《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Melt-Infiltrated SiC Matrix Composites

J.Am. Cera.Soe.,8951652-165802006 DOI:10.111.1551-2916.200600939x C 2006 The American Ceramic Society urna Creep and Stress-Strain Behavior after Creep for SiC Fiber Reinforced, Melt-Infiltrated SiC matrix Composites Gregory N Morschen Ohio Aerospace Institute, NASA Glenn Research Center, MS 106-5, Cleveland, Ohio 44135 Vijay v. Pujar Materials and Simulation Center, Goodrich Corporation, Brecksville, Ohio 44141 Silicon carbide fiber(Hi-Nicalon Type S, Nippon Carbon) re- combining CVI-SiC with Sic particulate slurry and molten sil- inforced silicon carbide matrix composites containing melt-in- icon(Si) infiltration(referred to hereafter as the MI matrix)has filtrated silicon were subjected to creep at 1315C at three resulted in the best combination of thermo-mechanical proper different stress conditions. for the lies for many targeted applications. In particular, the improved ture after 100 h of tensile creep, fast-fracture experiments were densification in these MI composites give improved matrix performed immediately following the creep test at the creep cracking properties relative to composites consisting of the con- temperature(1315 C)or after cooling to room temperature. All entional all-cvI matrix' There are two reasons for the im- specimens demonstrated excellent creep resistance and com- proved matrix cracking properties. First, the Si-infiltrant fills in pared well to the creep behavior published in the literature on most of the macro-porosity of the CVi-SiC matrix, thereby re- similar composite systems. Tensile results on the after-creep noving the stress concentrators in the cvi-sic matrix at cross- specimens showed that the matrix cracking stress actually in- over points in the woven structure, which are observed to be the creased. which is attributed to stress redistribution between site of matrix crack initiation at lower stresses 7.8 Second. the mposite constituents during tensile creep. volume expansion of Si upon cooling below its melting point and the lower thermal expansion coefficient of Si compared with SiC places the matrix under residual compression in the ma- L. Introduction trix relative to the fiber. However the mi matrix itself is ex pected to be less creepresistant than an all-CVI-SiC matrix ILICON carbide fiber reinforced silicon carbide matrix ceramic because of its reduced CVI-SiC content and the low creep re- matrix composites( CMCs), commonly referred to as SiC/ sistance of the Si infiltrant at temperatures near SiC CMCs, have been studied extensively over the last few dec-(1410.C) des for their potential as structural materials in different fields This study was aimed at understanding the tensile be. including advanced turbine engines for aerospace and pow ha avior of the Mi composite system at 1315.C in air, including generation applications. A primary reason for use of SiC fiber the effect of creep on the resulting stress-strain response and reinforced SiC CMCs is their ability to retain a high degree of ultimate properties of the system. There have been only limited structural performance at high temperatures because of their studies reported on inherent creep-resistance. Much work on SiC/SiC composite no studies on after eep behavior in SiC-SiC MI CMCS, and reep behavior of these materials. There creep has been performed for systems with lower-use tempera- have been. however, some studies on creep and after-creep be- ture fibers, e.g. NicalonM and Hi-NicalonM (Nippon Carbon havior on other ceramic composite, namely glass ceramic and Co., Tokyo, Japan)and with chemical vapor infiltrated(CvI) Si3N systems. Specifically, composite systems SiC matrices. In these studies, woven 2D 0/90 lay-up com which the fiber is more creep resistant than the matrix, it has were crept at temperatures ranging from 1000 to 1300C been shown that at creep stresses lower than the matrix cracking and argon,and creep stresses ranging from 30-180 stresses, the matrix relaxes during creep and transfers the load to the primary fiber direction. At the higher temperatures, the fiber It has been shown in these materials that the stress to 300C, and at applied tensile creep stresses above 100 cause matrix cracking in the after-creep specimens is actually MPa, rupture usually occurred after a few hours in these sam- higher, relative to that in the as-produced material. As de- ples with rupture strains upwards of 1%, thereby limiting their scribed in this paper, a similar phenomenon appears to be op use to lower temperatures. erative in the MI system studied he With the availability of more creep-resistant polycrystalline SiC fibers, e.g., Hi-Nicalon Type s(made by Nippon Carbon referred to here as HNS). .Tyranno SA(Ube Industries,To- kyo, Japan), Sylramic(formerly Dow Corning, and now made Il. Experimental Procedure by Col Ceramics, San Diego. CA), greater creep resistance Details about composite processing, constituent content, room can be achieved. This enables higher use-temperatures for Sic and some elevated temperature properties can be found in Mo- Sic composites based on these fibers. At the same time, signif- scher and Pujar for the three different panels that tensile spec icant progress has been made in matrix development to achieve imens for this study came from. The melt-infiltrated composites nearly full densification. The matrix consolidation approach all consisted of eight plies of stacked five harness satin (7.1 tow ends/cm)2D woven 0/90 HNS fabric, a CVI BN interphase, a editor VI-SiC matrix layer, followed by Sic slurry infiltration and molten Si infiltration. The three panels varied in their relative constituent contents, as shown in Table I. Note that panels labeled Al and a2 consisted of similar fractions of cvi SiC and luscript No. 21068 Received October 11, 2005: approved December 19, 2005. Si whereas panel B had a significantly lower fraction of CVI SiC thor to whom correspondence should be addressed. e-mail: vijay pujar(a goodrich. content and consequently a significantly greater fraction of si isenior Research Scientist residing at NASA Glenn Research Center. content

Creep and Stress–Strain Behavior after Creep for SiC Fiber Reinforced, Melt-Infiltrated SiC Matrix Composites Gregory N. Morscherz Ohio Aerospace Institute, NASA Glenn Research Center, MS 106-5, Cleveland, Ohio 44135 Vijay V. Pujarw Materials and Simulation Center, Goodrich Corporation, Brecksville, Ohio 44141 Silicon carbide fiber (Hi-Nicalon Type S, Nippon Carbon) re￾inforced silicon carbide matrix composites containing melt-in- filtrated silicon were subjected to creep at 13151C at three different stress conditions. For the specimens that did not rup￾ture after 100 h of tensile creep, fast-fracture experiments were performed immediately following the creep test at the creep temperature (13151C) or after cooling to room temperature. All specimens demonstrated excellent creep resistance and com￾pared well to the creep behavior published in the literature on similar composite systems. Tensile results on the after-creep specimens showed that the matrix cracking stress actually in￾creased, which is attributed to stress redistribution between composite constituents during tensile creep. I. Introduction SILICON carbide fiber reinforced silicon carbide matrix ceramic matrix composites (CMCs), commonly referred to as SiC/ SiC CMCs, have been studied extensively over the last few dec￾ades for their potential as structural materials in different fields including advanced turbine engines for aerospace and power generation applications. A primary reason for use of SiC fiber reinforced SiC CMCs is their ability to retain a high degree of structural performance at high temperatures because of their inherent creep-resistance. Much work on SiC/SiC composite creep has been performed for systems with lower-use tempera￾ture fibers, e.g., NicalonTM and Hi-NicalonTM (Nippon Carbon Co., Tokyo, Japan) and with chemical vapor infiltrated (CVI) SiC matrices.1–3 In these studies, woven 2D 0/90 lay-up com￾posites were crept at temperatures ranging from 10001 to 13001C in air2,3 and argon,1–3 and creep stresses ranging from 30–180 MPa in the primary fiber direction. At the higher temperatures, 12001–13001C, and at applied tensile creep stresses above 100 MPa, rupture usually occurred after a few hours in these sam￾ples with rupture strains upwards of 1%, thereby limiting their use to lower temperatures. With the availability of more creep-resistant polycrystalline SiC fibers, e.g., Hi-Nicalon Type S (made by Nippon Carbon, referred to here as HNS),4,5 Tyranno SA (Ube Industries, To￾kyo, Japan),5 Sylramic (formerly Dow Corning, and now made by COI Ceramics, San Diego, CA),5,6 greater creep resistance can be achieved. This enables higher use-temperatures for SiC/ SiC composites based on these fibers. At the same time, signif￾icant progress has been made in matrix development to achieve nearly full densification. The matrix consolidation approach combining CVI–SiC with SiC particulate slurry and molten sil￾icon (Si) infiltration (referred to hereafter as the MI matrix) has resulted in the best combination of thermo-mechanical proper￾ties for many targeted applications. In particular, the improved densification in these MI composites give improved matrix cracking properties relative to composites consisting of the con￾ventional all-CVI matrix.7 There are two reasons for the im￾proved matrix cracking properties. First, the Si-infiltrant fills in most of the macro-porosity of the CVI–SiC matrix, thereby re￾moving the stress concentrators in the CVI–SiC matrix at cross￾over points in the woven structure, which are observed to be the site of matrix crack initiation at lower stresses.7,8 Second, the volume expansion of Si upon cooling below its melting point and the lower thermal expansion coefficient of Si compared with SiC places the matrix under residual compression in the ma￾trix5,9 relative to the fiber. However, the MI matrix itself is ex￾pected to be less creepresistant than an all-CVI–SiC matrix because of its reduced CVI–SiC content and the low creep re￾sistance of the Si infiltrant at temperatures near its melting point (14101C). This study was aimed at understanding the tensile creep be￾havior of the MI composite system at 13151C in air, including the effect of creep on the resulting stress–strain response and ultimate properties of the system. There have been only limited studies reported on creep behavior in SiC–SiC MI CMCs, and no studies on after-creep behavior of these materials.10 There have been, however, some studies on creep and after-creep be￾havior on other ceramic composite, namely glass ceramic and Si3N4 matrix systems.11–13 Specifically, composite systems in which the fiber is more creep resistant than the matrix, it has been shown that at creep stresses lower than the matrix cracking stresses, the matrix relaxes during creep and transfers the load to the fiber.11 It has been shown in these materials that the stress to cause matrix cracking in the after-creep specimens is actually higher11,13 relative to that in the as-produced material. As de￾scribed in this paper, a similar phenomenon appears to be op￾erative in the MI system studied here. II. Experimental Procedure Details about composite processing, constituent content, room and some elevated temperature properties can be found in Mo￾rscher and Pujar5 for the three different panels that tensile spec￾imens for this study came from. The melt-infiltrated composites all consisted of eight plies of stacked five harness satin (7.1 tow ends/cm) 2D woven 0/90 HNS fabric, a CVI BN interphase, a CVI–SiC matrix layer, followed by SiC slurry infiltration and molten Si infiltration. The three panels varied in their relative constituent contents, as shown in Table I.5 Note that panels labeled A1 and A2 consisted of similar fractions of CVI SiC and Si whereas panel B had a significantly lower fraction of CVI SiC content and consequently a significantly greater fraction of Si content. Journal J. Am. Ceram. Soc., 89 [5] 1652–1658 (2006) DOI: 10.1111/j.1551-2916.2006.00939.x r 2006 The American Ceramic Society 1652 E. Lara-Curzio—contributing editor w Author to whom correspondence should be addressed. e-mail: vijay.pujar@goodrich. comz Senior Research Scientist residing at NASA Glenn Research Center. Manuscript No. 21068. Received October 11, 2005; approved December 19, 2005

May 2006 Creep and After-Creep Stress-Strain Behavior 1653 Table l. Physical Properties of the Composite Panels Tow ends per cm action CVI SiC fraction al-SiC(slurry) fraction Si fraction Panel density(g/cm) 0.2 0.14 2.83 7.1 0.21 0.21 0.13 2.74 B 0.33 0.14 2.79 sCatter in the thickness measurement was always less than 0.1 mm. CVI. chemical vapor infiltrated: SiC, silicon carbide: Si, silicon. Tensile test specimens were machined from the panels in the fore, the 0.002% offset strain stress was more useful for com- shape of a dog-bone (length= 152 mm; grip width= 12.6 mm; parison of all the loading curves. Selected specimens were cut gage width= 10 mm)with the 0 fibers aligned along the spec- and polished longitudinally and examined on an optical micro- imen length. Room temperature tensile tests were performe scope to determine the extent of matrix cracking Model 8562, Instron Inc, Canton, MA). Elevated temperature ensile tests were performed ing a screw-driven testing II. Results wrapped with wire( steel) screens in the grip region of the dog (1) Tensile Creep behavior bone test specimen (152 mm long tapered specimens: grip The tensile creep curves are plotted in Fig. I as total strain(creep width= 12.6 mm; gage width= 10 mm) in order to distribute strain plus loading strain) versus time for specimens from the grip pressure and alleviate stress concentrations from the hy three panels tested. In some cases, two specimens from the same draulic wedge grips(MTS 647). Room temperature strain meas- panel were tested per condition and some variation was urement was performed using a 25 mm gage length clip-on served for the two specimens. Table II lists the elastic moduli performed using a 25 mm gage length low-contact capacitance creep behavior of these composites for the most part exhibited a extensometer. For the elevated temperature tests, an MoSiz el ement. resistance-heated furnace was inserted in between the ed, 172 MPa, specimen ruptured and exhibited a rate of defor- two water-cooled grips. The furnace hot zone length was ap- mation increasing with time before failure. It is debatable proximately 15 mm. Room temperature tests were performed in whether"steady-state"creep is achieved. Figure 2 shows the load control, 4 kN/min. Elevated temperature tests(fast fracture strain rate versus strain for the creep experiments as determined nd creep) were loaded with a crosshead-displacement rate of by taking the instantaneous slope over 200 data points( 3.3 h) of the creep data. The composites do not appear to reach a 3.23 mm/min. The time to failure for the elevated temperature steady-state rate for 103 MPa stresses; however, for the 138 MPa tests was usually a few minutes under fast-fracture conditions similar to tests performed at room temperature. applied stresses a steady-state rate may have been achieved or at least nearly approached The 172 MPa specimen achieved a sec- ag, and reloading at predetermined loads to obtain hysteresis ondary-state strain rate; which perhaps was an inevitable arti- ops and determine residual stress in the specimen. Crack in- fact of a specimen experiencing an increasing creep rate before itation and propagation in the tensile specimens tested at room failure. The minimum strain rates and creep strains are listed in temperature was also monitored with wide band acoustic sen Table Il for each experiment. sors(B1025, Digital Wave Corp, Englewood, CO)attached di The stress-strain curves for several creep experiments as well rectly to the specimens. Two sensors were placed outside of the as the stress-strain curves for room temperature and 1315 fast fracture are plotted in Fig 3 for the composite panel from which gage section, and a third sensor was attached within the gage the most tests were performed. Several interesting observations section on the opposite side of the knife edges of the extenso- meter. This enabled accurate location of acoustic events. Only arise from the data in Fig. 3. First there is considerable non- the events that occurred in the gage section were used in the linearity in the 1315C fast-fracture curve at relatively low stress analysis (<100 MPa). For tests at room temperature, such non-linearity For elevated temperature tensile the furnace was first is almost always because of matrix cracking. At 1315C, this aised to 1315C and held at temperature for 10 min to insur could be because of either matrix cracking or composite creep temperature stability. The specimens were then loaded mon during loading or both. Note that the loading rate was relatively tonically to failure for fast-fracture tests or to the predetermined slow and as shown below much of this non-linear behavior is creep stress for the creep experiments. If specimens in the creep deed attributed to matrix creep although some minor micro- ived 100 h, they were rapidly unloaded to zero ress and immediately reloaded to failure at the cree ature under fast-fracture conditions. In some cases the mens were rapidly unloaded and immediately cooled to room temperature where a room temperature hysteresis to failure test A2-172 MPa was performed. In all cases, elastic moduli at room or elevat tures were determined using linear ssion of the 5- 50 MPa part of the loading curve. To determine the onset for B-138 MPa non-linear stress-strain (proportional limit), the 0.002% and 0.4 0.005% offset strain stresses were measured (ASTM C1275)and A2-138 MPa ompared for most of the fast-fracture tests on as-produced specimens, the initial loading curves of the creep tests, and the A1-103 MPa fast-fracture tests for the after-creep specimens. The 0.005% A1-103 MPa offset strain approach has been used at NASA Glenn Research Center for the Sylramic-iBN fiber reinforced MI composite sys- 0.1 B-103 MPa A2-103 MPa tem. 1 However as evident below the 0.005% offset strain stress was in some cases greater than the applied creep stress; there- 100 Time. hr Sylramic fibers were produced by Dow Coming(Midland, MD) foll a propn- etary in situ BN NASA-developed heat treatment. Fig 1. Tensile creep curves

Tensile test specimens were machined from the panels in the shape of a dog-bone (length 5 152 mm; grip width 5 12.6 mm; gage width 5 10 mm) with the 01 fibers aligned along the spec￾imen length. Room temperature tensile tests were performed using an electromechanical actuator universal testing machine (Model 8562, Instron Inc., Canton, MA). Elevated temperature tensile tests were performed using a screw-driven testing ma￾chine (Model 5569, Instron Inc.). The specimens ends were wrapped with wire (steel) screens in the grip region of the dog bone test specimen (152 mm long tapered specimens; grip width 5 12.6 mm; gage width 5 10 mm) in order to distribute grip pressure and alleviate stress concentrations from the hy￾draulic wedge grips (MTS 647). Room temperature strain meas￾urement was performed using a 25 mm gage length clip-on extensometer. Elevated temperature strain measurement was performed using a 25 mm gage length low-contact capacitance extensometer. For the elevated temperature tests, an MoSi2 el￾ement, resistance-heated furnace was inserted in between the two water-cooled grips. The furnace hot zone length was ap￾proximately 15 mm. Room temperature tests were performed in load control, 4 kN/min. Elevated temperature tests (fast fracture and creep) were loaded with a crosshead-displacement rate of 3.23 mm/min. The time to failure for the elevated temperature tests was usually a few minutes under fast-fracture conditions, similar to tests performed at room temperature. Room temperature tensile tests consisted of loading, unload￾ing, and reloading at predetermined loads to obtain hysteresis loops and determine residual stress in the specimen.14 Crack in￾itiation and propagation in the tensile specimens tested at room temperature was also monitored with wide band acoustic sen￾sors (B1025, Digital Wave Corp., Englewood, CO) attached di￾rectly to the specimens. Two sensors were placed outside of the gage section, and a third sensor was attached within the gage section on the opposite side of the knife edges of the extenso￾meter. This enabled accurate location of acoustic events. Only the events that occurred in the gage section were used in the analysis. For elevated temperature tensile tests, the furnace was first raised to 13151C and held at temperature for 10 min to insure temperature stability. The specimens were then loaded mono￾tonically to failure for fast-fracture tests or to the predetermined creep stress for the creep experiments. If specimens in the creep experiments survived 100 h, they were rapidly unloaded to zero stress and immediately reloaded to failure at the creep temper￾ature under fast-fracture conditions. In some cases the speci￾mens were rapidly unloaded and immediately cooled to room temperature where a room temperature hysteresis to failure test was performed. In all cases, elastic moduli at room or elevated temperatures were determined using linear regression of the 5– 50 MPa part of the loading curve. To determine the onset for non-linear stress–strain (proportional limit), the 0.002% and 0.005% offset strain stresses were measured (ASTM C1275) and compared for most of the fast-fracture tests on as-produced specimens, the initial loading curves of the creep tests, and the fast-fracture tests for the after-creep specimens. The 0.005% offset strain approach has been used at NASA Glenn Research Center for the Sylramic-iBNy fiber reinforced MI composite sys￾tem.15 However, as evident below, the 0.005% offset strain stress was in some cases greater than the applied creep stress; there￾fore, the 0.002% offset strain stress was more useful for com￾parison of all the loading curves. Selected specimens were cut and polished longitudinally and examined on an optical micro￾scope to determine the extent of matrix cracking. III. Results (1) Tensile Creep Behavior The tensile creep curves are plotted in Fig. 1 as total strain (creep strain plus loading strain) versus time for specimens from the three panels tested. In some cases, two specimens from the same panel were tested per condition and some variation was ob￾served for the two specimens. Table II lists the elastic moduli, creep, and retained strength properties for each experiment. The creep behavior of these composites for the most part exhibited a decreasing rate of deformation with time. Only the highly load￾ed, 172 MPa, specimen ruptured and exhibited a rate of defor￾mation increasing with time before failure. It is debatable whether ‘‘steady-state’’ creep is achieved. Figure 2 shows the strain rate versus strain for the creep experiments as determined by taking the instantaneous slope over 200 data points (B3.3 h) of the creep data. The composites do not appear to reach a steady-state rate for 103 MPa stresses; however, for the 138 MPa applied stresses a steady-state rate may have been achieved or at least nearly approached. The 172 MPa specimen achieved a sec￾ondary-state strain rate; which perhaps was an inevitable arti￾fact of a specimen experiencing an increasing creep rate before failure. The minimum strain rates and creep strains are listed in Table II for each experiment. The stress–strain curves for several creep experiments as well as the stress–strain curves for room temperature and 13151C fast fracture are plotted in Fig. 3 for the composite panel from which the most tests were performed. Several interesting observations arise from the data in Fig. 3. First, there is considerable non￾linearity in the 13151C fast-fracture curve at relatively low stress (o100 MPa). For tests at room temperature, such non-linearity is almost always because of matrix cracking. At 13151C, this could be because of either matrix cracking or composite creep during loading or both. Note that the loading rate was relatively slow and as shown below, much of this non-linear behavior is indeed attributed to matrix creep, although some minor micro- 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0 20 40 60 80 100 120 Time, hr Total Strain, % A1-103 MPa A2-138 MPa A2-138 MPa B-138 MPa A2-172 MPa A2-103 MPa A1-103 MPa B-103 MPa Fig. 1. Tensile creep curves. Table I. Physical Properties of the Composite Panels Panel Tow ends per cm Average f BN fraction CVI SiC fraction a-SiC (slurry) fraction Si fraction Panel density (g/cm3 ) A1 7.1 0.30 0.04 0.25 0.23 0.14 2.83 A2 7.1 0.35 0.04 0.21 0.21 0.13 2.74 B 7.1 0.33 0.04 0.14 0.25 0.20 2.79 wScatter in the thickness measurement was always less than 0.1 mm. CVI, chemical vapor infiltrated; SiC, silicon carbide; Si, silicon. y Sylramic fibers were produced by Dow Corning (Midland, MI) followed by a propri￾etary in situ BN NASA-developed heat treatment. May 2006 Creep and After-Creep Stress–Strain Behavior 1653

1654 Journal of the American Ceramic Society-Morscher and pujar Vol. 89. No. 5 Table Il. Tensile Creep Results Creep stress E(RT) E(1315)initial loading E(1315)after creep E(RT) after Minimum strain rate Creep strain Ultimate stress (GPa) (GPa) (GPa) I RT 1315FF 253+ 1315°C 2.80E-09 0.18 103 2.50E-09 267 A2 RT 二 232 41 1315FF 271t 1315°C 2.80E-10 1315°C 3.90E-090.3 1315°C 233 3.10E-09 0.23 1315°C 2.60E-080.44 270 1315°C 5.60E-10 1315°C 138 263 213 214 5.20E-090.36 247 Tested at 1315 C ' Did not fail in hot zone. Tested at room tempe Prior to the onset of tertiary creep. Each row corresponds to an individual tensile specimen. cracking is also evident. Second, the initial loading curves for the both the 103 and 138 MPa creep creep experiments exhibited a wide range of elastic moduli (Ta- fraction composites from panel a offset stress- ble ID). This could be because of variability in constituent con- es. Panel B with the higher Si cor higher offset tent from specimen to specimen or experimental error. Third, stress after creep; however, the increase was not as dramatic the elastic moduli measured from the fast-fracture curves after (only -20% higher) creep were slightly lower relative to moduli measured in the A specimen from panel A2 and a specimen from panel B that same specimens from the initial loading curves. This is discussed had survived 100 h, 138 MPa tensile creep were tested(unload- in more detail in the next section. Fourth, in all the specimens eload) to failure at room temperature. Figure 4(a) shows the that survived 100 h creep, no real loss in fast-fracture strength room temperature tensile stress-strain curves for the as-pro- was observed when compared with fast fracture of as-produced duced and after-creep A2 specimen, and Table IV compares the specimens at 1315.C. Fifth, the creep strain corresponded al- before-creep and after- creep properties. It is evident that the most exactly with the permanent displacement measured after "knee"in the after-creep curve is higher than the room temper- rapid u ture curve. Similarly, significant AE activity occurred at 140 MPa for the after-creep specimen, higher than the 115 MPa ob- served for the as-produced specimen (see dashed arrows in (2) Non-Linearity at Elevated and Room Temperature Fig 4(b). The stresses at which the 0.002% and 0.005% offset Table Ill lists the 0.002% and 0.005% offset-strain stresses de- strain curves intersect the stress-strain curve were 142 and 177 termined for tensile specimens tested at 1315C. Although there MPa, respectively, for the after-creep curve compared with 12 is some specimen-to-specimen variation in the measured offset and 147 MPa, respectively. for the as-produced curve(Fig 4(c) stresses, it is apparent that for panels Al and A2, which had Finally taking the average slope from the top portion of the relatively higher CVI SiC and lower Si content, the after-creep hysteresis loops(after Steen and Valles) one can determine a compressive stress for the crept specimen of approxi- in the intital loading of the creep experiment. This is true for 450 RT Fast Fracture 3.5E08 (hys loops remove 350 1315C Fast Fracture A1-103 MPa s 300(did not fail in hot A2-172 MPa E20E08 15E-08H B-138 MPa 00 138 Mpa; 100 hr; /1315C Fast Fracture 1315° C Creep 150 50E-09 方00E+00da 103Mpa;100hr; 040.50.60.7 1315° C Cree A2-138 MPa 10E-08 0.20.30.4 0.60.7 Total Strain. Fig. 2. Strain rate versus strain Fig 3. Stress-strain history for panel A2 specimens

cracking is also evident. Second, the initial loading curves for the creep experiments exhibited a wide range of elastic moduli (Ta￾ble II). This could be because of variability in constituent con￾tent from specimen to specimen or experimental error. Third, the elastic moduli measured from the fast-fracture curves after creep were slightly lower relative to moduli measured in the same specimens from the initial loading curves. This is discussed in more detail in the next section. Fourth, in all the specimens that survived 100 h creep, no real loss in fast-fracture strength was observed when compared with fast fracture of as-produced specimens at 13151C. Fifth, the creep strain corresponded al￾most exactly with the permanent displacement measured after rapid unloading. (2) Non-Linearity at Elevated and Room Temperature Table III lists the 0.002% and 0.005% offset-strain stresses de￾termined for tensile specimens tested at 13151C. Although there is some specimen-to-specimen variation in the measured offset stresses, it is apparent that for panels A1 and A2, which had relatively higher CVI SiC and lower Si content, the after-creep offset stresses are nearly doubled compared with those observed in the intital loading of the creep experiment. This is true for both the 103 and 138 MPa creep tests. The higher fiber volume fraction composites from panel A2 had the highest offset stress￾es. Panel B with the higher Si content also had a higher offset stress after creep; however, the increase was not as dramatic (only B20% higher). A specimen from panel A2 and a specimen from panel B that had survived 100 h, 138 MPa tensile creep were tested (unload– reload) to failure at room temperature. Figure 4(a) shows the room temperature tensile stress–strain curves for the as-pro￾duced and after-creep A2 specimen, and Table IV compares the before-creep and after-creep properties. It is evident that the ‘‘knee’’ in the after-creep curve is higher than the room temper￾ature curve. Similarly, significant AE activity occurred at 140 MPa for the after-creep specimen, higher than the 115 MPa ob￾served for the as-produced specimen (see dashed arrows in Fig. 4(b)). The stresses at which the 0.002% and 0.005% offset strain curves intersect the stress–strain curve were 142 and 177 MPa, respectively, for the after-creep curve compared with 125 and 147 MPa, respectively, for the as-produced curve (Fig. 4(c)). Finally, taking the average slope from the top portion of the hysteresis loops (after Steen and Valles14) one can determine a matrix compressive stress for the crept specimen of approxi- −1.0E−08 −5.0E−09 0.0E+00 5.0E−09 1.0E−08 1.5E−08 2.0E−08 2.5E−08 3.0E−08 3.5E−08 4.0E−08 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Strain, % Strain Rate, mm/mm/sec A1-103 MPa A2-138 MPa B-138 MPa A2-172 MPa MPa B-103 MPa rupture Fig. 2. Strain rate versus strain. Table II. Tensile Creep Results Panel Experiment Creep stress (MPa) E (RT) (GPa) E (1315) initial loading (GPa) E (1315) after creep (GPa) E (RT) after creep Minimum strain rate (s
1 ) Creep strain (%) Ultimate stress (MPa) A1 RT — 262 — — — — — 349 1315FF — — 189 — — — — 253w,z 13151C creep 103 — NA 209 — 2.80E–09 0.18 256w 13151C creep 103 — 223 209 — 2.50E–09 0.09 267w A2 RT — 232 — — — — 412 1315FF — — 182 — — — — 271w,z 13151C creep 103 — 225 198 — 2.80E
10 0.05 295w 13151C creep 138 — 184 157 — 3.90E
09 0.3 291w 13151C creep 138 — 203 — 233 3.10E
09 0.23 321y 13151C creep 172 — 177 — — 2.60E
08 0.44z — B RT — 270 — — — — — 362 13151C creep 103 — 217 209 — 5.60E
10 0.07 255w 13151C creep 138 263 213 — 214 5.20E
09 0.36 247y w Tested at 13151C. z Did not fail in hot zone. y Tested at room temperature. z Prior to the onset of tertiary creep. Each row corresponds to an individual tensile specimen. 0 50 100 150 200 250 300 350 400 450 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Total Strain, % Stress, MPa RT Fast Fracture (hys loops removed) 1315°C Fast Fracture (did not fail in hot zone) 138 Mpa;100 hr; 1315°C Creep 1315°C Fast Fracture after Creep Creep strain 103 Mpa;100 hr; 1315°C Creep Fig. 3. Stress–strain history for panel A2 specimens. 1654 Journal of the American Ceramic Society—Morscher and Pujar Vol. 89, No. 5

May 2006 Creep and After-Creep Stress-Strain Behavior 1655 Table lll. offset stress at 1315C and matrix Cracking observations Creep stress 1315C initial loading 1315C after creep Creep strain Panel Experim (MPa) (MPa) Observed matrix cracking due to cre 11315°C 1315°C 150/167 1315°C 82/>103 147/185 0.09 A21315°C 103/18 1315°C 6/188 1315°C 88/103 59/84 0.3 Some bridged surface cracks propagate 4 pI 0.23 Cracks only in surface 90 minicomposites 1315°C 0.44 Some bridged cracks containing local unbridged 95/>103 08/143 0.07 1315°C 138 90/107 36 Some bridged surface cracks propagate mately 60 MPa(Fig. 4a). This compressive stress was Cracks formed during after-creep fast-fracture were generally cantly higher than the 20 MPa residual stress obtained much finer, and could be observed only after plasma etching the as-polished sections. The difficulty in observing these cracks is matrix relaxes during creep so that upon stress removal, it was again because of the residual compressive stress in the matrix placed in even greater compression which had to be overcome in hich results in very little crack opening. The polished micro- order to cause matrix crack formation and growth. Micro- structures described here are as-polished sections which were not ructural observations of the specimen after fracture, as shown plasma etched in order to show the effects of creep on micro- below, are consistent with this notion, as only minor surface ructural change. micro-cracks were observed in this specimen. No matrix cracking attributed to creep could be discerned for The tensile spe n that had been crept for 138 MPa and imens from all three panels that were subject to creep at 103 100 h from panel B(high Si content, low CVI SiC content)was MPa. The specimen from panel Al that was subjected to tensile also tensile hysteresis tested at room temperature (Fig. 5(a)) reep at 138 MPa for 100 h and subsequently tested at room This specimen had the highest creep strain without failing during temperature did have relatively wide open, oxide- filled surface 100 h creep. The differences between the tensile stress-strain cracks emanating from the 90 tows at the surface of the com- havior of the specimen before creep and after creep was not posite; however, the matrix cracks did not appear to penetrate the same as for the panel A specimen (Table IV). For the panel b into the adjacent 0 fiber tow(Fig. 6). For the specimen from specimen no real difference was observed in offset-stress or AE panel A2 with a creep strain of 0.3% after 138 MPa creep which activity(Fig. 5(b)and only a mild improvement was observed as immediately fast fractured at 1315.C and the specimen from n residual compressive stress. Also the elastic modulus at room panel B subjected to 138 MPa tensile creep and tested at room temperature for this crept specimen was particularly low in temperature, similar surface 90 matrix cracks were observed to comparison with its as-produced modulus at room temperature ropagate several plies into the composite and were bridged by the specimen showed the specimen had bridged matrix cracks 0 fibers (Table IIf). The specimen from panel A2 that was crept at 172 MPa was the only specimen that showed fiber failure as- that had to occur during creep(next section). As matrix cracks lready existed before testing at room temperature after the cracks. However, with the exception of the failure crack, the creep test, it is not surprising that the elastic modulus was lower unbridged portion of the matrix cracks were localized and were and that no real improvement in matrix cracking stress was ob- all associated with the same 0 tow, three plies from one side of the specimen(Fig. 7). Six unbridged cracks were observed for the section polished and were spaced 2.3, 3.2, 2.6, 1.6, 1.2, and 1. 6 mm apart from one another moving away from the fracture 3) Matrix Cracking surface. Some of the unbridged matrix cracks did extend The matrix cracking behavior of several specimens is described er-bridged cracks to neighboring plies and/or extend to th failed specimens. Only the 172 MPa crept specimen had failed in appeared to extend a total of four plies, i. e, half the thickness pture. All of the other specimens were failed during the afte However, some of the cracks did not appear to propagate more creep tensile testing, which always occurred at a higher load thanthan one ply away from the unbridged bundle. The rest of the the creep stress. These latter specimens would be expected te matrix cracks were in between one and four plies in length. The have larger amounts of matrix cracks because of the fast-frac- polished surface in Fig. 7 corresponds to the region of the gage ture condition. It was generally easy to distinguish matrix cracks section that was 2 mm from one edge of the specimen. To de- formed during creep from those that were formed during the termine if any unbridged cracks existed in the interior of the after-creep fast-fracture test Matrix cracks formed during creep specimen, an additional 4 mm was sliced from the mounted were easy to observe in polished sections because they possessed polished section (i. e, 6 mm from the same edge or 4 mm from gnificant crack openings were generally accompanied by edge). None were observed oxide formation within the crack and along the pores, especially that significant fiber-bridged matrix cracking oc- those cracks that clearly extended to the composite surface. curred in specimens which had exhibited significant creep strains

mately 60 MPa (Fig. 4a). This compressive stress was signifi- cantly higher than the 20 MPa residual stress obtained for the as-produced specimen.7 These data support the notion that the matrix relaxes during creep so that upon stress removal, it was placed in even greater compression which had to be overcome in order to cause matrix crack formation and growth.11–13 Micro￾structural observations of the specimen after fracture, as shown below, are consistent with this notion, as only minor surface micro-cracks were observed in this specimen. The tensile specimen that had been crept for 138 MPa and 100 h from panel B (high Si content, low CVI SiC content) was also tensile hysteresis tested at room temperature (Fig. 5(a)). This specimen had the highest creep strain without failing during 100 h creep. The differences between the tensile stress–strain behavior of the specimen before creep and after creep was not the same as for the panel A specimen (Table IV). For the panel B specimen no real difference was observed in offset-stress or AE activity (Fig. 5(b)) and only a mild improvement was observed in residual compressive stress. Also the elastic modulus at room temperature for this crept specimen was particularly low in comparison with its as-produced modulus at room temperature and at temperature before creep (see Table II). Microscopy of the specimen showed the specimen had bridged matrix cracks that had to occur during creep (next section). As matrix cracks already existed before testing at room temperature after the creep test, it is not surprising that the elastic modulus was lower and that no real improvement in matrix cracking stress was ob￾served in this specimen. (3) Matrix Cracking The matrix cracking behavior of several specimens is described in Table III as determined from polished longitudinal sections of failed specimens. Only the 172 MPa crept specimen had failed in rupture. All of the other specimens were failed during the after￾creep tensile testing, which always occurred at a higher load than the creep stress. These latter specimens would be expected to have larger amounts of matrix cracks because of the fast-frac￾ture condition. It was generally easy to distinguish matrix cracks formed during creep from those that were formed during the after-creep fast-fracture test. Matrix cracks formed during creep were easy to observe in polished sections because they possessed significant crack openings and were generally accompanied by oxide formation within the crack and along the pores, especially those cracks that clearly extended to the composite surface. Cracks formed during after-creep fast-fracture were generally much finer, and could be observed only after plasma etching the as-polished sections.6 The difficulty in observing these cracks is again because of the residual compressive stress in the matrix which results in very little crack opening. The polished micro￾structures described here are as-polished sections which were not plasma etched in order to show the effects of creep on micro￾structural change. No matrix cracking attributed to creep could be discerned for specimens from all three panels that were subject to creep at 103 MPa. The specimen from panel A1 that was subjected to tensile creep at 138 MPa for 100 h and subsequently tested at room temperature did have relatively wide open, oxide-filled surface cracks emanating from the 901 tows at the surface of the com￾posite; however, the matrix cracks did not appear to penetrate into the adjacent 01 fiber tow (Fig. 6). For the specimen from panel A2 with a creep strain of 0.3% after 138 MPa creep which was immediately fast fractured at 13151C and the specimen from panel B subjected to 138 MPa tensile creep and tested at room temperature, similar surface 901 matrix cracks were observed to propagate several plies into the composite and were bridged by 01 fibers (Table III). The specimen from panel A2 that was crept at 172 MPa was the only specimen that showed fiber failure as￾sociated with matrix cracks, i.e., unbridged portions of matrix cracks. However, with the exception of the failure crack, the unbridged portion of the matrix cracks were localized and were all associated with the same 01 tow, three plies from one side of the specimen (Fig. 7). Six unbridged cracks were observed for the section polished and were spaced 2.3, 3.2, 2.6, 1.6, 1.2, and 1.6 mm apart from one another moving away from the fracture surface. Some of the unbridged matrix cracks did extend as fib￾er-bridged cracks to neighboring plies and/or extend to the sur￾face. The bridged portion of some of these matrix cracks appeared to extend a total of four plies, i.e., half the thickness. However, some of the cracks did not appear to propagate more than one ply away from the unbridged bundle. The rest of the matrix cracks were in between one and four plies in length. The polished surface in Fig. 7 corresponds to the region of the gage section that was 2 mm from one edge of the specimen. To de￾termine if any unbridged cracks existed in the interior of the specimen, an additional 4 mm was sliced from the mounted polished section (i.e., 6 mm from the same edge or 4 mm from the opposite edge). None were observed. It appears that significant fiber-bridged matrix cracking oc￾curred in specimens which had exhibited significant creep strains Table III. Offset Stress at 13151C and Matrix Cracking Observations Panel Experiment Creep stress (MPa) 0.002%/0.005% offset stress Creep strain (%) Observed matrix cracking due to creep 13151C initial loading (MPa) 13151C after creep (MPa) A1 13151C — 70 — — — 13151C creep 103 NA 150/167 0.18 — 13151C creep 103 82/4103 147/185 0.09 None A2 13151C — 103/118 — — — 13151C creep 103 67/97 166/188 0.05 None 13151C creep 138 88/103 159/184 0.3 Some bridged surface cracks propagate up to 4 plies 13151C creep 138 79/98 — 0.23 Cracks only in surface 901 minicomposites 13151C creep 172 85/103 — 0.44 Some bridged cracks containing local unbridged regions B 13151C creep 103 95/4103 108/143 0.07 None 13151C creep 138 90/107 — 0.36 Some bridged surface cracks propagate up to 4 plies Each row corresponds to an individual tensile specimen. May 2006 Creep and After-Creep Stress–Strain Behavior 1655

1656 Journal of the American Ceramic Society-Morscher and pujar Vol. 89. No. 5 (a) 400PanelA2After Panel B 138 Mpa 100 h Creep As-produced creep strain =0.23% AS-Produced 250 (hys loops removed Panel B After 138 MPa 100 100 h Creep creep strain= 0.36% 50 Compressive Stress 0.1 0.2 0.304 0.1 Strain. 0.8 0.8 u0.7 60.5 E0.4 0.produced 20.3 After-Creep 0 100150200250300350 Stress, MP 125MPa0.002% Fig. 5.(a)Room temperature unload-reload hysteresis curves to fail- ure for panel B specimens. Note, the hysteresi from the as-produced specimen for clarity.(b) Acoustic emission gener- ated from the gage section for both tensile specimens 1201 Panel A2 As-Produced 0.005% offset stress (>0.3%). This suggests most of the composite creep observed n these specimens was primarily because of the creep of the constituents in combination with some minor bridged crack 0. 002% offset wth 40 20 0.005%oset 0.06 For this HNS/MI composite system, excellent creep resistance % as observed when loaded up to 103 MPa at 1315C in the pri- Fig 4.(a) Room ary fiber direction. The creep strains measured in this study are s Note, the hysteresis loops have been re. comparable with those measured for the Sylramic-iBN MI com- oved from the as. men for clarity.(b)Acoustic emission posites when tested under similar conditions. For example, the generated from th tion for both tensile specimens. (c)The con- offset stresses for the as-produced specimen and 103 MPa for 100 h were 0.10%+0.05%. similar to the ested to failur creep strains measured in this study for the same conditions Table ID). In addition, composite creep appears to lead to load relaxation in the matrix and consequently residual compressive Table IV. Properties Associated with Non-Linearity in Room Temperature Tested Composites As-Produced and After 138 MPa, 100 Panela Panel b Mechanical property As-produced specimen After creep specimen s-produced specimen After creep specimen 0.002% offset 125 MPa 142 MPa 151 MPa 150 MPa 0.005% offset 147 MPa 177 MPa 169 MPa 177 MPa Significant AE 115 MP 140 MPa 160 MPa 160 MPa Matrix compression 20 MP 60 MPa 35 MPa 55 MPa Creep strain 0.23% 0.36% RT elastic modul 232 GPa 233 Gpa 270 Gpa Ultimate strength 412 MPa 321 MP 362 MPa 247 MPa n refers to a different tensile specimen from the same panel, not the same specimen before and after creep

( 0.3%). This suggests most of the composite creep observed in these specimens was primarily because of the creep of the constituents in combination with some minor bridged crack growth. IV. Discussion For this HNS/MI composite system, excellent creep resistance was observed when loaded up to 103 MPa at 13151C in the pri￾mary fiber direction. The creep strains measured in this study are comparable with those measured for the Sylramic-iBN MI com￾posites when tested under similar conditions. For example, the creep strains of Sylramic-iBN MI composites tested at 13151C and 103 MPa for 100 h were 0.10%70.05%,10 similar to the creep strains measured in this study for the same conditions (Table II). In addition, composite creep appears to lead to load relaxation in the matrix and consequently residual compressive 0 50 100 150 200 250 300 350 400 450 0 0.1 0.2 0.3 0.4 0.5 0.6 Strain, % Stress, MPa Panel A2 After 138 Mpa 100 h Creep creep strain = 0.23% Panel A2 As-Produced (hys loops removed) Residual Compressive Stress 0 20 40 60 80 100 120 140 160 0 0.02 0.04 0.06 0.08 Strain, % Stress, MPa Panel A2 As-Produced (E = 232 GPa) 0.002% offset 147 MPa 0.005% offset stress 0.005% offset 125 MPa 0.002% offset stress (a) (b) (c) 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 Norm Cum AE As￾produced After Creep 50 100 150 200 250 300 350 Stress, MPa Fig. 4. (a) Room temperature unload–reload hysteresis curves to fail￾ure for panel A2 specimens. Note, the hysteresis loops have been re￾moved from the as-produced specimen for clarity. (b) Acoustic emission generated from the gage section for both tensile specimens. (c) The con￾struction used to determine offset stresses for the as-produced specimen tested to failure. 0 50 100 150 200 250 300 350 400 0 0.1 0.2 0.3 0.4 0.5 0.6 Strain, % Stress, MPa Panel B As-produced (hys loops removed) Panel B After 138 MPa 100 h Creep creep strain= 0.36% Residual Compressive Stress (b) (a) 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 Norm Cum AE After-Creep As￾Produced 100 150 200 250 300 350 Stress, MPa Fig. 5. (a) Room temperature unload–reload hysteresis curves to fail￾ure for panel B specimens. Note, the hysteresis loops have been removed from the as-produced specimen for clarity. (b) Acoustic emission gener￾ated from the gage section for both tensile specimens. Table IV. Properties Associated with Non-Linearity in Room Temperature Tested Composites As-Produced and After 138 MPa, 100 H Creep Mechanical property Panel A Panel B As-produced specimenw After creep specimenw As-produced specimenw After creep specimenw 0.002% offset 125 MPa 142 MPa 151 MPa 150 MPa 0.005% offset 147 MPa 177 MPa 169 MPa 177 MPa Significant AE 115 MPa 140 MPa 160 MPa 160 MPa Matrix compression 20 MPa 60 MPa 35 MPa 55 MPa Creep strain — 0.23% — 0.36% RT elastic modulus 232 GPa 233 Gpa 270 Gpa 214 Gpa Ultimate strength 412 MPa 321 MPa 362 MPa 247 MPa w Each specimen refers to a different tensile specimen from the same panel, not the same specimen before and after creep. 1656 Journal of the American Ceramic Society—Morscher and Pujar Vol. 89, No. 5

May 2006 Creep and After-Creep Stress-Strain Behavior 1657 cracking strengths in a turbine blade com ature at relatively low stress conditions and for a short time before insertion into service. Improved matrix strengths would also result in applications involving stresses that are for most part thermally driven, as these can still alter the residual stress state in a similar fashion. Most importantly, the creep results show that matrix cracking properties actually im- rove with use, provided the matrix is more creep prone than the reinforcing fibers and can accommodate a high creep strain without cracking. Both attributes appear to exist for the as-pro- duced CVI-SiC matrix of this study as it was able to deform at least enough to accommodate a relatively high creep strain (>0.2% permanent strain on average) Finally, these results show that one should be cautious when nterpreting the stress for non-linearity(proportional limit)at 138 MPa creep; Fast-fracture at 0.1mm 1315C for these composite systems. Non-linearity is often as- oom temperature aft ociated with matrix cracking, but as it was observed in this study, much non-linearity early in the stress-strain curve was Fig. 6. Micrograph of surface 90 crack. because of creep. Faster loading rates may result in higher elastic dulus and different offset stresses the offset-stress stress in the matrix upon rapid unloading from the applied stress measurements of an after-creep specimen should not be affected ondition. This infers that a portion of the creep strain was by creep as the material has effectively experienced"creep-hard simply because of increased loading of the more creep-resistant ening. fibers and to some extent CVI-SiC portion of the matrix. It also infers that sufficient creep followed by rapid stress removal and specimen cool-down can lead to increased compressive stress on V. Conclusions the matrix and higher off-set stresses for matrix thru-thickness Woven Hi-Nicalon Type S, melt-infiltrated SiC matrix com- cracking under fast or slow fracture conditions at low temper- ites were shown to be relatively creep resistant because of the atures and under fast-fracture conditions near the creep tem- eep-resistant nature of the reinforcing fiber, consistent with perature. The need for rapid stress removal and specimen cool- reep behavior for similar SiC/SiC systems. Tensile creep at down follows from the possibility that upon stress removal at d stresses up to 138 MPa in the primary fiber direction the creep temperature, residual tensile stresses on the fibers ays survived the 100 h limit imposed in this study. Fast could force the matrix to creep in compression, thereby reduc- ing the tensile stresses on the fibers and corresponding com- fracture tensile testing of specimens after 103 or 138 MPa for pressive stresses on the matrix. However, creep rate in 100 h creep showed significantly higher offset stresses, or the compression of siliconized Sic is known to be over an order as-producte -linear stress-strain behavior when compared with ites. This was attributed to the stress relax- of magnitude less than in compression and may minimize the matrix recovery. Because of the limited number of specimens, tion of the matrix during tensile creep resulting in an increased ss to ci to use not enough tests were performed to determine the optimum attribute of composite system to raise matrix cracking stresses, a esign parameter for CMCs, or at least anticipate a stress relaxation will occur under reduced load at the creep tem- in matrix cracking stress for certain applications subjected to perature. However, it is apparent that the small amounts of creep strain for the 103 MPa creep condition can lead to signif- low applied stresses. icant increase in the off-set stresses of the o/e curve ( Tables II and Ill and Fig. 3) More detailed studies are needed to optimize the constituents References nd temperature-time-stress conditions to maximize the matrix Rospars, J. L. Chermant, and P. Ladeveze, ""On a First Creep Model for a nomenon could be used to enhance matrix cracking n applications where this is desirable. For example, Fatigue Behavior in Hi-NicalonTM-Fiber-Reinforced Silicon Carbide Composites at High Temperatures. "J.Am. Ceram. Soc., 81[1] 117-28(1999 3S. Zhu. M. Mizuno. y Kagawa and Y Mut onotonic Tension. Fati Creep Behavior of SiC- Fiber-Reinforced SiC-Matrix Composites: A Re Comp. Sci. Technol, 59 [6]833-51( 1999). R. Bunsell and M-H. Berger. "Fine Diameter Ceramic Fibres. J. Eur. SG.N. Morscher and V. V. Pujar. "Melt-Infiltrated SiC Com urbine gs of the 49th ASME IGTI Turbo Lane H M. Yun and J.A. DiCarlo, Comparison of the Tensile Creep, and Rupture ngth Properties of Stoichiometric SiC Fibers, Ceram. Eng. Sci. Proc., 20[31 G.N. Morscher and J. Z. Gyekenesi "Room Temperature Tensile Behavior and Damage Accumulation of Hi-Nicalon Reinforced SiC Matrix Composites, L. Guillaumat and J. Lamon, "Multi-Fissuration De Composites SiC/SiC. Matrix Cracking in 2D Woven Sic-Fibe Reinforced Melt-Infiltrated SiC Matrix Composites, Comp. Sci. Techno/ 64 Panel A2 IH. M. Yun, J. A. DiCarlo, R. T. Bhatt and J.B. Hurst, "Processing and 175 MPa Creep SiC Fiber for SiC/SiC Components. 1 mm 四g.Sci.Pro,24[3247-53(2003) olmes."Influence of Stress Ratio on the Elevated-Temperature Fa- le Fiber-Reinforced Silicon Nitride Composite, "J.An. ig. 7. Micrograph of unbridged cracks in interior 0 fiber tow ceran.Soc,74[163945(1991)

stress in the matrix upon rapid unloading from the applied stress condition. This infers that a portion of the creep strain was simply because of increased loading of the more creep-resistant fibers and to some extent CVI–SiC portion of the matrix. It also infers that sufficient creep followed by rapid stress removal and specimen cool-down can lead to increased compressive stress on the matrix and higher off-set stresses for matrix thru-thickness cracking under fast or slow fracture conditions at low temper￾atures and under fast-fracture conditions near the creep tem￾perature. The need for rapid stress removal and specimen cool￾down follows from the possibility that upon stress removal at the creep temperature, residual tensile stresses on the fibers could force the matrix to creep in compression, thereby reduc￾ing the tensile stresses on the fibers and corresponding com￾pressive stresses on the matrix. However, creep rate in compression of siliconized SiC is known to be over an order of magnitude less than in compression and may minimize the matrix recovery.16 Because of the limited number of specimens, not enough tests were performed to determine the optimum condition for maximizing matrix compression or whether matrix stress relaxation will occur under reduced load at the creep tem￾perature. However, it is apparent that the small amounts of creep strain for the 103 MPa creep condition can lead to signif￾icant increase in the off-set stresses of the s/e curve (Tables II and III and Fig. 3). More detailed studies are needed to optimize the constituents and temperature–time–stress conditions to maximize the matrix cracking stresses. Neverthless, this study shows that this phe￾nomenon could be used to enhance matrix cracking properties in applications where this is desirable. For example, the matrix cracking strengths in a turbine blade component could be im￾proved by subjecting the component to a higher than expected use temperature at relatively low stress conditions and for a short time before insertion into service.12 Improved matrix strengths would also result in applications involving stresses that are for most part thermally driven, as these can still alter the residual stress state in a similar fashion. Most importantly, the creep results show that matrix cracking properties actually im￾prove with use, provided the matrix is more creep prone than the reinforcing fibers and can accommodate a high creep strain without cracking. Both attributes appear to exist for the as-pro￾duced CVI–SiC matrix of this study as it was able to deform at least enough to accommodate a relatively high creep strain (40.2% permanent strain on average). Finally, these results show that one should be cautious when interpreting the stress for non-linearity (proportional limit) at 13151C for these composite systems. Non-linearity is often as￾sociated with matrix cracking, but as it was observed in this study, much non-linearity early in the stress–strain curve was because of creep. Faster loading rates may result in higher elastic modulus and different offset stresses. However, the offset-stress measurements of an after-creep specimen should not be affected by creep as the material has effectively experienced ‘‘creep-hard￾ening.’’ V. Conclusions Woven Hi-Nicalon Type STM, melt-infiltrated SiC matrix com￾posites were shown to be relatively creep resistant because of the creep-resistant nature of the reinforcing fiber, consistent with creep behavior for similar SiC/SiC systems. Tensile creep at ap￾plied stresses up to 138 MPa in the primary fiber direction al￾ways survived the 100 h limit imposed in this study. Fast￾fracture tensile testing of specimens after 103 or 138 MPa for 100 h creep showed significantly higher offset stresses, or the onset of non-linear stress–strain behavior when compared with as-produced composites. This was attributed to the stress relax￾ation of the matrix during tensile creep resulting in an increased stress to cause matrix cracking. It may be possible to use this attribute of composite system to raise matrix cracking stresses, a typical design parameter for CMCs, or at least anticipate a rise in matrix cracking stress for certain applications subjected to low applied stresses. References 1 C. Rospars, J. L. Chermant, and P. Ladeveze, ‘‘On a First Creep Model for a 2D SiCf–SiC Composite,’’ Mater. Sci. Eng. A, A250, 264–9 (1998). 2 S. Zhu, M. Mizuno, Y. Kagawa, J. Cao, Y. Nagano, and H. Kaya, ‘‘Creep and Fatigue Behavior in Hi-NicalonTM-Fiber-Reinforced Silicon Carbide Composites at High Temperatures,’’ J. Am. Ceram. Soc., 81 [1] 117–28 (1999). 3 S. Zhu, M. Mizuno, Y. Kagawa, and Y. Mutoh, ‘‘Monotonic Tension, Fatigue and Creep Behavior of SiC-Fiber-Reinforced SiC–Matrix Composites: A Re￾view,’’ Comp. Sci. Technol., 59 [6] 833–51 (1999). 4 A. R. Bunsell and M.-H. Berger, ‘‘Fine Diameter Ceramic Fibres,’’ J. Eur. Ceram. Soc., 20, 2249–60 (2000). 5 G. N. Morscher and V. V. Pujar, ‘‘Melt-Infiltrated SiC Composites for Gas Turbine Engine Applications’’; Proceedings of the 49th ASME IGTI Turbo Land, Sea and Air Conference, June 14–17, Vienna, Austria, 2004. Paper number: GT2004-53196. 6 H. M. Yun and J. A. DiCarlo, ‘‘Comparison of the Tensile Creep, and Rupture Strength Properties of Stoichiometric SiC Fibers,’’ Ceram. Eng. Sci. Proc., 20 [3] 259–72 (1999). 7 G. N. Morscher and J. Z. Gyekenesi, ‘‘Room Temperature Tensile Behavior and Damage Accumulation of Hi-Nicalon Reinforced SiC Matrix Composites,’’ Ceram. Eng. Sci. Proc., 19 [3] 241–9 (1998). 8 L. Guillaumat and J. Lamon, ‘‘Multi-Fissuration De Composites SiC/SiC,’’ Rev. Comp. Mater. Avances., 3, 159–71 (1993). 9 G. N. Morscher, ‘‘Stress-Dependent Matrix Cracking in 2D Woven SiC-Fiber Reinforced Melt-Infiltrated SiC Matrix Composites,’’ Comp. Sci. Technol., 64 [9] 1311–9 (2004). 10H. M. Yun, J. A. DiCarlo, R. T. Bhatt, and J. B. Hurst, ‘‘Processing and Structural Advantages of the Sylramic-iBN SiC Fiber for SiC/SiC Components,’’ Ceram. Eng. Sci. Proc., 24 [3] 247–53 (2003). 11J. W. Holmes, ‘‘Influence of Stress Ratio on the Elevated-Temperature Fa￾tigue of a Silicon Carbide Fiber-Reinforced Silicon Nitride Composite,’’ J. Am. Ceram. Soc., 74 [7] 1639–45 (1991). Fig. 6. Micrograph of surface 901 crack. Fig. 7. Micrograph of unbridged cracks in interior 01 fiber tow. May 2006 Creep and After-Creep Stress–Strain Behavior 1657

Journal of the American Ceramic Society--Morscher and Pujar Vol. 89. No 5 E. Lara-Curzio and c.m. russ. "On the matrix in ASTM STP 1309, Edited by M. G. Jenkins, et al. American Society for Testing Redistribution and Materials. West Conshohocken, PA. 1997. SS. Kalluri, A Calomino, and D N. Brewer, "An Assessment of variability in idaja. K. Jakus, J.E. Ritter, E. Lara-Curzio, T.R. Watkins, E.Y. Sun, the Average Tensile Properties of a Melt-Infiltrated SiC/SiC Composite, " Ceram. and J. J. Brennan, ""Creep-Induced Residual Strengthening in a Nicalon-Fiber Sci.Proe.25786(2004) Reinforced BMAS-Glass-Ceramic-Matrix Composite. "J. Am. Ceram. Soc., 82 obert. and T- Chuang, ""Dama ge-Enhancey 657-72101999) Creep in a Siliconized Silicon Carbide: Phenomenology. " J. Am. Ceram. Soc., 71 IM. Steen and J. L Valles. "Unloading-Reloading Sequences and the 7602-8(1988) of Mechanical Test Results for Continuous Fiber Ceramic Composites", P

12E. Lara-Curzio and C. M. Russ, ‘‘On the Matrix Cracking Stress and the Redistribution of Internal Stresses in Brittle–Matrix Composites,’’ Mater. Sci. Eng. A, 250 [2] 270–8 (1998). 13S. Widaja, K. Jakus, J. E. Ritter, E. Lara-Curzio, T. R. Watkins, E. Y. Sun, and J. J. Brennan, ‘‘Creep-Induced Residual Strengthening in a Nicalon-Fiber￾Reinforced BMAS-Glass-Ceramic-Matrix Composite,’’ J. Am. Ceram. Soc., 82 [3] 657–721 (1999). 14M. Steen and J. L. Valles, ‘‘Unloading-Reloading Sequences and the Analysis of Mechanical Test Results for Continuous Fiber Ceramic Composites’’; pp 49–65 in ASTM STP 1309, Edited by M. G. Jenkins, et al. American Society for Testing and Materials, West Conshohocken, PA, 1997. 15S. Kalluri, A. Calomino, and D. N. Brewer, ‘‘An Assessment of Variability in the Average Tensile Properties of a Melt-Infiltrated SiC/SiC Composite,’’ Ceram. Eng. Sci. Proc., 25 [4] 79–86 (2004). 16S. M. Wiederhorn, D. E. Roberts, and T-J Chuang, ‘‘Damage-Enhanced Creep in a Siliconized Silicon Carbide: Phenomenology,’’ J. Am. Ceram. Soc., 71 [7] 602–8 (1988). & 1658 Journal of the American Ceramic Society—Morscher and Pujar Vol. 89, No. 5

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