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SiBCN composites
Ceramics International 45 (2019) 22471–22478
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Effect of oxidation treatment on the mechanical properties of C/SiBCN composites
T
Yan Jiaa, Si'an Chena,*, Yong Lib, Xiaoyu Jia, Yanzi Goua, Haifeng Hua a b
Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, China Northwest Institute of Nuclear Technology, Xi'an 710024, China
A R T I C LE I N FO
A B S T R A C T
Keywords: C/SiBCN composites Mechanical properties Fracture Thermal applications
C/SiBCN composites are promising self-healing materials, which can be applied as structural materials in hightemperature oxidation environments in the aerospace industry; therefore, it is important to investigate their mechanical properties after oxidation treatment. In this study, C/SiBCN composites were oxidized at elevated temperature in 1200 °C-1700 °C in static air. The C/SiBCN composites demonstrated a desirable oxidation resistance at 1200 °C and 1600 °C. The flexural strength of the composites was retained at 168 MPa after static oxidation at 1200 °C for 30 min, owing to the low diffusion rate of oxygen into the composites. Even after an oxidation of 60 min, the strength was retained at 101 MPa. Furthermore, the strength of the composites was stable at approximately 130 MPa after oxidation at 1600 °C for 120 min, because of the formation of a dense oxide scale. The strength of C/SiBCN composites decreased at 1700 °C due to the interface reaction and oxidative damage of fibers. In 1200 °C-1500 °C, the carbon fiber in the composites was preferentially oxidized, and the strength of the composites decreased as the oxidation time increased.
1. Introduction Carbon fiber reinforced ceramic matrix composites (CFRCMCs) are promising for a variety of aerospace applications, owing to their low density, oxidation resistance, high strength, and excellent fracture toughness [1–4]. With the development of polymer-derived ceramics (PDCs), SiBCN ceramic derived from SiBCN precursor has attracted a lot of interest owing to their oxidation resistance properties, stability up to 2000 °C, and creep resistance [5–8]. Therefore, using continuous carbon fibers as reinforcement, C/SiBCN composites are expected to possess excellent self-healing properties and oxidation resistance because the glassy phase in high temperature can heal the cracks of the C/ SiBCN composites [9,10]. Considering that the C/SiBCN composite is required to withstand a high temperature and air environment in practical applications, it is important to determine the optimal service conditions of the composite; hence, it is essential to understand the evolution of mechanical properties and fracture behavior under hightemperature oxidation conditions. Only a few studies have been reported on C/SiBCN composites, and most of them have focused on the high temperature stability of C/ SiBCN composites under an argon environment. Lee et al. [11,12] prepared a C/SiBCN composite using precursor infiltration and pyrolysis (PIP) with a flexural strength of ~255 MPa. Because of the *
degradation of the SiBCN matrix, the flexural strength of the composite decreased to ~70 MPa at 1700 °C. Zhao et al. [13] fabricated a C/SiBCN composite using liquid precursors with a flexural strength of 265.2 MPa. Ding et al. [14,15] fabricated 3D C/SiBCN composites using PIP with a flexural strength of 371 MPa and flexural modulus of 31 GPa. The microstructure remained intact and the flexural strength of the composite dropped slightly to 85.2% at 1500 °C in argon. Further, the flexural strength of the composite degraded to 92 MPa at 1600 °C, and noticeable interface defects formed because of the carbothermal reduction reaction. In previous studies, the oxidation mechanism of the C/SiBCN composite was reported in detail [16]. The oxidation rate of the C/SiBCN composite was 5–7 μm2 h−1, and these composites were protected by the SiO2 oxide scale covered on the surface of the composites at and above 1600 °C. However, only a limited number of studies have been reported on the microstructure and mechanical properties of C/SiBCN composites at high temperature and air atmosphere, which are of crucial in realistic applications of in aircraft and spacecraft. Wang [17] reported that a C/ SiBCN composite exhibiting flexural strength of 334 MPa at room temperature maintained flexural strength of 198 MPa at 800 °C in static air. The flexural strength decreased to 173 MPa after oxidation at 1000 °C for 20 min. The mechanical properties of C/SiBCN composite at higher oxidation temperatures (i.e. 1500–1700 °C)—which is expected
Corresponding author. E-mail address: [email protected] (S. Chen).
https://doi.org/10.1016/j.ceramint.2019.07.269 Received 30 May 2019; Received in revised form 10 July 2019; Accepted 23 July 2019 Available online 24 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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Table 1 Properties of as-prepared C/SiBCN composites [16]. Properties
C/SiBCN composites −3
Density (g·cm ) Open porosity (%) Flexural strength (MPa) Flexural modulus (GPa)
1.61 6.05 324 56.9
to be the application temperature in the aerospace—have not been investigated. Therefore, the fracture behavior and mechanical properties of C/SiBCN composites at higher oxidation temperatures need to be investigated in detail. In this study, the fracture behavior and mechanical properties of the C/SiBCN composite after an oxidation treatment at 1200–1700 °C are reported, and the evolution of their mechanical properties are analyzed. Fig. 1. Morphologies of C/SiBCN specimens before and after oxidation.
2. Experimental section The C/SiBCN composites were fabricated by our group using PIP. The general properties of C/SiBCN composites are shown in Table 1. The preparation of C/SiBCN composites has been described elsewhere [18]. The C/SiBCN composites were machined into specimens of size of 3.0 mm × 4.0 mm × 60 mm. Oxidation experiments were performed in the range of 1200 °C1700 °C for 10 min–120 min in a muffle furnace under static air. The specimens were placed in a high purity (> 99.99%) alumina boat in a position where only two of their edges were in contact with the boat for minimizing the contact area between the specimen and crucible. They were then placed inside the muffle furnace at the test temperatures. After the specified oxidation time, the specimen was taken out of the muffle furnace and immediately cooled in air. The change in weight and dimension of the specimen was recorded with a precision of ± 0.001 g and ± 0.01 mm, respectively. In the remainder of the paper, the term “as-prepared” refers to untreated specimens; C12-10 represents that the C/SiBCN composite was oxidized at 1200 °C for 10 min, and so on. The microstructure of the composites after the three-point bending tests was characterized by scanning electron microscopy (SEM; TESCAN MAIA 3 XMH, Czech). The elemental analysis of specimens was performed using energy dispersive spectroscopy (EDS; FEI Quanta-200, USA). The phase compositions of specimens after the oxidation treatment were conducted by X-ray diffraction (XRD; Bruker D8 Advance, Cu Kα radiation, Germany). The TGA–DSC–MS curve of the composites was obtained by a thermogravimetric analysis and differential scanning calorimetry device with a mass spectrometer (METTLER TOLEDO, Switzerland; QMG 700, INFICON, Germany) in 30 °C-1600 °C in flowing air (the ramping rate: 10 °C·min−1). The open porosity and apparent density of C/SiBCN composite was measured using the Archimedes’ principle. The threepoint bending test was applied on specimens of 3.0 mm × 4.0 mm × 60 mm with 50 mm span and 0.5 mm min−1 crosshead speed to measure flexural strength and modulus (ASTM C1359-96). At least four specimens were used in each test.
1700 °C, the C/SiBCN specimens were severely oxidized, as shown in Fig. 1, and the weight loss rate was 79.4%. The TGA-DSC-MS curves of the C/SiBCN composites, and the TG curve of SiBCN ceramics in air [16] from ambient temperature to 1600 °C are shown in Fig. 2. A noticeable exothermic peak appears in 765 °C-1120 °C. At the same time, CO2 and a small amount of CO were released in this temperature range, accompanied by a noticeable weight loss (approximately 45.3%). However, similar weight loss did not occur in the SiBCN ceramics. Therefore, the exothermic peak was attributed to the oxidation of carbon fibers in the fresh cross section of C/SiBCN composites. As the oxidation temperature increased from 1120 °C to 1600 °C, the weight loss rate gradually increased to 50.7%. However, no significant amount of oxidizing gas was detected. The weight loss (5.4%) can be attributed to the slight oxidation of a small amount of residual carbon fibers and the SiBCN ceramic matrix. As the oxidation temperature increased from ambient temperature to 1600 °C, the carbon fibers in the C/SiBCN composites began to oxidize preferentially at 765 °C, while the SiBCN matrix did not show any noticeable oxidation. The weight loss rate of the C/SiBCN specimens after oxidation at 1200 °C-1700 °C for 10 min–120 min is shown in Fig. 3. The oxidation process of the composites was given by the following equations: The process of weight loss of composites is described by equations:
C(s) +
(1)
C(s) + O2 (g ) → CO2 (g )
(2)
B2 O3 (l) → B2 O3 (g )
(3)
SiO2 (l) → SiO2 (g )
(4)
SiC(s) + O2 (g ) → SiO (g ) + CO
(5)
The process of weight gain of composites is described by equations:
SiBCx Ny (s) +
3. Results and discussion 3.1. Morphology and weight loss of C/SiBCN composites The morphologies of C/SiBCN specimens before and after oxidation treatment are shown in Fig. 1. The color of specimens C12-120 and C14-120 changed from gray to black after oxidation. A glassy oxide scale formed and remained on the composites after oxidation treatment at 1600 °C, which avoided further oxidation of the composite [16]. However, after an extended oxidation treatment, such as 120 min, at
1 O2 (g ) → CO (g ) 2
7 + 2x O2 (g ) → SiO2 + 1/2B2 O3 (l) + xCO + y /2N2 4
(6)
2 2 2 SiC(s) + O2 (g ) → SiO2 (l) + CO 3 3 3
(7)
1 2 Si3 N4 (s) + O2 (g ) → SiO2 (l) + N2 3 3
(8)
4 2 2 BN(s) + O2 (g ) → B2 O3 (l) + N2 3 3 3
(9)
The weight loss rate increased gradually accompanied by reaction in Eqs. (1) and (2) with the increase in oxidation time at 1200 °C. However, as the oxidation time increased, the weight loss rate gradually
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Fig. 2. TGA–DSC–MS curve of C/SiBCN composites and TGA of SiBCN ceramics [16] in flowing air.
Fig. 3. wt loss of the C/SiBCN specimens after oxidation at 1200–1700 °C for 10–120 min.
decreased at 1400 °C, 1500 °C, and 1600 °C because of the weight-increasing oxidation of the matrix of the composites (reaction in Eqs. (6)–(9)). Compared with other oxidation temperatures, the weight loss rate was lower at 1600 °C because an oxide scale formed on the surface of the specimens by the sacrificial oxidation of the SiBCN ceramic matrix (reactions in Eqs. (6)–(9)) to prevent further oxidation of the composites. Specimen C17-120 was catastrophically oxidized, resulting in a weight loss rate of 79.4%, probably because the surface oxide scale was constantly volatilized and the composites were rapidly oxidized due to the direct contact with oxygen. 3.2. Mechanical properties of C/SiBCN composites after oxidation Fig. 4 and Fig. 5 show the open porosity, density, and mechanical properties of the specimens after oxidation treatment, respectively. After the oxidation treatment at 1200 °C, the flexural strength and modulus of the composites decreased gradually as the oxidation time increased (Fig. 5). The weight loss rate was only 14.5% in the C12-10 and C12-30 (Fig. 3) with the open porosity of approximately 20% (Fig. 4), indicating that only a small amount of fiber was oxidized. Therefore, the specimens C12-10 and C12-30 maintained high flexural strength retention rates of 65.4% and 51.9%, respectively (Fig. 5). With
Fig. 4. The open porosities and densities of C/SiBCN specimens after oxidation at 1200 °C–1700 °C for 10 min–120 min.
the increase in oxidation time, the oxidation rate of fibers was increased. The flexural strength and modulus of specimen C12-60 gradually decreased to 101 MPa and 18.5 GPa, respectively (Fig. 5) with the increase of weight loss rate (Fig. 3) and the open porosity (Fig. 4). A
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After oxidation at 1600 °C, an oxide scale formed on the surface of C/SiBCN composites, which healed the pores and cracks, and prevented the diffusion of oxygen into the interior of the composites [16]. Therefore, the flexural strength and flexural modulus of these specimens (approximately 130 MPa and 38 GPa, respectively) was significantly higher than those treated with oxidation at 1400 °C and 1500 °C, which indicated that C/SiBCN composites could maintain its mechanical properties via its self-healing behavior. The density of composites after oxidation at 1600 °C increased with the increase in the oxidation time (Fig. 4). This was because the oxidation reactions of the matrix (reaction in Eqs. (6)–(9)) were aggravated, resulting in a decrease in the weight loss rate of composites, and the open pores of composites were filled by the flow of molten oxides. The flexural strength the composites after oxidation at 1700 °C for 10 min–60 min decreased to 52.1 MPa–62.8 MPa, because the volatilization of the oxide scale led to the gradual oxidation of the fiber. After oxidation at 1700 °C for 120 min, the matrix and fibers of the C17-120 were violently oxidized, resulting in decreasing flexural strength and flexural modulus of C/SiBCN composites to 43 MPa and 6.4 GPa, respectively, (Fig. 5) with 79.4% weight loss (Fig. 3). 3.3. Reason for the evolution of mechanical properties
Fig. 5. The flexural strength and modulus of C/SiBCN specimens after oxidation treatment at 1200 °C–1700 °C for 10 min–120 min.
higher open porosity (> 30% shown in Fig. 4) facilitated the diffusion of oxygen and the oxidation rate of fibers increased as the oxidation time exceeded 60 min. Therefore, the flexural strength and modulus of C12–120 decreased sharply to 22.1 MPa and 7.49 GPa, respectively, with an open porosity of 48.8% (close to fiber volume fraction (44.4%)). The mechanical properties of the specimens after oxidation treatment at 1400 °C for 10 min–120 min were similar to those of treated with oxidation at 1500 °C. The strength retention rates of specimens C14-10 and C15-10 were 50.3% and 37.3% (Fig. 5), respectively, which were noticeably lower than that of C12-10 (65.4%). This is because the oxidation rate of the carbon fibers increased gradually as the oxidation temperature increased, which resulted in higher open porosity (Fig. 4). At the same time, the higher open porosity also increased the oxidation rate of fibers, because it facilitated the diffusion of oxygen. Therefore, the fractural strength and modulus of specimens C14-30 decreased sharply to 33.1 MPa and 21.1 GPa, respectively (Fig. 5). The fractural strength and modulus of C15-30 decreased to 10.4 MPa and 14.8 GPa, respectively (Fig. 5). The flexural strength of C14-120 and C15-120 was slightly higher than those of C14-60 and C15-60. It may be because the porosity of the composites decreased (Fig. 4) caused by the melting of the matrix, which resulted that the matrix loading capacity was increased.
Oxidation-induced damage of the mechanical properties of C/SiBCN composites mainly includes carbon fiber damage, SiBCN matrix damage and interphase damage. The carbon fiber content in the composites after oxidation has a strong effect on the mechanical properties of the composites, not only because the carbon fiber is the main load-bearing phase, but also because oxidized fibers form a large number of holes and cracks, which greatly damage the strength of the composites. The structural integrity of the matrix skeleton is destroyed in a high-temperature oxidizing environment, and this decreases the strength of the composites [19]. However, the oxidized matrix forms a molten glass phase, which contributes to the self-healing properties of C/SiBCN composites. As the oxidation temperature rise above 1200 °C, the residual thermal stress induced by the different expansion coefficients of the SiBCN matrix and the carbon fibers, results in the formation of microcracks at the interface (carbon, approximately 0 °C−1; SiBCN, 3.1 × 10−6–3.5 × 10−6 °C−1) [20], which affects the fracture behavior of the composites. The XRD patterns of the cross section of the specimens after oxidation treatment are shown in Fig. 6. The SiBCN ceramic matrix of the composites was amorphous before and after the high-temperature oxidation, therefore no characteristic diffraction peaks appeared. The peaks at 26° and 43° corresponded to carbon fibers in specimen C12-60; however, these peaks did not appeared in specimen C12-120, which indicated that the carbon fibers were oxidized completely. Therefore, the flexural strength of the C12-60 was maintained at 110 MPa and that of C12-120 is almost completely lost (22.1 MPa). The carbon fibers did not appear in specimens C14-30 and C15-60, which indicated that the fibers were oxidized faster at 1400 °C and 1500 °C than at 1200 °C. This is because the oxidation of carbon fibers was accelerated by the increase of the temperature. Therefore, the fractural strengths of the composites after oxidation at 1400 °C and 1500 °C decreased faster than those of the composites after oxidation at 1200 °C (Fig. 5). However, the carbon fibers were retained well in the C16-120 due to the formation of an oxide scale, resulting in the improved mechanical properties of C16120 (flexural strength: 130 MPa; flexural modulus: 38.0 GPa). SiC grains appeared in C17-30 and C17-60, which could be attributed to the interface reaction (2 C+ 2Si − N→ 2SiC + N2 ↑) and the crystallized of SiBCN matrix [14]. After oxidation at 1200 °C, the high temperature (1200 °C) induced the generation of microcracks at the interface, and oxygen diffused into the microcracks on the surface of composites and reacted with the fibers to form micropores. The microcracks and micropores were the path for the entrance of oxygen into the composites as well as the path for the
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Fig. 6. XRD of the cross-section of C/SiBCN specimens before and after oxidation at 1200 °C-1700 °C for 10–120 min.
Fig. 7. Micrographs of fracture surface of specimens (a) C12-10, (b) C12-30, (c) C12-60, and (d) C12-120.
exit of oxide gases out of the composites [21,22]. The oxidation of fibers inside specimen C12-10 was slow due to the hysteresis of oxygen diffusion through the microcracks and micropores. Carbon fibers in specimen C12-10 were remained as the main load bearing phase (Fig. 7(a–b)). Therefore, the fractural strength and modulus of specimen C12-10 was maintained at 212 MPa and 42.5 GPa, respectively.
Specimen C12-10 showed a typical fracture surface of the composites [23–26]. A certain extent of fiber-debonding and fiber pull-out were observed in specimen C12-10 (Fig. 7 (a)), indicating a relatively weak interface bonding. The fibers were gradually oxidized with an increase in oxidation time and a porous SiBCN matrix and a small amount of fiber pull-out
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Fig. 8. Micrographs of fracture surface of specimens (a) C14-10, (b) C14-30, (c) C14-60, (d) C14-120.
Fig. 9. Micrographs of fracture surfaces of specimens (a, b) C16-30, (c, d) C16-60, and (e, f) C16-120.
were observed in specimen C12-30 and C12-60 (Fig. 7(b–c)). Therefore, the flexural strength of specimen C12-30 and C12-60 decreased to 168 MPa and 101 MPa, respectively (Fig. 5). The fibers of the C12-120 were almost completely oxidized in Fig. 7 (d), which can also be seen in
the XRD pattern of specimen C12-120 (Fig. 6). Further, the structure of SiBCN ceramic matrix in the C12-120 collapsed, and the oval fiber holes were connected to irregular pores (Fig. 7 (d)). Therefore, the mechanical strength of the C12-120 decreased sharply (flexural strength:
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22.1 MPa; flexural modulus: 7.49 GPa). Specimen C14-10 did not show any noticeable obvious oxidation, similar to the as-prepared [18] and C12-10 specimens (Fig. 8 (a)). Therefore, the flexural strength and modulus of specimens C14-10 could retained at 163 MPa and 35.9 GPa, respectively. However, the fibers of specimens C14-30 were rapidly oxidized to form oval holes, while the SiBCN skeleton remained stable (Fig. 8 (b)). The flexural strength and modulus of specimens C14-30 decreased sharply to 33.1 MPa and 21.1 GPa, respectively (Fig. 5). The fibers were almost completely oxidized (Fig. 8 (c)). Furthermore, there were many cracks and holes in the SiBCN matrix of specimens C14-60 (Fig. 8 (c)), which resulted in a poor load bearing capacity and consequently decreased the mechanical properties of specimens C14-60 (fractural strength: 22.6 MPa; fractural modulus: 15.4 GPa). The SiBCN matrix of specimen C14-120 was oxidized and melted (Fig. 8 (d)). Some small pores were filled in C15-60, resulting in an increased density (1.02 g cm−3) and decreased porosity to 38.9% (Fig. 4). Some cracks in the SiBCN matrix were filled, which increased the load bearing capacity of the SiBCN matrix and resulted in enhanced mechanical properties of specimen C14-120 (Fig. 5). The fracture behavior and mechanical properties of the composite after oxidation at 1500 °C are similar to those of the composite after oxidation at 1400 °C, therefore, these are not described here. Fig. 9 shows the fracture surfaces of the composites after oxidation at 1600 °C. The composites exhibited significant self-healing properties at 1600 °C. As shown in Fig. 10, a dense oxide film was observed on the surface of the composites, healing the cracks in the composites. The fiber bundles of the composites after oxidation at 1600 °C showed little pull-out from the matrix of the composites, and the fracture surfaces were flat (Fig. 9), indicating a strong fiber–matrix interface bonding [27,28]. On the surface of the composites, an oxide scale formed by the oxidative melting of the SiBCN ceramic matrix, healing of the pores and cracks. The oxide layer contained Si and O elements (Fig. 10 (b)), which indicated that it was composed of SiO2. Borosilicate glass did not form on the composites due to the volatilization of B2O3 resulting from the decomposition of borosilicate glass at 1600 °C [29]. After oxidation at 1600 °C, the fibers and matrix of the composites did not show significant oxidative damage. The fibers remained intact in specimens C16-30, C16-60, and C16-120, because of the formation of a dense oxide scale on the composites that prevented further diffusion of oxygen into the composites (oxygen diffusion rate of SiO2: 3 × 10−14 g (m s)−1 [30]). Furthermore, the SiBCN matrix after oxidation at 1600 °C showed no obvious oxidative damage. Therefore, the flexural strength and modulus of specimen C16-120 could be maintained at 130 MPa and 38 GPa, respectively. Fig. 11 shows the fracture surfaces of the composite after oxidation at 1700 °C. Due to the difference in the expansion coefficients of the fibers and the matrix, a large number of cracks were generated at the
interface at 1700 °C. Therefore, specimens C17-10, C17-30, and C17-60 exhibited long fiber pull-out with an average length of more than 100 μm (Fig. 11 (b, d, g)). After oxidation at 1700 °C, the oxide scale became porous due to the volatilization of CO2, SiO2, and B2O3 (Fig. 11 (a)) [12]. Oxygen diffused along the porous surface oxide scales into the interior of specimens and the carbon fibers were oxidized. The fiber had a low oxidation rate at 1700 °C, because the oxygen diffusion rate through the porous oxide scale was slow. Therefore, the fibers were retained even in C17-60 (Fig. 11 (g–i)) with a slight increase in porosity (Fig. 4). Significant oxidative damage to fibers (Fig. 11 (e, h)) and interfacial reaction rings (Fig. 11 (f, i)) occurred in the C17-30 and C17-60, resulting in weakened fiber strength. The load bearing capacity of the SiBCN ceramic matrix was degraded because of an increased number of cracks and more isolated ceramic phases. Therefore, the flexural strengths of the specimens after oxidation at 1700 °C were significantly lower than those of the composites after oxidation at 1600 °C (Fig. 5).
4. Conclusion The fracture behavior and mechanical properties of the C/SiBCN composite were reported after oxidation at 1200 °C-1700 °C for different durations in static air. The elevated temperature induced the oxidation reaction of carbon fibers, and the oxidation rate gradually increased as the oxidation temperature increased from 1200 °C to 1500 °C. The flexural strength retention rate of the composites was 51.8% (168 MPa) after oxidation at 1200 °C for 30 min. Even for an oxidation time of 60 min, the flexural strength of the specimens were maintained at 101 MPa. After oxidation treatment at 1400 °C and 1500 °C for 30 min, the strength of the composites was nearly lost completely. The composites exhibited excellent self-healing properties at 1600 °C. The carbon fibers were appropriately retained after oxidation at 1600 °C, because of the formation of a dense oxide scale. The composites maintained a flexural strength of 130 MPa even after oxidation at 1600 °C for 120 min, exhibiting excellent long-term oxidation resistance. However, the mechanical properties of C/SiBCN composites decreased significantly at 1700 °C due to the oxidative damage of the fibers and the interface reaction between the carbon fibers and the matrix. The C/SiBCN composites demonstrate a desirable oxidation resistance at 1200 °C and 1600 °C for a long period of time (30 min and 120 min, respectively). More research focusing on the effects of fiber anti-oxidation coatings on the properties of C/SiBCN composites will be conducted in the future.
Fig. 10. SEM micrograph of the oxide scale (a) and the EDS result (b) of C/SiBCN specimens after oxidation at 1600 °C. 22477
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Fig. 11. Micrographs of fracture surface of specimens (a–c) C17-10, (d–f) C17-30, and (g–i) C17-60.
Acknowledgements This work was supported by the National Program for Key Basic Research Projects [grant numbers 2015CB655200] and the National Natural Science Foundation of China [grant numbers 51602346]. References [1] N.P. Padture, Advanced structural ceramics in aerospace propulsion, Nat. Mater. 15 (2016) 804–809. [2] Y. Ma, S. Wang, Z.H. Chen, Effects of high-temperature annealing on the microstructures and mechanical properties of Cf/SiC composites using polycarbosilane, Mater. Sci. Eng. A 528 (2011) 3069–3072. [3] J. Kong, M.J. Wang, J.H. Zou, L.N. An, Soluble and meltable hyperbranched polyborosilazanes toward high-temperature stable SiBCN ceramics, ACS Appl. Mater. Interfaces 7 (2015) 6733–6744. [4] C.K. Song, L.F. Cheng, Y.S. Liu, M.X. Zhao, F. Ye, Microstructure and electromagnetic wave absorption properties of RGO-SiBCN composites via PDC technology, Ceram. Int. 44 (2018) 18759–18769. [5] X.Y. Ji, S. Wang, C. Shao, H. Wang, High-temperature corrosion behavior of SiBCN fibers for aerospace applications, ACS Appl. Mater. Interfaces 10 (2018) 19712–19720. [6] E. Butchereit, K.G. Nickel, A. Muller, Precursor-derived Si-B-C-N ceramics: oxidation kinetics, J. Am. Ceram. Soc. 84 (2001) 2184–2188. [7] A. Viard, D. Fonblanc, M. Schmidt, A. Lale, D. Lopez-Ferber, M. Balestrat, M. Weinmann, F. Rossignol, Polymer derived Si-B-C-N ceramics: 30 Years of research, Adv. Eng. Mater. 20 (2018) 1800360. [8] R. Riedel, A. Kienzle, L. Ruwisch, W. Dressler, F. Aldinger, J. Bill, A silicoboron carbonitride ceramic stable to 2000 °C, Nature 382 (1996) 796–798. [9] C.J. Luo, Y.S. Tang, T. Jiao, J. Kong, High-temperature stable and metal-free electromagnetic wave-absorbing SiBCN ceramics derived from carbon-rich hyperbranched polyborosilazanes, ACS Appl. Mater. Interfaces 10 (2018) 28051–28061. [10] F. Liu, J. Kong, F. Ye, C.J. Luo, X.G. Luan, et al., High temperature self-healing SiBCN ceramics derived from hyperbranched polyborosilazanes, Adv. Compos. Hybrid Mater. 1 (2018) 506–517. [11] S.H. Lee, M. Weinmann, F. Aldinger, Processing and properties of C/Si-B-C-N fiberreinforced ceramic matrix composites prepared by precursor impregnation and pyrolysis, Acta Mater. 56 (2008) 1529–1538. [12] S.H. Lee, M. Weinmann, Cfiber/SiCfiller/Si-B-C-Nmatrix composites with extremely high thermal stability, Acta Mater. 57 (2009) 4374–4381. [13] H. Zhao, L. Chen, X. Luan, X. Zhang, J. Yun, T. Xu, Synthesis, pyrolysis of a novel liquid SiBCN ceramic precursor and its application in ceramic matrix composites, J. Eur. Ceram. Soc. 37 (2017) 1321–1329. [14] Q. Ding, D.W. Ni, Z. Wang, H.D. Wang, H.J. Zhou, Y.M. Kan, S.M. Dong, Mechanical properties and microstructure evolution of 3D Cf/SiBCN composites at elevated temperatures, J. Am. Ceram. Soc. 101 (2018) 4699–4707.
[15] Q. Ding, D.W. Ni, Z. Wang, S.M. Dong, H. Zhong, H.D. Wang, 3D Cf/SiBCN composites prepared by an improved polymer infiltration and pyrolysis, J. Adv. Ceram. 7 (2018) 266–275. [16] Y. Jia, S.A. Chen, X.Y. Ji, Y.Z. Gou, Y. Li, H.F. Hu, High-temperature oxidation behavior and oxidation mechanism of C/SiBCN composites in static air, Ceram. Int. 45 (2019) 12764–12772. [17] X.J. Wang, Z.B. Zhang, F. Zeng, Y.M. Li, C.H. Xu, Fabrication and properties of carbon fiber reinforced SiBCN matrix composites, Aerosp. Mater. Technol. 43 (2012) 47–50. [18] Y. Jia, X.Y. Ji, S.A. Chen, Y.Z. Gou, Y. Li, H.F. Hu, Preparation of C/SiBCN composites by precursor infiltration and pyrolysis using a novel precursor, Mater. Res. Express 6 (2019) 065612. [19] E. Butchereit, K.G. Nickel, Oxidation behavior of precursor dervied ceramics in the system Si-(B)-C-N, Key Eng. Mater. 175–176 (1999) 69–78. [20] S.H. Lee, F. Aldinger, Processing and thermal stability of a Cfiber/SiCfiller/Si-C-Nmatrix composite: effect of oligomer vaporization and the surface oxide layer of SiC filler, J. Mater. Sci. 44 (2009) 2101–2108. [21] J.T. Hou, S.R. Qiao, Y.B. Zhang, C.Y. Zhang, The influence of high temperature exposure to air on the damage to 3D-C/SiC composites, N. Carbon Mater. 24 (2009) 173–177. [22] X.G. Luan, L. Wang, Y. Zou, Q.Q. Zhang, R. Riedel, L.F. Cheng, Oxidation behavior of C/SiC-SiBCN composites at high temperature, J. Eur. Ceram. Soc. 39 (2019) 3003–3012. [23] Y.D. Wang, C. Han, D. Zheng, Y.P. lei, Large-scale, flexible and high-temperature resistant ZrO2/SiC ultrafine fibers with a radial gradient composition, J. Mater. Chem. A. 2 (2014) 9607–9612. [24] B. Wang, Y. Wang, Y. Lei, N. Wu, Y. Gou, C. Han, S. Xie, D. Fang, Hierarchically porous SiC ultrathin fibers mat with enhanced mass transport, amphipathic property and high-temperature erosion resistance, J. Mater. Chem. A. 2 (2014) 20873–20881. [25] F. Ye, L.T. Zhang, X.W. Yin, Y.J. Zhang, L. Kong, Y.S. Liu, L.F. Cheng, Dielectric and EMW absorbing properties of PDCs-SiBCN annealed at different temperatures, J. Eur. Ceram. Soc. 33 (2013) 1469–1477. [26] F. Ye, L.T. Zhang, X.W. Yin, Y.J. Zhang, L. Kong, Y.S. Liu, L.F. Cheng, Dielectric and microwave-absorption properties of SiC nanoparticle/SiBCN composite ceramics, J. Eur. Ceram. Soc. 34 (2014) 205–215. [27] Y. Mu, W.C. Zhou, H.Y. Wang, C. Wang, Y.C. Qing, Mechanical and dielectric properties of 2.5D SiCf/SiC–Al2O3 composites prepared via precursor infiltration and pyrolysis, Mater. Sci. Eng. A 596 (2014) 64–70. [28] K. Jian, Z.H. Chen, Q.S. Ma, H.F. Hu, W.W. Zheng, Effects of polycarbosilane infiltration processes on the microstructure and mechanical properties of 3D-Cf/SiC composites, Ceram. Int. 33 (2007) 905–909. [29] Y.P. Lei, Y.D. Wang, J.G. Xue, Y.C. Song, Influence of pyrolysis conditions on fabrication of polymer-derived BN fiber for wave transparent application, Compos. Part B 51 (2013) 254–259. [30] S.Y. Chen, G.Q. Jiang, J.J. Yu, D.B. Ou, Passive oxidation mechanism and transition temperature of silicon carbide, Aerosp. Mater. Technol. 39 (2009) 21–24.
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Effect of oxidation treatment on the mechanical properties of C/SiBCN composites
T
Yan Jiaa, Si'an Chena,*, Yong Lib, Xiaoyu Jia, Yanzi Goua, Haifeng Hua a b
Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, China Northwest Institute of Nuclear Technology, Xi'an 710024, China
A R T I C LE I N FO
A B S T R A C T
Keywords: C/SiBCN composites Mechanical properties Fracture Thermal applications
C/SiBCN composites are promising self-healing materials, which can be applied as structural materials in hightemperature oxidation environments in the aerospace industry; therefore, it is important to investigate their mechanical properties after oxidation treatment. In this study, C/SiBCN composites were oxidized at elevated temperature in 1200 °C-1700 °C in static air. The C/SiBCN composites demonstrated a desirable oxidation resistance at 1200 °C and 1600 °C. The flexural strength of the composites was retained at 168 MPa after static oxidation at 1200 °C for 30 min, owing to the low diffusion rate of oxygen into the composites. Even after an oxidation of 60 min, the strength was retained at 101 MPa. Furthermore, the strength of the composites was stable at approximately 130 MPa after oxidation at 1600 °C for 120 min, because of the formation of a dense oxide scale. The strength of C/SiBCN composites decreased at 1700 °C due to the interface reaction and oxidative damage of fibers. In 1200 °C-1500 °C, the carbon fiber in the composites was preferentially oxidized, and the strength of the composites decreased as the oxidation time increased.
1. Introduction Carbon fiber reinforced ceramic matrix composites (CFRCMCs) are promising for a variety of aerospace applications, owing to their low density, oxidation resistance, high strength, and excellent fracture toughness [1–4]. With the development of polymer-derived ceramics (PDCs), SiBCN ceramic derived from SiBCN precursor has attracted a lot of interest owing to their oxidation resistance properties, stability up to 2000 °C, and creep resistance [5–8]. Therefore, using continuous carbon fibers as reinforcement, C/SiBCN composites are expected to possess excellent self-healing properties and oxidation resistance because the glassy phase in high temperature can heal the cracks of the C/ SiBCN composites [9,10]. Considering that the C/SiBCN composite is required to withstand a high temperature and air environment in practical applications, it is important to determine the optimal service conditions of the composite; hence, it is essential to understand the evolution of mechanical properties and fracture behavior under hightemperature oxidation conditions. Only a few studies have been reported on C/SiBCN composites, and most of them have focused on the high temperature stability of C/ SiBCN composites under an argon environment. Lee et al. [11,12] prepared a C/SiBCN composite using precursor infiltration and pyrolysis (PIP) with a flexural strength of ~255 MPa. Because of the *
degradation of the SiBCN matrix, the flexural strength of the composite decreased to ~70 MPa at 1700 °C. Zhao et al. [13] fabricated a C/SiBCN composite using liquid precursors with a flexural strength of 265.2 MPa. Ding et al. [14,15] fabricated 3D C/SiBCN composites using PIP with a flexural strength of 371 MPa and flexural modulus of 31 GPa. The microstructure remained intact and the flexural strength of the composite dropped slightly to 85.2% at 1500 °C in argon. Further, the flexural strength of the composite degraded to 92 MPa at 1600 °C, and noticeable interface defects formed because of the carbothermal reduction reaction. In previous studies, the oxidation mechanism of the C/SiBCN composite was reported in detail [16]. The oxidation rate of the C/SiBCN composite was 5–7 μm2 h−1, and these composites were protected by the SiO2 oxide scale covered on the surface of the composites at and above 1600 °C. However, only a limited number of studies have been reported on the microstructure and mechanical properties of C/SiBCN composites at high temperature and air atmosphere, which are of crucial in realistic applications of in aircraft and spacecraft. Wang [17] reported that a C/ SiBCN composite exhibiting flexural strength of 334 MPa at room temperature maintained flexural strength of 198 MPa at 800 °C in static air. The flexural strength decreased to 173 MPa after oxidation at 1000 °C for 20 min. The mechanical properties of C/SiBCN composite at higher oxidation temperatures (i.e. 1500–1700 °C)—which is expected
Corresponding author. E-mail address: [email protected] (S. Chen).
https://doi.org/10.1016/j.ceramint.2019.07.269 Received 30 May 2019; Received in revised form 10 July 2019; Accepted 23 July 2019 Available online 24 July 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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Table 1 Properties of as-prepared C/SiBCN composites [16]. Properties
C/SiBCN composites −3
Density (g·cm ) Open porosity (%) Flexural strength (MPa) Flexural modulus (GPa)
1.61 6.05 324 56.9
to be the application temperature in the aerospace—have not been investigated. Therefore, the fracture behavior and mechanical properties of C/SiBCN composites at higher oxidation temperatures need to be investigated in detail. In this study, the fracture behavior and mechanical properties of the C/SiBCN composite after an oxidation treatment at 1200–1700 °C are reported, and the evolution of their mechanical properties are analyzed. Fig. 1. Morphologies of C/SiBCN specimens before and after oxidation.
2. Experimental section The C/SiBCN composites were fabricated by our group using PIP. The general properties of C/SiBCN composites are shown in Table 1. The preparation of C/SiBCN composites has been described elsewhere [18]. The C/SiBCN composites were machined into specimens of size of 3.0 mm × 4.0 mm × 60 mm. Oxidation experiments were performed in the range of 1200 °C1700 °C for 10 min–120 min in a muffle furnace under static air. The specimens were placed in a high purity (> 99.99%) alumina boat in a position where only two of their edges were in contact with the boat for minimizing the contact area between the specimen and crucible. They were then placed inside the muffle furnace at the test temperatures. After the specified oxidation time, the specimen was taken out of the muffle furnace and immediately cooled in air. The change in weight and dimension of the specimen was recorded with a precision of ± 0.001 g and ± 0.01 mm, respectively. In the remainder of the paper, the term “as-prepared” refers to untreated specimens; C12-10 represents that the C/SiBCN composite was oxidized at 1200 °C for 10 min, and so on. The microstructure of the composites after the three-point bending tests was characterized by scanning electron microscopy (SEM; TESCAN MAIA 3 XMH, Czech). The elemental analysis of specimens was performed using energy dispersive spectroscopy (EDS; FEI Quanta-200, USA). The phase compositions of specimens after the oxidation treatment were conducted by X-ray diffraction (XRD; Bruker D8 Advance, Cu Kα radiation, Germany). The TGA–DSC–MS curve of the composites was obtained by a thermogravimetric analysis and differential scanning calorimetry device with a mass spectrometer (METTLER TOLEDO, Switzerland; QMG 700, INFICON, Germany) in 30 °C-1600 °C in flowing air (the ramping rate: 10 °C·min−1). The open porosity and apparent density of C/SiBCN composite was measured using the Archimedes’ principle. The threepoint bending test was applied on specimens of 3.0 mm × 4.0 mm × 60 mm with 50 mm span and 0.5 mm min−1 crosshead speed to measure flexural strength and modulus (ASTM C1359-96). At least four specimens were used in each test.
1700 °C, the C/SiBCN specimens were severely oxidized, as shown in Fig. 1, and the weight loss rate was 79.4%. The TGA-DSC-MS curves of the C/SiBCN composites, and the TG curve of SiBCN ceramics in air [16] from ambient temperature to 1600 °C are shown in Fig. 2. A noticeable exothermic peak appears in 765 °C-1120 °C. At the same time, CO2 and a small amount of CO were released in this temperature range, accompanied by a noticeable weight loss (approximately 45.3%). However, similar weight loss did not occur in the SiBCN ceramics. Therefore, the exothermic peak was attributed to the oxidation of carbon fibers in the fresh cross section of C/SiBCN composites. As the oxidation temperature increased from 1120 °C to 1600 °C, the weight loss rate gradually increased to 50.7%. However, no significant amount of oxidizing gas was detected. The weight loss (5.4%) can be attributed to the slight oxidation of a small amount of residual carbon fibers and the SiBCN ceramic matrix. As the oxidation temperature increased from ambient temperature to 1600 °C, the carbon fibers in the C/SiBCN composites began to oxidize preferentially at 765 °C, while the SiBCN matrix did not show any noticeable oxidation. The weight loss rate of the C/SiBCN specimens after oxidation at 1200 °C-1700 °C for 10 min–120 min is shown in Fig. 3. The oxidation process of the composites was given by the following equations: The process of weight loss of composites is described by equations:
C(s) +
(1)
C(s) + O2 (g ) → CO2 (g )
(2)
B2 O3 (l) → B2 O3 (g )
(3)
SiO2 (l) → SiO2 (g )
(4)
SiC(s) + O2 (g ) → SiO (g ) + CO
(5)
The process of weight gain of composites is described by equations:
SiBCx Ny (s) +
3. Results and discussion 3.1. Morphology and weight loss of C/SiBCN composites The morphologies of C/SiBCN specimens before and after oxidation treatment are shown in Fig. 1. The color of specimens C12-120 and C14-120 changed from gray to black after oxidation. A glassy oxide scale formed and remained on the composites after oxidation treatment at 1600 °C, which avoided further oxidation of the composite [16]. However, after an extended oxidation treatment, such as 120 min, at
1 O2 (g ) → CO (g ) 2
7 + 2x O2 (g ) → SiO2 + 1/2B2 O3 (l) + xCO + y /2N2 4
(6)
2 2 2 SiC(s) + O2 (g ) → SiO2 (l) + CO 3 3 3
(7)
1 2 Si3 N4 (s) + O2 (g ) → SiO2 (l) + N2 3 3
(8)
4 2 2 BN(s) + O2 (g ) → B2 O3 (l) + N2 3 3 3
(9)
The weight loss rate increased gradually accompanied by reaction in Eqs. (1) and (2) with the increase in oxidation time at 1200 °C. However, as the oxidation time increased, the weight loss rate gradually
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Fig. 2. TGA–DSC–MS curve of C/SiBCN composites and TGA of SiBCN ceramics [16] in flowing air.
Fig. 3. wt loss of the C/SiBCN specimens after oxidation at 1200–1700 °C for 10–120 min.
decreased at 1400 °C, 1500 °C, and 1600 °C because of the weight-increasing oxidation of the matrix of the composites (reaction in Eqs. (6)–(9)). Compared with other oxidation temperatures, the weight loss rate was lower at 1600 °C because an oxide scale formed on the surface of the specimens by the sacrificial oxidation of the SiBCN ceramic matrix (reactions in Eqs. (6)–(9)) to prevent further oxidation of the composites. Specimen C17-120 was catastrophically oxidized, resulting in a weight loss rate of 79.4%, probably because the surface oxide scale was constantly volatilized and the composites were rapidly oxidized due to the direct contact with oxygen. 3.2. Mechanical properties of C/SiBCN composites after oxidation Fig. 4 and Fig. 5 show the open porosity, density, and mechanical properties of the specimens after oxidation treatment, respectively. After the oxidation treatment at 1200 °C, the flexural strength and modulus of the composites decreased gradually as the oxidation time increased (Fig. 5). The weight loss rate was only 14.5% in the C12-10 and C12-30 (Fig. 3) with the open porosity of approximately 20% (Fig. 4), indicating that only a small amount of fiber was oxidized. Therefore, the specimens C12-10 and C12-30 maintained high flexural strength retention rates of 65.4% and 51.9%, respectively (Fig. 5). With
Fig. 4. The open porosities and densities of C/SiBCN specimens after oxidation at 1200 °C–1700 °C for 10 min–120 min.
the increase in oxidation time, the oxidation rate of fibers was increased. The flexural strength and modulus of specimen C12-60 gradually decreased to 101 MPa and 18.5 GPa, respectively (Fig. 5) with the increase of weight loss rate (Fig. 3) and the open porosity (Fig. 4). A
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After oxidation at 1600 °C, an oxide scale formed on the surface of C/SiBCN composites, which healed the pores and cracks, and prevented the diffusion of oxygen into the interior of the composites [16]. Therefore, the flexural strength and flexural modulus of these specimens (approximately 130 MPa and 38 GPa, respectively) was significantly higher than those treated with oxidation at 1400 °C and 1500 °C, which indicated that C/SiBCN composites could maintain its mechanical properties via its self-healing behavior. The density of composites after oxidation at 1600 °C increased with the increase in the oxidation time (Fig. 4). This was because the oxidation reactions of the matrix (reaction in Eqs. (6)–(9)) were aggravated, resulting in a decrease in the weight loss rate of composites, and the open pores of composites were filled by the flow of molten oxides. The flexural strength the composites after oxidation at 1700 °C for 10 min–60 min decreased to 52.1 MPa–62.8 MPa, because the volatilization of the oxide scale led to the gradual oxidation of the fiber. After oxidation at 1700 °C for 120 min, the matrix and fibers of the C17-120 were violently oxidized, resulting in decreasing flexural strength and flexural modulus of C/SiBCN composites to 43 MPa and 6.4 GPa, respectively, (Fig. 5) with 79.4% weight loss (Fig. 3). 3.3. Reason for the evolution of mechanical properties
Fig. 5. The flexural strength and modulus of C/SiBCN specimens after oxidation treatment at 1200 °C–1700 °C for 10 min–120 min.
higher open porosity (> 30% shown in Fig. 4) facilitated the diffusion of oxygen and the oxidation rate of fibers increased as the oxidation time exceeded 60 min. Therefore, the flexural strength and modulus of C12–120 decreased sharply to 22.1 MPa and 7.49 GPa, respectively, with an open porosity of 48.8% (close to fiber volume fraction (44.4%)). The mechanical properties of the specimens after oxidation treatment at 1400 °C for 10 min–120 min were similar to those of treated with oxidation at 1500 °C. The strength retention rates of specimens C14-10 and C15-10 were 50.3% and 37.3% (Fig. 5), respectively, which were noticeably lower than that of C12-10 (65.4%). This is because the oxidation rate of the carbon fibers increased gradually as the oxidation temperature increased, which resulted in higher open porosity (Fig. 4). At the same time, the higher open porosity also increased the oxidation rate of fibers, because it facilitated the diffusion of oxygen. Therefore, the fractural strength and modulus of specimens C14-30 decreased sharply to 33.1 MPa and 21.1 GPa, respectively (Fig. 5). The fractural strength and modulus of C15-30 decreased to 10.4 MPa and 14.8 GPa, respectively (Fig. 5). The flexural strength of C14-120 and C15-120 was slightly higher than those of C14-60 and C15-60. It may be because the porosity of the composites decreased (Fig. 4) caused by the melting of the matrix, which resulted that the matrix loading capacity was increased.
Oxidation-induced damage of the mechanical properties of C/SiBCN composites mainly includes carbon fiber damage, SiBCN matrix damage and interphase damage. The carbon fiber content in the composites after oxidation has a strong effect on the mechanical properties of the composites, not only because the carbon fiber is the main load-bearing phase, but also because oxidized fibers form a large number of holes and cracks, which greatly damage the strength of the composites. The structural integrity of the matrix skeleton is destroyed in a high-temperature oxidizing environment, and this decreases the strength of the composites [19]. However, the oxidized matrix forms a molten glass phase, which contributes to the self-healing properties of C/SiBCN composites. As the oxidation temperature rise above 1200 °C, the residual thermal stress induced by the different expansion coefficients of the SiBCN matrix and the carbon fibers, results in the formation of microcracks at the interface (carbon, approximately 0 °C−1; SiBCN, 3.1 × 10−6–3.5 × 10−6 °C−1) [20], which affects the fracture behavior of the composites. The XRD patterns of the cross section of the specimens after oxidation treatment are shown in Fig. 6. The SiBCN ceramic matrix of the composites was amorphous before and after the high-temperature oxidation, therefore no characteristic diffraction peaks appeared. The peaks at 26° and 43° corresponded to carbon fibers in specimen C12-60; however, these peaks did not appeared in specimen C12-120, which indicated that the carbon fibers were oxidized completely. Therefore, the flexural strength of the C12-60 was maintained at 110 MPa and that of C12-120 is almost completely lost (22.1 MPa). The carbon fibers did not appear in specimens C14-30 and C15-60, which indicated that the fibers were oxidized faster at 1400 °C and 1500 °C than at 1200 °C. This is because the oxidation of carbon fibers was accelerated by the increase of the temperature. Therefore, the fractural strengths of the composites after oxidation at 1400 °C and 1500 °C decreased faster than those of the composites after oxidation at 1200 °C (Fig. 5). However, the carbon fibers were retained well in the C16-120 due to the formation of an oxide scale, resulting in the improved mechanical properties of C16120 (flexural strength: 130 MPa; flexural modulus: 38.0 GPa). SiC grains appeared in C17-30 and C17-60, which could be attributed to the interface reaction (2 C+ 2Si − N→ 2SiC + N2 ↑) and the crystallized of SiBCN matrix [14]. After oxidation at 1200 °C, the high temperature (1200 °C) induced the generation of microcracks at the interface, and oxygen diffused into the microcracks on the surface of composites and reacted with the fibers to form micropores. The microcracks and micropores were the path for the entrance of oxygen into the composites as well as the path for the
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Fig. 6. XRD of the cross-section of C/SiBCN specimens before and after oxidation at 1200 °C-1700 °C for 10–120 min.
Fig. 7. Micrographs of fracture surface of specimens (a) C12-10, (b) C12-30, (c) C12-60, and (d) C12-120.
exit of oxide gases out of the composites [21,22]. The oxidation of fibers inside specimen C12-10 was slow due to the hysteresis of oxygen diffusion through the microcracks and micropores. Carbon fibers in specimen C12-10 were remained as the main load bearing phase (Fig. 7(a–b)). Therefore, the fractural strength and modulus of specimen C12-10 was maintained at 212 MPa and 42.5 GPa, respectively.
Specimen C12-10 showed a typical fracture surface of the composites [23–26]. A certain extent of fiber-debonding and fiber pull-out were observed in specimen C12-10 (Fig. 7 (a)), indicating a relatively weak interface bonding. The fibers were gradually oxidized with an increase in oxidation time and a porous SiBCN matrix and a small amount of fiber pull-out
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Fig. 8. Micrographs of fracture surface of specimens (a) C14-10, (b) C14-30, (c) C14-60, (d) C14-120.
Fig. 9. Micrographs of fracture surfaces of specimens (a, b) C16-30, (c, d) C16-60, and (e, f) C16-120.
were observed in specimen C12-30 and C12-60 (Fig. 7(b–c)). Therefore, the flexural strength of specimen C12-30 and C12-60 decreased to 168 MPa and 101 MPa, respectively (Fig. 5). The fibers of the C12-120 were almost completely oxidized in Fig. 7 (d), which can also be seen in
the XRD pattern of specimen C12-120 (Fig. 6). Further, the structure of SiBCN ceramic matrix in the C12-120 collapsed, and the oval fiber holes were connected to irregular pores (Fig. 7 (d)). Therefore, the mechanical strength of the C12-120 decreased sharply (flexural strength:
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22.1 MPa; flexural modulus: 7.49 GPa). Specimen C14-10 did not show any noticeable obvious oxidation, similar to the as-prepared [18] and C12-10 specimens (Fig. 8 (a)). Therefore, the flexural strength and modulus of specimens C14-10 could retained at 163 MPa and 35.9 GPa, respectively. However, the fibers of specimens C14-30 were rapidly oxidized to form oval holes, while the SiBCN skeleton remained stable (Fig. 8 (b)). The flexural strength and modulus of specimens C14-30 decreased sharply to 33.1 MPa and 21.1 GPa, respectively (Fig. 5). The fibers were almost completely oxidized (Fig. 8 (c)). Furthermore, there were many cracks and holes in the SiBCN matrix of specimens C14-60 (Fig. 8 (c)), which resulted in a poor load bearing capacity and consequently decreased the mechanical properties of specimens C14-60 (fractural strength: 22.6 MPa; fractural modulus: 15.4 GPa). The SiBCN matrix of specimen C14-120 was oxidized and melted (Fig. 8 (d)). Some small pores were filled in C15-60, resulting in an increased density (1.02 g cm−3) and decreased porosity to 38.9% (Fig. 4). Some cracks in the SiBCN matrix were filled, which increased the load bearing capacity of the SiBCN matrix and resulted in enhanced mechanical properties of specimen C14-120 (Fig. 5). The fracture behavior and mechanical properties of the composite after oxidation at 1500 °C are similar to those of the composite after oxidation at 1400 °C, therefore, these are not described here. Fig. 9 shows the fracture surfaces of the composites after oxidation at 1600 °C. The composites exhibited significant self-healing properties at 1600 °C. As shown in Fig. 10, a dense oxide film was observed on the surface of the composites, healing the cracks in the composites. The fiber bundles of the composites after oxidation at 1600 °C showed little pull-out from the matrix of the composites, and the fracture surfaces were flat (Fig. 9), indicating a strong fiber–matrix interface bonding [27,28]. On the surface of the composites, an oxide scale formed by the oxidative melting of the SiBCN ceramic matrix, healing of the pores and cracks. The oxide layer contained Si and O elements (Fig. 10 (b)), which indicated that it was composed of SiO2. Borosilicate glass did not form on the composites due to the volatilization of B2O3 resulting from the decomposition of borosilicate glass at 1600 °C [29]. After oxidation at 1600 °C, the fibers and matrix of the composites did not show significant oxidative damage. The fibers remained intact in specimens C16-30, C16-60, and C16-120, because of the formation of a dense oxide scale on the composites that prevented further diffusion of oxygen into the composites (oxygen diffusion rate of SiO2: 3 × 10−14 g (m s)−1 [30]). Furthermore, the SiBCN matrix after oxidation at 1600 °C showed no obvious oxidative damage. Therefore, the flexural strength and modulus of specimen C16-120 could be maintained at 130 MPa and 38 GPa, respectively. Fig. 11 shows the fracture surfaces of the composite after oxidation at 1700 °C. Due to the difference in the expansion coefficients of the fibers and the matrix, a large number of cracks were generated at the
interface at 1700 °C. Therefore, specimens C17-10, C17-30, and C17-60 exhibited long fiber pull-out with an average length of more than 100 μm (Fig. 11 (b, d, g)). After oxidation at 1700 °C, the oxide scale became porous due to the volatilization of CO2, SiO2, and B2O3 (Fig. 11 (a)) [12]. Oxygen diffused along the porous surface oxide scales into the interior of specimens and the carbon fibers were oxidized. The fiber had a low oxidation rate at 1700 °C, because the oxygen diffusion rate through the porous oxide scale was slow. Therefore, the fibers were retained even in C17-60 (Fig. 11 (g–i)) with a slight increase in porosity (Fig. 4). Significant oxidative damage to fibers (Fig. 11 (e, h)) and interfacial reaction rings (Fig. 11 (f, i)) occurred in the C17-30 and C17-60, resulting in weakened fiber strength. The load bearing capacity of the SiBCN ceramic matrix was degraded because of an increased number of cracks and more isolated ceramic phases. Therefore, the flexural strengths of the specimens after oxidation at 1700 °C were significantly lower than those of the composites after oxidation at 1600 °C (Fig. 5).
4. Conclusion The fracture behavior and mechanical properties of the C/SiBCN composite were reported after oxidation at 1200 °C-1700 °C for different durations in static air. The elevated temperature induced the oxidation reaction of carbon fibers, and the oxidation rate gradually increased as the oxidation temperature increased from 1200 °C to 1500 °C. The flexural strength retention rate of the composites was 51.8% (168 MPa) after oxidation at 1200 °C for 30 min. Even for an oxidation time of 60 min, the flexural strength of the specimens were maintained at 101 MPa. After oxidation treatment at 1400 °C and 1500 °C for 30 min, the strength of the composites was nearly lost completely. The composites exhibited excellent self-healing properties at 1600 °C. The carbon fibers were appropriately retained after oxidation at 1600 °C, because of the formation of a dense oxide scale. The composites maintained a flexural strength of 130 MPa even after oxidation at 1600 °C for 120 min, exhibiting excellent long-term oxidation resistance. However, the mechanical properties of C/SiBCN composites decreased significantly at 1700 °C due to the oxidative damage of the fibers and the interface reaction between the carbon fibers and the matrix. The C/SiBCN composites demonstrate a desirable oxidation resistance at 1200 °C and 1600 °C for a long period of time (30 min and 120 min, respectively). More research focusing on the effects of fiber anti-oxidation coatings on the properties of C/SiBCN composites will be conducted in the future.
Fig. 10. SEM micrograph of the oxide scale (a) and the EDS result (b) of C/SiBCN specimens after oxidation at 1600 °C. 22477
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Fig. 11. Micrographs of fracture surface of specimens (a–c) C17-10, (d–f) C17-30, and (g–i) C17-60.
Acknowledgements This work was supported by the National Program for Key Basic Research Projects [grant numbers 2015CB655200] and the National Natural Science Foundation of China [grant numbers 51602346]. References [1] N.P. Padture, Advanced structural ceramics in aerospace propulsion, Nat. Mater. 15 (2016) 804–809. [2] Y. Ma, S. Wang, Z.H. Chen, Effects of high-temperature annealing on the microstructures and mechanical properties of Cf/SiC composites using polycarbosilane, Mater. Sci. Eng. A 528 (2011) 3069–3072. [3] J. Kong, M.J. Wang, J.H. Zou, L.N. An, Soluble and meltable hyperbranched polyborosilazanes toward high-temperature stable SiBCN ceramics, ACS Appl. Mater. Interfaces 7 (2015) 6733–6744. [4] C.K. Song, L.F. Cheng, Y.S. Liu, M.X. Zhao, F. Ye, Microstructure and electromagnetic wave absorption properties of RGO-SiBCN composites via PDC technology, Ceram. Int. 44 (2018) 18759–18769. [5] X.Y. Ji, S. Wang, C. Shao, H. Wang, High-temperature corrosion behavior of SiBCN fibers for aerospace applications, ACS Appl. Mater. Interfaces 10 (2018) 19712–19720. [6] E. Butchereit, K.G. Nickel, A. Muller, Precursor-derived Si-B-C-N ceramics: oxidation kinetics, J. Am. Ceram. Soc. 84 (2001) 2184–2188. [7] A. Viard, D. Fonblanc, M. Schmidt, A. Lale, D. Lopez-Ferber, M. Balestrat, M. Weinmann, F. Rossignol, Polymer derived Si-B-C-N ceramics: 30 Years of research, Adv. Eng. Mater. 20 (2018) 1800360. [8] R. Riedel, A. Kienzle, L. Ruwisch, W. Dressler, F. Aldinger, J. Bill, A silicoboron carbonitride ceramic stable to 2000 °C, Nature 382 (1996) 796–798. [9] C.J. Luo, Y.S. Tang, T. Jiao, J. Kong, High-temperature stable and metal-free electromagnetic wave-absorbing SiBCN ceramics derived from carbon-rich hyperbranched polyborosilazanes, ACS Appl. Mater. Interfaces 10 (2018) 28051–28061. [10] F. Liu, J. Kong, F. Ye, C.J. Luo, X.G. Luan, et al., High temperature self-healing SiBCN ceramics derived from hyperbranched polyborosilazanes, Adv. Compos. Hybrid Mater. 1 (2018) 506–517. [11] S.H. Lee, M. Weinmann, F. Aldinger, Processing and properties of C/Si-B-C-N fiberreinforced ceramic matrix composites prepared by precursor impregnation and pyrolysis, Acta Mater. 56 (2008) 1529–1538. [12] S.H. Lee, M. Weinmann, Cfiber/SiCfiller/Si-B-C-Nmatrix composites with extremely high thermal stability, Acta Mater. 57 (2009) 4374–4381. [13] H. Zhao, L. Chen, X. Luan, X. Zhang, J. Yun, T. Xu, Synthesis, pyrolysis of a novel liquid SiBCN ceramic precursor and its application in ceramic matrix composites, J. Eur. Ceram. Soc. 37 (2017) 1321–1329. [14] Q. Ding, D.W. Ni, Z. Wang, H.D. Wang, H.J. Zhou, Y.M. Kan, S.M. Dong, Mechanical properties and microstructure evolution of 3D Cf/SiBCN composites at elevated temperatures, J. Am. Ceram. Soc. 101 (2018) 4699–4707.
[15] Q. Ding, D.W. Ni, Z. Wang, S.M. Dong, H. Zhong, H.D. Wang, 3D Cf/SiBCN composites prepared by an improved polymer infiltration and pyrolysis, J. Adv. Ceram. 7 (2018) 266–275. [16] Y. Jia, S.A. Chen, X.Y. Ji, Y.Z. Gou, Y. Li, H.F. Hu, High-temperature oxidation behavior and oxidation mechanism of C/SiBCN composites in static air, Ceram. Int. 45 (2019) 12764–12772. [17] X.J. Wang, Z.B. Zhang, F. Zeng, Y.M. Li, C.H. Xu, Fabrication and properties of carbon fiber reinforced SiBCN matrix composites, Aerosp. Mater. Technol. 43 (2012) 47–50. [18] Y. Jia, X.Y. Ji, S.A. Chen, Y.Z. Gou, Y. Li, H.F. Hu, Preparation of C/SiBCN composites by precursor infiltration and pyrolysis using a novel precursor, Mater. Res. Express 6 (2019) 065612. [19] E. Butchereit, K.G. Nickel, Oxidation behavior of precursor dervied ceramics in the system Si-(B)-C-N, Key Eng. Mater. 175–176 (1999) 69–78. [20] S.H. Lee, F. Aldinger, Processing and thermal stability of a Cfiber/SiCfiller/Si-C-Nmatrix composite: effect of oligomer vaporization and the surface oxide layer of SiC filler, J. Mater. Sci. 44 (2009) 2101–2108. [21] J.T. Hou, S.R. Qiao, Y.B. Zhang, C.Y. Zhang, The influence of high temperature exposure to air on the damage to 3D-C/SiC composites, N. Carbon Mater. 24 (2009) 173–177. [22] X.G. Luan, L. Wang, Y. Zou, Q.Q. Zhang, R. Riedel, L.F. Cheng, Oxidation behavior of C/SiC-SiBCN composites at high temperature, J. Eur. Ceram. Soc. 39 (2019) 3003–3012. [23] Y.D. Wang, C. Han, D. Zheng, Y.P. lei, Large-scale, flexible and high-temperature resistant ZrO2/SiC ultrafine fibers with a radial gradient composition, J. Mater. Chem. A. 2 (2014) 9607–9612. [24] B. Wang, Y. Wang, Y. Lei, N. Wu, Y. Gou, C. Han, S. Xie, D. Fang, Hierarchically porous SiC ultrathin fibers mat with enhanced mass transport, amphipathic property and high-temperature erosion resistance, J. Mater. Chem. A. 2 (2014) 20873–20881. [25] F. Ye, L.T. Zhang, X.W. Yin, Y.J. Zhang, L. Kong, Y.S. Liu, L.F. Cheng, Dielectric and EMW absorbing properties of PDCs-SiBCN annealed at different temperatures, J. Eur. Ceram. Soc. 33 (2013) 1469–1477. [26] F. Ye, L.T. Zhang, X.W. Yin, Y.J. Zhang, L. Kong, Y.S. Liu, L.F. Cheng, Dielectric and microwave-absorption properties of SiC nanoparticle/SiBCN composite ceramics, J. Eur. Ceram. Soc. 34 (2014) 205–215. [27] Y. Mu, W.C. Zhou, H.Y. Wang, C. Wang, Y.C. Qing, Mechanical and dielectric properties of 2.5D SiCf/SiC–Al2O3 composites prepared via precursor infiltration and pyrolysis, Mater. Sci. Eng. A 596 (2014) 64–70. [28] K. Jian, Z.H. Chen, Q.S. Ma, H.F. Hu, W.W. Zheng, Effects of polycarbosilane infiltration processes on the microstructure and mechanical properties of 3D-Cf/SiC composites, Ceram. Int. 33 (2007) 905–909. [29] Y.P. Lei, Y.D. Wang, J.G. Xue, Y.C. Song, Influence of pyrolysis conditions on fabrication of polymer-derived BN fiber for wave transparent application, Compos. Part B 51 (2013) 254–259. [30] S.Y. Chen, G.Q. Jiang, J.J. Yu, D.B. Ou, Passive oxidation mechanism and transition temperature of silicon carbide, Aerosp. Mater. Technol. 39 (2009) 21–24.
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