Effect of graphene oxides on thermal degradation and crystallization behavior of poly(L-lactide)

Abstract

Graphene oxides (GO) were introduced into a poly(L-lactide) (PLLA) to prepare the PLLA/GO composites with different concentrations of the GOs. The main attention of the present work was focused on the thermal degradation and crystallization behaviors of a PLLA matrix induced by GOs. The measurements based on gel permeation chromatography (GPC), thermogravimetric analysis (TGA) and Fourier transform infrared spectroscopy (FTIR) clearly proved the thermal degradation of the PLLA matrix during the melt-processing procedure. Consequently, reduced complex viscosity of the composites was achieved. The crystallization behaviors of the PLLA matrix were comparatively investigated under different crystallization conditions including melt-crystallization (nonisothermal or isothermal crystallization occurring from the melt state) and cold-crystallization (crystallization occurring from an amorphous solid state) using differential scanning calorimetry (DSC), polarized optical microscopy (POM) and wide angle X-ray diffraction (WAXD). Largely enhanced crystallization ability of the PLLA matrix was achieved. The crystallization conditions and the GO content are key factors which influence the crystallization behavior of the PLLA matrix. This work proved that the stabilization of the PLLA matrix during the melt-processing procedure must be considered when designing and preparing the PLLA/GO materials.

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1. Introduction

As one of the biodegradable and biocompatible polymers, poly(L-lactide) (PLLA) has been one of the main subjects of many research reports for a long time due to its excellent physical properties, such as high tensile strength and tensile modulus, transparency and good processability, et al. Importantly, it can be obtained from natural resources, such as corn, sugarcane and sweet potato. The application of the PLLA greatly reduces the dependence on polymers which are obtained from petroleum resources. However, the application of the PLLA is also greatly limited by its poor crystallization ability. The common melt-extrusion, injection and compression procedures provide only the amorphous PLLA articles. This is because that PLLA material has little time to crystallize due to the relatively high cooling rate during the solidification process. Therefore, improving the crystallization ability of the PLLA is very significant whether from a viewpoint of reducing cost or from a viewpoint of maintaining mechanical performance.

The crystallization ability of the PLLA can be improved by adding a few amount of a plasticizer, which enhances the mobility of the PLLA molecular chains^1–4 and decreases the glass transition temperature (T_g). However, relatively high concentration of the plasticizer (usually more than 10 wt%) must be required.^5,6 Although the ductility of the plasticized PLLA material is improved, both the tensile strength and modulus are decreased dramatically. Another way to improve the crystallization ability of the PLLA is adding nucleating agent^7–12 or nanofiller which exhibits nucleation effect for semicrystalline polymers.^13–20 With the presence of the nucleating agent or the nanofiller, the nucleation activation energy is reduced, which promotes the crystallization of the PLLA at relatively higher melt temperatures. Although the nucleation density is increased greatly, the growth rate of crystallites is hardly enhanced. Contrarily, decreased growth rate of crystallites has been reported at high concentration of nanofiller. It is assumed that the decrease of the crystallization rate is mainly related to the reduced mobility of molecular chains resulted by the restriction of the nanofiller network structure.²¹ With respect to the role of plasticizer which enhances the mobility of chain segments and the role of nanofiller which increases the nucleation density, simultaneous addition of plasticizer and nanofiller into the PLLA is possibly an efficient way to improve the crystallization ability of the PLLA. This has been realized successfully in our previous work about the synergistic effects of polyethylene glycol (PEG) and carbon nanotubes (CNTs) on the PLLA crystallization have been confirmed.²² The crystallization rate of the PLLA was greatly enhanced in all conditions and bigger degree of crystallinity (X_c) was obtained. However, it should be stressed that the content of the plasticizer was 10 wt%, which inevitably results in the deterioration of tensile strength of the material even if CNTs are present in the matrix.²³ Therefore, improving the crystallization ability of the PLLA as much as possible and meanwhile maintaining the other physical properties at relatively high level is still an interesting challenge.

As one of the allotropes of carbon material, graphene oxides (GO) have similar chemical structure to the modified CNTs on one hand. On the other hand, there are many functional groups including carboxyl and hydroxyl groups on the surface of the GOs, and these polar groups facilitate the homogeneous dispersion of the GOs in the polar polymer matrix. Due to their large aspect ratio and the possible strong interfacial interaction between polymer matrix and GOs, GOs exhibit good nucleation effect during the crystallization of semicrystalline polymers. For example, He C. B. et al.²⁶ found that the incorporation of the GOs leaded to lower crystallization activation energy of stereocomplex and higher crystallinity in solution casting samples. They also found that the heterogeneous nucleation effect of the GOs was dependent on the crystallization conditions. Pinto A. M. et al.²⁴ also found that the incorporation of the 0.4 wt% GOs induced the cold crystallization of the solution casting PLLA/GO sample during the DSC heating process, while no crystallization was detected for the pure PLLA sample. Importantly, they reported that the presence of the GOs greatly enhanced the T_g of the PLLA. They also suggested that the good filler–matrix interaction was the main reason for the restricted chain segments mobility. However, Xu J.-Z. et al.²⁵ comparatively investigated the effects of the CNTs and the GOs on crystallization of the PLLA and they found that although CNTs and GOs exhibited heterogeneous nucleation effects for the PLLA, the induction ability of the GOs was weaker than that of the CNTs. With increasing of the CNT content, the induction period was shortened and the crystallization rate was increased for the PLLA/CNT composites, while the reverse situation was found for the PLLA/GO composites.

Although some work has been carried out to investigate the crystallization behavior of the PLLA/GO composites, the related reports are still few and much work needs to be done to understand the crystallization behavior of the PLLA in detail. Furthermore, the effects of the GOs on the melt-processing, the rheological properties and the crystallization behaviors of the PLLA/GO composites are still unclear. Furthermore, although the functional groups on the surface of the GOs possibly improve the interfacial interaction between the PLLA and the GOs, they also increase the possibility of the thermal degradation of the PLLA matrix under the combined effects of high temperature and shear stress during the melt-processing procedure. Therefore, in this work, we attempted to prepare the PLLA/GO composites via the common melt-processing procedure. The effects of the GOs on the thermal stability, the rheological properties and the crystallization behaviors of the PLLA matrix were investigated in detail.

2. Experimental part

2.1. Materials

PLLA (2002D, with D-isomer content of 4.3%, M_w of 2.53 × 10⁵ g mol⁻¹, melt flow rate (MFR) of 4–8 g/10 min (190 °C/2.16 kg), and density of 1.24 g cm⁻³) was purchased from NatureWorks®, USA. Graphite was obtained from Qingdao Heilong Graphite Co., Ltd.

2.2. Sample preparation

GOs were prepared according to Hummer's method²⁷ and the corresponding data of the GOs can be seen in our previous work.²⁸ After the surface modification, some functional groups, including carboxyl and hydroxyl groups, were introduced to the surface of the GOs. The weight fraction of functional groups of the GOs was about 34.59% which was measured by thermogravimetric analysis (TGA).

Samples were prepared through the two-step method: first, GOs and PLLA were dissolved in the DMF to prepare the solution of the GO/DMF and the PLLA/DMF, respectively; then the solution of the GO/DMF and the PLLA/DMF were compounded together with stirring at 80 °C for 2 h; and then the mixture was dried at 100 °C for enough time until the solvent was removed completely. Finally, the master batch of the PLLA/GO composite containing 5.0 wt% GOs was prepared. After that, the master batch was melt-compounded with the pure PLLA to prepare the composites. The melt-processing of all the composites was carried out in an internal mixer at a melt temperature of 185 °C and a rotate speed of 60 rpm. The mixing duration was 5 min. In this work, the weight fraction of the GOs in the composites was varied from 0.1 to 2.0 wt%. The sample notation of composite was defined as the PLLA/xGO, where x indicated the weight fraction of the GOs. For example, the PLLA/0.1GO represented that the concentration of the GOs in the composite was 0.1 wt%.

To prepare the sample for the measurement, the composite was first heated to melt completely. Some composites were quenched in the ice-water to obtain the amorphous samples and others were transferred into a hot-stage with the setting temperature of 115 °C and maintained at this temperature for 20 min to obtain the isothermal-crystallized samples. In addition, some amorphous samples were further treated at 115 °C for 20 min to prepare the annealed samples. To make a comparison, pure PLLA sample was prepared through the completely same procedures.

2.3. Scanning electron microscopy (SEM)

A scanning electron microscope (SEM, Fei Inspect FEI, the Netherlands) was used to investigate the dispersion of the GOs in the composites. The sample was first cryogenically fractured in liquid nitrogen, and then the fractured surface was coated with a thin layer of gold before the SEM characterization. An accelerating voltage of 5.0 kV was selected for the SEM characterization.

2.4. Gel permeation chromatography (GPC)

The number-average molecular weight (M̄_n), weight-average molecular weight (M̄_w) and the polydispersity d (d = M̄_w/M̄_n) of the representative composites obtained through melt-processing procedure were measured by a gel permeation chromatography (GPC) 1100 series (Aglient, USA) using a polystyrene as a standard. The GPC column type is super HM-H (6.0 mm × 15 mm × 2) and the flow rate is 6 ml min⁻¹. Before measurement, the PLLA/GO composite was first extracted by dichloromethane (CH₂Cl₂) and the GOs was removed, then the PLLA/CH₂Cl₂ solution was dried and the CH₂Cl₂ was removed. Subsequently, the PLLA sample was dissolved in chloroform at 40 °C. It should be stressed that the above sample preparation procedure for GPS measurement does not induce the variations of the molecular weight and molecular weight distribution because the solution temperature is relatively low and no chemical reaction occurs.

2.5. Thermogravimetric analysis (TGA)

The thermal stability of the GOs and the pure PLLA was comparatively investigated through a thermogravimetric analysis (TGA) measurement, which was carried out using the Q500 (TA Instrument, USA) in nitrogen atmosphere. Sample was heated from 30 °C to 450 °C at the heating rate of 40 °C min⁻¹ in the nitrogen atmosphere.

2.6. Fourier transform infrared spectroscopy (FTIR)

The chain scission of the pure PLLA and the PLLA/GO composites were characterized using a Fourier transform infrared (FTIR) AVATAR360 (Thermo Nicolet, USA) at an attenuated total reflection (ATR) mode. The resolution was 2 cm⁻¹ and the wavenumber range was from 600 to 4000 cm⁻¹. All the measurements were carried out at room temperature (23 °C).

2.7. Rheological measurement

The rheological measurement was carried out on a stress controlled rheometer AR2000ex (TA Instrument, USA). The sample disk was prepared through a compression molding way at the melt temperature of 185 °C and a pressure of 5 MPa. The thickness and diameter of the sample disk are 1 and 20 mm, respectively. During the rheological measurement process, the frequency sweep was performed from 0.05 to 300 rad per s. The melt temperature was set at 180 °C and the measurement was carried out in nitrogen atmosphere. For all the measurements, the samples were tested within the linear viscoelastic strain range.

2.8. Differential scanning calorimetry (DSC)

The crystallization and melting behaviors of samples were investigated using a differential scanning calorimetry (DSC) STA 449C Jupiter (Netzsch, Germany). For the quenched sample, it was first heated from 0 °C to 200 °C at a heating rate of 10 °C min⁻¹ and maintained at 200 °C for 3 min to erase any thermal history, then the sample was cooled to 0 °C at different cooling rates (2 and 5 °C min⁻¹), and then the sample was heated again to 200 °C at the heating rate of 10 °C min⁻¹. For the isothermally crystallized and/or the annealed samples, they were directly heated from 0 °C to 200 °C at the heating rate of 10 °C min⁻¹. All the DSC measurements were carried out in nitrogen atmosphere. The weight of each sample was about 8 mg. The degree of crystallinity (X_c-DSC) of the PLLA was calculated according to the following equation:where ΔH_m is the DSC measured value of fusion enthalpy, ΔH_cc is the cold crystallization enthalpy obtained during the DSC heating process, ΔH⁰_m is the fusion enthalpy of the completely crystalline PLLA, and ϕ is the weight fraction of the PLLA in the sample. Here, the value of ΔH⁰_m of the PLLA was 93 J g⁻¹.²⁹

2.9. Polarized optical microscopy (POM)

The crystalline morphology of sample was characterized using a polarized optical microscope (POM) XPN203E (Changfang Instruments, China) with a hot-stage. Firstly, a sample of about 5 mg was heated to 190 °C and pressed to obtain a slice with a thickness of about 20 μm. The slice was then maintained at 190 °C for 3 min to erase the thermal and mechanical history. Secondly, the slice was transferred to the hot-stage with a setting temperature of 115 °C and maintained at this temperature for 20 min. The crystalline morphology of the sample was taken images using a digital camera.

2.10. Wide angle X-ray diffraction (WAXD)

The crystalline structure of the PLLA matrix was investigated using a wide angle X-ray diffraction (WAXD) DX-1000 (China). The continuous scanning angle range used in this study was from 5° to 30° at 40 kV and 25 mA. The X_c-WAXD was calculated according to the following relation:where A_crystalline and A_amorphous are the fitted areas of the diffraction peaks of crystalline and amorphous, respectively. In the present work, the diffractions of (010), (110)/(200), (203) and (015) crystal planes were considered, and the sum of the fitted areas of these diffraction peaks represents the crystalline part, while the area below the diffraction peaks represents the amorphous part.

3. Results and discussion

3.1. Dispersion of the GOs and the PLLA/GO interaction

The dispersion state of the GOs in the PLLA matrix was first investigated by SEM and the results are shown in Fig. 1. There are at least two characteristics which can be deduced from the SEM images. First, although there is no serious GO agglomerates which consists of many random GO particles, the platelets of the GOs with multiple layers can be clearly seen. The width and the thickness of the GO sheets, which are measured directly from the SEM image, are in the range of 5–10 μm and 30–50 nm, respectively. The size of the GO sheets is larger than that of the GOs in the solution casting samples reported in the literature.³⁰ This implies that the dispersion of the GOs should be further improved. It is well known that the melt-processing using twin-screw extruder provides more intense shear stress to the melt, which facilitates the good dispersion of nanofiller. Therefore, it is expected that if the twin-screw extruder is used during the melt-processing procedure, the dispersion of the GOs can be further improved. The other strategy is improving the modification method of graphite so that GOs with fewer layers could be obtained. This facilitates the exfoliation and homogeneous dispersion of the GOs in the polymer matrix with the aid of shear stress. Second, large holes/cavities between GOs and PLLA can be seen in all the composites, indicating almost completely dewetted GO-PLLA interface. Furthermore, it can be seen that the more the GOs in the composites, the more the holes/cavities are observed. It is reported in the literature that GOs exhibit good affinity to PLLA due to the large amount of polar groups on the surface of the GOs, resulting in the good wetting of the GOs by the PLLA.^24,25 However, our results prove that the interfacial interaction between GOs and PLLA is weak. Probably, some functional groups degrade from the surface of the GOs and migrate into the matrix during the sample preparation process. When the sample is cryogenically fractured, GOs and PLLA matrix exhibit different abilities for deformation resistance and the stress transfer between GOs and PLLA becomes more difficult. This results in the initiation and propagation of crack at the interface, leading to the formation of large holes/cavities between GOs and PLLA. Therefore, it is speculated that the degradation of the functional groups of the GOs possibly leads to the thermal degradation of the adjacent PLLA matrix. This will be proved in the following section.

Fig. 1 SEM images show the dispersion of the GOs in the different PLLA composites. (a) PLLA/0.1GO, (b) PLLA/0.5GO, (c) PLLA/1GO and (d) PLLA/2GO.

3.2. Thermal degradation of PLLA during the melt processing

The molecular weight of the pure PLLA and the PLLA/GO composites after being melt-compounded were first measured using GPC. The results are shown in Table 1. It can be seen that addition of a few GOs induces the apparent change of the M̄_n. For example, for the PLLA/0.1GO sample, the M̄_n is decreased from 6.90 × 10⁴ g mol⁻¹ of the pure PLLA to 4.25 × 10⁴ g mol⁻¹. The PLLA/0.5GO sample exhibits the similar M̄_n to that of the PLLA/0.1GO sample, but the PLLA/2GO sample exhibits largely decreased M̄_n. This indicates that GOs at high concentration promote the thermal degradation of the PLLA matrix during the melt-processing procedure. This can be further proved by the dramatically decreased M̄_w in the PLLA/2GO composite. The M̄_w is decreased from 1.32 × 10⁵ g mol⁻¹ of the pure PLLA to 0.38 × 10⁵ g mol⁻¹ of the PLLA/2GO. Consequently, the PLLA/GO composites exhibit increased d as shown in Table 1.

Table 1 Number-average molecular weight (M̄_n), weight-average molecular weight (M̄_w) and the polydispersity (d) of molecular weight

Samples	M̄n × 10^−4 (g mol^−1)	M̄w × 10^−5 (g mol^−1)	d
PLLA	6.90	1.32	1.91
PLLA/0.1GO	4.25	1.37	3.23
PLLA/0.5GO	4.40	1.37	3.13
PLLA/2GO	1.09	0.38	3.48

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The thermal degradation of the PLLA matrix is possibly related to the thermal degradation of the GOs at the temperature range of 180–200 °C as shown in our previous work.²⁸ As shown in Fig. 2, which exhibits the TGA curves of the GOs and the pure PLLA. For the GOs, the weight loss at temperature lower than 100 °C is mainly related to the evaporation of absorbed water. However, at temperature higher than 150 °C, the weight loss of the GOs becomes more apparent. This can be attributed to the degradation of functional groups of the GOs. For the pure PLLA, there is no apparent weight loss at temperature lower than 280 °C. Therefore, it can be deduced that the thermal degradation of the PLLA matrix in the PLLA/GO composites is mainly related to the presence of the GOs. In this work, the weight fraction of the functional groups on the surface of the GOs is up to 34.59%. During the melt-processing procedure, the thermal degradation of the GOs is aggravated by the combined effects of high melt temperature and shear stress. Consequently, the thermal degradation of functional groups of the GOs initiates the thermal degradation of the PLLA matrix. The chain scission of the PLLA matrix can be proved further by Fig. 3, which shows the FTIR spectra of the pure PLLA and the PLLA/0.1GO and PLLA/1GO samples.

Fig. 2 Comparison of TGA curves of GOs and pure PLLA obtained at the heating rate of 40 °C min⁻¹.

Fig. 3 FTIR spectra of pure PLLA and PLLA/GO composites as indicated in the graphs.

For the pure PLLA, the characteristic absorption peaks at 1753, 1187, and 1086 cm⁻¹ are mainly attributed to the stretching vibration of >C=O groups, in-plane rocking vibration of –CH₃ groups, and symmetric stretching vibration of –COC– groups,³¹ respectively. Compared with the pure PLLA, the intensities of the three absorption peaks become smaller. Furthermore, largely deteriorated absorption bands are observed for the PLLA/1GO sample. This clearly indicates the reduction of >C=O, –CH₃ and –COC groups in the material, indicating the thermal degradation of the PLLA during the melt-compounding processing undoubtedly. Interestingly, besides the deteriorated absorption bands of >C=O, –CH₃ and –COC groups, one can observe the presence of some new absorption bands at 2919, 2850 cm⁻¹, 1640 and 1586 cm⁻¹ of the PLLA/GO composites. The former two absorption bands are attributed to the asymmetric stretching vibration and symmetric stretching vibration of –CH₂ groups while the latter two absorption bands are attributed to the symmetric stretching and asymmetric stretching of C=C groups. Specifically, the intensities of these absorption bands increase with increasing the content of GOs in the composites. It is well known that there are not –CH₂ and C=C groups in the molecular chains of the pure PLLA. Therefore, it can be deduced at least that both –CH₂ and C=C groups are introduced into the PLLA/GO material and the amount of these groups increases with increasing the content of the GOs.

With the aid of nuclear magnetic resonance (NMR) and mass spectrum (MS) for product identification, McNeill I. C. et al.³² proposed that the thermal degradation of the PLLA at lower temperature was a non-radical, backbiting ester interchange reaction involving the –OH chain ends, which resulted in the formation of lactide molecule or oligomeric ring with more than two repeat units. However, at relatively higher temperature, other two primary chain scission steps were envisaged, involving either alkyl-oxygen homolysis or carbonyl-oxygen homolysis, which resulted in the formation of carbon dioxide, lactide molecule or ring oligomer, and the primary product was methylketene (CH₃CH=C=O). However, Kopinke F.-D. et al.³³ proposed that the thermal degradation of PLLA was very complicated at temperature above 200 °C, involving at least five reaction pathways. Among these reaction pathways, the cis-elimination was believed one of the dominant pathway which resulted in the formation of acrylic acid and acyclic oligomers (CH₂=CHCOO(CH₃CHCOO)_mH). However, they did not provide the corresponding FTIR of the degraded product. Interestingly, in our work, the presence of CH₂ and C=C groups has been proved by FTIR spectra. Therefore, it is suggested that the thermal degradation of functional groups of the GOs at relatively low temperature induces the formation of macroradicals ((CH₃)ĊHCOOCH(CH₃)COOR), which come from the chain scission through alkyl-oxygen/carbonyl-oxygen homolysis of PLLA according to ref. 32. Then those macroradicals are enolized into (ROC(OH)=C(CH₃)OOR). Finally, the main production (CH₂=CHCOO(CH₃CHCOO)_mH) is obtained through the cis-elimination reaction according to ref. 33. That is also the reason why the FTIR spectra show the stronger absorption of C=C and CH₂ groups while the intensity of >C=O and –COC– groups is largely deteriorated.

3.3. Rheological properties

Studying on the rheological properties of polymeric composites in the molten state is crucial to obtain a fundamental understanding of the nature of the processability and the structure–property relationship of the composites. Importantly, rheological measurement also provides information about the microstructure of the nanofiller in the polymer melt.^34–36 Fig. 4 shows the rheological properties of the PLLA/GO composites. For making a comparison, the results of the pure PLLA are also shown. It is interesting to observe that with the incorporation of a few amount of the GOs (0.1 and 0.5 wt%), composites exhibit smaller storage modulus (G′, Pa) compared with the pure PLLA. Further increasing the content of the GOs induces the increase of G′. Furthermore, a plateau is present in the modulus curve at low frequency and it becomes more apparent at higher concentration of the GOs (2.0 wt%). However, at relatively higher frequency, all the composites exhibit smaller G′ compared with the pure PLLA. The similar phenomena are observed for the loss modulus (G′′, Pa). The analogous Cole–Cole plots of G′ versus G′′ are shown in Fig. 4c. It has been reported that both the shift and the change in the slope of the Cole–Cole plots indicate the significant microstructure changes of the composites with the presence of the fillers.^37–39 It has been suggested that both the platform at low frequency region in modulus curve and the deviation from a linear relationship between G′ and G′′ are indicative of the formation of a percolated rheological network structure of the filler.^40–42 Therefore, from Fig. 4a–c one can believe that the percolated rheological network structure of the GOs is also formed in the PLLA/GO composites when the GOs (1.0 and 2.0 wt%) are present in the composites. Fig. 4d shows the frequency dependence of loss tangent (tan δ, a ratio of G′′/G′) of the pure PLLA and the PLLA/GO composites. There are at least two features which can be concluded. First, it can be seen that with the incorporation of the GOs, the tan δ is dramatically decreased at relatively lower frequency. The higher the content of the GOs, the smaller the tan δ is. Second, at low concentration of the GOs, the tan δ increases with frequency until a moderate maximum is reached. Further increasing frequency results in the decrease of the tan δ. For the PLLA/2GO composite, one can also see that tan δ is independent upon the frequency at relatively lower frequency. Ko S. W.³⁴ suggested that the value of the tan δ could be used to reflect the dispersion of the nanofiller in the matrix. If tan δ is below 1, nanofillers have strongly associated relation. For example, nanofillers form the so called percolated rheological network structure in the matrix. In this work, the PLLA/2GO composite shows tan δ smaller than 1, which further indicates the formation of the percolated GO network structure.

Fig. 4 Rheological properties of different samples as indicated in the graphs. (a) storage modulus, (b) loss modulus, (c) Cole–Cole plots of storage modulus versus loss modulus and (d) tan δ as a function of the frequency in the different PLLA/GO composites. (e) Variation of complex viscosity of the pure PLLA and the PLLA/GO composites.

Fig. 4e shows the variation of the complex viscosity (η*, Pa s) of the pure PLLA and the composites. Pure PLLA exhibits the Newtonian fluid behavior with the invariant of the η* in the wide range of the frequency. The η* of material is greatly decreased with the presence of only 0.1 wt% GOs, and it further decreases with increasing GO content. For the PLLA/1GO composite, although it exhibits bigger η* compared with the PLLA/0.5GO composite, the value of the η* is still smaller than that of the pure PLLA. Even for the PLLA/2GO composite, it still exhibits smaller η* compared with the pure PLLA sample at relatively higher frequency. Furthermore, one can see that at relatively lower GO concentration, the composites exhibit the typical Newtonian fluid behavior. With increasing GOs content, one can observe a change from Newtonian fluid behavior to pseudoplastic fluid behavior, i.e. the η* of composites decreases with increasing of frequency. The more the GOs in the composites, the more apparent the pseudoplastic fluid behavior is.

The effect of the GOs on the rheological properties of polymer has been less researched. Zhang H. B.⁴³ investigated the effect of graphene sheets containing different oxygen content on the rheological properties of polymethylmethacrylate (PMMA) and found that almost all the composites exhibited higher modulus and viscosity compared with the pure PMMA. Filippone G. et al.⁴⁴ introduced the GOs into the polystyrene (PS) and found that high concentration of the GOs induced the apparent increase of the G′ at low frequency. However, they also showed that when the concentration of the GOs was only 1.9 wt%, the composite exhibited similar G′ and G′′ to those of the pure PS. However, Li Y. F. et al.⁴⁵ investigated the rheological properties of the polypropylene (PP) containing different contents of graphene nanoplatelet and found that even if the content of graphene nanoplatelet was increased up to 10 wt%, the composites exhibited much smaller G′, G′′ and η* compared with the pure PP. They believed that the decrease of the composites viscosity was resulted by the low surface friction and agglomeration of graphene nanoplatelet which promoted the occurrence of the interlayer slipperiness. However, they did not provide the variations of the molecular weight of the PP matrix. In this work, dramatically decreased molecular weight and increased polydispersity of PLLA matrix induced by the presence of the GOs are observed due to the aggravated thermal degradation of the PLLA during the melt-processing procedure. Therefore, it is believed that the rheological properties of the PLLA/GO composites are mainly related to the molecular weight of the PLLA matrix and the content of the GOs in the matrix. At relatively low GO content (0.1 and 0.5 wt%), the rheological properties of the composites are mainly determined by the PLLA matrix rather than by the dispersed GO particles. In this condition, the decrease of the molecular weight and the increase of the polydispersity of the matrix induce the decrease of the G′, G′′ and η*. Consequently, the reinforcement effect of the dispersed GO particles on the G′ becomes inconspicuous. However, at relatively high GO content (1 and 2 wt%), although the molecular weight of the PLLA matrix is further decreased, GOs form the rheological network structure in the matrix, which restricts the chain mobility of the PLLA matrix and induces the enhancement of the G′ at relatively low frequency. At relatively high frequency, the shear stress during the rheological measurement possibly destroys the rheological network structure of the GOs, and in this condition the rheological properties of the material is mainly determined by the PLLA matrix.

3.4. Crystallization behaviors

Melt crystallization. The melt crystallization behaviors of the pure PLLA and the PLLA/GO composites were comparatively investigated through nonisothermal and isothermal crystallization processes. Here, melt crystallization represents that the crystallization of the PLLA matrix occurs from the molten state and the molecular chains of the PLLA have enough mobility. This is completely different from the so called cold crystallization, in which the crystallization of the PLLA matrix occurs from the solid state and the motion of the PLLA molecular chains is greatly restricted although the temperature is higher than the T_g of the PLLA. In this work, the nonisothermal crystallization was carried out via two cooling rates, i.e. 5 and 2 °C min⁻¹. Fig. 5 shows the corresponding cooling curves and the subsequent heating curves of the different samples as indicated in the graphs. At the cooling rate of 5 °C min⁻¹ (Fig. 5a), pure PLLA exhibits only a linear curve without any exothermic phenomenon, indicating that no crystallization occurs during the cooling process. With the incorporation of the GOs in the matrix, the composites exhibit exothermic phenomena, indicating the occurrence of the melt crystallization. The intensity of the exothermic peak becomes stronger and the exothermic peak ascribing to the melt crystallization temperature of the PLLA (T_c) shifts to higher temperatures with increasing GO content. For example, the T_c is increased from 100.2 °C of the PLLA/0.1GO to 109.8 °C of the PLLA/2GO. The enhanced crystallization ability of the PLLA matrix is possibly related to the two factors, one factor is related to the decrease of the molecular weight of the PLLA matrix in the composites, which promotes the nucleation and growth of the PLLA crystallites, and the other factor is related to the nucleation effect of the GOs, which reduces the nucleation activation and promotes the crystallite growth on the surface of the GOs. Pan P. J.⁴⁶ found that the molecular weight greatly influenced the melt crystallization of the PLLA. With decreasing molecular weight, the T_c shifted to higher temperatures and the exothermic peak became much sharper. But it should be pointed out that when the M̄_n was decreased from 1.18 × 10⁵ g mol⁻¹ to 1.54 × 10⁴ g mol⁻¹, the T_c was increased only about 8 °C. Similarly, Pillin I. found⁴⁷ that when the M̄_n was decreased from 7.3 × 10⁴ g mol⁻¹ to 3.0 × 10⁴ g mol⁻¹, the T_c was increased from 94.9 °C to 97.7 °C. Obviously, the reduction of the molecular weight of the PLLA matrix in this work is much smaller than those reported in the literatures, but the variation of the T_c is much more apparent. Therefore, the enhanced crystallization ability of the PLLA matrix is resulted by the combined effects of molecular weight reduction and the nucleation effects of the GOs.

Fig. 5 (a and c) DSC cooling curves and (b and d) heating curves show the nonisothermal crystallization behaviors of the PLLA and the PLLA/GO composites and subsequent melting behaviors. Both the cooling rates and heating rates are indicated in the graphs.

Fig. 5b shows the subsequent heating curves of the samples obtained at the heating rate of 10 °C min⁻¹. Pure PLLA exhibits several transitions, including the glass transition at 59.1 °C, the exothermic phenomenon attributed to the cold crystallization behavior of the PLLA at 122.9 °C, and the endothermic phenomenon attributed to the melting of the PLLA crystallites at about 150.9 and 157.0 °C. As shown in Fig. 5a, pure PLLA does not crystallize during the cooling process. Therefore, the observed endothermic phenomenon originates from the fusion of the PLLA lamellae which is formed during the DSC heating process. Composites exhibit smaller T_g compared with the pure PLLA, and T_g decreases gradually with increasing GO content. The decrease of the T_g is mainly resulted by the decrease of the molecular weight. Furthermore, the cold crystallization temperature (T_cc) also decreases with the increase of the GO content. On the other hand, the intensity of the exothermic peak decreases gradually. For the PLLA/2GO sample, the exothermic peak disappears completely. Obviously, the shift of T_cc to lower temperatures is related to the nucleation effect of the pre-existed PLLA crystallites which promotes the occurrence of the cold crystallization at relatively lower temperature. The other possibility is related to the enhanced mobility of the molecular chains, which is induced by the decrease of the molecular weight. For the PLLA/2GO sample, the crystallization is completely finished at the cooling rate of 5 °C min⁻¹ and therefore, one can not observe the cold crystallization phenomenon during the subsequent heating process. Furthermore, all the composites exhibit the double endothermic peaks. The first one (T_m1) shifts to lower temperatures with increasing GO content while the second one (T_m2) maintains nearly invariant. Obviously, there are at least two kinds of PLLA crystallites which exhibit the different melting behaviors. This will be further discussed in the following.

When the cooling rate was decreased to 2 °C min⁻¹ (Fig. 5c), pure PLLA still exhibits inconspicuous melt crystallization phenomenon while more pronounced exothermic peaks are observed for all the composites. Similarly, the T_c shifts to higher temperatures with increasing GO content. The heating curves (Fig. 5d) exhibit both glass transition and melting behavior. However, there are some important features which should be pointed out. First, pure PLLA exhibits an endothermic peak at glass transition region rather than the common step change in the region, which indicates the melting of mesophase with certain molecular ordering during the glass transition.^48,49 Obviously, the molecular ordering is induced during the cooling process. Although no crystallization is detected, the relatively smaller cooling rate facilitates the formation of some molecular ordering structure. Second, the cold crystallization phenomenon is not observed for the PLLA/GO composites, indicating that the crystallization of PLLA in the composites is completely finished during the cooling process. Third, all the composites show the invariant T_m1 and T_m2. Making a comparison of melting temperatures as shown in Fig. 5b and d, it can be concluded that the different melting behaviors of the PLLA crystallites mainly originate from the different crystallization conditions which result in the formation of the different crystallites.

The isothermal crystallization of samples was carried out at melt temperature of 115 °C for 20 min. The crystallization morphologies were first characterized using POM. The results are shown in Fig. 6. For the pure PLLA, only sporadic spherulites are observed due to that the homogeneous nucleation of the PLLA is relatively difficult and the nucleation density is small. With the presence of only 0.1 wt% GOs, the number of the spherulites is greatly increased. This proves further the nucleation effect of the GOs for the PLLA crystallization. The more the GOs in the matrix, the more apparent the nucleation effect of the GOs is. Specifically for the PLLA/2GO sample, it is very difficult to differentiate the spherulites one by one. However, except PLLA/0.1GO and PLLA/0.5GO samples which show the similar spherulites size to that of the pure PLLA, both PLLA/1GO and PLLA/2GO samples show smaller spherulites size compared with the pure PLLA. This indicates that the enhanced molecular chain mobility does not make sure the increase of the spherulites diameter of the PLLA matrix. Most likely, the largely increased nucleation sites and the formation of the GO network structure at relatively high GO concentration mainly determine the size of spherulites in this work.

Fig. 6 POM images show the isothermal crystallization morphologies of the pure PLLA and the PLLA/GO composites obtained after being crystallized at 115 °C for 20 min. (a) PLLA, (b) PLLA/0.1GO, (c) PLLA/0.5GO, (d) PLLA/1GO and (e) PLLA/2GO.

Fig. 7 shows the crystalline structures and melting behaviors of isothermally crystallized samples. Pure PLLA exhibits a typical profile of amorphous polymers without any visible diffraction peaks possibly due to that the amount of PLLA crystals is too small to be detected by WAXD. POM images already show the crystalline morphologies of the pure PLLA. However, it should be pointed out that the spherulites in the pure PLLA are very few. The difference between the results of WAXD and POM is possibly resulted by the sample preparation ways. Anyhow, both WAXD and POM characterizations prove that at melt temperature of 115 °C, it is difficult for the occurrence of the pure PLLA crystallization. For all the composites, one can observe some characteristic diffraction peaks at 2θ = 14.9°, 16.5°, 19.0° and 22.3°, attributing to the diffractions of (010), (110)/(200), (203) and (015) crystal planes, respectively. This clearly proves that the crystalline structure of the PLLA matrix in the composites is induced. The crystallinity X_c-WAXD also proves the presence of crystalline structure. The X_c-WAXD of the pure PLLA is hardly calculated because the characteristic diffraction peaks are very inconspicuous. The PLLA/0.1GO shows the X_c-WAXD of 34.7%. Further increasing the content of the GOs induces the further enhancement of X_c-WAXD. As shown in Fig. 7b, although the sample was first isothermally crystallized at 115 °C for 20 min, pure PLLA still exhibits the cold crystallization phenomenon at T_cc of about 123.5 °C. Accordingly, the observed melting behavior is mainly related to the fusion of the PLLA crystallites formed during the DSC measurement through the cold crystallization process. All the composites exhibit only glass transition and melting behaviors. Similarly, the T_g decreases gradually with increasing content of GOs. Interestingly, it is observed that the PLLA/0.1GO and the PLLA/0.5GO samples exhibit single endothermic peak at 154.0 and 152.7 °C, while the PLLA/1GO and the PLLA/2GO samples exhibit double endothermic peaks at 150.0 °C (T_m1) and 153.7–155.8 °C (T_m2), respectively. Specifically, the intensity of the first endothermic peak (T_m1) increases while the second one (T_m2) decreases. Furthermore, the values of both T_m1 and T_m2 keep invariant. The observation is also different from the results obtained from Fig. 5b and it is similar to that obtained from Fig. 5d. Similarly, the crystallinity X_c was calculated. The data are 24.0%, 24.5%, 25.1% and 27.6% for the PLLA with different GO contents from 0.1 to 2.0 wt%, respectively. The increase of X_c further indicates the positive role of the GOs in enhancing the crystallization ability of the PLLA. In addition, the great difference of crystallinity between X_c-WAXD and X_c is mainly related to the different characterization methods.

Fig. 7 (a) WAXD profiles show the crystalline structures of the pure PLLA and the PLLA/GO composites and (b) DSC heating curves show the melting behaviors of the pure PLLA and the PLLA/GO composites. Samples were first crystallized at 115 °C for 20 min.

Cold crystallization. Previous results show that for the nonisothermal crystallization which is carried out at cooling rate of 2 °C min⁻¹ and the isothermal crystallization which is carried out at 115 °C for 20 min, the crystallization of the PLLA in the composites is finished completely and no cold crystallization occurs during the DSC heating process. In this condition, composites exhibit invariant T_m1 and T_m2. If the cold crystallization occurs during the DSC heating process, the T_m1 decreases gradually with increasing GO content while the T_m2 still maintains invariant. This indicates that the cold crystallization has a great effect on the formation of the PLLA crystallites. Therefore, the cold crystallization of the PLLA is further investigated through two different methods. The first method is studying the cold crystallization of the PLLA from completely amorphous state during the DSC heating process and the second one is studying the cold crystallization of the amorphous PLLA samples during the annealing treatment.

Fig. 8 shows the DSC heating curves of all the quenched samples studied in this work. When the sample is quenched in the ice-water, the relatively larger cooling rate is provided. Consequently, the crystallization of the PLLA does not occur and completely amorphous PLLA samples are obtained. It can be seen that all the samples exhibit the three transitions including the glass transition, the cold crystallization and the melting behaviors. T_g decreases from 50.4 °C of the pure PLLA to 46.2 °C of the PLLA/2GO. T_cc also decreases gradually from 113.6 °C of the pure PLLA to 96.7 °C of the PLLA/2GO. Similarly, double endothermic peaks are observed for all the samples. Interestingly, the intensity of the first endothermic peak decreases gradually and the T_m1 also decreases gradually with increasing GO content, while T_m2 maintains invariant. This is very similar to the results obtained from Fig. 5b.

Fig. 8 DSC heating curves of the pure PLLA and the PLLA/GO composites. Samples were first quenched in ice-water.

The effect of annealing treatment on crystallization behaviors of the PLLA is also investigated. From the variations of crystalline morphologies shown in Fig. 9 one can see that after being annealed at 115 °C for 20 min, all the samples show the apparent crystalline morphologies, indicating the occurrence of the cold crystallization during the annealing process. The presence of the GOs increases the nucleation density of cold crystallization; however, the effect of the molecular weight variations on the PLLA crystallization is still inconspicuous. This is consistent with the observations obtained from isothermal crystallization processes. Otherwise, the annealed PLLA/0.1GO and PLLA/0.5GO samples should exhibit larger crystallites compared with the annealed pure PLLA. The main factors which influence the crystallization behavior of the PLLA matrix are the nucleation effect of GOs and the crystallization conditions. The former is obviously related to the content of the GOs in the composites. Furthermore, making a comparison of crystalline morphologies between the annealed samples and the isothermally crystallized samples, there are at least two differences which should be stressed. The first one is that the annealed pure PLLA exhibits the formation of large quantity of crystallites, which is different from the isothermally crystallized PLLA sample in which only sporadic spherulites are observed. The second one is that the annealed samples exhibit much smaller spherulites size but larger nucleation density compared with the isothermally crystallized samples. This indicates at least that annealing treatment is more efficient in improving the nucleation ability of the PLLA. The melting behaviors of the annealed samples are shown in Fig. 10. The cold crystallization phenomenon is not observed, and only a single endothermic peak is observed at 154.3 °C for the pure PLLA. All the composites still exhibit the double endothermic peaks and the melt temperatures are 146.6–151.1 °C for T_m1 and 153.9–155.3 °C for T_m2, respectively. Similarly, T_m1 decreases with increasing GO content while T_m2 maintains nearly invariant.

Fig. 9 POM images show the crystallization morphology of the annealed samples. Amorphous samples were annealed at 115 °C for 20 min. (a) PLLA, (b) PLLA/0.1GO, (c) PLLA/0.5GO, (d) PLLA/1GO and (e) PLLA/2GO.

Fig. 10 DSC heating curves show the melting behaviors of the annealed pure PLLA and PLLA/GO composites.

Further understanding about crystallization and melting behaviors. To further understand the relationship between molecular chains mobility and the crystallization and melting behaviors of the pure PLLA and the PLLA/GO composites under the different conditions, the variations of T_g, T_c, T_cc, T_m1 and T_m2 versus the concentration of the GOs in the composites are illustrated in Fig. 11. It can be seen that in all the conditions, the T_g decreases with increasing GO content. This agrees well with the decrease of the molecular weight of the matrix. Furthermore, the quenched samples exhibit relatively smaller T_g compared with the samples which are cooled at the small cooling rates (5 and 2 °C min⁻¹). This is because that during the quenching process, chain segments motion is frozen in a very short time and almost all of chain segments are present in the nonequilibrium state with relatively larger free volume and smaller entanglement density. Therefore, during the subsequent heating process, chain segments of the quenched samples have greater ability to relax at relatively lower temperatures, resulting in smaller T_g accordingly. However, for the slowly cooled samples, the chain segments have enough time to adjust their conformation. In this condition, the free volume is relatively smaller and chain entanglement density is also bigger, resulting in higher T_g accordingly.

Fig. 11 (a) Variation of glass transition temperature (T_g) of different samples. (b) Variation of cold crystallization temperature (T_cc) and melt crystallization temperature (T_c) of different samples. (c) Variation of T_m1 and T_m2 obtained through cold crystallization and melt crystallization process, respectively.

As shown in Fig. 11b, one can see that for the melt crystallization process, the T_c increases with increasing GO content. This indicates undoubtedly the nucleation effect of the GOs for the PLLA melt crystallization. As expected, the smaller cooling rate results in relatively higher T_c. The T_cc decreases linearly with increasing GO content. Furthermore, it should be stressed that the quenched samples exhibit smaller T_cc compared with the samples which are first cooled at cooling rate of 5 °C min⁻¹. Generally, the primary PLLA crystallites formed during the cooling process promote the occurrence of the cold crystallization at relatively lower temperatures during the subsequent heating process. However, this work proves that the quenched samples with relatively bigger free volume and smaller chain entanglement density have greater ability to crystallize at lower temperatures during the DSC heating process.

The dependence of both T_m1 and T_m2 on the content of the GOs is illustrated in Fig. 11c. The melting temperatures are classified as two systems, System I is related to the fact that cold crystallization occurs either during the annealing process or during the DSC heating process and System II to the melt crystallization process without the occurrence of any cold crystallization. It can be seen that if cold crystallization occurs, T_m1 decreases apparently with increasing content of the GOs. However, if no cold crystallization occurs during the DSC heating process, T_m1 just changes in a very small range. Furthermore, it can be seen that the value of T_m1 of the System I is slightly smaller than that of the System II. This indicates that the T_m1 of the System I is related to the fusion of lamellae which is induced through cold crystallization process, while the T_m1 of the System II is related to the lamellae which is induced during the melt crystallization. The difference of T_m1 in System I and System II can be explained as follows. Although the cold crystallization temperature is much higher than the T_g of the PLLA matrix, the samples of the System I are still in the rubbery state when the cold crystallization occurs. In this condition, the mobility of chain segments is limited. Furthermore, the free volume is relatively smaller compared with the samples of the System II. On the other hand, the presence of the GOs most likely restricts the lamellar growth of the PLLA during the cold crystallization process, especially when the GOs exhibit the percolated network structure in the composites. Therefore, the lamellar growth of the PLLA matrix in the composites is difficult during the cold crystallization process, which results in lamellae formation with more defects or smaller thickness. For the System II, samples are cooled directly from the melt state, and therefore, both the chain segments mobility and the free volume of the PLLA are relatively larger, which facilitates the formation of lamellae with bigger thickness. Consequently, a relatively higher T_m1 is achieved. In this condition, the restriction effect of the GOs for the PLLA crystallization becomes inconspicuous. Furthermore, whether for the cold crystallization process or for the melt crystallization process, it is observed that the variation of the T_m2 is inconspicuous. Specifically, the System I and the System II exhibit the similarT_m2. This can be attributed to the melt–recrystallization–melt process occurred during the DSC heating process. Since the recrystallization of all the samples occurs at similar supercooling degree, they exhibit the similar PLLA lamellar structure. Consequently, the similar T_m2 is obtained for all the samples. Furthermore, it is also observed that some samples (pure PLLA and PLLA/0.1GO samples in Fig. 7b and pure PLLA in Fig. 10) exhibit single endothermic peak during the DSC heating scan. One possibility is that a relatively more perfect crystalline structure of the PLLA is induced during the crystallization process and no melt–recrystallization–melt process occurs during the subsequent DSC heating scan. The relative high temperature for the cold crystallization (during the DSC heating scan) and the enhanced chain mobility of the PLLA matrix are possibly the main reasons for the formation of the perfect crystalline structure.⁵⁰ The other possibility is that T_m1 overlaps with T_m2. This is possibly resulted by the inconspicuous melt–recrystallization–melt processing during the DSC heating scan.

4. Conclusions

The PLLA/GO composites containing different contents of the GOs have been prepared. The morphological characterization shows that the GOs exhibit the platelets with multiple layers. Large holes/cavities are observed between GOs and PLLA, indicating the weak interfacial interaction between GOs and PLLA. The presence of the GOs accelerates the thermal degradation of the PLLA matrix, which results in the decrease of the molecular weight. As a consequence, the storage modulus, loss modulus and complex viscosity of the melt are greatly decreased. However, GOs form the percolated network structure at relatively higher content, promoting the increase of modulus and complex viscosity of the melt at relatively low frequency. Studying on the crystallization behaviors shows that GOs exhibit great nucleation effect for the crystallization of the PLLA. Whether for the melt crystallization process or for the cold crystallization process, the presence of the GOs greatly improves the crystallization ability of the PLLA matrix. The effect of the molecular weight variations on the spherulites growth of the PLLA is inconspicuous. The content of the GOs and the crystallization conditions are the main factors which determine the crystalline structure of the PLLA matrix. Due to the relatively smaller chain segments mobility and the restriction of the GO sheet for the growth of the PLLA lamella, the primary crystallites of the PLLA induced by cold crystallization process are imperfect compared with those prepared from melt crystallization process. This work proves that GOs have a role of improving the crystallization ability of the PLLA matrix, but at the cost of the decrease of the PLLA molecular weight. Therefore, the stabilization of the PLLA matrix must be considered in designing the PLLA/GO materials.

Acknowledgements

Authors express their sincere thanks to the National Natural Science Foundation of China (51203129, 50973090), Distinguished Young Scholars Foundation of Sichuan (2012JQ0057) and the Fundamental Research Funds for the Central Universities (SWJTU12CX010, SWJTU11CX142, SWJTU11ZT10) for financial support.

References

1. Z. Kulinski and E. Piorkowska, Polymer, 2005, 46, 10290.
2. H. B. Li and M. A. Huneault, Polymer, 2007, 48, 6855.
3. L. Yu, H. S. Liu, K. Dean and L. Chen, J. Polym. Sci., Part B: Polym. Phys., 2008, 46, 2630.
4. W. C. Lai, W. B. Liau and T. T. Lin, Polymer, 2004, 45, 3073.
5. H. W. Xiao, W. Lu and J. T. Yeh, J. Appl. Polym. Sci., 2009, 113, 112.
6. H. W. Xiao, F. Liu, T. Jiang and J. T. Yeh, J. Appl. Polym. Sci., 2010, 117, 2980.
7. Y. H. Cai, S. F. Yan, J. B. Yin, Y. Q. Fan and X. S. Chen, J. Appl. Polym. Sci., 2011, 121, 1408.
8. Z. H. Xu, Y. H. Niu, L. Yang, W. Y. Xie, H. Li, Z. H. Gan and Z. G. Wang, Polymer, 2010, 51, 730.
9. S. C. Schmidet and M. A. Hillmyer, J. Polym. Sci., Part B: Polym. Phys., 2001, 39, 300.
10. H. Yamane and K. Sasai, Polymer, 2003, 44, 2569.
11. H. Tsuji, H. Takai and S. W. Saha, Polymer, 2006, 47, 3826.
12. J. Narita, M. Katagiri and H. Tsuji, Macromol. Mater. Eng., 2011, 296, 887.
13. T. Villmow, P. Pötschke, S. Pegel, L. Häussler and B. Kretzschmar, Polymer, 2008, 49, 3500.
14. Y. Y. Zhao, Z. B. Qiu and W. T. Yang, Compos. Sci. Technol., 2009, 69, 627.
15. Y. Y. Zhao, Z. B. Qiu, S. K. Yan and W. T. Yang, Polym. Eng. Sci., 2011, 51, 1564.
16. Y. T. Shieh, Y. K. Twu, C. C. Su, R. H. Lin and G. L. Liu, J. Polym. Sci., Part B: Polym. Phys., 2010, 48, 983.
17. S. M. Huang, J. J. Hwang, H. J. Liu and L. H. Lin, J. Polym. Sci., Part B: Polym. Phys., 2010, 117, 434.
18. N. Vasanthan, H. Ly and S. Ghosh, J. Phys. Chem. B, 2011, 115, 9556.
19. W. S. Chow and S. K. Lok, J. Therm. Anal. Calorim., 2009, 95, 627.
20. K. Das, D. Ray, I. Banerjee, N. R. Bandyopadhyay, S. Sengupta, A. K. Mohanty and M. Misra, J. Appl. Polym. Sci., 2010, 118, 143.
21. D. H. Xu and Z. G. Wang, Polymer, 2008, 49, 330.
22. Y. L. Li, H. Y. Wu, Y. Wang, L. Liu, L. Han, J. Wu and F. M. Xiang, J. Polym. Sci., Part B: Polym. Phys., 2010, 48, 520.
23. Y. L. Li, X. X. Li, F. M. Xiang, T. Huang, Y. Wang, J. Wu and Z. W. Zhou, Polym. Adv. Technol., 2011, 22, 1959.
24. Y. Sun and C. B. He, ACS Macro. Lett., 2012, 1, 709.
25. A. M. Pinto, J. Cabral, D. A. P. Tanaka, A. M. Mendes and F. D. Magalhaes, Polym. Int., 2013, 62, 33.
26. J. Z. Xu, T. Chen, C. L. Yang, Z. M. Li, Y. M. Mao, B. Q. Zeng and S. H. Benjamin, Macromolecules, 2010, 43, 5000.
27. N. I. Kovtyukhova, P. J. Ollivier, B. R. Martin, T. R. Mallouk, S. A. Chizhik, E. V. Buzaneva and A. D. Gorchinskiy, Chem. Mater., 1999, 11, 771.
28. J. H. Yang, C. X. Feng, J. Dai, N. Zhang, T. Huang and Y. Wang, Polym. Int., 2013, 62, 1085.
29. E. W. Fischer, H. J. Sterzel and G. Wegber, Kolloid-Zeitschrift und Zeitschrift für Polymere, 1973, 251, 980.
30. Y. Wang, Z. X. Shi, J. R. Yu, L. Chen, J. Zhu and Z. M. Hu, Carbon, 2012, 50, 5525.
31. J. Zhang, Y. Duan, H. Sato, H. Tsuji, N. Isao, S. Yan and Y. Ozaki, Macromolecules, 2005, 38, 8012.
32. I. C. McNeill and H. A. Leiper, Polym. Degrad. Stab., 1985, 11, 309.
33. F.-D. Kopinke, M. Remmler, K. Mackenzie, M. Milder and O. Wachsen, Polym. Degrad. Stab., 1996, 43, 329.
34. S. W. Ko, R. K. Gupta, S. N. Bhattacharya and H. J. Choi, Macromol. Mater. Eng., 2010, 295, 320.
35. B. Wang, T. Wan and W. Zeng, J. Appl. Polym. Sci., 2011, 121, 1032.
36. S. Singh, A. K. Ghosh, S. N. Maiti, S. Raha, R. K. Gupta and S. Bhattacharya, Polym. Eng. Sci., 2011, 55, 225.
37. C. D. Han and K. W. Lem, Polym. Eng. Rev., 1983, 2, 135.
38. P. Pötschke, T. D. Fornes and D. R. Paul, Polymer, 2002, 43, 3247.
39. C. McClory, T. McNally, M. Baxendale, P. Pötschke, W. Blau and M. Ruether, Eur. Polym. J., 2010, 46, 854.
40. R. A. Khare, A. R. Bhattacharyya, A. R. Kulkarni, M. Saroop and A. Biswas, J. Polym. Sci., Part B: Polym. Phys., 2008, 46, 2286.
41. H. K. F. Cheng, N. G. Sahoo, Y. Z. Pan, L. Li, S. H. Chan, J. H. Zhao and G. Chen, J. Polym. Sci., Part B: Polym. Phys., 2010, 48, 1203.
42. D. F. Wu, Y. Zhang, M. Zhang and W. Yu, Biomacromolecules, 2009, 10, 417.
43. H. B. Zhang, W. G. Zheng, Q. Yan, Z. G. Jiang and Z. Z. Yu, Carbon, 2012, 50, 5117.
44. G. Filippone, M. S. D. Luna, D. Acierno and P. Russo, Polymer, 2012, 53, 2699.
45. Y. F. Li, J. H. Zhu, S. Y. Wei, J. Ryu, L. Y. Sun and Z. H. Guo, Macromol. Chem. Phys., 2011, 212, 1951.
46. P. Pan, W. Kai, T. Dong and Y. Inoue, Macromolecules, 2007, 40, 6898.
47. I. Pillin, N. Montrelay, A. Bourmaud and Y. Grohens, Polym. Degrad. Stab., 2008, 93, 321.
48. G. Stoclet, R. Seguela, J. M. Lefebvre and C. Rochas, Macromolecules, 2010, 43, 7228.
49. R. H. Lv, B. Na, N. N. Tian, S. F. Zou, Z. J. Li and S. C. Jiang, Polymer, 2011, 52, 4979.
50. Y. Y. Shi, L. N. Shao, J. H. Yang, T. Huang, Y. H. Wang, N. Zhang and Y. Wang, Polym. Adv. Technol., 2013, 24, 42.

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