炭基體種類對炭/炭復合材料內(nèi)耗行為的影響

楊　威，　羅瑞盈，　侯振華，　張　鈾，　商海東，　郝名揚

(北京航空航天大學 物理科學與核能工程學院， 北京100191)

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炭基體種類對炭/炭復合材料內(nèi)耗行為的影響

楊威，羅瑞盈，侯振華，張鈾，商海東，郝名揚

(北京航空航天大學 物理科學與核能工程學院， 北京100191)

摘要：通過化學氣相沉積(CVI)和化學氣相沉積與先驅體轉化結合(CVI+PIP)的方法，制備了三種不同炭基組織結構的炭/炭復合材料。三種基體分別是光滑層基體(SLC)、粗糙層基體(RLC)和混合雙基體(DMC)(過度生長錐基體+呋喃樹脂炭基體)。對這三種復合材料樣品進行微觀組織結構和動態(tài)力學性能表征。結果表明，內(nèi)耗主要來源于炭基體缺陷的運動、纖維/基體界面的滑移和炭平面的滑移。復合材料的內(nèi)耗對于溫度和振幅變化非常敏感，但頻率的變化對復合材料的的內(nèi)耗影響不大。混合雙基體具有最高的缺陷密度和最高的內(nèi)耗，粗糙層基體具備較完美的炭平面和最低的內(nèi)耗。炭基體的微觀組織結構是影響內(nèi)耗的關鍵因素，由于光滑層基體、粗糙層基體和混合雙基體的微觀結構的區(qū)別，導致在不同基體中出現(xiàn)了不同的內(nèi)耗行為。在室溫狀態(tài)下，基體中缺陷和纖維/基體的界面的運動可能是影響內(nèi)耗的主要因素，隨著溫度的升高，內(nèi)耗的貢獻可能主要來源于炭平面的滑移，而且我們還發(fā)現(xiàn)動態(tài)模量與缺陷密度存在一定關聯(lián)。

關鍵詞：炭/炭復合材料; 致密化工藝; 力學性能; 內(nèi)耗

English edition available online ScienceDirect ( http:www.sciencedirect.comsciencejournal18725805 ).

1Introduction

Carbon/carbon (C/C) composites are considered to be potentially ideal high-temperature structural materials for advanced aero-engine applications owing to their low density, outstanding mechanical properties, high thermal conductivity and low thermal expansion coefficient (CTE)[1]. In addition, the engine weight can be effectively reduced, contributing to low fuel consumption. Especially at some engine parts, such as gasket and sealing rings, have to undergo high-speed rotational and dynamic load. Therefore, it is very essential to investigate the internal friction of C/C composites to match the requirement of engine components. Based on the research of internal friction, scientific and technological workers can design aeronautic and astronautic structural materials with satisfying internal friction to ensure that the component can be used reliability. Moreover, the internal friction analysis can be used as a non-destructive characterization method to evaluate the composite properties.

The matrix microcracks in fiber reinforced glass matrix composites are assessed by internal friction[2]. According to the study of unidirectional C/C composites, the internal friction of C/C composites decreases with the frequency from 0.01 to 1 Hz, and then it slightly increases to 5 Hz, where the value of the internal friction reach about 0.5×10-2-10×10-2[3]. But the internal friction of C/C composites increases with the frequency from 0.01 to 2 Hz, and the value of the internal friction is very small (only 3×10-3-8×10-3)[4]. In a recent report, the internal friction of C/C composites decreases with the frequency from 0.01 to 1 Hz, and then increases obviously, and the value of the internal friction is about 1×10-2-3×10-2[5]. In the work of C/C composites at elevated temperatures, before temperature reach 2 000 K, the internal friction of C/C composites is almost constant below 2 000 K and increase with temperature above 2 000 K[6]. The internal friction of C/C composites decreases with bulk density and increases with the volume fraction of fibers[7]. There is little research literature on the issue of using internal friction to characterize the matrix microstructure of C/C composites.

Various carbon matrices have exhibited excellent mechanical properties, however, these carbon matrices with different structures have hardly been considered in terms of internal friction behaviors of C/C composites in former studies. The aim of this work is to investigate internal friction behaviors of C/C composites with the three typical carbon matrices, and try to find the relationship between internal friction behaviors and the microstructure of carbon matrix. The C/C composites with satisfactory dynamic mechanical properties are obtained through CVI or PIP, and we can design aeronautic and astronautic structural materials with satisfying internal friction.

2Experimental

A quasi three dimensional needled polyacrylonitrile based carbon fiber felts were used as a preforms and the density of the preform was about 0.55 g/cm3. The size of preform wasΦ230×20 mm, and the carbon fiber preform was firstly heat-treated at 2 300 ℃ for 2 h. Chemical vapor infiltration using nitrogen-diluted propane and hydrogen-diluted methane was used to densify the preform to prepare C/C composites, named as SLC and RLC, respectively. The preform was firstly densified by chemical vapor infiltration using methane/hydrogen and carbon dioxide till the density up to 1.50-1.60 g/cm3, and then impregnated with furan resin and carbonized to yield another C/C composite, named as DMC. Finally, the three kinds of composites were heat-treated at 2 400 ℃ after the densities are about 1.72 g/cm3.

The dynamic mechanical properties were characterized with a dynamic mechanical analyzer (DMA800) by means of three-point bending forced vibration in air. The specimens were rectangular bars with a size of 60 mm×4 mm×2 mm, cut from the fabricated composites. The span was 40 mm. Loading direction was perpendicular to the cloth layer direction. The testing frequency was ranged from 0.1 to 50 Hz, and the amplitude was from 0.004% to 0.05% of the strain. The temperature was from 25 to 450 ℃ and the heating rate was 5 ℃/min. The microstructures of the three C/C composites were characterized by polarized light microscopy (PLM, Neophot 21). Then, the polished surfaces of the C/C composites were analyzed by Raman spectrometry (LabRAM Aramis) with a laser excitation wavelength of 532 nm. The powder samples were examined by X-ray diffraction (XRD, D/M-2200) in the 2θrange of 15 and 80° with monochromatic Cu Ka radiation. According to Bragg’s law,d002was obtained from the Equation (1):

(1)

Whereλis the wavelength of CuKαradiation andθis the diffraction angle in radians. Crystallite sizeLcis obtained from the Scherrer Equation (2):

(2)

WhereBis the half maximum intensity in radians of the (002) peak. Graphitization degreegis calculated from the Maire and Mering Equation (3):

(3)

In addition, the matrix morphologies of the C/C composites were observed by a scanning electron microscope (SEM, JSM-6700F).

3Results

3.1The microstructure of the C/C composites

Fig. 1 presents the microstructure of the three kinds of C/C composites SLC, RLC and DMC under polarized light. It can be seen that the C/C composites are composed of three parts: carbon fibers, matrix carbon and small pores. Carbon fibers are fabricated by polyacrylonitrile, exhibiting obviously optically isotropic. It is very clear that both SLC and RLC present single optical structure with regular extinction crosses, which represent smooth laminar (SL) and rough laminar (RL) pyrocarbon structure, respectively, while the DMC composite exhibits dual matrix including the pyrocarbon with the overgrowth cones, the resin carbon (RC) and the interface of pyrocarbon/RC. Both the interface of pyrocarbon/RC and the RC exhibit low anisotropy. This is because that dual matrix weakens the optical anisotropy and the uniformity of matrix.

Fig. 1　PLM images of the C/C composites: (a) SLC, (b) RLC and (c) DMC.

The XRD spectra of the C/C composites are shown in Fig. 2. All of the three kinds of composites exhibit a sharp (002) peak, and their physical properties are listed in Table 1. Since RL pyrocarbon is easy to be graphitized, RLC has the highestgandLcamong the three kinds of composites. In contrast, SL pyrocarbon is difficult to be graphitized, so the correspondinggandLcare rather low. Although graphitization of RC is also hard, DMC has the relatively high average values ofgandLc.

Fig. 2　XRD patterns of the three kinds of C/C composites.

The C/C composites with dual matrix[8]possess a highgvalue, which may be due to the fact that the interface stress of pyrocarbon/RC would thermally induce the stress graphitization of RC. This can be also confirmed by the Raman spectra of matrices for DMC as shown in Fig. 3c. The spectra exhibit two distinct peaks at approximately 1 350 and 1 580 cm-1, which correspond to theDband assigned to defects within the carbon lattice (edges, distorted graphene layers, et al.) andGband of the symmetry vibration mode for graphite, respectively. For further analyzing, the Raman spectra is fitted with Lorentzian functions forDandGbands. We primarily investigate the full width at half maximum of theDband (FWHMD) and the intensity ratio of the two bands (R=ID/IG). Because FWHMDis very sensitive to the low energy structural defects and it is recognized thatRis inversely proportional to the microcrystalline in-plane size and the ability of graphitization[9]. In DMC, the value of FWHMDis 64.69 cm-1in the case of the overgrowth cones, whereas it is only 42.67 cm-1for the interface of pyrocarbon/RC and 48.09 cm-1for RC. TheRare 1.63 for the overgrowth cones, 1.39 for the interface and 1.33 for RC. It is shown that the interface possesses the lowest defect density, theRof the pyrocarbon/RC interface is almost the same as that of the RC, which should be attributable to the stress graphitization. And it is also indicated that overgrowth cones have very high defect density and are difficult to be graphitized. For assessing the average FWHMDandRof the matrix based on the simple rule of mixture, we must use polished samples for a measurement of the fiber diameter, pyrocarbon layers thickness, the interface layer thickness and the matrix thickness, and these values are 7, 3.7, 2.3 and 12 μm, respectively. These statistic data are mean values of at least 25 measurements. The average FWHMDandRof DMC are 50.05 cm-1and 1.40, respectively. The FWHMDandRof the other composites are listed in Table 1 according to the Raman spectra shown in Fig. 3a and Fig. 3b. It is shown that the FWHMDof the composites has such a relationship: FWHMD(DMC)> FWHMD(SLC)> FWHMD(RLC), whereas theRof the composites is in the sequence ofR(SLC)>R(RLC)>R(DMC).The typical SEM images of different matrices in the composites are shown in Fig. 4.

The remarkable laminate structure with delamination or cleavage between sub-layers in RL pyrocarbon can be seen from Fig. 4a. Those microstructures are not found in SL (Fig. 4b) and RC (Fig. 4c) as both the morphology are plate like. However, the writhed sub-layers consisted of part delamination are found in the overgrowth cones (Fig. 4d) and the interface of pyrocarbon/RC (Fig. 4c). These micrographs agree well with the results of Raman spectra of the matrices.

Table 1　Physical properties of the composites.

Fig. 3　Raman spectra of the matrices:

3.2The internal friction and the dynamic modulus vs. frequency

As the frequency increased from 0.1 to 50 Hz, the internal friction of the C/C composites decline gradually (Fig. 5a), whereas the dynamic modulus of the C/C composites is nearly unchanged as the frequency rise from 0.1 to 10 Hz, and begin to reduce sharply when the frequency is beyond 20 Hz (Fig. 5b). In these circumstances, the internal friction of the composites decreases with frequency, as shown in Fig. 5a, which may be ascribed to the long relaxation time. On the whole, DMC possesses the highest internal friction among the three kinds of composites, and the internal friction of RLC is the lowest. The dynamic modulus of DMC is the lowest among the three kinds of composites, and that of RLC is the highest.

3.3The internal friction and the dynamic modulus vs. amplitude

The internal friction of all of the composites is almost the same at 0.004% of the strain, and begins to rise gradually when the strain increases from 0.004% to 0.05% (Fig. 6a). The relationship of internal friction is IF(DMC)> IF(SLC)> IF(RLC), whereas dynamic modulus are obviously reversed (Fig. 6b). The dynamic modulus of RLC and DMC decline gradually, whereas that of SLC increases steeply with the strain when the strain is less than 0.01%. The dynamic modulus of SLC decreases gradually with a further increase of the strain beyond 0.01%. These phenomena indicate that the internal friction and dynamic modulus of the C/C composites are sensitive to amplitude.

3.4The internal friction and the dynamic modulus vs. temperature

Compared with the internal friction of the C/C composites versus frequency and amplitude, the internal friction of the C/C composites versus temperature exhibits some special responding characteristics. The internal friction of SLC increases rapidly with temperature from 25 to 100 ℃, While the internal friction decreases slowly with temperature in the temperature range of 100 to 450 ℃. Meanwhile, When the temperature increases from RT to 450 ℃, the internal friction of both RLC and DMC decreases slowly and then begins to increase, and RLC shows a minimal value at about 300 ℃ (Fig. 7a). At 375 ℃, the maximum dynamic modulus of SLC is achieved. Meanwhile, the dynamic modulus of RLC increases slightly and that of DMC is nearly unchanged with the temperature from RT to 450 ℃(Fig. 7b). In principle, SLC possesses the highest internal friction among the three kinds of composites, and the internal friction of RLC is the lowest. The dynamic modulus of DMC is the lowest among the three kinds of composites, and that of RLC is the highest. These characteristics further confirm that the carbon matrix plays an important role in the internal friction characters of the C/C composites. The values of internal friction and dynamic modulus are listed in Table 2, in order to find the relationship between the internal friction and the structural parameter of carbon matrix.

Fig. 4　The typical SEM images of different matrices in the three kinds of C/C composites:

Fig. 5　Dynamic mechanical properties of the three kinds of C/C composites versus frequency

Fig. 6　Dynamic mechanical properties of the three kinds of C/C composites versus

Fig. 7　Dynamic mechanical properties of the three kinds of C/C composites vs.

Temperature(℃)SLCRLCDMCInternalfriction250.03750.02680.05061000.05420.02600.04712000.05030.02570.04233000.04050.01710.03824000.03730.02330.0402Dynamicmodulus(GPa)2513.313.911.810012.513.712.220012.513.712.030012.813.812.140013.113.912.0

4Discussion

The internal friction behaviors of the C/C composites are dominated by the properties of the carbon matrix, which dissipate energy under dynamic loading. Three kinds of carbon matrix movement under dynamic loading are sketched in Fig. 8. The graphene layers of RLC are easier to move and have low density of defects, but those of DMC are more difficult to move and have more defects that restrain the sliding of basal planes.

Fig. 8　Three kinds of carbon matrix

The internal friction in the C/C composites is usually explained by the dislocation mechanism (K-G-L theory)[10], which is primarily caused by reciprocating motion between movable poor pinning points or immovable strong pinning points. In the C/C composites, the graphene layers are easy to be slipped because of the weak Van der Waals forces between the layers under cyclic loading. Therefore, the increase of the basal carbon plane distance (d002) can enhance the energy dissipation by the sliding of basal planes and improve the internal friction of the C/C composites[11]. Thus, the relationship between the internal friction andd002of the composites is shown in Fig.9. There is an obvious correlation between the internal friction andd002at 100 to 300 ℃, and the internal friction of the C/C composites distinctly increases with thed002. The decrease of Van der Waals forces contribute to the sliding of carbon planes, and increase of the energy dissipation, but the correlation appears very complicated at RT due to a large difference of the internal friction of SLC and DMC. The typical laminate structure of RLC indicates a perfect graphene layer, and the complex plate like structure exists in SLC and DMC. Obviously, the carbon plane in perfect graphene structure is much easier to slide, but the existence of plate like structure restrain the sliding of carbon planes, leading to the reduction of internal friction. However, in fact, the internal friction of SLC and DMC is always higher than that of RLC since SLC and DMC have more defects than RLC, which indicate that the internal friction produced by the sliding of carbon planes only provide a low percentage to the overall internal friction.

Fig. 9　The relationship between the internal friction and d002.

According to the above key points, at RT, the correlation of the internal friction toLc,R-1and FWGMDare shown in Fig. 10. The correlation with the internal friction at RT toLcand R-1is not obvious (Fig. 10a), but there is an obvious correlation between the internal friction and FWGMDat RT(Fig. 10b). The internal friction of the C/C composites monotonously increases with the FWGMD, which represent the defect density of the matrix. The high defect density can form more dislocations. Under cyclic stress, the dislocations will be also dislocated to dissipate energy. The motion of the dislocations is able to increase the internal friction. Thus, the high FWGMDis able to bring more internal friction in the C/C composites. This opinion is similar to the view[12]about enhancement of internal friction of the C/C composites, by selective oxidation to increase the interface defects. These further confirm that the defect density plays a significant role in the internal friction of the C/C composites. And we also find that the dynamic modulus of the C/C composites decreases distinctly with FWGMD. It is because that the dynamic modulus of the three kinds of composites always has the opposite tend with the internal friction of the composites at RT. Therefore, the highest internal friction and highest dynamic modulus cannot be achieved simultaneously.

Fig. 9 and Fig. 10 represent simplified correlations, the other important microstructural features, like the interface of fiber/matrix, are not taken into account. The internal friction of the C/C composites depends on the properties of the carbon matrix and the contribution of the fiber/matrix interface is negligibly small[13]. But that the major cause of frictional heating in fiber-reinforced ceramic matrix composites is energy dissipation derived from micro-frictional sliding between reinforcing fibers and matrix along the debonding interface, and the friction work of one cycle generated in fiber-reinforced ceramic matrix composites subjected to cyclic loading[14]is estimated by the following Equation (4):

(4)

Wheredfis the fiber diameter,Δσis the fatigue stress range,vfis the fiber volume fraction,Efis the elastic modulus of the fibers,τdis the interfacial shear stress, and C stands for (1-vf)Em/Ef,Emthe elastic modulus of the matrix.

Fig. 10　The correlation of the internal friction at room temperature to Lc, R-1 and FWGMD.

In Eq. (4), the interfacial shear stressτdis mainly determined by the residual stress, which is caused by the mismatch of CTE between the fibers and matrix in manufacturing process. With increasing temperature, the interfacial thermal strain is reduced and the binding between the fibers and matrix is enhanced, restraining the sliding between the fibers and matrix. Therefore, with the increase of temperature, the internal friction of the C/C composite is reduced. The microstructure of the pyrocarbon has no evident effect on the CTE, and relationships between the CTE of the C/C composites (α) and temperature (T) can be expressed[15]by the following Equation (5):

α=b0+b1T

(5)

Whereb0reflects the effect of the composite porosity, structures of the pyrocarbon and performs,b1values are about 12×10-10-14×10-10/K2.

The variation of the residual stress in the interface with temperature may be similar for different carbon matrices. Hence, there is no significant variation on the internal friction for the three kinds of composites. In this work, the internal friction of the composites is mainly attributed to the different carbon matrices. Meanwhile, the internal friction caused by the sliding of the interface is also an important part to the overall internal friction at low temperature range, but it provides the same contribution to the overall internal friction for the three kinds of composites when the temperature is increased.

According to the above analysis, the overall internal friction is mainly originated from the motion of the dislocations, the sliding of the fiber/matrix interface and the sliding of the carbon planes. Moreover, the internal friction, which is caused by the dislocations derived from the defects in the matrix may have a large advantage over the other effects at RT. Because the internal friction of the interface with different carbon matrices have the same contribution to the overall internal friction, the internal friction produced by the sliding of carbon planes may have a large contribution to the overall internal friction when the temperature is increased. And the dynamic modulus is mainly determined by the defect density.

5Conclusions

Three different types of carbon matrices are obtained by three kinds of densification processes. The C/C composites with satisfactory dynamic mechanical properties were obtained through CVI or PIP. We prepare the matrix of DMC by a combination of CVI and PIP. DMC composite has the highest crystal-defect densities than the other composites. The matrix of SLC and RLC are prepared only through CVI using propane/nitrogen and methane/hydrogen, respectively. RLC indicated the perfect carbon plane and the lowest internal friction among these composites. It is found that internal friction is correlated to the microstructure of the composites. In some certain conditions, carbon matrix has mainly contribution to the internal friction in the C/C composites. We compare three kinds of C/C composites with different carbon matrices. The following conclusions are drawn from our study:

The internal friction of the C/C composites is more sensitive to temperature and amplitude, and more insensitive to frequency. The internal friction of composites has such a relationship: IF(DMC)>IF(SLC)>IF(RLC)at RT and from 300 to 450 ℃, whereas the relationship changes to IF(SLC)>IF(DMC)>IF(RLC)when the temperature increases from 50 to 300 ℃. The overall internal friction of the C/C composites is mainly originated from the motion of the dislocations, the sliding of the fiber/matrix interfaces and the sliding of the carbon planes. At room temperature, the internal friction caused by the dislocations and the interfaces is superior to the other effects, which can be increased obviously by increasing FWGMDin matrices. When the temperature increases, the internal friction produced by the sliding of carbon planes contributes mainly to the overall internal friction, and increases distinctly with an increase ofd002.

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Foundationitem: National Natural Science Foundation of China(21071011).

Authorintroduction: YANG Wei, Ph. D Candidate. E-mail: yangwei_vip@hotmail.com

Influence of the microstructure of the carbon matrices on the internal friction behavior of carbon/carbon composites

YANG Wei,LUO Rui-ying,HOU Zhen-hua,ZHANG You,SHAGN Hai-dong,HAO Ming-yang

(SchoolofPhysicsandNuclearEnergyEngineering,BeihangUniversity,Beijing100191,China)

Abstract:Three carbon/carbon composites with rough laminar, smooth laminar and dual matrix carbon were prepared by chemical vapor infiltration (CVI) using hydrogen-diluted methane, CVI using nitrogen-diluted propane, and two-step CVI using first methane/hydrogen and carbon dioxideand then furan resin impregnation and carbonization. The influence of the microstructure of the carbon matrix on the internal friction behavior of the composites was investigated. Results indicate that the microstructure of the carbon matrix plays an important role in the internal friction. The overall internal friction is related to the motion of dislocations, the sliding of the fiber/matrix interface and the sliding of the carbon planes. The internal friction of the composite is very sensitive to temperature and amplitude, but less sensitive to frequency. Among these composites, the dual matrix carbon has the highest density of crystal-defects and the highest internal friction while the rough laminar carbon has perfect carbon planes and the lowest internal friction.

Keywords:Carbon/carbon composites; Densification process; Mechanical characterization; Internal friction.

文章編號：1007-8827(2016)02-0159-08

中圖分類號:TQ342+.74

文獻標識碼:A

基金項目：國家自然科學基金(21071011).

通信作者：羅瑞盈，教授. E-mail： ryluo@buaa.edu.cn

作者簡介：楊威，博士研究生. E-mail： yangwei_vip@hotmail.com

Corresponding author:LUO Rui-ying, Professor. E-mail: ryluo@buaa.edu.cn

DOI:10.1016/S1872-5805(16)60009-4