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Effect of carbon fiber surface functional groups on the mechanical properties of carbon–carbon composites with HTT
Carbon 41 (2003) 1193–1203
Effect of carbon fiber surface functional groups on the mechanical properties of carbon–carbon composites with HTT S.R. Dhakate*, O.P. Bahl Carbon Technology Unit, Engineering Materials Division, National Physical Laboratory, Dr. K.S. Krishnan Marg, New Delhi 110012, India Received 18 March 2001; accepted 25 January 2003
Abstract The present investigation describes the quantitative measurement of surface functional groups present on commercially available different PAN based carbon fibers, their effect on the development of interface with resol-type phenol formaldehyde resin matrix and its effect on the physico–mechanical properties of carbon–carbon composites at various stages of heat treatment. An ESCA study of the carbon fibers has revealed that high strength (ST-3) carbon fibers possess almost 10% reactive functional groups as compared to 5.5 and 4.5% in case of intermediate modulus (IM-500) and high modulus (HM-45) carbon fibers, respectively. As a result, ST-3 carbon fibers are in a position to make strong interactions with phenolic resin matrix and HM-45 carbon fibers make weak interactions, while IM-500 carbon fibers make intermediate interactions. This observation is also confirmed from the pyrolysis data (volume shrinkage) of the composites. Bulk density and kerosene density more or less increase in all the composites with heat treatment up to 2600 8C. It is further observed that bulk density is minimum and kerosene density is maximum upon heat treatment at 2600 8C in case of ST-3 based composites compared to HM-45 and IM-500 composites. It has been found for the first time that the deflection temperature (temperature at which the properties of the material start to decrease or increase) of flexural strength as well as interlaminar shear strength is different for the three composites (A, B and C) and is determined by the severity of interactions established at the polymer stage. Above this temperature, flexural strength and interlaminar shear strength increase in all the composites up to 2600 8C. The maximum value of flexural strength at 2600 8C is obtained for HM-45 composites and that of ILSS for ST-3 composites. 2003 Published by Elsevier Science Ltd. Keywords: A. Carbon fibers; C. Electron spectroscopy; D. Functional groups, Mechanical properties
1. Introduction Carbon–carbon composites are used in a number of demanding applications such as space, defense, turbine blades, etc. [1,2]. Their performance is known to depend on the type of carbon fibers, matrix precursors, nature of bonding between fiber and matrix (fiber–matrix interface) and processing conditions [3,4]. The nature of the interface and its influence on physico–mechanical properties is extremely complex. Different types of interface may exist
*Corresponding author. Tel.: 191-11-574-6290; fax: 191-11572-6952. E-mail address: [email protected] (S.R. Dhakate).
in carbon–carbon composites between fiber and matrix, within fiber bundles, between the different microstructures which may exist within the matrix. It is well known that the fiber–matrix interaction depends upon the surface functional groups of the carbon fibers and the matrix precursor [3,4]. In particular, the mechanical properties of carbon–carbon composites are very sensitive to the bonding between fibers and the matrix and its stress transfer capability. In this respect, continuous research and development work is going on from the last 3–4 decades, and especially regarding the role of interface on controlling the overall performance of carbon–carbon composites [5–10]. A qualitative correlation between the amount of surface functional groups, the nature of the interface and composite properties has been reported by Fitzer et al. [7]. Also,
0008-6223 / 03 / $ – see front matter 2003 Published by Elsevier Science Ltd. doi:10.1016 / S0008-6223(03)00051-4
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Table 1 Properties of carbon fibers used (from manufacturer’s data sheet) Fiber type
Diameter (mm)
Density (g / cm 3 )
Tensile strength (MPa)
Tensile modulus (GPa)
Strain to failure (%)
ST-3 IM-500 HM-45
6.0 5.5 6.0
1.75 1.77 1.91
4200 4700 2200
250 300 440
1.68 1.57 0.5
there are plenty of data available on the analysis of surface functional groups of carbon fibers and their influence on the development of interface and mechanical properties of polymer matrix composites. In addition to this, some data are available in the literature describing the effect of heat treatment temperature (HTT) on the mechanical properties of polyarylacetylene (PAA) and furfuryl alcohol based carbon–carbon composites [11–13]. Some studies are also available on the influence of fiber surface functional groups and various types of surface treated carbon fibers on the development of interface with phenolic resin matrix [7,14,15]. In the present investigation, a systematic approach was adopted to understand the influence of carbon fiber surface functional groups on interface development by measuring the surface functional groups quantitatively and their influence on composite properties as a function of HTT.
2. Composite preparation and characterization Three types of commercially available PAN-based carbon fibers, manufactured by Toho Beslon Inc., and Torayca Industries Inc., Japan, were used as reinforcement. (a) ST-3 (high strength, likely HTT |1200–1500 8C). (b) IM-500 (intermediate modulus, likely HTT |1600– 1800 8C). (c) HM-45 (high modulus, likely HTT .2200 8C) Since these fibers are heat-treated to different temperatures, they exhibit different physical and mechanical properties (see Table 1) as well as different types of surface functional groups. The surface functional groups present were measured by ESCA, using an SSI 301 spectrometer employing monochromatic and focused Al Ka radiation (spot diameter 300 mm, 80 W; radiation energy 1486.6 eV) under a residual pressure of 5310 28 Torr (1 Torr5133.322 Pa). Unidirectional polymer composite samples (150 mm3 4.0 mm34.5 mm) were prepared using the wet winding and match mold die technique [16] with 4562% fiber volume. The resol type phenol formaldehyde resin was used as the matrix precursor for composite preparation.
The composites were coded as follows. (A) ST-3 carbon fiber composites. (B) IM-500 carbon fiber composites. (C) HM-45 carbon fiber composites. The polymer composites were heat-treated to 400– 2600 8C under an inert atmosphere. The heat-treated composites were characterized for volume shrinkage, density and mechanical properties during each stage of heat treatment. Flexural strength was measured by the three-point bending technique on a universal Instron testing machine (Model 4411, ASTM standard D-790-80) with a span length to depth ratio of 30:1. The interlaminar shear strength (ILSS) was measured using ASTM standard D-2344-74 with span length to depth ratio of 8:1. The kerosene density was measured by the kerosene pickup method using the Archimedes principle. The transverse coefficients of thermal expansion were measured using a thermo-mechanical analyzer (TMA) attached to a Mettler thermal system TA-3000, in the range 50–900 8C under an inert atmosphere. The optical micrographs of composites heat treated at 2600 8C were observed using polarized light microscopy.
3. Results and discussion
3.1. ESCA studies of carbon fibers The surface composition of carbon fibers obtained by
Table 2 Surface composition of carbon fibers Fiber type \element
Atomic % C (1s)
O (1s)
N (1s)
Si (2p)
ST-3 IM-500 HM-45
86.63 92.28 92.21
8.19 7.23 6.18
2.35 0.00 0.00
1.86 0.49 1.60
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Fig. 1. (a) ESCA spectra of (a) ST-3, (b) IM-500 and (c) HM-45 carbon fibers. (b) Deconvolution spectra of O 1s for (a) ST-3, (b) IM-500 and (c) HM-45 carbon fibers.
Table 3 Relative percentage of functional groups on the carbon fibers obtained by ESCA Functional group
Graphitic carbon (B. E., eV) Phenolic or hydroxyl (B. E., eV) Carbonyl (B. E., eV) Carboxylic (B. E., eV) Nitrogen containing (B. E., eV)
Relative atomic percentages of functional groups ST-3
IM-500
HM-45
69.9 (284.18) 18.76 (285.42) 7.01 (287.34) 4.31 (290.13) 2.35 (398.02, 400.43)
61.9 (284.06) 28.6 (285.08) 5.44 (287.90) 4.05 (290.48) –
90.88 (284.05, 284.83) – 4.3 (288.49) 4.8 (290.92) –
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the ESCA studies is presented in Table 2. As expected [17], the oxygen content on the surface decreases from ST-3 to HM-45 and the C / O ratio increases from ST-3 to HM-45, confirming the well known fact that the oxygencontaining functional groups, and consequently the reactivity of carbon fiber surface, decrease with increasing HTT. The nitrogen-containing functional groups are present only in ST-3 fibers, from the residual nitrogen present in the precursor fibers. Fig. 1a shows the ESCA spectra of the three carbon fibers. The three major peaks observed between binding energies 200 to 600 eV of carbon (C 1s), oxygen (O 1s) and nitrogen (N 1s) correspond to graphite-like carbon and various functional groups. The peak around 101 eV corresponds to silicon Si (2p). The C 1s peak at 284.24 eV is assigned to the carbon element only and between 284.05 and 284.83 eV to graphitic carbon, 285.08–285.42 eV to hydroxyl (.C–OH), 287.34–288.49 eV to carbonyl (.C=O) and 290.13–290.92 eV to carboxylic (–COOH) functional groups [18–20]. The chemical shift for all the functional groups of three different carbon fibers is different (Fig. 1b, Table 3), which may be due to the structural difference of the carbon fibers surfaces. The fibers used in the investigation are heat-treated at increasing temperature (ST-3 to HM-45) as described in Section 2. As a consequence, orientation of carbon layers preferentially improves parallel to fiber axis leading to different degrees of graphitization and therefore different electrical conductivities of the fibers. The ESCA spectrum for the ST-3 carbon fibers shows a very weak N 1s peak as compared to the C 1s and O 1s peaks, which arises from the residual nitrogen present in the fibers. Nitrogen-containing functional groups appear in ST-3 fibers at a binding energy of 398.02 eV, corresponding to the aromatic amines and piperidine structure, at 400.43 eV to aliphatic amines, nitrile and amides [21,22], while in the cases of IM-500 and HM-45 carbon fibers these groups are not detected. Deconvolution of these peaks gives relative percentages of functional groups present and these are compiled in Table 3. The high strength carbon fibers (ST-3) possess the highest number of surface functional groups and as a result the surface of these fibers may be more disordered possessing comparatively high active surface area. The ST-3 fibers possess carboxylic, phenolic and hydroxyl functional groups which are acidic in nature while carbonyl and some of nitrogen containing groups are basic in character [23–29]. The high modulus HM-45 carbon fiber, on the other hand, possesses a minimum number of surface functional groups and a better ordered graphite-like carbon fiber surface. Therefore such fibers should have the lowest active surface area. The intermediate modulus (IM-500) carbon fibers possess the greatest number of hydroxyl functional groups. The contribution of carboxylic groups is almost the same in all the three fibers while carbonyl groups decrease from ST-3 to HM-45 fibers.
3.2. Interactions of fibers with matrix The specific interactions are postulated to be Lewis acid–base type interactions or electron acceptor–donor interactions [30,31]. The polymer used here is a thermosetting phenolic resin which has acidic functional groups whereas fibers possess both acidic and basic functional groups. It is well known that carboxylic groups are responsible for strong interactions with a polymer matrix having basic functional groups [18,32]. Therefore, in the present case these groups are not likely to play an important role. It is found that ST-3 fibers possess maximum (9.46 relative percentage) functional groups which are basic in character and would make strong interactions, whereas HM-45 fibers possess minimum reactive functional groups (4.3 relative percentage) and would thus make weak interaction with the phenolic resin matrix. On the other hand, the IM-500 fibers possess a maximum amount of total functional groups (38.09 relative percentage) but the reactive functional groups are only 5.44 relative percentage (of carbonyl), which can make interactions with phenolic resin matrix intermediate between ST-3 and HM-45 carbon fibers [33].
3.3. Volume shrinkage vs. HTT of composites During heat treatment, carbon fiber reinforced polymer composites show changes in all the three directions of composites, i.e., length, width and thickness, which are due to the thermal degradation and shrinkage of polymer matrix [7,34]. In unidirectional composites, the changes taking place in the direction parallel to the fiber axis are controlled by the longitudinal thermal expansion of the carbon fibers themselves. The fiber thermal expansion (positive or negative) is very small [35]. As a consequence,
Fig. 2. Volume shrinkage observed with heat treatment temperature of composites.
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these composites exhibit very little change parallel to carbon fibers. However, appreciable changes in the width and thickness directions are observed. Fig. 2 shows the volume shrinkage of polymer composites with heat treatment temperature. The amount of shrinkage depends upon the type of polymer matrix used in the fabrication of composites and fiber–matrix interactions [9]. During heat treatment, chemical bonding between the fiber surface and the free functional groups of the resin undergoes rearrangement. Simultaneously, the molecular chains in the polymer matrix undergo pyrolysis and their rearrangement results in matrix shrinkage. Up to 400 8C the shrinkage pattern of all the composites does not show much difference, but with increasing HTT shrinkage increases linearly in all the composites. For composite A it is maximum during the whole range of temperature and minimum in case of composite C. Up to 800 8C, 15% shrinkage is noticed in the case of composite A, whereas in composites B and C it is 11–12%. The higher shrinkage in the case of composite A is attributed to strong fiber–matrix interactions. With increase in heat treatment temperature, shrinkage increases continuously in all the composites and it is |25% in composite A, 22% in composite B and 19% in composite C at 1800 8C. Above 1800 8C, the conversion of non-graphitic carbon into a graphite-like carbon structure (in the matrix), i.e., a reorientation of graphitic planes, starts taking place thus resulting in further composite shrinkage [11]. Further, with a continuous increase in heat treatment temperature to 2600 8C, the reorientation of graphitic planes takes place more extensively and the extent of shrinkage observed depends on the extent of reorientation. The incremental shrinkage between 2000 and 2600 8C is as high as 6–7% in composite A and only 4% in composite C. It is important to mention here that the initial fiber volume (456 2%) is kept the same in all three cases. This brings out very clearly the effect of reactive surface functional groups on the the fiber–matrix interactions. Consequently, during heat treatment the matrix
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shrinks onto the fibers when the interaction is maximum (as in composite A). Such composites should exhibit a maximum amount of shrinkage, as is observed. In composite C, fiber–matrix interactions are weak (see Section 3.2) and as a result there is minimum shrinkage.
3.4. Bulk density vs. HTT of composites The bulk or apparent density is the density of composites containing voids and porosity. Table 4 shows the changes in density observed with HTT of the composites. These changes in density may be due to two factors:
(i) Dimensional changes due to shrinkage during pyrolysis. (ii) Weight loss due to evolution of volatile products during pyrolysis.
The initial density of the polymer composites depends upon the density of carbon fibers used, its volume fraction (which is kept the same in all three composites) and compactness of the composites. The bulk density of polymer composites varied between 1.42 and 1.52 g / cm 3 . Upon heat treatment to 600 8C, a gradual density decrease is observed due to evolution of reaction products and formation of pores which results in volume expansion [36]. Above 600 8C, product evolution decreases to a large extent as pyrolysis of the resin matrix is almost complete and there is an increase in density up to 1400 8C which is due to structural changes. The small density decrease between 1400 and 1800 8C in composites A and B is due to the formation of closed microporosity [37,38] because of evolution of nitrogen-containing reaction products.
Table 4 Variation in bulk density of composites HTT (8C) 150 400 600 800 1000 1400 1800 2200 2600
Bulk density (g / cm 3 ) A
B
C
1.42 1.37 1.36 1.41 1.45 1.51 1.50 1.53 1.55
1.46 1.42 1.40 1.45 1.50 1.55 1.53 1.55 1.57
1.53 1.45 1.44 1.50 1.53 1.55 1.57 1.58 1.61
Fig. 3. Change in kerosene density of composites with stages of heat treatment.
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Table 5 Porosity in composites heat-treated at 2600 8C* Porosity (%)
Composite type A
B
C
Total Closed Open
31.4 17.5 13.9
30.5 20.2 10.3
28.8 20.8 8
* Determined from true (2.26 g / cm 3 ) and bulk densities.
Composites A and B were made from carbon fibers which are not heat treated to more than 1400–1800 8C and therefore, contain some amount of nitrogen containing functional groups on their surface. Composite C, on the other hand, does not show any decrease in the density since HM-45 carbon fibers (likely heat treatment temperature.2200 8C) are nitrogen free (Table 2). Further, on heat treatment to 2200 and 2600 8C, the bulk density gradually increases in all the composites, because of breakdown of microporosity due to structural rearrangement of carbon atoms. The final bulk density of composites varies between 1.55 and 1.61 g / cm 3 . The bulk density of all the composites increases more or less with HTT. In the case of glass like carbon it decreases with HTT due to the continuous volume expansion of pores [38]. In the composites the volume expansion is thought to be controlled by fiber–matrix interactions.
3.5. Kerosene density vs. HTT of composites In order to get an idea about the total open porosity in the composites, experiments were conducted to measure kerosene density since kerosene is known to wet all carbons is capable of entering pores of up to 50 nm in size. The kerosene density and bulk density give information about the open porosity present in the composites. Fig. 3 shows changes in the kerosene density of composites A, B and C with HTT. Initially at 1000 8C the density of all the composites is in the range 1.60–1.62 g / cm 3 . In the case of composites A and B, the density increases moderately between 1000 and 2000 8C, and then it registers a sharp increase from 1.64 to 1.80 g / cm 3 for composite A while it increases from 1.65 to 1.75 g / cm 3 for composite B. The sudden density increase in the case of composite A may be due to preferential graphitization (in the immediate vicinity of carbon fibers) of isotropic carbon derived from phenolic resin as well as the graphitization of carbon fibers [12] (see Fig. 7, composite A). However, in case of composites B and C kerosene density registers a continuous increase with the ultimate value of 1.75 g / cm 3 only at 2600 8C. This observation clearly brings out the role of fiber–matrix interaction in affecting
graphitization (and hence high density) of the matrix as well as the carbon fiber surface. Table 5 gives information about the total, open as well as closed porosity in the three composites heat treated at 2600 8C. Even though the same matrix has been used, open porosity is maximum (13.9%) in composite A and minimum (8%) in composite C. It is well known that when soft carbon is heat treated to temperatures of the order of 2200 8C and above, it becomes graphitized and, as a consequence of structural reorganization, a lot of porosity gets opened up [39]. However, in the present case, even though the maximum HTT (2600 8C) is the same in all the composites, the open porosity values are entirely different. In case of composite A, fiber–matrix interactions are strong this generates the maximum amount of stress, which is thought to induce stress graphitization in the carbon matrix (at the interface). As a result, a large amount of porosity (previously closed) gets opened up. Likewise, in composite C interactions are weak and hence minimum degree of stress graphitization is induced. As a consequence there is minimum open porosity in composite C. By comparing the data of bulk and kerosene density it is observed that the bulk density and kerosene density generally increase in all the composites with heat treatment to 2600 8C. It is further observed that bulk density is minimum and kerosene density is maximum upon heat treatment at 2600 8C in composite A, as compared to composites B and C. This is again thought to be due to the fact that, when carbon composites are heat treated above 2200 8C they become graphitized and, as a consequence of structural reorganization significant opening porosity is created. The extent of graphitization will be greater in ST-3 fiber-containing composite (e.g., due to frozen stresses). As a result, kerosene density of composite A is the highest even through its bulk density is minimum.
Fig. 4. Variation in CTE in transverse direction of composites A, B and C with stages of heat treatment.
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Fig. 5. Change in flexural strength of composites with stages of heat treatment.
3.6. Coefficient of thermal expansion ( CTE) vs. HTT of composites The thermal expansion measurement gives an idea of dimensional stability and the preferred orientation and graphitization of unidirectional carbon–carbon composites. The CTE parallel to the fiber axis is dominated by fiber thermal expansion [35]. In the transverse direction, however, the CTE contribution from the matrix is significant as well, as it would depend on the type of fibers and their surface characteristics. Therefore in the present study CTE measurements were carried out in the transverse direction only. Fig. 4 shows the CTE variation in the transverse direction. At 1000 8C the CTE of composite C is maximum that of composite A is minimum. The minimum value of CTE in composite A is attributed to the disordered structure of the carbon fibers. (In composite C the carbon fibers possess more ordered structure). The lower thermal expansion is a clear manifestation of the presence of a strong boundary restraint and anisotropic contraction of fibers and matrix. With increasing HTT, the CTE in all cases registers a gradual increase to almost 2000 8C, which is attributed to marginal improvements in the matrix structure. Above 2000 8C there is a sharp increase for composite A and moderate increases for composites B and C. The former is thought to be due to the onset of stress graphitization [11]. Stress graphitization clearly depends on the degree of frozen stresses in the composites which
are expected to be maximum in composite A and minimum in composite C. This would explain why composite A shows a maximum rise and composite C a minimum CTE rise between 2000 and 2600 8C [11]. This suggests that strong intercrystalline bonds in composite C prevent the full strain free c-axis due to minimum graphitization.
3.7. Flexural strength vs. HTT of composites Variations in flexural strength with HTT in the range 150 to 2600 8C are shown in Fig. 5. The strength of polymer composites made with high strength carbon fibers (composite A) is higher than that of composites made with intermediate and high modulus carbon fibers (composites B and C). The ST-3 fibers possess maximum reactive functional groups as compared to IM-500 and HM-45 carbon fibers (Section 3.1). As a result, ST-3 fibers are expected to have strongest interactions with the polymer matrix and thus possess a high stress transfer capability. With heat treatment, all the composites show a decrease in flexural strength up to a particular temperature and above it an upward trend is observed. In the case of composite A, flexural strength decreases suddenly from polymer stage to 600 8C, it increases thereafter to a maximum HTT of 2600 8C. On the other hand, flexural strength of composites B and C decreases from polymer stage to 1400 and 1000 8C, respectively, in the same manner as observed for composite A. The minima in flexural strength in compos-
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ites A, B and C are observed at 1800, 1400 and 1000 8C, respectively. In carbon fiber reinforced polymer composites, the flexural strength depends upon fiber–matrix bonding and mechanical properties of the fibers and the matrix precursor. The strength of phenolic resin decreases upon heat treatment to 600 8C [40]. Therefore, all the composites exhibit a steady decrease in strength upon heat treatment. Above 800 8C, the matrix becomes strong and composite failure is controlled by crack propagation via fiber–matrix interactions. In composites A and B, the fracture would initiate at the tip of prestressed microcracks perpendicular to the fiber surface [7]. However, in such composites the energy at the tip of the notches or cracks is expected to be smaller than the bonding energy between the fibers and the matrix. The crack, would therefore not stop at the fiber surface but pass straight through the composite causing it to fail at low loading and low flexural strength [40]. In composite C, the fiber–matrix bonding is quite weak and crack branching is expected to occur easily at the interface. The matrix crack propagates to the fiber surface, stops there, is eliminated by absorbing fracture energy or propagates further along the fiber–matrix interface. Hence, composite fracture by this mechanism requires a higher fracture energy and therefore composite C exhibits a high ultimate breaking strength [41]. On the other hand, the strain to failure of high modulus carbon fibers and carbonized matrix are nearly in the same range [6]. Therefore, composite C would fail at the strain level at which fibers or matrix reach the maximum obtainable stress. From Fig. 5 it is observed that the flexural strength, after reaching a minimum value, in each case registers an upward trend. The mechanical properties of brittle materials are highly notch sensitive. As described in our earlier
Fig. 7. Optical micrographs of composites A, B and C heattreated at 2600 8C.
Fig. 6. Change in ILSS of composites with stages of heat treatment.
study [42], the thermal stresses keep increasing upon heat treatment of composites and, as a result, the d spacing in all three composites increases, but only up to 2200 8C, falling thereafter because of stress relaxation [42]. However, the deflection temperature is different for each composite. Due to the generation of thermal stresses, many microcracks are expected in the matrix, which will be
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maximum in composite A, followed by composites B and C. As argued above, these microcracks act as centers of crack initiation and explains the varying deflection temperatures. For composites A, B and C these are |1800, 1400 and 1000 8C, respectively, and they indicate that the stress relaxation depends on the severity of fiber–matrix interactions established at the polymer stage. As the degree of interactions increases (strong fiber–matrix bonding), the higher should be the deflection temperature. Above the deflection temperature, an improvement in the flexural strength is attributed to modification of the interface exhibiting improved crack tip deflection [43,44]. This is more glaring in composite C which shows 50% higher strength (at 2600 8C) than composite A, even though the strength of carbon fibers in composite C is only 50% of the strength of carbon fibers in composite A.
3.8. ILSS vs. HTT of composites Fig. 6 shows the changes in interlaminar shear strength of composites A, B and C with HTT. At polymer stage (150 8C) composite A has the maximum ILSS followed by composites B and C. The maximum ILSS is a direct evidence of strong fiber–matrix interaction in composite A, while the minimum ILSS in composite C is evidence of weak interactions (see Section 3.2). The fibers in composite A possess maximum reactive functional groups and the fibers in composite C possess minimum reactive functional groups. Upon heat treatment, due to the evolution of volatile products and reorientation of matrix structure, the interfacial strength between fibers and matrix degrades, and as a result, ILSS decreases in all the composites. The extent of this decrease is higher in composite A and lower in composite C. This is directly related to the volume shrinkage observed during HTT (see Fig. 2). The higher the cross sectional shrinkage (composite A) the more the interfacial flaws, thus explaining the lower ILSS value in of composite A at 1000 8C. In all the composites, in the temperature region 1000–1800 8C, the value of ILSS is nearly the same but above 1800 8C it increases and is maximum in case of composite A. The maximum ILSS value at 2600 8C for all the composites is directly related to the severity of interactions established at the polymer stage which persist at higher heat treatment temperature. Lower ILSS values for composites B and C are also directly related to weak interactions established at the polymer stage which persist with HTT. These observations are in agreement with those of Dhami et al. [45]. It is also observed in the optical micrographs (Fig. 7) that, in the case of composite A heat treated at 2600 8C, the fiber and the matrix are in close proximity due to the strong fiber– matrix interaction even at 2600 8C. Composite A shows a lamellar type texture with strong extinction lines or a well
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defined columnar type texture which is due to the strong fiber–matrix interactions [12], while in composite B a lamellar type texture with less extinction lines is observed. In composite C a predominantly lamellar type texture is observed which is known to be due to the weak fiber– matrix interactions [46].
4. Conclusions It is observed in the present investigation that the type of fiber–matrix interactions (i.e., strong or weak) depends on the relative percentage of reactive functional groups present on the carbon fiber surface and not the total amount of functional groups present. In the case of high strength (ST-3) carbon fibers, almost 10% of the functional groups are reactive as compared to 5.5% and 4.5% in the case of intermediate (IM-500) and high modulus (HM-45) carbon fibers. Consequently, ST-3 carbon fibers are in a position to establish strongest interactions and HM-45 the weakest interactions with the matrix. The volume shrinkage occurring during heat treatment (pyrolysis) is maximum in the case of ST-3 and minimum in the case of HM-45 composites. The flexural strength and interlaminar shear strength in all the composites is maximum at the polymer stage and decreases with HTT up to a characteristic temperature (deflection temperature). To the best of our knowledge, this is the first report that the deflection temperature depends upon the severity of fiber–matrix interaction established at the polymer stage. Above the deflection temperature, flexural strength increases due to a modification of the interface to favor improved crack tip deflection. The coefficient of thermal expansion is confirmed to depend strongly upon fiber–matrix interactions and matrix structure.
Acknowledgements The authors are grateful to Dr. K. Lal, Director, National Physical Laboratory, New Delhi for his kind permission to publish these results. The authors also thank Dr. Anil K. Gupta, Head of Engineering Materials Division for the encouragement through out this investigation.
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[2] Fitzer E. The future of carbon–carbon composites. Carbon 1987;5:163–90. [3] Weisshaus H, Kenig S, Siegmann A. Effect of materials and processing on the mechanical properties of c–c composites. Carbon 1991;29:1203–20. [4] Dillon F, Thomas KM, Marsh H. The influence of matrix microstructure on mechanical properties of CFRC composites. Carbon 1993;31:1337–48. [5] Fitzer E, Geigle KH, Weiss R. Chemical interactions between carbon fiber surface and epoxy resins. Carbon 1980;18:389– 99. [6] Bradshaw WG, Vidoz AE. Fiber–matrix interactions in unidirectional carbon–carbon composites. Ceram Bull 1978;57:193–8. [7] Fitzer E, Geigle KH, Huettner W. Influence of carbon fiber surface treatment on the mechanical properties of carbon– carbon composites. Carbon 1980;18:265–70. [8] Manocha LM, Yasuda E, Tanabe Y, Kimura S. Effect of carbon fiber surface treatment on mechanical properties of c–c composites. Carbon 1988;26:333–7. [9] Manocha LM, Bahl OP, Singh YK. Mechanical behaviour of carbon–carbon composites made with surface treated carbon fibers. Carbon 1989;27:381–7. [10] Bahl OP, Dhami TL. In: Proceedings of the International Conference ‘‘Carbon and Carbonaceous Composite Materials’’, Malenovice, October 1995. [11] Zaldivar RJ, Rellick GS. Some observation on stress graphitization in carbon–carbon composites. Carbon 1991;29:1155–63. [12] Hishiyama Y, Inagaki M, Kimura S, Yamada S. Graphitization of carbon fiber / glassy carbon composites. Carbon 1974;12:249–58. [13] Manocha LM. The effect of heat treatment temperature on the properties of polyfurfuryl alcohol based carbon–carbon composites. Carbon 1994;32:213–23. [14] Chlopek J, Blzewicz S. Effect on processing variables on the properties of carbon–carbon composites. Carbon 1991;29:127–31. [15] Takano T, Uruno T, Kinjo T, Tlomak P, Ju CP. Structure and properties of unidirectionally reinforced PAN-resin based composites. J Mater Sci 1993;28:5610–9. [16] Dhakate SR, Parashar VK, Raman V, Bahl OP. Effect of titania (TiO 2 ) interfacial coating on mechanical properties of carbon–carbon composites. J Mater Sci Lett 2000;19(8):699–701. [17] Ishitani A. Application of X-ray photoelectron spectroscopy to surface analysis of carbon fiber. Carbon 1981;19:269–75. [18] Nakayama Y, Soeda F, Ishitani A. XPS study of the carbon fiber matrix interface. Carbon 1990;28:21–6. [19] Proctor A, Sherwood PMA, X-Ray photoelectron spectroscopic studies of carbon fiber surface II. The effect of electrochemical treatment. Carbon 1983;21:53–9. [20] Kozlowski C, Sherwood PMA. X-ray photoelectron spectroscopic studies of carbon fiber surface VIII. A comparison of type I and II fibers and their interactions with thin resin film. Carbon 1987;25:751–60. [21] Jones C, Samman E. The effect of low power plasmas on carbon fiber surfaces. Carbon 1990;28:509–14. [22] Jones C. Effect of electrochemical and plasma treatment on carbon fiber surfaces. Surf Interface Anal 1993;20:357–67.
[23] Puri BR. In: Walker PL, editor, Chemistry and physics of carbon, vol. 6, New York: Marcel Dekker, 1971, p. 191. [24] Bansal RC, Donnet JB, Stoeckli HF. In: Active carbon, New York: Marcel Dekker, 1988, p. 119. [25] Boehm HP, Voll M. Basische oberflachenoxide auf kohlenstoff—I. Adsorption von sauren. Carbon 1970;8:227–40. [26] Contescu A, Vass M, Contescu C, Putyera K, Schwarz JA. Acid buffering capacity of basic carbons revealed by their continuous pK distribution. Carbon 1998;36:247–58. [27] Boehm HP. Some aspects of the surface chemistry of carbon blacks and other carbons. Carbon 1994;32:759–69. [28] Papirer E, Li S, Donnet JB. Contribution to the study of basic groups on carbon. Carbon 1987;25:243–8. [29] Papirer E, Dentzer J, Li S, Donnet JB. Surface group on nitric acid oxidized carbon black sample determined by chemical thermodesorption. Carbon 1991;29:69–72. [30] Drago RS, Voel GC, Needham TE. A four-parameter equation for predicting enthalpies of adduct formation. J Am Chem Soc 1971;93:6014–26. [31] Drago RS, Parr LB, Chamberlain GS. Solent effects and their relationship to the E and C equation. J Am Chem Soc 1977;99:3203–9. [32] Krekel G, Huttinger KJ, Hoffman WP. The relevance of surface chemistry of carbon fiber in their adhesion to high temperature thermoplastic. Part II. Surface chemistry. J Mater Sci 1994;29:3461–8. [33] Dhakate SR, Bahl OP, Sahare PD. Oxidation behaviour of PAN based carbon fiber reinforced phenolic resin matrix composites. J Mater Sci Lett 2000;19(21):1959–61. [34] Jenkins GM, Kawamura K. Polymeric carbons. Cambridge University Press, 1978. [35] Morgan WC. Thermal expansion of graphite crystals. Carbon 1972;10:73–80. [36] Choe CR, Lee KH. Effect of processing parameters on the mechanical properties of carbonized phenolic resin matrix. Carbon 1992;30:247–9. [37] Kipling JJ, Sherwood JN, Shorter PV, Thompson NR. The pore structure and surface area of high temperature polymer carbons. Carbon 1964;1:321–8. [38] Mehrotra BN, Bragg RH, Rao AS. Effect of heat treatment temperature (HTT) on density, weight and volume of glasslike carbon (GC). J Mater Sci 1983;18:2671–8. [39] Savage G. Carbon–carbon composites. Chapman and Hall, 1993. [40] Lausevic Z, Marinkovic S. Mechanical properties and chemistry of carbonization of phenol formaldehyde resin. Carbon 1986;24:575–80. [41] Manocha LM. Changes in physical and mechanical properties of carbon fiber-reinforced polyfurfuryl alcohol composites during their pyrolysis to carbon–carbon composites. Composites 1988;19:311–9. [42] Dhakate SR, Bahl OP, Mathur RB, Dhami TL. Influence of interface on crystalline parameters, coefficient of thermal expansion and electrical conductivity in carbon–carbon composites. In: Carbon 2001 Conference, Kentucky, KY, 14–19 July 2001, Session 23, Paper 23.3. [43] Manocha LM, Bahl OP. Influence of carbon type and weave pattern on the development of 2D carbon–carbon composites. Carbon 1988;26:13–22. [44] McEnaney B, Mays T. Relation between microstructure and
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Effect of carbon fiber surface functional groups on the mechanical properties of carbon–carbon composites with HTT S.R. Dhakate*, O.P. Bahl Carbon Technology Unit, Engineering Materials Division, National Physical Laboratory, Dr. K.S. Krishnan Marg, New Delhi 110012, India Received 18 March 2001; accepted 25 January 2003
Abstract The present investigation describes the quantitative measurement of surface functional groups present on commercially available different PAN based carbon fibers, their effect on the development of interface with resol-type phenol formaldehyde resin matrix and its effect on the physico–mechanical properties of carbon–carbon composites at various stages of heat treatment. An ESCA study of the carbon fibers has revealed that high strength (ST-3) carbon fibers possess almost 10% reactive functional groups as compared to 5.5 and 4.5% in case of intermediate modulus (IM-500) and high modulus (HM-45) carbon fibers, respectively. As a result, ST-3 carbon fibers are in a position to make strong interactions with phenolic resin matrix and HM-45 carbon fibers make weak interactions, while IM-500 carbon fibers make intermediate interactions. This observation is also confirmed from the pyrolysis data (volume shrinkage) of the composites. Bulk density and kerosene density more or less increase in all the composites with heat treatment up to 2600 8C. It is further observed that bulk density is minimum and kerosene density is maximum upon heat treatment at 2600 8C in case of ST-3 based composites compared to HM-45 and IM-500 composites. It has been found for the first time that the deflection temperature (temperature at which the properties of the material start to decrease or increase) of flexural strength as well as interlaminar shear strength is different for the three composites (A, B and C) and is determined by the severity of interactions established at the polymer stage. Above this temperature, flexural strength and interlaminar shear strength increase in all the composites up to 2600 8C. The maximum value of flexural strength at 2600 8C is obtained for HM-45 composites and that of ILSS for ST-3 composites. 2003 Published by Elsevier Science Ltd. Keywords: A. Carbon fibers; C. Electron spectroscopy; D. Functional groups, Mechanical properties
1. Introduction Carbon–carbon composites are used in a number of demanding applications such as space, defense, turbine blades, etc. [1,2]. Their performance is known to depend on the type of carbon fibers, matrix precursors, nature of bonding between fiber and matrix (fiber–matrix interface) and processing conditions [3,4]. The nature of the interface and its influence on physico–mechanical properties is extremely complex. Different types of interface may exist
*Corresponding author. Tel.: 191-11-574-6290; fax: 191-11572-6952. E-mail address: [email protected] (S.R. Dhakate).
in carbon–carbon composites between fiber and matrix, within fiber bundles, between the different microstructures which may exist within the matrix. It is well known that the fiber–matrix interaction depends upon the surface functional groups of the carbon fibers and the matrix precursor [3,4]. In particular, the mechanical properties of carbon–carbon composites are very sensitive to the bonding between fibers and the matrix and its stress transfer capability. In this respect, continuous research and development work is going on from the last 3–4 decades, and especially regarding the role of interface on controlling the overall performance of carbon–carbon composites [5–10]. A qualitative correlation between the amount of surface functional groups, the nature of the interface and composite properties has been reported by Fitzer et al. [7]. Also,
0008-6223 / 03 / $ – see front matter 2003 Published by Elsevier Science Ltd. doi:10.1016 / S0008-6223(03)00051-4
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Table 1 Properties of carbon fibers used (from manufacturer’s data sheet) Fiber type
Diameter (mm)
Density (g / cm 3 )
Tensile strength (MPa)
Tensile modulus (GPa)
Strain to failure (%)
ST-3 IM-500 HM-45
6.0 5.5 6.0
1.75 1.77 1.91
4200 4700 2200
250 300 440
1.68 1.57 0.5
there are plenty of data available on the analysis of surface functional groups of carbon fibers and their influence on the development of interface and mechanical properties of polymer matrix composites. In addition to this, some data are available in the literature describing the effect of heat treatment temperature (HTT) on the mechanical properties of polyarylacetylene (PAA) and furfuryl alcohol based carbon–carbon composites [11–13]. Some studies are also available on the influence of fiber surface functional groups and various types of surface treated carbon fibers on the development of interface with phenolic resin matrix [7,14,15]. In the present investigation, a systematic approach was adopted to understand the influence of carbon fiber surface functional groups on interface development by measuring the surface functional groups quantitatively and their influence on composite properties as a function of HTT.
2. Composite preparation and characterization Three types of commercially available PAN-based carbon fibers, manufactured by Toho Beslon Inc., and Torayca Industries Inc., Japan, were used as reinforcement. (a) ST-3 (high strength, likely HTT |1200–1500 8C). (b) IM-500 (intermediate modulus, likely HTT |1600– 1800 8C). (c) HM-45 (high modulus, likely HTT .2200 8C) Since these fibers are heat-treated to different temperatures, they exhibit different physical and mechanical properties (see Table 1) as well as different types of surface functional groups. The surface functional groups present were measured by ESCA, using an SSI 301 spectrometer employing monochromatic and focused Al Ka radiation (spot diameter 300 mm, 80 W; radiation energy 1486.6 eV) under a residual pressure of 5310 28 Torr (1 Torr5133.322 Pa). Unidirectional polymer composite samples (150 mm3 4.0 mm34.5 mm) were prepared using the wet winding and match mold die technique [16] with 4562% fiber volume. The resol type phenol formaldehyde resin was used as the matrix precursor for composite preparation.
The composites were coded as follows. (A) ST-3 carbon fiber composites. (B) IM-500 carbon fiber composites. (C) HM-45 carbon fiber composites. The polymer composites were heat-treated to 400– 2600 8C under an inert atmosphere. The heat-treated composites were characterized for volume shrinkage, density and mechanical properties during each stage of heat treatment. Flexural strength was measured by the three-point bending technique on a universal Instron testing machine (Model 4411, ASTM standard D-790-80) with a span length to depth ratio of 30:1. The interlaminar shear strength (ILSS) was measured using ASTM standard D-2344-74 with span length to depth ratio of 8:1. The kerosene density was measured by the kerosene pickup method using the Archimedes principle. The transverse coefficients of thermal expansion were measured using a thermo-mechanical analyzer (TMA) attached to a Mettler thermal system TA-3000, in the range 50–900 8C under an inert atmosphere. The optical micrographs of composites heat treated at 2600 8C were observed using polarized light microscopy.
3. Results and discussion
3.1. ESCA studies of carbon fibers The surface composition of carbon fibers obtained by
Table 2 Surface composition of carbon fibers Fiber type \element
Atomic % C (1s)
O (1s)
N (1s)
Si (2p)
ST-3 IM-500 HM-45
86.63 92.28 92.21
8.19 7.23 6.18
2.35 0.00 0.00
1.86 0.49 1.60
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Fig. 1. (a) ESCA spectra of (a) ST-3, (b) IM-500 and (c) HM-45 carbon fibers. (b) Deconvolution spectra of O 1s for (a) ST-3, (b) IM-500 and (c) HM-45 carbon fibers.
Table 3 Relative percentage of functional groups on the carbon fibers obtained by ESCA Functional group
Graphitic carbon (B. E., eV) Phenolic or hydroxyl (B. E., eV) Carbonyl (B. E., eV) Carboxylic (B. E., eV) Nitrogen containing (B. E., eV)
Relative atomic percentages of functional groups ST-3
IM-500
HM-45
69.9 (284.18) 18.76 (285.42) 7.01 (287.34) 4.31 (290.13) 2.35 (398.02, 400.43)
61.9 (284.06) 28.6 (285.08) 5.44 (287.90) 4.05 (290.48) –
90.88 (284.05, 284.83) – 4.3 (288.49) 4.8 (290.92) –
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the ESCA studies is presented in Table 2. As expected [17], the oxygen content on the surface decreases from ST-3 to HM-45 and the C / O ratio increases from ST-3 to HM-45, confirming the well known fact that the oxygencontaining functional groups, and consequently the reactivity of carbon fiber surface, decrease with increasing HTT. The nitrogen-containing functional groups are present only in ST-3 fibers, from the residual nitrogen present in the precursor fibers. Fig. 1a shows the ESCA spectra of the three carbon fibers. The three major peaks observed between binding energies 200 to 600 eV of carbon (C 1s), oxygen (O 1s) and nitrogen (N 1s) correspond to graphite-like carbon and various functional groups. The peak around 101 eV corresponds to silicon Si (2p). The C 1s peak at 284.24 eV is assigned to the carbon element only and between 284.05 and 284.83 eV to graphitic carbon, 285.08–285.42 eV to hydroxyl (.C–OH), 287.34–288.49 eV to carbonyl (.C=O) and 290.13–290.92 eV to carboxylic (–COOH) functional groups [18–20]. The chemical shift for all the functional groups of three different carbon fibers is different (Fig. 1b, Table 3), which may be due to the structural difference of the carbon fibers surfaces. The fibers used in the investigation are heat-treated at increasing temperature (ST-3 to HM-45) as described in Section 2. As a consequence, orientation of carbon layers preferentially improves parallel to fiber axis leading to different degrees of graphitization and therefore different electrical conductivities of the fibers. The ESCA spectrum for the ST-3 carbon fibers shows a very weak N 1s peak as compared to the C 1s and O 1s peaks, which arises from the residual nitrogen present in the fibers. Nitrogen-containing functional groups appear in ST-3 fibers at a binding energy of 398.02 eV, corresponding to the aromatic amines and piperidine structure, at 400.43 eV to aliphatic amines, nitrile and amides [21,22], while in the cases of IM-500 and HM-45 carbon fibers these groups are not detected. Deconvolution of these peaks gives relative percentages of functional groups present and these are compiled in Table 3. The high strength carbon fibers (ST-3) possess the highest number of surface functional groups and as a result the surface of these fibers may be more disordered possessing comparatively high active surface area. The ST-3 fibers possess carboxylic, phenolic and hydroxyl functional groups which are acidic in nature while carbonyl and some of nitrogen containing groups are basic in character [23–29]. The high modulus HM-45 carbon fiber, on the other hand, possesses a minimum number of surface functional groups and a better ordered graphite-like carbon fiber surface. Therefore such fibers should have the lowest active surface area. The intermediate modulus (IM-500) carbon fibers possess the greatest number of hydroxyl functional groups. The contribution of carboxylic groups is almost the same in all the three fibers while carbonyl groups decrease from ST-3 to HM-45 fibers.
3.2. Interactions of fibers with matrix The specific interactions are postulated to be Lewis acid–base type interactions or electron acceptor–donor interactions [30,31]. The polymer used here is a thermosetting phenolic resin which has acidic functional groups whereas fibers possess both acidic and basic functional groups. It is well known that carboxylic groups are responsible for strong interactions with a polymer matrix having basic functional groups [18,32]. Therefore, in the present case these groups are not likely to play an important role. It is found that ST-3 fibers possess maximum (9.46 relative percentage) functional groups which are basic in character and would make strong interactions, whereas HM-45 fibers possess minimum reactive functional groups (4.3 relative percentage) and would thus make weak interaction with the phenolic resin matrix. On the other hand, the IM-500 fibers possess a maximum amount of total functional groups (38.09 relative percentage) but the reactive functional groups are only 5.44 relative percentage (of carbonyl), which can make interactions with phenolic resin matrix intermediate between ST-3 and HM-45 carbon fibers [33].
3.3. Volume shrinkage vs. HTT of composites During heat treatment, carbon fiber reinforced polymer composites show changes in all the three directions of composites, i.e., length, width and thickness, which are due to the thermal degradation and shrinkage of polymer matrix [7,34]. In unidirectional composites, the changes taking place in the direction parallel to the fiber axis are controlled by the longitudinal thermal expansion of the carbon fibers themselves. The fiber thermal expansion (positive or negative) is very small [35]. As a consequence,
Fig. 2. Volume shrinkage observed with heat treatment temperature of composites.
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these composites exhibit very little change parallel to carbon fibers. However, appreciable changes in the width and thickness directions are observed. Fig. 2 shows the volume shrinkage of polymer composites with heat treatment temperature. The amount of shrinkage depends upon the type of polymer matrix used in the fabrication of composites and fiber–matrix interactions [9]. During heat treatment, chemical bonding between the fiber surface and the free functional groups of the resin undergoes rearrangement. Simultaneously, the molecular chains in the polymer matrix undergo pyrolysis and their rearrangement results in matrix shrinkage. Up to 400 8C the shrinkage pattern of all the composites does not show much difference, but with increasing HTT shrinkage increases linearly in all the composites. For composite A it is maximum during the whole range of temperature and minimum in case of composite C. Up to 800 8C, 15% shrinkage is noticed in the case of composite A, whereas in composites B and C it is 11–12%. The higher shrinkage in the case of composite A is attributed to strong fiber–matrix interactions. With increase in heat treatment temperature, shrinkage increases continuously in all the composites and it is |25% in composite A, 22% in composite B and 19% in composite C at 1800 8C. Above 1800 8C, the conversion of non-graphitic carbon into a graphite-like carbon structure (in the matrix), i.e., a reorientation of graphitic planes, starts taking place thus resulting in further composite shrinkage [11]. Further, with a continuous increase in heat treatment temperature to 2600 8C, the reorientation of graphitic planes takes place more extensively and the extent of shrinkage observed depends on the extent of reorientation. The incremental shrinkage between 2000 and 2600 8C is as high as 6–7% in composite A and only 4% in composite C. It is important to mention here that the initial fiber volume (456 2%) is kept the same in all three cases. This brings out very clearly the effect of reactive surface functional groups on the the fiber–matrix interactions. Consequently, during heat treatment the matrix
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shrinks onto the fibers when the interaction is maximum (as in composite A). Such composites should exhibit a maximum amount of shrinkage, as is observed. In composite C, fiber–matrix interactions are weak (see Section 3.2) and as a result there is minimum shrinkage.
3.4. Bulk density vs. HTT of composites The bulk or apparent density is the density of composites containing voids and porosity. Table 4 shows the changes in density observed with HTT of the composites. These changes in density may be due to two factors:
(i) Dimensional changes due to shrinkage during pyrolysis. (ii) Weight loss due to evolution of volatile products during pyrolysis.
The initial density of the polymer composites depends upon the density of carbon fibers used, its volume fraction (which is kept the same in all three composites) and compactness of the composites. The bulk density of polymer composites varied between 1.42 and 1.52 g / cm 3 . Upon heat treatment to 600 8C, a gradual density decrease is observed due to evolution of reaction products and formation of pores which results in volume expansion [36]. Above 600 8C, product evolution decreases to a large extent as pyrolysis of the resin matrix is almost complete and there is an increase in density up to 1400 8C which is due to structural changes. The small density decrease between 1400 and 1800 8C in composites A and B is due to the formation of closed microporosity [37,38] because of evolution of nitrogen-containing reaction products.
Table 4 Variation in bulk density of composites HTT (8C) 150 400 600 800 1000 1400 1800 2200 2600
Bulk density (g / cm 3 ) A
B
C
1.42 1.37 1.36 1.41 1.45 1.51 1.50 1.53 1.55
1.46 1.42 1.40 1.45 1.50 1.55 1.53 1.55 1.57
1.53 1.45 1.44 1.50 1.53 1.55 1.57 1.58 1.61
Fig. 3. Change in kerosene density of composites with stages of heat treatment.
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Table 5 Porosity in composites heat-treated at 2600 8C* Porosity (%)
Composite type A
B
C
Total Closed Open
31.4 17.5 13.9
30.5 20.2 10.3
28.8 20.8 8
* Determined from true (2.26 g / cm 3 ) and bulk densities.
Composites A and B were made from carbon fibers which are not heat treated to more than 1400–1800 8C and therefore, contain some amount of nitrogen containing functional groups on their surface. Composite C, on the other hand, does not show any decrease in the density since HM-45 carbon fibers (likely heat treatment temperature.2200 8C) are nitrogen free (Table 2). Further, on heat treatment to 2200 and 2600 8C, the bulk density gradually increases in all the composites, because of breakdown of microporosity due to structural rearrangement of carbon atoms. The final bulk density of composites varies between 1.55 and 1.61 g / cm 3 . The bulk density of all the composites increases more or less with HTT. In the case of glass like carbon it decreases with HTT due to the continuous volume expansion of pores [38]. In the composites the volume expansion is thought to be controlled by fiber–matrix interactions.
3.5. Kerosene density vs. HTT of composites In order to get an idea about the total open porosity in the composites, experiments were conducted to measure kerosene density since kerosene is known to wet all carbons is capable of entering pores of up to 50 nm in size. The kerosene density and bulk density give information about the open porosity present in the composites. Fig. 3 shows changes in the kerosene density of composites A, B and C with HTT. Initially at 1000 8C the density of all the composites is in the range 1.60–1.62 g / cm 3 . In the case of composites A and B, the density increases moderately between 1000 and 2000 8C, and then it registers a sharp increase from 1.64 to 1.80 g / cm 3 for composite A while it increases from 1.65 to 1.75 g / cm 3 for composite B. The sudden density increase in the case of composite A may be due to preferential graphitization (in the immediate vicinity of carbon fibers) of isotropic carbon derived from phenolic resin as well as the graphitization of carbon fibers [12] (see Fig. 7, composite A). However, in case of composites B and C kerosene density registers a continuous increase with the ultimate value of 1.75 g / cm 3 only at 2600 8C. This observation clearly brings out the role of fiber–matrix interaction in affecting
graphitization (and hence high density) of the matrix as well as the carbon fiber surface. Table 5 gives information about the total, open as well as closed porosity in the three composites heat treated at 2600 8C. Even though the same matrix has been used, open porosity is maximum (13.9%) in composite A and minimum (8%) in composite C. It is well known that when soft carbon is heat treated to temperatures of the order of 2200 8C and above, it becomes graphitized and, as a consequence of structural reorganization, a lot of porosity gets opened up [39]. However, in the present case, even though the maximum HTT (2600 8C) is the same in all the composites, the open porosity values are entirely different. In case of composite A, fiber–matrix interactions are strong this generates the maximum amount of stress, which is thought to induce stress graphitization in the carbon matrix (at the interface). As a result, a large amount of porosity (previously closed) gets opened up. Likewise, in composite C interactions are weak and hence minimum degree of stress graphitization is induced. As a consequence there is minimum open porosity in composite C. By comparing the data of bulk and kerosene density it is observed that the bulk density and kerosene density generally increase in all the composites with heat treatment to 2600 8C. It is further observed that bulk density is minimum and kerosene density is maximum upon heat treatment at 2600 8C in composite A, as compared to composites B and C. This is again thought to be due to the fact that, when carbon composites are heat treated above 2200 8C they become graphitized and, as a consequence of structural reorganization significant opening porosity is created. The extent of graphitization will be greater in ST-3 fiber-containing composite (e.g., due to frozen stresses). As a result, kerosene density of composite A is the highest even through its bulk density is minimum.
Fig. 4. Variation in CTE in transverse direction of composites A, B and C with stages of heat treatment.
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Fig. 5. Change in flexural strength of composites with stages of heat treatment.
3.6. Coefficient of thermal expansion ( CTE) vs. HTT of composites The thermal expansion measurement gives an idea of dimensional stability and the preferred orientation and graphitization of unidirectional carbon–carbon composites. The CTE parallel to the fiber axis is dominated by fiber thermal expansion [35]. In the transverse direction, however, the CTE contribution from the matrix is significant as well, as it would depend on the type of fibers and their surface characteristics. Therefore in the present study CTE measurements were carried out in the transverse direction only. Fig. 4 shows the CTE variation in the transverse direction. At 1000 8C the CTE of composite C is maximum that of composite A is minimum. The minimum value of CTE in composite A is attributed to the disordered structure of the carbon fibers. (In composite C the carbon fibers possess more ordered structure). The lower thermal expansion is a clear manifestation of the presence of a strong boundary restraint and anisotropic contraction of fibers and matrix. With increasing HTT, the CTE in all cases registers a gradual increase to almost 2000 8C, which is attributed to marginal improvements in the matrix structure. Above 2000 8C there is a sharp increase for composite A and moderate increases for composites B and C. The former is thought to be due to the onset of stress graphitization [11]. Stress graphitization clearly depends on the degree of frozen stresses in the composites which
are expected to be maximum in composite A and minimum in composite C. This would explain why composite A shows a maximum rise and composite C a minimum CTE rise between 2000 and 2600 8C [11]. This suggests that strong intercrystalline bonds in composite C prevent the full strain free c-axis due to minimum graphitization.
3.7. Flexural strength vs. HTT of composites Variations in flexural strength with HTT in the range 150 to 2600 8C are shown in Fig. 5. The strength of polymer composites made with high strength carbon fibers (composite A) is higher than that of composites made with intermediate and high modulus carbon fibers (composites B and C). The ST-3 fibers possess maximum reactive functional groups as compared to IM-500 and HM-45 carbon fibers (Section 3.1). As a result, ST-3 fibers are expected to have strongest interactions with the polymer matrix and thus possess a high stress transfer capability. With heat treatment, all the composites show a decrease in flexural strength up to a particular temperature and above it an upward trend is observed. In the case of composite A, flexural strength decreases suddenly from polymer stage to 600 8C, it increases thereafter to a maximum HTT of 2600 8C. On the other hand, flexural strength of composites B and C decreases from polymer stage to 1400 and 1000 8C, respectively, in the same manner as observed for composite A. The minima in flexural strength in compos-
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ites A, B and C are observed at 1800, 1400 and 1000 8C, respectively. In carbon fiber reinforced polymer composites, the flexural strength depends upon fiber–matrix bonding and mechanical properties of the fibers and the matrix precursor. The strength of phenolic resin decreases upon heat treatment to 600 8C [40]. Therefore, all the composites exhibit a steady decrease in strength upon heat treatment. Above 800 8C, the matrix becomes strong and composite failure is controlled by crack propagation via fiber–matrix interactions. In composites A and B, the fracture would initiate at the tip of prestressed microcracks perpendicular to the fiber surface [7]. However, in such composites the energy at the tip of the notches or cracks is expected to be smaller than the bonding energy between the fibers and the matrix. The crack, would therefore not stop at the fiber surface but pass straight through the composite causing it to fail at low loading and low flexural strength [40]. In composite C, the fiber–matrix bonding is quite weak and crack branching is expected to occur easily at the interface. The matrix crack propagates to the fiber surface, stops there, is eliminated by absorbing fracture energy or propagates further along the fiber–matrix interface. Hence, composite fracture by this mechanism requires a higher fracture energy and therefore composite C exhibits a high ultimate breaking strength [41]. On the other hand, the strain to failure of high modulus carbon fibers and carbonized matrix are nearly in the same range [6]. Therefore, composite C would fail at the strain level at which fibers or matrix reach the maximum obtainable stress. From Fig. 5 it is observed that the flexural strength, after reaching a minimum value, in each case registers an upward trend. The mechanical properties of brittle materials are highly notch sensitive. As described in our earlier
Fig. 7. Optical micrographs of composites A, B and C heattreated at 2600 8C.
Fig. 6. Change in ILSS of composites with stages of heat treatment.
study [42], the thermal stresses keep increasing upon heat treatment of composites and, as a result, the d spacing in all three composites increases, but only up to 2200 8C, falling thereafter because of stress relaxation [42]. However, the deflection temperature is different for each composite. Due to the generation of thermal stresses, many microcracks are expected in the matrix, which will be
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maximum in composite A, followed by composites B and C. As argued above, these microcracks act as centers of crack initiation and explains the varying deflection temperatures. For composites A, B and C these are |1800, 1400 and 1000 8C, respectively, and they indicate that the stress relaxation depends on the severity of fiber–matrix interactions established at the polymer stage. As the degree of interactions increases (strong fiber–matrix bonding), the higher should be the deflection temperature. Above the deflection temperature, an improvement in the flexural strength is attributed to modification of the interface exhibiting improved crack tip deflection [43,44]. This is more glaring in composite C which shows 50% higher strength (at 2600 8C) than composite A, even though the strength of carbon fibers in composite C is only 50% of the strength of carbon fibers in composite A.
3.8. ILSS vs. HTT of composites Fig. 6 shows the changes in interlaminar shear strength of composites A, B and C with HTT. At polymer stage (150 8C) composite A has the maximum ILSS followed by composites B and C. The maximum ILSS is a direct evidence of strong fiber–matrix interaction in composite A, while the minimum ILSS in composite C is evidence of weak interactions (see Section 3.2). The fibers in composite A possess maximum reactive functional groups and the fibers in composite C possess minimum reactive functional groups. Upon heat treatment, due to the evolution of volatile products and reorientation of matrix structure, the interfacial strength between fibers and matrix degrades, and as a result, ILSS decreases in all the composites. The extent of this decrease is higher in composite A and lower in composite C. This is directly related to the volume shrinkage observed during HTT (see Fig. 2). The higher the cross sectional shrinkage (composite A) the more the interfacial flaws, thus explaining the lower ILSS value in of composite A at 1000 8C. In all the composites, in the temperature region 1000–1800 8C, the value of ILSS is nearly the same but above 1800 8C it increases and is maximum in case of composite A. The maximum ILSS value at 2600 8C for all the composites is directly related to the severity of interactions established at the polymer stage which persist at higher heat treatment temperature. Lower ILSS values for composites B and C are also directly related to weak interactions established at the polymer stage which persist with HTT. These observations are in agreement with those of Dhami et al. [45]. It is also observed in the optical micrographs (Fig. 7) that, in the case of composite A heat treated at 2600 8C, the fiber and the matrix are in close proximity due to the strong fiber– matrix interaction even at 2600 8C. Composite A shows a lamellar type texture with strong extinction lines or a well
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defined columnar type texture which is due to the strong fiber–matrix interactions [12], while in composite B a lamellar type texture with less extinction lines is observed. In composite C a predominantly lamellar type texture is observed which is known to be due to the weak fiber– matrix interactions [46].
4. Conclusions It is observed in the present investigation that the type of fiber–matrix interactions (i.e., strong or weak) depends on the relative percentage of reactive functional groups present on the carbon fiber surface and not the total amount of functional groups present. In the case of high strength (ST-3) carbon fibers, almost 10% of the functional groups are reactive as compared to 5.5% and 4.5% in the case of intermediate (IM-500) and high modulus (HM-45) carbon fibers. Consequently, ST-3 carbon fibers are in a position to establish strongest interactions and HM-45 the weakest interactions with the matrix. The volume shrinkage occurring during heat treatment (pyrolysis) is maximum in the case of ST-3 and minimum in the case of HM-45 composites. The flexural strength and interlaminar shear strength in all the composites is maximum at the polymer stage and decreases with HTT up to a characteristic temperature (deflection temperature). To the best of our knowledge, this is the first report that the deflection temperature depends upon the severity of fiber–matrix interaction established at the polymer stage. Above the deflection temperature, flexural strength increases due to a modification of the interface to favor improved crack tip deflection. The coefficient of thermal expansion is confirmed to depend strongly upon fiber–matrix interactions and matrix structure.
Acknowledgements The authors are grateful to Dr. K. Lal, Director, National Physical Laboratory, New Delhi for his kind permission to publish these results. The authors also thank Dr. Anil K. Gupta, Head of Engineering Materials Division for the encouragement through out this investigation.
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