A) continuous fiber ceramic composite

Composites: Part A 37 (2006) 2029–2040 www.elsevier.com/locate/compositesa

Effects of steam environment on high-temperature mechanical behavior of NextelTM720/alumina (N720/A) continuous fiber ceramic composite q M.B. Ruggles-Wrenn *, S. Mall, C.A. Eber, L.B. Harlan Department of Aeronautics and Astronautics, Air Force Institute of Technology, WPAFB, OH 45433-7765, USA Received 15 February 2005; received in revised form 30 November 2005; accepted 13 December 2005

Abstract Mechanical behavior of an oxide–oxide continuous fiber ceramic composite (CFCC) consisting of a porous alumina matrix reinforced with laminated, woven mullite/alumina fibers (NextelTM720) was investigated at 1200 and 1330 °C in laboratory air and in 100% steam environments. CFCC has no interface between the fiber and matrix, and relies on the porous matrix for flaw tolerance. Tension–tension fatigue behavior was studied for fatigue stresses ranging from 100 to 170 MPa at 1200 °C, and for fatigue stresses of 50 and 100 MPa at 1330 °C. Tensile creep behavior was examined for creep stresses ranging from 80 to 154 MPa at 1200 °C, and for creep stresses of 50 and 100 MPa at 1330 °C. At 1200 °C, the CFCC exhibited excellent fatigue resistance in laboratory air. The fatigue limit (based on a run-out condition of 105 cycles) was 170 MPa (88% UTS at 1200 °C). The material retained 100% of its tensile strength. Presence of steam caused noticeable degradation in fatigue performance at 1200 °C. Fatigue resistance at 1330 °C was poor. In creep tests, primary and secondary creep regimes were observed. Minimum creep rate was reached in all tests. At 1200 °C, creep rates were 108–105 s1 and maximum time to rupture was 255 h. At 1330 °C, creep rates were 107–105 s1 and maximum time to rupture was 87 h. Presence of steam accelerated creep rates and dramatically reduced creep life. Published by Elsevier Ltd. Keywords: A. Ceramic-matrix composites (CMCs); A. Fibres; A. Creep

1. Introduction Aerospace components require structural materials that have superior long-term mechanical properties and can be exposed to severe environmental conditions, such as high temperature, high pressure, or water vapor. Ceramic– matrix composites (CMCs), capable of maintaining excellent strength and fracture toughness at high temperatures continue to attract attention as candidate materials for q The views expressed are those of the authors and do not reflect the official policy or position of the United States Air Force, Department of Defense or the U.S. Government. * Corresponding author. Tel.: +937 255 3636x4641; fax: +937 656 4032. E-mail address: marina.ruggles-wrenn@afit.edu (M.B. RugglesWrenn).

1359-835X/$ - see front matter Published by Elsevier Ltd. doi:10.1016/j.compositesa.2005.12.008

aerospace applications. Compared to the conventional nickel-based superalloys, CMCs offer improved hightemperature performance at reduced weight. Advanced reusable space launch vehicles will likely incorporate fiber-reinforced CMCs in critical propulsion components. In these applications, CMCs will be subjected to mechanical loading in complex environments. For example, a typical service environment for a reusable rocket engine turbopump rotor includes hydrogen, oxygen and steam, at pressures >200 atm [1]. Ceramic–matrix composites are also being considered for aerospace turbine engine applications. Higher material operating temperatures and decreased cooling air requirement are the significant advantages that CMCs offer to the aerospace engine design community. However, these applications also require exposure to oxidizing environments. Recently CMCs have been

2030

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

demonstrated in various turbine components [2]. Many of these demonstration components have exhibited accelerated degradation after only a few hours in service environment. It is now widely recognized that the thermodynamic stability and oxidation resistance of CMCs have become important issues. Sintered structural ceramics are known to exhibit degradation in high-temperature environments. Non-oxide fiber/ non-oxide matrix composites generally show poor oxidation resistance [3,4]. The degradation involves oxidation of fibers and fiber coatings, and is accelerated by the presence of moisture [5–7]. Numerous studies addressed oxidation of SiC in moist environments [8–13]. Opila and Hann [9] and Opila [10,11] reported that the presence of water vapor increased the rate of SiO2 growth on SiC at high temperature, which led to accelerated rates of SiC recession. Degradation of BN fiber coatings in moist environments has also been a subject of extensive research [14–20]. Non-oxide fiber/oxide matrix composites or oxide fiber/non-oxide matrix composites do not exhibit high oxidation resistance either. For these materials, the high permeability constant for the diffusion of oxygen results in rapid oxygen permeation through the oxide matrix [21]. These considerations motivated the development of continuous fiber ceramic composites (CFCCs) based on environmentally stable oxide constituents [22–30]. The main advantage of CMCs over monolithic ceramics is their superior toughness, tolerance to the presence of cracks and defects, and non-catastrophic mode of failure. It is now well recognized that CFCCs can be designed to exhibit non-brittle fracture behavior and improved damage tolerance by introducing a weak fiber/matrix interface, which serves to deflect matrix cracks and to allow subsequent fiber pull-out [31–33]. It has recently been demonstrated that similar crack-deflecting behavior can also be achieved by means of a finely distributed porosity in the matrix instead of a separate interface between matrix and fibers [34]. This microstructural design philosophy implicitly accepts the formation of strong interfaces. It builds on the experience with porous interlayers as crack deflection paths [35,36] and extends the concept to utilize a porous matrix as a surrogate. The concept has been successfully demonstrated for oxide–oxide composites [22,26,30,37–41]. Resulting oxide/oxide CFCCs exhibit damage tolerance combined with inherent oxidation resistance. An extensive review of the mechanisms and mechanical properties of porous-matrix CMCs is given in [42]. The objective of this study is to investigate effects of steam environment on high-temperature mechanical behavior and durability of an oxide–oxide CFCC, consisting of a porous alumina matrix reinforced with the NextelTM720 fibers. Several previous studies examined high-temperature mechanical behavior of this material [2,43], and [44]. Unlike its counterpart reinforced with SiC fibers [45–48], the CFCC exhibited steady-state creep. Creep rates were 108–107 s1, similar to those expected from fibers alone. Zawada et al. [2] investigated the effect

of intermittent moisture exposure on the high-temperature fatigue durability of five different CMCs, including N720/ A. Fatigue testing of N720/A was performed at 1200 °C, fatigue stress levels were 6120 MPa. Cyclic loading and moisture exposure were applied alternately. Zawada et al. observed no degradation in fatigue performance or retained strength with intermittent moisture exposure. In the present study, fatigue and creep-rupture testing of N720/A specimens was conducted both in laboratory air and in 100% steam environment at high temperatures (1200 and 1330 °C). Applied stress levels used in both creep-rupture and fatigue tests (6170 MPa) were considerably higher than those employed in previous studies. As is seen in detail, effects of steam environment on fatigue and especially on creep performance cannot be neglected. 2. Experimental procedure 2.1. Material The composite studied was a commercially available material (N720/A, COI Ceramics, San Diego, CA) consisting of NextelTM720 fibers in a porous alumina matrix, supplied in a form of 2.8 mm thick plates. The plates consisted of 12 0°/90° woven layers, with a density of 2.78 g/cm3 and a fiber volume of approximately 44%. There was no fiber coating. The fiber fabric was infiltrated with the matrix in a sol–gel process. After drying with a ‘‘vacuum bag’’ technique under low pressure and low temperature, the composite was pressureless sintered [49]. Matrix porosity was 24%. Such porosity level renders the matrix sufficiently weak and gives the composite excellent damage tolerance during loading. Representative micrographs of the untested as-received material are shown in Fig. 1. Fig. 1(a) shows 0° and 90° fiber tows as well as numerous matrix cracks. In the case of untested material, most are shrinkage cracks formed during processing rather than matrix cracks generated during loading. Porous nature of the matrix is seen in Fig. 1(b). 2.2. Mechanical testing A servocontrolled MTS mechanical testing machine equipped with hydraulic water-cooled collet grips, a compact two-zone resistance-heated furnace, and two temperature controllers was used in all tests. An MTS TestStar digital controller was employed for input signal generation and data acquisition. Strain measurement was accomplished with an MTS high-temperature air-cooled uniaxial extensometer. For elevated temperature testing, thermocouples were bonded to the specimens to calibrate the furnace on a periodic basis. The furnace controller (using a non-contacting thermocouple exposed to the ambient environment near the test specimen) was adjusted to determine the power setting needed to achieve the desired temperature of the test specimen. Thus determined power setting was then used in actual tests. The power setting for testing

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

2031

Creep-rupture tests were conducted in load control in accordance with the procedure in ASTM standard C 1337. In each test, stress–strain data were recorded during the loading to the creep stress level and the actual creep period. Thus both total strain and creep strain could be calculated and examined. Creep-rupture tests were carried out at 1200 and 1330 °C, in laboratory air and in steam environments. Tension–tension fatigue tests were carried out in load control with an R ratio (minimum stress divided by the maximum stress) of 0.05 at a frequency of 1 Hz. Fatigue run-out was set to 105 cycles. The 105 cycle count value represents the number of loading cycles expected in aerospace applications at that temperature. Fatigue run-out limits were defined as the highest stress level, for which run-out was achieved. Cyclic stress–strain data were recorded throughout each test. In order to assess the damage development in the composite, stiffness degradation, changes in hysteresis energy density, as well as strain accumulation with fatigue cycles were examined. To determine retained strength and stiffness, specimens that achieved run-out were subjected to tensile test to failure at the temperature of the fatigue test. Fatigue tests were performed at 1200 and 1330 °C, in laboratory air and in steam environments. 2.3. Characterization Fracture surfaces of failed specimens were examined using SEM (Model 360FE, Leica) as well as optical microscopy. The SEM specimens were gold coated. 3. Results and discussion

in steam environment was determined by placing the specimen instrumented with thermocouples in 100% steam environment and repeating the furnace calibration procedure. Thermocouples were not bonded to the test specimens after the furnace was calibrated. Tests in steam environment employed an alumina susceptor (tube with end caps), which fits inside the furnace. The specimen gage section is located inside the susceptor, with the ends of the specimen passing through slots in the susceptor. Steam is introduced into the susceptor (through a feeding tube) in a continuous stream with a slightly positive pressure, expelling the dry air and creating 100% steam environment inside the susceptor. In all tests, a specimen was heated to test temperature in 25 min, and held at temperature for additional 15 min prior to testing. Dog bone shaped specimens of 152 mm total length with a 10-mm-wide gage section were used in all tests. Tensile tests were performed in stroke control with a constant displacement rate of 0.05 mm/s. Tensile tests were conducted at 23, 1200 and 1330 °C in laboratory air.

3.1. Monotonic tension Tensile stress–strain behavior at 23, 1200 and 1330 °C is shown in Fig. 2. The stress–strain curves obtained at 23 and 1200 °C are nearly linear to failure. Such linear behavior indicates that there is little additional matrix cracking 250

200

Stress (MPa)

Fig. 1. As-received material: (a) overview, optical microscope and (b) porous nature of the matrix is evident (SEM).

1200°C 23°C

150

1330°C

100

50

0 0.00

0.25

0.50

0.75

1.00

1.25

1.50

1.75

2.00

Strain (%) Fig. 2. Tensile stress–strain curves for NextelTM720/alumina ceramic composite at 23, 1200 and 1330 °C.

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

and that fiber-matrix debonding is insignificant. Material exhibits typical fiber-dominated composite behavior. Fiber fracture appears to be the dominant damage mode. At 23 °C, the ultimate tensile strength (UTS) was 169 MPa, elastic modulus, 60 GPa, and failure strain, 0.35%. At 1200 °C, the UTS, elastic modulus and failure strain were 192 MPa, 75 GPa and 0.38%, respectively. These results agree well with those reported by COI Ceramics [50]. The stress–strain behavior changes dramatically at 1330 °C. The stress–strain curve at 1330 °C is linear up to the proportional limit (74 MPa), where non-linear behavior sets in. At 1330 °C, the UTS, elastic modulus and failure strain were 120 MPa, 42 GPa and 1.7%, respectively. While the UTS and the elastic modulus are significantly lower than those at 1200 °C, failure strain increases almost tenfold. It is important to note that in all tension tests, as well as in all other tests reported herein, the failure occurred within the gage section of the extensometer. 3.2. Tension–tension fatigue Degradation of fatigue performance in high-temperature oxidizing environments remains among the key concerns that must be addressed before using CMCs in advanced aerospace applications. Therefore high-temperature fatigue tests, especially when conducted in steam environment are critical to assessing the durability of a given CMC. Tension–tension fatigue tests with a ratio, R of 0.05, were performed at 1200 and 1330 °C in air and in steam environments. Results are summarized in Table 1, where test temperature and environment are shown together with the maximum stress level and number of cycles to failure. Results are also presented in Fig. 3 as stress vs cycles to failure (S–N) curves for both temperatures and environments. At 1200 °C the in-air fatigue limit was 170 MPa (88% UTS at 1200 °C). This fatigue limit is based on a

250

200

Stress (MPa)

2032

Fatigue at 1330 °C Laboratory air Laboratory air Steam Steam a

Run-out.

100 125 150 170 100 125 150 170 50 100 50 100

Cycles to failure 120,199a 146,392a 167,473a 109,436a 100,780a 166,326a 11,782 202 97,282 1,519 25,852 347

1200°C, Steam

1330°C, Steam

150 UTS at 1330°C

100

0 1.E+00

1.E+01

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

Cycles (N) Fig. 3. Fatigue S–N curves for NextelTM720/alumina ceramic composite at 1200 and 1330 °C, in laboratory air and in steam environment.

run-out condition of 105 cycles, approximate number of loading cycles expected in aerospace applications at 1200 °C. It is believed that a more rigorous run-out condition would have resulted in a lower fatigue limit. Presence of steam (a highly oxidizing environment) causes noticeable degradation in fatigue performance. At 1200 °C, the insteam fatigue limit is only 125 MPa (65% UTS at 1200 °C). As seen in Fig. 3, increase in temperature from 1200 to 1330 °C results in significant degradation of the in-air fatigue performance. Even at the low fatigue stress level of 50 MPa (42% UTS at 1330 °C) the run-out was not achieved. As expected, steam environment even further degraded an already poor fatigue resistance. Of importance in cyclic fatigue is the reduction in stiffness (hysteresis modulus determined from the maximum and minimum stress–strain data points during a load cycle), reflecting the damage development during cycling. Change in modulus at 1200 °C is shown in Fig. 4, where normalized modulus (i.e. modulus normalized by the modulus obtained in the first cycle) is plotted vs fatigue cycles. It is noteworthy that although all in-air tests achieved

2.0

Normalized Modulus (E/Eo)

Fatigue at 1200 °C Laboratory air Laboratory air Laboratory air Laboratory air Steam Steam Steam Steam

Max stress (MPa)

1330°C, Air

50

Table 1 Summary of fatigue results for the N720/A composite at 1200 and 1330 °C, in laboratory air and steam environments Test environment

UTS at 1200°C

1200°C, Air

1.8

100 MPa, Air

100 MPa, Steam

1.6

125 MPa, Air

125 MPa, Steam

150 MPa, Air

150 MPa, Steam

170 MPa, Air

170 MPa, Steam

1.4

T = 1200°C f = 1 Hz R = 0.05

1.2 1.0 0.8 0.6 0.4 0.2 0.0 1.E+00

1.E+01

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

Cycles (N) Fig. 4. Normalized modulus vs fatigue cycles at 1200 °C in laboratory air and in steam environment.

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040 160

Stress (MPa)

120

Cycle 1025

Cycle 1

Cycle 10000

Cycle 103225

100

Cycle 2 80 60 40

Max Stress = 100 Mpa frequency = 1 Hz R = 0.05

20 0 0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

Strain (%) Fig. 6. Typical evolution of a stress–strain hysteresis loop with fatigue cycles.

presented in Fig. 7(a) and (b) for laboratory air and steam environments, respectively. It is seen that ratcheting takes place in all fatigue tests conducted in air at 1200 °C. Onset of ratcheting depends on the maximum fatigue stress. Earlier onset of ratcheting is observed in tests with higher maximum stress levels. In the 100 MPa test, there is little change in accumulated strain for up to 10,000 cycles, only

4.0 3.5 3.0

T = 1200°C Fatigue in Air

100 MPa 150 MPa 170 MPa

2.5 2.0 1.5 1.0 0.5 0.0 1.E+00

1.E+01

1.E+02

(a)

1.E+03

1.E+04

1.E+05

1.E+06

Cycles (N) 4.0

2.0 1.8

100 MPa, Air

1.6

100 MPa, Steam

1.4

50 MPa, Steam

T = 1330°C f = 1 Hz R = 0.05

3.5 3.0

Strain (%)

Normalized Modulus (E/Eo)

T = 1200°C Fatigue in Air

Cycle 125

140

Strain (%)

run-out, a decrease in normalized modulus with cycling was still observed. Modulus loss increased with increasing fatigue stress level. In air, normalized modulus was reduced by 5% in the 100 MPa test, 7% in the 125 MPa test, 8% in the 150 MPa test, and 17% in the 170 MPa test. Decrease in normalized modulus becomes more pronounced in steam environment. While in air the reduction in normalized modulus was limited to 17%, normalized modulus loss reached 30% in steam. In steam environment, normalized modulus loss in run-out fatigue tests was 10% in the 100 MPa test and 16% in the 125 MPa test. Normalized modulus in the 150 MPa test dropped by 17%, and in the 170 MPa test, by a significant 30% prior to failure at 202 cycles. Continuous decrease in modulus observed both in air and steam environments suggests progressive damage with continued cycling. Because the fatigue damage is still evolving at 105 cycles, the 105 fatigue limit does not meet the criteria of a true endurance fatigue limit proposed by Sorensen et al. [51] and may not be a true endurance fatigue limit. Normalized modulus evolution with cycling at 1330 °C is shown in Fig. 5. Modulus loss in the 100 MPa in-air test was 15%, noticeably greater than the 5% modulus loss in the corresponding 1200 °C test. This suggests accelerated damage growth, but may also be indicative of accelerated fiber degradation at higher temperature. Contrary to the expectations, presence of steam caused little additional modulus degradation. Modulus loss in steam was limited to 17%. Hysteresis loops for a 100 MPa test conducted in air at 1200 °C are presented in Fig. 6. Results in Fig. 6 are representative of the hysteresis loop evolution with cycling observed in all fatigue tests reported herein. It is seen that most extensive damage occurs on the first cycle. Afterwards hysteresis loops stabilize quickly. Results in Fig. 6 reveal that ratcheting, defined as progressive increase in accumulated strain with increasing number of cycles, continues throughout the test. Maximum and minimum cyclic strains as functions of cycle number for fatigue tests conducted at 1200 °C are

2033

1.2 1.0 0.8

T = 1200°C Fatigue in Steam

100 Mpa 125 MPa 150 MPa 170 MPa

2.5 2.0 1.5

0.6 1.0

0.4

0.5

0.2 0.0 1.E+00

1.E+01

1.E+02

1.E+03

1. E+04

1.E+05

0.0 1.E+00

1.E+06

Cycles (N) Fig. 5. Normalized modulus vs fatigue cycles at 1330 °C in laboratory air and in steam environment.

(b)

1.E+01

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

Cycles (N)

Fig. 7. Maximum and minimum strains as functions of cycle number at 1200 °C: (a) in laboratory air and (b) in steam environment.

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

4.0 3.5

Strain (%)

3.0

T = 1330°C

100 MPa, Air 100 MPa, Steam 50 MPa, Steam

2.5 2.0 1.5 1.0 0.5 0.0 1.E+00 1.E+01

1.E+02

1.E+03

1.E+04

1.E+05

1.E+06

Cycles (N) Fig. 8. Maximum and minimum strains as functions of cycle number at 1330 °C in laboratory air and in steam environment.

50

Hysteretic Energy Density (kJ/m3)

then does ratcheting begin. Conversely, in the 150 MPa test ratcheting begins after 1000 cycles, and in the 170 MPa test, after only 250 cycles. Earlier ratcheting is accompanied with higher strain accumulation. Maximum strains accumulated in the 100, 150 and 170 MPa tests were 0.6%, 1.7%, and 2.4%, respectively. As seen in Fig. 7(b), presence of steam accelerates ratcheting at 1200 °C. For a given maximum stress, specimens tested in steam exhibited a much earlier onset of ratcheting than those tested in air. In the 100 MPa test conducted in steam, ratcheting begins after 100 cycles, in the 150 MPa test, after 50 cycles, and in the 170 MPa test, after mere 10 cycles. Strain accumulations in the 100, 125, 150 and 170 MPa tests conducted in steam were 0.7%, 1.0%, 0.7%, and 0.8%, respectively. Strains accumulated in the 150 and 170 MPa tests in steam are considerably lower than those produced for the same fatigue stress levels in air. Generally, lower strain accumulation with cycling indicates that less damage has occurred, and that it is mostly limited to some additional matrix cracking. However, in the case of 150 and 170 MPa tests conducted in steam, low accumulated strains are more likely due to early bundle failures leading to specimen failure. In air environment, increase in test temperature results in accelerated ratcheting as well as in larger strain accumulations with cycling (see Fig. 8). In the 100 MPa test at 1330 °C, ratcheting begins immediately, strain is accumulated rapidly, reaching a significant 3.9% at failure. At 1330 °C, presence of steam results in earlier failure. Furthermore, in steam environment, shorter cyclic lives are accompanied with lower strain accumulations. The hysteresis energy density (HED) behavior is shown in Figs. 9 and 10 for 1200 and 1330 °C, respectively. The HED values at 1200 °C are fairly small, with the average of 10 kJ/m3. Most traditional composites with interfaces and classical fiber debonding typically produce HED values P80 kJ/m3 when fatigued above the proportional limit. It is seen that for each stress level tested at 1200 °C, the HED exhibits a significant decrease within the first 10 cycles. From this cycle number on there appears to be only

45 40 35 30 25

T = 1200°C f = 1 Hz R = 0.05

100 MPa, Air 150 MPa, Air 170 MPa, Air 100 MPa, Steam 125 MPa, Steam 150 MPa, Steam

20 15 10 5 0 1.E+00

1.E+01

1.E+02

1.E+03

1.E+04

1. E+05

1.E+06

Cycles (N) Fig. 9. Hysteresis energy density (HED) vs fatigue cycles at 1200 °C in laboratory air and in steam environment.

100

Hysteretic Energy Density (kJ/m3)

2034

100 MPa, Ai r 50 MPa, Steam 100 MPa, Steam

90 80

T = 1330°C f = 1 Hz R = 0.05

70 60 50 40 30 20 10 0 1.E+00

1. E+01

1.E+02

1. E+03

1.E+04

1. E+05

1.E+06

Cycles (N) Fig. 10. Hysteresis energy density (HED) vs fatigue cycles at 1330 °C in laboratory air and in steam environment.

a slight decrease in HED with continued cycling. However, upon closer examination the 1200 °C data reveals that at 104 cycles the HED becomes stable for all tests except the 150 MPa test conducted in steam. Among tests represented in Fig. 9, only the 150 MPa test in steam did not achieve run-out. In conventional composites, a decrease in HED with fatigue cycling is generally attributed to degradation of interfacial shear resistance at the fiber matrix interface. For brittle matrix composites, it was also observed [32] that continuous damage development, such as matrix cracking and fiber/matrix debonding, in a cyclically loaded specimen may have a significant effect on the HED behavior. The HED behavior at 1330 °C is qualitatively similar to that observed at 1200 °C. However, average HED values obtained in 100 MPa tests were higher at 1330 °C than at 1200 °C. The presence of steam appears to have little effect on the HED behavior at both temperatures investigated. Retained strength and stiffness of the fatigue specimens, which achieved fatigue run-out, are summarized in Table 2. Evaluation of retained properties is useful in assessing the damage state of the composite subjected to prior loading.

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

2035

Table 2 Retained properties of the N720/A specimens subjected to prior fatigue in laboratory air and in steam environment at 1200 °C Fatigue stress (MPa)

Retained strength (MPa)

Prior fatigue in laboratory air 100 194 125 199 150 199 170 192 Prior fatigue in steam environment 100 174 125 168

Strength retention (%)

Retained modulus (GPa)

Modulus retention (%)

Strain at failure (%)

P100 P100 P100 P100

56.6 54.9 43.4 40.7

74 73 72 67

0.44 0.44 0.53 0.51

90 88

47.6 52.0

84 80

0.40 0.43

It is seen that specimens tested in air exhibited no loss of tensile strength, irrespective of the fatigue stress level. However, considerable stiffness loss (28–33%) was observed. Stiffness degradation increases with increasing prior fatigue stress level. Full retention of tensile strength suggests that no fatigue damage occurred to the fibers. The reduction in stiffness is most likely due to additional matrix cracking. Conversely, prior fatigue in steam caused reduction of both strength and stiffness. Strength loss in steam was limited to 12% and stiffness loss, to 20%. In this case, the loss of strength may be associated with the environmental degradation of the fibers, while both fiber degradation and progressive matrix cracking may account for the loss of stiffness. The discrepancy between the retained modulus of a run-out specimen and the decrease in hysteresis modulus observed during fatigue testing most likely stems from different methods used to determine the retained and hysteresis moduli. Results in Table 2 demonstrate that fatigue in air did not cause reduction in strength. However, prior fatigue in steam environment resulted in noticeable strength loss, which cannot be neglected. Fiber degradation represents a possible source of composite degradation. NextelTM720 fibers consist of alumina grains with an approximate diameter of 0.1 lm distributed among larger (0.5 lm) mullite grains, consisting of many smaller subgrains [52]. Reported observations of the response of the fibers to thermal exposure are somewhat conflicting. Deleglise et al. [52] observed significant degradation only above 1400 °C for 5 h exposure times, while Milz et al. [53] reported severe degradation after 2 h at 1300 °C. Petry and Mah [54] report a small reduction in strength after 2 h at 1100 °C. The causes of degradation are not well understood; surface grooving, structural coarsening [54] and local impurity enrichment have been suggested [53]. Evidence in literature also suggests that, under specific conditions, reactions between the SiO2 in the fiber and moisture may occur. Wannaparhun et al. [55] showed that, at 1100 °C in water–vapor environment, SiO2 could be leached out of the NextelTM720 fiber. Exposure to water vapor resulted in the formation of volatile Si(OH)4 and was responsible for the loss of the mullite phase in the fiber. Furthermore, formation of volatile Si(OH)4 also resulted in surface recondensation of these silicon species with the Al2O3 matrix at the specimen surface, in turn causing an increase in aluminosilicate content at the surface of the

N720/A CMC (the same material as used in this study). Campbell et al. [56] exposed N720/A to a water–vapor environment for 1000 h at 1200 °C. Strength loss of 15% was observed after exposure. In the present study, a high fatigue limit (88% UTS) and 100% strength retention are observed for specimens tested in air. Presence of steam noticeably degrades fatigue performance of the material. In steam environment, fatigue limit is significantly lower (65% UTS) and strength retention is limited to 90%. It is noteworthy that strength losses similar to those observed by Campbell et al. [56] after 1000 h of noload exposure are seen after only 28 h (105 cycles at a frequency of 1 Hz) of fatigue cycling in steam at 1200 °C. The strength loss is strongly influenced by the loading conditions. In a given high-temperature environment, strength degradation increases with increasing fatigue load. 3.3. Creep rupture Results of the creep-rupture tests are summarized in Table 3, where test temperature and environment are shown together with the creep stress level and time to rupture. It is noteworthy that all specimens failed within the extensometer gage section. Creep curves obtained at 1200 and 1330 °C are presented in Figs. 11 and 12, respectively. Time scale in Figs. 11 and 12 is reduced in order to clearly show creep curves Table 3 Summary of creep-rupture results for the N720/A composite at 1200 and 1330 °C, in laboratory air and in steam environments Environment

Creep stress (MPa)

Time to rupture (s)

Creep at 1200 °C Air Air Air Air Steam Steam Steam Steam

80 100 125 154 80 100 125 154

917,573 147,597 15,295 968 165,777 8,966 869 98

Creep at 1330 °C Air Air Steam Steam

50 100 50 100

313,198 4,244 11,088 40

2036

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

Fig. 11. Creep strain vs time for NextelTM720/alumina ceramic composite at 1200 °C in laboratory air and in steam environment: (a) at 80 and 100 MPa, (b) at 125 and 154 MPa.

7.0

T = 1330°C

6.0

Strain (%)

5.0 4.0 3.0

100 MPa, Steam

100 MPa, Air

2.0

50 MPa, Steam

1.0

50 MPa, Air 0.0 0

5 00

1000

1500

2000

Time (s) Fig. 12. Creep strain vs time for NextelTM720/alumina ceramic composite at 1330 °C in laboratory air and in steam environment.

produced at higher stress levels. It is seen that all creep curves generated at 1200 °C in air environment exhibit primary and secondary creep regimes. Transition from primary to secondary creep occurs during the first 10% of creep life. Secondary creep appears to be nearly linear to failure. As the creep stress level increases from 80 MPa to

125 MPa, creep strain accumulation increases from 1% to 3.4%. However, for the creep stress of 154 MPa, creep strain accumulation reaches only about 0.6%. It is noteworthy that in all creep tests, the accumulated creep strain significantly exceeded failure strain obtained in tension test. It is important to recognize that the total strain incurred in a creep-rupture test represents a sum of two contributions: (1) that due to the initial loading up to the specific creep stress level, and (2) that accumulated during the actual creep period. For specimens tested at creep stresses 6125 MPa, close to 90% of total strain was accumulated during the creep period. However, for specimen tested at creep stress of 154 MPa, creep strain accounted only for 33% of the total strain. The creep curves produced in air environment at 1330 °C are qualitatively similar to those obtained at 1200 °C. Creep strain accumulation decreases from 5% to approximately 4% as the creep stress increases from 50 to 100 MPa. In both creep tests, close to 100% of the total strain was incurred during the creep period. Creep strain accounts for 98% of the total strain in the 50 MPa test, and for 94% of the total strain in the 100 MPa test. Results in Figs. 11 and 12 demonstrate that specimens tested in steam environment produced creep curves that are qualitatively similar to those produced in air. As expected, strains incurred during the initial loading to a given creep stress were not affected by the presence of steam. Conversely, creep strains accumulated in steam environment were significantly different from those accumulated in air. For both test temperatures and creep stress levels <100 MPa, specimens tested in steam accumulated more creep strain than specimens tested in air. At 1200 °C, creep strain produced in the 80 MPa test conducted in steam was 67% higher than that produced at the same creep stress in air. At 1330 °C and creep stress of 50 MPa, creep strain accumulated in steam was 20% higher than that accumulated in air. For creep stress levels P100 MPa, presence of steam resulted in lower creep strains and much lower creep lifetimes at both test temperatures. At 1200 °C, creep strains produced in steam environment for creep stress levels of 100, 125 and 154 MPa were, respectively, 53%, 73% and 92% lower than those produced at the same stress levels in air. At 1330 °C and creep stress of 100 MPa, creep strain accumulated in steam was 60% lower than that accumulated in air. Minimum creep rate was reached in all tests. Creep rate as a function of applied stress is presented in Fig. 13, where results of the present investigation are plotted together with the data from Wilson and Visser [57] for Nextel 720 fibers. To further facilitate comparison between the creep properties of the fibers and the composite, the Nextel 720 fiber data adjusted for Vf = 0.22 (volume fraction of the on-axis fibers in the N720/A composite) is also shown. For both temperatures, the minimum creep rates increase with increasing applied stress. At 1200 °C, as creep stress increases from 80 to 154 MPa creep rate increases by two orders of magnitude. Creep rates of the composite obtained

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040 1.E-02

Creep Rate (s-1)

1.E-03 1.E-04

1200°C, Air 1200°C, Steam 1330°C, Air 1330°C, Steam

N720 Fiber at Vf = 0.22 1200° C

1.E-05 1.E-06 1.E-07 N720 Fiber, 1200°C Wilson, 2001

1.E-08 1.E-09

10

100 Creep Stress (MPa)

1000

Fig. 13. Minimum creep rate as a function of applied stress at 1200 and 1330 °C in laboratory air and in steam environment. Data for Nextel 720 fibers (Wilson [57]) are also shown.

in air environment were close to the N720 fiber data adjusted for Vf = 0.22. For a given creep stress, creep rates in steam were approximately an order of magnitude higher than those in air. Fitting the experimental results (obtained in either air or steam) with a temperature-independent Norton–Bailey equation of the form e_ ¼ Arn yields the stress exponents that are considerably higher than that reported for the N720 fibers alone [57]. It is believed that the higher stress exponents are due to a contribution from the matrix. As expected, creep resistance decreases dramatically with increasing temperature. For the creep stress of 100 MPa, creep rate at 1330 °C was almost two orders of magnitude higher that at 1200 °C. The presence of steam further accelerates creep and degrades creep resistance. At 1330 °C, the presence of steam increases the minimum creep rate by a factor of 100. Stress-rupture behavior is summarized in Fig. 14, where creep stress is plotted vs time to rupture at 1200 and 1330 °C in air and in steam environments. At both temperatures, creep life decreases with increasing applied stress. For a 200 180

UTS at 1200°C

1200°C, Air 1200°C, Steam 1330°C, Air 1330°C, Steam

Stress (MPa)

160 140 120

UTS at 1330°C

100 80 60 40 20 0 1.E+00 1.E+01 1.E+02 1.E+03 1.E+04 1.E+05 1.E+06 1.E+07

Time (s) Fig. 14. Creep stress vs time to rupture at 1200 and 1330 °C in laboratory air and in steam environment.

2037

given creep stress, creep-rupture life at 1330 °C is considerably reduced compared to that at 1200 °C. With the creep run-out condition defined as 100 h, 80 MPa was the runout stress at 1200 °C. At 1330 °C the run-out was not achieved. Even at a low creep stress of 50 MPa the rupture time was 87 h < 100 h. The short creep lives produced at 1330 °C indicate that this oxide/oxide CMC should not be used under sustained loading at temperatures above 1200 °C. Presence of steam dramatically reduced creep lives at both test temperatures. At 1200 °C, reduction in creep life due to steam was at least 90% for applied stress levels P100 MPa, and 82% for the applied stress of 80 MPa. At 1330 °C, presence of steam reduced creep lives by 96–98%. It is recognized that NextelTM720 fiber has the best creep performance of any commercially available polycrystalline oxide fiber. The superior high-temperature creep performance of the NextelTM720 fibers results from the high content of mullite, which has a much better creep resistance than alumina [57]. Wannaparhum et al. [55] reported that exposure of the N720/A composite to water vapor at 1100 °C could result in an increase in Al2O3 content of the fiber because of the loss of SiO2 from its mullite phase. The loss of the mullite phase in the fiber may be the mechanism behind the higher creep rates and reduced creep resistance observed in steam. However, further experiments quantifying the mullite loss from the fiber exposed to water vapor at high temperature would be required to ascertain this. 3.4. Composite microstructure Fracture surface of a specimen tested in creep is shown in Fig. 15. It should be noted that the appearance of the fracture surface was not significantly affected by any of the following factors: test temperature (1200 vs 1330 °C), test environment (air vs steam), test type (tension vs creep vs tension–tension fatigue). Fracture surfaces of similar appearance were produced in all tests. Micrographs in Fig. 15 are typical and representative of fracture surfaces obtained in all tests in this study. It is seen in Fig. 15(a) that the fracture plane is not well defined. The fibers in the 0° tows in each cloth layer exhibit random failure producing fiber ‘‘pull-out’’. Fig. 15(b) shows that the 0° fiber tows break over a wide range of axial locations, in general spanning the entire width of the specimen. The locations of the fiber breaks within an individual tow also exhibit a broad distribution, typically 1 mm in length, as seen in Fig. 15(c). It is important to note that no matrix holes were observed on the fracture surface. In conventional CFCCs with ‘‘weak’’ interfaces, the fiber pull-out results in formation of matrix holes, where broken fibers slide out of the matrix. However, in the present composite, the pull-out of the fibers does not leave matrix sockets but causes fragmentation of intervening matrix in the region of strain localization. Some of the matrix debris and matrix still bonded to the fibers are seen in Fig. 15(d).

2038

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

Fig. 15. Fracture surface of a typical NextelTM720/alumina CMC specimen tested in creep at 1200 °C: (a) overall view showing general extent of fiber pullout, (b) fiber pullout across width of specimen, (c) fiber pullout within a 0° bundle, (d) matrix particles bonded to the fiber, (e) region of fairly coordinated fiber fracture in the 0° tow and (f) nearly planar fracture of the 90° fiber tows.

Fig. 15(e) shows that some fracture regions in the 0° tows exhibit flatter, more coordinated fracture topography. Close examination reveals that most of the fibers fracture on different planes, suggesting that a single crack front did not cause this fracture topography. However, one can see a few pairs of fibers (arrows) which exhibit planar fracture. These fiber pairs appear to have a common fracture origin where they touch. Haslam et al. [58] suggest that such common fracture origin is produced during fiber processing, when adjacent fibers in the bundle stick to each other and sinter together along their cylindrical axis. A typical fracture of the 90° fiber tows in a cloth layer is shown in Fig. 15(f). Here, the fracture surface topography can be characterized as nearly planar. 4. Concluding remarks The tension–tension fatigue behavior and the creep-rupture behavior of the NextelTM720/Alumina continuous

fiber ceramic composite were characterized in laboratory air and in 100% steam environment at 1200 and 1330 °C. 4.1. Fatigue behavior Tension–tension fatigue behavior of the N720/A CFCC was studied for fatigue stress levels of 100–170 MPa at 1200 °C, and for fatigue stress levels of 50 and 100 MPa at 1330 °C. Results suggest the following conclusions: (1) The N720/A composite exhibits excellent fatigue resistance in laboratory air at 1200 °C. The fatigue limit in air (based on a run-out condition of 105 cycles) is 170 MPa (88% UTS at 1200 °C). The material retains 100% of its tensile strength. However, considerable stiffness loss (30–50%) is observed. (2) Presence of steam causes noticeable degradation in fatigue performance at 1200 °C. The fatigue limit in steam environment is 125 MPa (65% UTS at

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

1200 °C). The material retains 90% of its tensile strength. Stiffness loss is limited to 37%. (3) Fatigue resistance at 1330 °C is poor, run-out was not achieved. (4) Ratcheting is observed in all fatigue tests. Presence of steam accelerates ratcheting at 1200 and 1330 °C.

4.2. Creep-rupture behavior Tensile creep behavior of the N720/A CMC was studied for creep stresses ranging from 80 to 154 MPa at 1200 °C, and for creep stresses of 50 and 100 MPa at 1330 °C. Results suggest the following conclusions: (1) The creep curves produced at 1200 and 1330 °C in laboratory air and in steam environment exhibit very short primary creep, which rapidly transitions into secondary creep. Secondary creep remains nearly linear to failure. In all creep tests, the accumulated creep strain significantly exceeds failure strain obtained in tension test. (2) Minimum creep rate was reached in all tests. Creep rates ranged from 108 to 105 s1 at 1200 °C, and from 107 to 105 s1at 1330 °C. At 1200 °C in air environment, creep rates of the composite are close to what may be expected from Nextel 720 fibers alone. The relationship between minimum creep rate and applied stress can be represented by a power law. Due to the contribution from the matrix, the stress exponent of the composite is higher than that reported for the fibers alone. Presence of steam significantly accelerates creep rates at both temperatures. (3) Creep-rupture lives ranged from 0.3 h (154 MPa test) to 255 h (80 MPa test) at 1200 °C, and from 1.2 h (100 MPa test) to 87 h (50 MPa test) at 1330 °C. Presence of steam dramatically reduces creep life. Reductions in creep life due to steam were 82–90% at 1200 °C, and 96–98% at 1330 °C.

References [1] Verrilli MJ, Opila EJ, Calomino A, Kiser JD. Effect of environment on the stress-rupture behavior of a carbon-fiber-reinforced silicon carbide ceramic matrix composite. J Am Ceram Soc 2004;87(8): 1536–42. [2] Zawada LP, Staehler J, Steel S. Consequence of intermittent exposure to moisture and salt fog on the high-temperature fatigue durability of several ceramic–matrix composites. J Am Ceram Soc 2003;86(8): 1282–91. [3] Prewo KM, Batt JA. The oxidative stability of carbon fibre reinforced glass-matrix composites. J Mater Sci 1988;23:523–7. [4] Mah T, Hecht NL, McCullum DE, Hoenigman JR, Kim HM, Katz AP, et al.. Thermal stability of SiC fibres (nicalon). J Mater Sci 1984;19:1191–201. [5] More KL, Tortorelli PF, Ferber MK, Keiser JR. Observations of accelerated silicon carbide recession by oxidation at high water–vapor pressures. J Am Ceram Soc 2000;83(1):211–3.

2039

[6] More KL, Tortorelli PF, Ferber MK, Walker LR, Keiser JR, Brentnall WD, et al. Exposure of ceramic and ceramic–matrix composites in simulated and actual combustor environments. In: Proceedings of international gas turbine and aerospace congress; 1999. Paper No. 99-GT-292. [7] Ferber MK, Lin HT, Keiser JR. Oxidation behavior of non-oxide ceramics in a high-pressure, high-temperature steam environment. In: Jenkins MG, Lara-Curzio E, Gonczy ST, editors. Mechanical, thermal, and environmental testing and performance of ceramic composites and components. ASTM STP 1392. American Society for Testing and Materials; 2000. p. 210–5. [8] Haynes JA, Lance MJ, Cooley KM, Ferber MK, Lowden RA, Stinton DP. CVD mullite coatings in high-temperature, high-pressure air–H2O. J Am Ceram Soc 2000;83(3):657–9. [9] Opila EJ, Hann Jr RE. Paralinear oxidation of SiC in water vapor. J Am Ceram Soc 1997;80(1):197–205. [10] Opila EJ. Oxidation kinetics of chemically vapor deposited silicon carbide in wet oxygen. J Am Ceram Soc 1994;77(3):730–6. [11] Opila EJ. Variation of the oxidation rate of silicon carbide with water vapor pressure. J Am Ceram Soc 1999;82(3):625–36. [12] Jacobson NS. Corrosion of silicon-based ceramics in combustion environments. J Am Ceram Soc 1993;76(1):3–28. [13] Thomas-Ogbuji L. A pervasive model of oxidation degradation in a SiC–SiC composite. J Am Ceram Soc 1998;81(11):2777–84. [14] More KL, Tortorelli PF, Lin HT, Lara-Curzio E, Lowden RA. Degradation mechanisms of BN interfaces in SiC/SiC composites in oxygen- and water-containing environments. Electrochem Soc Proc 1998;98–99:382–95. [15] Cofer CG, Economy J. Oxidative and hydrolytic stability of boron nitride—a new approach to improving the oxidation resistance of carbonaceous structures. Carbon 1989;33(4):389–95. [16] Matsuda T. Stability to moisture for chemically vapor deposited boron nitride. J Mater Sci 1989;24:2353–7. [17] Morscher GN, Bryant D, Tressler RE. Environmental durability of different BN interphases (for SiC/SiC) in H2O-containing atmospheres at intermediate temperatures. Ceram Eng Sci Proc 1997;18(3):525–33. [18] Thomas-Ogbuji L. Degradation SiC/BN/SiC composite in the burner rig. Ceram Eng Sci Proc 1998;19(4):257–64. [19] Jacobson N, Farmer S, Moore A, Sayir H. High-temperature oxidation of boron nitride: I: monolithic boron nitride. J Am Ceram Soc 1999;82(2):939–98. [20] Jacobson NS, Morscher GN, Bryant DR, Tressler RE. Hightemperature oxidation of boron nitride: II. boron nitride layers in composites. J Am Ceram Soc 1999;82(6):1473–82. [21] Hermes EE, Kerans RJ. Degradation of non-oxide reinforcement and oxide matrix composites. Mater Res Soc, Symp Proc 1988;125:73–8. [22] Szweda A, Millard ML, Harrison MG. Fiber-reinforced ceramic– matrix composite member and method for making, US Pat. No. 5 601 674; 1997. [23] Sim SM, Kerans RJ. Slurry infiltration and 3-D woven composites. Ceram Eng Sci Proc 1992;13(9–10):632–41. [24] Moore EH, Mah T, Keller KA. 3D composite fabrication through matrix slurry pressure infiltration. Ceram Eng Sci Proc 1994;15(4):113–20. [25] Lewis MH, Cain MG, Doleman P, Razzell AG, Gent J. Development of interfaces in oxide and silicate matrix composites. In: Evans AG, Naslain RG, editors. High-temperature ceramic–matrix composites II: manufacturing and materials development. American Ceramic Society; 1995. p. 41–52. [26] Lange FF, Tu WC, Evans AG. Processing of damage-tolerant, oxidation-resistant ceramic matrix composites by a precursor infiltration and pyrolysis method. Mater Sci Eng A 1995;A195: 145–50. [27] Lunderberg R, Eckerbom L. Design and processing of all-oxide composites. In: Evans AG, Naslain RG, editors. High-temperature ceramic–matrix composites II: manufacturing and materials development. American Ceramic Society; 1995. p. 95–104.

2040

M.B. Ruggles-Wrenn et al. / Composites: Part A 37 (2006) 2029–2040

[28] Mouchon E, Colomban P. Oxide ceramic matrix/oxide fiber woven fabric composites exhibiting dissipative fracture behavior. Composites 1995;26:175–82. [29] Morgan PED, Marshall DB. Ceramic composites of monazite and alumina. J Am Ceram Soc 1995;78(6):1553–63. [30] Tu WC, Lange FF, Evans AG. Concept for a damage-tolerant ceramic composite with strong interfaces. J Am Ceram Soc 1996;79(2):417–24. [31] Kerans RJ, Hay RS, Pagano NJ, Parthasarathy TA. The role of the fiber-matrix interface in ceramic composites. Am Ceram Soc Bull 1989;68(2):429–42. [32] Evans AG, Zok FW. Review: the physics and mechanics of fiberreinforced brittle matrix composites. J Mater Sci 1994;29:3857–96. [33] Kerans RJ, Parthasarathy TA. Crack deflection in ceramic composites and fiber coating design criteria. Composites: Part A 1999;30:521–4. [34] Levi CG, Yang JY, Dalgleish BJ, Zok FW, Evans AG. Processing and performance of an all-oxide ceramic composite. J Am Ceram Soc 1998;81:2077–86. [35] Davis JB, Lofvander JPA, Evans AG. Fiber coating concepts for brittle matrix composites. J Am Ceram Soc 1993;76(5):1249–57. [36] Mackin TJ, Yang JY, Levi CG, Evans AG. Environmentally compatible double coating concepts for sapphire fiber reinforced cTiAl. Mater Sci Eng A 1993;A161:285–93. [37] Hegedus AG. Ceramic bodies of controlled porosity and process for making same. US Pat. No. 5 0177 522, May 21; 1991. [38] Dunyak TJ, Chang DR, Millard ML. Thermal aging effects on oxide/ oxide ceramic–matrix composites. In: Proceedings of 17th conference on metal matrix, carbon, and ceramic matrix composites. NASA conference publication 3235, Part 2; 1993. p. 675–90. [39] Zawada LP, Lee SS. Mechanical behavior of CMCs for flaps and seals. In: Coblenz WS, editor. ARPA ceramic technology insertion program (DARPA). Defence advanced projects agency conference publication. Annapolis MD; 1994. p. 267–322. [40] Zawada LP, Lee SS. Evaluation of the fatigue performance of five CMCs for aerospace applications. In: Proceedings of the sixth international fatigue congress; 1996. p. 1669–74. [41] Lu TJ. Crack branching in all-oxide ceramic composites. J Am Ceram Soc 1996;79(1):266–74. [42] Zok FW, Levi CG. Mechanical properties of porous-matrix ceramic composites. Adv Eng Mater 2001;3(1–2):15–23. [43] Steel SG. Monotonic and fatigue loading behavior of an oxide/oxide ceramic matrix composite. MS Thesis. Air force institute of technology, WPAFB OH; 2000.

[44] Steel S, Zawada LP, Mall S. Fatigue behavior of a Nextel 720/ alumina (N720/A) composite at room and elevated temperature. Ceram Eng Sci Proc 2001;22(3):695–702. [45] Wu X, Holmes JW. Tensile creep and creep–strain recovery behavior of silicon carbide/calcium aluminosilicate matrix ceramic composites. J Am Ceram Soc 1993;76(10):2695–700. [46] Evans AG, Weber CH. Creep damage in SiC/SiC composites. Mater Sci Eng A 1996;A208:1–6. [47] Weber CH, Lofvander JPA, Evans AG. Creep anisotropy of a continuous-fiber-reinforced silicon carbide/calcium aluminosilicate composite. J Am Ceram Soc 1994;77(7):1745–52. [48] Weber CH, Kim KT, Heredia FE, Evans AG. High temperature deformation and rupture in SiC–C composites. Mater Sci Eng A 1995;196(1–2):25–31. [49] Jurf RA, Butner SC. Advances in oxide–oxide CMC. In: Proceedings of the 44th ASME gas turbine and aeroengine congress and exhibition. June 7–10; 1999. Indianapolis, Indiana. [50] COI Ceramics. Unpublished Data. [51] Sorensen BF, Holmes JW, Vanswijgenhoven EL. Does a true fatigue limit exist for continuous fiber-reinforced ceramic matrix composites?. J Am Ceram Soc 2002;85(2):359–65. [52] Deleglise F, Berger MH, Jeulin D, Bunsell AR. Microstructural stability and room temperature mechanical properties of the Nextel 720 Fibre. J Eur Ceram Soc 2001;21:569–80. [53] Milz C, Goering J, Schneider H. Mechanical and microstructural properties of NextelTM720 relating to its suitability for high temperature applications in CMCs. Ceram Eng Sci Proc 1999;21(9–10): 191–8. [54] Petry MD, Mah TI. The effect of thermal exposures on the strength of NextelTM 550 and 720 filaments. J Am Ceram Soc 1999;82(7):2801–7. [55] Wannaparhun S, Seal S, Desai VH, Varghese P, Campbell CX. A combined spectroscopic and thermodynamic investigation of Nextel720/alumina ceramic matrix composite in air and water vapor at 1100 °C. J Am Ceram Soc 2003;86(9):1628–30. [56] Campbell CX, Carelli EV, More KL, Varghese P, Seal S, Desai VH. Effect of high-temperature water vapor exposure on Nextel 720 in an alumina-matrix CMC. Siemens Westinghouse Power Corporation Technical Document TP-02076. Siemens Westinghouse Library in Orlando, Mail Code 00707-235. [57] Wilson DM, Visser LR. High performance oxide fibers for metal and ceramic composites. Composites: Part A 2001;32:1143–53. [58] Haslam JJ, Berroth KE, Lange FF. Processing and properties of an all-oxide composite with a porous matrix. J Eur Ceram Soc 2000;20:607–18.