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