. Ceram. Soc. 78 (9) (1995) 2456-2462. [32] Y. Zeng, X. Xiong, G. Li, Z. Chen, W. Sun, D. Wang, Microstructure and ablation behavior of carbon/carbon composites infiltrated with Zr–Ti, Carbon 54 (2013) 300-309. [33] E.W. Washburn, The Dynamics of Capillary Flow, Phys. Rev. 17 (3) (1921) 273-283. [34] Y. Wang, X. Zhu, L. Zhang, L. Cheng, Reaction kinetics and ablation properties of C/C–ZrC composites fabricated by reactive melt infiltration, Ceram. Int. 37 (4) (2011) 1277-1283. 24
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[35] J. Yang, O.J. Ilegbusi, Kinetics of silicon–metal alloy infiltration into porous carbon, Compos. Part A 31 (6) (2000) 617-625. [36] R. Asthana, Interface- and diffusion-limited capillary rise of reactive melts with a transient contact angle, Metall. Mater. Trans. A (2002) 2119-2128. [37] R. Asthana, Dissolutive capillary penetration with expanding pores and transient contact angles, J. Colloid Interface Sci. 231 (2000) 398-400. [38] R. Asthana, On reactive capillary flows with a dynamic contact angle, J. Mater. Processing Manuf. Sci., 10 (4) (2002) 187-197. [39] Y. Zeng, X. Xiong, D. Wang, L. Wu, Infiltration mechanism and factors influencing carbon/carbon– Zr–Ti–C composites prepared by liquid metal infiltration, J. Mater. Process Technol. 214 (12) (2014) 3150-3157. [40] Y. Chang, W. Sun, X. Xiong, Z. Chen, Y. Wang, Z. Hao, et al., A novel design of Al-Cr alloy surface sealing for ablation resistant C/C-ZrC-SiC composite, J. Eur. Ceram. Soc. 37 (2) (2017) 859-864. [41] T. Lin, M. Yang, P. He, C. Huang, F. Pan, Y. Huang, Effect of in situ synthesized TiB whisker on microstructure and mechanical properties of carbon–carbon composite and TiBw/Ti–6Al–4V composite joint, Mater. Des. 32 (8-9) (2011) 4553-4558. [42] C.A. Swenson, Recommended Values for the Thermal Expansivity of Silicon from 0 to 1000 K, J. Phys. Chem. Ref. Data 12 (2) (1983) 179-182. [43] T.S.R.C. Murthy, C. Subramanian, R.K. Fotedar, M.R. Gonal, P. Sengupta, S. Kumar, et al., Preparation and property evaluation of TiB2+TiSi2 composite, Int. J. Refract. Met. Hard Mater. 27 (3) (2009) 629-636. [44] Y. Zeng, X. Xiong, D. Wang, L. Wu, Residual thermal stresses in carbon/carbon–Zr–Ti–C composites and their effects on the fracture behavior of composites with different preforms, Carbon 81 (2015) 597-606. [45] C. Racault, F. Langlais, R. Naslain, Solid-state synthesis and characterization of the ternary phase Ti3SiC2, J. Mater. Sci. 29 (13) (1994) 3384-3392. [46] M. Jiang, H. Y. Xiao, H. B. Zhang, S. M. Peng, C. H. Xu, Z. J. Liu, et al., A comparative study of low energy radiation responses of SiC, TiC and ZrC, Acta Mater. 110 (2016) 192-199. [47] H. Wang, J. Cao, J. Feng, Brazing mechanism and infiltration strengthening of CC composites to TiAl alloys joint, Scr. Mater. 63 (8) (2010) 859-862. [48] K.K. Chawla, Composite Materials: Science and Engineering, Springer, New York (2012) 249-292. [49] J. Wang, K. Li, W. Li, H. Li, Z. Li, L. Guo, The preparation and mechanical properties of carbon/carbon composite joints using Ti–Si–SiC–C filler as interlayer, Mater Sci. Eng. A 574 (2013) 37-45.
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S0008-6223(18)30751-6
DOI:
10.1016/j.carbon.2018.08.021
Reference:
CARBON 13376
To appear in:
Carbon
Received Date: 27 May 2018 Revised Date:
22 July 2018
Accepted Date: 9 August 2018
Please cite this article as: Zongjing He, Chun Li, Junlei Qi, Yongxian Huang, Jicai Feng, Jian Cao, Preinfiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy, Carbon (2018), doi: 10.1016/j.carbon.2018.08.021 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Pre-infiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy Zongjing He, Chun Li, Junlei Qi*, Yongxian Huang, Jicai Feng, Jian Cao*
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State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
*
Corresponding authors. Tel: +86-451-86418882. E-mail: [email protected] (J. Cao); Tel: +86-451-86418882. E-mail: [email protected] (J. Qi)
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Pre-infiltration and brazing behaviors of Cf/C composites with high temperature Ti-Si eutectic alloy Zongjing He, Chun Li, Junlei Qi*, Yongxian Huang, Jicai Feng, Jian Cao*
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State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China
Abstract
A novel method is developed to join carbon fiber reinforced carbon composites (Cf/C
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composites) via a two-step approach using Ti14Si86 metal foil as the joining material. The
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infiltration behavior, joint microstructure, interface reaction, fracture characteristics and joining mechanism are investigated. The preparation of the almost compact deposit layer and the infiltration carbide can immediately heal the pores of Cf/C composites and form the gradient composite structure, which can effectively minimize the impact of the coefficient of
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thermal expansion (CTE) mismatch between the outer reaction layer and the carbon matrix and avoid the excessive filler infiltration in the brazing step. The brazing process reveals that
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the growing SiC interface and Ti-Si eutectic structure in the seam also behave as a graded CTE composite structure, which is beneficial for strengthening the joint. Compared with
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direct brazing, this method can obtain more reliable Cf/C composites joints due to strong interfacial bonding and the gradient structure. As a result, a favorable shear strength of 26 MPa is achieved within 10 min at a brazing temperature of 1400 °C. 1. Introduction Carbon fiber reinforced carbon composites are attractive materials in aerospace
*
Corresponding authors. Tel: +86-451-86418882. E-mail address: [email protected] (J. Cao). Tel: +86-451-86418882. E-mail address: [email protected] (J. L. Qi). 1
ACCEPTED MANUSCRIPT application such as space vehicle heat shields and aircraft brakes due to low thermal expansion and outstanding high-temperature mechanical properties at considerably reduced weight [1-5]. Along with extremely low neutron activation and low atomic number, great
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interest is drawn to the application of Cf/C composites in nuclear fusion reactors [6]. Cf/C composites are usually required to join to other substrates or to themselves to be fabricated into complex assemblies [7,8]. Among the current joining technologies, brazing has been
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extensively used due to its simplicity, high joint strength and excellent adaptability of sample
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configuration [9-12]. But the major problem in brazing Cf/C composites lies in its limited application at relatively low temperature [7,13,14]. It is essential to develop a new braze filler and improve high-temperature joint properties.
The judicious use of braze chemistry and application of innovative brazing strategies
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should combine reliable interfacial bonding and structure with graded CTE. Both Ti and Si elements show perfect wetting and adherence to the carbon material and equilibrium contact angles could reach near 0° in short brazing time [15,16]. Moreover, the reaction products
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between Ti, Si and carbon show excellent high temperature resistance [17]. In addition, Ti
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and Si are low-activation elements, which have shown great potential in thermonuclear fusion applications [6,18,19]. Dadras et al. [20] brazed Cf/C composites using TiSi2 and the achieved joint showed a maximum shear strength of 34.4 MPa when testing at 1164 °C, which indicates that Ti-Si alloy has potential applications in high temperature. With relatively low melting point and perfect fluidity, Ti-Si eutectic metal is much more beneficial to filler metal spreading and has exhibited good adhesion with silicon carbide [21]. However, the brazing of Cf/C composites with Ti-Si eutectic alloy has not been reported. In addition, the authors have 2
ACCEPTED MANUSCRIPT carried out some investigation of joining Cf/C composites. Because of the abundant porous structure of the low-density substrate, all attempts to directly braze the Cf/C joints have failed. Thus a necessary pretreatment of the Cf/C composites should be considered before brazing
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process. Chemical vapor infiltration (CVI), precursor infiltration and pyrolysis (PIP), slurry infiltration (SI) and reactive melt infiltration (RMI) have been applied to densify or modify
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fiber-reinforced ceramic-matrix composites (CMCs) or Cf/C composites [22-26]. SiC
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transition layer was prepared through CVI, chemical vapor deposition (CVD) and pack cementation processes to form a zigzag interface structure of Cf/C substrate to improve the poor wettability of LAS glass ceramic [27,28]. However, no practical methods can be found to moderate the residual stress in high-temperature brazed Cf/C joints. Compared with other
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technologies, RMI is an attractive modification method to introduce ceramics because of its time-saving preparation and relatively low cost [23,29,30]. Considering the relatively low density, small coefficient of thermal expansion as well as high working temperature in air of
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SiC [25,31], a novel method combined brazing with RMI was used to design a complex
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structure with graded CTE to mitigate residual stress among the joints. Furthermore, multiple effects of the simple method such as pinning the substrate and strengthening the Cf/C composites can also be expected. Taking advantage of the perfect wettability of the high-temperature Ti14Si86 eutectic metal and the porous structure of the substrate, a novel brazing method was developed to join Cf/C composites in argon gas atmosphere without extra pressure. The infiltration behavior and braze ability of Ti-Si for Cf/C composites were studied in this work. The joint 3
ACCEPTED MANUSCRIPT microstructure, shear fracture characteristics and interface bonding mechanism of the resulting joints were investigated in detail. 2. Experimental
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2.1. Materials Fig. 1(a) and (b) show the microstructure of the as-received Cf/C composites and Ti-Si eutectic alloy. The Cf/C composite used in this study had an open porosity of about 9%, and
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both the pores in the pyrocarbon and bundles could be observed, as shown in Fig. 1(a). The
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density and three point bending strength of the substrate were about 1.64 g/cm3 and 52.5±3 MPa, respectively. The Cf/C composites were cut into dimension of 5 mm × 5 mm × 5 mm for metallographic observation and 5 mm × 5 mm × 10 mm for shear tests. The surfaces perpendicular to the fiber planes were used as the joining planes. The eutectic filler alloy was
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prepared using the raw materials of sponge Ti (99.99 wt.% purity) and Si (99.99 wt.% purity) by vacuum arc-melting technology. As shown in Fig. 1(b), the alloy was homogeneous and mainly consisted of two phases, which were confirmed to be Si matrix phase (the gray phase)
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and TiSi2 intermetallic (the white phase) by X-ray diffraction (XRD) (Fig. 1(c)). The filler
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alloy was cut into 5 mm × 5 mm × 300 µm and 5 mm × 5 mm × 900 µm pieces, then ground by 800-grit silicon carbide paper. All the Cf/C composites and metal foils were ultrasonic cleaned in absolute ethanol for 5 min prior to brazing. 2.2. Pre-infiltration and brazing processes In this work, a novel two-step brazing method was designed to join the Cf/C composites. Fig. 1(d) and (e) show the schematics of the melt infiltration process and the brazing process. For the melt infiltration process, firstly, Ti-Si eutectic alloy foils of 300µm were placed on the 4
ACCEPTED MANUSCRIPT surface of the raw Cf/C composites (the area of 5 mm × 5 mm) used for brazing. Afterwards, the pre-treated assembly was heated to the holding temperature (1380 °C, 1400 °C, 1420 °C, 1440 °C) at a rate of 5 °C/min and held for 10 min, then slowly cooled down to room
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temperature. The brazing step was performed under the same heating cycle as the first step, and brazing assembly was like a sandwich structure with the filler alloy (300 µm) between two pre-infiltrated Cf/C substrates (PI-Cf/C). Particularly, the direct brazing of the Cf/C
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composites (using filler alloy of 900 µm) without per-infiltration process was served as a
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contrast experiment. The infiltration and brazing processes were carried out in argon atmospheric and no additional pressure was applied on each assembly except the dead weight.
2.3. Microstructural analysis and property measurement
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The microstructure of the PI-Cf/C samples and brazed joints was characterized by field emission scanning electron microscope (FE-SEM, Quanta 200FEG) coupled with energy dispersive spectroscope (EDS, INCA Energy 300) and transmission electron microscope
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(TEM, Talos F200X). The phases in the samples were identified by X-ray diffraction analysis
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(XRD, Philips X’Pert). The shear strength of the brazed joints at room temperature was measured through universal testing machine (Instron 5569) at a constant speed of 0.5mm/min, the schematic diagram is shown in Fig. 1(f). With the same method, the high temperature shear test was conducted at 1200 °C by radiation heating in the argon flow. Nano-indentation test (NanoIndenter G200) was also used to evaluate the mechanical behaviors of the brazed joints. The fracture characteristics was examined to determine the fracture propagation by SEM and XRD. Each value reported is the average of at least 10 tests. 5
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Fig. 1. The micrographs of (a) the Cf/C composites and (b) the Ti-Si alloy. (c) The XRD pattern of the Ti-Si alloy. The schematic diagrams of (d) the infiltration process, (e) the brazing process and (f) the shear test.
3. Results and discussion
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3.1. Analysis of infiltration behaviors
Fig. 2((a)-(c)) show the microstructure of the PI-Cf/C composites. As can be seen from
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Fig. 2(a), both the pores in bundles and among the pyrocarbon of the Cf/C substrate have been completely filled up with Ti-Si alloy [32], which might be attributed to the infiltration of
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melt. The transport of liquid melt through the pores is governed by capillarity and the resistance of the viscous forces, which can be represented by the modified Washburn model [33-35]:
2 3 B rt- At 2 3 A= 2D M
Mρ
C Si 一 Ti
ρ
Si 一 Ti C
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(3-1) (3-2)
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Cσcosθ 2µ
(3-3)
where h is the infiltration height, t is the annealing time, µ is the viscosity, σ is the surface tension, θ is the contact angle-about 0° according to sessile drop technique, r is the
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mean pore radius, C is the tortuosity factor, D is the effective diffusivity of carbon, Mi and ρi are the atomic weight and density of component i. Eq. (3-1) and (3-2) indicate that the infiltration behavior depends on two factors: chemical reaction and properties of the liquid
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melt. In this research, A and B can be assumed as constants in Eq. (3-1). Hence, the
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infiltration height (h) depends on the annealing time (t) and the pore dimension (r), and the trend graph can be drawn as Fig. 2(e). As shown in Fig. 2(e), the X-direction represents the annealing time at the infiltration temperature; and the Y-direction represents the infiltration height. It can be concluded that carbides formed at the small opening pores would seal the
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pores immediately at the beginning of infiltration process [35]. When the melt metal infiltrates into the bigger opening pores, not only carbide layer will be formed during the interface reaction, but also a flow of liquid can be remained through the center of pores. The
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combined effects of both the chemical reaction and the melt infiltration give rise to the results
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that the initial infiltration of the melt metal is rapid and then slow down with the annealing time increasing. The height of the melt in the large opening pores could be higher than that in the small opening pores. R. Asthana [36-38] reported a new reactive infiltration model characterized by pore shrinkage and decaying contact angle due to interphase formation, which could qualitatively explain the sudden drop in the penetration velocity with increasing distance in Si/porous carbon system. The dynamic model in the articles could better predict the actual penetration lengths for Si/C samples than the other analyses. And our function 7
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pyrocarbon [39]. The dimensions of residual spaces in the interspaces were about 100 µm. During the infiltration process, the molten metal of titanium-silicon was drawn into the residual spaces and reacted with the previously deposited pyrocarbon to form the SiC layer.
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The melt penetrated into the Cf/C composites and the penetration depth depends on the types
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and dimensions of the pores [3]. It can be found in Fig. 2(b) that the molten metal of titanium-silicon infiltrated along the carbon fiber bundles and reacted with carbon fiber to form silicon carbide. In contrast, the interspaces among the pyrocarbon were too large to be completely filled by the Ti-Si melt and residual unreacted Ti-Si alloy could be found in the
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center of the pores (Fig. 2(c)). Therefore, as shown in the schematic diagram of Fig. 2(d), the molten metal of Ti-Si penetrated into the pores of the substrate and formed a tight combination at the interface. From the insert in Fig. 2(a), it can be seen that the hole in the
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substrate was not filled up by the filler alloy when a dense deposit layer has already formed,
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which could prevent further infiltration of the braze filler in the second step. The interlaced infiltration across the whole substrate could effectively heal the pores of Cf/C composites [40], which could avoid excessive braze filler infiltration and strengthen the Cf/C composites. The coefficient of thermal expansion of SiC is 4.4×10-6 K-1 [19], higher than that of Cf/C composites (2×10-6 K-1) [41] and Si (3.5×10-6 K-1) [42], but lower than TiSi2 (10.4×10-6 K-1) [43]. This gradient structure could effectively minimize the impact of the mismatch in thermal expansion coefficient between the reaction layer and the carbon matrix, which could 8
ACCEPTED MANUSCRIPT help to improve the joint strength during the second step of brazing process. Zeng et al. found the release of residual thermal stress in Zr-Ti infiltrated porous Cf/C composites using Raman
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spectroscopy analysis [44].
Fig. 2. (a), (b), (c) The micrographs and (d) the schematic diagram of the PI-Cf/C composites by Ti-Si alloy.
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(e) The function curves of infiltration height as annealing time.
Fig. 3 shows the typical interfacial microstructure of the PI-Cf/C composites under different temperatures. Fig. 3(a) indicates that there were three different phases distributed in
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solidified melt. EDS results reveal that the three phases were Si (Si: 99.85 at.%, Ti: 0.15
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at.%), TiSi2 (Si: 66.27 at.%, Ti: 33.73 at.%) and SiC (Si: 48.96 at.%, C: 51.04 at.%), respectively. The existence of Si element reflected an inadequate reaction between Si from the Ti-Si eutectic foil and Cf/C composites at such a low temperature. When the melt infiltrating temperature was increased to 1400 °C or above, the content of Si and TiSi2 rapidly decreased (Fig. 3(b)) and the flaky phase was exposed. It should be Ti3SiC2 according to EDS results (Ti: 48.59 at.%, Si: 17.26 at.%, C: 34.15 at.%). As the temperature was increased to 1440 °C, the new spherical-like phase was identified which should be further confirmed. The 9
ACCEPTED MANUSCRIPT morphology changes during pre-infiltration process can also be visually observed by the inserts in Fig. 3. The well-combined metallization layer got thinning as the temperature was increased, reflecting the fine wettability and permeability of Ti14Si86 metal foil on Cf/C
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composites. The infiltration depth could be higher than low temperatures due to the changing of series of factors such as carbon diffusion, viscosity, wetting angle, surface tension [34] and more melt would infiltrate into porous structure of Cf/C composites. Uneven melt was
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detected at 1380 °C (insert in Fig. 3(a)), which is mainly due to the relatively low reaction
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temperature. Uniform and nearly compact deposit layer could be formed when the temperature was increased. However, it became loose and porous when the reaction temperature was further increased to 1440 °C because of the perfect wettability of the
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Ti14Si86 braze filler on Cf/C composites.
Fig. 3. Interface microstructure (magnification of the yellow box) of PI-Cf/C composites at (a) 1380 °C, (b) 1400 °C, (c) 1420 °C, (d) 1440 °C. (The insert images show the interface microstructure of the plane and depth direction.) 10
ACCEPTED MANUSCRIPT The interface phase evolution can also be reflected by XRD patterns. As shown in Fig. 4(a), for PI-Cf/C composites system, SiC and TiSi2 formed at the interface between filler metal and Cf/C composites at 1380 °C, indicating that Si has reacted with carbon. It shows
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that TiSi2 disappeared when the reaction temperature was raised up to 1420 °C. Si was almost exhausted at 1400 °C, which indicated a further interface reaction between Si and carbon as well as greater liquid infiltration. At the same time, Ti3SiC2 phase was detected
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caused by the reaction among Ti-Si-C system. Racault et al. [45] reported that the compound
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Ti3SiC2 in contact with graphite can decompose at high temperature as follows: Ti3SiC2=Si(g)+3TiC0.67(s)
(3-4)
Si(g)+C(s)=SiC(s)
(3-5)
TiC0.67(s)+0.33C(s)=TiC(s)
(3-6)
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As shown in Fig. 4(b), Ti3SiC2 began to decrease as the reaction temperature increased and the phases of TiC and TiC0.98 obviously appeared at about 1420 °C then continually increased. Therefore, it could be deduced that the spherical phases existed among Ti3SiC2
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were TiC and TiC0.98 (Fig. 3(d)). It was reported these carbides had more metal-like character,
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which was considered to be beneficial to the adhesion between the carbides and the matrix [46]. When the temperature was raised up to 1440 °C, the main phases were SiC and small amount of TiCx, indicating the reaction during melt infiltration was almost complete. In conclusion, an almost compact deposit layer was achieved through the pre-infiltration process, which provided a convenient way of densification and modification of the substrate. To enrich the phase analysis, Fig. 5 shows the detailed microstructure of the deposit layer of the PI-Cf/C composites observed by TEM. Bright field images and Selected Area Electron 11
ACCEPTED MANUSCRIPT Diffraction (SAED) patterns of TiSi2, Si and SiC are exhibited in Fig. 5(a), (b) and (d), respectively. The nanocrystal in Fig. 5(a) was confirmed to be Ti3SiC2 by High Resolution Transmission Electron Microscopy (HRTEM) image and its fast Fourier transform (FFT) (Fig.
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5(c)). Therefore, the TEM analysis agrees well with the above SEM and XRD results.
Fig. 4. (a) Full and (b) portion of XRD patterns of the deposit layer of the PI-Cf/C composites at various
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infiltrating temperatures.
Fig. 5. TEM investigation of the deposit layer of the PI-Cf/C composites: (a), (b) and (d): bright field images of TiSi2, Si, SiC and their SAED patterns, respectively; (c) HRTEM image and its FFT of the area labeled in (a). 12
ACCEPTED MANUSCRIPT 3.2. Joint microstructure The perfect wetting of Ti-Si system on Cf/C composites indicates that the Ti14Si86 foil is suitable to braze Cf/C composites. In order to evaluate the two-step brazing procedure of
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Ti14Si86 foil on Cf/C composites, the backscattered electron (BSE) image and the electron probe micro-analyzer (EPMA) maps of the brazed joints at 1380 °C are shown in Fig. 6. A fine contact brazing seam formed without any defects at 1380 °C. According to the spectrum
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analysis results in table 1, it could be deduced that eutectic microstructure of Si and TiSi2
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made up the brazing seam (Fig. 6(a)). Strip-like TiSi2 uniformly located in and strengthened the Si matrix, besides, there was also SiC phase dispersively distributing in the joining zone and Cf/C composites substrate. The dispersed SiC may be attributed to the SiC reaction layer formed during the pre-infiltration procedure. When the PI-Cf/C composites were brazed at
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1380 °C, lower than the melting point of SiC, partial SiC layer would be soured into debris by the melt of Ti-Si system owing to uncompact combination with Cf/C composites at low temperature. Finally, the dispersed SiC phase distributed across the joining zone, conductive
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to reducing CTE mismatch between brazing seam and Cf/C composites. Adjacent to Cf/C
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composites, a thin and discontinuous reaction layer of SiC already formed at 1380 °C. The maps of elements show that most of Si distributed all over the brazing seam (Fig. 6(b)) and Ti mainly dispersed in TiSi2 phase location and the interface between the brazing seam and the Cf/C composites (Fig. 6(c)). Besides, the rest of Si and Ti elements had infiltrated into the pores of Cf/C composites. The element of carbon only concentrated within the substrate and the dispersed SiC phase (Fig. 6(d)).
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Fig. 6. (a) The SEM image of the two-step brazed joints at 1380 °C; the EPMA maps of (b) Si, (c) Ti, and (d) C.
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Table 1 EDS chemical analysis (at.%) of different positions in Fig. 6(a) Si 64.91 98.97 49.33
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Phase Label A B C D
50.06
Ti 34.05 0.32 0.40
C 0.12 0.07 49.55
Possible Phases TiSi2 Si SiC
1.03
47.84
SiC
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Fig. 7 shows the microstructure of the brazed joints of the Cf/C composites at various temperatures. Similar to the micrograph of the brazed joint at 1380 °C (Fig. 6(a)), the brazing seam was mainly made up of eutectic Si matrix and distributed TiSi2 phase. Additionally, TiSi2 phases got refined and evident “nail” effect appeared as the brazing temperature increased (Fig. 7(a)-(d)). The thickness of brazing seam gradually reduced from about 270 µm (1380 °C) to 80 µm (1440 °C) and the reaction layer of SiC kept growing as temperature increased. During the brazing procedure, Ti-Si melt continually infiltrated 14
ACCEPTED MANUSCRIPT through the SiC layer designed in the first step to form the “nails” which could both strengthen the Cf/C joints and moderate the CTE mismatch [15,47]. Carbon atoms diffused through the interface to react with Ti-Si melt, which could further densify the SiC reaction
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layer. The insert in Fig. 7(b) shows that a dense and zigzag SiC layer which is more than 10µm has formed. The growing SiC layer compared with Fig. 3(b) illustrated a densification process of the reaction layer during the brazing step. Therefore, through the novel method
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which combined brazing process with pre-infiltration, the achieved brazed joint evolved into
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a graded composite structure which comprised modified and densified Cf/C substrate having closer CTE with metal filler, dense and continual SiC reaction layer as well as strip-like TiSi2 strengthened Si substrate. The gradual CTE transition effectively reduced the residual stress derived from the thermal mismatch among the brazed joints. The insert in Fig. 7(c) shows
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parallel interface cracks existed in the brazed joints at 1420 °C, which is detrimental to mechanical properties of the joints. A giant hole formed when the joint was brazed at 1440°C (insert in Fig. 7(d)), which can probably attribute to the excessive interfacial reaction and
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infiltration of the Ti-Si filler.
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Fig. 7. The microstructure of two-step brazed joints at (a) 1380 °C; (b) 1400 °C; (c) 1420 °C; (d) 1440 °C. (The insert image in (b) shows the interface and the insert images in (c) and (d) show the defects.)
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3.3. Mechanical properties and fracture characters of the joint Shear strength test and nano-indentation test were both used to evaluate the mechanical behaviors of brazed joints. Fig. 8 shows the shear strength of the two-step Cf/C joints tested at
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room temperature. When the brazing temperature increased from 1380 °C to 1400 °C, the
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average shear strength raised dramatically from 11 MPa to 26 MPa. The poor mechanical performance of the joint achieved at lower temperature was probably ascribed to uncompact reaction layer (Fig. 6(a)) and incomplete melting of the foils. Cf/C composites joined at 1400 ° C and 1420 ° C showed similar shear strength, indicating sufficient interfacial chemical reaction and strong bonding between the substrate and the brazing seam. In the research of X. Zhou [5], a contact pressure of 30 MPa was applied on Cf/C joints to obtain a maximum shear strength of 26.3 MPa. In this research, the developed pressureless method 16
ACCEPTED MANUSCRIPT also obtained similar joint strength. The mechanical properties of the phases in the joint were investigated by nano-indentation. As shown in Fig. 9, all the reaction products of SiC, TiSi2 and Si displayed larger modulus and hardness than the Cf/C composites [48]. When the
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brazing temperature was elevated, the growing SiC reaction layer contributed to reinforcing the whole Cf/C joints. Nevertheless, the flaws among the brazing seam such as longitudinal cracks (1420 °C) and bulky holes (1440 °C) supposedly caused negative impact on the
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mechanical properties of the Cf/C joints. Thus, both these two factors should be responsible
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for the shear strength of the joints. In the early stage, the mechanical performance was determined by the dense SiC reaction layer and strong combination among the graded CTE composite structure. But when temperature exceeded 1420 °C, the defects existed in the brazing seam may play a negative role. Besides, during the high temperature shear test, the
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joints brazed at 1400 °C for 10 min showed a maximum shear strength of 36 MPa, which indicated a promising application in high temperature. 30
20 15
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Shear Strength(MPa)
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1400
1420
1440
Brazing Temperature(°C) Fig. 8. The shear strength of the two-step Cf/C joints brazed at various temperatures.
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Fig. 9. (a) Load versus penetration depth in various phases. (b) Modulus and hardness of the phases.
SEM images of the fracture microstructure of the two-step Cf/C joints after shear
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strength test were presented in Fig. 10. The detailed feature information was expressed in high magnification and the overall perspectives were shown by the insert photographs. When the brazing temperature was 1380 °C, the SiC reaction layer seemed to be discontinuous and uncompact (Fig. 6(a)), making it easier for cracks initiating and propagating in the brazing
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seam near the interface. Hence, the phases of Si, TiSi2 and SiC were all exposed. The flat fracture appearance indicated low shear strength of the brazing joints. For the joints brazed at
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1400 °C and 1420 °C, strong adhesion formed between Cf/C substrate and brazing seam due to sufficient reaction and further densification of the SiC layer. Although with similar shear
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strength, the brazed joints at different temperatures showed different fracture behaviors. When the brazing temperature was 1400 °C, in addition to Si and SiC, pull-out TiSi2 and carbon fiber dispersed over the fracture surface, conductive to improving mechanical properties of the joints [14,49]. The mixed fracture paths ran across brazing seam, interface and Cf/C composites, which made the Cf/C brazed joints more reliable. However, as the brazing temperature was increased to 1420 °C, the cracks parallel to interface turned into failure initiation. Such failure mainly occurred in the interface and brazing seam without the 18
ACCEPTED MANUSCRIPT reinforcement effect of the carbon fibers (insert in Fig. 10(c)). With further increase of the temperature to 1440 °C, the phases exposed in the fracture surface did not change and the refinement of TiSi2 was apparently observed among the Si substrate. But the severe flaws
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within the brazed joints led to the bulk strength degradation despite the pull-out of TiSi2 and
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carbon fibers (insert in Fig. 10(d)).
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Fig. 10. Fracture morphology of the Cf/C joints brazed at different temperatures: (a) 1380 °C; (b) 1400 °C;
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(c) 1420 °C; (d) 1440 °C.
XRD was applied to analyze the phase distribution in the fracture surface. As shown in Fig. 11, the phases of Si, TiSi2 and SiC were all detected in each period from 1380 °C to 1440 °C, which was in accordance with the analysis results in Fig. 10.
19
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☆
▽
● --C ☆--SiC ◇--Si ▽--TiSi2
◇
Intensity(a.u.)
☆
● ◇
☆☆
☆ ▽ ●
▽ ● ◇
▽
◇ ▽
◇ ☆●
1440°C
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1400°C
1380°C
30
40
50
2Theta(degree)
60
70
80
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3.4. Reaction and joining mechanism
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Fig. 11. XRD pattern of the fracture surfaces of the joints brazed at various temperatures.
Fig. 12 reveals the reaction and joining mechanism of the two-step brazed Cf/C joints through a series of schematic diagrams. During the pre-infiltration step, the Ti-Si eutectic
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alloy exhibits excellent compatibility and wettability on the Cf/C composites. An almost tight and consecutive boundary zone mainly of SiC can be fabricated with the interaction among the elements of C, Si and Ti. Such modification treatment provides an economic and simple
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hole-sealing way to properly avoid the excessive liquid impregnation into the porous
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substrate, which highly benefits the reliable bonding of the Cf/C joints in the next brazing process. The designation of gradient structure of the near-surface Cf/C composites could reinforce the starting carbon matrix firstly, then effectively minimize the impact of the CTE mismatch between the outer reaction layer and the carbon matrix. The schematic of the PI-Cf/C sample and a Ti-Si foil are shown in Fig. 12(a). Next, in the non-pressure brazing procedure, the original unreacted phases of Si and TiSi2 attached to PI-Cf/C will dissolve into the liquid Ti-Si melt (Fig. 12(b)). The adherent 20
ACCEPTED MANUSCRIPT Ti3SiC2, TiC and TiC0.98 will also decompose in the surrounding of rich Si [5]. The carbon atoms pass through the micro-pore and micro-crack space of the designed SiC layer driven by thermodynamics force and react with Ti-Si melt to further densify the interface area. Other
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unreacted Si and Ti elements form the eutectic structure in the brazed joints. The gradual CTE transition structure has formed, which can effectively reduce the residual stress derived from the thermal mismatch (Fig. 12(c)). As the brazing temperature is elevated, the growing
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SiC reaction zone, refining TiSi2 strips, stronger nail effect as well as the graded CTE
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structure contribute to the reliable bonding of the brazed joints. However, serious reaction among the C-Ti-Si system under higher temperature leads to interface degradation and enormous filler exhaustion, causing defects such as cracks and voids in the joints. An optimal
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fracture mode of the joint should contain the interface, the substrate and the brazing seam.
Fig. 12. The joining mechanism of two-step brazed Cf/C joints: (a) PI-Cf/C and the filler alloy; (b) the filler alloy melt during the second step; (c) the brazed joint.
In contrast, the direct brazing experiment was also carried out at various parameters. However, the shear strengths of the joints obtained by this process were all less than 10 MPa. All the Cf/C shear strength samples were confronted the same problem: most of the filler lost caused by the uncontrolled infiltration of the whole porous Cf/C substrate, leaving 21
ACCEPTED MANUSCRIPT insufficient filler to fill the entire brazing seam. This may lead to great variation of the joint microstructure and incomplete bonding of the brazing joints. By the two-step brazing process, on the contrary, small amount of filler alloy reacted sufficiently with the Cf/C substrates and
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high temperature phases such as SiC, TiCx were produced in the first step (Fig. 3), which could serve as a barrier to prevent excess infiltration into the pores of Cf/C composites in the next step (insert in Fig. 2(a)). Therefore, strong bonding was achieved by the two-step
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method.
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As mentioned above, the key points to strong Cf/C brazed joints depend on both preparations of the nearly tight pre-infiltration layer and designation of the gradient-CTE composite joint structure with reliable bonding. Therefore, through the novel method which combined brazing process with pre-infiltration, the achieved brazed joint can effectively
4. Conclusion
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release thermal stress and strengthen the brazed joint and the intrinsic substrate.
In this work, reliable Cf/C joints have been successfully obtained with Ti-Si eutectic
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alloy through a novel two-step brazing method for the first time. The effects of the infiltration
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and brazing parameters on the phase compositions and mechanical behaviors of the joints were investigated. The joining mechanism was elaborated in detail. The infiltration structure effectively densified and strengthened the intrinsic Cf/C substrate firstly, and the PI-Cf/C further reacted with Ti-Si alloy to form a compact joint. Si and TiSi2 eutectic structure was retained in the brazing seam, and SiC phase was formed at the interface between Ti-Si structure and Cf/C composites. Therefore, both of the two steps are essential to the CTE-gradient bonding structure and excellent brazing joint of Cf/C composites. This research 22
ACCEPTED MANUSCRIPT provides a new high temperature braze filler (Ti-Si eutectic alloy) and technique to join the
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Cf/C composites.
Acknowledgment
The authors gratefully acknowledge the financial support from the National Natural
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Science Foundation of China under Grant Nos. 51622503, U1737205 and U1537206. References
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[1] S. Dietrich, J.M. Gebert, G. Stasiuk, A. Wanner, K.A. Weidenmann, O. Deutschmann, et al., Microstructure characterization of CVI-densified carbon/carbon composites with various fiber distributions, Compos. Sci. Technol. 72 (15) (2012) 1892-1900. [2] J. Yu, C. Zhou, H. Zhang, A micro-image based reconstructed finite element model of needle-punched C/C composite, Compos. Sci. Technol. 153 (2017) 48-61. [3] V. Casalegno, T. Koppitz, G. Pintsuk, M. Salvo, S. Rizzo, S. Perero, et al., Proposal for a modified non-active brazing alloy for joining CFC composites to copper, Compos. Part B 56 (2014) 882-888. [4] Y. Cai, X. Yin, S. Fan, L. Zhang, L. Cheng, Tribological behavior of three-dimensional needled ceramic modified carbon/carbon composites in seawater conditions, Compos. Sci. Technol. 87 (2013) 50-57. [5] X. Zhou, H. Yang, F. Chen, Y.-H. Han, J. Lee, S. du, et al., Joining of carbon fiber reinforced carbon composites with Ti3SiC2 tape film by electric field assisted sintering technique, Carbon 102 (2016) 106-115. [6] M. Salvo, P. Lemoine, M. Ferraris, M. Montorsi, Joining of carbon-carbon composites for thermonuclear fusion applications, J. Am. Ceram. Soc. 80 (1) (1997) 206-212. [7] X. Song, H. Li, V. Casalegno, M. Salvo, M. Ferraris, X. Zeng, Microstructure and mechanical properties of C/C composite/Ti6Al4V joints with a Cu/TiCuZrNi composite brazing alloy, Ceram. Int. 42 (5) (2016) 6347-6354. [8] M. Singh, R. Asthana, T.P. Shpargel, Brazing of carbon–carbon composites to Cu-clad molybdenum for thermal management applications, Mater. Sci. Eng. A 452-453 (2007) 699-704. [9] M. Singh, T.P. Shpargel, G.N. Morscher, R. Asthana, Active metal brazing and characterization of brazed joints in titanium to carbon–carbon composites, Mater. Sci. Eng. A 412 (1-2) (2005) 123-128. [10] W. Fu, S.P. Hu, X.G. Song, J.X. Li, J. Cao, J.C. Feng, et al., Wettability and bonding of graphite by Sn0.3Ag0.7Cu-Ti alloys, Carbon 121 (2017) 536-543. [11] S. Rizzo, S. Grasso, M. Salvo, V. Casalegno, M.J. Reece, M. Ferraris, Joining of C/SiC composites by spark plasma sintering technique, J. Eur. Ceram. Soc. 34 (4) (2014) 903-913. [12] Z. Wang, G. Wang, M. Li, J. Lin, Q. Ma, A. Zhang, et al., Three-dimensional graphene-reinforced Cu foam interlayer for brazing C/C composites and Nb, Carbon 118 (2017) 723-730. 23
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AC C
EP
TE D
M AN U
SC
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[13] N.V. Moutis, C. Jimenez, X. Azpiroz, T. Speliotis, C. Wilhelmi, S. Messoloras, et al., Brazing of carbon–carbon composites to Nimonic alloys, J. Mater Sci. 45 (1) (2009) 74-81. [14] J. Cao, C. Li, J. Qi, Y. Shi, J. Feng, Combustion joining of carbon–carbon composites to TiAl intermetallics using a Ti–Al–C powder composite interlayer, Compos. Sci. Technol. 115 (2015) 72-79. [15] M. Singh, R. Asthana, Characterization of brazed joints of CC composite to Cu-clad-Molybdenum, Compos. Sci. Technol. 68 (14) (2008) 3010-3019. [16] V. Casalegno, M. Salvo, M. Ferraris, Surface modification of carbon/carbon composites to improve their wettability by copper, Carbon 50 (6) (2012) 2296-2306. [17] X.R. Song, H.J. Li, X. Zeng, Brazing of C/C composites to Ti6Al4V using multiwall carbon nanotubes reinforced TiCuZrNi brazing alloy, J. Alloy Compd. 664 (2016) 175-180. [18] S. Pellegrino, L. Thomé, A. Debelle, S. Miro, P. Trocellier, Radiation effects in carbides: TiC and ZrC versus SiC, Nucl. Instrum. Methods Phys. Res. Sect. B 327 (2014) 103-107. [19] M. Li, X. Zhou, H. Yang, S. Du, Q. Huang, The critical issues of SiC materials for future nuclear systems, Scr. Mater. 143 (2018) 149-153. [20] P. Dadras, T.T. Ngai, G.M. Mehrotra, Joining of carbon-carbon composites using boron and titanium disilicide interlayers, J. Am. Ceram. Soc. 80 (1) (1997) 125-132. [21] J. Li, L. Liu, Y. Wu, Z. Li, W. Zhang, W. Hu, Microstructure of high temperature Ti-based brazing alloys and wettability on SiC ceramic, Mater. Des. 30 (2) (2009) 275-279. [22] N.M. Larson, F.W. Zok, In-situ 3D visualization of composite microstructure during polymer-to-ceramic conversion, Acta Mater. 144 (2018) 579-589. [23] R. Naslain, Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview, Compos. Sci. Technol. 64 (2) (2004) 155-170. [24] H.A. Bale, A. Haboub, A.A. MacDowell, J.R. Nasiatka, D.Y. Parkinson, B.N. Cox, et al., Real-time quantitative imaging of failure events in materials under load at temperatures above 1,600°C, Nat. Mater. 12 (2012) 40. [25] D.L. Poerschke, A. Braithwaite, D. Park, F. Lauten, Crystallization behavior of polymer-derived Si-O-C for ceramic matrix composite processing, Acta Mater. 147 (2018) 329-341. [26] Y. Zeng, D. Wang, X. Xiong, X. Zhang, P. J. Withers, W. Sun, et al., Ablation-resistant carbide Zr0.8Ti0.2C0.74B0.26 for oxidizing environments up to 3,000°C, Nature commun. 8 (2017) 15836. [27] Q. Fu, F. Zhao, H. Li, H. Peng, X. Nan, A Multi-interlayer LMAS Joint of C/C–SiC Composites and LAS Glass Ceramics, J. Mater. Sci. Technol. 31 (5) (2015) 467-472. [28] Q.G. Fu, H. Peng, H.J. Li, J. Wang, Q. Zhu, A gradient LMAS interlayer joint of SiC coated C/C composites to LAS glass ceramics, Ceram. Int. 40 (1) (2014) 2461-2466. [29] S. Chen, C. Zhang, Y. Zhang, H. Hu, Preparation and properties of carbon fiber reinforced ZrC–ZrB2 based composites via reactive melt infiltration, Compos. Part B 60 (2014) 222-226. [30] F.H. Gern, R. Kochendörfer, Liquid silicon infiltration: description of infiltration dynamics and silicon carbide formation, Compos. Part A 28 (4) (1997) 355-364. [31] H. Zhou, R.N. Singh, Kinetics model for the growth of silicon carbide by the reaction of liquid silicon with carbon, J. Am. Ceram. Soc. 78 (9) (1995) 2456-2462. [32] Y. Zeng, X. Xiong, G. Li, Z. Chen, W. Sun, D. Wang, Microstructure and ablation behavior of carbon/carbon composites infiltrated with Zr–Ti, Carbon 54 (2013) 300-309. [33] E.W. Washburn, The Dynamics of Capillary Flow, Phys. Rev. 17 (3) (1921) 273-283. [34] Y. Wang, X. Zhu, L. Zhang, L. Cheng, Reaction kinetics and ablation properties of C/C–ZrC composites fabricated by reactive melt infiltration, Ceram. Int. 37 (4) (2011) 1277-1283. 24
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[35] J. Yang, O.J. Ilegbusi, Kinetics of silicon–metal alloy infiltration into porous carbon, Compos. Part A 31 (6) (2000) 617-625. [36] R. Asthana, Interface- and diffusion-limited capillary rise of reactive melts with a transient contact angle, Metall. Mater. Trans. A (2002) 2119-2128. [37] R. Asthana, Dissolutive capillary penetration with expanding pores and transient contact angles, J. Colloid Interface Sci. 231 (2000) 398-400. [38] R. Asthana, On reactive capillary flows with a dynamic contact angle, J. Mater. Processing Manuf. Sci., 10 (4) (2002) 187-197. [39] Y. Zeng, X. Xiong, D. Wang, L. Wu, Infiltration mechanism and factors influencing carbon/carbon– Zr–Ti–C composites prepared by liquid metal infiltration, J. Mater. Process Technol. 214 (12) (2014) 3150-3157. [40] Y. Chang, W. Sun, X. Xiong, Z. Chen, Y. Wang, Z. Hao, et al., A novel design of Al-Cr alloy surface sealing for ablation resistant C/C-ZrC-SiC composite, J. Eur. Ceram. Soc. 37 (2) (2017) 859-864. [41] T. Lin, M. Yang, P. He, C. Huang, F. Pan, Y. Huang, Effect of in situ synthesized TiB whisker on microstructure and mechanical properties of carbon–carbon composite and TiBw/Ti–6Al–4V composite joint, Mater. Des. 32 (8-9) (2011) 4553-4558. [42] C.A. Swenson, Recommended Values for the Thermal Expansivity of Silicon from 0 to 1000 K, J. Phys. Chem. Ref. Data 12 (2) (1983) 179-182. [43] T.S.R.C. Murthy, C. Subramanian, R.K. Fotedar, M.R. Gonal, P. Sengupta, S. Kumar, et al., Preparation and property evaluation of TiB2+TiSi2 composite, Int. J. Refract. Met. Hard Mater. 27 (3) (2009) 629-636. [44] Y. Zeng, X. Xiong, D. Wang, L. Wu, Residual thermal stresses in carbon/carbon–Zr–Ti–C composites and their effects on the fracture behavior of composites with different preforms, Carbon 81 (2015) 597-606. [45] C. Racault, F. Langlais, R. Naslain, Solid-state synthesis and characterization of the ternary phase Ti3SiC2, J. Mater. Sci. 29 (13) (1994) 3384-3392. [46] M. Jiang, H. Y. Xiao, H. B. Zhang, S. M. Peng, C. H. Xu, Z. J. Liu, et al., A comparative study of low energy radiation responses of SiC, TiC and ZrC, Acta Mater. 110 (2016) 192-199. [47] H. Wang, J. Cao, J. Feng, Brazing mechanism and infiltration strengthening of CC composites to TiAl alloys joint, Scr. Mater. 63 (8) (2010) 859-862. [48] K.K. Chawla, Composite Materials: Science and Engineering, Springer, New York (2012) 249-292. [49] J. Wang, K. Li, W. Li, H. Li, Z. Li, L. Guo, The preparation and mechanical properties of carbon/carbon composite joints using Ti–Si–SiC–C filler as interlayer, Mater Sci. Eng. A 574 (2013) 37-45.
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