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Effect of high temperature treatment on the microstructure and elastoplastic properties of polyacrylonitrile-based carbon fibers
Carbon xxx (xxxx) xxx
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Carbon journal homepage: www.elsevier.com/locate/carbon
Effect of high temperature treatment on the microstructure and elastoplastic properties of polyacrylonitrile-based carbon fibers Fenghao Yang, Wen Liu, Maozhong Yi*, Liping Ran, Yicheng Ge, Ke Peng** State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 29 July 2019 Received in revised form 6 November 2019 Accepted 17 November 2019 Available online xxx
The effects of high temperature treatment (2000e2800 C) on the microstructure evolution and elastoplastic properties of polyacrylonitrile-based carbon fibers were investigated. It is found that the uneven shrinkage of the carbon fiber surfaces leads to formation of bulges whose length-to-width ratios increase with increasing heat treatment temperature. Because of the shrinkage stress, a closely stacked and curved graphene layer can be found in the sub-surface region of the carbon fibers. In the longitudinal section, the continuity of the graphite-like ribbons in the sub-surface region is worse than that in the skin region and the cross-linking regions between the graphite-like ribbons still contain a lot of defects after heat treatments. As the indentation depth increases, the axial and transverse elastic modulus of the carbon fibers decrease, which are related to the needle-like pores and the ordered graphite-like crystallites. The improvement of the slip between the graphene layers and the breakage of the graphite-like crystallites and the cross-linking regions are responsible for the decrease in axial elastic modulus with increasing heat treatment temperature. By contrast, the transverse elastic modulus is low and the excellent elasticity in the transverse direction is due to the curved graphene layers and the needle-like pores. © 2019 Published by Elsevier Ltd.
1. Introduction Polyacrylonitrile (PAN)-based carbon fiber reinforced carbon matrix (C/C) composites are the priority for aircraft braking systems and high temperature applications such as the nozzle and the throat liner of rockets due to their low density, outstanding tribological property and high temperature performances [1e5]. Because PAN-based carbon fibers are mostly continuous reinforcing structures, the properties of the carbon fibers such as tensile strength, tensile modulus and compressive strength have a great effect on the strength and ductility of C/C composites [6,7]. It is well known that C/C composites are subjected to complicated carbonization and graphitization processes during their preparation processing [5,6]. In inert and vacuum environments, C/C composites would be exposed in situations where the temperature exceeds 2000 C and even up to 3000 C [3,8]. These thermal conditions would cause a degradation of the properties of PAN-based carbon
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (K. Peng).
(M.
Yi),
[email protected]
fibers which is difficult to accurately evaluate in C/C composites. Thereby, the structure-property relationship of PAN-based carbon fibers heat-treated separately at high temperature is necessary to predict their properties in C/C composites. The structure of PAN-based carbon fibers is two-dimensional turbostratic graphene layers orientated preferentially along the fiber axis, and their skin-core heterogeneity has received extensive studies [9e15]. It has been found that the skin-core heterogeneity of the carbon fibers is mainly determined by the PAN precursor, preparation processing, heat treatment temperature and stretching conditions [16,17]. Compared with the graphene layers in the core region, the turbostratic graphene layers in the skin region easily transform into ordered graphene layers and the degree of the skincore heterogeneity increases with increasing heat treatment temperature [13,18]. Li et al. [19] have studied the skin-core heterogeneity and its effect on the tensile fracture mode, which showed that the fracture of the carbon fibers starts from the skin region. Our previous study found that the fracture sources of PAN-based carbon fibers are mainly located on the fiber surface or sub-surface region and the cracks prefer to propagate in the skin region [20]. Moreover, the carbon fiber surfaces with few voids or defects are beneficial to the fiber-matrix interlocking, but would decrease the
https://doi.org/10.1016/j.carbon.2019.11.055 0008-6223/© 2019 Published by Elsevier Ltd.
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tensile strength of the carbon fibers [21]. Therefore, the structural evolution in the sub-surface and skin regions is critical to reveal the properties of PAN-based carbon fibers. Nanoindentation, which can be used to characterize the elastoplastic properties of glass-like carbons [22e24], pyrocarbons [25] and carbon fibers [26e29], is a useful technique to determine the elastic modulus at nano- and micro-scales. The spherical tipped indentation onto glass-like carbons has been studied by Field and Swain et al. [22], suggesting that the glass-like carbons exhibit almost complete recoverable hysteresis response and the elastic modulus and yield stress of the glass-like carbons decrease with the development of the graphite structure [23]. However, unlike the microstructure of glass-like carbons, the graphite-like structure of PAN-based carbon fibers is anisotropic. Sun et al. [27] discussed the anisotropic behavior of single carbon fibers, which suggested that the axial elastic modulus and hardness are larger than the transverse elastic modulus and hardness. Maurin et al. [28] have analyzed the transverse modulus of high modulus carbon fibers and found that the transverse modulus has an inverse relationship with the tensile modulus. In addition, the elastic constants of single carbon fibers can also be estimated by the nanoindentation coupled with the modified Eshelby-mori-Tanaka model [29]. Given that the nanoindentation is depth-sensing and the heterogeneity of the microstructure of PAN-based carbon fibers has a close relationship with the heat treatment temperature, it is meaningful to investigate the effects of microstructure evolution on the elastoplastic properties of PAN-based carbon fibers by nanoindentation technique. In this paper, PAN-based carbon fibers were heat-treated at different high temperatures (2000e2800 C). The surface texture and the microstructure evolution of the carbon fibers in the subsurface and skin regions were studied. Nanoindentation experiments were performed in the transverse and axial directions to investigate the effects of heat treatment temperature on the elastoplastic properties of the carbon fibers. Finally, the relationship between the microstructure evolution and the elastoplastic properties of PAN-based carbon fibers was discussed. 2. Experimental 2.1. Sample preparation Commercially available 12 K PAN-based SYT55S carbon fibers (Zhongfu shenying Carbon Fiber Co., Ltd. Lianyungang, China) have been used. The cross-section of the as-received carbon fibers is circular (see Fig. 1). The density and diameter of the carbon fibers are 1.77 g/cm3 and 5.2 mm, respectively. Before heat treatments, the sizing agent was washed out by Soxhlet extraction method at 80 C for 48 h with a mixed solution of acetone and ethanol (volume
Fig. 1. Morphology of the cross-section of the as-received carbon fibers.
ration: 1:1). The carbon fibers were suspended in a graphite crucible and heated in a graphitization furnace (Advanced Corporation for Materials & Equipments Co., Ltd. Changsha, China). The heat treatment temperatures ranged from 2000 C to 2800 C. The heating rate was 5 C/min and the holding time at the desired temperature was 2 h under argon atmosphere. 2.2. Microstructure characterization The crystallite structure and the degree of graphitization were investigated by a Bruker D8 ADVANCE X-ray diffraction (XRD) with monochromatic Cu Ka radiation (l ¼ 0.154 nm). The acceleration voltage and emission current were 40 kV and 40 mA, respectively. The 2q value ranging from 15 to 60 was recorded under a scan speed of 1 /min. The standard silicon was added as an internal calibration to eliminate instrumental errors. The Ka2 was subtracted from the diffraction profile and the diffraction curve was fitted by a Voigt function to get the positions and full width at half maximums (FWHMs) of the (002), (100) peaks. The graphite-like crystallite parameters Lc and La, which represent the average thickness and the average length of the crystallites, respectively, were measured using Debye-Scherrer equation [30]:
Lc ¼
Kc l Ka l ; L ¼ bð002Þcosqð002Þ a bð100Þcosqð100Þ
(1)
The Scherrer parameters (Kc and Ka) in carbon materials depend on the form of layer stacking (turbostratic or AB stacking), and, more accurately, on the crystallite dimensions [31]. In the present study, Kc and Ka are taken as 0.90 and 1.84, respectively [31,32]. It should be noted that lower values of the scan speed contribute to more accurate crystallite parameters Lc and La [33]. According to the Bragg’s law, the interlayer spacing d(002) can be obtained and the degree of graphitization can be calculated by the following equation [34]:
gð%Þ ¼
0:3440 dð002Þ 100% 0:3440 0:3354
(2)
where g is the degree of graphitization (%), d(002) is the interlayer spacing (nm), 0.3440 is the interlayer spacing of the fully nongraphitized carbon [35], and 0.3354 is the interlayer spacing of the ideal graphite crystallite. Because equation (2) is an empirical formula, it is not applicable when the interlayer spacing of the turbostratic structure is larger than 0.3440 nm. Raman spectroscopy was obtained in a Raman microscope system (LabRAM ARAMIS) using an Ar ion laser as the excitation source (l ¼ 532 nm). A 50 microscope objective was used to focus the excitation laser beam on the carbon fiber surfaces and the laser spot diameter was 1 mm. The laser beam power was adjusted to no more than 5 mW using ND filters to avoid the laser heating effect. The Raman spectra were recorded in the range of 900e2000 cm1 and three spectra were measured for each carbon fiber. The spectra parameters such as band FWHM and band intensity were determined by a Voigt function. The carbon fiber surfaces were observed using a scanning electron microscope (SEM, FEI Helios Nanolab 600i). Focused ion beam machining (FIB, FEI Helios Nanolab 600i) was used to obtain the transverse and longitudinal slices of the carbon fibers, which were observed by a high resolution transmission electron microscopy (HRTEM, Tecnai G2 F20). In order to protect the edge of the slices, the carbon fiber surfaces were deposited a thin Platinum (Pt) layer. The selected-area electron-diffraction (SAED) patterns were recorded with a selected aperture of 110 nm in diameter. The orientation angle (OA) of the graphene layers was measured by the
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FWHM of the azimuthal intensity of the (002) arc reflections [18].
2.3. Nanoindentation test Nanoindentation experiments were performed on single carbon fiber surface in the transverse and axial directions using a CSM Indention Tester (UNHT) with a Berkovich diamond tip (see Fig. 2). The carbon fibers were embedded in poly(methyl methacrylate) and then polished parallel and perpendicular to the fiber axis. The initial depth of the indentation in the transverse direction was 500 nm to ensure that the indenter penetrated the carbon fiber surfaces accurately. In order to obtain elastic properties of the carbon fibers at different depths, the continuous multiple cyclic loading (CMC) mode was used. The indentation load was in the range of 0.1 mNe5 mN and 10 cycles were applied. The loading time and the unloading time were 30 s with a holding time of 10 s. The elastic modulus (EIT) was calculated from the load-displacement data based on the Oliver and Pharr theory [36]. The Poisson’s ratio of the carbon fibers was estimated to be 0.30 for calculation of the elastic modulus. The residual impression of the nanoindentation was analyzed by atomic force microscopy (AFM, Veeco NanoMan VS).
Fig. 3. XRD patterns of the as-received and heat-treated PAN-based carbon fibers. (A colour version of this figure can be viewed online.)
3.2. Surface texture 3. Results and discussion 3.1. X-ray diffraction analysis Fig. 3 shows the XRD patterns of the as-received and heattreated PAN-based carbon fibers and the corresponding results are listed in Table 1. A broad (002) diffraction peak of the graphitelike crystallites of the as-received carbon fibers can be seen at 25.31 and the interlayer spacing is about 0.3516, owing to the turbostratic stacking of the graphene layers. When the heat treatment temperature is higher than 2000 C, the (002) peak positions of the heat-treated carbon fibers shift from 25.95 to 26.14 and the interlayer spacing d(002) decreases significantly, indicating that the turbostratic stacking of the graphene layers gradually transforms into an ordered structure. Since the graphene layers are shifted, rotated and randomly oriented to each other, PAN-based carbon fibers are considered to be difficult to graphitize and the degree of graphitization only increases to 39.5% at 2800 C. The graphite-like crystallite parameters Lc and La increase with increasing heat treatment temperature. It is well known that there is a lot of amorphous carbon such as sp2 carbon clusters or sp2/sp3 carbons etc. distributes around the graphite-like crystallites [20,37]. During high temperature treatments, the transformation of amorphous carbon to ordered graphene layers, the rearrangement of the turbostratic graphene layers and the coalescence of graphite-like crystallites along the c-axis and a-axis are the main reasons for the growth of the graphite-like crystallites [38].
Fig. 2. Schematic diagram of the nanoindentation onto PAN-based carbon fibers in the transverse and axial directions. (A colour version of this figure can be viewed online.)
Raman spectroscopy is recognized as a suitable tool to provide quantitative information about the structure of carbon fiber surfaces at nano-scale, because the detection depth of Raman spectroscopy is about several tens of nanometers [37,39e41]. Fig. 4 shows the Raman spectra and the curve fitting results of the asreceived and heat-treated PAN-based carbon fibers. Three prominent bands related to the structure, including D1 band around 1360 cm1, G band (ideal graphitic lattice vibration) around 1580 cm1, and D2 band around 1620 cm1, can be found. The D1 and D2 bands are known to characterize the disordered carbon in hexagonal carbon layers or at the edges of crystallites. The overlapped D1 and G bands of the as-received carbon fibers are separated by taking into account another two bands, D4 and D3 bands, which derive from sp2-bonded amorphous carbon. After heat treatment at 2000 C, the sp2-bonded amorphous carbon almost disappears completely, and thus the D4 and D3 bands are not included in the fitting results. It is noticed that the FWHMs of the D1, G and D2 bands gradually decrease with increasing heat treatment temperature (see Fig. 4(e)). The integrated intensity ratio R ¼ ID1/IG can be used to estimate the degree of the organization of carbon materials [41], which decreases from 1.93 to 0.36 during heat treatments (see Fig. 4(f)). The results are consistent with the XRD analysis and indicate that the integrity and the ordering of the graphite-like crystallites in the sub-surface region have been improved. Fig. 5 shows the surface morphologies of the as-received and heat-treated PAN-based carbon fibers. The surface of the asreceived carbon fibers is relatively smooth and only a few longitudinal grooves can be seen. The surface morphology at higher magnification exhibits some closely stacked grain-like structure (see Fig. 5(a)), which is considered to be graphite-like crystallite clusters [17]. Because the graphite-like crystallite clusters consist of many turbostratic graphene layers, the rearrangement of the graphene layers would lead to the shrinkage of the carbon fiber surfaces. Given that the shrinkage of the edge of the graphite-like crystallite clusters may be the severest, a lot of bulges appear on the carbon fiber surfaces. The statistical distributions of the length-towidth ratio of the bulges and the corresponding results are shown in Fig. 6. It can be found that the average length-to-width ratios of the bulges increase from 3.02 to 5.10 with increasing heat
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Table 1 Graphite-like crystallite parameters, interlayer spacing d(002), degrees of graphitization of the as-received and heat-treated PAN-based carbon fibers. Heat treatment temperature/oC
Lc/nm
La/nm
d(002)/nm
Degree of graphitization/%
As-received 2000 2100 2300 2500 2700 2800
1.66 3.84 4.16 5.43 6.31 8.02 8.45
4.94 12.99 14.49 16.37 18.30 21.31 23.84
0.3516 0.3431 0.3427 0.3423 0.3419 0.3409 0.3406
e 10.5 15.1 19.8 24.4 36.0 39.5
Fig. 4. (a) Raman spectra and the curve fitting results of the as-received and heat-treated PAN-based carbon fibers: (b) As-received, (c) 2300 C, (d) 2800 C, (e) FWHMs of the D1, G and D2 bands and (f) integrated intensity ratio R (ID1/IG). (A colour version of this figure can be viewed online.)
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Fig. 5. Surface morphologies of the as-received and heat-treated PAN-based carbon fibers: (a) As-received, (b) 2300 C, (c) 2500 C, (d) 2800 C.
Fig. 6. (a) Statistical distributions of the length-to-width ratio of the bulges on the carbon fiber surfaces and (b) length-to-width ratios of the heat-treated PAN-based carbon fibers. (A colour version of this figure can be viewed online.)
treatment temperature. The increase in the length-to-width ratio of the bulges is due to the enhancement of the shrinkage caused by the growth and rearrangement of the graphene layers in the subsurface region of the carbon fibers.
3.3. Microstructure Fig. 7 shows the HRTEM images and SAED patterns in the transverse sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers. According to the starting position of the nanoindentation in the transverse direction (See Fig. 2), the skin region is about 500 nm from the carbon fiber surfaces. The diffraction (002) rings show that the OA values of the
graphene layers in the transverse sub-surface and skin regions of the carbon fibers are 180 , suggesting that the orientation of the graphene layers in the transverse section is random. As shown in Fig. 7(a) and (b), the graphene layers are short, disordered, bonded with each other by covalent bonds, which are under a thermodynamically unstable state. Considering that the molecular structure of the PAN precursor is helical and entangled, after heat treatment, some graphene layers that have similar directions of curvature would gather together to form onion-like structure. The core of the onion-like structure with disordered carbon atoms can be distinguished in the as-received carbon fibers. As the heat treatment temperature increases, the onion-like structure can be divided into three sections: a core section, a transition section and an outer
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Fig. 7. HRTEM images and SAED patterns in the transverse sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) as-received, (c) and (d) 2300 C, (e) and (f) 2800 C. (A colour version of this figure can be viewed online.)
section (see Fig. 7(c) and (d)). In the core section, the disordered carbon atoms become short and disordered graphene layers. Around the core section, the disordered graphene layers are orderly stacked to form graphite-like crystallites, which are marked with a yellow dotted rectangle, however, the graphene layers are relatively curved. In the outer section, some ordered graphene layers buckle along the transition section and connect the onion-like structure. Because the shrinkage stress would constrain the growth and rearrangement of the graphene layers, the size of the graphene layers in the sub-surface region is smaller than that in the skin region. When the carbon fibers are heat-treated at 2800 C, the growth and rearrangement of the graphene layers would be improved significantly, and thus the onion-like structure has transformed into fingerprint-like structure which contains a lot of long-ranged and orderly stacked graphene layers (see Fig. 7(f)). Many pores can be seen in the core of the fingerprint-like structure or at the edges of the crystallites. In the sub-surface region, the graphene layers are still short due to the increase in shrinkage stress, and a closely stacked layer about 20 nm in which the graphene layers are curved along the carbon fiber surface can be found. The width of the graphene layers in the transverse section increases with increasing heat treatment temperature, but their alignment is relatively tortuous, indicating that the graphene layers are subjected to a large amount of lattice strain. Owing to the lattice strain and strong covalent bond, a lot of structural defects such as crossover, bifurcation, stacking fault, dislocation, etc. appear. Fig. 8 shows the HRTEM images and SAED patterns in the longitudinal sub-surface and skin regions of the as-received and heat-
treated PAN-based carbon fibers. From the SAED patterns, it can be seen that the OA values both in the sub-surface and skin regions decrease with increasing heat treatment temperature, suggesting that the degree of the preferred orientation of the graphene layers is improved. In the sub-surface region of the as-received carbon fibers, the stacking and the orientation of graphene layers are disordered and the typical width of the disordered region is about several tens of nanometers (see Fig. 8(a)). In the skin region, the size of the graphite-like crystallites is very small and a lot of amorphous carbon and curved graphene layers distribute around these crystallites. After heat treatment at 2300 C, some graphite-like crystallites with preferred orientation appear in the sub-surface region, which could be inferred that the shrinkage stress has little effect on the growth of graphene layers parallel to the fiber axis. In the skin region, the disordered and curved graphene layers have transformed into ordered graphite-like crystallites because of the rearrangement of the graphene layers, but the graphene layers are still disordered and kinked between the graphite-like crystallites. As the heat treatment temperature increases to 2800 C, the stacking of the graphene layers becomes quite ordered and some graphitelike ribbons both in the sub-surface and skin regions appear, indicating that the continuity of the graphite-like crystallites has been improved significantly. Considering that the graphene layers are slightly bent in the longitudinal section, the lattice strain in the longitudinal section is much lower than that in the transverse section. A lot of structural defects also can be found in the graphitelike ribbons. Based on the HRTEM images, a schematic diagram of the microstructure of PAN-based carbon fibers heat-treated at 2800 C
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Fig. 8. HRTEM images and SAED patterns in the longitudinal sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) as-received, (c) and (d) 2300 C, (e) and (f) 2800 C. (A colour version of this figure can be viewed online.)
Fig. 9. (a) schematic diagram of the microstructure, morphologies of the needle-like pores: (b) transverse section and (c) longitudinal section, and (d) HRTEM images of the crosslinking region between the graphite-like ribbons of PAN-based carbon fibers heat-treated at 2800 C. (A colour version of this figure can be viewed online.)
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is shown in Fig. 9(a). It can be seen that the graphite-like crystallites are very tortuous in the transverse section, while straight and long in the longitudinal section. The morphologies of the graphite-like crystallites in the sub-surface region are different from those in the skin region. In the sub-surface region, owing to the shrinkage of the carbon fiber surfaces, the graphene layers are closely packed along the carbon fiber surface. Because the growth and rearrangement of the graphene layers would be suppressed, the longitudinal continuity of the graphite-like ribbons in the sub-surface region is worse than that in the skin region. During the rearrangement of the graphene layers, the needle-like pores appear around the graphite-like crystallites. Fig. 9(b) and (c) show the needle-like pores in the transverse and longitudinal sections of the carbon fibers, respectively. It can be seen that the transverse section of the needle-like pores is an irregular polygon. Because the shrinkage stress can compact the graphite-like crystallites, the size of the needle-like pores gradually increases from the sub-surface region to the skin region. Fig. 9(d) shows the cross-linking region between the graphite-like ribbons of the carbon fibers. Although the graphite-like ribbons have extended to the whole carbon fibers, the graphene layers are still deflected and kinked in the crosslinking region and the structural defects, needle-like pores and a small amount of cracks would aggregate in this region.
3.4. Elastoplastic properties The typical load-displacement curves of the nanoindentation in the axial and transverse directions are shown in Fig. 10(a) and (c),
respectively. Considering that the nanoindentation is very sensitive to surface roughness, the load-displacement curves of the first cycle are unconvinced and the corresponding results have been ignored. The hysteretic responses of PAN-based carbon fibers are significantly different between the axial and transverse directions, indicating that the carbon fibers exhibit anisotropic elastoplastic properties. In order to clarify the elastoplastic properties of the carbon fibers, the indentation ductility index D is calculated by Ref. [42]:
D¼
Wh Wh þ We
(3)
where, Wh is the hysteresis loop energy, defined as the area enclosed by the load-unloading curve. We is the elastic-stored energy, defined as the area under the unloading curve. In general, D is 1.00 for a purely plastic body without elastic recovery during unloading and 0.00 for a purely elastic body with a complete unloading recovery. As shown in Fig. 11, the transverse DIT-T of the carbon fibers ranges from 0.03 to 0.10, which is much lower than the axial DIT-A (0.15e0.36). Similar to the glass-like carbons [22,24], owing to the inter-planar slip of the graphene layers and localized dislocation motion, a yield-like behavior (‘plasticity’) of the carbon fibers would occur during nanoindentation. Therefore, it can be inferred that the carbon fibers exhibit almost complete elastic behavior in the transverse direction, while have a slight ‘plasticity’ in the axial direction. Moreover, the transverse DIT-T and axial DIT-A increase with increasing heat treatment temperature, suggesting
Fig. 10. Load-displacement curves of the nanoindentation: (a) axial direction and (c) transverse direction, and AFM images of the residual impressions of the nanoindentation onto PAN-based carbon fibers heat-treated at 2800 C: (b) axial direction and (d) transverse direction. (A colour version of this figure can be viewed online.)
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Fig. 11. Ductility index D of the nanoindentation in the axial and transverse directions.
that the ‘plasticity’ of the carbon fibers has been improved. As shown in Fig. 10(b) and (d), the residual impressions of the nanoindentation onto the carbon fibers heat-treated at 2800 C are very shallow and the residual penetration depths in the axial and transverse directions are 75.8 ± 1.6 nm and 87.8 ± 2.8 nm,
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respectively. A significant extrusion and cracks at the edge of the contact of the indenter can be found in the residual impressions. di et al. [29] found that the stress field penetrates along the Csana fiber axis under the axial compression, while extends perpendicular to the indentation direction under the transverse compression. Because of the difference of the stress field, a lot of cracks can be found in the transverse direction (see Fig. 10(d)). Furthermore, in the CMC mode of nanoindentation, repeated loading-unloading at the same position with an incremental load would induce crack initiation and propagation, resulting in accumulating damage to the carbon fiber surfaces. Consequently, the hysteresis loop energy (Wh) contains a little of energy consumed by ‘plasticity’ and a lot of energy consumed by crack initiation and propagation. The indentation load-depth and elastic modulus of the asreceived and heat-treated PAN-based carbon fibers derived from the load-displacement curves are shown in Fig. 12. As the heat treatment temperature increases, the maximum indentation depth gradually increases in the axial direction except for the carbon fibers heat-treated at 2300 C and 2700 C and the maximum indentation depth is about 389 ± 12 nm. The as-received carbon fibers exhibit the highest axial EIT-A, which decreases from 51.3 ± 1.4 GPa to 40.2 ± 1.0 GPa with increasing indentation depth (see Fig. 12(b)). As the heat treatment temperature increases, the axial EIT-A of the heat-treated carbon fibers decreases, suggesting that the elasticity of the carbon fibers is enhanced. The lowest axial EIT-A of the carbon fibers ranges from 33.6 ± 1.3 GPa to 22.7 ± 1.5 GPa after heat treatment at 2800 C.
Fig. 12. Indentation load-depth and elastic modulus of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) axial direction, (c) and (d) transverse direction. (A colour version of this figure can be viewed online.)
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F. Yang et al. / Carbon xxx (xxxx) xxx Table 2 Decline rates of the elastic modulus of the as-received and heat-treated PAN-based carbon fibers in the axial and transverse directions. Heat treatment temperature (oC)
Axial direction (102 GPa/nm)
Transverse direction (102 GPa/nm)
As-received 2000 2100 2300 2500 2700 2800
6.62 7.20 4.90 4.99 5.15 3.46 3.79
2.01 1.48 1.35 0.84 1.05 0.70 0.79
In the transverse direction, as the heat treatment temperature increases, the maximum indentation depth shows an upward trend and the maximum indentation depth is about 725 ± 7 nm which is nearly twice than that in the axial direction (see Fig. 12(c)). The transverse EIT-T of the carbon fibers decreases with increasing heat treatment temperature and indentation depth. The range of the transverse EIT-T of the as-received carbon fibers is from 16.9 ± 0.1 GPa to 9.6 ± 0.2 GPa, which is about one third of the axial EIT-A (see Fig. 12(d)). The transverse EIT-T of the carbon fibers decreases to 8.6 ± 0.1e4.2 ± 0.1 GPa when the heat treatment temperature increases to 2800 C. Based on the microstructure discussed above, it could be inferred that the decrease in elastic modulus along the indentation depth is closely related to the microstructure evolution. In order to clarify the relation, the decline rate is calculated based on the ratio of the change in elastic modulus to the corresponding change in indentation depth and the results are shown in Table 2. The accumulating damage in the form of cracks is an important reason for the decrease in elastic modulus along the indentation depth [22]. Because the cracks are easy to propagate along the graphite-like crystallites under the axial compression, the decline rates of the elastic modulus in the axial direction are larger than those in the transverse direction (see Table 2). As the heat treatment
temperature increases, the growth and rearrangement of the graphene layers would increase the size and order of the graphite-like crystallites, which could prevent the crack propagation effectively. Furthermore, the needle-like pores could also prevent the crack propagation of the nanoindentation in the transverse direction. It is clear that the microstructure evolution during heat treatments would reduce the effect of accumulating damage on the elastic modulus. Thus, the decline rates of the elastic modulus decrease after heat treatments and the carbon fibers heat-treated at 2700 C exhibit the minimum values. To explain the relationship between the microstructure evolution and the elastoplastic properties of PAN-based carbon fibers, a schematic mechanism of elastoplastic indentation onto as-received and heat-treated PAN-based carbon fibers at an early shallow stage is shown in Fig. 13. When the indenter penetrates the carbon fiber surfaces, the deformation of graphite-like crystallites and crosslinking regions between the graphite-like ribbons is vital to reveal the elastoplastic mechanism of PAN-based carbon fibers. In the as-received carbon fibers, the size of the graphite-like crystallites is small, and the amorphous carbon and turbostratic graphene layers are cross-linked by covalent bonds (see Fig. 13(a)), meaning that the breakage of the graphite-like crystallites and the crosslinking regions is difficult. Moreover, the slip of the turbostratic
Fig. 13. Schematic mechanism of elastoplastic indentation onto as-received and heat-treated PAN-based carbon fibers at an early shallow stage: (a) as-received and (b) 2800 C. (A colour version of this figure can be viewed online.)
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graphene layers is almost impossible. Thus, the as-received carbon fibers exhibit the highest axial EIT-A. After heat treatment at 2800 C, the ductility index D indicates that the ‘plasticity’ has been enhanced and the reasons may be explained as follows. First, the graphene layers would easily slip under compressive loads due to low energy of van der Waals bonds (see Fig. 13(b)). With increasing heat treatment temperature, the ordered stacking of the graphene layers means a significant slip of the graphene layers. Second, the crack would occur at the edges of the needle-like pores and propagate along the graphite-like crystallites, leading to the split of the crystallites. Finally, the graphite-like crystallites and the crosslinking regions would be easy to break due to the aggregation of the defects. Consequently, as the heat treatment temperature increases, the yield stress of the graphite-like crystallites and the cross-linking regions would decrease, leading to an improvement in ‘plasticity’ and a decrease in axial EIT-A. By contrast, in the transverse direction, because the turbostratic graphene layers are easy to bend and compress (see Fig. 13(a)), the transverse EIT-T is much lower than the axial EIT-A and the carbon fibers exhibit almost complete elasticity. When the turbostratic graphene layers transform into curved and continuous graphene layers, the bending deformation of the layers becomes easy and the needle-like pores also can provide the space for the elastic deformation (see Fig. 13(a)). Therefore, the transverse EIT-T gradually decreases with increasing heat treatment temperature. Moreover, the breakage of the graphite-like crystallites and the cross-linking regions is the main reason for the slight ‘plasticity’. 4. Conclusions (1) As the heat treatment temperature increases, the turbostratic graphene layers in the sub-surface region transform into an ordered graphite-like structure, leading to formation of many bulges on the carbon fiber surfaces. The average length-to-width ratio of the bulges increases from 3.02 to 5.10, owing to the growth and rearrangement of the graphene layers in the sub-surface region. (2) The shrinkage of the carbon fiber surfaces would constrain the growth and rearrangement of the graphene layers in the sub-surface region. The graphene layers are closely stacked and curved along the carbon fiber surface in the sub-surface region. In the longitudinal section, the disordered graphene layers have become ordered graphite-like ribbons whose continuity in the sub-surface region is worse than that in the skin region. The cross-linking regions between the graphitelike ribbons in the longitudinal direction still consist of a lot of kinked graphene layers, structural defects and needle-like pores after heat treatments. (3) The elastic modulus of PAN-based carbon fibers in the axial and transverse directions decreases with increasing indentation depth. The needle-like pores and the ordered graphite-like crystallites could prevent the crack propagation, leading to decrease in the decline rate of the elastic modulus along the indentation depth. Because of the improvement of the slip between the graphene layers and the breakage of the graphite-like crystallites and the crosslinking regions, the axial elastic modulus decreases with increasing heat treatment temperature. In the transverse direction, the carbon fibers exhibit almost complete elasticity, which strongly depends on the curved graphene layers and the needle-like pores. Declaration of competing interests The authors declare that they have no known competing
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Effect of high temperature treatment on the microstructure and elastoplastic properties of polyacrylonitrile-based carbon fibers Fenghao Yang, Wen Liu, Maozhong Yi*, Liping Ran, Yicheng Ge, Ke Peng** State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 29 July 2019 Received in revised form 6 November 2019 Accepted 17 November 2019 Available online xxx
The effects of high temperature treatment (2000e2800 C) on the microstructure evolution and elastoplastic properties of polyacrylonitrile-based carbon fibers were investigated. It is found that the uneven shrinkage of the carbon fiber surfaces leads to formation of bulges whose length-to-width ratios increase with increasing heat treatment temperature. Because of the shrinkage stress, a closely stacked and curved graphene layer can be found in the sub-surface region of the carbon fibers. In the longitudinal section, the continuity of the graphite-like ribbons in the sub-surface region is worse than that in the skin region and the cross-linking regions between the graphite-like ribbons still contain a lot of defects after heat treatments. As the indentation depth increases, the axial and transverse elastic modulus of the carbon fibers decrease, which are related to the needle-like pores and the ordered graphite-like crystallites. The improvement of the slip between the graphene layers and the breakage of the graphite-like crystallites and the cross-linking regions are responsible for the decrease in axial elastic modulus with increasing heat treatment temperature. By contrast, the transverse elastic modulus is low and the excellent elasticity in the transverse direction is due to the curved graphene layers and the needle-like pores. © 2019 Published by Elsevier Ltd.
1. Introduction Polyacrylonitrile (PAN)-based carbon fiber reinforced carbon matrix (C/C) composites are the priority for aircraft braking systems and high temperature applications such as the nozzle and the throat liner of rockets due to their low density, outstanding tribological property and high temperature performances [1e5]. Because PAN-based carbon fibers are mostly continuous reinforcing structures, the properties of the carbon fibers such as tensile strength, tensile modulus and compressive strength have a great effect on the strength and ductility of C/C composites [6,7]. It is well known that C/C composites are subjected to complicated carbonization and graphitization processes during their preparation processing [5,6]. In inert and vacuum environments, C/C composites would be exposed in situations where the temperature exceeds 2000 C and even up to 3000 C [3,8]. These thermal conditions would cause a degradation of the properties of PAN-based carbon
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (K. Peng).
(M.
Yi),
[email protected]
fibers which is difficult to accurately evaluate in C/C composites. Thereby, the structure-property relationship of PAN-based carbon fibers heat-treated separately at high temperature is necessary to predict their properties in C/C composites. The structure of PAN-based carbon fibers is two-dimensional turbostratic graphene layers orientated preferentially along the fiber axis, and their skin-core heterogeneity has received extensive studies [9e15]. It has been found that the skin-core heterogeneity of the carbon fibers is mainly determined by the PAN precursor, preparation processing, heat treatment temperature and stretching conditions [16,17]. Compared with the graphene layers in the core region, the turbostratic graphene layers in the skin region easily transform into ordered graphene layers and the degree of the skincore heterogeneity increases with increasing heat treatment temperature [13,18]. Li et al. [19] have studied the skin-core heterogeneity and its effect on the tensile fracture mode, which showed that the fracture of the carbon fibers starts from the skin region. Our previous study found that the fracture sources of PAN-based carbon fibers are mainly located on the fiber surface or sub-surface region and the cracks prefer to propagate in the skin region [20]. Moreover, the carbon fiber surfaces with few voids or defects are beneficial to the fiber-matrix interlocking, but would decrease the
https://doi.org/10.1016/j.carbon.2019.11.055 0008-6223/© 2019 Published by Elsevier Ltd.
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tensile strength of the carbon fibers [21]. Therefore, the structural evolution in the sub-surface and skin regions is critical to reveal the properties of PAN-based carbon fibers. Nanoindentation, which can be used to characterize the elastoplastic properties of glass-like carbons [22e24], pyrocarbons [25] and carbon fibers [26e29], is a useful technique to determine the elastic modulus at nano- and micro-scales. The spherical tipped indentation onto glass-like carbons has been studied by Field and Swain et al. [22], suggesting that the glass-like carbons exhibit almost complete recoverable hysteresis response and the elastic modulus and yield stress of the glass-like carbons decrease with the development of the graphite structure [23]. However, unlike the microstructure of glass-like carbons, the graphite-like structure of PAN-based carbon fibers is anisotropic. Sun et al. [27] discussed the anisotropic behavior of single carbon fibers, which suggested that the axial elastic modulus and hardness are larger than the transverse elastic modulus and hardness. Maurin et al. [28] have analyzed the transverse modulus of high modulus carbon fibers and found that the transverse modulus has an inverse relationship with the tensile modulus. In addition, the elastic constants of single carbon fibers can also be estimated by the nanoindentation coupled with the modified Eshelby-mori-Tanaka model [29]. Given that the nanoindentation is depth-sensing and the heterogeneity of the microstructure of PAN-based carbon fibers has a close relationship with the heat treatment temperature, it is meaningful to investigate the effects of microstructure evolution on the elastoplastic properties of PAN-based carbon fibers by nanoindentation technique. In this paper, PAN-based carbon fibers were heat-treated at different high temperatures (2000e2800 C). The surface texture and the microstructure evolution of the carbon fibers in the subsurface and skin regions were studied. Nanoindentation experiments were performed in the transverse and axial directions to investigate the effects of heat treatment temperature on the elastoplastic properties of the carbon fibers. Finally, the relationship between the microstructure evolution and the elastoplastic properties of PAN-based carbon fibers was discussed. 2. Experimental 2.1. Sample preparation Commercially available 12 K PAN-based SYT55S carbon fibers (Zhongfu shenying Carbon Fiber Co., Ltd. Lianyungang, China) have been used. The cross-section of the as-received carbon fibers is circular (see Fig. 1). The density and diameter of the carbon fibers are 1.77 g/cm3 and 5.2 mm, respectively. Before heat treatments, the sizing agent was washed out by Soxhlet extraction method at 80 C for 48 h with a mixed solution of acetone and ethanol (volume
Fig. 1. Morphology of the cross-section of the as-received carbon fibers.
ration: 1:1). The carbon fibers were suspended in a graphite crucible and heated in a graphitization furnace (Advanced Corporation for Materials & Equipments Co., Ltd. Changsha, China). The heat treatment temperatures ranged from 2000 C to 2800 C. The heating rate was 5 C/min and the holding time at the desired temperature was 2 h under argon atmosphere. 2.2. Microstructure characterization The crystallite structure and the degree of graphitization were investigated by a Bruker D8 ADVANCE X-ray diffraction (XRD) with monochromatic Cu Ka radiation (l ¼ 0.154 nm). The acceleration voltage and emission current were 40 kV and 40 mA, respectively. The 2q value ranging from 15 to 60 was recorded under a scan speed of 1 /min. The standard silicon was added as an internal calibration to eliminate instrumental errors. The Ka2 was subtracted from the diffraction profile and the diffraction curve was fitted by a Voigt function to get the positions and full width at half maximums (FWHMs) of the (002), (100) peaks. The graphite-like crystallite parameters Lc and La, which represent the average thickness and the average length of the crystallites, respectively, were measured using Debye-Scherrer equation [30]:
Lc ¼
Kc l Ka l ; L ¼ bð002Þcosqð002Þ a bð100Þcosqð100Þ
(1)
The Scherrer parameters (Kc and Ka) in carbon materials depend on the form of layer stacking (turbostratic or AB stacking), and, more accurately, on the crystallite dimensions [31]. In the present study, Kc and Ka are taken as 0.90 and 1.84, respectively [31,32]. It should be noted that lower values of the scan speed contribute to more accurate crystallite parameters Lc and La [33]. According to the Bragg’s law, the interlayer spacing d(002) can be obtained and the degree of graphitization can be calculated by the following equation [34]:
gð%Þ ¼
0:3440 dð002Þ 100% 0:3440 0:3354
(2)
where g is the degree of graphitization (%), d(002) is the interlayer spacing (nm), 0.3440 is the interlayer spacing of the fully nongraphitized carbon [35], and 0.3354 is the interlayer spacing of the ideal graphite crystallite. Because equation (2) is an empirical formula, it is not applicable when the interlayer spacing of the turbostratic structure is larger than 0.3440 nm. Raman spectroscopy was obtained in a Raman microscope system (LabRAM ARAMIS) using an Ar ion laser as the excitation source (l ¼ 532 nm). A 50 microscope objective was used to focus the excitation laser beam on the carbon fiber surfaces and the laser spot diameter was 1 mm. The laser beam power was adjusted to no more than 5 mW using ND filters to avoid the laser heating effect. The Raman spectra were recorded in the range of 900e2000 cm1 and three spectra were measured for each carbon fiber. The spectra parameters such as band FWHM and band intensity were determined by a Voigt function. The carbon fiber surfaces were observed using a scanning electron microscope (SEM, FEI Helios Nanolab 600i). Focused ion beam machining (FIB, FEI Helios Nanolab 600i) was used to obtain the transverse and longitudinal slices of the carbon fibers, which were observed by a high resolution transmission electron microscopy (HRTEM, Tecnai G2 F20). In order to protect the edge of the slices, the carbon fiber surfaces were deposited a thin Platinum (Pt) layer. The selected-area electron-diffraction (SAED) patterns were recorded with a selected aperture of 110 nm in diameter. The orientation angle (OA) of the graphene layers was measured by the
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FWHM of the azimuthal intensity of the (002) arc reflections [18].
2.3. Nanoindentation test Nanoindentation experiments were performed on single carbon fiber surface in the transverse and axial directions using a CSM Indention Tester (UNHT) with a Berkovich diamond tip (see Fig. 2). The carbon fibers were embedded in poly(methyl methacrylate) and then polished parallel and perpendicular to the fiber axis. The initial depth of the indentation in the transverse direction was 500 nm to ensure that the indenter penetrated the carbon fiber surfaces accurately. In order to obtain elastic properties of the carbon fibers at different depths, the continuous multiple cyclic loading (CMC) mode was used. The indentation load was in the range of 0.1 mNe5 mN and 10 cycles were applied. The loading time and the unloading time were 30 s with a holding time of 10 s. The elastic modulus (EIT) was calculated from the load-displacement data based on the Oliver and Pharr theory [36]. The Poisson’s ratio of the carbon fibers was estimated to be 0.30 for calculation of the elastic modulus. The residual impression of the nanoindentation was analyzed by atomic force microscopy (AFM, Veeco NanoMan VS).
Fig. 3. XRD patterns of the as-received and heat-treated PAN-based carbon fibers. (A colour version of this figure can be viewed online.)
3.2. Surface texture 3. Results and discussion 3.1. X-ray diffraction analysis Fig. 3 shows the XRD patterns of the as-received and heattreated PAN-based carbon fibers and the corresponding results are listed in Table 1. A broad (002) diffraction peak of the graphitelike crystallites of the as-received carbon fibers can be seen at 25.31 and the interlayer spacing is about 0.3516, owing to the turbostratic stacking of the graphene layers. When the heat treatment temperature is higher than 2000 C, the (002) peak positions of the heat-treated carbon fibers shift from 25.95 to 26.14 and the interlayer spacing d(002) decreases significantly, indicating that the turbostratic stacking of the graphene layers gradually transforms into an ordered structure. Since the graphene layers are shifted, rotated and randomly oriented to each other, PAN-based carbon fibers are considered to be difficult to graphitize and the degree of graphitization only increases to 39.5% at 2800 C. The graphite-like crystallite parameters Lc and La increase with increasing heat treatment temperature. It is well known that there is a lot of amorphous carbon such as sp2 carbon clusters or sp2/sp3 carbons etc. distributes around the graphite-like crystallites [20,37]. During high temperature treatments, the transformation of amorphous carbon to ordered graphene layers, the rearrangement of the turbostratic graphene layers and the coalescence of graphite-like crystallites along the c-axis and a-axis are the main reasons for the growth of the graphite-like crystallites [38].
Fig. 2. Schematic diagram of the nanoindentation onto PAN-based carbon fibers in the transverse and axial directions. (A colour version of this figure can be viewed online.)
Raman spectroscopy is recognized as a suitable tool to provide quantitative information about the structure of carbon fiber surfaces at nano-scale, because the detection depth of Raman spectroscopy is about several tens of nanometers [37,39e41]. Fig. 4 shows the Raman spectra and the curve fitting results of the asreceived and heat-treated PAN-based carbon fibers. Three prominent bands related to the structure, including D1 band around 1360 cm1, G band (ideal graphitic lattice vibration) around 1580 cm1, and D2 band around 1620 cm1, can be found. The D1 and D2 bands are known to characterize the disordered carbon in hexagonal carbon layers or at the edges of crystallites. The overlapped D1 and G bands of the as-received carbon fibers are separated by taking into account another two bands, D4 and D3 bands, which derive from sp2-bonded amorphous carbon. After heat treatment at 2000 C, the sp2-bonded amorphous carbon almost disappears completely, and thus the D4 and D3 bands are not included in the fitting results. It is noticed that the FWHMs of the D1, G and D2 bands gradually decrease with increasing heat treatment temperature (see Fig. 4(e)). The integrated intensity ratio R ¼ ID1/IG can be used to estimate the degree of the organization of carbon materials [41], which decreases from 1.93 to 0.36 during heat treatments (see Fig. 4(f)). The results are consistent with the XRD analysis and indicate that the integrity and the ordering of the graphite-like crystallites in the sub-surface region have been improved. Fig. 5 shows the surface morphologies of the as-received and heat-treated PAN-based carbon fibers. The surface of the asreceived carbon fibers is relatively smooth and only a few longitudinal grooves can be seen. The surface morphology at higher magnification exhibits some closely stacked grain-like structure (see Fig. 5(a)), which is considered to be graphite-like crystallite clusters [17]. Because the graphite-like crystallite clusters consist of many turbostratic graphene layers, the rearrangement of the graphene layers would lead to the shrinkage of the carbon fiber surfaces. Given that the shrinkage of the edge of the graphite-like crystallite clusters may be the severest, a lot of bulges appear on the carbon fiber surfaces. The statistical distributions of the length-towidth ratio of the bulges and the corresponding results are shown in Fig. 6. It can be found that the average length-to-width ratios of the bulges increase from 3.02 to 5.10 with increasing heat
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Table 1 Graphite-like crystallite parameters, interlayer spacing d(002), degrees of graphitization of the as-received and heat-treated PAN-based carbon fibers. Heat treatment temperature/oC
Lc/nm
La/nm
d(002)/nm
Degree of graphitization/%
As-received 2000 2100 2300 2500 2700 2800
1.66 3.84 4.16 5.43 6.31 8.02 8.45
4.94 12.99 14.49 16.37 18.30 21.31 23.84
0.3516 0.3431 0.3427 0.3423 0.3419 0.3409 0.3406
e 10.5 15.1 19.8 24.4 36.0 39.5
Fig. 4. (a) Raman spectra and the curve fitting results of the as-received and heat-treated PAN-based carbon fibers: (b) As-received, (c) 2300 C, (d) 2800 C, (e) FWHMs of the D1, G and D2 bands and (f) integrated intensity ratio R (ID1/IG). (A colour version of this figure can be viewed online.)
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Fig. 5. Surface morphologies of the as-received and heat-treated PAN-based carbon fibers: (a) As-received, (b) 2300 C, (c) 2500 C, (d) 2800 C.
Fig. 6. (a) Statistical distributions of the length-to-width ratio of the bulges on the carbon fiber surfaces and (b) length-to-width ratios of the heat-treated PAN-based carbon fibers. (A colour version of this figure can be viewed online.)
treatment temperature. The increase in the length-to-width ratio of the bulges is due to the enhancement of the shrinkage caused by the growth and rearrangement of the graphene layers in the subsurface region of the carbon fibers.
3.3. Microstructure Fig. 7 shows the HRTEM images and SAED patterns in the transverse sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers. According to the starting position of the nanoindentation in the transverse direction (See Fig. 2), the skin region is about 500 nm from the carbon fiber surfaces. The diffraction (002) rings show that the OA values of the
graphene layers in the transverse sub-surface and skin regions of the carbon fibers are 180 , suggesting that the orientation of the graphene layers in the transverse section is random. As shown in Fig. 7(a) and (b), the graphene layers are short, disordered, bonded with each other by covalent bonds, which are under a thermodynamically unstable state. Considering that the molecular structure of the PAN precursor is helical and entangled, after heat treatment, some graphene layers that have similar directions of curvature would gather together to form onion-like structure. The core of the onion-like structure with disordered carbon atoms can be distinguished in the as-received carbon fibers. As the heat treatment temperature increases, the onion-like structure can be divided into three sections: a core section, a transition section and an outer
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Fig. 7. HRTEM images and SAED patterns in the transverse sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) as-received, (c) and (d) 2300 C, (e) and (f) 2800 C. (A colour version of this figure can be viewed online.)
section (see Fig. 7(c) and (d)). In the core section, the disordered carbon atoms become short and disordered graphene layers. Around the core section, the disordered graphene layers are orderly stacked to form graphite-like crystallites, which are marked with a yellow dotted rectangle, however, the graphene layers are relatively curved. In the outer section, some ordered graphene layers buckle along the transition section and connect the onion-like structure. Because the shrinkage stress would constrain the growth and rearrangement of the graphene layers, the size of the graphene layers in the sub-surface region is smaller than that in the skin region. When the carbon fibers are heat-treated at 2800 C, the growth and rearrangement of the graphene layers would be improved significantly, and thus the onion-like structure has transformed into fingerprint-like structure which contains a lot of long-ranged and orderly stacked graphene layers (see Fig. 7(f)). Many pores can be seen in the core of the fingerprint-like structure or at the edges of the crystallites. In the sub-surface region, the graphene layers are still short due to the increase in shrinkage stress, and a closely stacked layer about 20 nm in which the graphene layers are curved along the carbon fiber surface can be found. The width of the graphene layers in the transverse section increases with increasing heat treatment temperature, but their alignment is relatively tortuous, indicating that the graphene layers are subjected to a large amount of lattice strain. Owing to the lattice strain and strong covalent bond, a lot of structural defects such as crossover, bifurcation, stacking fault, dislocation, etc. appear. Fig. 8 shows the HRTEM images and SAED patterns in the longitudinal sub-surface and skin regions of the as-received and heat-
treated PAN-based carbon fibers. From the SAED patterns, it can be seen that the OA values both in the sub-surface and skin regions decrease with increasing heat treatment temperature, suggesting that the degree of the preferred orientation of the graphene layers is improved. In the sub-surface region of the as-received carbon fibers, the stacking and the orientation of graphene layers are disordered and the typical width of the disordered region is about several tens of nanometers (see Fig. 8(a)). In the skin region, the size of the graphite-like crystallites is very small and a lot of amorphous carbon and curved graphene layers distribute around these crystallites. After heat treatment at 2300 C, some graphite-like crystallites with preferred orientation appear in the sub-surface region, which could be inferred that the shrinkage stress has little effect on the growth of graphene layers parallel to the fiber axis. In the skin region, the disordered and curved graphene layers have transformed into ordered graphite-like crystallites because of the rearrangement of the graphene layers, but the graphene layers are still disordered and kinked between the graphite-like crystallites. As the heat treatment temperature increases to 2800 C, the stacking of the graphene layers becomes quite ordered and some graphitelike ribbons both in the sub-surface and skin regions appear, indicating that the continuity of the graphite-like crystallites has been improved significantly. Considering that the graphene layers are slightly bent in the longitudinal section, the lattice strain in the longitudinal section is much lower than that in the transverse section. A lot of structural defects also can be found in the graphitelike ribbons. Based on the HRTEM images, a schematic diagram of the microstructure of PAN-based carbon fibers heat-treated at 2800 C
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Fig. 8. HRTEM images and SAED patterns in the longitudinal sub-surface and skin regions of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) as-received, (c) and (d) 2300 C, (e) and (f) 2800 C. (A colour version of this figure can be viewed online.)
Fig. 9. (a) schematic diagram of the microstructure, morphologies of the needle-like pores: (b) transverse section and (c) longitudinal section, and (d) HRTEM images of the crosslinking region between the graphite-like ribbons of PAN-based carbon fibers heat-treated at 2800 C. (A colour version of this figure can be viewed online.)
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is shown in Fig. 9(a). It can be seen that the graphite-like crystallites are very tortuous in the transverse section, while straight and long in the longitudinal section. The morphologies of the graphite-like crystallites in the sub-surface region are different from those in the skin region. In the sub-surface region, owing to the shrinkage of the carbon fiber surfaces, the graphene layers are closely packed along the carbon fiber surface. Because the growth and rearrangement of the graphene layers would be suppressed, the longitudinal continuity of the graphite-like ribbons in the sub-surface region is worse than that in the skin region. During the rearrangement of the graphene layers, the needle-like pores appear around the graphite-like crystallites. Fig. 9(b) and (c) show the needle-like pores in the transverse and longitudinal sections of the carbon fibers, respectively. It can be seen that the transverse section of the needle-like pores is an irregular polygon. Because the shrinkage stress can compact the graphite-like crystallites, the size of the needle-like pores gradually increases from the sub-surface region to the skin region. Fig. 9(d) shows the cross-linking region between the graphite-like ribbons of the carbon fibers. Although the graphite-like ribbons have extended to the whole carbon fibers, the graphene layers are still deflected and kinked in the crosslinking region and the structural defects, needle-like pores and a small amount of cracks would aggregate in this region.
3.4. Elastoplastic properties The typical load-displacement curves of the nanoindentation in the axial and transverse directions are shown in Fig. 10(a) and (c),
respectively. Considering that the nanoindentation is very sensitive to surface roughness, the load-displacement curves of the first cycle are unconvinced and the corresponding results have been ignored. The hysteretic responses of PAN-based carbon fibers are significantly different between the axial and transverse directions, indicating that the carbon fibers exhibit anisotropic elastoplastic properties. In order to clarify the elastoplastic properties of the carbon fibers, the indentation ductility index D is calculated by Ref. [42]:
D¼
Wh Wh þ We
(3)
where, Wh is the hysteresis loop energy, defined as the area enclosed by the load-unloading curve. We is the elastic-stored energy, defined as the area under the unloading curve. In general, D is 1.00 for a purely plastic body without elastic recovery during unloading and 0.00 for a purely elastic body with a complete unloading recovery. As shown in Fig. 11, the transverse DIT-T of the carbon fibers ranges from 0.03 to 0.10, which is much lower than the axial DIT-A (0.15e0.36). Similar to the glass-like carbons [22,24], owing to the inter-planar slip of the graphene layers and localized dislocation motion, a yield-like behavior (‘plasticity’) of the carbon fibers would occur during nanoindentation. Therefore, it can be inferred that the carbon fibers exhibit almost complete elastic behavior in the transverse direction, while have a slight ‘plasticity’ in the axial direction. Moreover, the transverse DIT-T and axial DIT-A increase with increasing heat treatment temperature, suggesting
Fig. 10. Load-displacement curves of the nanoindentation: (a) axial direction and (c) transverse direction, and AFM images of the residual impressions of the nanoindentation onto PAN-based carbon fibers heat-treated at 2800 C: (b) axial direction and (d) transverse direction. (A colour version of this figure can be viewed online.)
Please cite this article as: F. Yang et al., Effect of high temperature treatment on the microstructure and elastoplastic properties of polyacrylonitrile-based carbon fibers, Carbon, https://doi.org/10.1016/j.carbon.2019.11.055
F. Yang et al. / Carbon xxx (xxxx) xxx
Fig. 11. Ductility index D of the nanoindentation in the axial and transverse directions.
that the ‘plasticity’ of the carbon fibers has been improved. As shown in Fig. 10(b) and (d), the residual impressions of the nanoindentation onto the carbon fibers heat-treated at 2800 C are very shallow and the residual penetration depths in the axial and transverse directions are 75.8 ± 1.6 nm and 87.8 ± 2.8 nm,
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respectively. A significant extrusion and cracks at the edge of the contact of the indenter can be found in the residual impressions. di et al. [29] found that the stress field penetrates along the Csana fiber axis under the axial compression, while extends perpendicular to the indentation direction under the transverse compression. Because of the difference of the stress field, a lot of cracks can be found in the transverse direction (see Fig. 10(d)). Furthermore, in the CMC mode of nanoindentation, repeated loading-unloading at the same position with an incremental load would induce crack initiation and propagation, resulting in accumulating damage to the carbon fiber surfaces. Consequently, the hysteresis loop energy (Wh) contains a little of energy consumed by ‘plasticity’ and a lot of energy consumed by crack initiation and propagation. The indentation load-depth and elastic modulus of the asreceived and heat-treated PAN-based carbon fibers derived from the load-displacement curves are shown in Fig. 12. As the heat treatment temperature increases, the maximum indentation depth gradually increases in the axial direction except for the carbon fibers heat-treated at 2300 C and 2700 C and the maximum indentation depth is about 389 ± 12 nm. The as-received carbon fibers exhibit the highest axial EIT-A, which decreases from 51.3 ± 1.4 GPa to 40.2 ± 1.0 GPa with increasing indentation depth (see Fig. 12(b)). As the heat treatment temperature increases, the axial EIT-A of the heat-treated carbon fibers decreases, suggesting that the elasticity of the carbon fibers is enhanced. The lowest axial EIT-A of the carbon fibers ranges from 33.6 ± 1.3 GPa to 22.7 ± 1.5 GPa after heat treatment at 2800 C.
Fig. 12. Indentation load-depth and elastic modulus of the as-received and heat-treated PAN-based carbon fibers: (a) and (b) axial direction, (c) and (d) transverse direction. (A colour version of this figure can be viewed online.)
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F. Yang et al. / Carbon xxx (xxxx) xxx Table 2 Decline rates of the elastic modulus of the as-received and heat-treated PAN-based carbon fibers in the axial and transverse directions. Heat treatment temperature (oC)
Axial direction (102 GPa/nm)
Transverse direction (102 GPa/nm)
As-received 2000 2100 2300 2500 2700 2800
6.62 7.20 4.90 4.99 5.15 3.46 3.79
2.01 1.48 1.35 0.84 1.05 0.70 0.79
In the transverse direction, as the heat treatment temperature increases, the maximum indentation depth shows an upward trend and the maximum indentation depth is about 725 ± 7 nm which is nearly twice than that in the axial direction (see Fig. 12(c)). The transverse EIT-T of the carbon fibers decreases with increasing heat treatment temperature and indentation depth. The range of the transverse EIT-T of the as-received carbon fibers is from 16.9 ± 0.1 GPa to 9.6 ± 0.2 GPa, which is about one third of the axial EIT-A (see Fig. 12(d)). The transverse EIT-T of the carbon fibers decreases to 8.6 ± 0.1e4.2 ± 0.1 GPa when the heat treatment temperature increases to 2800 C. Based on the microstructure discussed above, it could be inferred that the decrease in elastic modulus along the indentation depth is closely related to the microstructure evolution. In order to clarify the relation, the decline rate is calculated based on the ratio of the change in elastic modulus to the corresponding change in indentation depth and the results are shown in Table 2. The accumulating damage in the form of cracks is an important reason for the decrease in elastic modulus along the indentation depth [22]. Because the cracks are easy to propagate along the graphite-like crystallites under the axial compression, the decline rates of the elastic modulus in the axial direction are larger than those in the transverse direction (see Table 2). As the heat treatment
temperature increases, the growth and rearrangement of the graphene layers would increase the size and order of the graphite-like crystallites, which could prevent the crack propagation effectively. Furthermore, the needle-like pores could also prevent the crack propagation of the nanoindentation in the transverse direction. It is clear that the microstructure evolution during heat treatments would reduce the effect of accumulating damage on the elastic modulus. Thus, the decline rates of the elastic modulus decrease after heat treatments and the carbon fibers heat-treated at 2700 C exhibit the minimum values. To explain the relationship between the microstructure evolution and the elastoplastic properties of PAN-based carbon fibers, a schematic mechanism of elastoplastic indentation onto as-received and heat-treated PAN-based carbon fibers at an early shallow stage is shown in Fig. 13. When the indenter penetrates the carbon fiber surfaces, the deformation of graphite-like crystallites and crosslinking regions between the graphite-like ribbons is vital to reveal the elastoplastic mechanism of PAN-based carbon fibers. In the as-received carbon fibers, the size of the graphite-like crystallites is small, and the amorphous carbon and turbostratic graphene layers are cross-linked by covalent bonds (see Fig. 13(a)), meaning that the breakage of the graphite-like crystallites and the crosslinking regions is difficult. Moreover, the slip of the turbostratic
Fig. 13. Schematic mechanism of elastoplastic indentation onto as-received and heat-treated PAN-based carbon fibers at an early shallow stage: (a) as-received and (b) 2800 C. (A colour version of this figure can be viewed online.)
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F. Yang et al. / Carbon xxx (xxxx) xxx
graphene layers is almost impossible. Thus, the as-received carbon fibers exhibit the highest axial EIT-A. After heat treatment at 2800 C, the ductility index D indicates that the ‘plasticity’ has been enhanced and the reasons may be explained as follows. First, the graphene layers would easily slip under compressive loads due to low energy of van der Waals bonds (see Fig. 13(b)). With increasing heat treatment temperature, the ordered stacking of the graphene layers means a significant slip of the graphene layers. Second, the crack would occur at the edges of the needle-like pores and propagate along the graphite-like crystallites, leading to the split of the crystallites. Finally, the graphite-like crystallites and the crosslinking regions would be easy to break due to the aggregation of the defects. Consequently, as the heat treatment temperature increases, the yield stress of the graphite-like crystallites and the cross-linking regions would decrease, leading to an improvement in ‘plasticity’ and a decrease in axial EIT-A. By contrast, in the transverse direction, because the turbostratic graphene layers are easy to bend and compress (see Fig. 13(a)), the transverse EIT-T is much lower than the axial EIT-A and the carbon fibers exhibit almost complete elasticity. When the turbostratic graphene layers transform into curved and continuous graphene layers, the bending deformation of the layers becomes easy and the needle-like pores also can provide the space for the elastic deformation (see Fig. 13(a)). Therefore, the transverse EIT-T gradually decreases with increasing heat treatment temperature. Moreover, the breakage of the graphite-like crystallites and the cross-linking regions is the main reason for the slight ‘plasticity’. 4. Conclusions (1) As the heat treatment temperature increases, the turbostratic graphene layers in the sub-surface region transform into an ordered graphite-like structure, leading to formation of many bulges on the carbon fiber surfaces. The average length-to-width ratio of the bulges increases from 3.02 to 5.10, owing to the growth and rearrangement of the graphene layers in the sub-surface region. (2) The shrinkage of the carbon fiber surfaces would constrain the growth and rearrangement of the graphene layers in the sub-surface region. The graphene layers are closely stacked and curved along the carbon fiber surface in the sub-surface region. In the longitudinal section, the disordered graphene layers have become ordered graphite-like ribbons whose continuity in the sub-surface region is worse than that in the skin region. The cross-linking regions between the graphitelike ribbons in the longitudinal direction still consist of a lot of kinked graphene layers, structural defects and needle-like pores after heat treatments. (3) The elastic modulus of PAN-based carbon fibers in the axial and transverse directions decreases with increasing indentation depth. The needle-like pores and the ordered graphite-like crystallites could prevent the crack propagation, leading to decrease in the decline rate of the elastic modulus along the indentation depth. Because of the improvement of the slip between the graphene layers and the breakage of the graphite-like crystallites and the crosslinking regions, the axial elastic modulus decreases with increasing heat treatment temperature. In the transverse direction, the carbon fibers exhibit almost complete elasticity, which strongly depends on the curved graphene layers and the needle-like pores. Declaration of competing interests The authors declare that they have no known competing
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