C60 as fine fillers to improve poly(phenylene sulfide) electrical conductivity and mechanical property

Electrical conductive poly(phenylene sulfide) (PPS)/fullerene (C60) composites were prepared by 1-chlornaphthalene blending method, and the interface effects of C60 and PPS on PPS/C60 properties were characterized. C60 is an excellent nanofiller for PPS, and 2 wt% PPS/C60 composite displayed the optimal conductivity which achieved 1.67 × 10−2 S/cm. However, when C60 concentration reached 2 wt%, the breaking strength and tensile modulus of PPS/C60 fiber achieved maximum 290 MPa and 605 MPa, and those values were 7.72 and 11.2 times as that of pure PPS. The excellent conductive and mechanical properties of PPS/C60 were attributed to the heterogeneous nucleation of C60 during PPS crystallization, formation of a large number of covalent bond by main C60-thiol adducts and minor C60-ArCl alkylation between C60 outer surface and PPS matrix. At same time, PPS/C60 thermal properties were also investigated.

, functionalized SWCNTs and MWCNTs [17][18][19][20] , nano particles of TiO 2 , ZnO, CuO and SiC 21 , nanoscale alumina particles 22 , graphite 23 , glass fiber 24 , metal inorganic salt 25 , nano-SiOx 26 , carbon and fiber 27 et al. However, it has been recently shown that the properties of PPS based composites hardly increase at low nanofiller loadings (up to 1 wt%) 14,28,29 . Because of van der Waals attraction between C 60 and their large surface area, C 60 tend to form agglomerates during mixing with PPS by melt blending. Therefore, it is difficult to use the conventional method to disperse C 60 in the PPS matrix.
Herein, C 60 was selected as fillers to improve PPS electrical conductivity and mechanical property, and the influence of the 1-chloronaphthalene solution mixing method and subsequent melt process on filler dispersion, C 60 interfacial adhesion with PPS matrix and composite properties were investigated to correlate the microscopic structure with macroscopic properties. Sample preparation. PPS (5 g) was dissolved in 100 ml of 1-chloronaphthalene at 205 °C under nitrogen atmosphere. After that, a certain amount of C 60 ranging from 0.5 to 10 wt% were loaded into PPS/1-chloronaphthalene solution at 205 °C under mechanical agitation. PPS/C 60 composites with 0.5 to 10 wt% nanofiller were obtained after removing 1-chloronaphthalene under vacuum condition. And then PPS and PPS/ C 60 composite fibers with diameters of 45~85 µm were prepared by melt spinning technology at 315 °C.
Characterization and Measurement. Fourier transform infrared spectroscopy (FT-IR) spectra were obtained using Bruker IFS66 at room temperature. Thermo gravimetric analyzer (TGA) analysis was carried out on a NETZSCH STA 409 TG analyzer, and the rate of temperature increase was at 10 °C/min. X-ray diffraction (XRD) data were obtained using an Elmer PHI-5600 instrument using a Mg Kα line as a radiation source and a D8 discover. The morphology of composites was characterized by a field emission scanning electron microscopy (FESEM, Hitachi 4800S, and Japans) and transmission electron micrographs (TEM, Hitachi H-7650 microscope). Differential scanning calorimetry (DSC) analysis was performed on a Perkin Elemer DSC-7 under nitrogen condition, and samples placed in aluminium pans were melted at 320 °C and kept at this temperature for 5 min to erase their thermal history. Subsequently, they were cooled from the melt to room temperature and then heated again up to 320 °C at a scan rate of 10 °C/min. From the DSC heating and cooling traces, peak melting temperature (T m ), heat of melting (ΔH m ), peak crystallization temperature (T c ) and heat of crystallize (ΔH c ) were obtained. The degree of crystallinity (X c ) was calculated from the following equation: Where ΔH c is the cold crystallization enthalpy from the DSC scan, W f is the weight fraction of C 60 in composites, and ΔH f is the melting enthalpy of 100% crystallized PPS which was taken as 105 J/g 23 . Fiber diameter data were obtained using a KEYENCE VHX-1000 microscope at room temperature. The breaking strength (δ t ), breaking elongation (ε) and tensile modulus (MPa) were measured on single fiber strength tester (China LLY-06). Each sample was tested ten times to evaluate the average value. The breaking strength was calculated from the following equation: Where F b is the maximum tension value, d is the fiber diameter. The breaking elongation was determined by the following equation: Where ε is the breaking elongation, L is the length of fiber elongation, L 0 is the initial length of the fiber before test. The fiber tensile modulus (E MPa) was determined by the following equation: Electrical conductivity of samples was measured by the four-point probe method using a Scientific Equipment device with a spacing probe S = 0.2 cm equipped with a DC precision power source (Model LCS-02) and a digital microvoltmeter (Model DMV-001). The powder sample was filled into a test slot, then applying a pressure about 18 MPa. The electrical resistivity was calculated from the following equation: Where S is the distance of adjacent probe, W is the sample thickness, D is the position correction factor, V is the test voltage, I is the test current. Each sample has been tested ten times, and the electrical resistivity is the average value of 10 measurement. The electrical conductivity of samples was calculated through the following equations: = Electrical conductivity 1/Electrical resistivity (6)

Results
After C 60 and PPS were dissolved in 1-chloronaphthalene, 1-chlornaphthalene evaporated under vacuum condition to obtain PPS/C 60 composites, and the dispersion and alignment of C 60 within the PPS matrix were characterized by TEM (Fig. 1). The dark and light areas correspond to nanofillers and PPS matix, respectively. The image of C 60 showed a homogeneous size distribution with average diameter around 300 nm, which revealed C 60 was easily agglomerated due to the strong van der Waals forces and intensive π-π stacking interactions among C 60 . However, the micrograph of 0.5-2 wt% C 60 /PPS composites indicated the aggregate C 60 diameter decrease, and no voids or discontinuities are detected between the C 60 outer surface and PPS matrix. The TEM micrograph of 0.5-2 wt% PPS/C 60 fibers indicated that C 60 was well dispersed into PPS matrix, however, when much more C 60 was incorporated into PPS (such as 10 wt% PPS/C 60 ), C 60 formed agglomerates inside the PPS matrix. TEM observation showed that C 60 were wrapped in PPS or covered by PPS layer and the heterogeneous dispersed bright dots with dimensions from 150-350 nm in 5 wt% and 10 wt% PPS/C60 composites were detected, indicating good adhesion between C 60 and PPS. Figure 2 A shows the SEM micrograph of PPS as a blank example, and the SEM images of Fig. 2B,C,D indicated that almost no apparent C 60 agglomerates were detected in 0.5 wt%, 1 wt% and 2 wt% PPS/C 60 composites, due to C 60 nanofillers further disaggregation within the polymer matrix during PPS/C 60 melt-process. C 60 nanofillers were randomly dispersed in the PPS matrix, occurring irregular crystal with dimensions from 20-50 nm when the content of C 60 is less than 2 wt%. C 60 nanofillers are difficult to disperse uniformly in PPS melt, therefore PPS can not well reinforced by C 60 nanofillers in melt blending technology. However, C 60 aggregates can be well dispersed in PPS matrix by solution blending method, in which the solvent plays a dual role in dispersing C 60 and preventing C 60 agglomeration. The heterogeneous dispersed bright dots with dimensions from 80-350 nm in 5 wt% and 10 wt% PPS/C 60 composites were detected, which was attributed to C 60 agglomerates. Similar results can also be found through the cross-sectional SEM images of the composites in Fig. 3. For PPS/C 60 , 2 wt% C 60 particles can be uniformly dispersed in PPS matrix, but further addition of 5 wt% C 60 would cause the agglomeration of C 60 , as can be seen in Fig. 3e and f. X-ray diffraction patterns of PPS/C 60 composites were displayed in Fig. 4. The two diffraction peaks at 19.2° and 20.8° are corresponding to the (110) and (200) crystalline planes of the orthorhombic structure of PPS 30 . And C 60 shows characteristic peaks at 2θ = 10.9°, 17.2°, 20.8°, 21.8°, 28.2°, 30.8° and 32.7° arising from the (110), (220), (310), (220), (330), (420) and (330) crystalline planes of the orthorhombic unit cell, respectively 31 . The C 60 agglomerate peaks were hardly visible in the diffractograms of 0.5-2wt% composites, which suggested that C 60 was well dispersed in PPS matrix. After 5 wt% C 60 was introduced into PPS, the serious C 60 agglomerate was clearly observed, and the crystal characteristic peaks of C 60 shifted to lower 2θ values bacause C 60 aggregates result in PPS lattice distortion 32 . Moreover, composite peaks become broader with reduced intensity, indicating the structural order decline which induced by the incorporation of C 60 . These observations are consistent with the behaviors of SWCNTs 33 and MWCNTs 34 , where the local order of the polymer decreased after the grafting reaction.
As shown in Fig. 5, the main bands at 565, 615, 1218, 1470 and 1680 cm −1 are characteristic absorption peak of C 60 35 . The phenyl groups of PPS exhibited the absorption peaks at 1628 and 1405 cm −1 , and the two bands at 1100 and 623 cm −1 were attributed to aromatic C-S stretching vibrations. After PPS was reinforced by C 60 , C 60 characteristic absorption peaks at 565 and 1475 cm −1 band were still observed, which demonstrated the C 60 successful incorporation into PPS matrix. However, C 60 band at 1218 cm −1 disappeared and a new peak at 1010 cm −1 appeared, which implied the well C 60 dispersion and C 60 -S formation (Fig. 6) 36,37 . Solvent could help the tangled PPS molecular chain creeping and stretching in 1-chloronaphthalene, which effectively contributes to the well dispersion of C 60 in PPS matrix. Although the alkylation or acylation was an effective functionalization method for fullerene 38,39 , herein a C 60 -thiol adducts by reacting C 60 with the end group SH of PPS was easier to take place 32 , and the covalent bond formation improved the C 60 -matrix interfacial adhesion 6,7 . It is general known that the formation of π-π stacking interactions can be characterized through the shift of -CH bond. It is clear that the C-H bond absorption peaks of PPS located at 808.7 and 818.8 cm −1 . In PPS/C 60 material, C-H vibration peak of PPS/C60 remain unchanged, which indicated that there only existed a weak π-π stacking interactions between C 60 and PPS matrix 36,40 .
In order to testify the existence of chain-extension reaction through the end-group reaction between PPS and C 60 , PPS and PPS/C 60 molecular weight were characterized by HTGPC ( Table 1). As C 60 content increased from 0.5 wt% to 2 wt%, PPS/C 60 composite molecular weight propagated from 3.30 × 10 4 g·mol −1 to 4.19 × 10 4  g·mol −1 , and polydispersity index (PDI) increased from 2.22 to 2.56. Only some PPS chains could interact with C 60 to lead to the longer polymer chains. However, when C 60 content reached 5 wt%, 5 wt% PPS/C 60 composite molecular weight declined to 3.71 × 10 4 g·mol −1 . The deviation could be attributed to much more the low molecular weight C 60 loading. A similar result was previously reported by Peng, K. J. & Liu, Y. L. 41 . HTGPC characterized results suggested some covalent bond formation between PPS and C 60 through the C 60 -thiol adducts and C 60 -ArCl alkylation.
The crystallization and melting behavior of composites were investigated (Fig. 7), and the calorimetric parameters derived from non-isothermal DSC scans were listed in Table 2. As it can be observed, 0.5 wt% C 60 had less influence on PPS crystallization temperature (Tc) increase (Fig. 7a), and 0.5 wt% PPS/C 60 exhibited T c of 217.2 °C with ΔH c being 42.98 J/g. However, for 2 wt% PPS/C 60 , T c shifted to 227.6 °C with maximum ΔH c being 48.23 J/g. AS for 5 wt% PPS/C 60 , T c increases to 223.8 °C with ΔH c decline to 42.42 J/g. The further increase of C 60 concentration in composites slowed the mobility and diffusion of PPS chains, which led to a significant decline of PPS crystallization temperature. Those behaviors observed are in agreement with that reported by Jeon et al. 42 . It was worthy of noting that 0.5 wt% PPS/C 60 could not well act as heterogeneous nucleating agent to accelerate PPS nucleation, which suggested that the intense restrictions on chain mobility are imposed by the C 60 -polymer chemical interactions.
T m shifted gradually to higher temperature with increasing C 60 content (See Fig. 7b and Table 2). 0.5 wt% PPS/ C 60 exhibited T m of 275.1 °C with ΔHm being 41.50 J/g. As for 2 wt% PPS/C 60 , T m shifted to 277.3 °C with maximum ΔHm being 44.10 J/g. While for 5 wt% PPS/C 60 , T m increased to 282.4 °C with ΔHm decline to 40.14 J/g. In the case of 2 wt% PPS/C 60 , both ΔH m and ΔH c achieve the maximum, and X c reaches the maximum value 46.87%.  The TG curves for the pure PPS matrix and composites under inert atmospheres were shown in Fig. 8, and their characteristic degradation parameters were summarized in Table 3. PPS displayed a single degradation stage that starts (T i ) at 502 °C and exhibited the maximum weight loss (T max ) rate at 541.3 °C. At 800 °C, the residual mass was about 52.2% of the initial weight. Clearly, the addition of 0.5-2 wt% C 60 filler led to an improvement in the thermal stability of PPS matrix, and a maximum T i increase (about 7 °C) was obtained at 2.0 wt% filler loading, and T max increment for 2.0 wt% PPS/C 60 was 5.5 °C. However, the significantly decline of T i and T max was detected at 10 wt% C 60 loading, but the residual mass at 800 °C increased to 60.39% of the initial weight. Such results should be attributed to different factors. Firstly, C 60 fillers are better dispersed within PPS matrix, which restricts chain mobility or diffusion to slow down the decomposition process 11 . Secondly, the covalent anchoring of PPS to C 60 leads to a strong enhancement in the thermal conductivity that facilitates heat dissipation within the composite 43 . The reason for the thermal stability decline at 5 and 10 wt% C 60 loading might be attributed to the appearance of C 60 agglomerates. Figure 9 shows single fiber morphology of each sample respectively. The characterized results suggested the diameter of fibers were 45.2 ± 0.8 μm, 40.5 ± 0.5 μm, 42.5 ± 1.0 μm, 64.3 ± 1.0 μm, 47.6 ± 1.0 μm and 87.2 ± 0.8 μm, respectively, changing with the content of C 60 . The mechanical behavior of PPS/C 60 fibers was investigated by single fiber strength tester technique which provides additional information about filler-matrix and filler-filler interactions. Figure 10 showed the breaking strength, breaking elongation and tensile modulus of PPS/C 60 fiber. The results of mechanical property study indicated that the concentration of C 60 had a greater influence on the mechanical performance of composites. As C 60 nanofiller content increase, the breaking strength and tensile modulus of composites firstly increased and then decreased. When C 60 concentration reached 2 wt%, the breaking  strength and tensile modulus of composites achieved maximum 290 MPa and 605 MPa, and those value were 7.72 and 11.2 times as that of pure PPS, respectively. The breaking elongation of PPS/C 60 composites always decreased with increasing C 60 content (Fig. 10b). The excellent mechanical properties of PPS/C 60 were attributed to the   heterogeneous nucleation of C 60 during PPS crystallization, the formation of a large number of covalent bond by C 60 -thiol adducts and π-π stacking interactions between C 60 surface and PPS matrix. However, the excessive addition of C 60 caused a significant reduction of breaking strength, e.g., the breaking strength of 10 wt% C 60 composite declined to 148 MPa. The excessive addition of C 60 reduced PPS crystallization degree and decreased the combination between PPS and C 60 , which were attributed to the phenomenon of C 60 aggregation. These results can also confirm by the SEM and TEM images of PPS/C 60 composites presented in Figs 1-2. PPS is an insulating material (σ 10 −16 S/cm), which limits its use in self-health monitoring, electro-actuation, etc 32 . Herein, C 60 was used as conductive fillers to improve PPS electrical conductivity and the electrical performance of PPS/C 60 composites were compared (Fig. 11). As C 60 nanofiller content increase, the electrical conductivity of  PPS/C 60 composites firstly increased and then decreased. When C 60 concentration reached 2 wt%, the electrical conductivity of composites achieved maximum 1.67 × 10 −2 S/cm, much higher than the value of pure PPS. The excellent electrical conductivity of 2 wt% PPS/C 60 composites was attributed to the well dispersed C 60 fillers, and the maximum conductive networks might be formed in the composite at appropriate C 60 content because of the conductive network formation 44 . C 60 fillers are well dispersed in the PPS matrix owing to the covalent bond formation by main C 60 -thiol adducts and mino C 60 -ArCl alkylation between C 60 surface and PPS. However, the excessive addition of C 60 nanofillers caused a significant reduction of electrical conductivity, e.g., the electrical conductivity of 10 wt% C 60 composite declined to 3.79 × 10 −3 S/cm, due to the phenomenon of C 60 aggregation.

Conclusions
PPS composites with well-dispersed C 60 had been prepared by solution co-blending method because solvent promotes the tangled PPS molecular chains creeping and stretching. The electrical conductivity value of 2 wt% achieved the maximum, and the excellent electrical conductivity of 2 wt% PPS/C 60 composites was mainly attributed to the covalent bond formation by main C 60 -thiol adducts and minor C 60 -ArCl alkylation between C 60 surface and PPS, but C 60 aggregation reduced composite electrical conductivity. Furthermore, 2 wt% C 60 could effectively increase PPS crystalli zation temperature, thermal stability and mechanical performance. However, the excessive C 60 loading reduced the PPS crystallization degree and caused C 60 re-aggregation, which led to poorer mechanical performance of PPS/C 60 composite. Figure 11. Room temperature electrical conductivity (r) of PPS/ C60 composites.