J.Am. Cera.Soc,84141787-94(2001) journal Influence of Strong Fiber/Coating Interfaces on the mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)/SiC Minicomposites Sebastien Bertrand, Rene Pailler, * and Jacques Lamon poratoire des Composites Thermostructuraux, UMR 5801(CNRS, SEP/SNECMA, UBL, CEA), 33600 Pessac. France Hi-Nicalon fiber-reinforced silicon carbide matrix minicom- PyC-based fiber coatings are the most common. Pyc displays a posites (Hi-Nicalon/SiC) with nanoscale multilayered (PyC layered microstructure, and it leads to high-strength-high- SiC)n fiber coatings(also referred to as interphases) have been toughness SiC/SiC composites. PyC, however, is not stable in an manufactured via pressure pulse chemical vapor infiltration oxidizing environment. Consequently, with a view to protect (P-CVI). Fiber/coating interfaces were strengthened by using treated fibers. The microstructures of the interphases as well ered Py C/SiC interphases has emerged as the propagation and deflection of cracks in the interfacial Multilayered PyC/SiC interphases in SiC/SiC composites have region were investigated by SEM and TEM. Interfacial shear been investigated by several authors. ,6., 7, 10-12 Composites with stress was estimated using various methods based on either the multilayered interphases display mechanical properties and life- width of hysteresis loops on unloading-reloading, crack spac- times at high temperatures that compare favorably with those of ing, or fitting of the force-deformation curve using their counterpart with a single-layer fiber coating. micromechanics-based model. Tensile behavior at room tem- The present paper inve tes the tensile mechanical behavior perature and lifetime in static fatigue in air at 700C were and lifetime of SiC/SiC minicomposites containing Hi-Nicalon related to the interphase/interface characteristics. fibers coated with nanoscale(PyC/SiC), multilayered interphases Strong fiber/coating interphases were obtained using treated fi- bers. Hi-Nicalon/SiC minicomposites reinforced with as-received L. Introduction fibers have been examined in a previous paper IL. Experimental ProceduR omposites with strong interfaces have been developed. The strong (1) Processing Conditions and Materials oating/fiber bond was obtained when the fibers had been prev Minicomposites are unidirectional composites reinforced with a ously treated. Features of the mechanical behavior of sic/sic single tow of fibers. The Hi-Nicalon tows consist of 500 filaments opposites with strong fiber/coating interfaces at room and el each having a diameter -13.5 um(+1.5 um). The minicompos vated temperatures have been examined in several papers ites were manufactured using either as-received-or as-treated Experiments as well as models have demonstrated that a strong Hi-Nicalon tows interface is beneficial to strength, toughness, lifetime, and creep he minicomposites have been produced via pressure pulse resistance. -By contrast, weak interfaces are shown to be chemical vapor infiltration(P-CVI). The P-CVI apparatus and the processing conditions were detailed. -The tows were mounted on SiC/SiC composites with strong fiber/coating bonds, SiC frames for deposition of the interphase and SiC matrix.The cracks are deflected within the coating cohesive failure tows were slightly twisted with a constant angle(I turn/5 cm)to short and branched multiple cracks. The associated ecrease the porosity and to increase the fiber volume fraction in short debonds and load transfers allow further cracking of the the minicomposites matrix via a scale effect related to the stressed volume of Two different interphases were deposited on the fibers(table D) uncracked matrix, leading to a higher density of matrix cracks (1)A(PyC/SiC)lo nanoscale multilayered coating consisting Sliding friction within the coating and multiple cracking of the of 10 Py C/SiC sequences. The thickness of each sublayer was matrix increase energy absorption, leading to toughening. Short e( PyC)=20 nm and e(SiC)=50 nm. PyC was deposited first on debonds and improved load transfers limit fiber overloading the fibers during matrix cracking, leading to strengthening of the composi (2)For comparison purposes, a single Pyc layer 100 nn The associated tensile stress-strain curve exhibits a wide curved domain, and the stress at matrix-cracking saturation is close The fiber volume fraction in the minicomposites was 40% ultimate failure. Values of the interfacial shear stress(T (+5%). The main properties of the minicomposite constituents are using various methods on SiC/SiC composites with sted in Table il. Additional data can be found elsewhere 2, based(PyC-based) fiber coating, range between 10 and weak interfaces, whereas they are 100-300 MPa interfaces. 3,4, 6 (2) Microstructural Characterization Surface analysis of the fibers was performed using electron spectroscopy(AES). The fiber/matrix interfacial F. Zok--contributing editor and longitudinal sections using SEM and HRTEM. Preparation of the thin foils was detailed elsewhere Manuscript No. 188914 Received November 29, 1999; approved November 2, Supported by SEP and CNRS through a grant given to S. Bertrand. Proprietary treatment(LCTS-SNECMA)
Influence of Strong Fiber/Coating Interfaces on the Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)n/SiC Minicomposites Sebastien Bertrand, Rene Pailler,* and Jacques Lamon* Laboratoire des Composites Thermostructuraux, UMR 5801 (CNRS, SEP/SNECMA, UB1, CEA), 33600 Pessac, France Hi-Nicalon fiber-reinforced silicon carbide matrix minicomposites (Hi-Nicalon/SiC) with nanoscale multilayered (PyC/ SiC)n fiber coatings (also referred to as interphases) have been manufactured via pressure pulse chemical vapor infiltration (P-CVI). Fiber/coating interfaces were strengthened by using treated fibers. The microstructures of the interphases as well as the propagation and deflection of cracks in the interfacial region were investigated by SEM and TEM. Interfacial shear stress was estimated using various methods based on either the width of hysteresis loops on unloading–reloading, crack spacing, or fitting of the force–deformation curve using a micromechanics-based model. Tensile behavior at room temperature and lifetime in static fatigue in air at 700°C were related to the interphase/interface characteristics. I. Introduction FIBER/MATRIX interfaces in the most advanced ceramic matrix composites consist of a thin coating layer (micrometer-scale) of one or several materials deposited on the fiber. Recently, SiC/SiC composites with strong interfaces have been developed. The strong coating/fiber bond was obtained when the fibers had been previously treated.1–3 Features of the mechanical behavior of SiC/SiC composites with strong fiber/coating interfaces at room and elevated temperatures have been examined in several papers.3–8 Experiments as well as models have demonstrated that a strong interface is beneficial to strength, toughness, lifetime, and creep resistance.3–9 By contrast, weak interfaces are shown to be detrimental. In those SiC/SiC composites with strong fiber/coating bonds, the matrix cracks are deflected within the coating (cohesive failure mode) into short and branched multiple cracks.3,10 The associated short debonds and load transfers allow further cracking of the matrix via a scale effect related to the stressed volume of uncracked matrix,4,5 leading to a higher density of matrix cracks. Sliding friction within the coating and multiple cracking of the matrix increase energy absorption, leading to toughening. Short debonds and improved load transfers limit fiber overloading during matrix cracking, leading to strengthening of the composite. The associated tensile stress–strain curve exhibits a wide curved domain, and the stress at matrix-cracking saturation is close to ultimate failure. Values of the interfacial shear stress (t), measured using various methods on SiC/SiC composites with pyrocarbonbased (PyC-based) fiber coating, range between 10 and 20 MPa for weak interfaces, whereas they are 100–300 MPa for strong interfaces.3,4,6 PyC-based fiber coatings are the most common. PyC displays a layered microstructure, and it leads to high-strength–hightoughness SiC/SiC composites.3 PyC, however, is not stable in an oxidizing environment. Consequently, with a view to protect PyC-based interphases against oxidation, the concept of multilayered PyC/SiC interphases has emerged. Multilayered PyC/SiC interphases in SiC/SiC composites have been investigated by several authors.2,6,7,10–12 Composites with multilayered interphases display mechanical properties and lifetimes at high temperatures that compare favorably with those of their counterpart with a single-layer fiber coating. The present paper investigates the tensile mechanical behavior and lifetime of SiC/SiC minicomposites containing Hi-Nicalon fibers coated with nanoscale (PyC/SiC)n multilayered interphases. Strong fiber/coating interphases were obtained using treated fibers.† Hi-Nicalon/SiC minicomposites reinforced with as-received fibers have been examined in a previous paper.12 II. Experimental Procedures (1) Processing Conditions and Materials Minicomposites are unidirectional composites reinforced with a single tow of fibers. The Hi-Nicalon tows consist of 500 filaments, each having a diameter ;13.5 mm (61.5 mm). The minicomposites were manufactured using either as-received12 or as-treated Hi-Nicalon tows.† The minicomposites have been produced via pressure pulsed chemical vapor infiltration (P-CVI). The P-CVI apparatus and the processing conditions were detailed.12 The tows were mounted on SiC frames for deposition of the interphase and SiC matrix. The tows were slightly twisted with a constant angle (1 turn/5 cm) to decrease the porosity and to increase the fiber volume fraction in the minicomposites. Two different interphases were deposited on the fibers (Table I): (1) A (PyC/SiC)10 nanoscale multilayered coating consisting of 10 PyC/SiC sequences. The thickness of each sublayer was e(PyC) 5 20 nm and e(SiC) 5 50 nm. PyC was deposited first on the fibers. (2) For comparison purposes, a single PyC layer 100 nm thick. The fiber volume fraction in the minicomposites was ;40% (65%). The main properties of the minicomposite constituents are listed in Table II. Additional data can be found elsewhere.12,13 (2) Microstructural Characterization Surface analysis of the fibers was performed using Auger electron spectroscopy (AES). The fiber/matrix interfacial region was examined on failed minicomposites using SEM and on cross and longitudinal sections using SEM and HRTEM. Preparation of the thin foils was detailed elsewhere.14 F. Zok—contributing editor Manuscript No. 188914. Received November 29, 1999; approved November 2, 2000. Supported by SEP and CNRS through a grant given to S. Bertrand. *Member, American Ceramic Society. † Proprietary treatment (LCTS-SNECMA). J. Am. Ceram. Soc., 84 [4] 787–94 (2001) 787 journal
788 Joumal of the American Ceramic Sociery--Bertrand et al. Vol 84. No 4 Table L. Description of Investigated Hi-Nicalon/SiC of 80N was applied using a dead weight that was progressively hung on the bottom grip via the displacement of the support at a constant speed. This force was -10% above the proportional limit (Table Ill). Acoustic emission showed that matrix cracking initi- atches Fibers, Interphase (nm) (nm) n ated below the proportional limit under a force around 50N. The first matrix cracks did not influence specimen compliance. At HN/(C/SIC)1o AR failure of the minicomposites, the chronometer was stopped by the HNT/(C/SIC)Io 50 10 falling dead weight, giving the lifetime HNT/C Il. Results and Discussion Pyc/siC b (1 Material Characterization Details on the microstructure of the nanoscale multilayered Table Il. Properties of Minicomposite Constituents (PyC/SiC)n fiber coatings have been provided. 2 AES depth profiles revealed the presence of an oxygen-enriched layer(15-50 Statistical nm)at the surface of the as-received Hi-Nicalon fibers (Table In This layer consisted of SiO, and free carbon. The first interfacial Constituents Composition (nm) (GPa) m (MPa) PyC sublayer deposited was bonded to this silicon/carbon/oxyge layer. Such sublayers have also been identified in composites Hi-Nicalon fibers Sio 15-502804.26 sponding fiber/matrix interactions were weak, and deflection of the Hi-Nicalon fibers(T) Free carbon 50-100 305 5.8 matrix cracks occurred at the fiber/interphase interfaces. 3, 10 PyC interphase Sic interphase 280 AEs depth profiles performed on treated Hi-Nicalon fibers SiC P-CVI matrix 4005.55.7 show that the surface of the fibers consists of a 50-100 nm thick TA R.= as-received, T. treated. Reference volume, I.=Im layer of free carbon(Fig. 1). The presence of such a superficial layer of free carbon increases the strength of the Py C coating/fiber interface in Nicalon/PyC/SiC composites ,,o The TEM micro- graph of Fig. I shows that the deposited Pyc is perfectly bonded ( Tensile Tests at Room Temperature to the fiber. By contrast, preexisting fiber/coating debonds are Uniaxial tension tests were performed at room temperature at a observed in the minicomposites reinforced with as-received Hi- Nicalon fibers I deformation rate of 50 um/mn using a machine and procedure metallic tubes that were then gripped into the testing machine. (2) Tensile Behavior at Room Temperature Gauge length was 20 mm. Load train compliance, Cs was Figure 2 shows that the force-deformation curves obtained for determined from tensile tests on fiber tows having various gauge HNT/(C/SiC)1o minicomposites reinforced with treated fibers engths(Cs= 0.3 um/N) exhibit typical features associated with strong fiber/matrix Unloading-reloading cycles were conducted on a few speci- bonds: ,3 mens of each batch to estimate the elastic modulus of the cracked (1) A wide curved domain up to ultimate failure attributed to minicomposites, residual strains at zero load, and T. After ultimate a high density of matrix cracks and short debonds, and failure, the test specimens were examined using SEM, and the (2) Rather narrow hysteresis loops, indicative of strong fiber/ crack spacings were measured. matrix interactions on unloading-reloading HN/(C/SiC)o minicomposites reinforced with as-received fi- (4 Static Fatigue Tests at 700C in Air bers display typical features associated with weak fiber/matrix The lifetime of the minicomposites under constant load was bonds, including measured in static fatigue at 700C in air. These temperature (1) A narrow curved domain reflecting longer debonds and a onditions were the worst, because the oxidation rate was high fo lower density of matrix cracks, and PyC and low for SiC. The minicomposite ends were glued within lumina tubes using an alumina-based ceramic adhesive. The gauge length(10 mm) was determined by the distance between the tubes. The tubes were gripped into the testing machine. The ses are used to describe the mechanical behavior of the minicomposites were positioned within the furnace hot zone where the temperature was uniform at 700C. The minicomposites were sites, as a result of the transverse cracks that locally reduce the stressed section to heated to the test temperature before loading. Then a constant force that of fibers Table Ill. Average Mechanical Properties of Minicomposites Tested Saturations FailureR Batches HN/(C/SIC)1o 351 71 0.12 135 0.30 19 0.83 HNT/(C/SiC)Io 0.79 34 750. 0.30 HNT/C 350 0.17 0.7 156 0.81 microcrack spacing in the internal matrix of the minicomposite. 'Microcrack spacing in the surface of the minicomposite. F
(3) Tensile Tests at Room Temperature Uniaxial tension tests were performed at room temperature at a deformation rate of 50 mm/mn using a machine and procedure detailed elsewhere.12 The minicomposite ends were glued within metallic tubes that were then gripped into the testing machine. Gauge length was 20 mm. Load train compliance, Cs, was determined from tensile tests on fiber tows having various gauge lengths (Cs 5 0.3 mm/N). Unloading–reloading cycles were conducted on a few specimens of each batch to estimate the elastic modulus of the cracked minicomposites, residual strains at zero load, and t. After ultimate failure, the test specimens were examined using SEM, and the crack spacings were measured. (4) Static Fatigue Tests at 700°C in Air The lifetime of the minicomposites under constant load was measured in static fatigue at 700°C in air. These temperature conditions were the worst, because the oxidation rate was high for PyC and low for SiC. The minicomposite ends were glued within alumina tubes using an alumina-based ceramic adhesive. The gauge length (10 mm) was determined by the distance between the tubes. The tubes were gripped into the testing machine. The minicomposites were positioned within the furnace hot zone where the temperature was uniform at 700°C. The minicomposites were heated to the test temperature before loading. Then a constant force of 80 N was applied using a dead weight that was progressively hung on the bottom grip via the displacement of the support at a constant speed. This force was ;10% above the proportional limit (Table III). Acoustic emission showed that matrix cracking initiated below the proportional limit under a force around 50 N.5 The first matrix cracks did not influence specimen compliance. At failure of the minicomposites, the chronometer was stopped by the falling dead weight, giving the lifetime. III. Results and Discussion (1) Material Characterization Details on the microstructure of the nanoscale multilayered (PyC/SiC)n fiber coatings have been provided.12 AES depth profiles revealed the presence of an oxygen-enriched layer (15–50 nm) at the surface of the as-received Hi-Nicalon fibers (Table II). This layer consisted of SiO2 and free carbon. The first interfacial PyC sublayer deposited was bonded to this silicon/carbon/oxygen layer. Such sublayers have also been identified in composites reinforced with as-received Nicalon fibers (NL 202). The corresponding fiber/matrix interactions were weak, and deflection of the matrix cracks occurred at the fiber/interphase interfaces.1,3,10 AES depth profiles performed on treated Hi-Nicalon fibers show that the surface of the fibers consists of a 50–100 nm thick layer of free carbon (Fig. 1). The presence of such a superficial layer of free carbon increases the strength of the PyC coating/fiber interface in Nicalon/PyC/SiC composites.1,3,10 The TEM micrograph of Fig. 1 shows that the deposited PyC is perfectly bonded to the fiber. By contrast, preexisting fiber/coating debonds are observed in the minicomposites reinforced with as-received HiNicalon fibers.12 (2) Tensile Behavior at Room Temperature‡ Figure 2 shows that the force–deformation curves obtained for HNT/(C/SiC)10 minicomposites reinforced with treated fibers exhibit typical features associated with strong fiber/matrix bonds:1,3 (1) A wide curved domain up to ultimate failure attributed to a high density of matrix cracks and short debonds, and (2) Rather narrow hysteresis loops, indicative of strong fiber/ matrix interactions on unloading–reloading. HN/(C/SiC)10 minicomposites reinforced with as-received fibers display typical features associated with weak fiber/matrix bonds, including (1) A narrow curved domain reflecting longer debonds and a lower density of matrix cracks, and ‡ Forces instead of stresses are used to describe the mechanical behavior of the minicomposites, because derivation of a stress from the applied force is not straightforward or appropriate. The stress state is not uniform within the minicomposites, as a result of the transverse cracks that locally reduce the stressed section to that of fibers. Table I. Description of Investigated Hi-Nicalon/SiC Minicomposites Batches Fibers† Interphase Interphase characteristics e(PyC)‡ (nm) e(SiC)‡ (nm) n§ HN/(C/SiC)10 A.R. (20/50)10 20 50 10 HNT/(C/SiC)10 T. (20/50)10 20 50 10 HN/C A.R. (100/0)1 100 0 1 HNT/C T. (100/0)1 100 0 1 † A.R. 5 as-received, T. 5 treated. ‡ Thickness per sublayer. § Number of PyC/SiC bilayers. Table II. Properties of Minicomposite Constituents12,13 Constituents† Superficial layer Young’s Modulus (GPa) Statistical parameters Composition Thickness (nm) m so ‡ (MPa) Hi-Nicalon fibers (A.R.) SiO2 15–50 280 4.2 6 Hi-Nicalon fibers (T.) Free carbon 50–100 305 5.8 26 PyC interphase 12–80 SiC interphase 280 SiC P-CVI matrix 400 5.5 5.7 † A.R. 5 as-received, T. 5 treated. ‡ Reference volume, Vo 5 1 m3 . Table III. Average Mechanical Properties of Minicomposites Tested Batches Ec (GPa) Proportional limit§ Saturation§ Failure§ ls (mm) FE (N) εE (%) FS (N) εs (%) FR (N) εR (%) HN/(C/SiC)10 351 71 0.12 135 0.30 195 0.83 50† 175‡ HNT/(C/SiC)10 360 73 0.07 ;FR ;εR 183 0.79 35† 67‡ HN/C 342 75 0.09 100 0.30 171 0.69 40† 95‡ HNT/C 350 75 0.17 140 0.7 156 0.81 30† 100‡ † Microcrack spacing in the internal matrix of the minicomposite. ‡ Microcrack spacing in the surface of the minicomposite. § F 5 force; ε 5 deformation. 788 Journal of the American Ceramic Society—Bertrand et al. Vol. 84, No. 4
April 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC) Sic 789 0.4 Ef Vf/Eo Fig 3. Relative elastic modulus versus applied E=elastic HN/C; D= HNT/C. eA= HN/(C/SiC)o; B uncracked minicom SiC)1o: C 010020030040050060070 indicates that the applied load is borne only by the fibers Therefore, fiber debonding is complete at this stage, and matrix cracking saturation has occurred. This minimum is not reached by 是100nm the modulus of HNT/(C/SiC)o minicomposites, indicating that saturation of matrix cracking has not occurred at ultimate failure. Figure 3 clearly shows that the behavior is significantly influ- Fig 1. TEM micrograph of fiber/coating interfacial region( batch C)and enced by the presence of treated fibers. Those minicomposites AES depth profile of treated Hi-Nicalon fibers. Pyc 1 shows deposited reinforced with as-received fibers display a significantly steep ublayer of PyC, "isotropic carbon"and "anisotropic carbon"indicate fiber modulus decrease. The minimum Err is observed at deformations superficial layer of free carbon. of -0 2%0.3%, which correspond to the strain at saturation indicated by the end of the curved domain of the force deforma tions curves(Fig. 2). For the HN/(C/SiC)o minicomposite, the minimum modulus becomes smaller than Ere This may be tainties in the data, including the modulus measurements By contrast, the minicomposites reinforced with treated fibers experience a gradual modulus decrease. The minimum is reached r larger deformations (20.7%), indicating a high strain at These trends suggest that debonding was more significant in those minicomposites reinforced with as-received fibers, which Strain(‰) implies the presence of weaker fiber/matrix bonds Fig. 2. Tensile force-deformation curves for Hi-Nicalon/SiC minicom- ( Matrix Cracking and Crack Deflection sites investigated in present paper. A= HN(C/SIC)o, B= HNT/(C/ SiC):C= HNC, D= HNT/C The crack spacing distance measured for the transverse crack (7)is always shorter in the internal matrix(Fig. 4). This effect seems to be related to tow g and may be attributed to the contribution of the radial compressive stress components that (2)Wide hysteresis loops, reflecting weaker fiber/matrix in teractions on unloading-reloading The above features can also be noticed on the force-deforma- tion curves obtained for those minicomposites with single-layer Pyc fiber coatings. However, a certain discrepancy is noticed in the force-deformation curves of some HNT/C minicomposites reinforced with treated fibers. The curved domain seems to be narrower, and the hysteresis loops seem to be wider than expected The elastic modulus pertinent to the cracked minicomposites is derived from the slope of the linear portion of the reloading curve a minimum tangent modulus). The tangent to this linear portion intercepts the origin in most cases. For the HNT/C minicompos- 80 ites, it intercepts the abscissa on the negative side, suggesting that the fibers tend to contract as a result of the presence of tensile 120 thermally induced residual stresses in the fibers. The permanent train at zero load includes contributions from misfit relief an sliding. The larger permanent elongations at zero load exhibited by Radial position(arbitrary unit those minicomposites with as-received fibers(Fig. 2) suggest the presence of weaker fiber/matrix interactions when compared with minicomposites reinforced with treated fibers. Figure 3 shows the typical trends in the elastic modulus during Surface of the Center of the Surface of the the tensile tests. For most minicomposites, the modulus decreas minicomposite minicomposite minicomposite a minimum value that coincides with the quantity Ep r(Er is the Fig. 4. Example of distribution of Is spacing distances measured at fiber Youngs modulus, and Ve the fiber volume fraction), which various locations in matrix of
(2) Wide hysteresis loops, reflecting weaker fiber/matrix interactions on unloading–reloading. The above features can also be noticed on the force–deformation curves obtained for those minicomposites with single-layer PyC fiber coatings. However, a certain discrepancy is noticed in the force–deformation curves of some HNT/C minicomposites reinforced with treated fibers. The curved domain seems to be narrower, and the hysteresis loops seem to be wider than expected. The elastic modulus pertinent to the cracked minicomposites is derived from the slope of the linear portion of the reloading curve (minimum tangent modulus).12 The tangent to this linear portion intercepts the origin in most cases. For the HNT/C minicomposites, it intercepts the abscissa on the negative side, suggesting that the fibers tend to contract as a result of the presence of tensile thermally induced residual stresses in the fibers. The permanent strain at zero load includes contributions from misfit relief and sliding. The larger permanent elongations at zero load exhibited by those minicomposites with as-received fibers (Fig. 2) suggest the presence of weaker fiber/matrix interactions when compared with minicomposites reinforced with treated fibers. Figure 3 shows the typical trends in the elastic modulus during the tensile tests. For most minicomposites, the modulus decreases to a minimum value that coincides with the quantity Ef Vf (Ef is the fiber Young’s modulus, and Vf the fiber volume fraction), which indicates that the applied load is borne only by the fibers. Therefore, fiber debonding is complete at this stage, and matrixcracking saturation has occurred. This minimum is not reached by the modulus of HNT/(C/SiC)10 minicomposites, indicating that saturation of matrix cracking has not occurred at ultimate failure. Figure 3 clearly shows that the behavior is significantly influenced by the presence of treated fibers. Those minicomposites reinforced with as-received fibers display a significantly steep modulus decrease. The minimum Ef Vf is observed at deformations of ;0.2%–0.3%, which correspond to the strain at saturation indicated by the end of the curved domain of the force deformations curves (Fig. 2). For the HN/(C/SiC)10 minicomposite, the minimum modulus becomes smaller than Ef Vf . This may be attributed to the presence of broken or bent fibers, and/or uncertainties in the data, including the modulus measurements. By contrast, the minicomposites reinforced with treated fibers experience a gradual modulus decrease. The minimum is reached for larger deformations ($0.7%), indicating a high strain at saturation. These trends suggest that debonding was more significant in those minicomposites reinforced with as-received fibers, which implies the presence of weaker fiber/matrix bonds. (3) Matrix Cracking and Crack Deflection The crack spacing distance measured for the transverse crack (ls) is always shorter in the internal matrix (Fig. 4). This effect seems to be related to tow twisting and may be attributed to the contribution of the radial compressive stress components that Fig. 1. TEM micrograph of fiber/coating interfacial region (batch C) and AES depth profile of treated Hi-Nicalon fibers. “PyC 1” shows deposited sublayer of PyC; “isotropic carbon” and “anisotropic carbon” indicate fiber superficial layer of free carbon. Fig. 2. Tensile force–deformation curves for Hi-Nicalon/SiC minicomposites investigated in present paper. A 5 HN/(C/SiC)10; B 5 HNT/(C/ SiC)10; C 5 HN/C; D 5 HNT/C. Fig. 3. Relative elastic modulus versus applied deformation: E 5 elastic modulus given by minimum tangent modulus, E0 5 elastic modulus of uncracked minicomposite. A 5 HN/(C/SiC)10; B 5 HNT/(C/SiC)10; C 5 HN/C; D 5 HNT/C. Fig. 4. Example of distribution of ls spacing distances measured at various locations in matrix of minicomposite (minicomposite HN/(C/SiC)10). April 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)n/SiC 789
790 Journal of the American Ceramic Society-Bertrand et al Vol 84. No 4 increase fiber/matrix interactions. These stresses are induced by the curved fibers that try to stretch under a tensile load The I data indicate that the fiber/matrix bond is strengthened in the minicomposites reinforced with treated fibers. Thus, the Is 1=E measured in the HNT/(C/SiC)Io minicomposites is significantl smaller than that in their counterpart with as-received fibers (HN/(C/SiC)o). Although the Is measured in the internal matrix b2 _(1+v)EE+(1-2)E falls within the same range for most minicomposites, that obtained for the minicomposites reinforced with treated fibers is the shortest where o is the applied stress in the unloading-reloading sequence (Table I) that corresponds to 84, o, the initial stress level at unloading, E Matrix cracking essentially involved transverse cracks the Youngs modulus of the minicomposite, R the fiber radius, v over, a few longitudinal cracks also were detected(Fig. 5) the Poisson's ratio (v=vm= v, Em the Youngs modulus of the cracks were not identified on the specimens inspected before matrix, Er that of fiber, and ve the fiber volume fraction. T was testing. They were created during the tensile tests derived from the SA-o data measured during the last unloading- Unlike the minicomposites with as-received fibers in which reloading sequence before ultimate failure of the minicomposites ignificant fiber/coating debonds were observed before testing, The number of matrix cracks in gauge length, N, was determined those minicomposites with treated fibers did not exhibit such from SEM inspection of the minicomposites after failure preexisting interface features. (2) The spacing distance of the matrix cracks at saturation. Figure 6 illustrates the significant differences observed in the Estimates of T are given by the following equations deflection of matrix cracks, depending on the fiber/coating bond In those minicomposites reinforced with as-received fibers, deflec- sRe tion of the matrix cracks occurs in the fiber/interphase interface and also in the matrix/interphase interface(Fig. 6). It is worth 24(1+ pointing out that the crack-opening displacement is rather large (0. 2 um). In those minicomposites reinforced with treated fibers, rrl matrix cracks are deflected within the coating in the PyC sublayers (Fig. 6). Crack branching can also be noticed on Fig. 6. Finally, the crack reaches the PyC sublayer bonded to the fiber surface. Unlike where o, is the applied stress at matrix-cracking saturation. in minicomposites with as-received fibers, the crack-openin 3) The force-deformation curves. A curve fitting procedur displacement is now very small (-10 nm been detailed and validated in previous works. s This model has based on a model of the tensile behavior was used e T adj (4 Interfacial Shear Stress (T) ment involved predictions of the force-deformation curves(Fig. results of tensile tests and crack examination, estimates of the T ters given in Table t properties and the flaw-strength parame- To assess the trends that have been identified based 7)from the cons were extracted from various data provided by the tensile tests The T estimates given in Table Iv evidence an unquestionable (1) The width of the hysteresis loops(8A)measured durin strengthening of the fiber/coating bond associated with the treated equation(established elsewhere for microcomposites 2. O/lowing fibers. The different methods agree satisfactorily in indicating this trend, despite certain discrepancies that can be noticed. The scatter in T measurements previously noticed does not affect this trend biN(I-a1Ve-Ro Therefore the results can be summarized as follows .t data are 2:E ( closer to 200 MPa for minicomposites reinforced with treated fibers, and they are close to 100 MPa for those reinforced with as-received fibers These trends compare satisfactorily with those established on Nicalon/SiC composites. This agreement can be attributed to the presence of a silicon/carbon/oxygen layer at the surface of as- received Nicalon and Hi-Nicalon fibers, and a free carbon layer at the surface of treated Nicalon and Hi-Nicalon fibers. which has been shown to dictate the fiber/coating bond (5) Lifetime in Static Fatigue at 700C n did not fail after 140 and 200 h. Further comparison of the data showed that the multilayered coating improved lifetime: from 2 to 20 h for those minicomposites reinforced with as-received fibers The influence might appear to be less clear for those minicompos ites reinforced with treated fibers because of scatter in the data 1.5pm for the HNT/(C/SiC)o minicomposites However, some HNT/ (C/SiC)o minicomposites were not broken after 200 and 140 h, was 114 h. This trend agreed with previous results reported elsewhere. 7, 8 The results suggest an effect of the fiber/coating bond. Strong nIcon matrix cracks. SEM images(Fig. 9)of the interfacial regions after the static fatigue tests show that PyC layers have disappeared Fig. 5. SEM images of longitudinal cracks detected in Hi-Nicalon/SiC However, some interesting differences can be noticed, depending minicomposite after tensile tests. on the batch, that can be related to the location of the debond crack
increase fiber/matrix interactions. These stresses are induced by the curved fibers that try to stretch under a tensile load. The ls data indicate that the fiber/matrix bond is strengthened in the minicomposites reinforced with treated fibers. Thus, the ls measured in the HNT/(C/SiC)10 minicomposites is significantly smaller than that in their counterpart with as-received fibers (HN/(C/SiC)10). Although the ls measured in the internal matrix falls within the same range for most minicomposites, that obtained for the minicomposites reinforced with treated fibers is the shortest (Table III). Matrix cracking essentially involved transverse cracks. Moreover, a few longitudinal cracks also were detected (Fig. 5). Such cracks were not identified on the specimens inspected before testing. They were created during the tensile tests. Unlike the minicomposites with as-received fibers in which significant fiber/coating debonds were observed before testing, those minicomposites with treated fibers did not exhibit such preexisting interface features. Figure 6 illustrates the significant differences observed in the deflection of matrix cracks, depending on the fiber/coating bond. In those minicomposites reinforced with as-received fibers, deflection of the matrix cracks occurs in the fiber/interphase interface and also in the matrix/interphase interface (Fig. 6). It is worth pointing out that the crack-opening displacement is rather large (0.2 mm). In those minicomposites reinforced with treated fibers, matrix cracks are deflected within the coating in the PyC sublayers (Fig. 6). Crack branching can also be noticed on Fig. 6. Finally, the crack reaches the PyC sublayer bonded to the fiber surface. Unlike in minicomposites with as-received fibers, the crack-opening displacement is now very small (;10 nm). (4) Interfacial Shear Stress (t) To assess the trends that have been identified based on the results of tensile tests and crack examination, estimates of the t were extracted from various data provided by the tensile tests: (1) The width of the hysteresis loops (dD) measured during unloading–reloading cycles, which is related to t by the following equation (established elsewhere for microcomposites15): t 5 b2N~1 2 a1Vf! 2 Rf 2V f 2 Em S sp 2 dDDS s sp DS1 2 s sp D (1) with a1 5 Ef Ec (2) b2 5 ~1 1 n! Em@Ef 1 ~1 2 2n!Ec# Ef@~1 1 n!Ef 1 ~1 2 n!Ec# where s is the applied stress in the unloading–reloading sequence that corresponds to dD, sp the initial stress level at unloading, Ec the Young’s modulus of the minicomposite, Rf the fiber radius, n the Poisson’s ratio (n5nm 5 nf ), Em the Young’s modulus of the matrix, Ef that of fiber, and Vf the fiber volume fraction. t was derived from the dD–s data measured during the last unloading– reloading sequence before ultimate failure of the minicomposites. The number of matrix cracks in gauge length, N, was determined from SEM inspection of the minicomposites after failure. (2) The spacing distance of the matrix cracks at saturation. Estimates of t are given by the following equations:16,17 t 5 ssRf 2VflsS1 1 EfVf EmVm D (3) t 5 ssRfVm 2Vfls (4) where ss is the applied stress at matrix-cracking saturation. (3) The force–deformation curves. A curve fitting procedure based on a model of the tensile behavior was used. This model has been detailed and validated in previous works.5,12 The t adjustment involved predictions of the force–deformation curves (Fig. 7) from the constituent properties and the flaw–strength parameters given in Table II. The t estimates given in Table IV evidence an unquestionable strengthening of the fiber/coating bond associated with the treated fibers. The different methods agree satisfactorily in indicating this trend, despite certain discrepancies that can be noticed. The scatter in t measurements previously noticed12 does not affect this trend. Therefore, the results can be summarized as follows: t data are closer to 200 MPa for minicomposites reinforced with treated fibers, and they are close to 100 MPa for those reinforced with as-received fibers. These trends compare satisfactorily with those established on Nicalon/SiC composites.3 This agreement can be attributed to the presence of a silicon/carbon/oxygen layer at the surface of asreceived Nicalon and Hi-Nicalon fibers, and a free carbon layer at the surface of treated Nicalon and Hi-Nicalon fibers, which has been shown to dictate the fiber/coating bond. (5) Lifetime in Static Fatigue at 700°C The lifetime data are shown in graphical form in Fig. 8. The lifetimes obtained for those minicomposites reinforced with treated fibers were unambiguously the longest. Some specimens did not fail after 140 and 200 h. Further comparison of the data showed that the multilayered coating improved lifetime: from 2 to ;20 h for those minicomposites reinforced with as-received fibers. The influence might appear to be less clear for those minicomposites reinforced with treated fibers because of scatter in the data for the HNT/(C/SiC)10 minicomposites. However, some HNT/ (C/SiC)10 minicomposites were not broken after 200 and 140 h, whereas the maximum lifetime of those HNT/C minicomposites was 114 h. This trend agreed with previous results reported elsewhere.7,8 The results suggest an effect of the fiber/coating bond. Strong bonds have been shown to reduce debonding and crack-opening displacement that limit the amount of oxygen migrating within the matrix cracks. SEM images (Fig. 9) of the interfacial regions after the static fatigue tests show that PyC layers have disappeared. However, some interesting differences can be noticed, depending on the batch, that can be related to the location of the debond crack Fig. 5. SEM images of longitudinal cracks detected in Hi-Nicalon/SiC minicomposite after tensile tests. 790 Journal of the American Ceramic Society—Bertrand et al. Vol. 84, No. 4
pril 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC) Sic 791 M (Fiber surface Sic 1 um (Fiber surface) Fiber Fibe Fig icrographs showing deflection of matrix microcracks in interfacial region of minicomposites: (a) SEM micrograph of a HN/(C/SIC)o minicomposite,(b)TEM micrograph of a HNT/(C/SiC)o minicomposite in the interfacial region. In the HN/(C/SiC)o minicomposites, on (a)250pT「·T·冂 the Pyc sublayer that lay on the fibers has been eliminated whereas, in the HNT/(C/SiC)o minicomposites, all the PyC layers exp in which matrix cracks are found to be deflected are affected V. Discussion 2150 It has been confirmed again that multilayered interphases are not detrimental to the mechanical behavior of sic/sic minicom- 100 posites, although these contain stiff sublayers of SiC and very thin (nanometer-scale)PyC sublayers The presence of rather strong fiber/coating interfaces in those minicomposites reinforced with treated Hi-Nicalon fibers has been Batcha revealed by a set of data, including features of the force deformation curves. location of the debond crack in the interfacial 002040.60.8 region, composition of the fiber surface, and estimates of T. Deformation (%) The force-deformation curves pertinent to those minicompos- ites reinforced with treated fibers exhibited the features previousl observed on Nicalon/SiC minicomposites and two-dimensional woven composites reinforced with treated fibers(NLM 202) including(i) a wide curved domain, (ii)a large stress/strain at saturation of matrix cracking that coincides with ultimate failures, (b)250 (iii) small residual deformations at zero load, and (iv) a gentle modulus decrease during tensile tests. The matrix cracks were deflected within each Pyc sublayer of the interphase in those minicomposites reinforced with treated fibers and at the fiber surface in those minicomposites reinforced 150 experiment vith as-received fibers. This debond patten was similar to that observed on Nicalon/SiC composites. 10 Deviation within the interphase resulted from the presence of a stronger bond between 100 the interphase and the fib In those minicomposites reinforced with as-received fibers, weakening of the fiber/interphase region has been evidenced usin SEM. The preexisting debond cracks were attributed to the lateral contraction of as-received Hi-Nicalon fibers during minicomposite Batch C rocessing. Yun et al. have shown that the lateral contraction of as-received Hi-Nicalon fibers during processing of composites can 0 xceed 1% at 1200C. The contraction was associated with fiber Deformation(%) shrinkage and with the transformation of the Sic amorphous phase nto a B-sic crystallized phase. Fig. 7. Comparison of predicted and experimental force deformation curves for SiC/SiC minicomposites(a) HN/(C/SiC)o and(b)HNT/(C/SiC)o The interfacial bond was characterized by the interfacial shear tress (T). As previously mentioned, all the methods indicated tha the fiber/coating bond was stronger in those minicomposites fibers and T > 200 MPa for those reinforced with treated fibers reinforced with treated fibers. However, the range of t data was In study, T 100 MPa for those Hi-Nicalon/SiC arrower than that observed on Nicalon/Sic two-dimensional InIc tes reinforced with as-received fibers. and T s 200 woven composites. In those Nicalon/SiC composites, the respe MPa e reinforced with treated fibers. A contribution of tive t values were different by more than I order of magnitude: twisting could be expected. Twisting generated radial compressive T 10 MPa for those composites reinforced with as-received stresses that enhanced fiber/matrix interactions and can tend to
in the interfacial region. In the HN/(C/SiC)10 minicomposites, only the PyC sublayer that lay on the fibers has been eliminated; whereas, in the HNT/(C/SiC)10 minicomposites, all the PyC layers in which matrix cracks are found to be deflected are affected. IV. Discussion It has been confirmed again that multilayered interphases are not detrimental to the mechanical behavior of SiC/SiC minicomposites, although these contain stiff sublayers of SiC and very thin (nanometer-scale) PyC sublayers. The presence of rather strong fiber/coating interfaces in those minicomposites reinforced with treated Hi-Nicalon fibers has been revealed by a set of data, including features of the force– deformation curves, location of the debond crack in the interfacial region, composition of the fiber surface, and estimates of t. The force–deformation curves pertinent to those minicomposites reinforced with treated fibers exhibited the features previously observed on Nicalon/SiC minicomposites and two-dimensional woven composites reinforced with treated fibers (NLM 202) including (i) a wide curved domain, (ii) a large stress/strain at saturation of matrix cracking that coincides with ultimate failures, (iii) small residual deformations at zero load, and (iv) a gentle modulus decrease during tensile tests. The matrix cracks were deflected within each PyC sublayer of the interphase in those minicomposites reinforced with treated fibers and at the fiber surface in those minicomposites reinforced with as-received fibers. This debond pattern was similar to that observed on Nicalon/SiC composites.3,10 Deviation within the interphase resulted from the presence of a stronger bond between the interphase and the fiber.6 In those minicomposites reinforced with as-received fibers, weakening of the fiber/interphase region has been evidenced using SEM. The preexisting debond cracks were attributed to the lateral contraction of as-received Hi-Nicalon fibers during minicomposite processing. Yun et al.20 have shown that the lateral contraction of as-received Hi-Nicalon fibers during processing of composites can exceed 1% at 1200°C. The contraction was associated with fiber shrinkage and with the transformation of the SiC amorphous phase into a b-SiC crystallized phase.21 The interfacial bond was characterized by the interfacial shear stress (t). As previously mentioned, all the methods indicated that the fiber/coating bond was stronger in those minicomposites reinforced with treated fibers. However, the range of t data was narrower than that observed on Nicalon/SiC two-dimensional woven composites. In those Nicalon/SiC composites, the respective t values were different by more than 1 order of magnitude: t ' 10 MPa for those composites reinforced with as-received fibers and t . 200 MPa for those reinforced with treated fibers.3 In the present study, t ' 100 MPa for those Hi-Nicalon/SiC minicomposites reinforced with as-received fibers, and t ' 200 MPa for those reinforced with treated fibers. A contribution of twisting could be expected. Twisting generated radial compressive stresses that enhanced fiber/matrix interactions and can tend to Fig. 7. Comparison of predicted and experimental force deformation curves for SiC/SiC minicomposites (a) HN/(C/SiC)10 and (b) HNT/(C/SiC)10. Fig. 6. Micrographs showing deflection of matrix microcracks in interfacial region of minicomposites: (a) SEM micrograph of a HN/(C/SiC)10 minicomposite, (b) TEM micrograph of a HNT/(C/SiC)10 minicomposite. April 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)n/SiC 791
792 Journal of the American Ceramic Society-Bertrand et al Vol 84. No 4 Table I. Interfacial Shear Stresses(T) Estimated Using Various Methods and Tow Youngs Modulus (M Batch Location HN/(C/SIC)1o Interior 105 78 HNT/(C/SIC)Io Interior Surface HNT/C Interior 150 trends have been observed on the Hi-Nicalon/SiC minicomposites 300 of the present study. The Hi-Nicalon/SiC minicomposites rein- forced with as-received fibers possessed the weaker fiber/matrix bonds. The interphase was deposited on a SiO,/PyC layer, which 250 is now well-known to lead to a weak fiber/coating interface. 3 Those Hi-Nicalon/SiC minicomposites reinforced with treated fibers appeared to exhibit features reflecting strong fiber/matrix 国200 bonds. The surface of the fibers was composed of an enriched carbon layer 50-100 nm thick, has been shown to give strong fiber/coating interfaces 150 Furthermore, the degree of organization of the carbon micro- structure also contributed to bonding. Droillard observed on two-dimensional Nicalon/SiC composites reinforced with as- received NL202 fibers that the PyC coating consisted of layers parallel to the fiber axis, whereas the PyC microstructure was rather disorganized in those composites reinforced with treated 50 Nicalon fibers. Similar observations were made on Hi-Nicalon minicomposites elsewhereas well as in the present work. In the minicomposites reinforced with treated fibers, the PyC sublayer exhibited a disordered microstructure (as shown in Fig. 2), an there was no discontinuity or interface between the fiber and the deposited PyC sublayer, which logically gave a strong fiber/ Its clearly indicate that longer in those minicomposites reinforced with treated fibers. This effect may be related to the presence of a stronger fiber/coating interface that limits debonding and crack-opening displacement. at 700C in air. Also given are lifetime data for Hi-Nicalon/BN/SIC The oxidized layers shown by Fig. 9 correspond exactly to the inicomposites18, 19 Arrows indicate those minicomposites that did not path of the deflected matrix cracks. In the minicomposites rein- forced with as-received fibers. the matrix cracks are deflected at the fiber/coating interface. Oxygen can diffuse instantaneously at the fiber surface and react with carbon in the sublayer on the fiber Increase these stresses in the minicomposites reinforced with and with the fiber itself. as-received fibers. This effect was supported by measurements of In those minicomposites reinforced with treated fibers, th the crack spacing distance that was smaller in the internal matrix matrix cracks are deflected first in the Pyc layers within th It could explain the large magnitude of T values determined on the coating, and finally in the PyC layer at the fiber surface. Oxygen ninicomposites reinforced with as-received fibers(90 MPa), reacts first with carbon in those sublayers within the multilayered although significant initial fiber/interphase debonding had been coating and ultimately with the carbon in the layer at the fiber surface. The oxygen propagation path is thus increased, whereas The T estimated for the minicomposites reinforced with treated the amount of oxygen is reduced, owing to the slight opening of fibers compared satisfactorily with those measured on twe the cracks. Thi yer ontributes to the dimensional woven Hi-Nicalon/(Py C/SiC)/SiC composites rein- forced with treated fibers for which t was -230 MPa. 8 treated fibers. Comparison of HNT/(C/SiC)o and HNT/C mini- The force-deformation curves predicted from the constituent composites indicates that this contribution is significant. However, properties and the flaw-strength statistical parameters(Table Il)fit it is worth emphasizing that this contribution is associated with the quite well the experimental ones(Fig. 7), which strongly supports presence of rather strong fiber/coating interfaces that determine the pertinence of the materials data and the model. The adjusted deviation of the matrix cracks, debonding, and crack opening. arameters agreed with those extracted using the other methods At high temperature, because of thermal expansion of the (Table IV). However, fitting was improved when tow Youngs composite constituents, the debond cracks and the matrix cracks moduli slightly larger than those measured using tensile tests on tend to close, or, at least, the crack-opening displacement tends to ingle filaments (Table IV) were introduced into the computations. decrease. Therefore, the effects associated with strong fiber This slight increase could be considered to reflect the tow coating interfaces that have been identified at room temperature stiffening associated with twisting. should be enhanced or. at least unaffected. As oxidation of the The composition of the superficial fiber layer has been shown to Py c sublayers proceeds, the influence of the above-mentioned determine the strength of the interphase/fiber bond in the Nicalon effects may decrease. The rate of this phenomenon is not known at (NL 202)SiC composites with Py C-based interphases. Simila this stage. However, comparison of either HN/C and HNT/C, or
increase these stresses in the minicomposites reinforced with as-received fibers. This effect was supported by measurements of the crack spacing distance that was smaller in the internal matrix. It could explain the large magnitude of t values determined on the minicomposites reinforced with as-received fibers (;90 MPa), although significant initial fiber/interphase debonding had been observed.12 The t estimated for the minicomposites reinforced with treated fibers compared satisfactorily with those measured on twodimensional woven Hi-Nicalon/(PyC/SiC)n/SiC composites reinforced with treated fibers for which t was ;230 MPa.8 The force–deformation curves predicted from the constituent properties and the flaw–strength statistical parameters (Table II) fit quite well the experimental ones (Fig. 7), which strongly supports the pertinence of the materials data and the model. The adjusted t parameters agreed with those extracted using the other methods (Table IV). However, fitting was improved when tow Young’s moduli slightly larger than those measured using tensile tests on single filaments (Table IV) were introduced into the computations. This slight increase could be considered to reflect the tow stiffening associated with twisting. The composition of the superficial fiber layer has been shown to determine the strength of the interphase/fiber bond in the Nicalon (NL 202)/SiC composites with PyC-based interphases.1,3 Similar trends have been observed on the Hi-Nicalon/SiC minicomposites of the present study. The Hi-Nicalon/SiC minicomposites reinforced with as-received fibers possessed the weaker fiber/matrix bonds. The interphase was deposited on a SiO2/PyC layer, which is now well-known to lead to a weak fiber/coating interface.1,3 Those Hi-Nicalon/SiC minicomposites reinforced with treated fibers appeared to exhibit features reflecting strong fiber/matrix bonds. The surface of the fibers was composed of an enriched carbon layer ;50–100 nm thick, which has been shown to give strong fiber/coating interfaces.1,3 Furthermore, the degree of organization of the carbon microstructure also contributed to bonding. Droillard2 observed on two-dimensional Nicalon/SiC composites reinforced with asreceived NL202 fibers that the PyC coating consisted of layers parallel to the fiber axis, whereas the PyC microstructure was rather disorganized in those composites reinforced with treated Nicalon fibers. Similar observations were made on Hi-Nicalon minicomposites elsewhere12 as well as in the present work. In the minicomposites reinforced with treated fibers, the PyC sublayer exhibited a disordered microstructure (as shown in Fig. 2), and there was no discontinuity or interface between the fiber and the deposited PyC sublayer, which logically gave a strong fiber/ coating bond. The experimental results clearly indicate that the lifetime is longer in those minicomposites reinforced with treated fibers. This effect may be related to the presence of a stronger fiber/coating interface that limits debonding and crack-opening displacement. The oxidized layers shown by Fig. 9 correspond exactly to the path of the deflected matrix cracks. In the minicomposites reinforced with as-received fibers, the matrix cracks are deflected at the fiber/coating interface. Oxygen can diffuse instantaneously at the fiber surface and react with carbon in the sublayer on the fiber and with the fiber itself. In those minicomposites reinforced with treated fibers, the matrix cracks are deflected first in the PyC layers within the coating, and finally in the PyC layer at the fiber surface. Oxygen reacts first with carbon in those sublayers within the multilayered coating and ultimately with the carbon in the layer at the fiber surface. The oxygen propagation path is thus increased, whereas the amount of oxygen is reduced, owing to the slight opening of the cracks. Thus, the multilayered interphase contributes to the lifetime improvement in those minicomposites reinforced with treated fibers. Comparison of HNT/(C/SiC)10 and HNT/C minicomposites indicates that this contribution is significant. However, it is worth emphasizing that this contribution is associated with the presence of rather strong fiber/coating interfaces that determine deviation of the matrix cracks, debonding, and crack opening. At high temperature, because of thermal expansion of the composite constituents, the debond cracks and the matrix cracks tend to close, or, at least, the crack-opening displacement tends to decrease. Therefore, the effects associated with strong fiber/ coating interfaces that have been identified at room temperature should be enhanced or, at least, unaffected. As oxidation of the PyC sublayers proceeds, the influence of the above-mentioned effects may decrease. The rate of this phenomenon is not known at this stage. However, comparison of either HN/C and HNT/C, or Table IV. Interfacial Shear Stresses (t) Estimated Using Various Methods and Tow Young’s Modulus Batch Location t (MPa) t# (MPa) Ef (GPa) Equation (1) Equation (3) Equation (4) Model of tensile behavior HN/(C/SiC)10 Interior 105 89 78 100 93 320 Surface 35 25 22 100 45 320 HNT/(C/SiC)10 Interior 216 117 103 350 197 350 Surface 113 61 54 350 144 350 HN/C Interior 120 81 71 100 93 320 Surface 50 34 29 100 53 320 HNT/C Interior 127 160 140 150 144 350 Surface 46 54 50 150 75 350 Fig. 8. Lifetime data for Hi-Nicalon SiC minicomposites in static fatigue at 700°C in air. Also given are lifetime data for Hi-Nicalon/BN/SiC minicomposites.18,19 Arrows indicate those minicomposites that did not fail. 792 Journal of the American Ceramic Society—Bertrand et al. Vol. 84, No. 4
April 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC) Sic Fiber Matrix um (b) 2 μm lu HN/(C/SiC)o and HNT/(C/SiC)o minicomposite lifetimes re- ing conditions. This rovided an upper bound for the flects the significant contribution of treated fibers perating on the The lifetimes obtained for the This contribution involves the previously mentioned effect of /(C/SiC)10 minicomposites compared satisfactorily with the fiber/coating interface. However, better resistance to oxidation obtained for the HN/BN/SiC minicomposites of the Hi-Nicalon fibers(slow crack growth) may also influence the lifetime. This influence will be assessed by an investigation of the oxidation resistance of fibers V. Conclusions For comparison purposes, lifetime data for HN/BN/SIC mini composites with BN interphases are also plotted in Fig. 8. These SiC/SiC minicomposites reinforced with treated Hi-Nica lata were reported elsewhere. 9The test conditions were similar fibers possess strong fiber/coating interfaces. When associated to those described in the present paper. However, the applied load with pyrocarbon/silicon carbide(PyC/SiC) nanoscale multilay was not given. This load was estimated from the stress on the fiber ered interphases deposited via pressure-pulsed chemical vapor if fully loaded, which was used by mosher to describe the infiltration(P-CVD), these strong interfaces lead to interesting
HN/(C/SiC)10 and HNT/(C/SiC)10 minicomposite lifetimes reflects the significant contribution of treated fibers. This contribution involves the previously mentioned effect of the fiber/coating interface. However, better resistance to oxidation of the Hi-Nicalon fibers (slow crack growth) may also influence the lifetime. This influence will be assessed by an investigation of the oxidation resistance of fibers. For comparison purposes, lifetime data for HN/BN/SiC minicomposites with BN interphases are also plotted in Fig. 8. These data were reported elsewhere.18,19 The test conditions were similar to those described in the present paper. However, the applied load was not given. This load was estimated from the stress on the fiber if fully loaded, which was used by Morsher18 to describe the loading conditions. This stress provided an upper bound for the stress operating on the fiber. The lifetimes obtained for the HNT/(C/SiC)10 minicomposites compared satisfactorily with those obtained for the HN/BN/SiC minicomposites. V. Conclusions SiC/SiC minicomposites reinforced with treated Hi-Nicalon fibers possess strong fiber/coating interfaces. When associated with pyrocarbon/silicon carbide (PyC/SiC)n nanoscale multilayered interphases deposited via pressure-pulsed chemical vapor infiltration (P-CVI), these strong interfaces lead to interesting Fig. 9. SEM images showing interfacial regions in minicomposites after static fatigue tests at 700°C in air: (a) HN/(C/SiC)10, (b) HNT/(C/SiC)10, (c) HN/C, and (d) HNT/C. April 2001 Influence of Interfaces on Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)n/SiC 793
Joumal of the American Ceramic Sociery--Bertrand et al. Vol 84. No 4 properties, including an improved lifetime in static fatigue in air at C. Droillard and J. Lamon, "Fracture of 2D Woven siC/iC cvi 700C. The effect of strong interfaces operates through the opposites with Multilayered Interphases, deviation of the crack within the PyC sublayers and limited Nonlinear Stress-Strain Behavior in Micro Influence of Interfacial pa- debonding and crack-opening displacement, which increase the s,"lm.J. fract.,82,297-316(1996) oxygen propagation path and decrease the quantity of oxygen Lissart and J, Lamon,"Damage and Failure in Ceramic Matrix Minicompos- migrating toward the fibers. ites: Experimental Study and Model, Acta Mater, 45 [3]1025-44(1997). Lamon, R. Naslain, E. La The mechanical behavior of the minicomposites under tension Besmann, "Properties of Multilayered Interphases in SiC/SiC CVI Comp exhibited the features previously observed on SiC/SiC composites Weak And"Strong Interfaces, " J. Am Ceram Soc., 81[9]2315-26(199 haracterized by a wide curved domain and large strains at Composites with Multilayered (Byc-sSicn, lntesolh sta ic Fatigue of 2D SicI/sic saturation of matrix cracking that were close to ultimate failure. As pointed out in previous papers, these features were not affected by trand, F. Germain, R Paillet, and J. Lamon,Thermomechanical Behavior of 2D-SiC/SiC Composites with Nanoscale Multilayered(PyC-SiC)m Interphases, the presence of multilayered fiber coatings, including stiff Sic Ad Compos. Lett, 8(61315-21(1999) sublayers. Deflection of the matrix cracks was observed in the Pyc er, and J. Lamon, "Interfacial Behavior the microstructure of the fiber surface that consisted of a thin layer for fficicont mdlilavered n terd a. o mite s ren sinte pre in Gs a g ng5 of free carbon. Such a microstructure was previously detected in SiC/SiC composites reinforced with treated Nicalon fibers (NL 202)and was shown to lead to strong fiber/coating bonds The t values estimated using various methods were comparable with(Pyrocarbon-Sic) Nanoscale Multilayered Interphases,"JAmCeramSoc Nicalon NL 202 fibers Bertrand, K. Pastor, F. Ronchail, and J. Lamon, "Determination of mechanical Statistical Properties of Hi-Nicalon Fibers and Tow acques,A. Guette, F. Langlais, and X. Bourrat,"Characterization of Predictions of the mechanical behavior of minicomposites from SiC/C(B)SiC Microcomposites by Transmission Electron Microscopy,J. Mater. constituent properties were in good agreement with experimental J. Lamon, F. Rebillat, and A. G. Evans, ""Microcomposite Test Procedure for results and allowed the extraction of consistent T values Evaluating the Soc,78[2]401-405(1995) Marshall, B. N. Cox, and A G. Evans, "The Mechanics of Matrix Cracking Acknowledgments eston, G. A. Cooper, and A, Kelly, ""Single and Multiple Fracture The authors wish to thank B Humez for help with mechanical testing, P. Forio for Properties of Fiber composites, Conference Proceedings of the National ph help with computations, x. Bourrat for help with TEM observations, S. Goujard for valuable discussion, and J. Forget and C. Dupouy for preparation of the manuscript. N. Morscher,"Tensile Stress Rupture of SiC,SiCm Minicomposites with Carbon and Boron Nitride Interphases at Elevated Temperatures in Air,J.Am Ceram. Soc.,80[S]2029-42(1997) ces S. Bertrand, O. Boisron, R. Pailler, J. Lamon, and R. Naslain, "(PyC/SiC), and (BN/SiC), Nanoscale-Multilayered Interphases by Pressure-Pulsed CVI, Pp. 321-24 R. Naslain, "Fiber-Matrix Interphases and Interfaces in Ceramic Matrix Compos- in Key Engineering Materials, Vols. 164-165. Trans Tech Publications, Aedermanns- ites Processed by CVI, Compos. Interfaces, 1 [3] 253-86(1993). Droillard,"Elaboration et Caracterisation de Composites a Matrice SiC et 2 H. M. Yun and J. A. DiCarlo, "High Temperature Contraction Behavior of Interphase Sequence C/SiC Ph D. Thesis No. 913. University of Bordeaux, Polymer-Derived SiC Fibers, "Ceram. Eng. Sci. Proc., 18 3]126(1997) Bordeaux, France, 1993 2S. Bertrand and R. Pailler, unpublished work
properties, including an improved lifetime in static fatigue in air at 700°C. The effect of strong interfaces operates through the deviation of the crack within the PyC sublayers and limited debonding and crack-opening displacement, which increase the oxygen propagation path and decrease the quantity of oxygen migrating toward the fibers. The mechanical behavior of the minicomposites under tension exhibited the features previously observed on SiC/SiC composites with strong fiber/coating bonds. The force–deformation curve was characterized by a wide curved domain and large strains at saturation of matrix cracking that were close to ultimate failure. As pointed out in previous papers, these features were not affected by the presence of multilayered fiber coatings, including stiff SiC sublayers. Deflection of the matrix cracks was observed in the PyC sublayers within the coating. This deflection pattern was related to the microstructure of the fiber surface that consisted of a thin layer of free carbon. Such a microstructure was previously detected in SiC/SiC composites reinforced with treated Nicalon fibers (NL 202) and was shown to lead to strong fiber/coating bonds. The t values estimated using various methods were comparable to those measured on SiC/SiC composites reinforced with treated Nicalon NL 202 fibers. Initial tow twisting exerted a certain influence that was reflected by the presence of a higher density of cracks in the internal matrix. Predictions of the mechanical behavior of minicomposites from constituent properties were in good agreement with experimental results and allowed the extraction of consistent t values. Acknowledgments The authors wish to thank B. Humez for help with mechanical testing, P. Forio for help with computations, X. Bourrat for help with TEM observations, S. Goujard for valuable discussion, and J. Forget and C. Dupouy for preparation of the manuscript. References 1 R. Naslain, “Fiber-Matrix Interphases and Interfaces in Ceramic Matrix Composites Processed by CVI,” Compos. Interfaces, 1 [3] 253–86 (1993). 2 C. Droillard, “Elaboration et Caracte´risation de Composites a` Matrice SiC et Interphase Se´quence´e C/SiC”; Ph.D. Thesis No. 913. University of Bordeaux, Bordeaux, France, 1993. 3 C. Droillard and J. Lamon, “Fracture Toughness of 2D Woven SiC/SiC CVI Composites with Multilayered Interphases,” J. Am. Ceram. Soc., 79, 849–58 (1996). 4 L. Guillaumat and J. Lamon, “Fracture Statistics Applied to Modelling the Nonlinear Stress–Strain Behavior in Microcomposites: Influence of Interfacial Parameters,” Int. J. Fract., 82, 297–316 (1996). 5 N. Lissart and J. Lamon, “Damage and Failure in Ceramic Matrix Minicomposites: Experimental Study and Model,” Acta Mater., 45 [3] 1025–44 (1997). 6 F. Rebillat, J. Lamon, R. Naslain, E. Lara-Curzio, M. K. Ferber, and T. M. Besmann, “Properties of Multilayered Interphases in SiC/SiC CVI Composites With ‘Weak’ And ‘Strong’ Interfaces,” J. Am. Ceram. Soc., 81 [9] 2315–26 (1998). 7 S. Pasquier, J. Lamon, and R. Naslain, “Tensile Static Fatigue of 2D SiC/SiC Composites with Multilayered (PyC–SiC)N Interphases at High Temperatures in Oxidizing Atmosphere,” Composites—Part A, 29A, 1157–64 (1998). 8 S. Bertrand, F. Germain, R. Pailler, and J. Lamon, “Thermomechanical Behavior of 2D-SiC/SiC Composites with Nanoscale Multilayered (PyC–SiC)n Interphases,” Adv. Compos. Lett., 8 [6] 315–21 (1999). 9 K. L. Rugg, R. E. Tressler, and J. Lamon, “Interfacial Behavior During Creep of Microcomposites at Elevated Temperature,” J. Eur. Ceram. Soc., 9, 2297–303 (1999). 10C. Droillard, J. Lamon, and X. Bourrat, “Strong Interface in CMCs, a Condition for Efficient Multilayered Interphases,” Mater. Res. Soc. Proc., 365, 371–76 (1995). 11F. Heurtevent, “Nanoscale (PyC–SiC)n Multilayered Interphases–Application as Interphase in Thermostructural Composites” (in Fr.); Ph.D. Thesis No. 1476. University of Bordeaux, Bordeaux, France, 1996. 12S. Bertrand, P. Forio, R. Pailler, and J. Lamon, “Hi-Nicalon/SiC Minicomposites with (Pyrocarbon–SiC)n Nanoscale Multilayered Interphases,” J. Am. Ceram., Soc., 82 [9] 2465–73 (1999). 13S. Bertrand, K. Pastor, F. Ronchail, and J. Lamon, “Determination of Mechanical and Statistical Properties of Hi-Nicalon Fibers and Tows”; to be published. 14S. Jacques, A. Guette, F. Langlais, and X. Bourrat, “Characterization of SiC/C(B)/SiC Microcomposites by Transmission Electron Microscopy,” J. Mater. Sci, 32, 2969 (1997). 15J. Lamon, F. Rebillat, and A. G. Evans, “Microcomposite Test Procedure for Evaluating the Interface Properties of Ceramic Matrix Composites,” J. Am. Ceram. Soc., 78 [2] 401–405 (1995). 16D. B. Marshall, B. N. Cox, and A. G. Evans, “The Mechanics of Matrix Cracking in Brittle-Matrix Fiber Composites,” Acta Metall., 33 [11] 2013 (1985). 17J. Aveston, G. A. Cooper, and A. Kelly, “Single and Multiple Fracture”; in Properties of Fiber Composites, Conference Proceedings of the National Physical Laboratory, Vol. 15. IPC Science and Technology Press Ltd., Surrey, U.K., 1971. 18G. N. Morscher, “Tensile Stress Rupture of SiCf/SiCm Minicomposites with Carbon and Boron Nitride Interphases at Elevated Temperatures in Air,” J. Am. Ceram. Soc., 80 [8] 2029–42 (1997). 19S. Bertrand, O. Boisron, R. Pailler, J. Lamon, and R. Naslain, “(PyC/SiC)n and (BN/SiC)n Nanoscale-Multilayered Interphases by Pressure-Pulsed CVI”; pp. 321–24 in Key Engineering Materials, Vols. 164–165. TransTech Publications, Aedermannsdorf, Switzerland, 1999. 20H. M. Yun and J. A. DiCarlo, “High Temperature Contraction Behavior of Polymer-Derived SiC Fibers,” Ceram. Eng. Sci. Proc., 18 [3] 126 (1997). 21S. Bertrand and R. Pailler; unpublished work. M 794 Journal of the American Ceramic Society—Bertrand et al. Vol. 84, No. 4