Effect of polystyrene long branch chains on melt behavior and foaming performance of poly(vinyl chloride)/graphene nanocomposites

Abstract

Several poly(vinyl chloride)-g-polystyrene graft copolymers (PVC-g-PS) with well defined molecular structures were synthesized via atom transfer radical polymerization (ATRP) from the structural defects of PVC. The effects of PS branch chains on the shear and extensional rheology as well as foaming properties were investigated. Compared to linear PVC, the introduction of PS branches results in increased complex viscosity, an elevated value of storage modulus at low shear frequencies, more pronounced shear-thinning behavior, more significant upshifted deviation from linear behaviour and a strain hardening phenomenon. Under the same foaming conditions, most of the resulting PVC-g-PS foams exhibit a closed cell structure, increased cell density and uniform cell size distribution while the linear PVC foam has serious cell coalescence. Moreover, graphene nanosheets could be well dispersed in the PVC-g-PS matrix due to the π–π stacking with PS relative to the PVC without PS branch chains. As expected, both the nucleation effect and increased melt viscosity from well-dispersed graphene sheets significantly improve the foaming behavior of PVC-g-PS/graphene nanocomposites, in comparison with the poor foamability of PVC/graphene composites due to the non-uniform dispersion of graphene.

---

1 Introduction

The relationship between structure and properties of polymers is a long-term issue in polymer science. The topological structure of polymer chains plays an important role in polymer performance, especially the melt behavior of polymers with long chain branch (LCB) structures.^1–9 For example, long chain branched polypropylene (LCBPP) with just a few LCBs can exhibit prominent strain hardening behavior, high melt strength and more obvious shear-thinning behavior.^2,3 These changes have a great effect on the foaming behavior of LCBPPs, such as higher volume expansion ratios, retarded cell coalescence, and a more homogeneous cell structure compared to linear PP.^6–9 Some recent reports showed that a short LCB (the molecular weight of the branch chain (M_b) is lower than its critical entanglement molecular weight (M_c)) also plays an important role in changing melt behavior, when the LCB is immiscible with the polymer backbone.^10–12 This can be explained according to the “tube model”.¹³ Before the relaxation of all LCB-polymer chains is realized, the branch chains should recoil to the region of the backbone. When the branch chain is immiscible with the backbone, the recoil of the branch chains is slow due to the backbone repulsion of the branch chain.

Recently the influence of nanofillers on melt behaviour of polymers has been a hot topic. The addition of nanoparticles changes rheological properties, strongly depending on the shape, the amount of nanoparticles, the dispersed state of nanoparticles and interaction between nanoparticles and polymer matrix.^14–17 For example, the presence of layered nanofillers (such as clay and graphene) in the polymer matrix significantly increases the melt viscosity of polymers at low shear frequency,^18,19 but the presence of spherical nanoparticles decreases the melt viscosity of polymers at low shear frequency,^20,21 when the M_w of polymer is larger than its M_c and R_p < R_g (R_p: the radius of nanoparticle; R_g: the radius of gyration of the polymer). However, most of researches about polymer nanocomposites are focused on linear polymer systems. The research on LCB-polymer nanocomposites is seldom.²² Also the presence of nanoparticles, including nanoclay, carbon nanofibers (CNFs), and carbon nanotubes (CNTs), plays a key role in melt processing, for instance, acting as nucleating agents in polymer foaming.^23–26 As a result, the cell density and size of polymer nanocomposites change dramatically compared to pure polymers. Recently, the scientists found graphene applied in polymer foaming as a nucleating agent can also receive attractive effect due to its high specific surface area and physical barrier effect to cell coalescence.^27,28

As demands for lightweight materials grow greatly, research on the polymeric foams has been stimulated.^29,30 Thermoplastic foams have great advantages, such as high impact strength, reduced weight, energy absorption and low thermal conductivity.^31–34 These unique properties make them ideal for a large number of applications. Commerical polyvinyl chloride (PVC) is a uniquely versatile linear polymer with the third largest tonnage after LDPE and PP, which is applied in many industry sectors.^35–39 In the fabrication of PVC foams, there is always a contradiction between the molecular weight and processability. It is difficult to shape PVC with high molecular weight; while the molecular weight was lower, it is hard to prepare PVC foam with satisfying quality due to its weak melt strength. Thus it is an important issue how to adjust melt behavior and increase the melt strength of PVC. To the best of our knowledge, how the branch structure influence the melt strength of PVC has not been investigated, furthermore, the branch structure and foaming behavior of PVC have not yet been correlated.

In this work, we demonstrated a novel way to improve the foaming behavior of PVC through introducing polystyrene (PS) long branch chains onto PVC backbone via ATRP route and adding graphene as nanofiller. The reason why PS branch chains are grafted onto PVC backbone results from the dual roles of PS chains. Here PS branch chain is immiscible with PVC backbone but compatible with graphene sheets. The natural incompatibility with polar polymer surface poses significant challenges to achieving complete exfoliation of graphene in PVC/graphene nanocomposites. The presence of PS branch chains onto PVC backbone can probably promote the exfoliation and dispersion of graphene in the matrix via the interaction between PS and graphene. Two fundamental issues were focused. On one hand, the effects of PS branches on rheological and foaming properties of PVC were studied. On the other hand, the effects of PS branches on morphology, rheological and foaming properties of PVC/graphene nanocomposites were studied. It is expected that the above research will promote fundamental understanding about the influences of long branch chains and the coexistence of long branch chains and nanofillers on rheological behavior and foam behavior of polymers.

2 Experimental section

2.1 Materials

Polyvinyl chloride (PVC) was purchased from Aldrich (M_w = 80 000 g mol⁻¹). Styrene (St), and N,N,N′,N′′,N′′-pentamethyldiethylenetriamine (PMDETA) were vacuum-distilled over calcium hydride before use. CuBr was purchased from Aldrich. Graphite powder, Di-n-octylo-phthalate (DOP) and dibutyltin dilaurate were purchased from Sinopharm Chemical Reagent Beijing Co., Ltd. DMF and THF (AR grade, from Beijing chemical works) were refluxed over metallic sodium with benzophenone and distilled under argon atmosphere prior to use. 98% H₂SO₄, 30% H₂O₂, and KMnO₄ were purchased from Shanghai Zhenxin Chemical Company. n-Octadecylamine (ODA) and hydroquinone were obtained commercially from Fuchen Chemical Reagents (Tianjin, China). All oxygen- and moisture-sensitive manipulations were carried out under dry and oxygen free argon atmosphere using standard Schlenk techniques.

2.2 Synthesis of functional graphene nanosheets

Graphene oxide was prepared with improved Hummers' method⁴⁰ and synthesis of FGS was according to the previously published procedure.⁴¹ In a typical process, a 9 : 1 mixture of concentrated H₂SO₄–H₃PO₄ (360 : 40 mL) was added to a mixture of graphite flakes (3.0 g, 1 wt equiv.) and KMnO₄ (18.0 g, 6 wt equiv.), The reaction was then heated to 50 °C and stirred for 12 h. The reaction was cooled to room temperature and poured onto ice (400 mL) with 30% H₂O₂ (3 mL). After multiple-wash process, the solid obtained was vacuum-dried at room temperature.

Graphene oxide (2.5 g) was dissolved and exfoliated in 800 mL deionized water via ultrasonication, the resulting suspension was mixed with the solution of ODA (3.5 g) in 150 mL ethanol in a three-neck flask. The mixture was refluxed with magnetic stirring for 20 h at 90 °C, then, 2.5 g hydroquinone was added, and the mixture was refluxed for another 20 h to reduce it to graphene nanosheet-ODA. The mixture was then purified with centrifugation and dried in an oven at 80 °C for 24 h. The content of ODA in the functional graphene nanosheets (FGS) is about 28 wt% measured by means of Elementar (Vario EL, Germany).

2.3 Preparation of PVC-g-PSx composites with FGS

The synthesis of PVC-g-PS graft copolymers was described in ESI.† PVC composites were prepared by solution blending. Typically, FGS was dispersed and exfoliated in THF by the ultrasonicator for 0.5 h. The resulting suspension was mixed with PVC-g-PSx, designed amounts of DOP (20 wt%) and dibutyltin dilaurate (5 wt%) which were dissolved in THF. After the mixture was homogenized at ambient temperature for 0.5 h, it was poured into a Teflon Petri dish and kept at 50 °C for film formation until its weight equilibrated. A series of PVC-g-PSx/FGS nanocomposite films (with 1 wt% FGS) were peeled off of the substrate, 1.5 mm thick circular disks by compression molding at 180 °C and 10 psi for 5 min.

2.4 Batch foaming

A stainless steel high-pressure vessel was used for batch foaming process. The samples with 1.5 mm thick circular disks and containing some amounts of plasticizer and thermal stabilizer were used in the foaming experiments. The high-pressure vessel was loaded with several samples. After the high-pressure vessel was purged with low-pressure CO₂, a given amount of CO₂ was pumped into the vessel. The vessel was heated to 40 °C and continuously charged with CO₂ to the pressure of 6 MPa. The samples were saturated under 6 MPa at 40 °C for 96 h to reach equilibrium adsorption of CO₂. Thereafter, the valve was rapidly opened to release the CO₂ in 2 s. Then the high-pressure vessel was opened up, and the saturated samples were removed from the pressure vessel and immediately foamed in a glycerin bath maintained at 110 °C for 3 min in order to reach steady-state foam structure. The bath temperature was controlled to 110 ± 0.1 °C, which will be referred to as the foaming temperature. After the foaming, all the samples were quenched in room-temperature water and dried under vacuum at 60 °C for 48 h. All the samples were foamed under the same conditions.

2.5 Characterization

¹H NMR spectra were recorded on a Bruker AV 400 MHz spectrometer with odichlorobenzene-d₄ as solvent. Relative molecular weight was determined by gel permeation chromatography (GPC) on TOSOH HLC 8220 GPC at 40 °C using THF as an eluent against linear polystyrene standards. The characteristic groups of PVC and PS were characterized by Fourier transform infrared spectroscopy (FT-IR, Bio-Rad FTS-135). Glass transition temperature (T_g) of the PVC-g-PSx samples was determined by differential scanning calorimetry (DSC) performed on a Metter Toledo star system (DSCI), using a heating rate of 10 °C min⁻¹ in the temperature range 0−180 °C under nitrogen. The above results were placed in ESI.† The rheological experiments of samples were performed on a strain-controlled rheometer (ARES-G2, TA Instruments) using a parallel plates geometry (25 mm diameter). The temperature was controlled by a forced convection oven (FCO) unit under a nitrogen gas purge to minimize the degradation of the samples. Oscillatory frequency sweeps were performed from 0.05 to 100 rad s⁻¹ with a strain at 1% in linear viscoelastic regime at 185 °C. The elongational viscosity of samples (size: 25 mm × 10 mm × 0.1 mm) was measured using a rheometer MCR 301 from Anton Paar instruments. Experimental temperature was 180 °C under a nitrogen gas purge. Extensional rate was set at 0.5 and 0.05 s⁻¹, using a UXF fixture accessory. Morphologies of the PVC-g-PSx/FGS composites were observed by transmission electron microscope (TEM, JEOL1011) at a 100 kV accelerating voltage. Ultrathin sections were cut using a Leica Ultracut and a glass knife under the frozen condition. The samples were collected on carbon-coated copper TEM grids. The cell morphology of the foamed samples was characterized by an XL 30 ESEM FEG scanning electron microscope (SEM). The foamed samples were fractured in liquid nitrogen and SEM images of the fractured surfaces were taken. The cell size and cell density were obtained by image analysis software. For statistical accuracy, the cell diameter (D) was the average size of at least 100 cells in the SEM image. Because the cells in cross section were mostly polygonal, the long axis was taken as cell diameter. The cell density (N₀), the number of cells per cubic centimeter of solid polymer, was calculated from eqn (1):⁴

N₀ = [n/A]^3/2ρ/ρ_f (1)

where n is the number of cells in the SEM micrograph; A is the area of the micrograph (cm²), calculated according to the scale bar; ρ and ρ_f are the mass densities of samples before and after foaming experiment, respectively, which can be measured by water displacement method.

3 Results and discussion

3.1 Effect of PS branch chains on the rheological properties and foaming behaviour of PVC-g-PS graft copolymers

The preparation of PVC-g-PS graft copolymers with labile chlorines of PVC has been reported previously,⁴² however, how PS branches influenced the physical properties of PVC has not yet been investigated, especially rheological properties. Fig. 1a shows the complex viscosity (η*) versus angular frequency (ω) at 185 °C for PVC-g-PSx. We compare the change of η* at the frequency of 0.1 rad s⁻¹, and the values are listed in Table 1. From Fig. 1a and Table 1, when PS was grafted on the backbone of PVC, the complex viscosity of PVC-g-PSx at low frequency region increases gradually with the length of LCB. Furthermore, the presence of LCB leads to more prominent shear-thinning phenomenon with the increasing length of LCB. For example, the PVC-g-PS18400 shows an obvious shear-thinning phenomenon compared to PVC and PVC-g-PS5100, although the length of PS branch chains (18 400 g moL⁻¹) is lower than the critical entanglement molecular weight (M_c) of PS (about 32 000 g mol⁻¹ (ref. 43 and 44)). Besides melt viscosity, the storage modulus is also sensitive to LCB. Fig. 1b shows the G′ versus ω for the PVC and PVC-g-PSx samples at 185 °C. Compared to PVC, the PVC-g-PSx samples exhibit higher G′ at low shear frequencies, and the terminal slope of G′ is less than that of PVC, especially PVC-g-PS18400 with PS branch chains of 18 400 g mol⁻¹. This shows that the presence of PS branch chains with the length of 18 400 g moL⁻¹ is long enough to dramatically change the rheological behavior of PVC. This is attributed to the following reasons. The retraction of PS branches in PVC-g-PSx melts is speculated to be impeded by two factors: one is the repulsion between PS branches and PVC backbones resulting from their incompatibility; the other is the interaction between PS branches (π–π interaction between PS branches). Therefore, the relaxation time of PVC-g-PS is much longer than that of linear PVC.

Fig. 1 (a) Complex viscosity vs. angular frequency, and (b) storage modulus vs. angular frequency for PVC-g-PS samples with different length of PS branches at 185 °C.

Table 1 Summary of the foams cell parameters

Samples	St content (mol%)	Branch length (g mol^−1)	GS content (%)	η0.1 (10^4 Pa s)	N0^a (10^8 cells per cm^3)	D^b (μm)
a Cell density.b Average cell diameter.
PVC	0	0	0	1.3	—	418
PVC-g-PS5100	15.7	5100	0	2.7	—	261
PVC-g-PS9300	28.6	9300	0	—	0.63	70
PVC-g-PS18400	56.6	18 400	0	35.0	2.19	36
PVC/FGS	0	0	1	2.2	—	607
PVC-g-PS5100/FGS	15.7	5100	1	11.7	0.52	75
PVC-g-PS9300/FGS	28.6	9300	1	24.8	1.35	45
PVC-g-PS18400/FGS	56.6	18 400	1	86.4	18.67	23

---

Elongational behavior of polymer melts plays an important role in the processing technique, such as foaming. The influence of PS branch chains on elongational flow properties of PVC was investigated. Fig. 2 shows the elongational viscosities for linear and typical branched samples. The elongational viscosity of a polymeric material in the molten state is strongly dependent on its structure, especially branched structure.⁴⁵ In comparison with PVC, PVC-g-PS shows noticeable elongational strain hardening. The faster the stretching rate, the more apparent the strain hardening is, despite the difference between branched samples.

Fig. 2 Elongational viscosities for PVC and PVC-g-PS samples with different length of PS branches at 0.5 s⁻¹ and 0.05 s⁻¹.

The significant improvement in rheological properties of PVC-g-PSx graft copolymers was anticipated to be able to improve their foaming properties. Compared to linear PVC, the change in rheological properties resulted in different foaming behaviors of PVC-g-PSx. For comparison, all the samples were foamed under the same conditions (foaming temperature: 110 °C, CO₂ saturation pressure: 6 MPa). Fig. 3 shows the cell morphologies of the foamed samples. PVC foam presents a large amount of open cell structure, unfoamed regions, and very non-uniform cell size distribution (Fig. 3a), indicating that the cell walls are not strong enough to bear the extensional force during cell growth due to its low melt strength. Similar phenomenon was also observed in PVC-g-PS5100 (Fig. 3b). In contrast, with continuing to increase the length of PS branches, most of cells in the PVC-g-PS foams appear in a closed form and open cells gradually disappeared (Fig. 3c and d). The calculated average cell sizes (D) and cell densities (N₀) of PVC-g-PS foams are collected in Table 1. The enhanced melt strain hardening of PVC-g-PS samples makes the cell walls stronger and prevents the cell coalescence. The PVC-g-PS18400 foam with the longest PS branches among the samples possessed a higher N₀ than the PVC-g-PS9300 foam did, and the cell size distribution of PVC-g-PS foam become narrower with increasing branch length. When the length of PS branch is 18 400 g mol⁻¹, the foam has an average cell size of 36 μm and cell density (N₀) of 2.19 × 10⁸ cells per cm³. This results from the different rheological behaviors. PVC-g-PS sample with longer branches shows higher melt viscosity and more obvious shear-thinning phenomenon. Therefore, under the same foaming conditions, the samples with these rheological properties can protect the foaming sample from forming large cell and even cell coalescence. On the other hand, the presence of PS branches shows nucleating effect on the foaming, resulting in the increased cell density.⁸

Fig. 3 The cell morphologies of (a) PVC, (b) PVC-g-PS5100, (c) PVC-g-PS9300 and (d) PVC-g-PS18400. Foaming temperature: 110 °C; CO₂ saturation pressure: 6.0 MPa.

3.2 Effect of PS branch chains on the dispersion of graphene in PVC-g-PS matrix

Owing to high theoretical properties and flat morphology of graphene, considerable interests in polymer/graphene nanocomposites have emerged to dramatically enhance performance at very low loadings, compared to conventional filler composites. However, the properties of polymer nanocomposites strongly depend on the dispersed state and exfoliated state of graphene platelets with high aspect ratios in polymer matrix. TEM measurements were carried out to compare the morphologies of PVC-g-PS/FGS composites with those of PVC/FGS composites. Fig. 4 shows the TEM images of PVC/FGS and PVC-g-PS/FGS composites. For the linear PVC/FGS mixture, numerous agglomerated particles (“tactoids”) were detected (Fig. 4a), indicating the aggregation of FGS platelets. Owing to the polarity mismatch between the non-polar FGS and polar PVC matrix, FGS was poorly dispersed in the linear PVC matrix. In contrast, the dispersed state of FGS in the PVC-g-PS matrix was totally different (Fig. 4b). The FGS was well dispersed in PVC-g-PS matrix, and uncorrelated nanosheets were clearly visible, confirming the uniformly dispersed without agglomeration. This behavior is mainly caused by the following reasons. When graphene is functionalized by octadecylamine, the long nonpolar octadecyl chains were grafted onto the graphene surface, the resultant FGS becomes hydrophobic. Also the presence of long nonpolar octadecyl chains reduces the interaction between graphene sheets. Furthermore, there are π–π interactions between FGS and PS. These factors resulted in the well dispersion of FGS in the PVC-g-PS matrix.

Fig. 4 TEM micrographs of PVC nanocomposite with 1 wt% FGS. (a) PVC/FGS; (b) PVC-g-PS/FGS.

3.3 Effect of PS branch length on rheological properties and foaming behaviour of PVC-g-PSx/FGS nanocomposites

Fig. 5 shows the plots of complex viscosity (η*) and storage modulus (G′) versus angular frequency (ω), at 185 °C for PVC/FGS and PVC-g-PSx/FGS systems. We approximate η* at the frequency of 0.1 rad s⁻¹, and the values are listed in Table 1. From Fig. 5a and Table 1, the complex viscosity of PVC/FGS composite is the lowest at low frequencies among the samples. When PS was grafted on the backbone of PVC, the complex viscosity of PVC-g-PSx/FGS nanocomposites at low frequency increased gradually with the increasing length of LCB. Furthermore, the presence of LCB led to more prominent shear-thinning phenomenon with the increasing length of LCB. Fig. 5b shows the G′ versus ω for PVC/FGS and PVC-g-PSx/FGS samples at 185 °C. The PVC/FGS showed a lower G′ in the terminal region. Comparatively, the PVC-g-PSx/FGS samples exhibited much higher G′ at low shear frequencies, and the terminal slopes of G′ were all less than that of PVC/FGS. The G′ of PVC-g-PSx/FGS samples at low shear frequencies was enhanced when the branch length increased. The extensional viscosities of the corresponding nanocomposites measured at 180 °C are presented in Fig. 6. The change trend is fairly similar with the samples without graphene (Fig. 2). For PVC/FGS, no strain hardening was observed. However, the samples modified with PS branch chains showed pronounced strain hardening, and this phenomenon was more notable as the stretching rate was increased.

Fig. 5 (a) Complex viscosity vs. angular frequency, and (b) storage modulus vs. angular frequency for PVC-g-PS/FGS samples at 185 °C.

Fig. 6 Elongational viscosities for PVC/FGS and PVC-g-PS/FGS samples with different length of PS branches at 0.5 s⁻¹ and 0.05 s⁻¹.

Fig. 7 shows the cell morphologies of PVC-g-PSx/FGS samples with the content of 1 wt% FGS. Similarly, all the samples were foamed under the same conditions as mentioned above. As shown in Fig. 7a, the introduction of FGS does not improve the foamability of PVC due to the non-uniform dispersion of FGS (Fig. 4a), and serious cell coalescence took place during foam processing. Thus PVC/FGS foam has a very low cell density, and the average cell sizes (D) is about 607 μm. In contrast, owing to the presence of PS branches, an obvious improvement in the foaming behavior of PVC-g-PSx/FGS foams is observed. These results demonstrate that the presence of PS chain promotes the positive effect of added FGS on the foaming behavior of PVC. The PVC-g-PS18400/FGS foam shows the minimum cell size of 23 μm and the maximum cell density with uniform cell size distribution, 1.87 × 10⁹ cells per cm³.

Fig. 7 The cell morphologies of (a) PVC/FGS, (b) PVC-g-PS5100/FGS, (c) PVC-g-PS9300/FGS and (d) PVC-g-PS18400/FGS. Foaming temperature: 110 °C; CO₂ saturation pressure: 6.0 MPa.

3.4 Discussion

Generally, the viscosity and the storage modulus of the composite melts will increase if the added plate-like nanofiller can be uniformly dispersed in polymer matrix. Compared to PVC, PVC/FGS composite shows no significant change in viscosity or storage modulus (Fig. 8a and b). However, there was a clear improvement in the η_0.1 and G′ when the PS branches was introduced on PVC backbone. Combining the above TEM results, we deem that PS branch chains in the PVC-g-PSx make FGS more uniformly dispersed due to the interaction between PS chains and graphene, which is one reason for the drastic variation of the rheological properties after the introduction of PS branch chains. In all the length range of PS branch chains (5100–18 400 g mol⁻¹), PVC-g-PSx/FGS nanocomposites exhibit a positive effect on the foaming behavior compared to PVC-g-PSx, where the resultant PVC-g-PSx/FGS foams possess uniform cell distribution and a gradually decreased size with increasing the length of PS chains. For example, the cell density of PVC-g-PS5100/FGS foam dramatically increases to 10⁷ cells per cm³ compared to PVC-g-PS5100, and an obvious decrease in cell size was observed (Fig. 3b vs. 7b). The changing trends of the cell size and cell density for PVC-g-PSx/FGS foams are the same as the samples without FGS (Fig. 3 vs. 7). The PVC-g-PS18400/FGS foam has the minimum cell size of 23 μm and the maximum cell density (1.87 × 10⁹ cells per cm³) with uniform cell size distribution (Fig. 7d).

Fig. 8 The comparison of (a) complex viscosity vs. angular frequency and (b) storage modulus vs. angular frequency for the samples with and without FGS.

Moreover, the foaming behavior is consistent with the rheological behavior, i.e., the higher the melt viscosity, the better the foaming behavior. These results demonstrate that the introduction of PS chain promote the positive effect of FGS on the foaming behavior of PVC. Possible reason for this phenomenon was attributed to the increased dispersion degree of FGS in the PVC matrix, which increased the viscosity, physical barrier effect of graphene to cell coalescence and the effective nucleating sites to induce cell nucleation. Therefore, PS long branch chains play the double functions in melt and foaming behaviors of PVC-g-PS/FGS nanocomposites.

4 Conclusion

The introduction of PS branch chains with different length onto PVC backbones is an efficient method to change rheological properties of PVC and dispersed state of graphene in PVC matrix. Compared to PVC, the presence of PS branch chains in the PVC-g-PS leads to increased η_0.1, elevated value of G′ at low shear frequencies, more pronounced shear thinning and strain hardening behavior. As a result, the foaming behavior is dramatically improved. The longer the length of PS branch, the more apparent this effect is. Furthermore the presence of PS branch chains promotes the dispersion degree of FGS in PVC-g-PS matrix due to the interaction between PS chains and graphene. Compared to PVC-g-PS, the shear rheological properties of PVC-g-PS/FGS nanocomposites change more significantly with increasing the branch length. Comparing PVC/FGS composite with PVC, PVC-g-PS/FGS nanocomposites exhibit better foaming behavior than PVC-g-PS, where the resultant foams possess uniform cell distribution and a decreased size with increasing the length of PS branch chains.

Acknowledgements

This work is financially supported by the National Natural Science Foundation of China for the Projects (51233005 and 51073149), the Ministry of Science and Technology of China (SQ2014AAJY1027) and the Science and Technology Development Plan of Jilin Province, China (20140203023GX).

Notes and references

1. P. Spitael and C. W. Macosko, Polym. Eng. Sci., 2004, 44, 2090–2100.
2. L. Wang, D. Wan, Z. Zhang, F. Liu, H. Xing, Y. Wang and T. Tang, Macromolecules, 2011, 44, 4167–4179.
3. L. Wang, Z. Jiang, F. Liu, Z. Zhang and T. Tang, Chin. J. Polym. Sci., 2014, 32, 333–349.
4. L. Wang, H. Yang, H. Tan, K. Yao, J. Gong, D. Wan, J. Qiu and T. Tang, Polymer, 2013, 54, 3641–3653.
5. H. E. Naguib, C. B. Park, U. Panzer and N. Reichelt, Polym. Eng. Sci., 2002, 42, 1481–1492.
6. F. Liu, D. Wan and T. Tang, Chin. J. Polym. Sci., 2014, 32, 51–63.
7. J. Stange and H. Munstedt, J. Cell. Plast., 2006, 42, 445–467.
8. L. Wang, D. Wan, J. Qiu and T. Tang, Polymer, 2012, 53, 4737–4757.
9. W. Zhai, H. Wang, J. Yu, J. Y. Dong and J. He, Polymer, 2008, 49, 3146–3156.
10. M. Gahleitner, Prog. Polym. Sci., 2001, 26, 895–944.
11. K. E. Mabrouk, J. S. Parent, B. I. Chaudhary and R. Cong, Polymer, 2009, 50, 5390–5397.
12. S. Li, M. Xiao, Y. Guan, D. Wei, H. Xiao and A. Zheng, Eur. Polym. J., 2012, 48, 362–371.
13. T. C. B. McLeish, Europhys. Lett., 1988, 6, 511.
14. G. G. Vogiatzis and D. N. Theodorou, Macromolecules, 2013, 47, 387–404.
15. A. Tuteja, P. M. Duxbury and M. E. Mackay, Macromolecules, 2007, 40, 9427–9434.
16. M. E. Mackay, T. T. Dao, A. Tuteja, D. L. Ho, B. Van Horn, H. C. Kim and C. J. Hawker, Nat. Mater., 2003, 2, 762–766.
17. S. B. Kharchenko, J. F. Douglas, J. Obrzut, E. A. Grulke and K. B. Migler, Nat. Mater., 2004, 3, 564–568.
18. M. J. Solomon, A. S. Almusallam, K. F. Seefeldt, A. Somwangthanaroj and P. Varadan, Macromolecules, 2001, 34, 1864–1872.
19. H. B. Zhang, W. G. Zheng, Q. Yan, Z. G. Jiang and Z. Z. Yu, Carbon, 2012, 50, 5117–5125.
20. A. Tuteja, M. E. Mackay, S. Narayanan, S. Asokan and M. S. Wong, Nano Lett., 2007, 7, 1276–1281.
21. F. H. Su, J. H. Yan and H. X. Huang, J. Appl. Polym. Sci., 2011, 119, 1230–1238.
22. H. Tan, D. Xu, D. Wan, Y. Wang, L. Wang, J. Zheng, F. Liu, L. Ma and T. Tang, Soft Matter, 2013, 9, 6282–6290.
23. W. Luo, Y. Li, H. Zou and M. Liang, RSC Adv., 2014, 4, 37302–37310.
24. B. Zhu, W. Zha, J. Yang, C. Zhang and L. J. Lee, Polymer, 2010, 51, 2177–2184.
25. J. Shen, C. Zeng and L. J. Lee, Polymer, 2005, 46, 5218–5224.
26. L. Sorrentino, M. Aurilia, L. Cafiero and S. Iannace, J. Appl. Polym. Sci., 2011, 122, 3701–3710.
27. H. Ma, H. Zhang, X. Li, X. Zhi, Y. Liao and Z. Yu, Ind. Eng. Chem. Res., 2014, 53, 4697–4703.
28. J. Ling, W. Zhai, W. Feng, B. Shen, J. Zhang and W. Zheng, ACS Appl. Mater. Interfaces, 2013, 5, 2677–2684.
29. H. Han, S. li and X. J. X. Kong, RSC Adv., 2014, 4, 33520–33529.
30. H. B. Zhang, Q. Yan, W. G. Zheng, Z. He and Z. Z. Yu, ACS Appl. Mater. Interfaces, 2011, 3, 918–924.
31. M. Sauceau, J. Fages, A. Common, C. Nikitine and E. Rodier, Prog. Polym. Sci., 2011, 36, 749–766.
32. H. W. Engels, H. G. Pirkl, R. Albers, R. W. Albach, J. Krause, A. Hoffmann, H. Casselmann and J. Dormish, Angew. Chem., Int. Ed., 2013, 52, 9422–9441.
33. W. Zhai, J. Yu, L. Wu, W. Ma and J. He, Polymer, 2006, 47, 7580–7589.
34. L. Chen, D. Rende, L. S. Schadler and R. Ozisik, J. Mater. Chem. A, 2013, 1, 3837–3850.
35. Y. Saeki and T. Emura, Prog. Polym. Sci., 2002, 27, 2055–2131.
36. K. Yao, J. Gong, J. Zheng, L. Wang, H. Tan, G. Zhang, Y. Lin, H. Na, X. Chen, X. Wen and T. Tang, J. Phys. Chem. C, 2013, 117, 17016–17023.
37. Y. Piette, A. Debuigne, C. Jerome, V. Bodart, R. Poli and C. Detrembleur, Polym. Chem., 2012, 3, 2880–2891.
38. A. Carretero-Genevrier, C. Boissiere, L. Nicole and D. Grosso, J. Am. Chem. Soc., 2012, 134, 10761–10764.
39. Y. Zhang, Y. Song and Q. Zheng, Chin. J. Polym. Sci., 2013, 31, 325–332.
40. D. C. Marcano, D. V. Kosynkin, J. M. Berlin, A. Sinitskii, Z. Sun, A. Slesarev, L. B. Alemany, W. Lu and J. M. Tour, ACS Nano, 2010, 4, 4806–4814.
41. G. Wang, X. Shen, B. Wang, J. Yao and J. Park, Carbon, 2009, 47, 1359–1364.
42. V. Percec and F. Asgarzadeh, J. Polym. Sci., Part A: Polym. Chem., 2001, 39, 1120–1135.
43. S. L. Shenoy, W. D. Bates, H. L. Frisch and G. E. Wnek, Polymer, 2005, 46, 3372–3384.
44. J. N. Yoo, L. H. Sperling, C. J. Glinka and A. Klein, Macromolecules, 1991, 24, 2868–2876.
45. Z. Zhang, D. Wan, H. Xing, Z. Zhang, H. Tan, L. Wang, J. Zheng, Y. An and T. Tang, Polymer, 2012, 53, 121–129.

---

Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c4ra09236h

---