《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_Ceramic fiber composites：Experimental analysis and modeling of mechanical properties

Availableonlineatwww.sciencedirect.com COMPOSITES scⅰ enceDirect CIENCE AND TECHNOLOGY ELSEVIER Composites Science and Technology 68 (2008)1165-l1 www.elsevier.com/locate/compscitech Ceramic fiber composites: Experimental analysis and modeling of mechanical properties Dietmar Koch, Kamen Tushtev, Georg Grathwohl Unicersity of Bremen, Department of Ceramic Materials and Components, D-28359 Bremen, Germany Received 16 May 2007: received in revised form 22 June 2007: accepted 29 June 2007 Available online 19 July 2007 Abstract Ceramic fiber reinforced ceramic matrix composites(CMC)are outstanding ceramics with high fracture toughness. This can be real ized if both brittle components of the composite, i.e., fibers and matrix are interacting with each other in an efficient way. Either a weak interface allowing debonding between fiber and matrix controls the fracture processes (WIC-CMC) or the matrix takes this role of a weak and more compliant component(WMC-CMC). An experimental test data base is presented for a WMC-type composite where the materials data are used to establish a model which describes the materials behavior in a macroscopic way. Inelastic deformation and materials damage processes are defined, measured and interpreted on the base of a continuum damage mechanics concept. The elas ic and inelastic response is then predictable up to failure as being dependent on the angle between fiber and loading directions of the specimen o 2007 Elsevier Ltd. All rights reserved Keywords: A. Ceramic-matrix composites( CMCs); B. Stress/strain curves; B. Mechanical properties; C Modelling: C Finite element analysis(FEA) 1. Introduction tolerance of cmc it has to be assured that the fibers remain intact and effective when cracks propagate. For this Continuous fiber reinforced ceramic matrix composites purpose the balance between the strength of the fibers and are mainly developed for specific applications at high tem- the crack resistance of the other microstructural compo- peratures and in oxidative atmosphere. The fibers as rein- nents, i. e. matrix and fiber matrix interface, has to be con forcing components offer higher strength and stiffness trolled in a way that the survival probability of the fibers is compared to the matrices which are, in contrast, character- not too low. This effect can be reached by the adjustment of ized by inferior properties as in most cases their microstruc- the CMC microstructure in two alternative ways tures exhibit microcracks, residual pores and gradients or inhomogeneities caused by the CMC fabrication process. The crack resistance of the fiber matrix interface is low- Typical processing routes are chemical vapor infiltration ered in order to allow debonding between fiber and (CVI), liquid polymer infiltration(LPI), liquid silicon infil matrix CMCs of this type are characterized by a weak tration (Lsi, directed metal oxidation (DiMOx), and cera interface or interphase. They are termed Weak Interface mic slurry impregnation(CSI) which lead to characteristic Composites, WIc properties of the resulting CMC. The interfacial bonding between fiber and matrix is of When CMCs are stressed crack initiation and propaga minor importance as the matrix is weak enough and tion generally originate from the matrix For enhanced flaw susceptible for multiple cracking while the fibers pro- vide strength and crack tolerance of the CMC. These composites are called Weak Matrix Composites, E-mail address: koch(ceramics. uni-bremen de (D. Koch) WMC 02663538/S. see front matter 2007 Elsevier Ltd. All rights reserved doi:10.1016j.compscitech.2007.06.029

Ceramic fiber composites: Experimental analysis and modeling of mechanical properties Dietmar Koch *, Kamen Tushtev, Georg Grathwohl University of Bremen, Department of Ceramic Materials and Components, D-28359 Bremen, Germany Received 16 May 2007; received in revised form 22 June 2007; accepted 29 June 2007 Available online 19 July 2007 Abstract Ceramic fiber reinforced ceramic matrix composites (CMC) are outstanding ceramics with high fracture toughness. This can be real￾ized if both brittle components of the composite, i.e., fibers and matrix are interacting with each other in an efficient way. Either a weak interface allowing debonding between fiber and matrix controls the fracture processes (WIC–CMC) or the matrix takes this role of a weak and more compliant component (WMC–CMC). An experimental test data base is presented for a WMC-type composite where the materials data are used to establish a model which describes the materials behavior in a macroscopic way. Inelastic deformation and materials damage processes are defined, measured and interpreted on the base of a continuum damage mechanics concept. The elas￾tic and inelastic response is then predictable up to failure as being dependent on the angle between fiber and loading directions of the specimens. 2007 Elsevier Ltd. All rights reserved. Keywords: A. Ceramic–matrix composites (CMCs); B. Stress/strain curves; B. Mechanical properties; C. Modelling; C. Finite element analysis (FEA) 1. Introduction Continuous fiber reinforced ceramic matrix composites are mainly developed for specific applications at high tem￾peratures and in oxidative atmosphere. The fibers as rein￾forcing components offer higher strength and stiffness compared to the matrices which are, in contrast, character￾ized by inferior properties as in most cases their microstruc￾tures exhibit microcracks, residual pores and gradients or inhomogeneities caused by the CMC fabrication process. Typical processing routes are chemical vapor infiltration (CVI), liquid polymer infiltration (LPI), liquid silicon infil￾tration (LSI), directed metal oxidation (DiMOx), and cera￾mic slurry impregnation (CSI) which lead to characteristic properties of the resulting CMC. When CMCs are stressed crack initiation and propaga￾tion generally originate from the matrix. For enhanced flaw tolerance of CMC it has to be assured that the fibers remain intact and effective when cracks propagate. For this purpose the balance between the strength of the fibers and the crack resistance of the other microstructural compo￾nents, i.e. matrix and fiber matrix interface, has to be con￾trolled in a way that the survival probability of the fibers is not too low. This effect can be reached by the adjustment of the CMC microstructure in two alternative ways: • The crack resistance of the fiber matrix interface is low￾ered in order to allow debonding between fiber and matrix. CMCs of this type are characterized by a weak interface or interphase. They are termed Weak Interface Composites, WIC. • The interfacial bonding between fiber and matrix is of minor importance as the matrix is weak enough and susceptible for multiple cracking while the fibers pro￾vide strength and crack tolerance of the CMC. These composites are called Weak Matrix Composites, WMC. 0266-3538/$ - see front matter 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2007.06.029 * Corresponding author. E-mail address: dkoch@ceramics.uni-bremen.de (D. Koch). www.elsevier.com/locate/compscitech Available online at www.sciencedirect.com Composites Science and Technology 68 (2008) 1165–1172 COMPOSITES SCIENCE AND TECHNOLOGY

D. Koch et al Composites Science and Technology 68(2008)1165-1172 In any case, the composites need a microstructural component which is weak enough to allow deformation, debonding and cracks to propagate without leading to at9oo°C state 0.20 pontaneous failure of the composite 2. Weak interface composites, WIC No debonding The microstructural design of CMC was primarily dri- 0.10 en by the development of a weak interface between fiber Initial and matrix in order to achieve crack deviation at the fiber debond ing matrix interface. With this microstructural model approach the mechanical behaviour of wic can be described accu tely. If an unidirectional composite is loaded in tensile mode the initial cracks propagate first in the matrix as the fibers are stronger and can reach a higher strain to fail- Indenter Displacement /mm Then the matrix crack propagates through the com- Fig. 2. Push-in-curves revealing the prevention of interfacial debonding in posite being bridged by the strong fibers which remain a SiC/DiMOx-CMc due to oxidation and formation of silica at 900C intact as the stress concentration at the interface does not resulting in brittle failure of the composite induce fiber failure but initial interfacial debonding. Using curve allowing slight increase of the critical ratio T/rEif the well-known relationship presented by He and Hutchin- on[I] the required (low) fracture energy of the interface r the relative Young's modulus of the composite deviates for initiation of these debonding processes can be calcu- from zero lated in relation to the fracture energy(surface energy) of If the condition of the critical fracture ratio the fiber rF and a ratio Tr< 0.25 must not be surpassed according to Fig. I is fulfilled, beyond ma acking for non-brittle behavior if fiber and matrix have similar strength WIC ceramics exhibit stress strain with a Young's moduli EF and EM, respectively. Typical represen- typical nonlinear characteristic arising from debonding, tatives of this type of composites are CMC with dense and multiple crack initiation, propagation and opening. The crystalline matrices as, e. g, the CVI derived SiC-matrix or load is increasingly transferred to the fibers and fiber fail- the DiMOx derived Al2O3-matrix. These CMCs are typical ure is initiated sequentially leading to the increase of the WIC materials for which fiber coating with the effect of fraction of failed fibers due to the statistic scattering of bon pyC, SiC, BN, and BCN layers or combinations of Applying the composites at high temperatures in oxida- these layers)is inevitable. The quantitative result is shown tive atmosphere the fiber coatings may be attacked due to in the He-Hutchinson-diagram( Fig 1) with the boundary environmental conditions. If oxidation of the interphase Brittle failure T>TE.krit due to Oxidation WMC 1008-06-04-020.00204060.81.0 Relative Youngs Modulus(EF-EM(E+EM Fig. I. Boundary curve according to He and Hutchinson for realization of non brittle behavior taking into consideration the critical relative fracture energy dependent on the stiffness ratio of fiber and matrix. Additionally, the effects of interfacial oxidation and matrix densification on failure behavior are

In any case, the composites need a microstructural component which is weak enough to allow deformation, debonding and cracks to propagate without leading to spontaneous failure of the composite. 2. Weak interface composites, WIC The microstructural design of CMC was primarily dri￾ven by the development of a weak interface between fiber and matrix in order to achieve crack deviation at the fiber matrix interface. With this microstructural model approach the mechanical behaviour of WIC can be described accu￾rately. If an unidirectional composite is loaded in tensile mode the initial cracks propagate first in the matrix as the fibers are stronger and can reach a higher strain to fail￾ure. Then the matrix crack propagates through the com￾posite being bridged by the strong fibers which remain intact as the stress concentration at the interface does not induce fiber failure but initial interfacial debonding. Using the well-known relationship presented by He and Hutchin￾son [1] the required (low) fracture energy of the interface CI for initiation of these debonding processes can be calcu￾lated in relation to the fracture energy (surface energy) of the fiber CF and a ratio CI/CF 6 0.25 must not be surpassed for non-brittle behavior if fiber and matrix have similar Young’s moduli EF and EM, respectively. Typical represen￾tatives of this type of composites are CMC with dense and crystalline matrices as, e.g., the CVI derived SiC-matrix or the DiMOx derived Al2O3-matrix. These CMCs are typical WIC materials for which fiber coating with the effect of lowering the fracture energy of the interface (e.g., pyrocar￾bon pyC, SiC, BN, and BCN layers or combinations of these layers) is inevitable. The quantitative result is shown in the He–Hutchinson-diagram (Fig. 1) with the boundary curve allowing slight increase of the critical ratio CI/CF if the relative Young’s modulus of the composite deviates from zero. If the condition of the critical fracture energy ratio according to Fig. 1 is fulfilled, beyond matrix cracking strength WIC ceramics exhibit stress strain curves with a typical nonlinear characteristic arising from debonding, multiple crack initiation, propagation and opening. The load is increasingly transferred to the fibers and fiber fail￾ure is initiated sequentially leading to the increase of the fraction of failed fibers due to the statistic scattering of fiber strength up to the final failure of the composite. Applying the composites at high temperatures in oxida￾tive atmosphere the fiber coatings may be attacked due to environmental conditions. If oxidation of the interphase -1.0 -0.8 -0.6 -0.4 -0.2 0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 Relative Fracture Energy ΓI / ΓF Relative Young's Modulus (EF -EM)/(EF +EM) Non-brittle failure Brittle failure ΓI > ΓF,krit due to reinfiltr . WMC due to Oxidation WIC Fig. 1. Boundary curve according to He and Hutchinson for realization of non brittle behavior taking into consideration the critical relative fracture energy dependent on the stiffness ratio of fiber and matrix. Additionally, the effects of interfacial oxidation and matrix densification on failure behavior are shown. Indenter Displacement / mm 0.0 0.2 0.4 0.6 0.8 1.0 1.2 Indenter Load / N 0.00 0.05 0.10 0.15 0.20 0.25 900˚C RT No debonding Initial debonding After oxidation at 900˚C Original state Silica layer After oxidation at 900°C Original state Silica layer Fig. 2. Push-in-curves revealing the prevention of interfacial debonding in a SiC/DiMOx–CMC due to oxidation and formation of silica at 900 C resulting in brittle failure of the composite. 1166 D. Koch et al. / Composites Science and Technology 68 (2008) 1165–1172

D. Koch et al. Composites Science and Technology 68(2008)1165-1172 1167 sic/sic,cⅥ +45°-45° 3 2 0 [12]风闾8]10冈冈图3图3[2 Porous matrix Dense matrix Tensile Test WMC WIC C/C. LPI 0°90° Fig. 4. WMC-WIC classification of different composites according to their stiffness ratio Eoo/E4so. Results from literature [2, 3, 5-8, 10, 12]and own measurements(marked with [x] Size effects are also not really significant and low scattering of strength is observed +45°-45° 3. Weak matrix composites, WMC Infiltration processes as LPl, Lsl, or CSI to provide the matrix in CMCs lead to microstructures which are charac- Strain[°%] terized by a fine porosity and therefore a low stifness. The significantly reduced matrix stifness and strength, com- Fig. 3. Representative stress versus strain curves of (a)WIC(CVI SiC/ pared to WIC, enables debonding processes and thus dam Si[2]and(b)WMC(LPIC/C) from axial(090° and diagonal(±45°) loaded tensile tests age tolerance even in the case of a strong fiber-matrix interface as the cracks which propagate through the matrix easily deviate close to the fiber surface through the matrix. layers occurs, the mechanical properties may change and This concept again corresponds to the theoretical analysis lead to an increase of the relative fracture energy of fiber of He and Hutchinson(Fig. 1)where the large difference and interface. Acceptable changes of the interfacial proper- between the Youngs moduli of the stiff fiber and the weak ties without provoking brittle behavior can be discussed matrix allows a much stronger bonding and a higher ratio using the borderline in Fig. 1. Hence, the interphase has between the fracture energies of interface and fiber. while to accomplish not only mechanical functions in order to the matrix fails at low stresses the composite can still be provide debonding: it also has to fulfill thermal and envi- loaded well above the matrix cracking strength as long as ronmental boundary conditions as e.g. sufficient oxidation the overall load can be transferred to the fibers. However resistance. Fig. 2 shows the effect of oxidation of the fiber as the redistribution of stresses from the fiber to the matrix matrix interface in a SiC/DiMOx composite as manifested does not take place in a significant manner final failure of by single fiber push-in tests. It shows that due to the forma- the composite occurs when the fibers do not fail locally tion of silica at the interface after oxidation at 900C initial restricted but in a large volume of the component. Thus, debonding is no longer possible at sufficiently low stresses. the mechanical behavior can no longer be described by a This leads to brittle failure of the composite as debonding micromechanical approach. Furthermore the mechanical is prevented behavior of WMC is dominated strongly by the properties It can be concluded that WIC with relatively strong and of the fibers and therefore the mechanical performanc stiff matrices and obligatory fiber coating provide high depends on their orientation and the loading direction fracture toughness values. Following the micromechanical As the matrix is not able to carry significant load, low mechanisms of debonding, these CMCs are relatively notch strength will be obtained under compression or loading insensitive with the highest achievable stress being fairly mode with an angle between fiber orientation and loading independent of the fiber and applied stress orientations. direction

layers occurs, the mechanical properties may change and lead to an increase of the relative fracture energy of fiber and interface. Acceptable changes of the interfacial proper￾ties without provoking brittle behavior can be discussed using the borderline in Fig. 1. Hence, the interphase has to accomplish not only mechanical functions in order to provide debonding; it also has to fulfill thermal and envi￾ronmental boundary conditions as e.g. sufficient oxidation resistance. Fig. 2 shows the effect of oxidation of the fiber matrix interface in a SiC/DiMOx composite as manifested by single fiber push-in tests. It shows that due to the forma￾tion of silica at the interface after oxidation at 900 C initial debonding is no longer possible at sufficiently low stresses. This leads to brittle failure of the composite as debonding is prevented. It can be concluded that WIC with relatively strong and stiff matrices and obligatory fiber coating provide high fracture toughness values. Following the micromechanical mechanisms of debonding, these CMCs are relatively notch insensitive with the highest achievable stress being fairly independent of the fiber and applied stress orientations. Size effects are also not really significant and low scattering of strength is observed. 3. Weak matrix composites, WMC Infiltration processes as LPI, LSI, or CSI to provide the matrix in CMCs lead to microstructures which are charac￾terized by a fine porosity and therefore a low stiffness. The significantly reduced matrix stiffness and strength, com￾pared to WIC, enables debonding processes and thus dam￾age tolerance even in the case of a strong fiber–matrix interface as the cracks which propagate through the matrix easily deviate close to the fiber surface through the matrix. This concept again corresponds to the theoretical analysis of He and Hutchinson (Fig. 1) where the large difference between the Young’s moduli of the stiff fiber and the weak matrix allows a much stronger bonding and a higher ratio between the fracture energies of interface and fiber. While the matrix fails at low stresses the composite can still be loaded well above the matrix cracking strength as long as the overall load can be transferred to the fibers. However, as the redistribution of stresses from the fiber to the matrix does not take place in a significant manner final failure of the composite occurs when the fibers do not fail locally restricted but in a large volume of the component. Thus, the mechanical behavior can no longer be described by a micromechanical approach. Furthermore the mechanical behavior of WMC is dominated strongly by the properties of the fibers and therefore the mechanical performance depends on their orientation and the loading direction. As the matrix is not able to carry significant load, low strength will be obtained under compression or loading mode with an angle between fiber orientation and loading direction. 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0 50 100 150 200 250 300 350 400 Tensile Test SiC/SiC, CVI +45°/ -45° 0°/ 90° Stress [MPa] Strain [%] 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0 50 100 150 200 250 300 350 400 Tensile Test C/C, LPI +45°/ -45° 0°/ 90° Stress [MPa] Strain [%] Fig. 3. Representative stress versus strain curves of (a) WIC (CVI SiC/ SiC) [2] and (b) WMC (LPI C/C) from axial (0/90) and diagonal (±45) loaded tensile tests. 0 1 2 3 4 5 6 E0 / E45 WHIPOX Nicalon/SiC Nicalon/SiC Nicalon/MAS C/SiC Nextel 610 /Mullite + Alumina Nextel 720 /Mullite + Alumina (Al C/C C/C, SIGRABOND C/C O2 3-ZrO2 )mc/Al O3 2 C/C [6] [5] [12] [x] [8] [7] [10] [x] [x] [3] [3] [2] WMC WIC Porous matrix Dense matrix Fig. 4. WMC–WIC classification of different composites according to their stiffness ratio E0/E45. Results from literature [2,3,5–8,10,12] and own measurements (marked with [x]). D. Koch et al. / Composites Science and Technology 68 (2008) 1165–1172 1167

D. Koch et aL Composites Science and Technology 68(2008)1165-1172 Table 1 ies of various CMC classified in Fig 4(as far as available from literature Material type Manufacturing method [6] Carbon T-300/Carbon, 16 layers, 0/900 56% fiber content LPL 4 infiltrations [5 Al2O3 Almax/ZrO2 AlO3, 6 layers, 0/900 33% fiber content minicomposite No detailed processing descripti Carbon Torayka M-40/Carbon, 15 layers, 0/90 50% fiber content Preformed yarn method LPI /C Sigrabond 1501 G, 24 layers, 0/90 open porosity 10-12% LPI Preformed 87x图2 [8 Carbon Torayka M40/C, 16 layers, 0 /90 50%fiber cont Nextel 720 /mullite and alumina, 12 layers, 0/900 39% fiber content matrix porosity 38% CSI up to [10] Nextel 610/80% mullite and 20% alumina 0/90 composite porosity 22-25% matrix porosity 38-42% CSI up to T800C/SiC 4 UD 0°/90° WHIPOX0°/90° Nicalon/MAS 12 0/900 40% fiber content fully dense MAS matrix Hot pressing Nicalon Sic/SiC 32% fiber content 8. 6% matrix porosity Nicalon SiC/SiC 0/90 plain weave 35% fiber content residual porosity 10-15% The characteristic difference between on-axis loading are infiltrated with the carbon matrix by several impregna (fibers are oriented in loading direction)and off-axis load- tion cycles with succeeding thermal treatments up to ing is manifested in Fig 3 for both composite types WIC 2000C. Finally, the resulting composite is characterized and WMC. Similar strength is measured for WIC under by an open porosity of 10-12% resulting from shrinkage both loading conditions, as discussed earlier on the base of micromechanical mechanisms On the contrary, in case of WMc the strength is strongly reduced under off-axis conditions because the fibers are not carrying the load sufficiently se I Longitudina The proposed concepts WIC and WMC are typical boundary examples with the real composites being situated somewhere in between. The properties of the composites not only depend on the manufacturing route, but also on d the chosen combination of fiber and matrix. In Fig. I it becomes obvious that with improved mechanical properties gD of the matrix the interfacial fracture energy plays a more important role and must be low enough in order to fulfill C 150/ →+1575 he boundary conditions of non brittle failure. This improvement of matrix properties can be reached by e. g reinfiltration cycles which is the regular case for the LPI process. Thus, if the matrix properties are enhanced the interfacial properties become more important. In Fig. 4 several composites from various sources are ranked con Strain (0%1900 and +45%/-459)as a measure of classification to b o WIC and WMC, respectively. Additionally, their proper- ties which were available from the literature are listed in q Table 1. It turns out that a typical WMC like C/C shows a high Eo/E4se ratio while a typical WIC like Nicalon/ +15%75° SiC is characterized by a Eoo/E4se ratio of 1. Other compos- 2 ites like oxide/oxide composites, LPI derived materials, and Sic fiber reinforced glass matrix composites are situated 090 between these boundary cases 方250 300 4. Materials and experiments ongitudinalTransverse Fundamental tests have been performed using a com- G」 mercially available C/C composite called Sigrabond 1501G(SGL Carbon, Germany). The material consists of 24 bidirectional reinforced layers of carbon fiber mats. Fig. 5. Typical stress strain curves of WMC composite C/C under(a) The composite is designed symmetrically and shows the tension and (b)compression with different angles between fiber orientation same properties in 00 and in 90 orientation. The fiber mats and loading direction

The characteristic difference between on-axis loading (fibers are oriented in loading direction) and off-axis load￾ing is manifested in Fig. 3 for both composite types WIC and WMC. Similar strength is measured for WIC under both loading conditions, as discussed earlier on the base of micromechanical mechanisms. On the contrary, in case of WMC the strength is strongly reduced under off-axis conditions because the fibers are not carrying the load sufficiently. The proposed concepts WIC and WMC are typical boundary examples with the real composites being situated somewhere in between. The properties of the composites not only depend on the manufacturing route, but also on the chosen combination of fiber and matrix. In Fig. 1 it becomes obvious that with improved mechanical properties of the matrix the interfacial fracture energy plays a more important role and must be low enough in order to fulfill the boundary conditions of non brittle failure. This improvement of matrix properties can be reached by e.g. reinfiltration cycles which is the regular case for the LPI process. Thus, if the matrix properties are enhanced the interfacial properties become more important. In Fig. 4 several composites from various sources are ranked con￾cerning their ratio of stiffness in on and off-axis orientation (0/90 and +45/45) as a measure of classification to WIC and WMC, respectively. Additionally, their proper￾ties which were available from the literature are listed in Table 1. It turns out that a typical WMC like C/C shows a high E0/E45 ratio while a typical WIC like Nicalon/ SiC is characterized by a E0/E45 ratio of 1. Other compos￾ites like oxide/oxide composites, LPI derived materials, and SiC fiber reinforced glass matrix composites are situated between these boundary cases. 4. Materials and experiments Fundamental tests have been performed using a com￾mercially available C/C composite called Sigrabond 1501G (SGL Carbon, Germany). The material consists of 24 bidirectional reinforced layers of carbon fiber mats. The composite is designed symmetrically and shows the same properties in 0 and in 90 orientation. The fiber mats are infiltrated with the carbon matrix by several impregna￾tion cycles with succeeding thermal treatments up to 2000 C. Finally, the resulting composite is characterized by an open porosity of 10–12% resulting from shrinkage Table 1 Properties of various CMC classified in Fig. 4 (as far as available from literature) Material type Manufacturing method [6] Carbon T-300/Carbon, 16 layers, 0/90 56% fiber content LPI, 4 reinfiltrations [5] Al2O3 Almax/ZrO2 + Al2O3, 6 layers, 0/90 33% fiber content minicomposite No detailed processing description [12] Carbon Torayka M-40 /Carbon, 15 layers, 0/90 50% fiber content Preformed yarn method LPI [x] C/C Sigrabond 1501 G, 24 layers, 0/90 open porosity 10–12% LPI [8] Carbon Torayka M40/C, 16 layers, 0/90 50% fiber content Preformed yarn method LPI [7] Nextel 720/mullite and alumina, 12 layers, 0/90 39% fiber content matrix porosity 38% CSI up to 1200 C [10] Nextel 610/80% mullite and 20% alumina 0/90 composite porosity 22–25% matrix porosity 38–42% CSI up to 1200 C [x] T800 C/SiC 4 UD layers 0/90 LPI [x] WHIPOX 0/90 open porosity 34% CSI [3] Nicalon/MAS 12 layers, 0/90 40% fiber content fully dense MAS matrix Hot pressing [3] Nicalon SiC/SiC 0/90 32% fiber content 8.6% matrix porosity CVI [2] Nicalon SiC/SiC 0/90 plain weave 35% fiber content residual porosity 10–15% CVI 0 50 100 150 200 250 300 350 400 -0.4 -0.2 0.0 0.2 0.4 Transverse Longitudinal +45°/-45° +15°/-75° +30°/-70° 0°/90° Strain [%] Stress [MPa] -400 -350 -300 -250 -200 -150 -100 -50 0 -0.4 -0.2 0.0 0.2 0.4 Longitudinal Transverse +45°/-45° +15°/-75° +10°/-80° 0°/90° Strain [%] Stress [MPa] ϕ σ σ σ ϕ σ Fig. 5. Typical stress strain curves of WMC composite C/C under (a) tension and (b) compression with different angles between fiber orientation and loading direction. 1168 D. Koch et al. / Composites Science and Technology 68 (2008) 1165–1172

D. Koch et al. Composites Science and Technology 68(2008)1165-1172 Matrix shear failure o Experimental data Fiber buckle failure Fiber tensile failure -100 Fig. 6. Summary of measured strength values and failure modes depending on fiber orientation and loading direction. induced cracks and pores within and between the fiber bun- shear stresses additionally reduce the overall compression dles. Further details are described elsewhere [ll]. trength In+45/45 orientation the specimen fails sim- The mechanical properties were investigated at room ilar to the tensile test with large strain to failure and temperatures under ambient atmosphere. The specimens extended nonlinear stress-strain behavior with various geometries were tested in tension, compres- The results from tension, shear and compression tests sion, and shear modes in a spindle testing machine(Zwick, are summarized in Fig. 6 showing 0I-T12 plane with I Germany)using different angles (0/90, +10/-800, +15/ and 2 representing the fiber orientation in the 2D rein- 75°,+30°/-60°,+45°-45°) between fiber orientation forced material. Depending on load and fiber orientation and loading direction. Strain was measured with strain gauges and with a laser based contactless strain measure- ment system. Complex loaded specimens as DEN (double end notch) coupons are tested in tensile mode in order to Stress- Strain Curve 2500 investigate the influence of stress concentrations on the mechanical behavior of the composites. Amor 2000 hensive description of the tests is found in [4, 9, 11, 13] 1500 5. Experimental results 0cQEu Depending on the angle between fiber orientation and 1000 loading direction significant changes of the stress strain o) 100 curves are observed in tensile as well as in compressive mode(Fig. 5). In 0/90 orientation the material behaves almost linear-elastic up to failure, the fibers that are ori ented in loading direction carry the load. Transversal strain 0.000.050.100.150200.250.300.350 is almost negligible due to the 90 fibers. Failure occurs Axial Strain [% when the fiber strength is reached resulting in large volume damage throughout the total gauge length. With increasing b35 Stress-Strain Curve 6000 angle between fiber orientation and loading direction(off axis loading) strength and Youngs modulus sharply decrease Due to the weak matrix damage occurs already at low stresses resulting in a reduced stifness of the com- 4000 posite. Under off-axis loading shear failure is always observed. The failure processes are not distributed over the whole gauge length but locally restricted. The fracture 2000 surface develops along the fiber axis under similar shear stresses and independent of the ofi-axis angle Under compression mode the material behaves in a sim ilar manner. In 0/90 orientation a linear-elastic behavior is observed, however, the specimens fail at much lower 00000.0020.0040.0060.0080.0100.012 stresses just above 200 MPa which is only half of tensile strength. The weak matrix is not able to prevent fiber buck ling which is also observed at specimens with +10/-80 tensile test and(b) pure shear test with associated acoustic emission and +15/-75 orientation. In these cases superimposed signals

induced cracks and pores within and between the fiber bun￾dles. Further details are described elsewhere [11]. The mechanical properties were investigated at room temperatures under ambient atmosphere. The specimens with various geometries were tested in tension, compres￾sion, and shear modes in a spindle testing machine (Zwick, Germany) using different angles (0/90, +10/80, +15/ 75, +30/60, +45/ 45) between fiber orientation and loading direction. Strain was measured with strain gauges and with a laser based contactless strain measure￾ment system. Complex loaded specimens as DEN (double end notch) coupons are tested in tensile mode in order to investigate the influence of stress concentrations on the mechanical behavior of the composites. Amore compre￾hensive description of the tests is found in [4,9,11,13]. 5. Experimental results Depending on the angle between fiber orientation and loading direction significant changes of the stress strain curves are observed in tensile as well as in compressive mode (Fig. 5). In 0/90 orientation the material behaves almost linear-elastic up to failure, the fibers that are ori￾ented in loading direction carry the load. Transversal strain is almost negligible due to the 90 fibers. Failure occurs when the fiber strength is reached resulting in large volume damage throughout the total gauge length. With increasing angle between fiber orientation and loading direction (off- axis loading) strength and Young’s modulus sharply decrease. Due to the weak matrix damage occurs already at low stresses resulting in a reduced stiffness of the com￾posite. Under off-axis loading shear failure is always observed. The failure processes are not distributed over the whole gauge length but locally restricted. The fracture surface develops along the fiber axis under similar shear stresses and independent of the off-axis angle. Under compression mode the material behaves in a sim￾ilar manner. In 0/90 orientation a linear-elastic behavior is observed, however, the specimens fail at much lower stresses just above 200 MPa which is only half of tensile strength. The weak matrix is not able to prevent fiber buck￾ling which is also observed at specimens with +10/80 and +15/75 orientation. In these cases superimposed shear stresses additionally reduce the overall compression strength. In +45/45 orientation the specimen fails sim￾ilar to the tensile test with large strain to failure and extended nonlinear stress–strain behavior. The results from tension, shear and compression tests are summarized in Fig. 6 showing r1–s12 plane with 1 and 2 representing the fiber orientation in the 2D rein￾forced material. Depending on load and fiber orientation Fig. 6. Summary of measured strength values and failure modes depending on fiber orientation and loading direction. 0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0 50 100 150 200 250 300 350 400 Stress-Strain Curve Stress [MPa] Axial Strain [%] 0 500 1000 1500 2000 2500 Sum of Acoustic Emission Signals Acoustic Emission Signals 0.000 0.002 0.004 0.006 0.008 0.010 0.012 0 5 10 15 20 25 30 35 Stress-Strain Curve Shear Stress [MPa] Shear Strain [%] 0 1000 2000 3000 4000 5000 6000 Acoustic Emission Signals Sum of Acoustic Emission Signals Fig. 7. Stress–strain curves with unloading reloading cycles of (a) on-axis tensile test and (b) pure shear test with associated acoustic emission signals. D. Koch et al. / Composites Science and Technology 68 (2008) 1165–1172 1169