Constructing 3D Graphene Networks in Polymer Composites for

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Constructing 3D Graphene Networks in Polymer Composites for Significantly Improved Electrical and Mechanical Properties Peng Wang,†,‡ Haodan Chong,† Jiajia Zhang,†,‡ and Hongbin Lu*,†,‡ †

State Key Laboratory of Molecular Engineering of Polymers, Department of Macromolecular Science, Collaborative Innovation Center of Polymers and Polymer Composites, Fudan University, 220 Handan Road, Shanghai 200433, China ‡ Shanghai Xiyin New Materials Corporation, 135 Guowei Road, Shanghai 200437, China S Supporting Information *

ABSTRACT: Graphene-based polymer composites with superior electrical and mechanical performance are highly desirable because of their wide range of applications. However, due to the mismatch between charge jumping and the load transfer of adjacent graphene sheets, it remains difficult to achieve significant, simultaneous improvements in electrical and mechanical properties of graphene−polymer composites. To overcome this issue, we here propose an effective strategy to constructed unique 3D conductive networks in which the compatibility of graphene and polymer can be improved by controlled decoration of few-defect graphene sheets, while segregated graphene networks retain good charge-jumping capability. The final composites exhibit an ultra-low electrical conductive percolation threshold of 0.032 vol % and an ultrahigh electrical conductivity of 60 S/m at only 2.45 vol %, superior to most of the reported results. They also reveal significantly improved thermodynamic properties, tensile strength, and toughness. We believe that such a simple, industrially feasible method contributes to boost the development of high-performance, functional graphene−polymer composites. KEYWORDS: few-defect graphene, graphene networks, controlled decoration, polymer composites, compatibility

1. INTRODUCTION Polymer composites with simultaneously improved electrical conductivity and mechanical properties are highly desired in many applications such as electromagnetic interference shielding,1,2 static electricity dissipation,3,4 sensor,5,6 and mechanically strengthened functional materials.7,8 For such composites, it is a critical step to construct 3D conductive particle networks. Compared with carbon nanotubes and carbon black, graphene has some advantages, including low cost, high aspect ratios, large specific surface areas, and outstanding electrical and mechanical properties, which make it an ideal filler for improving electrical and mechanical properties of polymer materials.9−12 Owing to huge aspect ratios, typically >103, significantly reduced electrical percolation thresholds have been demonstrated in graphene-based polymer composites.8,13−15 Nevertheless, it remains a key challenge to establish an eco-friendly, industrially viable route that can produce graphene-polymer composites with significantly improved electrical and mechanical properties.3,16 In polymer composites containing uniformly dispersed graphene sheets, the improvement in electrical conductivity is limited by the weak intersheet electron-tunneling transfer. Typical electrical conductivities are in order of magnitude of ∼1 S/m (polystyrene composite with 2.5 vol % of graphene8). Increasing the content of graphene is effective for enhancing © XXXX American Chemical Society

the electrical conductivities of composites, but drastically increased viscosities also bring some technical challenges in composite processing. This makes it a critical issue how to effectively construct electrically conductive particle networks in polymer that can maximize electrical and mechanical properties of polymer composites. Recently, constructing interconnected 3D graphene networks in polymer matrix has become a hot research topic, because it can significantly enhance the electrical and mechanical performance of polymer composites.10−12,15,17−22 3D graphene networks can be created by chemical vapor deposition on nickel foam templates that may be removed by etching12,17,18 and direct freeze-drying of graphene oxide hydrogels.15,19,20,22 Combining such 3D networks into polymer matrices, the resultant graphene-based polymer composites have exhibited superior electrical and mechanical performance. Jia et al. reported cellular-structured graphene foam (GF) reinforced epoxy composites, which exhibited an ultra-high electrical conductivity of 3 S/cm as well as 53 and 38% of increases in flexual modulus and strength with only 0.2 wt % GF,12 respectively. Wang et al. demonstrated 3D interconnected Received: May 24, 2017 Accepted: June 12, 2017 Published: June 12, 2017 A

DOI: 10.1021/acsami.7b07328 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Figure 1. (a) Schematic process for the preparation of m-GNS−PA6 composites, (b) AFM image of m-GNS sheet together with its thickness profile and a height profile along the blue line, (c) optical photograph of the m-GNS−PA6 composite sheets (0.5 mm in thickness) that were used for electrical conductivity measurement, and (d) good conductivity of m-GNS−PA6 composite sheet (1.45 vol % m-GNS) shown in a digital multimeter (the inset shows a conductive path containing the m-GNS−PA6 composite sheet).

good electrical properties.26 Covalent decoration of graphene contributes to the compatibility of the graphene network with polymer but always accompanying drastically reduced electrical properties,27 so that it is hard to reveal desired electrical conductivities in resultant composites. In this sense, how to balance the defect structure of graphene and the compatibility with polymer determines the performance optimization of composites to a large extent. We here demonstrate a simple, effective method to optimize the segregated graphene network structure so that the electrical and mechanical properties can be simultaneously improved. In reported studies, the segregated structures were mainly constructed in polymer matrices with high melt viscosity, such as ultra-high molecular polyethylene, cross-linked polymers, etc.28 The construction of segregated structures in semicrystalline polymer matrices with low melt viscosity had been rarely studied. Here, we create segregated graphene networks in a semicrystalline polyamide 6 (PA6) with low melt viscosity, which belongs to an important class of engineering plastics. To improve the electrical properties, we employed fewdefect graphene sheets to construct electrically percolated networks in PA6. Meanwhile, to optimize the mechanical performance, a limited amount of 3-aminopropyltriethoxysilane (APTES) with active functional groups (−NH2) were used to improve the compatibility of the percolated graphene network with PA6. Different from the reported methods, such a strategy is simple and industrially feasible, revealing outstanding advantages over those reported routes that adopt polymerdecorated or undecorated graphene sheets.

graphene aerogel reinforced epoxy resins with excellent electrical conductivities, high mechanical properties, and fracture toughness.19 Li et al. prepared thermally annealed anisotropic graphene aerogel (TAGAs) reinforced epoxy composites, exhibiting highly anisotropic mechanical and electrical properties and excellent electromagnetic interference (EMI) shielding efficiencies at low graphene loadings.20 Despite these advances, however, the template and freezedrying techniques involve multistep and energy-intensive operations, practically difficult to be applied in large-scale production of polymer composites. In contrast, constructing segregated graphene networks in polymer composites have been deemed to be a promising alternative to achieve highly electrically conductive, mechanically strong polymer composites.3 Segregated percolation networks may be created through mechanically mixing conductive fillers with polymer granules and be preserved in final composites after hot press, which improves the charge-transfer efficiency in composites, due to effective intersheet contacts.2,9,13,23 However, this also brings about additional issues, for example, impaired mechanical properties of composites because the particles aggregated at the surface of polymer granules introduce extra holes or free volume. These holes become defects or stress-concentrated points that largely decrease the strength and toughness of composites.3 To overcome these issues, double-percolation networks and polymer-grafted graphene sheets have been attempted and exhibited improved electrical and mechanical properties.4,24,25 Nevertheless, these improvements are still unsatisfactory. In practice, to optimize the electrical and mechanical properties of composites, the quality of graphene itself (electrical conductivity) and the intersheet contact resistance need to be considered simultaneously.3,16 Pristine graphene, without or with very few defects, has high electrical conductivities but tends to aggregate in composites, which usually impairs the mechanical properties of composites due to weak intersheet interactions (van del Waals force), despite

2. RESULTS AND DISCUSSION 2.1. Functionalization of Few-Defect Graphene. We employed chemically expanded graphite with a low extent of oxidization (LOEG), which has a specific surface area of >800 m2/g and uniform oxidation level in spacial, to prepare functionalized graphene sheets following our previously reported method.29 Different from reduced graphene oxide B


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ACS Applied Materials & Interfaces (RGO), the present method does not involve any chemical or thermal reduction process, but it still may effectively adjust the interface compatibility of graphene with polymer while preserving good graphene lattice structures and electrical conductivity. The chemical composition of LOEG was determined through X-ray photoelectron spectroscopy (XPS) (Figure S2) and the high-resolution core-level C 1s spectrum is given in Figure S2b. Its deconvolution shows the presence of three carbon bonds: C−C (284.5 eV), C−O(286.6 eV), CO (287.4 eV),30−32 in which the C−O and CO groups reveal 25.2% and 7.5% of contributions, respectively, to the whole peak area. A high C-to-O atomic ratio (3.1) indicates its low extent of oxidation, compared to traditional graphene oxide (typically C/O ratio is 2.1).30,31 We functionalize LOEG, rather than exfoliated graphene sheets, which enhances the production efficiency of modified graphene sheets (m-GNS), due to significantly improved washing and filtration processes. The Fourier transform infrared spectroscopy (FTIR) spectra of LOEG and m-LOEG and the corresponding thermogravimetric analysis (TGA) curves, as shown in Figure S3 in the Supporting Information, verified that APTES molecules were covalently grafted on the surface of LOEG. After the functionalization of LOEG, m-GNS sheets were exfoliated through 1 h of tip sonication in ethanol, and directly used to construct segregated electrically conductive networks in composites. The detailed description is presented in the Experimental Section. Figure 1a schematically shows the preparation process of m-GNS/PA6 composites. Mechanical mixing in ethanol enables m-GNS sheets to be adsorbed on the surface of PA6 granules through hydrogen bonding interaction. After drying, segregated conductive networks are formed throughout m-GNS/PA6 composites by hot compression. The resulting m-GNS/PA6 sheet (Figure 1c) exhibits good electrical conductivity (Figure 1d). The AFM image of m-GNS sheets (Figure 1b) shows that the thickness and lateral size of the m-GNS sheet are ∼2.64 nm and ∼4 um, respectively (Figure S4). Given the presence of the grafted APTES molecules, we conclude that after 1 h tip sonication, the m-LOEG had been exfoliated into single- or few-layer m-GNS sheets, which allows us to subsequently construct electrical conductive networks in composites. 2.2. Constructing Segregated Graphene Networks. PA6 powder and m-GNS were first mixed in ethanol by mechanical stirring at room temperature. After drying, m-GNS sheets were adsorbed on the surface of PA6 granules by hydrogen bond interactions that arise from the residual oxygencontaining groups, -NH2 groups of m-GNS and amide groups of PA6.16,19,33 In control experiments, different amounts of noncovalently modified GNS sheets (denoted as GNS), which were obtained by water-phase exfoliation of LOEG with the assistance of surfactant OP-15 (the AFM image of GNS sheets was shown in Figure S5), were also deposited on the surface of PA6 granules following the similar procedure. Field emission scanning electron microscopy (FESEM) images (Figure 2) unveil their significant difference in morphology. For the PA6 ganule coated with 0.19 vol % GNS (the content determination was seen in experimental section), a relative smooth surface (Figure 2a), in which some detached stacked GNS aggregates can be identified in the magnified image (Figure 2b, indicated by the red dot circle). In contrast, the PA6 granule containing m-GNS sheets reveals a coarse surface, where m-GNS sheets are loosely stacked (Figure 2c,d). Apparently, owing to the presence of APTES, these m-GNS sheets distributed over the whole surface of the PA6 granule failed to form layer-by-layer

Figure 2. FESEM images of (a, b) GNS−PA6 and (c, d) m-GNS− PA6 composite powder; (b, d) the amplification of panels a and c, respectively. The filler content is 0.19 vol %.

stacking similar to that of the GNS/PA6 granule. This suggests that although the covalent grafting decreased the number of oxygen-containing groups of GNS, the NH2 groups of APTES and the residual oxygen-containing groups can still ensure the effective adhesion between m-GNS and PA6 granules through relative strong hydrogen bonding interaction. In addition, the grafted APTES can prevent aggregation of GNS sheets, which is beneficial to improve the compatibility of graphene and PA6. To confirm the exfoliation of m-GNS sheets in composites and characterize their thickness, we have extracted m-GNS sheets from m-GNS/PA6 composite powder through dissolution, washing and centrifugation (see the method in Supporting Information). The collected m-GNS sheets (denoted as d-mGNS) were redispersed in alcohol by mild bath-sonication for 5 min prior to characterization. The AFM image of d-m-GNS sheet (Figure S6) shows a thickness of 3.14 nm, close to that of the neat m-GNS (2.64 nm). This indicates that APTES can indeed suppress the aggregation of graphene on PA6 granules. Given that two kinds of graphene-coated PA6 granules were prepared following the same procedure, their morphological difference would primarily originates from the grafted APTES molecules. Although their molecular size and number are limited, only 8 wt % of the m-GNS weight (Figure S2b, detailed analysis about the content of APTES grafted onto m-GNS sheets was also presented in Supporting Information), APTES molecules enable m-GNS sheets to be well-separated, uniformly deposited on the granule surface. This actually constitutes an important prerequisite to balance electrical and mechanical properties of segregated network-based composites. We constructed GNS- and m-GNS-based segregated networks in PA6 composites by hot press at 235 °C with a pressure of 13 MPa for 10 min (the hot press temperature is higher than the melting point of PA6, 218.5 °C that was obtained from the DSC curve in Figure S7). This operation can facilitate the flow of PA6 melt while preserving the 3D graphene networks preformed in composites. Optical observaC


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ACS Applied Materials & Interfaces tion exhibits the significant structural difference between two kinds of composites. As shown in Figure 3a,b, even at the solid

Figure 4. FESEM images of the fractured surface of (a, b) GNS−PA6 and (c, d) m-GNS−PA6 composites with segregated network structures; (b, d) amplification of panels a and c, respectively. The filler contents are 0.19 vol %. Marks A and B in panels a and d represent the graphene-rich region and PA6-rich regions, respectively.

but also large graphene aggregates (marked by red dot cycles) are clearly visible (Figure 4b). Such structural character indicates again that the noncovalently modified GNS is actually incompatible with PA6, which prevented polymer melt from infiltrating into the interlayer space between graphene sheets and also induced the formation of voids. In comparison, in the fracture surface of the m-GNS/PA6 composite, there exist hardly discernible, graphene-induced voids (Figure 4c), indicating the good compatibility between m-GNS and PA6. Especially, the structural character in the graphene-rich region (marked by two white dot lines in Figure 4d, constituting the wall of percolated networks) of the m-GNS/PA6 composite is significantly different from that of the GNS/PA6 composite (marked by a red dot cycle). For the former, the individual mGNS sheets (indicated by the red arrows in Figure 4d) are well separated by polymer while preserving the intersheet contact to some extent. No obvious voids or graphene aggregates are observed in the whole fracture surface. Such morphological feature implies the grafted APTES molecules not only inhibited restacking of graphene sheets in the process of hot press but also facilitated the infiltration of PA6 melt into the interlayer space of graphene sheets. Given the ability of amine groups of APTES to react with carboxyl groups of PA6 at high temperature, we speculate that some PA6 chains would probably be grafted on the surface of graphene, which in turn facilitates the infiltration of PA6 melt into the graphene-rich region. To verify this, we used formic acid to remove the ungrafted PA6 molecules, similar to the previously mentioned method, and the obtained sample is denoted as g-m-GNS/PA6. It was characterized with TGA in nitrogen atmosphere. Figure 5a presents its weight loss curve from 50 to 800 °C, in which the results of m-LOEG and the neat PA6 are also included for comparison. It is seen that a weight loss of ∼28 wt % occurred in the temperature range of 350−500 °C, which is ascribed to the thermal decomposition of the grafted PA6 chains.34 Figure 5b gives its FTIR spectrum, in which two characteristic bands at 1640 and 1539 cm−1 are attributed to the stretching vibration of the amide CO group and the bending vibration of the N− H bond or the stretching vibration of the amine C−N bond, respectively.34 These provide the evidence of PA6 molecules

Figure 3. Optical images of (a) GNS−PA6 and (b) m-GNS−PA6 composites; TEM images of (c) GNS−PA6 and (d) m-GNS−PA6 composites. The filler contents are 0.19 vol %.

content of 0.19 vol %, interconnected particle networks can still be observed in two composites. For noncovalently modified graphene-based composites (GNS/PA6 composite), however, some large aggregates exist (indicated by the red dot circles in Figure 3a) and the particle network reveals poor connectivity. This apparently is related to the initial stacking structure of GNS on the surface of PA6 grannules, which prevents the infiltration of PA6 into the interlayer space of the stacked GNS layer to some extent. This can be further confirmed from their TEM images, as shown in Figure 3c. Apparently, there exist some unoccupied holes or interstitial spaces within the wall of graphene networks, where few or almost no PA6 molecules are observed. In contrast, the m-GNS/PA6 composite reveals a clearer 3D network structure and good connectivity (Figure 3b), in which few large aggregates appear, implying that PA6 molecules can effectively infiltrate into the interstitial space of graphene wall of networks (Figure 3d). This not only reflects the good affinity of m-GNS to PA6 but also indicates that the stacking structure of graphene sheets on polymer granules has an important influence on the resulting network structures. Compared with the layer-by-layer stacking in GNS/PA6 granules, the loose, disordered stacking in m-GNS/PA6 granules is beneficial for infiltration of polymer melt and reduction of interstitial spaces or stress concentrated points. These are apparently critical for optimizing the comprehensive performance of composites containing segregated graphene networks. FESEM images of the fracture surface provide further evidence for the structural difference of two kinds of composites. To avoid structural variations occurred during preparing the samples, the compression molded samples were quickly fractured after immersing in liquid nitrogen for 10 min. As shown in Figure 4a, not only can 10−30 μm long voids (marked by green dot cycles) be seen in GNS/PA6 composite, D


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Figure 5. (a) TGA curves of m-LOEG, g-m-GNS−PA6, and neat PA6 and (b) FTIR spectra of m-LOEG, g-m-GNS−PA6, and neat PA6.

Figure 6. (a) Electrical conductivity as a function of filler content and (b) the fitted results of experimental data for GNS−PA6 and m-GNS−PA6 composites using the percolation law.

covalently grafting to the graphene surface. Given a limited amount of grafting chains, the modified graphene sheets can still preserve effective intersheet contact in composites, although they facilitated the infiltration of polymer melt and inhibited aggregation of graphene sheets. Next, we discuss the difference in electrical and mechanical properties by comparing two kinds of segregated graphene networks. 2.3. Ultra-low Percolation Threshold and High Electrical Conductivities. The huge aspect ratio of graphene (>103) provides an intriguing approach to achieve highly electrically conductive polymer composites with low filler contents. However, the reported results to date remain unsatisfactory, and typical percolation thresholds (ϕc) are in the range of 0.1−0.9 wt %3 for segregated graphene networks. For graphene/PA6 composites, the reported lowest ϕc and highest electrical conductivity are 0.39 wt %35 and 0.65 S/m at 3.0 wt %,36 respectively. The limited electrical properties primarily arise from several factors. (1) The presence of a large amount of defects reduces electrical conductivities of reduced GO (rGO); (2) the aggregation of rGO impairs the chargetransfer efficiency in 3D networks; (3) introduced polymer chains inhibit the charge transfer or jumping between adjacent graphene sheets to some extent, despite improved graphene− polymer compatibility.3,25 In contrast, the graphene sheets we used have few defects and a limited amount of grafting polymer chains, which is expected to afford a good solution to highly electrically conductive, mechanically strong graphene-nylon composites. We first determined the electrical conductivities of GNS and m-GNS films without PA6. Two films were prepared by vacuum filtration with a poly(tetrafluoroethylene) filter

membrane (0.45 um pore). The electrical conductivities are 1.85 × 104 and 4.5 × 103 S/m for GNS and m-GNS films, respectively (four-probe method). This indicates that the covalent decoration decreased the electrical conductivity of m-GNS, but an electrical conductivity of up to 103 S/m would be effective for improving the electrical performance of mGNS/PA6 composites. Figure 6a presents the electrical conductivities of two kinds of composites with different graphene contents, GNS−PA6 and m-GNS−PA6 composites. Both of them reveal more than 10 orders of magnitude of increases as the graphene content increases by 0.19 vol %, a typical percolation behavior. We use the classical percolation theory to describe the percolation behavior of PA6 composites, σ = σ0(ϕ − ϕc)t

(1)

where σ and σ0 are electrical conductivities of composite and filler, respectively, ϕ is the volume fraction of filler and t is a critical exponent.37−40 Fitting the above equation gives the ϕc values of two kinds of composites, 0.08 vol % for GNS/PA6 and 0.032 vol % for m-GNS/PA6 (Figure 6b), and the exponents (t) are 2.64 and 2.81 for GNS/PA6 and m-GNS/ PA6, respectively. Typically, t ≈ 1.6−2.0 and t ≈ 1.1−1.3 are for three-dimensional (3D) and two-dimensional (2D) conductive networks.3,40 Here, the t values of two kinds of graphene/PA6 composites are close to 2, suggesting the formation of 3D conductive networks. Nevertheless, they are actually higher than the theoretical prediction, which might arise from the specific structure and properties of GNS and mGNS.2,40 Notably, although there exist significant aggregation in E


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Figure 7. Storage modulus of (a) GNS−PA6 and (b) m-GNS−PA6 composites with different filler contents; tan δ of (c) GNS−PA6 and (d) mGNS−PA6 composites with different filler contents as a function of temperature.

GNS/PA6 composites, as shown in Figures 3 and 4, their ϕc is still lower than those of most graphene-based polymer composites(Table S1), including composites prepared by the latex self-assembly (0.15 vol %).39 This would be attributed to the high electrical conductivity of few-defect graphene and the small intersheet contact resistance. For m-GNS/PA6 composites, their ϕc is nearly 1 order of magnitude lower than the best value of the reported graphene−PA6 composites.35 Upon increasing graphene contents to 2.45 vol %, the m-GNS/PA6 composite reveals an electrical conductivity of 60 S/m, which is 3-fold higher than that of the GNS/PA6 composite and 2 orders of magnitude higher than those of reported results within similar graphene contents.36,41 2.4. Significant Mechanical Reinforcements. Previous studies had indicated that the microvoids along the segregated conductive channels always deteriorated the mechanical properties of composites with segregated structures,3,4,10,42 which has become a critical challenge that inhibits their practical applications. It is thus urgently necessary to establish an effective route that can produce graphene-based composites with excellent electrical and mechanical properties. In the following, we discuss the mechanical properties of GNS/PA6 and m-GNS/PA6 composites, and discuss possible mechanisms. Figure 7a,b shows the variation in storage moduli of GNS/ PA6 and m-GNS/PA6 composites with temperature, respectively, in which significant enhancements are observed. For GNS/PA6 composites, the storage modulus increases 36.8, 49.7 and 65.1% with the increase of graphene contents from 0.19, 0.58 to 1.45 vol % (Table 1), respectively. In comparison, mGNS/PA6 composites show larger improvements, for which the storage modulus increases 45.8, 72.2 and 92.0%, respectively (Table 1). The higher storage modulus for mGNS/PA6 composites primarily arises from better graphene dispersion and stronger interfacial interaction. The improved

Table 1. Thermomechanical Properties of GNS−PA6 and mGNS−PA6 Composites sample PA6 0.19 0.58 1.45 0.19 0.58 1.45

vol vol vol vol vol vol

% % % % % %

GNS GNS GNS m-GNS m-GNS m-GNS

Tg (°C)

storage modulus at 45 °C (MPa)

56.8 54.2 58.8 56.3 59.8 60.6 66.9

1789 2448 2678 2953 2609 3080 3435

graphene dispersion increases the specific surface area of graphene sheets so that graphene surfaces can take better effect in confining the motion of polymer chains or segments. The confinement effect on segment motion of a polymer can be characterized directly by the glass transition temperature (Tg). For GNS/PA6 composites, as shown in Figure 7c, no obvious Tg increase is observed regardless of what graphene content. Here, we determine the Tg of composites using the peak temperature of tan δ curves. For m-GNS/PA6 composites, however, the Tg of the composites increases 3, 4, 10 °C for 0.19, 0.58, and 1.45 vol % graphene contents, respectively, compared to that of the neat PA6. Increased confinement effect on segment motion reflects to some extent the contribution of interface interactions in composites. However, it should also be noted that for semicrystalline polymers like PA6, the crystal structure and crystallinity could likewise change the segment motion of polymer in amorphous domains, besides the confinement effect of filler surfaces.43,44 To clarify the possible influence of crystal structures, we further investigated the crystallization behavior of neat PA6 and graphene/PA6 composites using differential scanning calorimetry (DSC). As shown in Figure S7, the neat PA6 reveals a melting peak at 218.5 °C, which can be ascribed F


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Figure 8. Stress−strain curves of (a) GNS−PA6 and (b) m-GNS−PA6 composites with different filler contents.

Figure 9. Tensile strength (a), tensile modulus (b), and toughness (c) of GNS−PA6 and m-GNS−PA6 composites. Tensile strength is defined as the stress at the fracture point, tensile modulus is obtained from the slopes of linear areas in stress−strain curves, and toughness is calculated by integrating stress−strain curves.

arise from the formation of the new γ-form crystal in PA6 composites to some degree. To verify this, we calculated the crystallinities of all materials according to the DSC data. For the neat PA6, its crystallinity is 26.3%. The addition of GNS and mGNS sheets did not cause obvious changes in crystallinity. That is, for the GNS/PA6 composites, the crystallinities are 26.0, 26.2 and 26.4% for 0.19, 0.58, and 1.45 vol % GNS contents, respectively, while the crystallinities of m-GNS/PA6 composites are 26.4, 26.7 and 27.3% for 0.19, 0.58, and 1.45 vol % mGNS contents, respectively. These results suggest that the modulus improvements of PA6 composites would primarily stem from the interface effect of graphene. The greater improvements observed in m-GNS/PA6 composites reflect better graphene dispersion and stronger interfacial interaction, as opposed to GNS/PA6 composites. As aforementioned, the amide groups of APTES is able to react or form hydrogen bonds with PA6 during the hot press, which engendered strong interface interactions in m-GNS/PA6 composites. In comparison, the interface interaction is weaker in the GNS/PA6

to melting of the thermodynamically stable α-form crystal with hydrogen bonds between antiparallel chains.34 All the GNS/ PA6 and m-GNS/PA6 composites with different filler contents show similar DSC curves in which besides the main melting peak at 218.5 °C, a small shoulder peak at 209.5 °C can be ascribed to the γ-form crystal with hydrogen bonds between parallel chains.34 This means that the presence of graphene does not significantly change the α-crystal structure of PA6, but the appearance of the new γ-crystal indicates the nucleation effect of graphene in these composites.45 For the three groups of GNS/PA6 composites, they all show no obvious change in Tg compared with the neat PA6. In other words, formation of the new γ-form crystal in PA6 composites did not significantly affect the Tg of PA6 and the increase of Tg for the m-GNS/PA6 composites is thus ascribed to the strong interfacial interactions that restrict the mobility of polymer chains.33,43,44 Given that both crystal form and crystallinity can influence the storage modulus of the composites, the improvements in storage modulus of GNS/PA6 and m-GNS/PA6 composites may also G


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exhibit the best balance in electrical and mechanical performance among the PA-based composites and even comparable to the polymer composites reinforced by 3D graphene aerogels,20,22 indicating that the present strategy is beneficial for developing high-performance, functional graphene-based composites.

composites; that is, no obviously restricted chain motion occurred in the interface region (unchanged Tg). The representative stress−strain curves of the neat PA6 and polymer composites are shown in Figure 8a,b. The neat PA6 revealed a typical yield behavior with increasing stress during tension and the corresponding tensile modulus, tensile strength and toughness are 484 MPa, 59.7 MPa and 18.4 MPa m1/2, respectively (Figure 9). For GNS/PA6 composites, their tensile behavior changes significantly. No yielding was observed but the tensile strength and toughness decrease gradually from 60.8 to 40.3 MPa and 7.7 to 1.4 MPa m1/2 with increasing filler contents from 0.19 to 1.45 vol % (corresponding to a 32.5% decrease in tensile strength and a 92.4% decrease in toughness relative to those of the neat PA6). This indicates the incompatibilities between GNS and PA6 molecules. Because the polymer melt was unable to well infiltrate into the interlayer space between GNS, a large amount of microvoids or stress concentrated points resulted thereby and severely impaired their mechanical strength. For 0.58 vol % m-GNS−PA6 composite, the tensile strength surprisingly increased to 84.5 MPa (Figure 9), a 41.5% of increase relative to that of neat PA6, along with a typical yield behavior and an elongation at break up to about 40% (Figure 8b). This implies a 15.8% of increase in fracture toughness, relative to that of the neat PA6. When the m-GNS content increased to 1.45 vol %, a small decrease in tensile strength (4.6 MPa, corresponding to a 7.7% of decrease, relative to that of the neat PA6) was observed (Figure 9), in sharp contrast with the reported results.35 Apparently, the significant difference between two kinds of composites relates to their different conductive network structures (Figures 3, 4). The tensile properties of polymer composites depend largely on the load-transfer efficiency between polymer and fillers.33,34,46,47 For m-GNS/PA6 composites, the significantly improved mechanical properties mainly arise from two key points: (1) the good dispersion of mGNS sheets in the conductive channels and strong interface interactions, (2) the improved compatibility between m-GNS and PA6 that is critical for decreasing the number of microvoids or stress-concentrated points. The corresponding mechanistic models are illustrated in Figure 10. To elucidate the importance of the above two points, we also compare the properties of mGNS/PA6 composites with those of PA-based composites and typical 3D graphene/polymer composites reported in the literature. As shown in Table 2, m-GNS/PA6 composites

3. CONCLUSIONS We have proposed a new strategy to develop high-performance, functional graphene-based polymer composites through constructing well-controlled segregated graphene networks in polymer matrix. To chemically modify graphene sheets and to optimize the mechanical performance of graphene-polymer composites, some amount of oxygen-containing groups should be introduced to graphene sheets. To preserve the good electrical properties of graphene sheets and the subsequent mGNS sheets, the amount of introduced oxygen-containing groups should be limited. Based on those considerations, we employ LOEG. After the covalent modification of LOEG with APTES, the number of interstitial voids or stress-concentrated points in composites can be significantly decreased, so that highly electrically conductive, mechanically strong graphene/ PA6 composites resulted. This is distinctively different from the results reported previously and thus suggests a simple, industrially viable route to fabricate high-performance graphene-based composites. The limited amount of -NH2 groups grafted on GNS contribute to improve the compatibility of segregated graphene networks with PA6, and also strengthen the interface interactions of composites through the reaction between -NH2 groups and −COOH groups of PA6. This strategy exhibits good balance between electrical and mechanical properties of composites. The resulting composites showed a percolation threshold as low as 0.032 vol % and an electrical conductivity as high as 60 S/m at 2.45 vol % graphene content, which are superior to the results reported previously. Besides, m-GNS/PA6 composites revealed a 10 °C of increase in the glass transition temperature, and a 92.0% of increase in elastic storage modulus, relative to those of the neat PA6. The tensile strength and fracture toughness increased 41.5% and 15.8%, respectively, when the m-GNS content was 1.2 wt %. We believe that the present strategy affords a feasible solution for developing mechanically strong, tough, functional graphenebased polymer composites. 4. EXPERIMENTAL SECTION 4.1. Materials. Natural graphite flake (∼500 um) was purchased from Sigma. Concentrated sulfuric acid (H2SO4, ∼ 98%), hydrogen peroxide (H2O2, 30%), potassium permanganate (KMnO4) were purchased from Jiangsu Tongsheng Chemical Company. 3-aminopropyltriethoxysilane (APTES), octaphenyl polyoxyethyene (OP-15), ethanol were obtained from Sinopharm Chemical Reagent Company. PA6 powder (T-225, 100 mesh, supplied from Shanghai Zhenwei Composite Materials Co., Ltd.) with a melting point of 225 °C was used as the polymer matrix. All chemicals were used as received without further purification. 4.2. Preparation of Graphene Sheets and Modified GNS Sheets. We employed chemically expanded graphite with a low extent of oxidation (LOEG) as the starting material, which was prepared following our previous method.29 To prepare few-defect graphene, LOEG (2 g) was added into a 500 mL beaker containing deionized water (250 mL) and surfactant of OP-15 (2.5 g). After (tip) sonication for an hour, the GNS aqueous solution was obtained. For the preparation of GNS decorated with APTES (m-GNS), LOEG (2 g) was first added into a 500 mL three-neck flask containing APTES (15 g) and ethanol (250 mL) followed by refluxing the mixture under

Figure 10. Mechanistic model of reinforcement for (a) GNS- and (b) m-GNS-based composites. H


PA6

I

mechanical mixing and hot compression

interfacial polymerization latex self-assembly backfilling infiltration infiltration

melt mixing

PMMA is poly(methyl methacrylate).

PA6

m-GNS

a

PA6,6 PMMAa PMMAa epoxy epoxy

PA12

functionalized graphene SWCNT rGO graphene aerogel graphene foam graphene aerogels

functionalized MWCNTs graphene

melt compounding

PA6

PA12

solution blending

PA6

melt compounding

in situ intercalation polymerization melt mixing

PA6

PA6

in situ polymerization

PA6 PA6 PA6

RGO foliated graphite functionalized graphene functionalized graphene expanded graphite

graphene and carbon fiber (CF) MWCNTs

in situ polymerization in situ anionic ring-opening polymerization in situ polymerization in situ polymerization in situ polymerization

processing

PA6 PA6

polymer

GO GO

filler

126% (10 wt %) − 54.1% (0. 35 wt %)

− − − 12 °C (10 wt %) − 8.1 °C (0.35 wt %)

− − − 74% (10 wt %) −

92.0% (3 wt %)

− 30% (4.0 wt %) − − 67.4% (0.8 wt %)

62.8% (0.35 wt %)

−

−

−

− − − − − 41.5% (1.2 wt %)

− 15 °C (0.5 wt %) − 31 °C (0.2 wt %) ∼10 °C (0.8 wt %) 10.1 °C (3 wt %)

−

−

29% (0.1 wt %)

− − −16.7%

− − −

− − −

120% (0.1 wt %) 88.0% (1 wt %)

tensile strength (highest enhancement)

− −

Tg (highest enhancement)

− −

storage modulus (highest enhancement)

101% (3 wt %)

30% (2 wt %) − − − −

74.1% (0.35 wt %)

−

−

−

−

−

300% (0.1 wt %)

− − 156%

137% (0.1 wt %) 66.5% (1 wt %)

tensile modulus (highest enhancement)

Table 2. Comparison of the Mechanical and Electrical Performance of Polymer Composites with Different Nanofillers

−

0.3 vol %

1−5 wt %

0.5−1 wt %

3 vol %

0.75 vol %

−

1 wt % 0.75 vol % 0.19 vol %

− −

percolation threshold

∼ 0.1 (5 wt %) 1−2 wt % 64 (2.7 vol %) 0.15 vol % 0.859 (2.5 vol %) − 300 (0.2 wt %) 0.05 wt % 980 S/m (0.8 wt lower than 0.2 %) wt % 60 (2.45 vol %) 0.032 vol %

6.7 × 10−2 (1.38 vol %) −

1.43 × 10−2 (1 wt %) 0.12 (10 wt %)

16.7 (10 vol %)

0.01 (2 vol %)

0.65 (3 wt %) 0.1 (3 vol %) 6.84 × 10−2 (0.26 vol %) −

− −

highest conductivity (S/ m)

this work

55 39 22 12 20

33

54

53

52

51

50

49

36 41 35

34 48

ref

ACS Applied Materials & Interfaces Research Article


Research Article

ACS Applied Materials & Interfaces nitrogen atmosphere at 62 °C for 8 h (The chemical reaction occurs between APTES and LOEG is shown in Figure S1 in the Supporting Information). After the reaction, the product was washed with ethanol and deionized water three times to remove the residual modifier. Subsequently, the modified LOEG (m-LOEG) was added into a 500 mL beaker containing ethanol (250 mL) followed by tip sonication of an hour, giving the m-GNS suspension in ethanol. 4.3. Preparation of GNS−PA6 and m-GNS−PA6 Composites. PA6 powder (15 g) was added into a 500 mL beaker containing different amounts of m-GNS and ethanol (to convert the mass fraction of graphene sheets content to a volume fraction, the density of the graphene sheets and PA6 powder are taken as 2.227 and 1.05 g/cm3, respectively. The density of PA6 powder was obtained from Shanghai Zhenwei Composite Materials Co., Ltd.). After vigorous stirring for 2 h, the product was collected by filtration and dried in a vacuum oven at 80 °C for 24 h to remove the residual solvent. With a 15 min of preheating, the m-GNS coated PA6 particles were subsequently compressed into final composite samples under 13 MPa and 235 °C for 10 min. Finally, the m-GNS/PA6 composites with different amount of m-GNS were obtained. For comparison, the GNS/PA6 composites were also prepared in which the GNS aqueous solution was first washed by deionic water three times through centrifugation/ redispersion at 12000 rpm to remove the residual OP-15. Excess water was removed by successive centrifugation/redispersion using ethanol two times. The wet GNS was then added to a suitable amount of ethanol followed by tip sonication for 15 min to obtain the GNS suspension. After that, the preparation procedure for GNS/PA6 composites was the same as the m-GNS/PA6 composites. 4.4. Characterization. The morphology of PA6 powder and composites were characterized with Field-Emission Scanning Electron Microscope (FESEM, Ultra 55, Zeiss). The compression molded sheets of the composites were immersed in liquid nitrogen for 10 min, and then quickly fractured. Transmission electron microscopy (TEM) was performed on an FEI Tecnai G2 TF20 Twin, with an accelerating voltage of 200 kV. The resulting composites were cut with a diamond knife to ultrathin sheets using a Leica FC7-UC7 microtome and collected on copper grids. Optical microscopy was performed on a Leica DM2500P Polarizing Microscope. The composites were cut into films with thickness of 20 um using a microtome under ambient condition. Thermogravimetric analysis (TGA) was performed (Mettler Toledo TGA 1) from 50 to 800 °C at a heating rate of 10 °C min−1 under nitrogen atmosphere. Fourier transform infrared spectroscopy was conducted on a Nicolet 6700 spectrometer at room temperature over a frequency range of 500−4000 cm−1. The melting behavior of the composites and neat PA6 were determined using a TA DSC Q2000 differential scanning calorimeter (DSC), with a heating rate of 20 °C/min under nitrogen atmosphere. The data were collected from 125 to 300 °C. Atomic force microscope (AFM) images of the exfoliated GNS and m-GNS were acquired using Bruker-Multimode 8 in tapping mode by dropping diluted solutions onto freshly cleaved mica. The electrical conductivity of P-G composites was measured using the four-probe tester (ST2263) and Keithley 6571B Electrometer. Dynamic mechanical properties of the composites were measured using a dynamic mechanical analyzer (Mettler Toledo, DMA/ SDTA861e) in a stretch mode with a frequency of 1 Hz. The sample size was 10.5 × 6 × 0.5 mm3. The temperature was swept from 0 to 100 °C. The tensile properties of graphene-PA6 composites were measured using a universal testing machine (CMT4104, SANS Group, China) at room temperature. A load cell of 10000N was employed and the tensile rate imposed was 5 mm/min. All the samples were cut into the dumbbell shape with a razor blade. More than five tests were conducted for each sample, from which corresponding mean values and standard deviations were derived.

■

■

Figures showing a schematic of a chemical reaction of APTES, XPS and high-resolution spectra, TGA curves and FTIR spectra of LOEG and m-LOEG, FESEM image of the exfoliated m-GNS sheets, AFM images, DSC curves of GNS/PA6 and m-GNS/PA6 composites with different filler contents. Tables showing a comparison of percolation thresholds of graphenebased nanocomposites. (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Hongbin Lu: 0000-0001-7325-3795 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are grateful for the financial support by the 973 project (grant no. 2011CB605702), the National Science Foundation of China (grant no. 51173027), and the Shanghai Key Basic Research Project (grant no. 14JC1400600).

■

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b07328. J


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Constructing 3D Graphene Networks in Polymer Composites for

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