E≈S ournal of the European Ceramic Society 20(2000)589-597 Effect of coating deposition temperature on monazite coated fiber R.S. Haya, E Boakye, M.D. Petry lir Force Research Laboratory, Materials Directorate, Wright Patterson Air Force Base, OH 45433, USA 4401 Davi lia Road, Dayton, OH 45432 USA Accepted 13 August 1999 Abstract Monazite (apoA) was continuously coated on 3M Nextel 720 fiber tows with an ethanolic precursor using hexadecane for immiscible liquid displacement Coating deposition temperatures were varied from 900 to 1300.C. Fibers coated at 900C were heat-treated up to 100 h at 1200C. Coated fibers were characterized by analytical TEM, and tensile strengths were measured by structure evolution was complex and may have involved recrystallization of large defective grains into smaller grains and ther subsequent growth of these grains, along with coarsening of porosity. After 100 h at 1200 C there was significant roughening of the coating-fiber interface, with facetting of alumina grains in the fiber and some lanthanum segregation to these facetted boundaries Spheroidization of thin coatings was also observed. Tensile strength of coated fiber decreased with increasing deposition tempera- ture and with time at temperature after deposition. Possible reasons for the strength decrease are discussed lished by elsevier Science ltd Keywords: Aluminosilicate fibers; Coating: Grain growth; Interfaces; LaPO4: Mechanical properties 1. ntroduction itu between the fiber and matrix during processing. Sol and solution precursors have been used to coat fiber Improvement of intermediate and high temperature tows with monazite, as has continuous chemical vapor properties of ceramic matrix composites (CMCs) deposition(CVD). 10,15,16 Limited experiments demon- requires an oxidation resistant alternative to carbon or strate crack deflection and fiber pullout in tensile tests of boron nitride fiber-matrix interfaces. -6 One possible composites with CVD monazite coated fiber tows. 10 For alternative is monazite(LaPO4). Monazite is refractory liquid-phase precursors, crack deflection and debonding (mp 2072C)and thermochemically stable with many have only been demonstrated on composites with low other common refractory oxides such as alumina and volume fractions of dip coated, large diameter(100 um) mullite. Crack deflection and fiber pushout experi- single crystal alumina(Saphikon )monofilaments. The ments,as well as some limited mechanical tests on thickness uniformity of sol and solution derived coat composites suggest that monazite bonds weakly with ings on filaments in fiber tows is often poor, although other oxides 8-10 Some similar results were found for stoichiometry is easily controlled. 6 In contrast, limited xenotime(YPO4) and scheelite(CaWO4). 2 Monazite data suggest that CVD coatings have good thickness containing ceramics were demonstrated to be machine- uniformity, but are sometimes off stoichiometry. 15, 16 able, 3 which implies significant plasticity from some Filament tensile strength is often degraded during combination of cleavage microcracking, twinning, or fiber coating, 6. 17-19 or during exposure to various envir dislocation glide. 4 onments.20), 2I Strength increases after coating or envir Use of monazite as a fiber-matrix interface in MCs onmental exposure are also known 22.23 Large requires that it either be coated on fibers or formed in differences in tensile strength have been observed between fibers coated with different monazite pre- 4 Corresponding author Tel: 1-937 255 9825: fax: 1 937 6564296. strength dependence on coating temperature were also E- mail address: hayrs(@ ml wl wpafb af. mil(RS. Hay) found. Excessive degradation in fiber tensile strength 0955-2219/00/S- see front matter. Published by Elsevier Science Ltd PII:S0955-2219(99)00257-5
Eect of coating deposition temperature on monazite coated ®ber R.S. Haya,*, E. Boakyeb, M.D. Petryb a Air Force Research Laboratory, Materials Directorate, Wright Patterson Air Force Base, OH 45433, USA bUES Inc., 4401 Dayton-Xenia Road, Dayton, OH 45432, USA Accepted 13 August 1999 Abstract Monazite (LaPO4) was continuously coated on 3M NextelTM 720 ®ber tows with an ethanolic precursor using hexadecane for immiscible liquid displacement. Coating deposition temperatures were varied from 900 to 1300C. Fibers coated at 900C were heat-treated up to 100 h at 1200C. Coated ®bers were characterized by analytical TEM, and tensile strengths were measured by single ®lament tensile tests. The monazite precursor was characterized by X-ray, DTA/TGA, and mass spectrometry. Microstructure evolution was complex and may have involved recrystallization of large defective grains into smaller grains and then subsequent growth of these grains, along with coarsening of porosity. After 100 h at 1200C there was signi®cant roughening of the coating±®ber interface, with facetting of alumina grains in the ®ber and some lanthanum segregation to these facetted boundaries. Spheroidization of thin coatings was also observed. Tensile strength of coated ®ber decreased with increasing deposition temperature and with time at temperature after deposition. Possible reasons for the strength decrease are discussed. Published by Elsevier Science Ltd. Keywords: Aluminosilicate ®bers; Coating; Grain growth; Interfaces; LaPO4; Mechanical properties 1. Introduction Improvement of intermediate and high temperature properties of ceramic matrix composites (CMCs) requires an oxidation resistant alternative to carbon or boron nitride ®ber-matrix interfaces.1±6 One possible alternative is monazite (LaPO4). Monazite is refractory (mp 2072C) and thermochemically stable with many other common refractory oxides such as alumina and mullite.7±9 Crack de¯ection and ®ber pushout experiments, as well as some limited mechanical tests on composites suggest that monazite bonds weakly with other oxides.8±10 Some similar results were found for xenotime (YPO4) 11 and scheelite (CaWO4).12 Monazite containing ceramics were demonstrated to be machineable,13 which implies signi®cant plasticity from some combination of cleavage microcracking, twinning, or dislocation glide.14 Use of monazite as a ®ber-matrix interface in CMCs requires that it either be coated on ®bers or formed in situ between the ®ber and matrix during processing. Sol and solution precursors have been used to coat ®ber tows with monazite,6 as has continuous chemical vapor deposition (CVD).10,15,16 Limited experiments demonstrate crack de¯ection and ®ber pullout in tensile tests of composites with CVD monazite coated ®ber tows.10 For liquid-phase precursors, crack de¯ection and debonding have only been demonstrated on composites with low volume fractions of dip coated, large diameter (100 mm) single crystal alumina (Saphikon1) mono®laments.8 The thickness uniformity of sol and solution derived coatings on ®laments in ®ber tows is often poor, although stoichiometry is easily controlled.6 In contrast, limited data suggest that CVD coatings have good thickness uniformity, but are sometimes o stoichiometry.15,16 Filament tensile strength is often degraded during ®ber coating,6,17±19 or during exposure to various environments.20,21 Strength increases after coating or environmental exposure are also known.22,23 Large dierences in tensile strength have been observed between ®bers coated with dierent monazite precursors.6 Preliminary indications of a strong tensile strength dependence on coating temperature were also found. Excessive degradation in ®ber tensile strength 0955-2219/00/$ - see front matter. Published by Elsevier Science Ltd. PII: S0955-2219(99)00257-5 Journal of the European Ceramic Society 20 (2000) 589±597 * Corresponding author Tel.:1-937 255 9825; fax: 1 937 656 4296. E-mail address: hayrs@ml.wl.wpafb.af.mil (R.S. Hay)
al / Journal of the European Ceramic Society 20(2000)589-59 may cause an otherwise functional coating to appear tem thin foils were made of coated fiber cross-sec- non-functional. Knowledge of the conditions under tions as described elsewhere 35.36 Thin foils were which degradation occurs is therefore important for observed in either a JEOL 2000 FX operating at 200 interface evaluation, besides the obvious importance to kv, or in a Phillips CM 200 FEG operating at 200 k composite strength and failure mechanisms. 24-26 Isolation Energy dispersive spectroscopy(EDS)measurements of the causes of strength degradation may lead to were done in the Phillips CM 200 FEG with a 5 nm spot improved coating methods that minimize this degradation. size and a windowless detector This work describes characteristics of 3M Nextel 720 Coated filament tensile strengths and Weibull modulii fiber coated with monazite between 900 and 1300C. were measured using a 2.54 cm gauge length with 75 Characteristics of the 900oC coated fiber after heat- tests. For controls, tensile strengths were also measured treatment for up to 100 h at 1200oC are also described. for filaments that had been heat-treated without coat Changes in composition and microstructure of the ing, and for filaments that had been passed through the coating and fiber were observed by TEM. Precursor and fiber coater under conditions that mimicked a coating coating evolution were monitored by X-ray, TEM, and run, but without coating deposition. Further details of DTA/TGA. Precursor gas evolution was observed by tensile strength measurement are presented elsewhere. 3 DTA/TGA and mass spectrometry. Strength of coated and uncoated fibers were measured by single filament tensile tests. Changes in coating characteristics with 3. Results temperature were compared with changes in the coated filament tensile strength. The results are analyzed and 3. 1. X-ray, DTA/TGA, and mass spectrometry possible degradation mechanisms are discussed Precursor heat-treated for l h at 1200 C has monazite y peaks and small La3PO7 peaks, indicating a slight phosphate deficiency from the monazite stoichiometry (ig. 1). A weak exotherm at about 450C and a weak Lanthanum nitrate and phosphorous pentoxide were endotherm at about 950C were the only significant dissolved in dry ethanol with the appropriate stoichio- DTA features(Fig. 2). About 4.5% of the precursor metry to form 50 g/l of monazite. 6 These solutions had a mass was lost between 100 and 550 C. Above 550C density of 0.84 g/cm and a viscosity of 1. 39 mPa- S, as mass loss was less rapid, but an additional 1.5%was measured in a Brookfield programmable rheometer still lost between 550 and 1500C. This 1.5% loss cor- ( model DV-IID) at a shear rate of 1/300. The precursor responds to a volume of gas that is at least 10 to 50 was characterized after a I h heat- treatment at 1200c times larger than the volume of monazite it evolved by X-ray diffraction in a Rigaku Rotaflex Dif- from. A slight increase in mass loss at around 950C fractometer. Differential thermal analysis (DTA)and correlates with the temperature of the weak DTA thermogravimetric analysis (TGA) were done in a endotherm. By mass spectrometry, only HO, Nor Netzsch STA-409 at 10C/min up to 1500 C after the CO, and N2o or Co? were observed to evolve in sig precursor was dried for I h at 140 C. Mass spectro- nificant quantities above 600 C(Fig. 3). Lack of amu metry of gases evolved from the precursor was done in a resolution precluded distinguishing N2 from Co(amu Balzers QMs 420 at 5C/min up to 1050C. The mea- 28)or N20 from CO2(amu 44). Methane(CH4) and surement was done in argon so measurement overlap CH3 evolution at 500C roughly corresponds to the with atmospheric gases could be eliminated. A more detailed description of the mass spectrometry equip- ment is given elsewhere. 27 3M Nextel 720 alumina-mullite fibers 28-30 were con- tinuously coated with the monazite precursor in a ver- tical coater using hexadecane for immiscible liquid displacement. A coating speed of 1. 4 cm/s in an air atmosphere was used for all experiments. The furnace hot zone was about 8 cm in length, and total furnace length was 30 cm. The fiber coating apparatus and pro- cedures are described in more detail elsewhere 6.31-34 The fibers were desized in air at 1000C and 2.8 cm/s before coating. Coating runs were done at 900, 1000, 1 100, 1200, and 1300 C. Some fibers coated at 900C Fig. 1. X-ray diffraction pattern of monazite from precursor heat were then heat-treated at 1200%C for 0.2. 2. 10. or 100 h treated for I h at 1200C. All peaks correspond to monazite except in a furnace with MoSi, heating elements those highlighted in gray, which correspond to La
may cause an otherwise functional coating to appear non-functional. Knowledge of the conditions under which degradation occurs is therefore important for interface evaluation, besides the obvious importance to composite strength and failure mechanisms.24±26 Isolation of the causes of strength degradation may lead to improved coating methods that minimize this degradation. This work describes characteristics of 3M Nextel 720 ®ber coated with monazite between 900 and 1300C. Characteristics of the 900C coated ®ber after heattreatment for up to 100 h at 1200C are also described. Changes in composition and microstructure of the coating and ®ber were observed by TEM. Precursor and coating evolution were monitored by X-ray, TEM, and DTA/TGA. Precursor gas evolution was observed by DTA/TGA and mass spectrometry. Strength of coated and uncoated ®bers were measured by single ®lament tensile tests. Changes in coating characteristics with temperature were compared with changes in the coated ®lament tensile strength. The results are analyzed and possible degradation mechanisms are discussed. 2. Experiments Lanthanum nitrate and phosphorous pentoxide were dissolved in dry ethanol with the appropriate stoichiometry to form 50 g/l of monazite.6 These solutions had a density of 0.84 g/cm3 and a viscosity of 1.39 mPa.s, as measured in a Brook®eld programmable rheometer (model DV-III) at a shear rate of 1/300. The precursor was characterized after a 1 h heat-treatment at 1200C by X-ray diraction in a Rigaku Rota¯ex Diffractometer. Dierential thermal analysis (DTA) and thermogravimetric analysis (TGA) were done in a Netzsch STA-409 at 10C/min up to 1500C after the precursor was dried for 1 h at 140C. Mass spectrometry of gases evolved from the precursor was done in a Balzers QMS 420 at 5C/min up to 1050C. The measurement was done in argon so measurement overlap with atmospheric gases could be eliminated. A more detailed description of the mass spectrometry equipment is given elsewhere.27 3M Nextel 720 alumina±mullite ®bers28±30 were continuously coated with the monazite precursor in a vertical coater using hexadecane for immiscible liquid displacement. A coating speed of 1.4 cm/s in an air atmosphere was used for all experiments. The furnace hot zone was about 8 cm in length, and total furnace length was 30 cm. The ®ber coating apparatus and procedures are described in more detail elsewhere.6,31±34 The ®bers were desized in air at 1000C and 2.8 cm/s before coating. Coating runs were done at 900, 1000, 1100, 1200, and 1300C. Some ®bers coated at 900C were then heat-treated at 1200C for 0.2, 2, 10, or 100 h in a furnace with MoSi2 heating elements. TEM thin foils were made of coated ®ber cross-sections as described elsewhere.35,36 Thin foils were observed in either a JEOL 2000 FX operating at 200 kV, or in a Phillips CM 200 FEG operating at 200 kV. Energy dispersive spectroscopy (EDS) measurements were done in the Phillips CM 200 FEG with a 5 nm spot size and a windowless detector. Coated ®lament tensile strengths and Weibull modulii were measured using a 2.54 cm gauge length with 75 tests. For controls, tensile strengths were also measured for ®laments that had been heat-treated without coating, and for ®laments that had been passed through the ®ber coater under conditions that mimicked a coating run, but without coating deposition. Further details of tensile strength measurement are presented elsewhere.37 3. Results 3.1. X-ray, DTA/TGA, and mass spectrometry Precursor heat-treated for 1 h at 1200C has monazite X-ray peaks and small La3PO7 peaks, indicating a slight phosphate de®ciency from the monazite stoichiometry (Fig. 1). A weak exotherm at about 450C and a weak endotherm at about 950C were the only signi®cant DTA features (Fig. 2). About 4.5% of the precursor mass was lost between 100 and 550C. Above 550C mass loss was less rapid, but an additional 1.5% was still lost between 550 and 1500C. This 1.5% loss corresponds to a volume of gas that is at least 10 to 50 times larger than the volume of monazite it evolved from. A slight increase in mass loss at around 950C correlates with the temperature of the weak DTA endotherm. By mass spectrometry, only H2O, N2 or CO, and N2O or CO2 were observed to evolve in signi®cant quantities above 600C (Fig. 3). Lack of amu resolution precluded distinguishing N2 from CO (amu 28) or N2O from CO2 (amu 44). Methane (CH4) and CH3 evolution at 500C roughly corresponds to the Fig. 1. X-ray diraction pattern of monazite from precursor heattreated for 1 h at 1200C. All peaks correspond to monazite except those highlighted in gray, which correspond to La3PO7. 590 R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597
R.S. Hay et al. Journal of the European Ceramic Society 20(2000)589-597 DTA TGA Fig. 2. DTA/TGA of monazite precursor Scale for DTA trace is not 100nm ictures of typical coating anomalies: (a) Debonded n, poorly crystallize coating;(c) Broken filament to filament bridge. 3.3 The most obvious change in the coatings as deposi tion temperature increased from 900 to 1300C was a 100 300500700 decrease in pore volume fraction and an increase in pore size from 8 to 25 nm in diameter(Fig. 5). There were Temperature(C) other more subtle differences in coating microstructure. For coatings deposited at 900C, SAD patterns from Fig 3. Mass spectrometry of gas evolution from monazite precurso the coatings were roughly consistent with monazite, but diffraction spots were diffuse and streaked, and a fine <50 nm)grain size was inferred from the patterns Diffraction contrast did not change abruptly across DTA exotherm at about 450 C. The mass spectrometry grains as the sample was tilted; instead, diffuse diffraction specimen was black after the measurement, suggesting contrast swept irregularly across the coating, and iden- incomplete oxidation of carbonaceous species in the tification of specific grains was difficult. Evidently the argon environment. monazite coating grains were highly defective, but the defects could not be identified. For coatings deposited 3.2. Coatings-general at 1300C, diffraction patterns were sharp, with individual grains easily identifiable. Defects other than intra Coating thickness and morphology was similar to that granular porosity were not observed. The coating grain observed previously for this precursor. Coating cover- size was often much greater than the coating thickness age was continuous. Coating bridges between filaments thin poorly crystallized coatings, and debonded or 3.3.1. 1200C heat-treatments buckled"coatings were occasionally observed(Fig. 4) Coatings deposited at 900C and then heat-treated at Significant place to place variations in pore volume 1200oC had somewhat different microstructures. Two fraction and size were also observed. Coating thicknes hours at 1200.C changed the defective, porous, and averaged about 80 nm, with variations between 30 and poorly crystallized monazite grains(Fig. 5)to small( 120 common. Although a small amount of La PO, was nm diameter), well crystallized monazite grains with a observed by X-ray, no evidence of any phase other than small amount of porosity(Fig. 6). One hundred hours monazite was found in fiber coatings by tEM at 1200C caused grain growth to 100 nm diameter
DTA exotherm at about 450C. The mass spectrometry specimen was black after the measurement, suggesting incomplete oxidation of carbonaceous species in the argon environment. 3.2. Coatings Ð general Coating thickness and morphology was similar to that observed previously for this precursor.6 Coating coverage was continuous. Coating bridges between ®laments, thin poorly crystallized coatings, and debonded or ``buckled'' coatings were occasionally observed (Fig. 4). Signi®cant place to place variations in pore volume fraction and size were also observed. Coating thickness averaged about 80 nm, with variations between 30 and 120 common. Although a small amount of La3PO7 was observed by X-ray, no evidence of any phase other than monazite was found in ®ber coatings by TEM. 3.3. Coating temperature The most obvious change in the coatings as deposition temperature increased from 900 to 1300C was a decrease in pore volume fraction and an increase in pore size from 8 to 25 nm in diameter (Fig. 5). There were other more subtle dierences in coating microstructure. For coatings deposited at 900C, SAD patterns from the coatings were roughly consistent with monazite, but diraction spots were diuse and streaked, and a ®ne (<50 nm) grain size was inferred from the patterns. Diraction contrast did not change abruptly across grains as the sample was tilted; instead, diuse diraction contrast swept irregularly across the coating, and identi®cation of speci®c grains was dicult. Evidently the monazite coating grains were highly defective, but the defects could not be identi®ed. For coatings deposited at 1300C, diraction patterns were sharp, with individual grains easily identi®able. Defects other than intragranular porosity were not observed. The coating grain size was often much greater than the coating thickness. 3.3.1. 1200C heat-treatments Coatings deposited at 900C and then heat-treated at 1200C had somewhat dierent microstructures. Two hours at 1200C changed the defective, porous, and poorly crystallized monazite grains (Fig. 5) to small (40 nm diameter), well crystallized monazite grains with a small amount of porosity (Fig. 6). One hundred hours at 1200C caused grain growth to 100 nm diameter Fig. 2. DTA/TGA of monazite precursor. Scale for DTA trace is not given. Fig. 3. Mass spectrometry of gas evolution from monazite precursor. Fig. 4. TEM pictures of typical coating anomalies: (a) Debonded coating; (b) Thin, poorly crystallize coating; (c) Broken ®lament to ®lament coating bridge. R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597 591
.S. Hay et al. / Journal of the European Ceramic Society 20(2000)589-59 with even less porosity. The most obvious features after 3.4. Tensile strength 100 h at 1200oC were spheroidization of coatings less than 50 nm thick(Fig. 7), roughening of the monazite Coated filament tensile strength decreased from 2 coating-fiber interface( Figs. 6-8), and facetting of alu- GPa28 with both increased deposition temperature and mina grains in the fiber and near the coating along basal with time at 1200 C after deposition at 900 C ( Table 1 planes(Figs. 6-8). Lanthanum segregation at facetted Fig. 10). Filaments coated at 1300C had average tensile alumina-mullite interfaces near the coating was mea- strengths of 1. 21 GPa. Filaments heat-treated for 100 h sured by analytical TEM(Fig 8). Phosphorous segre- at 1200 C after coating at 900oC had only 0.83 GPa gation was not found and lanthanum was not found in average tensile strength. The actual average may have the interior of alumina grains. Wetting thin films were been slightly lower, because some filaments were too not observed along monazite-alumina, monazite-mul- weak to mount in grips for tensile testing, and their lite, or alumina-mullite interphase boundaries, or along contribution to the average would certainly have made the corresponding triple junctions(Fig. 9) it lower. Filaments coated at 1200C were slightly stronger than those coated at 1100 C, as were filaments 900℃C heat-treated for 2 h at 1200oC compared to 12 min at 1200.C. These observations were anomalous to the overall trend; it is possible that they are related to tran- sient flaw healing, but more data is required to sub stantiate this Control experiments suggest that much of the strength drop may be unrelated to the presence of a coating. Filaments heat-treated for 100 h at 1200C without a coating had only 1.04 GPa tensile strength (Table 1, Fig. 10). Similarly, filaments passed through the coating furnace at 1300C and 1200oC had only 1.5 GPa tensile strength, although filaments passed through water, water/HNO3, ethanol, and water/ octanol were stronger (Table I, Fig. 10). The strength drop between fibers coated at 1000 and 1100C was steeper than others, and may correlate with the slight 1300℃C Fig. 5. TEM pictures of coating deposited at(a)900oC,(b)I (c)1100 C, (d)1300C Simulated and observed diffraction patt 50 nm 900C coating and observed diffraction pattern from a 1300C are to the rig 200nm 2000mm Fig. 6. TEM pictures of coat posited at 900C and later heat- treated at 1200C for 2 h or 10 Fig. 7. TEM pictures of spheroidized coatings
with even less porosity. The most obvious features after 100 h at 1200C were spheroidization of coatings less than 50 nm thick (Fig. 7), roughening of the monazite coating±®ber interface (Figs. 6±8), and facetting of alumina grains in the ®ber and near the coating along basal planes (Figs. 6±8). Lanthanum segregation at facetted alumina±mullite interfaces near the coating was measured by analytical TEM (Fig. 8). Phosphorous segregation was not found, and lanthanum was not found in the interior of alumina grains. Wetting thin ®lms were not observed along monazite±alumina, monazite±mullite, or alumina±mullite interphase boundaries, or along the corresponding triple junctions (Fig. 9). 3.4. Tensile strength Coated ®lament tensile strength decreased from 2 GPa28 with both increased deposition temperature and with time at 1200C after deposition at 900C (Table 1, Fig. 10). Filaments coated at 1300C had average tensile strengths of 1.21 GPa. Filaments heat-treated for 100 h at 1200C after coating at 900C had only 0.83 GPa average tensile strength. The actual average may have been slightly lower, because some ®laments were too weak to mount in grips for tensile testing, and their contribution to the average would certainly have made it lower. Filaments coated at 1200C were slightly stronger than those coated at 1100C, as were ®laments heat-treated for 2 h at 1200C compared to 12 min at 1200C. These observations were anomalous to the overall trend; it is possible that they are related to transient ¯aw healing, but more data is required to substantiate this. Control experiments suggest that much of the strength drop may be unrelated to the presence of a coating. Filaments heat-treated for 100 h at 1200C without a coating had only 1.04 GPa tensile strength (Table 1, Fig. 10). Similarly, ®laments passed through the coating furnace at 1300C and 1200C had only 1.5 GPa tensile strength, although ®laments passed through water, water/HNO3, ethanol, and water/1- octanol were stronger (Table 1, Fig. 10). The strength drop between ®bers coated at 1000 and 1100C was steeper than others, and may correlate with the slight Fig. 5. TEM pictures of coating deposited at (a) 900C, (b) 1000C, (c) 1100C, (d) 1300C. Simulated and observed diraction pattern of 900C coating and observed diraction pattern from a 1300C coating are to the right. Fig. 6. TEM pictures of coating deposited at 900C and later heattreated at 1200C for 2 h or 100 h. Fig. 7. TEM pictures of spheroidized coatings. 592 R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597
.S. Hay et al. /Journal of the European Ceramic Society 20 (2000)589-597 2 nm t山Hd WAlNi w Mus Kev(x100) Fig 8. TEM picture of an area around a monazite coating -fiber interface analyzed by TEM. The EDS spectra of spots 1, 2, and 3 on the image are shown to the right. Spot 3 at an alumina-mullite interphase boundary shows trace lanthanum. increase in coating mass loss observed at about 950.C. Although the temperature interval does not exactly match, the heating rate was higher during coating than during the DTA/TGA measurements, so the corre- ponding mass loss during coating should be at a higher temperature. However, there are alternative explana tions for this strength drop that will be discussed in the next section Coating fibers with monazite at 900 C using this par- ticular precursor preserves most of the original tensile strength. However, this would not eliminate tensile strength degradation if composites with these coated fibers were processed or used at or above 1200.C, as shown by tensile strengths of 900C coated fiber heat treated at 1200C(Fig. 10). However, control experiments on uncoated fiber heat -treated at 1200 c have tensile strengths nearly as low as coated fibers. Beneath 1300C grain growth is insignificant in Nextel 720.38 The sensi Fig. 9. High resolution TEM micrograph of a monazite-alumina- tivity of tensile strength to heat-treatment or coating ng mullite triple junction at a coating-fiber interface. No evidence of a atmosphere (Table 1)suggests that degradation is most wetting phase at the interface or triple junction was found
increase in coating mass loss observed at about 950C. Although the temperature interval does not exactly match, the heating rate was higher during coating than during the DTA/TGA measurements, so the corresponding mass loss during coating should be at a higher temperature. However, there are alternative explanations for this strength drop that will be discussed in the next section. 4. Discussion Coating ®bers with monazite at 900C using this particular precursor preserves most of the original tensile strength. However, this would not eliminate tensile strength degradation if composites with these coated ®bers were processed or used at or above 1200C, as shown by tensile strengths of 900C coated ®ber heattreated at 1200C (Fig. 10). However, control experiments on uncoated ®ber heat-treated at 1200C have tensile strengths nearly as low as coated ®bers. Beneath 1300C grain growth is insigni®cant in Nextel 720.38 The sensitivity of tensile strength to heat-treatment or coating atmosphere (Table 1) suggests that degradation is most Fig. 8. TEM picture of an area around a monazite coating Ð ®ber interface analyzed by TEM. The EDS spectra of spots 1, 2, and 3 on the image are shown to the right. Spot 3 at an alumina±mullite interphase boundary shows trace lanthanum. Fig. 9. High resolution TEM micrograph of a monazite±alumina± mullite triple junction at a coating±®ber interface. No evidence of a wetting phase at the interface or triple junction was found. R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597 593
594 R.S. Hay et al./ Journal of the European Ceramic Society 20(2000)589-597 Table I Tensile strength measurements Deposition Environment Heat-treatment Heat-treatment Tensile Weibull temperature temperature time modulus (°C) (° (h) (GPa) 123 1000 Morante 6.85 Monazite Argon Monazite Air 0123 nnnn aaumuu rrrr Water/20% HNO, 111 Control Water/20% HNO/l-octanol 5.11 likely environmental assisted surface crack growth. Pre- heat-treatment atmosphere, and one from the monazite vious work showed that coating thickness did not affect precursor. It is not clear whether these effects are inde- tensile strength, and that different monazite precursors pendent or synergistic If degradation is predominantly used for fiber coating under identical conditions had environmental, hermetic matrices that seal fibers from much different tensile strengths, which was consistent the environment could preserve fiber strength in a with environmental effects specific to each precursor. CMC. For heat-treated fibers, atmospheric moisture or Flaws in weakly bonded fiber coatings should not func- silica from MoSiz heating elements are possible corro tion as fiber flaws, 9 which is consistent with lack of a sive chemical species coating thickness effect and the known debonding At 1200 C segregation of lanthanum from monazite properties of monazite. to interfaces in the fiber(Fig 8)may also cause strength Observations made here and elsewhere suggest that degradation. Lanthanum doped alumina has facetted there are separate effects on fiber strength, one from the and elongated grains. 3 Lanthanum doping also retards sintering and creep in alumina.39,40 Effects on strength have, to our knowl- 14 12 alumina basal planes causes easier intergranular fracture 3 along those planes analogous to cleavage in p-alumina and magnetoplumbite, 41.42 then significant weakening can be expected. We note that X-ray diffraction suggests our precursor was lanthanum rich(although TEM did not corroborate this). Therefore lanthanum activity may have been buffered to a relatively high value by la3 PO7 LaPO4, instead of a lower value fixed by LaPO4 LaP3Oo More measurements of lanthanum segregation as a function of temperature, time, and lanthanum Logit(s))at 1200 C activity at the two different buffers needs to be measured and a more definitive correlation with fiber tensile strength should be made Identification of the chemical species causing strength 100011 1200 l300 degradation remains problematic. Mass spectrometry sts that only N, or CO, H,O, and N,O or CO gases were evolved by the precursor above 700C(Fig. Fig 10. Tensile strength of coated fibers and control experiments vs coating temperature or heat-treatment time at 1200.C. Log(o) scale on 3). The black sample and argon atmosphere suggest that x-axis refers to solid line connecting squares. Coating temperature limited oxidation of carbon and partial reduction of the le on x-axis refers to dashed line connecting circles. Numbers for nitrates from the lanthanum precursor N2O control experiments are given in Table I more likely. The measured H20 could also be an artifact
likely environmental assisted surface crack growth. Previous work showed that coating thickness did not aect tensile strength, and that dierent monazite precursors used for ®ber coating under identical conditions had much dierent tensile strengths, which was consistent with environmental eects speci®c to each precursor.6 Flaws in weakly bonded ®ber coatings should not function as ®ber ¯aws,19 which is consistent with lack of a coating thickness eect and the known debonding properties of monazite. Observations made here and elsewhere6 suggest that there are separate eects on ®ber strength, one from the heat-treatment atmosphere, and one from the monazite precursor. It is not clear whether these eects are independent or synergistic. If degradation is predominantly environmental, hermetic matrices that seal ®bers from the environment could preserve ®ber strength in a CMC. For heat-treated ®bers, atmospheric moisture or silica from MoSi2 heating elements are possible corrosive chemical species. At 1200C segregation of lanthanum from monazite to interfaces in the ®ber (Fig. 8) may also cause strength degradation. Lanthanum doped alumina has facetted and elongated grains.39 Lanthanum doping also retards sintering and creep in alumina.39,40 Eects on strength have, to our knowledge, not been reported. If lanthanum segregation to alumina basal planes causes easier intergranular fracture along those planes analogous to cleavage in b-alumina and magnetoplumbites,41,42 then signi®cant weakening can be expected. We note that X-ray diraction suggests our precursor was lanthanum rich (although TEM did not corroborate this). Therefore lanthanum activity may have been buered to a relatively high value by La3PO7/ LaPO4, instead of a lower value ®xed by LaPO4/ LaP3O9. More measurements of lanthanum segregation as a function of temperature, time, and lanthanum activity at the two dierent buers needs to be measured, and a more de®nitive correlation with ®ber tensile strength should be made. Identi®cation of the chemical species causing strength degradation remains problematic. Mass spectrometry suggests that only N2 or CO, H2O, and N2O or CO2 gases were evolved by the precursor above 700C (Fig. 3). The black sample and argon atmosphere suggest that limited oxidation of carbon and partial reduction of the nitrates from the lanthanum precursor to N2 and N2O is more likely. The measured H2O could also be an artifact Table 1 Tensile strength measurements No. Deposition temperature ( C) Coating Environment Heat-treatment temperature ( C) Heat-treatment time (h) Tensile strength (GPa) Weibull modulus 1 900 Monazite Air 1.95 6.85 2 1000 Monazite Air 1.78 5.66 3 1100 Monazite Air 1.32 4.95 4 1200 Monazite Air 1.39 4.09 5 1300 Monazite Air 1.21 5.09 5a 1300 Monazite Argon 1.24 5.07 6 900 Monazite Air 1200 0.02 1.14 3.77 7 900 Monazite Air 1200 2 1.21 4.00 8 900 Monazite Air 1200 100 0.83 3.43 9 Control Air 1200 100 1.04 4.97 10 1200 Control Air 1.50 4.79 11 1300 Control Air 1.51 4.40 12 1300 Control Water 1.77 4.23 13 1300 Control Water/20% HNO3 1.59 4.08 14 1300 Control Water/20% HNO3/l-octanol 1.85 5.11 Fig. 10. Tensile strength of coated ®bers and control experiments vs. coating temperature or heat-treatment time at 1200C. Log(t) scale on x-axis refers to solid line connecting squares. Coating temperature scale on x-axis refers to dashed line connecting circles. Numbers for control experiments are given in Table 1. 594 R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597
R.S. Hay et al. Journal of the European Ceramic Society 20(2000)589-597 of wet tubing in the instrument since background runs The spheroidization of 50 nm thick coatings at consistently found this gas despite attempts to run in a 1200C is another concern(Fig. 7). Previous analysis dry environment. 43 This leaves N2 or N2O as the most and observations of the Rayleigh instability of poly- likely products of high temperature degassing, with crystalline thin films suggests that monazite coatings on H,O, CO, CO2, and undetected trace gases as additional other oxides may be prone to this instability. 44 45 Fc possibilities. If H2O or N2 were the principal high tem- example, the high monazite-alumina interphase bound- perature corrosive species, control experiments in air ary energy causes a large equilibrium contact angle B (72%nitrogen) and with water substituted for coating between monazite and alumina, which in turn allows precursor could be expected to degrade fiber strength in spheroidization of polycrystalline coatings at a smaller a manner similar to the coated fibers. Unfortunately, grain size(D) for a given coating thickness(a). Spher this was not the case. Water"coated"fibers had some oidization is predicted when: 44 of the highest strengths of any control experiment(Fig 10). Mass spectrometry of coating precursor was done 3 sin e in argon, but fiber coating was done in air, so the gases a 2-3cos 0+c0s38 evolved during coating might be more highly oxidized forms of carbon and nitrogen than those found by mass A more complex expression in later work5 yields simi- spectrometry. However, fibers coated in argon with this lar results. The grain size must be less than twice the precursor at 1300 C had 1.24 GPa tensile strength, a coating thickness if the film is to be stable with respect negligible increase over those coated in air to spheroidization at high 8. Coatings discussed in this The partial pressure of the evolved gases is another paper(Figs. 5-7)had unusually large grain sizes in concern. Partial pressures measured by mass spectro- comparison to coatings made from other precursors, so metry are quite low, because they are diluted by the they may be particularly prone to spheroidization argon atmosphere used for measurement. However, the Spheroidization kinetics of these coatings should partial pressure of a gas evolved in a closed coating pore depend on the surface diffusion coefficients of mon- may be much higher than atmospheric and depend on azite 46 However, when coated fibers are incorporated in the bursting strength of the pore and the driving force a dense matrix, the rate determining process must for gas formation. If these gas filled pores are adjacent change to interphase boundary diffusion control, which to the fiber surface, fiber surface flaws can be exposed to is generally slower than surface diffusion. 47 It is also high pressure gas from precursor decomposition. As the necessary to transport either matrix or fiber material to pore volume fraction decreases with coating tempera- accommodate monazite spheroidization, which could Ire(Fig. 5), it becomes more likely that the porosity is further retard spheroidization if diffusive mass transpor closed and that gas decomposition products are sealed in in the fiber or matrix is slow compared to monazite the coating. The large tensile strength drop between the 1000 and 1100 C coated fibers( Fig. 10)could, therefore, be related to densification and coarsening of coating por- 5. Summary and conclusions osity and the consequent transition from an open to hermetic coating. However, if the main cause of strength Lowering the deposition temperature of monazite degradation was environmental effects independent of the from 1300 to 900C for an ethanolic monazite precursor presence of a coating, such as those measured by control causes coatings from this particular pr to hay experiments, then it might also be argued that a hermetic higher pore volume fraction, smaller pore size, and coating should seal the fiber from this environment, and defectively crystallized grains with poorly defined grain the fiber should be stronger. More information is neces- size. The porosity is mostly eliminated and the grains sary to confidently explain these results recrystallize during heat-treatment at 1200.C. Thin As deposited coatings were porous, particularly at coatings spheroidized, which may be a problem with low temperature, but the porosity formed during coatings like monazite that have high substrate inter- deposition at 900 C was not stable at 1200 C. These phase-boundary energy. Spheroidization kinetics may coatings recrystallized to a small grain size at 1200oC. be retarded by incorporation of the coating in a dense Coatings deposited at 1300.C had a large grain size with matrix, and spheroidization can be prevented by stabi somewhat larger intragranular porosity. Heat-treat- lization of a small coating grain size in comparison to ments at 1200C were not done on these 1300C coated coating thick fibers. The large grain size and consequent lack of grain Fibers coated at 900oc did not lose much tensile boundary diffusion pathways might cause this intra- strength, but fibers coated at higher temperatures were granular porosity to be more resistant to ening more severely degraded. Control experiments that than the porosity present at 900C. Further experiments mimicked the deposition process in some cases were are necessary to establish the T-t path dependence of nearly as severely degraded as coated fibers. The high coating microstructure. strengths of fibers coated at low temperatures was not
of wet tubing in the instrument, since background runs consistently found this gas despite attempts to run in a dry environment.43 This leaves N2 or N2O as the most likely products of high temperature degassing, with H2O, CO, CO2, and undetected trace gases as additional possibilities. If H2O or N2 were the principal high temperature corrosive species, control experiments in air (72% nitrogen) and with water substituted for coating precursor could be expected to degrade ®ber strength in a manner similar to the coated ®bers. Unfortunately, this was not the case. Water ``coated'' ®bers had some of the highest strengths of any control experiment (Fig. 10). Mass spectrometry of coating precursor was done in argon, but ®ber coating was done in air, so the gases evolved during coating might be more highly oxidized forms of carbon and nitrogen than those found by mass spectrometry. However, ®bers coated in argon with this precursor at 1300C had 1.24 GPa tensile strength,6 a negligible increase over those coated in air. The partial pressure of the evolved gases is another concern. Partial pressures measured by mass spectrometry are quite low, because they are diluted by the argon atmosphere used for measurement. However, the partial pressure of a gas evolved in a closed coating pore may be much higher than atmospheric and depend on the bursting strength of the pore and the driving force for gas formation. If these gas ®lled pores are adjacent to the ®ber surface, ®ber surface ¯aws can be exposed to high pressure gas from precursor decomposition. As the pore volume fraction decreases with coating temperature (Fig. 5), it becomes more likely that the porosity is closed and that gas decomposition products are sealed in the coating. The large tensile strength drop between the 1000 and 1100C coated ®bers (Fig. 10) could, therefore, be related to densi®cation and coarsening of coating porosity and the consequent transition from an open to a hermetic coating. However, if the main cause of strength degradation was environmental eects independent of the presence of a coating, such as those measured by control experiments, then it might also be argued that a hermetic coating should seal the ®ber from this environment, and the ®ber should be stronger. More information is necessary to con®dently explain these results. As deposited coatings were porous, particularly at low temperature, but the porosity formed during deposition at 900C was not stable at 1200C. These coatings recrystallized to a small grain size at 1200C. Coatings deposited at 1300C had a large grain size with somewhat larger intragranular porosity. Heat-treatments at 1200C were not done on these 1300C coated ®bers. The large grain size and consequent lack of grain boundary diusion pathways might cause this intragranular porosity to be more resistant to coarsening than the porosity present at 900C. Further experiments are necessary to establish the T±t path dependence of coating microstructure. The spheroidization of 50 nm thick coatings at 1200C is another concern (Fig. 7). Previous analysis and observations of the Rayleigh instability of polycrystalline thin ®lms suggests that monazite coatings on other oxides may be prone to this instability.44,45 For example, the high monazite±alumina interphase boundary energy causes a large equilibrium contact angle between monazite and alumina, which in turn allows spheroidization of polycrystalline coatings at a smaller grain size (D) for a given coating thickness (a). Spheroidization is predicted when:44 D a 5 3 sin3 2 ÿ 3 cos cos3 1 A more complex expression in later work45 yields similar results. The grain size must be less than twice the coating thickness if the ®lm is to be stable with respect to spheroidization at high . Coatings discussed in this paper (Figs. 5±7) had unusually large grain sizes in comparison to coatings made from other precursors,6 so they may be particularly prone to spheroidization. Spheroidization kinetics of these coatings should depend on the surface diusion coecients of monazite.46 However, when coated ®bers are incorporated in a dense matrix, the rate determining process must change to interphase boundary diusion control, which is generally slower than surface diusion.47 It is also necessary to transport either matrix or ®ber material to accommodate monazite spheroidization, which could further retard spheroidization if diusive mass transport in the ®ber or matrix is slow compared to monazite. 5. Summary and conclusions Lowering the deposition temperature of monazite from 1300 to 900C for an ethanolic monazite precursor causes coatings from this particular precursor to have a higher pore volume fraction, smaller pore size, and defectively crystallized grains with poorly de®ned grain size. The porosity is mostly eliminated and the grains recrystallize during heat-treatment at 1200C. Thin coatings spheroidized, which may be a problem with coatings like monazite that have high substrate interphase-boundary energy. Spheroidization kinetics may be retarded by incorporation of the coating in a dense matrix, and spheroidization can be prevented by stabilization of a small coating grain size in comparison to coating thickness. Fibers coated at 900C did not lose much tensile strength, but ®bers coated at higher temperatures were more severely degraded. Control experiments that mimicked the deposition process in some cases were nearly as severely degraded as coated ®bers. The high strengths of ®bers coated at low temperatures was not R.S. Hay et al. / Journal of the European Ceramic Society 20 (2000) 589±597 595
R.S. Hay et al. Journal of the European Ceramic Society 20(2000)589-597 retained if those fibers were later heat-treated to higher 5. Sun, E. Y, Lin. H.-T. and Brennan, J.J., Intermediate-tempera temperatures. Again, degradation during heat-treatment ture environmental effects on boron nitride- coated silicon ca was nearly as severe in control experiments without a bide-fiber-reinforced glass-ceramic composites. J. Am. Ceram coating. Soc.,1997,80(31),609614 6. Boakye, E, Hay. R.S. and Petry. M. D. Continuous coating of Strength degradation at temperatures beneath 1200C oxide fiber tows using liquid precursors: monazite coatings on could not be caused by microstructural changes in the Nextel720.J.Am. Ceran.Soc,1999,82(9).232l-2331 fiber because there is no grain growth at those tem- 7. Morgan. P. E. D. and Marshall, D. B. Functional interfaces for peratures. An environmental cause was suggested to be oxide oxide composites. Mat Sci. Eng, 1993, A162 15-25. more likely. The chemical species in the environment or 8. Morgan, P.E. D. and Marshall, D. B, Ceramic composites of from coating degassing responsible for this degradation 9. Morgan, P. E D, Marshall, D. B and Housley, R.M,High were not successfully isolated. It is also possible that mperature stability of monazite-alumina composites. J. Ma strength degradation from the monazite coating was Sci.Eng.195,A195,215-222 related to sealing of corrosive gas decomposition pro- 10. Kanazawa, C. Johnson, S M. and Porter. J.R. Monazite coating ducts in the coating at relatively high pressure and promotes fiber pullout. J. Am. Ceram. Soc., 1997, 80(7), Back Cover I1. Kuo, D -H. and Riven, W. M, Characterization of yttrium therefore, related to transformation from open to closed phosphate and a yttrium phosphate/yttrium aluminate laminate. porosity in the coating Am. Ceran.Soc,1995,78(111),3121-3124 Lanthanum segregated from monazite to alumina 12. Goettler. R. W, Sambasivan, S. and Dravid, V.P., Isotropic mullite interphase boundaries at 1200C. Segregation omplex oxides as fiber coatings for oxide-oxide CFCC. Ceram. was associated with roughening of the fiber-coating Eng.Sci.Proc.1977,18(),279-286. 13. Davis, J. B, Marshall, D. B, Housley, R. M. and Morgan, P. E nterface and facetting of alumina grains in the fiber These microstructural changes may be at least partly responsible for low tensile strengths, but again it must 14. Hay, R.s. Davis, J.B., Marshall, D B and Morgan, P.E. D be emphasized that these tensile strengths were not much lower than those for control experiments on uncoated fibers More information on the kinetics and Hunt, A. T, Combustion chemical vapor deposition(CCVD) of LaPOa monazite and beta-alumina on alumina fibers for cerar lanthanum activity dependence for segregation are matrix composites. Mater. Sci. Eng. 4, 1998. 244, 91-96 desirable. It is clear that the tensile strength of Nextel 16. Chayka, P, personal communication. 720, with or without coating, is sensitive to processing 17. Hay, R.S., Petry, M. D. and Boakye, E.Fiber ure. furthe atings from sols and solutions Presented at the 20th Annual Conference on Composi investigation of this sensitivity is necessary to optimize processing conditions for this fiber, and to design new 18. Helmer, T, Peterlik, H and Kromp, K,Coating of carbon fibers with improved properties bers- the strength of the fibers. J. Am. Ceram. Soc., 1995 19. Parthasarathy, T.A., Folsom, C. A. and Zawada, L.P., bined effects of exposure to salt (NaCn) water and oxidatie Acknowledgements the strength of uncoated and BN- coated nicalon fibers. Ceran.Soc.,l998,81(7),1812-1818 Te thank S. Sambasivan for the monazite precursor, 20. Trumbauer, E.R. Hellmann, J.R Shelleman, D L and Ko Wilson of 3M for the Nextel 720 fiber. K. Von D. A. Effect of cleaning and abrasion induced damage on the Lehmden for thin section preparation and tensile testing. Weibull strength distribution of sapphire fiber. J. Am. Cer. Soc. 1994,77(8),2017-2024. J. Jones and P. Jero for mass spectrometry measure- 21. Inniss, D, Zhong, Q and Kurkjian, C R, Chemically corroded ments, and K. Keller for manuscript review stine silica fibers: blunt or sharp flaws. J. Am. Ceram. Soc., 1993,7612),3173-3177 22. Chang, x and Du, Y, Electrolytic treatment of continuous CVD References silicon carbide fibers. J. Am. Ceram. Soc., 1997, 80(10), 2754-2756. 23. Fernando, J.A Chawla, K. K, Ferber, M. K and Coffey, D, Effect of boron nitride coating on the tensile strength of Nextel 1. Mah. T. Mendiratta, M. G, Katz. A. P. Ruh, R. and Maz. 480fber. Mater.Sci.Eng.A,1992,154,103-108. diyasni, K. S, High-temperature mechanical behavior of fiber 24. Curtin. W. A. Strength versus gauge length in ceramic-matrix reinforced glass-ceramic-matrix composites. J. Am. Cer. Soc mposites. J. Am. Cer. Soc., 1994, 77(4). 1072-1074 1985,68(9),C248. 25. Curtin, w.A., In situ fiber strengths in ceramic-matri 2. Naslain, R, Dugne, 0, Guette, A, Sevely, J, Brosse, C. sites from fracture mirrors. J. Am. Cera. Soc.. 1994 4m4 Rocher. p and cotteret.. boron nitride i matrix composites. J. Am. Cer. Soc., 1991, 74(101), 2482-2488 26. Curtin. w.A. Ahn. B. K. and Takeda. N. d 3. Prewo. K. M. Brennan, J. J. and Starrett, S. Silicon carbide tough stress-strain behavior in unidirection amic matrIx opposites. Acta Mater., 1998, 46(10). 340 rated temperature. J. Mat. Sci., 1989, 24. 1373 27. Jones. J G.. Busbee. J. D. Jero. P. D. Kent. D. J. and 4. Zawada. L. P. and Wetherhold. R. C. The effects of thermal D. C. In situ control of interface coatings on fibers using atigue on a SIC fiber/aluminosilicate glass composite. J. Mat In: Composites and Functionally Graded Materials, ed. Sri Sci,1991,26,648-654 T.S.et al., ASME, Vol. 80, 1997, pp 379-384
retained if those ®bers were later heat-treated to higher temperatures. Again, degradation during heat-treatment was nearly as severe in control experiments without a coating. Strength degradation at temperatures beneath 1200C could not be caused by microstructural changes in the ®ber because there is no grain growth at those temperatures. An environmental cause was suggested to be more likely. The chemical species in the environment or from coating degassing responsible for this degradation were not successfully isolated. It is also possible that strength degradation from the monazite coating was related to sealing of corrosive gas decomposition products in the coating at relatively high pressure and, therefore, related to transformation from open to closed porosity in the coating. Lanthanum segregated from monazite to aluminamullite interphase boundaries at 1200C. Segregation was associated with roughening of the ®ber-coating interface and facetting of alumina grains in the ®ber. These microstructural changes may be at least partly responsible for low tensile strengths, but again it must be emphasized that these tensile strengths were not much lower than those for control experiments on uncoated ®bers. More information on the kinetics and lanthanum activity dependence for segregation are desirable. It is clear that the tensile strength of Nextel 720, with or without coating, is sensitive to processing conditions, environment, and temperature. Further investigation of this sensitivity is necessary to optimize processing conditions for this ®ber, and to design new ®bers with improved properties. Acknowledgements We thank S. Sambasivan for the monazite precursor, D. Wilson of 3M for the Nextel 720 ®ber, K. Von Lehmden for thin section preparation and tensile testing, J. Jones and P. Jero for mass spectrometry measurements, and K. Keller for manuscript review. References 1. Mah, T., Mendiratta, M. G., Katz, A. P., Ruh, R. and Mazdiyasni, K. S., High-temperature mechanical behavior of ®ber reinforced glass±ceramic±matrix composites. J. Am. Cer. Soc., 1985, 68(9), C248. 2. Naslain, R., Dugne, 0., Guette, A., Sevely, J., Brosse, C. R., Rocher, J. P. and Cotteret, J., Boron nitride interphase in ceramic matrix composites. J. Am. Cer. Soc., 1991, 74(101), 2482±2488. 3. Prewo, K. M., Brennan, J. J. and Starrett, S., Silicon carbide ®ber-reinforced glass ceramic composite tensile behavior at elevated temperature. J. Mat. Sci., 1989, 24, 1373. 4. Zawada, L. P. and Wetherhold, R. C., The eects of thermal fatigue on a SIC ®ber/aluminosilicate glass composite. J. Mat. Sci., 1991, 26, 648±654. 5. Sun, E. Y., Lin, H.-T. and Brennan, J. J., Intermediate-temperature environmental eects on boron nitride-coated silicon carbide-®ber-reinforced glass±ceramic composites. J. Am. Ceram. Soc., 1997, 80(31), 609±614. 6. Boakye, E., Hay, R. S. and Petry, M. D. Continuous coating of oxide ®ber tows using liquid precursors: monazite coatings on Nextel 720. J. Am. Ceram. Soc., 1999, 82(9), 2321±2331. 7. Morgan, P. E. D. and Marshall, D. B., Functional interfaces for oxide/oxide composites. Mat. Sci. Eng., 1993, A162, 15±25. 8. Morgan, P. E. D. and Marshall, D. B., Ceramic composites of monazite and alumina. J. Am. Cer. Soc., 1995, 78(61), 1553±1563. 9. Morgan, P. E. D., Marshall, D. B. and Housley, R. M., High temperature stability of monazite±alumina composites. J. Mat. Sci. Eng., 1995, A195, 215±222. 10. Kanazawa, C., Johnson, S. M. and Porter, J. R., Monazite coating promotes ®ber pullout. J. Am. Ceram. Soc., 1997, 80(7), Back Cover. 11. Kuo, D.-H. and Kriven, W. M., Characterization of yttrium phosphate and a yttrium phosphate/yttrium aluminate laminate. J. Am. Ceram. Soc., 1995, 78(111), 3121±3124. 12. Goettler, R. W., Sambasivan, S. and Dravid, V. P., Isotropic complex oxides as ®ber coatings for oxide±oxide CFCC. Ceram. Eng. Sci. Proc., 1977, 18(3), 279±286. 13. Davis, J. B., Marshall, D. B., Housley, R. M. and Morgan, P. E. D., Machinable ceramics containing rare-earth phosphates. J. Am. Ceram. Soc., 1998, 81(8), 2169±2175. 14. Hay, R. S., Davis, J. B., Marshall, D. B. and Morgan, P. E. D., Presented at the 1998 Annual Meeting of the Am. Ceram. Soc. 15. Hwang, T. J., Hendrick, M. R., Shao, H., Hornis, H. G. and Hunt, A. T., Combustion chemical vapor deposition (CCVD) of LaPO4 monazite and beta-alumina on alumina ®bers for ceramic matrix composites. Mater. Sci. Eng. A, 1998, 244, 91±96. 16. Chayka, P, personal communication. 17. Hay, R. S., Petry, M. D. and Boakye, E. Fiber strength with coatings from sols and solutions. Presented at the 20th Annual Conference on Composites and Advanced Ceramics, Cocoa Beach, 1996. 18. Helmer, T., Peterlik, H. and Kromp, K., Coating of carbon ®bers Ð the strength of the ®bers. J. Am. Ceram. Soc., 1995, 78(1), 133±136. 19. Parthasarathy, T. A., Folsom, C. A. and Zawada, L. 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