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carbon composites with vertically aligned carbon nanotubes: Providing direct and indirect reinforcements to the pyrocarbon matrix
Materials and Design 92 (2016) 120–128
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Three-dimensional carbon/carbon composites with vertically aligned carbon nanotubes: Providing direct and indirect reinforcements to the pyrocarbon matrix Lei Feng a, Kezhi Li a,⁎, Zhigang Zhao a, Hejun Li a, Leilei Zhang a, Jinhua Lu a, Qiang Song b,⁎ a b
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an 710072, China Center for Nano Energy Materials, Northwestern Polytechnical University, Xi'an 710072, China
a r t i c l e
i n f o
Article history: Received 27 September 2015 Received in revised form 30 November 2015 Accepted 8 December 2015 Available online 9 December 2015 Keywords: Carbon nanotubes Carbon/carbon composites Reinforcements Mechanical properties
a b s t r a c t Vertically aligned carbon nanotubes (CNTs) were grown in situ on the surface of carbon cloths coated by a thin SiO2 layer and then the hybrid cloths were stacked and densified by chemical vapor infiltration to obtain three-dimensional (3D) C/C composites. Effects of the length (5.2–21.8 μm) of aligned CNTs on the microstructures and mechanical properties of the composites were investigated. Results showed that aligned CNTs not only directly stiffened the matrix within the reach of CNTs, but also gave indirect reinforcement to the matrix out of the reach of CNTs by inducing the formation of small pyrocarbon grains that interlocked with each other. Both the direct and indirect reinforcements on the matrix could be dramatically increased by extending CNT length. Therefore, 3D C/C composites with 21.8 μm-long CNTs showed the most notably improvements of matrix-dominated mechanical properties: 63% and 275% improvements in out-of-plane and in-plane compressive strength; meanwhile, 13% improvements in fiber-dominated flexural strength. Compared with z-pinning and stitching, the use of vertically aligned CNTs would pave a meaningful way for effectively improving the global mechanical performance of woven-fabric C/C composites. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Carbon fiber (CF)-reinforced carbon (C/C) composites possess wide structural applications in aeronautics and astronautics [1,2]. Although excellent in-plane tensile properties have been achieved using various configurations of fiber preforms, the relatively weak compression and interlaminar properties of these composites remain major issues, which are mainly dependent on the low-strength brittle carbon matrix. Carbon nanotubes (CNTs) are regarded as the ideal secondary reinforcement, due to their excellent mechanical properties [3]. The use of nano-level CNTs in the micro-level CF preforms can provide the opportunity to improve the bonding state of fiber/matrix (F/M) interface and also to modify the microstructure of pyrocarbon (PyC) and then increase the structural continuity of C/C composites [4,5], improving their through-thickness mechanical properties. So far, CNTs with different morphologies, including fluff-like CNTs, curved CNTs, CNT thincoating, randomly-oriented CNTs, and so on, have been applied to reinforce the C/C composites [6–9]. Although some improvements in matrix-dominated mechanical strengths are achieved, the increments are always limited (30–115%), due to the poor dispersion morphologies
⁎ Corresponding authors. E-mail addresses: [email protected] (K. Li), [email protected] (Q. Song).
http://dx.doi.org/10.1016/j.matdes.2015.12.036 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
of CNTs and the resulting modest reinforcements. Recently, over 200% improvements in out-of-plane compressive strength and interlaminar shear strength are obtained for short CF felt-based C/C composites after grafting radially straight CNTs, the key result of which was that the radially straight CNTs provide substantial reinforcement to PyC [10]. Undoubtedly, the significance of these results consists in the fact that it is imperative to tailor the CNT morphology for maximizing the performance of CNT doped C/C composites. However, the tailoring CNT morphology should take the types and structures of CF preforms into full consideration. As for two-dimensional (2D) carbon clothbased preforms, the weakest link in mechanical performance is the gap between adjacent fiber plies due to the large width ranging from 10 μm to over 20 μm, and the macroscale surface that can be up to several square meters dependent on the size of C/C specimen. To improve the through-thickness mechanical properties of the composites, it is necessary to provide efficient reinforcements to the interlaminar PyC matrix. For that, the use of curved CNTs, CNT coating, or radially straight CNTs would be undesirable, due to the undesirable morphologies such as short length, curved body, low dispersion density and random orientation. Aligned CNTs with micron length prepared by catalytic chemical vapor deposition (CVD) can be used as continuous reinforcements to provide several times more enhancements to polymer composites than randomly distributed CNTs [11,12]. Besides, due to the unique
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dispersion morphology, aligned CNTs can also be used as matched reinforcements in woven-fabric composites for great improvements of interlaminar mechanical properties [13,14]. Unfortunately, the carbonized photoresist surface of CFs cannot support the growth of aligned CNTs due to the poor wettability of fibers within catalyst solutions (Fe, Ni and Co) [15]. Moreover, the high temperature and reactive conditions used for CNT growth can introduce defects on CF surface, which can significantly degrade in-plane mechanical properties of C/C composites. Here, a thin inorganic SiO2 interface layer was pre-coated on CF surface to facilitate the growth of vertically-aligned CNTs by injection CVD (ICVD) with minimum degradation to the underlying CFs in a woven cloth. And then, the CF cloths coated with aligned CNTs were stacked and densified though chemical vapor infiltration (CVI) to obtain threedimensional (3D) C/C composites. The aim of growing verticallyaligned CNTs on the surface of cloths is to provide enhanced mechanical properties to woven-fabric C/C composites along the thickness direction without altering the 2D stack design. Effect of the length of aligned CNTs on their reinforcing role were evaluated under compression and bending stresses.
2. Experimental 2.1. Growth of vertically aligned CNTs on cloths Polyacrylonitrile-based carbon cloths (T300, fiber diameter: 6–7 μm, 1000 fibers in a single bundle) with surface sizing removed were used as the preform materials. CFs were coated with a SiO2 interface layer by immersing cloths into toluene containing 5 vol.% silicon tetrachloride (SiCl4) and 5 vol.% tetraethoxy silane (TEOS), and followed by the hydrolysis and pyrolysis processes. The growth of CNTs on the surface-
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treated cloths was carried out using ICVD technique according to the procedure described in our previous work [16]. Hybrid cloths with CNTs of different lengths were obtained by controlling the growth time (10, 20 and 40 min). 2.2. Composite preparation Twenty-layers of CNT-coated cloths were stacked to build the multilayer 3D preforms and then were densified using isothermal CVI technique according to the procedure reported in Ref. [16]. C/C composites without and with CNTs grown for 10, 20 and 40 min were denoted as base, 3D-10, 3D-20 and 3D-40, respectively. The preparation procedure of the 3D composites is depicted in Fig. 1. 2.3. Mechanical property tests Tensile tests were carried out to investigate the effects of surface treatment and CNT growth processes on the mechanical strength of CFs. Tensile specimens were prepared by fixing both ends of fiber bundles onto two aluminum sheets using resin with a gauge length of 20 mm. The sizes of specimens for compression and bending tests were machined to be 7 mm × 6 mm × 4 mm and 55 mm × 7 mm × 4 mm, respectively. The support span for bending tests was 40 mm. The numbers of specimens for tensile, compression and bending testing were not less than 20, 10 and 7. All the tests were carried out on a universal testing machine (CMT5304) at a constant speed of 0.5 mm/min. 2.4. Characterization The chemical composition of CFs before and after the surface treatment was characterized by X-ray photoelectron spectroscopy (XPS),
Fig. 1. Overview of preparing 3D C/C composites with aligned CNTs.
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which was conducted on an ESCALAB 250Xi using AlKα radiation at a power of 300 W. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used to examine CNT growth morphology, fracture surfaces of the failure specimens and CNT microstructure. Microstructure of the PyC matrix was analyzed using polarized light microscopy (PLM). 3. Results and discussion 3.1. XPS analysis of the surface treatments of CFs Fig. 2 presents the XPS results of the CFs before and after the surface treatment. As seen, XPS survey spectrum of the CFs presents a strong C 1 s peak at 284.5 eV, along with a much weaker O 1 s peak at 532.5 eV ascribed to physically absorbed oxygen species. The distinguishing feature of the XPS spectra of the surface-treated CFs is the appearance of Si 2p (103.5 eV, seen inset) and Si 2 s (155.5 eV) peaks, together with concomitant increase in the O 1 s and decrease in the C 1 s peaks. The Si 2p (103.5 eV) is signal from SiO2 [17]. In addition, the C 1 s peak in the XPS spectra of the surface-treated CFs arises from the carbon substrate, indicating that the SiO2 interface layer is thinner than the XPS probe depth (b10 nm). Insets of Fig. 2 present the SEM images of CFs before (down) and after (up) the surface treatment. Similar to the CFs, the SiO2-coated CFs still show smooth with evenly distributing shallow folds along the direction of fiber axis. The thin nature of the SiO2 interface layer can ensure the efficient stress transfer from matrix to the underlying CFs [18]. 3.2. Characterization of CNTs grown on cloths The surface-treated cloths after growing CNTs on its surface by ICVD process are presented in Fig. 3a, c and e. Aligned CNTs grow uniformly and vertically on all of the exposed fibers on the surface of cloth. Another finding from the SEM images is that the CNTs become denser and longer with the increase of growth time. Side views (Fig. 3b, d and f) of the CNT forests provide accurate information of the CNT lengths which are 5.2, 13.1 and 21.8 μm, corresponding to the growth time of 10, 20 and 40 min (Table 1), respectively. Apparently, the length of the CNTs is increased with the increasing growth time. We define the yield of CNTs as weight gain of the cloth during the growth process. CNTs add 1.5% to the specimen weight with 10 min growth, and up to 3.6% and 5.4% (averaged over 10 specimens) with 20 and 40 min growth, respectively. From enlarged view of the CNTs with 40 min growth shown in inset of Fig. 3f, we can also get the spacing between nanotubes (from centre to centre) are 40–120 nm, and thus the area density of the CNTs on
Fig. 2. XPS spectra and SEM images of the CFs before (down) and after (up) the surface treatment.
the CF surface is estimated to be ~1.56 × 1010 tubes/cm2. This value is roughly equal to those of CNTs grown on silica by the ICVD process [19]. In contrast, to prior work where CNTs are grown on the surface of CFs without SiO2 coating [16,20], the CNTs here obtained at the same growth time are extremely long, dense and aligned. Typical TEM image shown in Fig. 3g reveals a typical structure of bamboo-like CNTs [21]. The CNTs have high quality with no by-products (e.g., amorphous carbon and catalyst particles). Statistics from over 50 CNTs suggest that the inner and outer diameters of CNTs range from 8 to 12 nm and 15–25 nm, respectively. Higher-resolution TEM image (Fig. 3h) shows clear graphene lattices in the tube wall, indicating a high crystallinity of the produced CNTs. 3.3. Tensile strength of fiber bundles Fig. 4 shows the tensile results of fiber bundles before and after the surface treatment and CNT growth processes. It can be seen that the surface-treated fiber bundles exhibit a small relative decrease in tensile strength compared with pristine CFs, probably attributed to the thermal degradation during high-temperature pyrolysis process. However, the most encouraging result is the large strength retention rate of CFs after growing CNTs. For the surface-treated fiber bundles after 40 min CNT growth, the strength retention rate is 88.5%. This value is much higher than those of fiber bundles (69.1% of strength retention rate) without surface treatment after the CNT growth with low growth temperature and roughly same growth time [20]. The large strength retention rate of CFs can be attributed to the SiO2 protection which may prevent the Fe catalyst diffusing into the carbon substrates and the trace oxygen attack [22]. It is expected that these 3D fabric reinforcements with less damaging to the CFs can result in excellent mechanical improvement in 3D C/C composites. 3.4. Characterization of 3D composites To investigate the effect of vertically aligned CNTs on the microstructure of PyC matrix, images of cross-sections for the base and 3D composites are captured using PLM (Fig. 5a–d) and SEM (Fig. 5e–h). From Fig. 5a, much large PyC grain size, long straight grain boundaries and evident matrix annular cracks can be clearly observed in the base composites. By contrast, the PyC matrix in 3D composites is highly integrated with no annular cracks, as shown in Fig. 5b–d. Another interesting observation from PLM images of 3D composites is that the matrix exhibits sandwich-like structures. The bottom and top layers next to fiber plies are PyC with CNTs (defined as PyC/CNT) having no optical activity due to the densely packed nanotubes [7], and the intermediate layer is pure PyC containing no CNTs. Note that the intermediate PyC has smaller grain size and short curved grain boundaries, largely different from those in base composites. The PyC matrix at its electron microscopy is shown in SEM images of Fig. 5e–h, where PyC next to ply surface exhibits rough containing CNTs and aloof from fibers is smooth without CNTs but with interlocked grain boundaries. These characteristics observed in PLM and SEM images become much more distinct and clearer with the prolonged CNT length. Based on the PLM and SEM observations, corresponding models regarding the microstructure of PyC matrix of the base and 3D composites are established, as shown in Fig. 6. In the case of base composite (Fig. 6a), large PyC grains both in size and thickness are rather susceptible to cracking caused by thermal stress during the fabrication process, leading to the occurrence of matrix annular cracks [23]. Moreover, long straight grain boundaries preferentially aligned roughly either along or parallel to the fiber plies are more likely to form among these large PyC grains. After introducing aligned CNTs, two conclusions can be drawn from the change of matrix microstructure. Firstly, the vertically aligned nanotubes extending into the surrounding matrix as continuous reinforcers significantly stiffen the matrix. With the increase in CNT length, this direct reinforcement on PyC matrix can extend to the area aloof
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Fig. 3. SEM images (a–f) of SiO2-coated CFs grafted by aligned CNTs with different lengths controlled by the growth time: (a, b) 10 min; (c, d) 20 min; (e, f) 40 min. Inset of (f) is an enlarged view of the CNT forests. TEM images (g, h) of the produced CNTs.
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Table 1 Comparison of CNT length (L), volume fractions (V) of CFs and CNTs, density and open porosity of base, 3D-10, 3D-20 and 3D-40 composites. Composites
LCNT (μm)
VCF (%)
VCNT (%)
Density (g/cm3)
Open porosity (%)
Base 3D-10 3D-20 3D-40
– 5.2 13.1 21.8
38.1 34.7 33.8 32.2
– 1.2 3.0 4.2
1.62 1.63 1.60 1.59
12.8 12.6 14.1 14.9
from plies because of the expansion of 3D CNT forests into distant space. Secondly, CNTs also play an indirect reinforcing role in reinforcing the PyC matrix out of the reach of nanotubes (here, it is called as “indirect reinforcement”). Concretely, the top surface of CNT forests provides abundant nano nucleation sites, which can induce the formation of spherical or cone-shaped PyC on the surface of PyC/CNT [24]. With the proceeding of CVI process, these PyC grains contact with each other forming interlocked grain boundaries. Furthermore, the expansion of 3D CNT forests can efficiently reduce the growth space for the single PyC both in thickness and size, resulting in smaller PyC grains with shorter grain boundaries among them. These refined PyC have high cohesion and fracture resistance, which are expected to inhibit the growth and propagation of cracks under loading. Therefore, compared with CNTs reinforced polymer or ceramic composites wherein the CNTs can only modify the matrix within the reach of nanotubes, the great significance of aligned CNTs used in C/C composites is their ability to not only directly reinforce the matrix within the reach of nanotubes, but also to give indirect reinforcement to the matrix aloof from nanotubes. The microstructure models of 3D composites with CNTs of different lengths are clearly described in Fig. 6b–d. Clearly, both the direct and indirect reinforcements on the PyC matrix can be significantly increased by extending the CNT length. 3.5. Mechanical properties of 3D composites Fig. 7 shows the compressive stress–strain curves obtained for all the composite specimens. It is apparent that the out-of-plane compressive strength (OCS, Fig. 7a) and in-plane compressive strength (ICS, Fig. 7b) of the C/C composites are significantly increased after the growth of vertically aligned CNTs on cloths, especially for the composites containing long nanotubes. The curves also display a clear behavior in the specimens' response as seen by the difference in slope: specimens from 3D specimens containing long CNTs show a higher slope than the base specimen and 3D specimens containing short CNTs. This observation suggests that the former composites are stiffer than the latter composites, i.e., higher out-of-plane compressive modulus (OCM) and inplane compressive modulus (ICM). The results listed in Table 2 show that 3D-10, 3D-20 and 3D-40 have improvements in OCS of 33%, 42%
and 63%, respectively. Compared with the base composites, 3D-10, 3D-20 and 3D-40 composites also have 13%, 19% and 24% improvements in OCM, respectively. The measured ICS of the 3D-10, 3D-20 and 3D-40 composites show 152%, 206% and 275% enhancements compared with base composites, respectively. The percentage improvements in the ICM of the 3D-10, 3D-20 and 3D-40 composites compared with those for the base specimens are 90%, 110% and 151%, respectively. For the base composites, upon 90° loading of specimen, massive destructive cracks are formed which propagate along the length and height of specimen and then break the base specimens into small pieces, whereas for the 3D composites the fracture only penetrates the height of specimens forming roughly 45° angle with the thickness direction resulting in fewer and bigger fragments. Further investigation is performed based on SEM observations into the failure specimen to acquire the information regarding reinforcement mechanism. The observation from the low-magnification SEM images of out-of-plane compressive fracture surface shows that compared with the completely disintegrated base specimens (Fig. 8a), the crushed PyC is significantly decreased and even absent in the fragments of 3D specimens (for illustrative purpose, only image for representative specimen of 3D-20 is presented, as shown in Fig. 8c). Examination of the fracture surface at higher magnification provides detailed information regarding the morphology of the fractured PyC. As stated in Fig. 8b, the surface of fractured PyC in base specimens looks flat and exhibits brittle fracture character. Thus, it can be speculated that, when the base specimens are subjected to out-of-plane stress, the matrix annular cracks and straight grain boundaries (labeled by arrows in Fig. 8b) provide the main channels for the destructive crack propagation along the length and height of the composites, respectively. But comparatively, the fractured PyC of 3D composites appears very integrated and rugged (Fig. 8d). These observations indicate that a strongly-enhanced PyC matrix change the mode of failure of the composites to a higher level of reinforcement. Furthermore, sidewalls and ends of CNTs exposed to the fractured PyC of 3D composites can be clearly observed, as shown in Fig. 8e and f, which can be considered as another important factor for the increase of OCS and OCM. On the one hand, the existence of CNT forests sharply increase the interface area between PyC matrix and reinforcements, and thus damage is more difficult to occur in 3D composites by virtue of additional surface energy and friction work because of the newly formed surface area [25]. On the other hand, the fracture of densely-packed CNTs also dissipates great energy, which can further endow the high fracture strength of the 3D composites. Fig. 9 presents the SEM images of failed composite specimens after the in-plane compression tests. When the base specimens are subjected to 0° loading stress, overall shear failure occurs characterized by delamination failures, as displayed in Fig. 9a. Enlarged SEM images reveal that the destructive cracks spread in the PyC matrix (Fig. 9b) or along the
Fig. 4. Tensile stress–strain curves (a) and corresponding strengths (b) of fiber bundles before and after surface treatment (STCF) and CNT growth processes (STCF-X, X represents the growth time).
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Fig. 5. PLM (a–d) and SEM (e–h) images of the base and 3D composites: (a, e) base; (b, f) 3D-10; (c, g) 3D-20; (d, h) 3D-40.
Fig. 6. Structure models of the base and 3D composites: (a) base; (b) 3D-10; (c) 3D-20; (d) 3D-40.
interface between matrix and outer surface of fiber bundles (Fig. 9c) until they absolutely open the plies. For the spreading of destructive cracks in PyC matrix of base composites, long straight PyC grain boundaries orientated roughly parallel to the fiber plies and the matrix
annular cracks provide the main channels for the destructive crack growth and subsequent long-distance propagation. After grafting CNT forests on the CFs, the delamination failures of 3D composites under in-plane stress are alleviated and even inhibited, especially for the
Fig. 7. Compressive stress–strain curves of the base and 3D composites: (a) out-of-plane compression and (b) in-plane compression.
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Table 2 Compressive and flexural properties of the base, 3D-10, 3D-20 and 3D-40 composites (± values represent standard deviation). Out-of-plane compression
In-plane compression
Flexure
Composites Base 3D-10 3D-20 3D-40
Strength (MPa)
Modulus (GPa)
Strength (MPa)
Modulus (GPa)
Strength (MPa)
Modulus (GPa)
158 ± 8 210 ± 11 225 ± 14 258 ± 16
3.12 ± 0.18 3.54 ± 0.11 3.72 ± 0.16 3.86 ± 0.21
48 ± 10 121 ± 9 147 ± 11 180 ± 18
5.12 ± 0.65 9.72 ± 0.54 10.75 ± 0.87 12.86 ± 0.81
135 ± 9 143 ± 13 148 ± 11 153 ± 9
22.1 ± 1.4 24.6 ± 1.2 25.4 ± 1.1 26.6 ± 1.6
composites containing long CNTs. As illustrated in Fig. 9d selected from 3D-20 specimens, where the failure initiates from the top edge of the composite and develops roughly along the diagonal of test specimen. Enlarged SEM images demonstrate that the stiffer matrix impedes the extending of cracks which want to develop along straight path
(Fig. 9e), and the resultant crack deflection leads to the fracture of fiber bundles (Fig. 9f). Great energy can be consumed during crack deflection thus resulting in the remarkable enhancement in in-plane compressive performance [26]. It is difficult to find the CNT ends and sidewalls that are observed in the out-of-plane compressive fracture
Fig. 8. Different-magnification SEM images of the out-of-plane compressive fracture surfaces of the base (a and b) and 3D-20 (c–f) composites.
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Fig. 9. SEM images of the representative failure specimens from the base (a–c) and 3D-20 (d–f) composites after the in-plane compression test (red dotted arrows represent the extending paths of the destructive cracks). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
surface of the 3D composites. Actually, compared with the stronglystiffened PyC/CNT, the destructive cracks are more likely to spread in intermediate PyC layer. For the 3D-10 composites containing short CNTs, the absence of annular cracks in intermediate PyC layer is the main reason for the improvements of ICS and ICM. With the extending of CNT forests in 3D composites, the intermediate PyC is significantly decreased in width and meanwhile has more convoluted and interlocking grain boundaries, as a result, the growth and propagation of destructive cracks in the PyC matrix of long CNT-reinforced 3D composites becomes increasingly difficult such that crack deflection occurs, resulting in the higher ICS and ICM. However, the stiffer carbon matrix may not be the only contributor to ICS and ICM enhancements, since they can also be attributed to the increased lateral support for the load-bearing plies provided by vertically aligned CNTs that may prevent the CFs from microbuckling, a critical composite failure mode associated with fibers under in-plane compressive stress [27]. The percentage enhancements in ICS and ICM obtained for 3D-40 composites are much higher than those obtained by grafting curved and radially CNTs on CFs [10,16,20]. Therefore, from the above analysis, one can comprehend that highdensity CNT forests with long extending length which not only significantly stiffen the surrounding matrix at greater depth (i.e., direct reinforcement), but also influence the peripheral matrix at higher level (i.e., indirect reinforcement), and thus largely increase the compressive properties.
Fig. 10. Flexural stress–strain curves of the base and 3D composites.
More surprisingly, the dramatic enhancement of the matrixdominated mechanical performance does not compromise the inplane properties. The flexural stress–strain curves for the test specimens are presented in Fig. 10, in which all the 3D composites show high failure flexural strength (FS) and modulus (FM) compared with the base specimens. From Table 2, the FS of 3D-10, 3D-20 and 3D-40 are improved by 6%, 10% and 13%, although the fiber volume fractions are decreased by 8.9%, 11.3% and 15.5% (Table 1), respectively. Compared with the base composites, 3D-10, 3D-20 and 3D-40 also have 11%, 15%, and 20% enhancements in FM. Another observation in Fig. 10 is that all the composite specimens show typical brittle fracture, evidenced by the sudden failure as the stress goes up to the peak of curves, which suggests that, the introduction of CNT forests increases the mechanical properties of C/C composites but not significantly changes their fracture behavior. Besides, it needs to be emphasized that the little damage to the CFs attributed to the SiO 2 protection is an important factor for the improvement of FS and FM, since the direct growth of CNTs on the CF surface leads to a serious decrease in flexural properties of C/C composites due to the excessive strength loss of CFs by iron etching and thermal degradation reported in our previous work [20].
4. Conclusions The 3D composites introduced here is the first real instance where the CNTs are used in high density, vertically aligned orientation and controllable length to effectively improve the mechanical performance of C/C composite. Vertically aligned CNTs grown on the surface-treated carbon cloths not only directly stiffen the matrix within the CNT reach, but also give indirect reinforcement to the matrix out of the CNT reach by inducing the formation of small PyC gains that interlock with each other. Both the direct and indirect reinforcements on the PyC matrix can be significantly increased by extending the CNT length. Our work demonstrates that introducing vertically aligned CNTs in the thickness direction of woven-fabric C/C composites is an effective strategy to largely increase longitudinal and transverse compressive properties without sacrificing the in-plane performance. To further understand this new type composite and to enlarge its application, further research should be done on the multifunctional properties, such as thermal conductivity, thermal diffusivity and thermal expansion behaviors.
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Acknowledgments This work has been supported by the Key Grant Project of the Chinese Ministry of Education (Grant No. 313047), the FP7-International Research Staff Exchange Scheme—Advanced Biomaterials for Regenerative Medicine, the Global Innovation Initiative—Nanostructured Materials for the Control of Contaminants Detrimental to Health, the Fundamental Research Funds for the Central universities under Grant No. 3102014JCQ01030 and “111” Project of China (B08040), the Research Fund for the Doctoral Program of Higher Education of China (Grant No. 20126102110013) and the Natural Science Foundation of China (Grant Nos. 51432008, 51202194, 51272212 and 51272213). References [1] E. Fitzer, L.M. Manocha, Carbon Reinforcements and Carbon/Carbon Composites, Christiane, Berlin, 1998. [2] G. Savage, Carbon/Carbon Composites, Springer, London, 1993. [3] H. Qian, E.S. Greenhalgh, M.S.P. Shaffer, A. Bismarck, Carbon nanotube-based hierarchical composites: a review, J. Mater. Chem. 20 (2010) 4751–4762. [4] Q.M. Gong, Z. Li, X.W. Zhou, J.J. Wu, Y. Wang, J. Liang, Synthesis and characterization of in situ grown carbon nanofiber/nanotube reinforced carbon/carbon composites, Carbon 43 (2005) 2426–2429. [5] Q.M. Gong, Z. Li, X.D. Bai, D. Li, J. Liang, The effect of carbon nanotubes on the microstructure and morphology of pyrolytic carbon matrices of C–C composites obtained by CVI, Compos. Sci. Technol. 65 (2005) 1112–1119. [6] J.S. Li, R.Y. Luo, Study of the mechanical properties of carbon nanofiber reinforced carbon/carbon composites, Compos. A Appl. Sci. 39 (2008) 1700–1704. [7] P. Xiao, X.F. Lu, Y.Q. Liu, L.L. He, Effect of in situ grown carbon nanotubes on the structure and mechanical properties of unidirectional carbon/carbon composites, Mater. Sci. Eng. A 528 (2011) 3056–3061. [8] H.L. Li, H.J. Li, J.H. Lu, K.Z. Li, C. Sun, D.S. Zhang, Mechanical properties enhancement of carbon/carbon composites by in situ grown carbon nanofibers, Mater. Sci. Eng. A 547 (2012) 138–141. [9] K.Z. Li, L. Li, H.J. Li, Q. Song, J.H. Lu, Q.G. Fu, Electrophoretic deposition of carbon nanotubes onto carbon fiber felt for production of carbon/carbon composites with improved mechanical and thermal properties, Vacuum 104 (2014) 105–110. [10] Q. Song, K.Z. Li, H.L. Li, H.J. Li, C. Ren, Grafting straight carbon nanotube radially onto carbon fibers and their effect on the mechanical properties of carbon/carbon composites, Carbon 50 (2012) 3943–3960. [11] L. Ci, J. Suhr, V. Pushparaj, X. Zhang, P.M. Ajayan, Continuous carbon nanotube reinforced composites, Nano Lett. 8 (2008) 2762–2766. [12] K.G. Dassios, S. Musso, C. Galiotis, Compressive behavior of MWCNT/epoxy composite mats, Compos. Sci. Technol. 72 (2012) 1027–1033.
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Three-dimensional carbon/carbon composites with vertically aligned carbon nanotubes: Providing direct and indirect reinforcements to the pyrocarbon matrix Lei Feng a, Kezhi Li a,⁎, Zhigang Zhao a, Hejun Li a, Leilei Zhang a, Jinhua Lu a, Qiang Song b,⁎ a b
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an 710072, China Center for Nano Energy Materials, Northwestern Polytechnical University, Xi'an 710072, China
a r t i c l e
i n f o
Article history: Received 27 September 2015 Received in revised form 30 November 2015 Accepted 8 December 2015 Available online 9 December 2015 Keywords: Carbon nanotubes Carbon/carbon composites Reinforcements Mechanical properties
a b s t r a c t Vertically aligned carbon nanotubes (CNTs) were grown in situ on the surface of carbon cloths coated by a thin SiO2 layer and then the hybrid cloths were stacked and densified by chemical vapor infiltration to obtain three-dimensional (3D) C/C composites. Effects of the length (5.2–21.8 μm) of aligned CNTs on the microstructures and mechanical properties of the composites were investigated. Results showed that aligned CNTs not only directly stiffened the matrix within the reach of CNTs, but also gave indirect reinforcement to the matrix out of the reach of CNTs by inducing the formation of small pyrocarbon grains that interlocked with each other. Both the direct and indirect reinforcements on the matrix could be dramatically increased by extending CNT length. Therefore, 3D C/C composites with 21.8 μm-long CNTs showed the most notably improvements of matrix-dominated mechanical properties: 63% and 275% improvements in out-of-plane and in-plane compressive strength; meanwhile, 13% improvements in fiber-dominated flexural strength. Compared with z-pinning and stitching, the use of vertically aligned CNTs would pave a meaningful way for effectively improving the global mechanical performance of woven-fabric C/C composites. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Carbon fiber (CF)-reinforced carbon (C/C) composites possess wide structural applications in aeronautics and astronautics [1,2]. Although excellent in-plane tensile properties have been achieved using various configurations of fiber preforms, the relatively weak compression and interlaminar properties of these composites remain major issues, which are mainly dependent on the low-strength brittle carbon matrix. Carbon nanotubes (CNTs) are regarded as the ideal secondary reinforcement, due to their excellent mechanical properties [3]. The use of nano-level CNTs in the micro-level CF preforms can provide the opportunity to improve the bonding state of fiber/matrix (F/M) interface and also to modify the microstructure of pyrocarbon (PyC) and then increase the structural continuity of C/C composites [4,5], improving their through-thickness mechanical properties. So far, CNTs with different morphologies, including fluff-like CNTs, curved CNTs, CNT thincoating, randomly-oriented CNTs, and so on, have been applied to reinforce the C/C composites [6–9]. Although some improvements in matrix-dominated mechanical strengths are achieved, the increments are always limited (30–115%), due to the poor dispersion morphologies
⁎ Corresponding authors. E-mail addresses: [email protected] (K. Li), [email protected] (Q. Song).
http://dx.doi.org/10.1016/j.matdes.2015.12.036 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
of CNTs and the resulting modest reinforcements. Recently, over 200% improvements in out-of-plane compressive strength and interlaminar shear strength are obtained for short CF felt-based C/C composites after grafting radially straight CNTs, the key result of which was that the radially straight CNTs provide substantial reinforcement to PyC [10]. Undoubtedly, the significance of these results consists in the fact that it is imperative to tailor the CNT morphology for maximizing the performance of CNT doped C/C composites. However, the tailoring CNT morphology should take the types and structures of CF preforms into full consideration. As for two-dimensional (2D) carbon clothbased preforms, the weakest link in mechanical performance is the gap between adjacent fiber plies due to the large width ranging from 10 μm to over 20 μm, and the macroscale surface that can be up to several square meters dependent on the size of C/C specimen. To improve the through-thickness mechanical properties of the composites, it is necessary to provide efficient reinforcements to the interlaminar PyC matrix. For that, the use of curved CNTs, CNT coating, or radially straight CNTs would be undesirable, due to the undesirable morphologies such as short length, curved body, low dispersion density and random orientation. Aligned CNTs with micron length prepared by catalytic chemical vapor deposition (CVD) can be used as continuous reinforcements to provide several times more enhancements to polymer composites than randomly distributed CNTs [11,12]. Besides, due to the unique
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dispersion morphology, aligned CNTs can also be used as matched reinforcements in woven-fabric composites for great improvements of interlaminar mechanical properties [13,14]. Unfortunately, the carbonized photoresist surface of CFs cannot support the growth of aligned CNTs due to the poor wettability of fibers within catalyst solutions (Fe, Ni and Co) [15]. Moreover, the high temperature and reactive conditions used for CNT growth can introduce defects on CF surface, which can significantly degrade in-plane mechanical properties of C/C composites. Here, a thin inorganic SiO2 interface layer was pre-coated on CF surface to facilitate the growth of vertically-aligned CNTs by injection CVD (ICVD) with minimum degradation to the underlying CFs in a woven cloth. And then, the CF cloths coated with aligned CNTs were stacked and densified though chemical vapor infiltration (CVI) to obtain threedimensional (3D) C/C composites. The aim of growing verticallyaligned CNTs on the surface of cloths is to provide enhanced mechanical properties to woven-fabric C/C composites along the thickness direction without altering the 2D stack design. Effect of the length of aligned CNTs on their reinforcing role were evaluated under compression and bending stresses.
2. Experimental 2.1. Growth of vertically aligned CNTs on cloths Polyacrylonitrile-based carbon cloths (T300, fiber diameter: 6–7 μm, 1000 fibers in a single bundle) with surface sizing removed were used as the preform materials. CFs were coated with a SiO2 interface layer by immersing cloths into toluene containing 5 vol.% silicon tetrachloride (SiCl4) and 5 vol.% tetraethoxy silane (TEOS), and followed by the hydrolysis and pyrolysis processes. The growth of CNTs on the surface-
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treated cloths was carried out using ICVD technique according to the procedure described in our previous work [16]. Hybrid cloths with CNTs of different lengths were obtained by controlling the growth time (10, 20 and 40 min). 2.2. Composite preparation Twenty-layers of CNT-coated cloths were stacked to build the multilayer 3D preforms and then were densified using isothermal CVI technique according to the procedure reported in Ref. [16]. C/C composites without and with CNTs grown for 10, 20 and 40 min were denoted as base, 3D-10, 3D-20 and 3D-40, respectively. The preparation procedure of the 3D composites is depicted in Fig. 1. 2.3. Mechanical property tests Tensile tests were carried out to investigate the effects of surface treatment and CNT growth processes on the mechanical strength of CFs. Tensile specimens were prepared by fixing both ends of fiber bundles onto two aluminum sheets using resin with a gauge length of 20 mm. The sizes of specimens for compression and bending tests were machined to be 7 mm × 6 mm × 4 mm and 55 mm × 7 mm × 4 mm, respectively. The support span for bending tests was 40 mm. The numbers of specimens for tensile, compression and bending testing were not less than 20, 10 and 7. All the tests were carried out on a universal testing machine (CMT5304) at a constant speed of 0.5 mm/min. 2.4. Characterization The chemical composition of CFs before and after the surface treatment was characterized by X-ray photoelectron spectroscopy (XPS),
Fig. 1. Overview of preparing 3D C/C composites with aligned CNTs.
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which was conducted on an ESCALAB 250Xi using AlKα radiation at a power of 300 W. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used to examine CNT growth morphology, fracture surfaces of the failure specimens and CNT microstructure. Microstructure of the PyC matrix was analyzed using polarized light microscopy (PLM). 3. Results and discussion 3.1. XPS analysis of the surface treatments of CFs Fig. 2 presents the XPS results of the CFs before and after the surface treatment. As seen, XPS survey spectrum of the CFs presents a strong C 1 s peak at 284.5 eV, along with a much weaker O 1 s peak at 532.5 eV ascribed to physically absorbed oxygen species. The distinguishing feature of the XPS spectra of the surface-treated CFs is the appearance of Si 2p (103.5 eV, seen inset) and Si 2 s (155.5 eV) peaks, together with concomitant increase in the O 1 s and decrease in the C 1 s peaks. The Si 2p (103.5 eV) is signal from SiO2 [17]. In addition, the C 1 s peak in the XPS spectra of the surface-treated CFs arises from the carbon substrate, indicating that the SiO2 interface layer is thinner than the XPS probe depth (b10 nm). Insets of Fig. 2 present the SEM images of CFs before (down) and after (up) the surface treatment. Similar to the CFs, the SiO2-coated CFs still show smooth with evenly distributing shallow folds along the direction of fiber axis. The thin nature of the SiO2 interface layer can ensure the efficient stress transfer from matrix to the underlying CFs [18]. 3.2. Characterization of CNTs grown on cloths The surface-treated cloths after growing CNTs on its surface by ICVD process are presented in Fig. 3a, c and e. Aligned CNTs grow uniformly and vertically on all of the exposed fibers on the surface of cloth. Another finding from the SEM images is that the CNTs become denser and longer with the increase of growth time. Side views (Fig. 3b, d and f) of the CNT forests provide accurate information of the CNT lengths which are 5.2, 13.1 and 21.8 μm, corresponding to the growth time of 10, 20 and 40 min (Table 1), respectively. Apparently, the length of the CNTs is increased with the increasing growth time. We define the yield of CNTs as weight gain of the cloth during the growth process. CNTs add 1.5% to the specimen weight with 10 min growth, and up to 3.6% and 5.4% (averaged over 10 specimens) with 20 and 40 min growth, respectively. From enlarged view of the CNTs with 40 min growth shown in inset of Fig. 3f, we can also get the spacing between nanotubes (from centre to centre) are 40–120 nm, and thus the area density of the CNTs on
Fig. 2. XPS spectra and SEM images of the CFs before (down) and after (up) the surface treatment.
the CF surface is estimated to be ~1.56 × 1010 tubes/cm2. This value is roughly equal to those of CNTs grown on silica by the ICVD process [19]. In contrast, to prior work where CNTs are grown on the surface of CFs without SiO2 coating [16,20], the CNTs here obtained at the same growth time are extremely long, dense and aligned. Typical TEM image shown in Fig. 3g reveals a typical structure of bamboo-like CNTs [21]. The CNTs have high quality with no by-products (e.g., amorphous carbon and catalyst particles). Statistics from over 50 CNTs suggest that the inner and outer diameters of CNTs range from 8 to 12 nm and 15–25 nm, respectively. Higher-resolution TEM image (Fig. 3h) shows clear graphene lattices in the tube wall, indicating a high crystallinity of the produced CNTs. 3.3. Tensile strength of fiber bundles Fig. 4 shows the tensile results of fiber bundles before and after the surface treatment and CNT growth processes. It can be seen that the surface-treated fiber bundles exhibit a small relative decrease in tensile strength compared with pristine CFs, probably attributed to the thermal degradation during high-temperature pyrolysis process. However, the most encouraging result is the large strength retention rate of CFs after growing CNTs. For the surface-treated fiber bundles after 40 min CNT growth, the strength retention rate is 88.5%. This value is much higher than those of fiber bundles (69.1% of strength retention rate) without surface treatment after the CNT growth with low growth temperature and roughly same growth time [20]. The large strength retention rate of CFs can be attributed to the SiO2 protection which may prevent the Fe catalyst diffusing into the carbon substrates and the trace oxygen attack [22]. It is expected that these 3D fabric reinforcements with less damaging to the CFs can result in excellent mechanical improvement in 3D C/C composites. 3.4. Characterization of 3D composites To investigate the effect of vertically aligned CNTs on the microstructure of PyC matrix, images of cross-sections for the base and 3D composites are captured using PLM (Fig. 5a–d) and SEM (Fig. 5e–h). From Fig. 5a, much large PyC grain size, long straight grain boundaries and evident matrix annular cracks can be clearly observed in the base composites. By contrast, the PyC matrix in 3D composites is highly integrated with no annular cracks, as shown in Fig. 5b–d. Another interesting observation from PLM images of 3D composites is that the matrix exhibits sandwich-like structures. The bottom and top layers next to fiber plies are PyC with CNTs (defined as PyC/CNT) having no optical activity due to the densely packed nanotubes [7], and the intermediate layer is pure PyC containing no CNTs. Note that the intermediate PyC has smaller grain size and short curved grain boundaries, largely different from those in base composites. The PyC matrix at its electron microscopy is shown in SEM images of Fig. 5e–h, where PyC next to ply surface exhibits rough containing CNTs and aloof from fibers is smooth without CNTs but with interlocked grain boundaries. These characteristics observed in PLM and SEM images become much more distinct and clearer with the prolonged CNT length. Based on the PLM and SEM observations, corresponding models regarding the microstructure of PyC matrix of the base and 3D composites are established, as shown in Fig. 6. In the case of base composite (Fig. 6a), large PyC grains both in size and thickness are rather susceptible to cracking caused by thermal stress during the fabrication process, leading to the occurrence of matrix annular cracks [23]. Moreover, long straight grain boundaries preferentially aligned roughly either along or parallel to the fiber plies are more likely to form among these large PyC grains. After introducing aligned CNTs, two conclusions can be drawn from the change of matrix microstructure. Firstly, the vertically aligned nanotubes extending into the surrounding matrix as continuous reinforcers significantly stiffen the matrix. With the increase in CNT length, this direct reinforcement on PyC matrix can extend to the area aloof
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Fig. 3. SEM images (a–f) of SiO2-coated CFs grafted by aligned CNTs with different lengths controlled by the growth time: (a, b) 10 min; (c, d) 20 min; (e, f) 40 min. Inset of (f) is an enlarged view of the CNT forests. TEM images (g, h) of the produced CNTs.
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Table 1 Comparison of CNT length (L), volume fractions (V) of CFs and CNTs, density and open porosity of base, 3D-10, 3D-20 and 3D-40 composites. Composites
LCNT (μm)
VCF (%)
VCNT (%)
Density (g/cm3)
Open porosity (%)
Base 3D-10 3D-20 3D-40
– 5.2 13.1 21.8
38.1 34.7 33.8 32.2
– 1.2 3.0 4.2
1.62 1.63 1.60 1.59
12.8 12.6 14.1 14.9
from plies because of the expansion of 3D CNT forests into distant space. Secondly, CNTs also play an indirect reinforcing role in reinforcing the PyC matrix out of the reach of nanotubes (here, it is called as “indirect reinforcement”). Concretely, the top surface of CNT forests provides abundant nano nucleation sites, which can induce the formation of spherical or cone-shaped PyC on the surface of PyC/CNT [24]. With the proceeding of CVI process, these PyC grains contact with each other forming interlocked grain boundaries. Furthermore, the expansion of 3D CNT forests can efficiently reduce the growth space for the single PyC both in thickness and size, resulting in smaller PyC grains with shorter grain boundaries among them. These refined PyC have high cohesion and fracture resistance, which are expected to inhibit the growth and propagation of cracks under loading. Therefore, compared with CNTs reinforced polymer or ceramic composites wherein the CNTs can only modify the matrix within the reach of nanotubes, the great significance of aligned CNTs used in C/C composites is their ability to not only directly reinforce the matrix within the reach of nanotubes, but also to give indirect reinforcement to the matrix aloof from nanotubes. The microstructure models of 3D composites with CNTs of different lengths are clearly described in Fig. 6b–d. Clearly, both the direct and indirect reinforcements on the PyC matrix can be significantly increased by extending the CNT length. 3.5. Mechanical properties of 3D composites Fig. 7 shows the compressive stress–strain curves obtained for all the composite specimens. It is apparent that the out-of-plane compressive strength (OCS, Fig. 7a) and in-plane compressive strength (ICS, Fig. 7b) of the C/C composites are significantly increased after the growth of vertically aligned CNTs on cloths, especially for the composites containing long nanotubes. The curves also display a clear behavior in the specimens' response as seen by the difference in slope: specimens from 3D specimens containing long CNTs show a higher slope than the base specimen and 3D specimens containing short CNTs. This observation suggests that the former composites are stiffer than the latter composites, i.e., higher out-of-plane compressive modulus (OCM) and inplane compressive modulus (ICM). The results listed in Table 2 show that 3D-10, 3D-20 and 3D-40 have improvements in OCS of 33%, 42%
and 63%, respectively. Compared with the base composites, 3D-10, 3D-20 and 3D-40 composites also have 13%, 19% and 24% improvements in OCM, respectively. The measured ICS of the 3D-10, 3D-20 and 3D-40 composites show 152%, 206% and 275% enhancements compared with base composites, respectively. The percentage improvements in the ICM of the 3D-10, 3D-20 and 3D-40 composites compared with those for the base specimens are 90%, 110% and 151%, respectively. For the base composites, upon 90° loading of specimen, massive destructive cracks are formed which propagate along the length and height of specimen and then break the base specimens into small pieces, whereas for the 3D composites the fracture only penetrates the height of specimens forming roughly 45° angle with the thickness direction resulting in fewer and bigger fragments. Further investigation is performed based on SEM observations into the failure specimen to acquire the information regarding reinforcement mechanism. The observation from the low-magnification SEM images of out-of-plane compressive fracture surface shows that compared with the completely disintegrated base specimens (Fig. 8a), the crushed PyC is significantly decreased and even absent in the fragments of 3D specimens (for illustrative purpose, only image for representative specimen of 3D-20 is presented, as shown in Fig. 8c). Examination of the fracture surface at higher magnification provides detailed information regarding the morphology of the fractured PyC. As stated in Fig. 8b, the surface of fractured PyC in base specimens looks flat and exhibits brittle fracture character. Thus, it can be speculated that, when the base specimens are subjected to out-of-plane stress, the matrix annular cracks and straight grain boundaries (labeled by arrows in Fig. 8b) provide the main channels for the destructive crack propagation along the length and height of the composites, respectively. But comparatively, the fractured PyC of 3D composites appears very integrated and rugged (Fig. 8d). These observations indicate that a strongly-enhanced PyC matrix change the mode of failure of the composites to a higher level of reinforcement. Furthermore, sidewalls and ends of CNTs exposed to the fractured PyC of 3D composites can be clearly observed, as shown in Fig. 8e and f, which can be considered as another important factor for the increase of OCS and OCM. On the one hand, the existence of CNT forests sharply increase the interface area between PyC matrix and reinforcements, and thus damage is more difficult to occur in 3D composites by virtue of additional surface energy and friction work because of the newly formed surface area [25]. On the other hand, the fracture of densely-packed CNTs also dissipates great energy, which can further endow the high fracture strength of the 3D composites. Fig. 9 presents the SEM images of failed composite specimens after the in-plane compression tests. When the base specimens are subjected to 0° loading stress, overall shear failure occurs characterized by delamination failures, as displayed in Fig. 9a. Enlarged SEM images reveal that the destructive cracks spread in the PyC matrix (Fig. 9b) or along the
Fig. 4. Tensile stress–strain curves (a) and corresponding strengths (b) of fiber bundles before and after surface treatment (STCF) and CNT growth processes (STCF-X, X represents the growth time).
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Fig. 5. PLM (a–d) and SEM (e–h) images of the base and 3D composites: (a, e) base; (b, f) 3D-10; (c, g) 3D-20; (d, h) 3D-40.
Fig. 6. Structure models of the base and 3D composites: (a) base; (b) 3D-10; (c) 3D-20; (d) 3D-40.
interface between matrix and outer surface of fiber bundles (Fig. 9c) until they absolutely open the plies. For the spreading of destructive cracks in PyC matrix of base composites, long straight PyC grain boundaries orientated roughly parallel to the fiber plies and the matrix
annular cracks provide the main channels for the destructive crack growth and subsequent long-distance propagation. After grafting CNT forests on the CFs, the delamination failures of 3D composites under in-plane stress are alleviated and even inhibited, especially for the
Fig. 7. Compressive stress–strain curves of the base and 3D composites: (a) out-of-plane compression and (b) in-plane compression.
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Table 2 Compressive and flexural properties of the base, 3D-10, 3D-20 and 3D-40 composites (± values represent standard deviation). Out-of-plane compression
In-plane compression
Flexure
Composites Base 3D-10 3D-20 3D-40
Strength (MPa)
Modulus (GPa)
Strength (MPa)
Modulus (GPa)
Strength (MPa)
Modulus (GPa)
158 ± 8 210 ± 11 225 ± 14 258 ± 16
3.12 ± 0.18 3.54 ± 0.11 3.72 ± 0.16 3.86 ± 0.21
48 ± 10 121 ± 9 147 ± 11 180 ± 18
5.12 ± 0.65 9.72 ± 0.54 10.75 ± 0.87 12.86 ± 0.81
135 ± 9 143 ± 13 148 ± 11 153 ± 9
22.1 ± 1.4 24.6 ± 1.2 25.4 ± 1.1 26.6 ± 1.6
composites containing long CNTs. As illustrated in Fig. 9d selected from 3D-20 specimens, where the failure initiates from the top edge of the composite and develops roughly along the diagonal of test specimen. Enlarged SEM images demonstrate that the stiffer matrix impedes the extending of cracks which want to develop along straight path
(Fig. 9e), and the resultant crack deflection leads to the fracture of fiber bundles (Fig. 9f). Great energy can be consumed during crack deflection thus resulting in the remarkable enhancement in in-plane compressive performance [26]. It is difficult to find the CNT ends and sidewalls that are observed in the out-of-plane compressive fracture
Fig. 8. Different-magnification SEM images of the out-of-plane compressive fracture surfaces of the base (a and b) and 3D-20 (c–f) composites.
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Fig. 9. SEM images of the representative failure specimens from the base (a–c) and 3D-20 (d–f) composites after the in-plane compression test (red dotted arrows represent the extending paths of the destructive cracks). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
surface of the 3D composites. Actually, compared with the stronglystiffened PyC/CNT, the destructive cracks are more likely to spread in intermediate PyC layer. For the 3D-10 composites containing short CNTs, the absence of annular cracks in intermediate PyC layer is the main reason for the improvements of ICS and ICM. With the extending of CNT forests in 3D composites, the intermediate PyC is significantly decreased in width and meanwhile has more convoluted and interlocking grain boundaries, as a result, the growth and propagation of destructive cracks in the PyC matrix of long CNT-reinforced 3D composites becomes increasingly difficult such that crack deflection occurs, resulting in the higher ICS and ICM. However, the stiffer carbon matrix may not be the only contributor to ICS and ICM enhancements, since they can also be attributed to the increased lateral support for the load-bearing plies provided by vertically aligned CNTs that may prevent the CFs from microbuckling, a critical composite failure mode associated with fibers under in-plane compressive stress [27]. The percentage enhancements in ICS and ICM obtained for 3D-40 composites are much higher than those obtained by grafting curved and radially CNTs on CFs [10,16,20]. Therefore, from the above analysis, one can comprehend that highdensity CNT forests with long extending length which not only significantly stiffen the surrounding matrix at greater depth (i.e., direct reinforcement), but also influence the peripheral matrix at higher level (i.e., indirect reinforcement), and thus largely increase the compressive properties.
Fig. 10. Flexural stress–strain curves of the base and 3D composites.
More surprisingly, the dramatic enhancement of the matrixdominated mechanical performance does not compromise the inplane properties. The flexural stress–strain curves for the test specimens are presented in Fig. 10, in which all the 3D composites show high failure flexural strength (FS) and modulus (FM) compared with the base specimens. From Table 2, the FS of 3D-10, 3D-20 and 3D-40 are improved by 6%, 10% and 13%, although the fiber volume fractions are decreased by 8.9%, 11.3% and 15.5% (Table 1), respectively. Compared with the base composites, 3D-10, 3D-20 and 3D-40 also have 11%, 15%, and 20% enhancements in FM. Another observation in Fig. 10 is that all the composite specimens show typical brittle fracture, evidenced by the sudden failure as the stress goes up to the peak of curves, which suggests that, the introduction of CNT forests increases the mechanical properties of C/C composites but not significantly changes their fracture behavior. Besides, it needs to be emphasized that the little damage to the CFs attributed to the SiO 2 protection is an important factor for the improvement of FS and FM, since the direct growth of CNTs on the CF surface leads to a serious decrease in flexural properties of C/C composites due to the excessive strength loss of CFs by iron etching and thermal degradation reported in our previous work [20].
4. Conclusions The 3D composites introduced here is the first real instance where the CNTs are used in high density, vertically aligned orientation and controllable length to effectively improve the mechanical performance of C/C composite. Vertically aligned CNTs grown on the surface-treated carbon cloths not only directly stiffen the matrix within the CNT reach, but also give indirect reinforcement to the matrix out of the CNT reach by inducing the formation of small PyC gains that interlock with each other. Both the direct and indirect reinforcements on the PyC matrix can be significantly increased by extending the CNT length. Our work demonstrates that introducing vertically aligned CNTs in the thickness direction of woven-fabric C/C composites is an effective strategy to largely increase longitudinal and transverse compressive properties without sacrificing the in-plane performance. To further understand this new type composite and to enlarge its application, further research should be done on the multifunctional properties, such as thermal conductivity, thermal diffusivity and thermal expansion behaviors.
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Acknowledgments This work has been supported by the Key Grant Project of the Chinese Ministry of Education (Grant No. 313047), the FP7-International Research Staff Exchange Scheme—Advanced Biomaterials for Regenerative Medicine, the Global Innovation Initiative—Nanostructured Materials for the Control of Contaminants Detrimental to Health, the Fundamental Research Funds for the Central universities under Grant No. 3102014JCQ01030 and “111” Project of China (B08040), the Research Fund for the Doctoral Program of Higher Education of China (Grant No. 20126102110013) and the Natural Science Foundation of China (Grant Nos. 51432008, 51202194, 51272212 and 51272213). References [1] E. Fitzer, L.M. Manocha, Carbon Reinforcements and Carbon/Carbon Composites, Christiane, Berlin, 1998. [2] G. Savage, Carbon/Carbon Composites, Springer, London, 1993. [3] H. Qian, E.S. Greenhalgh, M.S.P. Shaffer, A. Bismarck, Carbon nanotube-based hierarchical composites: a review, J. Mater. Chem. 20 (2010) 4751–4762. [4] Q.M. Gong, Z. Li, X.W. Zhou, J.J. Wu, Y. Wang, J. Liang, Synthesis and characterization of in situ grown carbon nanofiber/nanotube reinforced carbon/carbon composites, Carbon 43 (2005) 2426–2429. [5] Q.M. Gong, Z. Li, X.D. Bai, D. Li, J. Liang, The effect of carbon nanotubes on the microstructure and morphology of pyrolytic carbon matrices of C–C composites obtained by CVI, Compos. Sci. Technol. 65 (2005) 1112–1119. [6] J.S. Li, R.Y. Luo, Study of the mechanical properties of carbon nanofiber reinforced carbon/carbon composites, Compos. A Appl. Sci. 39 (2008) 1700–1704. [7] P. Xiao, X.F. Lu, Y.Q. Liu, L.L. He, Effect of in situ grown carbon nanotubes on the structure and mechanical properties of unidirectional carbon/carbon composites, Mater. Sci. Eng. A 528 (2011) 3056–3061. [8] H.L. Li, H.J. Li, J.H. Lu, K.Z. Li, C. Sun, D.S. Zhang, Mechanical properties enhancement of carbon/carbon composites by in situ grown carbon nanofibers, Mater. Sci. Eng. A 547 (2012) 138–141. [9] K.Z. Li, L. Li, H.J. Li, Q. Song, J.H. Lu, Q.G. Fu, Electrophoretic deposition of carbon nanotubes onto carbon fiber felt for production of carbon/carbon composites with improved mechanical and thermal properties, Vacuum 104 (2014) 105–110. [10] Q. Song, K.Z. Li, H.L. Li, H.J. Li, C. Ren, Grafting straight carbon nanotube radially onto carbon fibers and their effect on the mechanical properties of carbon/carbon composites, Carbon 50 (2012) 3943–3960. [11] L. Ci, J. Suhr, V. Pushparaj, X. Zhang, P.M. Ajayan, Continuous carbon nanotube reinforced composites, Nano Lett. 8 (2008) 2762–2766. [12] K.G. Dassios, S. Musso, C. Galiotis, Compressive behavior of MWCNT/epoxy composite mats, Compos. Sci. Technol. 72 (2012) 1027–1033.
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