E噩≈S Journal of the European Ceramic Society 20(2000)1505-1514 Degradation at 1200C of a SiC coated 2D-Nicalon/C/SIC composite processed by SICFILLR method C. Badini a,*. P Fino a G. Ubertallia F. Taricco b Dipartimento di scienza dei Materiali e Ingegneria Chimica, Politecnico di Torino, Corso Duca degli Abruzzi, 24-10129 Turin, ital Fiat avio. Turin. ita Received 28 June 1999; received in revised form 25 October 1999; accepted 4 November 1999 Abstract The thermal stability of a 2D-Nicalon/C/SiC composite was studied through the variation of both mechanical properties and microstructure occurring during heat treating. The composite was processed by infiltration of Sic preforms according to sIC improve the thermal stability a CVI layer was deposited on the carbon interphase and the specimen surfaces were CVD covered by an external sic seal coating about 165 um thick. The aging tests were carried out at 1200C in air or in non oxidizing environment (vacuum). Other specimens were thermally cycled between 25 and 1150 C. Three point bending tests and Charpy impact measure- ments were performed before and after these treatments. The composite microstructure was investigated by scanning electron microscope(SEM), electron probe microanalysis(EPMA), X-ray diffraction(XRD), reflectance infrared spectroscopy(FTIR) and ce area BET measurements. The as-processed material showed a modulus of rupture(MOR) of 483 MPa and appreciable toughness. These characteristics were retained after aging(200 h at 1200oC) under vacuum. Air thermal treatments caused heavy loss of strength and increase of brittleness. Strong oxidation occurred during these last treatments at both the carbon interlayer and the matrix, while the Sic external sample coating was not oxidized. The oxygen needed for composite bulk oxidation flowed through the Sic coating due to the occasional presence of very few structural defects. C 2000 Elsevier Science Ltd. All rights Keywords: Composites; Interfaces: SiC; Thermal shock resistance 1. Introduction about 3. 2 g/cm). Finally, the mechanical properties of Sic/SiC composite can be retained at high temperatures Long fiber-reinforced ceramic matrix composites have and under severe service environments. Generally been proposed as advanced materials suitable for struc- speaking, most of ceramics and ceramic/ceramic com tural applications. In particular, in the last years many posites show better oxidation resistance than metal efforts have been devoted to the development of SiCo/ alloys. However, many industrial applications (i.e Sic composites. These composites show some attractive the aerospace field) require the development of new properties and advantages over traditional ceramics oxidation resistant materials able to work in extreme higher tensile and flexural strength (provided by the conditions(temperature above 1000 C and oxidizing continuous fiber reinforcement). enhanced fracture toughness and impact resistance (chiefly achieved by In these conditions also the Sic/Sic composites can tailoring the fiber/matrix interface characteristics). Fur- suffer degradation. In fact, at temperatures higher than thermore, their specific strength and modi uus are 1000C both the polymer-derived(Nicalon) and the greater than those of many other structural materials chemically vapor-deposited(CVD) SiC fibers undergo (metal alloys or ceramic)because both fiber and ceramic damaging through decomposition and consequent eva matrix are made of Sic(theoretical density for Sic poration of gaseous species. The vaporizing compounds responsible for degradation chiefly are CO, in the case 4 Corresponding author. Fax: 39-11-564 of CVD fibers, and Sio plus CO for Nicalon fibers. @athena polito. it(C The oxygen needed for the formation of Sio and CO 0955-2219/00/S- see front matter C 2000 Elsevier Science Ltd. All rights reserved PII:S0955-2219(00)000297
Degradation at 1200C of a SiC coated 2D-Nicalon/C/SiC composite processed by SICFILL1 method C. Badini a,*, P. Fino a , G. Ubertalli a , F. Taricco b a Dipartimento di Scienza dei Materiali e Ingegneria Chimica, Politecnico di Torino, Corso Duca degli Abruzzi, 24-10129 Turin, Italy bFiat Avio, Turin, Italy Received 28 June 1999; received in revised form 25 October 1999; accepted 4 November 1999 Abstract The thermal stability of a 2D-Nicalon/C/SiC composite was studied through the variation of both mechanical properties and microstructure occurring during heat treating. The composite was processed by in®ltration of SiC preforms according to SICFILL1 method. The material toughness was enhanced by a carbon interphase put between the ®bers and the matrix. In order to improve the thermal stability a CVI layer was deposited on the carbon interphase and the specimen surfaces were CVD covered by an external SiC seal coating about 165 mm thick. The aging tests were carried out at 1200C in air or in non oxidizing environment (vacuum). Other specimens were thermally cycled between 25 and 1150C. Three point bending tests and Charpy impact measurements were performed before and after these treatments. The composite microstructure was investigated by scanning electron microscope (SEM), electron probe microanalysis (EPMA), X-ray diraction (XRD), re¯ectance infrared spectroscopy (FTIR) and surface area BET measurements. The as-processed material showed a modulus of rupture (MOR) of 483 MPa and appreciable toughness. These characteristics were retained after aging (200 h at 1200C) under vacuum. Air thermal treatments caused heavy loss of strength and increase of brittleness. Strong oxidation occurred during these last treatments at both the carbon interlayer and the matrix, while the SiC external sample coating was not oxidized. The oxygen needed for composite bulk oxidation ¯owed through the SiC coating due to the occasional presence of very few structural defects. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: Composites; Interfaces; SiC; Thermal shock resistance 1. Introduction Long ®ber-reinforced ceramic matrix composites have been proposed as advanced materials suitable for structural applications. In particular, in the last years many eorts have been devoted to the development of SiCf/ SiC composites. These composites show some attractive properties and advantages over traditional ceramics: higher tensile and ¯exural strength (provided by the continuous ®ber reinforcement), enhanced fracture toughness and impact resistance (chie¯y achieved by tailoring the ®ber/matrix interface characteristics). Furthermore, their speci®c strength and modulus are greater than those of many other structural materials (metal alloys or ceramic) because both ®ber and ceramic matrix are made of SiC (theoretical density for SiC about 3.2 g/cm3 ). Finally, the mechanical properties of SiC/SiC composite can be retained at high temperatures and under severe service environments. Generally speaking, most of ceramics and ceramic/ceramic composites show better oxidation resistance than metal alloys. However, many industrial applications (i.e. in the aerospace ®eld) require the development of new oxidation resistant materials able to work in extreme conditions (temperature above 1000C and oxidizing atmosphere). In these conditions also the SiC/SiC composites can suer degradation. In fact, at temperatures higher than 1000C both the polymer-derived (Nicalon) and the chemically vapor-deposited (CVD) SiC ®bers undergo damaging through decomposition and consequent evaporation of gaseous species. The vaporizing compounds responsible for degradation chie¯y are CO, in the case of CVD ®bers, and SiO plus CO for Nicalon ®bers.1,2 The oxygen needed for the formation of SiO and CO 0955-2219/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(00)00029-7 Journal of the European Ceramic Society 20 (2000) 1505±1514 * Corresponding author. Fax: +39-11-564-4699. E-mail address: badini@athena.polito.it (C. Badini)
1506 C. Badini et al. Journal of the European Ceramic Society 20 (2000)1505-1514 may come from both the fiber itself (which contains a Both these reactions, involving the formation of gas non negligible percentage of this element) and an oxi eous carbon monoxide. can be slowed down if an dizing environment. The heat treatment of Nicalon atmosphere of CO or a SiC seal-coating(deposited by fibers performed over 1200 C in air also causes a further Cvd on the surface of the composite specimen) are SiC fiber oxidation, which results in the formation of a adopted. Anyway, a progressive weakening of the SiO2 surface layer. On the other hand, according to interfacial bond between fiber and matrix(and the con Mah et al., this silica layer prevents the reaction pro- sequent strength decrease) was also observed during duct evaporation and slows down the loss of fiber aging carried out in the conditions described above. The strength, which, for this reason, occurs more quickly iC coating is able to retain after a treatment of 80 h at under vacuum than in air. Furthermore, a Sic grain 1200oC, about the 70% of the untreated material tensile growth happens at the fiber core during the thermal strength. More prolonged thermal treatments at 1200oC treatment and concurs to decrease the fiber mechanical or at higher temperatures result in a more marked trength degradation due to either grain growth of Sic inside the The thermal stability of a SiCr/SiC composite greatly fibers or formation of a Sio2+C layer according to the depends on that of the SiC fibers, even though the following reaction degradation mechanism of this kind of composites is affected by some changes occurring at the fiber/matrix 2CO(g)+ SiC interface and by the presence of protective sealing coat- CO(g+ SiO(g SiOz(s)+ C(s) ings deposited on the composite surface as well In order to understand the phenomena happening at According to literature, an external coating of siC is the fiber/matrix interface, it is to be considered that the also suitable for avoiding the oxidative degradation of a SiC fibers are frequently coated with a CVD carbon SiC/C/SiC composite caused by a thermal treatment at layer, referred to as the interphase, prior to the infiltra- 1200C carried out under air. tion of the matrix, with the purpose of allowing fiber In conclusion, literature data put in evidence that an matrix debonding under stress. Debonding is one of the external SiC coating enhances the SiC/Sic thermal sta mechanisms responsible for an increase of fracture bility, hindering both composite oxidation and decom work, which results in an higher strain and in an position reaction of SiC fibers enhanced toughness. However, Filippuzzi et al. put in This paper deals with the thermal stability of a 2D- evidence that this interphase of carbon is prone to Nicalon /C/SiC composite processed by a new infiltra undergo oxidation, giving gaseous carbon monoxide tion way(SICFILL method developed by F.N. S pA and leaving the fiber surface free to react with oxygen Boscomarengo, Alessandria, Italy) ? and coated with and to form a SiO2 layer CVD (or Cvi)-SiC protective layers. The thermal stability in the temperature range of The changes occurring in both microstructure and 1000-1300oC of a 2D-Nicalon/C/SiC composite under mechanical features during composite aging, performed different aging environments(vacuum, atmosphere of in different conditions, have been investigated argon or carbon monoxide) has been well investigated by Labrugere et al. 4. 5 These authors studied the beha vior of a composite produced(by SEP) according to 2. Experimental three steps: preparation of a 2D preform by pressing together fabrics of Nicalon fibers; deposition of a pyr- 2. 1. Materials and methods ocarbon layer (less than I um thick) on the fiber surface isothermal/isobaric chemical vapor infiltration (ICVI) The composite material under investigation was fab of the preform to obtain an in-situ SiC matrix ricated by a hybrid method involving CVI, CVD and During the aging treatment of this composite per- polymer impregnation-pyrolysis(PIP), starting from a formed under argon or vacuum, several phenomena 2D-preform(obtained by lay-up of 100x 100 mm fabrics ccur. Firstly, fibers undergo decomposition through of CG Nicalon fibers). The production process included the reaction the following steps SiO2xCI-x- SiOg+(2x-1)CO(g+(2-3x)Ceg deposition of a pyrocarbon layer(about 0.3 um thick)on the fiber surface by chemical vapor infill A subsequent reaction causes the destruction of the tration(CVD), carried out at 1100C using a CH4/ yrocarbon layer and the growth of large Sic crystals H, gaseous mixture on the fiber surface. deposition(by CVi) of a second surface layer of Sic (about 2 um thick), performed at about siog)+2Cs→SCs)+CO 1200C using a methyltrichlorosilane/hydrogen MTS/H2) mixture
may come from both the ®ber itself (which contains a non negligible percentage of this element) and an oxidizing environment. The heat treatment of Nicalon ®bers performed over 1200C in air also causes a further SiC ®ber oxidation, which results in the formation of a SiO2 surface layer. On the other hand, according to Mah et al.,1 this silica layer prevents the reaction product evaporation and slows down the loss of ®ber strength, which, for this reason, occurs more quickly under vacuum than in air. Furthermore, a SiC grain growth happens at the ®ber core during the thermal treatment and concurs to decrease the ®ber mechanical strength. The thermal stability of a SiCf/SiC composite greatly depends on that of the SiC ®bers, even though the degradation mechanism of this kind of composites is aected by some changes occurring at the ®ber/matrix interface and by the presence of protective sealing coatings deposited on the composite surface as well. In order to understand the phenomena happening at the ®ber/matrix interface, it is to be considered that the SiC ®bers are frequently coated with a CVD carbon layer, referred to as the interphase, prior to the in®ltration of the matrix, with the purpose of allowing ®ber/ matrix debonding under stress. Debonding is one of the mechanisms responsible for an increase of fracture work, which results in an higher strain and in an enhanced toughness. However, Filippuzzi et al.3 put in evidence that this interphase of carbon is prone to undergo oxidation, giving gaseous carbon monoxide and leaving the ®ber surface free to react with oxygen and to form a SiO2 layer. The thermal stability in the temperature range of 1000±1300C of a 2D-Nicalon/C/SiC composite under dierent aging environments (vacuum, atmosphere of argon or carbon monoxide) has been well investigated by LabrugeÁre et al.4,5 These authors studied the behavior of a composite produced (by SEP) according to three steps: preparation of a 2D preform by pressing together fabrics of Nicalon ®bers; deposition of a pyrocarbon layer (less than 1 mm thick) on the ®ber surface; isothermal/isobaric chemical vapor in®ltration (ICVI) of the preform to obtain an in-situ SiC matrix. During the aging treatment of this composite performed under argon or vacuum, several phenomena occur. Firstly, ®bers undergo decomposition through the reaction: SiO2xC1ÿx ! SiO g 2x ÿ 1CO g 2 ÿ 3xC g A subsequent reaction causes the destruction of the pyrocarbon layer and the growth of large SiC crystals on the ®ber surface: SiO g 2C s ! SiC s CO g Both these reactions, involving the formation of gaseous carbon monoxide, can be slowed down if an atmosphere of CO or a SiC seal-coating (deposited by CVD on the surface of the composite specimen) are adopted. Anyway, a progressive weakening of the interfacial bond between ®ber and matrix (and the consequent strength decrease) was also observed during aging carried out in the conditions described above. The SiC coating is able to retain, after a treatment of 80 h at 1200C, about the 70% of the untreated material tensile strength. More prolonged thermal treatments at 1200C or at higher temperatures result in a more marked degradation due to either grain growth of SiC inside the ®bers or formation of a SiO2+C layer according to the following reactions: 2CO g SiC s ! SiO2 s 3C s CO g SiO g ! SiO2 s C s According to literature,6 an external coating of SiC is also suitable for avoiding the oxidative degradation of a SiC/C/SiC composite caused by a thermal treatment at 1200C carried out under air. In conclusion, literature data put in evidence that an external SiC coating enhances the SiCf/SiC thermal stability, hindering both composite oxidation and decomposition reaction of SiC ®bers. This paper deals with the thermal stability of a 2DNicalon/C/SiC composite processed by a new in®ltration way (SICFILL1 method developed by F.N. S.p.A., Boscomarengo, Alessandria, Italy)7 and coated with CVD (or CVI)-SiC protective layers. The changes occurring in both microstructure and mechanical features during composite aging, performed in dierent conditions, have been investigated. 2. Experimental 2.1. Materials and methods The composite material under investigation was fabricated by a hybrid method involving CVI, CVD and polymer impregnation-pyrolysis (PIP), starting from a 2D-preform (obtained by lay-up of 100100 mm fabrics of CG Nicalon ®bers). The production process included the following steps: . deposition of a pyrocarbon layer (about 0.3 mm thick) on the ®ber surface by chemical vapor in®ltration (CVI), carried out at 1100C using a CH4/ H2 gaseous mixture; . deposition (by CVI) of a second surface layer of SiC (about 2 mm thick), performed at about 1200C using a methyltrichlorosilane/hydrogen (MTS/H2) mixture; 1506 C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514
C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 preform liquid infiltration under vacuum by a (SEM)and electron probe microanalysis(EPMA, Jeol slurry composed of polycarbosilane(PCS), xylene Superprobe JXA-8600), infrared spectroscopy, BET and crystalline p-SiC nanopowders produced by measurements laser assisted synthesis at XRD measurements were carried out on the sample solvent evaporation at room temperature and pyr- surface and repeated after progressive mechanical olysis of PCS at 1 C under inert atmosphere removal of external parts(with definite thickness)of the further six densification steps(carried out by PIP samples. Grazing angle XRD were performed on the process again, but without the addition of Sic sample surfaces and repeated after mechanical removal particles to the slurry) suitable for achieving a of a surface layer about 40 um thick. density of the composite plate of about 2.2 g/cm EPMA analyses were done on the transversal sections of the composite specimens: silicon, carbon and oxygen The composite plates were then machined by a cutting were analyzed. Wavelength dispersion spectrometers machine equipped with a diamond saw in order to (WDS), equipped with thallium acid phthalate crystal obtain 50x10x4 mm' samples. These samples were d =25.757 A)for Si, layered dispersion element coated by CVD(MTS/H2 precursors, temperature of multilayer(2d=60 A)for O and layered dispersion ele 1200C)obtaining a seal layer of SiC about 165 um ment C multilayer (2d=99.1 A)for C, were used to thick and a final density of 2.56 g/cm detect the K line of these different elements Standards Both this last coating and the previously CVI-depos- of pure SiC(Sigma-Aldrich) and pure quartz were used ited on the carbon interphase were placed with the aim for calibration. After calibration the analysis of each of of avoiding the composite oxidation as well as the eva- these standards gave reproducible results, differing no poration of fiber decomposition products more than I at% from the nominal standard composi The composite samples were submitted to the follow- tion. The sample analyses were repeated several times different thermal treatments. (and the results were averaged) in each of the different thermal aging at 1200 C in air for periods up to pecimen parts: fiber core, external part of the fibers 200h; dual C/SiC interphase, matrix and external coating of SiC. Also, the thin fiber coatings were well distinguish thermal aging for 200 h at 1200 C under vacuum able in the sample sections of the untreated samples, (to this purpose the samples were sealed under vacuum in a silica tube) however, the electron probe(about I um in size)cannot thermal cycling (1000 cycles) between 25 and be completely contained inside the thin layer of pyr carbon. For this reason, the results of the analysis 1150 C(each cycle was performed by keeping the performed on this fiber coating were affected by the samples in a tubular furnace for 20 min, taking the samples out of the furnace and leaving them to presence of the fiber and the Sic interphase, neigh bouring with the analyze Diffuse reflectance Fourier transformed infrared Bars of as-processed composite were cut in several spectroscopy (DR-FTIR; Bruker IFS66 instrument slices (in the parallel and transversal directions with equipped with MCT-Cryodetector) was used with the respect to the bar major axis) and the section surfaces ulm of checking the presence of Si-o bonds(arising were observed by SEM in order to test the homogeneity from SiC oxidation) on the sample surface; this analysis of the external sic coating was repeated after removal of a 40 um thick layer from The mechanical behavior of the as-prepared compo- the external Sic sample coating site was compared with that of the treated samples by Following the results obtained with the methods bending tests The flexural th was described above, it was considered necessary to perform measured, as average of three tests, using a Sintech 10d further experiments aimed at assessing the capability of equipment, with support span of 40 mm and crosshead the Sic coating of avoiding the gas penetration inside speed of 1 mm/min the material. Indirect indications about this coating Some specimens, of the as-processed material and of characteristic were obtained by measuring the surface the composite aged in extreme conditions, were also area of the composite bars by BET adsorption iso- ubmitted to Charpy test. An instrumented Izod- therms of nitrogen at 77 K( Carlo Erba Sorptomatic arpy equipment(ATS-FAAR) with a 3. 5 kg hammer 1800 instrument) was used to obtain the impact-load/time curves and to calculate the work of fracture. The microstructure of 2. 2. Experimental results and discussion untreated and aged samples was studied by different techniques: X-ray diffraction(XRD) and grazing-angle 2.2.1. Mechanical tests XRD(Philips diffractometer equipped with a PW3020 The flexural strength and the Charpy impact resis- goniometer for grazing angle measurements, Cu Ka), tance of the composite samples in the thermal treated optical microscopy, scanning electron microscopy and untreated conditions are compared in Table I
. preform liquid in®ltration under vacuum by a slurry composed of polycarbosilane (PCS), xylene and crystalline b-SiC nanopowders produced by laser assisted synthesis at ENEA; . solvent evaporation at room temperature and pyrolysis of PCS at 1100C under inert atmosphere; . further six densi®cation steps (carried out by PIP process again, but without the addition of SiC particles to the slurry) suitable for achieving a density of the composite plate of about 2.2 g/cm3 . The composite plates were then machined by a cutting machine equipped with a diamond saw in order to obtain 50104 mm3 samples. These samples were coated by CVD (MTS/H2 precursors, temperature of 1200C) obtaining a seal layer of SiC about 165 mm thick and a ®nal density of 2.56 g/cm3 . Both this last coating and the previously CVI-deposited on the carbon interphase were placed with the aim of avoiding the composite oxidation as well as the evaporation of ®ber decomposition products. The composite samples were submitted to the following dierent thermal treatments: . thermal aging at 1200C in air for periods up to 200 h; . thermal aging for 200 h at 1200C under vacuum (to this purpose the samples were sealed under vacuum in a silica tube); . thermal cycling (1000 cycles) between 25 and 1150C (each cycle was performed by keeping the samples in a tubular furnace for 20 min, taking the samples out of the furnace and leaving them to cool in stationary air) Bars of as-processed composite were cut in several slices (in the parallel and transversal directions with respect to the bar major axis) and the section surfaces were observed by SEM in order to test the homogeneity of the external SiC coating. The mechanical behavior of the as-prepared composite was compared with that of the treated samples by three point bending tests. The ¯exural strength was measured, as average of three tests, using a Sintech 10D equipment, with support span of 40 mm and crosshead speed of 1 mm/min. Some specimens, of the as-processed material and of the composite aged in extreme conditions, were also submitted to Charpy test. An instrumented IzodCharpy equipment (ATS-FAAR) with a 3.5 kg hammer was used to obtain the impact-load/time curves and to calculate the work of fracture. The microstructure of untreated and aged samples was studied by dierent techniques: X-ray diraction (XRD) and grazing-angle XRD (Philips diractometer equipped with a PW3020 goniometer for grazing angle measurements, Cu Ka), optical microscopy, scanning electron microscopy (SEM) and electron probe microanalysis (EPMA, JeolSuperprobe JXA-8600), infrared spectroscopy, BET measurements. XRD measurements were carried out on the sample surface and repeated after progressive mechanical removal of external parts (with de®nite thickness) of the samples. Grazing angle XRD were performed on the sample surfaces and repeated after mechanical removal of a surface layer about 40 mm thick. EPMA analyses were done on the transversal sections of the composite specimens: silicon, carbon and oxygen were analyzed. Wavelength dispersion spectrometers (WDS), equipped with thallium acid phthalate crystal (2d=25.757 AÊ ) for Si, layered dispersion element 1 multilayer (2d=60 AÊ ) for O and layered dispersion element C multilayer (2d=99.1 AÊ ) for C, were used to detect the Ka line of these dierent elements. Standards of pure SiC (Sigma-Aldrich) and pure quartz were used for calibration. After calibration the analysis of each of these standards gave reproducible results, diering no more than 1 at% from the nominal standard composition. The sample analyses were repeated several times (and the results were averaged) in each of the dierent specimen parts: ®ber core, external part of the ®bers, dual C/SiC interphase, matrix and external coating of SiC. Also, the thin ®ber coatings were well distinguishable in the sample sections of the untreated samples, however, the electron probe (about 1 mm in size) cannot be completely contained inside the thin layer of pyrocarbon. For this reason, the results of the analysis performed on this ®ber coating were aected by the presence of the ®ber and the SiC interphase, neighbouring with the analyzed area. Diuse re¯ectance Fourier transformed infrared spectroscopy (DR-FTIR; Bruker IFS66 instrument, equipped with MCT-Cryodetector) was used with the aim of checking the presence of Si±O bonds (arising from SiC oxidation) on the sample surface; this analysis was repeated after removal of a 40 mm thick layer from the external SiC sample coating. Following the results obtained with the methods described above, it was considered necessary to perform further experiments aimed at assessing the capability of the SiC coating of avoiding the gas penetration inside the material. Indirect indications about this coating characteristic were obtained by measuring the surface area of the composite bars by BET adsorption isotherms of nitrogen at 77 K (Carlo Erba Sorptomatic 1800 instrument). 2.2. Experimental results and discussion 2.2.1. Mechanical tests The ¯exural strength and the Charpy impact resistance of the composite samples in the thermal treated and untreated conditions are compared in Table 1. C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514 1507
iery20(2000)1505-1514 Table I treated in air gave a very different response(Fig Mechanical properties of untreated and thermal treated composite curve C). In this last case, the material failed in a brittle samples mode, after much lower deformation and stress(cross- Material Bending harpy head displacement of about 0. 1 mm and stress of 71 strength fracture work MPa only) MOR MPa] E[GPa][kJ/m The thermal treatments performed in air result in a marked decrease of modulus of rupture(MOR), but do not cause Youngs modulus variations After200hatl200°C After 200 h at 1200%C. the residual flexural strength under vacuum After 200 h at 1200C in air 71 was about only the 15% of that showed by the as-fab ricated composite. Thermal cycling carried out between (1000 cycles) 25 and 1150C (1000 cycles) caused a similar effect on he composite mechanical behavior The decrease of MOR strictly depended on the pro gress of oxidation phenomena, because samples aged A characteristic stress/displacement curve obtained by for 200 h at 1200C under vacuum maintained a high bending the untreated composite is reported in Fig. 1 MOR value(469 MPa) (curve A). This graph shows that the composite under- The embrittlement of the composite samples aged in went a severe deformation during flexural test. In the air at 1150-12000C was confirmed by Charpy tests. The first part of the test, the stress increase caused a pro- work of fracture measured by Charpy method greatly gressive crosshead displacement (up to about 0.6 mm), depended on the sample orientation with respect to the afterward, the load suddenly fell to about 150 MPa and impacting hammer: the fracture energy was greater remained practically constant meanwhile the displace- when the composite bar was put with its major section ment increased up to more than 1.2 mm. Then the test thickness (10 mm) perpendicularly to the falling ham was interrupted because deformations as large as these mer. The results reported in Table I refer to this sample were not compatible with the sample-holder geometry. orientation. The untreated composites showed an aver At the test end the specimens were heavily bent, but not age work of fracture(81 kJ/m2)much greater than that hared in two parts, the pull-out of fibers providing the of the samples thermally treated under air (ranging connection of the composite bar at the bending center. between 3. 6 and 4.2 kJ/m). Furthermore, the untreated The specimens aged under vacuum showed a similar composite specimens underwent a deformation process behavior(Fig. 1, curve B), while the samples thermal before fracture(for a period of 0.9 ms), contrary to the Displacement [mmI Fig 1. Three-point bending curve 2000 posite specimens: A=untreated sample; B= sample after 200 h of aging under vacuum at 1200C; C
A characteristic stress/displacement curve obtained by bending the untreated composite is reported in Fig. 1 (curve A). This graph shows that the composite underwent a severe deformation during ¯exural test. In the ®rst part of the test, the stress increase caused a progressive crosshead displacement (up to about 0.6 mm), afterward, the load suddenly fell to about 150 MPa and remained practically constant meanwhile the displacement increased up to more than 1.2 mm. Then the test was interrupted because deformations as large as these were not compatible with the sample-holder geometry. At the test end the specimens were heavily bent, but not shared in two parts, the pull-out of ®bers providing the connection of the composite bar at the bending center. The specimens aged under vacuum showed a similar behavior (Fig. 1, curve B), while the samples thermal treated in air gave a very dierent response (Fig. 1, curve C). In this last case, the material failed in a brittle mode, after much lower deformation and stress (crosshead displacement of about 0.1 mm and stress of 71 MPa only). The thermal treatments performed in air result in a marked decrease of modulus of rupture (MOR), but do not cause Young's modulus variations. After 200 h at 1200C, the residual ¯exural strength was about only the 15% of that showed by the as-fabricated composite. Thermal cycling carried out between 25 and 1150C (1000 cycles) caused a similar eect on the composite mechanical behavior. The decrease of MOR strictly depended on the progress of oxidation phenomena, because samples aged for 200 h at 1200C under vacuum maintained a high MOR value (469 MPa). The embrittlement of the composite samples aged in air at 1150±1200C was con®rmed by Charpy tests. The work of fracture measured by Charpy method greatly depended on the sample orientation with respect to the impacting hammer: the fracture energy was greater when the composite bar was put with its major section thickness (10 mm) perpendicularly to the falling hammer. The results reported in Table 1 refer to this sample orientation. The untreated composites showed an average work of fracture (81 kJ/m2 ) much greater than that of the samples thermally treated under air (ranging between 3.6 and 4.2 kJ/m2 ). Furthermore, the untreated composite specimens underwent a deformation process before fracture (for a period of 0.9 ms), contrary to the Table 1 Mechanical properties of untreated and thermal treated composite samples Material Bending strength Charpy fracture work MOR [MPa] E [GPa] [kJ/m2 ] Untreated 483 58 81 After 200 h at 1200C under vacuum 469 58 ± After 200 h at 1200C in air 71 60 3.6 Cyled between 25 and 1150C (1000 cycles) 44 59 4.2 Fig. 1. Three-point bending curves of composite specimens: A=untreated sample; B= sample after 200 h of aging under vacuum at 1200C; C= sample after 200 h of aging in air at 1200C. 1508 C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514
C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 air aged samples which immediately broke in a brittle debonding and pull-out(Fig 3), while the surface frac manner(Fig. 2) ture of samples aged at 1200C was completely flat(Fig The Charpy fracture surface of this two kinds of 4). On the other hand, the comparison of the transverse composite specimens was observed by scanning electron sections of these specimens, obtained cutting them some microscope. The fracture of untreated samples was mm apart from the fracture surface, put in evidence that characterized by important phenomena of fiber aging affects the interphase morphology(Figs. 5-7) This treatment also caused the fiber damage in several points(Fig. 6). This fiber damage, which likely can occur either during mechanical test or sample cutting also. is a clear evidence of the fiber embrittlement 2.22 Microstructure characterization 2.2.2.1. X-ray diffraction. The XRD patterns of the as- fabricated composite show that all the stronger peaks pertain to silicon carbide. However, after ablation of the more external part of the specimens, the XRD spectra present some changes In Fig. 8A the spectrum of the sample surface(pattern"a")and those recorded after Ilal wloa progressive removal by polishing of surface layers 200 0.9 Time (ns ⊥n、A△A△△A几△△A△A mm301kU 930E1 0558/01 sE Fig. 4. Flat Charpy fracture surface of a specimen after thermal Fig. 2. Charpy load/ time curves of composite specimens: a=untreated treatment(200 h) at 1200.C in air sample; b= sample aged in air for 200 h at 1200C 1mm38.1kU93E1g55181sE Fig 3. Charpy fracture surface of untreated specimen: fiber pull-out Fig. 5. Section of untreated composite ba
air aged samples which immediately broke in a brittle manner (Fig. 2). The Charpy fracture surface of this two kinds of composite specimens was observed by scanning electron microscope. The fracture of untreated samples was characterized by important phenomena of ®ber debonding and pull-out (Fig. 3), while the surface fracture of samples aged at 1200C was completely ¯at (Fig. 4). On the other hand, the comparison of the transverse sections of these specimens, obtained cutting them some mm apart from the fracture surface, put in evidence that aging aects the interphase morphology (Figs. 5±7). This treatment also caused the ®ber damage in several points (Fig. 6). This ®ber damage, which likely can occur either during mechanical test or sample cutting also, is a clear evidence of the ®ber embrittlement. 2.2.2. Microstructure characterization 2.2.2.1. X-ray diffraction. The XRD patterns of the asfabricated composite show that all the stronger peaks pertain to silicon carbide. However, after ablation of the more external part of the specimens, the XRD spectra present some changes. In Fig. 8A the spectrum of the sample surface (pattern ``a'') and those recorded after progressive removal by polishing of surface layers 200 Fig. 2. Charpy load/time curves of composite specimens: a=untreated sample; b=sample aged in air for 200 h at 1200C. Fig. 3. Charpy fracture surface of untreated specimen: ®ber pull-out. Fig. 5. Section of untreated composite bar. Fig. 4. Flat Charpy fracture surface of a specimen after thermal treatment (200 h) at 1200C in air. C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514 1509
1510 C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 and 400 um thick(patterns"b"and"c")are compared. 2.2. 2 2. Infrared spectroscopy. In order to confirm the These XRD patterns show that the SiC surface coating outcomes of XRD measurements regarding the forma- processed at 1200oC by CVD is better crystallized than tion of silica inside the composite samples aged in oxi- he Sic matrix produced by polycarbosilane pyrolysis at dative environment, infrared spectroscopy was used 1 100C. In fact, the intensities of the more important The infrared spectrum of the untreated sample surface peaks of B-SiC (JCPDS-ICDD card 29-1192)progres- did not show bands between 1000 and 1250 cm, that is sively become weaker and their shape widens with the in the characteristic range for silica bands. However, the increase of the distance from the sample surface; in the spectra of samples treated at 1200C or cycled in air meantime faint peaks likely belonging to a-SiC (ICDD showed bands in this range only when the analysis was card 42-1091) grow in the XRD pattern. The XRD of carried out on the sample surface, while these signals the sample surface shows few small, not identified disappeared after sample polishing(removal of a mate- flexes, that disappear as soon as the surface is rial layer about 40 um thick). These IR measurements polished. Fig 8B shows the XRD patterns of the sam- strengthen the XRD results, putting in evidence that the ples treated in air for 200 h at 1200oC. In these spectra a SiC coating oxidation occurs close to the sample surface new peak placed just below 20=22, which can be rea- only sonably attributed to cristobalite(ICDD card 39-1425) was observed. The intensity of this peak decreases 2.2.2.3. Electron probe microanalysis. The SEM-WDS moving from the surface toward the inner of the speci- analyses allow us to compare the chemical composition men, becoming close to zero at a distance of about 200 of composite samples showing a non-brittle behavior um from the surface. Afterward, its intensity grows in (composite specimens untreated or treated under the specimen core, at 400 um from the surface or more. vacuum)with that of composite bars aged in oxidizing This particular trend(observed for samples thermally environment(kept at 1200oC or cycled between 25 and cycled too) was confirmed by grazing-angle XRD mea l150°C) surements. The grazing-angle XRD patterns of the The microanalysis results collected in Table 2 show sample aged for 200 h at 1200oC are reported in Fig 9. that the atomic percent of C, Si and O, measured in the Pattern"a"refers to the sample surface and pattern"b" different part of the composite(fibers, matrix, inter- to the same specimen after removal of a layer 40 um phases and external coating), may appreciably change thick: the ablation resulted in a slight decrease in the after thermal treatment. intensity of the peak placed at 20=22. Furthermore, if Some of these composite parts display a non-homo- the XRD measure of the polished specimen(Fig. 9, geneous composition; for this reason, the composition pattern"b") is repeated adopting a Bragg-Brentano of both core and external part (labeled with the symbol geometry, that is using more penetrating X-rays, the ,)of Nicalon fibers and Sic bar coating are sepa spectrum"c"is obtained. This last spectrum shows that rately reported in Table 2. The changes in the material the peak at 20=22 is almost vanished composition can be better appreciated by considering On the basis of XRD measurements, it can be inferred the C/Si and O/Si atomic ratios, reported as histograms that oxidation occurred both on the sample surface and for untreated and thermal aged samples in Figs. 10 and in the sample core, but not (or not so much)in internal part of the Sic coating deposited on the com The SiC fiber coating and the sample coating of the posite bar surface untreated composite show a C/Si atomic ratio close to b 1\um Fig. 6. Fiber damage in the composite aged in air for 200 h at 1200C. Fig. 7. Fiber coating(a) before and(b)after thermal treatment
and 400 mm thick (patterns ``b'' and ``c'') are compared. These XRD patterns show that the SiC surface coating processed at 1200C by CVD is better crystallized than the SiC matrix produced by polycarbosilane pyrolysis at 1100C. In fact, the intensities of the more important peaks of b-SiC (JCPDS-ICDD card 29-1192) progressively become weaker and their shape widens with the increase of the distance from the sample surface; in the meantime faint peaks likely belonging to a-SiC (ICDD card 42-1091) grow in the XRD pattern. The XRD of the sample surface shows few small, not identi®ed re¯exes, that disappear as soon as the surface is polished. Fig. 8B shows the XRD patterns of the samples treated in air for 200 h at 1200C. In these spectra a new peak placed just below 2y =22, which can be reasonably attributed to cristobalite (ICDD card 39-1425), was observed. The intensity of this peak decreases moving from the surface toward the inner of the specimen, becoming close to zero at a distance of about 200 mm from the surface. Afterward, its intensity grows in the specimen core, at 400 mm from the surface or more. This particular trend (observed for samples thermally cycled too) was con®rmed by grazing-angle XRD measurements. The grazing-angle XRD patterns of the sample aged for 200 h at 1200C are reported in Fig. 9. Pattern ``a'' refers to the sample surface and pattern ``b'' to the same specimen after removal of a layer 40 mm thick: the ablation resulted in a slight decrease in the intensity of the peak placed at 2y=22. Furthermore, if the XRD measure of the polished specimen (Fig. 9, pattern ``b'') is repeated adopting a Bragg-Brentano geometry, that is using more penetrating X-rays, the spectrum ``c'' is obtained. This last spectrum shows that the peak at 2y =22 is almost vanished. On the basis of XRD measurements, it can be inferred that oxidation occurred both on the sample surface and in the sample core, but not (or not so much) in the internal part of the SiC coating deposited on the composite bar surface. 2.2.2.2. Infrared spectroscopy. In order to con®rm the outcomes of XRD measurements regarding the formation of silica inside the composite samples aged in oxidative environment, infrared spectroscopy was used. The infrared spectrum of the untreated sample surface did not show bands between 1000 and 1250 cm-1 , that is in the characteristic range for silica bands. However, the spectra of samples treated at 1200C or cycled in air showed bands in this range only when the analysis was carried out on the sample surface, while these signals disappeared after sample polishing (removal of a material layer about 40 mm thick). These IR measurements strengthen the XRD results, putting in evidence that the SiC coating oxidation occurs close to the sample surface only. 2.2.2.3. Electron probe microanalysis. The SEM-WDS analyses allow us to compare the chemical composition of composite samples showing a non-brittle behavior (composite specimens untreated or treated under vacuum) with that of composite bars aged in oxidizing environment (kept at 1200C or cycled between 25 and 1150C). The microanalysis results collected in Table 2 show that the atomic percent of C, Si and O, measured in the dierent part of the composite (®bers, matrix, interphases and external coating), may appreciably change after thermal treatment. Some of these composite parts display a non-homogeneous composition; for this reason, the composition of both core and external part (labeled with the symbol ``*'') of Nicalon ®bers and SiC bar coating are separately reported in Table 2. The changes in the material composition can be better appreciated by considering the C/Si and O/Si atomic ratios, reported as histograms for untreated and thermal aged samples in Figs. 10 and 11. The SiC ®ber coating and the sample coating of the untreated composite show a C/Si atomic ratio close to Fig. 6. Fiber damage in the composite aged in air for 200 h at 1200 Fig. 7. Fiber coating (a) before and (b) after thermal treatment. C. 1510 C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514
C Badini et al. Journal of the European Ceraic Society 20(2000)1505-1514 I and less than 1% of oxygen, while the composition more well distinguishable, even though an intermediate of Nicalon fibers is consistent with the presence of B- zone between fibers and matrix can be seen by SEM Sic. Sio.c. and free carbon. 8 (Fig. 7) The analysis of the sample aged under vacuum at In Table 2 as well as in Fig. ll, we have labeled the 1200C puts in evidence only small composition changes zone closer to the Nicalon fiber surface as"inner reac- with respect to the as-fabricated composite. The tion layerand the adjacent more external one as vacuum treated composite shows a more uniform oxy-"outer reaction layer". Actually, the intermediate zone gen content in the Nicalon fibers and a maximum oxy- at the fiber/matrix interface is strongly oxidized, its gen percentage in the matrix of 4%. Furthermore, a oxygen atomic percent ranging between 45 and 6% carbon depletion was observed in the more external part Conversely, the carbon content decreases here to low of the SiC sample coatin values: down to 26% in the inner reaction zone grown The aging in air at 1200C and the thermal cycling on the primary carbon interphase. The oxygen atomic cause much more significant variations in the sample composition. In these last samples the interphases(ori- samples(between 7 and 9%). The surface of the Sic ginally of C and Sic) close to the fiber surface are no coating sealing the composite bars is oxie wing a=a-Sic;β=阝sic;S=(SiO2);*= unidentified (a) (b) XRD patterns of (A) the untreated material and(B)of composite specimen after 200 h in air at 1200C. Patterns: a= sample surface; ter removal of a surface layer 200 um thick; c=after removal of a surface layer 400 um thick
1:1 and less than 1% of oxygen, while the composition of Nicalon ®bers is consistent with the presence of bSiC, SiOxCy and free carbon.8 The analysis of the sample aged under vacuum at 1200C puts in evidence only small composition changes with respect to the as-fabricated composite. The vacuum treated composite shows a more uniform oxygen content in the Nicalon ®bers and a maximum oxygen percentage in the matrix of 4%. Furthermore, a carbon depletion was observed in the more external part of the SiC sample coating. The aging in air at 1200C and the thermal cycling cause much more signi®cant variations in the sample composition. In these last samples the interphases (originally of C and SiC) close to the ®ber surface are no more well distinguishable, even though an intermediate zone between ®bers and matrix can be seen by SEM (Fig. 7). In Table 2 as well as in Fig. 11, we have labeled the zone closer to the Nicalon ®ber surface as ``inner reaction layer'' and the adjacent more external zone as ``outer reaction layer''. Actually, the intermediate zone at the ®ber/matrix interface is strongly oxidized, its oxygen atomic percent ranging between 45 and 6%. Conversely, the carbon content decreases here to low values: down to 26% in the inner reaction zone grown on the primary carbon interphase. The oxygen atomic percent also increases in the matrix of the air treated samples (between 7 and 9%). The surface of the SiC coating sealing the composite bars is oxidized, showing Fig. 8. XRD patterns of (A) the untreated material and (B) of composite specimen after 200 h in air at 1200C. Patterns: a=sample surface; b=after removal of a surface layer 200 mm thick; c=after removal of a surface layer 400 mm thick. C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514 1511
1512 C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 = aSiC Fig 9. Grazing-angle XRD patterns("a"and"b")and Bragg-Brentano XRD pattern (c")of a sample aged in air for 200 h at 1200C: a=graz. ing-angle XRD of the sample surface; b=grazing-angle XRD after removal of a surface layer about 40 um thick; c= Bragg-Brentano XRD after removal of a surface layer about 40 um thick an oxygen percentage greater than that of the coating oxidized than the composite matrix or the fiber inter core. However, it is surprising that only near the coating faces surface the oxygen content is greater than 1% at and that it quickly decreases with depth inside this coating. 2. 4. Microstructure and gas permeability of Sic seal In conclusion, it is clear that the external coating is less coating. The sections of the as-processed composite (observed by sEm)show that the external Sic coating is free from cracks, while only occasionally pores can b Table 2 detected inside this protective layer. However, the Electron probe microanalysis thickness of this layer is well far to be constant, because Fiber interphase Interphase it becomes very thin in some sample parts. Conversely, of sic cracks placed across the Sic layer were observed after Untreated composite(ar%) thermal treatment at 1200C and, more frequently, after cycling between 25 and 1150C Si45-40*36 These results show that the Sic coating looses its 0<1-5.5*<1 integrity during rapid temperature changes, becoming Composite thermal treated for 200 h at 1200 C under vacuum (at%) unsuitable for the prevention of oxygen diffusion 44-41 towards the composite bulk. This behavior is likely Si47-46*35 55-57* foreseeable in the case of samples submitted to thermal cycling, but not for specimens treated in isothermal Fiber Outer conditions which experience only one heating step up to reaction zone reaction zon 1200C. Furthermore, for these last composite samples. Composite thermal treated for 200 h at 1200C in air(ar%) cracks could arise from cooling at the test end or from 54-44* sami ple cutting and polishing (due to the increased 45.5-55* material brittleness) 1-2* As the results reported above seemed quite puzzling, Composite cycled between 25 and 1150C(1000 cycles)(ar% the gas permeability of the as-processed Sic coating, 53-47* apparently free from both cracks running through this SiC layer and open porosity, was further investigated by BET measurements. This test gives an indication about composition of the more external part of the zone. the material porosity through the measurement of the
an oxygen percentage greater than that of the coating core. However, it is surprising that only near the coating surface the oxygen content is greater than 1% at and that it quickly decreases with depth inside this coating. In conclusion, it is clear that the external coating is less oxidized than the composite matrix or the ®ber interfaces. 2.2.2.4. Microstructure and gas permeability of SiC seal coating. The sections of the as-processed composite (observed by SEM) show that the external SiC coating is free from cracks, while only occasionally pores can be detected inside this protective layer. However, the thickness of this layer is well far to be constant, because it becomes very thin in some sample parts. Conversely, cracks placed across the SiC layer were observed after thermal treatment at 1200C and, more frequently, after cycling between 25 and 1150C. These results show that the SiC coating looses its integrity during rapid temperature changes, becoming unsuitable for the prevention of oxygen diusion towards the composite bulk. This behavior is likely foreseeable in the case of samples submitted to thermal cycling, but not for specimens treated in isothermal conditions which experience only one heating step up to 1200C. Furthermore, for these last composite samples, cracks could arise from cooling at the test end or from sample cutting and polishing (due to the increased material brittleness). As the results reported above seemed quite puzzling, the gas permeability of the as-processed SiC coating, apparently free from both cracks running through this SiC layer and open porosity, was further investigated by BET measurements. This test gives an indication about the material porosity through the measurement of the Fig. 9. Grazing-angle XRD patterns (``a'' and ``b'') and Bragg±Brentano XRD pattern (``c'') of a sample aged in air for 200 h at 1200C: a=grazing-angle XRD of the sample surface; b=grazing-angle XRD after removal of a surface layer about 40 mm thick; c=Bragg±Brentano XRD after removal of a surface layer about 40 mm thick. Table 2 Electron probe microanalysis Fiber Interphase of C Interphase of SiC Matrix SiC coating Untreated composite (at%) C 54±55*a 63 48 50 51 Si 45±40* 36 51 48 48 O <1±5.5* <1 <1 2 <1 Composite thermal treated for 200 h at 1200C under vacuum (at%) C 51±54* 64 51.5 49 44-41* Si 47±46* 35 48 47 55±57* O 1 <1 <1 4 <1 Fiber Inner reaction zone Outer reaction zone Matrix SiC coating Composite thermal treated for 200 h at 1200C in air (at%) C 54 26 47 47 54±44* Si 46 29 47 44 45.5±55* O <1 45 6 9 <1±2* Composite cycled between 25 and 1150C (1000 cycles) (at%) C 51 33 41 40 53±47* Si 48 48 46 53 47±53* O <1 18.5 13 7 <1* a *=composition of the more external part of the zone. 1512 C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514
C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 total amount of gas(nitrogen)that can be adsorbed on area and the presence of close surface porosity only the pore surface. Probably the gas penetrates inside the composite bars BET measurements show that the as-fabricated com- where it is adsorbed on the walls of the internal pores posite bars have a high specific surface area(up to about On the basis of BEt analyses the presence of few coat 10 m/g), which is not explainable with their geometrical ing defects, not easy to find by SEM, has to be inferred Untreated 25 Sample treated Sample under vacuum F FS C S M L LE F Fs C S M Untreated Sample treated 0,3 0,1 F FS C S M L LE F FS C S M L LE F=fiber; Fs= fiber surface; C= carbon layer; S= Sic layer; M=matrix; L= sic surface coating; LE= SiC surface coating, external part Fig. 10. C/Si and O/Si ratios measured by WDS inside different parts of composite specimens untreated or aged at 1200C under vacur Sample aged 200 h1200° in air upto1150°c 恒‘ 匿變 F FS RI RO M L LE Sample cycled h1200° in air upto1150°c 2 F FS RI RO M L LE F FS RI RO M L LE F= fiber: fs= fiber surface: Ri= inner reaction zone: ro outer reaction zone ng Fig. l1. C/Si and O/Si ratios measured by WDS inside different parts of composite specimens aged at 1200C in air(200 h) or thermally cycled
total amount of gas (nitrogen) that can be adsorbed on the pore surface. BET measurements show that the as-fabricated composite bars have a high speci®c surface area (up to about 10 m2 /g), which is not explainable with their geometrical area and the presence of close surface porosity only. Probably the gas penetrates inside the composite bars where it is adsorbed on the walls of the internal pores. On the basis of BET analyses the presence of few coating defects, not easy to ®nd by SEM, has to be inferred. Fig. 10. C/Si and O/Si ratios measured by WDS inside dierent parts of composite specimens untreated or aged at 1200C under vacuum. Fig. 11. C/Si and O/Si ratios measured by WDS inside dierent parts of composite specimens aged at 1200C in air (200 h) or thermally cycled. C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514 1513
1514 C. Badini et al. Journal of the European Ceramic Society 20(2000)1505-1514 These defects make the composite core prone to Anyway, the results concerning the oxidation undergo oxidation, even though oxygen would diffuse mechanism put in evidence the difficulty to obtain a very slowly towards a compact SiC protective layer surface coating suitable for avoiding oxygen diffusion in the working conditions that this class of composite can experience in service (thermal shocks coupled with 3. Conclusions mechanical stresses). Further difficulties in carrying out effective oxidation barriers should arise from the A 2D-Nicalon/ C/SiC composite, cheaper than other joint points between SiC/Sic components and other composites of this class because of the use of common materials. Probably the improvement of SiC/Sic oxi ceramic grade fibers and due to the adoption of infil dation resistance would be more easily pursued tration process instead of CVI for the matrix fabrica- through a suitable modification of the fiber/matrix tion, was produced and investigated. This material interfaces shows high flexural strength and appreciable toughness he composite has a high thermal stability at 1200Cin non-oxidizing environment, but suffers from degrada- Acknowledgements ion when aged in air at this temperature The study of the degradation mechanism allows to The authors wish to thank Ing. A. Ortona of F. N draw the following conclusions Company (Bosco Marengo, Alessandria, Italy), who delivered the composite, for the useful discussion the oxidation of the carbon interphase (put between fibers and matrix) and the formation of silica are responsible for the embrittlement and the References strength decrease of the composite the oxygen contained in the Nicalon fibers scarcely I. Mah, T, Necht N. L. McCullum, D. E, Hoenigman. J. R. reacts with the carbon interphase (as shown by the Kim, H. M., Katz, A. P. and Lipsitt, H. A, Thermal stability of SiC fibres(Nicalon).J Mater. Sci. 1984, 19. 1191-120 aging tests under vacuum), probably also due to 2. Johnson. S. M.. Brittain. R. D. Lamoreaux. R. H. and Row the presence of a second interphase which cliffe. D. J. Degradation mechanism of silicon carbide fibers. J prevents the emission of gaseous reaction products A. Ceram Soc. 1988. 71. C132-C-135. the oxygen needed for the detrimental process ppuzzi, L Camus, G and Naslain, R, Oxidation mechanism comes from the environment, flowing towards the and kinetics of a ID-SiC/C/SiC composite material: I, an experi- mental approach. J. Am. Ceram. Soc., 1994, 77, 459-466 Sic barrier deposited by CVd on the sample sur 4. Labrugere, G, Guette, A. and Naslain, R, Effect of ageing eatments at high temperatures on the microstructure and this Sic coating in the as-processed specimens ems to be free of cracks. while these last form mechanical behaviour of 2D-Nicalon( C o 1. 17, 623-640 under vacuum or argon. J. Eur. Ceram. Soc. 199 fter thermal aging or after thermal cycling 5. Labrugere G. Guillaumat. L. Guette A. and Naslain R. Effect f ageing treatments at high to ures on the microstructur BET measurements suggest that the as-deposited and mechanical behaviour of 2D-Nicalon/C/SiC composites: 2 Sic layer shows a non-negligible gas permeability, ageing under CO and influence of a SiC seal-coating. J. Eur. probably caused by very few defects only occa- Ceran.Soc.1997,17,641-657 sionally present and, for this reason, difficult to be 6. Cavalier. J. C. Lacom be. A. and Rouges. J. M. Ceramic matrix observed by SEM; composites: new materials with very high performances. In Pro- eedings of the International Congress Developments in the Science furthermore, as after thermal treatments the oxy and Technology of Composite Materials, ed. A. R. Bunsell gen percent in the more internal part of the Sic P. Lamicq and A. Messiah. Elsevier Applied Science. London, coating is near to zero and, generally, very lower 1989,pp.99-10 than that observed in the composite bulk, the 7. Casadio. S. Donato. A. Nannetti. A. Ortona. A. and Rescio occurrence of a mechanism for oxygen migration M.R., Liquid infiltration and pyrolysis of SiC matrix composite Is Ceramic Transactions. 1995. 58. 193 (different from solid state diffusion inside the SiC Montioux, M. and Oberlin, A, Understanding lattice) has to be inferred fibre. J. Eur. Ceram Soc. 1993.. 95-103
These defects make the composite core prone to undergo oxidation, even though oxygen would diuse very slowly towards a compact SiC protective layer. 3. Conclusions A 2D-Nicalon/C/SiC composite, cheaper than other composites of this class because of the use of common ceramic grade ®bers and due to the adoption of in®ltration process instead of CVI for the matrix fabrication, was produced and investigated. This material shows high ¯exural strength and appreciable toughness. The composite has a high thermal stability at 1200C in non-oxidizing environment, but suers from degradation when aged in air at this temperature. The study of the degradation mechanism allows to draw the following conclusions: . the oxidation of the carbon interphase (put between ®bers and matrix) and the formation of silica are responsible for the embrittlement and the strength decrease of the composite; . the oxygen contained in the Nicalon ®bers scarcely reacts with the carbon interphase (as shown by the aging tests under vacuum), probably also due to the presence of a second SiC interphase which prevents the emission of gaseous reaction products; . the oxygen needed for the detrimental process comes from the environment, ¯owing towards the SiC barrier deposited by CVD on the sample surface; . this SiC coating in the as-processed specimens seems to be free of cracks, while these last form after thermal aging or after thermal cycling; . BET measurements suggest that the as-deposited SiC layer shows a non-negligible gas permeability, probably caused by very few defects only occasionally present and, for this reason, dicult to be observed by SEM; . furthermore, as after thermal treatments the oxygen percent in the more internal part of the SiC coating is near to zero and, generally, very lower than that observed in the composite bulk, the occurrence of a mechanism for oxygen migration (dierent from solid state diusion inside the SiC lattice) has to be inferred. Anyway, the results concerning the oxidation mechanism put in evidence the diculty to obtain a surface coating suitable for avoiding oxygen diusion in the working conditions that this class of composite can experience in service (thermal shocks coupled with mechanical stresses). Further diculties in carrying out eective oxidation barriers should arise from the joint points between SiC/SiC components and other materials. Probably the improvement of SiC/SiC oxidation resistance would be more easily pursued through a suitable modi®cation of the ®ber/matrix interfaces. Acknowledgements The authors wish to thank Ing. A. Ortona of F. N. Company (Bosco Marengo, Alessandria, Italy), who delivered the composite, for the useful discussion. References 1. Mah, T., Necht, N. L., McCullum, D. E., Hoenigman, J. R., Kim, H. M., Katz, A. P. and Lipsitt, H. A., Thermal stability of SiC ®bres (Nicalon). J. Mater. Sci., 1984, 19, 1191±1201. 2. Johnson, S. M., Brittain, R. D., Lamoreaux, R. H. and Rowclie, D. J., Degradation mechanism of silicon carbide ®bers. J. Am. Ceram. Soc., 1988, 71, C132±C±135. 3. Filippuzzi, L., Camus, G. and Naslain, R., Oxidation mechanism and kinetics of a 1D-SiC/C/SiC composite material: I, an experimental approach. J. Am. Ceram. Soc., 1994, 77, 459±466. 4. LabrugeÁre, G., Guette, A. and Naslain, R., Eect of ageing treatments at high temperatures on the microstructure and mechanical behaviour of 2D-Nicalon/C/SiC composites: 1 ageing under vacuum or argon. J. Eur. Ceram. Soc., 1997, 17, 623±640. 5. LabrugeÁre, G., Guillaumat, L., Guette, A. and Naslain, R., Eect of ageing treatments at high temperatures on the microstructure and mechanical behaviour of 2D-Nicalon/C/SiC composites: 2 ageing under CO and in¯uence of a SiC seal-coating. J. Eur. Ceram. Soc., 1997, 17, 641±657. 6. Cavalier, J. C., Lacombe, A. and Rouges, J. M., Ceramic matrix composites: new materials with very high performances. In Proceedings of the International Congress Developments in the Science and Technology of Composite Materials, ed. A. R. Bunsell, P. Lamicq and A. Messiah. Elsevier Applied Science, London, 1989, pp. 99±100. 7. Casadio, S., Donato, A., Nannetti, A., Ortona, A. and Rescio, M. R., Liquid in®ltration and pyrolysis of SiC matrix composite materials. Ceramic Transactions, 1995, 58, 193. 8. Le Coustumer, P., Montioux, M. and Oberlin, A., Understanding Nicalon ®bre. J. Eur. Ceram. Soc., 1993, 11, 95±103. 1514 C. Badini et al. / Journal of the European Ceramic Society 20 (2000) 1505±1514