- Email: [email protected]
polystyrene nanocomposite by in situ microemulsion polymerization
Journal of Colloid and Interface Science 350 (2010) 530–537
Contents lists available at ScienceDirect
Journal of Colloid and Interface Science www.elsevier.com/locate/jcis
A facile approach to the fabrication of graphene/polystyrene nanocomposite by in situ microemulsion polymerization Archana S. Patole a, Shashikant P. Patole b, Hyuck Kang a, Ji-Beom Yoo b, Tae-Ho Kim a, Jeong-Ho Ahn a,* a b
Department of Polymer Science and Engineering, Sungkyunkwan University, Suwon 440-746, Republic of Korea SKKU Advanced Institute of Nanotechnology (SAINT), Sungkyunkwan University, Suwon 440-746, Republic of Korea
a r t i c l e
i n f o
Article history: Received 21 November 2009 Accepted 13 January 2010 Available online 18 January 2010 Keywords: Polystyrene Graphene In situ microemulsion polymerization Functionalization Nanoparticles Morphology Dispersion Composite Thermal properties Electrical properties Transmission electron microscopy
a b s t r a c t This paper reports a large scale production route for polystyrene (PS) nanoparticle-functionalized graphene sheets using water based in situ microemulsion polymerization. The higher surface area of the graphene basal plane and the better proximity of the reactant species in in situ microemulsion polymerization were used to functionalize the graphene sheets using PS nanoparticles. The thermal properties of the PS were improved with the incorporation of graphene in the composite. The modified graphene exhibited good compatibility and interactions with the host PS matrix to form conducting PS films. Ó 2010 Published by Elsevier Inc.
1. Introduction Graphene have attracted considerable interest over the last few years on account of its extraordinary electrical, thermal, and mechanical properties arising from its unique structure [1]. This two dimensional sp2 bonded carbon structure has potential applications in many technological fields, such as high frequency analog electronics [2], single molecule sensors [3], nanocomposites [4], batteries [5], supercapacitors [6], liquids crystal display [7], and hydrogen storage [8]. One possible way of exploiting their properties for real world applications would be to incorporate graphene sheets in a composite material [9,10]. The manufacture of such composites requires that graphene sheets be produced on a sufficient scale but be incorporated and distributed homogeneously into various matrices. Over the last few years, many groups have reported either top down or bottom up approaches to synthesize graphene, including chemical vapor deposition (CVD) [11], mechanical exfoliation (repeated peeling) of graphite [12], chemical intercalation and exfoliation of graphite [13,14] and thermal expansion of graphite [15]. Of these routes, the thermal expansion
* Corresponding author. Fax: +82 31 292 8790. E-mail address: [email protected] (J.-H. Ahn). 0021-9797/$ - see front matter Ó 2010 Published by Elsevier Inc. doi:10.1016/j.jcis.2010.01.035
of graphite is an easy way of obtaining a bulk amount of graphene that allows it to be incorporated in various composite materials. The use of polymers structurally similar to the matrix polymer to functionalize graphene is a desirable approach for developing polymeric graphene composites. This ensures that the functionalized graphene is compatible with the polymer matrix to avoid potential microscopic phase separation in the composites. Therefore, an ideal polymeric graphene composite might be prepared using solubilized graphene that are functionalized with the matrix polymer. Polystyrene (PS) is one of the most widely used commercial polymers. The incorporation of functionalized graphene sheets in a PS matrix results in composite thin films that are semiconducting and exhibit an ambipolar field effect [10]. These composites can be deposited uniformly over large areas in the form of thin films from solution onto which the devices, such as thin film transistors, can be fabricated without extensive lithography. The dispersion of graphene into a PS matrix has stimulated considerable interest among researchers [9,10]. However, the lack of reactivity of the graphene basal plane needs to be solved. A perfect graphene does not have dangling bonds on the basal plane for chemical bond formation. The chemical potential of the basal plane is lower than the edges or defects. Therefore, graphene composites allow grafting on the edges and defects only. The dispersion and agglomeration of graphene is
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
another problem while forming the composite materials in bulk amounts. The hydrophobic nature and high specific surface area in graphene causes poor dispersion in a variety of organic solvents. Van der Waals forces cause graphene to agglomerate and restack, which destroys the properties of the individual graphene sheet. The functionalization of graphene sheets is an easy way of overcoming this problem. Until now, several strategies were proposed based on non-covalent and covalent bonding between inorganic and organic molecules on the surface of grapheme [9,16–19]. However, to date, there is no report on PS nanoparticle-functionalized graphene sheets for PS-graphene composite. In our previous work we presented effective in situ synthesis and characteristics of polystyrene nanoparticle covered multiwall carbon nanotube (MWCNT) composite by microemulsion polymerization [20]. This paper reports the use of microemulsion polymerization to functionalize graphene sheets by PS nanoparticles. This polymerization method has some merit over other emulsion methods because it produces thermodynamically stable polymer particles tens of nm in size, and uses the advantage of surfactants through polymerization [20–22]. The higher surface area in graphene and in situ microemulsion enhances the anchoring of PS nanoparticles on graphene through p–p interactions. The functionality of these nanoparticles affords better compatibility and interactions with the host PS matrix, thereby imparting improved thermal properties and enhanced conductivity. It is believed that this rational approach will provide a better method for functionalizing graphene in various host matrices in order to improve its applications in the real world. 2. Materials and methods 2.1. Materials Commercially available expandable graphite (EG) with >99.9% purity was used to obtain a few layers of graphene sheets. Styrene, sodium dodecyl sulfate (SDS), azobisisobutyronitrile (AIBN), PS beads (Mw-280,000) and 1-pentanol were purchased from Sigma Aldrich and used without further purification. Tetrahydrofuran (THF) and methyl alcohol as solvents were purchased from Aldrich Co. and used as received. 2.2. Graphene synthesis In a coexisting research direction, before selecting the material for experiments, we explored the rapid thermal expansion of EG. Initially, 50 mg of as-purchased EG were loaded in the CVD chamber using an alumina boat with a lid. The CVD chamber was evacuated to <0.01 Torr using a rotary pump. The ramping rate was set for 800 °C/min using a tungsten lamp heater and programmable heater controller. Thermal expansion was carried out at various temperatures (700, 800, 900 and 1000 °C) for two minutes. The samples were then cooled to room temperature and removed from the CVD chamber for further experimentation. This scheme can be used for the mass production of graphene depending on the CVD reactor size.
531
containing 200 ml deionized water, 4.5 ml pentanol and 1.5 g SDS for 1 h. Pentanol was used as a hydrophobe and SDS was used as the surfactant. After dispersion, the aqueous graphene mixture was charged into a 500 ml four-neck glass reactor equipped with a condenser, dropping funnel, stirrer and nitrogen inlet. The reactor was then placed in an ice bath. Styrene (5 g) and AIBN (0.05 g), which were the monomer and initiator, respectively, were mixed through a funnel prior to polymerization. The solution was deoxygenated with oxygen-free nitrogen for 20 min by bubbling nitrogen gas. The mixture was then sonicated at 0 °C for 4 h. After homogeneous mixing, the ice bath was replaced with an oil bath for heating. Polymerization was then carried out by increasing the oil bath temperature to 85 °C for 4 h under an inert nitrogen blanket. After polymerization, the grey colored emulsion obtained was washed with methanol and water to remove the excess surfactants, hydrophobe and residual monomer from the reaction mixture. The procedure was repeated more than three times. The self-assembled graphene/PS nanoparticle composite was then dried under vacuum at 60 °C for 24 h, and used for further characterization. As a comparison, the pristine polymer nanoparticles were prepared using the same procedure in the absence of graphene. 2.4. Preparation of conducting PS films PS nanoparticle-functionalized graphene was used to prepare the PS films. In typical experiments, a 2.5 wt.% graphene/PS nanoparticle sample was selected as a filler to make the composite films. The host matrix was prepared by dissolving 5 mg of PS beads in 20 ml of THF. 3–20 wt.% filler was added to this host matrix. The solution was then cast on a glass plate and dried at room temperature to form the free standing, 600 lm thick and uniform films. Fig. 1 summarizes the process steps. 2.5. Characterizations The morphology of graphene, pristine PS nanoparticles and graphene/PS nanoparticles was examined by scanning electron microscopy (FESEM; JSM6700F, JEOL). The samples were dispersed in methanol and drop coated onto a copper grid for high resolution transmission electron microscopy (HR-TEM; JEOL 300 kV) to explore the structural properties. The spectroscopic analyses were carried out using a Fourier transform infrared spectrometer (FTIR-Nicolet IR 200) and a micro-Raman spectrometer (Invia Basic, Renishaw Co. England) using an Ar-ion laser at 640 nm as the excitation light source. Thermo-gravimetric/differential thermal analysis (TG/DTA) of graphene, PS nanoparticles and graphene/PS nanoparticles were carried out using a thermobalance (TGA 2050) from room temperature to 800 °C, at a rate of 10 °C/min in a continuous nitrogen flow. The dynamic differential scanning calorimetry (DSC) experiments were carried out with a thermal analyzer (DSC 2920) at heating rates of 5 °C/min in a continuous nitrogen flow. The sheet resistance (Rs) of the composite films was measured using a 4-point probe resist meter (AIT CMTSR1000 N).
2.3. Ultrasonically initiated in situ microemulsion polymerization 3. Results and discussion An ultrasonic irradiation device (Sonomaster Model ULH-700S) was equipped with a standard aluminium horn with a replaceable tip diameter of 13 mm and temperature controller. The nominal energy output of the probe was set to 300 W. The PS nanoparticle-functionalized graphene was prepared using the following procedure. First, approximately 0.15–1 g of graphene, which was thermally expanded at 1000 °C was sonicated in a beaker
Fig. 2a shows a photograph of the as-purchased EG powder and thermally expanded graphene. The temperature used for thermal expansion was 1000 °C. The bulk volume of the expanded graphene was much higher than that of the EG powder. It is evident that the EG expanded at 1000 °C. The bulk volume of thermally expanded graphene increased with increasing heating temperature.
532
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 1. Multi step process for producing the conducting graphene–PS films.
Fig. 2b shows a SEM image of EG, which consists of several graphene stacks. Here, the main challenge was to separate these stacks to obtain the individual graphene flakes. Van der Waals forces keep these stacks together. The heating rate of 800 °C/min used for expansion is more than sufficient to overcome this force to obtain a few layer graphene sheets [15]. The samples heated to 700 and 800 °C showed a much larger number of graphene layers than the samples expanded at 900 and 1000 °C. Fig. 2c–d shows SEM images of the graphene flakes expanded at 1000 °C. In Fig. 2c, the folded graphene flake had the appearance of a crumbled piece of paper. A higher magnification image shows the folded edge of the typical graphene flakes (Fig. 2d). A few layers of graphene (1–10 layers) were observed by HR-TEM. Fig. 2e shows a one layer graphene flake at its edge and Fig. 2f shows two layers of graphene at its edge. A higher number of graphene layers (up to 35 graphene layers) were observed in the case of the 700 and 800 °C expanded samples. The samples expanded at 1000 °C were chosen for further experimentation. In situ polymerization was carried out to functionalize the graphene flakes with the PS nanoparticles. The main motivation behind the in situ reaction was to increase the proximity of the reactant species and use a higher surface area of graphene flakes to anchor the PS nanoparticles. Previously, in the case of MWCNT, a higher surface area was reported to provide better anchoring of PS nanoparticles to the outer walls of the MWCNTs [20]. A similar analogy was used for the present study, rather in the case of graphene, a higher surface area provides better proximity for the reactant species to anchor the PS nanoparticles. Fig. 3a–c shows SEM images of the dried composite. The images show that the composite material consists of randomly aggregated, thin, crumpled graphene flakes closely associated with each other and a polystyrene nanoparticle covered on it, forming a disordered solid. The
individual graphene flake size was approximately 5–10 lm. The high resolution images (Fig. 3b–c) confirmed that several PS nanoparticles were anchored on both sides of the graphene flakes. At the resolution limit of our instrument (see Fig. 3c–d), the data suggests, but does not prove, the presence of individual flakes in this material. The absence of charging during SEM imaging indicates that the network of the graphene based sheets and the individual sheets were electrically conducting. This qualitative conclusion was further confirmed by the electrical measurements on graphene/PS nanoparticle composite thin films. Fig. 3d–f shows TEM images of the PS nanoparticles-covered graphene flakes. The PS nanoparticles (with average diameter of 37.5 nm) were distributed randomly over the graphene flakes. The number of PS nanoparticles on the graphene flakes depends on the amount of starting graphene flakes; the number of PS nanoparticles decreases with increasing number of graphene flakes. Fig. 3d shows a TEM image of the 3 wt.% composite and Fig. 3e shows a TEM image of the 20 wt.% composite. HR-TEM analysis confirmed that the individual graphene flakes consists of 1–10 graphene sheets, which is consistent with the thermally expanded samples at 900 and 1000 °C (see Fig. 2e–f). A closer view (see Fig. 3f) shows the attachment of PS nanoparticles on graphene or incorporation into the graphene sheet. It appears that Brownian motion of the emulsified monomer droplets and graphene inside water causes collisions between these two species, which provides further anchoring sites for juvenile PS nanoparticles to attach to the graphene surface. It was assumed that the hydrophilic ends of SDS on graphene and on the emulsified droplets caused repulsion, which was overcome by the kinetic energy of Brownian motion of these species. During the collision, SDS at the contact region might have been displaced or removed, and provided more stable sites for anchoring the PS nanoparticles on the graphene surface.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
533
Fig. 2. (a) Digital camera photograph of the as-purchased EG and thermally exfoliated graphene flakes. SEM images showing, (b) the graphite stacks in the EG, (c) crumbled graphene flake and (d) edge of the graphene flake. HR-TEM images showing, (e) single layer and (f) two layers graphene.
The in situ reaction and larger surface area of the graphene sheet provided good adhesion at the contact regions. This adhesion can be attributed to some interaction between the phenyl groups in PS and graphene through p–p stacking [23,24]. The HR-TEM image in Fig. 3f confirms the hypothesis. Fig. 4 shows the proposed hypothesis which was applicable to the higher surface area MWCNTs [20]. In the first step, graphene was functionalized by the anionic SDS surfactant. The graphene was isolated from the water by SDS, keeping the hydrophobic end towards the surface of the graphene. Similarly, the emulsified droplets were formed by the SDS coverage, maintaining the hydrophilic ends of SDS towards water. During polymerization, collisions between the graphene and emulsified monomer due to Brownian motion caused the selective etching of SDS. This provides anchoring sites on the graphene for nanoparticles. There may be possibility that the anchor of the PS particles is due to the weak Van der Waals force originating from the random contact of carbon atoms and PS chain. The attachment of PS nanoparticles to the graphene sheet becomes a probabilistic event due to collisions, and does not cause the selective anchoring of all PS nanoparticles, the unan-
chored PS latex and SDS were removed by successive washing in methanol. Finally, the PS nanoparticles-covered graphene composite was obtained after successive washing. FTIR spectral analysis was performed to confirm the chemical structure of the coating of PS nanoparticles on the graphene sheet. Fig. 5a shows the FTIR spectra of the pristine graphene, pristine PS nanoparticles and graphene/PS nanoparticle composite. Although no visible peaks were observed in graphene, there were several peaks in the pristine PS nanoparticles and composites. Prominent peaks in PS nanoparticles and composite at approximately 2820– 2967 cm 1, 2994–3043 cm 1, 1500 cm 1 and 700–1400 cm 1 indicate aliphatic C–H stretching, aromatic C–H stretching, aliphatic CH2 and different conformation sensitive vibration modes of polystyrene, respectively. This confirms that the polymerization reaction produced PS from styrene. A new peak at approximately 3036 cm 1 appeared in case of composite, which could be attributed to the C–H stretching of aromatic ring. There was no significant change in the peak positions apart from the peak intensity, indicating that the PS nanoparticles in pristine and composite forms have the same structural properties. The anchoring of the
534
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 3. SEM and TEM images of the graphene sheets functionalized by PS nanoparticles. In figure (a) the micron size of the sheets are revealed. (b and c) The higher magnification images show that the PS nanoparticles are attached to both sides of the graphene sheets. TEM images of (d) 3 wt.% and (e) 20 wt.% samples. (f) HR-TEM image of the attachment of PS nanoparticles on the graphene flake. The interface region in (f) is marked with an arrow.
Fig. 4. Step wise functionalization of graphene sheets with PS nanoparticles.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 5. (a) FTIR spectra of the pristine graphene, pristine PS nanoparticles and graphene/PS nanoparticles. (b) Raman spectra of pristine graphene and graphene/PS nanoparticles.
nanoparticles occurs at the surface of the graphene sheet, which maintains the inner structure of PS nanoparticle. However, the possibility of a structural change on the surface of the graphene sheet cannot be excluded. Raman spectroscopy is a powerful nondestructive tool for characterizing carbonaceous materials, particularly for distinguishing ordered and disordered crystal structures of carbon. The major Raman features of graphene and graphite are the so called G (1580 cm 1) and 2D (2670 cm 1) bands. The G band originates from the in-plane vibration of sp2 carbon atoms and is a doubly degenerate (TO and LO) phonon mode (E2g symmetry) at the Brillouin zone center [25]. The 2D band originates from a two phonon double resonance Raman process [26]. The 2D band is the obvious difference between the Raman features of graphene and graphite. For graphene, the 2D band can be fitted with a sharp and symmet-
535
ric peak while that of graphite can be fitted with two peaks [27]. The supporting material (see Fig. S1) shows that in the case of graphene, the 2D band sharpens and becomes blue shifted. In addition to the differences in the 2D band, the intensity of the D band decreases almost linearly (as shown in Supporting Fig. S1) with decreasing graphene thickness. A comparison of the Raman spectra of graphene and the graphene/PS nanoparticle composite (Fig. 5b) revealed broadening of the 2D band in graphene due to the creation of defects. When the D band of graphene and composite Raman (Fig. 5b) were compared, the peak at 1354 cm 1 appeared after introducing polystyrene nanoparticles on the top of graphene, which is likely because the PS nanoparticles anchored on the grapheme basal plane, which increased the disorder in graphene. The intensity ratio I(D)/I(G) was increased significantly after polystyrene modification. Thermo-gravimetric analysis (TGA) was performed on graphene, PS nanoparticles and composites at all concentrations to obtain additional confirmation of both the graphene content and the structure of the composite, as well as to determine the effects of the graphene on the thermal stability of the composite. TGA was carried out under a nitrogen atmosphere. The resulting curves are shown in Fig. 6a. In the case of graphene, an initial 8% weight loss at <350 °C revealed the oxidation of amorphous carbon present in the sample. Pristine PS nanoparticles began to decompose at 380 °C and degraded completely at 401 °C. In the case of the pristine PS nanoparticles and composites, the weight loss below 200 °C was attributed to the decomposition of the lower molecular weight PS and the weight loss above 200 °C was attributed to the decomposition of higher molecular weight PS. However, the mass that remained constant above 400 °C was assigned to graphene. This method can determine the actual graphene content in the composites. The degradation temperature of the composites increased with graphene content. A maximum increase of 16 °C was observed for the 20 wt.% composite (refer Table 1). The first derivative of the TGA curve (DTA curves shown in Fig. 6b) clearly shows the variation in weight with time (dW/dT) as a function of temperature. The DTA peaks indicate the temperature of the maximum reactive velocity. Table 1 lists the maximum reactive velocity of the DTA curve of pristine polystyrene and graphene composites. The pristine PS nanoparticles reach the maximum reactive velocity at 401 °C and the value was lowest of all the curves obtained. The temperature at the maximum reactive velocity increased with increasing graphene content, and for the 20 wt.% composite it was increased by 15 °C compared to the pristine PS nanoparticles. This indicates that the thermal stability of the PS nanoparticles was improved by the addition of graphene. Previous studies on CNT/PS nanoparticle composites reported that the introduction of CNTs to the PS system can enhance the thermal stability of CNT–PS composites due to an interaction between the outer walls of the CNTs and PS lattices. Moreover, the formation of a barrier of CNT inhibits mass transfer and provides thermal insulation to shield the underlying polymer from the heat source [28,29]. Fig. 6c shows the glass transition temperature (Tg), which defines a pseudo second order phase transition. The incorporation of PS nanoparticles on graphene sheets resulted in an increase in Tg. In the case of pristine PS nanoparticles, the Tg was observed at 96 °C and was shifted to 113 °C for the 20 wt.% composite (see Table 1). Such an effect on the Tg obtained may be due to motion restrictions of the polymer chains at the nanoparticle–graphene interface [30]. A region of a strongly bonded polymer chains might have formed near the graphene surface. In this region, the polymer chains exhibited different behavior than that in the nanoparticle. The strong packing hinders chain segmental mobility. Therefore, more energy is needed to allow the first thermal transition, shifting Tg to a higher temperature. In the present case, similar effect might
536
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537 Table 1 Thermal properties of nanocomposites.
a b
Sample
Tdeg (°C)a
Graphene content (wt.%)a
Tg (°C)b
PSNP 3 wt.% 5 wt.% 10 wt.% 15 wt.% 20 wt.%
401 410 413 415 417 417
– 5.28 7.0 19.0 22.0 31.0
96 106 107 108 110 113
Determined by TGA. Determined by DSC.
Fig. 6. Thermal properties of the graphene/PS nanoparticles composite. (a) TGA, (b) DTA and (c) DSC analysis. The net amount of graphene estimated from TGA curves above 450 °C and the shift in glass transition temperature from DSC are listed in Table 1.
have occurred due to in situ polymerization. The observed increase in Tg with increasing graphene amount suggests good interfacial compatibility between the PS and graphene sheet. The ultimate goal of making graphene/PS nanoparticle composites was to develop a lightweight but highly conducting material. The PS nanoparticle-functionalized graphene flakes were used to make PS films. The details of the process steps involved in making the conducting PS films are discussed in the experimental section. The graphene/PS nanoparticle-enforced PS films were both
Fig. 7. Digital camera photograph shows (a) functionalized graphene embedded PS films. The number indicates a content of 2.5 wt.% functionalized graphene in the film. The inset shows the slurry in THF used for making the PS films. (b–e) Photographs show the flexibility of the composite films. (f) Sheet resistance of the composite films as a function of filler amount.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
conducting and flexible. Fig. 7a shows a photograph of the different wt.% composite films. Fig. 7b–e highlight the flexibility of the film. Fig. 7f shows the sheet resistance of these films as a function of the graphene content. The sheet resistance decreased several orders of magnitude with increasing graphene concentration due to percolation. For a 3 wt.% graphene content, the sheet resistance was 36.67 MX/h, which increased to 65.46 X/h for a 20 wt.% graphene content. These properties are superior to those reported elsewhere [9]. The sheet resistances differed slightly between the top and bottom of the film. This might be due to the sedimentation effect while forming the films by a wet cast method. No anisotropy effect was observed in these samples. It should be noted that at a 3 wt.% graphene content, the conductivity of the samples (4.5 10 5 S m 1) already satisfied the antistatic criterion (10 6 S m 1) for thin films and rapidly increased over the 20 wt.% range. The electrical properties of these composites compare well with the best values reported for nanotube–polymer composites thus far [31]. In the present study, the approach behind dispersing the PS nanoparticle-functionalized graphene in a PS host matrix was to achieve good dispersion and compatibility between the graphene and PS matrix (see Supporting Fig. S2). Therefore, it is expected that this straightforward, large-scale approach for the preparation, functionalization and incorporation of graphene sheets into polymer matrices will lead to further developments of a broad new class of materials with enhanced properties, and even allow the introduction of new functionalities to polymer composites in bulk quantities and reduces production cost. 4. Summary A large scale production route for PS nanoparticle-functionalized graphene sheets using water based in situ microemulsion polymerization was developed. SEM and HR-TEM showed the successful anchoring of PS nanoparticles on the surface of graphene. The thermal properties of the PS were improved with the incorporation of graphene in the composite. The modified graphene showed good compatibility and interactions with the host PS matrix to form conducting PS films. The reported scheme for fabricating the PS composite thin films from graphene and a commodity plastic highlights the potential for low-cost, macroscale thin film electronics. Acknowledgments This work was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2010-
537
0015035) and Ministry of Knowledge and Economy (Project No. 10037449). The authors also appreciate the project and equipment support from Gyeonggi Province through the GRRC program in Sungkyunkwan University. One of the authors (S.P.P.) is grateful to the Sungkyunkwan University for awarding him a Postdoctoral Research Fellowship. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.jcis.2010.01.035. References [1] A.K. Geim, K.S. Novoselov, Nat. Mater. 6 (2007) 183–191. [2] N. Tombros, C. Jozsa, M. Popinciuc, H.T. Jonkman, B.J.V. Wees, Nature 1448 (2007) 571–575. [3] J.M. Carlsson, Nat. Mater. 6 (2007) 801–802. [4] T. Ramanathan, A.A. Abdala, S. Stankovich, D.A. Dikin, M. Herrera-Alonso, R.D. Pinar, D.H. Adamson, H.C. Schnipp, X. Chen, R.S. Ruoff, S.T. Nguyen, Nat. Nanotechnol. 3 (2008) 327–331. [5] M. Liang, L. Zhi, J. Mater. Chem. 19 (2009) 5871–5878. [6] S.R.C. Vivekchand, C.S. Rout, K.S. Subrahmanyam, A. Govindaraj, C.N.R. Rao, J. Chem. Sci. 120 (2008) 9–13. [7] J. Zhu, Nat. Nanotechnol. 3 (2008) 528–530. [8] S. Patchkovskii, J.S. Tse, S.N. Yurchenko, L. Zhechkov, T. Heine, G. Seifert, PNAS 102 (2005) 10439–10444. [9] N.A. Kotov, Nature 442 (2006) 254–255. [10] G. Eda, M. Chhowalla, Nano Lett. 9 (2009) 814–818. [11] A.N. Obraztsov, Nat. Nanotechnol. 4 (2009) 212–213. [12] J.C. Meyer, C.O. Girit, M.F. Crommie, A. Zettl, Nature 454 (2008) 319–322. [13] C.T. Vincent, J.A. Matthew, Y. Yang, B.K. Richard, Nat. Nanotechnol. 4 (2009) 25–29. [14] S. Park, R. Ruoff, Nat. Nanotechnol. 4 (2009) 217–224. [15] M.J. McAllister, J. Li, D.H. Adamson, H.C. Schniepp, A.A. Abdala, J. Liu, M. Herrera-Alonso, D.L. Milius, Chem. Mater. 19 (2007) 4396–4404. [16] G. Lu, S. Mao, S. Park, R.S. Ruoff, J. Chen, Nano Res. 2 (2009) 192–200. [17] G. Maurin, F. Henn, B. Simon, J.F. Colomer, J.B. Nagy, Nano Lett. 2 (2001) 75–79. [18] S. Niyogi, E. Bekyarova, M.E. Itkis, J.L. McWilliams, M.A. Hamon, R.C. Haddon, J. Am. Chem. Soc. 128 (2006) 7720–7721. [19] D. Li, M.B. Muller, S. Gilje, R.B. Kaner, G.G. Wallence, Nat. Nanotechnol. 3 (2008) 101–105. [20] A.S. Patole, S.P. Patole, J.B. Yoo, J.H. Ahn, T.H. Kim, J. Polym. Sci. Part B: Polym. Phys. 47 (2009) 1523–1529. [21] S.S. Atik, J.K. Thomas, J. Am. Chem. Soc. 103 (1981) 4279–4280. [22] A.S. Patole, S.P. Patole, M.H. Sung, T.H. Kim, Elasto. Comp. 44 (2009) 34–40. [23] Z. Zhang, J. Zhang, P. Chen, B. Zhang, J. He, G.H. Hu, Carbon 44 (2006) 692–698. [24] A. Star, T.R. Han, J.P. Gabriel, K. Bradley, G. Gruner, Nano Lett. 3 (2003) 1421– 1423. [25] M.A. Pimenta, G. Dresselhaus, M.S. Dresselhaus, L.G. Cançado, A. Jorio, R. Saito, Phys. Chem. Chem. Phys. 9 (2007) 1276–1290. [26] C. Thomsen, S. Reich, Phys. Rev. Lett. 85 (2000) 5214–5217. [27] Z. Ni, Y. Wang, T. Yu, Z. Shen, Nano Res. 1 (2008) 273–291. [28] L. Jiang, L. Gao, J. Sun, J. Colloid Interface Sci. 260 (2003) 89–94. [29] M. Zhang, L. Su, L. Mao, Carbon 44 (2006) 276–283. [30] H. Oh, F. Peter, Nat. Mater. 8 (2009) 139–142. [31] P. Xiao, M. Xiao, K. Gong, Polymer 42 (2001) 4813–4816.
Contents lists available at ScienceDirect
Journal of Colloid and Interface Science www.elsevier.com/locate/jcis
A facile approach to the fabrication of graphene/polystyrene nanocomposite by in situ microemulsion polymerization Archana S. Patole a, Shashikant P. Patole b, Hyuck Kang a, Ji-Beom Yoo b, Tae-Ho Kim a, Jeong-Ho Ahn a,* a b
Department of Polymer Science and Engineering, Sungkyunkwan University, Suwon 440-746, Republic of Korea SKKU Advanced Institute of Nanotechnology (SAINT), Sungkyunkwan University, Suwon 440-746, Republic of Korea
a r t i c l e
i n f o
Article history: Received 21 November 2009 Accepted 13 January 2010 Available online 18 January 2010 Keywords: Polystyrene Graphene In situ microemulsion polymerization Functionalization Nanoparticles Morphology Dispersion Composite Thermal properties Electrical properties Transmission electron microscopy
a b s t r a c t This paper reports a large scale production route for polystyrene (PS) nanoparticle-functionalized graphene sheets using water based in situ microemulsion polymerization. The higher surface area of the graphene basal plane and the better proximity of the reactant species in in situ microemulsion polymerization were used to functionalize the graphene sheets using PS nanoparticles. The thermal properties of the PS were improved with the incorporation of graphene in the composite. The modified graphene exhibited good compatibility and interactions with the host PS matrix to form conducting PS films. Ó 2010 Published by Elsevier Inc.
1. Introduction Graphene have attracted considerable interest over the last few years on account of its extraordinary electrical, thermal, and mechanical properties arising from its unique structure [1]. This two dimensional sp2 bonded carbon structure has potential applications in many technological fields, such as high frequency analog electronics [2], single molecule sensors [3], nanocomposites [4], batteries [5], supercapacitors [6], liquids crystal display [7], and hydrogen storage [8]. One possible way of exploiting their properties for real world applications would be to incorporate graphene sheets in a composite material [9,10]. The manufacture of such composites requires that graphene sheets be produced on a sufficient scale but be incorporated and distributed homogeneously into various matrices. Over the last few years, many groups have reported either top down or bottom up approaches to synthesize graphene, including chemical vapor deposition (CVD) [11], mechanical exfoliation (repeated peeling) of graphite [12], chemical intercalation and exfoliation of graphite [13,14] and thermal expansion of graphite [15]. Of these routes, the thermal expansion
* Corresponding author. Fax: +82 31 292 8790. E-mail address: [email protected] (J.-H. Ahn). 0021-9797/$ - see front matter Ó 2010 Published by Elsevier Inc. doi:10.1016/j.jcis.2010.01.035
of graphite is an easy way of obtaining a bulk amount of graphene that allows it to be incorporated in various composite materials. The use of polymers structurally similar to the matrix polymer to functionalize graphene is a desirable approach for developing polymeric graphene composites. This ensures that the functionalized graphene is compatible with the polymer matrix to avoid potential microscopic phase separation in the composites. Therefore, an ideal polymeric graphene composite might be prepared using solubilized graphene that are functionalized with the matrix polymer. Polystyrene (PS) is one of the most widely used commercial polymers. The incorporation of functionalized graphene sheets in a PS matrix results in composite thin films that are semiconducting and exhibit an ambipolar field effect [10]. These composites can be deposited uniformly over large areas in the form of thin films from solution onto which the devices, such as thin film transistors, can be fabricated without extensive lithography. The dispersion of graphene into a PS matrix has stimulated considerable interest among researchers [9,10]. However, the lack of reactivity of the graphene basal plane needs to be solved. A perfect graphene does not have dangling bonds on the basal plane for chemical bond formation. The chemical potential of the basal plane is lower than the edges or defects. Therefore, graphene composites allow grafting on the edges and defects only. The dispersion and agglomeration of graphene is
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
another problem while forming the composite materials in bulk amounts. The hydrophobic nature and high specific surface area in graphene causes poor dispersion in a variety of organic solvents. Van der Waals forces cause graphene to agglomerate and restack, which destroys the properties of the individual graphene sheet. The functionalization of graphene sheets is an easy way of overcoming this problem. Until now, several strategies were proposed based on non-covalent and covalent bonding between inorganic and organic molecules on the surface of grapheme [9,16–19]. However, to date, there is no report on PS nanoparticle-functionalized graphene sheets for PS-graphene composite. In our previous work we presented effective in situ synthesis and characteristics of polystyrene nanoparticle covered multiwall carbon nanotube (MWCNT) composite by microemulsion polymerization [20]. This paper reports the use of microemulsion polymerization to functionalize graphene sheets by PS nanoparticles. This polymerization method has some merit over other emulsion methods because it produces thermodynamically stable polymer particles tens of nm in size, and uses the advantage of surfactants through polymerization [20–22]. The higher surface area in graphene and in situ microemulsion enhances the anchoring of PS nanoparticles on graphene through p–p interactions. The functionality of these nanoparticles affords better compatibility and interactions with the host PS matrix, thereby imparting improved thermal properties and enhanced conductivity. It is believed that this rational approach will provide a better method for functionalizing graphene in various host matrices in order to improve its applications in the real world. 2. Materials and methods 2.1. Materials Commercially available expandable graphite (EG) with >99.9% purity was used to obtain a few layers of graphene sheets. Styrene, sodium dodecyl sulfate (SDS), azobisisobutyronitrile (AIBN), PS beads (Mw-280,000) and 1-pentanol were purchased from Sigma Aldrich and used without further purification. Tetrahydrofuran (THF) and methyl alcohol as solvents were purchased from Aldrich Co. and used as received. 2.2. Graphene synthesis In a coexisting research direction, before selecting the material for experiments, we explored the rapid thermal expansion of EG. Initially, 50 mg of as-purchased EG were loaded in the CVD chamber using an alumina boat with a lid. The CVD chamber was evacuated to <0.01 Torr using a rotary pump. The ramping rate was set for 800 °C/min using a tungsten lamp heater and programmable heater controller. Thermal expansion was carried out at various temperatures (700, 800, 900 and 1000 °C) for two minutes. The samples were then cooled to room temperature and removed from the CVD chamber for further experimentation. This scheme can be used for the mass production of graphene depending on the CVD reactor size.
531
containing 200 ml deionized water, 4.5 ml pentanol and 1.5 g SDS for 1 h. Pentanol was used as a hydrophobe and SDS was used as the surfactant. After dispersion, the aqueous graphene mixture was charged into a 500 ml four-neck glass reactor equipped with a condenser, dropping funnel, stirrer and nitrogen inlet. The reactor was then placed in an ice bath. Styrene (5 g) and AIBN (0.05 g), which were the monomer and initiator, respectively, were mixed through a funnel prior to polymerization. The solution was deoxygenated with oxygen-free nitrogen for 20 min by bubbling nitrogen gas. The mixture was then sonicated at 0 °C for 4 h. After homogeneous mixing, the ice bath was replaced with an oil bath for heating. Polymerization was then carried out by increasing the oil bath temperature to 85 °C for 4 h under an inert nitrogen blanket. After polymerization, the grey colored emulsion obtained was washed with methanol and water to remove the excess surfactants, hydrophobe and residual monomer from the reaction mixture. The procedure was repeated more than three times. The self-assembled graphene/PS nanoparticle composite was then dried under vacuum at 60 °C for 24 h, and used for further characterization. As a comparison, the pristine polymer nanoparticles were prepared using the same procedure in the absence of graphene. 2.4. Preparation of conducting PS films PS nanoparticle-functionalized graphene was used to prepare the PS films. In typical experiments, a 2.5 wt.% graphene/PS nanoparticle sample was selected as a filler to make the composite films. The host matrix was prepared by dissolving 5 mg of PS beads in 20 ml of THF. 3–20 wt.% filler was added to this host matrix. The solution was then cast on a glass plate and dried at room temperature to form the free standing, 600 lm thick and uniform films. Fig. 1 summarizes the process steps. 2.5. Characterizations The morphology of graphene, pristine PS nanoparticles and graphene/PS nanoparticles was examined by scanning electron microscopy (FESEM; JSM6700F, JEOL). The samples were dispersed in methanol and drop coated onto a copper grid for high resolution transmission electron microscopy (HR-TEM; JEOL 300 kV) to explore the structural properties. The spectroscopic analyses were carried out using a Fourier transform infrared spectrometer (FTIR-Nicolet IR 200) and a micro-Raman spectrometer (Invia Basic, Renishaw Co. England) using an Ar-ion laser at 640 nm as the excitation light source. Thermo-gravimetric/differential thermal analysis (TG/DTA) of graphene, PS nanoparticles and graphene/PS nanoparticles were carried out using a thermobalance (TGA 2050) from room temperature to 800 °C, at a rate of 10 °C/min in a continuous nitrogen flow. The dynamic differential scanning calorimetry (DSC) experiments were carried out with a thermal analyzer (DSC 2920) at heating rates of 5 °C/min in a continuous nitrogen flow. The sheet resistance (Rs) of the composite films was measured using a 4-point probe resist meter (AIT CMTSR1000 N).
2.3. Ultrasonically initiated in situ microemulsion polymerization 3. Results and discussion An ultrasonic irradiation device (Sonomaster Model ULH-700S) was equipped with a standard aluminium horn with a replaceable tip diameter of 13 mm and temperature controller. The nominal energy output of the probe was set to 300 W. The PS nanoparticle-functionalized graphene was prepared using the following procedure. First, approximately 0.15–1 g of graphene, which was thermally expanded at 1000 °C was sonicated in a beaker
Fig. 2a shows a photograph of the as-purchased EG powder and thermally expanded graphene. The temperature used for thermal expansion was 1000 °C. The bulk volume of the expanded graphene was much higher than that of the EG powder. It is evident that the EG expanded at 1000 °C. The bulk volume of thermally expanded graphene increased with increasing heating temperature.
532
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 1. Multi step process for producing the conducting graphene–PS films.
Fig. 2b shows a SEM image of EG, which consists of several graphene stacks. Here, the main challenge was to separate these stacks to obtain the individual graphene flakes. Van der Waals forces keep these stacks together. The heating rate of 800 °C/min used for expansion is more than sufficient to overcome this force to obtain a few layer graphene sheets [15]. The samples heated to 700 and 800 °C showed a much larger number of graphene layers than the samples expanded at 900 and 1000 °C. Fig. 2c–d shows SEM images of the graphene flakes expanded at 1000 °C. In Fig. 2c, the folded graphene flake had the appearance of a crumbled piece of paper. A higher magnification image shows the folded edge of the typical graphene flakes (Fig. 2d). A few layers of graphene (1–10 layers) were observed by HR-TEM. Fig. 2e shows a one layer graphene flake at its edge and Fig. 2f shows two layers of graphene at its edge. A higher number of graphene layers (up to 35 graphene layers) were observed in the case of the 700 and 800 °C expanded samples. The samples expanded at 1000 °C were chosen for further experimentation. In situ polymerization was carried out to functionalize the graphene flakes with the PS nanoparticles. The main motivation behind the in situ reaction was to increase the proximity of the reactant species and use a higher surface area of graphene flakes to anchor the PS nanoparticles. Previously, in the case of MWCNT, a higher surface area was reported to provide better anchoring of PS nanoparticles to the outer walls of the MWCNTs [20]. A similar analogy was used for the present study, rather in the case of graphene, a higher surface area provides better proximity for the reactant species to anchor the PS nanoparticles. Fig. 3a–c shows SEM images of the dried composite. The images show that the composite material consists of randomly aggregated, thin, crumpled graphene flakes closely associated with each other and a polystyrene nanoparticle covered on it, forming a disordered solid. The
individual graphene flake size was approximately 5–10 lm. The high resolution images (Fig. 3b–c) confirmed that several PS nanoparticles were anchored on both sides of the graphene flakes. At the resolution limit of our instrument (see Fig. 3c–d), the data suggests, but does not prove, the presence of individual flakes in this material. The absence of charging during SEM imaging indicates that the network of the graphene based sheets and the individual sheets were electrically conducting. This qualitative conclusion was further confirmed by the electrical measurements on graphene/PS nanoparticle composite thin films. Fig. 3d–f shows TEM images of the PS nanoparticles-covered graphene flakes. The PS nanoparticles (with average diameter of 37.5 nm) were distributed randomly over the graphene flakes. The number of PS nanoparticles on the graphene flakes depends on the amount of starting graphene flakes; the number of PS nanoparticles decreases with increasing number of graphene flakes. Fig. 3d shows a TEM image of the 3 wt.% composite and Fig. 3e shows a TEM image of the 20 wt.% composite. HR-TEM analysis confirmed that the individual graphene flakes consists of 1–10 graphene sheets, which is consistent with the thermally expanded samples at 900 and 1000 °C (see Fig. 2e–f). A closer view (see Fig. 3f) shows the attachment of PS nanoparticles on graphene or incorporation into the graphene sheet. It appears that Brownian motion of the emulsified monomer droplets and graphene inside water causes collisions between these two species, which provides further anchoring sites for juvenile PS nanoparticles to attach to the graphene surface. It was assumed that the hydrophilic ends of SDS on graphene and on the emulsified droplets caused repulsion, which was overcome by the kinetic energy of Brownian motion of these species. During the collision, SDS at the contact region might have been displaced or removed, and provided more stable sites for anchoring the PS nanoparticles on the graphene surface.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
533
Fig. 2. (a) Digital camera photograph of the as-purchased EG and thermally exfoliated graphene flakes. SEM images showing, (b) the graphite stacks in the EG, (c) crumbled graphene flake and (d) edge of the graphene flake. HR-TEM images showing, (e) single layer and (f) two layers graphene.
The in situ reaction and larger surface area of the graphene sheet provided good adhesion at the contact regions. This adhesion can be attributed to some interaction between the phenyl groups in PS and graphene through p–p stacking [23,24]. The HR-TEM image in Fig. 3f confirms the hypothesis. Fig. 4 shows the proposed hypothesis which was applicable to the higher surface area MWCNTs [20]. In the first step, graphene was functionalized by the anionic SDS surfactant. The graphene was isolated from the water by SDS, keeping the hydrophobic end towards the surface of the graphene. Similarly, the emulsified droplets were formed by the SDS coverage, maintaining the hydrophilic ends of SDS towards water. During polymerization, collisions between the graphene and emulsified monomer due to Brownian motion caused the selective etching of SDS. This provides anchoring sites on the graphene for nanoparticles. There may be possibility that the anchor of the PS particles is due to the weak Van der Waals force originating from the random contact of carbon atoms and PS chain. The attachment of PS nanoparticles to the graphene sheet becomes a probabilistic event due to collisions, and does not cause the selective anchoring of all PS nanoparticles, the unan-
chored PS latex and SDS were removed by successive washing in methanol. Finally, the PS nanoparticles-covered graphene composite was obtained after successive washing. FTIR spectral analysis was performed to confirm the chemical structure of the coating of PS nanoparticles on the graphene sheet. Fig. 5a shows the FTIR spectra of the pristine graphene, pristine PS nanoparticles and graphene/PS nanoparticle composite. Although no visible peaks were observed in graphene, there were several peaks in the pristine PS nanoparticles and composites. Prominent peaks in PS nanoparticles and composite at approximately 2820– 2967 cm 1, 2994–3043 cm 1, 1500 cm 1 and 700–1400 cm 1 indicate aliphatic C–H stretching, aromatic C–H stretching, aliphatic CH2 and different conformation sensitive vibration modes of polystyrene, respectively. This confirms that the polymerization reaction produced PS from styrene. A new peak at approximately 3036 cm 1 appeared in case of composite, which could be attributed to the C–H stretching of aromatic ring. There was no significant change in the peak positions apart from the peak intensity, indicating that the PS nanoparticles in pristine and composite forms have the same structural properties. The anchoring of the
534
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 3. SEM and TEM images of the graphene sheets functionalized by PS nanoparticles. In figure (a) the micron size of the sheets are revealed. (b and c) The higher magnification images show that the PS nanoparticles are attached to both sides of the graphene sheets. TEM images of (d) 3 wt.% and (e) 20 wt.% samples. (f) HR-TEM image of the attachment of PS nanoparticles on the graphene flake. The interface region in (f) is marked with an arrow.
Fig. 4. Step wise functionalization of graphene sheets with PS nanoparticles.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
Fig. 5. (a) FTIR spectra of the pristine graphene, pristine PS nanoparticles and graphene/PS nanoparticles. (b) Raman spectra of pristine graphene and graphene/PS nanoparticles.
nanoparticles occurs at the surface of the graphene sheet, which maintains the inner structure of PS nanoparticle. However, the possibility of a structural change on the surface of the graphene sheet cannot be excluded. Raman spectroscopy is a powerful nondestructive tool for characterizing carbonaceous materials, particularly for distinguishing ordered and disordered crystal structures of carbon. The major Raman features of graphene and graphite are the so called G (1580 cm 1) and 2D (2670 cm 1) bands. The G band originates from the in-plane vibration of sp2 carbon atoms and is a doubly degenerate (TO and LO) phonon mode (E2g symmetry) at the Brillouin zone center [25]. The 2D band originates from a two phonon double resonance Raman process [26]. The 2D band is the obvious difference between the Raman features of graphene and graphite. For graphene, the 2D band can be fitted with a sharp and symmet-
535
ric peak while that of graphite can be fitted with two peaks [27]. The supporting material (see Fig. S1) shows that in the case of graphene, the 2D band sharpens and becomes blue shifted. In addition to the differences in the 2D band, the intensity of the D band decreases almost linearly (as shown in Supporting Fig. S1) with decreasing graphene thickness. A comparison of the Raman spectra of graphene and the graphene/PS nanoparticle composite (Fig. 5b) revealed broadening of the 2D band in graphene due to the creation of defects. When the D band of graphene and composite Raman (Fig. 5b) were compared, the peak at 1354 cm 1 appeared after introducing polystyrene nanoparticles on the top of graphene, which is likely because the PS nanoparticles anchored on the grapheme basal plane, which increased the disorder in graphene. The intensity ratio I(D)/I(G) was increased significantly after polystyrene modification. Thermo-gravimetric analysis (TGA) was performed on graphene, PS nanoparticles and composites at all concentrations to obtain additional confirmation of both the graphene content and the structure of the composite, as well as to determine the effects of the graphene on the thermal stability of the composite. TGA was carried out under a nitrogen atmosphere. The resulting curves are shown in Fig. 6a. In the case of graphene, an initial 8% weight loss at <350 °C revealed the oxidation of amorphous carbon present in the sample. Pristine PS nanoparticles began to decompose at 380 °C and degraded completely at 401 °C. In the case of the pristine PS nanoparticles and composites, the weight loss below 200 °C was attributed to the decomposition of the lower molecular weight PS and the weight loss above 200 °C was attributed to the decomposition of higher molecular weight PS. However, the mass that remained constant above 400 °C was assigned to graphene. This method can determine the actual graphene content in the composites. The degradation temperature of the composites increased with graphene content. A maximum increase of 16 °C was observed for the 20 wt.% composite (refer Table 1). The first derivative of the TGA curve (DTA curves shown in Fig. 6b) clearly shows the variation in weight with time (dW/dT) as a function of temperature. The DTA peaks indicate the temperature of the maximum reactive velocity. Table 1 lists the maximum reactive velocity of the DTA curve of pristine polystyrene and graphene composites. The pristine PS nanoparticles reach the maximum reactive velocity at 401 °C and the value was lowest of all the curves obtained. The temperature at the maximum reactive velocity increased with increasing graphene content, and for the 20 wt.% composite it was increased by 15 °C compared to the pristine PS nanoparticles. This indicates that the thermal stability of the PS nanoparticles was improved by the addition of graphene. Previous studies on CNT/PS nanoparticle composites reported that the introduction of CNTs to the PS system can enhance the thermal stability of CNT–PS composites due to an interaction between the outer walls of the CNTs and PS lattices. Moreover, the formation of a barrier of CNT inhibits mass transfer and provides thermal insulation to shield the underlying polymer from the heat source [28,29]. Fig. 6c shows the glass transition temperature (Tg), which defines a pseudo second order phase transition. The incorporation of PS nanoparticles on graphene sheets resulted in an increase in Tg. In the case of pristine PS nanoparticles, the Tg was observed at 96 °C and was shifted to 113 °C for the 20 wt.% composite (see Table 1). Such an effect on the Tg obtained may be due to motion restrictions of the polymer chains at the nanoparticle–graphene interface [30]. A region of a strongly bonded polymer chains might have formed near the graphene surface. In this region, the polymer chains exhibited different behavior than that in the nanoparticle. The strong packing hinders chain segmental mobility. Therefore, more energy is needed to allow the first thermal transition, shifting Tg to a higher temperature. In the present case, similar effect might
536
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537 Table 1 Thermal properties of nanocomposites.
a b
Sample
Tdeg (°C)a
Graphene content (wt.%)a
Tg (°C)b
PSNP 3 wt.% 5 wt.% 10 wt.% 15 wt.% 20 wt.%
401 410 413 415 417 417
– 5.28 7.0 19.0 22.0 31.0
96 106 107 108 110 113
Determined by TGA. Determined by DSC.
Fig. 6. Thermal properties of the graphene/PS nanoparticles composite. (a) TGA, (b) DTA and (c) DSC analysis. The net amount of graphene estimated from TGA curves above 450 °C and the shift in glass transition temperature from DSC are listed in Table 1.
have occurred due to in situ polymerization. The observed increase in Tg with increasing graphene amount suggests good interfacial compatibility between the PS and graphene sheet. The ultimate goal of making graphene/PS nanoparticle composites was to develop a lightweight but highly conducting material. The PS nanoparticle-functionalized graphene flakes were used to make PS films. The details of the process steps involved in making the conducting PS films are discussed in the experimental section. The graphene/PS nanoparticle-enforced PS films were both
Fig. 7. Digital camera photograph shows (a) functionalized graphene embedded PS films. The number indicates a content of 2.5 wt.% functionalized graphene in the film. The inset shows the slurry in THF used for making the PS films. (b–e) Photographs show the flexibility of the composite films. (f) Sheet resistance of the composite films as a function of filler amount.
A.S. Patole et al. / Journal of Colloid and Interface Science 350 (2010) 530–537
conducting and flexible. Fig. 7a shows a photograph of the different wt.% composite films. Fig. 7b–e highlight the flexibility of the film. Fig. 7f shows the sheet resistance of these films as a function of the graphene content. The sheet resistance decreased several orders of magnitude with increasing graphene concentration due to percolation. For a 3 wt.% graphene content, the sheet resistance was 36.67 MX/h, which increased to 65.46 X/h for a 20 wt.% graphene content. These properties are superior to those reported elsewhere [9]. The sheet resistances differed slightly between the top and bottom of the film. This might be due to the sedimentation effect while forming the films by a wet cast method. No anisotropy effect was observed in these samples. It should be noted that at a 3 wt.% graphene content, the conductivity of the samples (4.5 10 5 S m 1) already satisfied the antistatic criterion (10 6 S m 1) for thin films and rapidly increased over the 20 wt.% range. The electrical properties of these composites compare well with the best values reported for nanotube–polymer composites thus far [31]. In the present study, the approach behind dispersing the PS nanoparticle-functionalized graphene in a PS host matrix was to achieve good dispersion and compatibility between the graphene and PS matrix (see Supporting Fig. S2). Therefore, it is expected that this straightforward, large-scale approach for the preparation, functionalization and incorporation of graphene sheets into polymer matrices will lead to further developments of a broad new class of materials with enhanced properties, and even allow the introduction of new functionalities to polymer composites in bulk quantities and reduces production cost. 4. Summary A large scale production route for PS nanoparticle-functionalized graphene sheets using water based in situ microemulsion polymerization was developed. SEM and HR-TEM showed the successful anchoring of PS nanoparticles on the surface of graphene. The thermal properties of the PS were improved with the incorporation of graphene in the composite. The modified graphene showed good compatibility and interactions with the host PS matrix to form conducting PS films. The reported scheme for fabricating the PS composite thin films from graphene and a commodity plastic highlights the potential for low-cost, macroscale thin film electronics. Acknowledgments This work was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2010-
537
0015035) and Ministry of Knowledge and Economy (Project No. 10037449). The authors also appreciate the project and equipment support from Gyeonggi Province through the GRRC program in Sungkyunkwan University. One of the authors (S.P.P.) is grateful to the Sungkyunkwan University for awarding him a Postdoctoral Research Fellowship. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.jcis.2010.01.035. References [1] A.K. Geim, K.S. Novoselov, Nat. Mater. 6 (2007) 183–191. [2] N. Tombros, C. Jozsa, M. Popinciuc, H.T. Jonkman, B.J.V. Wees, Nature 1448 (2007) 571–575. [3] J.M. Carlsson, Nat. Mater. 6 (2007) 801–802. [4] T. Ramanathan, A.A. Abdala, S. Stankovich, D.A. Dikin, M. Herrera-Alonso, R.D. Pinar, D.H. Adamson, H.C. Schnipp, X. Chen, R.S. Ruoff, S.T. Nguyen, Nat. Nanotechnol. 3 (2008) 327–331. [5] M. Liang, L. Zhi, J. Mater. Chem. 19 (2009) 5871–5878. [6] S.R.C. Vivekchand, C.S. Rout, K.S. Subrahmanyam, A. Govindaraj, C.N.R. Rao, J. Chem. Sci. 120 (2008) 9–13. [7] J. Zhu, Nat. Nanotechnol. 3 (2008) 528–530. [8] S. Patchkovskii, J.S. Tse, S.N. Yurchenko, L. Zhechkov, T. Heine, G. Seifert, PNAS 102 (2005) 10439–10444. [9] N.A. Kotov, Nature 442 (2006) 254–255. [10] G. Eda, M. Chhowalla, Nano Lett. 9 (2009) 814–818. [11] A.N. Obraztsov, Nat. Nanotechnol. 4 (2009) 212–213. [12] J.C. Meyer, C.O. Girit, M.F. Crommie, A. Zettl, Nature 454 (2008) 319–322. [13] C.T. Vincent, J.A. Matthew, Y. Yang, B.K. Richard, Nat. Nanotechnol. 4 (2009) 25–29. [14] S. Park, R. Ruoff, Nat. Nanotechnol. 4 (2009) 217–224. [15] M.J. McAllister, J. Li, D.H. Adamson, H.C. Schniepp, A.A. Abdala, J. Liu, M. Herrera-Alonso, D.L. Milius, Chem. Mater. 19 (2007) 4396–4404. [16] G. Lu, S. Mao, S. Park, R.S. Ruoff, J. Chen, Nano Res. 2 (2009) 192–200. [17] G. Maurin, F. Henn, B. Simon, J.F. Colomer, J.B. Nagy, Nano Lett. 2 (2001) 75–79. [18] S. Niyogi, E. Bekyarova, M.E. Itkis, J.L. McWilliams, M.A. Hamon, R.C. Haddon, J. Am. Chem. Soc. 128 (2006) 7720–7721. [19] D. Li, M.B. Muller, S. Gilje, R.B. Kaner, G.G. Wallence, Nat. Nanotechnol. 3 (2008) 101–105. [20] A.S. Patole, S.P. Patole, J.B. Yoo, J.H. Ahn, T.H. Kim, J. Polym. Sci. Part B: Polym. Phys. 47 (2009) 1523–1529. [21] S.S. Atik, J.K. Thomas, J. Am. Chem. Soc. 103 (1981) 4279–4280. [22] A.S. Patole, S.P. Patole, M.H. Sung, T.H. Kim, Elasto. Comp. 44 (2009) 34–40. [23] Z. Zhang, J. Zhang, P. Chen, B. Zhang, J. He, G.H. Hu, Carbon 44 (2006) 692–698. [24] A. Star, T.R. Han, J.P. Gabriel, K. Bradley, G. Gruner, Nano Lett. 3 (2003) 1421– 1423. [25] M.A. Pimenta, G. Dresselhaus, M.S. Dresselhaus, L.G. Cançado, A. Jorio, R. Saito, Phys. Chem. Chem. Phys. 9 (2007) 1276–1290. [26] C. Thomsen, S. Reich, Phys. Rev. Lett. 85 (2000) 5214–5217. [27] Z. Ni, Y. Wang, T. Yu, Z. Shen, Nano Res. 1 (2008) 273–291. [28] L. Jiang, L. Gao, J. Sun, J. Colloid Interface Sci. 260 (2003) 89–94. [29] M. Zhang, L. Su, L. Mao, Carbon 44 (2006) 276–283. [30] H. Oh, F. Peter, Nat. Mater. 8 (2009) 139–142. [31] P. Xiao, M. Xiao, K. Gong, Polymer 42 (2001) 4813–4816.