Excellent mechanical properties of carbon fiber semi-aligned electrospun carbon nanofiber hybrid polymer composites

Abstract

The present investigation reports the excellent mechanical properties of carbon fiber fabric hybrid polymer composites obtained by integrating continuous semi-aligned carbon nanofiber layers in the interlaminar region between carbon fiber fabric plies. Different weight fractions (0 to 2.0 wt%) of continuous nanofiber layers were sandwiched with carbon fiber fabric preform to prepare polymer matrix hybrid nanocomposites by a compression molding technique. It was found that layers with 1.1 wt% content of semi-aligned nanofibers significantly enhance the bending strength by 175% and the modulus by 200%. The greater interfacial adhesion of the high surface area of the nanofibers significantly increases the interlaminar shear strength by 190%. The improvement in the mechanical properties of the polymer composites is due to the modification of the matrix properties, crack deflection and debonding. The high surface area and voids that form around the carbon nanofibers can absorb higher fracture energy. However, nanofiber content greater than 1.1 wt% at the interlaminar region of the composites has negative effects on the overall mechanical properties of the hybrid composites. The thick nanofiber layers are believed to be acting as micron size carbon fibers, which results in a downward trend in the bending properties. The present strategy for the preparation of high performance carbon fiber hybrid polymer composites utilizing electrospun continuous carbon nanofibers at the interlaminar region provides a promising approach for improving the properties of composites for different applications.

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1. Introduction

Carbon fiber reinforced polymer composites have been extensively used in different applications, such as wind energy, marine current turbine blades, the automobile, aerospace sectors, owing to their high specific strength, stiffness and excellent corrosion resistance. In these composites, fiber is the load-bearing component; the matrix dissipates loads to the fibers, maintains the fiber orientation and protects the fibers from environmental conditions. The composites are prepared through impregnation of fibers (high strength and high modulus) with polymeric resin matrices. In practice, prepregs consisting of unidirectional or woven fiber fabrics are usually prepared prior to preparation of the composites.¹ In unidirectional carbon fiber polymer matrix composites, the strength is dominated by the fiber volume content.² Meanwhile, in laminated composites, the carbon fibers dominate the in-plane mechanical properties, which are typically sufficiently high for applications, whereas the resin matrix dominates the out of plane properties, which are significantly lower than the in plane properties.³ This is due to the poor matrix and fiber-matrix interfacial properties. In addition, the strain to failure of the matrix and carbon fibers plays a vital role in realizing the full potential of carbon fiber in the composites.⁴

It is known that the laminate composite structure is extremely susceptible to crack initiation and propagation along the laminar interfaces in various failure modes. Delamination is one of the most prevalent life-limiting crack growth modes in laminate composites, as it may cause severe reductions in in-plane strength and stiffness, potentially leading to catastrophic failure of the whole structure.^5,6 To address this, several techniques have been devised to improve delamination resistance,^7–10 such as designing 3D fabric architectures, transverse stitching or pinning the fabrics, fiber hybridization, toughening the matrix resin, and placing interleaves made of tough resin materials in the interplay regions of the laminates. These methods improve the inter-laminar properties, but at the cost of the in-plane mechanical properties.¹¹ Therefore, it is necessary to find an effective technique to improve the out of plane properties without compromising the other mechanical and fracture properties of the composites.

Since the discovery of carbon nanotubes, the unique and fascinating properties of nanoscale materials have triggered tremendous motivation among scientists. Polymer composites reinforced with nanoscale fillers have attracted growing interest among the scientific community. Polymer composites reinforced with nanotubes, nanofibers and nanoparticles in polymeric matrices are expected to possess excellent mechanical properties. However, these composites present several technological issues such as inadequate dispersion, alignment, and low content of the nano-reinforcements, as well as poor interfacial bonding strength and load transfer at the interfaces. To date, the strength of the nanotubes or nanofibers has not yet translated into the bulk composites, i.e., the improvement of the mechanical properties of nanocomposites is not as significant as predicted compared to continuous carbon fiber reinforced composites.^12–18 It is reported that nanoscale reinforcements could distinguishably enhance the toughness and damage tolerance of conventional structural composites.^19–22 It is predicted theoretically and validated experimentally that hybrid fiber-reinforced composites with uniformly distributed nano-reinforcement fillers possess greatly enhanced mechanical properties.²³ The hybrid or multiscale approach in composites has been often motivated by nature, which shows many examples of interesting biocomposites with multiple hierarchical levels of reinforcement. For example, bone is a fascinating multiscale hierarchical composite consisting of macro, meso and micro as well as nano-characteristics, all of which are responsible for imparting unique mechanical characteristics such as high stiffness and high toughness.

Therefore, one promising approach is based upon the incorporation of nano-reinforcement fillers between micron size carbon fibers and polymer composite prepregs to form hybrid nanocomposites. However, the mechanical properties of many hybrid multi-scale composites are still not as high as expected due to the technological challenge of uniform dispersion of the nanoscale fillers in highly viscous resins. Hence, it is important to develop a process to prepare carbon fiber laminated polymer composites with uniformly dispersed nano-reinforcements in the interfacial regions in order to exploit the carbon fiber properties.

In this direction, very little effort has been made; Dzenis et al.^21,22 incorporated a polymer nanofiber interface in advanced aerospace based carbon fiber reinforced polymer composites, and it was established that >1 wt% of the electrospun polymer nanofibers improved the static and fatigue peel and shear interlaminar fracture toughness. Khan et al.²⁴ introduced carbon nanofiber bucky paper in between unidirectional carbon fiber prepreg layers. The interlaminar shear strength and fracture toughness of multiscale composites containing 10 wt% carbon nanofiber bucky papers were enhanced by 31% and 104% compared to those of the composite without bucky paper at the interface. Recently, Chen et al.¹⁹ used electrospun carbon nanofibers to modify epoxy resin for use as a matrix for the development of hybrid multi-scale composites by vacuum assisted resin transfer molding. They reported that the addition of 0.3% carbon nanofibers in epoxy resin increased the impact absorption energy by 79.1%, the interlaminar shear strength by 42.2% and the flexural strength by 13.6%.

However, in the biological system of human bone, the macro, meso and micro phases are all aligned in a continuous form. Therefore, in the present investigation, semi aligned continuous carbon nanofibers with different volume contents in the form of a carbon nanofiber layer at the interlaminar region of carbon fiber fabric and polymer composites are integrated by sandwiching with carbon fabric layers; the influence of this structure on the mechanical properties of carbon fiber hybrid polymer composites is investigated.

2. Experimental

2.1 Synthesis of electrospun nanofibers

Chopped poly-acrylonitrile (PAN) co-polymer (with 6% monomer methyl methacrylate) microfibers of diameter 12.5 μm were used as the starting source. N,N-Dimethylformamide (DMF) of 99% purity (B. P. = 157 °C) procured from Fisher Scientific, India was used as the solvent. The 12 wt% PAN copolymer–DMF solution was prepared by mixing with ultra-sonication for 50 min in a sonication bath, followed by magnetic stirring of the solution for 5 to 10 h at room temperature to obtain a clear, transparent and homogenous solution.

The PAN-copolymer solution was electrospun using an electro-spinning apparatus (ESPIN) procured from the Physics Instrument Company, Chennai, India.^25,26 The polymer solution was loaded into a syringe equipped with a stainless steel capillary metal-hub needle. The solution was electrospun at a flow rate of 0.2 mL h⁻¹ with an applied voltage of 15 kV, a needle tip to collector distance of 20 cm and a cylindrical collector speed of 2000 rpm; the spinning time was 12 h.

The electrospun PAN nanofibers mat was peeled off from the aluminum foil and then placed on the surface of a graphite plate (4 × 25 cm²). Fixed weights were placed over the ends of the nanofiber mat so that tension existed to a certain degree during the stabilization process. The stabilization was carried out in oxidizing atmosphere (air) at a temperature of 320 °C and a heating rate of 2 °C min⁻¹, followed by holding at 320 °C for 1 h to allow complete stabilization of the PAN-based polymer nanofibers. The stabilized nanofiber mat was carbonized at a temperature of 1000 °C and a heating rate of 5 °C min⁻¹ from room temperature to 1000 °C in an inert (high purity nitrogen gas) atmosphere, followed by holding at 1000 °C for 1 h. The changes which took place during the conversion of polymer nanofibers to carbon nanofibers are described in an earlier report.²⁷

2.2 Composite preparation process

In order to prepare the carbon fiber fabric epoxy resin composite, epoxy resin LY 556, diglycidyl ether of bisphenol A (DGEBA), procured from Huntsman, USA, and the hardener, triethylenetetramine (TETA), procured from Central Drug House, India, were mixed in a ratio of 100 : 12.5 by wt and sonicated for 10 min to eliminate air bubbles. Carbon fiber cloth (Twill Woven T-300 carbon fiber of density 1.78 g cm³, tensile strength 3.5 GPa, tensile modulus 230 GPa and number of filaments per tow of carbon fibers 3000) was procured from Toray, Japan. Carbon fiber fabric sheets of the desired size (10 cm × 10 cm) were impregnated with a solution of epoxy resin LY556 and TETA hardener to obtain the resin impregnated carbon fabric prepreg. Initially, the carbon fiber fabric reinforced polymer composites with 30 ± 2 vol% were prepared by a compression molding technique in the laboratory. The stacked prepregs were placed in die moulds of size 10 cm × 10 cm on the hot plate of a hydraulic press; when a temperature of 50 °C was attained, a load of 100 kg was applied on the die mould and the mould was left at a temperature of 100 °C for 1 h for complete curing.

Semi-aligned carbon nanofiber (CNF) sheets of the desired size were sandwiched between the carbon fabric prepregs. The sandwich structure was left for 1 h at atmospheric conditions so that excess resin poured out and a prepreg was formed. The prepregs were placed in a die mould of size 10 cm × 10 cm on the hot plate of the hydraulic press; after reaching a temperature of 50 °C, a load of 100 kg was applied on the die mould and the mould was left on the hot plate of hydraulic press at a temperature of 100 °C for 1 h for complete curing.

The same procedure (as per above) was applied to make successive hybrid composites with different CNF (0.6, 0.9, 1.1, 1.3 and 1.9 wt%) layers. The different wt% of the CNF layers were sandwiched at the interlaminar regions in between the carbon fiber fabric prepreg to obtain composites with varying CNF nanofiber contents. Fig. 1 shows the schematic of the stacking of the semi-aligned electrospun carbon nanofiber layers between the twill woven carbon fiber fabric layers.

Fig. 1 Schematic of the stacking of semi-aligned electrospun carbon nanofiber layers between twill woven carbon fiber fabric for preparation of the composites.

2.3 Characterizations

The surface morphology of the carbon nanofibers and the fracture surface of the hybrid polymer nanocomposite were observed by scanning electron microscope (SEM, VP-EVO, MA-10, Carl-Zeiss, UK) operating at an accelerating potential of 10.0 kV. The surface morphology of the carbon nanofibers was characterized by transmission electron microscopy (TEM) using a Tecnai G20 S-TWIN 300 kV instrument. The carbon fibers and carbon nanofibers were characterized by X-ray diffraction to determine their structural differences (XRD, D-8 Advanced Bruker diffractometer) using CuK_α radiation (λ = 1.5418 Å). The mechanical properties of the carbon fiber fabric epoxy and the carbon fiber fabric-carbon nanofiber hybrid polymer composites were measured by an Instron universal testing machine, model 5967, at room temperature. The bending strength measurements were performed on samples of length 60 mm, width 15 mm and thickness 2.0 mm that were cut from the prepared composites. The three point bending strength was measured with a span length of 30 mm at a cross head speed of 1 mm min⁻¹, and the short beam shear strength was measured at a span length to thickness ratio of 4 and a cross head speed of 0.5 mm min⁻¹.

The coefficient of thermal expansion (CTE) of both types of composites was measured with a thermo-mechanical analyzer (TMA-Q400, TA Instruments) in nitrogen atmosphere in the temperature range of 50 to 150 °C @ 10 °C min⁻¹.

3. Results and discussion

Fig. 2(a) shows the SEM micrograph of the carbon nanofiber sheet derived from electrospun 12 wt% PAN solution followed by stabilization in oxidizing atmosphere at a temperature of 320 °C and a heating rate of 2 °C min⁻¹, and carbonization at a temperature of 1000 °C in inert atmosphere at a heating rate of 5 °C min⁻¹. The diameter of the carbon nanofibers is in the range of 400 to 500 nm, and the fibers are semi-aligned almost in one direction. The surface of the nanofibers is almost smooth. Fig. 2(b) shows the TEM micrograph of CNF, which shows a turbostratic structure because it was only heat treated to 1000 °C. The surface of CNF shows non-aligned layers of carbon, which can be useful for intercalation in the composite.

Fig. 2 (a) SEM micrograph of carbon nanofibers, (b) TEM micrograph of carbon nanofibers.

Fig. 3 shows the XRD patterns of PAN based T-300 carbon fibers of diameter ∼7 μm and carbon nanofibers derived from the electrospun technique with diameters of 400 to 500 nm; this reveals that the diameter of CNF is one order of magnitude less than that of the carbon fibers. The XRD spectra provide details of the carbon structure in both types of carbon fibers. In the case of the T-300 PAN based carbon fibers, the 002 peak appears at 2θ = 25.29°, corresponding to a d-spacing value of 0.352 nm, while in the case of the carbon nanofibers, the 002 peak appears at 2θ = 24.42°, corresponding a to d-spacing value of 0.365 nm. This demonstrates that CNF has more turbostratic structure than carbon fiber; this is because the commercial carbon fibers were heat treated at a higher temperature. Thus, nanofibers with turbostratic structures and high surface area can be suitable for interacting with the polymer matrix in the composites.

Fig. 3 XRD patterns of (a) T-300 carbon fibers and (b) carbon nanofibers.

Fig. 4 shows the bulk densities of the carbon fiber polymer composites with increasing carbon nanofiber content. Initially, in the case of the carbon fiber fabric polymer composites, the bulk density is 1.25 g cm³, while on incorporation of CNF at the interlaminar regions of the composites, the bulk density continuously increases; it reaches a maximum of 1.45 g cm³ at 1 wt% of CNF content. However, at higher CNF contents, the bulk density of the hybrid polymer composites decreases to 1.29 g cm³ at 1.9 wt% of CNF. Enhancement in the bulk density can reduce voids and porosity in the composites, which has a positive effect on the mechanical properties of fiber reinforced polymer composites.

Fig. 4 Variation in bulk density of nanocomposites with increasing CNF content.

The coefficient of thermal expansion (CTE) can give an idea of the shrinkage that occurs in fiber reinforced polymer composites with increasing temperature at a defined load. It also indicates the bonding between the reinforcing constituent in the carbon fiber polymer and the carbon fiber–carbon nanofiber hybrid composites. It is found that the linear CTE in the case of composites without a nanofiber interface is 235 μm m⁻¹ °C. The value of CTE is dominated by the contribution of epoxy polymer shrinkage in the temperature range of 30 to 150 °C. Meanwhile, after incorporating the continuous nanofiber layer at the interlaminar region in the carbon fiber fabric polymer composites, the value of CTE decreases, which may be due to the improvement in the polymeric matrix properties. This is due to the intermingling of CNF, which suppresses the shrinkage in the matrix region; as a result, in the 0.9 wt% CNF layer incorporated composite, the CTE value decreases to 220 μm m⁻¹ °C. However, in the 1.1 wt% nanofiber incorporated composite, the value of CTE further decreases to 189 μm m⁻¹ °C; this is due to the increased contraction during shrinkage of the composite. Above 1.1 wt% of the CNF layer at the interlaminar region in the composites, the CTE value increases to 249 μm m⁻¹ °C. This demonstrates that a higher nanofiber content does not suppress shrinkage, which ultimately influences the mechanical properties.

Fig. 5 shows the bending strength of polymer composites with increasing content of CNF in the form of layers sandwiched with the carbon fabric preform. Bending strength is a material parameter that indicates the capability of the material to resist failure under external loading in the form of compression and tensile force. Initially, the bending strength of the conventional carbon fiber fabric polymer composite is 398 MPa; on sandwiching with different weight contents of carbon nanofiber sheets, the strength of the composite increases significantly. The mechanical properties of fiber reinforced polymer composites are generally governed by the volume fraction of reinforcing fibers. For the 0.6 wt% carbon nanofiber layer at the interlaminar region, the bending strength of the fiber reinforced composite increases to 449 MPa; for the 0.9 wt% CNF layer, the bending strength increases to 528 MPa, while at 1.1 wt% of carbon nanofibers, the strength increases to 730 MPa. As the nanofiber content increases further, the strength decreases continuously; for the 1.9 wt% nanofiber layer, the strength approaches a value of 594 MPa, which is still much higher than that of the fiber reinforced polymer composite without nanofiber layers (398 MPa). The increase in strength of the polymer composites after continuous CNF incorporation at the interlaminar region is due to the lower content of CNF; the nanofibers intermingled with the epoxy resin and thus improved the properties of the epoxy resin.

Fig. 5 Variation in bending strength with increasing CNF content in the composites.

When load is applied, for the composite with a 0.6 wt% CNF interlaminar layer, the nanofibers break from the matrix and dissipate the strain energy, preventing the failure of the composites and leading to a greater value of flexural strength. However, a higher content of nanofibers (0.9 and 1.1 wt%) layer at the interlaminar region can assist load transfer from the matrix to the carbon fibers. As reported in an earlier section, the bulk density of the composites increases, which also plays a role in increasing the bending strength. A higher value of bulk density with nanofiber interlaminar layers can reduce the number of defects and flaws in the composites, and the decrease in CTE value can also be helpful in improving the bending strength of the composites. The nanofibers intermingled in the matrix can form a void around the nanofibers due to the mismatch in the coefficients of thermal expansion. This results in crack deflection and induces crack pinning. These voids around the CNF can absorb higher fracture energies due to the high surface area of the carbon nanofibers. As the nanofiber content of the interlaminar layers in the composites (above 1.1 wt%) increases, the thickness of the CNF layer increases. This results in a downward trend in the out of plane properties. This is due to the thick nanofiber layer, which is believed to be acting in the same role as micron carbon fibers, resulting in strong interaction between the nanofiber layers. This shows that a higher content of continuous CNF as an interlaminar layer does not influence the properties of composites to a great extent. Thus, the optimum nanofiber layer content in the present investigation is not greater than 1.1 wt%.

Fig. 6 shows the variation in the modulus with increasing CNF content in the carbon fiber reinforced polymer nanocomposites. The modulus of the polymer composites is initially in the range of 20 GPa but increases as the content of CNF at the interlaminar region increases above 0.6 wt%. This demonstrates that a lower amount of CNF is not effective in changing the fracture behavior of the composites but that more than 0.6 wt% of CNF is effective in increasing the modulus, due to the change in the mode of fracture behavior. With 1.1 wt% of CNF, the modulus is ≥40 MPa, which is the maximum among all the cases. Meanwhile, as the CNF content further increases, the modulus decreases. This may be due to the percolation threshold of carbon nanofibers with high surface area at the interlaminar region in the carbon fiber polymer composites.

Fig. 6 Variation in modulus with increasing CNF content in the composites.

Fig. 7 shows the variation in the interlaminar shear strength of the polymer nanocomposites with increasing CNF content. The interlaminar shear strength describes the resistance of the composites against interlaminar shear failure, which generally indicates free delamination in angle-ply polymer composites. Initially, in the polymer composite without CNF, the ILSS is 27 MPa; upon addition of CNF to the composites, the ILSS increases continuously up to 55 MPa with a CNF content of 1.1 wt%; as the CNF content further increases, the ILSS decreases to 42 MPa. The low value of the ILSS in this study is associated with the lower fiber content (30 ± 2 vol%) in the polymer matrix composites.

Fig. 7 ILSS with increasing CNF content in the fiber reinforced polymer composites.

It is well known that the interlaminar shear strength of a fiber reinforced polymer composite depends on both the fiber-matrix interfacial bond and the shear properties of the matrix material.²⁸ The lower continuous CNF content of the layer at the interlaminar region acted as a reinforcement to the matrix rich interface and enhances the shear properties of the matrix by imparting resistance to matrix cracking. In the case of higher content (1.1 wt%), the continuous CNF layer around the carbon fiber fabrics increases the interfacial surface due to the high aspect ratio of the nanofibers. The higher interfacial adhesion surface increases energy absorption during debonding or crack propagation. Therefore, there is significant enhancement in the interlaminar shear strength. The CNF between the matrix and interfacial region can also controls the shrinkage in the polymer composites, which also prevents delamination at the interfacial region. However, higher contents of high aspect ratio CNF have a negative effect on the ILSS. Above 1.1 wt% of CNF, the ILSS decreases continuously with increasing CNF content. At 1.9 wt% of CNF, the ILSS value decreases to 42 MPa.

In the case of laminated carbon fiber polymer composites, shear failure occurs via the shear stress that is typically transferred from layer to layer though the polymer matrix. Thus, delamination is one of the primary mechanisms responsible for failure of the composites.^29,30 To overcome the problem of delamination in laminated fiber composites, the properties of the polymeric matrix and fiber-matrix interface must be improved. Fig. 8(a) shows the optical micrograph of the carbon fabric polymer composite after failure, in which stress is transferred from the interlaminar region by crack propagation.

Fig. 8 Optical micrographs of composites after failure (a) without and (b) with carbon nanofiber interlaminar layers (1.1 wt%) in the composites.

The fracture surfaces of the laminated carbon fiber composites can provide important information about the interfacial bonding strength between the fabric and resin matrix as well as the fabric/nanofibers and resin. Fig. 9 shows the fracture surfaces of carbon fiber composites observed by SEM with and without carbon nanofiber interlaminar layers. Fig. 9(a and b) show SEM images of conventional carbon fiber fabric polymer composites from different directions; in both the micrographs, the epoxy resin is in contact with the carbon fiber surface, and as a result, the carbon fibers have a rough surface due to remnants of carbon fibers in the resin. Moreover, the cross section of the fibers in the composites shows that the carbon fibers are well bonded with the epoxy resin. In the case of carbon fabric polymer composites with different weight fractions of CNF layers at the interlaminar region (Fig. 9(c–f)), it can be observed that the CNF are intermingled with the epoxy resin and are broken during the failure mechanism (Fig. 9(c)); the high surface area of the nanofibers enhanced the properties of the polymer matrix. This resulted in improved properties of the composites. The nanofiber content (0.6 and 0.9 wt%) and the fracture surface demonstrated that the CNF are broken and that some of them emerged from the polymer matrix phase, which results in a rough fracture surface. This shows that the presence of CNF can deflect microcracks and can increase resistance to the propagation of cracks in the composite. In Fig. 9(e and f), images of the composite with a higher nanofiber content (1.1 wt%) shows that the carbon fibers and matrix are in close contact with each other through the nanofibers at the interlaminar region. The diameter of the nanofibers is one order of magnitude less than that of the carbon fibers; thus, the nanofibers can easily mix with the epoxy resin. The intermingling of the nanofibers with epoxy resin can improve the properties of the epoxy resin as well as control the shrinkage of the composites. This is demonstrated by the bulk density and CTE data. Upon failure of the composite, debonding ultimately occurred and nanofibers were pulled out of the matrix; this resulted in voids and angular gaps around some of the nanofibers. This is attributed to the increase in resistance to crack propagation due to high surface energy dissipation. It is also evident from the optical micrograph of 1.1 wt% nanofiber interlaminar layer based composites, which shows the micro cracking in the composites (Fig. 8). The strengthening mechanism by bridging the microcracks occurs in the interfacial region. The crack bridging promotes the redistribution of stresses between the high surface area carbon fibers as well.

Fig. 9 SEM images of fracture surfaces of composites without and with carbon nanofiber interlaminar layers. (a and b) Fracture surface of carbon fabric composites, (c) 0.6 wt%, (d) 0.9 wt%, (e and f) 1.1 wt% carbon nanofibers at the interlaminar regions in the polymer composites.

As observed, the bending strength decreases as the wt% of the carbon nanofibers layer decreases from 1.3 to 1.9. As the content of nanofiber layers increases, the ultimate surface area in the nanofiber composites increases; this does not occur at the interfacial region, which results in a thick layer of nanofibers turning towards the carbon fibers, and as a result, the effectiveness of CNF continuously decreases. This shows that a higher content of continuous CNF as an interlaminar layer does not influence the properties of composites to a great extent. This clearly demonstrated that the percolation limit of the nanofibers in the present investigation is a 1.1 wt% equivalent layer of nanofibers in the carbon fiber fabric polymer composites. This clearly shows that a thick layer of CNF reduces the nano-interfacing effect.

In this direction, very little effort has been made. Chen et al.³¹ used random electrospun carbon nanofiber sheets sandwiched between carbon fiber fabrics to develop hybrid multiscale epoxy composites. One sheet of carbon nanofiber was placed between two carbon fiber fabric sheets to prepare the composites. They reported that the interlaminar shear strength increased by 86% and the bending strength increased by 11% compared to a control composite with carbon nanofibers. In another course of investigation by Chen et al.,³² PAN based nanofibers were directly electrospun onto conventional T-300 carbon fiber fabric for different (5, 10, 20 and 30 min) time intervals. Stabilization and carbonization of the PAN based nanofibers on carbon fiber fabric were carried out. A hybrid multiscale epoxy reinforced with electrospun carbon nanofiber–carbon fiber fabric was developed by a vacuum resin transfer moulding technique. It was reported that the flexural strength of the hybrid composites increased from 376 MPa to 465 MPa (an increase of 24%) and the interlaminar shear strength increased from 27.5 MPa to 88 MPa (an increase of 220%) while the modulus increased from 12.1 to 24.8 GPa (an increase of 105%) for the optimum 10 min collection time of PAN nanofiber on the carbon fiber fabrics.

In another study by Chen et al.,¹⁹ electrospun carbon nanofibers were used to modify epoxy resin for use as a matrix for the development of hybrid multiscale composites by vacuum assisted resin transfer moulding. It was reported that the addition of 0.3% of carbon nanofibers in epoxy resin could increase the impact absorption energy by 79.1%, the interlaminar shear strength by 42.2% and the flexural strength by 13.6%. Dzenis et al.^21,22 incorporated a polymer nanofiber interface in advanced aerospace based carbon fiber reinforced polymer composites, and they established that >1 wt% of electrospun polymer nanofibers improved the static and fatigue peel and the shear interlaminar fracture toughness. Khan et al.²⁴ introduced carbon nanofiber bucky paper between unidirectional carbon fiber prepreg layers. The interlaminar shear strength and fracture toughness of the multiscale composites containing 10 wt% carbon nanofiber bucky paper were enhanced by 31% and 104%, respectively, compared to a composite without bucky paper at the interface.

However, in the present investigation, semi aligned carbon nanofiber sheets of different thicknesses were applied at the interfacial region, and it was found that the bending strength increased by 175%, the modulus increased by 200% and the interlaminar shear strength increased by 190% at 1.1 wt% carbon nanofiber content. The enhancement in the bending strength in the present study is much higher than in previous reports^31,32 to the best of our knowledge. This may be due to the fact that a higher content of the almost-aligned carbon nanofibers at the interfacial regions played a significant role in load transferring and surface energy dissipation. This clearly indicates that the incorporation of aligned nanofibers at the interfacial region can be very constructive for greatly improving the bending properties of carbon fiber polymer composites.

4. Conclusions

To improve the out of plane properties (bending and interlaminar shear strength) of carbon fiber fabric reinforced polymer composites, a nano-interfacing approach is adapted to develop hybrid composites in the present investigation. The assumption behind this approach is that improvement of interfacial bonding will suppress crack propagation through the interlaminar regions by applying a continuous carbon nanofiber layer with a high surface area at the interlaminar regions in carbon fiber fabric composites. Layers with different weight fractions of electrospun continuous carbon nanofibers are sandwiched with the carbon fiber fabric preform to make the composites by a compression moulding technique. It is revealed that the carbon nanofiber layer at the interlaminar region improved the properties of the polymeric matrix and the load bearing capability of the composites. The layer with 1.1 wt% content of continuous carbon nanofibers significantly enhances the bending strength by 175% and the modulus by 200%. The high surface area nanofibers significantly increase the interlaminar shear strength by 190%. It was found that percolation threshold of the high surface area carbon nanofibers is 1.1 wt% in the present study; below and above 1.1 wt% nanofiber contents in the composites, the out of plane properties of the carbon fiber composites decrease, but the values of bending strength and interlaminar shear strength are higher compared to those of conventional fiber composites.

The present strategy of fabricating high performance carbon fiber polymer composites utilizing electrospun continuous carbon nanofibers at the interlaminar region provides an emerging path for improving the out of plane properties of lightweight carbon fiber polymer composites for different applications.

Acknowledgements

The authors are highly grateful to Director, CSIR-NPL, and Head, Material Physics and Engineering Division, for his kind permission to publish the results. The authors would like to thank Dr R. P. Pant for providing XRD spectra of the carbon fibers and Mr Jai Tawale for providing SEM characterization of the carbon fiber composites. The authors would like to thank the Department of Science and Technology for financial support in the form of an India-Japan international collaborative project (DST/INT/JSPS-199/2015).

References

1. E. Bkyarova, E. T. Thostenson, A. Yu, H. Kim, A. Gao, J. Tang, H. T. Hahn, T. W. Chou, M. E. Itkis and R. C. Haddon, Langmuir, 2007, 23, 3970–3974.
2. S. R. Dhakate, R. B. Mathur and T. L. Dhami, Carbon Sci., 2002, 3, 1–6.
3. J. Zhu, A. Imam and R. Crane, Compos. Sci. Technol., 2007, 67, 1509–1517.
4. C. R. Thomas, Essentials of carbon–carbon composites, The Royal Society of Chemistry, 1993.
5. J. K. Kim and Y. W. Mai, Engineered interfaces in fiber reinforced composites, Oxford Elsevier, 1998.
6. J. K. Kim and Y. W. Mai, Compos. Sci. Technol., 1991, 41, 333–378.
7. K. Dransfield, K. C. Baillie and Y. W. Mai, Compos. Sci. Technol., 1994, 50, 305–317.
8. J. K. Kim, Key Eng. Mater., 1998, 141(143), 149–168.
9. A. P. Mourutz, Composites, Part A, 2007, 38, 2383–2397.
10. M. Hoja, S. Matsuda, M. Tanaka, S. Ochiai and A. Murakami, Compos. Sci. Technol., 2006, 66, 665–675.
11. C. A. Steeves and N. A. Fleck, Int. J. Solids Struct., 2006, 43, 3197–3212.
12. M. Alexandre and P. Dubois, Mater. Sci. Eng., R, 2000, 28, 1–63.
13. E. T. Thostenson, Z. F. Ren and T. W. Chou, Compos. Sci. Technol., 2001, 61, 1899–1912.
14. J. Njuguna and K. Pielichowski, Adv. Funct. Mater., 2003, 5, 769–778.
15. E. T. Thostenson, C. Y. Li and T. W. Chou, Compos. Sci. Technol., 2005, 65, 491–516.
16. J. N. Coleman, U. Khan and Y. K. Gunoko, Adv. Mater., 2006, 18, 689–706.
17. S. C. Tjoing, Mater. Sci. Eng., R, 2006, 53, 73–197.
18. T. W. Chou, L. M. Gao, E. T. Thostenson, Z. G. Zhang and J. H. Byun, Compos. Sci. Technol., 2010, 70, 1–19.
19. Q. Chen, W. Wu, Y. Zhao, M. Xi, T. Xu and H. Fong, Composites, Part B, 2014, 58, 43–53.
20. E. T. Thostenson, Appl. Phys. Lett., 2009, 97, 073111.
21. Y. A. Dzenis and D. H. Reneker, US Pat., 6265333, 2001.
22. Y. Dzenis, Science, 2008, 319, 419–420.
23. X. F. Wu, Fracture of advanced polymer composites with nanostructured interfaces: fabrication, characterization and modeling, VDM Verlag Publishing House, Germany, 2009.
24. S. U. Khan and J. K. Kim, Carbon, 2012, 50, 5265–5277.
25. A. Sharma, A. Gupta, G. Rath, A. Goyal, R. B. Mathur and S. R. Dhakate, J. Mater. Chem. B, 2013, 1(27), 3410–3418.
26. R. Sharma, N. Singh, A. Gupta, S. Tiwari, S. K. Tiwari and S. R. Dhakate, J. Mater. Chem. A, 2014, 2, 16669–16677.
27. S. R. Dhakate, A. Gupta, A. Chaudhari, J. Tawale and R. B. Mathur, Synth. Met., 2011, 161, 411–419.
28. J. K. Kim and Y. W. Mai, Engineered interfaces in fiber reinforced composites, Elsevier, Oxford, 1998.
29. J. Li, M. L. Sham and J. K. Kim, Compos. Sci. Technol., 2007, 67, 296–305.
30. M. F. Hibbs, M. K. Tse and W. L. Bradley, Interlaminar fracture toughness and real time fracture mechanism of some toughened graphite/epoxy composites, in Toughened composites, ed. J. Johnston Norman, ASTM STP, 1987, vol. 937, pp. 115–130.
31. Q. Chen, L. Zhang, A. Rahman, Z. Zhou, X. F. Wu and H. Fong, Composites, Part A, 2011, 42, 2036–2042.
32. Q. Chen, Y. Zhao, Z. Zhou, A. Rahman, X. F. Wu, W. Wu, T. Xu and H. Fong, Composites, Part B, 2013, 44, 1–7.

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