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SiBCN composite with CuPd-V filler alloy
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Joining of Cf /SiBCN composite with CuPd-V filler alloy Wen-Wen Li a , Bo Chen a , La-Mei Cao b , Wei Liu b , Hua-Ping Xiong a,∗ , Yao-Yong Cheng a a b
Welding and Plastic Forming Division, Beijing Institute of Aeronautical Materials, Beijing 100095, China Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, China
a r t i c l e
i n f o
Article history: Received 20 March 2017 Received in revised form 10 May 2017 Accepted 6 June 2017 Available online xxx Keywords: Cf /SiBCN composite Wettability Joint strength Interfacial reaction
a b s t r a c t Two compositions of CuPd-V system filler alloy were designed for joining the Cf /SiBCN composite. Their dynamic wettability on the Cf /SiBCN composite was studied with the sessile drop method. The CuPd-8V alloy exhibited a contact angle of 57◦ after holding at 1170 ◦ C for 30 min, whereas for CuPd-13V alloy, a lower contact angle of 28◦ can be achieved after heating at 1200 ◦ C for 20 min. Sound Cf /SiBCN joints were successfully produced using the latter filler alloy under the brazing condition of (1170-1230)◦ C for 10 min. The results showed that the active element V strongly diffused to the surface of Cf /SiBCN composite, with the formation of V2 C/VN reaction layer. The microstructure in the central part of the joint brazed at 1200 ◦ C was characterized by the V2 C/VN particles distributing scatteringly in CuPd matrix. The corresponding joints showed the maximum three-point bend strength of 82.4 MPa at room temperature. When the testing temperature was increased to 600 ◦ C, the joint strength was even elevated to 108.8 MPa. Furthermore, the joints exhibited the strength of 92.4 MPa and 39.8 MPa at 800 ◦ C and 900 ◦ C, respectively. © 2018 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction As we know, ceramics based on silicon nitride and carbide has superior high-temperature strength and oxidation resistance. But silicon nitride decomposes at about 1400 ◦ C in vacuum and 1775 ◦ C in 0.1 MPa nitrogen, limiting the high-temperature range of its technological uses [1,2]. And for SiC ceramic, when the temperature exceeded 1500 ◦ C, recrystallization and grain coarsening would cause significant performance deterioration. However, a boroncontaining silicon nitride/carbide ceramic (Si3.0 B1.0 C4.3 N2.0 ) would not degrade at temperatures up to 2000 ◦ C even in nitrogen-free environments [3]. And other studies have reported that amorphous silicoboron carbonitride ceramic frequently resists thermal decomposition even at temperatures exceeding 1800 ◦ C [4,5]. Compared with SiC or Si3 N4 ceramics, the oxidation resistance at 1700 ◦ C of Si-B-C-N quaternary ceramic is the most outstanding [6]. Cf /SiBCN composite, which consists of reinforcing carbon fibers within the Si-B-C-N ceramic matrix, should have superior properties compared with any other single materials [7,8]. As a result, the Cf /SiBCN composite has been a promising material for high-temperature applications, including aeronautical turbines,
∗ Corresponding author. E-mail address: [email protected] (H.-P. Xiong).
advanced rocket propulsion thrust chambers, aircraft nose cone, leading edge and etc. [9]. The practical use of the composite requires development of appropriate joining techniques. In many cases, the joining of the composite to itself or to metals is also needed, and the joints must retain the good high-temperature property. Nevertheless, the research on the Si-B-C-N ceramic has just started over the last two decades, and the reports about its joining are rather lacking. AuCuTi filler alloy has been used to join the 2Si-B-3C-N ceramic with the shear strength of 138 MPa [10]. However, generally, the service temperature of AgCuTi system brazing alloys is limited to 500 ◦ C. In the previous work, Cf /SiBCN composite was brazed with Ni-based filler alloys with the bend strength of 62.9 MPa [11]. But regretfully, no further study on the high-temperature strength was conducted. According to the previous references [12,13], in amorphous Si-B-C-N at high temperature the phases of SiC and Si3 N4 easily crystallized to precipitate. Therefore, it is likely possible to join Cf /SiBCN composite with the filler metals, which can be used for brazing SiC and Si3 N4 ceramics or Cf /SiC composite. To our knowledge, vanadium can be used as an active element for ceramic joining [14–17], and new progress has been made in the development of V-containing filler alloys for Si3 N4 ceramic joining [18,19]. And the possibility for joining AlN ceramic by using AuPdCoNi-V [20] filler alloy was also verified, with V2 N reaction layer formed at the surface of AlN ceramic. For the joining of Cf /SiC composite, V played the interfacial activity to form the VC0.75 reaction layer and the room-
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Table 1 Composition, liquidus temperature, heating condition in wetting experiment and contact angle of two filler alloys on Cf /SiBCN. Filler alloy
Composition (wt %)
Liquidus temperature (◦ C)
Heating condition
Contact angle (deg.)
CuPd-8V CuPd-13V
Cu-(35–42)Pd-8V Cu-(35–45)Pd-13V
1134.4 1155.6
1170 ◦ C/30 min 1200 ◦ C/20 min
57 28
Fig. 1. Wetting morphologies of two kinds of filler alloys on Cf /SiBCN composite: (a, c) CuPd-8V (1170 ◦ C/30 min); (b, d) CuPd-13V (1200 ◦ C/20 min).
temperature strength of 135 MPa can be maintained up to 800 ◦ C [21]. According to the previous report, firstly, two CuPd-V filler alloys were contrasted in their dynamic wettability on Cf /SiBCN composite substrate, and on this basis the joining of the Cf /SiBCN composite was studied using one of the filler alloys. The effects of brazing temperature on the joint microstructure and strength were investigated. In addition, the high-temperature jointing strength was also studied. 2. Experimental procedure Two compositions of CuPd-V system filler alloys were designed, as listed in Table 1. The two filler alloys can be fabricated into brazing foil with the thickness of 100 m by alternating roomtemperature rolling process and vacuum annealing treatment, attributing to the good plasticity. The liquidus temperatures of two filler alloys were measured by differential thermal analysis (DTA) method, and dynamic wettability of the two filler alloys on Cf /SiBCN composite was studied with the sessile drop method, and the results are listed in Table 1. The used Cf /SiBCN composite substrate was carbon fiber reinforced amorphous SiBCN matrix composite, with about 10 vol. % porosity and about 50% carbon fibers. This material is prepared by polymer infiltration pyrolysis (PIP) process [22]. For dynamic wetting experiment, five CuPd-V foils with a diameter of 3 mm were prepared, and then 5-layer foils were laminated by low-electric-current spot welding. Afterwards the
laminated filler plate was placed on the Cf /SiBCN substrate (10 mm × 10 mm × 3 mm). The chamber of the furnace was evacuated to (3.1–3.5) × 10−3 Pa at room temperature, and when heated to 900 ◦ C high-purity (99.999%) argon was put into the chamber with a pressure of 0.1 atm. Then the assembly was heated to the wetting temperature at a heating rate of 10 ◦ C/min. During the heating process, the morphologies of the molten droplet were recorded dynamically by taking photos at intervals. The contact angles were thus calculated from the droplet images based on the Laplace equation. The accuracy of the contact angle measurements should be within ±1◦ . The composite bars to be joined were 3 × 4 × 20 mm3 and the total length of the butt joint was 40 mm for three-point bend test. Three brazing temperatures of 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, were used in the brazing experiment, but the dwelling time was fixed at 10 min. The heating rate of the brazing experiment was 10 ◦ C/min. After brazing the samples were cooled to 500 ◦ C at a cooling rate of 5 ◦ C/min, followed by furnace cooling to room temperature. During the heat process, the vacuum of furnace was kept between (3.0–7.0) × 10−3 Pa. The joint strengths at room temperature and elevated temperatures were determined by the three-point bend test. The joint microstructures were analyzed by a scanning electron microscope (SEM) equipped with an X-ray wave-length dispersion spectrometer (XWDS) and an X-ray energy-dispersive spectrometer (EDS). After the bend test the morphology and the typical microzones on the fracture surface were characterized by SEM and EDS. Furthermore, the X-ray diffraction (XRD) spectrometer with
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Fig. 2. Change of contact angle of CuPd-8V filler alloy with dwelling time at 1170 ◦ C.
CuK˛ radiation at 30 kV and 20 mA was used for determination of the phases on the fracture surface, with the scanning range between 10◦ and 90◦ . 3. Results and discussion The wetting morphologies of the two filler alloys are shown in Fig. 1. After holding at 1170 ◦ C for 30 min the contact angle of CuPd8V alloy on the Cf /SiBCN composite was 57◦ . For the filler alloy of CuPd-13V, the contact angle showed a lower value of 28◦ after holding at 1200 ◦ C for 20 min. It is not strange that the higher concentration of active element V corresponds to the better wettability. Fig. 2 presents the change of contact angle of CuPd-8V filler alloy with holding time at 1170 ◦ C. The results show that this filler alloy does not melt until holding at 1170 ◦ C for 15 min. Afterwards, the contact angle decreases gradually, and reaches an equilibrium value of 57◦ with the dwelling time of 30 min (Fig. 2). The wetting kinetics of CuPd-13V filler alloy is presented in Fig. 3. It is found that when heated to 1100 ◦ C, the contact angle was about 83◦ (Fig. 3(a)). Before heating up to 1170 ◦ C, the contact angle shows almost no change. As temperature is further increased, the contact angle is decreased greatly. When heated to 1200 ◦ C, the contact angle is decreased to 43.9◦ (Fig. 3(a)). The holding time of 20 min is given to the sample at 1200 ◦ C. Unlike the high spreading rate at the range from 1170 ◦ C to 1200 ◦ C, it shows slow wetting kinetics when holding at 1200 ◦ C (Fig. 3(b)). The contact angle is decreased slowly until to an equilibrium value of 28◦ with the total holding time of 20 min, as shown in Fig. 1(b) and (d). Microstructure of the Cf /SiBCN composite joint brazed with CuPd-8V filler alloy at 1170 ◦ C for 10 min is shown in Fig. 4. It is clear that two reaction layers V(C,N) with a thickness of about
3
1 m are formed at the surface of Cf /SiBCN composite. Two kinds of solid solutions CuPd and Cu3 Pd coexisted in the central part of the joint. Though good bonding was achieved the joint only exhibited a low bend strength of 58.1 MPa at room temperature. The too thin reaction layer would cause the weak interfacial bonding. Compared with the CuPd-8V filler alloy, CuPd-13V alloy showed better wettability on the Cf /SiBCN composite. The main reason was that when the concentration of element V was low, the activity of element V in the interfacial reaction with Cf /SiBCN was insufficient (the thickness of the reaction layer was too thin). As a consequence, it was difficult to achieve a strong bonding. In order to ensure desirable wettability and thus improve the bonding quality at the joining interface, this work focused on the CuPd-13V filler alloy for the Cf /SiBCN joining. Fig. 5 presents the backscattered electron images of the Cf /SiBCN joints brazed with CuPd-13V filler alloy with different brazing temperatures. It is deduced that evident reactions between the CuPd-13V filler alloy and the Cf /SiBCN composite occurred as the two reaction layers (RL) had been formed at the surfaces of the joined composite. Furthermore, with the increase of the brazing temperature, the thickness of the reaction layer was increased, for example at brazing temperature of 1170 ◦ C, the average thickness of the reaction layer (RL “1” and “2”) was only 2.5 m. When brazed at 1200 ◦ C, reaction layer (RL“3” and “4”) showed a thickness of 3.5 m, and that was increased to 6.7 m (RL“5” and “6”) at the brazing temperature of 1230 ◦ C. According to the EDS analysis results listed in Table 2, the reaction layer (labeled “1”–“6”) was mainly composed of V-(C,N) compounds, enriching with elements V, C and N. The active element V diffused strongly to the surface of the Cf /SiBCN composite, as shown in the elemental area distribution maps of the joint brazed at 1200 ◦ C (Fig. 6(c)). The XRD patterns of the fracture surface verified that the interfacial reaction products consisted of V2 C and VN compounds (Fig. 7). In the central part of the joint brazed at 1170 ◦ C (Fig. 5(a)), there were two phases labeled “7” and “8”, which were enriched with elements Cu and Pd. Firstly a grey phase of Cu3 Pd (microzone “7”) was composed of 62.42 at. % Cu and 19.92 at. % Pd. Secondly, a grayish white phase marked as microzone “8” can be deduced as CuPd phase in which the concentration of elements Cu and Pd was 46.17 at. % and 38.04 at. % Pd, respectively. From Fig. 5(b) and (c), when rising the brazing temperature, besides the increase of the thickness of the reaction layer, the microstructure in the central part of the joint was also changed to some extent. At first, both the microzones “9” and “10” in Fig. 5(b), which can be confirmed as CuPd phase by the EDS analysis results, indicated that only one kind of solid solution existed. As shown in Fig. 6(a) and (b), the elements Cu and Pd segregated in the central part of the joint. Similarly, microzone “11” in Fig. 5(c) for brazing temperature of 1230 ◦ C was also the CuPd phase. Therefore with the
Fig. 3. Wetting kinetics of the CuPd-13V alloy on Cf /SiBCN composite: (a) change of contact angle with temperature, (b) change of contact angle with holding time at 1200 ◦ C.
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Fig. 4. Microstructure of the Cf /SiBCN composite joint brazed with CuPd-8V filler alloy at 1170 ◦ C for 10 min: (a) at a low magnification, and (b) at a high magnification.
Fig. 5. Microstructure of the Cf /SiBCN joint brazed with CuPd-13V filler alloy at 1170 ◦ C (a), 1200 ◦ C (b), and 1230 ◦ C (c).
increase of the brazing temperature, the inter-diffusion between elements Cu and Pd became strong so that only CuPd phase existed within the joint. The XRD analysis result (Fig. 7) confirmed that both CuPd and Cu3 Pd phases co-existed at the fracture surface of the joint brazed at 1170 ◦ C (Fig. 7(a)). Whereas, the diffraction peaks of Cu3 Pd phase disappeared at the brazing temperature of 1200 ◦ C (Fig. 7(b)).
For three joints brazed at 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, there were black particles distributing scatteringly in the central part of the joint marked as “12”, “13” and “14” in Fig. 5. Fig. 6(c) signified that V segregated here and V2 C/VN compounds should have been formed according to the EDS results and the XRD patterns. Therefore, during the brazing process, active element V would firstly diffuse from the CuPd-13V filler alloy to the surface of the
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Table 2 EDS analysis results for the typical microzones in Fig. 5. Microzone
1 2 3 4 5 6 7 8 9 10 11 12 13 14
Composition (at. %)
Deduced phase
Cu
Pd
V
Si
B
C
N
0.32 0.24 0.35 0.28 0.28 0.05 62.42 46.17 39.08 39.78 33.54 3.29 3.67 0.49
0.29 0.22 0.31 0.20 0.69 0.07 19.92 38.04 34.50 34.21 45.25 3.32 3.62 2.34
43.93 46.81 45.72 51.82 42.71 42.32 1.14 1.41 5.10 1.58 0.41 41.08 43.51 38.05
2.99 0.18 1.07 0.31 0.12 0.12 1.41 7.61 6.13 6.37 6.50 0.39 0.29 0.34
0.22 / 1.10 / 4.84 5.67 6.91 1.70 7.43 8.82 5.24 1.21 2.03 5.02
18.02 29.06 29.14 33.38 40.57 39.59 5.47 3.75 4.81 5.48 5.40 34.69 35.62 35.67
34.23 23.49 22.31 14.01 10.79 12.18 2.73 1.32 2.95 3.76 3.66 16.02 11.26 18.09
V-(C,N)
Cu3 Pd CuPd
V-(C,N)
Fig. 6. Elemental area distribution maps for Fig. 5(b): (a) Cu, (b) Pd, (c) V, (d) Si, (e) B, (f) C, and (g) N.
Fig. 7. XRD pattern of the fracture surface after bend test: (a) brazing temperature of 1170 ◦ C, (b) brazing temperature of 1200 ◦ C.
Cf /SiBCN composite, and react with Cf /SiBCN through the following formulae: 2V + C = V2 C
(1)
V + N = VN
(2)
Fukai et al. [23] studied the SiC/V joint at (1200-1400)◦ C, and reported that element V reacted with SiC to form V2 C and silicides. In this work, V2 C was detected as the main reaction product within the joint (Fig. 7). In the previous study, when joining Si3 N4 ceramic and Cf /SiC composite using Pd- and V- containing filler alloy, Pd2 Si compounds were observable in the corresponding joints [19,21]. However, here, Pd2 Si compound is absent for the joints brazed at three different temperatures. Compared with the previous Cf /SiC composite, the concentration of Si in the amorphous 3Si-1B-4C-2N
ceramic was remarkably lower than that in the SiC ceramic matrix, and this might be the main reason for the absence of Pd-silicides within the Cf /SiBCN composite joint. According to the V-N binary phase diagram, both VN and V2 N are the stable phases in system of V-N [24]. In this study, only VN compound was as the reaction product. In addition, based on the calculation (f G (VN) = −434.3 + 0.1688T (kJ/mol)) the free energy of formation for VN was −185.7 kJ/mol at 1200 ◦ C [18], indicating that the VN compound can be formed spontaneously. Fig. 8 presents the joint appearances for the three kinds of the joints brazed at 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, respectively, and the correspondingly room-temperature strength of the joints is shown in Fig. 9. The room-temperature strength of the joint brazed at 1170 ◦ C was 57.8 MPa, but it was increased to 82.4 MPa at the brazing temperature of 1200 ◦ C. For the joint brazed at 1170 ◦ C, the
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Fig. 8. Joint appearances for the three kinds of the joints brazed at 1170 ◦ C (a), 1200 ◦ C (b) and 1230 ◦ C (c).
Fig. 9. Room-temperature strength of the joints brazed at different temperature.
thickness of the reaction layer was only 2.5 m (Fig. 5(a)), whereas that was increased to 3.5 m (Fig. 5(b)) when brazing at 1200 ◦ C. If the reaction layer was too thin, the joint strength would decrease due to the insufficient reactivity [25]. In addition, in the case of 1200 ◦ C the V-(C,N) compounds distributed more uniformly in the central part of the joint, with effective strengthening of the joint.
With the brazing temperature up to 1230 ◦ C, the thickness of the reaction layer was increased to 6.7 m, nearly twice of that at 1200 ◦ C. And the obvious cracks appeared within the reaction layer. Consequently, the joint strength was decreased to 62.9 MPa. The relationship between the reaction layer thickness and the defects produced at the reaction region was studied quantitatively by Nakao et al. [26]. They pointed out that when the reaction layer exceeded the optimum thickness, some defects such as microvoids and cracks might be produced at the reaction layer. In general, the mechanical and physical properties of the reaction layer (V2 C,VN) were very different from those of Cf /SiBCN composite; for example, the coefficient of thermal expansion (CTE) for VN was (9–10) × 10−6 K−1 , whereas that of SiBCN matrix ceramic to be joined was about 3.18 × 10−6 K−1 [27]. Undoubtedly, when the reaction layer was relatively thick [28], the residual stresses within the joint would accumulate to cause the decrease of the joint strength. The fractured surface morphologies of the joint brazed at 1200 ◦ C are shown in Fig. 10, and the corresponding EDS analysis results for typical microzones are listed in Table 3. The fracture surface presented double-layered characterization, and the cracking path was evident, and the Cf /SiBCN base material was exposed on the two layered surface, as shown in Fig. 10(a) and (b). For the mag-
Fig. 10. Fracture surface of the Cf /SiBCN joint brazed at 1200 ◦ C (a), the magnification image for area S1 (b), the magnification image for area S2 (c).
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Table 3 EDS analysis result for the typical microzones in Fig. 10. Microzone
1 2
Composition (at. %)
Deduced phase
Cu
Pd
V
Si
B
C
N
31.25 1.54
20.99 0.95
22.04 36.16
3.52 4.22
3.49 5.18
8.46 40.28
10.26 11.67
Fig. 11. Model of the crack propagation when subjected to the bend test.
CuPd/V-(C,N) V-(C,N)
can be regarded as ductile buffer to release the residual stress through their plastic deformation [30]. As a result, at the high testing temperatures the joint strength can be elevated to some extent. Furthermore, the V-C and V-N were typical transition metal carbide/nitride, which have high melting point, high hardness, high strength at high temperatures [31,32]. Therefore, the high strength of the joints should also be attributed to the strengthening effect by the V2 C/VN particles which distributed scatteringly in the central part of the joint. But when the temperature increased to 900 ◦ C, the joint strength was decreased sharply due to the softening of the CuPd phase. 4. Conclusions
Fig. 12. Effect of test temperature on three point bend strength of the Cf /SiBCN composite joint brazed at 1200 ◦ C.
nification of selected area S2, both the CuPd phase and the V-(C,N) compound were detected. It is reasonable to deduce that when subjected to the bend test, the crack initiated at or near the V-(C,N) compound layer at the surface of the Cf /SiBCN composite. On one hand, the formation of the V-(C,N) layer played an important role in the wetting and joining of the Cf /SiBCN composite. However, on the other hand, it seems that the V-(C,N) reaction layer was also the weak link of the Cf /SiBCN-Cf /SiBCN joint. Afterwards, the crack propagated through the brazing seam to the other interface between the Cf /SiBCN composite and the filler alloy, followed by the final failure. The model of the crack propagation within the joint is illustrated in Fig. 11. As shown in Fig. 12, the newly designed CuPd-13V filler alloy can provide the joint strength of 82.4 MPa at room temperature. When the testing temperature was increased to 600 ◦ C, the joint strength was even elevated to 108.8 MPa. Though further increase of the testing temperature to 800 ◦ C resulted in a decrease of the joint strength to 92.4 MPa, it was still at the same level of the room-temperature strength. However, when the testing temperature reached 900 ◦ C, the joint strength was sharply decreased to 39.8 MPa. It is well known that, in the joining of ceramic to ceramic, interfacial thermal mismatch exists between the brazing alloy and the ceramic. In the case of Cf /SiBCN-Cf /SiBCN joint, the V-(C,N) reaction layer can not effectively release the thermal stresses within the joint. Only a soft central layer can secure the stresses relax [29]. At the elevated temperatures of (600-800)◦ C, CuPd phases
(1) CuPd-8V filler alloy exhibited a contact angle of 57◦ after 1170 ◦ C for 30 min, and the CuPd-13V filler alloy showed a lower contact angle of 28◦ on Cf /SiBCN after heating at 1200 ◦ C for 20 min. The higher concentration of active element V corresponded to the better wettability. (2) Element V played an active role in the interfacial reactions, with V2 C and VN compounds formed at the surface of Cf /SiBCN composite. In the central part of the joint brazed with CuPd-13V filler alloy, V2 C/VN compounds distributed scatteringly in the CuPd solid solution. When brazed at 1170 ◦ C, the thickness of V-(C,N) reaction layer was only 2.5 m. With the brazing temperature increased to 1200 ◦ C and 1230 ◦ C, the thickness of the reaction layer was increased to 3.5 m and 6.7 m, respectively. (3) The joint brazed with CuPd-13V filler alloy showed the maximum strength of 82.4 MPa at room temperature, whereas that brazed with CuPd-8V filler was 58.1 MPa. For the joints brazed with CuPd-13V filler alloy, the joint strength was even elevated to 108.8 MPa when the testing temperature increased to 600 ◦ C, and the room-temperature strength value can be maintained up to 800 ◦ C. Acknowledgements This research work was financially supported by the National Natural Science Foundation of China (Grant Nos. 59905022, 50475160 and 51275497), and the Aeronautical Science Foundation of China (Grant No. 2008 ZE21005). References [1] H.D. Batha, E.D. Whitney, J. Am. Ceram. Soc. 56 (1973) 365–373. [2] R. Raj, J. Am. Ceram. Soc. 76 (1993) 2147–2174. [3] R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill, F. Aldinger, Nature 382 (1996) 796–798. [4] M. Jansen, B. Jaschke, T.J. Jaschke, Bond 101 (2002) 137–191. [5] A. Zern, J. Mayer, N. Janakiraman, M. Weinmann, J. Bill, M. Rühle, J. Eur. Ceram. Soc. 22 (2002) 1621–1629. [6] M. Weinmann, J. Schuhmacher, H. Kummer, S. Prinz, J.Q. Peng, H.J. Seifert, M. Christ, K. Muller, J. Bill, F. Aldinger, Chem. Mater. 12 (2000) 623–632. [7] X. Hernandez, C. Jimenez, K. Mergia, P. Yialouris, S. Messoloras, V. Liedtke, C. Wilhelmi, J. Barcena, J. Mater. Eng. Perform. 23 (2014) 3069–3076. [8] M. Weinmann, T.W. Kamphowe, J. Schuhmacher, K. Muller, F. Aldinger, Chem. Mater. 12 (2000) 2112–2122. [9] P. Baldus, M. Jansen, D. Sporn, Science 285 (1999) 699–703. [10] R. Pan, T.S. Lin, P. He, H.M. Wei, Z.H. Yang, Y.X. Shen, J. Chin. Ceram. Soc. 43 (2015) 1719–1724 (in Chinese).
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Joining of Cf /SiBCN composite with CuPd-V filler alloy Wen-Wen Li a , Bo Chen a , La-Mei Cao b , Wei Liu b , Hua-Ping Xiong a,∗ , Yao-Yong Cheng a a b
Welding and Plastic Forming Division, Beijing Institute of Aeronautical Materials, Beijing 100095, China Science and Technology on Advanced High Temperature Structural Materials Laboratory, Beijing Institute of Aeronautical Materials, Beijing 100095, China
a r t i c l e
i n f o
Article history: Received 20 March 2017 Received in revised form 10 May 2017 Accepted 6 June 2017 Available online xxx Keywords: Cf /SiBCN composite Wettability Joint strength Interfacial reaction
a b s t r a c t Two compositions of CuPd-V system filler alloy were designed for joining the Cf /SiBCN composite. Their dynamic wettability on the Cf /SiBCN composite was studied with the sessile drop method. The CuPd-8V alloy exhibited a contact angle of 57◦ after holding at 1170 ◦ C for 30 min, whereas for CuPd-13V alloy, a lower contact angle of 28◦ can be achieved after heating at 1200 ◦ C for 20 min. Sound Cf /SiBCN joints were successfully produced using the latter filler alloy under the brazing condition of (1170-1230)◦ C for 10 min. The results showed that the active element V strongly diffused to the surface of Cf /SiBCN composite, with the formation of V2 C/VN reaction layer. The microstructure in the central part of the joint brazed at 1200 ◦ C was characterized by the V2 C/VN particles distributing scatteringly in CuPd matrix. The corresponding joints showed the maximum three-point bend strength of 82.4 MPa at room temperature. When the testing temperature was increased to 600 ◦ C, the joint strength was even elevated to 108.8 MPa. Furthermore, the joints exhibited the strength of 92.4 MPa and 39.8 MPa at 800 ◦ C and 900 ◦ C, respectively. © 2018 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction As we know, ceramics based on silicon nitride and carbide has superior high-temperature strength and oxidation resistance. But silicon nitride decomposes at about 1400 ◦ C in vacuum and 1775 ◦ C in 0.1 MPa nitrogen, limiting the high-temperature range of its technological uses [1,2]. And for SiC ceramic, when the temperature exceeded 1500 ◦ C, recrystallization and grain coarsening would cause significant performance deterioration. However, a boroncontaining silicon nitride/carbide ceramic (Si3.0 B1.0 C4.3 N2.0 ) would not degrade at temperatures up to 2000 ◦ C even in nitrogen-free environments [3]. And other studies have reported that amorphous silicoboron carbonitride ceramic frequently resists thermal decomposition even at temperatures exceeding 1800 ◦ C [4,5]. Compared with SiC or Si3 N4 ceramics, the oxidation resistance at 1700 ◦ C of Si-B-C-N quaternary ceramic is the most outstanding [6]. Cf /SiBCN composite, which consists of reinforcing carbon fibers within the Si-B-C-N ceramic matrix, should have superior properties compared with any other single materials [7,8]. As a result, the Cf /SiBCN composite has been a promising material for high-temperature applications, including aeronautical turbines,
∗ Corresponding author. E-mail address: [email protected] (H.-P. Xiong).
advanced rocket propulsion thrust chambers, aircraft nose cone, leading edge and etc. [9]. The practical use of the composite requires development of appropriate joining techniques. In many cases, the joining of the composite to itself or to metals is also needed, and the joints must retain the good high-temperature property. Nevertheless, the research on the Si-B-C-N ceramic has just started over the last two decades, and the reports about its joining are rather lacking. AuCuTi filler alloy has been used to join the 2Si-B-3C-N ceramic with the shear strength of 138 MPa [10]. However, generally, the service temperature of AgCuTi system brazing alloys is limited to 500 ◦ C. In the previous work, Cf /SiBCN composite was brazed with Ni-based filler alloys with the bend strength of 62.9 MPa [11]. But regretfully, no further study on the high-temperature strength was conducted. According to the previous references [12,13], in amorphous Si-B-C-N at high temperature the phases of SiC and Si3 N4 easily crystallized to precipitate. Therefore, it is likely possible to join Cf /SiBCN composite with the filler metals, which can be used for brazing SiC and Si3 N4 ceramics or Cf /SiC composite. To our knowledge, vanadium can be used as an active element for ceramic joining [14–17], and new progress has been made in the development of V-containing filler alloys for Si3 N4 ceramic joining [18,19]. And the possibility for joining AlN ceramic by using AuPdCoNi-V [20] filler alloy was also verified, with V2 N reaction layer formed at the surface of AlN ceramic. For the joining of Cf /SiC composite, V played the interfacial activity to form the VC0.75 reaction layer and the room-
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Table 1 Composition, liquidus temperature, heating condition in wetting experiment and contact angle of two filler alloys on Cf /SiBCN. Filler alloy
Composition (wt %)
Liquidus temperature (◦ C)
Heating condition
Contact angle (deg.)
CuPd-8V CuPd-13V
Cu-(35–42)Pd-8V Cu-(35–45)Pd-13V
1134.4 1155.6
1170 ◦ C/30 min 1200 ◦ C/20 min
57 28
Fig. 1. Wetting morphologies of two kinds of filler alloys on Cf /SiBCN composite: (a, c) CuPd-8V (1170 ◦ C/30 min); (b, d) CuPd-13V (1200 ◦ C/20 min).
temperature strength of 135 MPa can be maintained up to 800 ◦ C [21]. According to the previous report, firstly, two CuPd-V filler alloys were contrasted in their dynamic wettability on Cf /SiBCN composite substrate, and on this basis the joining of the Cf /SiBCN composite was studied using one of the filler alloys. The effects of brazing temperature on the joint microstructure and strength were investigated. In addition, the high-temperature jointing strength was also studied. 2. Experimental procedure Two compositions of CuPd-V system filler alloys were designed, as listed in Table 1. The two filler alloys can be fabricated into brazing foil with the thickness of 100 m by alternating roomtemperature rolling process and vacuum annealing treatment, attributing to the good plasticity. The liquidus temperatures of two filler alloys were measured by differential thermal analysis (DTA) method, and dynamic wettability of the two filler alloys on Cf /SiBCN composite was studied with the sessile drop method, and the results are listed in Table 1. The used Cf /SiBCN composite substrate was carbon fiber reinforced amorphous SiBCN matrix composite, with about 10 vol. % porosity and about 50% carbon fibers. This material is prepared by polymer infiltration pyrolysis (PIP) process [22]. For dynamic wetting experiment, five CuPd-V foils with a diameter of 3 mm were prepared, and then 5-layer foils were laminated by low-electric-current spot welding. Afterwards the
laminated filler plate was placed on the Cf /SiBCN substrate (10 mm × 10 mm × 3 mm). The chamber of the furnace was evacuated to (3.1–3.5) × 10−3 Pa at room temperature, and when heated to 900 ◦ C high-purity (99.999%) argon was put into the chamber with a pressure of 0.1 atm. Then the assembly was heated to the wetting temperature at a heating rate of 10 ◦ C/min. During the heating process, the morphologies of the molten droplet were recorded dynamically by taking photos at intervals. The contact angles were thus calculated from the droplet images based on the Laplace equation. The accuracy of the contact angle measurements should be within ±1◦ . The composite bars to be joined were 3 × 4 × 20 mm3 and the total length of the butt joint was 40 mm for three-point bend test. Three brazing temperatures of 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, were used in the brazing experiment, but the dwelling time was fixed at 10 min. The heating rate of the brazing experiment was 10 ◦ C/min. After brazing the samples were cooled to 500 ◦ C at a cooling rate of 5 ◦ C/min, followed by furnace cooling to room temperature. During the heat process, the vacuum of furnace was kept between (3.0–7.0) × 10−3 Pa. The joint strengths at room temperature and elevated temperatures were determined by the three-point bend test. The joint microstructures were analyzed by a scanning electron microscope (SEM) equipped with an X-ray wave-length dispersion spectrometer (XWDS) and an X-ray energy-dispersive spectrometer (EDS). After the bend test the morphology and the typical microzones on the fracture surface were characterized by SEM and EDS. Furthermore, the X-ray diffraction (XRD) spectrometer with
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Fig. 2. Change of contact angle of CuPd-8V filler alloy with dwelling time at 1170 ◦ C.
CuK˛ radiation at 30 kV and 20 mA was used for determination of the phases on the fracture surface, with the scanning range between 10◦ and 90◦ . 3. Results and discussion The wetting morphologies of the two filler alloys are shown in Fig. 1. After holding at 1170 ◦ C for 30 min the contact angle of CuPd8V alloy on the Cf /SiBCN composite was 57◦ . For the filler alloy of CuPd-13V, the contact angle showed a lower value of 28◦ after holding at 1200 ◦ C for 20 min. It is not strange that the higher concentration of active element V corresponds to the better wettability. Fig. 2 presents the change of contact angle of CuPd-8V filler alloy with holding time at 1170 ◦ C. The results show that this filler alloy does not melt until holding at 1170 ◦ C for 15 min. Afterwards, the contact angle decreases gradually, and reaches an equilibrium value of 57◦ with the dwelling time of 30 min (Fig. 2). The wetting kinetics of CuPd-13V filler alloy is presented in Fig. 3. It is found that when heated to 1100 ◦ C, the contact angle was about 83◦ (Fig. 3(a)). Before heating up to 1170 ◦ C, the contact angle shows almost no change. As temperature is further increased, the contact angle is decreased greatly. When heated to 1200 ◦ C, the contact angle is decreased to 43.9◦ (Fig. 3(a)). The holding time of 20 min is given to the sample at 1200 ◦ C. Unlike the high spreading rate at the range from 1170 ◦ C to 1200 ◦ C, it shows slow wetting kinetics when holding at 1200 ◦ C (Fig. 3(b)). The contact angle is decreased slowly until to an equilibrium value of 28◦ with the total holding time of 20 min, as shown in Fig. 1(b) and (d). Microstructure of the Cf /SiBCN composite joint brazed with CuPd-8V filler alloy at 1170 ◦ C for 10 min is shown in Fig. 4. It is clear that two reaction layers V(C,N) with a thickness of about
3
1 m are formed at the surface of Cf /SiBCN composite. Two kinds of solid solutions CuPd and Cu3 Pd coexisted in the central part of the joint. Though good bonding was achieved the joint only exhibited a low bend strength of 58.1 MPa at room temperature. The too thin reaction layer would cause the weak interfacial bonding. Compared with the CuPd-8V filler alloy, CuPd-13V alloy showed better wettability on the Cf /SiBCN composite. The main reason was that when the concentration of element V was low, the activity of element V in the interfacial reaction with Cf /SiBCN was insufficient (the thickness of the reaction layer was too thin). As a consequence, it was difficult to achieve a strong bonding. In order to ensure desirable wettability and thus improve the bonding quality at the joining interface, this work focused on the CuPd-13V filler alloy for the Cf /SiBCN joining. Fig. 5 presents the backscattered electron images of the Cf /SiBCN joints brazed with CuPd-13V filler alloy with different brazing temperatures. It is deduced that evident reactions between the CuPd-13V filler alloy and the Cf /SiBCN composite occurred as the two reaction layers (RL) had been formed at the surfaces of the joined composite. Furthermore, with the increase of the brazing temperature, the thickness of the reaction layer was increased, for example at brazing temperature of 1170 ◦ C, the average thickness of the reaction layer (RL “1” and “2”) was only 2.5 m. When brazed at 1200 ◦ C, reaction layer (RL“3” and “4”) showed a thickness of 3.5 m, and that was increased to 6.7 m (RL“5” and “6”) at the brazing temperature of 1230 ◦ C. According to the EDS analysis results listed in Table 2, the reaction layer (labeled “1”–“6”) was mainly composed of V-(C,N) compounds, enriching with elements V, C and N. The active element V diffused strongly to the surface of the Cf /SiBCN composite, as shown in the elemental area distribution maps of the joint brazed at 1200 ◦ C (Fig. 6(c)). The XRD patterns of the fracture surface verified that the interfacial reaction products consisted of V2 C and VN compounds (Fig. 7). In the central part of the joint brazed at 1170 ◦ C (Fig. 5(a)), there were two phases labeled “7” and “8”, which were enriched with elements Cu and Pd. Firstly a grey phase of Cu3 Pd (microzone “7”) was composed of 62.42 at. % Cu and 19.92 at. % Pd. Secondly, a grayish white phase marked as microzone “8” can be deduced as CuPd phase in which the concentration of elements Cu and Pd was 46.17 at. % and 38.04 at. % Pd, respectively. From Fig. 5(b) and (c), when rising the brazing temperature, besides the increase of the thickness of the reaction layer, the microstructure in the central part of the joint was also changed to some extent. At first, both the microzones “9” and “10” in Fig. 5(b), which can be confirmed as CuPd phase by the EDS analysis results, indicated that only one kind of solid solution existed. As shown in Fig. 6(a) and (b), the elements Cu and Pd segregated in the central part of the joint. Similarly, microzone “11” in Fig. 5(c) for brazing temperature of 1230 ◦ C was also the CuPd phase. Therefore with the
Fig. 3. Wetting kinetics of the CuPd-13V alloy on Cf /SiBCN composite: (a) change of contact angle with temperature, (b) change of contact angle with holding time at 1200 ◦ C.
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Fig. 4. Microstructure of the Cf /SiBCN composite joint brazed with CuPd-8V filler alloy at 1170 ◦ C for 10 min: (a) at a low magnification, and (b) at a high magnification.
Fig. 5. Microstructure of the Cf /SiBCN joint brazed with CuPd-13V filler alloy at 1170 ◦ C (a), 1200 ◦ C (b), and 1230 ◦ C (c).
increase of the brazing temperature, the inter-diffusion between elements Cu and Pd became strong so that only CuPd phase existed within the joint. The XRD analysis result (Fig. 7) confirmed that both CuPd and Cu3 Pd phases co-existed at the fracture surface of the joint brazed at 1170 ◦ C (Fig. 7(a)). Whereas, the diffraction peaks of Cu3 Pd phase disappeared at the brazing temperature of 1200 ◦ C (Fig. 7(b)).
For three joints brazed at 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, there were black particles distributing scatteringly in the central part of the joint marked as “12”, “13” and “14” in Fig. 5. Fig. 6(c) signified that V segregated here and V2 C/VN compounds should have been formed according to the EDS results and the XRD patterns. Therefore, during the brazing process, active element V would firstly diffuse from the CuPd-13V filler alloy to the surface of the
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Table 2 EDS analysis results for the typical microzones in Fig. 5. Microzone
1 2 3 4 5 6 7 8 9 10 11 12 13 14
Composition (at. %)
Deduced phase
Cu
Pd
V
Si
B
C
N
0.32 0.24 0.35 0.28 0.28 0.05 62.42 46.17 39.08 39.78 33.54 3.29 3.67 0.49
0.29 0.22 0.31 0.20 0.69 0.07 19.92 38.04 34.50 34.21 45.25 3.32 3.62 2.34
43.93 46.81 45.72 51.82 42.71 42.32 1.14 1.41 5.10 1.58 0.41 41.08 43.51 38.05
2.99 0.18 1.07 0.31 0.12 0.12 1.41 7.61 6.13 6.37 6.50 0.39 0.29 0.34
0.22 / 1.10 / 4.84 5.67 6.91 1.70 7.43 8.82 5.24 1.21 2.03 5.02
18.02 29.06 29.14 33.38 40.57 39.59 5.47 3.75 4.81 5.48 5.40 34.69 35.62 35.67
34.23 23.49 22.31 14.01 10.79 12.18 2.73 1.32 2.95 3.76 3.66 16.02 11.26 18.09
V-(C,N)
Cu3 Pd CuPd
V-(C,N)
Fig. 6. Elemental area distribution maps for Fig. 5(b): (a) Cu, (b) Pd, (c) V, (d) Si, (e) B, (f) C, and (g) N.
Fig. 7. XRD pattern of the fracture surface after bend test: (a) brazing temperature of 1170 ◦ C, (b) brazing temperature of 1200 ◦ C.
Cf /SiBCN composite, and react with Cf /SiBCN through the following formulae: 2V + C = V2 C
(1)
V + N = VN
(2)
Fukai et al. [23] studied the SiC/V joint at (1200-1400)◦ C, and reported that element V reacted with SiC to form V2 C and silicides. In this work, V2 C was detected as the main reaction product within the joint (Fig. 7). In the previous study, when joining Si3 N4 ceramic and Cf /SiC composite using Pd- and V- containing filler alloy, Pd2 Si compounds were observable in the corresponding joints [19,21]. However, here, Pd2 Si compound is absent for the joints brazed at three different temperatures. Compared with the previous Cf /SiC composite, the concentration of Si in the amorphous 3Si-1B-4C-2N
ceramic was remarkably lower than that in the SiC ceramic matrix, and this might be the main reason for the absence of Pd-silicides within the Cf /SiBCN composite joint. According to the V-N binary phase diagram, both VN and V2 N are the stable phases in system of V-N [24]. In this study, only VN compound was as the reaction product. In addition, based on the calculation (f G (VN) = −434.3 + 0.1688T (kJ/mol)) the free energy of formation for VN was −185.7 kJ/mol at 1200 ◦ C [18], indicating that the VN compound can be formed spontaneously. Fig. 8 presents the joint appearances for the three kinds of the joints brazed at 1170 ◦ C, 1200 ◦ C and 1230 ◦ C, respectively, and the correspondingly room-temperature strength of the joints is shown in Fig. 9. The room-temperature strength of the joint brazed at 1170 ◦ C was 57.8 MPa, but it was increased to 82.4 MPa at the brazing temperature of 1200 ◦ C. For the joint brazed at 1170 ◦ C, the
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Fig. 8. Joint appearances for the three kinds of the joints brazed at 1170 ◦ C (a), 1200 ◦ C (b) and 1230 ◦ C (c).
Fig. 9. Room-temperature strength of the joints brazed at different temperature.
thickness of the reaction layer was only 2.5 m (Fig. 5(a)), whereas that was increased to 3.5 m (Fig. 5(b)) when brazing at 1200 ◦ C. If the reaction layer was too thin, the joint strength would decrease due to the insufficient reactivity [25]. In addition, in the case of 1200 ◦ C the V-(C,N) compounds distributed more uniformly in the central part of the joint, with effective strengthening of the joint.
With the brazing temperature up to 1230 ◦ C, the thickness of the reaction layer was increased to 6.7 m, nearly twice of that at 1200 ◦ C. And the obvious cracks appeared within the reaction layer. Consequently, the joint strength was decreased to 62.9 MPa. The relationship between the reaction layer thickness and the defects produced at the reaction region was studied quantitatively by Nakao et al. [26]. They pointed out that when the reaction layer exceeded the optimum thickness, some defects such as microvoids and cracks might be produced at the reaction layer. In general, the mechanical and physical properties of the reaction layer (V2 C,VN) were very different from those of Cf /SiBCN composite; for example, the coefficient of thermal expansion (CTE) for VN was (9–10) × 10−6 K−1 , whereas that of SiBCN matrix ceramic to be joined was about 3.18 × 10−6 K−1 [27]. Undoubtedly, when the reaction layer was relatively thick [28], the residual stresses within the joint would accumulate to cause the decrease of the joint strength. The fractured surface morphologies of the joint brazed at 1200 ◦ C are shown in Fig. 10, and the corresponding EDS analysis results for typical microzones are listed in Table 3. The fracture surface presented double-layered characterization, and the cracking path was evident, and the Cf /SiBCN base material was exposed on the two layered surface, as shown in Fig. 10(a) and (b). For the mag-
Fig. 10. Fracture surface of the Cf /SiBCN joint brazed at 1200 ◦ C (a), the magnification image for area S1 (b), the magnification image for area S2 (c).
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Table 3 EDS analysis result for the typical microzones in Fig. 10. Microzone
1 2
Composition (at. %)
Deduced phase
Cu
Pd
V
Si
B
C
N
31.25 1.54
20.99 0.95
22.04 36.16
3.52 4.22
3.49 5.18
8.46 40.28
10.26 11.67
Fig. 11. Model of the crack propagation when subjected to the bend test.
CuPd/V-(C,N) V-(C,N)
can be regarded as ductile buffer to release the residual stress through their plastic deformation [30]. As a result, at the high testing temperatures the joint strength can be elevated to some extent. Furthermore, the V-C and V-N were typical transition metal carbide/nitride, which have high melting point, high hardness, high strength at high temperatures [31,32]. Therefore, the high strength of the joints should also be attributed to the strengthening effect by the V2 C/VN particles which distributed scatteringly in the central part of the joint. But when the temperature increased to 900 ◦ C, the joint strength was decreased sharply due to the softening of the CuPd phase. 4. Conclusions
Fig. 12. Effect of test temperature on three point bend strength of the Cf /SiBCN composite joint brazed at 1200 ◦ C.
nification of selected area S2, both the CuPd phase and the V-(C,N) compound were detected. It is reasonable to deduce that when subjected to the bend test, the crack initiated at or near the V-(C,N) compound layer at the surface of the Cf /SiBCN composite. On one hand, the formation of the V-(C,N) layer played an important role in the wetting and joining of the Cf /SiBCN composite. However, on the other hand, it seems that the V-(C,N) reaction layer was also the weak link of the Cf /SiBCN-Cf /SiBCN joint. Afterwards, the crack propagated through the brazing seam to the other interface between the Cf /SiBCN composite and the filler alloy, followed by the final failure. The model of the crack propagation within the joint is illustrated in Fig. 11. As shown in Fig. 12, the newly designed CuPd-13V filler alloy can provide the joint strength of 82.4 MPa at room temperature. When the testing temperature was increased to 600 ◦ C, the joint strength was even elevated to 108.8 MPa. Though further increase of the testing temperature to 800 ◦ C resulted in a decrease of the joint strength to 92.4 MPa, it was still at the same level of the room-temperature strength. However, when the testing temperature reached 900 ◦ C, the joint strength was sharply decreased to 39.8 MPa. It is well known that, in the joining of ceramic to ceramic, interfacial thermal mismatch exists between the brazing alloy and the ceramic. In the case of Cf /SiBCN-Cf /SiBCN joint, the V-(C,N) reaction layer can not effectively release the thermal stresses within the joint. Only a soft central layer can secure the stresses relax [29]. At the elevated temperatures of (600-800)◦ C, CuPd phases
(1) CuPd-8V filler alloy exhibited a contact angle of 57◦ after 1170 ◦ C for 30 min, and the CuPd-13V filler alloy showed a lower contact angle of 28◦ on Cf /SiBCN after heating at 1200 ◦ C for 20 min. The higher concentration of active element V corresponded to the better wettability. (2) Element V played an active role in the interfacial reactions, with V2 C and VN compounds formed at the surface of Cf /SiBCN composite. In the central part of the joint brazed with CuPd-13V filler alloy, V2 C/VN compounds distributed scatteringly in the CuPd solid solution. When brazed at 1170 ◦ C, the thickness of V-(C,N) reaction layer was only 2.5 m. With the brazing temperature increased to 1200 ◦ C and 1230 ◦ C, the thickness of the reaction layer was increased to 3.5 m and 6.7 m, respectively. (3) The joint brazed with CuPd-13V filler alloy showed the maximum strength of 82.4 MPa at room temperature, whereas that brazed with CuPd-8V filler was 58.1 MPa. For the joints brazed with CuPd-13V filler alloy, the joint strength was even elevated to 108.8 MPa when the testing temperature increased to 600 ◦ C, and the room-temperature strength value can be maintained up to 800 ◦ C. Acknowledgements This research work was financially supported by the National Natural Science Foundation of China (Grant Nos. 59905022, 50475160 and 51275497), and the Aeronautical Science Foundation of China (Grant No. 2008 ZE21005). References [1] H.D. Batha, E.D. Whitney, J. Am. Ceram. Soc. 56 (1973) 365–373. [2] R. Raj, J. Am. Ceram. Soc. 76 (1993) 2147–2174. [3] R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill, F. Aldinger, Nature 382 (1996) 796–798. [4] M. Jansen, B. Jaschke, T.J. Jaschke, Bond 101 (2002) 137–191. [5] A. Zern, J. Mayer, N. Janakiraman, M. Weinmann, J. Bill, M. Rühle, J. Eur. Ceram. Soc. 22 (2002) 1621–1629. [6] M. Weinmann, J. Schuhmacher, H. Kummer, S. Prinz, J.Q. Peng, H.J. Seifert, M. Christ, K. Muller, J. Bill, F. Aldinger, Chem. Mater. 12 (2000) 623–632. [7] X. Hernandez, C. Jimenez, K. Mergia, P. Yialouris, S. Messoloras, V. Liedtke, C. Wilhelmi, J. Barcena, J. Mater. Eng. Perform. 23 (2014) 3069–3076. [8] M. Weinmann, T.W. Kamphowe, J. Schuhmacher, K. Muller, F. Aldinger, Chem. Mater. 12 (2000) 2112–2122. [9] P. Baldus, M. Jansen, D. Sporn, Science 285 (1999) 699–703. [10] R. Pan, T.S. Lin, P. He, H.M. Wei, Z.H. Yang, Y.X. Shen, J. Chin. Ceram. Soc. 43 (2015) 1719–1724 (in Chinese).
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Please cite this article in press as: W.-W. Li, et al., J. Mater. Sci. Technol. (2018), https://doi.org/10.1016/j.jmst.2018.01.001