/. Am Cern. So, 87 ITI In4-12C0KH1 urna Effect of a Boron Nitride Interphase That Debonds between the Interphase and the Matrix in SiC/SiC Composites Gregory N. Morscher Ohio Aerospace Institue, Cleveland, Ohio 44142 Hee Mann Yun*, Cleveland State University. Cleveland Ohio 441 15 James A. DiCarlo NASA Glenn Research Center. Cleveland Ohio 44135 Linus Thomas-Ogbuji* QSS Group, Inc, Cleveland, Ohio 44135 ally, the debonding and sliding interface enabling fibe interfaces as well as oxidation of the fiber surface (Fig. I(a). The it for Sic-tiber-reinforeed Sic-matrix composites with liquid boria reaction product reacts with the SiC fiber to form phase. Recently, composites have been fabricated where inter- oxidation of the SiC. Also. B,O, reacts with water vapor in the occur between the BN interp atmosphere to form volatile B-containing hydrated species result and the matrix. This results in two major improvements in mechanical properties. First, significantly higher failure phenomena result in a solid oxidation product (glass) that strongly strength with no loss in ultimate strength properties of the matrix itself and causes subsequent composite embrittlement(Fig composites. Second, significantly longer stress-rupture times at l(an higher stresses were observed in air at 815 C. In addition, no One proposal to curtail this type of rapid oxidative process that oss in mechanical properties was observed for composites that leads to composite embrittlement would be for the debonding and did not possess a thin carbon layer between the fiber and the sliding interface to be some distance away from the reinforcing interphase when subjected to burner-rig exposure. Two pri- fibers, For SiC/SiC composites this has been attempted with mary factors were hypothesized for the occurrence of debond C/SiC multilayers as the"interphase" and more recently with ng and sliding between the BN interphase and the SiC matrix BN/SiC multilayers. In theory, debonding and sliding would a weaker interface at the BN/matrix interface than the fi. me of the outer layers, prohibiting or complicating the ber/BN interface and a residual tensile/shear stress-state at the diffusion of oxidizing species to the inner fiber/interphase region BN/matrix interface of melt-infiltrated composites. Also, the occurrence of outside debonding was believed to occur during demonstrated for stress-rupture of minicomposites with multilayer C/SiC coating infiltration For SiC/SiC composites with BN interphases. if the debonding and sliding layer was between the BN and the matrix, a simil benefit proposed for the multilayer approach could be achieved I. Introduction Oxidation of the Bn would occur from the "outside"of the BN F OR woven SiC/SiC composites with BN interphases, the typical inwards toward the fiber. The resulting boria oxidation product interface where debonding and sliding occur is between the would react with the SiC matrix to eventually form a borosilicate fiber and the BN interphase, We refer to this phenomenon as glass that would act as a"sealant"slowing diffusion of oxidizing "inside debonding. "Unfortunately, the inside debonding of the species to the BN. In order for the fibers to be fused together or to interphase exacerbates the environmental durability problem of the matrix, oxidation of the entire thickness of the bN would have SiC/SiC composites with BN interphases at intermediate temper- to occur(Fig. 1(b). This may take a considerable amount of time considering the effects of sealing and the reduced surface area of spheres. "2 When matrix cracks are formed, the environment has "inside"debonding case(Figs. I(a)and (b). Therefore, the major BN interphase preferentially at both the fiber/BN and BN/CVI SiC benefit expected from an outside-debonded interphase in SiC/SiC composites would be improved intermediate-temperature mechan cal properties, e.g., stress-rupture. in oxidizing environments. Such behavior has been demonstrated and will be described and R. Naslain--contributing editor discussed in this work I. Experimental Procedure npt No. 186784 Received August 2: approved April 23. 203. the NASA UEET program. Sic-tiber-reinforced melt-infiltrated SiC-matrix composite pan els that exhibited outside debonding were fabricated from 2D. ntist al NASA Glenn Rescarch Center. Cleveland, OH woven, balanced, 5 harness satin, 0/90 fabric, by General Electric
January 200-4 Effect of a BN Interphase Thar Debonds between the Interphase and the Matrix in SiC/SiC Composites a: Inside Debonding Oxidation SiO+B, O, b: Outside Debonding Fig. I. Schematic representation of oxidation of the interphase for (a) debonding and sliding between the fiber and the BN interphase, i c,inside debonding, and (b) between the BN interphase and the matrix, i.e. "outside debonding he composite fabri machine (Instron Model 8562 Instron, Ltd. Canton, MA). Modal cation process involves the following steps: chemical vapor ic emission (AE) was monitored during the room- infiltration (CVI) of a stacked (152 mm x 229 mm) 2D-woven fabric with BN, CVI SiC infiltration, SiC particle slurry infiltra laced rature tests with two wide- band (50 kHz to 2.0 MHz) sensors outside the tapered region of the tensile bar. The AE tion, and final liquid Si infiltration. The occurrence of outside waveforms were recorded and digitized using a fracture wave debonding was initially a processing aberration, but has since been detector(FWD, Digital Wave Corp, Englewood, CO). The AE under study to optimize and control its occurrence. Outside data were filtered using the location software provided by the debonding was observed for over 20 different SiC/SiC composite FWD manufacturer. after the tensile test, to separate out the AE panels fabricated with Sylramic(Dow Corning. Midland, MI hat occurred outside the gauge section fibers, Hi-Nicalon type S(Nippon Carbon, Tokyo, Japan, referred Intermediate-temperature stress-rupture tests were performed to as HNS in the following), and Sylramic-iBN ( treated Sylramic on dogbone specimens using a different universal-testing machine fibers that possess an in sitm BN coating"). Most of the panels (Instron Model 4502 Instron, Ltd.. Canton MA) in air at 815 as were fabricated with Sylramic-iBN or Sylramic fibers and ranged in Ref. 3. Specimens were tabbed with graphite-epoxy composite in fiber volume fraction in the loading direction from 0. 13 to 0.2 tabs. The test specimens were gripped with water-cooled hydraulic (i.e, total fiber volume fraction of 0.26 to 0.40). Table I lists some grips, and a very low load (100 N) was applied to account for f the variations in the physical characteristics of composite thermal expansion of the material during heating. The specimens panels were exposed to elevated temperature using a resistance-heated furnace( MoSi, elements ). Although the fumace was 75 mm long, intermediate-temperature tensile testing. Room-temperature ten the hot zone region was only about 15 mm. When the furnace sile testing was performed on at least two dogbone specimens from reached the desired temperature, 815 C, the load was raised to the each panel. Dogbone specimens, 152 mm long, were cut so that the rupture stress where it was held until failure auge section was 10 mm wide and the grip section was 12.5 mm Specimens from some panels were also subjected to an atr wide Both monotonic and load/unload/reload hysteresis tensile spheric pressure burner-rig under zero-stress exposure at 815C tests were performed at room temperature using a universal-testing i.e., uncracked, and then tensile tested at room temperature to Table I. Physical and Mechanical Properties of Some of the SiC/SiC Composites Tested T(MPa ocation of Estimated trum EIGPat a,n (MPar 7.1/8 0.17 HNS-inside 7.1/8 0.38 SYL--outsid 7.1/6 0.13 224 5.0/8 SYL-mixe 0.19 246 0.33 SYL-insid 8.7/8 0.2 64 SYL-inside 0.17 270 0.31 70 SYL-iBN outside 0.2 SYL-iBN outsid 0.17 220 39 SYL-iBN mixed 0.19 >476 SYLiBn inside 8.7/8 73 SYL-iBN insid 7.9/8 0.2 248 502 0.42 SYI-ibn inside 5.0/8 0.12 279 284 0.2 63 Tow ends per centimeter
106 Journal of the American Ceramic SocietyMorscher et al. Vol, 87. No. I determine the retained strength. The low-pressure burner rig(1.0 measure of the residual stress can be approximated from the atm)uses a high-velocity (Mach 0.3) flame and is designed to intersection of the average slopes of the hysteresis loops for simulate the combustion environments of turbine engines stresses higher than approximately half the peak stress of the Fracture surfaces of the failed composites were examined with hysteresis loop(Fig. 2), 5-60 MPa for the inside-debondi a field emission scanning electron microscope(FESEM), Hitachi composite and -35 MPa for the outside-debonding composite Model S-4700. A fiber push-in technique.4 was performed on Figure 3 shows typical stress-strain curves(hysteresis loops polished sections of untested panels to determine the interfacial removed) for the same architecture MI composites with"outside" shear stress of the sliding interface. At least 20 different fibers and"inside"debonding. In general. although similar in ultimate were tested for each specimen. Finally, the interphase region of strength, two differences between outside- and inside-debonding some specimens was examined using Auger electron spectroscopy composites were evident for room-temperature stress-strain be- (AES) and transmission electron microscopy (TEM), For AES. havior:"outside-debondingcomposites had (1) lower elastic mall slivers of composite material were fractured in bending in situ under vacuum to prevent the fracture surface from contami moduli (Table I)and (2)a higher strain at a given applied stress including higher strains to failure (Table I and Fig. 3). However. nation. Depth profiles were then performed at regions where the one panel, which exhibited a mixture of inside and outside BN layer adhered to the matrix and at other regions where the BN debonding, was an exception and had a high elastic modulus(246 layer adhered to the fiber. Figure 4 shows examples of composite fracture surfaces after IL. Results room-temperature tensile failure. Some bundle pullout was ob- served for both types of composites: however, individual fiber (1) Room-Temperature Tensile Stress-Strain Behavior lout was significantly longer for outside-debonding composites Typical unload-reload tensile hysteresis stress-strain curves (Figs. 4(a) and (b)) than for inside-debonding composites(Figs nd AE activity are shown in Fig. 2 for MI SYL-iBN/SiC c)and (d)). Note the adherence of the BN layer to the fibers for omposite that displays inside and outside debonding. It was he outside-debonding composites(Fig 4(b)) compared with the observed that the first detectable AE that occurs in the gauge outside-debonding composites (Fig. 4(d). It would be ideal if section occurs at 110 t 20 MPa for both inside- and outside debonding outside the Bn interphase occurred for each fiber debonding composites. Also note that on unloading the material independently from one another (e. g, Fig. 1). However, because of tiffens, indicating that the matrix is in residual compression. A he close packing of fibers in woven bundles, debonding between SYL-IBN 14 35087emBf=02 E=280 GPa 岳 0 04日 Strain, 500 SYL-BN 1.4 8. epcm: 8 ply, f0.2 E=216 GP 350 300 250 150 50 0.1 0.2 0.3 0.4 0.5 0.6 Strain, Fig. 2. Tensile load-unload-reload hysteresis curves for (a) inside-debonding and (b) outside-debonding SYL-iBN SiC/SiC composites, Also plotted is the normalized cumulative AE energy. Squares are stress-strain model for best-fit interfacial shear stress
January 2004 Effect of a BN Interphase That Debonds between the Interphase and the Matrit in SiC/SiC Composites 600 was estimated from the measured final crack density of failed composites multiplied by the normalized cumulative AE energy 500 Outside Debond ng ( Fig. 2), assuming the latter represented the stress-dependent Inside Debonding distribution of matrix cracks, which has been demonstrated for Inside Debonding similar systems. Therefore. the only variable not known was T. E f=0.18 which was adjusted to best fit the predicted stress-strain curve to d30018py the experimental stress-strain curve. For the case where the sliding B SYL-iBN lengths overlap, Ahn and Curtin showed that if the cracks are utside Debonding f=0.17 ll equally spaced, the composite strain could then be modeled by 100 18 epi e=o/(,+au/Er-a(o+oM4E S(c)p I (for p26) nm BN layer on the fiber surface. The rupture life depends on the time it takes to bond nearest-neighbor fibers together, which takes where o is the applied stress. o,, is the residual (thermal)stress in more time with increasing separation distance. In addition, the debonding interface for inside-debonding SYL-iBN actually oc- subscripts m, f, and c refer to matrix. fiber. and composite. curs between the in situ BN and the CVI-deposited BN. In other respectively, and pe is the matrix crack density. The first part of the words, for inside-debonding SYL-iBN composites, the debonding equation corresponds to the elastic strain response of an uncracked and sliding interface was some distance(-100 nm)away from the composite and the second part of the equation corresponds to the ber surface, which contained SiC. extra strain (displacement) of the fibers at and away from a For both fiber composite systems possessing an outside- through-thickness matrix crack dictated by the sliding length debonding interface. further improvements in intermediate- temperature stress-rupture life were observed (Fig. 6). For SYL 2 debonding composites, stress-rupture improved by over 250 MPa where in fiber stress (-50 MPa for an /=0.2 composite). For a=(1-f)E=∥E outside-debonding SYL-iBN composites in comparison to inside debonding SYL- BN composites, there was over an order of and ris the fiber radius, f is the fiber volume fraction in the loading magnitude in time improvement at high stresses and --200 MP direction, and T is the interfacial shear strength, Ec and o,were improvement in fiber stress (40 MPa for an/=0,2 composite determined from the stress-strain curves. Er is 380 GPa and E at lower stresses near the run-out condition. It should be noted that was determined from the rule-of-mixtures. The stress-dependent p hese high-stress conditions for stress-rupture are significantly
Journal of the American Ceramie Society-Morscher er al Vol 87. No. 1 2014 Fib BN Fig. 4. FESEM images of fracture surfaces of SYL- iBN composites showing outside debonding (a, b) and inside debonding (e, d). higher than the stresses for matrix cracks to penetrate the load- contact. the thinner areas of Bn were oxidized and fiber-to-fiber bearing fibers (determined from the onset of hysteresis loop fusion occurred for rupture times greater than 80 h. There were activity,-175 MPa for/= 0.2 composites used in this study ) In several regions of significant fiber pullout throughout the cross other words, the SYL-iBN composites are significantly cracked at section of the fracture surface. the stress-rupture conditions of this study, even for specimens that did not fail after long periods of time. L A few specimens(SYL and SYL-iBN) were precracked at room mperature and compared with the rupture behavior of pristine Examination of the rupture specimen fracture surfaces con- specimens from the same panel (Fig. 8). It was evident that firmed the survival of most of the BN around the fibers in the inside-debonding SYL composites with nominally good matrix crack even though significant oxidation had occurred in the properties were significantly poorer in rupture behavior with matrix crack (Fig. 7). However, at regions of near fiber-to-fiber precracking as has been observed in another study. On the other 6001 77 GPa 50050 f=0.18 f 400 Strain. f=0.17 = Change in slope at 70 MPa not associated with occurrence of ae 0.1 030 Strain, Fig. 5. Room-temperature tensile stress-strain curves for a number of outside-debonding composites with difterent fiber volume tractions
January 2004 Effect of a BN Interphase Thar Debonds between the Interphase and the Matrix in SiC/SiC Composites 180 nside Debonding Outside Debonding+350 SYL-iBN 1600 295o00g苏 L1200 Outside Debonding一 T o et cireles a SYL-BN 150 SYL as-produced 600 Time to fail. h The data are plotted as stress on the fibers, Le composite stress divided by f. The composite stress for an/=0.2 is plotted on the right atlsat815C. hand. the outside-debonding specimen with SYL- iBN fibers that for (1) Hi-Nicalon due to a carbon-rich layer that occurs for MI was precracked did not fail after 330 h compared with the pristine composites after fiber and composite processing, (2) Hi-Nicalon specimen which failed after -190 h. With so few data it is not s due to a carbon-rich laver on the fiber surface after fiber possible to conclude that precracked outside-debonding specimens processing,and(3) Sylramiccomposites when a sizing is used re superior to pristine specimens in rupture. However, this is the on the fibers that was not burned off completely before BN first time a precracked SiC/SiC specimen outperformed a nonprec terphase deposition. SYL and SYL-iBN composite racked specimen from the same panel at a stress where time to siz izing that has low char yield are unaffected by burner-rig failure of the nonprecracked specimen was greater than 10 h exposure. Figure 9 compares the burner-rig degradation(or lack Sylramic, SYL-iBN, and HNS outside-debonding composites thereof). Also shown is an example of outside-debonding SYL were subjected to burner-rig exposure at 815C for-100 h with no iBN before and after burner- rig exposure. No significant strength determine the retained strength properties. Whereas the rupture with outside debonding and complete sizing removal. Burner-rig zero-stress burner- rig experiment has proved to be an effective test ing were often observed to stiffen and fracture at slightly lower to evaluate the ability of an undamaged composite material to ultimate strain (Fig 9, Sylramic composite not shown). However withstand severe intermediate-temperature oxidation through the exposed (as-machined)edges of the composite specimen. It has rig exposure due to the presence of a carbon layer on the fiber been found that if carbon exists on the surface of any fiber type, th surface. HNS outside-debonding composites were also observed to SiC/SiC MI composites will be significantly degraded after stiffen slightly after burner-nig exposure Stiffening does not occur burner-rig exposure. This type of degradation has been observed Fig. 7. SEM micrograph from the fracture surface of outside debonding composite after stress-rupture atter-I00 h at 8I5C
110 Journal of the American Ceramic Society Morscher et al Vol. 87. No. I () Improved Mechanical Properties 350 MPa Precrack The increased strain to failure of outside-debonding composites can be attributed to the lower T of the BN-CvI SiC interface over that of the fiber/BN interface. However, if decreased and global oad sharing exists, one would expect the ultimate strength properties to decrease. The converse has been observed for 9190310MPa Precrack strongly bonded interfaces compared with weakly bonded inter- Inside Debonding faces in CVI SiC matrix composites where the higher T interface composites exhibit higher ultimate strengths. In this study, the lower T composites did not lose strength and in some cases were 150 stronger than comparable high-T composites, Two explanations 0 400 can account for this. First, it is possible that high-T composites Time, h exhibit local load sharing and the lower-T composites exhibit global load sharing. If the high-T composites exhibit local load Fig. 8. Stress-rupture of as-received and precracked SYL and SYL-IBN sharing, stress concentrations would develop for load-bearing SiC/SiC composites at 815 C in air fibers surrounding individual or groups of broken fibers in a matrix crack. This would result in lower composite ultimate strengths than expected based on global load sharing. Second, global load sharing may occur for both low- and high-T composites, and as Xia Some fracture surfaces were examined from burner-rig-exposed and Curtin have theorized, fibers with an adhered coating would SYL and SYL-iBN composites. The composites exhibited long be effectively stronger than fibers without a coating because the SiO2-containing layer was often observed on the surface of the BN coating, even for low-modulus coatings such as C or BN. In fact, in between the BN and the matrix throughout the cross section their model would predict approximately the same composite (Fig. 10). Evidently, oxidation occurred through the BN/CVI SiC trength for outside debonding composites with a T= 10 MPa and interface region at the exposed cut edge into the interior of the inside-debonding composites with a T= 70 MPa composite, HNS composites exhibited a flat fracture surface and The improved intermediate-temperature rupture life of outside- trong bonding of fibers as has been reported for other systems debonding composites occurs in the manner put forward in Fig with carbon layers that exist at the fiber surface. 2 I(b)(Fig. 7). Even after 100 h at 815C in a bridged matrix crack a significant portion of the Bn remained as a barrier between the ()Analysis of the BN-CVI SiC Interface oxidation reaction product and the fiber surface for the majority of AES depth profiles were conducted on specimen surfaces that fiber circumference. However, thinner regions of BN separating were fractured in the AES chamber for several specimens exhib- nearest-neighbor fibers were oxidized and appear to have led to the th time-dependent strength degradation. This would explain why the layers adhered to the CVI SiC matrix were performed(not shown). Sylramic composites with outside debonding are poorer in A mild enrichment of C appeared to exist at the BN-CVI SiC stress-rupture than the Sylramic-iBN composites(Fig. 7). The ver no difference in the amount of C enrichment at the BN-CVI little or no BN interphase in between, whereas the latter possesses SiC interface for inside- and outside-debonding composites could the in situ bn layers on the fiber surface that enable greater be discerned given the error in the AES measurement (-10%0) protection of the fibers as well as some degree of fiber separation. Representative TEM micrographs of inside- and outside. Finally, for typical inside-debonding composites that fail at sig. debonding specimens are shown in Fig. 11. Carbon maps of the nificantly shorter lives and lower stresses at the same temperature. ame region are also shown. There does appear to be some C no BN was detectable in the oxidized portion of a fracture surface enrichment at the BN-CVI SiC interface for the outside-debonding (see e composites and little if any C enrichment for the inside-debonding interphase region of inside-debonding composites would be com composites. However. this cannot be considered conclusive evi- ly oxidized, with the fibers strongly bonded to the matrix dence of a c-laver. (2) Why Outside Debonding? Two potential mechanisms are considered for outside debond- '.e for these MI composite systems: (I) a weaker BN-CVI SiC 450 400 SYL-IBN SYL-BN at the interface to cause debonding of the weak interface probably atter bu as-produced on cooling after infiltration of molten Si. The outside-debonding 350 composites possess a lower T than inside-debonding composites 300 (Table I); presumably, the debond energy of the BN/CVI-SiC g20 interface is also lower than the debond energy of the fiber/BN 2), which presumably forces the fibers into residual tension. This 100 is due to free Si. The volume expansion of Si from the liquid to after burner-nig solid state is%. Therefore, expansion of the Si phase takes 50 place during cooling of the composite from its fabrication temper ature for Ml(1400 C depending on the additives to the Si). This 00.10.20.30.40.50.6 places the Si in compression. Si also has a lower thermal expansion coefficient than SiC.-3×10 C compared with-4.5× in. 10/C. respectively. Therefore, on further cooling. the Si is placed in further compression. The crack closure effect(Fig. 2) side debonding. The HNS stress-strain curves are offset on the strain axis fiber/interphase bundles taken together are in residual compression or clarity necessitating residual tension in the fibers. Locally, the residual
January 2004 Effect of a BN Interphase That Debonds betwveen the Interphase and the Matrit in SiC/SiC Composites 0.601201.802.40 0.601.201.8024 601.201.80240 101703009460k11.3mm×10.0kSE(L)02212001 500um Fig. 10. SEM micrograph and EDS spectra for an oxide layer on the outside of the BN, the BN interphase, and the SiC fiber surface of an outside-debonding SYL-iBN SIC/SiC composite fracture surface after 815 C burner-rig exposure and tensile testing at room temperature tress states would be expected to be quite complex. Nevertheless existence of a"gap between the BN and the CVI-SiC, i.e,,an the interphase and interfaces between the fibers and matrix should already debonded interface before testing be subjected to residual tensile and shear stress. This could create the scenario where, if the strength of those interfaces were weak enough, the interface could debond during cooling of the compos- ite or result in a stress state ahead of an approaching crack that could lead to preferential debonding of the BN-CVI SiC interface MI SiC/SiC composites with BN interphases that exhibited rather than the fiber-BN interface In this regard, since MI systems interface debonding and sliding at the BN-CVI SiC interface will inherently have residual compression in the matrix, they may showed significantly higher strain capabilities and intermediate be an ideal composite system to enable outside debonding mperature stress-rupture life over conventional composites that Ithough outside-debonding type of behavior has been observed in exhibit interface debonding and sliding at the BN-fiber interface SiC/BN/CVI SiC minicomposites with tailored BN interfaces. Higher strain to failure was attributed to lower interfacial shear Regarding the observance of a weaker BN-matrix interface for stress at the BN-CVI SiC interface. Improved intermediate outside-debonding CMCs, the presence of carbon either as a thin mperature properties were attributed to the protection from the layer outside of the BN or in an enriched form appears to be the oxidizing environment due to the adherence of the BN layer to the most likely factor, even though the detection of carbon enrichment fiber surface which is not the situation for the inside-debonding is not compelling. One other possible explanation is that differ- composites. Thus the environment does not have direct access to low-temperature-deposited BN. BN shrinkage. the formation of a oxidative process of strongly bonding fibers to nearest-neighbor been observed for fiber/BN/CVI SiC preforms that have been served after 100 h burner- rig exposure, which typically occurs heat-treated to higher temperatures. Nevertheless, oxidation that when carbon exists at the fiber/BN interface occurs between the BN and the CVI-SiC matrix during burner-rig The cause of outside debonding was believed to be due to(1)a xposure clearly implies either the presence of a C layer, as was weaker BN/CVI-SiC interface than BN/fiber i the case for the earlier-mentioned composite systems where a thin caused by the presence of C at the BN/CVI-SiC it C layer existed between the fiber and the BN, 522- or the residual tensile/shear stress at the BN-CVI SiC
Journal of the American Ceramie Soociety-Morscher et al Vol. 87. No. I Sic BN nn 200mm Fig. Il. TEM micrographs (top) and carbon map(bottom, labeled"C")for SYL-iBN SiC/SiC composites showing (a) inside- debonding composite and (b) sufficient to cause interface debonding during cooldown after L U I. I Ogbuji. "A Pervasive Mode of Oxidation Degradation in a SiC-Sic composite processing. The residual stress stale in SiC fiber/MI Composite.JAm Cer, 81 |111 2777(1998 J. I. Eldridge. "Desktop Fiber Push-Out Apparatus, NASA Technical Memo composites is primarily caused by the infiltration of molten Si in hs or of Woven mand 1. I. Eldridge."Constiten Effects on the Strews-Strain 534 December 191) he final step of matrix processing. The volume expansion of the Si expansion coefficient for Si compared with SiC results in residual o/an or of woven Melt-Infiltrated SiC Composites."ICF 10, Intemational Congress iquid to solid phase transformation coupled with a lower thermal re. Elsevier Science. Oxford. U. K. in press American Society for Testing and Materials, West Conshohocken, PA. 1997. A w. Pryee and P. A Smith, "Matrix Cracking in Unithrectienal Ceramic Matri Composites Under Quasi-Static and Cyclie Loading. Acta Metall. Mater. 41 141 1269-81(1993 References Iw. A. Curtin, B, K. Ahn, and N, Takeda. "Modeling Brittle and. Tough Stress- Strain Behavior in Undirectional Ceramie Matrix Compnites, Artu Marer Rupiure of a Woven Hi-Nicokon. BN- Interphase, SiC-Matrix Composite in Air IG. N. Morscher. "Modal Acoustic Emission Source Determination in Silicun Carbide Matrix Composites: pp. 383-ND in Rewiri of Progress in Quantiti N. S. Jacobson. G. N. Morscher. D. R, Bryant, and R. E Tressler, "High- Nondestructive Fiwlutiont. Conference Proceedings 509. Vol 19A. Edited by D. Temperature Oxidation of Boron Nitride: Il. Boron Nitride Layers in Composites Thompson and D. E, Chimenti American Institute of Physics, Melville, NY, 200LI. 1Am. ceram.S,82间61473-82199 Bohlen,"Fiher Coatings for Ceramic Matrix m Ceramic Matrix Compostes,/, Mech Phvs. Solis, 45 121177-209(19971 Composites, Cerm Eng Srl Prox. 13 17-8123-56(19 R. Naslain, "The Concept of Layered Interphases in SiC/SiC": pp. 23-39 in MG. N. Morscher and 1. Hunt."Suress-Rupture and Stress-Relaxation of SiCSic Composites at Intermediate Temperature, Cent Eng. Sci. Proc. 22 131539-46(2001) Manufacturing and Materials Development, Edited by A. G. Evans and R. Naslain G. N. Morscter and J, D. Cawley, "Intermediate Temperuture Strength Degra- Amencan Ceramie Society. Westerville, OH. I ation in SiC/Sic PE Rebilla. J. Lamon, R. Naslain E Lar Curno. M. K. Ferber, and T M. Besmann - H M, Yun. J A DiCarlo. L T Oubuji, and Y L Chen, "Tensile Behavior of Prperties of Multilayered Interphase in SiC/S C Chemcal- vaper-Infiltrateu Compe As-Fabneated and Burmer. Rig Expomed SiC/SIC Composites with Hi-Nicalon Type-S ites with Weak and"Strong Interfaces"/. Ann. Ceran. So, 81 1912315-26(1998) Fibers, Cerim. Eng. Sri. PrINe. in press. "S. Betrand, O, Boison. R. Pailler, J, Lamon, and R. Naslain, "(PyC/SiC) and L U I. T Ogbuji, D, R. Wheeler, and T R: McCue. "Process-Induced Carbon (BNSiC) Nanoseale-Multitayered interphases by Pressure-Pulsed CVI Key Eng ubs-Layer in SIC/BN/SIC Composites, Charactenization and Consequences. "Ceram Mater.164-165.16465(1999 EngS.Pmw,22131379-87(2001 S. Pasquier. J, Lamon and R, Nitslain, "Static Fatigue nF D SIC/SIC Composite A Cunin. ""Theory of Mechanical Properness of Ceramie-Matrix Compow with Multilayered (PyC-SIC) Interphases at High Temperatures in Air, Key Eng JAm. Ceran.So,7412837-4519 "s.Bertrand.R. Pallet, and l. Lama l ifetime of Hi-Nicalon/ PyC/siC)./Si Composites: Localized Load-Sharing and Associated Size Effects." Int. /. Solidy imicemposites,/ A. Ceram. Soe &4 HI Z. Xia and w.A. Curtin, "Design of Fiber/Coating Systems for High Strength in D Brewer, "HSREPM Combustor Materials Development Progran, Muter. Set Ceramic Matrix Composites, Cerat Eng Sa. Prmc, 22 131371-78(201) Eng AA 84-91(19y) G N Morscher, H. Y. Yun, and F. I Hurwitz. " High Temperature Si-Doped BN "hL P Yum and 1, s CAdet "C iC Fiber s- ce Ten. sep and zp In ntss js e: A. Les s c iaC C c vincent. H vi eent. sd i' be in rscsic 259-7201999 Minicompoites with Stnucture-Giraded BN Interphases. "J. Eur. Cera. Srac, 20 G. N. Morscher. "Modal Acoustic Emission of Damage Accumulation in 929-38(2000 Woven SiC/SiC Composite. Comp. Sci. Tec. 59. 687-97(19 R Bhat private commun
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