Availableonlineatwww.sciencedirect.com SCIENCE DIRECT● E噩≈S ELSEVIER Journal of the European Ceramic Society 25(2005)1629-1636 www.elsevier.com/locate/jeurceramsoc Delayed failure of ceramic matrix composites in tension at elevated temperatures Sung R Choi*, Narottam P. Bansal, Michael J. Verrilli Received 22 February 2004; received in revised form 3 May 2004; accepted 23 May 2004 Available online 14 august 2004 Abstract Ultimate tensile strength of five different continuous fiber-reinforced ceramic matrix composites(CMCs), including SiC /BSAS(two dimensional(2D),2 types), SiC/MAS(2D), SiC,/SiC(2D), and C/SiC (2D, 2 types), was determined as a function of test rate at 1 100-1200C in air. All five CMCs exhibited a significant dependency of ultimate tensile strength on test rate such that the ultimate tensile strengt decreased with decreasing test rate. The dependency of ultimate tensile strength on test rate, the applicability of preload technique, and the predictability of life from one loading configuration(constant stress-rate loading) to another(constant stress loading) all suggested that the overall, phenomenological delayed failure of the CMCs would be governed by a power-law type of slow crack growth C2004 Elsevier Ltd. All rights reserved Keywords: Composites; Strength; Stress-rupture testing: Slow crack growth, Lifetime; SiC/SIC 1. Introduction This paper, as a continuation of the previous study describes delayed failure behavior of five different fiber- The successful development and design of continuous reinforced CMCs at 1 100-1200 C in air, including three Sic fiber-reinforced ceramic matrix composites(CMCs)are de pendent on thorough understanding of their basic properties fiber-reinforced CMCs(SiCr/BSAS(2D, 2 types), SiCr/MAS (D)and SiCe/sic(2D woven)) and one carbon fiber- such as deformation, fracture, and delayed failure(slow crack reinforced CMC(C, /SiC(2D woven, 2 types ).Ultimate ten- rowth, fatigue, or damage evolution/accumulation) behav- sile strength of each composite was determined as a function ior. Particularly, complete evaluation and characterization of of test rate in constant stress-rate testing and its rate depen environmental conditions is a prerequisite to ensure acura- dency was analyzed with a power-law type of slow crack life prediction of structural components stand the governing failure mechanism(s)of the composites In a previous study, ultimate tensile strength of three Finally, the results of elevated-temperature constant stress Sic fiber-reinforced CMCs(SiC/CAs, SiCe/MAS-5, and (stress rupture")testing were obtained for SiC:/BSAS and SiCr/SiC)at 1100-1200oC in air was found to be a strong compared with those of constant stress-rate testing to ve function of test rate. This rate dependency of ultimate tensile ify the overall failure mechanism(s)of the composite and to strength, in conjunction with the additional results of both accelerated and stress rupture testing, was found to be at- establish constant stress-rate testing as a means of life pre- diction test methodology for CmCs tributed to a power-law type of slow crack growth or damage evolution/accumulation that described adequately the phe- nomenological time-dependent behavior of the CMCs 2. Experimental procedure Corresponding author. Tel: +1 216 433 8366; fax: +1 216 433 8366 E-mail address: sung. r choi @grc. nasa. gov(SR Choi). Five different CMCs-four SiC fiber-reinforced and one NASA Resident Principle Scientist carbon fiber-reinforced-were used in this study, includ- 0955-2219/S-see front matter c 2004 Elsevier Ltd. All rights reserved doi: 10.1016/j-jeurceramsoc. 2004.05.024
Journal of the European Ceramic Society 25 (2005) 1629–1636 Delayed failure of ceramic matrix composites in tension at elevated temperatures Sung R. Choi∗,1, Narottam P. Bansal, Michael J. Verrilli National Aeronautics and Space Administration, John H. Glenn Research Center, Cleveland, OH 44135, USA Received 22 February 2004; received in revised form 3 May 2004; accepted 23 May 2004 Available online 14 August 2004 Abstract Ultimate tensile strength of five different continuous fiber-reinforced ceramic matrix composites (CMCs), including SiCf/BSAS (two dimensional (2D), 2 types), SiCf/MAS (2D), SiCf/SiC (2D), and Cf/SiC (2D, 2 types), was determined as a function of test rate at 1100–1200 ◦C in air. All five CMCs exhibited a significant dependency of ultimate tensile strength on test rate such that the ultimate tensile strength decreased with decreasing test rate. The dependency of ultimate tensile strength on test rate, the applicability of preload technique, and the predictability of life from one loading configuration (constant stress-rate loading) to another (constant stress loading) all suggested that the overall, phenomenological delayed failure of the CMCs would be governed by a power-law type of slow crack growth. © 2004 Elsevier Ltd. All rights reserved. Keywords: Composites; Strength; Stress–rupture testing; Slow crack growth; Lifetime; SiC/SiC 1. Introduction The successful development and design of continuous fiber-reinforced ceramic matrix composites (CMCs) are dependent on thorough understanding of their basic properties such as deformation, fracture, and delayed failure (slow crack growth, fatigue, or damage evolution/accumulation) behavior. Particularly, complete evaluation and characterization of delayed failure behavior of CMCs under specified loadingenvironmental conditions is a prerequisite to ensure accurate life prediction of structural components. In a previous study,1 ultimate tensile strength of three SiC fiber-reinforced CMCs (SiCf/CAS, SiCf/MAS-5, and SiCf/SiC) at 1100–1200 ◦C in air was found to be a strong function of test rate. This rate dependency of ultimate tensile strength, in conjunction with the additional results of both accelerated and stress rupture testing, was found to be attributed to a power-law type of slow crack growth or damage evolution/accumulation that described adequately the phenomenological time-dependent behavior of the CMCs. ∗ Corresponding author. Tel.: +1 216 433 8366; fax: +1 216 433 8366. E-mail address: sung.r.choi@grc.nasa.gov (S.R. Choi). 1 NASA Resident Principle Scientist. This paper, as a continuation of the previous study,1 describes delayed failure behavior of five different fiberreinforced CMCs at 1100–1200 ◦C in air, including three SiC fiber-reinforced CMCs (SiCf/BSAS (2D, 2 types), SiCf/MAS (1D) and SiCf/SiC (2D woven)) and one carbon fiberreinforced CMC (Cf/SiC (2D woven, 2 types)). Ultimate tensile strength of each composite was determined as a function of test rate in constant stress-rate testing and its rate dependency was analyzed with a power-law type of slow crack growth. Preload testing was also carried out to better understand the governing failure mechanism(s) of the composites. Finally, the results of elevated-temperature constant stress (“stress rupture”) testing were obtained for SiCf/BSAS and compared with those of constant stress-rate testing to verify the overall failure mechanism(s) of the composite and to establish constant stress-rate testing as a means of life prediction test methodology for CMCs. 2. Experimental procedure Five different CMCs—four SiC fiber-reinforced and one carbon fiber-reinforced—were used in this study, includ- 0955-2219/$ – see front matter © 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.jeurceramsoc.2004.05.024
S.R. Choi et al. /Journal of the European Ceramic Society 25(2005)1629-1636 ing Nicalon or Hi-Nicalon M SiC crossply(2D)fiber- Cooled Wedge Grip reinforced barium strontium aluminosilicate(designated Test Specimen Nicalon/BSAS and Hi-Nicalon/BSAS), Nicalon'unidirec tionally (ID) reinforced magnesium aluminosilicate(des- ignated SiCr/MAS), Nicalon"M plain-woven(2D)silicon carbide(designated SiCr/SiC: "enhanced"), and T-300 car- bon fiber-reinforced plain-woven(2D)silicon carbide(des- ignated C/SiC: "standard"and"enhanced ). The matrices Extensometer of the composites, except for Cr/SiC, were reinforced by Nicalon m or hi-nicalon m fibers with a fiber volume frac Induction Coil tion of about 0. 40. The unidirectional, crossply or plain- woven laminates of the SiC fiber-reinforced composites were typically 12-18 plies thick with a nominal thickness of around Cooled Wedge Grip 3-3.5 mm, depending on material. The Cr/Sic composite had a total of 29 plies, a fiber volume fraction of0. 46, and a nomi- nal laminate thickness of3 mm. The enhanced Sice/SiC com- Fig. 1. Schematics of experimental setup used in tensile testing for ceramic posite was modified from its standard matrix by a proprietary matrix composites at elevated temperatures in air process to increase the oxidation resistance of the composite Sic was also chemically vapor deposited on the composite sile testing was performed in accordance with ASTM Test panels to cover the residual porosity. The enhanced Cf/SiC composite had a boron carbide that was introduced into the Preload or accelerated tensile testing, primarily appl composite prior to deposition of pyrolytic carbon interface to protect the carbon fibers from oxidation Both Nicalon/bSAs conducted at 1 100 or 1200oC the lowest test rate of and Hi-Nicalon/BSAS were fabricated at NASA Glenn Re- 0.005 MPa/s in an attempt to better understand the govern- search Center, 2 SiCr/MAS by Corning, Inc. Corning, NY), 3 ng failure mechanism(s)of the composites. A predetermined SiC /Sic by duPont Company(Newark, DE), 3 and C/Sic preload, corresponding to an% of ultimate tensile strength by Honeywell Advanced Composites, Inc.(Newark, DE). 4 of each composite that was determined at 0.005 MPa/s with cessing can be found elsewhere. 2y omposites and their pro- no preload, was applied quickly (100 MPa/s) to the test Detailed information regarding the specimen prior to testing, and monotonic tensile testing at The dogboned tensile test specimens measuring 152mm 0.005 MPa/s started and continued until the test specimen (length) by 12.7 mm(width) were machined from the com- failed. The corresponding ultimate tensile strength was de- posite laminates, with the gage section of about 30mm termined. One test specimen was used in preload testing for long, 10 mm wide, and 3.0-3.5 mm thick(as-furnished) each composite The Cr/Sic test specimens were supplied with a notch ma- Constant stress("stress rupture") tensile testing was con- chined(in depth=2.5 mm and root radius 1.2 mm)at ducted at 100C in air for the Nicalon/BSAS (wIth two ne side of gage section at the longitudinal center of each different batches"A"and"B") composite using test speci test specimen. After machining, the Ce/Sic test specimens mens with the same geometry and the same test frame and were seal coated with SiC by the chemical vapor infiltration equipment that were employed in constant stress-rate tensile testing. The limited availability of test materials confined the Monotonic tensile testing was conducted in air at 1100oC testing to three to four test specimens, depending on batch for Nicalon/BSAS, Hi-Nicalon/BSAS and SiC/MAS-5 and Two to three different constant stresses were applied to test at 1200C for SiC /SiC and Cr/SiC, using a servohydraulic specimens and corresponding times to failure were deter test frame(Model 8501, Instron, Canton, MA). A schematic test setup is shown in Fig. 1. Each test specimen, located in- side of a SiC susceptor via two upper and lower water-cooled hydraulic grips, was induction-heated by radiation through 3. Results and discussion 15-kw power supply. Two high-temperature extensometers were placed on edges of each test specimen to measure ten- 3.. Ultimate tensile strength sile strain. Detailed descriptions on test setup and induction heating equipment were found in a previous study. 'A total The results of ultimate tensile strength as a function of of three different test rates in force control, corresponding test rate determined for the aforementioned CMCs are pre- to stress rates of 5, 0.16, and 0.005 MPa/s, were employed sented in Fig. 2, where ultimate tensile strength was plot for a given composite. This test method, when applied to ad- ted as a function of applied stress rate using log-log scales vanced monolithic ceramics is called constant stress rate o Each solid line in the figure represents the best-fit regres- dynamic fatigue"testing. Typically, one to three test speci- sion based on the log(ultimate tensile strength) versus log mens were tested at each test rate for a given composite. Ten-(applied stress rate)relation. The strength data determined
1630 S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 ing NicalonTM or Hi-NicalonTM SiC crossply (2D) fiberreinforced barium strontium aluminosilicate (designated Nicalon/BSAS and Hi-Nicalon/BSAS), NicalonTM unidirectionally (1D) reinforced magnesium aluminosilicate (designated SiCf/MAS), NicalonTM plain-woven (2D) silicon carbide (designated SiCf/SiC: “enhanced”), and T-300 carbon fiber-reinforced plain-woven (2D) silicon carbide (designated Cf/SiC: “standard” and “enhanced”). The matrices of the composites, except for Cf/SiC, were reinforced by NicalonTM or Hi-NicalonTM fibers with a fiber volume fraction of about 0.40. The unidirectional, crossply or plainwoven laminates of the SiC fiber-reinforced composites were typically 12–18 plies thick with a nominal thickness of around 3–3.5 mm, depending on material. The Cf/SiC composite had a total of 29 plies, a fiber volume fraction of 0.46, and a nominal laminate thickness of 3 mm. The enhanced SiCf/SiC composite was modified from its standard matrix by a proprietary process to increase the oxidation resistance of the composite. SiC was also chemically vapor deposited on the composite panels to cover the residual porosity. The enhanced Cf/SiC composite had a boron carbide that was introduced into the composite prior to deposition of pyrolytic carbon interface to protect the carbon fibers from oxidation. Both Nicalon/BSAS and Hi-Nicalon/BSAS were fabricated at NASA Glenn Research Center,2 SiCf/MAS by Corning, Inc. (Corning, NY),3 SiCf/SiC by DuPont Company (Newark, DE),3 and Cf/SiC by Honeywell Advanced Composites, Inc. (Newark, DE).4 Detailed information regarding the composites and their processing can be found elsewhere.2–4 The dogboned tensile test specimens measuring 152 mm (length) by 12.7 mm (width) were machined from the composite laminates, with the gage section of about 30 mm long, 10 mm wide, and 3.0–3.5 mm thick (as-furnished). The Cf/SiC test specimens were supplied with a notch machined (in depth = 2.5 mm and root radius = 1.2 mm) at one side of gage section at the longitudinal center of each test specimen. After machining, the Cf/SiC test specimens were seal coated with SiC by the chemical vapor infiltration method.4 Monotonic tensile testing was conducted in air at 1100 ◦C for Nicalon/BSAS, Hi-Nicalon/BSAS and SiCf/MAS-5 and at 1200 ◦C for SiCf/SiC and Cf/SiC, using a servohydraulic test frame (Model 8501, Instron, Canton, MA). A schematic test setup is shown in Fig. 1. Each test specimen, located inside of a SiC susceptor via two upper and lower water-cooled hydraulic grips, was induction-heated by radiation through a 15-kW power supply. Two high-temperature extensometers were placed on edges of each test specimen to measure tensile strain. Detailed descriptions on test setup and induction heating equipment were found in a previous study.3 A total of three different test rates in force control, corresponding to stress rates of 5, 0.16, and 0.005 MPa/s, were employed for a given composite. This test method, when applied to advanced monolithic ceramics, is called constant stress rate or “dynamic fatigue” testing. Typically, one to three test specimens were tested at each test rate for a given composite. TenFig. 1. Schematics of experimental setup used in tensile testing for ceramic matrix composites at elevated temperatures in air. sile testing was performed in accordance with ASTM Test Standard C 1359.5 Preload or accelerated tensile testing, primarily applied to monolithic ceramics in order to save test time,6 was also conducted at 1100 or 1200 ◦C using the lowest test rate of 0.005 MPa/s in an attempt to better understand the governing failure mechanism(s) of the composites. A predetermined preload, corresponding to an 80% of ultimate tensile strength of each composite that was determined at 0.005 MPa/s with no preload, was applied quickly (∼100 MPa/s) to the test specimen prior to testing, and monotonic tensile testing at 0.005 MPa/s started and continued until the test specimen failed. The corresponding ultimate tensile strength was determined. One test specimen was used in preload testing for each composite. Constant stress (“stress rupture”) tensile testing was conducted at 1100 ◦C in air for the Nicalon/BSAS (with two different batches “A” and “B”) composite using test specimens with the same geometry and the same test frame and equipment that were employed in constant stress-rate tensile testing. The limited availability of test materials confined the testing to three to four test specimens, depending on batch. Two to three different constant stresses were applied to test specimens and corresponding times to failure were determined. 3. Results and discussion 3.1. Ultimate tensile strength The results of ultimate tensile strength as a function of test rate determined for the aforementioned CMCs are presented in Fig. 2, where ultimate tensile strength was plotted as a function of applied stress rate using log–log scales. Each solid line in the figure represents the best-fit regression based on the log (ultimate tensile strength) versus log (applied stress rate) relation. The strength data determined
S.R. Choi et al. /Journal of the European Ceramic Sociery 25(2005)1629-1636 Hi-Nicalon/BSAS (2-D b) 200 n=13 80%P一△ BATCH B 60 0410310210110 10410310210110010110 Applied stress rate, o[MPa/s] Applied stress rate o [MPa/s] SiC!/MAS(1-D SiC+SiC (2-D; Enhanced Tension/1200°c 5 E 700F Tension/1100@C 80%PL 9 13101010101 10410310210110°101102 Applied stress rate, d [MPa/s CsIc(2-D) 300 Tension/1200c 80%PL Applied stress rate, o[MPa/s Fig. 2. Results of ultimate tensile strength as a function of applied stress rate determined for(a)NicalonBSAS, (b) Hi-Nicalon/BSAS, (c)SiCr/MAS, (d) SiCr/SiC, and (e)Cr/Sic ceramic matrix composites at elevated temperatures in air. The solid lines represent the best-fit regression lines based on Eq (1). The delayed failure(or slow crack growth) parameter n was also included for each material with an 80% preload were also included. The decrease in Nicalon/BSAS exhibited higher strength than Nicalon/BSAS ultimate tensile strength with decreasing stress rate, which at test rates >0.16 MPa/s, the former exhibited a greater sus- represents a susceptibility to delayed failure(or slow crack ceptibility to delayed failure than the latter, viewed from growth or damage accumulation), was significant for all the the degree of strength degradation with decreasing test rate omposite materials tested, consistent with the previous re- Also note that the enhanced Cr/SiC composite showed lower sults determined with other CMCs such as ID SiCf/CAs strength at high test rates but greater resistance to delayed calcium aluminosilicate ), 2D SiC:/MAS-5, and 2D woven failure than the standard counterpart, thus achieving more SiCr/SiC (standard). The strength degradation was about improved resistance to strength-environmental degradation 3, 59, 48, and 28%, respectively, for Nicalon/BSAS(batch by the boron carbide enhancement A), Hi-Nicalon/BSAS, SiCr/MAS, and SiCr/SiC, when the Typical examples of fracture pattern of each composite stress rate decreased from the highest(=5 MPa/s)to the low- tested at the highest(=S MPa/s)and lowest(=0.005 MPa/s) est(=0.005 MPa/s). The corresponding strength degradation test rates are shown in Fig. 3. The mode of fracture for Cr/Sic was 61 and 30%, respectively, for the standard for both Nicalon/BSAS and Hi-Nicalon/BSAS composites and enhanced composites. It is noteworthy that although Hi showed fiber pullout with zigzag matrix cracking through the
S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 1631 Fig. 2. Results of ultimate tensile strength as a function of applied stress rate determined for (a) Nicalon/BSAS, (b) Hi-Nicalon/BSAS, (c) SiCf/MAS, (d) SiCf/SiC, and (e) Cf/SiC ceramic matrix composites at elevated temperatures in air. The solid lines represent the best-fit regression lines based on Eq. (1). The delayed failure (or slow crack growth) parameter n was also included for each material. with an 80% preload were also included. The decrease in ultimate tensile strength with decreasing stress rate, which represents a susceptibility to delayed failure (or slow crack growth or damage accumulation), was significant for all the composite materials tested, consistent with the previous results determined with other CMCs such as 1D SiCf/CAS (calcium aluminosilocate), 2D SiCf/MAS-5, and 2D woven SiCf/SiC (standard).1 The strength degradation was about 43, 59, 48, and 28%, respectively, for Nicalon/BSAS (batch A), Hi-Nicalon/BSAS, SiCf/MAS, and SiCf/SiC, when the stress rate decreased from the highest (=5 MPa/s) to the lowest (=0.005 MPa/s). The corresponding strength degradation for Cf/SiC was 61 and 30%, respectively, for the standard and enhanced composites. It is noteworthy that although HiNicalon/BSAS exhibited higher strength than Nicalon/BSAS at test rates ≥0.16 MPa/s, the former exhibited a greater susceptibility to delayed failure than the latter, viewed from the degree of strength degradation with decreasing test rate. Also note that the enhanced Cf/SiC composite showed lower strength at high test rates but greater resistance to delayed failure than the standard counterpart, thus achieving more improved resistance to strength–environmental degradation by the boron carbide enhancement. Typical examples of fracture pattern of each composite tested at the highest (=5 MPa/s) and lowest (=0.005 MPa/s) test rates are shown in Fig. 3. The mode of fracture for both Nicalon/BSAS and Hi-Nicalon/BSAS composites showed fiber pullout with zigzag matrix cracking through the
S.R. Choi et al. /Journal of the European Ceramic Society 25(2005)1629-1636 b) Fig 3. Fracture patterns for(a) Nicalon/BSAS, (b )SiCr/MAS, (c)SiCr/SiC (enhanced), and(d)Cr/SiC (enhanced")ceramic matrix composites tested in tension at elevated temperatures in air. The upper and lower pictures for a given composite material indicate the specimens tested at the lowest (=0.005 MPa/ and the highest(=5 MPa/s)test rates, respectively specimen-thickness direction. A change in surface(matrix) The strength dependency on test rate exhibited by these color of test specimens from dark grey to white was more pro- composites at elevated temperatures(see Fig. 2)is very simi nounced at 0.005 MPa/s than 5 MPa/s, an evidence of more lar to that commonly observed in advanced monolithic ceram aggressive high-temperature reaction/oxidation involved cs at elevated temperatures. The strength degradation with the lower test rate. attributed to increased test time. fracture decreasing stress rate in monolithic ceramics is known to oc. patterns for the SiCr/MAS composite indicated some fiber cur by a slow crack growth process(also known as delayed pullout with jagged faceted matrix cracking often propagat- failure, subcritical crack growth or fatigue)and is expressed ng along the test-specimen length One specimen tested at as follows the fast test rate of 5 MPa/s failed close to the transition and grip regions. No significant difference in the mode of frac (1) ture was observed between SiCr/SiC (enhanced)and Cr/SiC (standard or enhanced), where almost all the specimens tested where n and D are slow crack growth parameters, and of at either a high or low test rate exhibited relatively flat, straight (MPa)and o(MPa/s)are fracture strength and applied stress fracture surfaces with little fiber pullout, termed brittle frac- rate, respectively. Eq- (1)is based on the conventional power ture. A similar brittle mode of fracture was also observed law crack velocity formulation as expressed previously for the 2D standard SiCf/Sic composite. Black KI to bluish discoloration in the heated region of tested spec U=A mens was obvious for either standard or enhanced Cr/sic. accompanying a weight loss after testing due to oxidation: where u, Kl, and Kle are crack velocity, mode I stress intensity he more weight loss occurred at the lower test rate, and vice factor and fracture toughness, respectively. A is also called versa. Detailed oxidation and stress rupture/life behaviors of slow crack growth parameter. The parameter D is associated this Cr/Sic composite system have been explored previously with n, A, KIc, and inert strength of a material. The parameters in a low partial pressure of oxygen environment.. 7,8 nand D in Eq (1)can be obtained by a linear regression
1632 S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 Fig. 3. Fracture patterns for (a) Nicalon/BSAS, (b)SiCf/MAS, (c) SiCf/SiC (“enhanced”), and (d) Cf/SiC (“enhanced”) ceramic matrix composites tested in tension at elevated temperatures in air. The upper and lower pictures for a given composite material indicate the specimens tested at the lowest (=0.005 MPa/s) and the highest (=5 MPa/s) test rates, respectively. specimen-thickness direction. A change in surface (matrix) color of test specimens from dark grey to white was more pronounced at 0.005 MPa/s than 5 MPa/s, an evidence of more aggressive high-temperature reaction/oxidation involved at the lower test rate, attributed to increased test time. Fracture patterns for the SiCf/MAS composite indicated some fiber pullout with jagged faceted matrix cracking often propagating along the test-specimen length. One specimen tested at the fast test rate of 5 MPa/s failed close to the transition and grip regions. No significant difference in the mode of fracture was observed between SiCf/SiC (enhanced) and Cf/SiC (standard or enhanced), where almost all the specimens tested at either a high or low test rate exhibited relatively flat, straight fracture surfaces with little fiber pullout, termed brittle fracture. A similar brittle mode of fracture was also observed previously for the 2D standard SiCf/SiC composite.1 Black to bluish discoloration in the heated region of tested specimens was obvious for either standard or enhanced Cf/SiC, accompanying a weight loss after testing due to oxidation: the more weight loss occurred at the lower test rate, and vice versa. Detailed oxidation and stress rupture/life behaviors of this Cf/SiC composite system have been explored previously in a low partial pressure of oxygen environment.4,7,8 The strength dependency on test rate exhibited by these composites at elevated temperatures (see Fig. 2) is very similar to that commonly observed in advanced monolithic ceramics at elevated temperatures. The strength degradation with decreasing stress rate in monolithic ceramics is known to occur by a slow crack growth process (also known as delayed failure, subcritical crack growth or fatigue) and is expressed as follows:9–11 log σf = 1 n + 1 log σ˙ + log D (1) where n and D are slow crack growth parameters, and σf (MPa) and σ˙ (MPa/s) are fracture strength and applied stress rate, respectively. Eq. (1) is based on the conventional powerlaw crack velocity formulation as expressed: v = A KI KIc n (2) where v, KI, and KIc are crack velocity, mode I stress intensity factor and fracture toughness, respectively. A is also called slow crack growth parameter. The parameter D is associated with n, A,KIc, and inert strength of a material. The parameters n and D in Eq. (1) can be obtained by a linear regression anal-
S.R. Choi et al. /Journal of the European Ceramic Sociery 25(2005)1629-1636 633 ysis from slope and intercept, respectively, when log(fracture ength)is plotted as a function of log(applied stress rate) Constant stress-rate(or called dynamic fatigue)testing based on Eq (1)has been established as ASTM Test Methods to determine slow crack growth parameters of advanced mono- Sic/CAS lithic ceramics at ambient and elevated temperatures. 0, II Notwithstanding the limited number of test specimens used, the data fit to Eq. (1) was very reasonable for the cur rent CMCs with the coefficients of correlation in regression 98 all greater than 0.930. This implies that delayed failure the composites could be described by the power-law type of slow crack growth formulation, Eq(2). With this in mind 10010110210 the apparent delayed-failure parameters n and D for the composites were determined based on Eq(1)using the ex- Applied stress rate, d[MPa/s perimental data shown in Fig. 2(with the units of of in MPa Fig 4. Ultimate tensile strength as a function of test rates determined for us determined were n= ID SiC./ CAS at 1100oC, 2D SiC /MAS-5 at 1100C, and woven SiC /Sic 13 and D= 127, n=7 and D= 188, n=13 and D= 367, ( "standard") at 1200.C in air from previous studies and n= 20 and D= 160, respectively, for Nicalon/BSAS (batch A), Hi-Nicalon/BSAS, SiCr/MAS-5, and SiCr/SiC the degree of strength degradation with respect to test rate. As The prime was used here for composites to distinguish them oxidation prevails into the material system, porosity increases from monolithic ceramic counterparts. The apparent param- and the effective number and sizes of load-bearing fibers de- eters for the Cr/sic composite were n=6 and D =196 crease. This oxidation-induced damage would be considered and n=18 and D=166, respectively, for the standard and to be equivalent to crack-like flaws growing through matri- enhanced versions. The value of n represents a measure of ces and fibers from a fracture-mechanics point of view.The susceptibility to delayed failure, and is typically categorized equivalent crack propagates under a driving force(K1)based in brittle materials such that the susceptibility is very high for on Eq (2)so that the resulting strength follows in accordance n=20, intermediate for n=30-50, and very low for n>50. with Eq(1). Note that oxidation-induced damage increases Hence, the current composites exhibited a significant suscep- with decreasing test rate since more time is available for ox- tibility to delayed failure as compared with monolithic coun- idation at lower test rate. and vice versa. In other words at terparts such as silicon nitrides and silicon carbides which are faster test rates, an equivalent crack has little time to grow. typically in the range of n>20 at temperatures >1200C. 12 resulting in higher strength; whereas at lower test rates the Similar results showing greater susceptibility of CMCs to de- crack has longer time to grow appreciably, thereby yielding layed failure at elevated temperatures were also found from lower strength. However, oxidation was not likely a unique the previous study in ID SICf/CAS, 2D SICr/MAS, and 2D source of final fracture responsible in the Cr/SiC material system, as will be discussed in Section 3.2 from n=6-18, as depicted in Fig. 4 Unlike the other CMCs, the Cr/Sic composite, as men- tioned before, was subjected to significant oxidation of car- 3.2. Preload tests results bon fibers, resulting in material loss. Therefore, strength degradation was increased, attributed to decreasing fiber vol The results of preload tests carried out at 0.005 MPa/s with ume fraction and subsequently increasing porosity, as the test an 80% preload are presented in Fig.5, where ultimate tensile rate decreased. The strength degradation due to oxidation fo strength was plotted against preload factor(a=0.8)for each the Cr/Sic composite has been described based on the results composite. The preload factor a is defined such that a preload of stress rupture through a finite difference model on oxy ensile stress(op)applied to a test specimen prior to testing is gen concentrations and carbon consumption, 13 Although this normalized with respect to the ultimate tensile strength (ae) stress-oxidation model would give a better physical explana- tion, the phenomenological power-law formulation used in this study still provides a simple, convenient way to quantify (3) A The number of test specimens used in this study, one to three at each tes For advanced monolithic ceramics whose delayed failure rate,would be insufficient to determine reliable delayed failure paramete is governed by the power-law slow crack growth formulation However, considering a small scatter in ultimate strength of (Eq(2), it has been shown that fracture strength is a func- CMCs, typically with coefficients of variation <5%, the current delayed ion of preload factor and slow crack growth parameter n as failure parameters provided will not be changed too much even with a large follows 6. 10,11 number of test specimens; hence, the variation of the n and D to the number of test specimen would be expected to be minimal and statistically insignif =o(1+a+1)1(m+1)
S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 1633 ysis from slope and intercept, respectively, when log (fracture strength) is plotted as a function of log (applied stress rate). Constant stress-rate (or called dynamic fatigue) testing based on Eq. (1) has been established as ASTM Test Methods to determine slow crack growth parameters of advanced monolithic ceramics at ambient and elevated temperatures.10,11 Notwithstanding the limited number of test specimens used, the data fit to Eq. (1) was very reasonable for the current CMCs with the coefficients of correlation in regression all greater than 0.930. This implies that delayed failure of the composites could be described by the power-law type of slow crack growth formulation, Eq. (2). With this in mind, the apparent delayed–failure parameters n and D for the composites were determined based on Eq. (1) using the experimental data shown in Fig. 2 (with the units of σf in MPa and σ˙ in MPa/s). The parameters1 thus determined were n = 13 and D = 127, n = 7 and D = 188, n = 13 and D = 367, and n = 20 and D = 160, respectively, for Nicalon/BSAS (batch A), Hi-Nicalon/BSAS, SiCf/MAS-5, and SiCf/SiC. The prime was used here for composites to distinguish them from monolithic ceramic counterparts. The apparent parameters for the Cf/SiC composite were n = 6 and D = 196, and n = 18 and D = 166, respectively, for the standard and enhanced versions. The value of n represents a measure of susceptibility to delayed failure, and is typically categorized in brittle materials such that the susceptibility is very high for n ≤ 20, intermediate for n = 30–50, and very low for n > 50. Hence, the current composites exhibited a significant susceptibility to delayed failure as compared with monolithic counterparts such as silicon nitrides and silicon carbides which are typically in the range of n > 20 at temperatures ≥1200 ◦C.12 Similar results showing greater susceptibility of CMCs to delayed failure at elevated temperatures were also found from the previous study1 in 1D SiCf/CAS, 2D SiCf/MAS, and 2D woven SiCf/SiC (standard) composites with n values ranging from n = 6–18, as depicted in Fig. 4. Unlike the other CMCs, the Cf/SiC composite, as mentioned before, was subjected to significant oxidation of carbon fibers, resulting in material loss. Therefore, strength degradation was increased, attributed to decreasing fiber volume fraction and subsequently increasing porosity, as the test rate decreased. The strength degradation due to oxidation for the Cf/SiC composite has been described based on the results of stress rupture through a finite difference model on oxygen concentrations and carbon consumption.13 Although this stress-oxidation model would give a better physical explanation, the phenomenological power-law formulation used in this study still provides a simple, convenient way to quantify 1 The number of test specimens used in this study, one to three at each test rate, would be insufficient to determine reliable delayed failure parameters. However, considering a relatively small scatter in ultimate strength of many CMCs, typically with coefficients of variation ≤5%, the current delayedfailure parameters provided will not be changed too much even with a large number of test specimens; hence, the variation of the n and D to the number of test specimen would be expected to be minimal and statistically insignificant as well. Fig. 4. Ultimate tensile strength as a function of test rates determined for 1D SiCf/CAS at 1100 ◦C, 2D SiCf/MAS-5 at 1100 ◦C, and woven SiCf/SiC (“standard”) at 1200 ◦C in air from previous studies.1 the degree of strength degradation with respect to test rate. As oxidation prevails into the material system, porosity increases and the effective number and sizes of load-bearing fibers decrease. This oxidation-induced damage would be considered to be equivalent to crack-like flaws growing through matrices and fibers from a fracture-mechanics point of view. The equivalent crack propagates under a driving force (KI) based on Eq. (2)so that the resulting strength follows in accordance with Eq. (1). Note that oxidation-induced damage increases with decreasing test rate since more time is available for oxidation at lower test rate, and vice versa.7 In other words, at faster test rates, an equivalent crack has little time to grow, resulting in higher strength; whereas, at lower test rates, the crack has longer time to grow appreciably, thereby yielding lower strength. However, oxidation was not likely a unique source of final fracture responsible in the Cf/SiC material system, as will be discussed in Section 3.2. 3.2. Preload tests results The results of preload tests carried out at 0.005 MPa/s with an 80% preload are presented in Fig. 5, where ultimate tensile strength was plotted against preload factor (α = 0.8) for each composite. The preload factor α is defined such that a preload tensile stress (σp) applied to a test specimen prior to testing is normalized with respect to the ultimate tensile strength (σf) with no preload6,10,11 α = σp σf (3) For advanced monolithic ceramics whose delayed failure is governed by the power-law slow crack growth formulation (Eq. (2)), it has been shown that fracture strength is a function of preload factor and slow crack growth parameter n as follows6,10,11 σfp = σf(1 + αn+1) 1/(n+1) (4)
634 S.R. Choi et al. /Journal of the European Ceramic Society 25(2005)1629-1636 O Nicalon/BSAS W200 △H- Nicalon/BSAS Theoretical 0 20 Prelaod factor, a x100 [9 Prelaod factor, a x100 [% C/sic (2-D) ension1200°c Tension1200°c Theoretical Theoretical 100 10 020406080100 Prelaod factor, a x100 [ Fig. 5. Results of preload tests, plotted with ultimate tensile strength as a function of preload factor for(a) Nicalon/BSAS and Hi-Nicalon/BSAS, (b)SiCr/MAs, (c)SiCr/SiC (enhanced), and(d)Cr/SiC ("standard"and"enhanced")ceramic matrix composites at elevated temperatures in air. A theoretical prediction based on Eq (4)is included for comparison for each composite material where ofp is ultimate tensile strength with a preload stress or test time was greater than 80% of fracture stress or a is in the range of0 s a < 1. It is noted from Eq.(4) total test time that ultimate tensile strength under preload is more sensi- The insignificant difference in ultimate tensile strength of tive to higher preload factor a and lower n value, because the Cr/sic composite between two preloads of 0 and 80% of much augmented delayed failure occurring under these is particularly noteworthy. The total test time at no preload conditions was about 5-7 h for the Cr/SiC composite(standard and en- Each solid line in Fig. 5 indicates the theoretical predic- hanced), while the respective total test time with an 80% tion based on Eq(4), together with the estimated value of preload was about 1-2h. Despite this appreciable difference n and the ultimate tensile strength with no preload. The pre- in test time(or exposure time for oxidation) between the two diction, despite a limited number of test specimens used preloads, the resulting strength difference was minimal. This in good agreement with the experimental data except for the suggests that oxidation would not be a sole source of the SiCH/MAS composite. Note that Eq(4)was derived based composite failure as well as of the rate dependency. A further on the power-law, slow crack growth formulation, Eq.(2) tudy using increased number of test specimens would reveal Therefore, the reasonable applicability of the preload anal- more detail aspects of failure mechanism(s) involved in the sis to the current composites suggests that delayed failure Cr/SiC composite. However, it is important to state at this process of these composites would be the one governed by point that apart from detailed understanding of a complex the power-law type of slow crack growth, as expressed in oxidation kinetics associated with the C/Sic composite, the from the figure that the overall difference in ultimate tensile results of both ultimate tensile strength and preload tes E s Eq.(2). This is consistent with the observations of the pre- composite failure can be described phenomenologically vious preload studies using other CMCs. It is also noted the simple power-law formulation of Eq(2), based on the trength between two preloads (a=0 and 80%)was insignif- icant, resulting in a considerable saving (80%)of test time 3.3. Constant stress(stress rupture)tests by applying the 80% preload. This indicates that any signifi cant crack growth that would control ultimate tensile strength A summary of the results of constant stress(stress rupture of a composite did not occur even when the applied stress to testing for the Nicalon/BSAS composite(batches A and B)at test specimen during test reached up to 80%of its fracture 1100C is presented in Fig. 6, where time to failure was plot tress. Conversely, the crack growth or damage to control fi- ted against applied stress in log-log scales. a decrease intime nal catastrophic failure would have occurred when applied to failure with increasing applied stress, which represents a
1634 S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 Fig. 5. Results of preload tests, plotted with ultimate tensile strength as a function of preload factor for (a) Nicalon/BSAS and Hi-Nicalon/BSAS, (b) SiCf/MAS, (c) SiCf/SiC (“enhanced”), and (d) Cf/SiC (“standard” and “enhanced”) ceramic matrix composites at elevated temperatures in air. A theoretical prediction based on Eq. (4) is included for comparison for each composite material. where σfp is ultimate tensile strength with a preload and α is in the range of 0 ≤ α < 1. It is noted from Eq. (4) that ultimate tensile strength under preload is more sensitive to higher preload factor α and lower n value, because of much augmented delayed failure occurring under these conditions. Each solid line in Fig. 5 indicates the theoretical prediction based on Eq. (4), together with the estimated value of n and the ultimate tensile strength with no preload. The prediction, despite a limited number of test specimens used, is in good agreement with the experimental data except for the SiCf/MAS composite. Note that Eq. (4) was derived based on the power-law, slow crack growth formulation, Eq. (2). Therefore, the reasonable applicability of the preload analysis to the current composites suggests that delayed failure process of these composites would be the one governed by the power-law type of slow crack growth, as expressed in Eq. (2). This is consistent with the observations of the previous preload studies using other CMCs.1 It is also noted from the figure that the overall difference in ultimate tensile strength between two preloads (α = 0 and 80%) was insignificant, resulting in a considerable saving (∼80%) of test time by applying the 80% preload. This indicates that any signifi- cant crack growth that would control ultimate tensile strength of a composite did not occur even when the applied stress to test specimen during test reached up to 80% of its fracture stress. Conversely, the crack growth or damage to control fi- nal catastrophic failure would have occurred when applied stress or test time was greater than 80% of fracture stress or total test time. The insignificant difference in ultimate tensile strength of the Cf/SiC composite between two preloads of 0 and 80% is particularly noteworthy. The total test time at no preload was about 5–7 h for the Cf/SiC composite (standard and enhanced), while the respective total test time with an 80% preload was about 1–2 h. Despite this appreciable difference in test time (or exposure time for oxidation) between the two preloads, the resulting strength difference was minimal. This suggests that oxidation would not be a sole source of the composite failure as well as of the rate dependency. A further study using increased number of test specimens would reveal more detail aspects of failure mechanism(s) involved in the Cf/SiC composite. However, it is important to state at this point that apart from detailed understanding of a complex oxidation kinetics associated with the Cf/SiC composite, the composite failure can be described phenomenologically by the simple power-law formulation of Eq. (2), based on the results of both ultimate tensile strength and preload tests. 3.3. Constant stress (stress rupture) tests A summary of the results of constant stress (stress rupture) testing for the Nicalon/BSAS composite (batches A and B) at 1100 ◦C is presented in Fig. 6, where time to failure was plotted against applied stress in log–log scales. A decrease in time to failure with increasing applied stress, which represents a
S.R. Choi et al. /Journal of the European Ceramic Sociery 25(2005)1629-1636 635 ing configuration (constant stress-rate testing)to another Nicalon/BSAS(2-D) (constant stress)for a selected composite all support that delayed failure of the composites was controlled by the BATCH 'A power-law type of slow crack growth(or damage evolu- tion/accumulation), confirmed with not only the current 80} BATCH日 CMCs but the previous CMCs. It is noted that despite many 0000 differences in their processing, architecture, microstructure, and interface, compared with monolithic ceramics, CMCs still exhibit delayed failure, similar in principle to mono- lithic counterparts. This is simply due to the fact that macro- 310410510° topically, a CMC is a composite composed of two or more delayed-failure susceptible monolithic materials(co Time to failure, t, [s] stituents of fibers, matrices, interfaces, etc. ) The vulnera- bility to delayed failure would be greater in composite be- Fig.6.Results of constant stress("stress rupture")testing for Nicalon/BSAS cause of its more likelihood of chance to environmental composite(batches""(open triangle)and"B(closed triangle)at 1100C in air. The solid lines represent the predictions made based on Eq (5)from exposure by its inherently more open, porous microstruc the results of constant stress-rate testing( Fig. 2) tures, as compared to dense monolithic counterparts. The vulnerability of a composite, of course, will be increased ible to delaved susceptibility to delayed failure, was evident for the com- failure posite. The mode of fracture was similar to that in constant A subsequent importance drawn from the results of this stress-rate testing showing some fiber pullout with jagged work is that constant stress-rate testing, commonly utilized matrix cracking through the specimen-thickness direction. in monolithic ceramics, could be applicable to CMCs to de- At lower applied stresses, however, the fracture surfaces of termine their delayed-failure(or life prediction) parameters, both batches were somewhat flat with decreased fiber pull- at least for a short range of lifetimes, consistent with the pre- out, a change in fracture mode more likely to brittle fracture. vious observation. The merits of constant stress-rate testing The lines in Fig. 6 indicate life prediction from the constant are enormous in terms of test simplicity and test economy stress-rate data. The prediction, primarily applied to brittle (short test time and less data scatter)over other stress rup. monolithic materials, was made using the following relation ture or cyclic fatigue testing. A simplistic, phenomenolog based on the power-law formulation of Eq(2). ical law of delayed failure was only explored in this study to develop lifetime prediction testing and methodology for (5) CMCs, without accounting for detailed failure mechanisms associated with matrix/fiber interaction, matrix cracking and ts effect on slow crack growth, delayed failure of sustain- where tf and o are time to failure and applied constant stress ing fibers, and creep-associated deformation. 4-I8A micr respectively. Use of Eq.(5)together with n and d de- scopic level of study on this subject is thus needed. Finally termined in constant stress-rate testing allows one to pre- the results of this work suggest that care must be exer dict life under constant stress(stress rupture)loading. Even cised when characterizing elevated-temperature strength of though the small number of test specimens was used here composite materials. This is due to the fact that elevated due to limited material availability, the prediction was in rea- temperature strength has a relative meaning if a material ex- sonable agreement with experimental data at least for the hibits rate dependency: the strength simply depends on which Nicalon/BSAS composite. This indicates that the governing test rate one chooses(see Fig. 2). Therefore, use of at least failure law of the Nicalon/BSAS composite was very simi- two test rates(high and low) is recommended to better chan lar either in constant stress-rate or in constant stress loading. acterize high-temperature strength behavior of a composite Since the prediction(Eq (5))was made based on Eq (2), the material governing delayed-failure mechanism of the Nicalon/BSAS would be the one controlled by the power-law type of slow crack growth(Eq (2). Other CMCs also showed reasonable 4. Conclusions agreement in life between constant stress-rate and constant- stress loading configurations Elevated-temperature ultimate tensile strength of five different continuous fiber-reinforced ceramic composites 3.4. Implications including Nicalon/BSAS (2D), Hi-Nicalon/BSAS (2D), SiCr/MAS (ID), SiCr/SiC(2D woven: enhanced), and As seen in the preceding sections 3.1-3.3. the stren Cr/sic (2D woven: standard and enhanced), exhibited a dependency on test rate, the applicability of the preload tech- strong dependency on test rate, consistent with the behav- nique, and the reasonable life prediction from one load- ior observed in other CMCs as well as in many advanced
S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 1635 Fig. 6. Results of constant stress (“stress rupture”) testing for Nicalon/BSAS composite (batches “A” (open triangle) and “B” (closed triangle)) at 1100 ◦C in air. The solid lines represent the predictions made based on Eq. (5) from the results of constant stress-rate testing (Fig. 2). susceptibility to delayed failure, was evident for the composite. The mode of fracture was similar to that in constant stress-rate testing showing some fiber pullout with jagged matrix cracking through the specimen-thickness direction. At lower applied stresses, however, the fracture surfaces of both batches were somewhat flat with decreased fiber pullout, a change in fracture mode more likely to brittle fracture. The lines in Fig. 6 indicate life prediction from the constant stress-rate data. The prediction, primarily applied to brittle monolithic materials, was made using the following relation based on the power-law formulation of Eq. (2). 1 tf = Dn +1 n + 1 σ−n (5) where tf and σ are time to failure and applied constant stress, respectively. Use of Eq. (5) together with n and D determined in constant stress-rate testing allows one to predict life under constant stress (stress rupture) loading. Even though the small number of test specimens was used here due to limited material availability, the prediction was in reasonable agreement with experimental data at least for the Nicalon/BSAS composite. This indicates that the governing failure law of the Nicalon/BSAS composite was very similar either in constant stress-rate or in constant stress loading. Since the prediction (Eq. (5)) was made based on Eq. (2), the governing delayed–failure mechanism of the Nicalon/BSAS would be the one controlled by the power-law type of slow crack growth (Eq. (2)). Other CMCs also showed reasonable agreement in life between constant stress-rate and constantstress loading configurations.1 3.4. Implications As seen in the preceding Sections 3.1–3.3, the strength dependency on test rate, the applicability of the preload technique, and the reasonable life prediction from one loading configuration (constant stress-rate testing) to another (constant stress) for a selected composite all support that delayed failure of the composites was controlled by the power-law type of slow crack growth (or damage evolution/accumulation), confirmed with not only the current CMCs but the previous CMCs.1 It is noted that despite many differences in their processing, architecture, microstructure, and interface, compared with monolithic ceramics, CMCs still exhibit delayed failure, similar in principle to monolithic counterparts. This is simply due to the fact that macroscopically, a CMC is a composite composed of two or more delayed–failure susceptible monolithic materials (constituents of fibers, matrices, interfaces, etc.). The vulnerability to delayed failure would be greater in composite because of its more likelihood of chance to environmental exposure by its inherently more open, porous microstructures, as compared to dense monolithic counterparts. The vulnerability of a composite, of course, will be increased if the constituents added are highly susceptible to delayed failure. A subsequent importance drawn from the results of this work is that constant stress-rate testing, commonly utilized in monolithic ceramics, could be applicable to CMCs to determine their delayed–failure (or life prediction) parameters, at least for a short range of lifetimes, consistent with the previous observation.1 The merits of constant stress-rate testing are enormous in terms of test simplicity and test economy (short test time and less data scatter) over other stress rupture or cyclic fatigue testing. A simplistic, phenomenological law of delayed failure was only explored in this study to develop lifetime prediction testing and methodology for CMCs, without accounting for detailed failure mechanisms associated with matrix/fiber interaction, matrix cracking and its effect on slow crack growth, delayed failure of sustaining fibers, and creep-associated deformation.14–18 A microscopic level of study on this subject is thus needed. Finally, the results of this work suggest that care must be exercised when characterizing elevated-temperature strength of composite materials. This is due to the fact that elevatedtemperature strength has a relative meaning if a material exhibits rate dependency: the strength simply depends on which test rate one chooses (see Fig. 2). Therefore, use of at least two test rates (high and low) is recommended to better characterize high-temperature strength behavior of a composite material. 4. Conclusions Elevated-temperature ultimate tensile strength of five different continuous fiber-reinforced ceramic composites, including Nicalon/BSAS (2D), Hi-Nicalon/BSAS (2D), SiCf/MAS (1D), SiCf/SiC (2D woven: enhanced), and Cf/SiC (2D woven: standard and enhanced), exhibited a strong dependency on test rate, consistent with the behavior observed in other CMCs as well as in many advanced
636 S.R. Choi et al. /Journal of the European Ceramic Society 25(2005)1629-1636 monolithic ceramics at elevated temperatures. The rate de- 6.(a)Choi, S.R. and Gyekenyesi, J. P, Fatigue strength as a function pendency of ultimate tensile strength, the applicability of reloading in dynamic fatigue testing of glass and ceramics. Trans the preload technique, and the predictability of life from one MME.J. Eng. Gas Turbines Power, 1997, 119(3), 493-499 loading configuration(constant stress rate)to another(con- b)Choi, S.R. and Salem, J. A, Effect of preloading on fatigue rength in dynamic fatigue testing of ceramic materials at elevated stant stress)for the Nicalon/BSAS composite suggested that temperatures. Ceram. Eng. Sci. Proc 1995, 16(4),87-94 the overall failure law of the composites would be governed 7. Halbig, M. C, The influence of temperature, stress, and environment by a power-law type of slow crack growth(or damage evo- on the oxidation and life of C/SiC composites. Ceram. Eng. Sci. Proc lution/accumulation ) It was further confirmed that constant stress-rate testing could be utilized as a means of life predic- 8. Calomino, A, verrilli, M.J. and Thomas, D. J, Stress/life behavion of C/SiC composites in a low partial pressure of oxygen environment. tion test methodology for composites when short lifetimes lI-Stress rupture life and residual strength relationship. Ceram. Eng are expected and when ultimate tensile strength is used as a Sc.PPoc.,2002,23,443-451. ilure criterion 9. Evans, A.G., Slow crack growth in brittle materials under dynamic 10. ASTM C 1368. Standard test method fo wth parameters of advanced ceramics by constant stress-rate flex- Acknowledge ng at ambient atureIn Anmual Book of AsTM Standards This work was supported in part by Higher Operating Tem- 1. ASTM C 1465. Standard test method for determination of slow crack perature Propulsion Components(HOTPC) Program and the owth parameters of advanced ceramics by constant stress-rate fl Ultra-Efficient Engine Technology (UEET) Program, NASA ural testing at elevated temperatures. In Annual Book of ASTM Stan- dards(Vol 15.01). ASTM, West Conshohocken, PA, 2001 lenn Research Center, Cleveland, OH, USA. The authors 12. Choi, S. R, Gyekenyesi, J. P. et al., Ultra-fast fracture strength of are grateful to R. Pawlik for experimental work during the advanced structural ceramics at elevated temperatures: an approach to course of this research high-temperature"inert "strength. In Fracture Mechanics of Ceramics 13, ed. R. C. Bradt. Kluwer Academic/Plenum Publishers, New 13. Halbig, M. C, Modeling the oxidation of carbon fibres in a C/SiC References omposite under stressed oxidation. Ceram. Eng. Sci. Proc., 2002 23(3),427-434 1. Choi, S.R and Gyekenyesi, J. P, Effect of load rate on tensile 14. Sorenson, B F and Holmes, J w, Effect of loading rate on the mono- ngth of various CFCCs at elevated temperatures: an a tonic tensile behavior of a continuous-fiber-reinforced glass-ceramic prediction testing. Ceram. Eng. Sci. Proc., 2001, 22(3), 597-606 matrix composite. J. Am. Ceram Soc, 1996, 79(2), 313-320 2. Bansal, N P and Setlock, J. A, Fabrication of fiber-reinforced celsian 15. Curtin, w.A. and Halverson, H. G, High Temperature Deforma- matrix composites Composites Part 4, 2001, 32, 1021-1029 tion and Failure in Oxide/oxide Composites, HITEMP Review 1999 3. Worthen. D. Thermomechanical Fatigue Behavior of Three re Engine Materials Technology Project CFCC's. NAS CR-195441. NASA Glenn Research Center, Cleve. (NASA/CP-1999-208915/ol 2, Paper 48). NASA Glenn Research enter. Cleveland. OH. 1999 4. Verrilli, M. J, Calomino, A and Thomas, D. J. Stress/life behavior of 16. Lewinsohn, C. A, Henager Jr, C. H. and Jones, R. H, Environmen a C/SiC composite in a low partial pressure of oxygen environment: tally induced time-dependent failure mechanisms in CFCCs at elevated I-Static strength and stress rupture database. Ceram. Eng. Sci. Proc. temperatures. Ceram. Eng. Sci. Proc., 1998, 19(4), 11-18 2002,23(3),435-442. 17. Henager, C. H. and Jones, R. H, Subcritical crack growth in CVI 5. ASTM C 1359, Standard test method for monotonic tensile streng silicon-carbide reinforced with Nicalon fibers-experiment and model esting of continuous fiber-reinforced ceramics with solid JAm. Ceram.Soc,1994,77(9),2381-2394 mperatures. In An- 18. Spearing, S. M, Zok, F. W. and Evans, A. G, Stress-corrosion crack mal Book of ASTM Standards (ol 15.01). ASTM, West Con- ng in a unidirectional ceramic-matrix composite. J. Am. Ceram. Soc. shohocken. PA. 2001 1994,77(2),562-570
1636 S.R. Choi et al. / Journal of the European Ceramic Society 25 (2005) 1629–1636 monolithic ceramics at elevated temperatures. The rate dependency of ultimate tensile strength, the applicability of the preload technique, and the predictability of life from one loading configuration (constant stress rate) to another (constant stress) for the Nicalon/BSAS composite suggested that the overall failure law of the composites would be governed by a power-law type of slow crack growth (or damage evolution/accumulation). It was further confirmed that constant stress-rate testing could be utilized as a means of life prediction test methodology for composites when short lifetimes are expected and when ultimate tensile strength is used as a failure criterion. Acknowledgments This work was supported in part by Higher Operating Temperature Propulsion Components (HOTPC) Program and the Ultra-Efficient Engine Technology (UEET) Program, NASA Glenn Research Center, Cleveland, OH, USA. The authors are grateful to R. Pawlik for experimental work during the course of this research. References 1. Choi, S. R. and Gyekenyesi, J. P., Effect of load rate on tensile strength of various CFCCs at elevated temperatures: an approach to life-prediction testing. Ceram. Eng. Sci. Proc., 2001, 22(3), 597–606. 2. Bansal, N. P. and Setlock, J. A., Fabrication of fiber-reinforced celsian matrix composites. Composites Part A, 2001, 32, 1021–1029. 3. Worthem, D. W., Thermomechanical Fatigue Behavior of Three CFCC’s, NASA CR-195441. NASA Glenn Research Center, Cleveland, OH, 1995. 4. Verrilli, M. J., Calomino, A. and Thomas, D. J., Stress/life behavior of a C/SiC composite in a low partial pressure of oxygen environment: I—Static strength and stress rupture database. Ceram. Eng. Sci. Proc., 2002, 23(3), 435–442. 5. ASTM C 1359, Standard test method for monotonic tensile strength testing of continuous fiber-reinforced advanced ceramics with solid rectangular cross-section specimens at elevated temperatures. In Annual Book of ASTM Standards (Vol 15.01). ASTM, West Conshohocken, PA, 2001. 6. (a) Choi, S. R. and Gyekenyesi, J. P., Fatigue strength as a function of preloading in dynamic fatigue testing of glass and ceramics. Trans. ASME. J. Eng. Gas Turbines Power, 1997, 119(3), 493–499; (b) Choi, S. R. and Salem, J. A., Effect of preloading on fatigue strength in dynamic fatigue testing of ceramic materials at elevated temperatures. Ceram. Eng. Sci. Proc., 1995, 16(4), 87–94. 7. Halbig, M. C., The influence of temperature, stress, and environment on the oxidation and life of C/SiC composites. Ceram. Eng. Sci. Proc., 2002, 23, 419–426. 8. Calomino, A., Verrilli, M. J. and Thomas, D. J., Stress/life behavior of C/SiC composites in a low partial pressure of oxygen environment. II—Stress rupture life and residual strength relationship. Ceram. Eng. Sci. Proc., 2002, 23, 443–451. 9. Evans, A. G., Slow crack growth in brittle materials under dynamic loading conditions. Int. J. Fracture, 1974, 10(2), 251–259. 10. ASTM C 1368, Standard test method for determination of slow crack growth parameters of advanced ceramics by constant stress-rate flexural testing at ambient temperature. In Annual Book of ASTM Standards (Vol 15.01). ASTM, West Conshohocken, PA, 2001. 11. ASTM C 1465, Standard test method for determination of slow crack growth parameters of advanced ceramics by constant stress-rate flexural testing at elevated temperatures. In Annual Book of ASTM Standards (Vol 15.01). ASTM, West Conshohocken, PA, 2001. 12. Choi, S. R., Gyekenyesi, J. P. et al., Ultra-fast fracture strength of advanced structural ceramics at elevated temperatures: an approach to high-temperature “inert” strength. In Fracture Mechanics of Ceramics, Vol 13, ed. R. C. Bradt. Kluwer Academic/Plenum Publishers, New York, 2002, pp. 27–46. 13. Halbig, M. C., Modeling the oxidation of carbon fibres in a C/SiC composite under stressed oxidation. Ceram. Eng. Sci. Proc., 2002, 23(3), 427–434. 14. Sorenson, B. F. and Holmes, J. W., Effect of loading rate on the monotonic tensile behavior of a continuous-fiber-reinforced glass–ceramic matrix composite. J. Am. Ceram. Soc., 1996, 79(2), 313–320. 15. Curtin, W. A. and Halverson, H. G., High Temperature Deformation and Failure in Oxide/oxide Composites, HITEMP Review 1999: Advanced High Temperature Engine Materials Technology Project (NASA/CP—1999-208915/Vol 2, Paper 48). NASA Glenn Research Center, Cleveland, OH, 1999. 16. Lewinsohn, C. A., Henager Jr., C. H. and Jones, R. H., Environmentally induced time-dependent failure mechanisms in CFCCs at elevated temperatures. Ceram. Eng. Sci. Proc., 1998, 19(4), 11–18. 17. Henager, C. H. and Jones, R. H., Subcritical crack growth in CVI silicon-carbide reinforced with Nicalon fibers—experiment and model. J. Am. Ceram. Soc., 1994, 77(9), 2381–2394. 18. Spearing, S. M., Zok, F. W. and Evans, A. G., Stress–corrosion cracking in a unidirectional ceramic-matrix composite. J. Am. Ceram. Soc., 1994, 77(2), 562–570