Mechanical properties of polymer composites reinforced by functionalized graphene prepared via direct exfoliation of graphite flakes in styrene

Abstract

Using a convenient sonochemical method, we have prepared polystyrene (PS) functionalized graphenes (FGs) by direct exfoliation of graphite flakes in the monomer of styrene. This material could be dispersed in toluene and subsequently formed into composites with PS by solution casting followed by a compression molding method. The FGs are well dispersed in the PS matrix as evidenced by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The addition of a small amount of FGs significantly increased the glass transition temperature (T_g) of the PS matrix. It's worth noting that compared to pure PS, a 103.7% increase in tensile strength and a 407% improvement of Young's modulus were achieved by addition of only 0.5 wt% FGs due to the effective load transfer between the FGs and the PS matrix. The storage modulus was also considerably enhanced upon addition of FGs. The PS/FG composites show superior mechanical properties as compared to their counterparts reinforced by carbon fillers such as multi-walled carbon nanotubes (MWNT) and thermally reduced graphene (TrG). The presence of PS on the FG surfaces and thus a strong interaction between the FGs and PS is essential for the property enhancements. Moreover, the Young's modulus of the PS/FG composites was compared with the predictions of the well-established Halpin–Tsai model. The FGs are very effective as a reinforcing filler for graphene-based polymeric composites.

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Introduction

Graphene, which is composed of a single layer of sp²-hybridized carbon atoms arranged in a hexagonal lattice, has received tremendous attention in both scientific and industrial areas. It is the strongest material ever measured and has other remarkable qualities, including high thermal stability and electron mobility as well as large surface area and high aspect ratio that elevates its potential for use as a filler in high-performance polymer nanocomposites.^1,2 When incorporated appropriately, graphene can dramatically enhance physical properties of polymer composites at low loadings.³ The strengthening effect of graphene to various polymers has recently been reported, e.g., polyvinyl alcohol,⁴ polycaprolactone,⁵ polyurethane and polystyrene (PS).⁶ Most graphene/polymer composites have been developed using both reduced and unreduced graphite oxide as the synthesis of graphene oxide (G-O) is a cost-effective and scalable process. The exfoliation of graphite oxide (GO) can be readily done by mild sonication and oxygen-containing functional groups on GO provide rich chemistry for further surface modification. Yu et al.⁷ synthesized functionalized graphene oxide (FGO)/polyurethane acrylate nanocomposites by UV curing technology. The incorporation of functionalized graphene oxide effectively enhanced the thermal stability and mechanical properties of polyurethane acrylate. The glass transition temperature (T_g) of the composites was slightly enhanced as compared to pure polyurethane acrylate and the initial degradation temperature of the composite with 1.0 wt% FGO was increased. Qiu and Wang et al. reported the exfoliation of GO by π–π interactions using porphyrin.⁸ Sheng et al.⁹ reported a facile and eco-friendly method for preparing polymer nanocomposites with grafted GO homogeneously dispersed in poly(methyl methacrylate) (PMMA) matrix, then melt blending the grafted GO with PMMA matrix, and simultaneous in situ thermal reduction of GO. Interestingly, the storage modulus of the nanocomposites increases by 45% and the T_g increases by 7.5 °C at 1.5 wt% grafted GO loading. Although the thermal and mechanical properties of the composites can be effectively enhanced, the preparation of both GO and subsequent surface modifications is a multistep process, being troublesome and time-consuming.^10–12 The oxidation process significantly breaks the intrinsic structure of graphene and thus deteriorates its mechanical performance. It is necessary to find a convenient approach to prepare graphene from direct exfoliation of natural graphite flakes via sonication in organic solvents, surfactants/water solutions, or other media.

Recently, it is reported that by using a reactive medium, styrene as the solvent, the combined mechanochemical effects of high intensity ultrasound, in a single step, induce exfoliation of graphite to produce functionalized graphenes (FGs).¹³ The surface tension (35 dyn cm⁻¹ at 0 °C) of styrene match the surface energy of graphite, thus facilitating its exfoliation. Meanwhile, the reactive vinyl groups capable of polymerization undergo sonochemical reactions during the powerful sonication to produce radical functionalization of graphene sheets. Thus, ultrasonic irradiation could mechanochemically exfoliate graphite into graphene sheets combined with FGs with PS chains. The convenient sonochemical route of preparing PS-functionalized graphene sheets in styrene enables its homogenous dispersion in common organic solvents. Herein, taking advantage of this possibility, we fabricated PS/FGs by means of solution casting and subsequent compression molding.

Recent investigations indicate that graphene can considerably increase the low-strain stress of elastomers (low-modulus polymers), but it is much less effective at reinforcing thermoplastics (more rigid matrix materials). Our results reveal that by incorporating FGs into PS matrix appropriately, the obtained PS/FGs not only exhibit dramatic enhancement in mechanical properties but also show improved thermal stability at low FGs loading. Due to the presence of PS chains on graphene, the FGs outperform multi-walled carbon nanotubes (MWNT) and thermally reduced graphene (TrG) as a reinforcing filler in the PS matrix. The Young's modulus of the PS/FGs composites was also evaluated with the predictions of the well-established Halpin–Tsai model. At low filler content, PS/FGs composites offer significant reinforcement. However, the mechanical properties of PS/FGs composites degrade as the content of FGs is beyond a certain value due to the aggregation of FGs.

Experimental

Materials

Natural graphite powder (graphite) was obtained from Qingdao Ruisheng Graphite Co. Ltd. (purity 99.99%, particle size 40 μm). Multi-walled carbon nanotubes (MWNT-2040, 99% purity) with a diameter of 20 to 40 nm and a length of 4–8 μm were purchased from Shenzhen Nano Hong Kong Co., LTD. PS (PG-32, M_w = 20 000) were obtained from Shanghai Aladdin Chemical Co., LTD. Styrene (>99.0%) were purchased from Sinopharm Chemical Reagent Co., LTD.

Preparation of TrG

Thermally reduced graphene was synthesized according to the previously described method.^14,15 0.8 g of GO (GO was prepared according to the method described by Hummers with some modifications¹⁴) was placed into a long quartz tube that was sealed at one end, and the other end was closed with a rubber stopper. Then, a nitrogen inlet was inserted through the rubber stopper. The sample was flushed with nitrogen for 5 min, and the quartz tube was quickly inserted into a muffle tube furnace preheated to 1050 °C and held in the furnace for ten minutes under nitrogen flow. According to the investigation by Schniepp et al.¹⁶ the surface areas of dry TrG powders determined by the Brunauer–Emmett–Teller (BET) method were in the range of 600–1500 m² g⁻¹, but still lower than theoretical surface area of graphene because of the overlap of the exfoliated sheets.

Preparation of FGs

Following exactly the same experimental procedure as in the literature,¹³ the preparation began with natural graphite flakes (20 mg) in freshly distilled styrene (40 mL), which was then sonicated with high intensity ultrasound using an ultrasonic cell disruptor JY 92-Π, 3 cm² Ti tip at 300 W at 0 °C for 4 h under Ar flow. The obtained mixture was then subjected to centrifugation at 1000 rpm for 5 min. The black supernatant was vacuum-filtered through a 0.22 μm PVDF membrane and washed with tetrahydrofuran (THF). The collected black solids were re-suspended in THF and then filtered and washed with excess THF again. This dispersion–filtration–washing cycle was repeated at least 3 times to remove any residual styrene or PS. Thermal gravimetric analysis (TGA) indicates that 17.2% PS is contained in FGs, consistent with the results (18.0 wt%) obtained by Xu et al.¹³

Preparation of PS/FGs composite samples

In a typical experiment, certain amount of FGs was sonicated in toluene for about 60 min. PS pellets were dissolved in toluene. Then, the FGs suspension was added to the PS solution and the resulting mixture was stirred for two hours. After ultrasonication for 60 min, the mixture was poured into large amount of methanol to coagulate the product, following by vacuum-drying at 60 °C until the weight was constant. The dried composites were further compression-molded above their melting temperature to prepare composite films for mechanical testing. The FGs loadings were varied from 0 wt% to 1 wt%. The composites with 0.5 wt% loading of MWNT and TrG were fabricated following the same procedure.

Characterization

The functionalized graphenes were analyzed by Fourier transform infrared spectrometer (FTIR) (TENSOR27, Bruker Co.). The FGs an graphite samples were deposited on the silicon wafer and the Raman spectra was measured using French HORIBA Jobin Yvon company's Lab RAM type Raman analyzer. Tensile properties were measured using rectangular test pieces on a tensile testing machine (AI-700M, Taiwan High Speed Rail Technology Co., LTD). The measurements were performed at a crosshead speed of 1 mm min⁻¹ at room temperature. The average values were obtained from at least five tests of each sample. The dynamic mechanical properties of the PS/FGs composite were measured using a dynamic mechanical analyzer (DMA) (DMA-Q800, TA instrument) in the double cantilever deformation mode at a frequency of 1 Hz. The temperature was swept from room temperature to 180 °C at a heating rate of 3 °C min⁻¹. Thermogravimetric analysis (TGA) was conducted on TA Instruments Q600-SDT at a heating rate of 10 °C min⁻¹ under N₂. For transmission electron microscopy (TEM) analysis, compression-molded composites were microtomed to slices of 50–90 nm thickness using an ultra45 diamond blade (Leica Ultracut UCT, Reichert Inc.) and dropped onto 300 mesh copper TEM grids (Ted Pella). The TEM images were acquired on a JEOL 2010F at 120 keV. Scanning electron microscopy (SEM) samples were prepared by sputter coating freeze-fractured compression-molded samples using a 6 nm coating of iridium metal deposited in an Ar atmosphere. Images were acquired using a JSM-6700F SEM (Japan Electronics Co., Ltd) at 5 keV and 10⁻⁵ torr.

Results and discussion

Due to the physical and chemical effects arising from acoustic cavitation, ultrasound could find important applications in a wide range of materials and chemical synthesis. Taking advantage of this simple technique, Xu et al.¹³ recently reported that FGs were prepared through ultrasonic induced exfoliation of graphite flakes in styene, and simultaneously PS chains were grafted on graphene surfaces through sonochemically initiated radical polymerization of styrene. The TEM observation, FTIR and Raman spectra of FGs confirm that graphene was functionalized by PS chains and exfoliated to single-layer or few-layer graphene (Fig. S1–3, see ESI†). The content of polystyrene in the functionalized graphene is about 10.2 wt%, as determined by comparing the weight loss (TGA) of FGs and graphite (Fig. S4†). The obtained FGs were easily dispersed in organic solvents such as THF, toluene and chloroform, forming single- and few-layer graphene.¹³ The stably dispersed FGs in solution provide a prerequisite for fabricating PS/FGs composites by solvent blending. In this work, the dispersion of the FGs in toluene was used for preparing PS/FGs composites.

The dispersion of nanofillers in polymer matrix is essential for improving the mechanical, thermal and other properties of the composites. The solution phase mixing has been widely used to prepare various kinds of polymer–graphene composites because it is a simple method to uniformly disperse graphene into polymer matrix.¹⁷ The morphologies of freeze-fractured unfilled PS and the sample with 0.5 wt% of FGs were probed by scanning electron microscopy (SEM). As can be seen in Fig. 1a and b, the fracture surface of neat PS is flat and smooth. By contrast, the fracture surface (Fig. 1c–d) of the PS/FGs composite is rough and hierarchical.¹⁸ The rough fractured surface can be observed more clearly in the high resolution SEM image (Fig. 1d), which could be defined as the FGs protruding out of the fracture surfaces.¹⁸

Fig. 1 Cross-sectional SEM images of the freshly-fractured surfaces of PS (a and b) and PS/FGs with various FGs loading (wt%), (c and d) 0.5 wt% PS/FGs, (b and d) high magnification, (a and c) low magnification.

To study the exfoliation of the flakes and the level of sheet dispersion achieved in the composites, TEM images of microtomed sections of the composites are shown in Fig. 2.

Fig. 2 TEM of (a) single-layer FG and (b) few-layer FGs in PS (0.5 wt%) with SAED confirming single-layer (c) and few-layer FGs (d), respectively.

The FGs were indeed present in the composites as exfoliated sheets. Single-(Fig. 2a) or few-layer (Fig. 2b) graphenes were sparsely dispersed in the matrix, with a distribution of lateral dimensions of hundreds of nanometers to several microns. The ultrasonication used to exfoliate graphite flakes fragments the sheets, and is responsible for the small lateral dimensions of many of the platelets. The SAED patterns of the region marked in Fig. 2a and b are shown in Fig. 2c and d, respectively. We observe a hexagonal pattern due to the presence of graphene sheets. The spots of the hexagonal patterns are labeled using the Miller–Bravais (hkil) notation. The different intensities between (1210) spots and (0110) spots can be used to identify the single-layer or few-layer graphene. As we can see from Fig. 2c, the (0110) spots of center ring is more obvious than the (1210) spots of outer ring, indicating that the FGs in the PS matrix are single-layer.¹⁹ In the case of Fig. 2d, the (1210) spots is clearly more intense than (0110), confirming few-layer FGs in the PS matrix.²⁰ That is, FGs are well dispersed in the PS matrix. The significant exfoliation of the FGs was possibly due to the interfacial π–π interaction between graphene nanosheets and styrene.⁸

The representative stress–strain curves of neat PS and PS/FGs composites are depicted in Fig. 3 and the mechanical properties of the composites are summarized in Table 1. Tensile strength is significantly improved at all filler loadings. Compared with neat PS (16.2 MPa), the tensile strength of the PS/FGs, containing only 0.1 wt% FGs increases by 129.6% and reaches maximum value (37.2 MPa). Such enhancements in the tensile strength can be attributed to the homogeneous dispersion of the FGs in the PS matrix and efficient load transfer between the nanofillers and PS matrix. Here, the PS chains on the surface of FGs provide stronger interaction with the PS matrix. As a result, a more effective load transfer across the graphene–PS interface is obtained, resulting in a significant increase of tensile mechanical properties. As can be seen in Table 1, Young's modulus initially increases with the increasing content of the filler. The maximum improvement (407% improvement with respect to neat PS) is found at 0.5 wt% loading where we also observed 103.7% enhancement in tensile strength. However, Young's modulus decreases at higher filler loadings. The reduction in mechanical properties observed at higher graphene content is usually associated with filler aggregation. The appreciable improvements in strength and modulus at low filler loadings are significant as it would facilitate a dramatic reduction in the mass of plastics needed for many structural applications. Upon addition of FGs, elongation at break decreases since strong interactions with the matrix in conjunction with high aspect ratio values substantially restrict the mobility of the polymer chains.

Fig. 3 Stress–strain curves for PS/FGs composites.

Table 1 Summary of tensile properties of pure PS and PS/FGs composites

Sample	Young's modulus/MPa	Tensile strength/MPa	Elongation at break/%
PS	327.8 ± 12.3	16.2 ± 1.7	1.8 ± 0.3
PS/FGs 0.1	516.3 ± 11.6	37.2 ± 1.5	1.9 ± 0.1
PS/FGs 0.2	961.5 ± 15.9	25.5 ± 2.3	1.4 ± 0.2
PS/FGs 0.5	1662 ± 20.3	33.0 ± 2.2	1.3 ± 0.1
PS/FGs 1.0	771.9 ± 18.7	34.8 ± 1.3	1.4 ± 0.1

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It is noted that such mechanical improvement is large compared to the results reported in the literature. Graphene sheets have been attempted to improve the mechanical properties of polymers, however, only moderate enhancements were observed in rigid matrix materials-based systems. For instance, Ramanathan et al.²¹ reported that around 20% and 80% of increases in tensile strength and Young's modulus relative to the pristine PMMA with 1.0 wt% graphene. Qiu et al.²² revealed that 10% increases in tensile strength of the epoxy composite with the addition of TrG at 0.54 vol%. In this regard, poor particle dispersion and interface interaction are presumably responsible for the limited performance improvement. The mechanical enhancements were observed in GO–poly(vinyl alcohol) (PVA) systems, where the PVA nanocomposites with 0.7 wt% GO sheets revealed 62% and 76% increases in tensile strength and Young's modulus.²³ The superior hydrophilicity of graphene oxide is conducive to their dispersion and interaction with PVA. When 0.9 wt% of graphene sheets were added to PS, around 70% and 57% of increases in tensile strength and Young's modulus relative to the pristine PS were reported. These are in contrast with the pronounced property enhancements observed in the present PS/FGs system, in which both dispersion and interfacial interaction (thus load transfer efficiency) were mediated by PS chains covalently bonded to the graphene surface.

Halpin–Tsai model is a well-known composites theory for the prediction of elastic modulus of unidirectional or randomly distributed filler-reinforced composites. The Halpin–Tsai model is based on two assumptions: the particle and matrix are linearly elastic, isotropic and firmly bonded, and the particle–particle interactions are not explicitly considered. It relates composite moduli to the volume fraction of the filler, the relative moduli of the constituents and the reinforcement geometry (e.g. the aspect ratio) in a straightforward manner. Herein, to study the dispersion status of FGs in PS matrix, the Halpin–Tsai equation is employed to simulate the modulus of the PS/FGs composites. The modulus E_c are given by relative moduli of the constituents and the reinforcement geometry (e.g. the aspect ratio) in a straightforward manner. Herein, to study the dispersion status of FGs in PS matrix, the Halpin–Tsai equation is employed to simulate the modulus of the PS/FGs composites. The modulus E_c are given by

where E_c represents the tensile modulus of the composites with randomly distributed FGs, E_g and E_m are the modulus of the FGs and neat PS. The parameters of W, L, t and W_g refer to the width, length, thickness, and weight fraction of volume fraction of FGs, respectively. ξ is the shape factor depending on the filler geometry, ϕ_g is the volume fraction of FGs. ρ_g is the FGs density, and ρ_m is the matrix density. E_g, the Young's modulus of FGs, is taken as 1060 GPa which may approach that of graphene and is used here.¹² The modulus of neat PS, E_m, is 327.8 MPa from the experimental data. The statistical average width, length and thickness of FG sheet are about 1.0 μm, 1.5 μm and 2.0 nm, as determined from TEM images.^1,2 The density of neat PS is 1.05 g cm⁻³ and the density of FGs is 2.1 g cm⁻³ as measured by a pycnometry method in our work. Substituting these parameters into eqn (1)–(5), E_c is calculated under the hypothesis that FG sheets are randomly distributed throughout the composites. As shown in Fig. S5,† it was found that the experimental data for the composites are larger than the theoretical simulation at low FGs content presumably due to improved interface adhesion of FGs with PS (the contribution of an effective interphase zone to the reinforcement effect). However, the mechanical properties of PS/FGs composites were decreased as the content of FGs is beyond a certain value due to the aggregation of FGs.

We also compare the tensile properties of polystyrene reinforced by MWNT and TrG. As can be seen in Fig. 4 and Table 2, FGs show an obvious advantage over TrG in strengthening PS composites. This is due to the fact that PS adsorbed on graphene during the sonochemical preparation process results in an increased compatibility and thus an enhanced interaction between FGs and PS. Although the presence of PS on FGs lead to good stability and solubility in common organic solvents, such a sonochemical treatment do not extensively affect the graphitic structure of FGs.¹³ By contrast, the presence of defects on TrG (also stated below) and the reduction of the lateral size of the sheets during the thermal treatment also represent disadvantages for mechanical reinforcement.^24,25

Fig. 4 Stress–strain curves for PS composites with different fillers at 0.5 wt% loading.

Table 2 Summary of tensile properties of pure PS and PS composites with different fillers at filler loading of 0.5 wt%

Sample	Tensile modulus/MPa	Tensile strength/MPa	Elongation at break/%
PS	327.8 ± 15.6	16.2 ± 2.1	1.8 ± 0.3
PS/FGs	1662 ± 21.9	32.5 ± 2.2	1.3 ± 0.1
PS/MWNT	1054.5 ± 20.1	27.3 ± 2.3	1.4 ± 0.1
PS/TrG	617.1 ± 14.3	22.2 ± 1.9	1.2 ± 0.2

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It is interesting to discuss the reinforcement effect of carbon nanotube (CNT) filled PS nanocomposites. Similarly, the result did not reveal remarkable mechanical improvements compared with FGs due to the entanglement of CNT and relatively poor load transfer efficiency (due to the absence of PS chains on nanotube surface).²⁴ However, MWNT is superior to TrG in property enhancement of PS composites. Two reasons are presumably responsible for the difference between MWNT and TrG. One of them is that there exist more defects in TrG. Although the reduction by high-temperature pyrolysis can partially restore the aromatic structure of graphene oxide, the residual defects still decrease the mechanical strength of reduced graphene sheets.²⁵ The other is the difference in aspect ratio of two particles, which gives rise to an increase of critical loading to establish a particle network in matrices.^26,27 The aspect ratio of MWNT is several hundreds while the aspect ratio of TrG sheets is 20–40 if assuming 1 nm thick monolayer graphene sheets. Smaller aspect ratios require higher particle loadings to achieve identical degree of mechanical enhancement.^26,27 As a result, it is plausible for MWNT to exhibit a prominent mechanical reinforcement at low filler content than that of TrG. The dynamic mechanical properties of neat PS and PS/FGs composites are shown in Fig. 5. The samples were characterized in the constant frequency temperature scan (1 Hz, 3 °C min⁻¹) to determine the effect of the elastic and damping behavior of the nanocomposites. Fig. 5a shows the effect of the FGs concentration on storage modulus. The change in the modulus reveals the change in rigidity, and ultimately the strength of the samples. As expected, the results from DMA experiments on the PS/FGs composites indicate that the storage modulus of the nanocomposites increases compared with neat PS, indicating reinforcement (see Fig. 5a). An increase in storage modulus with the increasing FGs content at low filler loading is followed by a decline at higher loading (1.0 wt% PS/FGs), consistent with the tendency from tensile testing. The improvements in the storage modulus suggest that FGs acts as effective reinforcements in the polymer matrix by transferring the load from the polymer to these fillers. It is notable that a relative increase in the storage modulus was observed with a maximum value of approximately 41% corresponding to 0.5 wt% of FGs loading at 40 °C. An obvious increase in storage modulus of the PS/FGs composites could be explained by the uniform dispersion of FGs sheets in the PS matrix and the strong interaction between the FGs and PS matrix. The main interaction force between the FGs and PS matrix is van der Waals force such as π–π interactions.⁸

Fig. 5 (a) Storage modulus and (b) tan δ curves of PS/FGs composites.

Fig. 5b shows that the presence of FGs in the composites leads to a move to a higher temperature in loss modulus.²⁸ T_g determined from the tan δ peak increases with the content of FGs, indicating that the segmental mobility of the PS chains during glass transition was significantly limited and obstructed by the presence of FGs. The maximum shift of T_g is 9.2 °C (from 87.7 to 96.9 °C) at 1.0 wt% FGs loading. The origin of T_g shifts has been attributed to the presence of these so-called ‘interphase’ polymer, which arises due to the interaction of the PS chains with FGs surface.²⁹ Furthermore, FGs act as a resistance to the viscous flow of the polymer chain in the glass transition region. Interestingly, these fillers show improved elastic properties of the composites in the T_g region, suggesting good compatibility between the matrix polymer and fillers in the composites at elevated temperatures. Percolation of this network of interphase polymer could then manifest the large T_g shift of the bulk composite.²⁹

Layered nanofillers usually increase the thermal stability of a polymer due to the physical barrier effect which retards the diffusion of degradation products, gases and heat.³⁰ The thermal stability of the PS/FGs composites was characterized by TGA under non-oxidative condition. As shown in Fig. 6, the degradation temperature of PS/FGs increases remarkably in comparison with pure PS. The onset degradation temperature (T_onset) of PS/FGs increased by 9 °C with only 0.1 wt% of FGs loading and the midpoint degradation temperature (T_midpoint) of PS/FGs composites increased by 10 °C. The appreciable improvement in T_onset at very low FGs loading indicate that the FGs were well-dispersed in the PS matrix and well-interacted with the PS chains. A small amount of FGs is sufficient to create inflammable jammed network that retards transport of the degradation products.³¹ Insets in Fig. 6a shows that the onset degradation temperature (T_onset) and the midpoint degradation temperature (T_midpoint) level off beyond 0.1 wt% FGs loading. Fig. 6b shows the derivative thermogravimetric analysis (DTG) of TGA, which gives the peak temperature at the rapidest degradation rate (T_rap). The DTG of pure PS showed a single large peak at 406 °C. In contrast, all of the FGs/PS nanocomposites of composites displayed increased T_rap of about 21.2 °C higher than that of the pure PS. The enhanced thermal stability for the PS nanocomposites relative to that for the neat PS may be ascribed to the nanoscale dispersion of the FGs and the barrier of the degraded gases.^32,33

Fig. 6 (a) TGA and (b) differential thermogravimetry (DTG) curves of PS/FGs as a function of FGs weight fraction. The inset is onset and midpoint decomposition temperature of PS/FGs composites.

Conclusion

We have used functionalized graphene sheets prepared through sonochemical exfoliation of graphite in styrene as a filler in polystyrene-based composites. Composites were successfully fabricated by a solution casting method followed by compression molding. We found that great improvements in the mechanical properties of polystyrene were obtained with small amount of FGs loading. Compared with neat PS, a 103.7% increase in tensile strength and a 407% improvement of Young's modulus were achieved by the addition of only 0.5 wt% FGs. Appreciable increases in the dynamic modulus were observed by DMA where we also observed a shifts of 9.2 °C in T_g at 1.0 wt% loading versus neat polystyrene. In terms of tensile property, FGs/polystyrene composites significantly outperform their counterparts reinforced by other carbon fillers such as MWNT and TrG. An enhancement in the thermal properties was also obtained with low FGs loading. Due to the presence of PS chains on graphene surfaces, the FGs could be well dispersed in the PS matrix and the strong interaction between the FGs and PS matrix forms, which account for the significant property enhancements. The Young's modulus predicted by the conventional Halpin–Tsai model underestimated the experimental data since the contribution of an effective interphase zone to the reinforcement effect was not considered. Such an interphase is manifest given high surface area of graphene and improved interface adhesion of FGs with PS. The results indicate that functionalized graphene prepared through simple sonochemical route behaves as a very effective filler for application in high performance polymeric composites.

Acknowledgements

The work was funded in part by the National Natural Science Foundation of China (no. 51073082 and 51373088) and Open Project of Collaborative Innovation Center of Green Tyres & Rubber, Qingdao University of Science & Technology (No. 2015GTR0021).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra24479c

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