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mechanical properties of short carbon fiber-reinforced natural graphite flake composites with mesophase pitch as the binder
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Microstructure and thermal/mechanical properties of short carbon fiber-reinforced natural graphite flake composites with mesophase pitch as the binder Yun Zhao a,b, Zhanjun Liu a,*, Huiqi Wang Zechao Tao a, Quangui Guo a, Lang Liu a a b c
a,b
, Jingli Shi
a,c,* ,
Jincai Zhang a,
Key Laboratory of Carbon Materials, Institute of Coal Chemistry, Chinese Academy of Sciences, Taiyuan 030001, China Graduate University of Chinese Academy of Sciences, Beijing 100049, China School of Materials Science and Engineering, Tianjin Polytechnic University, Tianjin 300387, China
A R T I C L E I N F O
A B S T R A C T
Article history:
Short carbon fiber reinforced graphite blocks (SFGs) were fabricated from a mixture of
Received 1 April 2012
mesophase pitch, natural graphite flakes and short carbon fibers by hot-pressing at
Accepted 7 November 2012
2773 K. The effect of fiber content on the structure and thermal/mechanical properties
Available online 15 November 2012
of the SFGs was investigated. It was found that introducing the fibers lowered the densification, and also changed the pore structure and pore size distribution. Compared with the pristine block, all the SFGs earned improved in-plane thermal conductivity and mechanical strength. The formation of a heat flow network and the increase of crystalline sizes made a synergistic effect on the promotion of in-plane thermal conductivity. In-plane thermal conductivity reached the maximum when the fiber content was 6 wt.%. The increase of mechanical strength was mainly attributed to the pull-out of fibers from the matrix. The bend and compressive strength in the direction perpendicular to graphite layers reached the maximum values of 39.6 MPa and 65.5 MPa for fiber content of 8 wt.%, respectively. 2012 Elsevier Ltd. All rights reserved.
1.
Introduction
Owing to low thermal expansion and density, high thermal conductivity, resistance to corrosive environment as well as excellent mechanical properties at elevated temperature, carbon/graphite materials have been used for engineering materials at high temperature [1], thermal management devices such as heat exchangers and heat sinks [2,3] and cathode blocks in aluminum electrolysis cell [4]. As well-known, the ideal graphite crystal has an extremely high in-plane thermal conductivity and the value could be as high as 4180 Wm1 K1 according to molecular dynamics simulations [5]. However,
the thermal conductivity of common poly-crystalline graphite is as low as 70–150 Wm1 K1 at room temperature. It has therefore much room to develop graphite blocks with higher thermal conductivity. Many attempts have been made to improve the thermal conductivity and mechanical properties of graphite materials. Typically, pitch-based carbon fibers/carbon composites can obtain high in-plane thermal conductivity and mechanical strength [6]. However, the composites are expensive and the manufacturing process is complex, which limited their wide application. Some studies indicated that doping catalytic elements (Si, Ti, Zr, etc.) into the carbon matrix could
* Corresponding authors. Address: Key Laboratory of Carbon Materials, Institute of Coal Chemistry, Chinese Academy of Sciences, Taiyuan 030001, China. Fax: +86 351 4083952 (Z. Liu). E-mail address: [email protected] (Z. Liu). 0008-6223/$ - see front matter 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.carbon.2012.11.013
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provide an effective pathway to improve the thermal conductivity of graphite materials [7–9]. It was suggested that the effective increase of thermal conductivity was attributed to the large amounts of these catalytic elements, and the elements were difficult to achieve a uniform distribution in the carbon matrix due to the large difference of density between the additives and the carbon matrix. In addition, precursor carbon also plays a key role in the thermal conductivity of the resultant graphite [10]. Commonly, the precursor carbon to synthetic graphite was composed of the filler of coke and the binder of pitch. Compared with other fillers (e.g., carbon black, petroleum coke and needle coke), natural graphite (NG) flake was perfect precursor for fabricating graphite materials with excellent thermal properties because of their high degree of graphitization and preferential crystalline orientation. Most recently, Yuan et al. [11] prepared a graphite block with in-plane thermal conductivity of 522 Wm1 K1 and bend strength of 7.7 MPa from 86 wt.% NGs and 14 wt.% mesophase pitch by a low temperature hot-pressing process. Also, Liu et al. [12] prepared a graphite block from NGs and mesophase pitch by a high temperature hot-pressing method. The in-plane thermal conductivity was up to 704 Wm1 K1, and the bend strength was 21.1 MPa. However, these results indicated that the bend strength of the NG-derived blocks was poor, which restricted their applications in the thermal management devices where strength was a consideration. There was some work [13–15] suggesting that the graphite blocks cooperated with short carbon fibers could achieve the significant improvement of mechanical properties. The moderate bonding between fiber and matrix can induce the cracks deflecting and propagating along the fiber/matrix interfaces instead of along the matrix directly [15]. In the present work, short fiber reinforced graphite blocks (SFGs) with different fiber content were prepared by a facile mixing of NGs, mesophase pitch and short mesophase pitch-based carbon fibers and the following hot-pressing process at 2773 K. The purpose was to evaluate the potential of short fibers to improve the mechanical properties while maintaining or even increasing the inplane thermal conductivity of the graphite matrix. In addition, this work was directed at investigating the correlation between fiber content and microstructure as well as thermal and mechanical strength of the as-obtained graphite blocks.
2.
Experimental
2.1.
Raw materials
The raw materials consisted of fillers of NGs, binder of mesophase pitch, and reinforcement of short mesophase pitchbased carbon fibers. NGs were supplied by Shandong Graphite Co. Ltd., China, and the mean particle size and purity was 270 lm and 99.8%, respectively. The Mitsubishi naphthalene-based mesophase pitch was used as the binder, and its properties were listed in [12]. Short mesophase pitch-based carbon fibers were offered by our lab, which were prepared through mesophase pitch spinning with a round nozzle, oxidation stabilization at 533 K and carbonization at 1173 K in nitrogen atmosphere [16].
2.2.
Mixing procedure of the raw materials
To minimize the damage and breakage of short fibers during the mixing procedure, a disperse system was introduced that methylcellulose was the dispersant and distilled water was the solvent. The dispersant of 0.7 wt.% methylcellulose was firstly dissolved into 99.3 wt.% distilled water, and then the short fibers were added into the blending solvent. The average length of the fibers used was in the range 3–4 mm. After an adequate stirring dispersion of the fibers, the mesophase pitch powder (150 lm) and NGs were added into the dispersal system in sequence at a proper stirring speed. With continuous stirring, the mixing system came into being homogeneous slurry gradually. Then the mixing slurry was desiccated at 393 K for 3 h to remove the solvent thoroughly. Each proportion of fibers, mesophase pitch powder and NGs was involved in such process to produce the corresponding mixture. In addition, a graphite block without fiber addition named pristine graphite was also prepared for a convenient comparison.
2.3.
Preparation of the graphite blocks
The dried mixture was firstly compacted in a graphite mould, and then pressed uniaxially with a hot-pressing device to obtain the resultant composites, which were cylindrical blocks of A55 mm · 15 mm. One-step hot-pressing method was used to prepare the graphite blocks. The hotpressing device was highly efficient and it integrated the molding, carbonization and graphitization into one process. The hot-pressing temperature was raised at a heating rate of 300 K/h to 2773 K, and then dwelled for 30 min at the ultimate temperature. An infrared radiation thermometer was used to determine the temperature over 1273 K. The hot-pressing pressure was provided by an oil pressure pump, and the ultimate pressure was 25 MPa for all samples. In this study, the fiber content was respectively set at 2, 4, 6, 8, and 10 wt.% in the raw materials. The composites were accordingly named as SFG2, SFG4, SFG6, SFG8 and SFG10.
2.4.
Properties and microstructure characterization
The specimens for test were mechanically cut from the graphite blocks, polished and ultrasonically washed in ethyl alcohol for 15 min. The density of the samples was calculated by measuring the weight and dimension. The bend strength (40 mm · 10 mm · 10 mm) and compressive strength (10 mm · 10 mm · 10 mm) of samples were tested by INSTRON-5500R Electronic Universal Materials Testing System. The thermal diffusivities of all the specimens (10 mm · 10 mm · 3 mm) were determined at room temperature using a laser-flash diffusivity instrument (Nano-Flash-Apparatus, LFA 447, NETZSCH). The thermal conductivity was then calculated by the following equation: k = q · Cp · a, where k was the thermal conductivity, q was density, Cp was the specific heat capacity, and a is the thermal diffusivity of the sample. Three specimens were used for each test. X-ray diffractometer (Rigaku-D/max-cB, Cu Ka, k = 0.15406 nm, 40 kV, 100 mA) was used to obtain the X-ray
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diffraction (XRD) patterns and determine the primary structure. The crystalline parameters (d002, La and Lc) of the composites were calculated by the formulas in Ref. [17], where d002 was the graphite inter-layer spacing, La was the average crystalline diameter and Lc was the average crystalline thickness. Open porosity and pore size distribution of the samples were analyzed by a mercury porosimeter instrument (AutoPore IV 9500). Fractured micrographs of the composites were observed through JEOL JSM-6700E scanning electron microscopy (SEM). Transmission electron microscope (TEM) specimen was carefully prepared by the conventional process of slicing, grinding and ion-milling. Then the prepared foil was examined on a Technai F30 field-emission TEM at 300 kV.
3.
Results and discussion
3.1. SFGs
Effect of fiber content on the microstructure of the
3.1.1.
Morphology and pore structure
Fig. 1 displayed the typical images of the SFGs and the pristine graphite. It can be seen that both the short fibers and NGs oriented in the direction perpendicular to hot-pressing. The fibers were distributed along the graphite layers of the NGs randomly. The 2D structure was developed during the hot-pressing process due to the large aspect ratio of the fibers and NGs. Further, the addition of the short fibers increased the anisotropy of the matrix, and the reinforced type was regarded as 2D reinforcement. In the densification
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process, NGs were firstly wrapped by the binder of pitch, and then oriented and arranged layer by layer under the hot-pressure, as shown in Fig. 1a. Introducing the short fibers into system disturbed the arrangement of graphite layers, and also hindered the adhesion between the NGs and matrix. In addition, the shrinkage of the matrix during heat treatment was restrained. It resulted in the change of pore structure of the matrix carbon, and brought about new cavities and defects around the fibers, as shown in Fig. 1b. It was found that a homogeneous distribution of fibers was observed with fiber content of 6 wt.%, which had a positive effect on developing a 2D structure network, as shown in Fig. 1c. However, the excessive addition resulted in the agglomeration of the fibers. As shown in Fig. 1d, the agglomeration hindered the cohesion of the adjacent NGs greatly, and induced the generation of defects and the disorder of graphite layers. Pore structure greatly affects the thermal/mechanical properties of graphite materials [12]. Fig. 2 showed the pore size distribution of the blocks. It could be seen that the addition of short fibers disturbed the pristine pore structure. The pore size distribution got broader and the average pore size got larger gradually with increasing fiber content. However, when the fiber content increased to 6 wt.%, the average pore size began to decrease and the pore size distribution got narrower. Compared to the block with 6 wt.% fiber content, the block with 8 wt.% fiber content had more uniform pore size distribution, even though the average pore size got larger. As the fiber content increased to 10 wt.%, the pore structure
Fig. 1 – SEM images of the composites (a) pristine graphite, (b) SFG2, (c) SFG6 and (d) SFG10 (perpendicular to the hot-pressing direction).
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the binder content between 6 wt.% and 8 wt.% was most appropriate, our dates indicated that it was probably true.
3.1.2. Crystalline parameters and stress-induced graphitization
Fig. 2 – Pore size distribution of the pristine graphite and SFGs.
of the matrix was destroyed greatly, and the pore diameter exhibited disordered distribution. It was well known that, the relative content of the binder is a key factor in the pore structure of graphite materials [12]. In this work, the binder content in each starting mixture was set at 25 wt.% despite the difference of the fiber content. Since the short fibers had a higher specific surface area than NGs, the increase of the proportion between fibers and NGs reduced the relative content of the binder. Assume that binder content between 6 wt.% and 8 wt.% was just appropriate to achieve a compact block, it could be concluded that the binder content more than 8 wt.% or less than 6 wt.% was insufficient or excessive, respectively, which could lead to non-compact structure and excessive pores. Although we had no direct evidence that
The crystalline parameters of the SFGs as a function of fiber content were shown in Fig. 3, where significant influence of fiber content was demonstrated. In order to have a convenient comparison, pristine graphite without fiber addition was also listed. Compared with the pristine graphite, the d002 of SFGs increased in all cases. Similar to the results in Ref. [18], an increasing tendency of the d002 was observed with fiber fraction. This could be attributed to the following two reasons: first was that the fibers had larger d002 compared with NG (shown in Table S1), and the second was that the fiber addition hindered the effective combination between the NGs and the matrix. As far as the crystalline sizes was concerned, an increasing tendency was observed with fiber fraction increasing from 0 to 6 wt.%, even though the crystalline sizes of the fibers were similar (shown in Table S1) to that of the pristine graphite. The La and Lc reached 147 and 55 nm at 6 wt.% fiber content, respectively. The increase of the crystalline sizes could be explained by the stress-induced graphitization which occurred at the interfaces between fibers and matrix [17–22]. The build-up of the stress originated from the large volume shrinkage of matrix precursor while short fibers had little changes of volume during heat-treatment process [18]. In our system, since the short fibers were wrapped by binder pitch, the stress arose during carbonization and then promoted the graphitization and increased crystalline dimension of the pitch carbon around the fibers. Stress graphitization could improve the compatibility of the interfaces between short fibers and pitch carbon, and it was beneficial for the increase of thermal/mechanical properties. It should be noted that the crystalline sizes began to decrease when the fiber content was more than 6 wt.%. The weakening of stress-induced graphitization might be an important cause. The degree of the stress-induced graphitization mainly depends on the combination between the fibers and the matrix [22]. As mentioned above, with the increasing fiber content, the condition of the interface between fiber/matrix varied and the combination during carbonization was different from each other. It seemed that the fiber content lower
Fig. 3 – Changes in crystalline parameters (d002, La and Lc) of the composites as a function of fiber content.
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Fig. 4 – TEM images of the cross-section of SFG6. (a) Bright-field images and SAED patterns of carbon fiber (CF), interface and NG, respectively. (b) High-magnification image of the interface between the fiber and pitch carbon.
than 6 wt.% was relatively proper for superior combination, while that higher than 6 wt.% began to induce a weaker combination due to the reduction of relative content of the binder. As a result, the crystalline sizes began to decrease with the increasing fiber content higher than 6 wt.%. In order to investigate the interfaces between fibers and pitch carbon, TEM observation for SFG6 was investigated, as shown in Fig 4. Fig. 4a showed the cross-section image of SFG6 in the direction perpendicular to graphite layers. Interfaces with compact structure were observed (indicated with white arrows), and strong bonding between fiber and interface without any defects was also revealed. The selected-area electron diffraction (SAED) patterns were obtained from the marked regions on the cross-section related to the carbon fiber, interface and NG, respectively. The pattern for fiber presented narrow and clear diffraction rings, which could be indexed as polycrystalline graphite, and no broadening for (0 0 2) ring was observed. It meant that the fibers in SFG6 had high crystallinity and well-orientation of the basal planes [23]. It also can be seen from the pattern obtained at the interface that some strips of sharp dots aligned uniformly in a same direction. This revealed that the interface had a highly-orientated layered structure [24], which was better than that of the matrix away from the fiber (shown in Fig. S1) but inferior to that of NG (the white circle 3 of SAED pattern) in SFG6. Fig. 4b showed the high-magnification micrograph of the interface between the fiber and matrix. Well-ordered
graphene layers and high degree of graphitization were observed, which demonstrated that the stress-induced graphitization has occurred. An induced graphitization region was developed around the surface of carbon fiber, where the growth of the graphitic crystallites was parallel to the fiber axis. In addition, many nanopores came into being among the induced layers due to the volatilization of micromolecules during hot-pressing process.
3.2.
Effects of fiber content on the properties of the SFGs
3.2.1.
Density and thermal conductivity
Table 1 listed the main properties of the SFGs. Compared to the pristine graphite, the density of the SFGs decreased, and the open porosity increased in all cases. The depression of densification after fiber addition was because that the existence of fibers hindered the cohesion of the graphite layers, which led to many cavities in SFGs. Moreover, the lower density of fibers would be taken into consideration for the decreasing densification of the SFGs. The thermal conductivity of the SFGs was also listed in Table 1, where the significant influence of the fiber content was demonstrated. Compared to the pristine graphite, all the SFGs had higher in-plane thermal conductivity. A reducing tendency of in-plane thermal conductivity was observed when the fiber content was less or higher than 6 wt.%. The in-plane thermal conductivity up to 518 Wm1 K1 was obtained for SFG6. It was well known that, for carbon materials,
Table 1 – Basic properties of the blocks. Materials
Density (g/cm3)
Open porosity (%)
Thermal conductivity (Wm1 K1) ||a
Pristine graphite SFG2 SFG4 SFG6 SFG8 SFG10 a b
1.93 1.89 1.87 1.89 1.88 1.85
Parallel to graphite layers. Perpendicular to graphite layers.
11.3 14.5 16.4 14.8 15.2 14.8
328 ± 7 386 ± 10 420 ± 11 518 ± 15 428 ± 13 419 ± 14
?b 35.9 ± 0.8 31.7 ± 0.7 23.0 ± 0.6 22.4 ± 0.7 20.4 ± 0.6 18.0 ± 0.8
Compressive strength (MPa) ?b 25.7 ± 0.8 31.0 ± 1.1 40.1 ± 1.5 46.0 ± 1.6 65.5 ± 2.6 29.3 ± 1.2
||a 15.3 ± 0.5 18.2 ± 0.6 22.1 ± 0.7 24.6 ± 1.0 29.9 ± 1.0 17.8 ± 0.7
Bend strength (MPa) ?b 20.8 ± 0.7 22.4 ± 0.7 28.2 ± 1.1 31.8 ± 1.3 39.6 ± 1.5 25.4 ± 1.0
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Fig. 5 – Development of the 2D heat flow network in SFGs as a function of the fiber content.
the larger the crystalline size, the higher the thermal conductivity. It should be noted from Table 1 that, the in-plane thermal conductivity of SFGs as a function of fiber content was just the same tendency as the crystalline sizes. This verified that there were some interrelations between thermal conductivity and crystalline sizes of the SFGs. Apart from the improving crystalline sizes, the formation of the heat flow network would be in consideration for the increase of thermal conductivity of SFGs. As far as the pristine graphite was concerned, the inferior conductive property mainly resulted from the pitch-derived graphite, which had a poor degree of graphitization, random crystalline orientation and defects. These hindered the transfer of phonons in the pristine graphite, and restrained the improvement of thermal conductivity. The pitch fibers after heat treatment of 2773 K had a higher orientation of crystalline grains, fewer defects and higher thermal conductivity [25], which played an important role in the improvement of thermal conductivity of the SFGs. Considering the 2D structure of the SFGs discussed above, a model for the heat conduction was proposed to explain the increase of thermal conductivity when short fibers were introduced into the SFGs, as shown in Fig. 5. Based on the high orientation along graphite layers and large aspect ratio, the fibers added in SFGs connected the adjacent NGs, and enlarged the heat flow region. Many heat flow regions were merged into larger regions with the increase of fiber content, and then developed into a heat flow network, which led to the prominent improvement of in-plane thermal conductivity. The mechanism for the enhancement of in-plane thermal conductivity for SFGs seemed obscure. The synergistic effect of multi-factors might affect the thermal conductivity. According to the above analysis, the following four factors would be contributed to the improved thermal conductivity. First, the short fibers themselves had a high thermal conductivity along the fiber axis. The thermal conductivity of the fibers used in this work was about 580 Wm1 K1 after graphitization at 2773 K, which measured by a four probe method like that in Ref. [16] and the equation between thermal conductivity (k) and electrical resistance (q): k = 440,000/(q + 258)295 [26]. Second, Larger crystalline sizes led to higher thermal conductivity of the blocks. The introduction of the fibers increased the crystalline sizes, which depended on the stress-induced
graphitization. Third, the formation of the 2D heat flow network accelerated the in-plane transfer of phonons and then improved the thermal transfer ability. Finally, the pore structure would be taken into account. As was stated above, more homogeneous pore size distribution and smaller pore size in SFG6 contributed to the higher thermal conductivity. It could also be seen from Table 1 that, the thermal conductivity perpendicular to graphite layers was much lower than the corresponding values in the parallel direction. A decreasing tendency of thermal conductivity was observed as fiber content increased. The anisotropic degree of thermal conductivity was prominently increased with fiber fraction. Anisotropy existed commonly in graphite material due to its special hexagonal net structure. The weak Van der Waals forces among graphite layers resulted in poor thermal transfer properties. As far as the SFGs were concerned, addition of short fibers could increase the interfacial heat resistance in the direction perpendicular to the graphite layers. It might be due to the low thermal conductivity in the radial direction of the fibers, which was deduced from the sheet-like structure of the fibers with high orientation along the fiber axis (shown in Fig. S2). The increase of thermal anisotropy is beneficial to reduce the temperature gradient in the basal plane of graphite and control the conduction of heat flow when suffering heat sink [27].
3.2.2.
Mechanical properties
The bend and compressive strength of the blocks were also listed in Table 1. The bend strength perpendicular to graphite layers increased with fiber content from 0 to 8 wt.%, and the value reached 39.6 MPa for SFG8, increased by 90.6% compared with the pristine graphite. However, an abrupt decrease was observed for SFG10. As for the compressive strength in the direction perpendicular to graphite layers, similar tendency was exhibited and the maximum value, 65.5 MPa, increased by 155% than that of the pristine graphite, was obtained also at the 8 wt.% fiber content. The mechanical properties exhibited an increasing anisotropy, but it was not obvious than that of thermal conductivity, and this could due to the anisotropic structure of the SFGs. From the values of compressive strength in Table 1, it could be seen that both the ones perpendicular and parallel to graphite layers had the same tendency of
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Fig. 6 – (a) SEM images of the interface between fiber and matrix in SFG8 after mechanical tests; (b) Magnified image of the white rectangle region in (a).
variation as a function of fiber fraction. It meant that the mechanical properties in both directions might be dominated by the similar mechanism of reinforcement. The mechanical strength of graphite was closely related to their microstructure. It was well known that more porosity and defects in the materials easily led to low mechanical strength. In addition, the interface characteristics between fibers and matrix also played key roles in the mechanical properties of SFGs [28,29]. Fractured micrographs of SFG8 after the bend strength test were investigated. Fig. 6 showed the interface characteristics between the fibers and the matrix in SFG8. It can be seen from Fig. 6a that the obvious pull-out of fibers was observed, which indicated that the fiber and the matrix had a weak combination. Fig. 6b revealed the weak fiber–matrix bonding, where the crack branching occurred across the interfaces. It could be deduced that the bonding energy between fibers and matrix was smaller than both the energy at the tip of defects and fracture energy of the fiber. Therefore, the composite failure would be mainly controlled by the crack propagation linked with the fiber–matrix interfaces [15]. The micro-cracks propagated forward parallel to the load direction, which initiated at the tip of defects or pores in the matrix. When met the fiber–matrix interfaces, the cracks stopped at the fiber surfaces and were eliminated by absorbing fracture energy or continue propagated along the fiber–matrix interfaces. The fiber pull-out effect would play a dominant role in the composite fracture mechanism. The higher the fiber content, the more the interfaces, and the more energy was consumed for the further propagation of the cracks. What should be noted is that, unlike the clean fiber/ polymer or fiber/cement surface after failure [30,31], some matrix graphite with high orientation along the fiber axis adhered to the surface of the fiber (shown in Fig. 6b). This indicated that some chemical bonding between the fibers and the matrix were developed during the hot-pressing process. Furthermore, it could be concluded that the combination between the fiber and the matrix was stronger than that of the matrix itself. The coarse surface showed that more energy could be consumed than the clean one when the pull-out occurred. Therefore, the fiber addition in SFGs resulted in the improved mechanical strength. The special characteristic of the pull-out surface also reflected the poor mechanical properties of the matrix, and the introduction of the fibers could strengthen parts of the matrix which was around the fibers. The high orientation of
matrix along the fiber axis showed the well-compatibility between the fibers and the binder. The behavior of orientation might result from the stress-induced graphitization, since the observation of the SEM in Fig. 6b was similar to the results of the TEM (shown in Fig. 4b). Considering that the binder pitch was increasingly insufficient for wrapping fibers with fiber fraction, the excessive addition of short fibers would result in inferior interface combination. Furthermore, excessive fiber addition not only deteriorated the pore structure, but also caused the aggregation of fibers, which might lead to the abrupt decrease of the mechanical strength of SFG10.
4.
Conclusions
The goal of the present work was to improve the thermal conductivity and mechanical strength of graphite blocks. The effect of the short mesophase pitch carbon-fiber addition on the microstructure and thermal/mechanical properties of the NG-derived graphite blocks has been investigated. In the results discussed above, some important insights have been gained in this system: (1) Compared to the pristine graphite, the d002 of the SFGs increased with the increasing fiber content. The crystalline sizes were improved, which could due to the stress-induced graphitization. For the fiber content of 6 wt.%, the La and Lc reached 147 and 55 nm, respectively. (2) The fiber content played a key role in the thermal conductivity of SFGs. A tendency of reduced in-plane thermal conductivity of the graphite blocks was observed when the fiber content was less or higher than 6 wt.%. At a fiber content of 6 wt.%, the in-plane thermal conductivity achieved the maximum value of 518 Wm1 K1. The improvement of thermal conductivity relied on the increase of crystalline sizes and the formation of the 2D heat flow network. (3) Compared to the pristine graphite, all the SFGs had higher compressive and bend strength. The mechanical strength of the SFGs were reduced when the fiber content was less or higher than 8 wt.%. At a fiber content of 8 wt.%, the bend and compressive strength reached 39.6 and 65.5 MPa, respectively. The increase of the mechanical strength might be due to the pull-out effect of fibers. (4) The fiber addition increased the anisotropy of matrix both in microstructure and thermal/mechanical properties due to the orientation behavior of the fibers after hotpressing. Considering the poor mechanical properties of
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NG-derived graphite, the introduction of short fibers with a low fraction in this system was demonstrated to be a simple and effective approach to improve the mechanical properties while increasing the in-plane thermal conductivity prominently.
Acknowledgments The work was financially supported by the National Basic Research Program (973 Program, No. 2011CB605802) and National Natural Science Foundation (50902136) of China. The authors are grateful to Prof. Jinren Song for his helpful suggestion and encouragement.
Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.carbon. 2012.11.013.
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5 3 ( 2 0 1 3 ) 3 1 3 –3 2 0
Available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/carbon
Microstructure and thermal/mechanical properties of short carbon fiber-reinforced natural graphite flake composites with mesophase pitch as the binder Yun Zhao a,b, Zhanjun Liu a,*, Huiqi Wang Zechao Tao a, Quangui Guo a, Lang Liu a a b c
a,b
, Jingli Shi
a,c,* ,
Jincai Zhang a,
Key Laboratory of Carbon Materials, Institute of Coal Chemistry, Chinese Academy of Sciences, Taiyuan 030001, China Graduate University of Chinese Academy of Sciences, Beijing 100049, China School of Materials Science and Engineering, Tianjin Polytechnic University, Tianjin 300387, China
A R T I C L E I N F O
A B S T R A C T
Article history:
Short carbon fiber reinforced graphite blocks (SFGs) were fabricated from a mixture of
Received 1 April 2012
mesophase pitch, natural graphite flakes and short carbon fibers by hot-pressing at
Accepted 7 November 2012
2773 K. The effect of fiber content on the structure and thermal/mechanical properties
Available online 15 November 2012
of the SFGs was investigated. It was found that introducing the fibers lowered the densification, and also changed the pore structure and pore size distribution. Compared with the pristine block, all the SFGs earned improved in-plane thermal conductivity and mechanical strength. The formation of a heat flow network and the increase of crystalline sizes made a synergistic effect on the promotion of in-plane thermal conductivity. In-plane thermal conductivity reached the maximum when the fiber content was 6 wt.%. The increase of mechanical strength was mainly attributed to the pull-out of fibers from the matrix. The bend and compressive strength in the direction perpendicular to graphite layers reached the maximum values of 39.6 MPa and 65.5 MPa for fiber content of 8 wt.%, respectively. 2012 Elsevier Ltd. All rights reserved.
1.
Introduction
Owing to low thermal expansion and density, high thermal conductivity, resistance to corrosive environment as well as excellent mechanical properties at elevated temperature, carbon/graphite materials have been used for engineering materials at high temperature [1], thermal management devices such as heat exchangers and heat sinks [2,3] and cathode blocks in aluminum electrolysis cell [4]. As well-known, the ideal graphite crystal has an extremely high in-plane thermal conductivity and the value could be as high as 4180 Wm1 K1 according to molecular dynamics simulations [5]. However,
the thermal conductivity of common poly-crystalline graphite is as low as 70–150 Wm1 K1 at room temperature. It has therefore much room to develop graphite blocks with higher thermal conductivity. Many attempts have been made to improve the thermal conductivity and mechanical properties of graphite materials. Typically, pitch-based carbon fibers/carbon composites can obtain high in-plane thermal conductivity and mechanical strength [6]. However, the composites are expensive and the manufacturing process is complex, which limited their wide application. Some studies indicated that doping catalytic elements (Si, Ti, Zr, etc.) into the carbon matrix could
* Corresponding authors. Address: Key Laboratory of Carbon Materials, Institute of Coal Chemistry, Chinese Academy of Sciences, Taiyuan 030001, China. Fax: +86 351 4083952 (Z. Liu). E-mail address: [email protected] (Z. Liu). 0008-6223/$ - see front matter 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.carbon.2012.11.013
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provide an effective pathway to improve the thermal conductivity of graphite materials [7–9]. It was suggested that the effective increase of thermal conductivity was attributed to the large amounts of these catalytic elements, and the elements were difficult to achieve a uniform distribution in the carbon matrix due to the large difference of density between the additives and the carbon matrix. In addition, precursor carbon also plays a key role in the thermal conductivity of the resultant graphite [10]. Commonly, the precursor carbon to synthetic graphite was composed of the filler of coke and the binder of pitch. Compared with other fillers (e.g., carbon black, petroleum coke and needle coke), natural graphite (NG) flake was perfect precursor for fabricating graphite materials with excellent thermal properties because of their high degree of graphitization and preferential crystalline orientation. Most recently, Yuan et al. [11] prepared a graphite block with in-plane thermal conductivity of 522 Wm1 K1 and bend strength of 7.7 MPa from 86 wt.% NGs and 14 wt.% mesophase pitch by a low temperature hot-pressing process. Also, Liu et al. [12] prepared a graphite block from NGs and mesophase pitch by a high temperature hot-pressing method. The in-plane thermal conductivity was up to 704 Wm1 K1, and the bend strength was 21.1 MPa. However, these results indicated that the bend strength of the NG-derived blocks was poor, which restricted their applications in the thermal management devices where strength was a consideration. There was some work [13–15] suggesting that the graphite blocks cooperated with short carbon fibers could achieve the significant improvement of mechanical properties. The moderate bonding between fiber and matrix can induce the cracks deflecting and propagating along the fiber/matrix interfaces instead of along the matrix directly [15]. In the present work, short fiber reinforced graphite blocks (SFGs) with different fiber content were prepared by a facile mixing of NGs, mesophase pitch and short mesophase pitch-based carbon fibers and the following hot-pressing process at 2773 K. The purpose was to evaluate the potential of short fibers to improve the mechanical properties while maintaining or even increasing the inplane thermal conductivity of the graphite matrix. In addition, this work was directed at investigating the correlation between fiber content and microstructure as well as thermal and mechanical strength of the as-obtained graphite blocks.
2.
Experimental
2.1.
Raw materials
The raw materials consisted of fillers of NGs, binder of mesophase pitch, and reinforcement of short mesophase pitchbased carbon fibers. NGs were supplied by Shandong Graphite Co. Ltd., China, and the mean particle size and purity was 270 lm and 99.8%, respectively. The Mitsubishi naphthalene-based mesophase pitch was used as the binder, and its properties were listed in [12]. Short mesophase pitch-based carbon fibers were offered by our lab, which were prepared through mesophase pitch spinning with a round nozzle, oxidation stabilization at 533 K and carbonization at 1173 K in nitrogen atmosphere [16].
2.2.
Mixing procedure of the raw materials
To minimize the damage and breakage of short fibers during the mixing procedure, a disperse system was introduced that methylcellulose was the dispersant and distilled water was the solvent. The dispersant of 0.7 wt.% methylcellulose was firstly dissolved into 99.3 wt.% distilled water, and then the short fibers were added into the blending solvent. The average length of the fibers used was in the range 3–4 mm. After an adequate stirring dispersion of the fibers, the mesophase pitch powder (150 lm) and NGs were added into the dispersal system in sequence at a proper stirring speed. With continuous stirring, the mixing system came into being homogeneous slurry gradually. Then the mixing slurry was desiccated at 393 K for 3 h to remove the solvent thoroughly. Each proportion of fibers, mesophase pitch powder and NGs was involved in such process to produce the corresponding mixture. In addition, a graphite block without fiber addition named pristine graphite was also prepared for a convenient comparison.
2.3.
Preparation of the graphite blocks
The dried mixture was firstly compacted in a graphite mould, and then pressed uniaxially with a hot-pressing device to obtain the resultant composites, which were cylindrical blocks of A55 mm · 15 mm. One-step hot-pressing method was used to prepare the graphite blocks. The hotpressing device was highly efficient and it integrated the molding, carbonization and graphitization into one process. The hot-pressing temperature was raised at a heating rate of 300 K/h to 2773 K, and then dwelled for 30 min at the ultimate temperature. An infrared radiation thermometer was used to determine the temperature over 1273 K. The hot-pressing pressure was provided by an oil pressure pump, and the ultimate pressure was 25 MPa for all samples. In this study, the fiber content was respectively set at 2, 4, 6, 8, and 10 wt.% in the raw materials. The composites were accordingly named as SFG2, SFG4, SFG6, SFG8 and SFG10.
2.4.
Properties and microstructure characterization
The specimens for test were mechanically cut from the graphite blocks, polished and ultrasonically washed in ethyl alcohol for 15 min. The density of the samples was calculated by measuring the weight and dimension. The bend strength (40 mm · 10 mm · 10 mm) and compressive strength (10 mm · 10 mm · 10 mm) of samples were tested by INSTRON-5500R Electronic Universal Materials Testing System. The thermal diffusivities of all the specimens (10 mm · 10 mm · 3 mm) were determined at room temperature using a laser-flash diffusivity instrument (Nano-Flash-Apparatus, LFA 447, NETZSCH). The thermal conductivity was then calculated by the following equation: k = q · Cp · a, where k was the thermal conductivity, q was density, Cp was the specific heat capacity, and a is the thermal diffusivity of the sample. Three specimens were used for each test. X-ray diffractometer (Rigaku-D/max-cB, Cu Ka, k = 0.15406 nm, 40 kV, 100 mA) was used to obtain the X-ray
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diffraction (XRD) patterns and determine the primary structure. The crystalline parameters (d002, La and Lc) of the composites were calculated by the formulas in Ref. [17], where d002 was the graphite inter-layer spacing, La was the average crystalline diameter and Lc was the average crystalline thickness. Open porosity and pore size distribution of the samples were analyzed by a mercury porosimeter instrument (AutoPore IV 9500). Fractured micrographs of the composites were observed through JEOL JSM-6700E scanning electron microscopy (SEM). Transmission electron microscope (TEM) specimen was carefully prepared by the conventional process of slicing, grinding and ion-milling. Then the prepared foil was examined on a Technai F30 field-emission TEM at 300 kV.
3.
Results and discussion
3.1. SFGs
Effect of fiber content on the microstructure of the
3.1.1.
Morphology and pore structure
Fig. 1 displayed the typical images of the SFGs and the pristine graphite. It can be seen that both the short fibers and NGs oriented in the direction perpendicular to hot-pressing. The fibers were distributed along the graphite layers of the NGs randomly. The 2D structure was developed during the hot-pressing process due to the large aspect ratio of the fibers and NGs. Further, the addition of the short fibers increased the anisotropy of the matrix, and the reinforced type was regarded as 2D reinforcement. In the densification
315
process, NGs were firstly wrapped by the binder of pitch, and then oriented and arranged layer by layer under the hot-pressure, as shown in Fig. 1a. Introducing the short fibers into system disturbed the arrangement of graphite layers, and also hindered the adhesion between the NGs and matrix. In addition, the shrinkage of the matrix during heat treatment was restrained. It resulted in the change of pore structure of the matrix carbon, and brought about new cavities and defects around the fibers, as shown in Fig. 1b. It was found that a homogeneous distribution of fibers was observed with fiber content of 6 wt.%, which had a positive effect on developing a 2D structure network, as shown in Fig. 1c. However, the excessive addition resulted in the agglomeration of the fibers. As shown in Fig. 1d, the agglomeration hindered the cohesion of the adjacent NGs greatly, and induced the generation of defects and the disorder of graphite layers. Pore structure greatly affects the thermal/mechanical properties of graphite materials [12]. Fig. 2 showed the pore size distribution of the blocks. It could be seen that the addition of short fibers disturbed the pristine pore structure. The pore size distribution got broader and the average pore size got larger gradually with increasing fiber content. However, when the fiber content increased to 6 wt.%, the average pore size began to decrease and the pore size distribution got narrower. Compared to the block with 6 wt.% fiber content, the block with 8 wt.% fiber content had more uniform pore size distribution, even though the average pore size got larger. As the fiber content increased to 10 wt.%, the pore structure
Fig. 1 – SEM images of the composites (a) pristine graphite, (b) SFG2, (c) SFG6 and (d) SFG10 (perpendicular to the hot-pressing direction).
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the binder content between 6 wt.% and 8 wt.% was most appropriate, our dates indicated that it was probably true.
3.1.2. Crystalline parameters and stress-induced graphitization
Fig. 2 – Pore size distribution of the pristine graphite and SFGs.
of the matrix was destroyed greatly, and the pore diameter exhibited disordered distribution. It was well known that, the relative content of the binder is a key factor in the pore structure of graphite materials [12]. In this work, the binder content in each starting mixture was set at 25 wt.% despite the difference of the fiber content. Since the short fibers had a higher specific surface area than NGs, the increase of the proportion between fibers and NGs reduced the relative content of the binder. Assume that binder content between 6 wt.% and 8 wt.% was just appropriate to achieve a compact block, it could be concluded that the binder content more than 8 wt.% or less than 6 wt.% was insufficient or excessive, respectively, which could lead to non-compact structure and excessive pores. Although we had no direct evidence that
The crystalline parameters of the SFGs as a function of fiber content were shown in Fig. 3, where significant influence of fiber content was demonstrated. In order to have a convenient comparison, pristine graphite without fiber addition was also listed. Compared with the pristine graphite, the d002 of SFGs increased in all cases. Similar to the results in Ref. [18], an increasing tendency of the d002 was observed with fiber fraction. This could be attributed to the following two reasons: first was that the fibers had larger d002 compared with NG (shown in Table S1), and the second was that the fiber addition hindered the effective combination between the NGs and the matrix. As far as the crystalline sizes was concerned, an increasing tendency was observed with fiber fraction increasing from 0 to 6 wt.%, even though the crystalline sizes of the fibers were similar (shown in Table S1) to that of the pristine graphite. The La and Lc reached 147 and 55 nm at 6 wt.% fiber content, respectively. The increase of the crystalline sizes could be explained by the stress-induced graphitization which occurred at the interfaces between fibers and matrix [17–22]. The build-up of the stress originated from the large volume shrinkage of matrix precursor while short fibers had little changes of volume during heat-treatment process [18]. In our system, since the short fibers were wrapped by binder pitch, the stress arose during carbonization and then promoted the graphitization and increased crystalline dimension of the pitch carbon around the fibers. Stress graphitization could improve the compatibility of the interfaces between short fibers and pitch carbon, and it was beneficial for the increase of thermal/mechanical properties. It should be noted that the crystalline sizes began to decrease when the fiber content was more than 6 wt.%. The weakening of stress-induced graphitization might be an important cause. The degree of the stress-induced graphitization mainly depends on the combination between the fibers and the matrix [22]. As mentioned above, with the increasing fiber content, the condition of the interface between fiber/matrix varied and the combination during carbonization was different from each other. It seemed that the fiber content lower
Fig. 3 – Changes in crystalline parameters (d002, La and Lc) of the composites as a function of fiber content.
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Fig. 4 – TEM images of the cross-section of SFG6. (a) Bright-field images and SAED patterns of carbon fiber (CF), interface and NG, respectively. (b) High-magnification image of the interface between the fiber and pitch carbon.
than 6 wt.% was relatively proper for superior combination, while that higher than 6 wt.% began to induce a weaker combination due to the reduction of relative content of the binder. As a result, the crystalline sizes began to decrease with the increasing fiber content higher than 6 wt.%. In order to investigate the interfaces between fibers and pitch carbon, TEM observation for SFG6 was investigated, as shown in Fig 4. Fig. 4a showed the cross-section image of SFG6 in the direction perpendicular to graphite layers. Interfaces with compact structure were observed (indicated with white arrows), and strong bonding between fiber and interface without any defects was also revealed. The selected-area electron diffraction (SAED) patterns were obtained from the marked regions on the cross-section related to the carbon fiber, interface and NG, respectively. The pattern for fiber presented narrow and clear diffraction rings, which could be indexed as polycrystalline graphite, and no broadening for (0 0 2) ring was observed. It meant that the fibers in SFG6 had high crystallinity and well-orientation of the basal planes [23]. It also can be seen from the pattern obtained at the interface that some strips of sharp dots aligned uniformly in a same direction. This revealed that the interface had a highly-orientated layered structure [24], which was better than that of the matrix away from the fiber (shown in Fig. S1) but inferior to that of NG (the white circle 3 of SAED pattern) in SFG6. Fig. 4b showed the high-magnification micrograph of the interface between the fiber and matrix. Well-ordered
graphene layers and high degree of graphitization were observed, which demonstrated that the stress-induced graphitization has occurred. An induced graphitization region was developed around the surface of carbon fiber, where the growth of the graphitic crystallites was parallel to the fiber axis. In addition, many nanopores came into being among the induced layers due to the volatilization of micromolecules during hot-pressing process.
3.2.
Effects of fiber content on the properties of the SFGs
3.2.1.
Density and thermal conductivity
Table 1 listed the main properties of the SFGs. Compared to the pristine graphite, the density of the SFGs decreased, and the open porosity increased in all cases. The depression of densification after fiber addition was because that the existence of fibers hindered the cohesion of the graphite layers, which led to many cavities in SFGs. Moreover, the lower density of fibers would be taken into consideration for the decreasing densification of the SFGs. The thermal conductivity of the SFGs was also listed in Table 1, where the significant influence of the fiber content was demonstrated. Compared to the pristine graphite, all the SFGs had higher in-plane thermal conductivity. A reducing tendency of in-plane thermal conductivity was observed when the fiber content was less or higher than 6 wt.%. The in-plane thermal conductivity up to 518 Wm1 K1 was obtained for SFG6. It was well known that, for carbon materials,
Table 1 – Basic properties of the blocks. Materials
Density (g/cm3)
Open porosity (%)
Thermal conductivity (Wm1 K1) ||a
Pristine graphite SFG2 SFG4 SFG6 SFG8 SFG10 a b
1.93 1.89 1.87 1.89 1.88 1.85
Parallel to graphite layers. Perpendicular to graphite layers.
11.3 14.5 16.4 14.8 15.2 14.8
328 ± 7 386 ± 10 420 ± 11 518 ± 15 428 ± 13 419 ± 14
?b 35.9 ± 0.8 31.7 ± 0.7 23.0 ± 0.6 22.4 ± 0.7 20.4 ± 0.6 18.0 ± 0.8
Compressive strength (MPa) ?b 25.7 ± 0.8 31.0 ± 1.1 40.1 ± 1.5 46.0 ± 1.6 65.5 ± 2.6 29.3 ± 1.2
||a 15.3 ± 0.5 18.2 ± 0.6 22.1 ± 0.7 24.6 ± 1.0 29.9 ± 1.0 17.8 ± 0.7
Bend strength (MPa) ?b 20.8 ± 0.7 22.4 ± 0.7 28.2 ± 1.1 31.8 ± 1.3 39.6 ± 1.5 25.4 ± 1.0
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Fig. 5 – Development of the 2D heat flow network in SFGs as a function of the fiber content.
the larger the crystalline size, the higher the thermal conductivity. It should be noted from Table 1 that, the in-plane thermal conductivity of SFGs as a function of fiber content was just the same tendency as the crystalline sizes. This verified that there were some interrelations between thermal conductivity and crystalline sizes of the SFGs. Apart from the improving crystalline sizes, the formation of the heat flow network would be in consideration for the increase of thermal conductivity of SFGs. As far as the pristine graphite was concerned, the inferior conductive property mainly resulted from the pitch-derived graphite, which had a poor degree of graphitization, random crystalline orientation and defects. These hindered the transfer of phonons in the pristine graphite, and restrained the improvement of thermal conductivity. The pitch fibers after heat treatment of 2773 K had a higher orientation of crystalline grains, fewer defects and higher thermal conductivity [25], which played an important role in the improvement of thermal conductivity of the SFGs. Considering the 2D structure of the SFGs discussed above, a model for the heat conduction was proposed to explain the increase of thermal conductivity when short fibers were introduced into the SFGs, as shown in Fig. 5. Based on the high orientation along graphite layers and large aspect ratio, the fibers added in SFGs connected the adjacent NGs, and enlarged the heat flow region. Many heat flow regions were merged into larger regions with the increase of fiber content, and then developed into a heat flow network, which led to the prominent improvement of in-plane thermal conductivity. The mechanism for the enhancement of in-plane thermal conductivity for SFGs seemed obscure. The synergistic effect of multi-factors might affect the thermal conductivity. According to the above analysis, the following four factors would be contributed to the improved thermal conductivity. First, the short fibers themselves had a high thermal conductivity along the fiber axis. The thermal conductivity of the fibers used in this work was about 580 Wm1 K1 after graphitization at 2773 K, which measured by a four probe method like that in Ref. [16] and the equation between thermal conductivity (k) and electrical resistance (q): k = 440,000/(q + 258)295 [26]. Second, Larger crystalline sizes led to higher thermal conductivity of the blocks. The introduction of the fibers increased the crystalline sizes, which depended on the stress-induced
graphitization. Third, the formation of the 2D heat flow network accelerated the in-plane transfer of phonons and then improved the thermal transfer ability. Finally, the pore structure would be taken into account. As was stated above, more homogeneous pore size distribution and smaller pore size in SFG6 contributed to the higher thermal conductivity. It could also be seen from Table 1 that, the thermal conductivity perpendicular to graphite layers was much lower than the corresponding values in the parallel direction. A decreasing tendency of thermal conductivity was observed as fiber content increased. The anisotropic degree of thermal conductivity was prominently increased with fiber fraction. Anisotropy existed commonly in graphite material due to its special hexagonal net structure. The weak Van der Waals forces among graphite layers resulted in poor thermal transfer properties. As far as the SFGs were concerned, addition of short fibers could increase the interfacial heat resistance in the direction perpendicular to the graphite layers. It might be due to the low thermal conductivity in the radial direction of the fibers, which was deduced from the sheet-like structure of the fibers with high orientation along the fiber axis (shown in Fig. S2). The increase of thermal anisotropy is beneficial to reduce the temperature gradient in the basal plane of graphite and control the conduction of heat flow when suffering heat sink [27].
3.2.2.
Mechanical properties
The bend and compressive strength of the blocks were also listed in Table 1. The bend strength perpendicular to graphite layers increased with fiber content from 0 to 8 wt.%, and the value reached 39.6 MPa for SFG8, increased by 90.6% compared with the pristine graphite. However, an abrupt decrease was observed for SFG10. As for the compressive strength in the direction perpendicular to graphite layers, similar tendency was exhibited and the maximum value, 65.5 MPa, increased by 155% than that of the pristine graphite, was obtained also at the 8 wt.% fiber content. The mechanical properties exhibited an increasing anisotropy, but it was not obvious than that of thermal conductivity, and this could due to the anisotropic structure of the SFGs. From the values of compressive strength in Table 1, it could be seen that both the ones perpendicular and parallel to graphite layers had the same tendency of
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Fig. 6 – (a) SEM images of the interface between fiber and matrix in SFG8 after mechanical tests; (b) Magnified image of the white rectangle region in (a).
variation as a function of fiber fraction. It meant that the mechanical properties in both directions might be dominated by the similar mechanism of reinforcement. The mechanical strength of graphite was closely related to their microstructure. It was well known that more porosity and defects in the materials easily led to low mechanical strength. In addition, the interface characteristics between fibers and matrix also played key roles in the mechanical properties of SFGs [28,29]. Fractured micrographs of SFG8 after the bend strength test were investigated. Fig. 6 showed the interface characteristics between the fibers and the matrix in SFG8. It can be seen from Fig. 6a that the obvious pull-out of fibers was observed, which indicated that the fiber and the matrix had a weak combination. Fig. 6b revealed the weak fiber–matrix bonding, where the crack branching occurred across the interfaces. It could be deduced that the bonding energy between fibers and matrix was smaller than both the energy at the tip of defects and fracture energy of the fiber. Therefore, the composite failure would be mainly controlled by the crack propagation linked with the fiber–matrix interfaces [15]. The micro-cracks propagated forward parallel to the load direction, which initiated at the tip of defects or pores in the matrix. When met the fiber–matrix interfaces, the cracks stopped at the fiber surfaces and were eliminated by absorbing fracture energy or continue propagated along the fiber–matrix interfaces. The fiber pull-out effect would play a dominant role in the composite fracture mechanism. The higher the fiber content, the more the interfaces, and the more energy was consumed for the further propagation of the cracks. What should be noted is that, unlike the clean fiber/ polymer or fiber/cement surface after failure [30,31], some matrix graphite with high orientation along the fiber axis adhered to the surface of the fiber (shown in Fig. 6b). This indicated that some chemical bonding between the fibers and the matrix were developed during the hot-pressing process. Furthermore, it could be concluded that the combination between the fiber and the matrix was stronger than that of the matrix itself. The coarse surface showed that more energy could be consumed than the clean one when the pull-out occurred. Therefore, the fiber addition in SFGs resulted in the improved mechanical strength. The special characteristic of the pull-out surface also reflected the poor mechanical properties of the matrix, and the introduction of the fibers could strengthen parts of the matrix which was around the fibers. The high orientation of
matrix along the fiber axis showed the well-compatibility between the fibers and the binder. The behavior of orientation might result from the stress-induced graphitization, since the observation of the SEM in Fig. 6b was similar to the results of the TEM (shown in Fig. 4b). Considering that the binder pitch was increasingly insufficient for wrapping fibers with fiber fraction, the excessive addition of short fibers would result in inferior interface combination. Furthermore, excessive fiber addition not only deteriorated the pore structure, but also caused the aggregation of fibers, which might lead to the abrupt decrease of the mechanical strength of SFG10.
4.
Conclusions
The goal of the present work was to improve the thermal conductivity and mechanical strength of graphite blocks. The effect of the short mesophase pitch carbon-fiber addition on the microstructure and thermal/mechanical properties of the NG-derived graphite blocks has been investigated. In the results discussed above, some important insights have been gained in this system: (1) Compared to the pristine graphite, the d002 of the SFGs increased with the increasing fiber content. The crystalline sizes were improved, which could due to the stress-induced graphitization. For the fiber content of 6 wt.%, the La and Lc reached 147 and 55 nm, respectively. (2) The fiber content played a key role in the thermal conductivity of SFGs. A tendency of reduced in-plane thermal conductivity of the graphite blocks was observed when the fiber content was less or higher than 6 wt.%. At a fiber content of 6 wt.%, the in-plane thermal conductivity achieved the maximum value of 518 Wm1 K1. The improvement of thermal conductivity relied on the increase of crystalline sizes and the formation of the 2D heat flow network. (3) Compared to the pristine graphite, all the SFGs had higher compressive and bend strength. The mechanical strength of the SFGs were reduced when the fiber content was less or higher than 8 wt.%. At a fiber content of 8 wt.%, the bend and compressive strength reached 39.6 and 65.5 MPa, respectively. The increase of the mechanical strength might be due to the pull-out effect of fibers. (4) The fiber addition increased the anisotropy of matrix both in microstructure and thermal/mechanical properties due to the orientation behavior of the fibers after hotpressing. Considering the poor mechanical properties of
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CARBON
5 3 ( 2 0 1 3 ) 3 1 3 –3 2 0
NG-derived graphite, the introduction of short fibers with a low fraction in this system was demonstrated to be a simple and effective approach to improve the mechanical properties while increasing the in-plane thermal conductivity prominently.
Acknowledgments The work was financially supported by the National Basic Research Program (973 Program, No. 2011CB605802) and National Natural Science Foundation (50902136) of China. The authors are grateful to Prof. Jinren Song for his helpful suggestion and encouragement.
Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.carbon. 2012.11.013.
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