Micro- and nanostructure of the carbon matrix of infiltrated carbon fiber felts

PERGAMON

Carbon 39 (2001) 215–229

Micro- and nanostructure of the carbon matrix of infiltrated carbon fiber felts b ¨ B. Reznik a , *, D. Gerthsen a , K.J. Huttinger a

¨ Elektronenmikroskopie, Universitat ¨ Karlsruhe, Kaiserstr. 12, 76128 Karlsruhe, Germany Laboratorium f ur b ¨ Chemische Technik, Universitat ¨ Karlsruhe, Kaiserstr. 12, 76128 Karlsruhe, Germany Institut f ur Received 14 January 2000; accepted 5 May 2000

Abstract The structural properties of carbon / carbon-composites fabricated by chemical vapor infiltration (CVI) were studied by polarized light microscopy (PLM), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) on a micrometer and nanometer scale. The types of carbon bonds were estimated by electron-energy-loss spectroscopy (EELS). Using a methane / hydrogen gas mixture at a temperature of 11008C two different methane partial pressures were applied. The carbon fibers are surrounded by ring-shaped layers with different optical reflectance. The SEM analyses of fracture surfaces revealed differences in the micromechanical behavior depending on the matrix morphology. Particular emphasis was put on the distinction of individual forms of pyrolytic carbons with similar optical behavior, which reveals significant structural differences in detailed SEM and TEM analyses.  2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Carbon / Carbon composites, Pyrolytic carbon; B. Infiltration; C. Electron microscopy, Electron-energy-loss spectroscopy

1. Introduction Carbon / carbon (C / C)-composites consisting of carbon fibers in a pyrolytic carbon matrix have attracted particular interest due to their high strength and light weight. However, the industrial application of infiltrated carbon fiber felts is still restricted due to their relatively high porosity and structure inhomogenity. Therefore it is very difficult to predict the mechanical properties of the felts. Since the chemical vapor infiltration (CVI) is frequently applied for the C / C-composite production [1,2], a thorough characterization of the micro- and nanostructure is required to establish a correlation between the structural and mechanical properties of infiltrated felts and the fabrication parameters. The understanding of the relationship between the CVI conditions and the structural properties will finally allow the design of composites with material properties appropriate for the desired application. The influence of the deposition parameters on the structure of the composites can be illustrated on the basis of results received by polarized light microscopy (PLM) *Corresponding author. Tel.: 149-0721-608-3720; fax: 1490721-608-3721. E-mail address: [email protected] (B. Reznik).

on a micrometer scale. The behavior of the optical textures is qualitatively correlated with the orientation and crystallinity of graphite lamellae. According to Bokros [3] and Pierson and Liebermann [4], laminar deposits exhibit a preferential orientation of the pyrocarbon layers parallel to the growth surface. A distinction is made between smooth laminar (SL) and rough laminar (RL) forms of carbon, which induce well-defined or numerous irregular extinction crosses in the PLM. In contrast, structures which possess little, if any, optical activity are called isotropic (ISO). Granoff et al. [5] investigated carbon fiber felts densified by means of thermal-gradient chemical vapor infiltration (CVI) using methane as the source gas. The materials with a SL matrix showed higher flexural strength than those with a RL matrix. The main source of a low strength is attributed to the presence of cracks in the heat-treated and highly-oriented pyrocarbons. Oh and Lee [6,7] analyzed the effects of the matrix structure on the mechanical properties of carbon composites prepared by thermal-gradient chemical vapor deposition. Pyrolytic carbon matrices were deposited within carbon mats at temperatures between 11008C and 14008C and from 10 to 70% propane concentration. The lower fracture stress and lower elastic modulus of isotropic and columnar (RL-like) carbons are caused by the low matrix density and the presence of

0008-6223 / 01 / $ – see front matter  2001 Elsevier Science Ltd. All rights reserved. PII: S0008-6223( 00 )00116-0

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matrix cracks, respectively. In addition, it was shown that the decomposition of propane at a temperature of about 12008C induces the formation of a duplex matrix consisting of a transition structure with SL and ISO forms of pyrocarbon. These composites are characterized by a relatively high Young’s modulus and fracture stress in comparison to matrices containing only one form of pyrocarbon. Similar observations were made by Benzinger ¨ and Huttinger [8,9] with respect to the fabrication process and properties of infiltrated carbon felts that were obtained by low temperature (T¯11008C) isothermal CVI using a methane / hydrogen gas mixture as a reaction gas. It was shown that the felts with a triplex matrix consisting of ISO, SL, and RL pyrocarbons produce a higher flexural strength and a higher Young’s modulus compared to composites containing mostly the RL pyrocarbon in the matrix. In the present study, a detailed structural characterization of infiltrated carbon fiber felts was performed to study the effects of the variation of the total reactor pressure on the microstructure and on the micromechanical properties. Particular efforts have been undertaken to distinguish the structure of individual forms of pyrocarbon with similar and very low optical activities, which could only be accomplished by the combination of polarized light microscopy, transmission electron microscopy (TEM) and scanning electron microscopy (SEM). In addition, the properties of interfaces between different pyrocarbon layers and between the matrix and the fiber are of prominent interest, because they are considered to influence the mechanical properties of the composites. Information about the types of bonds (p and s bonds) was obtained by electronenergy-loss spectroscopy (EELS), which can be carried out with high spatial resolution in a transmission electron microscope.

2. Experimental techniques

2.1. Chemical vapor infiltration The infiltration of the fiber felts was performed at a temperature of 11008C in a hot wall reactor. The PANfibers had a typical mean diameter of 12 mm and were randomly oriented in the felt before the infiltration. The gas mixture consisted of CH 4 and H 2 with a partial pressure ratio of 7:1 at a total pressure of 20 kPa (specimen felt I) and 30 kPa (specimen felt II). The total size of the composite was 20320335 mm 3 . Further details of the infiltration procedure and the sample properties are de¨ scribed by Benzinger and Huttinger [8,9].

2.2. Metallography In order to obtain comparable PLM, TEM and SEM samples, the following preparation steps were performed. At first a 3 mm disk was cut from the central part of the

Fig. 1. Schematic drawing of the sample locations for the PLM, TEM and SEM investigations. The dashed line shows the fracture direction.

composite with an ultrasonic cutter (Fig. 1) which was used for the PLM observations and for the TEM foil preparation. The remaining piece of the composite was broken approximately along the dashed line to obtain fracture surfaces for the SEM studies. Metallographic examinations were carried out by PLM using a Leitz DM RX microscope. The samples were mounted in epoxy resin and polished. A SiO 2 polishing paste with a 0.5 mm grain size was used in the final stage of the polishing procedure. In order to characterize the optical activity, the extinction angle ae was measured according to the procedure described by Jacques et al. [10].

2.3. Fractography Freshly fractured surfaces were examined by SEM in a LEO 1530 microscope with a Schottky field-emission gun without evaporating conductive layers prior to the investigation. With respect to the varying conditions of loading and deformation in the infiltrated felt and the random orientation of the fibers, the fracture surfaces of fibers and the matrix layers are oriented in different directions. Those surfaces were preferentially chosen for the fracture analysis, where the failure took place perpendicular to the fiber axis.

2.4. Transmission electron microscopy For the TEM specimen preparation, slices with a thickness of 200 mm were cut from the composite and further reduced in thickness by mechanical grinding and polishing to 80 mm. As already pointed out by Bruneton et al. [11], pyrocarbon layers may display different thinning rates during the Ar 1 -ion milling. The ions peel off preferentially the poorly organized structures in contrast to the slow thinning rate of the structures with a high degree of parallel ordering of the (002)-basal planes. To reduce the influence of the preferential ion milling, the samples

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were carefully dimpled down to a few microns or even to perforation. Therefore, the duration of the Ar 1 -ion bombardment and consequently the inhomogeneous thinning of the different carbons could be minimized. Typically, two argon guns operating at 1 kV, a current of 1 mA under an angle of 48 were used for the final cleaning of the sample from both sides. The TEM was carried out in a Philips CM 200 FEG / ST electron microscope operated at 200 kV. The instrument is characterized by a Scherzer resolution of 0.24 nm and an information limit of 0.15 nm.

2.5. Electron-energy-loss spectroscopy Information regarding the type of chemical bonding was obtained by parallel electron-energy-loss spectroscopy (PEELS) analyses of the fine structure of the carbon K-edge which distinctly differs for sp 2 - and sp 3 -coordinated carbons [12]. The analyses were carried out in a LEO EM 912 Omega transmission electron microscope at an electron energy of 120 keV, which is equipped with an Omega electron-energy spectrometer integrated into the projection lens system. The cathode was undersaturated to obtain spectra with the highest possible energy resolution. Under those conditions, the energy resolution was estimated to be better than 1.2 eV according to the measured full width at half maximum (FWHM) of the zero-loss peak. The PEELS was performed in the image mode of the microscope at a magnification of 30 0003. The spatial resolution is determined by the diameter of spectrometer entrance aperture and corresponds to an analyzed area of about 2003200 nm. The spectra were taken with a 10243 1024 pixel charge-coupled-device (CCD) camera with an acquisition time of typically 500 ms. The background subtraction was performed using a power-law function. Amorphization or structure changes were not observed under the applied illumination conditions. Polycrystalline graphite was used as a reference material. The powder-like material was placed on a holy-carbon supporting film. The diffraction pattern was characterized by Debye-Scherrer rings which minimizes the polarization factor of the EELS signal [12].

3. Experimental results

3.1. Metallography PLM images of felt I and felt II are shown in Fig. 2(a,b). Cross-sectioned and longitudinal sections of fibers are observed which are surrounded by concentric pyrocarbon layers. Three layers with different optical textures and thicknesses can be distinguished in felt I which are indicated in Fig. 2(a). The black rings, which are additionally observed around some fibers in Fig. 2(a) result from the debonding between the matrix and the fiber, which is most likely induced by the specimen preparation. It the

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case of felt II (Fig. 2(b)), the presence of up to five layers can be recognized around each fiber which are marked by the appropriate numbers. The extinction angles ae obtained by PLM are listed in Table 1 for the different layers in felt I and felt II. The optical textures are deduced according to the classification of Duppel et al. [14]. A homogeneous contrast (ae ,48) is typical for ISO carbon, while SL (128#ae ,188) and RL (188$ae ) textures are characterized by an inhomogeneous reflectivity. An intermediate optical activity (48#ae ,128) corresponds to the dark laminar (DL) carbons. The layer thicknesses were measured by PLM and also by SEM due to the resolution limit of the light microscopy. Only a rough estimation could be made for the extinction angle of the 1 mm thin layer 2 in felt I. To compare the two samples with respect to changes of the carbon microstructure with progressive infiltration, we will concentrate our structural analysis on the first three layers.

3.2. Fractography Fig. 3 shows low magnification SEM images of fracture surfaces of single fibers and the surrounding matrix in a cross-section view for felt I and felt II. The fracture surfaces of the fibers in Fig. 3(a) and Fig. 3(b) show the same features indicating similar loading conditions during the fracture which is a prerequisite for obtaining comparable results. According to Roulin-Moloney [15], the fracture occurs in both cases from the top to the bottom of the fiber which exhibits a smooth mirror-like zone and a hackle zone where bifurcation takes place producing the rough region, which is schematically sketched in the left insert of Fig. 3(a). The same concentric arrangement of layers as in Fig. 2 can be recognized. The fiber in felt I (Fig. 3(a)) is surrounded by a ring with a smooth fracture surface and a thickness of approximately 2 mm which corresponds to the ISO1 layer. The layer ISO2 is visible in the enlarged section of the image (right insert). The rough, layered fracture surface represents the third layer with the RL optical texture. An intensive fragmentation in concentric sublayers takes place in this layer. The fracturing of the matrix and carbon fibers occurs approximately in the same plane perpendicular to the fiber axis. In felt II (Fig. 3(b)), the fiber is partly covered by a thin matrix layer (DL according to Table 1) with a smooth fracture surface. The surface roughness increases in the second (SL) and third (RL) layers. The roughness is reduced in the fourth (SL) layer and rises again in the fifth (RL) layer. In contrast to felt I, the fracture surface level of the matrix changes from one layer to another. The adhesion between the matrix and the fiber can be assessed by the fiber pullout which occurs in the fracture experiments. Taking into account the influence of the different loading conditions that locally prevail in the felts, significant differences between the two samples could not be observed in overview SEM images.

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Fig. 2. Polarized light micrographs of felt I (a) and felt II (b). The numbers and the corresponding optical textures of pyrolytic layers are: (a) 1-ISO1, 2-ISO2, 3-RL; (b) 1-DL, 2 and 4-SL, 3 and 5-RL.

The laminar-type morphology of the pyrolytic carbon is resolved in more detail at a higher magnification in Fig. 4 for felt I and Fig. 5 for felt II. Globular-like fragments grouped in bands and oriented parallel to the fiber surface, whose orientation is indicated by the dashed line, can be

recognized in the case of layer ISO1 of felt I in Fig. 4(a). The surface of the layer ISO2 (Fig. 4(b)) with a similar optical reflectance exhibits a more elongated, coarser morphology. Fig. 4(c) shows the layered structure of the RL carbon. A smooth fracture surface is observed between

B. Reznik et al. / Carbon 39 (2001) 215 – 229 Table 1 Optical texture (OT), extinction angle a and thickness t of pyrocarbon layers a Layer, (OT) Felt I 1, (ISO1) 2, (ISO2) 3, (RL) Felt II 1, (DL) 2, (SL) 3, (RL) 4, (SL) 5, (RL)

a [degrees]

t [mm] (PLM)

t [mm] (SEM)

,4 ,4 .18

2.1 |1 |30

2.5 1.5 |30

4–12 12–18 .18 12–18 .18

2 7 9.7 4.5 8.8

2 6.5 8 4 9

a The thicknesses measured by PLM are accurate within 61 mm. The error for the SEM thicknesses is 60.2 mm.

steps with distances between 40 and 80 nm. Cracks are additionally found, which extend parallel to the fiber surface. In contrast to the layer ISO1 of felt I, the fracture surface of the DL layer of felt II depicted in Fig. 5(a) looks completely different. The surface morphology appears to be induced by flakes with sizes in the order of some 10 nm, which are oriented parallel to the fiber surface. The fracture surface of the SL layer is shown in Fig. 5(b), which is similar to that of the DL layer. However, a higher degree of texture parallel to the fiber surface is observed, which could be induced by a more pronounced alignment of the flakes. High-magnification SEM images perpendicular to the flakes (not shown here) indicate, that the flake thickness is in the order of a few nanometers. The RL layer of felt II (Fig. 5(c)) is characterized by a lower crack density and a rougher fracture surface compared to the RL layer of felt I. The step distances can be as low as 20 nm which contributes to the rough appearance of the fracture surface. The RL layers of both felts (Figs. 4(c) and 5(c)) are characterized by a pronounced laminar structure without any indication of globular- or grain-like fragments.

3.3. Transmission electron microscopy The structure of the different forms of carbon and the interfaces was also studied by TEM on a submicron- and nanometer scale. Low magnification TEM images are presented in Fig. 6(a) for felt I and Fig. 6(b) for felt II, which contain a cross-section view of a fiber surrounded by different layers. Three layers labeled 1 to 3 can be distinguished in felt I whose widths are similar to the layers displayed in Figs. 2(a) and 3(a). The layers ISO1 and ISO2 can be clearly distinguished in the image. Assuming comparable sample thicknesses around the TEM specimen edge, the intensity of the image can be qualitatively interpreted in terms of the material density which increases

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from ISO1 to ISO2 and the RL layer. Cracks parallel to the curved surface of the fiber can be recognized in particular inside the RL layer. Debonding does not preferentially occur at the interfaces between different layers. The cracks may be introduced by the TEM specimen preparation. However, the location and direction of the cracks can be taken as indicators for the strength of the different carbon layers and the interfaces. The arrangement of the DL, SL and RL layers in felt II is visible in Fig. 6(b). The measured layer thicknesses correspond to those in Figs. 2(b) and 3(b). The interface between the fiber and the first (DL) layer of felt II is displayed at an intermediate magnification in Fig. 7. Contrast modulations are observed with typical dimensions around 100 nm which can be interpreted in terms of a wavy, laminar-like arrangement of the layers. The observed morphology is consistent with the structural features of the fracture surface visible in Fig. 5(a), where structural units (flakes) of similar sizes are found. A selected area diffraction (SAD) pattern is inserted in Fig. 7. The intensity distribution of the ring with the smallest diameter corresponds to the orientation distribution of the (002)-basal planes. The increased intensity perpendicular to the fiber surface is an indicator of a weak texture, i.e. slight preferential orientation of the (002) planes parallel to the fiber surfaces. A crack is located in the carbon matrix at some distance from the interface, whose opening follows the laminar-like morphology of the matrix. This observation is indicative of a strong adhesion between fiber and matrix while debonding takes place more commonly inside the layer. Fig. 8 shows the microstructure of the matrix of felt I in the region of the layers ISO1, ISO2 and RL. The distinction between the two layers ISO1 and ISO2 can be easily accomplished by the SAD patterns inserted in Fig. 8. A homogeneous angular intensity would be expected for isotropic carbon. However, a weak texturing can already be recognized in the layer ISO1. The texture increases in the layer ISO2. A quantification of the texture degree can be obtained by the measurement of the orientation angle (OA) along the azimutal intensity distribution of ringshaped (002) reflections in SAD patterns as described by Tressaud et al. [16]. An intensity profile along a circle with a radius 1 /d 0002 is extracted from the diffraction patterns which negatives were digitized by a CCD camera with 14 bit gray scale resolution. The OA corresponds to the FWHM of the intensity maxima along the circle, which yields an OA of 458 for RL, 1058 for ISO1 and 868 for ISO2 correspondingly. The diffraction pattern of the RL layer is typical for a high degree of parallel alignment of the (002) planes. The RL layer contains numerous cracks, which are oriented parallel to the (002) planes. Cracking is also observed inside the ISO1 layer while the interfaces between the different carbon layers do not contain any cracks. The nanostructure of the interfacial region between the

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Fig. 3. Representative SEM micrographs of the fracture surface of a single fiber surrounded by the carbon matrix of (a) felt I infiltrated at Ptot 520 kPa and (b) felt II infiltrated at Ptot 530 kPa. Layers with different surface topography are marked by numbers which correspond to the layers identified by PLM. The left insert in (a) presents a schematic drawing of the regions on the fracture surface of carbon fiber which corresponds to the description of [21]. The right insert in (a) displays an enlarged section showing the ISO2 layer.

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Fig. 4. High-resolution SEM images of the fracture surface of felt I of the (a) ISO1, (b) ISO2 and (c) RL layers. The dashed line indicates the orientation of the fiber surface for all images.

Fig. 5. High-resolution SEM images of the fracture surface of felt II of the (a) DL, (b) SL and (c) RL layers. The dashed line marks the orientation of the fiber surface for all images.

ISO1 and ISO2 carbon layers of felt I is depicted in Fig. 9. The dashed line indicates the orientation of the fiber surface. The inserts are diffractograms of areas corresponding to the layers ISO1 and ISO2, which are obtained by the Fourier transformation of appropriately selected image regions. Apart from the superimposed information

regarding the imaging conditions, the diffractograms yield information about the orientation distribution of the basal planes similar to the SAD patterns. The increased intensity of the (002) Debye Scherrer ring perpendicular to the fiber surface again evidences the microtexture in the ISO carbon. The interface region, which extends over a few

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Fig. 6. TEM overview images of (a) felt I and (b) felt II showing a carbon fiber surrounded by pyrolytic carbon layers with different thicknesses.

nm, can be only roughly localized due to the very smooth transition between the different pyrocarbons. Fig. 10 shows a high magnification image of the interface region between highly and weakly textured carbons (the SL and RL layers of felt II) with the SAD patterns. The microcracks, which are located parallel to the graphite basal planes in the RL layer, are most likely induced by the specimen preparation. However, independent from the factors affecting the crack generation and propagation, it can be recognized, that the cracks do not extend through the interface into the weakly textured carbon. The wavy-like, curled structure of the SL carbon appears to suppress the crack propagation.

3.4. Electron-energy-loss spectroscopy EELS spectra of the carbon K-edge from the different layers of felt II displayed in Fig. 11 yield information about the bond types. The spectrum at the bottom of Fig.

11 was taken using a polycrystalline graphite reference sample which exhibits a fine structure at energy losses DE.293 eV similar to the previous EELS data of graphite [13]. The presence of the p* prepeak (labeled p) at 286 eV is characteristic for p bonds in sp 2 -coordinated carbon while the main resonance peak of the s bonds lies at the s* edge at 293 eV (labeled s). The comparison of the spectra of the different layers (labeled 1 to 3 in Fig. 11) with the graphite shows that all layers are crystallized in the sp 2 -rich form. The main trend in all spectra is the loss of fine structure at DE.296 eV and a fluctuation of the relative intensities of the p* and s* peaks. An additional feature in the spectra of the different layers is the variation of the height of the minimum between the s* and p* peaks at DE5(28761) eV which will be further discussed in Section 4.2. Assuming that the intensity of carbon K-edge represents properly the density of anti-bonding states located within the electron probe volume it is possible to determine the relative amounts of p- and s-bonded carbon from the relative intensities of the p and s features in the near-edge structure. To estimate the fraction of p and s bonds, it is assumed that the ratio of integrated areas under the p* peak from 283 to 287 eV and the s* peak from 290 to 297 eV is proportional to Np /Ns , the ratio of the density of p and s states [17]. This ratio is normalized with respect to the value determined for polycrystalline graphite supposed to be fully sp 2 in character and with a Np /Ns ratio of 1 / 3. The graphite sample as well as the DL and SL pyrocarbons have a polycrystalline character as revealed by diffraction patterns (Fig. 7–10). Therefore, we assume that the polarization factor can be ignored in the case of DL and SL pyrocarbons. In contrast, the measured intensity Ip /Is for RL carbon should depend on the orientation of the electron beam with respect to the texture. Based on this assumption we receive for DL and SL carbons Np /Ns ¯2 / 5 and for the RL carbon Np /Ns ¯1 / 3 correspondingly.

4. Discussion The structure of carbon fiber felts, which were infiltrated at 11008C and total reactor pressures Ptot of 20 kPa and 30 kPa, was characterized by PLM, SEM and TEM on different scales ranging from the micrometer to the subnanometer range. The pyrolytic carbon is deposited in ring-shaped layers around the fibers whose texture and crystallinity changes abruptly on a light-optical scale depending on the infiltration duration and the total reactor pressure. On the basis of the experimental results, the following topics will addressed in the discussion: (a) the techniques for the structural characterization of the different forms of pyrocarbon and in particular the distinction of pyrocarbons with similar optical activity, (b) the information contained in the EELS spectra, (c) the microstructure of the carbon matrix and the influence of the total

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Fig. 7. TEM image of an area close to the fiber / matrix interface in felt II taken at an intermediate magnification.

reactor pressure and (d) the influence of the microstructure on the micromechanical properties.

4.1. Structural characterization of different forms of pyrolytic carbon Only a coarse distinction of the different forms of pyrolytic carbons is possible by the PLM, which is in addition limited by the spatial resolution of the light microscopy. A more sensitive discrimination requires the application of TEM, where SAD patterns or diffractograms yield more detailed data on the texture which is determined by the degree of alignment of the (002) planes parallel to the fiber surface. A quantification of the texturing can be obtained by the measurement of the orientation angle. The benefits of this technique are in particular obvious in Fig. 8, where the slight preferential orientation (laminar structure) of the basal planes in the layers ISO1 and even more pronounced in ISO2 are visualized, while the PLM suggests an isotropic structure, i.e. a random orientation of the (002) planes. High-resolution SEM images (Figs. 4 and 5) of fracture surfaces contribute additional valuable information regarding the distinction of pyrocarbon forms with similar optical activity. Distinct differences are again revealed for the layers ISO1 and ISO2 of felt I in Fig. 4(a,b), where the ISO1 layer exhibits globular structural units in contrast to

the more elongated, coarser structural units of ISO2. Another example for the sensitivity of the high-resolution SEM with respect to small structural differences is given by the RL layers of felt I and felt II. A smooth fracture surface with a high crack density is found for felt I in contrast to felt II with a rough fracture surface induced by small step distances and a lower crack density. It can be finally stated that the combination of PLM, TEM and SEM significantly increases the sensitivity towards small structural differences which cannot be discriminated by PLM. It is worth noting, that the applied sample preparation has allowed to find similar areas for PLM, SEM and TEM investigations. Therefore, the layer thicknesses and types of pyrocarbons have always been comparable.

4.2. Electron-energy-loss spectroscopy Of special interest is the characterization of possible differences of the ratio between p and s bonds in the different pyrocarbon layers which prevail in all cases in the sp 2 -rich form according to Fig. 11. However, the measured Np /Ns ratio of 2 / 5 in SL and DL carbon requires a detailed discussion because Np /Ns is expected to rise only up to a maximum value of 1 / 3 in completely sp 2 -coordinated crystalline carbon. The comparison of the EELS and HRTEM data dem-

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Fig. 8. TEM image of the pyrocarbon matrix of felt I taken at an intermediate magnification. The inserts show the SAD diffraction patterns of the ISO1, ISO2 and RL layers with the appropriate orientation angles. The dashed line indicates the orientation of the fiber surface.

onstrates some tendencies of structure evolution in pyrocarbons: (1) the higher concentration of p states (Np /Ns ¯2 / 5) corresponds to the highly disordered DL and SL carbon (Figs. 7 and 10); (2) the lower concentration of p states (Np /Ns ¯1 / 3), i.e. a graphite-like concentration, corresponds to the RL carbon which con-

sists turbostratic-oriented graphite layers (Fig. 10). The significant loss of the fine structure in the s* region of the spectra of the DL and SL carbon compared to the graphite reference sample is indicative of bond length variations. Those can be attributed to a high density of defects and possibly vacancies which are expected to induce partially

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Fig. 9. High-resolution image of the interface region between the layers ISO1 and ISO2 of felt I. The inserted diffractograms were obtained by the Fourier transformation of regions which were selected from the ISO1 and ISO2 regions. The dashed line indicates the orientation of the fiber surface.

occupied (dangling) s bonds at the neighboring carbon atoms. In analogy to diamond-like films [12,18] it can be assumed that uncoordinated s bonds contribute to defect states in the region of the p density of states and therefore to the p* signal in the EELS spectra. Another point to consider is the possibility that hydrogen could be contained in the deposited pyrocarbon. A signal at DE5(28761) eV between the s* and the p* resonance is typically observed in hydrocarbons [18–20] and was assigned to electronic states of the C–H bonds. At

the same energy loss, the 1s→p* transition occurs in CO with a natural line width of 0.1 eV [19]. However, the presence of C–O bonds at 11008C does not appear to be possible. One may therefore tentatively attribute the varying height of the minimum between the s* and p* signals to varying hydrogen contents. The highest intensity at DE5287 eV corresponds to the first deposited (DL) pyrocarbon layer indicating that a hydrogen-rich phase could be formed during the initial stage of the infiltration process.

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Fig. 10. Lattice fringes image and corresponding SAD diffraction patterns of the interface region between the SL and RL layers in felt II.

In summary, the Np /Ns ratio of approximately 2 / 5 in the DL and SL carbon, which exceeds the maximum Np /Ns ratio in crystalline graphite, is mainly assigned to the high density of defects and vacancies in the highly disordered DL and SL structure with possibly some influence of varying hydrogen concentrations.

4.3. Microstructure of the carbon matrix The pyrolytic carbon is deposited in ring-shaped layers around the fibers in both samples. Significant changes of the carbon structure are observed in the carbon matrix with progressing infiltration, which typically leads to an increase of the texture during the initial and intermediate deposition stages. The variation of Ptot and the methane partial pressure PCH 4 respectively also strongly affects the microstructure of the carbon matrices. Only a weak texture is observed in felt I in the layers ISO1 and ISO2 with a total thickness of only 2 to 3 mm. The matrix of felt I is dominated by RL carbon which grows directly on top of

the ISO layers. In contrast, an oscillating behavior of the texture in felt II is observed (DL, SL, RL, SL, RL). However, all carbons are characterized by a laminar morphology even during the initial stages of deposition. The interfaces between the different layers are abrupt on a light-optical scale. HRTEM images show that the interfaces can extend over regions in the order of some 10 nm (Figs. 9 and 10). An explanation for the observed sequence of layers was ¨ already presented by Huttinger [21] and Benzinger and ¨ Huttinger [8,9] which is based on the increasing A p /Vp – ratio (A p : internal surface area of pores, Vp : volume of pores) during the infiltration. At a small A p /Vp –ratio (small surface area, large pore volume) in the beginning of the infiltration, the carbon deposition is proposed to occur from large aromatic or even polyaromatic hydrocarbon molecules which are formed in the gas phase. According to Pearson and Liebermann [4], weakly textured carbons are the consequence from the deposition of large hydrocarbon molecules. The formation of large molecules is prevented with progressing duration of the infiltration for an increas-

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infiltration. A pronounced layer structure and steps with different distances on the fracture surface of felt I and felt II are typical for RL carbons, which could reflect growth alterations during the later infiltration stages. The influence of PCH 4 cannot be understood in terms of the above model comparing the initial stages of the infiltration of felt I and II. Larger molecules or clusters are expected at a higher PCH 4 (felt II) due to an increased probability of gas phase reactions. Consequently, carbons with a less pronounced texture are expected, which is in contradiction to the experimental observations. The formation of the oscillating texture in the matrix of felt II is also difficult to explain in the context of the considerations above. Bruneton et al. [11] obtained carbon matrices with a similar microstructure after the rapid pyrolysis of cyclohexane, which was attributed to local temperature instabilities. However, due to the isothermal character of CVI experiments the oscillating character of the textures could be rather related to an altering diffusion of gas molecules in the pore-like channels of felts. More experiments are in progress to understand the structure formation in more detail.

4.4. Influence of the microstructure on the micromechanical properties

Fig. 11. PEELS spectra in the energy loss region of the carbon K-shell excitation edge of the different pyrocarbons in felt II.

ing A p /Vp –ratio. The deposition then occurs mainly from light, linear carbon molecules (C 1 - and C 2 -species) which leads to the formation of carbon with a RL structure. Experimental data about the size and type of molecules in the gas phase are at present not available. However, some information can be extracted from the high-resolution SEM images (Fig. 4 and Fig. 5) which reveal globular- or flake-like structural units in ISO, DL or SL carbons. The lateral size of these structural units can be up to some 10 nm indicating that large clusters could already be present in the gas phase during the early stages of

The micromechanical behavior is determined by the arrangement, thickness and texture of the layers in the carbon matrix and the adhesion between the carbon matrix and the fiber. An idea of the influence of the matrix microstructure on the mechanical properties can be obtained by the SEM analysis of fracture surfaces and by the investigation of the density, size and location of cracks in TEM specimens. A good adhesion between fiber and matrix was observed in both felts which is in particular visualized by Fig. 7 for felt II (comparable results exist for felt I), where the crack was generated in the matrix and not directly at the fiber / matrix interface. This also illustrates that SEM does not necessarily allows the exact localization of the debonding which can occur either directly at the fiber / matrix interface or in the matrix. The comparable adhesion between the fiber and matrix could be the main cause for the relatively small difference of the mechanical properties (flexural strength, Young’s modulus and strain-to-failure) measured by Benzinger and ¨ Huttinger [8,9]. Nevertheless, the fracture behavior of the two composites is not completely identical, which can be ascribed to the microstructure of the carbon matrix. The matrix of the felt I mostly consists of highly textured pyrocarbon (Fig. 2(a)) in contrast to the texture in felt II which oscillates between RL and SL carbons (Fig. 2(b)). It is common to both samples, that the highest crack densities are observed in the RL layers (e.g. Fig. 6), where a low crack propagation resistance exists parallel to the highly ordered graphite

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layers. The intensive fragmentation in the form of ringshaped sublayers (Fig. 3(a,b)) and the sharp edges of cracks inside of RL layers are typical for brittle failure. An important observation is contained in Fig. 10 where the crack propagation is impeded at the interface between highly and weakly textured carbon. Therefore, crack deflection is expected at the interfaces between highly and weakly carbon layers of felt II, which results in different height levels on the fracture surface of felt II (Fig. 3(b)) in contrast to felt I (Fig. 3(a)). An indication of multidirectional cracking (crack branching) along the structural units in the weakly textured carbons is obtained by the high-resolution SEM images (Figs. 4 and 5) and the TEM image Fig. 7, which show highly developed fracture surfaces and rough crack edges. According to Bokros [22], Chawla [23] and McEnaney et al. [24], this can lead to an enhanced toughness because the stress required to drive a number of cracks is higher than that one for driving a single crack. Depending on the sizes and shapes of the structural units, the resistance towards the crack propagation could be different in the weakly textured carbons of felt I and felt II. Two other factors could lead to a variation of the mechanical properties. The folds below the fiber in the SL layer in Fig. 3(b) indicate, that a portion of fracture energy can be dissipated in stress-induced deformation in SL carbons. Slight texture differences between the RL carbons in felt I and felt II may influence the crack generation and propagation in the RL layers which is visualized by Figs. 4(c) and 5(c) where different step and crack densities are observed.

5. Summary A thorough study of the microstructure of infiltrated carbon fiber felts was performed by PLM, SEM and TEM. The felts were infiltrated at 11008C using a hydrogen / methane mixture of 7:1 at total reactor pressures of 20 kPa and 30 kPa. It was shown that even small differences between the micro- and nanostructure of the different forms of pyrocarbons can be revealed by combining PLM, TEM and SEM. The measurement of the extinction angles by PLM only allows a coarse distinction between the different pyrocarbons. More detailed information about the texture can be obtained by the analysis of SAD patterns and the measurement of the (0002)-reflection opening angle. Additional valuable information about the structural units is obtained by high-resolution SEM on fracture surfaces. EELS spectra of felt II suggest that all carbon layers consist of sp 2 -coordinated carbon atoms. The high Np /Ns ratio of approximately 2 / 5 in the DL and SL carbon, which may attributed to the high density of defects and vacancies in disordered graphitic structure as well as to the influence of hydrogen.

The carbon is deposited in ring-shaped layers with an increasing degree of texture during the initial and intermediate infiltration stages, which can be understood in terms of the increasing A p /Vp -ratio resulting in smaller molecules in the gas phase with progressing infiltration duration. Some questions remain to be solved regarding the understanding of the oscillating texture of felt II during the final infiltration stage and the details of the structure formation depending on the methane partial pressure. High-resolution SEM analyses of the fracture surfaces reveal the size and shape of the structural units of the weakly textured carbon layers which could be a fingerprint of the sizes of the clusters present in the gas phase. Taking into account that there is a comparable good adhesion between the fiber and matrix in both felts, the differences between the mechanical properties could be mainly related to the matrix architecture (sequence of the layers with different widths and texture). Brittle fracture is deduced for the RL carbon layers containing a high density of cracks with a low crack propagation resistance. The number of interfaces between highly and weakly textured carbon layers, where the crack propagation is impeded, could be a dominant factor toward the toughness enhancement of pyrolytic carbon matrices.

Acknowledgements This research was performed in the Collaborative Research Center (Sonderforschungsbereich) 551 financed by the Deutsche Forschungsgemeinschaft. We thank Prof. A. Oberlin for the stimulating discussions. The texture measurements were carried out with the help of V. De Pauw and M. Guellali which is greatly appreciated.

References [1] Savage G. In: Carbon–carbon composites, London: Chapman & Hall, 1993, pp. 103–20. [2] Fitzer E, Manocha LM. In: Carbon reinforcements and carbon / carbon composites, Berlin: Springer-Verlag, 1998, pp. 65–90. [3] Bokros JC. Deposition, structure, and properties of pyrolytic carbon. In: Walker PL, editor, Chemistry and physics of carbon, vol. 5, New York: Dekker, 1969, pp. 1–118. [4] Pierson HO, Lieberman ML. Carbon 1975;13:159–66. [5] Granoff B, Pierson HO, Schuster DM. Carbon 1973;11:177– 87. [6] Oh S-M, Lee J-Y. Carbon 1988;26:769–76. [7] Oh S-M, Lee J-Y. Carbon 1988;26:763–8. ¨ [8] Benzinger W, Huttinger KJ. Carbon 1999;37:941–6. ¨ [9] Benzinger W, Huttinger KJ. Carbon 1999;37:1311–22. [10] Jacques S, Guette A, Bourrat X, Langlais F, Guimon C, Labrugere C. Carbon 1996;34:1135–43. [11] Bruneton E, Narcy B, Oberlin A. Carbon 1995;35:1599–611.

B. Reznik et al. / Carbon 39 (2001) 215 – 229 [12] Bruley J, Williams DB, Cuomo JJ, Pappas PP. J Microscopy 1995;180:22–32. [13] Rosenberg RA, Love PJ, Rehn V. Phys Rev B 1986;33:4034– 7. [14] Duppel P, Bourrat X, Pailler R. Carbon 1995;33:1193–204. [15] Roulin-Moloney AC, editor, Fractography and failure mechanisms of polymers and composites, London: Elsevier Applied Science, 1989, pp. 233–90. [16] Tresaud A, Chambon M, Gupta V, Flandrois S, Bahl OP. Carbon 1995;33:1339–45. [17] Lossy D, Pappas L, Ronnen AR, James PD. J Appl Phys 1995;77:4750–6. ¨ ¨ [18] Fink J, Muller-Heinzerling T. Pfluger J Solid State Comm 1983;47:687–91.

229

¨ J, Robinson CJ, Jark W. Phys Rev B [19] Comelli G, Stohr 1988;38:7511–9. ¨ J, Sette F. J Chem [20] Hitchcock AP, Beaulieu S, Steel T, Stohr Phys 1984;80:3927–35. ¨ [21] Huttinger KJ. Adv Mater 1998;4:151–8. [22] Bokros JC, Price RJ. Carbon 1966;3:503–19. [23] Chawla KK. In: Ceramic matrix composites, London: Chapman & Hall, 1993, pp. 291–340. [24] McEnaney B, Mays TJ. Relationships between microstructure and mechanical properties in carbon–carbon composites. In: Thomas CR, editor, Essentials of carbon–carbon composites, Cambridge: Royal Society of Chemistry, 1993, pp. 143–74.