Strong and ductile graphene oxide reinforced PVA nanocomposites

Accepted Manuscript Strong and ductile graphene oxide reinforced PVA nanocomposites Swarnima Kashyap, Swadesh K. Pratihar, Shantanu K. Behera PII:

S0925-8388(16)31495-5

DOI:

10.1016/j.jallcom.2016.05.162

Reference:

JALCOM 37682

To appear in:

Journal of Alloys and Compounds

Received Date: 4 February 2016 Revised Date:

2 May 2016

Accepted Date: 16 May 2016

Please cite this article as: S. Kashyap, S.K. Pratihar, S.K. Behera, Strong and ductile graphene oxide reinforced PVA nanocomposites, Journal of Alloys and Compounds (2016), doi: 10.1016/ j.jallcom.2016.05.162. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Strong and ductile graphene oxide reinforced PVA nanocomposites

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Swarnima Kashyap, Swadesh K Pratihar, Shantanu K Behera* Department of Ceramic Engineering Centre for Development of Nanomaterials National Institute of Technology, Rourkela, Odisha 769008 INDIA

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Abstract

Graphene oxide (GO) and its derivatives have been widely used as reinforcement in polymer matrices for enhanced mechanical and electrical properties. While graphene has higher elastic

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modulus and tensile strength than that of graphene oxide, it is the latter’s hydrophilic nature that provides a distinct advantage as a filler in the aqueous processing of polymer nanocomposites. In this work, we have tried to disseminate the effects of GO and reduced graphene oxide (rGO) as fillers in polyvinyl alcohol (PVA) matrix. At a loading of only 0.3 wt.% GO in PVA, we observed ~150% increment both in elastic modulus and tensile strength, which is unprecedented. The property enhancement was attributed to the homogeneous distribution of fillers, and strong interfacial interactions (hydrogen bonds) between the fillers and the matrix. Composites fabricated after in-situ

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reduction of GO fillers had elastic modulus comparable to that of pure PVA, but had considerably improved tensile strength. The failure strain of rGO based composites was much higher than that of

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both pure PVA and GO reinforced PVA composites, thus showing enhanced ductility. The microstructure of the PVA-rGO composites exhibited alignment of the fillers in the plane of the

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polymer film, with interfacial bonds between the fillers and matrix only at the rGO sheet edges. The distribution of GO fillers in the PVA-GO composites, however, was random with no preferred orientation. A higher degree of interfacial interactions and the homogeneous distribution of the fillers

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led to enhanced elastic modulus and strength, along with considerable ductility for the PVA-GO composites. Keywords

Graphene; PVA; Nanocomposites; Interfacial bonding; Mechanical properties.

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Corresponding author: [email protected]

Department of Ceramic Engineering, NIT Rourkela, Odisha 769008, INDIA Tel: 91-661-246-2214, Fax: 91-661-246-2201

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1. Introduction Graphene exhibits exceptional thermal, mechanical and electrical properties [1] for which it has gained tremendous scientific attention in the recent years. Its thermal conductivity (~5300 Wm-1K1

) and tensile elastic modulus (~1060 GPa) are amongst the best, and its tensile strength is comparable

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to that of carbon nanotubes [2,3]. Besides, its large theoretical specific surface area (2630 m2g−1) [4], high intrinsic mobility (200,000 cm2v−1s−1) [5,6], and optical transmittance (∼97.7%) merit attention for exploitation in many functional applications [7,8]. One of the most promising areas of this

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material is in polymer nanocomposites, where the exceptionally high elastic modulus and tensile strength of graphene or graphene derivatives [9] manifest in significant improvement in the mechanical properties of the composites. It is logical to expect that the interface-controlled properties

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are functions of the structure of graphene (or its derivatives) in a polymer nanocomposite. Exfoliation of graphene oxide produced by thermal or sonication methods provide morphology with varying lateral dimensions [10,11], which influences the degree of interfacial interactions and mechanical interlocking. Additionally, interfacial chemistry of the fillers also plays an important role in the mechanical and electrical properties of reinforced polymer composites. According to the LerfKlinowski model [12], graphene oxide (GO) produced by Hummers or modified Hummers method creates graphene sheets with hydroxyl and epoxy groups present in higher concentrations in the basal

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plane, and carboxylic acid groups at the edges of the sheets [13]. The presence of these functional groups makes GO hydrophilic, enabling its easy dispersion in polar solvents to form a stable colloid.

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This property of GO provides excellent opportunity for exploitation as reinforcement in aqueous based polymer composite fabrication.

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Polyvinyl alcohol (PVA) is a hydroxyl rich, non-toxic, biocompatible, and water-soluble polymer system, which can be processed by aqueous methods. It is an important engineering polymer

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with various applications in membrane technology [14], fuel cells [15], drug delivery [16], and shape memory effects [17]. It has often been considered as a model polymer matrix to study the reinforcement behaviour of GO based fillers. Several authors have reported the strong influence of filler loading and the synthesis routes on the mechanical behaviour with moderate to more than 100% increase in tensile strength with filler loadings in PVA composites having of 2-3%fillers [18,19,20,21,22,23]. Mo et al. fabricated GO based PVA composites with tensile strength of 280 MPa and elastic modulus of 13.5 GPa, albeit with 50 wt% loading of fillers [24]. These reports have shown mechanical property improvement with filler content that remained in the levels of a few percent or more. Progressively, however, comparable levels of mechanical property enhancements have reported for matrices with lower levels of filler loading, eg. 0.5-0.7 wt% with approaches, such as in-situ reduction of the GO, and covalent functionalization of GO filler [25,26,27,28]. Additionally, synergistic effect of two types of fillers (multiwalled carbon nanotubes and graphene oxide nanosheets) has recently exhibited enhanced tensile strength with 1.25 wt% fillers [29].

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ACCEPTED MANUSCRIPT A few general trends can be observed despite the scatter reported in the literature on the mechanical properties of PVA based composites. Increase in the tensile elastic modulus has generally resulted in the lowering of failure strain [24,30,31,32,33]. The ductile to brittle transition has recently been related to the dehydration of water in the composite and its interfacial bonding [34]. More

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importantly, the difference in the effect of graphene oxide and graphene on the mechanical properties in PVA (or any polymer system) has not been clearly spelt out. Additionally, since the word graphene is often used to describe functionalised or partially reduced graphene oxide (rGO), careful reading of

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the original papers is often necessary to determine the exact type of filler employed [35]. The purpose of this study is to explore the possibility of fabricating composites that possess high elastic modulus along with considerable failure strain (which is difficult to synthesize but desirable for engineering

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applications). We have also addressed the subtle difference in the effect of GO and rGO on the properties of the composites by fabricating composites with 0.3 wt.% of fillers. GO was used via solution blending technique with water as the processing solvent. Subsequently, the GO loaded polymer blend was reduced in-situ to prepare rGO loaded composite films. Tensile properties of the films have been measured followed by a series of chemical and structural characterization techniques.

2. Experimental

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Graphite fine powder (Loba Chemie, 99% pure, particle size 50µm), potassium permanganate (Loba Chemie), sulphuric acid (98%, Merck), hydrochloric acid (98%, Merck), nitric acid (Merck),

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PVA with molecular weight 1,15,000 and degree of polymerization 1700-1800, were procured from commercial sources and used without any further purification. Composite fabrication and characterization methods are outlined as follows. Synthesis of graphene oxide was carried out by the

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modified Hummers method with graphite as the precursor. In a typical experiment, 2 g of graphite powder was mixed with concentrated H2SO4 (46 ml) in a three necked round bottom flask kept in an

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ice bath. The temperature of the ice bath was well controlled and not allowed to rise above 10 °C. The cooled mixture was stirred for 30 minutes, after which the oxidant KMnO4 (6 g) was gradually added to it. This mixture was stirred further for 30 more minutes. Subsequently, the ice bath was removed and the mixture was heated to bring the temperature up to 35 °C, and was kept at this temperature for the next 1 hour with vigorous stirring. Subsequently, a large amount of deionised (DI) water (92 ml) was added to the mixture, which increased the reaction temperature to around 98 °C. Controlled addition of water is an important step here to prevent explosion of the reactants. The mixture was further stirred for about an hour, with the temperature of the reaction maintained at 98 °C, following which additional DI water (280 ml) was added to the mixture. Then the reaction was terminated by the addition of 30% H2O2 (50 ml) to oxidize the unreacted graphite which turned the slurry bright yellow in colour. The GO suspension was then washed with 1:10 HCl aqueous solution (500 ml) in order to remove metal ions. The paste collected from the filter paper was dried at 60 °C. The agglomeration powder was re-dispersed into deionized water and washed rigorously 5-7 times with

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ACCEPTED MANUSCRIPT more deionized water until the pH was nearly 7. The paste collected on the filter paper was dispersed in water by ultrasonication, and recollected by a research centrifuge operating at 4000 rpm separating out any unexfoliated GO. The GO platelets were obtained by drying the filtrate at 60 °C in air. Two sets of PVA composites were fabricated by conventional solution casting method. The fabrication

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procedure can be briefly outlined as follows. 2 g of PVA was dissolved in 30 ml DI water at 95°C under reflux to form an aqueous solution. Subsequently, GO powder (at 0.3 wt.% of PVA) was sonicated in 20 ml DI water to form a yellowish brown colloidal solution. The as obtained solutions

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were mixed together under vigorous stirring for about 2 hours at 60 °C to form a colloid. The PVA solution is comparatively thicker than the GO colloid, which makes the mixing of the two components difficult. Therefore, vigorous stirring of the aqueous GO mixed PVA colloid is required to ensure

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complete and uniform mixing. Finally, 10 ml of this homogeneous GO mixed PVA colloid was poured onto a Teflon petri dish, and kept in a vacuum dryer at 40 °C for 48 h. The dried films (hereafter, PVA-GO) were then peeled off the substrate for further characterizations. For the second type of films, 0.2 ml of hydrazine hydrate (35%wt) was added to the GO-PVA blend, and was refluxed for 24 hours at a temperature of 80 °C. The brown polymer blend gradually turned black. 10 ml of the reduced blend was then cast onto a Teflon petri dish, and dried in vacuum at 40 °C until weight equilibrium. The cast films (hereafter, PVA-rGO) were peeled off the petri dish for further

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characterization.

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The mechanical properties of the films were measured using a universal testing machine (TA.XT2i Texture Analyzer, Stable Micro Systems Ltd, Surrey, UK) at room temperature. A load cell of 500N was used with a tensile loading rate of 5 mm min-1 according to ASTM D638 standards. All

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of the samples were cut with a razor blade into dog bone shapes of 5 mm width and 20 mm gage length, and fixed into regular cardboard windows of (30 mm x 30 mm dimension) to hold the films in

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the tensile tester teeth and to ensure uniform stretching. Five samples were tested for each film type, and their mean value and standard deviation were recorded. Cross sectional imaging of the PVAbased composites was carried out by scanning electron microscopy (FEI Nova NanoSEM, FEI, Eindhoven, Netherlands). X-Ray diffraction studies (Rigaku Ultima IV) were conducted with a CuKα radiation in the scanning range of 5 to 50° of 2Ɵ. Fourier transform Infrared spectroscopy (FTIR) was recorded (Alpha-E, Bruker, Germany) with a scan resolution of 2 cm-1 in the range from 4000 to 500 cm-1, using the attenuated total reflectance (ATR) technique. The thermal behaviour of the films was investigated via differential scanning calorimetry (DSC, Mettler-DSC822e TA, USA) on square shaped coupons (~10 mg) over 30-350 °C with a heating rate of 5 °C min-1.

3. Results and discussions The X-ray diffractogram of graphite flakes exhibited a sharp diffraction peak at 26.6°, corresponding to the (002) reflections, and d-spacing value of 0.336 nm (Fig.1). The diffractogram of GO powders exhibited a peak at a 2θ value of 10.9°, which corresponds to d-spacing of 0.78 nm. Such

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complete exfoliation of the graphene layers, and no restacking of the layers after drying. Although the reduction method followed was done in-situ for the PVA-based composites, the individual reduction of GO powders in solution by hydrazine indicated that the process has been successful. UV-Vis

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spectroscopy of the aqueous colloids of GO and rGO was performed to check if the reduction process has been successful. The absorption peak was found at 230 nm for the GO colloid, where as it has shifted to around 260 nm for the rGO colloid (cf. Fig. 2). This shift is the characteristic feature of a

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reduced form of graphene oxide, which is attributed to the restoring of electronic conjugation of the C=C bonds in graphene [37]. The zeta potential measurements of the rGO colloid system was observed in the range of -30 to -35 mV, which is consistent with a stable aqueous dispersion. The scanning electron micrograph (SEM) of the synthesized GO platelets (Fig. 3) exhibited highly crumpled morphology. The lateral dimensions were about 9-10 µm in size as calculated from SEM images. The thickness of the GO layers was found out to be ~8 nm from Scherrer’s formula from the XRD patterns.

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Since property enhancement of the polymers by reinforcing the matrix with a lower amount

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of graphene based filler is the overall objective, we first report the results of mechanical properties. The uniaxial tensile properties of the pristine as well as reinforced PVA composite films are presented in (Fig. 4), and the values are provided in Table I. The pristine PVA films showed a elastic modulus

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of 2.32 GPa, and tensile strength of 25 MPa, which are typical values for PVA (although the same polymer with a different molecular weight can have different tensile properties). The reinforced

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composites, on the other hand, showed considerably improved elastic modulus and tensile strength values for a low filler loading of only 0.3 wt. %. For the PVA-GO composites, the elastic modulus was 5.82 GPa, which is ~250 % of pure PVA values (i.e. increment of 150%). The tensile strength was 63 MPa, which also is almost 250% of that of PVA films. The PVA-rGO films, however, exhibited elasticity of 2.55 GPa, which is comparable to that of PVA. The tensile strength for PVArGO was 39 MPa, which is higher than that of pristine PVA films, and lower than that of PVA-GO films. In summary, the GO fillers improve elastic modulus as well as tensile strength of the PVA composites, whereas upon hydrazine reduction (where GO becomes rGO fillers) the elastic modulus reduces back to values similar to that of pure PVA. Improvement in tensile strength in the PVA-rGO composites nevertheless is much better than that of pure PVA. Based on the results, one might argue that the reduction process with the use of hydrazine can cause changes in the polymeric matrix itself other than the targeted reduction of the GO platelets. Hence, we have refluxed the pure PVA solution with hydrazine hydrate for 24 hours, followed by solution casting, film fabrication, and mechanical

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ACCEPTED MANUSCRIPT testing. These films (PVA-H) exhibited tensile strength of 23 MPa, similar to the pure PVA values, and elastic modulus of 1.73 GPa, which is somewhat lower than that of pure PVA. It follows then that the hydrazine reduction step, although doesn’t affect the tensile strength, lowers the elastic modulus (by ~25%). The fact that the PVA-rGO composites showed similar elastic modulus to that of pure

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PVA indicates that the decrease in elastic modulus in the PVA matrix, as a result of hydrazine reduction, is offset by filler (rGO) induced increase in elastic modulus. Further, interesting behaviour on the ductility of the samples can be observed by the failure strain values (Table I & Fig. 4). The

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pure PVA and PVA-H films failed at an average strain of 149% and 147%, respectively. The PVAGO films failed at an average strain of 170% exhibiting higher ductility than pure PVA. Interestingly, the PVA-rGO films failed at a much higher strain of 208% showing enhanced ductility of the

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composites.

The properties obtained in the current set of experiments are significantly improved in PVA based systems. Such improved mechanical properties with only 0.3 wt.% of filler loading without any functionalization of the fillers has rarely been reported [18-33]. Additionally, some interesting features stand out from the data set. Generally, an increase in elasticity is observed with a decrease in the failure strain [24,38,39]. However, the current results from PVA-GO composites exhibit improved elastic modulus, tensile strength, as well as high ductility, which is a desired combination of

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properties. In addition, PVA-rGO composites exhibited improved tensile strength with a profound

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increase in ductility. In the subsequent sections, the composites are evaluated with a series of characterization tools, and the behaviour is explained.

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Physical properties of polymeric materials are dependent on the crystallinity of the polymer and its glass transition temperature. The XRD analysis of pristine PVA films, as can be seen in Fig. 5, shows a peak at 19.6° 2θ, followed by a few weak signals at higher angles. These features are

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indicative of a semi-crystalline polymer. The PVA-GO composites also exhibited similar features with only one major peak with only slightly reduced intensity to that of pure PVA (Fig. 5). No peak corresponding to graphene or graphitic carbon layers was discernible, which could be due to the low volume fraction of the fillers in the composite. The purpose of using a low amount of filler content in the current experiments was to understand if crystallinity of the polymers increases elastic modulus. In our case a significant improvement in strength and elastic modulus has been observed (PVA vs PVA-GO composites), without major change in the crystallinity of the polymer. It may thus be inferred that incorporation of GO in the PVA matrix does not change polymer crystallinity significantly, at least not at the level of 0.3 wt.% of filler content. With higher amount of filler, however, polymer crystallinity and tensile elastic modulus may increase, where part of the improvement may be ascribed to increased crystallinity of the polymer. The diffractogram for the PVA-rGO composite, on the other hand, had the major peak (at 19.6°) less intense than that of pure PVA and PVA-GO films (Fig. 5), indicating reduced crystallinity. The diffraction pattern of the PVA-

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for such reduced crystallinity is not clear yet.

The results of glass transition (Tg) and melting temperatures (Tm) from DSC analysis of the films are listed in Table II. The GO or rGO based PVA films showed moderate variation in the glass

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transition temperature (Fig. 6). The PVA-GO composites exhibited distinctly higher Tg (114 °C) than that of PVA films (109 °C). Such change could be ascribed to the hydrogen bonding between oxygen

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containing functional groups on the GO layers and the polymer matrix. In addition, the crumpled morphology of the GO fillers may also limit segmental mobility of the polymeric chains, as has been discussed in the literature [40]. The PVA-rGO composite exhibited Tg of 110 °C which is similar to that of PVA composite. Interestingly, the rGO effect at 0.3 wt.% loading did not seem to cause major change in the Tg, indicating that the hydrogen bonding effects are reduced. The Tm of all of the composites were ~225 °C, consistent with the behaviour reported in the literature [41]. FTIR spectroscopy of the polymer composites (Fig. 7) exhibited standard absorption peaks

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for the PVA system. The appearance of absorption peaks at ~1080 cm-1 and ~1140 cm-1 and their relative intensity is indicative of the semi-crystalline nature of the polymer composites [42]. The

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absorption peaks for pure PVA matrix and the PVA-H specimens were similar with distinct and broad –OH stretch around 3200 cm-1, which can be attributed to the presence of alcohols for intermolecular

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H-bonds in the polymer. The PVA-GO films exhibited similar features with a slight shift of the peak to lower wavenumbers, indicating strong interfacial interactions between the matrix and filler. The PVA-rGO films, however, exhibited a shallower –OH peak as compared to that of the pristine PVA

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film. The peak shift here was between that of PVA and PVA-GO composites. Thus the extent of intermolecular H-bonding may be between that of pure PVA and PVA-GO composites. Due to the reduced intermolecular H-bonds as observed in the spectra, we conclude that hydrazine reduction process lowers interfacial interaction between the matrix and the filler. The morphological studies of the nanocomposites reveal some new information that has not been reported in the literature. The cross section of the PVA film indicated features typical of a glassy polymeric system with no appreciable contrast. In the PVA-GO films, however, outlines of the filler can easily be observed (indicated by arrows in Fig.8a). In addition, the GO layers were found to be randomly oriented, with no preferred direction of alignment. While the enhanced tensile strength and modulus of the PVA-GO composites are explained by strong interfacial bonding, the slightly increased failure strain of these films can be attributed to the random alignment of the fillers. It is possible that on application of stress, the randomly oriented GO fillers progressively align themselves

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reduced PVA-rGO composites not all of the functional groups of graphene oxide are reduced. The rGO fillers are not the ideal graphene sheets. Some of the functional groups, in particular the epoxy groups on the basal plane of GO sheets, are reduced in the hydrazine reduction process. The groups on

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the edge of the sheets (according to the Lerf-Klinowski model) are –COOH and –OH groups that tend to remain even after hydrazine treatment [43]. The absence of functional groups on the basal plane may have facilitated orientation of the rGO sheets in the plane of the film. The groups at the edge of

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the sheets can still form hydrogen bonds with the polymer matrix.

Load transfer depends on interfacial interactions of the reinforcements and the matrix. GO contains oxygen containing functional groups that can form strong bonds at the interface, thus resulting in efficient load transfer. Factors such as strong interactions at the interface generally tend to restrict molecular chain movement resulting in a lower failure strain. Enhanced homogeneity of the dispersion of the fillers indicate formation of more such interfaces (between reinforcements and the matrix), and thus formation of more interfacial bonds. These features in the PVA-GO composites

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exhibited enhanced tensile properties. Slippage of the rGO fillers can occur upon loading in the PVA-

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rGO composites, due to the lack of interfacial interactions between the basal planes of rGO and the PVA matrix. Such a process can lead to high strain of the film. However, owing to the strong interfacial bonds between the edge groups in rGO sheets and the polymer matrix the PVA-rGO

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composites exhibit better tensile strength as compared to that of pure PVA. A simplified model for each of the composite microstructure is presented in Fig. 9 with only emphasis on the presence of

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epoxy, hydroxyl, and carboxylic functional groups. The PVA-GO composites with bonding interactions between the polymer matrix and the edge as well as basal planes of the GO sheets are shown in Fig. 9a. The PVA-rGO composites with interfacial interactions only at the filler edges are shown in Fig. 9b. Although speculative, the model is based on the micrography of the PVA-GO and PVA-rGO composites.

4. Summary PVA composites with minimal GO loading of 0.3 wt.% exhibited ~150% increment in tensile elastic modulus as well as strength, and higher failure strain than those of pure PVA films. The improvement in the mechanical properties was ascribed to stronger interfacial bonding between the fillers and the matrix, and the wrinkled morphology of the fillers. The comparatively higher failure strain of the PVA-GO composites along with a higher elastic modulus resulted from the randomly oriented morphology of the fillers that align themselves parallel to the loading direction. The PVArGO composites showed elastic modulus comparable to that of pure PVA. In these composites, the

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ACCEPTED MANUSCRIPT decrease in elastic modulus as an effect of hydrazine on the polymer matrix was offset by an equivalent increase in modulus due to rGO fillers. PVA-rGO composites also exhibited better tensile strength than the PVA films, but comparatively lower than those of PVA-GO composites. The failure strain was ~60% higher than that of pure PVA composites. The improved strength of the PVA-rGO

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composites resulted from the strong interfacial bonds between the matrix and the edges of the rGO sheets. The substantially higher ductility for the PVA-rGO composites can be ascribed to the rGO fillers aligned in the plane of the film with limited or no interfacial bonds in the basal planes, which

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facilitates easy gliding of the molecular chains in the polymer. With these results tailoring of PVA

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composites concerning elastic modulus, tensile strength, and ductility is possible.

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Figures Captions Figure 2: UV-vis spectroscopy of GO and rGO colloids.

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Figure 1: X-ray diffractograms of graphite, graphene oxide, and reduced graphene oxide

Note the wrinkled morphology of the graphene oxide sheets.

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Figure 3: Field emission scanning electron micrograph of the synthesized graphene oxide platelets.

Figure 4: Tensile stress-strain curves of the polymer nanocomposites. The curves shown here are one

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of the many runs conducted for each sample. Refer to Table I for average values.

Figure 5: X-ray diffractograms of the polymer nanocomposites. The scale in Y-axis is the same for all of the composites.

Figure 6: Differential scanning calorimetry traces of the composites indicating shift in the glass transition temperature.

Figure 7: FTIR spectroscopy of the polymer composites.

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Figure 8: Scanning electron micrographs of the cross sections of the polymer composites; (a) PVA-

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GO composites, exhibiting randomly oriented GO fillers, and (b) PVA-rGO composites, exhibiting reduced GO fillers aligned parallel to the film plane. Figure 9: Schematic showing the morphology of the reinforced PVA composites: (a) PVA-GO

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composites with randomly oriented GO platelets with bonding interactions all around the GO sheets,

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and (b) PVA-rGO composites with interfacial interactions only in the edges.

Tables captions

Table I: Tensile properties of the polymer nanocomposites Table II: Thermal properties of the polymer nanocomposites

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E (GPa) 2.32±0.3

σf (MPa) 25.3±3

εf (%) 149±7

PVA-GO

5.82±0.6

63±5

170±5

PVA-rGO

2.55±0.3

39±5

208±4

PVA-H

1.73±0.5

23±3

147±5

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Material PVA

Table II: Thermal properties of the polymer nanocomposites

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Material PVA PVA-GO PVA-rGO

Tg (°C) 109 114 110

Tm (°C) 225 225 224

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Intensity (a.u)

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Graphite

Reduced Graphene Oxide

20

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Graphene Oxide

30 2

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50

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Graphite

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GO

Graphite

250

300

350

400

450

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W avelength (nm)

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rGO

GO

D

200

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Absorbance

rGO

500

550

600

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RI PT

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60

PVA-GO

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40

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Stress [ ] (MPa)

50

30 20

PVA-rGO

PVA-H

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40

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10

80

PVA film

120 Strain [ ] (%)

160

200

240

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Intensity [a.u.]

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PVA-rGO

PVA-GO

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PHA-H

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PVA

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30 2

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PVA-rGO

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-0.5

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-1

Heat Flow [mW mg ]

0.0

-1.0

PVA

PVA-GO

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-1.5

80

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60

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100

120

Temperature [ C]

140

160

d

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PVA-H

c

PVA-rGO

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Transmittance

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b

PVA-GO

a

4000

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PVA film

3500

3000

2500

2000 -1

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Wavenumber (cm )

1500

1000

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Highlights: o PVA-GO (0.3 wt%) exhibited 150% increment in tensile strength over PVA.

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o Composite with in-situ reduced fillers showed excellent ductility (220% strain). o PVA-GO films had randomly oriented fillers in the matrix with strong interfacial.

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o PVA-rGO composites had preferred orientation of the fillers in the plane of the film.

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o Such interesting mechanical properties with only 0.3 wt% fillers are new information.