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Article Cite This: ACS Omega 2019, 4, 1228−1237
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Higher-Order Structure in Amorphous Poly(ethylene terephthalate)/ Graphene Nanocomposites and Its Correlation with Bulk Mechanical Properties Shigeru Aoyama,†,§ Issam Ismail,*,†,∥ Yong Tae Park,†,⊥ Christopher W. Macosko,† and Toshiaki Ougizawa*,‡ †
ACS Omega 2019.4:1228-1237. Downloaded from pubs.acs.org by 31.40.211.195 on 01/22/19. For personal use only.
Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States ‡ Department of Materials Science and Engineering, Tokyo Institute of Technology, 2-12-1-S8-33, O-okayama, Meguro-ku, Tokyo 152-8552, Japan S Supporting Information *
ABSTRACT: Graphene of two different aspect ratios, Af, was melt mixed with poly(ethylene terephthalate) (PET) to form amorphous PET/graphene composites with less than 5% crystallinity. The higher-order structure and mechanical properties of poly(ethylene terephthalate) (PET) in these composites were investigated using techniques such as differential scanning calorimetry and dynamic mechanical analysis, whereas transmission electron microscopy, melt rheology, and electrical conductivity were used to study the graphene dispersion. A decrease in heat capacity changes, ΔCp, of PET in nanocomposites at the glass transition temperature, Tg, without Tg change suggests that a rigid amorphous fraction (RAF) of PET was formed at the PET/graphene interface. The stiffening effect of graphene below 1 wt % loading is quite small in the glassy state region and independent of the Af of graphene. Above 2 wt %, graphene forms a mechanical percolation network with the RAF of PET and the PET chains are geometrically restricted by the incorporation of graphene with a high Af, resulting in an unexpectedly higher modulus of nanocomposites both below and above Tg. properties of PET4,5 as well as its wear resistance.6 Fu et al.7 reported that the intermediate phase (i.e., RAF) enhances the tenacity of PET fiber. Rastogi et al.2 reported unexpectedly low yield stress during compression due to the formation of small crystals and because of the presence of a large RAF content. Lin et al.4 reported unexpectedly higher O2 solubility in crystalline PET due to the formation of RAF. The introduction and blending of nanoparticles as fillers is a commonly used technique to improve the mechanical and gasbarrier properties of polymers, including PET, as well as their thermal and dimensional stability. An especially potent class of nanofillers is that of two-dimensional (2D) nanoplatelets, as these have high aspect ratios, which enable them to provide a greater improvement to polymeric properties than their spherical counterparts. 8,9 Modified clay 10−19 and graphene20−33 are commonly utilized as 2D nanoplatelet fillers, with graphene recently attracting wide interest due to its combination of 2D morphology and excellent mechanical and electrical properties.20−33 Graphene has been shown to enhance the mechanical and gas-barrier properties of nano-
1. INTRODUCTION Poly(ethylene terephthalate) (PET) is a semicrystalline thermoplastic characterized by good mechanical properties and easy moldability as well as a high glass transition temperature (Tg) and chemical resistance. The preceding characteristics result in PET being a high-performance polymer that is useful in numerous industrial products in the forms of films, fibers, and bottles.1 PET films are typically produced by extrusion of PET resin, following which additional processing steps are used, such as annealing and tentering. These steps play a significant part in directing the formation of higher-order structures in PET, including chain orientation and crystallites, which results in improvements to the physicochemical film properties. The higher-order structure of solid PET is usually represented by a three-phase model, which includes a crystalline phase, a rigid amorphous fraction (RAF), and a movable amorphous fraction (MAF).2 The RAF represents the fraction of the amorphous phase that possesses partial order and is responsible for reducing the mobility of PET chains. The RAF is usually positioned between the MAF and the crystalline phase, which means that it has no contribution to the jump in heat capacity at the glass transition. These higherorder structures also affect the mechanical2,3 and gas-barrier © 2019 American Chemical Society
Received: November 25, 2018 Accepted: January 8, 2019 Published: January 15, 2019 1228
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Figure 1. Representative TEM images of (a) PET/G1 and (b) PET/G2 with 2 wt % of graphene and (c) interparticle distances, λ, of PET/G1 and PET/G2.30
nanocomposites was found to be greater than that of neat PET. In particular, PET/G2 demonstrated a higher η* than PET/G1, reflecting the higher aspect ratio of the former, as typically observed in polymer nanocomposites of plateletshaped fillers.36 At higher frequencies, the η* of the nanocomposites are decreased and are closer to that of neat PET; this is because the contribution of the polymeric matrix is more significant at such frequencies.36 In previous studies, it was demonstrated that PET/clay nanocomposites exhibit a lower η* than neat PET at higher frequencies, which was correlated to the high levels of polymer matrix degradation.36,37 However, the η*of both the PET/G1 and PET/G2 nanocomposites in this study are equal to or more than that for neat PET at all frequencies, indicating that the decomposition of PET is not enhanced by the addition of graphene and that the molecular weight of PET in all of the mentioned nanocomposites is thus similar to that of neat PET.37 The measured G′ of the PET/graphene nanocomposites further corroborates the preceding observation; contrary to the more strongly degenerating PET/clay nanocomposites,36,37 the G′ of the PET/graphene nanocomposites examined in this study approaches that of the neat PET as a lower limit at higher frequencies, never falling beneath the G′ of the neat PET (Figure S1c,d). Additionally, it may be observed that at lower frequencies, G′ rises as a function of graphene loading, becoming frequency independent above a certain critical loading, which is a typical solid-like frequency response displayed by polymeric composite melts containing 2D fillers.30 The dispersion quality of the nanocomposite samples was studied qualitatively and quantitatively using transmission electron microscopy (TEM) and melt rheology, respectively, as in our previous studies.30 Af, which is a ratio of diameter D to thickness h, was estimated for the two graphenes used in this study to be 14 for PET/G1 and 63 for PET/G2 based on the dispersion rheology of the melts.30 Representative TEM images and interparticle distances, λ, of PET/G1 and PET/G2 are shown in Figure 1 (the TEM images based on which λ was calculated are shown in Figures S2 and S3 in the Supporting Information and the detailed procedure for λ estimation is reported elsewhere30). Both G1
composites as well as their electrical conductivity and dimensional stability.25−29 It was previously shown by Zhang et al.25 that the increase in electrical conductivity of melt-mixed PET/graphene nanocomposites compared with pure PET is significantly more than that obtained via melt-mixed PET/ graphite composites. Moreover, it was shown by Feng et al.26 that PET/graphene nanocomposites produced via in situ polymerization resulted in a much larger enhancement in electrical conductivity compared with melt blending (approximately 4 orders of magnitude). Furthermore, Bandla et al.27 and Li et al.28 blended graphene with PET and observed an increase in the Young’s modulus and dynamic storage modulus of the resulting nanocomposites. However, all of the preceding studies only discussed the properties of the mentioned composites (particularly, mechanical properties) from the perspective of the graphene-in-PET dispersibility. There have been no studies investigating the higher-order structure of PET in PET/graphene nanocomposites. The objective of this investigation is to study the higherorder structure in PET/graphene nanocomposites and its effect on their mechanical properties. The study was conducted by selecting two graphenes, each with its own aspect ratio, Af, and melt mixing them as nanofillers with PET. The higherorder structure of PET in the resulting nanocomposites was investigated by differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA), and its effect on the mechanical properties of PET/graphene nanocomposites with different graphene loading was then discussed.
2. RESULTS AND DISCUSSION 2.1. Rheological Properties and Dispersion Level of PET-Based Nanocomposites. Melt rheology was used to characterize the nanocomposites in terms of (1) the level of graphene-in-PET dispersibility and (2) the degradation level of the PET matrix, if any, as a result of graphene blending. The complex viscosity, η*, and storage modulus, G′, of the nanocomposite melt are demonstrated in Figure S1. The η* of the neat PET melt exhibits a pseudo-Newtonian behavior, whereas those of both nanocomposites showed significant shear thinning, as typically exhibited by such nanocomposites (Figure S1a,b).36,37 At low frequencies, the η* of the 1229
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preceding equation and using Φper = 0.023 (4.0 wt %) for PET/G1 and Φper = 0.003 (0.62 wt %) for PET/G2, Af values for PET/G1 and PET/G2 were estimated to be 20 and 128, respectively (summary in Table 1). This indicates that G2 can form an electrical percolation network more easily than G1.
(the lower aspect ratio and thicker stacks) and G2 (the higher aspect ratio and thinner stacks) appear to have dispersed well in the PET matrix, even though some G1 aggregated stacks can be observed in the images. G2 dispersed in the composites in the form of few-layered stacks, which appear to be aligned in roughly the same direction, with the sheet edges twisted and wrinkled. The λ of PET/G2 was measured to be approximately ≤1 μm for the lower graphene loadings and became ∼100 nm for loadings greater than 0.5 wt %. It was also observed that λ in PET/G1 was greater than that for PET/G2, given the similar graphene loadings. Presumably due to its lower aspect ratio, it took twice as much loading of G1 (compared with G2) for its λ in the PET/G1 composites to match that of G2 in the PET/G2 composites.30 2.2. Electrical Properties. The electrical resistance of the PET/graphene composites was measured as a corroborative method of studying the graphene dispersion level. The surface resistances of the PET/graphene composites are shown in Figure 2. The melt blending of both G1 and G2 was observed
Table 1. Aspect Ratios of Nanofillers in PET-Based Nanocomposites from Various Characterization Methods PET/G1 PET/G2
electrical resistance
14 63
20 128
It is worthy to mention that the electrical percolation measurements consistently seem to overestimate the aspect ratios of G1 and G2 compared with the rheological percolation measurements, as seen in Table 1. This is thought to be because electrical and rheological percolation are phenomenologically different; electrical percolation occurs when free electrons become capable of conducting electric current by hopping between neighboring (adjacent) graphene sheets due to proximity. On the other hand, rheological percolation is the result of a combined effect of (1) molecular restriction of the matrix polymer by the platelet graphene fillers, which results in chain confinement effects, and (2) the interaction between neighboring graphene sheets. This chain confinement effect, absent from the picture in electrical percolation, is thought to be particularly strong in PET/graphene nanocomposites such as our own, owing to the relatively large aspect ratio of graphene sheets, which are effective at chain confinement due to the large sheer ratio of the sheet surface area to chain length. The phenomenological difference between the two mechanisms is thus believed to be responsible for the discrepancy between the two sets of aspect ratio estimates. 2.3. Thermal Properties and Higher-Order Structure of PET in Nanocomposites. DSC was measured on the nanocomposites to study the higher-order structure of PET. The DSC curves for the first heating scan are depicted in Figure 3 (complete profiles and Tg of nanocomposites are depicted in Figures S4 and S5 in the Supporting Information). Figure 4 shows the higher-order structures such as crystallinity and the RAF contents of PET in the amorphous PET/ graphene nanocomposites per the three-phase model. The first
Figure 2. Electrical resistances of PET/G1 and PET/G2 nanocomposites. Dashed lines are exponential approximation lines.
to decrease the nanocomposite surface resistance significantly below that of neat PET, with the mentioned reduction in surface resistance increasing as a function of graphene loading. The electrical percolation threshold Φper of G2 in PET/G2 composites was lower than that for G1 in PET/G1 composites, reflecting that G2 has a higher aspect ratio than G1 in the PET matrix, which is consistent both with earlier observations from rheology and with the qualitative TEM image evidence, and is moreover consistent with characterization results from our previous work on the same.30 To convert the weights of the graphenes and PET to volumes and enable the calculation of volume fractions, the density of graphene was considered as 2.28 g cm−3 (as typically reported in the literature),23,24,30 whereas 1.335 g cm−3 was taken as the density for pure PET.30,38 From exponential approximation lines of reduced surface resistance area, we estimate Φper of graphene in PET to be 0.023 (4.0 wt %) for G1 and 0.003 (0.62 wt %) for G2. Based on the estimated percolation thresholds and the assumption of graphene sheets being monodisperse 2D circular disks, Af was estimated using eq 117,39−41 A f = D/h = 3Φsphere /2Φper
melt rheology30
(1)
where Φsphere = 0.29, which is the onset of percolation for interpenetrating, randomly packed spheres.42 Based on the
Figure 3. Representative DSC profiles of neat PET and PET/ graphene nanocomposites (first heating scan at 10 °C min−1). 1230
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Figure 4. (a) Crystallinity and (b) RAF contents of PET/graphene nanocomposites.
heating scan, starting from room temperature, showed the following for all nanocomposite samples: (1) stepwise endothermic changes for glass transition, (2) exothermic peaks for cold crystallization, and (3) endothermic peaks for crystal melting. The peak heat flow for cold crystallization of all nanocomposites shifted to lower temperatures compared with that of neat PET, indicating that graphene plays the role of a nucleating agent for crystallization in all nanocomposite samples. Furthermore, the peak shapes of PET-based nanocomposites became wider, indicating that the nanocomposite cold crystallization process takes a longer time to complete than for neat PET.43 In spite of the variability observed across samples in terms of the cold crystallization behavior, it was observed for all of the samples that the crystallinity level is less than 5% (meaning that these samples are all nearly amorphous, as depicted in Figure 4a). Since the RAF of PET forms when the volume fraction of crystallinity is above 0.07,4 the amount of RAF in the amorphous neat PET can be regarded as zero. No significant Tg change (Figure S5 in Supporting Information) is observed in any sample, which is a typical phenomenon of nanocomposites produced by melt blending, indicating no covalent bonding or strong hydrogen bonding in our systems.44 Despite the crystallinity of all samples being similar and quite small, ΔCp of nanocomposites at Tg became smaller than that for neat PET (Figure 3), suggesting that graphene induces the formation of RAF. The RAF contents in PET/graphene nanocomposites increased with an increase in graphene concentration (Figure 4b; i.e., MAF concentrations are decreased (Figure S6), and, notably, G2 with a higher aspect ratio created more RAF than G1 for the same loading. Ma et.al.34 reported that the RAF in amorphous PET/SiO2 nanocomposites was caused by the restriction of the mobility of PET chains by the interaction between PET and SiO2. The interfacial area per unit volume, (Sv)p−c, of graphene in the PET/graphene nanocomposites was measured from the TEM images using Basu’s method, which assumes that the 2D graphene nanofillers can all be morphologically approximated as circular disks.30,45 The relationship between (Sv)p−c and RAF contents is shown in Figure 5. The RAF content increased in the PET/graphene nanocomposites with an increase in (Sv)p−c. The fact that PET/G1 and PET/G2 show the same dependence on area indicates that RAF is indeed controlled by surface interactions. In our previous study,30 the PET/G1 and PET/G2 nanocomposites were studied using Raman spectroscopy and the peak locations of the ring C−C band in the spectra were shown to be shifted to lower wavenumbers, which increased with the graphene
Figure 5. Relationship between the interfacial area per unit volume, (Sv)p−c, and RAF contents of PET/graphene nanocomposites. The detailed procedure for analysis of (Sv)p−c is reported elsewhere.30
loading used in the nanocomposites, with PET/G2 experiencing a greater shift than PET/G1 (Figure S7). The latter observations suggest that PET can interact with graphene by the aromatic PET/graphene interactions. Therefore, it can be inferred that the RAF of PET in nanocomposites is mainly formed at the PET/graphene interface and that the PET chain interfacial mobility is hindered even in the amorphous state.2 Normally, the Tg of PET increases with an increase in the crystallinity due to the mobility restriction of PET chains by the formation of RAF at the interface between crystals and the MAF.4 However, no such Tg change is observed in our systems. Therefore, the RAF in the nanocomposites is fundamentally different from the RAF in crystallized PET. Molecular chains of PET in nanocomposites are confined by π−π interactions between graphene and PET, but the restriction and confinement of these chains by the addition of graphene is relatively weaker than for chains in crystallized PET. 2.4. Glass Transition Behavior and Mechanical Properties. DMA was used to measure the mechanical properties of the PET/graphene nanocomposites. The dependencies of the dynamic storage moduli, E′, and tan δ of the nanocomposites on temperature are shown in Figure 6. The E′ values of all samples were relatively constant below 70 °C (the glassy state region), and then dropped significantly over the glass transition range, and finally increased due to cold crystallization above 100 °C (Figure 6a,b). The peak of tan δ occurs due to the chain-segmental relaxation phenomena taking place during the glass−rubber 1231
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Figure 6. Dynamic storage moduli, E′, for (a) PET/G1 and (b) PET/G2 and tan δ for (c) PET/G1 and (d) PET/G2. Measurements were conducted at 1 Hz at 2 °C min−1.
amorphous regions (i.e., MAF). On the other hand, the peak values of tan δ of PET/G2 decreased with an increase in graphene loading even though E″ was increased. Our previous study30 suggests that narrower interparticle distances of G2 in PET/G2 cause a stronger restriction of PET chains than G1 by geometrical confinement. Considering that RAF is mainly formed in the interfacial region between PET and graphene, the decrease in the tan δ of PET/G2 accompanied with the increase in E″ indicates that G2 in the PET matrix, especially around 2 wt %, restricts the molecular movement of the amorphous PET chains both in the interfacial region (i.e., in RAF) and in the amorphous region (i.e., MAF).46 The moduli at 40 °C (glassy state) and 100 °C (rubbery state) were normalized by the moduli of neat PET and plotted against the nanofiller loading (Figure 8). At 40 °C, PET exhibits a glassy behavior and so the moduli values at this temperature are significantly greater than at 100 °C. The addition of G1 increased the moduli of PET slightly and nearly linearly. The E′ of PET/G1 with 2 wt % of graphene was only 4% higher than that of neat PET at 40 °C. However, a different reinforcement effect was exhibited for G2. Below 1 wt % of graphene loading, the addition of G2 caused no significant difference in the trend of E′ compared with G1, whereas an unexpectedly higher modulus of PET/G2 was seen at 2 wt % G2, with the E′ of PET/G2 nanocomposites being 21% more than that of PET at 40 °C. At 100 °C, PET displays a rubbery behavior, thus leading to a larger reinforcement effect being observed in nanocomposites at this temperature compared with 40 °C. The loading of G1 to the PET by melt blending increased the modulus of PET nearly linearly. The E′ of PET/G1 with 2 wt % of graphene was 20% more than that of PET at 100 °C. However, the modulus of PET/G2 below 1 wt % graphene loading became
transition in nanocomposites of PET, and these peak positions in DMA are normally taken to correspond to Tg. The DMA results (Figure 6) also show that the Tg values for the PET/ graphene composites remain nearly constant despite the incorporation of graphene, which is consistent with the Tg trends obtained via DSC (see Figure S5 in the Supporting information). Tan δ peak values are plotted against the nanofiller loading in Figure 7 and the E′ and E″ at the peak
Figure 7. Tan δ peak values of PET/graphene nanocomposites. Measurements were conducted at 1 Hz at 10 °C min−1. Error bars of PET/G1 are smaller than the size of points.
tan δ are plotted in Figure S8. Although the RAF content increased with an increase in graphene loading (Figures 4b and 5), the peak values of tan δ of PET/G1 with a smaller aspect ratio remained nearly constant. Both E′ and E″ of PET/G1 also remained constant (Figure S8). This indicates that the restriction of molecular chains of PET/G1 is limited to the interfacial region (i.e., in RAF) and does not affect other 1232
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Figure 8. Normalized dynamic storage moduli, E′/E0, of PET/graphene nanocomposites at (a) 40 °C and (b) 100 °C.
Figure 9. Suggested structures of PET/G1 and PET/G2 at (a) below 1% and (b) 2 wt % of graphene loading.
composites are influenced by the morphology and physicochemical properties of the nanofillers, the evenness of nanofiller dispersion throughout the nanocomposite, and the matrix−nanofiller interfacial adhesion. Our previous studies suggest that G2 in PET is wrinkled and twisted30 and that the PET/graphene interfacial adhesion is relatively lower than for other composites in the glassy state.50 Therefore, the lower stress transfer efficiency by the lower PET/graphene interfacial adhesion in the glassy state,50 coupled with the lower alignment of fillers compared with the ideal structure, both contribute to low stiffening effects for G2 below 1% in the glassy state region. Above Tg, PET softens and the matrix modulus becomes much lower than that of the filler modulus. Additionally, the interfacial adhesion in both PET/graphene and PET/clay
higher than that of PET/G1, and a remarkable enhancement of E′ was observed at 2 wt % of graphene loading. The E′ of PET/G2 with 2 wt % of graphene was 370% greater than that of PET at 100 °C. Di Lorenzo et al.47 reported that the overall rigid fraction (i.e., the total of crystallinity and RAF contents, Xc + Xra) correlates with the moduli of polymeric materials in the glassy state and that the confinement effect enhances the modulus of the PET matrix.34 However, there is no significant difference between G1 and G2 below 1% at 40 °C even though the Af of G2 and the RAF contents of PET/G2 are higher than those for PET/G1. This lowered stiffening effect below 1% is partially due to the fact that the reinforcing effect decreases when the value of the matrix modulus approaches that of the material used as a filler.23,48,49 Additionally, the properties of nano1233
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Scheme 1. Representative Processing of the Thin Amorphous Nanocomposites Sample
stiffness enhancement effects caused by the addition of the two graphenes are dictated by simple mixing rules. However, below Tg, the matrix/filler interfacial adhesion grows weaker in the glassy state with the result being a drop in the stress transfer efficiency.52 The reduced stress transfer efficiency, in turn, means that there is a smaller increase in the modulus, which is independent of Af. On the other hand, above Tg, the difference in the filler size significantly affects the mentioned reinforcement effects due to the improved stress transfer efficiency resulting from the enhanced interfacial adhesion that occurs in the rubbery state. Therefore, the PET/G2 composites (which have a higher Af) showed higher moduli than PET/G1 composites in the rubbery state region. Above 2 wt %, the restriction of molecular chains of PET/ G1 is still limited to the interfacial region (i.e., in RAF) mainly due to wider interparticle distances than in PET/G2 as shown in Figure 9b. Therefore, the stiffness effects of G1 are still dictated by simple mixing rules, resulting in the modulus of PET/G1 increasing slightly and nearly linearly at a graphene loading of 2 wt %. On the other hand, the formation of hybrid mechanical percolation networks of G2 with the RAF of PET geometrically restricts the mobility of PET chains in PET/G2 due to the higher aspect ratio and the narrower interparticle distance. Integrated confinement effects of RAF and the geometrical restriction enhance the modulus of the PET matrix.34,47 Therefore, the combined effects of the mechanical percolation network of graphene with the RAF of PET at the interfaces and the confined PET matrix cause a drastic enhancement of the moduli of PET/G2 both below and above Tg.
nanocomposites is markedly stronger than in the glassy region.50 This is most likely due to the PET viscoelasticity in the rubbery state: a commonly observed phenomenon in pressure-sensitive adhesives above Tg.51 The big difference in the moduli between the polymer and the nanofiller coupled with the enhanced interfacial adhesion above Tg causes a larger reinforcement effect in the nanocomposites, resulting in PET/ G2 composites with a higher Af displaying higher moduli than the PET/G1 composites in the rubbery state. Above 1 wt % of graphene loading, the unexpectedly higher modulus of PET/G2 with 2 wt % of graphene at both 40 and 100 °C may be justified by the formation of rigid filler networks.52,53 As mentioned above, the RAF contents in PET/ G2 with 2 wt % graphene increased up to 22% and formed at interfacial regions. Our previous study suggests that PET chains in PET/G2 are more confined compared with PET/ G1.30 Therefore, the unexpectedly higher modulus of PET/G2 with 2 wt % of graphene indicates the formation of mechanical percolation networks of graphene with the RAF of PET at interfacial regions with confined PET matrixes. This mechanical percolation threshold was estimated between 1 and 2 wt %, which matches with the percolation thresholds of the molten dispersions (∼1.11 wt %) from our previous study.30 2.5. Internal Structure and Its Correlation with Bulk Mechanical Properties. The preceding results led us to form the following summative view of the mechanical properties of PET/graphene nanocomposites (as summarized in Figure 9): the RAF of PET in nanocomposites is formed at the PET/ graphene interface by the aromatic interactions between the matrix and the filler, both of which contain a π−π conjugation. Molecular chains of PET in nanocomposites are clearly confined compared with neat PET, but the restriction of these chains by the addition of graphene is relatively weaker than in crystallized PET. Below 1 wt % of graphene loading, the restriction of PET chains is mostly limited to the interfacial region (i.e., RAF) as shown in Figure 9a. Therefore, the
3. CONCLUSIONS PET was melt mixed with two graphene nanofillers (G1 and G2), each with its own aspect ratio. From the electrical properties, the percolation threshold of G2 in PET/G2 composites was estimated to be lower than that of G1 in PET/G1 composites, indicating that G2 has a larger aspect 1234
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(ARES, TA Instruments) at 270 °C under an N2 blanket. Detailed rheological measurement parameters are reported elsewhere.30 4.4. TEM Imaging and Interparticle Distance Measurement. TEM images were acquired using a Tecnai T12 TEM from FEI using 120 kV as the accelerating voltage. Detailed sample preparation for TEM and imaging conditions and parameters are reported elsewhere.30 4.5. DSC Measurement. DSC for the samples was measured on a Q1000 device from TA Instruments under an inert N2 blanket. The amorphous sheet samples were initially heated to 280 °C at 10 °C min−1 and held at 280 °C for 1 min (first heating scan). To enable the investigation of the higherorder structure, the glass transition parameters (i.e., glass transition temperature, Tg, and heat capacity change, ΔCp), cold crystallization parameters (i.e., heat of fusion, ΔHcc), and crystal melting parameters (i.e., melting temperature, Tm, and heat of fusion, ΔHm) were obtained from the first heating scan. Because the RAF of PET typically forms when the volume fraction of crystallinity is more than 0.07,4 we assume that amorphous neat PET has an RAF of 100%,4 which means that the change of ΔCp in nanocomposites is assumed to reflect the change of MAF, Xma, with the incorporation of graphene. The Xma in nanocomposites was calculated by eq 234,35
ratio than G1 in the PET matrix, which is consistent with observations from melt rheology. From the DSC study, PET in amorphous nanocomposites showed a decreased heat capacity at Tg with increasing graphene concentration, suggesting the presence of a significant RAF of PET in the mentioned composites even in the amorphous state. The RAF contents in the nanocomposites also increased with an increase in the graphene loading and correlated with the interfacial area between PET and graphene, but no Tg change was observed. The preceding observations indicate that the RAF is primarily formed in the interfacial region between PET and graphene and that the restriction effect on these chains resulting from the addition of graphene is relatively weaker than in crystallized PET. From the DMA study, below 1 wt % of graphene loading, the stiffening effects of G1 and G2 were observed to be quite small and showed no significant differences in the glassy region due to the wrinkled structure and the nonperfect alignment of the graphene in PET as well as the lower PET/graphene interfacial adhesion. On the other hand, a significant reinforcement effect was observed above Tg due to the enhanced interfacial adhesion. Moreover, PET/G2 composites showed higher moduli than PET/G1 composites in the rubbery state region, an observation that was attributed to the higher Af of G2 in the former. Above 2 wt % loading, hybrid mechanical percolation networks of G2 and RAF of PET suppress the molecular movements of PET chains more effectively than G1 because of the larger specific interfacial area and narrower interparticle distance of the former. The combined effect of the hybrid mechanical percolation network and confined PET matrix drastically enhances the modulus of PET/G2 both below and above Tg.
X ma = ΔCp,c/ΔCp,0 × 1/(1 − w) × 100
(2)
where ΔCp,0 is the ΔCp for amorphous neat PET, ΔCp,c is the ΔCp for nanocomposite samples, and w is the weight fraction of the nanofillers. The percent amount of RAF, Xra, in the nanocomposites was calculated by eq 334,35 X ra = 100 − Xc − X ma
4. EXPERIMENTAL SECTION 4.1. Materials. PET pellets were obtained from Toray Plastics America (North Kingstown, RI). These pellets were crushed into a powder of particle size 30 mesh using a Polyvision instrument (Manchester, PA). Following the grinding, the intrinsic viscosity of the PET was found to be 0.61 dL g−1 using ortho-chlorophenol. As nanofillers, two batches of graphene (G1 and G2) of different aspect ratios were used as received. G1 was obtained from XG Science (Product code: xGnP-C750, multilayer, diameter < 2 μm), whereas G2 was obtained from Angstron Materials (Product code: N002-PDR, 12 h prior to melt mixing. 4.2. Preparation of Nanocomposites. The PET and graphene were melt mixed to produce nanocomposites according to a previously published method described in Scheme 1.30 PET/graphene mixtures of 5.5 g were fed into a recirculating twin-screw extruder (Microcompounder, DACA Instruments) using different graphene loadings (0−12 wt % for G1 and 0−2 wt % for G2) at 280 °C using N2 as the purge gas. The PET and graphene were blended at 360 rpm for 8 min and cooled afterward via extrusion into an ice/water bath. The nanocomposites produced by this method were dried in vacuo at 120 °C for >12 h. Thin sections (120−180 μm) were prepared for characterization (TEM, DSC, and DMA) by pressing the mentioned samples between fiber-reinforced Teflon sheets at 270 °C and 1−1.5 MPa for 2 min followed by rapid quenching in ice water to prevent crystallization. 4.3. Melt Rheology. The rheology of the nanocomposites was investigated using a strain-controlled rotational rheometer
(3)
4.6. Electrical Properties. The surface resistance of the nanocomposite films was measured using an 11-probe DC resistance meter (Prostat-801). 4.7. DMA Measurement. DMA for the samples was measured on a Rheometrics Solid Analyzer, RSA-G2 (TA Instruments, New Castle, DE). The specimens for DMA measurements were sectioned from nanocomposite sheets into strips of 6 mm × 40 mm of rectangular shape. Temperature ramp measurements were performed from 25 to 180 °C at 2 °C min−1 and a frequency of 1 Hz. 4.8. Raman Spectroscopy. Raman measurements on the samples were obtained using an α 300R confocal Raman microscope equipped with a UHTS200 spectrometer and a DV401 charged-coupled device detector from WITec (Ulm, Germany). Detailed measurement conditions are reported elsewhere.30
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsomega.8b03280. TEM images of amorphous PET/graphene nanocomposites showing dispersion levels from our previous study, DSC profiles for the composites, Tg of the mentioned composites as a function of filler content as obtained by DSC and DMA, mobile amorphous contents of composites as a function of filler content, and dynamic storage and loss moduli and tan δ of composites as obtained by DMA (PDF) 1235
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected]. Tel: +971-2-607-5835. Fax: +971-2-607-5200 (I.I.). *E-mail:
[email protected]. Tel: +81 3 5734 2423. Fax: +81 3 5374 2423 (T.O.). ORCID
Issam Ismail: 0000-0002-2574-4095 Yong Tae Park: 0000-0001-9296-6790 Christopher W. Macosko: 0000-0002-2892-3267 Present Addresses ⊥
Department of Mechanical Engineering, Myongji University, 1st Engineering Bldg., 116 Myongji-ro, Cheoin-gu, Yongin, Gyeonggi-do 449-728, Korea (Y.T.P.). ∥ Khalifa University of Science and Technology, Sas Al Nakhl Campus, Abu Dhabi 127788, UAE (I.I.). § Films & Film Products Research Laboratories, Toray Industries, Inc., 1-1-1 Sonoyama, Otsu, Shiga 520-8558, Japan (S.A.). Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This research was supported by Toray Industries through their membership in the University of Minnesota Industrial Partnership for Research in Interfacial and Materials Engineering (IPRIME). Parts of this work were performed in the University of Minnesota Characterization Facility, which received partial support from the NSF through the MRSEC program.
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