C Pergamon Acta mater.4902001)3727-3738 www.elsevier.com/locate/actamat SUBCRITICAL CRACK GROWTH IN CVI SIC!SIC COMPOSITES AT ELEVATED TEMPERATURES: EFFECT OF FIBER CREEP RATE C H. HENAGER Jr, C.A. LEWINSOHN+ and R. H. JONES Pacific Northwest National Laboratory,$ Materials Sciences/P8-15, 902 Battelle Blvd, PO box 999, Richland. WA 99352-0999 US Received 3 April 2000, received in revised form 9 July 2001: accepted 9 July 2001) Abstract-Subcritical crack-growth studies in SiC!SiC composites were conduc reinforced with Hi-Nicalon fibers over a broad temperature range materials reinforced with Nicalon-CG fibers. Composites with a 0/90 plain weave and carbon interphase were tested in argon from 1 173 to 1473 K. Crack growth data obtained ronments are onsistent with a proposed fiber-creep-controlled crack-growth mechanism Measured crack-growth activation energies and ti perature exponents in argon agree with fiber creep-activation energies and nonlinear reep equations for both fiber types. Estimates of local strains during crack growth are in reasonable agreement with estimated fiber creep strains for the given times and temperatures. The increased creep resistance of Hi-Nicalon fibers is reflected in reduced crack-growth rates for composites containing those fibers. 200/ Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Ceramics-structural; Composites; Mechanical properties-creep; High temperature; Fracture 1 INTRODUCTION stand critical time-dependent deformation mech- Continuous fiber ceramic matrix composites( CFCCs) anisms are more reliable than unreinforced ceramics [I], but Previous studies(this group and others)have are prone to time-dependent failures from stable crack shown that creep of bridging fibers at elevated tem- [2-71 or from mechanical embrittlement(unstable nation for the observed creep or slow crack growth crack growth)caused by environmental exposure 18, of these continuous-fiber composites where the Sic 9. In particular, we are interested in understanding matrix is more creep resistant than the fibers [2, 3 time-dependent properties for these materials in gas- 3-26]. These results support the hypothesis that cooled advanced fission and fusion reactor environ creeping fibers transfer stress back to the matrix ments [10-12). Such environments are pristine in causing further matrix cracking, a loss of matrix stiff terms of oxygen content, and composites with carbon- ness, and increased loading of the crack-bridging fib- based interphases appear attractive. Since life-predic- ers, ultimately leading to failure. Recent results rel- tion methodologies for these materials would neces- evant to this work include experimental creep or sarily include time-dependent crack growth as an crack-growth tests on SiC/SiC composites [18, 19, important failure mechanism, it is essential to under- 23-26] and models of time-dependent crack growth [27-31]A complete discussion of the models appears in a companion paper [32] and will not be I To whom all correspondence should be addressed E-mail address: chuck. henager(@pnl. gov(C. H. Hen- Evans and Weber [18] documented increased com- pliances due to matrix cracking and also observed supported by Associated Western Univer- fiber sliding stresses at 1473 K that were almost one west Division(AWU NW) under Grant order of magnitude smaller than at room temp rature DE-FC -75522, DE-FG07-93ER-75912, or DE- Wilshire et al. [23] compared composite creep rates AC06-76RLO1830 le to matrix cracking to account for observed com 1359-6454/01/S20.00@ 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. P:S1359-6454(01)00276-2
Acta mater. 49 (2001) 3727–3738 www.elsevier.com/locate/actamat SUBCRITICAL CRACK GROWTH IN CVI SiCf/SiC COMPOSITES AT ELEVATED TEMPERATURES: EFFECT OF FIBER CREEP RATE C. H. HENAGER Jr†, C. A. LEWINSOHN‡ and R. H. JONES Pacific Northwest National Laboratory,§ Materials Sciences/P8-15, 902 Battelle Blvd, PO box 999, Richland, WA 99352-0999, USA ( Received 3 April 2000; received in revised form 9 July 2001; accepted 9 July 2001 ) Abstract—Subcritical crack-growth studies in SiCf/SiC composites were conducted with composites reinforced with Hi-Nicalon fibers over a broad temperature range for comparison to earlier studies on materials reinforced with Nicalon-CG fibers. Composites with a 0/90 plain weave architecture and carbon interphase were tested in argon from 1173 to 1473 K. Crack growth data obtained in inert environments are consistent with a proposed fiber-creep-controlled crack-growth mechanism. Measured crack-growth activation energies and time–temperature exponents in argon agree with fiber creep-activation energies and nonlinear creep equations for both fiber types. Estimates of local strains during crack growth are in reasonable agreement with estimated fiber creep strains for the given times and temperatures. The increased creep resistance of Hi-Nicalon fibers is reflected in reduced crack-growth rates for composites containing those fibers. 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Ceramics—structural; Composites; Mechanical properties—creep; High temperature; Fracture & fracture toughness 1. INTRODUCTION Continuous fiber ceramic matrix composites (CFCCs) are more reliable than unreinforced ceramics [1], but are prone to time-dependent failures from stable crack growth occurring in inert and oxidizing environments [2–7] or from mechanical embrittlement (unstable crack growth) caused by environmental exposure [8, 9]. In particular, we are interested in understanding time-dependent properties for these materials in gascooled advanced fission and fusion reactor environments [10–12]. Such environments are pristine in terms of oxygen content, and composites with carbonbased interphases appear attractive. Since life-prediction methodologies for these materials would necessarily include time-dependent crack growth as an important failure mechanism, it is essential to under- † To whom all correspondence should be addressed. E-mail address: chuck.henager@pnl.gov (C. H. Henager Jr) ‡ Research supported by Associated Western Universities, Inc., Northwest Division (AWU NW) under Grant DE-FG06-89ER-75522, DE-FG07-93ER-75912, or DEFG06-92RL-12451 with the US Department of Energy. § Pacific Northwest National Laboratory is operated for the US Department of Energy by Battelle under Contract DE-AC06-76RLO 1830. 1359-6454/01/$20.00 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6454(01)00276-2 stand critical time-dependent deformation mechanisms. Previous studies (this group and others) have shown that creep of bridging fibers at elevated temperatures in inert environments is one possible explanation for the observed creep or slow crack growth of these continuous-fiber composites where the SiC matrix is more creep resistant than the fibers [2, 3, 13–26]. These results support the hypothesis that creeping fibers transfer stress back to the matrix, causing further matrix cracking, a loss of matrix stiffness, and increased loading of the crack-bridging fibers, ultimately leading to failure. Recent results relevant to this work include experimental creep or crack-growth tests on SiCf/SiC composites [18, 19, 23–26] and models of time-dependent crack growth [27–31]. A more complete discussion of the models appears in a companion paper [32] and will not be included here. Evans and Weber [18] documented increased compliances due to matrix cracking and also observed fiber sliding stresses at 1473 K that were almost one order of magnitude smaller than at room temperature. Wilshire et al. [23] compared composite creep rates at 1573 K to fiber creep rates to demonstrate the degree of load transfer to the fibers that must occur due to matrix cracking to account for observed com-
3728 HENAGER et al. SUBCRITICAL CRACK GROWTH: PART I posite creep rates. They observed extensive matrix exponents if the understanding produced by the pre- cracking due to fiber creep and pointed out that this vious studies is correct. present authors have observed [4, 33). Zhu et al. [2 performed creep tests in air and argon at 1573 K on 2. EXPERIMENTAL APPROACH Hi-Nicalon-fiber SiC/SiC and Nicalon-CG-fiber The experimental approach has been described SiC/siC and made many of the same observations. more detail elsewhere [3] and will only be summar They did measure a slower creep rate and longer ized here. Subcritical crack growth was obtained by time-to-rupture for the Hi-Nicalon-fiber composite loading single-edge-notched beam(SENB)specimens attributed to a greater creep resistance of Hi-Nicalon (typically 50x55x4 mm)in 1/4 four-point bending in fibers compared to Nicalon-CG fibers [34]. Tressler a fully articulated silicon-carbide fixture with an outer al. [26] crept Hi-Nicalon-fiber SiC/SiC microcom- span length of 40 mm. We incorporated the specimen posites at 1473 to 1673 K and demonstrated a model and fixture in a vertically oriented mullite tube con that used explicit fiber and matrix creep data to tained within a high-temperature furnace mounted to account for creep deformation with and without an electromechanically controlled mechanical test matrix cracks. They observed transient creep curves frame. t The SENB specimens contained an initial for all materials and conditions in their study. A rule- notch with a depth-to-width ratio(all) of approxi- of-mixtures creep model was used for the material mately 0. 2. The notch was made by a high-speed dia- without matrix cracks, while a different approach mond saw and was typically 3.9x10-4 m wide at the suitable for a crack bridging fiber in tension was tip of the notch and 5.9x10-4 m wide at the mouth adopted for the cracked microcomposites No attempt was made to sharpen the notch. Speci- However, few studies have compared creep or mens were heated to the test temperature at a rate of crack-growth rates in materials with similar com- about 0. 25 K/s and were allowed to equilibrate for posite architecture but different fibers [13, 25, 35]. 1200 s at temperature. Then a load calculated [38]to Studies by the present authors [2-4, 13] have demon- provide an initial applied stress intensity of 9 to 10 strated that thermal activation energies for crack MPa m 2 was applied to the sample and held for the growth in inert environments are consistent with acti- duration of the test. This stress intensity was chosen vation energies measured for these same fibers in sin- to induce some initial crack extension and crack e-fiber creep or stress relaxation tests [34, 36]. bridging and it falls between that required for matrix However, this strong evidence for a crack growth cracking and the peak load fracture toughness, Ko, mechanism controlled by fiber creep would be bols- reported in Table 1. In addition, we periodically tered by supporting evidence in similar materials with unloaded to 95% of the constant applied load and then different fibers. Additional data would provide an reloaded; at the time of the initial loading and 5, 25 opportunity for testing our ability to model crack and 50 h after the initial loading to generate hysteresis growth in a specific specimen geometry using loops(for the Hi-Nicalon materials only) detailed fiber properties where the fiber type was the The atmosphere inside the mullite tube was con- most significant variable trollable and maintained at atmospheric pressure Moreover, since these composites are being con- (1.01x10 Pa). We used gettered argon, initially sidered for use in high-temperature, gas-cooled reac- 99.999% pure, for testing, with an oxygen content tor concepts [10-12], we are interested in lifetimes in reduced to less than 0.01 Pa by passing the gas ert environments or in low oxygen concentrations. through a titanium-gettering furnace. The deflection Many of the studies discussed above were performed of the specimen midpoint was measured by using ar in air, which is a very degrading environment, parti- alumina pushrod, also containing a thermocouple, cularly for CFCCs with a carbon-based fiber-matrix attached to a strain-gauge extensometer. The dis- interphase. This teaches us little about life prediction placements were corrected for differences between and failure mechanisms in environments where these the load-point and midpoint and for the compliance materials are seriously being considered for use. of the test apparatus 3] Therefore, we began the present study using similar An SiC/SiC CFCC containing Hi-NicalonTM fibers 0/90-woven, chemical vapor infiltration(CVI)-SiC was examined in this study and compared to the matrix CFCCs reinforced with Hi-Nicalon fibers to results of previous studies [2-4, 13] on materials con- compare with the previous Nicalon-CG fiber CFCCs taining "Ceramic-grade" Nicalon TM(Nicalon-CG) 24, 13]. This study is the first to critically compare fibers. Data obtained using the previous Nicalon-CG the effect of fiber creep characteristics on crack- materials were extended over a temperature range growth kinetics and activation energies in an inert from 1173 to 1398 K. The Hi-Nicalon materials were environment. Hi-Nicalon fibers have higher creep fabricated from two-dimensional, plain weave fiber activation energies and a reduced creep rate compared to Nicalon-CG fibers [34, 36, 37]. Composites made with Hi-Nicalon fibers should exhibit crack-growth ates and activation energies consis on fiber creep rates, activation energies, and creep t Instron1125,Instron Corp,Canton,MA,USA
3728 HENAGER et al.: SUBCRITICAL CRACK GROWTH: PART I posite creep rates. They observed extensive matrix cracking due to fiber creep and pointed out that this cracking facilitates environmental ingress, as the present authors have observed [4, 33]. Zhu et al. [25] performed creep tests in air and argon at 1573 K on Hi-Nicalon-fiber SiCf/SiC and Nicalon-CG-fiber SiCf/SiC and made many of the same observations. They did measure a slower creep rate and longer time-to-rupture for the Hi-Nicalon-fiber composite attributed to a greater creep resistance of Hi-Nicalon fibers compared to Nicalon-CG fibers [34]. Tressler et al. [26] crept Hi-Nicalon-fiber SiCf/SiC microcomposites at 1473 to 1673 K and demonstrated a model that used explicit fiber and matrix creep data to account for creep deformation with and without matrix cracks. They observed transient creep curves for all materials and conditions in their study. A ruleof-mixtures creep model was used for the material without matrix cracks, while a different approach suitable for a crack bridging fiber in tension was adopted for the cracked microcomposites. However, few studies have compared creep or crack-growth rates in materials with similar composite architecture but different fibers [13, 25, 35]. Studies by the present authors [2–4, 13] have demonstrated that thermal activation energies for crack growth in inert environments are consistent with activation energies measured for these same fibers in single-fiber creep or stress relaxation tests [34, 36]. However, this strong evidence for a crack growth mechanism controlled by fiber creep would be bolstered by supporting evidence in similar materials with different fibers. Additional data would provide an opportunity for testing our ability to model crack growth in a specific specimen geometry using detailed fiber properties where the fiber type was the most significant variable. Moreover, since these composites are being considered for use in high-temperature, gas-cooled reactor concepts [10–12], we are interested in lifetimes in inert environments or in low oxygen concentrations. Many of the studies discussed above were performed in air, which is a very degrading environment, particularly for CFCCs with a carbon-based fiber-matrix interphase. This teaches us little about life prediction and failure mechanisms in environments where these materials are seriously being considered for use. Therefore, we began the present study using similar 0/90-woven, chemical vapor infiltration (CVI)-SiC matrix CFCCs reinforced with Hi-Nicalon fibers to compare with the previous Nicalon-CG fiber CFCCs [2–4, 13]. This study is the first to critically compare the effect of fiber creep characteristics on crackgrowth kinetics and activation energies in an inert environment. Hi-Nicalon fibers have higher creep activation energies and a reduced creep rate compared to Nicalon-CG fibers [34, 36, 37]. Composites made with Hi-Nicalon fibers should exhibit crack-growth rates and activation energies consistent with Hi-Nicalon fiber creep rates, activation energies, and creep exponents if the understanding produced by the previous studies is correct. 2. EXPERIMENTAL APPROACH The experimental approach has been described in more detail elsewhere [3] and will only be summarized here. Subcritical crack growth was obtained by loading single-edge-notched beam (SENB) specimens (typically 50×5.5×4 mm) in 1/4 four-point bending in a fully articulated silicon-carbide fixture with an outer span length of 40 mm. We incorporated the specimen and fixture in a vertically oriented mullite tube contained within a high-temperature furnace mounted to an electromechanically controlled mechanical test frame.† The SENB specimens contained an initial notch with a depth-to-width ratio (a/W) of approximately 0.2. The notch was made by a high-speed diamond saw and was typically 3.9×104 m wide at the tip of the notch and 5.9×104 m wide at the mouth. No attempt was made to sharpen the notch. Specimens were heated to the test temperature at a rate of about 0.25 K/s and were allowed to equilibrate for 1200 s at temperature. Then a load calculated [38] to provide an initial applied stress intensity of 9 to 10 MPa m1/2 was applied to the sample and held for the duration of the test. This stress intensity was chosen to induce some initial crack extension and crack bridging and it falls between that required for matrix cracking and the peak load fracture toughness, KQ, reported in Table 1. In addition, we periodically unloaded to 95% of the constant applied load and then reloaded; at the time of the initial loading and 5, 25, and 50 h after the initial loading to generate hysteresis loops (for the Hi-Nicalon materials only). The atmosphere inside the mullite tube was controllable and maintained at atmospheric pressure (1.01×105 Pa). We used gettered argon, initially 99.999% pure, for testing, with an oxygen content reduced to less than 0.01 Pa by passing the gas through a titanium-gettering furnace. The deflection of the specimen midpoint was measured by using an alumina pushrod, also containing a thermocouple, attached to a strain-gauge extensometer. The displacements were corrected for differences between the load-point and midpoint and for the compliance of the test apparatus [3]. An SiCf/SiC CFCC containing Hi-Nicalon fibers was examined in this study and compared to the results of previous studies [2–4, 13] on materials containing “Ceramic-grade” Nicalon (Nicalon-CG) fibers. Data obtained using the previous Nicalon-CG materials were extended over a temperature range from 1173 to 1398 K. The Hi-Nicalon materials were fabricated from two-dimensional, plain weave fiber † Instron 1125, Instron Corp., Canton, MA, USA
HENAGER et al. SUBCRITICAL CRACK GROWTH: PART I 3729 Table 1. Physical properties of the materials tested(variability given as 95% confidence intervals) Composite designation CG-C Hi-C Fiber architecture 2D plain weave fiber mats 2D plain weave fiber mats thermal cvi at-1050°C othermal cvi at-1050° eramic-grade nical Fiber vol ameter(um) 14±2.3 ol. fractio Composite coating CVD SIC CVD SIC Modulus (GPa) 122±21 172±23 Fracture toughness, Ko(MPa m")5 nterface sliding stress(MPay 8.4+6.6 72.1±33.5 aWe do not have access to all the processing details for these vendor-supplied materials. However, the processing of these CVI-SiC/SiC The modulus is calculated from the measured from the linear portion of loading during flexural testing of SENB specimens at elevated temperatur ture compliance. "Calculated using peak load as described in ASTM E399 mEasured by fiber push-in test at room temperature as described in [43] mats.t Before matrix infiltration, a 1-um-thick carbon 3. RESULTS interphase layer was deposited on the fibers by chemi- 3. 1. Displacement-time cures each bend-bar were coated with a 2-um-thick layer The displacement-time curve for a representative of silicon carbide, deposited by CVD, to provide oxi- CG-C specimen tested at a constant load of 630 N at dation resistance and to protect the surfaces from 1373 K for 8x10 s in argon is a nonlinear curve( Fig damage. These materials will be abbreviated as"Hi- 1), suggestive of a transient creep curve, but indicate the fiber type and the interphase com- accompanied by subcritical, time-dependent crack position(Table 1) growth in a multiply-cracked damage zone(see The materials studied previously were fabricated in below) that extends from the notch root much as a a similar manner. t The Nicalon-CG fibers were also mode I crack. The data shown in Fig. I are the midp- oated with either a I-um-thick carbon interphase int displacement of the notched bar as a function of layer, a 0. 4-um-thick carbon/boron nitride(C/BN) time, t, and are compared to a functional fit of the interphase layer, or a I-Hm-thick multi-layer form nterphase layer consisting of carbon/boron carbide/boron nitride (C/B,C/BN), with the carbon f= alt exp(-bIDIe layer next to the fiber. These materials will be abbreviated"CG-C. The time-independent mechan- ical properties of the CG-C materials in argon were corresponds to a Sherby-Dorn creep equation entical, within experimental uncertainties, in a,b, and c are independent fitting parameters pendent of the interphase chemistry and thickness Composite properties and identification details are listed in Table 0.004 We tested Hi-C specimens at 1373, 1423, 1448 - fit to data nd 1473 K in argon ed 10 CG-C specimens tested at 1173, 1273, 1298, 1323, 5 1348,1373,and1398K gon. The Hi-C com-§610 sites were tested at temperatures higher than the CG-C materials because Hi-Nicalon fibers have greater thermal stability compared to Nicalon-CG 2I0 8105 t Composite fabrication performed by DuPont Lanxide Fig. 1. Displacement-time curve for a CG-C material(Table Corp, Wilmington, DE, USA 1)at 600 N constant load(corresp Composite fabrication performed by RCL, Whittier, m 2 )test at 1373 K in argon. Experi ve (solid line) CA, USA(RCI is no longer existent). is compared to equation(1)
HENAGER et al.: SUBCRITICAL CRACK GROWTH: PART I 3729 Table 1. Physical properties of the materials testeda (variability given as 95% confidence intervals) Composite designation CG-C Hi-C Fiber architecture 2D plain weave fiber mats 2D plain weave fiber mats Processing conditions Isothermal CVI at 1050°C Isothermal CVI at 1050°C Fiber type Ceramic-grade Nicalon Hi-Nicalon Fiber coating thickness (µm) 1.0±0.1 1.0±0.1 Fibers/tow 420 321 Fiber diameter (µm) 15.0±0.5 14±2.3 Fiber vol. fraction 40 40 Composite coating CVD SiC CVD SiC Porosity 20±5% 6±1.4% Modulus (GPa)b 122±21 172±23 Fracture toughness, KQ (MPa m1/2) c 17.5 22.4±0.1 Interface sliding stress (MPa)d 8.4±6.6 72.1±33.5 a We do not have access to all the processing details for these vendor-supplied materials. However, the processing of these CVI-SiC/SiC materials is rather standardized. b The modulus is calculated from the specimen compliance measured from the linear portion of loading during flexural testing of SENB specimens at elevated temperature and corrected for fixture compliance. c Calculated using peak load as described in ASTM E399. d Measured by fiber push-in test at room temperature as described in [43]. mats.† Before matrix infiltration, a 1-µm-thick carbon interphase layer was deposited on the fibers by chemical vapor deposition (CVD). The outer surfaces of each bend-bar were coated with a 2-µm-thick layer of silicon carbide, deposited by CVD, to provide oxidation resistance and to protect the surfaces from damage. These materials will be abbreviated as “HiC” to indicate the fiber type and the interphase composition (Table 1). The materials studied previously were fabricated in a similar manner.‡ The Nicalon-CG fibers were also coated with either a 1-µm-thick carbon interphase layer, a 0.4-µm-thick carbon/boron nitride (C/BN) interphase layer, or a 1-µm-thick multi-layer interphase layer consisting of carbon/boron carbide/boron nitride (C/B4C/BN), with the carbon layer next to the fiber. These materials will be abbreviated “CG-C.” The time-independent mechanical properties of the CG-C materials in argon were identical, within experimental uncertainties, independent of the interphase chemistry and thickness. Composite properties and identification details are listed in Table 1. We tested Hi-C specimens at 1373, 1423, 1448, and 1473 K in argon and compared the results with CG-C specimens tested at 1173, 1273, 1298, 1323, 1348, 1373, and 1398 K in argon. The Hi-C composites were tested at temperatures higher than the CG-C materials because Hi-Nicalon fibers have greater thermal stability compared to Nicalon-CG fibers. † Composite fabrication performed by DuPont Lanxide Corp., Wilmington, DE, USA. ‡ Composite fabrication performed by RCI, Whittier, CA, USA (RCI is no longer existent). 3. RESULTS 3.1. Displacement–time curves The displacement–time curve for a representative CG-C specimen tested at a constant load of 630 N at 1373 K for 8×105 s in argon is a nonlinear curve (Fig. 1), suggestive of a transient creep curve, but accompanied by subcritical, time-dependent crack growth in a multiply-cracked damage zone (see below) that extends from the notch root much as a mode I crack. The data shown in Fig. 1 are the midpoint displacement of the notched bar as a function of time, t, and are compared to a functional fit of the form: f a[t exp(b/T)]c (1) which corresponds to a Sherby–Dorn creep equation where a, b, and c are independent fitting parameters Fig. 1. Displacement–time curve for a CG-C material (Table 1) at 600 N constant load (corresponding to a Ka = 9.6 MPa m1/2) test at 1373 K in argon. Experimental curve (solid line) is compared to equation (1) (dashed line).
3730 HENAGER et al. SUBCRITICAL CRACK GROWTH: PART I 2.50 0.000l 三 ∠1398K 1473K 2 81 1510 g610 51 2 10 1273K CG-C 1373K 110421043104410451 Fig. 2. Displacement-time curves for a Nicalon-CG (CG-C) Fig 4. Displacement-time curves for identical Hi-C SENI range of 1273 he in 25 K temperature steps for 4XI0- four CG-C specimens tested at 1373 K in argon are included nt-time data were adjusted for each temperat and only the time-depen and T is the temperature in Kelvin. The 630 N load corresponds to an initial applied stress intensity (K) 9.6 MPa m/ at the initial normalized notch length f alW=0. 17(a=0.93 mm). We explored the tem perature dependence of this subcritical cracking for (a) the CFCC materials by performing identical tests over a range of temperatures(Figs 2-4). Initially, one specimen of the CG-C material was tested by taking it to progressively higher temperatures in argon, each exposure lasting for 4x10s(Fig. 2), from 1273 to 2 500 1398 K in 25 K increments and at a constant load of 630 N. The specimen was unloaded during tempera- 4001 ture increases and allowed to come to thermal equilib- rium at the new, higher temperature before it was Total plastic reloaded. Similar temperature-dependent data for the train at zero load CG-C material were obtained by performing identical 1373 K in argon where the specimens were held for elastic. ot tests on three separate specimens at 1 173, 1273, and 1030.000120.0016 longer times at constant loads corresponding to initial (b) K values of 10 MPa m(Fig. 3). The loading dis-70 placements of the samples are not shown--only the hours in Ar: min in O: le-dependent displacements under constant load are plotted in Figs 2-4 400 300 3105 l373K 2105010"21034105 11041.210-41.4104 273K Mid-Point Displacement (m) 1173K Fig. 5.(a)Load-displacement data showing unloading measured in flexure at 0. 5. 25. and 5 10 1105 1.510 2 10 ment on Hi-C material at 1448K in argon. Note that there is essentially no change in the appearance of the loops after 50 Fig. 3. Displacement-time curves for three identical CG-C cated(see Table 2).(b) Similar load-di The data were adjusted so that zero time coincides with th beginning of the time-dependent displacements. The loads on to 202 Pa O,(added each specimen correspond to a K,=10 MPa mn and loop widening, which is attributed to interphase recession
3730 HENAGER et al.: SUBCRITICAL CRACK GROWTH: PART I Fig. 2. Displacement–time curves for a Nicalon-CG (CG-C) SENB specimen tested at 600 N in argon over a temperature range of 1273 to 1398 K in 25 K temperature steps for 4×104 s per step. The displacement–time data were adjusted to zero for each temperature step and only the time-dependent displacements are shown. and T is the temperature in Kelvin. The 630 N load corresponds to an initial applied stress intensity (Ka) of 9.6 MPa m1/2 at the initial normalized notch length of a/W = 0.17 (a = 0.93 mm). We explored the temperature dependence of this subcritical cracking for the CFCC materials by performing identical tests over a range of temperatures (Figs 2–4). Initially, one specimen of the CG-C material was tested by taking it to progressively higher temperatures in argon, each exposure lasting for 4×104 s (Fig. 2), from 1273 to 1398 K in 25 K increments and at a constant load of 630 N. The specimen was unloaded during temperature increases and allowed to come to thermal equilibrium at the new, higher temperature before it was reloaded. Similar temperature-dependent data for the CG-C material were obtained by performing identical tests on three separate specimens at 1173, 1273, and 1373 K in argon where the specimens were held for longer times at constant loads corresponding to initial K values of 10 MPa m1/2 (Fig. 3). The loading displacements of the samples are not shown—only the time-dependent displacements under constant load are plotted in Figs 2–4. Fig. 3. Displacement–time curves for three identical CG-C SENB specimens tested at 1373, 1273, and 1173 K in argon. The data were adjusted so that zero time coincides with the beginning of the time-dependent displacements. The loads on each specimen correspond to a Ka = 10 MPa m1/2. Fig. 4. Displacement–time curves for identical Hi-C SENB specimens tested at 1373 to 1473 K in argon. The curves for four CG-C specimens tested at 1373 K in argon are included for comparison. Fig. 5. (a) Load–displacement data showing unloading– reloading hysteresis loops measured in flexure at 0, 5, 25, and 50 h, respectively, during a subcritical-crack-growth experiment on Hi-C material at 1448 K in argon. Note that there is essentially no change in the appearance of the loops after 50 h of crack growth. Elastic, inelastic, and plastic strains are indicated (see Table 2). (b) Similar load–displacement data showing unloading–reloading hysteresis loops at 0, 5, 25 h in argon at 1448 K but immediately followed by exposure during testing to 202 Pa O2 (added to argon gas feed). Note the slope decrease and loop widening, which is attributed to interphase recession
HENAGER et al: SUBCRITICAL CRACK GROWTH: PART I Table 2. SCG specimens total strain and time data for representative CG-C and Hi-C materials Specimen ID Test temperature(K) Test duration(s) ent Strain at load Strain at unloading displacement"(m) 戀 1.0x10-4 2.0×10-2 1473 12×10 1.82×105 1373 2.73×105 5.6×10 5.5×10-3 ent at load excluding elastic and inelastic loading displacements(Fig. 5a). sing four-point bending relation between outer fiber strain and specimen deflection but allowing the displacement only in the see equation(7)and Table 3). Subcritical-crack-growth (SCG) data for Hi-C immediately followed by naterials are similar to those of the previously tested (pO2=202 Pa) for which the hysteresis loops appear CG-C materials [2-4, 13]. The temperature depen- distinctly different(Fig. 5b). Although the fidelity of dence of cracking in the Hi-C materials was investi- these hysteresis curves lacks the precision required gated by testing separate and identical specimens at to obtain detailed information regarding interface slip 1373, 1423, 1448, and 1473 K(Fig. 4) at constant transfer(because of surrounding elastic unloading of ads corresponding to initial Ka values of 10 MPa the SENB specimen), it is apparent that the interface m. The displacement-time curves for several CG- mechanical properties are unchanged when compared C specimens tested at 1373 K are included for com- to oxidation-induced interface removal mechanism parison with the Hi-C materials 42,43] recoverable, strain during the scG testing (Table 2), 3.2. Subcritical crack growth damage sone both under load and unloaded (for selected The resultant cracking observed under these con- specimens). It was found that the total plastic strain ditions for CG-C materials is shown in Fig. 6a in an under load was equal to the total plastic strain after optical micrograph of a polished SENB cross-section, unloading for the specimens where an accurate value where the section was taken from the center of the for the inelastic loading strain due to matrix cracking SENB bar by cutting the bar lengthwise parallel to the on loading above the matrix cracking stress was crack propagation direction and normal to the plane obtained. In addition, we investigated the possibility containing the cracks. Multiple cracks are seen in the that other inelastic deformation mechanisms were CVI-SiC matrix material. but little fiber breakage is responsible for the observed behavior by qualitatively observed, even at the root of the notch where the analyzing unloading-reloading hysteresis loops that crack opening displacement would be the largest were performed periodically during the crack-growth schematic of the cracked damage zone is shown in experiments on the Hi-C materials(Fig 5a). The load Fig. 6b. A single Hi-C specimen was also sectioned versus displacement data for the hysteresis loops and showed similar multiple cracking (uncorrected for machine compliance) in inert Several specimens were tested, unloaded while at environments exhibited characteristics observed by temperature, cooled with no applied load, and then others [39-41]. In contrast to this, we also show hys- sectioned for optical microscopy of the cracks and of teresis loops for a specimen tested in argon but the damage zone. Data from these nens, In microscopy of sectioned SCG specime emperature Test duration Initial crack Number of cracks"(n) Dam ength(m) length(m) CG-C-I ■睡 Hi-C-2 2.0×105 1,12×10-3 265×10-3 1.45×10-3 3.6×10-4 Initial number emanating from notch and final number at maximum damage zone extent final crack length)
HENAGER et al.: SUBCRITICAL CRACK GROWTH: PART I 3731 Table 2. SCG specimens total strain and time data for representative CG-C and Hi-C materials Specimen ID Test temperature (K) Test duration (s) Total time-dependent Strain at loadb Strain at unloading displacementa (m) CG-C-1 1373 1.16×104 1.32×105 2.6×103 – CG-C-2 1373 7.45×104 2.82×105 5.6×103 – CG-C-3 1373 9.62×104 3.7×105 7.5×103 – CG-C-4 1373 6.22×105 9.2×105 1.85×102 – CG-C-5 1373 7.65×105 1.0×104 2.0×102 – Hi-C-1 1473 1.81×105 1.04×104 1.7×102 – Hi-C-2 1448 2.0×105 7.3×105 1.2×102 1.2×102 Hi-C-3 1423 1.82×105 4.4×105 7.3×103 – Hi-C-4 1373 2.73×105 3.35×105 5.6×103 5.5×103 a Specimen mid-point displacement at load excluding elastic and inelastic loading displacements (Fig. 5a). b Strain computed at notch root using four-point bending relation between outer fiber strain and specimen deflection but allowing the displacement to occur only in the damage zone (see equation (7) and Table 3). Subcritical-crack-growth (SCG) data for Hi-C materials are similar to those of the previously tested CG-C materials [2–4, 13]. The temperature dependence of cracking in the Hi-C materials was investigated by testing separate and identical specimens at 1373, 1423, 1448, and 1473 K (Fig. 4) at constant loads corresponding to initial Ka values of 10 MPa m1/2. The displacement–time curves for several CGC specimens tested at 1373 K are included for comparison with the Hi-C materials. Data were obtained for the total “plastic,” or nonrecoverable, strain during the SCG testing (Table 2), both under load and unloaded (for selected specimens). It was found that the total plastic strain under load was equal to the total plastic strain after unloading for the specimens where an accurate value for the inelastic loading strain due to matrix cracking on loading above the matrix cracking stress was obtained. In addition, we investigated the possibility that other inelastic deformation mechanisms were responsible for the observed behavior by qualitatively analyzing unloading–reloading hysteresis loops that were performed periodically during the crack-growth experiments on the Hi-C materials (Fig. 5a). The load versus displacement data for the hysteresis loops (uncorrected for machine compliance) in inert environments exhibited characteristics observed by others [39–41]. In contrast to this, we also show hysteresis loops for a specimen tested in argon but Table 3. Summary of optical microscopy of sectioned SCG specimens Test Temperature Test duration Initial crack Final crack Number of cracksa (n) Damage zone Mean crack specimen (K) (s) length (m) length (m) width (m) spacing (m) Initial Final CG-C-1 1373 1.16×104 1.0×103 1.18×103 1 – – CG-C-2 1373 7.45×104 0.82×103 1.62×103 9 5 1.24×103 2.5×104 CG-C-3 1373 9.62×104 0.93×103 1.95×103 8 3 1.14×103 3.8×104 CG-C-4 1373 6.22×105 0.76×103 2.44×103 10 4 1.33×103 3.3×104 CG-C-5 1373 7.65×105 0.93×103 2.75×103 10 5 1.24×103 2.5×104 Average –––– 9 4 1.24×103 3.0×104 CG-C Hi-C-2 1448 2.0×105 1.12×103 2.65×103 4 1.45×103 3.6×104 a Initial number emanating from notch and final number at maximum damage zone extent (final crack length). immediately followed by testing in oxygen (pO2 = 202 Pa) for which the hysteresis loops appear distinctly different (Fig. 5b). Although the fidelity of these hysteresis curves lacks the precision required to obtain detailed information regarding interface slip transfer (because of surrounding elastic unloading of the SENB specimen), it is apparent that the interface mechanical properties are unchanged when compared to oxidation-induced interface removal mechanisms [42, 43]. 3.2. Subcritical crack growth damage zone The resultant cracking observed under these conditions for CG-C materials is shown in Fig. 6a in an optical micrograph of a polished SENB cross-section, where the section was taken from the center of the SENB bar by cutting the bar lengthwise parallel to the crack propagation direction and normal to the plane containing the cracks. Multiple cracks are seen in the CVI–SiC matrix material, but little fiber breakage is observed, even at the root of the notch where the crack opening displacement would be the largest. A schematic of the cracked damage zone is shown in Fig. 6b. A single Hi-C specimen was also sectioned and showed similar multiple cracking. Several specimens were tested, unloaded while at temperature, cooled with no applied load, and then sectioned for optical microscopy of the cracks and of the damage zone. Data from these specimens, in