《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_j.1551-2916.2007.01887.x[1]

J.Am. Ceran.Soc,90o3185-3193(2007 DOL: 10.1111 C 2007 The American No claim to original US government works urna In-Plane Cracking behavior and Ultimate Strength for 2D Woven and Braided Melt-Infiltrated SiC/SiC Composites Tensile Loaded in Off-Axis Fiber directions gregory n. morscher Ohio Aerospace Institute, Cleveland, Ohio Hee mann yun Matech GSM. Irvine. Califor NASA Glenn Research Center. Cleveland, ohio The tensile mechanical properties of ceramic matrix composites referred to as proportional limit stress), and fiber-pullout mech- CMC in directions off the primary axes of the reinforcing fi- anisms that lead to graceful failure and high ultimate tensile bers are important for the architectural design of CMc com- strength (UTS). This is the case because the high-modulus fi nents that are subjected to multiaxial stress states. In this bers are oriented to carry much of the tensile load applied to the study, two-dimensional (2D)woven melt-infiltrated (MD) sic composite before and after the dfls point, which is caused by Sic composite panels with balanced fiber content in the 0 ai matrix cracks 90 directions were tensile loaded in-plane in the 0 direction and (TTMc)in the CMC. However, when loaded in a direction at a at 45 to this direction. In addition, a 2D triaxially braided MI significant angle to the primary fiber axes, large reductions in Sic/SiC composite panel with a higher fiber content in the key CMC design properties such as low DFLS and UTS can +67 bias directions compared with the axial direction was occur. Whereas on-axis properties are strongly dependent on tensile loaded perpendicular to the axial direction tows (i.e 23 fiber properties, off-axis properties are strongly dependent on from the bias fibers). Stress-strain behavior, acoustic emission, matrix properties, particularly on their stiffness and load-carry nd optical microscopy were used to quantify stress-dependent ing ability, which are typically related to their porosity content matrix cracking and ultimate strength in the panels. It was ob- For example, when CMC with highly porous oxide matrices served that both off-axis-loaded panels displayed higher com- were tested off-axis, highly nonlinear stress-strain behavior and posite onset stresses for through-thickness matrix cracking elatively weak strengths were observed because the porous ma- he 2D-woven 0/90 panels loaded in the primary 0 direction. trix could not carry significant load But when the matrix stiff- d load-c be attributed to higher effective fiber fractions in the loading porous matrix, higher dFl stresses and ultimate strengths were direction, which in turn reduces internal stresses on weak regions obtained, but still not as high as the on-axis value ite tows oriented normal For CMC structural applications, high stresses for TTMC are the loading direction and or critical flaws in the matrix for enerally desired in all directions. This not only allows mainte- posite stress. Both off-axis-oriented panels also nance of composite modulus and thermal conductivity to high showed relatively good ultimate tensile strength when compared stresses but also results in greater composite life for CMC with with other off-axis-oriented composites in the literature, both on nonoxide constituents that can be degraded by environmental an absolute strength basis as well as when normalized by the permeability into TTMC. To achieve these high cracking stress- average fiber strength within the composites. Initial implications es, high-stiffness low-porosity matrices are generally preferred are discussed for constituent and architecture design to improve hich can also help to improve the CMc creep-rupture resis- the directional cracking of SiC/SiC CMC components with MI tance and thermal conductivity. One such CMC system is the matrices melt-infiltrated (MD) SiC matrix system reinforced by the syl- ramic-iBN SiC fiber that is near-stoichiometric in composition and contains a thin in sittt-grown BN layer on its surface. This L. Introducti Sic/SiC composite system is typically processed by taking a woven or braided fiber-preform, coating the fibers with another HE tensile mechanical properties of continuous fiber-rein- thin bn layer by chemical vapor infiltration(CVD, and then forced ceramic matrix composites( CMC) vary according to forming an initial SiC matrix by CV. The remaining matrix the orientation of fibers with respect to the loading direction porosity is then filled with a slurry of SiC particles. After drying CMC mechanical properties are typically measured for tensile the slurry-infiltrated preform, final matrix formation is by mel stresses in a direction parallel to one of the primary fiber axes, infiltration (MD) of liquid Si, which fills nearly all of the large hich generally results in desirable linear stress-st ores in the structure, leaving 5% closed porosity a high stress for deviation from linearity or(DFLS)(also The primary objective of this study was to measure and an- alyze the off-axis in-plane tensile stress-strain behavior of five F. Zok--contributing editor thin-walled panels with the MI Sylramic-iBNSIC system con taining two different basic fiber architectures. A and B. Four panels with architecture A, consisting of a two-dimensional (2D)-woven fabric lay-up with balanced fiber content in the 0 anuscript No. 22879 Received March 5. 2007; approved June 2.2007. uthor to whom correspondence should be addressed. e-mail: gmorscher(a and 90 directions, were tensile loaded in the 0 direction(panels grc.nasa.go Al-3)and at 45. to this direction(panel A4). Another panel 3185

In-Plane Cracking Behavior and Ultimate Strength for 2D Woven and Braided Melt-Infiltrated SiC/SiC Composites Tensile Loaded in Off-Axis Fiber Directions Gregory N. Morscherw Ohio Aerospace Institute, Cleveland, Ohio Hee Mann Yun Matech GSM, Irvine, California James A. DiCarlo NASA Glenn Research Center, Cleveland, Ohio The tensile mechanical properties of ceramic matrix composites (CMC) in directions off the primary axes of the reinforcing fi- bers are important for the architectural design of CMC com￾ponents that are subjected to multiaxial stress states. In this study, two-dimensional (2D)-woven melt-infiltrated (MI) SiC/ SiC composite panels with balanced fiber content in the 01 and 901 directions were tensile loaded in-plane in the 01 direction and at 451 to this direction. In addition, a 2D triaxially braided MI SiC/SiC composite panel with a higher fiber content in the 7671 bias directions compared with the axial direction was tensile loaded perpendicular to the axial direction tows (i.e., 231 from the bias fibers). Stress–strain behavior, acoustic emission, and optical microscopy were used to quantify stress-dependent matrix cracking and ultimate strength in the panels. It was ob￾served that both off-axis-loaded panels displayed higher com￾posite onset stresses for through-thickness matrix cracking than the 2D-woven 0/90 panels loaded in the primary 01 direction. These improvements for off-axis cracking strength can in part be attributed to higher effective fiber fractions in the loading direction, which in turn reduces internal stresses on weak regions in the architecture, e.g., minicomposite tows oriented normal to the loading direction and/or critical flaws in the matrix for a given composite stress. Both off-axis-oriented panels also showed relatively good ultimate tensile strength when compared with other off-axis-oriented composites in the literature, both on an absolute strength basis as well as when normalized by the average fiber strength within the composites. Initial implications are discussed for constituent and architecture design to improve the directional cracking of SiC/SiC CMC components with MI matrices. I. Introduction THE tensile mechanical properties of continuous fiber-rein￾forced ceramic matrix composites (CMC) vary according to the orientation of fibers with respect to the loading direction. CMC mechanical properties are typically measured for tensile stresses in a direction parallel to one of the primary fiber axes, which generally results in desirable linear stress–strain behavior, a high stress for deviation from linearity or (DFLS) (also referred to as proportional limit stress), and fiber-pullout mech￾anisms that lead to graceful failure and high ultimate tensile strength (UTS).1 This is the case because the high-modulus fi- bers are oriented to carry much of the tensile load applied to the composite before and after the DFLS point, which is caused by the development of transverse through-thickness matrix cracks (TTMC) in the CMC. However, when loaded in a direction at a significant angle to the primary fiber axes, large reductions in key CMC design properties such as low DFLS and UTS can occur.2 Whereas on-axis properties are strongly dependent on fiber properties, off-axis properties are strongly dependent on matrix properties, particularly on their stiffness and load-carry￾ing ability, which are typically related to their porosity content. For example, when CMC with highly porous oxide matrices were tested off-axis, highly nonlinear stress–strain behavior and relatively weak strengths were observed because the porous ma￾trix could not carry significant load.3 But when the matrix stiff￾ness and load-carrying ability were increased via sintering the porous matrix, higher DFL stresses and ultimate strengths were obtained, but still not as high as the on-axis value. For CMC structural applications, high stresses for TTMC are generally desired in all directions. This not only allows mainte￾nance of composite modulus and thermal conductivity to high stresses but also results in greater composite life for CMC with nonoxide constituents that can be degraded by environmental permeability into TTMC. To achieve these high cracking stress￾es, high-stiffness low-porosity matrices are generally preferred, which can also help to improve the CMC creep-rupture resis￾tance and thermal conductivity. One such CMC system is the melt-infiltrated (MI) SiC matrix system reinforced by the Syl￾ramic-iBN SiC fiber that is near-stoichiometric in composition and contains a thin in situ-grown BN layer on its surface.4 This SiC/SiC composite system is typically processed by taking a woven or braided fiber-preform, coating the fibers with another thin BN layer by chemical vapor infiltration (CVI), and then forming an initial SiC matrix by CVI. The remaining matrix porosity is then filled with a slurry of SiC particles. After drying the slurry-infiltrated preform, final matrix formation is by melt infiltration (MI) of liquid Si, which fills nearly all of the large pores in the structure, leaving B5% closed porosity. The primary objective of this study was to measure and an￾alyze the off-axis in-plane tensile stress–strain behavior of five thin-walled panels with the MI Sylramic–iBN/SiC system con￾taining two different basic fiber architectures, A and B. Four panels with architecture A, consisting of a two-dimensional (2D)-woven fabric lay-up with balanced fiber content in the 01 and 901 directions, were tensile loaded in the 01 direction (panels A1–3) and at 451 to this direction (panel A4). Another panel F. Zok—contributing editor w Author to whom correspondence should be addressed. e-mail: gmorscher@ grc.nasa.gov Manuscript No. 22879. Received March 5, 2007; approved June 2, 2007. Journal J. Am. Ceram. Soc., 90 [10] 3185–3193 (2007) DOI: 10.1111/j.1551-2916.2007.01887.x r 2007 The American Ceramic Society No claim to original US government works 3185

3186 Journal of the American Ceramic Society-Morscher et al. VoL. 90. No. 10 Table I. Individual Specimen Properties From Different Composite Panels Thickness, Total fiber Test angle to Panel 2D architecture fraction ( primary fibers (GPa)(MPa) in matrix(MPa) Al 8.7 epcm 0/90, 8 ply, 5HS, single-tow 0.39 261410 7.9 epcm 090.8 ply, 5HS, single-tow A3 3.95 epcm(2)epi 0/90, 8 ply, 5HS, double-tow 2.1 000 50 A4 8.7 epcm 45/45. 8 ply, 5HS, single-tow Bl 0/+67 Braid, 4 layer, tri-axial braid, double-tow Fraction of fibers in the axial direction=0.06. Fraction of fibers in each of the +67"bias directions =0.13 epcm, ends per centimeter: UTS, ultimate tensile strength; with architecture B(panel Bl), consisting of a 2D triaxially was symmetrically located between the two primary fiber direc- braided architecture [0/+67] with balanced fiber content in the tions that contained equal volume fractions +67 bias directions and reduced fiber content in the axial di- Tensile unload-reload hysteresis tests were performed on the rection, was tensile loaded perpendicular to the axial directio dogbone specimens using a universal testing machine(Model (i.e, 23 from the bias fibers). The primary room-temperature 8562: Instron Ltd, Canton, MA) with AE monitoring as de- properties of interest were the elastic modulus, the DFLs as scribed in Morscher.6, 7 A clip-on strain gage(25.4 mm gage, measured by two off-set methods, the UTS, and the matrix- 0.25% strain) was used to measure strain in the gage section. acking behavior as monitored by acoustic emission(AE). Ini Panels Al, A2, A3, and bl were tested in load control at tial implications are discussed for architecture design to model 4 kN/min (- 200 MPa/min depending on composite thickness) and improve the directional cracking strengths of Sic/SiC CMc This loading rate is relatively slow compared with typical mono- components with MI matrices tonic fast-fracture tests and has been determined to be a good ate for ae acquisition A Fracture Wave Detector by Digital Wave Corporation II. Experimental Procedure Englewood, CO) was used to monitor AE waveforms. Three For this study, four panels with architecture A were fabricated wide-band (b1025, Digital Wave Corporation) AE sensors were by cutting 150 mm x 225 mm plies from a 2D-woven 0/90 mounted on the specimens. Two sensors were placed above and Sylramic SiC fabric with a five-harness satin weave and bal below the gage section approximately 50 mm apart from one anced number of tow ends per centimeter(epcm, i.e., the num- another. The third sensor was placed between the other two ber of fiber tows per centimeter in the woven cloth when looking sensors in the middle of the gage section. aE waveforms on all at the weave edge-on)in the 0 and 90 directions. Each single three channels were captured simultaneously when any of the tow consisted of x800 fibers with a 10 um average diameter. three sensors was triggered, i.e., the channels were synchronized For panels Al and A4, the fabric had 8.7 single-tow epcm; but This allowed for easy separation of events. Only events that the plies were cut along the 0 and 90 directions for panel Al triggered the middle sensor were used in the analysis, i.e,onl and at 450 to the orthogonal directions for panel A4. For panels each of the a panels and two specimens from panel B were A2 and a3, the plies were cut along the oand 90 directions; b the fabric had 7.9 single-tow epcm for panel A2, and 3.95 dou tested in this way for this study (see Table D) ble-tow epcm for panel A3. For each A-type panel, eight plies were then stacked and converted to Sylramic-iBN fiber at NASA Glenn. The 2D Sylramic-iBN stacks were then sent to Orthogonal fibers GE Composites, Newark, DE, for MI SiC/SiC processing For panel Bl, the architecture was first formed by creating a four-layer 0/+67 tri-axial braid on a 50-mm-diameter tubular andre. Approximately 23% of the fibers were in the axial di- rection and 77% were in the bias direction at an angle of 67 to the axial fibers. Two as-produced Sylramic fiber tows were Tensile Axis ar architecture was then removed from the mandrel and laid flat to form a 75 mm x -150 mm rectangular preform which was converted to the Sylramic-iBN fibers at NASA and then into a Sic/sic panel with typical Mi processing at GE Tensile 150-mm-long dogbone specimens with a contoured (a gage section(12.7 mm width in grip region and 10 mm width in age region) were machined from each panel. Architecture thickness, and total fiber volume fraction for all tested speci- IS mens from the five panels are listed in Table I. For panels a A2, and A3 specimens, the testing direction was along the pri- nary or 0 direction; but for panel A4, testing was at 45 to the Tensile 0o and 90 directions of the original fabric as shown in Fig. 1(a) Axis For panel B specimens, Fig. I(b) shows that the testing direction was perpendicular to the axial or 0%fiber direction, or along the hoop"direction of the original tubular architecture. Thus, for both off-axis panels of this study, the tensile loading direction Axial Fibers(below surface) FThe Syiramic fibers of this study were originally produced ochester. NH. Both he Sylramic and Syiramic-iBN SiC fibers are currently produced at ATK COI Ceramics, Fig 1. Photographs of composite surface sh ber orientations San Diego. CA and tensile axis for(a)[+45] panel A4 and (b)[0/+67] braid panel Bl

with architecture B (panel B1), consisting of a 2D triaxially braided architecture [0/767] with balanced fiber content in the 7671 bias directions and reduced fiber content in the axial di￾rection, was tensile loaded perpendicular to the axial direction (i.e., 231 from the bias fibers). The primary room-temperature properties of interest were the elastic modulus, the DFLS as measured by two off-set methods, the UTS, and the matrix￾cracking behavior as monitored by acoustic emission (AE). Ini￾tial implications are discussed for architecture design to model and improve the directional cracking strengths of SiC/SiC CMC components with MI matrices. II. Experimental Procedure For this study, four panels with architecture A were fabricated by cutting 150 mm
225 mm plies from a 2D-woven 0/90 Sylramic SiC fabricz with a five-harness satin weave and bal￾anced number of tow ends per centimeter (epcm, i.e., the num￾ber of fiber tows per centimeter in the woven cloth when looking at the weave edge-on) in the 01 and 901 directions. Each single tow consisted of B800 fibers with a 10 mm average diameter. For panels A1 and A4, the fabric had 8.7 single-tow epcm; but the plies were cut along the 01 and 901 directions for panel A1 and at 451 to the orthogonal directions for panel A4. For panels A2 and A3, the plies were cut along the 01and 901 directions; but the fabric had 7.9 single-tow epcm for panel A2, and 3.95 dou￾ble-tow epcm for panel A3. For each A-type panel, eight plies were then stacked and converted to Sylramic-iBNz fiber at NASA Glenn.5 The 2D Sylramic-iBN stacks were then sent to GE Composites, Newark, DE, for MI SiC/SiC processing.4 For panel B1, the architecture was first formed by creating a four-layer 0/767 tri-axial braid on a 50-mm-diameter tubular mandrel. Approximately 23% of the fibers were in the axial di￾rection and 77% were in the bias direction at an angle of B671 to the axial fibers. Two as-produced Sylramic fiber tows were combined in the axial and bias directions. A 75 mm-length of the tubular architecture was then removed from the mandrel and laid flat to form a 75 mm
B150 mm rectangular preform, which was converted to the Sylramic-iBN fibers at NASA and then into a SiC/SiC panel with typical MI processing at GE composites. Tensile 150-mm-long dogbone specimens with a contoured gage section (12.7 mm width in grip region and 10 mm width in gage region) were machined from each panel. Architecture, thickness, and total fiber volume fraction for all tested speci￾mens from the five panels are listed in Table I. For panels A1, A2, and A3 specimens, the testing direction was along the pri￾mary or 01 direction; but for panel A4, testing was at 451 to the 01 and 901 directions of the original fabric as shown in Fig. 1(a). For panel B specimens, Fig. 1(b) shows that the testing direction was perpendicular to the axial or 01 fiber direction, or along the ‘‘hoop’’ direction of the original tubular architecture. Thus, for both off-axis panels of this study, the tensile loading direction was symmetrically located between the two primary fiber direc￾tions that contained equal volume fractions. Tensile unload–reload hysteresis tests were performed on the dogbone specimens using a universal testing machine (Model 8562; Instron Ltd., Canton, MA) with AE monitoring as de￾scribed in Morscher.6,7 A clip-on strain gage (25.4 mm gage, 0.25% strain) was used to measure strain in the gage section. Panels A1, A2, A3, and B1 were tested in load control at 4 kN/min (B200 MPa/min depending on composite thickness). This loading rate is relatively slow compared with typical mono￾tonic fast-fracture tests and has been determined to be a good rate for AE acquisition. A Fracture Wave Detector by Digital Wave Corporation (Englewood, CO) was used to monitor AE waveforms. Three wide-band (B1025, Digital Wave Corporation) AE sensors were mounted on the specimens. Two sensors were placed above and below the gage section approximately 50 mm apart from one another. The third sensor was placed between the other two sensors in the middle of the gage section. AE waveforms on all three channels were captured simultaneously when any of the three sensors was triggered, i.e., the channels were synchronized. This allowed for easy separation of events. Only events that triggered the middle sensor were used in the analysis, i.e., only events that occurred in the gage section. One specimen from each of the A panels and two specimens from panel B were tested in this way for this study (see Table I). Table I. Individual Specimen Properties From Different Composite Panels Panel 2D architecture Thickness, mm Total fiber fraction (f) Test angle to primary fibers E (GPa) UTS (MPa) Residual stress in matrix (MPa) A1 8.7 epcm 0/90,8 ply, 5HS, single-tow 2.3 0.39 01 261 410 60 A2 7.9 epcm 0/90, 8 ply, 5HS, single-tow 2.0 0.39 01 250 463 50 A3 3.95 epcm (2) epi 0/90, 8 ply, 5HS, double-tow 2.1 0.39 01 202 444 50 A4 8.7 epcm 45/45,8 ply, 5HS, single-tow 2.4 0.36 451 233 242 40 B1 0/167 Braid, 4 layer, tri-axial braid, double-tow 1.8 0.32w 231 240 338 60 1.8 0.32 231 260 366 60 w Fraction of fibers in the axial direction 5 0.06. Fraction of fibers in each of the 1671 bias directions 5 0.13. epcm, ends per centimeter; UTS, ultimate tensile strength; 2D, two-dimensional. Fig. 1. Photographs of composite surface showing fiber orientations and tensile axis for (a) [745] panel A4 and (b) [0/767] braid panel B1. z The Sylramic fibers of this study were originally produced by Dow Corning, Midland, MI, and were woven into fabric at Albany International Techniweave, Rochester, NH. Both the Sylramic and Sylramic-iBN SiC fibers are currently produced at ATK COI Ceramics, San Diego, CA. 3186 Journal of the American Ceramic Society—Morscher et al. Vol. 90, No. 10

October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 500 7.9 epcm [Orgo From Table l, it can be seen that in-plane E values are similar 8.7 epcm [o/90 for all panels with the exception of the double-tow 5HS-woven composite(panel A3). This suggests that for the Sic/SiC system 350[0/467] double studied, the in-plane elastic modulus is not strongly dependent 3 95 epem [090 on 2D fiber architectures or on tensile-loading direction. For in- 300 f=0.39 plane UTS, the 2D-woven 0/90 panels aligned in the primary fiber axes are of course the strongest because they were test 8.7 epem [+45] parallel to the fiber direction. Nevertheless, the braided panel with over three-quarters of the fibers oriented 23 from the loading-axis displayed a high in-plane UTS; however, the [+45 panel displayed a relatively low UTS. But perhaps the most 50 striking mechanical results were the high in-plane DFLs values a3 for both the braided and the [+] panel 040.506 These high DFLS results correlate(Fig. 4 and Table In) with higher stress ranges over which high-energy ae events were re- corded in the off-axis panels. In Fig. 4. the aE activity versus Fig. 2. Stress-strain curves of off-axis and orthogonally aligned com- applied CMC stress is plotted as normalized cumulative AE en- psites with the hysteresis loop removed for clarity. An example of the ergy, which is the cumulative AE energy of each AE event up ysteresis loops for 0/+67 braid (I)is shown in the inset. a given event divided by the total AE energy of all the eve o For the MI siC/SiC system, it has been shown that this type of plot represents a good relative distribution of transverse or direction in order to ns were cut and polished along the TTMC. In addition, multiplying the final matrix crack density measured from polished sections after failure, by the normalized The polished specimens were plasma etched(CFa at 500 cumulative AE energy is a very good estimate of the actual for 30 min) in order to enhance the matrix cracks in the CvI sic stress-dependent matrix crack density versus applied stress. The of the matrix. matrix cracks were counted over lengths of neasured matrix crack densities at failure are shown in Table ll approximately 10 mm in order to obtain a matrix crack density. and the estimated matrix crack density with stress in Fig. 5. The difference in the shape of the matrix crack distribution between e double-tow and single-tow 2D-woven composites is hypoth esized to be due to the greater concentration and longer lengt II Results of back-to-back 90 minicomposites (see Section IV). Two The room-temperature tensile stress-strain curves for the off- axis-oriented SiC/SiC panels A4 and Bl and for the orthogonal ifferent off-axis-loaded specimens. The [+45] specimen has a oriented 2D-woven SiC/SiC panels Al-3 are shown in Fig. 2 ery steep curve, whereas the braided specimens were similar in Hysteresis loops used for residual stress determination have shape to the single-tow 2D composites. Neither of the off-axis been eliminated from these curves except for a braided panel composites appears to have reached matrix crack saturation be BI specimen, which is given as an example in Fig. 2. The values ause a decrease in the rate of ae activity was never achieved at of initial elastic modulus e. UTs. and residual stress on matrix her stresses. It is also interesting to note that at a composit stress of 200 MPa, the off-axis-tested braided specimens had are listed in Table I and are consistent with other composite little or no cracks: however, the double-tow 2D-woven compo panels fabricated with the same fiber types and fiber architec tures , As shown in Fig. 3, the offset strain construct method ites had -5 TTMC/mm. was also used to determine DFLs. It consists of drawing a line For this study, several aE criteria were used to evaluate stress levels near initiation of matrix cracking as these values ith the same slope as the initial elastic modulus, but offset by some amount of positive strain, where the dFls would be de- t upper design limits beyond which CMC mod- termined by the intersection of that curve with the stress-strain and axial thermal conductivity irreversibly decrease and the curve. Typical offset-strain values used are 0.005% and bility for adverse environmental effects for Sic/SiC com- 0.002%. 0 As indicated in Table II. both values were dete ites exists. These stress levels. which are indicated in Fig. 6 n this study and Table Il, are associated with(I)the first AE event, (2) the first loud ae event. which is defined as an ae event with an energy of at least one-tenth the highest energy event not corresponding to final failure of the composite, and ( 3)the effective AE onset stress, which is determined by extrapolation of the steep slope of the normalized cumulative AE energy with increasing stress back down to the zero energy axis(see Fig. 6 For the braided specimens, there was an initial increase in slope (AE activity) followed by a further increase in slope, which con- tinued at the higher rate until failure. For AE onset stress, the Stress-strain curve initial moderate AE slope(Fig. 6) was used because it was con- 0. 005% offset firmed that high-energy AE events were occurring in this regime It is evident in Table II that except for the first AE event, the 0. 002% offset other two stress measures for matrix cracking are significantl higher for the off-axis specimens than for the on-axis specimens. Underlying mechanisms and technical significance for these stress levels will be discussed in the following sections Strain. The effect of architecture and orientation on matrix cracking(or Fig3. Deviation from linearity-stress construction for a braided spec- DFLS) and UTs are key properties that need to be understood imen. Note, only the low strain portion of the stress-strain curve is for applications using MI SiC/SiC composites under multiaxial plotted ress states. In the following sections, the results shown in Figs

The tested specimens were cut and polished along the loading direction in order to measure transverse matrix cracks optically. The polished specimens were plasma etched (CF4 at 500 Watts for 30 min) in order to enhance the matrix cracks in the CVI SiC of the matrix. Matrix cracks were counted over lengths of approximately 10 mm in order to obtain a matrix crack density.7 III. Results The room-temperature tensile stress–strain curves for the off￾axis-oriented SiC/SiC panels A4 and B1 and for the orthogonal￾oriented 2D-woven SiC/SiC panels A1–3 are shown in Fig. 2. Hysteresis loops used for residual stress determination have been eliminated from these curves except for a braided panel B1 specimen, which is given as an example in Fig. 2. The values of initial elastic modulus E, UTS, and residual stress on matrix are listed in Table I and are consistent with other composite panels fabricated with the same fiber types and fiber architec￾tures.4,7 As shown in Fig. 3, the offset strain construct method8 was also used to determine DFLS. It consists of drawing a line with the same slope as the initial elastic modulus, but offset by some amount of positive strain, where the DFLS would be de￾termined by the intersection of that curve with the stress–strain curve. Typical offset-strain values used are 0.005%9 and 0.002%.10 As indicated in Table II, both values were determined in this study. From Table I, it can be seen that in-plane E values are similar for all panels with the exception of the double-tow 5HS-woven composite (panel A3). This suggests that for the SiC/SiC system studied, the in-plane elastic modulus is not strongly dependent on 2D fiber architectures or on tensile-loading direction. For in￾plane UTS, the 2D-woven 0/90 panels aligned in the primary fiber axes are of course the strongest because they were tested parallel to the fiber direction. Nevertheless, the braided panel with over three-quarters of the fibers oriented 231 from the loading-axis displayed a high in-plane UTS; however, the [745] panel displayed a relatively low UTS. But perhaps the most striking mechanical results were the high in-plane DFLS values for both the braided and the [745] panels. These high DFLS results correlate (Fig. 4 and Table II) with higher stress ranges over which high-energy AE events were re￾corded in the off-axis panels. In Fig. 4, the AE activity versus applied CMC stress is plotted as normalized cumulative AE en￾ergy, which is the cumulative AE energy of each AE event up to a given event divided by the total AE energy of all the events. For the MI SiC/SiC system, it has been shown7 that this type of plot represents a good relative distribution of transverse or TTMC. In addition, multiplying the final matrix crack density, measured from polished sections after failure, by the normalized cumulative AE energy is a very good estimate of the actual stress-dependent matrix crack density versus applied stress.7 The measured matrix crack densities at failure are shown in Table II and the estimated matrix crack density with stress in Fig. 5. The difference in the shape of the matrix crack distribution between the double-tow and single-tow 2D-woven composites is hypoth￾esized to be due to the greater concentration and longer lengths of back-to-back 901 minicomposites (see Section IV). Two different matrix crack distributions are evident for the two different off-axis-loaded specimens. The [745] specimen has a very steep curve, whereas the braided specimens were similar in shape to the single-tow 2D composites. Neither of the off-axis composites appears to have reached matrix crack saturation be￾cause a decrease in the rate of AE activity was never achieved at higher stresses. It is also interesting to note that at a composite stress of 200 MPa, the off-axis-tested braided specimens had little or no cracks; however, the double-tow 2D-woven compos￾ites had B5 TTMC/mm. For this study, several AE criteria were used to evaluate key stress levels near initiation of matrix cracking as these values typically represent upper design limits beyond which CMC mod￾uli and axial thermal conductivity irreversibly decrease and the possibility for adverse environmental effects for SiC/SiC com￾posites exists.11 These stress levels, which are indicated in Fig. 6 and Table II, are associated with (1) the first AE event, (2) the first loud AE event, which is defined as an AE event with an energy of at least one-tenth the highest energy event not corresponding to final failure of the composite, and (3) the effective AE onset stress,7 which is determined by extrapolation of the steep slope of the normalized cumulative AE energy with increasing stress back down to the zero energy axis (see Fig. 6). For the braided specimens, there was an initial increase in slope (AE activity) followed by a further increase in slope, which con￾tinued at the higher rate until failure. For AE onset stress, the initial moderate AE slope (Fig. 6) was used because it was con- firmed that high-energy AE events were occurring in this regime. It is evident in Table II that except for the first AE event, the other two stress measures for matrix cracking are significantly higher for the off-axis specimens than for the on-axis specimens. Underlying mechanisms and technical significance for these stress levels will be discussed in the following sections. IV. Analysis and Discussion The effect of architecture and orientation on matrix cracking (or DFLS) and UTS are key properties that need to be understood for applications using MI SiC/SiC composites under multiaxial stress states. In the following sections, the results shown in Figs. Fig. 2. Stress–strain curves of off-axis and orthogonally aligned com￾posites with the hysteresis loop removed for clarity. An example of the hysteresis loops for 0/767 braid (1) is shown in the inset. 0 50 100 150 200 250 0 0.02 0.04 0.06 0.08 0.1 Strain, % Stress, MPa Stress-strain curve 0.002% offset 0.005% offset Fig. 3. Deviation from linearity-stress construction for a braided spec￾imen. Note, only the low strain portion of the stress–strain curve is plotted. October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3187

3188 Journal of the American Ceramic Society-Morscher et al. VoL. 90. No. 10 Table Il. Matrix Cracking Properties of Composite Specimens DFLS,0002% DFLS.0.005% First AE First loud ae AE onset before AE onset stress- Final matrix crack Panel offset(MPa) offset(MPa stress(MPa stress(MPa) stress(MPa otal #f loud events density(#/mm) Orthogonal oriented composites 147 174 190 4-141 l30 173 182 1-141 135 157 2-134 9.0 Off-axis oriented composites 10 4.0 4.9 195 231 193 210 3-134 7. 9 epcm [±45 7.9 epcm [/9o 8.7 epcm 0/901 8.7 epcm double-tot [0/90] 3.95 epcm Double Tow [0/901 p 2 [O/+67] braid 8.7 epcm, 00 Fig 4. Acoustic emission activity versus composite stress Fig. 5. Estimated matrix crack density with stress based on acoustic 2 and 4 and in Table I and Il will be mechanistically analyzed and compared with other data in the literature to better ur dividual 90" minicomposites and the size of two 90% minicom- tand their scientific and technical significan osites that happen to be adjacent to one another. Whenever this back-to-back tow circumstance exists. the effective width of an unbridged tunnel crack would be approximately twice the (1) Matrix Cracking in 0/90 2D-Woren Composites Tested in back-to-back individual 90 minicomposites is much more com- osite. This characteristic of the o direction non in the double-tow woven 3.95 epcm composite panel com- A variety of microscopic studies of [o/90]2D-woven CMC spec- pared with single-tow 7.9 epcm woven panels. Not only are there mens have shown that at lower stresses. initial transverse matrix nore regions of multiple 90@ minicomposites but also the length cracks are usually either"tunnel"microcracks, which occur in of these regions(distance the tow is woven over four tows before the 90 minicomposites oriented perpendicular to the 00loadir axis, and or nonsteac ady-state microcracks that are partially it is woven under the fifth tow in the five-harness satin archi- bridged due to sufficient fiber traction in the matrix crack tecture)would be approximately twice the length of single-tow to stop matrix crack propagation through-thickness. At woven composites for the five- harness satin weave because the higher stresses, these microcracks propagate through-thickness epcm of the double-tow CMC was one-half that of the single or link up with other microcracks to form TTMC over a range tow CMC. As a result, as shown in Fig 4, the double-tow woven AE methodologies have been successfully used not only uantify but also to understand and model the occurrence and tress-strain dependence of microcrack and TTMc behavior. Initial low-energy events generally correspond to tunnel micro crack formation in 90 minicomposites perpendicular to the loading direction. High-energy events, those in the upper loga- rithmic decade of energy, correspond to either large microcracks 1st Loud AE event raid &[0/90] tribution for TTMC is controlled by(I) the size distribution of 1st AE even 90 minicomposites perpendicular to the load-bearing 0 mini 5 0.31 Braid &(0r-9o) composites and (2)the in situ stress in the region of the com- 0 posite outside the load-bearing 0 minicomposite, i.e., the portion of the composite composed of 90 minicomposites and the MI matrix. For panels fabricated from the random lay-up of standard single-tow-woven fabric plies, the size distribution of 90 minicomposites can crudely vary between the size of Stress. MPa ' For the CMC of this study, a minicom of a single multi fiber tow, the CvI Fig. 6. Acoustic emission(AE) events and onset stress determination. The arrows below the x-axis indicate ae onset stress

2 and 4 and in Table I and II will be mechanistically analyzed and compared with other data in the literature to better under￾stand their scientific and technical significance. (1) Matrix Cracking in 0/90 2D-Woven Composites Tested in the 01 Direction A variety of microscopic studies of [0/90] 2D-woven CMC spec￾imens have shown that at lower stresses, initial transverse matrix cracks are usually either ‘‘tunnel’’ microcracks,12 which occur in the 901 minicompositesy oriented perpendicular to the 01 loading axis, and/or nonsteady-state microcracks that are partially bridged due to sufficient fiber traction in the matrix crack wake to stop matrix crack propagation through-thickness. At higher stresses, these microcracks propagate through-thickness or link up with other microcracks to form TTMC over a range of stress levels. AE methodologies have been successfully used not only to quantify but also to understand and model the occurrence and stress–strain dependence of microcrack and TTMC behavior. Initial low-energy events generally correspond to tunnel micro￾crack formation in 901 minicomposites perpendicular to the loading direction. High-energy events, those in the upper loga￾rithmic decade of energy, correspond to either large microcracks and/or TTMC.13 For 2D-woven 0/90 MI SiC/SiC panels tested in the 01 direction, it has been demonstrated that the stress dis￾tribution for TTMC is controlled by (1) the size distribution of 901 minicomposites perpendicular to the load-bearing 01 mini￾composites and (2) the in situ stress in the region of the com￾posite outside the load-bearing 01 minicomposite, i.e., the portion of the composite composed of 901 minicomposites and the MI matrix.7,14 For panels fabricated from the random lay-up of standard single-tow-woven fabric plies, the size distribution of 901 minicomposites can crudely vary between the size of in￾dividual 901 minicomposites and the size of two 901 minicom￾posites that happen to be adjacent to one another. Whenever this back-to-back tow circumstance exists, the effective width of an unbridged tunnel crack would be approximately twice the crack width of a single minicomposite. This characteristic of back-to-back individual 901 minicomposites is much more com￾mon in the double-tow woven 3.95 epcm composite panel com￾pared with single-tow 7.9 epcm woven panels. Not only are there more regions of multiple 901minicomposites but also the length of these regions (distance the tow is woven over four tows before it is woven under the fifth tow in the five-harness satin archi￾tecture) would be approximately twice the length of single-tow woven composites for the five-harness satin weave because the epcm of the double-tow CMC was one-half that of the single￾tow CMC. As a result, as shown in Fig. 4, the double-tow woven Table II. Matrix Cracking Properties of Composite Specimens Panel DFLS, 0.002% offset (MPa) DFLS, 0.005% offset (MPa) First AE stress (MPa) First loud AE stress (MPa) AE onset stress (MPa) No. of loud events before AE onset stress— total # loud events Final matrix crack density (#/mm) Orthogonal oriented composites A1 147 174 132 170 190 4–141 — A2 130 173 100 159 182 1–141 10.3 A3 135 176 128 138 157 2–134 9.0 Off-axis oriented composites A4 210 225 56 197 220 1–65 4.0 B1 232 259 83 187 215 2–95 4.9 195 231 135 193 210 3–134 — AE, acoustic emmision. 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 0 100 200 300 400 500 Stress, MPa Norm Cum AE Energy 7.9 epcm [0/90] [0/+67] braid 8.7 epcm, [+45] 3.95 epcm double-tow [0/90] 8.7 epcm [0/90] Fig. 4. Acoustic emission activity versus composite stress. 0 2 4 6 8 10 12 0 100 200 300 400 500 Composite Stress, MPa Estimated Crack Density, #/mm 7.9 epcm; [0/90] 3.95 epcm Double Tow [0/90] 8.7 epcm [0/90] [0/+67] braid 8.7 epcm, [+45] Fig. 5. Estimated matrix crack density with stress based on acoustic emission. 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 0 100 200 300 400 500 Stress, MPa Norm Cum AE Energy 8.7 epcm [0/90] 1st AE event Braid & [0/90] 1st Loud AE event Braid & [0/90] [0/+67] braid Fig. 6. Acoustic emission (AE) events and onset stress determination. The arrows below the x-axis indicate AE onset stress. y For the CMC of this study, a minicomposite consists of a single multi fiber tow, the CVI BN interphase coating, and the initial CVI SiC matrix coating associated with tow. 3188 Journal of the American Ceramic Society—Morscher et al. Vol. 90, No. 10

October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3189 orthogonal this study D[o+67] double tow braid (hoop direction)-this study Fiber Fraction in Loading Direction emission(AE) matrix cracking(onset)stress for two- osites with fibers tows oriented perpendicular to the loading direction. panel displays the lowest onset stress and the smallest stress dis- tribution for TTMC in the on-axis panels In an effort to model architectural effects on the onset of TTMC in 2D-woven 0/90 SiC/SiC composites tested in the 0o direction, a previous study has examined the important micro- tructural factors affecting the internal stresses on the 90 mini composites since these appear to be the most critical flaws within the matrix that eventually cause TTMC. Two important facto 1 mm were identified: (1)a built-in residual stress on the matrix and (2) the applied composite stress as modified by the load shared by the 0%minicomposites. As such, one can estimate the stress on the matrix region containing the 90 minicomposites(=mini- matrix) by the following simple rule of mixtures relationship (oc +Oth)Ec-fmini em ominimatrix 1) Fig 8. Micrograph of a two-dimensional Ec posite after room temperature tensile failure urce.Most notably, there was also a very na Here,σ is the composite stress, Oth is the residual tribution for TTMC(Fig. 4), which would corres ess in the matrix Ec is the measured composite ela population(high Weibull modulus)of flaws and tic modulus from the oe curve (Table I): mini is the volume that propagate through-thickness at and slightly above the ma ically -2 times the fiber content in the 0 direction for the com-(panel A4) displayed a higher onset stress for TTMC than its can be reduced by increasing a compressive residual stress on the cracking stress and the narrow stress-distribution for matrix matrix and/or by increasing mini. As shown in Table l, one ben acking for this system is primarily due to a lack of minicom- efit of the mi sic/SiC system is a compressive residual stress on posites perpendicular to the direction of stress so that higher the as-fabricated matrix, but this stress did not vary much with tresses are required for tunnel crack propagation at an ang the 2D architectures of this study. The exact source of this stress 45 to the loading axis. There is also the possibility that the is not completely understood, but probably can urce of matrix cracking are surface flaws or pores in the ma silicon content in the matrix and to the composite fabrication trix and not the +45-oriented minicomposites Matrix cracks conditions. On the other hand, fiber content in the or primary were observed to be perpendicular to the loading direction as MI SiC/SiC system, which in turn should increase the composite stress for the onset of ttmc. that this latter mechanism can indeed be utilized is seen in Fig. 7, which plots as open circles the Ae onset cracking strength as a function of primary for the (3) Matrix Cracking in 01+67 2D-Braided Composite Tested 0o-loaded 0/90 MI SiC/SiC composites of this study and for a in the Hoop Direction For the triaxially braided specimen, approximately 77% of the fibers or a total fiber fraction of 0. 26 were oriented 23 from the loading axis. This fraction is effectively higher than that of the (2) Matrix Cracking in 0/90 2D-Woren Composites Tested 2D-woven composites loaded in the primary fiber direction, a rhe45° Direction thus appears to be one cause for the higher matrix cracking The stress-distribution for matrix cracking in the [+45]off-axis stress for the braided composite. As described above for the 0/90 panel A4 differs considerably from those of the [o/90] compos- composites, matrix cracks emanate from the region outside the ites tested in a 0 or fiber direction first ae events show that load-bearing minicomposites, e.g., from the minicomposites ori some small microcracks were formed at low stresses in the [t45 ented 90 to the loading axis. The same type of analysis can be composite (Table ID), the source of which has not been deter- applied here to the braided composites to determine whether the mined. There was considerable porosity in the matrix region stress ranges in the TTMC flaw-source regions(axial tow mini- tween the outer two plies ( Fig 8), which could have been one composites, see Fig. 1)of the matrix are similar to standard

panel displays the lowest onset stress and the smallest stress dis￾tribution for TTMC in the on-axis panels. In an effort to model architectural effects on the onset of TTMC in 2D-woven 0/90 SiC/SiC composites tested in the 01 direction, a previous study7 has examined the important micro￾structural factors affecting the internal stresses on the 901 mini￾composites since these appear to be the most critical flaws within the matrix that eventually cause TTMC. Two important factors were identified: (1) a built-in residual stress on the matrix and (2) the applied composite stress as modified by the load shared by the 01 minicomposites. As such, one can estimate the stress on the matrix region containing the 90 minicomposites ( 5 mini￾matrix) by the following simple rule of mixtures relationship7 : sminimatrix ¼ ðsc þ sthÞ Ec Ec fminiEmini 1 fmini
(1) Here, sc is the applied composite stress, sth is the residual stress in the matrix (Table I); Ec is the measured composite elas￾tic modulus from the se curve (Table I); fmini is the volume fraction of the 01 minicomposites in the loading direction (typ￾ically B2 times the fiber content in the 01 direction for the com￾posites of this study); and Emini is the effective modulus of these 01 minicomposites. Thus, the stress on the 90 minicomposites can be reduced by increasing a compressive residual stress on the matrix and/or by increasing fmini. As shown in Table I, one ben￾efit of the MI SiC/SiC system is a compressive residual stress on the as-fabricated matrix, but this stress did not vary much with the 2D architectures of this study. The exact source of this stress is not completely understood, but probably can be related to the silicon content in the matrix and to the composite fabrication conditions. On the other hand, fiber content in the 01 or primary fiber direction, fprimary, can be increased to a large degree for the MI SiC/SiC system, which in turn should increase the composite stress for the onset of TTMC. That this latter mechanism can indeed be utilized is seen in Fig. 7, which plots as open circles the AE onset cracking strength as a function of fprimary for the 01-loaded 0/90 MI SiC/SiC composites of this study and for a previous study.7 (2) Matrix Cracking in 0/90 2D-Woven Composites Tested in the 451 Direction The stress-distribution for matrix cracking in the [745] off-axis panel A4 differs considerably from those of the [0/90] compos￾ites tested in a 01 or fiber direction. First AE events show that some small microcracks were formed at low stresses in the [745] composite (Table II), the source of which has not been deter￾mined. There was considerable porosity in the matrix region between the outer two plies (Fig. 8), which could have been one source. Most notably, there was also a very narrow stress-dis￾tribution for TTMC (Fig. 4), which would correspond to a large population (high Weibull modulus) of flaws and/or microcracks that propagate through-thickness at and slightly above the ma￾trix cracking stress. The important fact is that this composite (panel A4) displayed a higher onset stress for TTMC than its theoretically equivalent 0/90 composite (panel A1) loaded along its primary fiber axis. It is suggested that the higher matrix cracking stress and the narrow stress-distribution for matrix cracking for this system is primarily due to a lack of minicom￾posites perpendicular to the direction of stress so that higher stresses are required for tunnel crack propagation at an angle 451 to the loading axis. There is also the possibility that the source of matrix cracking are surface flaws or pores in the ma￾trix and not the 7451-oriented minicomposites. Matrix cracks were observed to be perpendicular to the loading direction as has been reported before.2 (3) Matrix Cracking in 0/767 2D-Braided Composite Tested in the Hoop Direction For the triaxially braided specimen, approximately 77% of the fibers, or a total fiber fraction of 0.26, were oriented 231 from the loading axis. This fraction is effectively higher than that of the 2D-woven composites loaded in the primary fiber direction, and thus appears to be one cause for the higher matrix cracking stress for the braided composite. As described above for the 0/90 composites, matrix cracks emanate from the region outside the load-bearing minicomposites, e.g., from the minicomposites ori￾ented 901 to the loading axis. The same type of analysis can be applied here to the braided composites to determine whether the stress ranges in the TTMC flaw-source regions (axial tow mini￾composites, see Fig. 1) of the matrix are similar to standard 100 120 140 160 180 200 220 240 0.1 0.15 0.2 0.25 0.3 Fiber Fraction in Loading Direction AE Matrix Cracking Stress, MPa 2D orthogonal [7] 2D orthogonal - this study 2D double tow orthogonal [7] 2D double tow orthogonal this study [0/+67] double tow braid (hoop direction) - this study Fig. 7. Effect of fiber fraction in the loading direction on acoustic emission (AE) matrix cracking (onset) stress for two-dimensional com￾posites with fibers tows oriented perpendicular to the loading direction. Fig. 8. Micrograph of a two-dimensional woven [745]-oriented com￾posite after room temperature tensile failure. October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3189

3190 Journal of the American Ceramic Society-Morscher et al. VoL. 90. No. 10 2D-woven 0/90 composites oriented in one of the orthogonal directions Assuming the applicability of Eq (1), there is a question as to model for [0/90] single-tow 2D woven data(ref. 14) how to approximate Emini because the fibers are oriented at an angle to the loading axis. One limit would be to assume that the ould act as if they were in parallel, si 3.95 epcm [o/9o there was the same fraction of minicomposites +23 as there are 8 uble-tow woven 7.9 epcm[o/90]single- 23 from the tensile axis The other extreme would be to as- sume that the modulus of a minicomposite at an angle to the 3 loading axis would behave similar to a unidirectional ply ori- 8.7 epcm [o/9o] single- ented at an angle and could therefore be estimated from the properties of an anisotropic lamina cos"(e)( 1-22)si2()cs() Braid Upper Bound +-sin4(0)(2) where subscript I refers to the 0 orientation of a minicomposite, subscript 2 refers to the transverse direction or to the 90orien- tation of a minicomposite, and 0 refers to the orientation of the 0 loading direction off the 0 fiber axis. G1 is the bulk modulu 200 and v12 is Poissons ratio. Minimatrix stress, MPa To determine the effect of orientation, ea was determined for the 23 fiber orientation case relative to the value of E,(360 GPa Fig 9. Estimated matrix crack density versus minimatrix stress for the braided composite ) E, is unknown for these minicom- posites, but a conservative estimate would be 100 GPa, which portant to discuss the response of the different types of CMCs to would be on the low -end of elastic moduli estimated for 90o off-axis loading CMC mechanical behavior has been defined as matrix dom- minicomposites in CVI SiC matrix composites. V12 was as- inated(Class ID)and fiber dominated( Class lID). 2.17 For matrix- umed to be 0. 15 and G1 was estimated from the simple iso tropic relationship E1/2(1+v12). Based on these simple dominated composites(.g, SiC fiber-reinforced glass ceramic ssumptions Ee/E1=0.943, a minor effect. Figure 9 shows the or Sic matrix composites), initial loads are shared by an un- timated matrix crack density versus minimatrix stress(Eq (D)) cracked matrix and by the fibers. Therefore, elastic properties for both the cases in comparison with the 2D composites. The lower bound"assumes that the minicomposites behave in the same way as parallel minicomposites and Emini is determined from the rule of mixtures based on the fractional content of fi- ber, BN, and CVI SiC. The"upper bound"refers to a reduced Emini from Eq. (2). There is only a small difference in minimatrix ess for the two extremes and they compare well with the composites loaded in the orthogonal direction. From this i would appear that matrix cracking in the braided composites is controlled by the stress acting on the axial minicomposites ori- ented perpendicular to the loading axis in the same manner a the 2D-woven 0/90 composites with single tows. However, in order to increase fiber volume fraction, the 0+67 braided com- sites were braided with two(double) tows woven together for both the axial and bias fibers. But there were a smaller fraction of minicomposites oriented perpendicular to the loading axis for the braided composites compared with the 2D-woven 0/90 com- osites. Also, the minicomposites oriented perpendicular to the loading direction of the braided composites do not contact one another(Fig. 10)and do not form back-to-back minicomposites as is the case for the 2D-woven 0/90 double-tow composites These two factors are probably the reason for the cracking be- havior of the braided composite being more in line with that of tow woven composites and not the double-tow woven composites. Figure 7 compares the effect of effective fiber frac- tion bridges a transverse matrix crack in the loadi tion on AE onset stress for all the MI SiC/SiC composites studied with 2D-woven and braided architectures, which possess minicomposites oriented perpendicular to the loading direction i.e., the likely source for the initiation of matrix cracks. (4) Ultimate Strength for Off-Axis Panels The ultimate strengths of the panels loaded off-axis were gen- erally less than those for panels tested in a primary fiber direc tion. This result is to be expected since the fibers were not aligned in the direction of applied stress and are subject to local bending and shear within a matrix crack. At this point, it is im- The fractional content of fiber BN. and CvI SiC We ere very similar to the data already re- ported in Morscher'fmini for the bias fibers of the braided composites was 0.54 compared Fig 10. Micrograph of a 0/+67 braid composite after roor with 0.36 for the orthogonal composit temperature tensile failur

2D-woven 0/90 composites oriented in one of the orthogonal directions. Assuming the applicability of Eq. (1), there is a question as to how to approximate Emini because the fibers are oriented at an angle to the loading axis. One limit would be to assume that the bias minicomposites would act as if they were in parallel, since there was the same fraction of minicomposites 1231 as there are 231 from the tensile axis. The other extreme would be to as￾sume that the modulus of a minicomposite at an angle to the loading axis would behave similar to a unidirectional ply ori￾ented at an angle, and could therefore be estimated from the properties of an anisotropic lamina15 1 Ey ¼ 1 E1 cos4 ðyÞ 1 G12 2n12 E1
sin2 ðyÞ cos2 ðyÞ þ 1 E2 sin4 ðyÞ (2) where subscript 1 refers to the 01 orientation of a minicomposite, subscript 2 refers to the transverse direction or to the 901orien￾tation of a minicomposite, and y refers to the orientation of the loading direction off the 01 fiber axis. G12 is the bulk modulus and n12 is Poissons ratio. To determine the effect of orientation, Ey was determined for the 231 fiber orientation case relative to the value of E1 (360 GPa for the braided compositez ). E2 is unknown for these minicom￾posites, but a conservative estimate would be 100 GPa, which would be on the low-end of elastic moduli estimated for 901 minicomposites in CVI SiC matrix composites.16 n12 was as￾sumed to be 0.15 and G12 was estimated from the simple iso￾tropic relationship E1/2 (11n12). Based on these simple assumptions Ey/E1 5 0.943, a minor effect. Figure 9 shows the estimated matrix crack density versus minimatrix stress (Eq. (1)) for both the cases in comparison with the 2D composites. The ‘‘lower bound’’ assumes that the minicomposites behave in the same way as parallel minicomposites and Emini is determined from the rule of mixtures based on the fractional content of fi- ber, BN, and CVI SiC. The ‘‘upper bound’’ refers to a reduced Emini from Eq. (2). There is only a small difference in minimatrix stress for the two extremes and they compare well with the composites loaded in the orthogonal direction. From this it would appear that matrix cracking in the braided composites is controlled by the stress acting on the axial minicomposites ori￾ented perpendicular to the loading axis in the same manner as the 2D-woven 0/90 composites with single tows. However, in order to increase fiber volume fraction, the 0767 braided com￾posites were braided with two (double) tows woven together for both the axial and bias fibers. But there were a smaller fraction of minicomposites oriented perpendicular to the loading axis for the braided composites compared with the 2D-woven 0/90 com￾posites. Also, the minicomposites oriented perpendicular to the loading direction of the braided composites do not contact one another (Fig. 10) and do not form back-to-back minicomposites as is the case for the 2D-woven 0/90 double-tow composites.7 These two factors are probably the reason for the cracking be￾havior of the braided composite being more in line with that of the single-tow woven composites and not the double-tow woven composites. Figure 7 compares the effect of effective fiber frac￾tion that bridges a transverse matrix crack in the loading direc￾tion on AE onset stress for all the MI SiC/SiC composites studied with 2D-woven and braided architectures, which possess minicomposites oriented perpendicular to the loading direction, i.e., the likely source for the initiation of matrix cracks. (4) Ultimate Strength for Off-Axis Panels The ultimate strengths of the panels loaded off-axis were gen￾erally less than those for panels tested in a primary fiber direc￾tion. This result is to be expected since the fibers were not aligned in the direction of applied stress and are subject to local bending and shear within a matrix crack. At this point, it is im￾portant to discuss the response of the different types of CMCs to off-axis loading. CMC mechanical behavior has been defined as matrix dom￾inated (Class II) and fiber dominated (Class III).2,17 For matrix￾dominated composites (e.g., SiC fiber-reinforced glass ceramic or SiC matrix composites), initial loads are shared by an un￾cracked matrix and by the fibers. Therefore, elastic properties 0 2 4 6 8 10 12 0 50 100 150 200 250 300 Minimatrix stress, MPa Estimated Crack Density, #/mm model for [0/90] single-tow 2D woven data (ref. 14) 7.9 epcm [0/90] single￾tow woven (this study) 8.7 epcm [0/90] single￾tow woven (this study) 3.95 epcm [0/90] double-tow woven - this study - ref. 7 Braid Lower Bound Braid Upper Bound Fig. 9. Estimated matrix crack density versus minimatrix stress. Fig. 10. Micrograph of a [0/767] braid composite after room temperature tensile failure. z The fractional content of fiber, BN, and CVI SiC was 0.32, 0.07, and 0.25, respectively. The fractional contents of the other composites were very similar to the data already re￾ported in Morscher7 fmini for the bias fibers of the braided composites was 0.54 compared with B0.36 for the orthogonal composites. 3190 Journal of the American Ceramic Society—Morscher et al. Vol. 90, No. 10

October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3191 Table Ill. Ultimate Strength Properties of 10/901 Composite It is difficult to compare all of the off-axis UTs data on an in the Literature Tested in Orthogonal and 45/45 Orientations absolute basis since composites vary considerably in fiber vol- ume fractions, in situ strength properties of the fibers, interfacial 0)UTS(0)E45)UTs(45) architecture f=2/o(GPa)(MPa)(GPa)(GPa) on fiber modulus and diameter. However, when comparing the Matrix-dominated composites UTS of [0/90] composites oriented in the [+45]-direction, the Sylramic-MI- 5 HS 0.34268344258200 matrix-dominated SiC composites, whether reinforced by 5 HS 0.34210351169158 Nicalon, Hi-Nicalon, Sylra or Sylramic-ibN SiC fibers NicCⅥISiC2 Weave~0.4265255NA210 (see Tables I and II), exhibited the best off-axis UTS values Nic-CVI SIC 0.36240248183192 (158-242 MPa) compared with the fiber-dominated composites NIC-CAS 0/90 laminate 0.4 136 215 NA (all< 100 MPa) regardless of the fiber type. To estimate the general effect of composite-type and off-axis testing direction on effective fiber strength or load-carrying abil Fiber-dominated composites ity, one can eliminate the fiber tensile strength and volume 0.460320NA80 fraction effects from the strength data by normalizing the 8 HS 0.36 70 252 34 76 off-axis UTS data of Tables I, Ill, and IV using the following C-CVI SiC2 8 HS 0.36873246274 relationship 8 HS 0.361013052170 ALO3 /Mullite-8HS ~04~972105052 Normalized effective fiber strength MI, melt-infiltrated: UTS. ultimate tensile strength. Sr(off-axis)/Sr(o)=(cult off-axis/out o)(o/Jfofr-a 3) where Sr(off-axis) is the average strength of a fiber at failure that are heavily influenced by the elastic properties of the matrix. is oriented at an angle to the loading direction and Sno is the Nonlinearity in the stress-strain curve is caused by transverse average strength of a fiber at failure when oriented in the direc- matrix cracks(cracks perpendicular to the applied load) due te tion of loading. This can be determined by dividing the UTS of terface debonding and fiber sliding resulting in fiber pullout site by the fraction of fibers. ult o and fo refers, re- within the matrix crack and matrix crack opening. Transverse spectively, to the UTS and fraction of fibers when the fibers are matrix cracking occurs for both on-axis and off-axis loading of aligned with the loading direction, i.e., for a [o/90] composite fo half the total volume fraction of fibers g and off-ax For fiber-dominated composites(e.g, porous oxide/oxide refers, respectively, to the UTS and effective load-bearing fra C/C, or C/SiC composites), the matrix carries little load due tion of fibers when the fibers are aligned at an angle to the to the fact that the matrix is porous or is heavily microcracked in loading direction. It can be assumed for the [0/90) composites the as-processed condition, and the mechanical properties are hen loaded symmetric to the primary fiber directions, e.g controlled by the fiber(architecture)response to loading. When [+45] that both fiber axes are carrying load, so that fofr-axis=2 loaded on-axis, elastic moduli are typically only slightly larger fo. For nonsymmetric loading, it can be assumed that the fibers han f(0)x Ef, and there is only minor nonlinearity to the stress in the axis at the greatest angle to the loading direction will strain curve since the matrix is highly porous and/ or a high fracture first, debond, and or pull apart, leaving the other pri damage condition already exists in the matrix. When loaded in ary axis fibers to control UTS, so that fff-axis equals fo of the the off-axis. there is significant nonlinear emaining fibers. For braided composites, specimens tested in behavior beginning near zero stress due to the absence of he hoop direction were compared with specimens allel fibers in the loading direction. More damage is created in same panel tested in the axial direction(C/epoxy- the composite due to shear band formation, fiber-rearran or to [0/90 specimens tested in the 0 direction(Syl-i ment, inter-ply delamination, and fiber"scissoring. This results Table in high ultimate strains as long as the fibers can withstand the Based on these assumptions and Eq ( 3), Fig. ll shows the tensile, bending, and shear forces, but low ultimate strengths due normalized fiber strength results as a function of test angle usi to these combined forces on the fibers. This low off-axis Uts the UtS data from Tables l, Ill, and IV. Again the four matrix- behavior for these fiber-dominated composites may be a limita- dominated Sic-based matrix composites where the elastic mod- tion for certain applications. ulus of the matrix was relatively high exhibit the best relative The Sylramic-iBN composites are an excellent example of retention of fiber strength. For the fiber-dominated or low-mod- matrix-dominated composites and of how off-axis properties ulus matrix composites, C/epoxy, Al2O3/ mullite, Nic/C, C/C can be maximized which may be tant for future applica and C/SiC, there was excellent correlation at an angle of 45, but tions. To demonstrate this, the uts data from the composite poorer off-axis strengths than the matrix-dominated compo tested in this study (see table i) can be con with UTS data ites. The two matrix-dominated glass-CMC were in between the in the literature for a variety of fiber-dominated and matrix- fiber-dominated and Sic-based matrix-dominated composites dominated 2D-woven, braided, and cross-plied ceramic and probably due to the lower modulus of the glass-ceramic matrix polymer composites that were tested in a primary fiber direc- compared with CVI and MI SiC matrices. It is significant that tion as well as in an off-axis direction. Descriptions of these the braided and [0/90] C/epoxy data from different studies over composites and their ultimate properties are listed in Table Ill the range of angles tested correlate with one another, justifying for 0/90 composites tested in the 0 and 45/45 direction and in some of the crude assumptions made above Tables IV for 0/90 and braided composites tested at multiple For fiber orientations 25-45 off the loading axis. the relative ngles. These composites included 2D plain-woven carbon fiber stress-carrying ability of the Syl-iBN fibers in high-modulus SiC- reinforced epoxy, 2D triaxially braided carbon fiber-reinforced based matrices ranged from 50% to 30% of the fiber strength epoxy, 2D eight-harness satin-woven Al2Ox-fiber-reinforced respectively, when oriented in the primary 0 direction. Though porous oxide matrix (mullite) composites, CG Nicalon the relative stress-carrying ability decreased with increasing an Nippon Carbon, Tokyo, Japan, referred to as"Nic")fiber- gle, it was still far superior to the fiber-dominated systems(50% einforced carbon matrix. - CG Nicalon fiber-reinfe greater relative stress at 25 and 100% greater relative stress at matrix composites, CG Nicalon-fiber reinforced matrix c posites, an carbon fiber-reinforced carbo Note that in Eq. (3). the ellipsoid character of off-axis oriented fibers in the pla composites (Nippon Carbon, Japan; referred to as"HN)an fiber-reinforced MI composites

are heavily influenced by the elastic properties of the matrix. Nonlinearity in the stress–strain curve is caused by transverse matrix cracks (cracks perpendicular to the applied load) due to interface debonding and fiber sliding resulting in fiber pullout within the matrix crack and matrix crack opening. Transverse matrix cracking occurs for both on-axis and off-axis loading of composites. For fiber-dominated composites (e.g., porous oxide/oxide, C/C, or C/SiC composites), the matrix carries little load due to the fact that the matrix is porous or is heavily microcracked in the as-processed condition, and the mechanical properties are controlled by the fiber (architecture) response to loading. When loaded on-axis, elastic moduli are typically only slightly larger than f(01)
Ef, and there is only minor nonlinearity to the stress strain curve since the matrix is highly porous and/or a high damage condition already exists in the matrix. When loaded in the off-axis, there is significant nonlinearity in the stress–strain behavior beginning near zero stress due to the absence of par￾allel fibers in the loading direction. More damage is created in the composite due to shear band formation, fiber-rearrange￾ment, inter-ply delamination, and fiber ‘‘scissoring.’’ This results in high ultimate strains as long as the fibers can withstand the tensile, bending, and shear forces, but low ultimate strengths due to these combined forces on the fibers. This low off-axis UTS behavior for these fiber-dominated composites may be a limita￾tion for certain applications. The Sylramic-iBN composites are an excellent example of matrix-dominated composites and of how off-axis properties can be maximized, which may be important for future applica￾tions. To demonstrate this, the UTS data from the composites tested in this study (see Table I) can be compared with UTS data in the literature for a variety of fiber-dominated and matrix￾dominated 2D-woven, braided, and cross-plied ceramic and polymer composites that were tested in a primary fiber direc￾tion as well as in an off-axis direction. Descriptions of these composites and their ultimate properties are listed in Table III for 0/90 composites tested in the 0 and 45/45 direction and in Tables IV for 0/90 and braided composites tested at multiple angles. These composites included 2D plain-woven carbon fiber￾reinforced epoxy,18 2D triaxially braided carbon fiber-reinforced epoxy,19 2D eight-harness satin-woven Al2O3-fiber-reinforced porous oxide matrix (mullite) composites,20 CG Nicalon (Nippon Carbon, Tokyo, Japan, referred to as ‘‘Nic’’) fiber￾reinforced carbon matrix,2,21 CG Nicalon fiber-reinforced glass matrix composites,2,22 CG Nicalon-fiber reinforced CVI SiC matrix composites,2,21,23 carbon fiber-reinforced carbon matrix composites,17,21 and 2D five-harness satin-woven Hi-Nicalon (Nippon Carbon, Japan; referred to as ‘‘HN’’) and Sylramic fiber-reinforced MI composites.24 It is difficult to compare all of the off-axis UTS data on an absolute basis since composites vary considerably in fiber vol￾ume fractions, in situ strength properties of the fibers, interfacial sliding stress, and fiber-bending stiffness, which are dependent on fiber modulus and diameter. However, when comparing the UTS of [0/90] composites oriented in the [745]-direction, the matrix-dominated SiC composites, whether reinforced by Nicalon, Hi-Nicalon, Sylramic, or Sylramic-iBN SiC fibers (see Tables I and II), exhibited the best off-axis UTS values (158–242 MPa) compared with the fiber-dominated composites (allo100 MPa) regardless of the fiber type. To estimate the general effect of composite-type and off-axis testing direction on effective fiber strength or load-carrying abil￾ity, one can eliminate the fiber tensile strength and volume fraction effects from the strength data by normalizing the off-axis UTS data of Tables I, III, and IV using the following relationshipJ : Normalized effective fiber strength ¼ Sfðoff-axisÞ=Sfð0Þ ¼ ðsult off-axis=sult 0Þðf0=f off-axisÞ (3) where Sf(offaxis) is the average strength of a fiber at failure that is oriented at an angle to the loading direction and Sf(0) is the average strength of a fiber at failure when oriented in the direc￾tion of loading. This can be determined by dividing the UTS of the composite by the fraction of fibers. sult 0 and f0 refers, re￾spectively, to the UTS and fraction of fibers when the fibers are aligned with the loading direction, i.e., for a [0/90] composite f0 is half the total volume fraction of fibers. sult offaxis and f
offaxis refers, respectively, to the UTS and effective load-bearing frac￾tion of fibers when the fibers are aligned at an angle to the loading direction. It can be assumed for the [0/90] composites when loaded symmetric to the primary fiber directions, e.g., [745], that both fiber axes are carrying load, so that foffaxis
5 2 f0. For nonsymmetric loading, it can be assumed that the fibers in the axis at the greatest angle to the loading direction will fracture first, debond, and/or pull apart, leaving the other pri￾mary axis fibers to control UTS, so that foffaxis
equals f0 of the remaining fibers. For braided composites, specimens tested in the hoop direction were compared with specimens from the same panel tested in the axial direction (C/epoxy—Table IV) or to [0/90] specimens tested in the 0 direction (Syl-iBN MI— Table I). Based on these assumptions and Eq. (3), Fig. 11 shows the normalized fiber strength results as a function of test angle using the UTS data from Tables I, III, and IV. Again the four matrix￾dominated SiC-based matrix composites where the elastic mod￾ulus of the matrix was relatively high exhibit the best relative retention of fiber strength. For the fiber-dominated or low-mod￾ulus matrix composites, C/epoxy, Al2O3/mullite, Nic/C, C/C, and C/SiC, there was excellent correlation at an angle of 451, but poorer off-axis strengths than the matrix-dominated compos￾ites. The two matrix-dominated glass-CMC were in between the fiber-dominated and SiC-based matrix-dominated composites probably due to the lower modulus of the glass–ceramic matrix compared with CVI and MI SiC matrices. It is significant that the braided and [0/90] C/epoxy data from different studies over the range of angles tested correlate with one another, justifying some of the crude assumptions made above. For fiber orientations 251–451 off the loading axis, the relative stress-carrying ability of the Syl-iBN fibers in high-modulus SiC￾based matrices ranged from 50% to 30% of the fiber strength, respectively, when oriented in the primary 01 direction. Though the relative stress-carrying ability decreased with increasing an￾gle, it was still far superior to the fiber-dominated systems (50% greater relative stress at 251 and 100% greater relative stress at Table III. Ultimate Strength Properties of [0/90] Composites in the Literature Tested in Orthogonal and 45/45 Orientations Composite Fiber architecture f 5 2f0 E(0) (GPa) UTS(0) (MPa) E(45) (GPa) UTS(45) (GPa) Matrix-dominated composites Sylramic-MI24 5 HS 0.34 268 344 258 200 Hi-Nic-MI24 5 HS 0.34 210 351 169 158 Nic-CVI SiC2 Plain Weave B0.4 265 255 NA 210 Nic-CVI SiC21 8 HS 0.36 240 248 183 192 NiC-CAS2 0/90 laminate B0.4 136 215 NA 95 Fiber-dominated composites Nic-C2 — B0.4 60 320 NA 80 Nic-C21 8 HS 0.36 70 252 34 76 C-CVI SiC21 8 HS 0.36 87 324 62 74 C-C21 8 HS 0.36 101 305 21 70 Al2O3/Mullite20 8 HS B0.4 B97 210 50 52 MI, melt-infiltrated; UTS, ultimate tensile strength. J Note that in Eq. (3), the ellipsoid character of off-axis oriented fibers in the plane of failure was not taken into account. A more exhaustive and complicated analysis would re￾quire not only the angle, but the nature of fiber/matrix sliding which is not well understood,2 the matrix crack opening and the degree of fiber straightening. From an engineering and design standpoint, the simple analysis performed here is considered to be more straightfor￾ward and useful. October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3191

3192 Journal of the American Ceramic Society-Morscher et al. VoL. 90. No. 10 Table IV. Ultimate Strength Properties of Composites in the (5) Implications for Architecture Design of MI SiC/SiC Literature Tested in Multiple Directions From the observations of this study, one can summarize several Composite architecture Orientation ffas(GPa)(MPa) onstituent and architectural guidelines that can be applied to Matrix-dominated composites uture designs of components fabricated with ceramic compos- Nic-MAS 0/90 laminate 00.185130385 ites in general and nonoxide Syl-iBN MI SiC/SiC composites in 300.185120147 articular where high off-axis strengths are required. It is as- 450.37110157 sumed that the design goals will be to achieve as high a matrix Nic-CVI23 2 D woven0/9000.2220190 racking stress as possible as well as a high UTS along the prin- 200.2210170 cipal stress directions within the components 450.4210170 First, the matrix constituent should display a high stiffness Fiber-dominated composites and high strain capability by utilizing a high-modulus compo- C- epoxy 00.25NA461 sition, such as SiC, and a fabrication approach that results in as 150.25NA274 low a porosity as possible, such as MI SiC. For the Sic/sic composites of this study, the high-modulus low-porosity MI 300.25NA143 450.5NA12 matrix allows the composite elastic modulus to be fairly inde pendent of architecture and in-plane testing direction. Further oxy braid0/±45 00.1540.8417 more, as is the case with the MI process, the 0/±60 axial 11 31.6318 fabrication process should(if possible)result in a residual com- 0/±60hoop300.4649.9400 pressive stress on the matrix critical flaws or weak portion of the 0/±45hoop 5 0.42 19.8 165 fiber architecture after final composite fabrication(see Table D) CVI, chemical vapor infiltration: UTS. ultimate tensile strength. In contrast to residual stresses caused by thermal expansion mismatch between the fiber and the matrix the residual stress of the mi process appears to be independent of temperature up to at least 81 45). This is an important design consideration for composite Second, the fiber constituent should be as strong as possible and architecture selection when appreciable off-axis loading ap n its as-produced condition and should retain a high fraction of plications are pursued this strength after composite fabrication. The on-axis UTS val Note that the Nic/CVI SiC ,composites had the highest ues for the Sylramic-iBN MI composites of this study(see Table elative off-axis strength. As discussed above, the eq. (3)anal D)and other studieswhen normalized by the fiber volume ysis does not account for such material properties as interfacial fraction are the highest( 2400 MPa)displayed to date by ding stress and fiber bending stiffness, which will affect the any woven Sic-based fiber in an MI composite. For off-axis mechanics of fiber fracture when aligned at an angle within a UTS, high fiber strength is also important for obtaining high transverse matrix crack. The absolute off-axis strengths of the fiber load-carrying ability during bending within matrix cracks. iC/CVI SiC composites were similar to the HN/MI and In addition, the fiber should have as high a modulus as possible Sylramic/MI and less than the Syl-iBN/MI composites. Also in order to shift composite loads away from the matrix flaws or Culto of the Nic/CVI SiC composites were poor, indicating a low rom weak portions of the fiber architecture and onto the load- n sI ter processing) fiber strength in comparison with the bearing fiber. Also, in combination with the interphase coatin othe es, possibly resulting in higher than ex- the fiber surface conditions should be such as to provide hig pected normalized off-axis fiber strength results. Another pos- interfacial shear to inhibit large crack-bridging fiber lengths that bility is that the beneficial effects of a lower fiber modulus and will both statistically reduce fiber strength within the cracks interfacial sliding stress are becoming evident for the nicalon auge length effect)and allow more fiber bending for off-axis composites. Nevertheless, it is suggested that the two lines in ading. Due to the relatively high surface roughness of the Fig. Il represent, albeit crudely, off-axis UTS for matrix-dor Iramic-ibn fiber and due to the stiffness of the combined inated and fiber-dominated CMC. These relationships can be iBN/BN interphase coating, interfacial shear strengths in Sylra used by designers as a"rule-of-thumb"when considering the mic- iBN/BN/SiC systems are approximately 70 Mpa, much type of composite and fiber architecture to be used in a com- than observed with other smoother fiber types and or ponent with off-axis stresses. carbon interfacial coatings when processed in the same manner and with similar thickness 13 Third, the type of fiber architecture and orientation of fiber must be judiciously selected in relation to the directions and magnitudes of the principal stresses within the CMC compo- Matrix Dominated SiC/SiC CMC nent. a primary guideline shown in this study is to effective fiber volume fractions as high as possible in these pi a 0. pal stress directions, both to reduce stress on matrix flaws and/ r weak 90 minicomposites and to increase the stress and strain for ultimate failure. However as shown here. the conventional 0.61 Dominated Nic/MAS [22] approach of putting the primary fiber axes directly along the 0.5 principal stress directions may not be required. For example for Nic/CVI the 0/90 panel aligned at 45 to the primary fiber axes, the ae Epoxy PW[181 onset cracking stress increased from 190 to 220 MPa due in part to the removal of minicomposites being perpendicular to the boxy braid [191 oading direction. But UTS values decreased significantly from Al203/mullite[20] 410 to 242 MPa because of fiber strength loss within open ma C/C& C/CVI SiC [21] rix cracks. However, for the braided panel aligned at a smaller 3 to the primary fiber axes, onset stresses remained high at 220 MPa and UTS only degraded to 350 MPa. Generally, it ngle of loaded-fibers off axis, degrees is thought that high us values are desirable for co Fig. 11. Effect of testing direction and composite type on the relative materials, but for nonoxide ceramic composites, such as MI SiC/ trength retention of the fibers for the composite data in Tables l, l SiC, structural life degrades in a complex manner above the cracking stress due to environment ingress through the open

451). This is an important design consideration for composite and architecture selection when appreciable off-axis loading ap￾plications are pursued. Note that the Nic/CVI SiC2,21,23 composites had the highest relative off-axis strength. As discussed above, the Eq. (3) anal￾ysis does not account for such material properties as interfacial sliding stress and fiber bending stiffness, which will affect the mechanics of fiber fracture when aligned at an angle within a transverse matrix crack. The absolute off-axis strengths of the NiC/CVI SiC composites were similar to the HN/MI and Sylramic/MI and less than the Syl-iBN/MI composites. Also, sult0 of the Nic/CVI SiC composites were poor, indicating a low in situ (after processing) fiber strength in comparison with the other SiC/SiC composites, possibly resulting in higher than ex￾pected normalized off-axis fiber strength results. Another pos￾sibility is that the beneficial effects of a lower fiber modulus and interfacial sliding stress are becoming evident for the Nicalon composites. Nevertheless, it is suggested that the two lines in Fig. 11 represent, albeit crudely, off-axis UTS for matrix-dom￾inated and fiber-dominated CMC. These relationships can be used by designers as a ‘‘rule-of-thumb’’ when considering the type of composite and fiber architecture to be used in a com￾ponent with off-axis stresses. (5) Implications for Architecture Design of MI SiC/SiC Components From the observations of this study, one can summarize several constituent and architectural guidelines that can be applied to future designs of components fabricated with ceramic compos￾ites in general and nonoxide Syl-iBN MI SiC/SiC composites in particular where high off-axis strengths are required. It is as￾sumed that the design goals will be to achieve as high a matrix cracking stress as possible as well as a high UTS along the prin￾cipal stress directions within the components. First, the matrix constituent should display a high stiffness and high strain capability by utilizing a high-modulus compo￾sition, such as SiC, and a fabrication approach that results in as low a porosity as possible, such as MI SiC. For the SiC/SiC composites of this study, the high-modulus low-porosity MI matrix allows the composite elastic modulus to be fairly inde￾pendent of architecture and in-plane testing direction. Further￾more, as is the case with the MI process, the composite fabrication process should (if possible) result in a residual com￾pressive stress on the matrix critical flaws or weak portion of the fiber architecture after final composite fabrication (see Table I). In contrast to residual stresses caused by thermal expansion mismatch between the fiber and the matrix, the residual stress of the MI process appears to be independent of temperature up to at least 8151C.25 Second, the fiber constituent should be as strong as possible in its as-produced condition and should retain a high fraction of this strength after composite fabrication. The on-axis UTS val￾ues for the Sylramic-iBN MI composites of this study (see Table I) and other studies7 when normalized by the fiber volume fraction are the highest (B2400 MPa) displayed to date by any woven SiC-based fiber in an MI composite. For off-axis UTS, high fiber strength is also important for obtaining high fiber load-carrying ability during bending within matrix cracks. In addition, the fiber should have as high a modulus as possible in order to shift composite loads away from the matrix flaws or from weak portions of the fiber architecture and onto the load￾bearing fiber. Also, in combination with the interphase coating, the fiber surface conditions should be such as to provide high interfacial shear to inhibit large crack-bridging fiber lengths that will both statistically reduce fiber strength within the cracks (gauge length effect) and allow more fiber bending for off-axis loading. Due to the relatively high surface roughness of the Sylramic-iBN fiber and due to the stiffness of the combined iBN/BN interphase coating, interfacial shear strengths in Sylra￾mic-iBN/BN/SiC systems are approximately 70 Mpa,7 much higher than observed with other smoother fiber types and/or carbon interfacial coatings when processed in the same manner and with similar thickness.13,26 Third, the type of fiber architecture and orientation of fibers must be judiciously selected in relation to the directions and magnitudes of the principal stresses within the CMC compo￾nent. A primary guideline shown in this study is to achieve effective fiber volume fractions as high as possible in these prin￾cipal stress directions, both to reduce stress on matrix flaws and/ or weak 901 minicomposites and to increase the stress and strain for ultimate failure. However, as shown here, the conventional approach of putting the primary fiber axes directly along the principal stress directions may not be required. For example for the 0/90 panel aligned at 451 to the primary fiber axes, the AE onset cracking stress increased from 190 to 220 MPa due in part to the removal of minicomposites being perpendicular to the loading direction. But UTS values decreased significantly from 410 to 242 MPa because of fiber strength loss within open ma￾trix cracks. However, for the braided panel aligned at a smaller 231 to the primary fiber axes, onset stresses remained high at B220 MPa and UTS only degraded to B350 MPa. Generally, it is thought that high UTS values are desirable for composite materials, but for nonoxide ceramic composites, such as MI SiC/ SiC, structural life degrades in a complex manner above the cracking stress due to environment ingress through the open Table IV. Ultimate Strength Properties of Composites in the Literature Tested in Multiple Directions Composite Fiber architecture Orientation f0 or f
off-axis E (GPa) UTS (MPa) Matrix-dominated composites Nic-MAS22 0/90 laminate 0 0.185 130 385 30 0.185 120 147 45 0.37 110 157 Nic-CVI23 2D woven 0/90 0 0.2 220 190 20 0.2 210 170 45 0.4 210 170 Fiber-dominated composites C - epoxy PW18 Plain Weave 0 0.25 NA 461 15 0.25 NA 274 30 0.25 NA 143 45 0.5 NA 127 C - epoxy braid19 0/745 axial 0 0.15 40.8 417 0/760 axial 0 0.11 31.6 318 0/760 hoop 30 0.46 49.9 400 0/745 hoop 45 0.42 19.8 165 CVI, chemical vapor infiltration; UTS, ultimate tensile strength. 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 0 10 20 30 40 50 Angle of loaded-fibers off axis, degrees Norm. Effective Fiber Strength Al2O3/mullite [20] Nic/C [2] C/epoxy PW [18] Syl-iBN/MI (this study) Syl-iBN/MI Braid (this study) Syl/MI [22] Nic/MAS [22] Nic/CAS [2] HN/MI [22] Nic/CVI [2] C/epoxy braid [19] Nic/CVI [21] C/C & C/CVI SiC [21] Fiber Axis Loading Axis angle Fiber Dominated CMC Matrix Dominated SiC/SiC CMC Nic/CVI [23] Fig. 11. Effect of testing direction and composite type on the relative strength retention of the fibers for the composite data in Tables I, III, and IV. 3192 Journal of the American Ceramic Society—Morscher et al. Vol. 90, No. 10

October 2007 Tensile Mechanical Properties of Ceramic Matrix Composites 3193 matrix cracks. In addition, above the cracking stress, composite Acknowledgments properties such as modulus and thermal conductivity also de- grade in complex and unpredictable ways. Thus, today most would like to thank Wayne Steffier of Hyper- Therm Composites or provision of the excellent property database of the various CMCs that CMC designers of nonoxide CMc components that require ng service life find it more desirable to achieve as high a crack ing stress as possible in the principal stress directions, but still References retain some comfortable level of ultimate strength. As shown in his study, this goal can best be achieved by increasing the effec ceedings of the the properties o f fibre composites Conference. pp. 15-24. IPC Sci- attempt to minimize the size and volume fraction of minicom- Several Ceramic-Matrix Composites,"JAm ceran posites oriented perpendicular to these directions. For thos cases where perpendicular minicomposites cannot be avoided Aging on the Mechanical Properties of a Porous-Matrix Ceramic Composite, lorscher and R. T. Bhatt. "Sic mechanistic-based mini-matrix approach can be helpful in un- derstanding and modeling matrix cracking for a variety of 2D Chap and 3D fiber architectures. It should also be noted that the PH M. Yun, J. Z. Gyekenyesi, Y.LChen,DR.Wheeler, and J.A.DiCarlo effects of creep rupture on fiber orientation need to be deter- Fibers. "Ceran. Eng. Sci. Proc. 22.521-3(2001 Reinforced by Treated Sylramic Sic mined and considered as well before deciding on a fiber archi- tecture and fiber orientation ven SiC/SiC Composite, "Comp. Sci. Technol, 59, 687-9(1999 N. Morscher. ""Stress-Dependent Matrix CI Reinforced Melt-Infiltrated SiC Matrix Composites, Comp. Sci. Tecno, 64, V. Summary and Conclusions Test Specimens at Ambient Temperature. ASTM, West Conshohocken, PA, (2000) A 2D-woven 0/90 and a 2D-braided [0/+67 Sylramic-iBN-re- s. Kalluri, A. Calomino, and D N. Brewer, "An Assessment of Variability in gle symmetrically located between the two primary fiber di- Eng. Sei Pr lensile pr nforced MI SiC matrix panel when tensile tested in-plane at an the A 用 (2004 Mel-Infiltrated SiC/SiC Composite,"Ceram rections were both shown to exhibit high stress for the onset of N. Morscher and V. Pujar, ""Melt-Infiltrated SiC Composites for Gas Tur- TTMC. The matrix cracking and deflection from linearity stress- Land, Sea, and Air, Ju 17, 2004, Vienna, Austria, Paper no. GT2004-5423 es were actually higher for the off-axis loading condition com- pared with 2D-woven [0/90] panels loaded in one of the primary Sic/SiC Composites. eram.Soe,214-152777880200 fiber directions. For both panels, this improved cracking behav. B. N. Cox and D. B. Marshall. "Crack Initiation in Fiber-Reinforced Brittle Laminates, " J. An. Ceram. Soc. 79 [5]1181-88(1996). ior can be explained in part by a higher effective fraction of fi- 3G. N. Morscher. "Modal Acoustic Emission Source Determination in Silicon bers in the loading direction, which reduces internal stress on Review te stress. For the woven ondestructive Evaluation. Edited by D. O. Thompson, and D. E Chimenti CP panel, there also existed an absence of weak minicomposit IG. N. Morscher"Matrix Cracking in Four Different 2D SiC/SiC Composite oriented perpendicular to the loading direction. For the braided ystems. Published in the 35th International SAMPE panel, which did have a low fraction of axial minicomposit ites were well separated from one another, which prevented the ig n morsch deling the Elastic Modulus of 2D Woven CVI SiC occurrence of‘back-to-back”90° minicomposites that are the source of low-strength matrix cracks in 2D-woven [0/90)com- sis of Fiber-Reinforced posites with similar tow size. AE measurements on the braided 407-18 in ASM Engineered Materials Handbook. Vol 21: Composites. ASM also showed that the onset stress for ttmc and the K. Naik. P. S. Shemekar, and M. v Ire Behavior of woven cracking distribution with increasing stress behaved effectively in the same way as 2D-woven [0/90]composites when loaded in °p. L. Falzon and l. Herzberg,“ Mechanical of 2-D Braided Car- a primary fiber direction. Not only does this confirm that the weak minicomposites perpendicular to the loading direction are Properties of an All-Oxide Ceramic Composite, " J. Am. Ceram. cal sources of it should also enable So,820272l-30(1999) matrix"approach?to model and improve matrix cracking and Wayne Steffier Composite Data Sheets, Hyper- mposites Inc (2006 DFL stresses of CMC components with braided and other 2C. S. Lynch and A. G. Evans, "Effects of OfT-Ax the Tensile bu architectures Lamon. and Allix "Model of the Non In terms of in-plane UTS, the Syl-iBN/MI SiC composite Behavior of 2D SiC-SiC Chemical Vapor Infiltration Composites, J.Amm. Ceran. were found to be superior in both on-axis and off-axis behavior n comparison with other 2D ceramic composites in the NASAs Enabling Propulsion Materials Program literature with either fiber-dominated or matrix-dominated me- chanical behavior. The off-axis behavior can be attributed pri- 2G. N. Morscher and J. Z. Gyekenesi, "Room Temperature Tensile Behavior marily to the relatively high modulus and strength of the mI on Reinforced SiC Matrix Composite matrix, which carries significant load and to the high strength Ceram. Eng. Sci. Proc., 19[3]241-9(1998) of the Sylramic-iBN fiber that is retained during composite G. N. Morscher, H. M. Yun, and J. A. DiCarlo, " Matrix J.Am. Cerami.Soc,88l]146-53(2005

matrix cracks. In addition, above the cracking stress, composite properties such as modulus and thermal conductivity also de￾grade in complex and unpredictable ways. Thus, today most CMC designers of nonoxide CMC components that require long service life find it more desirable to achieve as high a crack￾ing stress as possible in the principal stress directions, but still retain some comfortable level of ultimate strength. As shown in this study, this goal can best be achieved by increasing the effec￾tive fiber fraction in these directions, but in doing so, one should attempt to minimize the size and volume fraction of minicom￾posites oriented perpendicular to these directions. For those cases where perpendicular minicomposites cannot be avoided, this and other studies7,27 have shown that the fairly robust mechanistic-based mini-matrix approach can be helpful in un￾derstanding and modeling matrix cracking for a variety of 2D and 3D fiber architectures. It should also be noted that the effects of creep rupture on fiber orientation need to be deter￾mined and considered as well before deciding on a fiber archi￾tecture and fiber orientation. V. Summary and Conclusions A 2D-woven 0/90 and a 2D-braided [0/767] Sylramic-iBN-re￾inforced MI SiC matrix panel when tensile tested in-plane at an angle symmetrically located between the two primary fiber di￾rections were both shown to exhibit high stress for the onset of TTMC. The matrix cracking and deflection from linearity stress￾es were actually higher for the off-axis loading condition com￾pared with 2D-woven [0/90] panels loaded in one of the primary fiber directions. For both panels, this improved cracking behav￾ior can be explained in part by a higher effective fraction of fi- bers in the loading direction, which reduces internal stress on critical matrix flaws for a given composite stress. For the woven panel, there also existed an absence of weak minicomposites oriented perpendicular to the loading direction. For the braided panel, which did have a low fraction of axial minicomposites loaded perpendicular to the loading direction, the minicompos￾ites were well separated from one another, which prevented the occurrence of ‘‘back-to-back’’ 901 minicomposites that are the source of low-strength matrix cracks in 2D-woven [0/90] com￾posites with similar tow size. AE measurements on the braided panel also showed that the onset stress for TTMC and the cracking distribution with increasing stress behaved effectively in the same way as 2D-woven [0/90] composites when loaded in a primary fiber direction. Not only does this confirm that the weak minicomposites perpendicular to the loading direction are the critical sources of, it should also enable use of the ‘‘mini￾matrix’’ approach7 to model and improve matrix cracking and DFL stresses of CMC components with braided and other architectures. In terms of in-plane UTS, the Syl-iBN/MI SiC composites were found to be superior in both on-axis and off-axis behavior in comparison with other 2D ceramic composites in the literature with either fiber-dominated or matrix-dominated me￾chanical behavior. The off-axis behavior can be attributed pri￾marily to the relatively high modulus and strength of the MI matrix, which carries significant load, and to the high strength of the Sylramic-iBN fiber that is retained during composite processing. Acknowledgments We would like to thank Wayne Steffier of Hyper-Therm Composites Incor￾porated for provision of the excellent property database of the various CMCs that they fabricate. References 1 J. 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