Introduction

Fullerene/polymer composites have recently attracted considerable attention from the materials research community1,2,3. It has been shown that the incorporating fullerene or its derivatives into a polymer matrix could significantly improve its photovoltaic, electromechanical and thermomechanical properties4, 5. The literature already identified several techniques to blend fullerene with polymers3, and the properties of fullerene/polymer composites are closely related to the morphologies of the polymers formed during melt or solution processing3, 4. Chemical functionalization of fullerenes is often performed to increase the miscibility of fullerenes with host polymers6, but the chemical functionalization method makes fullerenes lost certain original precious properties, which resulted in chemical attachment of fullerenes to a polymer is not always the most effective method for getting high performance composite materials, despite a number of fullerene-attached polymers have been successfully synthesized in the past7.

PPS is widely used in practice due to its excellent properties, such as chemical resistance, low coefficient of friction, mechanical behaviors, dimensional stability and electrical property8,9,10. To improve the performance of PPS, the nanofiller-reinforced composites require homogenous filler dispersion and good interfacial adhesion with the host matrix. Attaching functional groups onto the filler surface has been proven to be an effective approach to prepare polymer composites, and another effective method is wrapping fillers by organic molecules to provide the π-π stacking interactions between PPS matrix and nanofiller sidewalls. The conventional fillers included single-walled carbon nanotubes (SWCNTs)11, 12, inorganic fullerene-like tungsten disulfide (IF-WS2)13, SWCNT-IF-WS2 14,15,16, functionalized SWCNTs and MWCNTs17,18,19,20, nano particles of TiO2, ZnO, CuO and SiC21, nanoscale alumina particles22, graphite23, glass fiber24, metal inorganic salt25, nano-SiOx26, carbon and fiber27 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 C60 and their large surface area, C60 tend to form agglomerates during mixing with PPS by melt blending. Therefore, it is difficult to use the conventional method to disperse C60 in the PPS matrix.

Herein, C60 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, C60 interfacial adhesion with PPS matrix and composite properties were investigated to correlate the microscopic structure with macroscopic properties.

Experimental

Materials and reagents

C60 powder (>99.9 wt/wt purity) was purchased from Puyang Yongxin Fullerene Co., Ltd, and used as received. PPS (Mw~3.3 × 104, Mw/Mn~1.4 × 104, d25 °C~1.35 g/cm3, Tg~90 °C, Tm~280 °C) was synthesized from Na2S and 1,4-dichlorobenzene in N-methylpyrrole, and the end groups of synthesized PPS contained –SH and -PhCl. 1-chlornaphthalene (95%, Fluka) was purchased from J& K Chemical Co., Ltd and purified by distillation under reduced pressure before use.

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 C60 ranging from 0.5 to 10 wt% were loaded into PPS/1-chloronaphthalene solution at 205 °C under mechanical agitation. PPS/C60 composites with 0.5 to 10 wt% nanofiller were obtained after removing 1-chloronaphthalene under vacuum condition. And then PPS and PPS/C60 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:

$${X}_{c}( \% )=\frac{{\rm{\Delta }}{H}_{c}}{{\rm{\Delta }}{H}_{f}(1-{W}_{f})}\times 100 \% $$
(1)

Where ΔH c is the cold crystallization enthalpy from the DSC scan, \({W}_{f}\) is the weight fraction of C60 in composites, and ΔH f is the melting enthalpy of 100% crystallized PPS which was taken as 105 J/g23.

Fiber diameter data were obtained using a KEYENCE VHX-1000 microscope at room temperature. The breaking strength (\({\delta }_{t}\)), breaking elongation (\(\varepsilon \)) 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:

$${\delta }_{t}(MPa)=4\frac{{F}_{b}}{\pi {d}^{2}}\times {10}^{-6}$$
(2)

Where F b is the maximum tension value, d is the fiber diameter. The breaking elongation was determined by the following equation:

$$\varepsilon ( \% )=\frac{L-{L}_{0}}{{L}_{0}}$$
(3)

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:

$$E=\frac{{\delta }_{t}}{\varepsilon }=\frac{4{F}_{b}{L}_{0}}{\pi {d}^{2}(L-{L}_{0})}\times {10}^{-4}$$
(4)

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:

$${\rm{Electrical}}\,{\rm{resistivity}}={\rm{2}}\,{\rm{\pi }}S\times (W/S)\times D\times (V/I)$$
(5)

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:

$${\rm{Electrical}}\,{\rm{conductivity}}=1/\mathrm{Electrical}\,{\rm{resistivity}}$$
(6)

Results

After C60 and PPS were dissolved in 1-chloronaphthalene, 1-chlornaphthalene evaporated under vacuum condition to obtain PPS/C60 composites, and the dispersion and alignment of C60 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 C60 showed a homogeneous size distribution with average diameter around 300 nm, which revealed C60 was easily agglomerated due to the strong van der Waals forces and intensive π-π stacking interactions among C60. However, the micrograph of 0.5–2 wt% C60/PPS composites indicated the aggregate C60 diameter decrease, and no voids or discontinuities are detected between the C60 outer surface and PPS matrix. The TEM micrograph of 0.5–2 wt% PPS/C60 fibers indicated that C60 was well dispersed into PPS matrix, however, when much more C60 was incorporated into PPS (such as 10 wt% PPS/C60), C60 formed agglomerates inside the PPS matrix. TEM observation showed that C60 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 C60 and PPS.

Figure 1
figure 1

TEM micrographs of PPS/C60 powder with different weight fractions of C60 (a) Pure PPS; (b) 0.5 wt%; (c) 1 wt%; (d) 2 wt%; (e) 5 wt%; (f) 10 wt%.

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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 C60 agglomerates were detected in 0.5 wt%, 1 wt% and 2 wt% PPS/C60 composites, due to C60 nanofillers further disaggregation within the polymer matrix during PPS/C60 melt-process. C60 nanofillers were randomly dispersed in the PPS matrix, occurring irregular crystal with dimensions from 20–50 nm when the content of C60 is less than 2 wt%. C60 nanofillers are difficult to disperse uniformly in PPS melt, therefore PPS can not well reinforced by C60 nanofillers in melt blending technology. However, C60 aggregates can be well dispersed in PPS matrix by solution blending method, in which the solvent plays a dual role in dispersing C60 and preventing C60 agglomeration. The heterogeneous dispersed bright dots with dimensions from 80–350 nm in 5 wt% and 10 wt% PPS/C60 composites were detected, which was attributed to C60 agglomerates. Similar results can also be found through the cross-sectional SEM images of the composites in Fig. 3. For PPS/C60, 2 wt% C60 particles can be uniformly dispersed in PPS matrix, but further addition of 5 wt% C60 would cause the agglomeration of C60, as can be seen in Fig. 3e and f.

Figure 2
figure 2

SEM micrographs of cryogenically fractured melt-processed PPS/C60 with different weight fractions of C60 (a) PPS; (b) 0.5 wt%; (c) 1 wt%; (d) 2 wt%; (e) 5 wt%; (f) 10 wt%.

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Figure 3
figure 3

Cross-sectional SEM micrographs of cryogenically fractured melt-processed PPS/C60 with different weight fractions of C60 (a) PPS; (b) 0.5 wt%; (c) 1 wt%; (d) 2 wt%; (e) 5 wt%; (f) 10 wt%.

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X-ray diffraction patterns of PPS/C60 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 PPS30. And C60 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, respectively31. The C60 agglomerate peaks were hardly visible in the diffractograms of 0.5–2wt% composites, which suggested that C60 was well dispersed in PPS matrix. After 5 wt% C60 was introduced into PPS, the serious C60 agglomerate was clearly observed, and the crystal characteristic peaks of C60 shifted to lower 2θ values bacause C60 aggregates result in PPS lattice distortion32. Moreover, composite peaks become broader with reduced intensity, indicating the structural order decline which induced by the incorporation of C60. These observations are consistent with the behaviors of SWCNTs33 and MWCNTs34, where the local order of the polymer decreased after the grafting reaction.

Figure 4
figure 4

XRD spectra of C60, PPS and PPS/C60 composite powder.

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As shown in Fig. 5, the main bands at 565, 615, 1218, 1470 and 1680 cm−1 are characteristic absorption peak of C60  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 C60, C60 characteristic absorption peaks at 565 and 1475 cm−1 band were still observed, which demonstrated the C60 successful incorporation into PPS matrix. However, C60 band at 1218 cm−1 disappeared and a new peak at 1010 cm−1 appeared, which implied the well C60 dispersion and C60-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 C60 in PPS matrix. Although the alkylation or acylation was an effective functionalization method for fullerene38, 39, herein a C60-thiol adducts by reacting C60 with the end group SH of PPS was easier to take place32, and the covalent bond formation improved the C60-matrix interfacial adhesion6, 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/C60 material, C-H vibration peak of PPS/C60 remain unchanged, which indicated that there only existed a weak π-π stacking interactions between C60 and PPS matrix36, 40.

Figure 5
figure 5

FT-IR spectra of C60, PPS and PPS/C60 composite powder.

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Figure 6
figure 6

The chemical functionalization process of fullerenes and PPS/C60 interfacial adhesion changes during solution and melt processing.

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In order to testify the existence of chain-extension reaction through the end-group reaction between PPS and C60, PPS and PPS/C60 molecular weight were characterized by HTGPC (Table 1). As C60 content increased from 0.5 wt% to 2 wt%, PPS/C60 composite molecular weight propagated from 3.30 × 104 g·mol−1 to 4.19 × 104 g·mol−1, and polydispersity index (PDI) increased from 2.22 to 2.56. Only some PPS chains could interact with C60 to lead to the longer polymer chains. However, when C60 content reached 5 wt%, 5 wt% PPS/C60 composite molecular weight declined to 3.71 × 104 g·mol−1. The deviation could be attributed to much more the low molecular weight C60 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 C60 through the C60-thiol adducts and C60-ArCl alkylation.

Table 1 Molecular weight and polydispersity index of different PPS.
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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% C60 had less influence on PPS crystallization temperature (Tc) increase (Fig. 7a), and 0.5 wt% PPS/C60 exhibited T c of 217.2 °C with ΔH c being 42.98 J/g. However, for 2 wt% PPS/C60, T c shifted to 227.6 °C with maximum ΔH c being 48.23 J/g. AS for 5 wt% PPS/C60, T c increases to 223.8 °C with ΔH c decline to 42.42 J/g. The further increase of C60 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/C60 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 C60–polymer chemical interactions.

Figure 7
figure 7

DSC thermograms obtained at a rate of 10 °C/min: (a) DSC cooling curves of PPS/C60; (b) DSC heating curves of PPS/C60.

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Table 2 DSC traces for PPS and PPS/C60.
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T m shifted gradually to higher temperature with increasing C60 content (See Fig. 7b and Table 2). 0.5 wt% PPS/C60 exhibited T m of 275.1 °C with ΔHm being 41.50 J/g. As for 2 wt% PPS/C60, T m shifted to 277.3 °C with maximum ΔHm being 44.10 J/g. While for 5 wt% PPS/C60, T m increased to 282.4 °C with ΔHm decline to 40.14 J/g. In the case of 2 wt% PPS/C60, 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 (Ti) at 502 °C and exhibited the maximum weight loss (Tmax) 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% C60 filler led to an improvement in the thermal stability of PPS matrix, and a maximum Ti increase (about 7 °C) was obtained at 2.0 wt% filler loading, and Tmax increment for 2.0 wt% PPS/C60 was 5.5 °C. However, the significantly decline of Ti and Tmax was detected at 10 wt% C60 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, C60 fillers are better dispersed within PPS matrix, which restricts chain mobility or diffusion to slow down the decomposition process11. Secondly, the covalent anchoring of PPS to C60 leads to a strong enhancement in the thermal conductivity that facilitates heat dissipation within the composite43. The reason for the thermal stability decline at 5 and 10 wt% C60 loading might be attributed to the appearance of C60 agglomerates.

Figure 8
figure 8

TG curves of PPS and PPS/C60).

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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 C60. The mechanical behavior of PPS/C60 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/C60 fiber. The results of mechanical property study indicated that the concentration of C60 had a greater influence on the mechanical performance of composites. As C60 nanofiller content increase, the breaking strength and tensile modulus of composites firstly increased and then decreased. When C60 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/C60 composites always decreased with increasing C60 content (Fig. 10b). The excellent mechanical properties of PPS/C60 were attributed to the heterogeneous nucleation of C60 during PPS crystallization, the formation of a large number of covalent bond by C60-thiol adducts and π-π stacking interactions between C60 surface and PPS matrix. However, the excessive addition of C60 caused a significant reduction of breaking strength, e.g., the breaking strength of 10 wt% C60 composite declined to 148 MPa. The excessive addition of C60 reduced PPS crystallization degree and decreased the combination between PPS and C60, which were attributed to the phenomenon of C60 aggregation. These results can also confirm by the SEM and TEM images of PPS/C60 composites presented in Figs 1–2.

Figure 9
figure 9

Micrographs of PPS/C60 with different weight fractions of C60 (a) PPS; (b) 0.5 wt%; (c) 1 wt%; (d) 2 wt%; (e) 5 wt%; (f) 10 wt%.

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Figure 10
figure 10

Mechanical properties of PPS/C60 fiber with different addition amount of C60: (a) breaking strength; (b) breaking elongation; (c) tensile modulus.

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PPS is an insulating material (σ 10−16 S/cm), which limits its use in self-health monitoring, electro-actuation, etc32. Herein, C60 was used as conductive fillers to improve PPS electrical conductivity and the electrical performance of PPS/C60 composites were compared (Fig. 11). As C60 nanofiller content increase, the electrical conductivity of PPS/C60 composites firstly increased and then decreased. When C60 concentration reached 2 wt%, the electrical conductivity of composites achieved maximum 1.67 × 10−2S/cm, much higher than the value of pure PPS. The excellent electrical conductivity of 2 wt% PPS/C60 composites was attributed to the well dispersed C60 fillers, and the maximum conductive networks might be formed in the composite at appropriate C60 content because of the conductive network formation44. C60 fillers are well dispersed in the PPS matrix owing to the covalent bond formation by main C60-thiol adducts and mino C60-ArCl alkylation between C60 surface and PPS. However, the excessive addition of C60 nanofillers caused a significant reduction of electrical conductivity, e.g., the electrical conductivity of 10 wt% C60 composite declined to 3.79 × 10−3S/cm, due to the phenomenon of C60 aggregation.

Figure 11
figure 11

Room temperature electrical conductivity (r) of PPS/C60 composites.

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Conclusions

PPS composites with well-dispersed C60 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/C60 composites was mainly attributed to the covalent bond formation by main C60-thiol adducts and minor C60-ArCl alkylation between C60 surface and PPS, but C60 aggregation reduced composite electrical conductivity. Furthermore, 2 wt% C60 could effectively increase PPS crystalli zation temperature, thermal stability and mechanical performance. However, the excessive C60 loading reduced the PPS crystallization degree and caused C60 re-aggregation, which led to poorer mechanical performance of PPS/C60 composite.

Table 3 Thermogravimetric data under nitrogen atmosphere.
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