MATERIALS HIENGE& ENGIEERING ELSEVIER Materials Science and Engineering A 454-455(2007)590-601 www.elsevier.com/locate/msea Compressive creep behavior of an oxide-oxide ceramic composite with monazite fiber coating at elevated temperatures P.R Jackson a M.B. Ruggles-Wrenn a, *. S.S. Baek D K.A. Keller C, Department of Aeronautics and Astronautics, Air Force Institute of Technology, Wright-Patterson Air Force Base, OH 45433-7765 USA UES, Inc, 4401 Dayton Xenia Road, Dayton, Oh, oath Korea Received 26 June 2006: received in revised form 13 November 2006: accepted 23 November 2006 The compressive creep behavior of a N610/LaPO/Al2O3 composite was investigated at 900 and 1100C. The composite consists of a poro alumina matrix reinforced with Nextel M610 fibers coated with monazite in a symmetric cross-ply (0/90 /0/90% ),orientation. The compressive tress-strain behavior was investigated and the compressive properties were measured. Compressive creep behavior was examined for creep stresses in the -50to-95 MParange. Minimum creep rate was reached in all tests At 900C, both monazite-containing and control(N61O/Al2O3)specimens produced creep strains<-0.05%. At 1100C, compressive creep strains approached-9%, and compressive creep strain rates, -10-7s-I Creep run-out defined as 100 h at creep stress, was achieved in all tests. Composite microstructure, damage and failure mechanisms, as well as effects of variation in microstructure on mechanical response were examined. Differences in processing and consequently in the composite microstructure had a significant effect on compressive response of the ceramic-matrix composite(CMC) o 2006 Elsevier B. v. All rights reserved Keywords: Ceramic-matrix composites( CMCs): Oxides; Fibres; Coatings; Creep: High-temperature properties; Mechanical properties; Fractography 1. Introduction the CMCs into aerospace turbine engine applications, such as ombustor walls [3-5]. Because these applications require expo- Advances in aerospace propulsion technologies have raised sure to oxidizing environments, the thermodynamic stability and the demand for structural materials that have superior long-term oxidation resistance of CMCs are vital issues. mechanical properties and retained properties under high tem- The main advantage of CMCs over monolithic ceramics perature, high pressure and varying environmental factors, such their superior toughness, tolerance to the presence of cracks as moisture [1]. Ceramic-matrix composites(CMCs), capable and defects, and non-catastrophic mode of failure. It is widely of maintaining excellent strength and fracture toughness at high accepted that in order to avoid brittle fracture behavior in CMCs temperatures are prime candidate materials for such aerospace and improve the damage tolerance, a weak fiber/matrix inter- applications. Additionally, the lower densities of CMCs and their face is needed, which serves to deflect matrix cracks and to higher use temperatures, together with a reduced need for cool- allow subsequent fiber pullout [6-9. Historically, following ing air, allow for improved high-temperature performance when the development of SiC fibers, fiber coatings such as C or BN compared to conventional nickel-based superalloys[2]. Con- have been employed to promote the desired composite behav current efforts in optimization of the CMCs and in design of the ior. However, the non-oxide fiber/non-oxide matrix composites combustion chamber are expected to accelerate the insertion of generally show poor oxidation resistance [10, 11], particularly at intermediate temperatures(800C). These systems are sus- ceptible to embrittlement due to oxygen entering through the The views expressed are those of the authors and do not reflect the official matrix cracks and then reacting with the interphase and the policy or position of the United States Air Force, Department of Defense or the U.S. Government fibers [12-15]. The degradation, which involves oxidation of Corresponding author. Tel: +1937 255 3636x4641: fax: +1 9376567621. fibers and fiber coatings, is typically accelerated by the pres- E-mail address: marina. ruggles-wrenn@ afit. edu(M B. Ruggles-w ence of moisture [16-22]. Using oxide fiber/non-oxide matrix I Under USAF Contract # F33615-01-C-5214 or non-oxide fiber/oxide matrix composites generally does not 0921-5093/S-see front matter 2006 Elsevier B v. All rights reserved doi:10.1016/1msea.2006.11.131
Materials Science and Engineering A 454–455 (2007) 590–601 Compressive creep behavior of an oxide–oxide ceramic composite with monazite fiber coating at elevated temperatures P.R. Jackson a, M.B. Ruggles-Wrenn a,∗, S.S. Baek b, K.A. Keller c,1 a Department of Aeronautics and Astronautics, Air Force Institute of Technology, Wright-Patterson Air Force Base, OH 45433-7765, USA b Agency for Defense Development, Daejeon, South Korea c UES, Inc., 4401 Dayton Xenia Road, Dayton, OH 45433-7817, USA Received 26 June 2006; received in revised form 13 November 2006; accepted 23 November 2006 Abstract The compressive creep behavior of a N610/LaPO4/Al2O3 composite was investigated at 900 and 1100 ◦C. The composite consists of a porous alumina matrix reinforced with NextelTM610 fibers coated with monazite in a symmetric cross-ply (0◦/90◦/0◦/90◦)s orientation. The compressive stress–strain behavior was investigated and the compressive properties were measured. Compressive creep behavior was examined for creep stresses in the−50 to−95 MPa range. Minimum creep rate was reached in all tests. At 900 ◦C, both monazite-containing and control (N610/Al2O3) specimens produced creep strains ≤ −0.05%. At 1100 ◦C, compressive creep strains approached −9%, and compressive creep strain rates, −10−7 s−1. Creep run-out defined as 100 h at creep stress, was achieved in all tests. Composite microstructure, damage and failure mechanisms, as well as effects of variation in microstructure on mechanical response were examined. Differences in processing and consequently in the composite microstructure had a significant effect on compressive response of the ceramic–matrix composite (CMC). © 2006 Elsevier B.V. All rights reserved. Keywords: Ceramic–matrix composites (CMCs); Oxides; Fibres; Coatings; Creep; High-temperature properties; Mechanical properties; Fractography 1. Introduction Advances in aerospace propulsion technologies have raised the demand for structural materials that have superior long-term mechanical properties and retained properties under high temperature, high pressure and varying environmental factors, such as moisture [1]. Ceramic–matrix composites (CMCs), capable of maintaining excellent strength and fracture toughness at high temperatures are prime candidate materials for such aerospace applications. Additionally, the lower densities of CMCs and their higher use temperatures, together with a reduced need for cooling air, allow for improved high-temperature performance when compared to conventional nickel-based superalloys [2]. Concurrent efforts in optimization of the CMCs and in design of the combustion chamber are expected to accelerate the insertion of 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.: +1 937 255 3636x4641; fax: +1 937 656 7621. E-mail address: marina.ruggles-wrenn@afit.edu (M.B. Ruggles-Wrenn). 1 Under USAF Contract # F33615-01-C-5214. the CMCs into aerospace turbine engine applications, such as combustor walls[3–5]. Because these applications require exposure to oxidizing environments, the thermodynamic stability and oxidation resistance of CMCs are vital issues. 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 widely accepted that in order to avoid brittle fracture behavior in CMCs and improve the damage tolerance, a weak fiber/matrix interface is needed, which serves to deflect matrix cracks and to allow subsequent fiber pullout [6–9]. Historically, following the development of SiC fibers, fiber coatings such as C or BN have been employed to promote the desired composite behavior. However, the non-oxide fiber/non-oxide matrix composites generally show poor oxidation resistance [10,11], particularly at intermediate temperatures (∼800 ◦C). These systems are susceptible to embrittlement due to oxygen entering through the matrix cracks and then reacting with the interphase and the fibers [12–15]. The degradation, which involves oxidation of fibers and fiber coatings, is typically accelerated by the presence of moisture [16–22]. Using oxide fiber/non-oxide matrix or non-oxide fiber/oxide matrix composites generally does not 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.11.131
PR Jackson et al. / Materials Science and Engineering A 454-455(2007)590-601 bstantially improve the high-temperature oxidation resistance Nextel M720/alumina composite and reported excellent fatigue 6. The need for environmentally stable composites moti- resistance and 100% retention of tensile strength at 1200C vated the development of CMCs based on environmentally stable The fatigue limit(based on a run-out condition of 10 cycles) oxide constituents [24-32] was 170 MPa(88% UTS at 1200C). However, the creep pe c More recently it has been demonstrated that similar crack- formance at 1200 C was poor. Creep run-out(defined as 100h filecting behavior can also be achieved by means of a finely at creep stress)was achieved only at stress levels below 50% distributed porosity in the matrix instead of a separate interface UTS between matrix and fibers [33]. This microstructural design phi- Because creep was shown to be considerably more damag losophy implicitly accepts the strong fiber/matrix interface. It ing than cyclic loading to oxide-oxide CMCs with porous matrix builds on the experience with porous interlayers as crack deflec- [50,51], high-temperature creep resistance remains among the tion paths 34,35] and extends the concept to utilize a porous key issues that must be addressed before using these materials matrix as a surrogate. The concept has been successfully demon- in advanced aerospace applications. The objective of this study strated for oxide-oxide composites [24, 28, 32, 36-39]. Resulting is to investigate the effects of monazite fiber coating on com- oxide/oxide Cmcs exhibit damage tolerance combined with pressive creep resistance of NextelM610/alumina(N610/A) inherent oxidation resistance. However, due to the strong bond- composite with a porous matrix. Keller et al. [52] investigated ing between the fiber and matrix, a minimum matrix porosity is the effectiveness of monazite coatings in Nextel M610/alumina needed for this concept to work [40]. An extensive review of the porous matrix composites after long-term exposure at and mechanisms and mechanical properties of porous matrix CMCs 1200.C Coated fiber samples exhibited better tensile strength is given in [41]. retention after 1000 h at 1200C when compared to the control For a dense(>90%)matrix composite, an interfacial coat-(uncoated fiber) material. Ruggles-Wrenn et al. [53]reported ing is needed for crack deflection. An extensive review of the that the use of monazite coating in a NextelM610/alumina research on oxide coatings for oxide and non-oxide compos- porous matrix CMC resulted in improved tensile creep resistance ites has been given by Kerans et al. [42]. The development at 900C. This effort investigates the compressive creep-rupture of oxide-oxide composites that rely on a weak fiber/matrix behavior of the NextelTM610/monazite/alumina(N610/M/A) interface for crack deflection prompted research into oxidation- composite developed at the Air Force Research Labora- resistant fiber coatings that are chemically stable with the tory(AFRL/MLLN), Materials and Manufacturing Directorate omposite constituents. Monazite (Lapo4) and scheelite are Compressive creep-rupture tests were conducted at 900 and among the various oxidation-resistant coating materials that 1100 C for compressive creep stresses ranging from -50 have been investigated. Numerous studies examined composites to -95 MPa. Composite microstructure, damage and failure gan and Marshall [31, 43, mechanisms as well as effects of variation in microstruc Morgan et al. 144] and Chawla et al. [45] showed that due to ture on mechanical response were examined. While differences the chemical compatibility of monazite with alumina at high in processing and consequently in the composite microstruc temperature, monazite was a good candidate for a weak inter- ture did not have a significant effect on tensile respons face material for alumina-based composites. Since then, multiple of the CMC, effects on the compressive properties were investigations into the production of monazite coatings and its dramatic use with different fiber/matrix combinations [46-48] have been carried out. Degradation of fiber strength caused by the coating 2. Material and experimental arrangements nd long-term, high-temperature heat treatments was identified as the key problem with monazite coatings [48]. Boakye et al. The NextelTM610/monazite/alumina and Nextel TM610/ [49]explored the effects of different liquid precursors on coating alumina composites were processed as described elsewhere characteristics and tensile strength of coated fibers, and devel- [ 53]. Both the monazite-containing billets and the control oped monazite coating that did not significantly degrade the fiber (N610/A)billets consisted of eight layers in a symmetric cross- ply orientation of [(0/90))2s. Billet properties, namely fiber Porous matrix oxide/oxide CMCs exhibit several behavior volume fraction (Vr) and density, are summarized in Table 1 trends that are distinctly different from those exhibited by tradi- where properties of N610/M/A and n610/a billets used in tional CMCs with a fiber/matrix interface. For these materials, prior work [53] are also included. Samples produced using fatigue is significantly more damaging than creep. On the other the monazite-coated NextelTM610 fiber exhibited bulk den hand, recent investigations into the high-temperature mechan- sities of 2.48-2.68 g/em,, with 35 vol. composite porosity ical behavior of porous matrix oxide-oxide CMCs [50,51] and 45-50 vol %o matrix porosity (including large microcracks). revealed that creep loading was considerably more damaging Micrographs of the as-processed material shown in Fig. I reveal than fatigue. Zawada et al. [50] have shown that a porous shrinkage and sintering cracks that occurred during the cooling matrix Nextel M610/aluminosilicate composite exhibited high stage of the composite processing. Fig. 1(a)shows extensive sur fatigue limit, long fatigue life and near 100% strength reten- face microcracking, while interlaminar matrix cracks are seen tion at 1000oC. However, creep lives were short, indicating in Fig. 1(b). The N610/M/A specimens were cut from 10 dif- low creep resistance and limiting the use of that CMC to tem- ferent billets, and N610/A specimens from 5 different billets peratures below 1000C. Ruggles-Wrenn et al. [51] examined Specimen numbers contain reference to the billet number. For high-temperature fatigue and creep behaviors of a porous matrix example, number Bl-1 refers to the specimen 1 from billet 1
P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 591 substantially improve the high-temperature oxidation resistance [23]. The need for environmentally stable composites motivated the development of CMCs based on environmentally stable oxide constituents [24–32]. More recently it has been demonstrated that similar crackdeflecting behavior can also be achieved by means of a finely distributed porosity in the matrix instead of a separate interface between matrix and fibers [33]. This microstructural design philosophy implicitly accepts the strong fiber/matrix interface. It builds on the experience with porous interlayers as crack deflection paths [34,35] and extends the concept to utilize a porous matrix as a surrogate. The concept has been successfully demonstrated for oxide–oxide composites[24,28,32,36–39]. Resulting oxide/oxide CMCs exhibit damage tolerance combined with inherent oxidation resistance. However, due to the strong bonding between the fiber and matrix, a minimum matrix porosity is needed for this concept to work [40]. An extensive review of the mechanisms and mechanical properties of porous matrix CMCs is given in [41]. For a dense (>90%) matrix composite, an interfacial coating is needed for crack deflection. An extensive review of the research on oxide coatings for oxide and non-oxide composites has been given by Kerans et al. [42]. The development of oxide–oxide composites that rely on a weak fiber/matrix interface for crack deflection prompted research into oxidationresistant fiber coatings that are chemically stable with the composite constituents. Monazite (LaPO4) and scheelite are among the various oxidation-resistant coating materials that have been investigated. Numerous studies examined composites containing monazite coatings. Morgan and Marshall [31,43], Morgan et al. [44] and Chawla et al. [45] showed that due to the chemical compatibility of monazite with alumina at high temperature, monazite was a good candidate for a weak interface material for alumina-based composites. Since then, multiple investigations into the production of monazite coatings and its use with different fiber/matrix combinations [46–48] have been carried out. Degradation of fiber strength caused by the coating and long-term, high-temperature heat treatments was identified as the key problem with monazite coatings [48]. Boakye et al. [49] explored the effects of different liquid precursors on coating characteristics and tensile strength of coated fibers, and developed monazite coating that did not significantly degrade the fiber strength. Porous matrix oxide/oxide CMCs exhibit several behavior trends that are distinctly different from those exhibited by traditional CMCs with a fiber/matrix interface. For these materials, fatigue is significantly more damaging than creep. On the other hand, recent investigations into the high-temperature mechanical behavior of porous matrix oxide–oxide CMCs [50,51] revealed that creep loading was considerably more damaging than fatigue. Zawada et al. [50] have shown that a porous matrix NextelTM610/aluminosilicate composite exhibited high fatigue limit, long fatigue life and near 100% strength retention at 1000 ◦C. However, creep lives were short, indicating low creep resistance and limiting the use of that CMC to temperatures below 1000 ◦C. Ruggles-Wrenn et al. [51] examined high-temperature fatigue and creep behaviors of a porous matrix NextelTM720/alumina composite and reported excellent fatigue resistance and 100% retention of tensile strength at 1200 ◦C. The fatigue limit (based on a run-out condition of 105 cycles) was 170 MPa (88% UTS at 1200 ◦C). However, the creep performance at 1200 ◦C was poor. Creep run-out (defined as 100 h at creep stress) was achieved only at stress levels below 50% UTS. Because creep was shown to be considerably more damaging than cyclic loading to oxide–oxide CMCs with porous matrix [50,51], high-temperature creep resistance remains among the key issues that must be addressed before using these materials in advanced aerospace applications. The objective of this study is to investigate the effects of monazite fiber coating on compressive creep resistance of NextelTM610/alumina (N610/A) composite with a porous matrix. Keller et al. [52] investigated the effectiveness of monazite coatings in NextelTM610/alumina porous matrix composites after long-term exposure at 1100 and 1200 ◦C. Coated fiber samples exhibited better tensile strength retention after 1000 h at 1200 ◦C when compared to the control (uncoated fiber) material. Ruggles-Wrenn et al. [53] reported that the use of monazite coating in a NextelTM610/alumina porous matrix CMC resulted in improved tensile creep resistance at 900 ◦C. This effort investigates the compressive creep–rupture behavior of the NextelTM610/monazite/alumina (N610/M/A) composite developed at the Air Force Research Laboratory (AFRL/MLLN), Materials and Manufacturing Directorate. Compressive creep–rupture tests were conducted at 900 and 1100 ◦C for compressive creep stresses ranging from −50 to −95 MPa. Composite microstructure, damage and failure mechanisms, as well as effects of variation in microstructure on mechanical response were examined. While differences in processing and consequently in the composite microstructure did not have a significant effect on tensile response of the CMC, effects on the compressive properties were dramatic. 2. Material and experimental arrangements The NextelTM610/monazite/alumina and NextelTM610/ alumina composites were processed as described elsewhere [53]. Both the monazite-containing billets and the control (N610/A) billets consisted of eight layers in a symmetric crossply orientation of [(0◦/90◦)]2s. Billet properties, namely fiber volume fraction (Vf) and density, are summarized in Table 1, where properties of N610/M/A and N610/A billets used in prior work [53] are also included. Samples produced using the monazite-coated NextelTM610 fiber exhibited bulk densities of 2.48–2.68 g/cm3, with 35 vol.% composite porosity and 45–50 vol.% matrix porosity (including large microcracks). Micrographs of the as-processed material shown in Fig. 1 reveal shrinkage and sintering cracks that occurred during the cooling stage of the composite processing. Fig. 1(a) shows extensive surface microcracking, while interlaminar matrix cracks are seen in Fig. 1(b). The N610/M/A specimens were cut from 10 different billets, and N610/A specimens from 5 different billets. Specimen numbers contain reference to the billet number. For example, number B1-1 refers to the specimen 1 from billet 1
PR Jackson et al. /Materials Science and Engineering A 454-455(2007)590-601 Table 1 (Model 360FE, Leica)as well as optical microscopy. The SEM Summary of billet properties for the N61O/M/A and the N6l0VA composit specimens were coated with gold or car Billet Fiber volume fraction(%) Density (g/cm) N610/monazite/alumina composite 3. Test procedures All tests were conducted in laboratory air environment at 900 43.0 and 1100C. In all tests, a specimen was heated to the test tem- 44 perature at a rate of 1C/s, and held at temperature for additional BIO 429 BIla 883 30 min prior to testing. Monotonic tension and monotonic com- pression tests were performed in stroke control with a constant B12 displacement rate of 0.05 mm/s In compressive creep-rupture B13 tests specimens were loaded to the creep stress level at the stress N610/alumina composite 839939 rate magnitude of 20 MPa/s Creep run-out was defined as either 100 h at a given creep stress or 50h at creep stress if creep strain rate magnitude remained below 10-9s-I. In each test. stress-strain data were recorded during the loading to the creep stress level as well as during the creep period. Thus, both total B strain and creep strain could be calculated and examined. To B19 2.82 determine the retained tensile(compressive)strength and mod- a billets used in prior work [53]. ulus, specimens that achieved creep run-out were subjected to tensile(compressive) tests to failure at the temperature of the Billets were cut into 16 mm x 126 mm flat rectangular coupons, creep test. In some cases one specimen was tested per test con- which were machined into dog bone-shaped tensile specimens dition. The authors recognize that this is a limited set of data with a 10 mm x 18 mm gage section, and into 20 mm x 126 mm However, this scoping research serves to identify the tempera straight-sided compression specimens. Diamond-grit grinding ture range where the use of monazite coating results in improved was used for billets B1-B8, and the abrasive water-jet machin- creep resistance. Furthermore, results of this exploratory effort ing, for billets B9-B19 can be used to determine whether a more rigorous investigation a servocontrolled mrs mechanical testing machine of the effectiveness of monazite coating in this CMC or in a equipped with hydraulic water-cooled collet grips, a compact different material system should be undertaken wo-zone resistance-heated furnace, and two temperature con- trollers were used in all tests. An mrs TestStar ii digital 4. Results and discussion controller was employed for input signal generation and data acquisition Strain measurement was accomplished with an MTs 4.1. Monotonic tension and monotonic compression high-temperature air-cooled uniaxial extensometer. For elevated ature testing, thermocouples were bonded to the speci- The N610/M/A and N610/A specimens were tested in ten- mens using alumina cement(Zircar) to calibrate the furnace on sion at 900C. In addition, the monazite-containing composite a periodic basis. The furnace controller(using a non-contacting was subjected to tensile test to failure at 23C. Tensile test thermocouple exposed to the ambient environment near the test results are summarized in Table 2, where elastic modulus, ulti- specimen)was adjusted to determine the power setting needed mate tensile strength (UTS)and failure strain are presented for to achieve the desired temperature of the test specimen. The each test temperature. Results from prior work [53] are also determined power setting was then used in actual tests included In compression, the N601/M/A specimens were tested ture surfaces of failed specimens were examined usi at 23, 900 and 1100C, and the N61O/A specimens, at 900 10m Fig. 1. Optical micrographs of the as-processed material showing shrinkage cracks:(a) extensive surface microcracking and (b) interlaminar matrix cracks
592 P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 Table 1 Summary of billet properties for the N610/M/A and the N610/A composites Billet Fiber volume fraction (%) Density (g/cm3) N610/monazite/alumina composite B1a 44.3 2.55 B2a 39.7 2.48 B4a 40.2 2.49 B5a 43.0 2.50 B9a 44.7 2.58 B10a 42.9 2.58 B11a 41.0 2.53 B12 30.0 2.68 B13 29.1 2.55 B14 26.4 2.58 B15 29.0 2.53 N610/alumina composite B3a 51.6 2.95 B6a 54.8 2.99 B17 35.0 2.73 B18 36.0 2.79 B19 34.5 2.82 a Billets used in prior work [53]. Billets were cut into 16 mm × 126 mm flat rectangular coupons, which were machined into dog bone-shaped tensile specimens with a 10 mm × 18 mm gage section, and into 20 mm × 126 mm straight-sided compression specimens. Diamond-grit grinding was used for billets B1–B8, and the abrasive water-jet machining, for billets B9–B19. A servocontrolled MTS mechanical testing machine equipped with hydraulic water-cooled collet grips, a compact two-zone resistance-heated furnace, and two temperature controllers were used in all tests. An MTS TestStar II 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 using alumina cement (Zircar) 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. The determined power setting was then used in actual tests. Fracture surfaces of failed specimens were examined using SEM (Model 360FE, Leica) as well as optical microscopy. The SEM specimens were coated with gold or carbon. 3. Test procedures All tests were conducted in laboratory air environment at 900 and 1100 ◦C. In all tests, a specimen was heated to the test temperature at a rate of 1 ◦C/s, and held at temperature for additional 30 min prior to testing. Monotonic tension and monotonic compression tests were performed in stroke control with a constant displacement rate of 0.05 mm/s. In compressive creep–rupture tests specimens were loaded to the creep stress level at the stress rate magnitude of 20 MPa/s. Creep run-out was defined as either 100 h at a given creep stress or 50 h at creep stress if creep strain rate magnitude remained below 10−9 s−1. In each test, stress–strain data were recorded during the loading to the creep stress level as well as during the creep period. Thus, both total strain and creep strain could be calculated and examined. To determine the retained tensile (compressive) strength and modulus, specimens that achieved creep run-out were subjected to tensile (compressive) tests to failure at the temperature of the creep test. In some cases one specimen was tested per test condition. The authors recognize that this is a limited set of data. However, this scoping research serves to identify the temperature range where the use of monazite coating results in improved creep resistance. Furthermore, results of this exploratory effort can be used to determine whether a more rigorous investigation of the effectiveness of monazite coating in this CMC or in a different material system should be undertaken. 4. Results and discussion 4.1. Monotonic tension and monotonic compression The N610/M/A and N610/A specimens were tested in tension at 900 ◦C. In addition, the monazite-containing composite was subjected to tensile test to failure at 23 ◦C. Tensile test results are summarized in Table 2, where elastic modulus, ultimate tensile strength (UTS) and failure strain are presented for each test temperature. Results from prior work [53] are also included. In compression, the N601/M/A specimens were tested at 23, 900 and 1100 ◦C, and the N610/A specimens, at 900 Fig. 1. Optical micrographs of the as-processed material showing shrinkage cracks: (a) extensive surface microcracking and (b) interlaminar matrix cracks
PR Jackson et al. / Materials Science and Engineering A 454-455(2007)590-601 2 mary of tensile properties for the N610/M/A and the N610/A composites Specimen Temperature Tensile modulus Tensile Failure strength(MPa) strain(%) B13-3 0. 3(54) l80(117b B13-1900 107 0.19 B2.1a1100 76(56) 57(115) N610/alumina composite 129(73) 17(665) 0.09 Tension B17-190 75(62) 0.07 T=23° B3-2a1100 116(65b) 05(59) 0.11 0 0.05 Data from Ruggles-Wrenn et al. [53] ABS STRAIN (% b Adjusted for V:=0.29 and 1100 C. Modulus, strength and failure strain obtained in compression tests at various temperatures are summarized in Table 3. Tensile and compressive stress-strain curves obtaine for N610/A and N610/M/A composites at 23, 900 and 1100oC are shown in Fig. 2(a-c), respectively. Note that in the case of w compression, stress magnitude versus strain magnitude curves are presented. In order to facilitate comparison between results n 100 obtained for specimens with different fiber volume fractions results in Table 2 as well as all tensile data in Fig. 2 were adjusted nsion for Vr=0.29. Results of the monotonic tensile tests in this study T=900° are consistent with those obtained in prior work [53], where the monotonic tensile behavior of the two composites is described ABS STRAIN (% Compressive failure in fiber-reinforced composites is erally associated with microbuckling of the fibers [54-57] T=1100°c Flexural stresses in a fiber due to in-phase buckling lead to the formation of kink zones. which can cause fracture in brittle longitudinal compressive damage and fracture typically involve 9 200 axial splitting of the matrix, buckling of the fibers, and kink g 160F Compression fibers [58, 59]. In the case of the 0/90 cross-ply composites, for a 0/900 cross-ply CMC, compressive failure initiated with a nucleation of axial cracks between adjacent fibers in the 90 ATension plies. These cracks gradually form shear zones, which induce Ruggles-Wrenn, 2006 0 ply flexure and cause buckling and kinking of the 0 fibers, leading to local fiber fracture and subsequent composite failure For porous matrix composites, the matrix is exceptionally weak (c) ABS STRAIN (% Fig. 2. Monotonic stress-strain curves for N610/M/A and N610/A composites at:(a)23C.(b)900 C and(c)1100C. Tensile data from Ruggles-Wrenn et ummary of compressive properties for the N610/M/A and the N610/A al. [53] are also shown. All tensile data are adjusted for V(=0.29 composites Specimen Temperature Compressive Compressive Failure and the fibers bear most of the load once the oo bundles buckle modulus(GPa) strength(MPa) strain(%) profuse matrix microcracking takes place, resulting in the loss of N610/monazite/alumina composite fiber stabilization and consequently the loss of the composites B12-1 -113 0.19 load-bearing capacity. Composite failure is then reached Bl4-1 11 -0.18 At all temperatures investigated, compressive stress-strain B14-2 -0.17 curves of N610/M/A are nearly linear to failure, indicating that B14-3 -0.16 compression damage and fracture occur in close succession. N610/alumina composite At 23 and 900C compressive m B8-8 72 0.29 B18-71100 0.59 monazite-containing CMC are similar to the corresponding ten- sile values. However, at 1100C the compressive strength is
P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 593 Table 2 Summary of tensile properties for the N610/M/A and the N610/A composites Specimen Temperature ( ◦C) Tensile modulus (GPa) Tensile strength (MPa) Failure strain (%) N610/monazite/alumina composite B13-3 23 64 105 0.20 B9-1a 900 83 (54b) 180 (117b) 0.31 B13-1 900 55 107 0.19 B2-1a 1100 76 (56b) 157 (115b) 0.34 N610/alumina composite B3-1a 23 129 (73b) 117 (66b) 0.09 B17-1 900 75 (62b) 64 (53b) 0.07 B3-2a 1100 116 (65b) 105 (59b) 0.11 a Data from Ruggles-Wrenn et al. [53]. b Adjusted for Vf = 0.29. and 1100 ◦C. Modulus, strength and failure strain obtained in compression tests at various temperatures are summarized in Table 3. Tensile and compressive stress–strain curves obtained for N610/A and N610/M/A composites at 23, 900 and 1100 ◦C are shown in Fig. 2(a–c), respectively. Note that in the case of compression, stress magnitude versus strain magnitude curves are presented. In order to facilitate comparison between results obtained for specimens with different fiber volume fractions, results in Table 2 as well as all tensile data in Fig. 2 were adjusted for Vf = 0.29. Results of the monotonic tensile tests in this study are consistent with those obtained in prior work [53], where the monotonic tensile behavior of the two composites is described in detail. Compressive failure in fiber-reinforced composites is generally associated with microbuckling of the fibers [54–57]. Flexural stresses in a fiber due to in-phase buckling lead to the formation of kink zones, which can cause fracture in brittle fibers [58,59]. In the case of the 0◦/90◦ cross-ply composites, longitudinal compressive damage and fracture typically involve axial splitting of the matrix, buckling of the fibers, and kink banding or shear banding [60–62]. Lankford [60] reported that for a 0◦/90◦ cross-ply CMC, compressive failure initiated with nucleation of axial cracks between adjacent fibers in the 90◦ plies. These cracks gradually form shear zones, which induce 0◦ ply flexure and cause buckling and kinking of the 0◦ fibers, leading to local fiber fracture and subsequent composite failure. For porous matrix composites, the matrix is exceptionally weak Table 3 Summary of compressive properties for the N610/M/A and the N610/A composites Specimen Temperature ( ◦C) Compressive modulus (GPa) Compressive strength (MPa) Failure strain (%) N610/monazite/alumina composite B12-1 23 74 −113 −0.19 B14-1 900 47 −110 −0.18 B14-2 900 58 −103 −0.17 B14-3 1100 63 −97 −0.16 N610/alumina composite B18-8 900 72 −230 −0.29 B18-7 1100 73 −244 −0.59 Fig. 2. Monotonic stress–strain curves for N610/M/A and N610/A composites at: (a) 23 ◦C, (b) 900 ◦C and (c) 1100 ◦C. Tensile data from Ruggles-Wrenn et al. [53] are also shown. All tensile data are adjusted for Vf = 0.29. and the fibers bear most of the load. Once the 0◦ bundles buckle, profuse matrix microcracking takes place, resulting in the loss of fiber stabilization and consequently the loss of the composite’s load-bearing capacity. Composite failure is then reached. At all temperatures investigated, compressive stress–strain curves of N610/M/A are nearly linear to failure, indicating that compression damage and fracture occur in close succession. At 23 and 900 ◦C compressive modulus and strength of the monazite-containing CMC are similar to the corresponding tensile values. However, at 1100 ◦C the compressive strength is
PR Jackson et al. /Materials Science and Engineering A 454-455(2007)590-601 slightly lower than the UTS. Note that at 1100C, in contrast to Table 4 the mostly linear compressive behavior, the tensile stress-strain Summary of compressive creep-rupture results for the NIO/M/A and the curve departs from linearity. Such non-linear behavior is indica- N610/A composites tive of progressive matrix cracking and crack deflection, the Specimen Temperature Creep stress Creep Time to mechanisms likely responsible for a considerably larger tensile rupture(s) failure strain and a somewhat higher tensile strength N610/monazite/alumina composite is seen in Fig. 2(b) that at 900C, B1461100 360.0003 stress-strain behavior of the uncoated fiber composite is also B14 360.0003 nearly linear to failure. However, at 1100C, the compres- B14-41100 360.0003 sive stress-strain curve of N610/A contains two parts. The first is a linear part which extends to a stress of approxi- mately-157 MPa The second part is also approximately linear with a noticeably lower slope. This stress-strain behavior indi- B19.2 1100 cates probable composite damage at the change in slope. It is B19-5 900 500650万 0.03 180.0003 3600004 -2.65 3600004 360.0002 0.03 180.000 likely that as the compressive stress approaches -157 MPa shear B196 0.01 180.000 cracking in the 90 fiber bundles and ply delamination take place. 0.05 180.000 The damage relieves constraints acting on the 0o bundles allow Run-out ing them to deform more readily by buckling, and leading to a lower slope in the stress-strain curve. When the compressive °C( see Fig.3(b)) stress reaches -240 MPa, the composite uckling and exhibit primary and secondary creep regimes. Unlike in the shear fracture of the 0o bundles. At 900 and 1100C, the com- case of tension, in compression transition from primary to sec pressive modulus and strength of N610/A considerably exceed ondary creep occurs much later in creep life. For N610/M/A the corresponding tensile values. It appears that the processing- composite, primary creep transitions into secondary creep after related cracks in both matrix and 0 fiber bundles close up during 5h at the stress of-50MPa. At stresses <-65 MPa, primary compression producing a higher modulus. It can be further con- jectured that in tension the same processing-related cracks would start propagating at lower stress levels, leading to lower UTS It is noteworthy that while the use of the monazite coat- N610/Monazite/Alumina T=1100°c ing served to improve the high-temperature tensile strength, the 4.0 gth of the monazite -contai considerably lower than that of the uncoated fiber CMC. The 3.0 addition of the monazite coating resulted in v54% loss in com- o pressive strength at 900C, and in -60% loss in compressive u2.0 100 MP strength at 1100°C 20 MPa It is important to note that in all monotonic tension and com- 61.0 pression tests, as well as in all other tests reported herein, the 40 MPa failure occurred within the gage section of the extensometer. 25050075010001250150017502000 4.2. Creep-rupture 10.0 Results of the compressive creep-rupture tests are summa- N610/Monazite/Alumina T=1100°c rized in Table 4, where creep strain accumulation and rupture 9 N610/Alumina time are shown for each test temperature and creep stress z level. Creep curves obtained at 1100 and 900C are shown in Figs. 3 and 4, respectively. Tensile creep data from prior work [53]is included in Figs. 3 and 4 for comparison Results of the u 65 MP. Tensile creep curves obtained for the n610/M/A at 1100%C 54.0 tensile creep tests of the two composites appear in Ref. [ 53] (see Fig 3(a)) exhibit primary, secondary and tertiary creep regimes. Primary creep rapidly transitions to secondary creep. For stresses 2 100 MPa, transition from secondary to tertiary creep occurs during the first third of the creep life. Creep strain 100000200000300000400000500000 and creep life decrease with increasing applied stress. Note that (b) TIME (s) the tensile creep strains accumulated in all tests conducted at 1100oC significantly exceed the failure strain obtained in the Fig. 3. Creep curves for NlO/MA and Nlo/A composites at 1100C: (a) tension test. The tensile creep run-out(set to 100h) was not Time scale in(a)is selected to clearly show creep curves obtained at stress levels above 100 MPa
594 P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 slightly lower than the UTS. Note that at 1100 ◦C, in contrast to the mostly linear compressive behavior, the tensile stress–strain curve departs from linearity. Such non-linear behavior is indicative of progressive matrix cracking and crack deflection, the mechanisms likely responsible for a considerably larger tensile failure strain and a somewhat higher tensile strength. It is seen in Fig. 2(b) that at 900 ◦C, the compressive stress–strain behavior of the uncoated fiber composite is also nearly linear to failure. However, at 1100 ◦C, the compressive stress–strain curve of N610/A contains two parts. The first is a linear part which extends to a stress of approximately −157 MPa. The second part is also approximately linear with a noticeably lower slope. This stress–strain behavior indicates probable composite damage at the change in slope. It is likely that as the compressive stress approaches −157 MPa shear cracking in the 90◦ fiber bundles and ply delamination take place. The damage relieves constraints acting on the 0◦ bundles allowing them to deform more readily by buckling, and leading to a lower slope in the stress–strain curve. When the compressive stress reaches −240 MPa, the composite fails by buckling and shear fracture of the 0◦ bundles. At 900 and 1100 ◦C, the compressive modulus and strength of N610/A considerably exceed the corresponding tensile values. It appears that the processingrelated cracks in both matrix and 0◦ fiber bundles close up during compression producing a higher modulus. It can be further conjectured that in tension the same processing-related cracks would start propagating at lower stress levels, leading to lower UTS. It is noteworthy that while the use of the monazite coating served to improve the high-temperature tensile strength, the compressive strength of the monazite-containing composite was considerably lower than that of the uncoated fiber CMC. The addition of the monazite coating resulted in ∼54% loss in compressive strength at 900 ◦C, and in ∼60% loss in compressive strength at 1100 ◦C. It is important to note that in all monotonic tension and compression tests, as well as in all other tests reported herein, the failure occurred within the gage section of the extensometer. 4.2. Creep–rupture Results of the compressive creep–rupture tests are summarized in Table 4, where creep strain accumulation and rupture time are shown for each test temperature and creep stress level. Creep curves obtained at 1100 and 900 ◦C are shown in Figs. 3 and 4, respectively. Tensile creep data from prior work [53] is included in Figs. 3 and 4 for comparison. Results of the tensile creep tests of the two composites appear in Ref. [53]. Tensile creep curves obtained for the N610/M/A at 1100 ◦C (see Fig. 3(a)) exhibit primary, secondary and tertiary creep regimes. Primary creep rapidly transitions to secondary creep. For stresses ≥ 100 MPa, transition from secondary to tertiary creep occurs during the first third of the creep life. Creep strain and creep life decrease with increasing applied stress. Note that the tensile creep strains accumulated in all tests conducted at 1100 ◦C significantly exceed the failure strain obtained in the tension test. The tensile creep run-out (set to 100 h) was not achieved. Table 4 Summary of compressive creep–rupture results for the N610/M/A and the N610/A composites Specimen Temperature ( ◦C) Creep stress (MPa) Creep strain (%) Time to rupture (s) N610/monazite/alumina composite B14-6 1100 −50 −1.05 360,000a B14-5 1100 −65 −4.53 360,000a B14-4 1100 −75 −7.79 360,000a B15-2 900 −50 −0.03 180,000a N610/alumina composite B19-4 1100 −50 −1.53 360,000a B19-3 1100 −65 −2.65 360,000a B19-2 1100 −75 −6.95 360,000a B19-5 900 −50 −0.03 180,000a B19-6 900 −75 −0.01 180,000a B19-7 900 −95 −0.05 180,000a a Run-out. Compressive creep curves obtained at 1100 ◦C (see Fig. 3(b)) exhibit primary and secondary creep regimes. Unlike in the case of tension, in compression transition from primary to secondary creep occurs much later in creep life. For N610/M/A composite, primary creep transitions into secondary creep after ∼5 h at the stress of −50 MPa. At stresses ≤ −65 MPa, primary Fig. 3. Creep curves for N610/M/A and N610/A composites at 1100 ◦C: (a) tensile creep, data from Ruggles-Wrenn et al. [53] and (b) compressive creep. Time scale in (a) is selected to clearly show creep curves obtained at stress levels above 100 MPa
PR Jackson et al. / Materials Science and Engineering A 454-455(2007)590-601 1E+00 ■N610MA.1100"cT 口N610/MA.11o0℃c, Compression T=900° 态1Ea月·wmm:解 0°c, Compression Ruggles-Wrenn, 2006 0.15 80 MP 0.10 N610A 1E07 0.05 810 a1E° Creep rate magnitude 105 80000 00000 1E-11 ABS CREEP STRESS (MPa) T=900°c for N610/M/A and N610A ceramic composites at 900 and 1100C. from Ruggles-Wrenn et al. [53] are also shown. All tensile data are adjusted for V=0.29. creep strain accumulations were low and steady-state creep rate magnitudes remained below 10-10s-I. It is seen that the com- pressive creep strain magnitudes accumulated in all tests at 900C are an order of magnitude lower than the failure strain magnitude obtained in the compression test. Note that a run- out was achieved in all compressive creep tests conducted at 900C. Conversely, tensile creep run-out was achieved only for 50000 100000 150000 200000 N610/M/A specimens tested at creep stresses s 120 MPa. Minimum creep rate was reached in all tests. Creep strain rate magnitude as a function of applied stress magnitude is Fig 4. Creep curves for NIO/M/Aand NlO/A composites at"C: (a)tensile shown in Fig. 5, where results of previous work [53]are also creep, data from Ruggles-Wrenn et al. [53] and(b) compressive creep included. To facilitate the comparison between the creep prop- erties of specimens with different fiber volume fractions, all creep persists during the first 50h of the creep test. Similar tensile data in Fig. 5 were adjusted for Vf=0. 29. As expected, observations can be made for N610/A. All specimens tested in tensile creep strain rates increase with increasing temperature oppressive creep at 1100C accumulated extensive amounts As demonstrated in prior work [53], at 1100C tensile creep of creep strain. For both composites, compressive creep strain rates of N610/M/A are what may be expected from N610 fibers increases with the magnitude of applied stress. At-50 MPa, the alone. Results in Fig. 5 also show that at 1100C the tensile uncoated fiber composite produced a slightly higher creep strain creep rates are at least two orders of magnitude higher than than the monazite-containing CMC. However, at creep stress the compressive creep rate magnitudes produced at the same levels <-65 MPa, N610/M/A accumulated larger compressive applied stress magnitude. It is recognized that different failure creep strains than N61O/A Compressive creep strain magni- mechanisms are associated with tensile and compressive creep. tudes accumulated in all tests at 1100C significantly exceed Fibers do not play as vital a role in the response of the compos- the failure strain magnitude obtained in the compression test. ite material under compressive creep as they do under tensile All compressive creep tests conducted at 1100C achieved run- creep. While compressive creep rate magnitudes also increase out. Both the monazite-containing composite and the uncoated with rising temperature, the increase is not as dramatic as that in fiber CMC survived 100h of compressive creep at stress levels the case of tensile creep In addition, it is seen that compressive ranging from -50 to -75 MPa. creep rates of the monazite-containing composite are similar At 900C tensile as well as compressive creep curves to those produced by the uncoated fiber CMC. At 900C,ten- obtained for both composites(see Fig. 4)exhibit primary and sile as well as ssive creep strain rate magnitudes of both secondary creep regimes. In both tension and compression, tran- N610/M/A and N61O/A( with the exception of the tensile creep sition from primary to secondary creep occurs early in creep life. rate at 150 MPa)are s 10-8s- In tension, secondary creep continues to failure. Compressive Stress-rupture behavior is summarized in Fig. 6, where creep creep curves in Fig 4(b )indicate that secondary creep is likely to stress magnitude is plotted versus time to rupture at 900 and persist for the duration of the creep life. In both tension and com- 1100C for both composites. Tensile creep-rupture life of the ression, increasing magnitude of creep stress appears to have N610/M/A increases considerably with decreasing temperature little effect on creep strain magnitude, which remains <.05%. At 1100C, tensile creep life was 50,432s(14 h)at the low Compressive creep tests were interrupted after 50h because stress of 40 MPa, and mere 75 s at 120 MPa. At 900C, the
P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 595 Fig. 4. Creep curves for N610/M/A and N610/A composites at 900 ◦C: (a) tensile creep, data from Ruggles-Wrenn et al. [53] and (b) compressive creep. creep persists during the first 50 h of the creep test. Similar observations can be made for N610/A. All specimens tested in compressive creep at 1100 ◦C accumulated extensive amounts of creep strain. For both composites, compressive creep strain increases with the magnitude of applied stress. At −50 MPa, the uncoated fiber composite produced a slightly higher creep strain than the monazite-containing CMC. However, at creep stress levels ≤ −65 MPa, N610/M/A accumulated larger compressive creep strains than N610/A. Compressive creep strain magnitudes accumulated in all tests at 1100 ◦C significantly exceed the failure strain magnitude obtained in the compression test. All compressive creep tests conducted at 1100 ◦C achieved runout. Both the monazite-containing composite and the uncoated fiber CMC survived 100 h of compressive creep at stress levels ranging from −50 to −75 MPa. At 900 ◦C tensile as well as compressive creep curves obtained for both composites (see Fig. 4) exhibit primary and secondary creep regimes. In both tension and compression, transition from primary to secondary creep occurs early in creep life. In tension, secondary creep continues to failure. Compressive creep curves in Fig. 4(b) indicate that secondary creep is likely to persist for the duration of the creep life. In both tension and compression, increasing magnitude of creep stress appears to have little effect on creep strain magnitude, which remains ≤ 0.05%. Compressive creep tests were interrupted after 50 h because Fig. 5. Minimum creep rate magnitude as a function of applied stress magnitude for N610/M/A and N610/A ceramic composites at 900 and 1100 ◦C. Tensile data from Ruggles-Wrenn et al. [53] are also shown. All tensile data are adjusted for Vf = 0.29. creep strain accumulations were low and steady-state creep rate magnitudes remained below 10−10 s−1. It is seen that the compressive creep strain magnitudes accumulated in all tests at 900 ◦C are an order of magnitude lower than the failure strain magnitude obtained in the compression test. Note that a runout was achieved in all compressive creep tests conducted at 900 ◦C. Conversely, tensile creep run-out was achieved only for N610/M/A specimens tested at creep stresses ≤ 120 MPa. Minimum creep rate was reached in all tests. Creep strain rate magnitude as a function of applied stress magnitude is shown in Fig. 5, where results of previous work [53] are also included. To facilitate the comparison between the creep properties of specimens with different fiber volume fractions, all tensile data in Fig. 5 were adjusted for Vf = 0.29. As expected, tensile creep strain rates increase with increasing temperature. As demonstrated in prior work [53], at 1100 ◦C tensile creep rates of N610/M/A are what may be expected from N610 fibers alone. Results in Fig. 5 also show that at 1100 ◦C the tensile creep rates are at least two orders of magnitude higher than the compressive creep rate magnitudes produced at the same applied stress magnitude. It is recognized that different failure mechanisms are associated with tensile and compressive creep. Fibers do not play as vital a role in the response of the composite material under compressive creep as they do under tensile creep. While compressive creep rate magnitudes also increase with rising temperature, the increase is not as dramatic as that in the case of tensile creep. In addition, it is seen that compressive creep rates of the monazite-containing composite are similar to those produced by the uncoated fiber CMC. At 900 ◦C, tensile as well as compressive creep strain rate magnitudes of both N610/M/A and N610/A (with the exception of the tensile creep rate at 150 MPa) are ≤ 10−8 s−1. Stress–rupture behavior is summarized in Fig. 6, where creep stress magnitude is plotted versus time to rupture at 900 and 1100 ◦C for both composites. Tensile creep–rupture life of the N610/M/A increases considerably with decreasing temperature. At 1100 ◦C, tensile creep life was 50,432 s (∼14 h) at the low stress of 40 MPa, and mere 75 s at 120 MPa. At 900 ◦C, the
596 PR Jackson et al. /Materials Science and Engineering A 454-455(2007)590-601 the as-processed material. Both specimens retained 100% of their compressive strength. Furthermore, prior creep appears 250日·M61oMA.o0c. Tensi N610/A, 900C, Compression to have increased compressive modulus. The pre-crept speci ◆N610A,900°c, Tensi mens exhibited higher stiffness values. This indicates that the processing-related cracks may have closed up during com- 9 50 FN610MA, UTS at 110-C pressive creep producing a higher compressive modulus. To evaluate the effects of compressive creep on tensile strength and stiffness, a monazite-containing specimen that achieved a run-out in a-50 MPa compressive creep test at 900C was subjected to a tensile test to failure at that temperature The pre-crept specimen produced a modulus of 68 GPa and strength of 162 MPa, retaining better than 100% of its tensile 100010000100000100000 strength and exhibiting a noticeable increase in modulus. Ten- Time(s) sile stress-strain behavior of the specimen subjected to prior Fig. 6. Creep stress magnitude vs time to rupture for N610/M/A and N610vA compressive creep remained qualitative similar to that of the as- ramic composites at 900 and 1100.C. Tensile data from Ruggles-Wrenn et processed material( see Fig. 8). Retained tensile strength and al. [53] are also shown modulus of the two specimens that reached tensile creep run- out are given in Ref. [53]. Both specimens retained over 90% Table 5 of their tensile strength and at least 90% of their modulus. Prior Retained compressive compressive creep at 1100C ies of the N610/M/A specimens subjected to prior tensile creep had no qualitative effect on tensile stress-strain rength(MPa) modulus(GPa) failure(%) 4.3. Composite microstructure B145 0.13 10 -0.12 Fracture surfaces of the N610/M/A specimens tested in com- pression at 900 and 1100C are shown in ig. 9(d). Brushy 120MPa tensile creep test achieved a run-out. Furthermore, fracture surfaces indicative of fibrous fracture are produced at addition of the monazite coating is seen to significantly improve both temperatures. The optical micrographs in Fig. 9(a and the tensile creep life at 900C. Compressive creep life appears b)reveal the"stepwise"topography of the fracture surfaces to be relatively independent of temperature or applied stress. All obtained at 900oC. The specimen exhibits a fairly large dam- compressive creep tests achieved a run-out. Furthermore, the use age zone of 30-35 mm in length. Extensive delamination is of the monazite coating had little effect on compressive creep readily seen in Fig 9(b). Failure surfaces produced at 1100C life at both temperatures investigated. (see Fig 9(c and d)) are similar to those obtained at 900C, Retained compressive strength and modulus of the speci- although the specimen tested at 1100 C produced somewhat mens that achieved a run-out in -65 and-75 MPa compressive smaller damage zones creep tests at 1100C are summarized in Table 5. Compres- Fracture surfaces of the N610/A specimens tested in sive stress-strain curves obtained for the N610/M/A specimens monotonic compression at 900 and 1100C are presented subjected to prior compressive creep at 1100C are presented Fig I0(a-d). The contrast between the goooC fracture st in Fig. 7 together with the compressive stress-strain curve for faces of the two composites(Figs. 9(a and b)and 10(a and 200 106hat65 T=1100°c T=900°c 0 h at-50 MPa 02 h at-75 MPa Enu Enu As-Processed on doo As-Processed 50 N610/Monazite/Alumina 0 0.05 0.10 0.0 0.4 ABS STRAIN (% STRAIN (% Fig. 7. Effects of prior essive creep at 1100C on compressive Fig 8. Effects of prior compressive creep at 900C on tensile stress-strain tress-strain behavior of N610/M/A ceramic composite. behavior of N610/M/A ceramic composite Results from [53] are also included
596 P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 Fig. 6. Creep stress magnitude vs. time to rupture for N610/M/A and N610/A ceramic composites at 900 and 1100 ◦C. Tensile data from Ruggles-Wrenn et al. [53] are also shown. Table 5 Retained compressive properties of the N610/M/A specimens subjected to prior compressive creep at 1100 ◦C Specimen Creep stress (MPa) Retained strength (MPa) Retained modulus (GPa) Strain at failure (%) B14-5 −65 −113 89 −0.13 B14-4 −75 −102 72 −0.12 120 MPa tensile creep test achieved a run-out. Furthermore, addition of the monazite coating is seen to significantly improve the tensile creep life at 900 ◦C. Compressive creep life appears to be relatively independent of temperature or applied stress. All compressive creep tests achieved a run-out. Furthermore, the use of the monazite coating had little effect on compressive creep life at both temperatures investigated. Retained compressive strength and modulus of the specimens that achieved a run-out in −65 and −75 MPa compressive creep tests at 1100 ◦C are summarized in Table 5. Compressive stress–strain curves obtained for the N610/M/A specimens subjected to prior compressive creep at 1100 ◦C are presented in Fig. 7 together with the compressive stress–strain curve for Fig. 7. Effects of prior compressive creep at 1100 ◦C on compressive stress–strain behavior of N610/M/A ceramic composite. the as-processed material. Both specimens retained 100% of their compressive strength. Furthermore, prior creep appears to have increased compressive modulus. The pre-crept specimens exhibited higher stiffness values. This indicates that the processing-related cracks may have closed up during compressive creep producing a higher compressive modulus. To evaluate the effects of compressive creep on tensile strength and stiffness, a monazite-containing specimen that achieved a run-out in a −50 MPa compressive creep test at 900 ◦C was subjected to a tensile test to failure at that temperature. The pre-crept specimen produced a modulus of 68 GPa and strength of 162 MPa, retaining better than 100% of its tensile strength and exhibiting a noticeable increase in modulus. Tensile stress–strain behavior of the specimen subjected to prior compressive creep remained qualitative similar to that of the asprocessed material (see Fig. 8). Retained tensile strength and modulus of the two specimens that reached tensile creep runout are given in Ref. [53]. Both specimens retained over 90% of their tensile strength and at least 90% of their modulus. Prior tensile creep had no qualitative effect on tensile stress–strain behavior. 4.3. Composite microstructure Fracture surfaces of the N610/M/A specimens tested in compression at 900 and 1100 ◦C are shown in Fig. 9(d). Brushy fracture surfaces indicative of fibrous fracture are produced at both temperatures. The optical micrographs in Fig. 9(a and b) reveal the “stepwise” topography of the fracture surfaces obtained at 900 ◦C. The specimen exhibits a fairly large damage zone of 30–35 mm in length. Extensive delamination is readily seen in Fig. 9(b). Failure surfaces produced at 1100 ◦C (see Fig. 9(c and d)) are similar to those obtained at 900 ◦C, although the specimen tested at 1100 ◦C produced somewhat smaller damage zones. Fracture surfaces of the N610/A specimens tested in monotonic compression at 900 and 1100 ◦C are presented Fig. 10(a–d). The contrast between the 900 ◦C fracture surfaces of the two composites (Figs. 9(a and b) and 10(a and Fig. 8. Effects of prior compressive creep at 900 ◦C on tensile stress–strain behavior of N610/M/A ceramic composite. Results from [53] are also included
PR Jackson et al. / Materials Science and Engineering A 454-455(2007)590-601 (b) 5mm 4 mm 5 mm mm Fig 9. Fracture surfaces of N601/M/A specimens tested in compression at: (a)900C, (b)900oC-side view, (c)1100 C and(d)1100 C-side view. b))is striking. The fracture surface of the N610/M/A spec- related fiber fracture was observed in all cases. However, imen exhibits a large damage zone, significant delamination, the pre-crept specimens produced somewhat smaller damage and extensive uncorrelated fiber fracture with the length of fiber zones and less extensive delamination than the as-processed "brushes"reaching 8 mm. Conversely, the fracture surface of material the N61O/A specimen does not have a"brushy"appearance and Additional monotonic tension and compression tests were shows only a short damage zone of approximately 6 mm. In conducted at 900C using specimens from billet B15. Tensile compression at 900C, the uncoated fiber CMC responds more strength and modulus values were close to those reported for like a monolithic ceramic than a composite As seen in Fig. 10(c other billets in Table 3, indicating that no fiber degradation and d), the fracture surface topography of the N610/A speci- occurred during processing of this billet. Compressive stiff- mens becomes more serrated at 1100C. Some delamination ness values (ranging from 48 to 58 GPa) obtained for billet and a larger damage zone(approaching 17 mm in length) are B15 were also close to those reported in Table 3 for billet observed B14. However, compressive strength values of B15 specimens Fracture surfaces produced in compression tests to failure were surprisingly low. While specimens from billet B14 pro- conducted on N61O/M/A specimens that had achieved com- duced compressive strength values of -110 and -103 MPa, the pressive creep run-out at 1100C were also examined. Prior strength values obtained for the B15 specimens were between creep had little effect on the fracture surface appearance. Uncor- 6 and -58 MPa Such unusually low compressive strength
P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 597 Fig. 9. Fracture surfaces of N601/M/A specimens tested in compression at: (a) 900 ◦C, (b) 900 ◦C—side view, (c) 1100 ◦C and (d) 1100 ◦C—side view. b)) is striking. The fracture surface of the N610/M/A specimen exhibits a large damage zone, significant delamination, and extensive uncorrelated fiber fracture with the length of fiber “brushes” reaching ≈8 mm. Conversely, the fracture surface of the N610/A specimen does not have a “brushy” appearance and shows only a short damage zone of approximately 6 mm. In compression at 900 ◦C, the uncoated fiber CMC responds more like a monolithic ceramic than a composite. As seen in Fig. 10(c and d), the fracture surface topography of the N610/A specimens becomes more serrated at 1100 ◦C. Some delamination and a larger damage zone (approaching 17 mm in length) are observed. Fracture surfaces produced in compression tests to failure conducted on N610/M/A specimens that had achieved compressive creep run-out at 1100 ◦C were also examined. Prior creep had little effect on the fracture surface appearance. Uncorrelated fiber fracture was observed in all cases. However, the pre-crept specimens produced somewhat smaller damage zones and less extensive delamination than the as-processed material. Additional monotonic tension and compression tests were conducted at 900 ◦C using specimens from billet B15. Tensile strength and modulus values were close to those reported for other billets in Table 3, indicating that no fiber degradation occurred during processing of this billet. Compressive stiffness values (ranging from 48 to 58 GPa) obtained for billet B15 were also close to those reported in Table 3 for billet B14. However, compressive strength values of B15 specimens were surprisingly low. While specimens from billet B14 produced compressive strength values of −110 and −103 MPa, the strength values obtained for the B15 specimens were between −36 and −58 MPa. Such unusually low compressive strength
PR Jackson et al. /Materials Science and Engineering A 454-455(2007)590-601 (a) (b) 5 mm 4 mm mm 4 m Fig 10. Fracture surfaces of N601/A specimens tested in compression at: (a)900C.( b)900C-side view, (c)1100"C and (d)1100C-side view values suggested that composite microstructure of billet B15 into the 90 fiber layer, where they self-arrest. A magnified may be different from that of billet B14. Optical micrographs view of the fracture surface detail in Fig. 13(b) shows sepa- of the as-processed material from billets B15(Fig. 11)and B14 rate cracks initiating between the laminae, propagating along (Fig. 12)reveal significant differences in microstructure. The angled paths into the 90 fiber layer, and finally self-arresting cross-section of the b 15 material in Fig. lI shows thick matrix Note that no matrix-rich areas are visible in micrographs in bands located between the fibrous layers. In contrast, micro- Fig 13. A much different failure mechanism is observed in the graph of the as-processed material from B 14 in Fig 12 shows specimen from billet B15(see Fig. 14). Cracks initiate within only a few isolated matrix-rich areas and good infiltration of the and propagate directly through the 90 fiber layers. Unlike the matrix material into the 900 fiber I short self-arresting cracks seen in the B14 specimens, cracks The thick matrix-rich layers found in billet B15 together with seen in Fig. 14(a) propagate to considerable lengths before eas devoid of matrix in the fibrous layers suggest a process dissipating their energy. A higher magnification view of the frac control problem. Further examination of the fracture surfaces of ture surface of the B15 specimen in Fig. 14(b)shows a crack pecimens from billets B14 and B15 reveals that the presence propagating through the 90 fiber layer. Apparent differences of the matrix-rich regions and poor infiltration of the matrix in microstructure resulted in different failure mechanisms in material into the fibrous layers indeed caused early compr N610/M/A specimens from billets B14 and B15, and caused sive failures and low compressive strength. Fracture surfaces of early compression failures of the B15 specimens. Poor infil- N610/M/A specimens from billets B14 and B15 presented in tration of the matrix material into the fibrous layers in billet Figs. 13 and 14, respectively, suggest different failure mech- B15 may have been caused by either fiber bridging due to the nisms. It is seen in Fig. 13(a) that in B14 specimens the application of fiber coating or by high viscosity of the matrix cracks initiate between the laminae then propagate at an angle slurry
598 P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 Fig. 10. Fracture surfaces of N601/A specimens tested in compression at: (a) 900 ◦C, (b) 900 ◦C—side view, (c) 1100 ◦C and (d) 1100 ◦C—side view. values suggested that composite microstructure of billet B15 may be different from that of billet B14. Optical micrographs of the as-processed material from billets B15 (Fig. 11) and B14 (Fig. 12) reveal significant differences in microstructure. The cross-section of the B15 material in Fig. 11 shows thick matrix bands located between the fibrous layers. In contrast, micrograph of the as-processed material from B14 in Fig. 12 shows only a few isolated matrix-rich areas and good infiltration of the matrix material into the 90◦ fiber layer. The thick matrix-rich layers found in billet B15 together with areas devoid of matrix in the fibrous layers suggest a process control problem. Further examination of the fracture surfaces of specimens from billets B14 and B15 reveals that the presence of the matrix-rich regions and poor infiltration of the matrix material into the fibrous layers indeed caused early compressive failures and low compressive strength. Fracture surfaces of N610/M/A specimens from billets B14 and B15 presented in Figs. 13 and 14, respectively, suggest different failure mechanisms. It is seen in Fig. 13(a) that in B14 specimens the cracks initiate between the laminae then propagate at an angle into the 90◦ fiber layer, where they self-arrest. A magnified view of the fracture surface detail in Fig. 13(b) shows separate cracks initiating between the laminae, propagating along angled paths into the 90◦ fiber layer, and finally self-arresting. Note that no matrix-rich areas are visible in micrographs in Fig. 13. A much different failure mechanism is observed in the specimen from billet B15 (see Fig. 14). Cracks initiate within and propagate directly through the 90◦ fiber layers. Unlike the short self-arresting cracks seen in the B14 specimens, cracks seen in Fig. 14(a) propagate to considerable lengths before dissipating their energy. A higher magnification view of the fracture surface of the B15 specimen in Fig. 14(b) shows a crack propagating through the 90◦ fiber layer. Apparent differences in microstructure resulted in different failure mechanisms in N610/M/A specimens from billets B14 and B15, and caused early compression failures of the B15 specimens. Poor infiltration of the matrix material into the fibrous layers in billet B15 may have been caused by either fiber bridging due to the application of fiber coating or by high viscosity of the matrix slurry
PR Jackson et al. / Materials Science and Engineering A 454-455(2007)590-601 Matrix r 90° Fibers 0° Fibers Fig. 12. As-processed N601//A composite from billet B14. Few isolated matrix-rich areas and good matrix infi iber layerare apparent. 5. Concluding remarks 5.l.M 1 The compressive stress-strain behavior of N610/A and N610/M/A composites was investigated and the compressive properties measured at room and elevated temperatures. At temperatures<1100C, the compressive stress-strain behavior of the monazite-containing composite is nearly linear to fail ure; compressive modulus and strength magnitudes are close Fig. 11. As-processed N601/M/A composite from billet B15. Thick matrix to the corresponding tensile values obtained in prior work bands between the fibrous layers are clearly visible. (a) Optical micrograph [53]. Compressive behavior of the N6IO/A composite remains and(b) backscatter SEM image showing matrix-rich areas between fibrous lay. approximately linear to failure at 900C. However, at 1100C, ers and regions devoid of matrix in the 90 fiber layer. Bright white streaks are the compressive stress-strain curve consists of two nearly lin- ear portions with a decrease in slope occurring as the stress Fig. 13. Fracture surfaces of a N6IO//A specimen from billet B 14.(a) Multiple cracks initiate between the layers and (b) several short cracks(arrows)propagate into the 90 fiber layer and self-arrest
P.R. Jackson et al. / Materials Science and Engineering A 454–455 (2007) 590–601 599 Fig. 11. As-processed N601/M/A composite from billet B15. Thick matrix bands between the fibrous layers are clearly visible. (a) Optical micrograph and (b) backscatter SEM image showing matrix-rich areas between fibrous layers and regions devoid of matrix in the 90◦ fiber layer. Bright white streaks are the monazite coating. Fig. 12. As-processed N601/M/A composite from billet B14. Few isolated matrix-rich areas and good matrix infiltration into the 90◦ fiber layer are apparent. 5. Concluding remarks 5.1. Monotonic compression The compressive stress–strain behavior of N610/A and N610/M/A composites was investigated and the compressive properties measured at room and elevated temperatures. At temperatures ≤ 1100 ◦C, the compressive stress–strain behavior of the monazite-containing composite is nearly linear to failure; compressive modulus and strength magnitudes are close to the corresponding tensile values obtained in prior work [53]. Compressive behavior of the N610/A composite remains approximately linear to failure at 900 ◦C. However, at 1100 ◦C, the compressive stress–strain curve consists of two nearly linear portions with a decrease in slope occurring as the stress Fig. 13. Fracture surfaces of a N610/M/A specimen from billet B14. (a) Multiple cracks initiate between the layers and (b) several short cracks (arrows) propagate into the 90◦ fiber layer and self-arrest