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Toughening efficiency and mechanism of carbon fibre epoxy matrix composites by PEK-C
Composite Structures 229 (2019) 111431
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Toughening efficiency and mechanism of carbon fibre epoxy matrix composites by PEK-C
T
Jiawei Yaoa,b, Kangmin Niua, , Yifan Niub, Teng Zhangb ⁎
a b
School of Materials Science and Engineering, University of Science & Technology Beijing, Beijing 100083, China Sino-European Institute of Aviation Engineering, Civil Aviation University of China, Tianjin 300300, China
ARTICLE INFO
ABSTRACT
Keywords: Polymer-matrix composites (PMCs) Fracture toughness Interleaving Atomic force microscopy (AFM) Phase structure
Polyaryletherketone with Cardo (PEK-C) was used as matrix additive, film and particle interleaves to improve the fracture toughness of carbon fibre epoxy matrix composites. The toughening efficiency and mechanism of the three methods were studied. The PEK-C film interlayer was more efficient and stable in the interlaminar toughening. The flexural strength was not compromised with the incorporation of PEK-C due to the limited addition. The enhanced interlaminar toughening was mainly attributed to the two-phase structure of micrometer scale formed by PEK-C “peel” and the rounded epoxy matrix wrapped by PEK-C. The nanoscale sea-island substructure in the rounded epoxy matrix was observed by atomic force microscopy (AFM), which might be also correlated with the toughness improvement.
1. Introduction Carbon fibre reinforce plastics (CFRPs) have been extensively employed in a wide range of industries, especially in aeronautical industry due to their outstanding specific mechanical properties. Thermoset epoxies exhibit many advantages such as good mechanical properties and high heat resistance, and are widely used as matrices for CFRPs. However, the high degree of cross-linking density after curing results in their inherent brittleness and low fracture toughness. In addition, weak out-of-plane mechanical properties due to the anisotropic nature of unidirectional CFRPs results in lower impact resistance [1,2]. These weaknesses may lead to the ultimate failure during the service and the limitation of their applications. The incorporation of thermoplastic resin into CFRPs has been reported to be a prevalent method of improving the fracture toughness and impact resistance [3–10]. There are two ways to do that, method one is by the dispersion of thermoplastic resin, such as polyaryletherketone (PEK)[4,5], polyethersulfone (PES) [6,7], polyetheretherketone (PEEK) [8,9], or Polysulfone (PSF) [10,11] in epoxy matrix. The related researches began in the 1990s and demonstrated that adding of thermoplastic resin could enhance fracture toughness. Another newer method is interleaving, which can be described by adding discrete layers of thermoplastic resin in the form of particles [12,13], films [14,15] and fibres [16–18]. Gao et al. [13] evaluated thermoplastic particle interlayer toughening effect properties of three
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types of CFRPs. Their results showed that the interlayer zones formed by particles and matrix could significantly suppress delamination and mode II delamination resistance were significantly enhanced. In the studies of Yao et al. [15], thermoplastic resin films of three different thickness were used to interleave CFRPs. The results indicated that GIC and interlaminar shear strength increased maximally at a film thickness of 10 μm due to the fact that PEK-C and epoxy resin formed a dualphase structure during the hot pressing curing process. However, the flexural strength and flexural modulus decreased with the increase of the film thickness. Zhang et al. [16] studied the influences of PEK-C nanofibre diameter and interlayer thickness on the properties of CFRPs and results showed that interlaminar fracture toughness has been developed by interleaving PEK-C nanofibres with the weight loading as low as 0.4%. Finer nanofibres resulted in more stable crack propagation and better mechanical performance under flexure loading. In brief, these methods have been widely proven to be effective to improve the fracture toughness. Nevertheless, these studies were conducted on different material systems and it is still confused about the toughening efficiency and mechanism of these methods. In the latest researches, nano-particles, such as graphene oxide particles or carbon nanotubes, have been used to be dispersed in epoxy matrix or to be added into interlayers, which could achieve a significant rise in toughness by a quite small amount (< 1%) [19–22]. Compared to these fillers, thermoplastic resin is cheaper and easier to manufacture for industries. Therefore, in the present work, the thermoplastic resin
Corresponding author. E-mail address: [email protected] (K. Niu).
https://doi.org/10.1016/j.compstruct.2019.111431 Received 4 April 2019; Received in revised form 15 July 2019; Accepted 11 September 2019 Available online 13 September 2019 0263-8223/ © 2019 Elsevier Ltd. All rights reserved.
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Table 1 Amount of PEK-C and thickness of samples. CFRP samples
Control Film interleaved Particle interleaved Matrix toughened
Fig. 1. Schematic of fabrication process of prepregs.
PEK-C was blended in epoxy resin or prepared as interlayers in form of particle and film to enhance the fracture toughness. Such modified epoxy resin or interlayers were used to manufacture CFRPs. The effects of different toughening methods on the fracture toughness, impact resistance and mechanical properties of CFRPs were investigated.
Amount of PEK-C 0 0.33 mg/cm2 0.33 mg/cm2 The total amount is equal to that of film interleaved samples.
Thickness of samples (mm) Bending 2.11 2.07 2.37 2.21
± ± ± ±
0.03 0.03 0.10 0.03
Impact 2.96 2.96 3.03 3.05
± ± ± ±
0.14 0.06 0.17 0.14
Mode I 3.17 3.16 3.25 3.26
± ± ± ±
0.07 0.06 0.11 0.10
amount of epoxy resin was dispersed into the DCM by a mechanical mixer, until completely dissolved. The unidirectional carbon fibres T700 were immersed into the epoxy resin solution and pulled out at a constant speed, then placed at room temperature for 24 h to evaporate DCM. The prepregs without PEK-C were fabricated for control and interleaved CFRP samples. The prepregs with PEK-C added into epoxy resin were also fabricated for epoxy matrix toughened CFRP samples.
2. Experimental work 2.1. Processing of composites 2.1.1. Materials The PEK-C powders used in this study were produced by Xuzhou Engineering Plastics Factory with the particle size of about 180 μm. Tetrahydrofuran (THF) and Dichloromethane (DCM) purchased from Tianjin Kemiou Chemical Reagent Co., Ltd were used to dissolve PEK-C and epoxy resin respectively. The epoxy resin (Hansort® 6240) was provided by Hansort Advanced Materials Co., LTD. The glass transition temperature was 130 °C~140 °C. The curing temperature recommended by producer was 120 °C~130 °C. The unidirectional T700 carbon fibre fabrics were supplied by Yixing Huifeng Carbon Fibre Technology Co., LTD.
2.1.2.2. Preparation of CFRP composites. The prepregs prepared by the above method were cut into small pieces. The total number of plies and fibre orientation were set according to the test standards for the control and toughened CFRP samples, as depicted in Fig. 2. For mode I interlaminar fracture toughness test, the laminates [0]18 were fabricated with unidirectional T700/epoxy prepregs. The release film (PTFE) was added at the mid-plane to form the pre-crack of 65 mm. The film interlayer was directly placed at the mid-plane for the film interleaving laminates. The PEK-C particles were scattered using a sieve at the mid-plane for the particle interleaving laminates. For bending test, 12 plies of unidirectional prepregs were stacked with interlayers placed between each ply. For the low velocity impact test, the laminates [+45/0/−45/90]2s were fabricated with interlayers placed between each ply. For all the tests, control samples and epoxy matrix toughened samples were also fabricated. The stacked laminates were cured in an autoclave and the pressure, temperature and holding time were set up to 0.6 MPa, 120 °C and 120 min, respectively. Based on the areal density of PEK-C film, a careful calculation was conducted to ensure the equal amount of PEK-C added in CFRPs. The added amount of PEK-C and thickness of samples after curing for different tests were displayed in Table 1.
2.1.2. Preparation of interlayer films PEK-C powders were dissolved in THF at the weight fraction of 10% at room temperature. PEK-C film was formed by dip-coating method on top of glass substrate at a defined dragging rate of 0.2 cm/s and was then placed in the oven for 20 min to evaporate THF. The film thickness was restricted to 10 μm to avoid the decrease about in-plane properties caused by excessive amount of thermoplastic resin added [15]. The thickness of film obtained was 8 ~ 10 μm measured by micrometers and the areal density of PEK-C film was averagely 0.33 mg/cm2 calculated by dividing the masse by the area of film. 2.1.2.1. Impregnation of carbon fibres with epoxy resin. A schematic of the fabrication process of prepregs was shown in Fig. 1. A calculated
Fig. 2. Preparation process of CFRP composites.
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modulus and strength of the CFRPs according to the ISO14125-1998 test standard [26]. The sample size was 100 mm × 15 mm. The test were conducted at room temperature with a constant displacement rate of 2 mm/min. Five samples were tested for each value. Scanning electron microscope (SEM, Hitachi S-3400N) was applied to investigate the fracture and lateral surfaces of samples after mode I interlaminar fracture toughness tests and the lateral surfaces of samples after bending tests. The accelerating voltage was set as 15 kV. Gold layer was sputter-coated on the surface before tests. The phase images were studied by atomic force microscope (AFM) tapping mode. Fig. 4 gives the schematic setup of phase imaging by AFM tapping mode [15,27]. The cantilever was modulated by a sinusoidal wave and phase contrast caused by differences in surface adhesion and viscoelasticity was monitored by the detector. Silicon tapping cantilever (OTESPA) was used with a radius of curvature of about 2 nm. Tests were performed by Bruker Dimension Icon AFM. The amplitude setpoint was chosen as 40 mV. Height and phase images were recorded simultaneously. Each test was performed at room temperature and repeated five or more times.
Fig. 3. Experimental set-up for mode I interlaminar fracture toughness test.
2.2. Characterization and testing
3. Results and discussion
Mode I interlaminar facture toughness, GIC were measured on a universal testing machine (Instron 5982) according to the ASTM D5528-01 [23]. The loadcell of 5 kN was used. The data acquisition frequency was set as the default setting (10 Hz). The experimental device was shown in Fig. 3. The size of samples was 125 mm × 20 mm. The loading rate was 2 mm/min. The modified beam theory (MBT) method was used to calculate GIC.:
GIC =
3.1. Mode I interlaminar fracture toughness R-curves of all the CFRPs were shown in Fig. 5. R-curve displayed the relation of GIC with the length of crack, which characterized the initiation and propagation of a delamination. From the R-curve, it was obvious that fracture toughness increased with the crack length and eventually reached the plateau. The rise behavior could be explained by the fibre bridging mechanism. It was clear that higher crack growth resistance was observed in toughened composites. Up to crack length of 70 mm, there was no large distinction between control and toughened composites. However, the contribution of PEK-C became apparent after crack length of 70 mm. The initial GIC value of control samples was only 102.5 J/m2 due to its bad interfacial adhesion, while that of matrix toughened, particle interleaved and film interleaved samples were
3P 2b (a + | |)
where P is the load, δ is load point displacement, b is the specimen width, a is the delamination length, and Δ is the corrective factor for crack length calculated for each test based on standard ASTM D552801. Five samples were tested for each value. Low velocity Impact testing was conducted on an impact testing machine (CEAST 9350) according to ASTM D7136/D7136M standard [24]. The loadcell of 45 kN and DAQ systems (CEAST DAS 64 k) were used. The sampling frequency was 1000 kHz. Four rubber tips linked with toggle clamps were used to restrain the corners of specimen during impact testing. The specimen size was 150 mm × 100 mm. The impact energy of the CFRP laminates was set as E = 19 J. Five tests were conducted for each value. Based on JB/T4008-1999 standard [25], ultrasonic inspection (Cscan, C505030) was used to identify the damage, such as delamination, fibre breakage, local punctures or indentation caused by impact. Three-point bending test was used to determine the bending
Fig. 5. Mode I interlaminar fracture toughness of all the CFRPs.
Fig. 4. Schematic setup of phase imaging by AFM tapping mode.
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Fig. 6. Lateral surface images of all the CFRPs: (a) control sample; (b) matrix toughened samples; (c), (d) particle interleaved samples; (e), (f) film interleaved samples.
increased by 64.2%, 82.3% and 71.6% respectively. The film interleaved composites showed high and stable GIC value along the crack propagation. The highest GIC was obtained in film interleaved composites, revealing 472.78 J/m2 at crack length of 85 mm. A large variation of fracture toughness in particle interleaved composite could be attributed to uneven distribution of PEK-C particles on the mid-plane caused by scattering using a sieve by hand. The lateral surfaces of samples after mode I interlaminar fracture toughness were investigated by SEM to track the crack propagation path. Fig. 6(a) showed the magnification of crack route for the samples without toughening, from which, the failure mechanism of fibre bridging was observed and matrix damage in the neighbouring layer was rare. Compared with Fig. 6(a), the crack propagation moved into the neighbouring layer irregularly and damage of matrix was observed for the toughened composites (Fig. 6(b)–(f)). Fibre strip breakage and delamination could be particularly observed for the particle and film interleaved composites in Fig. 6(c) and (e). It should be noted that the
delamination phenomenon in the neighbouring layer of the film interleaved samples was extraordinarily obvious. With an expectation to further investigate the improvement mechanism of mode I interlaminar fracture toughness, the fracture surfaces were examined. Fig. 7 showed fracture surface images of control and toughened composites. The epoxy matrix of control samples was thin in the interlaminar region. The failure mechanism was mainly the matrix breakage and the fibre-matrix debonding. The matrix between fibres exhibited brittle fracture features. The smooth fracture surfaces implied low resistance to crack propagation. (Fig. 7(a, b)). After adding PEK-C into epoxy matrix, the failure mechanism was consistent with that of control samples (Fig. 7(c)). However, it was obvious that fracture surface was transformed from brittle features to ductile features. Submitted to a curing process, the second phase PEK-C and epoxy matrix interdiffused due to the enhanced molecular chain under high temperature and then the phase separation began while the epoxy resin crosslinked to form the network structure [28,29]. The phase
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Fig. 7. Fracture surface images of all the CFRPs: (a), (b) control samples; (c)-(e) matrix toughened samples; (f)-(h) particle interleaved samples; (i), (j) film interleaved samples.
Fig. 8. The force versus time curves for all the specimens: (a) force versus time curves; (b) the smoothed curves.
separation phenomena appeared but was rare in the interlaminar region for matrix toughened samples (Fig. 7(d)). Fig. 7(e) displayed the representative two-phase structure formed by epoxy resin and PEK-C [29,30]. The two-phase structure was composed of the PEK-C “peel” and the epoxy matrix wrapped by the “peel”. When the crack extended to the two-phase structure, the epoxy was pulled out from the “peel” and the smooth fracture surface revealed the brittle fracture. What’s more, the “peel” was pulled during the crack growth and exhibited
Table 2 Fmax, DFmax, Elastic energy for all the samples. Fmax (N) Control Resin toughened Particle interleaved Film interleaved
5512 5942 5871 5961
± ± ± ±
32 85 64 22
DFmax (mm) 2.88 3.18 3.14 3.06
± ± ± ±
0.12 0.08 0.17 0.23
Elastic energy (J) 7.0 8.5 7.7 8.0
± ± ± ±
0.6 0.7 0.5 0.3
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Fig. 11. Bending properties of all the samples.
lead to the deviation of cracks and the formation of the secondary cracks in the interlaminar region, even the delamination of matrix in the weak region and the debonding between fibres and matrix due to weak adhesion of neighbouring layers, which could dissipate more energy and enhance further the interlaminar fracture toughness of composites.
Fig. 9. Energy versus time curves of all the samples.
superior ability of ductile deformation. The ductile deformation and pulled-out of epoxy toughened by PEK-C consumed more energy and contributed to the improvement of fracture toughness. The clean fibre surface of control and matrix toughened samples were observed, demonstrating weak interfacial adhesion between the matrix and the fibres (Fig. 7(a–c)). The fracture surfaces of particle interleaving composites were displayed in Fig. 7(f–h). Fig. 7(f) showed a rough surface and the interlaminar region was covered with interlayer, which revealing the high resistance to the crack propagation. The phase separation of PEK-C and epoxy matrix was also observed. Cavities left by the pulled-out epoxy and plastic deformation caused by PEK-C were clearly displayed in Fig. 7(g, h). However, the thickness difference of the PEK-C “peel” was found between Fig. 7(g) and (h) and the “peel” was thicker in Fig. 7(h). The thicker “peel” was easily formed in the region of more PEK-C resin [29,30]. Therefore, the different thickness was attributed to the uneven distribution of PEK-C particles. What’s more, the difference of thickness made clear the fluctuation of GIC with the length of crack for particle interleaved laminates found in the Section 3.1. Similar toughening features were observed in film interleaved samples (Fig. 7(i, j)), but were less. This could be because PEK-C film was thin (< 10 μm) and uniform distribution, resulting in the easier diffusion of PEK-C from interlaminar region into neighbouring layers. During the crack propagation in the mode I interlaminar fracture toughness test, in case that the crack moved to the area of the twophase structure, higher energy dissipation and fracture toughness might
3.2. Low velocity impact The force versus time curves for all the specimens at impact energy of 19 J were shown in Fig. 8. The first phase (Between [0–2.5 ms]) was similar, revealing that the PEK-C addition affected less the elastic behavior of specimens. The maximum load (Fmax) and the displacement corresponding to the maximum load (DFmax) obtained from the testing could be used to evaluate the impact resistance [31]. For the toughened laminate composites, the improvement of impact resistance was seen in terms of the higher Fmax and higher DFmax. Fmax and DFmax increased from 5512 N and 2.88 mm for the control specimen to 5942 N and 3.18 mm, to 5874 N and 3.14 mm and to 5961 N and 3.06 mm, respectively, shown in Table 2. What’s more, the significant drop of load after Fmax for the control specimen in Fig. 8(b) indicated that damage propagation in the untoughened composite plate was more severe. Composites toughened by three different methods revealed almost the same performance in terms of impact resistance due to the same added amount. Fig. 9 showed the energy versus time curves of all the samples. The maximum value of each curve represents the associated impact energies. The constant value reached at the final phase of impact event gives the total energy absorbed by composited specimens. The
Fig. 10. C-scan results of delaminated areas of all the CFRPs: (a) control sample; (b) matrix toughened samples; (c) particle interleaved samples; (d) film interleaved samples.
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difference between them is considered as elastic energy, retained in the impactor and used to rebouned the impactor from the nonperforated samples [31]. From Fig. 9 and Table 2 it could be seen that the elastic energy in toughened samples was slightly larger, showing a higher impact resistance. Fig. 10 showed C-scan results of delaminated areas of all the CFRPs. The control CFRP laminate showed a damaged area of 281.28 mm2 calculated by gray-scale value based on MATLAB. With adding PEK-C to the CFRP, the damage area was reduced. A drop of damage area by 28.9% was observed for particle interleaved composites. PEK-C and two-phase structure formed between PEK-C and epoxy resin could absorb the impact energy and further restrain the crack growth in matrix and between the layers.
3.3. Three-point bending Fig. 11 showed the bending properties of all the CFRPs. The results showed that the flexural strength was not compromised with the incorporation of PEK-C in the CFRP composites due to the limited addition of PEK-C [15]. The bending strength were 1141, 1157, 1125 and 1099 MPa for control, matrix toughened, particle interleaved and film interleaved samples respectively. The modulus was decreased by 16% for film interleaved composites due to the inherent lower rigidness of thermoplastic resin than epoxy matrix, which could weaken the rigidness of interleaved laminates. Nonetheless, a distinct increase by 38% in particle interleaved samples was observed. The flexural modulus was sensitive to sample thickness. The thickness of particle interleaved
Fig. 12. The lateral faces of all the samples after bending tests: (a), (b) control samples; (c), (d) matrix toughened samples; (e)-(h) particle interleaved samples; (i)-(k) film interleaved samples.
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3.4. Phase image by AFM tapping mode The phase images of all the samples scanned by AFM Tapping mode were displayed in Fig. 13. Fig. 13(a) showed the profile in epoxy matrix for laminates without toughening and obviously there exist only one phase. The two-phase morphology of sea-island was observed for toughened composites (Fig. 13(b)–(d)). The island of PEK-C was distributed in the epoxy matrix. Submitted to the curing process, PEK-C was firstly dissolved in epoxy matrix and then the phase separation occurred while the epoxy resin began crosslinking to form the network structure. Compared to the SEM images of two-phase structure in Fig. 7, the scanned results by AFM tapping mode were different. In Fig. 7, the epoxy matrix was segregated with PEK-C “peel” and the size of inner region was about 1–2 μm, or even larger. Fig. 13 revealed that there existed the nanoscale sea-island substructure, from which, it could be speculated that the island phase of PEK-C was formed in the round epoxy matrix wrapped by PEK-C “peel”. The formation of phase structure and the morphology as described were summarized and sketched schematically in Fig. 14. During the curing process, the kinetic ability of molecular chain of epoxy resin was much more enhanced than that of PEK-C resin at 120 °C due to the obvious difference in glass transition temperature of epoxy resin (about 130 °C) and PEK-C resin (above 230 °C). A large amount of epoxy resin diffused into PEK-C resin as shown in Fig. 14(b). The representative two-phase structures of micrometer scale were formed and composed of PEK-C “peel” and the inner region of epoxy resin, which were scattered in the epoxy-rich matrix. The limited amount of PEK-C diffused, resulting in the distribution of nanoscale island phase of PEK-C in the sea phase of epoxy matrix shown in Fig. 14(c). It was estimated that the enhanced interlaminar toughening was mainly attributed to the representative twophase structures of micrometer scale; in addition, it might be attributed to the sea-island substructure of nanometer scale, which need more research to be done.
Fig. 13. The phase images of all the samples scanned by AFM Tapping mode: (a) control sample; (b) matrix toughened samples; (c) particle interleaved samples; (d) film interleaved samples.
samples was changed to 2.37 mm from 2.11 mm for control samples. The rise in modulus could attribute to the augmentation of sample thickness caused by too large PEK-C particle size (about 180 μm provided by producer). The lateral faces of samples after bending tests were analyzed using SEM images as displayed in Fig. 12. It was obvious that the failures were caused by fibre breakage and crack propagation in the interlaminar region. The cracks were smooth, indicating the low crack propagation resistance for samples without toughening. The large and long cracks were formed due to the weak interfacial adhesion, which leaded to the shortest path of fibre breakage for samples without toughening. The neat fibre surface demonstrated weak interfacial adhesion between the matrix and the fibres, which was coincide with that of samples for mode I interlaminar fracture toughness (Fig. 12(a) and (b)). With toughening, the resin “burrs” after plastic deformation caused generally by shear stress in the regions of delamination were observed (Fig. 12(d), (g), and (j)). Fig. 12(f) showed that the fibres and matrix of particle interleaved composites in the inner ply was consistent with that of composites without toughening and the fracture surface displayed the brittle features. After interleaving, the deviation of crack propagation path appeared in the interlaminar region. The crack path from the interleaf to the neighbouring ply indeed increased the fracture toughness (Fig. 12(h) and (k)).
4. Conclusion The toughening efficiency and mechanism of the three methods, including the matrix toughed by PEK-C, PEK-C film interleaving and particle interleaving, were studied in this work. High fracture toughness and impact resistance were achieved for toughened CFRP composites. The PEK-C film interlayer was more efficient and stable in the fracture toughness than the others. The flexural strength was not compromised with the incorporation of PEK-C due to the limited addition of PEK-C. The enhanced toughening was mainly attributed to the two-phase structure of micrometer scale formed by epoxy matrix and PEK-C resin. There existed nanoscale sea-island substructure according to the AFM scanning results, which might be also correlated with the toughness
Fig. 14. Schematic of formation and morphology of phase structure of PEK-C and epoxy matrix.
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improvement. Characterizing the substructure was difficult due to the nanometer size and more efforts should be taken to further develop it.
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Acknowledgement The authors would acknowledge financial support from the Fundamental Research Funds for the Central Universities (Grant No. 3122018Z007). Data availability The processed data required to reproduce these findings are available to download from [DOI: 10.17632/8mdjrmdsfc.1]. References [1] Cartie DDR, Irving PE. Effect of resin and fibre properties on impact and compression after impact performance of CFRP. Compos Part A Appl Sci Manuf 2002;33:483–93. [2] Cantwell WJ, Morton J. The impact resistance of composite materials-a review. Composites 1991;22:347–62. [3] Shivakumar K, Panduranga R. Interleaved polymer matrix composites-a review. AIAA/ASME/ASCE/AHS/ASC Struct, Struct Dyn, Mater Conf 2013. [4] Yi XS, An XF. Effect of interleaf sequence on impact damage and residual strength in a graphite/epoxy laminate. J Mater Sci 2002;39:3253–5. [5] Zhang J, Niu H, Zhou J, Wang XG, Lin T. Synergistic effects of PEK-C/VGCNF composite nanofibres on a trifunctional epoxy resin. Compos Sci Technol 2011;71:1060–7. [6] Yang JP, Chen ZK, Feng QP, Deng YH, Liu Y, Ni QQ, et al. Cryogenic mechanical behaviors of carbon nanotube reinforced composites based on modified epoxy by poly(ethersulfone). Compos Part B Eng 2012;43:22–6. [7] Lee SE, Jeong E, Lee MY, Lee MK, Lee YS. Improvement of the mechanical and thermal properties of polyethersulfone-modified epoxy composites. J Ind Eng Chem 2016;33:73–9. [8] Francis B, Thomas S, Jose J, Ramaswamy V, Lakshmana R. Hydroxyl terminated poly (ether ether ketone) with pendent methyl group toughened epoxy resin: miscibility, morphology and mechanical properties. Polymer 2005;46:12372–85. [9] Francis B, Thomas S, Asari GV, Ramaswamy RS, Seno J, Rao VL. Synthesis of hydroxyl-terminated poly (ether ether ketone) with pendent tert-butyl groups and its use as a toughener for epoxy resins. J Polym Sci Pol Phys 2006;44:541–56. [10] Gyun YN, Gu WY, Chul KS. Toughening of epoxy composite by dispersing polysulfone particle to form morphology spectrum. Polym Bull 2005;52:365–72. [11] Li G, Li P, Yu Y, Jia XL, Zhang S, Yang XP, et al. Novel carbon fibre/epoxy composite toughened by electrospun polysulfone nanofibres. Mater Lett 2008;62:511–4. [12] Gilbert EN, Hayes BS, Seferis JC. Interlayer toughened unidirectional carbon prepreg system: effect of preformed particle morphology. Compos Part A Appl Sci Manuf 2006;34:245–52.
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Toughening efficiency and mechanism of carbon fibre epoxy matrix composites by PEK-C
T
Jiawei Yaoa,b, Kangmin Niua, , Yifan Niub, Teng Zhangb ⁎
a b
School of Materials Science and Engineering, University of Science & Technology Beijing, Beijing 100083, China Sino-European Institute of Aviation Engineering, Civil Aviation University of China, Tianjin 300300, China
ARTICLE INFO
ABSTRACT
Keywords: Polymer-matrix composites (PMCs) Fracture toughness Interleaving Atomic force microscopy (AFM) Phase structure
Polyaryletherketone with Cardo (PEK-C) was used as matrix additive, film and particle interleaves to improve the fracture toughness of carbon fibre epoxy matrix composites. The toughening efficiency and mechanism of the three methods were studied. The PEK-C film interlayer was more efficient and stable in the interlaminar toughening. The flexural strength was not compromised with the incorporation of PEK-C due to the limited addition. The enhanced interlaminar toughening was mainly attributed to the two-phase structure of micrometer scale formed by PEK-C “peel” and the rounded epoxy matrix wrapped by PEK-C. The nanoscale sea-island substructure in the rounded epoxy matrix was observed by atomic force microscopy (AFM), which might be also correlated with the toughness improvement.
1. Introduction Carbon fibre reinforce plastics (CFRPs) have been extensively employed in a wide range of industries, especially in aeronautical industry due to their outstanding specific mechanical properties. Thermoset epoxies exhibit many advantages such as good mechanical properties and high heat resistance, and are widely used as matrices for CFRPs. However, the high degree of cross-linking density after curing results in their inherent brittleness and low fracture toughness. In addition, weak out-of-plane mechanical properties due to the anisotropic nature of unidirectional CFRPs results in lower impact resistance [1,2]. These weaknesses may lead to the ultimate failure during the service and the limitation of their applications. The incorporation of thermoplastic resin into CFRPs has been reported to be a prevalent method of improving the fracture toughness and impact resistance [3–10]. There are two ways to do that, method one is by the dispersion of thermoplastic resin, such as polyaryletherketone (PEK)[4,5], polyethersulfone (PES) [6,7], polyetheretherketone (PEEK) [8,9], or Polysulfone (PSF) [10,11] in epoxy matrix. The related researches began in the 1990s and demonstrated that adding of thermoplastic resin could enhance fracture toughness. Another newer method is interleaving, which can be described by adding discrete layers of thermoplastic resin in the form of particles [12,13], films [14,15] and fibres [16–18]. Gao et al. [13] evaluated thermoplastic particle interlayer toughening effect properties of three
⁎
types of CFRPs. Their results showed that the interlayer zones formed by particles and matrix could significantly suppress delamination and mode II delamination resistance were significantly enhanced. In the studies of Yao et al. [15], thermoplastic resin films of three different thickness were used to interleave CFRPs. The results indicated that GIC and interlaminar shear strength increased maximally at a film thickness of 10 μm due to the fact that PEK-C and epoxy resin formed a dualphase structure during the hot pressing curing process. However, the flexural strength and flexural modulus decreased with the increase of the film thickness. Zhang et al. [16] studied the influences of PEK-C nanofibre diameter and interlayer thickness on the properties of CFRPs and results showed that interlaminar fracture toughness has been developed by interleaving PEK-C nanofibres with the weight loading as low as 0.4%. Finer nanofibres resulted in more stable crack propagation and better mechanical performance under flexure loading. In brief, these methods have been widely proven to be effective to improve the fracture toughness. Nevertheless, these studies were conducted on different material systems and it is still confused about the toughening efficiency and mechanism of these methods. In the latest researches, nano-particles, such as graphene oxide particles or carbon nanotubes, have been used to be dispersed in epoxy matrix or to be added into interlayers, which could achieve a significant rise in toughness by a quite small amount (< 1%) [19–22]. Compared to these fillers, thermoplastic resin is cheaper and easier to manufacture for industries. Therefore, in the present work, the thermoplastic resin
Corresponding author. E-mail address: [email protected] (K. Niu).
https://doi.org/10.1016/j.compstruct.2019.111431 Received 4 April 2019; Received in revised form 15 July 2019; Accepted 11 September 2019 Available online 13 September 2019 0263-8223/ © 2019 Elsevier Ltd. All rights reserved.
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Table 1 Amount of PEK-C and thickness of samples. CFRP samples
Control Film interleaved Particle interleaved Matrix toughened
Fig. 1. Schematic of fabrication process of prepregs.
PEK-C was blended in epoxy resin or prepared as interlayers in form of particle and film to enhance the fracture toughness. Such modified epoxy resin or interlayers were used to manufacture CFRPs. The effects of different toughening methods on the fracture toughness, impact resistance and mechanical properties of CFRPs were investigated.
Amount of PEK-C 0 0.33 mg/cm2 0.33 mg/cm2 The total amount is equal to that of film interleaved samples.
Thickness of samples (mm) Bending 2.11 2.07 2.37 2.21
± ± ± ±
0.03 0.03 0.10 0.03
Impact 2.96 2.96 3.03 3.05
± ± ± ±
0.14 0.06 0.17 0.14
Mode I 3.17 3.16 3.25 3.26
± ± ± ±
0.07 0.06 0.11 0.10
amount of epoxy resin was dispersed into the DCM by a mechanical mixer, until completely dissolved. The unidirectional carbon fibres T700 were immersed into the epoxy resin solution and pulled out at a constant speed, then placed at room temperature for 24 h to evaporate DCM. The prepregs without PEK-C were fabricated for control and interleaved CFRP samples. The prepregs with PEK-C added into epoxy resin were also fabricated for epoxy matrix toughened CFRP samples.
2. Experimental work 2.1. Processing of composites 2.1.1. Materials The PEK-C powders used in this study were produced by Xuzhou Engineering Plastics Factory with the particle size of about 180 μm. Tetrahydrofuran (THF) and Dichloromethane (DCM) purchased from Tianjin Kemiou Chemical Reagent Co., Ltd were used to dissolve PEK-C and epoxy resin respectively. The epoxy resin (Hansort® 6240) was provided by Hansort Advanced Materials Co., LTD. The glass transition temperature was 130 °C~140 °C. The curing temperature recommended by producer was 120 °C~130 °C. The unidirectional T700 carbon fibre fabrics were supplied by Yixing Huifeng Carbon Fibre Technology Co., LTD.
2.1.2.2. Preparation of CFRP composites. The prepregs prepared by the above method were cut into small pieces. The total number of plies and fibre orientation were set according to the test standards for the control and toughened CFRP samples, as depicted in Fig. 2. For mode I interlaminar fracture toughness test, the laminates [0]18 were fabricated with unidirectional T700/epoxy prepregs. The release film (PTFE) was added at the mid-plane to form the pre-crack of 65 mm. The film interlayer was directly placed at the mid-plane for the film interleaving laminates. The PEK-C particles were scattered using a sieve at the mid-plane for the particle interleaving laminates. For bending test, 12 plies of unidirectional prepregs were stacked with interlayers placed between each ply. For the low velocity impact test, the laminates [+45/0/−45/90]2s were fabricated with interlayers placed between each ply. For all the tests, control samples and epoxy matrix toughened samples were also fabricated. The stacked laminates were cured in an autoclave and the pressure, temperature and holding time were set up to 0.6 MPa, 120 °C and 120 min, respectively. Based on the areal density of PEK-C film, a careful calculation was conducted to ensure the equal amount of PEK-C added in CFRPs. The added amount of PEK-C and thickness of samples after curing for different tests were displayed in Table 1.
2.1.2. Preparation of interlayer films PEK-C powders were dissolved in THF at the weight fraction of 10% at room temperature. PEK-C film was formed by dip-coating method on top of glass substrate at a defined dragging rate of 0.2 cm/s and was then placed in the oven for 20 min to evaporate THF. The film thickness was restricted to 10 μm to avoid the decrease about in-plane properties caused by excessive amount of thermoplastic resin added [15]. The thickness of film obtained was 8 ~ 10 μm measured by micrometers and the areal density of PEK-C film was averagely 0.33 mg/cm2 calculated by dividing the masse by the area of film. 2.1.2.1. Impregnation of carbon fibres with epoxy resin. A schematic of the fabrication process of prepregs was shown in Fig. 1. A calculated
Fig. 2. Preparation process of CFRP composites.
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modulus and strength of the CFRPs according to the ISO14125-1998 test standard [26]. The sample size was 100 mm × 15 mm. The test were conducted at room temperature with a constant displacement rate of 2 mm/min. Five samples were tested for each value. Scanning electron microscope (SEM, Hitachi S-3400N) was applied to investigate the fracture and lateral surfaces of samples after mode I interlaminar fracture toughness tests and the lateral surfaces of samples after bending tests. The accelerating voltage was set as 15 kV. Gold layer was sputter-coated on the surface before tests. The phase images were studied by atomic force microscope (AFM) tapping mode. Fig. 4 gives the schematic setup of phase imaging by AFM tapping mode [15,27]. The cantilever was modulated by a sinusoidal wave and phase contrast caused by differences in surface adhesion and viscoelasticity was monitored by the detector. Silicon tapping cantilever (OTESPA) was used with a radius of curvature of about 2 nm. Tests were performed by Bruker Dimension Icon AFM. The amplitude setpoint was chosen as 40 mV. Height and phase images were recorded simultaneously. Each test was performed at room temperature and repeated five or more times.
Fig. 3. Experimental set-up for mode I interlaminar fracture toughness test.
2.2. Characterization and testing
3. Results and discussion
Mode I interlaminar facture toughness, GIC were measured on a universal testing machine (Instron 5982) according to the ASTM D5528-01 [23]. The loadcell of 5 kN was used. The data acquisition frequency was set as the default setting (10 Hz). The experimental device was shown in Fig. 3. The size of samples was 125 mm × 20 mm. The loading rate was 2 mm/min. The modified beam theory (MBT) method was used to calculate GIC.:
GIC =
3.1. Mode I interlaminar fracture toughness R-curves of all the CFRPs were shown in Fig. 5. R-curve displayed the relation of GIC with the length of crack, which characterized the initiation and propagation of a delamination. From the R-curve, it was obvious that fracture toughness increased with the crack length and eventually reached the plateau. The rise behavior could be explained by the fibre bridging mechanism. It was clear that higher crack growth resistance was observed in toughened composites. Up to crack length of 70 mm, there was no large distinction between control and toughened composites. However, the contribution of PEK-C became apparent after crack length of 70 mm. The initial GIC value of control samples was only 102.5 J/m2 due to its bad interfacial adhesion, while that of matrix toughened, particle interleaved and film interleaved samples were
3P 2b (a + | |)
where P is the load, δ is load point displacement, b is the specimen width, a is the delamination length, and Δ is the corrective factor for crack length calculated for each test based on standard ASTM D552801. Five samples were tested for each value. Low velocity Impact testing was conducted on an impact testing machine (CEAST 9350) according to ASTM D7136/D7136M standard [24]. The loadcell of 45 kN and DAQ systems (CEAST DAS 64 k) were used. The sampling frequency was 1000 kHz. Four rubber tips linked with toggle clamps were used to restrain the corners of specimen during impact testing. The specimen size was 150 mm × 100 mm. The impact energy of the CFRP laminates was set as E = 19 J. Five tests were conducted for each value. Based on JB/T4008-1999 standard [25], ultrasonic inspection (Cscan, C505030) was used to identify the damage, such as delamination, fibre breakage, local punctures or indentation caused by impact. Three-point bending test was used to determine the bending
Fig. 5. Mode I interlaminar fracture toughness of all the CFRPs.
Fig. 4. Schematic setup of phase imaging by AFM tapping mode.
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Fig. 6. Lateral surface images of all the CFRPs: (a) control sample; (b) matrix toughened samples; (c), (d) particle interleaved samples; (e), (f) film interleaved samples.
increased by 64.2%, 82.3% and 71.6% respectively. The film interleaved composites showed high and stable GIC value along the crack propagation. The highest GIC was obtained in film interleaved composites, revealing 472.78 J/m2 at crack length of 85 mm. A large variation of fracture toughness in particle interleaved composite could be attributed to uneven distribution of PEK-C particles on the mid-plane caused by scattering using a sieve by hand. The lateral surfaces of samples after mode I interlaminar fracture toughness were investigated by SEM to track the crack propagation path. Fig. 6(a) showed the magnification of crack route for the samples without toughening, from which, the failure mechanism of fibre bridging was observed and matrix damage in the neighbouring layer was rare. Compared with Fig. 6(a), the crack propagation moved into the neighbouring layer irregularly and damage of matrix was observed for the toughened composites (Fig. 6(b)–(f)). Fibre strip breakage and delamination could be particularly observed for the particle and film interleaved composites in Fig. 6(c) and (e). It should be noted that the
delamination phenomenon in the neighbouring layer of the film interleaved samples was extraordinarily obvious. With an expectation to further investigate the improvement mechanism of mode I interlaminar fracture toughness, the fracture surfaces were examined. Fig. 7 showed fracture surface images of control and toughened composites. The epoxy matrix of control samples was thin in the interlaminar region. The failure mechanism was mainly the matrix breakage and the fibre-matrix debonding. The matrix between fibres exhibited brittle fracture features. The smooth fracture surfaces implied low resistance to crack propagation. (Fig. 7(a, b)). After adding PEK-C into epoxy matrix, the failure mechanism was consistent with that of control samples (Fig. 7(c)). However, it was obvious that fracture surface was transformed from brittle features to ductile features. Submitted to a curing process, the second phase PEK-C and epoxy matrix interdiffused due to the enhanced molecular chain under high temperature and then the phase separation began while the epoxy resin crosslinked to form the network structure [28,29]. The phase
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Fig. 7. Fracture surface images of all the CFRPs: (a), (b) control samples; (c)-(e) matrix toughened samples; (f)-(h) particle interleaved samples; (i), (j) film interleaved samples.
Fig. 8. The force versus time curves for all the specimens: (a) force versus time curves; (b) the smoothed curves.
separation phenomena appeared but was rare in the interlaminar region for matrix toughened samples (Fig. 7(d)). Fig. 7(e) displayed the representative two-phase structure formed by epoxy resin and PEK-C [29,30]. The two-phase structure was composed of the PEK-C “peel” and the epoxy matrix wrapped by the “peel”. When the crack extended to the two-phase structure, the epoxy was pulled out from the “peel” and the smooth fracture surface revealed the brittle fracture. What’s more, the “peel” was pulled during the crack growth and exhibited
Table 2 Fmax, DFmax, Elastic energy for all the samples. Fmax (N) Control Resin toughened Particle interleaved Film interleaved
5512 5942 5871 5961
± ± ± ±
32 85 64 22
DFmax (mm) 2.88 3.18 3.14 3.06
± ± ± ±
0.12 0.08 0.17 0.23
Elastic energy (J) 7.0 8.5 7.7 8.0
± ± ± ±
0.6 0.7 0.5 0.3
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Fig. 11. Bending properties of all the samples.
lead to the deviation of cracks and the formation of the secondary cracks in the interlaminar region, even the delamination of matrix in the weak region and the debonding between fibres and matrix due to weak adhesion of neighbouring layers, which could dissipate more energy and enhance further the interlaminar fracture toughness of composites.
Fig. 9. Energy versus time curves of all the samples.
superior ability of ductile deformation. The ductile deformation and pulled-out of epoxy toughened by PEK-C consumed more energy and contributed to the improvement of fracture toughness. The clean fibre surface of control and matrix toughened samples were observed, demonstrating weak interfacial adhesion between the matrix and the fibres (Fig. 7(a–c)). The fracture surfaces of particle interleaving composites were displayed in Fig. 7(f–h). Fig. 7(f) showed a rough surface and the interlaminar region was covered with interlayer, which revealing the high resistance to the crack propagation. The phase separation of PEK-C and epoxy matrix was also observed. Cavities left by the pulled-out epoxy and plastic deformation caused by PEK-C were clearly displayed in Fig. 7(g, h). However, the thickness difference of the PEK-C “peel” was found between Fig. 7(g) and (h) and the “peel” was thicker in Fig. 7(h). The thicker “peel” was easily formed in the region of more PEK-C resin [29,30]. Therefore, the different thickness was attributed to the uneven distribution of PEK-C particles. What’s more, the difference of thickness made clear the fluctuation of GIC with the length of crack for particle interleaved laminates found in the Section 3.1. Similar toughening features were observed in film interleaved samples (Fig. 7(i, j)), but were less. This could be because PEK-C film was thin (< 10 μm) and uniform distribution, resulting in the easier diffusion of PEK-C from interlaminar region into neighbouring layers. During the crack propagation in the mode I interlaminar fracture toughness test, in case that the crack moved to the area of the twophase structure, higher energy dissipation and fracture toughness might
3.2. Low velocity impact The force versus time curves for all the specimens at impact energy of 19 J were shown in Fig. 8. The first phase (Between [0–2.5 ms]) was similar, revealing that the PEK-C addition affected less the elastic behavior of specimens. The maximum load (Fmax) and the displacement corresponding to the maximum load (DFmax) obtained from the testing could be used to evaluate the impact resistance [31]. For the toughened laminate composites, the improvement of impact resistance was seen in terms of the higher Fmax and higher DFmax. Fmax and DFmax increased from 5512 N and 2.88 mm for the control specimen to 5942 N and 3.18 mm, to 5874 N and 3.14 mm and to 5961 N and 3.06 mm, respectively, shown in Table 2. What’s more, the significant drop of load after Fmax for the control specimen in Fig. 8(b) indicated that damage propagation in the untoughened composite plate was more severe. Composites toughened by three different methods revealed almost the same performance in terms of impact resistance due to the same added amount. Fig. 9 showed the energy versus time curves of all the samples. The maximum value of each curve represents the associated impact energies. The constant value reached at the final phase of impact event gives the total energy absorbed by composited specimens. The
Fig. 10. C-scan results of delaminated areas of all the CFRPs: (a) control sample; (b) matrix toughened samples; (c) particle interleaved samples; (d) film interleaved samples.
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difference between them is considered as elastic energy, retained in the impactor and used to rebouned the impactor from the nonperforated samples [31]. From Fig. 9 and Table 2 it could be seen that the elastic energy in toughened samples was slightly larger, showing a higher impact resistance. Fig. 10 showed C-scan results of delaminated areas of all the CFRPs. The control CFRP laminate showed a damaged area of 281.28 mm2 calculated by gray-scale value based on MATLAB. With adding PEK-C to the CFRP, the damage area was reduced. A drop of damage area by 28.9% was observed for particle interleaved composites. PEK-C and two-phase structure formed between PEK-C and epoxy resin could absorb the impact energy and further restrain the crack growth in matrix and between the layers.
3.3. Three-point bending Fig. 11 showed the bending properties of all the CFRPs. The results showed that the flexural strength was not compromised with the incorporation of PEK-C in the CFRP composites due to the limited addition of PEK-C [15]. The bending strength were 1141, 1157, 1125 and 1099 MPa for control, matrix toughened, particle interleaved and film interleaved samples respectively. The modulus was decreased by 16% for film interleaved composites due to the inherent lower rigidness of thermoplastic resin than epoxy matrix, which could weaken the rigidness of interleaved laminates. Nonetheless, a distinct increase by 38% in particle interleaved samples was observed. The flexural modulus was sensitive to sample thickness. The thickness of particle interleaved
Fig. 12. The lateral faces of all the samples after bending tests: (a), (b) control samples; (c), (d) matrix toughened samples; (e)-(h) particle interleaved samples; (i)-(k) film interleaved samples.
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3.4. Phase image by AFM tapping mode The phase images of all the samples scanned by AFM Tapping mode were displayed in Fig. 13. Fig. 13(a) showed the profile in epoxy matrix for laminates without toughening and obviously there exist only one phase. The two-phase morphology of sea-island was observed for toughened composites (Fig. 13(b)–(d)). The island of PEK-C was distributed in the epoxy matrix. Submitted to the curing process, PEK-C was firstly dissolved in epoxy matrix and then the phase separation occurred while the epoxy resin began crosslinking to form the network structure. Compared to the SEM images of two-phase structure in Fig. 7, the scanned results by AFM tapping mode were different. In Fig. 7, the epoxy matrix was segregated with PEK-C “peel” and the size of inner region was about 1–2 μm, or even larger. Fig. 13 revealed that there existed the nanoscale sea-island substructure, from which, it could be speculated that the island phase of PEK-C was formed in the round epoxy matrix wrapped by PEK-C “peel”. The formation of phase structure and the morphology as described were summarized and sketched schematically in Fig. 14. During the curing process, the kinetic ability of molecular chain of epoxy resin was much more enhanced than that of PEK-C resin at 120 °C due to the obvious difference in glass transition temperature of epoxy resin (about 130 °C) and PEK-C resin (above 230 °C). A large amount of epoxy resin diffused into PEK-C resin as shown in Fig. 14(b). The representative two-phase structures of micrometer scale were formed and composed of PEK-C “peel” and the inner region of epoxy resin, which were scattered in the epoxy-rich matrix. The limited amount of PEK-C diffused, resulting in the distribution of nanoscale island phase of PEK-C in the sea phase of epoxy matrix shown in Fig. 14(c). It was estimated that the enhanced interlaminar toughening was mainly attributed to the representative twophase structures of micrometer scale; in addition, it might be attributed to the sea-island substructure of nanometer scale, which need more research to be done.
Fig. 13. The phase images of all the samples scanned by AFM Tapping mode: (a) control sample; (b) matrix toughened samples; (c) particle interleaved samples; (d) film interleaved samples.
samples was changed to 2.37 mm from 2.11 mm for control samples. The rise in modulus could attribute to the augmentation of sample thickness caused by too large PEK-C particle size (about 180 μm provided by producer). The lateral faces of samples after bending tests were analyzed using SEM images as displayed in Fig. 12. It was obvious that the failures were caused by fibre breakage and crack propagation in the interlaminar region. The cracks were smooth, indicating the low crack propagation resistance for samples without toughening. The large and long cracks were formed due to the weak interfacial adhesion, which leaded to the shortest path of fibre breakage for samples without toughening. The neat fibre surface demonstrated weak interfacial adhesion between the matrix and the fibres, which was coincide with that of samples for mode I interlaminar fracture toughness (Fig. 12(a) and (b)). With toughening, the resin “burrs” after plastic deformation caused generally by shear stress in the regions of delamination were observed (Fig. 12(d), (g), and (j)). Fig. 12(f) showed that the fibres and matrix of particle interleaved composites in the inner ply was consistent with that of composites without toughening and the fracture surface displayed the brittle features. After interleaving, the deviation of crack propagation path appeared in the interlaminar region. The crack path from the interleaf to the neighbouring ply indeed increased the fracture toughness (Fig. 12(h) and (k)).
4. Conclusion The toughening efficiency and mechanism of the three methods, including the matrix toughed by PEK-C, PEK-C film interleaving and particle interleaving, were studied in this work. High fracture toughness and impact resistance were achieved for toughened CFRP composites. The PEK-C film interlayer was more efficient and stable in the fracture toughness than the others. The flexural strength was not compromised with the incorporation of PEK-C due to the limited addition of PEK-C. The enhanced toughening was mainly attributed to the two-phase structure of micrometer scale formed by epoxy matrix and PEK-C resin. There existed nanoscale sea-island substructure according to the AFM scanning results, which might be also correlated with the toughness
Fig. 14. Schematic of formation and morphology of phase structure of PEK-C and epoxy matrix.
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improvement. Characterizing the substructure was difficult due to the nanometer size and more efforts should be taken to further develop it.
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