J. An. Ceram Soc..87[101967-197602004) ournal Zirconia-Silica-Carbon Coatings on Ceramic Fibers Emmanuel E Boakye, *T Randall S. Hay, *f M. Dennis Petry, and Triplicane A Parthasarathy* f UES, Inc, Dayton, Ohio 45432 Materials and Manufacturing Directorate, Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio 45433 Precursors for zircon-carbon mixtures were made to coat ZrSiO4 forms from ZrO, and Sio, at -1400-1600C.,For fibers for ceramic-matrix composites. Precursors were char- NextelM 720 fiber coatings, fib er strength degra acterized using XRD, TGA, and DTA. Zircon formed from with grain growth in the fiber requires ZrSiO4 to form at vanadium-or lithium-doped precursors after heat treatments 1350%C are neces- fiber tensile strengths were measured. Although zircon formed sary for high purity. 3 Seeding and fluxes can significantly in powders, only tetragonal-zirconia-silica mixed phases formed in fiber coatings at 1200.C in air. Loss of vanadium air has been studied extensively, 239. 0.45-47and there is one study oxide flux to the fibers may have caused the lack of conversion for a nitrogen atmosphere, but there is little information for other to zircon. The strengths of the coated fibers were severely atmospheres. It has been claimed that ZrSiOa should be stable with degraded after heat treatment at 21000C in air, but not in carbon at <1527 C, based on thermodynamic calculations, but carbon-coated fibers made using similar methods. Mecha- reduction of ZrSiO4 at 1450C Ran be produced by carbothermal argon. The coated fibers were compared with zirconia- other work shows that Zro isms for fiber strength degradation are discussed. In this work, attempts to synthesize precursors for porous ZrSiO4 fiber coatings are reported. The precursors use carbon as a fugitive phase to hold porosity open during coating deposition. Conditions that may promote ZrSiO4 formation in the presence of carbon are examined. Fibers were coated with these precursors. here has been extensive research on development of conversion to ZrSio, was attempted, and strengths of the coated oxidation-resistant fiber-matrix interfaces for ceramic-matrix composites(CMCs). Recent results demonstrate the utility of tensile strengths of the coated fibers are discussed composites are another widely used approach. Here, a discrete coatings were previously discussed in a minicomposite study of intrinsic weakness of the porous matrix does not allow buildup of a matrix crack-tip stress that is sufficient to break fibers. A weakness of this approach is that matrix-dominated mechanical properties of porous-matrix CMCs tend to be poor Fretting and Il. Experimental Procedures wear at contact and attachment points can severely degrade porous matrices.A modified approach eserve matn Fiber coatings are typically 100 nm thick on 12 um fila- dominated properties is use of a porous fiber-matrix interface in a ments. 0-33 For a homogenous coating, the oxide--carbon disper dense-matrix CMC. -8 Candidate materials for a porous fiber sion should be homogeneous at a "20 nm scale. Separate matrix interface must be themochemically stable with the fiber and precipitation of oxides and carbon from precursors introduces the matrix, and they must be sufficiently refractory so that porosity extensive filament-to-filament variation in coatings. One approach designed into the interface material does not coarsen or densify is to bond the carbon and oxide precursors through electrostatic omposite processing and use and steric interactions. After the matrix is processed, carbon SiO4) decomposes to SiO and ZrO, at-16000- can be oxidized to leave a porous-oxide fiber-matrix interface. 1700C,,and it has a coefficient of thermal expansion(CTE)of The following chemicals were used: zirconyl nitrate hydrate 4.3X 10/C, which is a good match with SiC.,4 Conse- ammonium vanadate. lithium nitrate,(Aldrich Chemical Co uently, ZrSiOa has been used as a matrix in CMCs with SiC milwaukee, Wn), tetraethoxysilane(Alfa Chemical Co., Ward Hil fibers.- The low diffusion and creep rates of ZrSiO4 suggest it MA), and poly(acrylic acid)(PA; Fischer Scientific Co.,Pitts should resist densification and pore coarsening.-- Except for burgh, PA). Water was purified by deionization of distilled water HF, it is chemically resistant to acids. These properties make with a nanopure ultrapure system (Model D4744, Barnstead/ ZrSiO4 an attractive candidate for a porous fiber-matrix interface. Thermolyne Corp, Dubuque, IA) (2)ZNS Precursor F. w. Zok--contributing editor Zirconyl nitrate hydrate(ZN, 25 g) was dissolved in 500 mL of absolute ethanol and refluxed at 45C for 30 min. Tetraethoxy lane (TEoS, 22.5 added and the mixture was refluxed at anuscript No 10557. Received September 25. 2003; approved May 3, 2004. 55°Cfor24h.A g)was a m vanadate or lithium nitrate was added as a dopant to mber, American Ceramic Society ZrSiOa formation temperat concentration was 1.2 g/L, except where stated otherwise. In some ir Force Research Laboratory. formulations, TEOS was prehydrolyzed for 2 h at pH 2, with an
Zirconia–Silica–Carbon Coatings on Ceramic Fibers Emmanuel E. Boakye,* ,† Randall S. Hay,* ,‡ M. Dennis Petry,† and Triplicane A. Parthasarathy* ,† UES, Inc., Dayton, Ohio 45432 Materials and Manufacturing Directorate, Air Force Research Laboratory, Wright–Patterson Air Force Base, Ohio 45433 Precursors for zircon– carbon mixtures were made to coat fibers for ceramic-matrix composites. Precursors were characterized using XRD, TGA, and DTA. Zircon formed from vanadium- or lithium-doped precursors after heat treatments at >900°C in air, but it did not form at 1200°–1400°C in argon when large amounts of carbon were added. Some precursors were used to coat NextelTM 720 and Hi-NicalonTM fibers. The coatings were characterized using SEM and TEM, and coatedfiber tensile strengths were measured. Although zircon formed in powders, only tetragonal-zirconia–silica mixed phases formed in fiber coatings at 1200°C in air. Loss of vanadium oxide flux to the fibers may have caused the lack of conversion to zircon. The strengths of the coated fibers were severely degraded after heat treatment at >1000°C in air, but not in argon. The coated fibers were compared with zirconia– carbon-coated fibers made using similar methods. Mechanisms for fiber strength degradation are discussed. I. Introduction THERE has been extensive research on development of oxidation-resistant fiber–matrix interfaces for ceramic-matrix composites (CMCs).1 Recent results demonstrate the utility of monazite (LaPO4, NdPO4) for this interface.2–5 Porous-matrix composites are another widely used approach.6 –10 Here, a discrete phase does not exist at the fiber–matrix interface; instead, the intrinsic weakness of the porous matrix does not allow buildup of a matrix crack-tip stress that is sufficient to break fibers. A weakness of this approach is that matrix-dominated mechanical properties of porous-matrix CMCs tend to be poor. Fretting and wear at contact and attachment points can severely degrade porous matrices.8 A modified approach that can preserve matrixdominated properties is use of a porous fiber–matrix interface in a dense-matrix CMC.11–18 Candidate materials for a porous fiber– matrix interface must be themochemically stable with the fiber and the matrix, and they must be sufficiently refractory so that porosity designed into the interface material does not coarsen or densify during composite processing and use. Zircon (ZrSiO4) decomposes to SiO2 and ZrO2 at 1600°– 1700°C,19,20 and it has a coefficient of thermal expansion (CTE) of 4.3 106 /°C, which is a good match with SiC.21,22 Consequently, ZrSiO4 has been used as a matrix in CMCs with SiC fibers.22–24 The low diffusion and creep rates of ZrSiO4 suggest it should resist densification and pore coarsening.25–29 Except for HF, it is chemically resistant to acids.30 These properties make ZrSiO4 an attractive candidate for a porous fiber–matrix interface. ZrSiO4 forms from ZrO2 and SiO2 at 1400°–1600°C.31,32 For NextelTM 720 fiber coatings, fiber strength degradation associated with grain growth in the fiber requires ZrSiO4 to form at 1300°C.33 However, near-stoichiometric SiC fibers survive much higher temperatures in argon without such degradation.34 –37 Reacting ZrO2 with various SiO2 allotropes has little effect on ZrSiO4 formation.32 Sol– gel precursors decrease the formation temperature to 1200°C,32 but temperatures 1350°C are necessary for high purity.38 Seeding and fluxes can significantly decrease ZrSiO4 formation temperature.39 – 42 ZrSiO4 formation in air has been studied extensively,32,39,40,43– 47 and there is one study for a nitrogen atmosphere,42 but there is little information for other atmospheres. It has been claimed that ZrSiO4 should be stable with carbon at 1527°C, based on thermodynamic calculations,22 but other work shows that ZrO2 can be produced by carbothermal reduction of ZrSiO4 at 1450°C.48 In this work, attempts to synthesize precursors for porous ZrSiO4 fiber coatings are reported. The precursors use carbon as a fugitive phase to hold porosity open during coating deposition.13,49 Conditions that may promote ZrSiO4 formation in the presence of carbon are examined. Fibers were coated with these precursors, conversion to ZrSiO4 was attempted, and strengths of the coated fibers were measured. Characteristics of the fiber coatings and tensile strengths of the coated fibers are discussed and compared with those of ZrO2– carbon-coated fibers. Some features of these coatings were previously discussed in a minicomposite study of porous ZrO2–SiO2 fiber coatings.15 II. Experimental Procedures (1) General Fiber coatings are typically 100 nm thick on 12 m filaments.50 –53 For a homogenous coating, the oxide– carbon dispersion should be homogeneous at a 20 nm scale. Separate precipitation of oxides and carbon from precursors introduces extensive filament-to-filament variation in coatings. One approach is to bond the carbon and oxide precursors through electrostatic and steric interactions.13,49 After the matrix is processed, carbon can be oxidized to leave a porous-oxide fiber–matrix interface. The following chemicals were used: zirconyl nitrate hydrate, ammonium vanadate, lithium nitrate, (Aldrich Chemical Co., Milwaukee, WI), tetraethoxysilane (Alfa Chemical Co., Ward Hill, MA), and poly(acrylic acid) (PA; Fischer Scientific Co., Pittsburgh, PA). Water was purified by deionization of distilled water with a nanopure ultrapure system (Model D4744, Barnstead/ Thermolyne Corp., Dubuque, IA). (2) ZNS Precursor Zirconyl nitrate hydrate (ZN, 25 g) was dissolved in 500 mL of absolute ethanol and refluxed at 45°C for 30 min. Tetraethoxysilane (TEOS, 22.5 g) was added, and the mixture was refluxed at 55°C for 24 h. Ammonium vanadate or lithium nitrate was added as a dopant to decrease ZrSiO4 formation temperature. Their concentration was 1.2 g/L, except where stated otherwise. In some formulations, TEOS was prehydrolyzed for 2 h at pH 2, with an F. W. Zok—contributing editor Manuscript No. 10557. Received September 25, 2003; approved May 3, 2004. Supported by the Air Force Office of Scientific Research. *Member, American Ceramic Society. † UES, Inc. ‡ Air Force Research Laboratory. J. Am. Ceram. Soc., 87 [10] 1967–1976 (2004) 1967 journal
Joumal of the American Ceramic Sociery-Boakye et al. Vol 87. No. 10 Table L. Coated Nextel M 720 Fiber Strength for Various Coating Precursors, Heat Treatment Temperatures, Filament strength(GPa). Weibull modulus Argon, as coated Air. I h Coating temperature at-treat temperature Heat-treat temperature Precursor O0°C 1100° 1200°C 600°C 00°C ZNS-C 195,631.88,6.7 69,4.7 1.83.6 1.12.5.3 ZNS-C 93,5.4 2.35*,6.0 ZNS-C 0.19*,1.3 ZP. 1.50,4.8 STrength values are an average from 50 tests. Precursor weight loss(percent: see Fig. 4)above 1000C also is shown. indicates single-filament strengths computed from tow tests(Eq.(I); the Weibull modulii accompanying these values are those measured thanol: TEOS molar ratio of 5.8 and water: TEOS molar ratio of (5) ZP Precursor harge of hydrous Sioz and hydrolyzed ZN are -2 and-10-ll, ethanol and refluxed at 60%C for 4 h. PA (15 g) was added as the mixture was -1. At a pH of 1, Pa electrosterically interacted with 125 g/L and the ZrO PA weight ratio was 40: 60. Precursor the hydrated Sioz and the zirconyl nitrate hydrate to form a stable without Pa is referred to as ZP, and precursor with PA is referred to as ZP-C The ZrO2: SiO2: PA weight ratio was 27 ZrO,-SiO,carbon concentration of 40 iscosities of 1.26, 2.44, and 5.32 cP(1 cP= 1 X 10 N-s/m2) (6) Precursor Characterization respectively. Concentration refers to the ZrO2-SiO2-carbon yield after heat treatment at 140 C for 72 h. Precursor without Pa is viscosities were measured with a programmable referred to as ZNS, and precursor with PA is referred to as ZNS-C rheometer (Model DV-lll, Brookfield Engineering Laboratories Stoughton, MA)at a shear rate of 1/300s. Differential thermal analysis (DTA)and thermogravimetric analysis (TGA)were con- (3) ZES Precursor ducted (Model STA-409, Netzsch, Bayern, Germany). Powder The double alkoxide method was used to test the effects of samples for DTA and TGa were heat-treated at 140C for 18 h and molecular-scale mixing on ZrSiO formation. Zirconium ethoxide X-ray powder diffractometry (XRD)analyses were conducted (ZEOS, 10 g)and TEOS(8.6 g) were added to 1100 cm of Powder samples were heat-treated at 9000-1400C for I h (Model absolute ethanol in a dry box. Ethanol was distilled before use. The Rotaflex, Rigaku Co., Tokyo, Japan). The carbon concentration mixture was refluxed at 78 C for 18 h. For a concentration of 18 was evaluated from TGA curves for each preheat temperature. The g/L, 200 cm of the mixture was extracted, and ammonium ZrSiOa concentration was calculated from the ratio of the peak vanadate and dilute HNO, were added and hydrolyzed at 50 C for intensities for 20= 26.98(ZrSiO4) and the sum of the peak 24 h at pH 2. The water: (TEOS +ZEOS)molar ratio was 20. The intensities of 20=30.17(1-ZrO2), 20=2819%(m-ZrO2), and ZrO2-SiO carbon concentration was limited to 15 g/L by the ZrSiO4 Crystallite sizes were estimated for 1-ZrO2, m-ZrO2, and solubility of ZEOS in ethanol. The ZrO,: SiO,: PA weight ratio was ZrSiO4 using the Scherer formula 27: 13: 60. PA was added as the carbon precursor. Precursor without PA is referred to as ZES, and precursor with PA is referred to as ZES-C (7) Fiber Coating Nextel 720(AlOx-mullite)and Hi-Nicalon-STM(SiC)fiber (4) ZN Precursor tows were coated using the ZNS-C precursor Characteristics of Zirconyl nitrate hydrate(ZN, 25 g) was dissolved in 200 mL of these fibers have been reviewed elsewhere 33,36,58-60 ZNS-C water and refluxed at 60oC for 4 h. PA(15 g) was added. The precursor concentrations of 40-160 g/L were used. The ZES ZrO2-C sol concentration was 125 g/L and the Zro, PA weight precursor was not used, because it had relatively low concentration ratio was 40: 60. Precursor without Pa is referred to as zN. and and high viscosity. Nextel 720 also was coated using the ZP-C and precursor with Pa is referred to as ZN-C ZN-C precursors, using precursor concentrations of 125 g/L. A continuous vertical coater that used hexadecane for immiscible luid displacement was used for fiber coating 6.62 For Nextel 720, the hexadecane layer thickness was varied from 0 to 20 cm to Table Il. Al,O3 Grain Size (Major Axis) study the effect on coating thickness and uniformity. Nextel 720 after Heat Treatment at various coatings were heat-treated in-line at 1000-1300%C and Hi- Te peratures and Times for Nextel M 720 Nicalon-S coatings were heat-treated in-line at 1000%-1600C. al Coated with ZNS-C and zP-C Precursors" at a speed of 1. 4 cm/s in argon. Coated Nextel 720 and Hi- Heat treatment log grain size(nm) Nicalon-S were given further heat treatments in air and argon for 1-100 h at temperatures from 600 to 1200C (Tables I and D) 1.80±0.16 The fiber coatings were characterized using scanning electron microscopy(SEM: Model FEG, Leica, Buffalo, NY) and trans- 000/100 1.80±0.12 mission electron microscopy (TEM: Model CM 200 FEG, Philli 200/100 2.05±0.1 Eindhoven, The Netherlands). Coating phases were determined ZP-O 1000/1 195±0.16 using analytical TEM from energy-dispersive spectroscopy(EDS) in spot mode(5 nm) and from ring patterns obtained by selected area diffraction spectroscopy (SADS)of the fine-grained coatings
ethanol:TEOS molar ratio of 5.8 and water:TEOS molar ratio of 10. PA was added as the carbon precursor. The point of zero charge of hydrous SiO2 and hydrolyzed ZN are 2 and 10 –11, respectively.54 –56 The pH of the zirconyl nitrate hydrate–TEOS mixture was 1. At a pH of 1, PA electrosterically interacted with the hydrated SiO2 and the zirconyl nitrate hydrate to form a stable colloidal dispersion. The ZrO2:SiO2:PA weight ratio was 27:13:60. Precursors with ZrO2–SiO2– carbon concentration of 40, 80, and 160 g/L had viscosities of 1.26, 2.44, and 5.32 cP (1 cP 1 10–3 Ns/m2 ), respectively. Concentration refers to the ZrO2–SiO2– carbon yield after heat treatment at 140°C for 72 h. Precursor without PA is referred to as ZNS, and precursor with PA is referred to as ZNS-C. (3) ZES Precursor The double alkoxide method was used to test the effects of molecular-scale mixing on ZrSiO4 formation. Zirconium ethoxide (ZEOS, 10 g) and TEOS (8.6 g) were added to 1100 cm3 of absolute ethanol in a dry box. Ethanol was distilled before use. The mixture was refluxed at 78°C for 18 h. For a concentration of 18 g/L, 200 cm3 of the mixture was extracted, and ammonium vanadate and dilute HNO3 were added and hydrolyzed at 50°C for 24 h at pH 2. The water:(TEOS ZEOS) molar ratio was 20. The ZrO2–SiO2– carbon concentration was limited to 15 g/L by the solubility of ZEOS in ethanol. The ZrO2:SiO2:PA weight ratio was 27:13:60. PA was added as the carbon precursor. Precursor without PA is referred to as ZES, and precursor with PA is referred to as ZES-C. (4) ZN Precursor Zirconyl nitrate hydrate (ZN, 25 g) was dissolved in 200 mL of water and refluxed at 60°C for 4 h. PA (15 g) was added. The ZrO2–C sol concentration was 125 g/L and the ZrO2:PA weight ratio was 40:60. Precursor without PA is referred to as ZN, and precursor with PA is referred to as ZN-C. (5) ZP Precursor Zirconium propoxide (ZP, 26 g) was dissolved in 200 mL of ethanol and refluxed at 60°C for 4 h. PA (15 g) was added as the carbon precursor. The concentration of the ZrO2–C precursor was 125 g/L and the ZrO2:PA weight ratio was 40:60. Precursor without PA is referred to as ZP, and precursor with PA is referred to as ZP-C. (6) Precursor Characterization Precursor viscosities were measured with a programmable rheometer (Model DV-III, Brookfield Engineering Laboratories, Stoughton, MA) at a shear rate of 1/300 s1 . Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) were conducted (Model STA-409, Netzsch, Bayern, Germany). Powder samples for DTA and TGA were heat-treated at 140°C for 18 h and X-ray powder diffractometry (XRD) analyses were conducted. Powder samples were heat-treated at 900°–1400°C for 1 h (Model Rotaflex, Rigaku Co., Tokyo, Japan). The carbon concentration was evaluated from TGA curves for each preheat temperature. The ZrSiO4 concentration was calculated from the ratio of the peak intensities for 2 26.98° (ZrSiO4) and the sum of the peak intensities of 2 30.17° (t-ZrO2), 2 28.19° (m-ZrO2), and ZrSiO4. Crystallite sizes were estimated for t-ZrO2, m-ZrO2, and ZrSiO4 using the Scherer formula.57 (7) Fiber Coatings Nextel 720 (Al2O3–mullite) and Hi-Nicalon-STM (SiC) fiber tows were coated using the ZNS-C precursor. Characteristics of these fibers have been reviewed elsewhere.33,36,58 – 60 ZNS-C precursor concentrations of 40 –160 g/L were used. The ZES precursor was not used, because it had relatively low concentration and high viscosity. Nextel 720 also was coated using the ZP-C and ZN-C precursors, using precursor concentrations of 125 g/L. A continuous vertical coater that used hexadecane for immiscible liquid displacement was used for fiber coating.61,62 For Nextel 720, the hexadecane layer thickness was varied from 0 to 20 cm to study the effect on coating thickness and uniformity. Nextel 720 coatings were heat-treated in-line at 1000°–1300°C and HiNicalon-S coatings were heat-treated in-line at 1000°–1600°C, all at a speed of 1.4 cm/s in argon. Coated Nextel 720 and HiNicalon-S were given further heat treatments in air and argon for 1–100 h at temperatures from 600° to 1200°C (Tables I and II). The fiber coatings were characterized using scanning electron microscopy (SEM; Model FEG, Leica, Buffalo, NY) and transmission electron microscopy (TEM; Model CM 200 FEG, Phillips, Eindhoven, The Netherlands). Coating phases were determined using analytical TEM from energy-dispersive spectroscopy (EDS) in spot mode (5 nm) and from ring patterns obtained by selected area diffraction spectroscopy (SADS) of the fine-grained coatings. Table I. Coated Nextel™ 720 Fiber Strength for Various Coating Precursors, Heat Treatment Temperatures, and Atmospheres† Precursor Precursor weight loss (%) Filament strength (GPa), Weibull modulus Argon, as coated Argon, 1 h temperature Air, 1 h Coating temperature Heat-treat temperature Heat-treat temperature 1000°C 1100°C 1200°C 1000°C 1200°C 600°C 1000°C ZNS-C 1.95, 6.3 1.88, 6.7 1.69, 4.7 ZP-C 1.83, 6.2 ZN-C 1.12, 5.3 ZNS-C 1.93, 5.4 2.05*, 5.0 2.35*, 6.0 ZNS-C 0.13 0.19*, 1.3 ZP-C 1.34 1.55*, 5.4 ZN-C 0.90 0.93*, 5.2 No coating 1.50, 4.8 † Strength values are an average from 50 tests. Precursor weight loss (percent; see Fig. 4) above 1000°C also is shown. * indicates single-filament strengths computed from tow tests (Eq. (1)); the Weibull modulii accompanying these values are those measured for tows. Table II. Al2O3 Grain Size (Major Axis) after Heat Treatment at Various Temperatures and Times for Nextel™ 720 Coated with ZNS-C and ZP-C Precursors† Heat treatment log grain size(nm) As received 1.80 0.16 ZNS-C 1000/1 1.80 0.13 1000/100 1.80 0.12 1200/100 2.05 0.12 ZP-C 1000/1 1.95 0.16 † All grain-size distributions were lognormal, and the standard deviations of the distributions are shown. 1968 Journal of the American Ceramic Society—Boakye et al. Vol. 87, No. 10
October 2004 Zirconia-Silica-Carbon Coatings on Ceramic Fibers 1969 TEM samples of coated fibers were made using published meth- ods.6364 The effect of fiber coating and subsequent heat treatment on Nextel 720 grain size was measured from TEM micrographs using published methods, and fiber grain-boundary chemistry was probed using analytical TEM in EDS spot mode Coated-filament tensile strengths were measured using a 2.54 7100 dir cm gauge length and 50 filament tests, and, for some strengths also were measured using fiber-tow testing 65,66 300A filament strength(o was too low to be measured, only fiber-tow 1200A testing was done. Single-filament tensile strengths (o) were calculated from tow tensile strengths(o) using a dry bundle failure model modified to include slack in filaments in the 空人0 tows.67,68 Single-filament and tow tensile strengths were measured 300A for 1 1 Nextel 720 and Nextel 610 fiber tows each with a different fiber coating and/or heat treatment. The results suggested that the model should be modified to account for slack in filaments in the tow bundle(Fig. 1). The following relationship between o and o est fits the data G:=0.750em) g=1.340(em) where m is the measured Weibull modulus of single filaments for each fiber tow(typically about six opq II Results and discussion 宏巴A10 (1) Zircon Formation Heat treatment of the vanadium-doped ZNS precursor in air at 900 and 1000C formed 78%o and 86%0 ZrSio4, respectively and the remainder was [-ZrO,(Fig. 2(a)). ZrSiOa was the major rystalline phase(100%)after heat treatment in air at 1100C (Fig. 2(b)). Other studies have shown that ZrsiOa precursors 100d doped with vanadium and heat-treated in air show similar behavior In argon at <1300C, ZrSiO did not form with vanadium (Fig. 2(c), but ZNS precursors doped with lithium formed 4 at 1200.C(Fig. 2(d)). For vanadium-doped ZNS, heat treatment in argon at 1300.C formed only f-ZrO2,(Fig. 2(c)), and Fig. 2. Powder XRD pattens of ZrSiO4 and ZrsiO-carbon precursors heat treatment at 1400.C formed 67% ZrSiO. 33% I-ZrO, and heat-treated in air and argon at various temperatures traces of m-ZrO2 and cristobalite(Fig. 2(e). Vanadium-doped ZNS with prior heat treatment of 1200-1300oC in argon and urther heat treatment in air at 1000C formed only Zrsio similar phase evolution In argon at 1400.C, ZNS-C formed Zrc and SiC, with only a trace of ZrSiOa, and ZES-C formed no (2) Zircon-Carbon Fe ZrSiO,(Figs. 20) and(k). At 1300C, ZNS-C formed I-ZrO, in d heat teaCe a nd areo- a preoopsors were fored wand the huam amorphous SiO,; ZES-C formed t-ZrO, and m-ZrO, in amorphous SiO,(Figs. 2()and(m). ZNS-C and ZES-C formed I-ZrO, after heat treatment in argon at 1200C(Figs. 2(n)and(o)) Increased ammonium vanadate concentrations(4.8 g/)did not enhance ZrSioa formation. Replacing vanadium with lithium in the ZNS-C precursor was not as effective as in the ZNs precursor. Only 5%o and 3%o ZrsiO4 formed at 1200 and 1300oC,respec- tively, in argon(Figs. 2(p) and(q)) 8 (3) DTA/TGA The ZNS, ZNS-C, ZES, and ZES-C precursors experienced mass loss between 25.and 200C, which corresponded to an endothermic reaction from loss of water of hydration(Fig. 3) Weight loss and an endothermic reaction at 350C(C-1 in Fig 3) 0 were assigned to loss of nitrates and low-molecular-weight carbo- 0.5 1.5 Predicted Tow Strength(a) naceous species Weight loss and an exothermic reaction(C-2 in Fig 3)at 520C for the ZNS and zeS precursors were assigned to carbon combustion. The exothermic reactions at 875 and 1 150%C correlated with the crystallization of I-ZrO, and ZrSiOa, observed Fig. 1. Linear best fit of measured fiber-tow strengths to fiber-tow using XRd strengths calculated from single-filament strengths using the dry-bundle The small weight loss at higher temperatures was more easily model: a, Ko(me), using a Weibull modulus m measured indepen- observed when TGa weight loss was plotted logarithmically after dently for filaments in each fiber tow (typically about six). Best fit oc normalization to the final solids weight (+.1%, TGA precision) at for K=0.75 1500C (Fig. 4). The ZNS precursor contained 2 wt%o volatile
TEM samples of coated fibers were made using published methods.63,64 The effect of fiber coating and subsequent heat treatment on Nextel 720 grain size was measured from TEM micrographs using published methods,33 and fiber grain-boundary chemistry was probed using analytical TEM in EDS spot mode. Coated-filament tensile strengths were measured using a 2.54 cm gauge length and 50 filament tests, and, for some samples, strengths also were measured using fiber-tow testing.65,66 If filament strength ( f ) was too low to be measured, only fiber-tow testing was done. Single-filament tensile strengths ( f ) were calculated from tow tensile strengths ( t ) using a dry bundle failure model modified to include slack in filaments in the tows.67,68 Single-filament and tow tensile strengths were measured for 11 Nextel 720 and Nextel 610 fiber tows, each with a different fiber coating and/or heat treatment. The results suggested that the model should be modified to account for slack in filaments in the tow bundle (Fig. 1). The following relationship between f and t best fits the data: t 0.75 f(em) –1/m (1a) or f 1.34 t(em) 1/m (1b) where m is the measured Weibull modulus of single filaments for each fiber tow (typically about six). III. Results and Discussion (1) Zircon Formation Heat treatment of the vanadium-doped ZNS precursor in air at 900° and 1000°C formed 78% and 86% ZrSiO4, respectively, and the remainder was t-ZrO2 (Fig. 2(a)). ZrSiO4 was the major crystalline phase (100%) after heat treatment in air at 1100°C (Fig. 2(b)). Other studies have shown that ZrSiO4 precursors doped with vanadium and heat-treated in air show similar behavior.69,70 In argon at 1300°C, ZrSiO4 did not form with vanadium dopants (Fig. 2(c)), but ZNS precursors doped with lithium formed ZrSiO4 at 1200°C (Fig. 2(d)). For vanadium-doped ZNS, heat treatment in argon at 1300°C formed only t-ZrO2, (Fig. 2(c)), and heat treatment at 1400°C formed 67% ZrSiO4, 33% t-ZrO2, and traces of m-ZrO2 and cristobalite (Fig. 2(e)). Vanadium-doped ZNS with prior heat treatment of 1200°–1300°C in argon and further heat treatment in air at 1000°C formed only ZrSiO4. (2) Zircon–Carbon Formation The ZNS-C and ZES-C precursors were doped with vanadium and heat-treated in argon at 1200°–1400°C for 1 h, and they had similar phase evolution. In argon at 1400°C, ZNS-C formed ZrC and SiC, with only a trace of ZrSiO4, and ZES-C formed no ZrSiO4 (Figs. 2(j) and (k)). At 1300°C, ZNS-C formed t-ZrO2 in amorphous SiO2; ZES-C formed t-ZrO2 and m-ZrO2 in amorphous SiO2 (Figs. 2(l) and (m)). ZNS-C and ZES-C formed t-ZrO2 after heat treatment in argon at 1200°C (Figs. 2(n) and (o)). Increased ammonium vanadate concentrations (4.8 g/L) did not enhance ZrSiO4 formation. Replacing vanadium with lithium in the ZNS-C precursor was not as effective as in the ZNS precursor. Only 5% and 3% ZrSiO4 formed at 1200° and 1300°C, respectively, in argon (Figs. 2(p) and (q)). (3) DTA/TGA The ZNS, ZNS-C, ZES, and ZES-C precursors experienced mass loss between 25° and 200°C, which corresponded to an endothermic reaction from loss of water of hydration (Fig. 3). Weight loss and an endothermic reaction at 350°C (C-1 in Fig. 3) were assigned to loss of nitrates and low-molecular-weight carbonaceous species. Weight loss and an exothermic reaction (C-2 in Fig. 3) at 520°C for the ZNS and ZES precursors were assigned to carbon combustion. The exothermic reactions at 875° and 1150°C correlated with the crystallization of t-ZrO2 and ZrSiO4, observed using XRD. The small weight loss at higher temperatures was more easily observed when TGA weight loss was plotted logarithmically after normalization to the final solids weight (0.1%, TGA precision) at 1500°C (Fig. 4). The ZNS precursor contained 2 wt% volatile Fig. 1. Linear best fit of measured fiber-tow strengths to fiber-tow strengths calculated from single-filament strengths using the dry-bundle model: t K f (me) –1/m, using a Weibull modulus m measured independently for filaments in each fiber tow (typically about six). Best fit occurs for K 0.75. Fig. 2. Powder XRD patterns of ZrSiO4 and ZrSiO4– carbon precursors heat-treated in air and argon at various temperatures. October 2004 Zirconia–Silica–Carbon Coatings on Ceramic Fibers 1969
1970 Joumal of the American Ceramic Society-Boakye et al. Vol 87. No. 10 100 ZNS 60 ZES-C ZNS-C ZNS.C ZNS ES-C 20040060080010001200 40060080010001200 Temperature(C) Fig 3. DTA/TGA of ZNS, ZNS-C, ZES, and ZES-C precursors species at 1000oC. The ZNS-C precursor had -1% weight gain above refractory VO, phase should have formed, which would have been 900C, which possibly was from formation and subsequent oxidation ineffective as a flux of trace amounts of carbides. SiO, may have locally sea Unlike vanadium-doped ZNS, lithium-doped ZNS formed Zr- rom oxidation at -900-1200oC and, consequently, promoted local SiO4 in argon at 1200C(Fig. 2(d). However, in argon, lithium- carbothermal reduction of oxides to carbides. It is surprising that the doped ZNS-C formed only a trace of ZrSiO4 at 1200oC (Fig. 2(m)) TGAs of the ZNS and ZNS-C precursors were dissimilar at high and formed no ZrSiO4 at 1300C (Fig. 2(1)). Lithium also emperatures. The only difference between the two precursors was the enhanced ZrsiO4 formation by a liquid-phase flux mechanism, but PA added to ZNS-C, which should have burned off above 600oC. lithium formed only one stable oxide, which was unaffected by an When Sio, was absent from the precursors, as for ZN-C and ZP-C, argon atmosphere. Unlike the V2Os-doped precursors, the Li,o the weight loss was similar to that for the ZNS-C precursor below flux should have been effective; therefore, other explanations 550C; above 550C, weight loss was similar to that for the ZNs lack of ZrSiO4 formation must be sought. The slightly higher precursor(Fig. 4), as was expected formation rate at the lower temperature is difficult to explai Comparison of the weight loss from the ZNS and ZNS-C kinetically and suggests a thermodynamic cause precursors showed that the ZNS-C precursor yielded --55 wt%(36 The stability of ZrsiO4 in the presence of carbon was tested Dl%) ZrSiO and 45 wt%(64 vol%)carbon (4) ZrSio, Stability In the absence of vanadium or lithium dopa ZrSiO4(s)+C(s)-ZrO(s)+ Sio(g)+Co(g) conversion to Zrsio4 occurred at 1400.C for ZNS pre ZrSiO4 was stable in argon, but, in the presence of carbon, it Vanadium-doped ZNS recurse ely reacted to air at 21100C(Fig. 2(b)), but, in argon, the precursors did not partially decomposed to m-Zro2(Figs. 2(u)and (v). This result react until 1400C. Vanadium has been suggested to function as a uid-Phase flux through formation of melted V2Os,but, in argon and particularly in the presence of carbon, the much more (a) epoxy 100 Vextel20 10 (c) Nextel 20 2( nm1 Temperature℃ Fig. 5. TEM images of ZNS-C-derived coatings on Nextel 720 Coating was done at 1000.C in argon. Coatings are an intimate mechanical Fig 4. TGA of ZNS, ZNS-C, ZN-C, and ZP-C precursors. Weight loss is mixture of I-ZrO,, SiO2, and carbon (a)Coating pliability is evident on normalized to final solids yield at 1500C, and it is plotted logarithmically cracks fo filaments when they are close together. degradation occurs. Circled points show weight loss for precursors above during TEM specimen preparation. (c)Continuous coating coverage is 1000C, and they are used in Fig. 14 achieved. althou are sometimes thin
species at 1000°C. The ZNS-C precursor had 1% weight gain above 900°C, which possibly was from formation and subsequent oxidation of trace amounts of carbides. SiO2 may have locally sealed carbon from oxidation at 900°–1200°C and, consequently, promoted local carbothermal reduction of oxides to carbides. It is surprising that the TGAs of the ZNS and ZNS-C precursors were dissimilar at high temperatures. The only difference between the two precursors was the PA added to ZNS-C, which should have burned off above 600°C. When SiO2 was absent from the precursors, as for ZN-C and ZP-C, the weight loss was similar to that for the ZNS-C precursor below 550°C; above 550°C, weight loss was similar to that for the ZNS precursor (Fig. 4), as was expected. Comparison of the weight loss from the ZNS and ZNS-C precursors showed that the ZNS-C precursor yielded 55 wt% (36 vol%) ZrSiO4 and 45 wt% (64 vol%) carbon. (4) ZrSiO4 Stability In the absence of vanadium or lithium dopants, complete conversion to ZrSiO4 occurred at 1400°C for ZNS precursors.31,41 Vanadium-doped ZNS precursors completely reacted to ZrSiO4 in air at 1100°C (Fig. 2(b)), but, in argon, the precursors did not react until 1400°C. Vanadium has been suggested to function as a liquid-phase flux through formation of melted V2O5, 69 but, in argon and particularly in the presence of carbon, the much more refractory VO2 phase should have formed, which would have been ineffective as a flux. Unlike vanadium-doped ZNS, lithium-doped ZNS formed ZrSiO4 in argon at 1200°C (Fig. 2(d)). However, in argon, lithiumdoped ZNS-C formed only a trace of ZrSiO4 at 1200°C (Fig. 2(m)) and formed no ZrSiO4 at 1300°C (Fig. 2(l)). Lithium also enhanced ZrSiO4 formation by a liquid-phase flux mechanism, but lithium formed only one stable oxide, which was unaffected by an argon atmosphere. Unlike the V2O5-doped precursors, the Li2O flux should have been effective; therefore, other explanations for lack of ZrSiO4 formation must be sought. The slightly higher formation rate at the lower temperature is difficult to explain kinetically and suggests a thermodynamic cause. The stability of ZrSiO4 in the presence of carbon was tested. ZrSiO4 and 46:54 ZrSiO4:carbon (PA) mixtures were heat-treated in argon at 1300°C for 2 h. The ZrSiO4 decomposition reaction is ZrSiO4(s) C(s) 3 ZrO2(s) SiO(g) CO(g) (2) ZrSiO4 was stable in argon, but, in the presence of carbon, it partially decomposed to m-ZrO2 (Figs. 2(u) and (v)). This result Fig. 3. DTA/TGA of ZNS, ZNS-C, ZES, and ZES-C precursors. Fig. 4. TGA of ZNS, ZNS-C, ZN-C, and ZP-C precursors. Weight loss is normalized to final solids yield at 1500°C, and it is plotted logarithmically to emphasize weight loss at high temperatures, where fiber strength degradation occurs. Circled points show weight loss for precursors above 1000°C, and they are used in Fig. 14. Fig. 5. TEM images of ZNS-C-derived coatings on NextelTM 720. Coating was done at 1000°C in argon. Coatings are an intimate mechanical mixture of t-ZrO2, SiO2, and carbon. (a) Coating pliability is evident on cracks formed during thin-section preparation. (b) Coatings bridge adjacent filaments when they are close together, and sometimes crack or debond during TEM specimen preparation. (c) Continuous coating coverage is achieved, although coatings are sometimes thin. 1970 Journal of the American Ceramic Society—Boakye et al. Vol. 87, No. 10
October 2004 Zirconia-Silica-Carbon Coatings on Ceramic Fibers was consistent with the results for -doped ZNS-C(Figs ostructure development of ZNS-C-derived 2(p) and (q)) and suggested that was unstable in the fiber coatings formed in argon on Nextel 720 and Hi-Nicalon -s presence of carbon above 1250c was a lower tempera- was similar in most respects to that in powders. Coatings were ture than previously suggested. 224& At 1400 C, ZrC formed(Figs deposited as small(--5 nm)t-ZrO2 particles in an amorphous Sio 20)and (k). The uniformity and small particle size of the powders and carbon matrix at 1000%-1200C(Fig 5). Coating thickness and coatings derived from the ZNs-C precursors may have ranged from 10 to >100 nm. The coatings were pliable; they influenced reduction maintained continuity over Mode ll cracks that developed at fiber surfaces during thin-section preparation(Fig. 5(a)). At 1300 (5) Fiber Coatings 1400C, the t-ZrO2 particle size was slightly larger(5-10 nm) (A)Phase and Microstructure: Coating morphology was and, at 1500%C, it was -15 nm(Fig. 6). Some coatings formed at similar to that observed for other systems.so Every filament was 1300C had an outer layer of -20 nmt-ZrOz particles and an inner oated but the thickness varied. The median coating thickness was 50 nm for Nextel 720 fibers using a 65 g/L precursor. The 1.2 present at temperatures as low as 1300.C(Fig. 7). Coatings g/L of ammonium vanadate should have yielded a V2O5 flux deposited at 1600C consisted of "50 nm ZrC mosaic crystals, concentration of 2.5 wt% relative to ZrSiO4 for this precursor with evidence of graphitic carbon, SiO2, and trace SiC from EDS Coating bridges between filaments(Fig. 5(b)and crust around the fiber-tow perimeter were common for higher-concentration and 1600%c and electron diffraction patterns(Fig. 6). At 15000 the carbon was turbostratic. Residence time at precursors, particularly for the 160 g/L precursor with viscosity of emperature was only a few seconds during coating; therefore, ZrC 5.3 cP. This was consistent with previous work on monazite formation required higher temperatures to form during coating coatings, where crust and other coating imperfections were more than it did during I h heat treatments of powders(Fig. 2). TE prevalent when precursor viscosity was >3 cP. Coating mor vanadium flux was not detected using EDS in the argon-deposited phology and thickness were not affected by the hexadecane layer coatings, except for occasional high-contrast particles with a high thickness vanadium content which were inferred to be vanadium carbide 了400 10001200 1c8 5001000& ZrC gra shite 10 nm Fig. 6. TEM images of SiO,, some turbostratic carbon, and I-ZrO2 are present after coating at 1400%and 1500C pze turbostatic graphitic carbon, and some amorphous
was consistent with the results for lithium-doped ZNS-C (Figs. 2(p) and (q)) and suggested that ZrSiO4 was unstable in the presence of carbon above 1250°C, which was a lower temperature than previously suggested.22,48 At 1400°C, ZrC formed (Figs. 2(j) and (k)). The uniformity and small particle size of the powders and coatings derived from the ZNS-C precursors may have influenced reduction.71–74 (5) Fiber Coatings (A) Phase and Microstructure: Coating morphology was similar to that observed for other systems.50 Every filament was coated, but the thickness varied. The median coating thickness was 50 nm for Nextel 720 fibers using a 65 g/L precursor. The 1.2 g/L of ammonium vanadate should have yielded a V2O5 flux concentration of 2.5 wt% relative to ZrSiO4 for this precursor. Coating bridges between filaments (Fig. 5(b)) and crust around the fiber-tow perimeter were more common for higher-concentration precursors, particularly for the 160 g/L precursor with viscosity of 5.3 cP. This was consistent with previous work on monazite coatings, where crust and other coating imperfections were more prevalent when precursor viscosity was 3 cP.50 Coating morphology and thickness were not affected by the hexadecane layer thickness. The phase and microstructure development of ZNS-C-derived fiber coatings formed in argon on Nextel 720 and Hi-Nicalon-S was similar in most respects to that in powders. Coatings were deposited as small (5 nm) t-ZrO2 particles in an amorphous SiO2 and carbon matrix at 1000°–1200°C (Fig. 5). Coating thickness ranged from 10 to 100 nm. The coatings were pliable; they maintained continuity over Mode II cracks that developed at fiber surfaces during thin-section preparation (Fig. 5(a)). At 1300°– 1400°C, the t-ZrO2 particle size was slightly larger (5–10 nm), and, at 1500°C, it was 15 nm (Fig. 6). Some coatings formed at 1300°C had an outer layer of 20 nm t-ZrO2 particles and an inner layer of 5 nm particles (Fig. 7). Evidence of graphitic carbon was present at temperatures as low as 1300°C (Fig. 7). Coatings deposited at 1600°C consisted of 50 nm ZrC mosaic crystals, with evidence of graphitic carbon, SiO2, and trace SiC from EDS spot analysis and electron diffraction patterns (Fig. 6). At 1500° and 1600°C, the carbon was turbostratic. Residence time at temperature was only a few seconds during coating; therefore, ZrC formation required higher temperatures to form during coating than it did during 1 h heat treatments of powders (Fig. 2). The vanadium flux was not detected using EDS in the argon-deposited coatings, except for occasional high-contrast particles with a high vanadium content, which were inferred to be vanadium carbide. Fig. 6. TEM images of ZNS-C-derived coatings on Hi-NicalonTM-S SiC fiber. Coatings were applied at 1400°, 1500°, and 1600°C in argon. Amorphous SiO2, some turbostratic graphitic carbon, and t-ZrO2 are present after coating at 1400° and 1500°C. ZrC, turbostratic graphitic carbon, and some amorphous SiO2 are present after coating at 1600°C. October 2004 Zirconia–Silica–Carbon Coatings on Ceramic Fibers 1971
1972 Joumal of the American Ceramic Society-Boakye et al. Vol 87. No. 10 t动m SNexte 720 graphitic carbon Nextel 20 on NexteJTM1720 Fig. 7. TEM images of ZNS-C-derived coatings on Nextel 720 Coatings were applied at 1300.C in argon Coatings were an intimate mechanical mixture of 5 nm t-ZrO2, SiOz, and carbon. In a few places, an overlayer of 20 nm t-ZrO2 particles was present. In some places, the carbon was graphitic. Vanadium from the flux was only occasionally detected at locally high concentrations in particles, which were inferred to be vanadium carbide. Coatings deposited at 1000C on Nextel 720 and heat-treated Grain growth was not measurable in the ZNS-C coated fibers for I h in argon at 1200C showed spatially nonuniform coarsen- after I or 100 h at 1000C in air(Fig. 10). Lognormal Al,O3 ing of the I-ZrO2 particles to >10 nm, with no trace of ZrSio4 grain-size distributions with an inverse logaverage of 63 nm(long formation(Fig 8). In most coatings, the Sio2-ZrO2 dispersion was dimension) was found in both cases, which was identical to that in uniform. However, in isolated locations on some coatings, the as-received fiber. ,Slight grain growth to 89 nm was detected in dispersion was nonuniform on a local scale, with 5 nm 1-ZrO ZP-C-coated fiber after I h at 1000C in air(Table m). Significant articles dispersed over 30-60 nm Sio, spheres after I h at rain growth was evident in the ZNS-C-coated Nextel 720 after 1000C in argon, and 20 nm t-ZrO2 particles dispersed over 100 h at 1200oC, with noticeable faceting of the Al,O3 grains(Fig 100-300 nm SiO, spheres after I h at 1000 C in air(Fig 9). There 11). A lognormal Al O3 grain-size distribution with an inverse was no obvious cause of the sporadic nonuniformity in the logaverage of 112 nm was found. It is tempting to attribute this di grain growth to the V2O3 flux, but the effect of this flux on grain Although heat treatments of vanadium-doped powders formed growth in the Al2O3, Al2O3-mullite, or mullite systems has, to our ZrSio, at 900%C for I h in air heat treatments of 100 h at 1000 knowledge, not been studied. and 1200C in air did not form ZrSio, in ZNS-C-derived fiber (B) Fiber Strength: The as-received Nextel 720 tensile coatings On Nextel 720 fibers, dispersions of 20 nm t-ZrO2 grai strength of 2 GPa was retained after it was coated with the in dense, amorphous SiO, formed after I and 100 h at 1000.C ZNS-C precursor at 1000C in argon, but there was a slight (Fig. 10). ith log mean(nm) of 1.28+ 0.25 and 1.33 0.22 These I-ZrO2 grains had a lognormal particle-siz strength decrease after it was coated at 1200C(Fig. 13 and Table distribution I). Coated-fiber strength was independent of precursor concentra- after I and 100 h, respectively. Little or no coarsening occurred at tion. Strength was not degraded after heat treatment in argon at 1000%-1200C for 1 h. Coated-fiber strength also was retained amorphous SiO, after 100 h at 1200C(Fig. 11). The metastability after 1 h in air at 600 C and I h in argon at 1000C, but, after I h of the I-ZrO, phase for small particle sizes was frequently in air at 1000C, the fiber strength was severely degraded; the fiber observed and has been extensively discussed elsewhere. 75 Vana- had only -10% of the as-received strength. ZP-C-coated Nextel um was not detected in the coating using EDS. Lack of ZrSio4 720 was degraded only slightly after it was coated at 1000oC in formation was attributed to loss of V,Os flux, perhaps to the fibers argon, but was not degraded nearly as much as ZNS-C coated However, vanadium was not detected along fiber grain boundaries Nextel 720 after 1 h in air at 1000 C. However, ZN-C-coated or triple junctions using EDS with analytical TEM. The 2.5 wt% Nextel 720 strength was degraded by almost 50% after it was V2Os flux in a 100 nm thick coating would have been diluted by coated in argon at 1000 C, but it underwent only little further a factor of 30 if it had partitioned equally with the fiber, and it degradation after heat treatment in air at 1000C for I h would not have been detectable unless it concentrated as precipi- Hi-Nicalon fibers coated with the ZNS-C precursor and heat tates. ZP-C-derived coatings(no SiO,) formed porous 15-20 nm treated in air at 1000oC for I h were so weak that they could no I-ZrO, after I h at 1000C in air(Fig. 12) be tested for tow-bundle strength. This contrasted with results
Coatings deposited at 1000°C on Nextel 720 and heat-treated for 1 h in argon at 1200°C showed spatially nonuniform coarsening of the t-ZrO2 particles to 10 nm, with no trace of ZrSiO4 formation (Fig. 8). In most coatings, the SiO2–ZrO2 dispersion was uniform. However, in isolated locations on some coatings, the dispersion was nonuniform on a local scale, with 5 nm t-ZrO2 particles dispersed over 30 – 60 nm SiO2 spheres after 1 h at 1000°C in argon, and 20 nm t-ZrO2 particles dispersed over 100 –300 nm SiO2 spheres after 1 h at 1000°C in air (Fig. 9). There was no obvious cause of the sporadic nonuniformity in the dispersion. Although heat treatments of vanadium-doped powders formed ZrSiO4 at 900°C for 1 h in air, heat treatments of 100 h at 1000° and 1200°C in air did not form ZrSiO4 in ZNS-C-derived fiber coatings. On Nextel 720 fibers, dispersions of 20 nm t-ZrO2 grains in dense, amorphous SiO2 formed after 1 and 100 h at 1000°C (Fig. 10). These t-ZrO2 grains had a lognormal particle-size distribution with log mean(nm) of 1.28 0.25 and 1.33 0.22 after 1 and 100 h, respectively. Little or no coarsening occurred at 1000°C. Dispersions of 200 nm m-ZrO2 grains formed in dense, amorphous SiO2 after 100 h at 1200°C (Fig. 11). The metastability of the t-ZrO2 phase for small particle sizes was frequently observed and has been extensively discussed elsewhere.75 Vanadium was not detected in the coating using EDS. Lack of ZrSiO4 formation was attributed to loss of V2O5 flux, perhaps to the fibers. However, vanadium was not detected along fiber grain boundaries or triple junctions using EDS with analytical TEM. The 2.5 wt% V2O5 flux in a 100 nm thick coating would have been diluted by a factor of 30 if it had partitioned equally with the fiber, and it would not have been detectable unless it concentrated as precipitates. ZP-C-derived coatings (no SiO2) formed porous 15–20 nm t-ZrO2 after 1 h at 1000°C in air (Fig. 12). Grain growth was not measurable in the ZNS-C coated fibers after 1 or 100 h at 1000°C in air (Fig. 10). Lognormal Al2O3 grain-size distributions with an inverse logaverage of 63 nm (long dimension) was found in both cases, which was identical to that in as-received fiber.33,53 Slight grain growth to 89 nm was detected in ZP-C-coated fiber after 1 h at 1000°C in air (Table II). Significant grain growth was evident in the ZNS-C-coated Nextel 720 after 100 h at 1200°C, with noticeable faceting of the Al2O3 grains (Fig. 11). A lognormal Al2O3 grain-size distribution with an inverse logaverage of 112 nm was found. It is tempting to attribute this grain growth to the V2O3 flux, but the effect of this flux on grain growth in the Al2O3, Al2O3–mullite, or mullite systems has, to our knowledge, not been studied. (B) Fiber Strength: The as-received Nextel 720 tensile strength of 2 GPa was retained after it was coated with the ZNS-C precursor at 1000°C in argon, but there was a slight strength decrease after it was coated at 1200°C (Fig. 13 and Table I). Coated-fiber strength was independent of precursor concentration. Strength was not degraded after heat treatment in argon at 1000°–1200°C for 1 h. Coated-fiber strength also was retained after 1 h in air at 600°C and 1 h in argon at 1000°C, but, after 1 h in air at 1000°C, the fiber strength was severely degraded; the fiber had only 10% of the as-received strength. ZP-C-coated Nextel 720 was degraded only slightly after it was coated at 1000°C in argon, but was not degraded nearly as much as ZNS-C coated Nextel 720 after 1 h in air at 1000°C. However, ZN-C-coated Nextel 720 strength was degraded by almost 50% after it was coated in argon at 1000°C, but it underwent only little further degradation after heat treatment in air at 1000°C for 1 h. Hi-Nicalon fibers coated with the ZNS-C precursor and heattreated in air at 1000°C for 1 h were so weak that they could not be tested for tow-bundle strength. This contrasted with results Fig. 7. TEM images of ZNS-C-derived coatings on NextelTM 720. Coatings were applied at 1300°C in argon. Coatings were an intimate mechanical mixture of 5 nm t-ZrO2, SiO2, and carbon. In a few places, an overlayer of 20 nm t-ZrO2 particles was present. In some places, the carbon was graphitic. Vanadium from the flux was only occasionally detected at locally high concentrations in particles, which were inferred to be vanadium carbide. 1972 Journal of the American Ceramic Society—Boakye et al. Vol. 87, No. 10
October 2004 Zirconia-Silica-Carbon Coatings on Ceramic Fibers 1973 1000C, Targon coating 01 5nm 20nm 000c,如ha Nextel 200nm d7000 C, ih air. atings )O nm 100 nm 200nm Fig. 8. TEM images of ZNS-C-derived coatings on Nextel720 Datings were applied at 1000.C in argon and heat-treated for I h at Fig 9. TEM images of ZNS-C-derived coatings on Nextel I-ZrO2, SiO,, and carbon, but some 1-ZrO2 particles coarsened to >10 how a breakdown in the intimate mechanical mixture between ((c)and(d)). There was no trace of ZrSiO4 formation SiO, Coatings were applied at 1000.C in argon Coatings in(a) heat-treated for I h at 1000C for I h at 1000C in air. In all cases, smaller f-ZrO2 particles surrou arger spheres of SiO2. Spheres were significantly larger after reported for Tyranno ZMIM SiC fiber, which was coated with dense ZrSiOa at 1200.C without degrading fiber strength using a metal citrate-polymerized complex method. However, the fiber was not heat-treated after the coating was applied; therefore, direct was tentatively attributed to oxidation of trace carbide. Actual comparison with our results was difficult. weight loss from trapped gaseous decomposition products of th For all precursors and fibers, burnoff of carbon at 1000c ZNS-C precursor may have been masked by this oxidation weight degraded fiber strength. This has been observed for other oxide_ g fiber grain growth could not account for lower strength from carbon coatings. One possible explanation for the effect of carbon is that surface-active carbon in a coating scavenges surface-active coating decomposition products that would other- fibers at 1000C, whereas slight growth was observed for ZP-C- wise cause stress corrosion of the fibers coated fibers (Table ID). Monazite-coated fibers exhibited more Previous work on monazite-(LaPOa-) coated Nextel 720 fiber grain growth than that observed for various ZrO2-containing showed that coated fibers were strong when coating precursors that precursors in this study. No correlation between fiber grain growth had low weight loss above the coating temperature were used. It and fiber strength was observed for monazite-coated fibers, per- owth was insufficient to create critical flaws active agent that caused the stress corrosion responsible for larger than those that formed by stress corrosion. Lack of a fiber zns Ch degradation zpe relaten Sextet wezo iber s trener h tar grain-zgrewcth el fce was consistent with that observed for treatment at 1000C for I h in air and the weight loss observed for All monazite precursors studied earlier contained a nitrate as the these precursors above 1000.C is shown in Fig. 14. The results lanthanum source.. This reinforced the observations made have been superimposed on results observed earlier for monazite- previously for monazite, where fibers coated with precursors for coated Nextel 720. The trends were exactly opposite to th which nitrate was washed-out had higher strengths. In this stud observed for monazite-coated fibers: Fibers coated with the pre- the only precursor that did not contain nitrate was ZP-C. Relative cursor with the highest weight loss above 1000C, ZP-C, had the to most nitrate-containing precursors for ZrO, and monazite, fibers highest strength, and fibers coated with the precursor with the oated with this precursor were strong after exposure to high lowest weight loss above 1000 C, ZNS-C, had the lowest strengt temperatures in air, which was consistent with the idea that a However, as discussed earlier, ZNS-C had a slight weight gain nitrogen-containing gaseous species was primarily responsible for (1%)between 900and 1300C, rather than weight loss, which stress corrosion at high temperatures
reported for Tyranno ZMITM SiC fiber, which was coated with dense ZrSiO4 at 1200°C without degrading fiber strength using a metal citrate-polymerized complex method.76 However, the fiber was not heat-treated after the coating was applied; therefore, direct comparison with our results was difficult. For all precursors and fibers, burnoff of carbon at 1000°C degraded fiber strength. This has been observed for other oxide– carbon coatings.14,77 One possible explanation for the effect of carbon is that surface-active carbon in a coating scavenges surface-active coating decomposition products that would otherwise cause stress corrosion of the fibers. Previous work on monazite- (LaPO4-) coated Nextel 720 showed that coated fibers were strong when coating precursors that had low weight loss above the coating temperature were used.51 It was suggested that a nitrogen-containing species was the surfaceactive agent that caused the stress corrosion responsible for strength degradation.51 The relationship between fiber strength for ZNS-C-, ZN-C-, and ZP-C-coated Nextel 720 fibers after heat treatment at 1000°C for 1 h in air and the weight loss observed for these precursors above 1000°C is shown in Fig. 14. The results have been superimposed on results observed earlier for monazitecoated Nextel 720.51,53 The trends were exactly opposite to those observed for monazite-coated fibers: Fibers coated with the precursor with the highest weight loss above 1000°C, ZP-C, had the highest strength, and fibers coated with the precursor with the lowest weight loss above 1000°C, ZNS-C, had the lowest strength. However, as discussed earlier, ZNS-C had a slight weight gain (1%) between 900° and 1300°C, rather than weight loss, which was tentatively attributed to oxidation of trace carbide. Actual weight loss from trapped gaseous decomposition products of the ZNS-C precursor may have been masked by this oxidation weight gain. Fiber grain growth could not account for lower strength from ZNS-C-derived coatings; no grain growth was observed in these fibers at 1000°C, whereas slight growth was observed for ZP-Ccoated fibers (Table II). Monazite-coated fibers exhibited more fiber grain growth than that observed for various ZrO2-containing precursors in this study. No correlation between fiber grain growth and fiber strength was observed for monazite-coated fibers, perhaps because grain growth was insufficient to create critical flaws larger than those that formed by stress corrosion.53 Lack of a fiber grain-growth effect was consistent with that observed for monazite-coated fibers. All monazite precursors studied earlier contained a nitrate as the lanthanum source.51,53 This reinforced the observations made previously for monazite, where fibers coated with precursors for which nitrate was washed-out had higher strengths. In this study, the only precursor that did not contain nitrate was ZP-C. Relative to most nitrate-containing precursors for ZrO2 and monazite, fibers coated with this precursor were strong after exposure to high temperatures in air, which was consistent with the idea that a nitrogen-containing gaseous species was primarily responsible for stress corrosion at high temperatures. Fig. 8. TEM images of ZNS-C-derived coatings on NextelTM 720. Coatings were applied at 1000°C in argon and heat-treated for 1 h at 1200°C in argon. Coatings were an intimate mechanical mixture of 5 nm t-ZrO2, SiO2, and carbon, but some t-ZrO2 particles coarsened to 10 nm ((c) and (d)). There was no trace of ZrSiO4 formation. Fig. 9. TEM images of ZNS-C-derived coatings on NextelTM 720 that show a breakdown in the intimate mechanical mixture between t-ZrO2 and SiO2. Coatings were applied at 1000°C in argon. Coatings in (a)–(c) were heat-treated for 1 h at 1000°C in argon. Coatings in (d) were heat-treated for 1 h at 1000°C in air. In all cases, smaller t-ZrO2 particles surrounded larger spheres of SiO2. Spheres were significantly larger after heat treatment in air. October 2004 Zirconia–Silica–Carbon Coatings on Ceramic Fibers 1973
1974 Joumal of the American Ceramic Society-Boakye et al. Vol 87. No. 10 1000C,1h epoxy m-sircouia 50 nm s.Nextel 1000C,100h epoxy 。 Nexte720 200nm Nextel 720 Fig. 11. TEM images of ZNS-C coatings on Nextel 720 that and transformation to a twinned sImms show coarsening of ZrO, particles monoclinic phase. Coatings were at 1000@C in argon and heat- reated for 1 and 100 h at 1200.C where it may promote grain growth and degrade fiber strength Finely dispersed multiphase fiber coatings of I-ZrO2-SiOx-carbon bon phase is rapidly filled by SiO,, and dense I-4 oatings form. Consequently, such (-ZrO2-SiO2 coatings are n weak and. therefore, are not useful as fiber-matrix interfaces in eramIc-matrix composites A more promising approach to synthesis of porous ZrSiO4-fiber coatings might involve formation of a coating green body that is a uniform dispersion of fine ZrSiO and carbon powder. Matrix densification and subsequent carbon burnout would have to be done below the temperature at which ZrSiOa is carbothermal 10m reduced, and the Zrsio and carbon particle sizes would have to be significantly smaller than coating thicknesses, typically 100 nm. Another possibility is use of a flux for ZrSiO4 formation that Fig 10. TEM images of ZNS-C-derived coatings on NextelTM 720 that functions in a reducing atmosphere, without affecting fibers show densification of the coating and oxidation of carbon in air that yields a dispersion of t-ZrO2 particles in SiOz. Coatings were applied at 1000C in argon and heat-treated for I and 100 h at 1000.C in ai In earlier studies, it was noted that surface-active decomposition products were more easily trapped in coatings that were less porous and more likely to be locally hermetic. The SiOz that formed from the ZNS-C precursor at high temperatures in air formed a dense, t Zirconia hermetic coating(Fig. 10)that may have trapped even very small amounts of surface-active species at very high local partial pressures. This may have accounted for the unusually low strengths of Nextel 720 and Hi-Nicalon-s with ZNS-C-derived coatings after high- temperature exposure to air. However, very small amounts of Sio2 might have been expected to seal fiber surface flaws and increase fiber strength, if behavior was analogous to that observed for trace AlPO lass in monazite-coated fibers .5 V. Summary and Conclusions Methods using fluxes that readily form ZrSiO Nextel720 fiber 1200-1400C do not do so in the presence of carbon. The 20 nm ommonly used V,Os liquid-phase flux is reduced in the presence
In earlier studies, it was noted that surface-active decomposition products were more easily trapped in coatings that were less porous and more likely to be locally hermetic.51 The SiO2 that formed from the ZNS-C precursor at high temperatures in air formed a dense, hermetic coating (Fig. 10) that may have trapped even very small amounts of surface-active species at very high local partial pressures. This may have accounted for the unusually low strengths of Nextel 720 and Hi-Nicalon-S with ZNS-C-derived coatings after hightemperature exposure to air. However, very small amounts of SiO2 might have been expected to seal fiber surface flaws and increase fiber strength, if behavior was analogous to that observed for trace AlPO4 glass in monazite-coated fibers.51 IV. Summary and Conclusions Methods using fluxes that readily form ZrSiO4 in air at 1200°–1400°C do not do so in the presence of carbon. The commonly used V2O5 liquid-phase flux is reduced in the presence of carbon and, therefore, is ineffective. In an open system, ZrSiO4 can be carbothermally reduced. Flux methods also are inappropriate for fiber coatings because flux can be easily lost to the fibers, where it may promote grain growth and degrade fiber strength. Finely dispersed multiphase fiber coatings of t-ZrO2–SiO2– carbon can be deposited in argon from appropriate precursors. However, on heat treatment in air at 1000°C, porosity from the fugitive carbon phase is rapidly filled by SiO2, and dense t-ZrO2–SiO2 coatings form. Consequently, such t-ZrO2–SiO2 coatings are not weak and, therefore, are not useful as fiber–matrix interfaces in ceramic-matrix composites. A more promising approach to synthesis of porous ZrSiO4-fiber coatings might involve formation of a coating green body that is a uniform dispersion of fine ZrSiO4 and carbon powder. Matrix densification and subsequent carbon burnout would have to be done below the temperature at which ZrSiO4 is carbothermally reduced, and the ZrSiO4 and carbon particle sizes would have to be significantly smaller than coating thicknesses, typically 100 nm. Another possibility is use of a flux for ZrSiO4 formation that Fig. 10. TEM images of ZNS-C-derived coatings on NextelTM 720 that functions in a reducing atmosphere, without affecting fibers. show densification of the coating and oxidation of carbon in air that yields a dispersion of t-ZrO2 particles in SiO2. Coatings were applied at 1000°C in argon and heat-treated for 1 and 100 h at 1000°C in air. Fig. 11. TEM images of ZNS-C-derived coatings on NextelTM 720 that show coarsening of ZrO2 particles in SiO2 and transformation to a twinned monoclinic phase. Coatings were applied at 1000°C in argon and heattreated for 1 and 100 h at 1200°C in air. Fig. 12. TEM images of ZP-C-derived coatings on NextelTM 720 that show t-ZrO2 particles without SiO2. Coatings were applied at 1000°C in argon and heat-treated for 1 h at 1000°C in air. 1974 Journal of the American Ceramic Society—Boakye et al. Vol. 87, No. 10
October 2004 Zirconia-Silica-Carbon Coatings on Ceramic Fibers 1975 2.5 2.2 四5苏鸟 ZN-C 1.5 0.2? ZNS-C 15 Precursor Wt Loss% above t Fig 14. Strengths of ZrOx-SiO-carbon- and ZrO carbon-coated Nex. telM 720 after heat treatment for I h at 1000 C in air. Strengths are plotted 0.5 a function of coating precursor weight loss above 1000C(see Fig. 3). Results are mposed on results for monazite-coated Nextel 720 from an earlier study(Ref. 51). The weight loss value for ZNS-C is questionable because of weight gain that may be from oxidation of trace carbide. Long-Term Exposure at High Temperature, "J. Am. Ceram. Soc., 86(21 325-32 C Kaya, E. G. Butler, A Selcuk, A R. Boccaccini, and M. H. Lewis, "Mullite 40060080010001200 Properties, J. Eur. Ceram. Soc., 22, 2333-42(2002). E. D. Morgan and D. B. Marshall, "Ceramic Composites of Monazite and Temperature(C) eram.Soe.,7861553-63(1995) tel 720 fibers after coating and heat treatment in air and roated Nex. Oxide-Oxide Composites,"J. Eur. Ceram Soc, 20(51583-87(200Containing and Performance of an All-Oxide Ceramic Composite. "J Am. Cera Soc., 81 [8] 2077-86(1998) C. G. Levi and F. w. Zok. "Effects e on the Mechanical Properties of a Porous-Matrix. All-Oxide Ceramic Composi Hi-Nicalon s and Nextel720 fiber strengths are not "Performance of Four Ceramic-Matrix Com- significantly affected by coating as long as carbon remains in the posite Divergent Flap Inserts Following Ground Testing on an F110 Turbofan coating. Coated-fiber strengths degrade when carbon is removed by heat treatment in air at 21000C. Strength degradation of 上VcmH1123w面 of thermalA Hi-Nicalon s is so severe that thorough characterization of relationships between strength, temperature, and microstructures low-C. Tu. F. F. Lange. and A. G. Evans, "Concept for a Damage-Tolerant cannot be done Ceramic Composite with"" Interfaces, J. Am. Ceram. Soc., 79 [2]417-24 Degradation of Nextel 720 fibers is more severe when SiO, is present and when coating precursors contain a nitrate. A small Interface. Ceram. Eng. Sci. Proc. 16(41497-505 (1995). ant Fiber Coating for CMC containing coatings. bu with monazite-coated fibers, this1"E. BPakite, ". Mater. Res, 12 (5)1287-96(1997. ber Coating in SiC-Si, amount of fiber grain growth may be associated with Zro rain growth does not correlate with strength degradation. Degra R. S. Hay, and M. D. Petry, "Mixed Carbon-Aluminum Oxide dation may be caused by high-temperature stress corrosion from a Geochemisty of Oxides, 0onshandr stides and elated Ma en al. Edited by l. voight. surface-active decomposition product of the coating precursors T Wood, B Bunker, and B. casey. Materials Research Socicty. Pittsburgh, PA 1997 that contains nitrogen. This also is consistent with previous conclusions for monazite-coated fibers. This gas may be trapped and Oxide Coatings on Continuous Ceramic Fibers": Pp. 377-82 in Ceramic Matrix at higher local partial pressures by dense coatings, such as those aterials. M: 365. Edited by R. A. Lowden, M. K. Ferber that contain SiO,, and, consequently, degradation may be more J.R. Hellmann, and S G. DiPetro. Materials Research Society. Pittsburgh, PA, 1995 serious. It is possible that surface-active carbon in a coating may scavenge this gas and, consequently, decrease the fiber strength Porous Zio=SiO, and Monazite Coatings using Nextel ni 720 Fiber- Reinton d IeM. 3. 0Brien and B. W. Sheldon. "Porous Alumina Coating with Ta opposites. "J Am. Ceran. Soc., 82(1235 (1999) References J.B. Davis, AKristoffersson, E Carlstrom, and W.I.Clegg."Fabrication and tion in Ceramic Laminates with Porous Interlayers, "J. Am. Ceram Soc IR. J. Kerans, R. S. Hay, T. A. Parthasarathy, and M. K. Cinibulk, 8310]2369-74(2000 Design for Oxidation-Resistant Ceramic Composites, J. Am. Ceran. Soc., 85 [11] K.A. Keller. T Mah, T.A. Parthasarathy, EE Boakye, P. Mogilevsky, and M.K. Kanno, "Thermodynamic and Crystallographic Discussion of the Formation Cinibulk,"Effectiveness of Monazite Coatings in Oxide/Oxide Composites after and Dissociation of Zircon J. Mater. Sci. 24. 2415-20(1989)
Hi-Nicalon STM and NextelTM 720 fiber strengths are not significantly affected by coating as long as carbon remains in the coating. Coated-fiber strengths degrade when carbon is removed by heat treatment in air at 1000°C. Strength degradation of Hi-Nicalon S is so severe that thorough characterization of relationships between strength, temperature, and microstructures cannot be done. Degradation of Nextel 720 fibers is more severe when SiO2 is present and when coating precursors contain a nitrate. A small amount of fiber grain growth may be associated with ZrO2- containing coatings, but, as with monazite-coated fibers,53 this grain growth does not correlate with strength degradation. Degradation may be caused by high-temperature stress corrosion from a surface-active decomposition product of the coating precursors that contains nitrogen. This also is consistent with previous conclusions for monazite-coated fibers.51 This gas may be trapped at higher local partial pressures by dense coatings, such as those that contain SiO2, and, consequently, degradation may be more serious. It is possible that surface-active carbon in a coating may scavenge this gas and, consequently, decrease the fiber strength degradation. References 1 R. J. Kerans, R. S. Hay, T. A. Parthasarathy, and M. K. Cinibulk, “Interface Design for Oxidation-Resistant Ceramic Composites,” J. Am. Ceram. Soc., 85 [11] 2599–632 (2002). 2 K. A. Keller, T. Mah, T. A. Parthasarathy, E. E. Boakye, P. Mogilevsky, and M. K. Cinibulk, “Effectiveness of Monazite Coatings in Oxide/Oxide Composites after Long-Term Exposure at High Temperature,” J. Am. Ceram. Soc., 86 [2] 325–32 (2003). 3 C. Kaya, E. G. Butler, A. Selcuk, A. R. Boccaccini, and M. H. 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Soc., 83 [7] 1727–38 (2000). 9 E.A. V. Carelli, H. Fujita, J. Y. Yang, and F. W. Zok, “Effects of Thermal Aging on the Mechanical Properties of a Porous-Matrix Composite,” J. Am. Ceram. Soc., 85 [3] 595– 602 (2002). 10W.-C. Tu, F. F. Lange, and A. G. Evans, “Concept for a Damage-Tolerant Ceramic Composite with ‘Strong’ Interfaces,” J. Am. Ceram. Soc., 79 [2] 417–24 (1996). 11L. U. J. T. Ogbuji, “A Porous, Oxidation-Resistant Fiber Coating for CMC Interface,” Ceram. Eng. Sci. Proc., 16 [4] 497–505 (1995). 12L. U. J. T. Ogbuji, “Evaluation of a Porous Fiber Coating in SiC–Si3N4 Minicomposite,” J. Mater. Res., 12 [5] 1287–96 (1997). 13E. Boakye, R. S. Hay, and M. D. Petry, “Mixed Carbon–Aluminum Oxide Coatings from Aqueous Sols and Solutions”; pp 363– 68 in Aqueous Chemistry and Geochemistry of Oxides, Oxyhydroxides, and Related Materials. Edited by J. Voight, T. Wood, B. Bunker, and B. Casey. Materials Research Society, Pittsburgh, PA, 1997. 14R. S. Hay, M. D. Petry, K. A. Keller, M. K. 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