.Am. Conan Soc.851711815-22(2002 Journal Developing Interfacial Carbon-Boron-Silicon Coatings for Silicon Nitride-Fiber-Reinforced Composites for Improved Oxidation Resistance Kiyoshi Sato, " Hiroki Morozumi, Osamu Funayama, Hiroshi Kaya, and Takeshi Isoda Tonen General Sekiya. 'Iruma-gun. Saitama 356-8500, Japan C-B-Si coatings were formed on a Si,N, fiber using chemical deposition( CVD). That work has demonstrated the advantage of vapor deposition and embedded in a Si-N-C matrix using oxidation resistance on single-filament-reinforced composites polymer impregnation and pyrolysis. The boron-containing only--a microcomposite layer was anticipated to form borosilicate glass and seal A BN and oxide layer has been investigated elsewhere as a oxygen-diffusion passes. Two types of C-B-Si coatings were substitute for the carbon layer. h-BN has a layered crystal sted on the fiber-matrix interface, and they improved the structure, similar to graphite, which causes slip and debondin adhesion between the fiber and matrix in many composites, such as sandwiched between two graphitelike carbon sublayers. The Sic- fiber-reinforced SiclI and SiC-fiber-reinforced glass. 2The second coating was a graphitelike carbon layer containing a small amount of boron and silicon. The carbon (sub)layer of oxidation starting temperature of Bn is only 100 K higher than both coatings weakened the fiber-matrix bonding, giving the arbon, but Bn is expected to form B,O3. which seals the composites a high nexural strength(l1 GPa). The composites oxygen-diffusion passes. A composite with a BN layer shows retained 60%6-70% of their initial strength, even after oxida- oxidation resistance to >1100 K, but it experiences embrittlement tion at 1523 K for 100 h. The mechanism for improved oxidation at intermediate temperatures, 900-1100 K, because of the active resistance was discussed through the microstructure of the oxidation of BN. 3. interface, morphology of the fracture surface, and oxygen An interfacial oxide coating offers the considerable advantage distribution on a cross section of the oxidized composi of causing no oxidation degradation. 5-17 However, an oxide interface is not adequate for a non-oxide fiber, because the oxide layer diffuses oxygen to the fiber. Recently, all-oxide composites F IBER-REINFORCED ceramic composites show promise for over oxidizing atmosphere. 9-22 However, among the disadvantages of coming the brittleness of monolithic ceramics. -Control of all-oxide composites is high-temperature creep above -1500 K, the bonding between fiber and matrix is important to giving which results from an intrinsic property of ordinal oxides. Thus, the problem of improving the oxidation resistance of ceramic a matrix crack from propagating to a fiber. 5 A carbon laver has composites has remained unsolved been applied at the fiber-matrix interface for this purpose, but study focused on a non-oxide composite, the carbon layer oxidizes above -700 K in air, limiting the the high strength of such a composite at high temperature improves oxidation resistance of the composite. " .7 Two approaches have oxidation resistance. A C-B-Si coating was produced using CVD been investigated as a solution to this problem: (i) preventing and applied at the fiber-matrix interface. We anticipated that () oxidation of the carbon layer; and (ii) replacing the carbon layer the C-B-Si layer would form a borosilicate glass, sealing the with an oxidation-resistant layer. oxygen-diffusion passes, as had the BN. and boron-containing Various studies have investigated preventing oxidation of the carbon layers, and (ii)controlling the B: Si ratio in the C-B-Si arbon layer. The addition of boron in a matrix or an interface layer would change the softening temperature and viscosity of the layer, to form a B, O, or borosilicate-glass seal at low tempera- borosilicate glass, improving the performance of the seal. The Ires, has been studied. The oxygen-diffusion passes that must be ternary system C-B-Si was selected to form a carbon-rich phase. sealed are the matrix cracks and interfacial gaps resulting from the which would cause fiber-matrix debonding. A Si-B-N layer could oxidation loss of the carbon layer. The addition of a boride powder to the matrix partially prevents oxidation of the carbon interface, coating was more difficult than a carbon coating to fabricate u CVD. A olymer-impregnation and pyrolysis(PIP)method" boron-containing carbon on a Sic fiber using chemical vapor was used in the present study to fabricate the composite. The dvantage of the PIP method was that it could make use of shaping techniques from the industrial process for fiber-reinforced plastics T.A. Parthasarathy-contributing editor However, improving the oxidation resistance of PIP composites was difficult because of open pores and matrix cracks, formed by pyrolysis shrinkage of the polymer, that diffused oxygen into the composite. The present method of providing oxidation resistance Manuscript No. 188719 Received February 25, 2000: approved September 10. to PIP composites should be effective on other porous composites, ported by the Ministry of Economy, Trade and Industry, and work conducted such as reaction-bonded composites and pore-free composites. such as chemical vapor infiltrated composites Member, American Ceramic Society addressed. Now with Advanced This paper first describes the properties of a C-B-Si 946. Japan. a Si,N fiber. Results for the investigation of compos Nissan Motor Co, Ltd, Atsugi. C-B-Si interfacial layer are described in terms of the fib Formerly Tonen Corporation interface microstructure, mechanical properties, and
1816 Journal of the American Ceramic Society-Sato et al Vol. 85. No. 7 Bobbin Exhaust Precursor gas Exhaust Cⅲi chambe Deposition Feeder chamber Winder N2 N2 obb (seal gas (seal gas) (seal gas Fig. 1. Schematic drawing of the apparatus for fiber coating resistance of the composites Properties of the SiC-fiber-reinforced reinforcement and methy lhydrosilazane(MHS: NN710, Tonen composite produced when a C-B-Si layer is applied to a SiC fiber Corp, Tokyo, Japan)as a matrix sor. Details of the fabrication process and properties of MHS are described else- where 30.31 MHs was a random copolymer, composed of IL. Experimental Procedure -SiH, NH] -[]- units and converted to an amor phous Si-N-C by pyrolysis at 1200 K in a nitrogen-gas or inert (1) Preparation and Analysis of Coatings atmosphere. The chemical composition of the pyrolyzed MHS was Two fibers were used as reinforcements in the present study: () (in mol%)43 silicon, 38 nitrogen, 18 carbon, and I oxygen, with an amorphous Si,Na fiber272* developed by Tonen Corporation and (ii)a commercial SiC fiber( Hinicalon, Nippon Carbon Co, forced composites were fabricated by the following process: () performed in nitrogen gas at 1623 K. Unidirectionally rein- Tokyo, Japan). The SiaN, fiber was a strand composed of 1000 single filaments. The diameter of the filament was 10 um. The brication of unidirectionally fiber-aligned prepreg: (ii) stacking chemical composition of the fiber was (in mass %)60 silicon, 37 and curing of the prepreg sheets;(ii) pyrolysis of a cured sample gen,<I carbon, and <3 oxygen. The fabrication method and at 1623 K; and (iv) densification of the sample by seven cycles of basic character of the fiber are described elsewhere 27, 28 mpregnation and pyrolysis, Details of the process are described The C-B-Si coating was formed using the CVD elsewhere. 30 described in Fig. 1, which consisted of a strand feeder. The mechanical properties of the composites were evaluated in chamber, deposition chamber, and strand winder. The cle terms of flexural strength and interlaminar shear strength(ILSS)at deposition chambers were tubular furnaces with an inner diameter room temperature. The test pieces, which were cut from the of 60 mm and lengths of I and 3 m respectively. A strand of fiber omposite panels, measured 4 mm x 40 mm x 3 mm for the was fed from a bobbin to the cleaning chamber at 33 mm/s and flexural test and 4 mm x 12 mm x 2 mm for the ILSS test. The heated to 1073 K under a nitrogen- gas atmosphere, where the longitudinal direction of the test pieces was aligned with the fiber sizing on the fiber thermally decomposed. The desized strand wa orientation. The flexural strength was measured using the three- oint bend test. with a 30 mm span and a 8.3 um/s testing speed under atmospheric pressure, in the deposition chamber. Nitrogen The span and testing speed of the ILSS test were 8 mm and 8.3 was used as the carrier gas. Finally the coated strand was resized um/'s, respectively. Both tests were completed five times. The with a polyether and wound onto another bobbin in the winder volume content of fiber(V) was calculated from the dimensions of The single-filament strengths of the uncoated fiber and the the composite panel and the amount of fiber used. The bulk density oated fiber were tested, according to ASTM standard. at room was calculated from the weight and dimensions of the test piece temperature (25 mm gauge length, 8.3 um/s testing speed. used for the flexural-strength test. The true density of the compo Cale. Cross-section areas of the filaments, needed for strength ite was measured using a pycnometer at 303 K, using n-butanol as the medium, on a sample crushed under No. 80 mesh For the oxidation test on the composites, the test pieces measuring tron spectroscopy (AES: Model No. PHI650. ULVAC, Ltd. 4 mm X 40 mm X 3 mm were placed inside a tube fumace, heated to a given temperature at a rate of o 167 K/s, and maintained at that current of 3 nA, was used to examine the elemental depth profiles temperature for a given time under a dry -air flow of 74 mmol/s. Some of the coatings. The etching rate, determined using argon sputter samples were exposed at the same temperature and time. under a trogen-gas flow of 74 mmol/s, as references The deposits in the deposition chamber were analyzed The microstructure of the fiber-matrix interface was investi investigate the synthesis mechanism of the multilayered gated using transmission electron microscopy (TEM: Model No Graphite sheets were placed inside the chamber, along its JEM3010, JEOL) using the following techniques: (i) observation wall, before the coating operation. The graphite on a bright-field image; (ii)crystal-structure analysis by nanobeam removed from the furnace after the operation was and electron diffraction(NBED) with a 5 nm beam: and (ii)elemental analyzed using electron probe microanalysis(EPMA analysis by energy dispersive spectroscopy(EDS), with a 5 nm JXA8600MX, JEOL. Tokyo, Japan)and x-ray diffractometry beam. AES depth profiles were investigated on the surfaces of a (XRD; Model No RINT 1400 Rigaku Co., Ltd, Tokyo, Japan) after fracture. The measuring conditions for AES were the same as (2) Preparation and Analysis of the Composite those for the coated fiber Composite panels measuring 100 mm X 100 mm x 3 mm were fabricated using the PIP process, with the coated fiber as a III. Results Two types of fiber coating, coating I and Il. were effective in American the improvement of the oxidation resistance of composites. The iety for Testing and Materials. West Conshohocken, PA deposition conditions of the coatings are shown in Table L. A
July 20( C-B-Si Coatings for S N Fiber-Reinforced Composites for Improved Oxidation Resistance 817 Table 1. Fabrication Conditions and Thicknesses of C-B-Si Coatings Coating condition Thiekness of Gas-flow rate(mmols) Temperature Si,Na Coating Bo I0.29 08000 30-80 R0.000 00-200 02040.60.81.0 Sputter time (ks) eference condition. R, included a carbon coating, the properties of which were described in a previous paper. Figure 2 shows depth Fiber profiles of coatings I and IL Coating thickness, determined by AES alysis, is shown in Table I. Coating I had a boron-containing sublayer Coating II conta ned a small amount of boron. However, the AES depth profile had the following problems, so that only relative changes in elemental composition were valid: (i)disagree ment of the composition of the Si,Na fiber with the chemical analytical value: 27.2 and(ii) miscalculation of 15 mol% of boron on the Si,N, fiber, nevertheless nondetection of a boron peak in the AES spectra. These disagreements were explained by inaccu- races of the default-calculation parameters and misreading of the (b) Sputter t,6.8.0 pectrum background by the calculation program. Therefore, the depth profile on Fig. 2 was revised using the AES spectra as follows. (1) For coating I. the surface consisted of carbon only, and Fig. 2. Depth profile of apparent clemental composition on (a) coating I ron,silicon, and nitrogen were detected at depths of 15. 30, and and (b) coating II using AES 60 nm, respectively. Boron had the maximum value, at a depth of 60 nm. The concentrations of silicon and nitrogen increased gradually and reached saturation at a depth of 120 nm, which he fiber side to the matrix side. The L2 presented the image of a responded to a coating-fiber interface crystal lattice with 0. 26 nm interlayer spacing. The NBED pattern ty,(2) For coating Il, the surface consisted of carbon, and boron, of the L2 layer was obscure spots, suggesting that L2 was licon, and nitrogen were detected at depths of 15, 30, and 30 nm. composed of disordered crystallites. The NBED patterns of bright respectively. Boron did not mark the obvious maximum value. layers LI and L3 lacked a lattice image and showed obscure ring Silicon and nitrogen increased gradually, reaching saturation at a which polarized the brightness toward the direction of the fiber depth of 80 nm surface. Therefore, LI and L3 were composed of crystallites The deposits in the CvD chamber after the preparation of very low crystallinity, oriented toward the fiber surface. EDS was ating I were examined using EPMA and XRD. The deposits used on the fiber-matrix interface to obtain elemental information ke th Boron, although definitely contained in the sample, was not ubstrate graphite negligible on EPMA. Figure 3 shows changes in detected because of limitations in the present equipment. The peak ak intensity for each element along the longitudinal direction of the chamber. The temperature distribution in the chamber wa times and one-forth, respectively, that on the Si-N-C matrix: measured, using a thermocouple inserted into the chamber, under therefore, the interface consisted mainly of carbon and a small the flow of a carrier gas only. No deposit was observed below -0 8 amount of silicon m(0. 8 m upstream from the center of the furnace), At-0 8 n boron and carbon were first detected. The boron attained maximum value at-0.75 m. The carbon intensity exhibited a minimum at the boron maximum and then increased gradually on the downstream side. Silicon was detected from -0 7 m and 1500 exhibited a maximum at-03 m Peaks of B,C were detected from -0.7 to-06 m from XRD analysis of the deposits. Very weak 1.0 peaks of 3C-SiC were detected from-05 to-0.1 m. The existence of carbon in the deposit was not confirmed using XRD, Because 0.8 the intrusion depth of the X-rays was greater than the thickness of B the deposits, the diffraction peaks of the substrate graphite over 0.6 lapped that of the deposit. The single-filament strength of the coated n 0 Table Il. The strength of the as-fabricated fiber was scattered among the fabrication lots of the fiber. Thus, a direct comparison c 0.2 of the strengths of the coated fibers was inadequate: the strength retention ratio was used. The retention ratio of each fiber was 7%-121% (2) Interface between Fiber and Matrix Position from center of furnace (m) e composite reinforced with Si,Na fiber coated with coating I is described here as composite I. Figure 4(a) shows the TEM Fig 3. Change of relative intensity of boron, carbon, and silicon peaks by image of the fiber-matrix interface of non-oxidized composite I EMPA analysis on the deposits in the coating chamber after fabrication of The interface of this composite had a layered structure, composed coating I, Relative intensity was the ratio of the peak intensity of an element on the deposit to that on the simple body of the element under the of a bright layer 5 nm thick(LI), a dark layer 10-15 nm thick same measuring conditions. Origin of the horizontal axis is positioned on (L2), and a bright layer 10-15 nm thick(L3), in that order, from the center of the coating chamber. Negative direction is the upstream side
18 Journal of the American Ceramic Socieny-Sato er al. Vol. 85. No. 7 Table l. Strength and Elastic Modulus of Row Fibers and Coated Fibers As-fabricated fiber Coated fiher Strength(GPa E(GPa) Strength(GPa) Couting SiN 2.0 0.50 182 0.70 0.54 4617 1.90 0.6 7053 106 169 0.65 Figure 4(b) shows a TEM image of the fiber-matrix interface of carbon. Silicon and nitrogen were detected at a depth from 15 nm. non-oxidized composite Il. The interface was a monolayer 20 nm increased slowly, and reached saturation at a depth of 50 nm, where pattern under NBED. Under EDS, the peak intensities of carbon similar to those in Fig. 5(b); the one exception had a boron-containing and one-third that on the Si-N-C matrix; therefore, the interface surface layer 80-110 nm thick, which consisted of carbon andae ad a and silicon on the interface were, respectively, about three times urface layer. In the case of composite Il. the pullout fiber consisted of carbon and a small amount of silicon amount of boron. The surface layer of the fractured matrix consisted Figure 4(c) shows the fiber-matrix interface of composite I mainly of carbon and was 45 vol%),nonbrittle intensity of oxygen was about twice that on LI and L2: on the fracture, and high strength. Table IV shows the strengths of the matrix(M)and the fiber(F). no oxygen was detected composites after oxidation. Composites I and I retained high strengtH coa igure 5 shows the AES spectra of the fractured surfaces of 0.8-0.6 GPa, and maintained their nonbrittle fracture, even after composite L. The AES spectra(Fig. 5) are shown instead of AES oxidation at 1523 K for 100 h. The fracture surfaces of oxidized depth profiles to illustrate the depth distribution of boron, because the composites I and Il had a region 0. 1-0. 4 mm wide around their AES depth profile cannot show boron content, as explained for Fig. 2. periphery, in which fibers failed along a matrix crack plane. The Figure 5(a) is the typical spectra of the surface of a pullout fiber. The center region of the fracture surface showed many pullout fibers. fiber surface consisted mainly of carbon and a small amount of Reference composite R showed a large degradation in strength. 0.2 nitrogen and boron. Silicon was detected from a depth of 15 nm. GPa, and britle fracture after oxidation. No pullout fiber as Silicon and nitrogen gradually increased, reaching saturation at a observed, Heat exposure in a nitrogen-gas flow at 1523 K for 10 h depth of 80 nm because of the proximity of the fiber. Boron was caused no change of the strengths for all composites: therefore, the detected at depths from 15 to 80 nm. Four pullout fibers were strength degradation using the oxidation test resulted from the examined, and all revealed similar results. Figure 5(b) is the typical oxidation AES spectra of the fracture surface of the matrix from which a fiber Figure 6 shows the oxygen-concentration map of a cross section of had pulled out. The surface of the fractured matrix consisted of composites L, IL, and R after oxidation at 1523 K for 100 h. The matrix B L3 H/L2 L1 (a) =10nm(b) 5n 10nm Fig. 4. TEM image of the fiber-matrix interface of the as-fabricated composites (a) I and (b) Il and ()of the oxidized composite I at 1523 K for 10 h. M and F indicate matrix and fiber. respectively. L, LI, L2 and L3 indicate sublayers on the interface. O and B indicate an oxidized sublayer and a bubble
July 2002 C-B-Si Coatings for S,N,- Fiber-Reinforced Composites for Improved Oxidation Resistance 1819 C t=0. 9ks t=0. 6ks AYY t=0.3ks r t=Oks 12 250 30 405 Kinetic energy(ev) N t=0. 9ks p t=0. 6ks p t=0.3ks D teoks 250 300 480 530 155 200 355 Kinetic energy (ev) Fig. 5. Depth change of the first derivative AES spectra on the surfaces of a (a) pullout fiber and(b)matrix from which a fiber was debonded of the fractured composite I. Time interval of each spectra measurement was 60 s, which corresponded to-15 nm in depth composite R, of which the interface was a referential carbon layer, To evaluate the effect of C-B-Si coating for another fiber, fully oxidized to its center region. On the other hand, oxidation of oating I was applied on SiC fiber. The results are shown on omposites I and ll was limited to a region 0. 1-0. 4 mm wide around Tables II-IV. The mechanical properties of the coated fiber and the periphery of the sample. The oxidized region corresponded as-fabricated composite were equal to the Si, N4 fiber. The strength roughly to the region that showed the flat surface after fracture. The of oxidized composite at 1523 K for 100 h was 51% of as- oxidation of composites L, Il, and R proceeded microscopically on the fabricated composite, The strength after oxidation was lower than parts of the matrix adjacent to the fiber-matrix interfaces or the matrix the Si, N fiber-reinforced composite, but higher than the case cracks, and no oxidation occurred on the fibers where carbon coating was used Table Ill. Mechanical Properties of Fabricated Composites Flexural strength(GPa) ILSS (MPa 107 16 Vr is volume faction of fiber, SD is standard deviation. ' ILSS is interlaminar shear strength
1820 Journal of the American Ceramic Society-Sato et al. Table IV. Flexural Strength of Composites after Oxidation EDS, which was applied simultaneously with TEM, was insuffi- fter oxidation at 1523 K for Afler oxidation at 1523 K for ient for lightweight elements, such as boron If it is confirmed that the fiber-matrix interface became coated with almost no change Flexural strength Flexural strength during the PIP composite fabrication, the chemical composition ol (GP the coating applies to the fiber-matrix interface. To confirm the Fiber Coating Average SD rate(S) Average SD rate(%) SD Rate nton difference between the coating and the fiber-matrix interface, the 0.83 surface of the fiber-matrix interface as follows. First the boron containing sublayer, the second sublayer of the coating, was found 0.20003 0.590.03550.510.0348 on the pullout fiber only. Second, the carbon-rich layer. the third sublayer of the coating, was found on the pullout fiber and the SD is standar deviation, Brittle fracture was show fractured matrix surface. These results showed that (i) the layered structure of the coating changed little during the PIP process and that(ii)the fracture of the fiber-matrix interface proceeded in the IV. Discussion outer layer of the interface, which corresponded to the third (1) Structure of Coatings and Fiber-Matrix Interface sublayer of the coating. The difference between the coating The coating was prepared using the CVD reactor, in which thickness of 40-120 nm(Table D) revealed by AES analysis, and fibers were fed continuously. Therefore, it was expected that the the interface thickness of 30-50 nm( Fig. 4), obtained from TEM elemental depth profile of the coating would roughly agree with observation, would present a problem if the coating were to he chemic omposition of the deposit. which changed along the became the fiber-matrix interface with no changes. However the horizontal direction of the CVD reactor. Actually, the depth profile thickness obtained using two TEM observations was considered within the range of eight measurements using AES deposit( Fig 3)on the concentration change of boron and carbon Therefore. the sublayers of the coating and the fiber-matrix AES and EPMA results showed coating I consisted of three interface corresponded to each other as follows. L3 was the third sublayers(from the fiber side ) (i) first sublayer without boron; (ii) sublayer of the coating, a carbon layer with a small amount of econd sublayer of boron, silicon, and carbon; and (iii) third silicon and a graphitelike structure L2 was the second sublayer of sublayer mainly of carbon. The weak peaks of B C and SiC. the coating, composed of crystallites, with a 0.26 nm lattice sublayer: therefore, boron and silicon were determined to be The L2 interlayer spacing was similar to the 0.25 nm of the closest contained as BA C and SiC. packing plane of a-or B-SiC, the 0. 26 nm of the(104)plane of TEM observation of the fiber-matrix interface of composite I B.C. and the 0.24 nm of the(021) plane of B, C. B,C, and SiC that revealed a three-layered structure. However, the chemical compo were detected using XRD on the deposits in the coating apparatus sition of each sublayer was uncertain, because the sensitivity of as mentioned before. Therefore, the crystallite seemed to be SiC 4ze05 4ze05 4ze05 42=02 (b) (d)-20m R(carbon coating). (c)composite L, and (d) composite II using EPMA. (a)SEM image of the composite
July 2002 C-B-Si Coatings for S N,Fiber-Reinforced Composites for Improved Oxidation Resistance 1821 and/or B, C; however, the crystal phase could not be identified supplied from the outside. More et al. showed a borosilicate- because of the obscurity of the electron diffraction pattern of L2. glass layer with bubbles on the fiber-matrix interface of the uncertain. LI was estimated to be a graphitelike carbon, because of was similar to that of the present oxidized composite(Fig 4(c)) the similarity of its TEM image and electron diffraction pattern to These similarities also suggested the formation of borosilicate hose for L3 glass at the interface. In the case of coating Il. TEM observation and EDS analysis of the interface showed that the fiber-matrix interface of this com- V. Conclusions site was a monolayer 20) nm thick, consisting mainly of car ith a small amount of silicon, and having a weak crystal Newly developed C-B-Si interfacial coatings were applied at orientation parallel to the fiber surface. AES analysis of the coated the fiber-matrix interfaces of a Si, N,-fiber-reinforced cor fiber and the fracture surface of the interface rewt as in coating fiber was coated with the C-B-Si layer using CVD and embedded coating structure remained even after the PlP process, L. When the information of the AES analysis was added to the in the Si-N-C matrix by a PIP process. Two types of C-B-Si TEM observation, the interface was determined to be a graphite coatings enhanced the oxidation resistance of the PIP composites. like carbon layer containing a small amount of boron and silicon although the matrix had many cracks, resulting from pyrolysis with an outer carbon-rich sublayer shrinkage of the precursor, that allowed the easy permeation of The fracture of the fiber-matrix interface of composites I and ll oxygen. The first coating, coating L, formed a multilayered proceeded on the matrix side of the interface layer, which fiber-matrix interface, which consisted of three sublayers: a responded to the outer sublayer of the fiber coating. The crystalline sublayer containing boron, silicon, and carbon was debonding on the outer surface of the coating was necessary for the sandwiched between two graphitelike carbon layers. The second coating to be applied on the PIP composite If debonding occurred coating, coating Il, formed a morphologically monolayered inter- at the inner part of the fiber coating, the fiber coating was lost from face, which consisted of a graphitelike carbon layer containing a he fiber surface after cyclic impregnation of PIP. small amount of boron and silicon, Debonding between the fiber and the matrix occurred at the carbon (sub )layer for both of the composites and gave the composites a flexural strength as high as (2) Mechanism for Improving Oxidation Resistance 1.1 GPa. The composites retained 77%(coating I) and 60% The high strength of the composites was obtained by weakened (coating Il) of their original strength, even after oxidation at 1523 fiber-matrix bonding by the carbon sublayer for the monolayered K for 360 ks. Coating I was also effective in the improvement of nterface(composite ID) and the multilayered interface( composite the oxidation resistance of a Sic- fiber-reinforced composite I), as shown by AES analysis of the fracture surfaces. The carbon The mechanism by which oxidation resistance was improved is sublayer prevented the propagation of a matrix crack through the hypothesized as follows. The carbon(sub )layer was easily oxi- fiber as a conventional carbon interface. -5 dized near the surface of the composite. Simultaneously, the The mechanism for improved oxidation resistance apparently boron- and silicon-containing (sub )layer (the center crystalline involves the microstructure of the interface, the morphology of the coating II) formating I and the graphitelike carbon layer itself fo fracture surface, and the oxygen distribution on a cross section of the matrix cracks. As a result, the 0. 1-0.3 nm wide periphery of the oxidized composite. The carbon sublayer adjacent to the surface of the sample oxidizes easily if the sample is exposed the composite showed brittle fracture, caused by the hard bond under an oxidizing atmosphere at high temperature. The advantage between the fiber and the matrix with the borosilicate glass, but th of the C-B-Si interface over a conventional carbon interface is the inside of the composite was unoxidized and showed many long formation of borosilicate glass. The center crystalline sublayer pullout fibers on its fracture surface boron-containing graphitelike carbon layer supplies boron in the the mechanism of oxidation resistance, were directly detec supplies boron and silicon in the case of coating L, and the No borosilicate layers, the existence of which would have case of coating II. Boron-rich borosilicate glass melts at700 K, nterface by microanalytical techniques, such as TEM an ear the starting temperature for the oxidation of carbon. Borosil because of equipment limitations for the detection of boron. However, cate glass seals the matrix cracks and the fiber-matrix gaps the proposed mechanism adequately explained the morphology of the fracture surfaces and the oxygen concentration distributions of the resulting from oxidation loss of the carbon sublayers and prevents he permeation of oxygen into the composite. cross sections of the oxidized composites If few matrix cracks exist in a composite, oxygen permeation is pped within the thin oxygen-sealing layer around the composite References However, a composite fabricated using the PIP process has many E. Fitzer and R iber-Reinforced Silicon Carbide, " Anm Ceran Soc matrix cracks. Therefore, full suppression of oxygen permeation uires a thick oxygen-sealing layer around the composite. In the K. M. Prew,J and G, K. Layden. "Fiber-Reinforced Glasses and oxygen-sealing layer, the borosilicate glass bonds matrix cracks nance Applications," An. Ceram. Soe, BulL, 65 [21 305-13(1986) nd fiber-matrix interfaces. This hard bonding of the interfaces K, M. Prewo,""Fiber-Reinforced Ceramics: New Opportunities for Co results in a flat fracture surface at the periphery of the composite On the other hand, all the C-B-Si interface layers remain unox "R I, Kerans, R. S. Hay. N. J. Pagano, and T. A. Parthasarathy. "The Rol dized at the inside of the composite, causing much fiber pullout on 429-42(1989 iber-Matrix Interface in Ceramic Composites, "Am. Ceran. Soc. Bull,, 68 121 H. C Cao. E. Bischoff. O. Sbaizero, M. Rhule, A G, Evans, D. B. Marshall, and No borosilicate glass, the direct evidence of an antio J.J. Brennan.“Eiec mechanism. was detected using TEM or EDs in the J. Am Ceru. Soc. 73 [61 1691-99(1990 K.M. Prewo,""Fatigue and Stress Rupture of Silicon Carbide-Fiber-Reinforeed ever, the phenomenon is well supported by the hypothesis that L. Filipuzzi, G. Camus R. Naslain, and J. Thebault "Oxidation Mechanisms and borosilicate glass seals the oxygen-diffusion passes In an attempt ID-SiC/C/SiC Composite Materials: L, An Experimental Approach. to determine whether borosilicate glass and B,O, could form by Int. Soc77121455 the oxidation of the other boron-containing interfacial layer, NT, E. Steyer, F. W. Zok, and D. P, Walls, "Stress Rupture of an Enhanced Nicalon/Silicon Carbide Composite at Intermediate Temperatures, "J. Am. Cet Sheldon et al. 5 conducted thermodynamic calculations on the A 8]2140-46(1998 ystem isting of a siC fiber. a Bn interface, and an SiC agano. J. Cab, Y matrix. Lee et al. 7 observed fiber-matrix interfaces for the same Fatigue Behavior in an Enhanced Sic/SiC Composite at High Temperature. "/Am system. Those researchers showed that B,O, and/or borosilicate wm.So,8lp92269-77(998 Guette, F, Langlais, R. Naslain, and S, Goujatd,"High glass formed when the amount of oxygen was high or oxygen was Temperature Lifetime in Air of SiC/C(B)SiC Microeomposites Prepared by
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