Materials Science and Engineering A 498(2008)430-436 Contents lists available at Science Direct Materials Science and Engineering A ELSEVIER journalhomepagewww.elsevier.com/locate/msea Preparation and oxidation resistance of 2D C/Sic composites modified by partial boron carbide self-sealing matrix Yongsheng Liu, Litong Zhang, Aifei Cheng, Xin gang Luan Wenbin Yang, Neihua Zhang, Yongdong Xu, Qingfeng Zeng National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, Xi'an 710072, Shaanxi, People's Republic of china ARTICLE INFO ABSTRACT The C/SiC composites with partial boron carbide(BCx) self-healing matrix were prepared by isothermal pril 2008 or infiltration technique(ICVi). SEM results showed that the multilayer SiC-BCx matrix exist revised form 16 August 2008 18 August 2008 ber bundles, nevertheless there were still SiC matrix among monofilaments. The mechanical prop- room temperature were tested and compared with conventional C/SiC composites. The composites hold better strain to fracture than the conventional C/SiC composites, although th mechanical properties of the two kinds of composites were similar. The pull-out of multilayer 1 carbide as the main factor of the increase of strain to fracture. The high residual strengths were demor after oxidized at static air and simulation combustion environments. The improvement for oxidation resis- tances of the modified composites were contributed to the self-healing of crack and pore during oxidation athermal chemical vapor infiltration in static air and simulation combustion environments ared with conventional C/Sic composites. Oxidation resistance Crown Copyright@ 2008 Published by Elsevier B V. All rights reserved 1. Introduction self-healing matrix, the A400, A410 and A500 materials were devel- oped [ 14], which were tested as seals of F100-Pw-229 engine [15]. ontinuous carbon fiber reinforced silicon carbide(c/sic) com- The application capabilities of A410 and A500 for long-term time posites are potential candidates for a variety of applications in the were demonstrated. aerospace field including rocket nozzles, aeronautic jet engines and Chemical vapor infiltration( Cvi)has been demonstrated to be a aircraft braking systems [1-3. However the oxidation of carbon very effective and matured enough preparation method to fabricate fiber and interface limits long-term applications of C/SiC compos- ceramic matrix composites, for example C/Sic [16, Sic/Sic[17 or ites in high-temperature oxidizing environments. C/Si3N4 [18 composites with ultra-pure and controllable matrix In order to enhance oxidation resistance, boron-containing thickness. materials were widely used to protect C/SiC composites. Boron In the present work, we have attempted to modify the CSic ,B-c[5. B4C [6 and Si-B-c [7. 8 were also applied to the composites using partial boron carbide self-sealing matrix by ICVL xidation protection coa mposites. These Firstly, the microstructure of the modified composites is discussed. boron-containing materials were all used as intermediate layer Secondly, the mechanical properties of the modified composites in multilayer self-healing coatings, with Sic layer as internal and are tested and compared with conventional C/Sic composites. external layer. The oxidation resistances of these C/SiC composites Finally, the oxidation resistance of the composites in static air and coated with the multilayer self-healing coatings had been improved simulation combustion environment are also demonstrated, and static air Multilayer self-healing matrices [9-13 were also devel- compared with C/Sic composites coated with multilayer coatings. oped to improve the oxidation resistance of SiC matrix composites B-C and Si-B-C were used to modify Sic matrix by CNRS-SEP and University of Bordeaux-l in France. The oxidation testes showed 2. Experimental procedures that the self-healing matrix modification can improve the oxI- 2.1. Preparation of specimens dation resistance of SiC matrix composites Based on multilayer Firstly, the 2d preforms were fabricated from laminated car- bon cloth, which was molded by graphite mold in our laboratory. luthor.Tel:+862988486068823;fax:+86 Secondly, pyrolytic carbon(Pyc) interface was deposited on the fiber using C3H6 precursor at 870C for I h at a reduced pressure 0921-5093s-see front matter. Crown Copyright o 2008 Published by Elsevier B V. All rights reserved. doi:10.1016 J. msea200808022
Materials Science and Engineering A 498 (2008) 430–436 Contents lists available at ScienceDirect Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea Preparation and oxidation resistance of 2D C/SiC composites modified by partial boron carbide self-sealing matrix Yongsheng Liu∗, Litong Zhang, Laifei Cheng, Xin’gang Luan Wenbin Yang, Weihua Zhang, Yongdong Xu, Qingfeng Zeng National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, Xi’an 710072, Shaanxi, People’s Republic of China article info Article history: Received 9 April 2008 Received in revised form 16 August 2008 Accepted 18 August 2008 Keywords: Modification Boron carbide C/SiC Multilayer matrices Isothermal chemical vapor infiltration Oxidation resistance abstract The C/SiC composites with partial boron carbide (BCx) self-healing matrix were prepared by isothermal chemical vapor infiltration technique (ICVI). SEM results showed that the multilayer SiC–BCx matrix exist among fiber bundles, nevertheless there were still SiC matrix among monofilaments. The mechanical properties at room temperature were tested and compared with conventional C/SiC composites. The modified composites hold better strain to fracture than the conventional C/SiC composites, although the other mechanical properties of the two kinds of composites were similar. The pull-out of multilayer matrices was the main factor of the increase of strain to fracture. The high residual strengths were demonstrated after oxidized at static air and simulation combustion environments. The improvement for oxidation resistances of the modified composites were contributed to the self-healing of crack and pore during oxidation in static air and simulation combustion environments, compared with conventional C/SiC composites. Crown Copyright © 2008 Published by Elsevier B.V. All rights reserved. 1. Introduction Continuous carbon fiber reinforced silicon carbide (C/SiC) composites are potential candidates for a variety of applications in the aerospace field including rocket nozzles, aeronautic jet engines and aircraft braking systems [1–3]. However the oxidation of carbon fiber and interface limits long-term applications of C/SiC composites in high-temperature oxidizing environments. In order to enhance oxidation resistance, boron-containing materials were widely used to protect C/SiC composites. Boron [4], B–C [5], B4C [6] and Si–B–C [7,8] were also applied to the oxidation protection coatings of C/SiC and C/C composites. These boron-containing materials were all used as intermediate layer in multilayer self-healing coatings, with SiC layer as internal and external layer. The oxidation resistances of these C/SiC composites coated with themultilayer self-healing coatings had been improved in static air.Multilayer self-healingmatrices [9–13] were also developed to improve the oxidation resistance of SiC matrix composites. B–C and Si–B–C were used to modify SiC matrix by CNRS-SEP and University of Bordeaux-1 in France. The oxidation testes showed that the self-healing matrix modification can improve the oxidation resistance of SiC matrix composites. Based on multilayer ∗ Corresponding author. Tel.: +86 29 8848 6068 823; fax: +86 29 8849 4620. E-mail addresses: liuys99067@163.com, yongshengliu@nwpu.edu.cn (Y. Liu). self-healing matrix, the A400, A410 and A500 materials were developed [14], which were tested as seals of F100-PW-229 engine [15]. The application capabilities of A410 and A500 for long-term time were demonstrated. Chemical vapor infiltration (CVI) has been demonstrated to be a very effective and matured enough preparation method to fabricate ceramic matrix composites, for example C/SiC [16], SiC/SiC [17] or C/Si3N4 [18] composites with ultra-pure and controllable matrix thickness. In the present work, we have attempted to modify the C/SiC composites using partial boron carbide self-sealing matrix by ICVI. Firstly, the microstructure of the modified composites is discussed. Secondly, the mechanical properties of the modified composites are tested and compared with conventional C/SiC composites. Finally, the oxidation resistance of the composites in static air and simulation combustion environment are also demonstrated, and compared with C/SiC composites coated with multilayer coatings. 2. Experimental procedures 2.1. Preparation of specimens Firstly, the 2D preforms were fabricated from laminated carbon cloth, which was molded by graphite mold in our laboratory. Secondly, pyrolytic carbon (PyC) interface was deposited on the fiber using C3H6 precursor at 870 ◦C for 1 h at a reduced pressure 0921-5093/$ – see front matter. Crown Copyright © 2008 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.08.022
Y Liu et al. Materials Science and Engineering A 498(2008)430-436 500 Pa, yielding a thickness of 200 nm. Thirdly, two layers of (a) SiC matrices were infiltrated among monofilaments. Subsequently the two layers of Sic and two layers of BCx were infiltrated by alternate ICVI in different reactors. Each layer of Sic matrix was chieved at 1000C for 80 h at reduced pressure of 2 kPa by using methyltrichlorosilane(MTS, CH3 SiCl3) with a H2: MTS molar ratio of 10. This was achieved by bubbling hydrogen gas through the MTS. The argon diluent was used to slow down the chemical reaction rate during deposition. BCx matrix deposition condi- tions were as follows: temperature 950C, pressure 1 kPa, time Oh, boron trichloride(BCl3 299.99 vol% and iron99.95 vol %)flow 10 ml min-I, hydro- gen(H2 >99.999 vol %)flow 60 ml min-, Argon(Ar>99.99 vol%) ow 60 ml min-I. Then, the as-received composite was machined nd polished, and 3 mm x 4 mm x 30 mmsubstrates were obtained Finally, the specimens were coated with three layers CVD SiC coat ing. The conditions for CVD Sic coating were the same as the sic sa17cmm×06 matrix except for the deposition time, which was 30 ha time. On the other hand, in order to compare the mechanical properties, the C/ Sic (b)F. composites were prepared under the same deposition conditions, which had six layers Sic matrix and three layers Sic coatings 2. 2. Characterization and measurements of the composites The morphologies of the BCx layer and the modified com posites were observed with scanning electron microscope(SEM, JSM6700F). The flexural strength of the composite specimens before and after oxidation was measured by a three-point bend- and the loading rate was O5 nure. The span dimension was 20 mm ing method at room tempel Fracture toughness was measured by the single-edge notch beam method with samples of 3 mm x 5 mm x 40 mm. Five samples vere tested. The notch was produced by electro-discharge machin- ing with a depth and a width of 2.5 and 0.2 mm, respectively. the span dimension was 30 mm and the loading rate was 0.5 mm min- m x2.00k The value of fracture toughness(Klc) was calculated by using of the American ASTME 399-74 expression (1) 后)=2(0)2-4829)+2189() 382(1)+23829( where Y is the geometrical factor for an edge crack in a three-point bend beam, calculated by Eq (2): Pc is the fracture load; C is the notch depth; S is the span dimension, H and B are the thickness and broadness of the sample, respectively. Fracture work was calculated by the following formula [19 here ac is characteristic area of fracture curve which refers to the area under load-displacement curve above 90% stress: H and B are Fig. 1. Morphologies of the multilayer matrices of the as-fabricated modified the thickness and broadness of the sample, respectively composites:(a) multilayer matrices among fiber bundles; (b)Sic matrix among mono-filament; (c)EDAX graph of multilayer matrices. The monotonic tensile tests were performed at room temper ature to determine the tensile properties of the material, namel the ultimate tensile stress on a servo-hydraulic machine(Model INSTRON 1196 from INSTRON Ltd, England). Five samples were Tensile strength and tensile modulus was calculated by the fol- tested. The load increased at a constant rate of 0.2 mm min-1 lowing formulas: to fracture of the specimens. The displacement, load were moni- red Strain to failure was measured through strain gage(INSTRON CATALOGUE 2620-601) BH (4)
Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 431 of 500 Pa, yielding a thickness of 200 nm. Thirdly, two layers of SiC matrices were infiltrated among monofilaments. Subsequently the two layers of SiC and two layers of BCx were infiltrated by alternate ICVI in different reactors. Each layer of SiC matrix was achieved at 1000 ◦C for 80 h at reduced pressure of 2 kPa by using methyltrichlorosilane (MTS, CH3SiCl3) with a H2:MTS molar ratio of 10. This was achieved by bubbling hydrogen gas through the MTS. The argon diluent was used to slow down the chemical reaction rate during deposition. BCx matrix deposition conditions were as follows: temperature 950 ◦C, pressure 1 kPa, time 20 h, boron trichloride (BCl3 ≥ 99.99 vol.% and iron ≤ 10 ppm) flow 20 ml min−1, methane (CH4 ≥ 99.95 vol.%) flow 10 ml min−1, hydrogen (H2 ≥ 99.999 vol.%) flow 60 ml min−1, Argon (Ar ≥ 99.99 vol.%) flow 60 ml min−1. Then, the as-received composite was machined and polished, and 3 mm × 4 mm × 30 mm substrates were obtained. Finally, the specimens were coated with three layers CVD SiC coating. The conditions for CVD SiC coating were the same as the SiC matrix except for the deposition time, which was 30 h a time. On the other hand, in order to compare the mechanical properties, the C/SiC composites were prepared under the same deposition conditions, which had six layers SiC matrix and three layers SiC coatings. 2.2. Characterization and measurements of the composites The morphologies of the BCx layer and the modified composites were observed with scanning electron microscope (SEM, JSM6700F). The flexural strength of the composite specimens before and after oxidation was measured by a three-point bending method at room temperature. The span dimension was 20 mm and the loading rate was 0.5 mm min−1. Five samples were tested. Fracture toughness was measured by the single-edge notch beammethod with samples of 3 mm × 5 mm × 40 mm. Five samples were tested. The notch was produced by electro-discharge machining with a depth and a width of 2.5 and 0.2 mm, respectively. The span dimension was 30 mm and the loading rate was 0.5 mm min−1. The value of fracture toughness (K1c) was calculated by using of the American ASTME 399-74 expression: K1c = PC B S H3/2 f C H (1) f C H = 2.9 C H 1/2 − 4.62.9 C H 3/2 + 21.82.9 C H 5/2 −37.62.9 C H 7/2 + 38.72.9 C H 9/2 (2) where Y is the geometrical factor for an edge crack in a three-point bend beam, calculated by Eq. (2); PC is the fracture load; C is the notch depth; S is the span dimension, H and B are the thickness and broadness of the sample, respectively. Fracture work was calculated by the following formula [19]: W = Ac BH (3) where Ac is characteristic area of fracture curve, which refers to the area under load–displacement curve above 90% stress; H and B are the thickness and broadness of the sample, respectively. The monotonic tensile tests were performed at room temperature to determine the tensile properties of the material, namely, the ultimate tensile stress on a servo-hydraulic machine (Model INSTRON 1196 from INSTRON Ltd., England). Five samples were tested. The load increased at a constant rate of 0.2 mm min−1 up to fracture of the specimens. The displacement, load were monitored. Strain to failure was measured through strain gage (INSTRON CATALOGUE 2620-601). Fig. 1. Morphologies of the multilayer matrices of the as-fabricated modified composites: (a) multilayer matrices among fiber bundles; (b) SiC matrix among mono-filament; (c) EDAX graph of multilayer matrices. Tensile strength and tensile modulus was calculated by the following formulas: T = P BH (4)
Y Liu et aL/ Materials Science and Engineering A 498(2008 )430-436 Mechanic properties of the as-fabricated modified composites and C/Sic composites Composite Flexural strength(MPa) Fracture toughness Work of fracture Tensile strength(MPa) Tensile modulus(GPa Strain to failure(%) (MPam) (km-2) Modified composites 4873* 31.7 21.25±2.31 53±1.23 263.35±2432 106.67±1154 C/SiC composites 506.5±276 19.35±1.97 258.13±2675 0.63±00 ET L△P 3. 2. Mechanical properties of the modified composites BH△L where or is tensile strength; P is tensile fracture load. H and b are The mechanical properties of the modified composites and C/Sic the thickness and broadness of the sample, respectively Er is ten- composites are shown in Table 1. It is obvious that the frac- sile modulus: L is test length in sample, which is 30 mm; AP and tensile strength of the modified composites are similar as that of C/Sic composites, though the fle d increasing value, which is corresponded with the linear part ure strength and tensile modulus of the modified composites are on load-displacement curve. AL is displacement which is corre- slightly lower than that of C/SiC composites. The strain to failure of sponded with△P. the modified composites has been improved. The load-displacement curve and fracture morphologies of the modified composites were 23. Oxidation tests shown in Fig. 2. Fig. 2(a)shows that the modified composites had large displacement and high residual strength. The fiber bundle le The static oxidation tests were conducted in a MoSiz furnace and multilayer matrices pull-out were found as shown in Fig. 2(b). in static air environments at temperatures ranging from 700 to There were three pull-out stages to multilayer matrices, and the 1300C for 10h. Three specimens put in an alumina tube with a improved toughness of the modified composites mainly was due purity of 99.99% were used for each experimental condition. They to the multilayer matrices pull-out according to Fig. 2. were introduced into the heating furnace at the desired temper- The better strain to fracture of the modified composites can be ature. The mass of the specimens were recorded after they were explained using micro-crack progress models as shown in Fig 3.The oxidized for 0. 2, 5, and 10h at the given temperature(700, 1000 deflexion of the micro-crack would occur between the interphase electronic balance(sensitivity =0.01 mg) layer matrix to the next layer matrix. Then, the path of the crack The specimens were also tested in a high temperature simula- tion combustion environment coupled with a creep stress for 10, (a)500 and 60 h, respectively. Three sample ronment contained with 12 kPa 0,, 8 kPa H,O and 80 kPa Ar. The otal pressure was 100 kPa and the gas velocity was 0.53 m/s. The mperature of the combustion gases were 700 C detected by a platinum-rhodium thermocouple during testing since this temper ature is pest temperature for C/SiC composites. The stress, which are applied by hydraulic servo frame(INSRTON 8872), were 100 MPa. The fracture faces were observed by SEM 3. Results and discussion 3.1. Morphologies of the modified composites Fig 1 shows that the morphologies of the polished cross-section Displacement/mm of the composites. The( Sic-BCx)multilayer matrices among fiber bundles are shown in Fig. 1(a). Fig. 1(b) shows that the matrix among monofilaments was silicon carbide matrix. The element compositions of monolayer matrix in(Sic-BCx)multilayer matri- ces are shown in Fig. 1(c). It is obvious that the grey matrix layer is Sic matrix, and the dark matrix is BCx matrix. The distribution of different matrix in the fiber preform was correlative with infiltration process. The Sic matrix among monofil nents is derived from the initial two-times infiltration, according to the following reactior CH3SiCl3+H2- SiC H2+HCl pore among monofilaments were fully filled matrix, and the bCx matrix could not occur among monofilaments. Whereafter the BCx and sic matrix appeared among fiber bundles due to the alterative bCx and sic infiltration, according to reaction (7)and reaction(1): Fig. 2. The load-displ rve and fracture morphology of the as-fabricated BCl3+xCH4 +H2- BCx+H2+HCI (7) modified composites: (a)load-displacement curve: (b)fracture morphology
432 Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 Table 1 Mechanic properties of the as-fabricated modified composites and C/SiC composites Composites Flexural strength (MPa) Fracture toughness (MPa m1/2) Work of fracture (kJ m−2) Tensile strength (MPa) Tensile modulus (GPa) Strain to failure (%) Modified composites 487.3 ± 31.7 21.25 ± 2.31 9.53 ± 1.23 263.35 ± 24.32 106.67 ± 11.54 0.87 ± 0.07 C/SiC composites 506.5 ± 27.6 19.35 ± 1.97 7.36 ± 0.96 258.13 ± 26.75 107.25 ± 10.67 0.63 ± 0.06 ET = L BH P L (5) where T is tensile strength; P is tensile fracture load. H and B are the thickness and broadness of the sample, respectively; ET is tensile modulus; L is test length in sample, which is 30 mm; P is load increasing value, which is corresponded with the linear part on load–displacement curve. L is displacement which is corresponded with P. 2.3. Oxidation tests The static oxidation tests were conducted in a MoSi2 furnace in static air environments at temperatures ranging from 700 to 1300 ◦C for 10 h. Three specimens put in an alumina tube with a purity of 99.99% were used for each experimental condition. They were introduced into the heating furnace at the desired temperature. The mass of the specimens were recorded after they were oxidized for 0, 2, 5, and 10 h at the given temperature (700, 1000 and 1300 ◦C), respectively. Sample weights were measured using an electronic balance (sensitivity = 0.01 mg). The specimens were also tested in a high temperature simulation combustion environment coupled with a creep stress for 10, 25 and 60 h, respectively. Three samples were tested. The test environment contained with 12 kPa O2, 8 kPa H2O and 80 kPa Ar. The total pressure was 100 kPa and the gas velocity was 0.53 m/s. The temperature of the combustion gases were 700 ◦C detected by a platinum–rhodium thermocouple during testing since this temperature is pest temperature for C/SiC composites. The stress, which are applied by hydraulic servo frame (INSRTON 8872), were 100 MPa. The fracture faces were observed by SEM. 3. Results and discussion 3.1. Morphologies of the modified composites Fig. 1 shows that the morphologies of the polished cross-section of the composites. The (SiC–BCx) multilayer matrices among fiber bundles are shown in Fig. 1(a). Fig. 1(b) shows that the matrix among monofilaments was silicon carbide matrix. The element compositions of monolayer matrix in (SiC–BCx) multilayer matrices are shown in Fig. 1(c). It is obvious that the grey matrix layer is SiC matrix, and the dark matrix is BCx matrix. The distribution of different matrix in the fiber preform was correlative with infiltration process. The SiCmatrix amongmonofilaments is derived from the initial two-times infiltration, according to the following reaction: CH3SiCl3 + H2 → SiC + H2 + HCl (6) So, the pore among monofilaments were fully filled by SiC matrix, and the BCx matrix could not occur among monofilaments. Whereafter the BCx and SiC matrix appeared among fiber bundles due to the alterative BCx and SiC infiltration, according to reaction (7) and reaction (1): BCl3 + xCH4 + H2 → BCx + H2 + HCl (7) 3.2. Mechanical properties of the modified composites The mechanical properties of the modified composites and C/SiC composites are shown in Table 1. It is obvious that the fracture toughness, fracture work and tensile strength of the modified composites are similar as that of C/SiC composites, though the flexure strength and tensile modulus of the modified composites are slightly lower than that of C/SiC composites. The strain to failure of the modified composites has been improved. The load–displacement curve and fracture morphologies of the modified composites were shown in Fig. 2. Fig. 2(a) shows that the modified composites had large displacement and high residual strength. The fiber bundle and multilayer matrices pull-out were found as shown in Fig. 2(b). There were three pull-out stages to multilayer matrices, and the improved toughness of the modified composites mainly was due to the multilayer matrices pull-out according to Fig. 2. The better strain to fracture of the modified composites can be explained usingmicro-crack progressmodels as shown inFig. 3. The deflexion of the micro-crack would occur between the interphase of the multilayer matrices, when the crack progresses from one layer matrix to the next layer matrix. Then, the path of the crack Fig. 2. The load–displacement curve and fracture morphology of the as-fabricated modified composites: (a) load–displacement curve; (b) fracture morphology.
Y Liu et al/ Materials Science and Engineering A 498(2008)430-436 1000°C Multilayer matrices Oxidation temperature (C) Fig 3. Micro-crack deflexion model in the SiC-BCx multilayer matrices of the mod- ted for 10h ngth of the modified composites and C/SiC before and 2B2Cs)+(172)0(g)03B2O3()+4CO2(g) 10 B2O(1) (10) siO2()+B2O2()∞B2O3,x5i02() (11) 1000 B2O3· aSiO(1) B203(g)+xSiO2(s) “002 2C(s)+O2(g)302co(g) The reactions( 8)and (9) led to weight gain. The other reactions led 006 to weight loss At 700 C, B2O3 liquid would be produced at 700C due to the reaction(8), which led to weight gain. At the same time reactions(10)and(13) would occur which led to weight loss. A previous report 2] indicated that the B2O3 would flow at 630.C Oxidation time /h The liquid flowed through the cracks in the composites, and led to Fig 4. Weight change of modified composites after oxidized for 10h the partially sealing of the crack since the formed speed of B203 from the reaction(9)and flowing speed of B2O3 liquid were not enough faster. The weight change was collective results of reac could be prolonged, and the energy of the micro-crack tions(9)and(13). Therefore, the larger weight loss than that after ogress would be consumed. Therefore the strain to fracture of xidation at 1000 and 1300C and the weight loss occurred at the whole oxidation process based on the same cause. At 1000C, the reactions(8)-(11)and(13)would be faster. Large 3.3. Oxidation resistance of the composites in static( amount of B203 would be formed at this temperature. At the same ime, Sioz would be formed. B203 would accelerate the oxida tion of CVD Sic by forming a melt, according to B2O3-SiOz system Fig 4 shows the weight change of the modified composite after phase diagram [ 21]. The thermal expansion of CVD Sic would also oxidation at 700, 1000 and 1300C. There are similar trend of prompt the crack sealing On the other hand, B2O3 volatilization weight change for the modified composite at different oxidation would also occur at the same time. The weig an co formao s was also temperature, which can be described as: the weight change n-collective results of B2O3 formation, Sioz formation, CO near with oxidation time prolonging, and the weight loss increases and B2O3 volatilization. Based on the above causes, the weight for the initial 2h, then decrease for the next 8 h. The weight changes was less than that of composites after oxidation at 1000C, which after oxidation for 2, 5, and 10 h, respectively, were included in showed that the sealing of the micro-cracks are better than that of Table 2. The maximum weight losses all occurred after oxidation for composites after oxidation at 700C. 2h. At the same time, the weight gain ratio increased with increas- At 1300C, the weight change curve was similar as that at ing temperature. According to Ref [20], the BCx can be written as 1000 C. But the weight losses were less than that at 1000 C dur B3C2. For the modified composites, the oxidation reactions in air ing the initial 2 h. And the weight gains were larger than that at 1000C during the last 4h. This phenomenon can be contributed to the faster formation of B2O3 and B2O3 SiO2 liquid glass. Due to 2SiC(s)+302(g)→2SiO2(1)+2C0(g) 3) the small weight changes at all oxidation temperatures, we can con- clude that the C/(Sic-BCx)n composite have low weight loss after Table 2 oxidized from 700 to 1300.C The maximum, minimum and final weight changes of modified composites after The residual flexural strength of C/SiC and the modified compos oxidized at different temperature in air during 10h ites were compared as shown in Fig. 5. The residual strength of the temperature°c) 1000 modified composite stayed nearly constant after oxidation for 10h change after 2h 0.015 0.053 at 700C, then a little strength increase(102.7% retained strength) hange after 5 h oxidation(%) -0.016 after oxidation at 1000 and 1300.C The residual strength of the change after 10 h oxidation(%)-0021 0.015 C/SiC composites showed a strength loss(73% retained strength)
Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 433 Fig. 3. Micro-crack deflexion model in the SiC–BCx multilayer matrices of the modified composites. Fig. 4. Weight change of modified composites after oxidized for 10 h. progress could be prolonged, and the energy of the micro-crack progress would be consumed. Therefore, the strain to fracture of the composites was improved. 3.3. Oxidation resistance of the composites in static air environment Fig. 4 shows the weight change of the modified composite after oxidation at 700, 1000 and 1300 ◦C. There are similar trend of weight change for the modified composite at different oxidation temperature, which can be described as: the weight change is nonlinear with oxidation time prolonging, and the weight loss increases for the initial 2 h, then decrease for the next 8 h. The weight changes after oxidation for 2, 5, and 10 h, respectively, were included in Table 2. The maximum weight losses all occurred after oxidation for 2 h. At the same time, the weight gain ratio increased with increasing temperature. According to Ref. [20], the BCx can be written as B3C2. For the modified composites, the oxidation reactions in air from 700 to 1300 ◦C are as follows: 2SiC(s) + 3O2(g)800 ◦C −→ 2SiO2(l) + 2CO(g) (8) Table 2 The maximum, minimum and final weight changes of modified composites after oxidized at different temperature in air during 10 h Oxidation temperature (◦C) 700 1000 1300 The weight change after 2 h oxidation (%) −0.074 −0.015 −0.053 The weight change after 5 h oxidation (%) −0.011 0 −0.016 The weight change after 10 h oxidation (%) −0.021 0.015 0.042 Fig. 5. Residual flexural strength of the modified composites and C/SiC before and after oxidized for 10 h. 2B3C2(s) + (17/2)O2(g)600 ◦C −→ 3B2O3(l) + 4CO2(g) (9) B2O3(l)600−1000 ◦C −→ B2O3(g) (10) SiO2(l) + B2O3(l)1000 ◦C −→ B2O3 · xSiO2(l) (11) B2O3 · xSiO2(l)≥1000 ◦C −→ B2O3(g) + xSiO2(s) (12) 2C(s) + O2(g)400 ◦C −→ 2CO(g) (13) The reactions (8) and (9) led to weight gain. The other reactions led to weight loss. At 700 ◦C, B2O3 liquid would be produced at 700 ◦C due to the reaction (8), which led to weight gain. At the same time, reactions (10) and (13) would occur which led to weight loss. A previous report [2] indicated that the B2O3 would flow at 630 ◦C. The liquid flowed through the cracks in the composites, and led to the partially sealing of the crack since the formed speed of B2O3 from the reaction (9) and flowing speed of B2O3 liquid were not enough faster. The weight change was collective results of reactions (9) and (13). Therefore, the larger weight loss than that after oxidation at 1000 and 1300 ◦C and the weight loss occurred at the whole oxidation process based on the same cause. At 1000 ◦C, the reactions (8)–(11) and (13) would be faster. Large amount of B2O3 would be formed at this temperature. At the same time, SiO2 would be formed. B2O3 would accelerate the oxidation of CVD SiC by forming a melt, according to B2O3–SiO2 system phase diagram [21]. The thermal expansion of CVD SiC would also prompt the crack sealing. On the other hand, B2O3 volatilization would also occur at the same time. The weight change was also collective results of B2O3 formation, SiO2 formation, CO formation and B2O3 volatilization. Based on the above causes, the weight loss was less than that of composites after oxidation at 1000 ◦C, which showed that the sealing of the micro-cracks are better than that of composites after oxidation at 700 ◦C. At 1300 ◦C, the weight change curve was similar as that at 1000 ◦C. But the weight losses were less than that at 1000 ◦C during the initial 2 h. And the weight gains were larger than that at 1000 ◦C during the last 4 h. This phenomenon can be contributed to the faster formation of B2O3 and B2O3·SiO2 liquid glass. Due to the small weight changes at all oxidation temperatures, we can conclude that the C/(SiC–BCx)n composite have low weight loss after oxidized from 700 to 1300 ◦C. The residual flexural strength of C/SiC and the modified composites were compared as shown in Fig. 5. The residual strength of the modified composite stayed nearly constant after oxidation for 10 h at 700 ◦C, then a little strength increase (102.7% retained strength) after oxidation at 1000 and 1300 ◦C. The residual strength of the C/SiC composites showed a strength loss (73% retained strength)
Y Liu et aL/ Materials Science and Engineering A 498(2008 )430-436 Micro-crack self-sealing '5. 0oum' Pores self-sealing 50k6.6mmx2009720D6 c Fig. 7. Cross-section morphologies of the modified composite after oxidation for Fig. 6. Surface morphologies of the modified composite after oxidation for 10h: 10h:(a)at 700 C: (b)at 1000 C: (c)at 1300C. at dation from 700 to 1300C for 10 h as shown in Figs. 6 and 7. The after oxidation at 700C, and little strength loss after oxidation at surface micro-crack of the composites was sealed by glass after 1000and1300°C oxidation at 700C. however, there was no obvious glass exist on The high residual strength of the modified composites can be the other position of the surface as shown in Fig. 6(a). The fibe contributed to the good oxidation protection of the (Sic-BCxn mul- and(SiC-BCxn multilayer matrices were not attacked as shown tilayer matrices for carbon fiber and Pyc interface from 700 to in Fig. 7(a). After oxidation at 1000C. a large amount of glass 1300C for 10 h in static air, which can be demonstrated by the appeared and the pore between fiber bundles were filled as shown morphologies of the composites surface and cross-section after oxi- in Fig. 6(b). The fibers and(sic-BCxn multilayer matrices were still
434 Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 Fig. 6. Surface morphologies of the modified composite after oxidation for 10 h: (a) micro-crack self-healing at 700 ◦C; (b) pore self-healing at 1000 ◦C; (c) glass at 1300 ◦C. after oxidation at 700 ◦C, and little strength loss after oxidation at 1000 and 1300 ◦C. The high residual strength of the modified composites can be contributed to the good oxidation protection of the (SiC–BCx)n multilayer matrices for carbon fiber and PyC interface from 700 to 1300 ◦C for 10 h in static air, which can be demonstrated by the morphologies of the composites surface and cross-section after oxiFig. 7. Cross-section morphologies of the modified composite after oxidation for 10 h: (a) at 700 ◦C; (b) at 1000 ◦C; (c) at 1300 ◦C. dation from 700 to 1300 ◦C for 10 h as shown in Figs. 6 and 7. The surface micro-crack of the composites was sealed by glass after oxidation at 700 ◦C, however, there was no obvious glass exist on the other position of the surface as shown in Fig. 6(a). The fibers and (SiC–BCx)n multilayer matrices were not attacked as shown in Fig. 7(a). After oxidation at 1000 ◦C, a large amount of glass appeared and the pore between fiber bundles were filled as shown in Fig. 6(b). The fibers and (SiC–BCx)n multilayer matrices were still
Y Liu et al. Materials Science and Engineering A 498(2008)430-436 Oxidation of pcr bundles Oxidation of Oxidation of fiber Micro-crack self-sealn 馨 Fig 8. Cross-section morphologies of the modified composite after tensile tests at 700C and simulation combustion environment for different time: (a)10h; (b)25 h,(c) nd(d)60h not attacked as shown in Fig. 7(b). After oxidation at 1300C. the ronment (12%02 /8%H20/80%Ar ), the samples were tested with at the same time the bursting glass foam were obsergunt of glass, 100 MPa and at 700oC. The high residual strengths were retained d as shown is shown in Table 3 after tested at different time. After tested for in Fig. 6(c). The multilayer matrices were seriously oxidized, how- 60h, the residual strength was still 219. 24 MPa, and the strength er the fibers were also not obvious oxidized as shown in Fig. 7(c). retained ratio was 83. 25%. Compared with the conventional c/sic The above results showed that the oxidation of the composites composites, the fracture life of the modified composite was obvi ccur at the near surface when the temperature is 700 and 1000.C, ously prolonged. therefore the fibers and ( sic-BCx)n multilayer matrices were not The improvement of the oxidization resistance and fractur seriously attacked which led to high residual strength After oxi- life were explained according to the cross-section morphologies dized at1300°c, the fi ere not obviously attacked although as shown in Fig. 8. After tested for 10h, very small amount of the inner(SiC-BCxnm multilayer matrices were seriously oxidized fiber were oxidized, however, a large amount of fiber were not Therefore the residual strength was still high oxidized as shown in Fig. 8(a), therefore the strength retained ratio achieved 95.44%. after tested for 25h a few fiber were oxi- 3.4. Oxidation resistance of the composites in simulation dized, however, a large amount of fiber were still not oxidized as shown in Fig 8(b). therefore the strength retained ratio achieved 93.55%. The oxidation of multilayer matrices was found. After In order to demonstrate to the oxidation protection of tested for 60h, more fiber were oxidized, the multilayer matri- (Sic-BCx)n multilayer matrices in simulation combustion envi- ces were also oxidized as shown in Fig 8(c). therefore the strength Residual tensile strength of the modified composites after tested at 700 C and simulation combustion environment for different Test time(h) Modified composites(strength retained ratio 26335±24.322(100%)25124±22672(9544%) 246.36±2135(93.55%)21924±20462(8325%) C/SiC composites 25813±26.75(100%) Fractured after tested 6.7 h
Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 435 Fig. 8. Cross-section morphologies of the modified composite after tensile tests at 700 ◦C and simulation combustion environment for different time: (a) 10 h; (b) 25 h; (c) and (d) 60 h. not attacked as shown in Fig. 7(b). After oxidation at 1300 ◦C, the surface of composites was covered with a large amount of glass, at the same time the bursting glass foam were observed as shown in Fig. 6(c). The multilayer matrices were seriously oxidized, however the fibers were also not obvious oxidized as shown in Fig. 7(c). The above results showed that the oxidation of the composites occur at the near surface when the temperature is 700 and 1000 ◦C, therefore the fibers and (SiC–BCx)n multilayer matrices were not seriously attacked which led to high residual strength. After oxidized at 1300 ◦C, the fibers were not obviously attacked although the inner (SiC–BCx)n multilayer matrices were seriously oxidized. Therefore the residual strength was still high. 3.4. Oxidation resistance of the composites in simulation combustion environment In order to demonstrate to the oxidation protection of the (SiC–BCx)n multilayer matrices in simulation combustion environment (12%O2/8%H2O/80%Ar), the samples were tested with 100 MPa and at 700 ◦C. The high residual strengths were retained as shown in Table 3 after tested at different time. After tested for 60 h, the residual strength was still 219.24 MPa, and the strength retained ratio was 83.25%. Compared with the conventional C/SiC composites, the fracture life of the modified composite was obviously prolonged. The improvement of the oxidization resistance and fracture life were explained according to the cross-section morphologies as shown in Fig. 8. After tested for 10 h, very small amount of fiber were oxidized, however, a large amount of fiber were not oxidized as shown in Fig. 8(a), therefore the strength retained ratio achieved 95.44%. After tested for 25 h, a few fiber were oxidized, however, a large amount of fiber were still not oxidized as shown in Fig. 8(b), therefore the strength retained ratio achieved 93.55%. The oxidation of multilayer matrices was found. After tested for 60 h, more fiber were oxidized, the multilayer matrices were also oxidized as shown in Fig. 8(c), therefore the strength Table 3 Residual tensile strength of the modified composites after tested at 700 ◦C and simulation combustion environment for different Composites Test time (h) As-fabricated 10 25 60 Modified composites (strength retained ratio) 263.35 ± 24.32a (100%) 251.24 ± 22.67a (95.44%) 246.36 ± 21.35a (93.55%) 219.24 ± 20.46a (83.25%) C/SiC composites 258.13 ± 26.75 (100%)a Fractured after tested 6.7 h a Residual strength (MPa)
Y Liu et aL/ Materials Science and Engineering A 498(2008 )430-436 retained ratio only 83. 25% The crack sealing were found after oxi- References dized for 60 h as shown in Fig. 8(d). the above results showe that the modified composites have better oxidation resistance in 1I R. Naslain, A Guette E Rebillat, et al. J. Solid State Chem. 177(2004)449-456 simulation combustion environment than that of conventiona Sci Eng. A 35 [4] Y.S. Liu, LF 4. Conclusions 172-177. [5] S. Wu, LF. Cheng. W.B. Yang Y.S. Liu, LT Zhang, Y.D. Xu, Appl. Compos Mater. We have demonstrated the fabrication of modified compos- [ 6S. Goujard, L Vandenbulcke. J Mater. Sci. 29(1994)6212-6220 ites with good mechanical properties and oxidation resistance. ai sGou! S. Goujard, L Vandenbulcke, H Tawil, Thin Solid Films 252(1994)120-130. The Sic-BCx multilayer matrices The modified composites provide similar mechanical proper- oux, Solid State ionics 141-142 (2001)541-54 ties and better strain to fracture than C/SiC composites due to [101 F Lamouroux, S Bertrand, R. Pailler, R. Naslain, Key Eng Mater. 164-165(1999) he pull-out of multilayer matrices though slight lower flexure [111 E Lamouroux, S Bertrand, R Pailler, R N aslan, M Cataldi, Compos.Sci. Technol strength. On the and, the modified composites provide 59(1999)1073-1085 etter oxidation e than C/Sic composites due to the [121 J.P. viricelle, P Goursat, D. Bahloul-Hourlier, Compos. Sci TechnoL. 61(2001) 607-614 self-healing ability of oxide glass which resulted from oxidation [13] L Quemard, E. Rebillat, A Guette, H Tawil, C. Louchet-Pouillerie, ]. Eur. Ceram of bc Soc.27(2007)2085-2094 [14] FA Christin, IntJ AppL. Ceram. TechnoL. 2(2)(2005)97-104 Acknowledgments [15] E bouillon eedings of GT2005, ASME Turbo Expo 2005: Power for Land, Sea and Air, Reno-Tahoe, Nevada, USA, GT2005-68428, June 6-9, 2005. The authors ack e su t of the ch 001)2079-2085 Foundation for Natural Sciences under Contracts No. 90405015. 181 i S LiL, D E heng, I zhang. .D. x u, . Liu, norg, Mater: 20 (5)(2005)979 was also supported by the doctorate Foundation of northwester Polytechnical University(CX200505) [211 T]. Rocket, W.R. Foster. J Am. Ceram. Soc. 48(1965)78-85
436 Y. Liu et al. / Materials Science and Engineering A 498 (2008) 430–436 retained ratio only 83.25%. The crack sealing were found after oxidized for 60 h as shown in Fig. 8(d). The above results showed that the modified composites have better oxidation resistance in simulation combustion environment than that of conventional composites. 4. Conclusions We have demonstrated the fabrication of modified composites with good mechanical properties and oxidation resistance. The SiC–BCx multilayer matrices located among fiber bundles. The modified composites provide similar mechanical properties and better strain to fracture than C/SiC composites due to the pull-out of multilayer matrices though slight lower flexure strength. On the other hand, the modified composites provide better oxidation resistance than C/SiC composites due to the self-healing ability of oxide glass which resulted from oxidation of BCx. Acknowledgments The authors acknowledge the support of the Chinese National Foundation for Natural Sciences under Contracts No. 90405015, No.50672076, No.50642039, No.50802076 and the NSFC Distinguished Young Scholar under Contract No.50425208. This work was also supported by the Doctorate Foundation of Northwestern Polytechnical University (CX200505). References [1] R. Naslain, A. Guette, F. Rebillat, et al., J. Solid State Chem. 177 (2004) 449–456. [2] R. Naslain, Compos. Sci. Technol. 64 (2004) 155–170. [3] J. Schulte-Fischedick, J. Schmidt, R. Tamme, U. Kröner, J. Arnold, B. Zeiffer, Mater. Sci. Eng. A 386 (2004) 428–434. [4] Y.S. Liu, L.F. Cheng, L.T. Zhang, S.J. Wu, D. Li, Y.D. Xu, Mater. Sci. Eng. A 466 (2007) 172–177. [5] S.J. Wu, L.F. Cheng, W.B. Yang, Y.S. Liu, L.T. Zhang, Y.D. Xu, Appl. Compos. Mater. 13 (2006) 397–406. [6] S. Goujard, L. Vandenbulcke, J. Mater. Sci. 29 (1994) 6212–6220. [7] S. Goujard, L. Vandenbulcke, Thin Solid Films 245 (1994) 86–97. [8] S. Goujard, L. Vandenbulcke, H. Tawil, Thin Solid Films 252 (1994) 120–130. [9] R. Naslain, R. Pailler, X. Bourrat, S. Bertrand, F. Heurtevent, P. Dupel, F. Lamouroux, Solid State Ionics 141–142 (2001) 541–548. [10] F. Lamouroux, S. Bertrand, R. Pailler, R. Naslain, Key Eng. Mater. 164–165 (1999) 365–368. [11] F. Lamouroux, S. Bertrand, R. Pailler, R. Naslain, M. Cataldi, Compos. Sci. Technol. 59 (1999) 1073–1085. [12] J.P. Viricelle, P. Goursat, D. Bahloul-Hourlier, Compos. Sci. Technol. 61 (2001) 607–614. [13] L. Quemard, F. Rebillat, A. Guette, H. Tawil, C. Louchet-Pouillerie, J. Eur. Ceram. Soc. 27 (2007) 2085–2094. [14] F.A. Christin, Int. J. Appl. Ceram. Technol. 2 (2) (2005) 97–104. [15] E. Bouillon, G. Ojard, Z. Ouyang, L. Zawada, G. Habarou, C. Louchet, et al., Proceedings of GT2005, ASME Turbo Expo 2005: Power for Land, Sea and Air, Reno-Tahoe, Nevada, USA, GT2005-68428, June 6–9, 2005. [16] S.Q. Guo, Y. Kagawa, J. Am. Ceram. Soc. 84 (9) (2001) 2079–2085. [17] F. Christin, Adv. Eng. Mater. 4 (2002) 903–912. [18] Y.S. Liu, L.F. Cheng, L.T. Zhang, Y.D. Xu, Y. Liu, J. Inorg. Mater. 20 (5) (2005) 979 (in Chinese). [19] Y.D. Xu, L.F. Cheng, L.T. Zhang, J. Chin. Ceram. Soc. 30 (2) (2002) 184–188 (in Chinese). [20] Y.S. Liu, Ph.D. Thesis, Northwestern Polytechnical University, 2008 (in Chinese). [21] T.J. Rocket, W.R. Foster, J. Am. Ceram. Soc. 48 (1965) 78–85
