Development of Advanced Elastomeric Conductive Nanocomposites

Article pubs.acs.org/Macromolecules

Development of Advanced Elastomeric Conductive Nanocomposites by Selective Chemical Affinity of Modified Graphene Horacio J. Salavagione,*,† Susana Quiles-Díaz,† Patricia Enrique-Jimenez,‡ Gerardo Martínez,† Fernando Ania,‡ Araceli Flores,‡ and Marián A. Gómez-Fatou† †

Departamento de Física de Polímeros, Elastómeros y Aplicaciones Energéticas, Instituto de Ciencia y Tecnología de Polímeros (ICTP-CSIC), Juan de la Cierva 3, 28006 Madrid, Spain ‡ Departamento de Física Macromolecular, Instituto de Estructura de la Materia (IEM-CSIC), Serrano 119, 28006 Madrid, Spain S Supporting Information *

ABSTRACT: The preparation of conductive nanocomposites of poly(styrene-b-ethylene-co-butylene-b-styrene) (SEBS) triblock copolymers with chemically modified graphene is reported. Graphene has been functionalized with short polyethylene brushes by thiol−ene click chemistry in order to enhance the interactions of the filler with the ethylenic phase of the block copolymer. The presence of graphene modifies the morphology of the SEBS domains, reducing their size and promoting an isotropic distribution of cylinders arranged in hexagonal packing, as shown by SAXS, AFM, and TEM. The elastomeric character of SEBS persists in the nanocomposites with filler content below 1 vol %, retaining high elongation at break. In addition, the percolation threshold for electrical conductivity is found to be slightly above 0.7 vol %, reaching conductivity values of ∼10−3 S cm−1 for ∼2.5 vol % of filler. These materials represent a new class of flexible conductors that open promising opportunities for applications like actuators or bodily motion sensors.

■

INTRODUCTION Among the wide pool of applications of graphene-based materials, polymeric nanocomposites constitute one of the areas of major growth and short-term impact on real applications.1−5 The nature of the polymeric matrices that can be reinforced with graphene and the strategies used to incorporate this filler are numerous and diverse.1−5 The key to obtain superior materials lies in the efficient functionalization of graphene to achieve good dispersion and strong interphases between the components, the covalent graphene−polymer coupling being one of the most successful routes.4,6−8 Alternatively, the functionalization of graphene with short polymer brushes to improve its dispersion on similar polymeric matrices represents another useful strategy.9−12 Although to a lesser extent, this approach has also been employed with copolymers very different in nature.13−22 The combination of elastomeric matrices with graphene is very exciting since electromechanical materials find niche applications in crucial fields like flexible electronics,23 actuators,24,25 body motion sensing,26,27 and monitoring structural deterioration in materials.28 Thermoplastic elastomers like poly(styrene-b-ethylene-co-butylene-b-styrene) (SEBS) triblock copolymers constitute a special family since they display the physical properties of rubbers but can also be processed as thermoplastics. In spite of their potential, investigation on the incorporation of graphene into a SEBS matrix is limited to a handful of examples, and advances to date are rather modest.21,29−31 In particular, to the best of our © XXXX American Chemical Society

knowledge, conductivity values in the range required for practical applications have not been achieved. The styrenic phase has been commonly used as the target to incorporate graphene in SEBS, aiming at the establishment of π−π graphene/styrene interactions.21,28,31 However, this route has only yielded improvements in mechanical properties.21,29,32 These studies have shown that the conductive material (graphene) exhibits more affinity for domains with higher electron density (styrenic) that favor electron mobility while the electron-deficient domains (ethylenic phase) remain unaltered hindering electron transport. Hence, promoting the dispersion of graphene in the electron-poor domains could substantially increase the mobility of carriers with the consequent impact on the electrical conductivity. Consequently, directing graphene toward the ethylenic phase of SEBS seems to be a promising strategy to prepare nanocomposites with high conductivity values using a small amount of filler. Other nanofillers like multiwalled carbon nanotubes (MWCNT), montmorillonite (MMT), and silver nanoparticles have been incorporated into block copolymers, and changes in the morphology and phase distribution were induced.33−36 In the former case, MWCNT were selectively located in ABC block terpolymers of styrene, butadiene, and methyl methReceived: March 8, 2016 Revised: June 6, 2016

A

DOI: 10.1021/acs.macromol.6b00490 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules acrylate,33,34 while in the second example MMT modified with polymer brushes was incorporated to poly(styrene-block-tertbutyl acrylate).35 However, these copolymers are very different from SEBS because they contain polar acrylate blocks and unsaturated groups in the “hydrophobic” blocks with higher electron density than SEBS. Therefore, conferring electrical conductivity to SEBS constitutes a significant advance. Recently, we addressed the study of nanocomposites of graphene and high-density polyethylene (HDPE), employing a route based on a thiol−ene click reaction between graphene and short-chain polyethylene (PE) followed by two-step processing with HDPE.37,38 The final materials displayed very low percolation threshold and high absolute dc conductivity values. The same strategy can be extended to a wide range of polymers having similar structures. In this work we report the preparation of nanocomposites of SEBS and graphene by using graphene modified with short polyethylene brushes as filler. Considering the composition of SEBS, on the basis of styrenic and ethylenic domains, we have used the PE-modified graphene to specifically enhance the interactions with the ethylenic phase. As a result of the controlled chemical affinity of graphene, the nanocomposites with small amounts of graphene display good electrical conductivity, while the elastomeric character is maintained.

■

Table 1. Details of the Different Nanocomposites Studied sample

GPE (wt %) (graphene content)a

SEBS−GPE1 SEBS−GPE2 SEBS−GPE3 SEBS−GPE4 SEBS−GPE5 a

1 1.5 2 5 10

(0.8) (1.0) (1.6) (4.2) (9.1)

GPE (vol %) 0.5 0.7 1.0 2.5 5.2

Measured by TGA after heating up to 800 °C.

to be ≈135 × 1200 μm2. A 2D SAXS detector was placed at a distance of 6020 mm from the sample that was calibrated against silver behenate. Portions of the films were positioned perpendicularly to the incident beam. A Linkam hot stage (THMS600) was used to heat the samples above room temperature. The 2D SAXS patterns were analyzed with the help of the fit2D program.39 Profiles of the scattered intensity, I, as a function of the wave vector q (q = 4π sin θ/λ) were obtained by integrating I over a small portion of the azimuthal angle. Atomic force microscopy (AFM) measurements of thin films prepared by drop-casting from o-DCB solutions were conducted on a Multimode Nanoscope IVa, Digital Instrument/Veeco operated in tapping mode under ambient conditions. Silicon tapping mode cantilevers with a spring constant of 40 N.m−1 and a resonance frequency of 300 kHz were used. The morphology of the nanocomposites was also examined by transmission electron microscopy (TEM). Images were obtained with a Philips Tecnai 20 microscope. Ultrathin sections, 50−100 nm in thickness, were cryogenically microtomed with a diamond knife at ∼−60 °C and supported on copper TEM grids. Some of these sections were used without further preparation to evaluate the graphene dispersion in the matrix, while others were stained with ruthenium tetroxide to visualize the phase morphology of the block copolymers. Thermogravimetric analysis was carried out using a TA Instruments Q50 thermobalance in the range of 50−800 °C at a heating rate of 10 °C min−1, under an inert atmosphere (nitrogen, 60 cm3 min−1). The dc-conductivity measurements were carried out on films dried under vacuum for 24 h. The measurements were made using a fourprobe setup equipped with a dc low-current source (LCS-02) and a digital microvoltmeter (DMV-001) from Scientific Equipment & Services. The temperature of the samples was regulated with an accuracy of ±0.5 °C using a PID controlled oven from Scientific Equipment & Services. The conductivity σ was calculated as the inverse of the resistivity ρ:

EXPERIMENTAL SECTION

The SEBS employed was provided by Dynasol (Madrid, Spain). This elastomer contains 30 wt % of styrene units and the following molecular weight characteristics as determined by GPC: weightaverage molecular weight, Mw = 85 000 g/mol; polydispersity index, Mw/Mn = 1.45, where Mn is the number-average molecular weight. Graphene in powder form (N002-PDR) was supplied by Anstrong Materials Inc. (USA). Modification of Graphene. Graphene was click-functionalized according to a previously reported procedure.37 Briefly, 0.5 g of PE terminated with −SH groups was dissolved in 50 mL of anhydrous oDCB, under a nitrogen atmosphere. Then, a thermal initiator (AIBN 1.72 g, 0.01 mol) and graphene (0.5 g) were added to the polymer solution. The mixture was heated at 70 °C and stirred overnight. The solid product was collected by filtration and washed with abundant amounts of methanol and hot toluene to remove excess initiator and nonreacted polymer, respectively. Soxhlet extraction in hot toluene was also conducted in order to remove free polymer; however, no differences were observed after the extraction. The modified-graphene product was denominated GPE. The actual composition of GPE was estimated by TGA, where the mass loss due to the elimination of the PE brushes corresponds to 18%. Hence, GPE contains 82 wt % of graphene. Preparation of Nanocomposites. The nanocomposites were prepared by mixing in o-DCB appropriate amounts of SEBS and GPE to obtain nanocomposites with 1, 1.5, 2.0, 5, and 10 wt % of GPE. A nanocomposite with 2.0 wt % of unmodified graphene was also prepared for comparison (SEBS-G). The volume fraction of GPE was calculated assuming density values of 0.94 and 1.91 g cm−3 for SEBS and GPE, respectively. The latter density was estimated considering the composition of GPE together with the density of the starting graphene (∼2.2 g cm−3) and that of the modifier PE brushes (0.985 g cm−3), both of them provided by the suppliers. The estimated volume fractions are shown in Table 1. Characterization. The dispersion of the fillers in the SEBS matrix was examined by scanning electron microscopy (SEM) using a SU8000 Hitachi scanning electron microscope. The nanocomposite samples were cryofractured from drop-casted films. Small-angle X-ray scattering (SAXS) patterns were obtained using synchrotron radiation at the Non-Crystalline Diffraction beamline (BL11-NCD) of ALBA (Barcelona, Spain). A wavelength of λ = 0.0999 nm was used. The spot size at the sample location is estimated

σ=

1 ρ

(1)

where

⎛V ⎞ ρ = 4.5324t ⎜ ⎟f1 f2 ⎝I⎠

(2)

with t being the thickness of the sample, I the applied current, V the measured voltage, f1 the finite thickness correction for thick samples on an insulating bottom boundary, and f 2 the finite width correction. Tensile properties of the composites were measured with an Instron 4204 tensile tester at room temperature and 50 ± 5% relative humidity, using a crosshead speed of 50 mm min−1 and a load cell of 100 N. Five specimens for each type of composite were tested to ensure reproducibility.

■

RESULTS AND DISCUSSION The important influence of the inclusion of nanofillers on the microstructure of block copolymers has been previously highlighted.33−36,40 Controlling the distribution of these nanofillers can provide a versatile tool to tailor their properties. It is well-known that the triblock copolymer SEBS undergoes microphase separation into self-assembled ordered morphologies. Therefore, the incorporation of nanoparticles, in particular B


Article

Macromolecules

surfaces of nanocomposites with 1.0 vol % of modified (SEBS− GPE3) and unmodified graphene (SEBS−G) are compared in Figure 2. The influence of the covalent attachment of the

graphene, and its effect on the polymer properties need to be analyzed considering the dissimilar chemical nature of each domain. Unlike homopolymers, in this case the key issue to be considered is the specific location of the filler in the matrix and the specific interactions with each domain. It is not clear whether graphene has a preference for any particular domain or whether this is conditioned by the graphene loading. Furthermore, the chemical functionalization of graphene may change the affinity for a given domain. In this study we have investigated the effect of modified graphene on the nano/ microscale morphology and on the structure of the block copolymer and evaluated its impact on selected bulk properties of the nanocomposites. In order to investigate the degree of dispersion of the nanofiller in the polymer matrix, relevant to the optimization of the properties, cryofractured surfaces of all samples were analyzed by SEM. Figure 1 shows the SEM images of SEBS/

Figure 2. SEM images of fractured surface of SEBS/GPE nanocomposites with 2 wt % unmodified graphene (SEBS−G) (A) and 2 wt % (1 vol %) modified graphene (SEBS−GPE3) (B). Both images are collected using the same magnification (scale bar corresponds to 5 μm).

polyethylene brushes on the graphene dispersion can be clearly observed in the SEM images. Agglomerates were found in composites with unmodified graphene, whereas a more homogeneous dispersion was displayed in the nanocomposite with GPE. The influence of modified graphene on the microdomain structure of SEBS was investigated by AFM. Figure 3 shows AFM phase images of SEBS and SEBS/GPE nanocomposites. The microphase separation of the triblock copolymer was clearly observed in neat SEBS, showing cylindrical domains of

Figure 1. SEM images of the fractured surfaces of SEBS and SEBS/ GPE nanocomposites: SEBS (A), SEBS−GPE1 (B), SEBS−GPE2 (C), SEBS−GPE3 (D), SEBS−GPE4 (E), SEBS−GPE5 (F). All images are collected using the same magnification (scale bar corresponds to 5 μm).

GPE nanocomposites with different modified-graphene content. At low filler content (below 1.0 vol % GPE), a homogeneous graphene distribution is observed. At filler contents of 1.0 vol % GPE or higher, the formation of graphene islands evenly distributed throughout the polymeric matrix can be discerned. These islands of partially wrinkled graphene laminates increase in size when the filler content increases, displaying values of aproximately 1 and 5 μm for samples with 1.0 and 2.5 vol % of GPE, respectively. Closer inspection of these islands by TEM in unstained samples suggests that graphene sheets are well dispersed and separated by the presence of the polymer (see Figure S1 in the Supporting Information). To examine the effect of the polymeric brushes of GPE on improving the interactions with the SEBS matrix, cryofractured

Figure 3. AFM phase images of SEBS and SEBS/GPE nanocomposites: SEBS (A), SEBS−GPE1 (B), SEBS−GPE3 (C), SEBS− GPE4 (D), SEBS−GPE5 (E). Field of view 1 μm × 1 μm. C


Article

Macromolecules PS blocks with an average diameter of ≈20 nm. The nanocomposites also present a microphase-separated morphology with a similar cylindrical structure. Some local orientation of the PS cylinders can be distinguished in Figure 3. For example, Figure 3B shows PS cylinders predominantly ordered parallel to the free surface, and Figure 3D illustrates a hexagonal arrangement of cylinders perpendicular to the scanned surface. Figure 3 suggests that as the amount of modified graphene increases, a more disordered self-assembled microphaseseparated structure is observed. This effect is clearly observed in Figure 3E which shows a clear distortion of the cylindrical packing due to the effect of the very high content of GPE (5.2 vol %). These results are in agreement with the SAXS and TEM observations that will be described below. In order to investigate in detail the morphological features of the nanocomposites, SAXS and TEM measurements were performed. Figure 4 shows representative 2D SAXS patterns for

distribution of each maximum, indicating some local preferential orientation. In addition, the SAXS diffraction maxima of this material appear as slightly eccentric ellipses and are not circular in shape. A similar eccentricity has been reported for SEBS with the same composition subjected to small deformations (5−10% strain).41 The local character of the orientation anisotropy and eccentricity was demonstrated by performing similar experiments at different locations on the same SEBS film. It revealed preferential orientation directed along diverse azimuthal angles and a range of ellipsoid eccentricities. These particular features detected on the diffraction patterns of SEBS are progressively lost as graphene is incorporated into the matrix in such a way that full isotropic rings can be distinguished in the nanocomposite with 2.5 vol % GPE. In addition, SEBS exhibits a number of diffracted SAXS maxima that become clearly distinct by plotting the scattered intensity, integrated over a small range of azimuthal angles (±5°) along the minor axis of the ellipsoid (at ≈55°), as a function of q, as shown in Figure 4A for q below 0.5 nm−1. This intensity profile can be separated into a number of diffraction peaks as shown in the Supporting Information (see Figure S3). Four maxima at q = 0.204, 0.220, 0.250, and 0.409 nm−1 can be discerned and are indicated in Figure 4A by arrows. In contrast, all the nanocomposites exhibit only two maxima in this q range. As an example, the intensity profile of the SEBS/GPE material with 1 vol % of filler is illustrated above Figure 4D, and the deconvolution into two principal maxima at q = 0.215 nm−1 and at q = 0.372 nm−1 is included in Figure S3. It is noteworthy that these peaks appear at relative q positions of 1:√3 and suggest a planar hexagonal packing in agreement with the hexagonal arrangement of cylinders. This statement is further confirmed by the appearance of a small maxima, out of the q range shown in Figure 4, at q = 0.569 nm−1 (q relative position of √7). In contrast, the SEBS diffraction maxima shown in Figure 4A are consistent with an oblique planar unit cell with a = 32.4 nm, b = 30.0 nm, and γ = 72°. It is worth remarking the local character of this specific unit cell: experiments at different locations have yielded oblique unit cells with comparable a and b parameters and γ values in the range 60°−76°. In other words, SEBS can locally exhibit hexagonal packing (γ = 60°) or oblique (60° < γ < 80°), the latter appearing to arise from the deformation of the hexagonal symmetry. Electron microscopy observations are in agreement with these statements. Figure 5 shows the TEM images of SEBS and two of the nanocomposites. Dark regions correspond to PS domains stained with ruthenium tetroxide, and the bright areas are related to the unstained ethylenic phase. In the case of SEBS, an hexagonal arrangement of cylinders can be found at certain locations, as marked with an arrow in Figure 5A. This hexagonal packing can appear locally distorted, possibly due to tensions arising during the formation of the polymer film, as shown in Figure 5B. In this case, the packing of cylinders can be appropriately represented by an oblique unit cell. Moreover, the unit cell of Figure 5B adopts a ≈ 28 nm, b ≈ 21 nm, and γ ≈ 80°, which is in the range of the unit cell determined by SAXS. Figures 5C and 5D show, as an example, a clear hexagonal packing for SEBS−GPE3 and SEBS−GPE5, respectively, an arrangement which is also commonly found for the rest of nanocomposites, as already suggested by the SAXS data. In addition, comparison of Figures 5A and 5D reveals a decrease in domain size and packing order with the introduction of 5.2 vol % of graphene. This result is in agreement with SAXS studies that allowed the recognition of a progressive

Figure 4. SAXS patterns for neat SEBS (A), SEBS−GPE1 (B), SEBS− GPE2 (C), SEBS−GPE3 (D), SEBS−GPE4 (E), and SEBS−GPE5 (F). The plots of I vs q for SEBS and SEBS−GPE3 are included on top of (A) and (D) 2D patterns and were obtained by integrating the scattered intensity over a range of azimuthal angles (±5°) around the minor axis of the ellipsoid (at ≈55° for both materials) as indicated in the figure. Arrows indicate the position of the main diffraction maxima.

SEBS and the SEBS/GPE nanocomposites. In all cases, a number of diffracted maxima appear that arise from the packing arrangement of the morphological domains. In the first place, Figure 4 shows a broadening of the diffraction maxima as graphene content increases, suggesting a reduced size and/or perfection of the morphological domains, especially for graphene contents above 2.5 vol %. Second, for neat SEBS, one can distinguish a small modulation on the intensity D


Article

Macromolecules

Figure 5. TEM micrographs of SEBS and the SEBS/GPE nanocomposites: SEBS (A, B), SEBS−GPE3 (C), SEBS−GPE5 (D). The arrows guide the reader in locating hexagonal packing. The square in (B) defines the unit cell (oblique) that best describes the morphological arrangement of cylinders. The scale bar represents 500 nm for micrographs A, C, and D and 50 nm for B.

deterioration of the cylindrical packing as the graphene content increased that is more significant above 2.5 vol %. In fact, such changes are difficult to distinguish with TEM for the sample with only 1 vol % due to the local character of the microscopy studies (Figure 5C). In contrast, the SAXS measurements provide an average vision of the sample morphology over a significantly larger scale (spot size of 1.2 mm by 135 μm). Finally, it is important to remark that graphene laminates were difficult to visualize by TEM, especially in those nanocomposites with very low nanofiller content. As mentioned above, the distribution of graphene laminates was clearly observed in unstained samples where the phase-separated morphology of the copolymer was not resolved. Only in extremely thin areas located at the edges of stained sections and with high nanofiller content (SEBS−GPE4) were the graphene layers observed in the ethylenic phase (see Figure S2 in the Supporting Information). The thermal stability of all nanocomposites was investigated by TGA under a nitrogen atmosphere. The thermogravimetric (TG) and first derivative thermogravimetric (DTG) results are shown in Figure 6 and Table 2. The table includes the characteristic degradation temperatures for all samples. Data for modified graphene (GPE) and SEBS nanocomposites with 1.0 vol % of unmodified graphene (G) are also presented in Table 2 for comparison. The increase in all degradation temperatures measured for the 0.5 vol % GPE nanocomposite as compared with those of SEBS demonstrates enhancement in thermal stability of the former. With increasing graphene content (up to 1.0 vol % GPE) only the temperatures corresponding to 10% weight loss (T10) and the maximum rate of weight loss (Tmax) were observed to increase. At the highest contents (2.5 and 5.2 vol % GPE) both the initial degradation temperature, Ti, and T10 decrease, and a shoulder on the DGT curves at lower temperature is clearly observed (Figure 6B). The TGA curves display a single step at lower composition while two steps are discerned at higher content. Several effects must be considered in order to explain this behavior. It has already been proposed that well dispersed graphene can act as a barrier in polymer matrices and effectively hinder the diffusion of the degradation

Figure 6. TG (A) and DTG (B) curves under a nitrogen atmosphere at a rate of 10 °C min−1 for the different samples.

Table 2. Characteristic Degradation Temperatures Obtained from TGAa sample

Ti (°C)

T10 (°C)

Tmax (°C)

SEBS SEBS−GPE1 SEBS−GPE2 SEBS−GPE3 SEBS−GPE4 SEBS−GPE5 SEBS−G GPE

404 408 400 397 383 373 406 290b

426 429 430 427 418 414 430 385

456 457 464 470 474 482 462 447

a

Ti: initial degradation temperature obtained at 2% weight loss. T10: temperature corresponding to 10% weight loss. Tmax: maximum degradation rate temperature. bSee ref 37.

products slowing down the decomposition process. A key factor is the dispersion of graphene in the polymer matrix. A homogeneous dispersion of the carbon filler in the polymer will favor the barrier effect and will stabilize the material as is the case for SEBS/GPE nanocomposites with low filler content. However, at higher content the nanocomposites showed (by SEM) the presence of agglomerates (Figure 1) reducing the influence of the filler on the thermal stability. Another factor is that the PE brushes from GPE start to degrade at lower temperature. As the amount of GPE increases, it is manifested by a broadening on the lower temperature side of the peak of E


Article

Macromolecules

where σf is the conductivity of the filler, ϕ is the filler volume fraction, ϕc is the percolation threshold, and t is the universal critical exponent. Plotting log σ as a function log(ϕ − ϕc) yields a linear dependence, with fitting parameters t = 3.2 ± 0.4 and σf = 103.2±0.2 S cm−1 (inset in Figure 7A), values that are very similar to those obtained for nanocomposites of PS and graphene.43 On the other hand, the critical exponent t is expected to depend on the system dimensionality, and some calculations suggest values of t ≈ 1.33 and t ≈ 2 for 2D and 3D systems, respectively.45 However, higher values are frequently estimated, and theories like the “Swiss cheese model”46 and tunneling effects47 have been proposed to justify them. On the other hand, some authors have argued that the higher the value of t, the higher the percolation threshold,45 and this statement seems to be consistent with our results. For instance, the t and ϕc values in our work (t = 3.2, ϕc = 0.8 vol %) are both higher than those obtained in PS/graphene (t = 2.7, ϕc = 0.1 vol %).43 This looks reasonable as SEBS is composed of two welldifferentiated domains and in one of them (ethylenic) the electronic mobility is not assisted by π-electrons. Despite the presence of the ethylenic phase, the conductivity values for SEBS/GPE nanocomposites with very small amounts of graphene are higher than that for an antistatic material (1 × 10−8 S cm−1), becoming much better for higher amounts of graphene (∼1 × 10−6 S cm−1 for 1.0 vol % of GPE). The latter value is markedly higher than the conductivity reported to date for graphene/SEBS nanocomposites (1.2 × 10−11 S cm−1 for 4 wt % of graphene), in which the matrix is SEBS grafted with maleic anhydride (SEBS-g-MAH).30 Moreover, the good conductivity values obtained in the nanocomposites reported here corroborate that the polyethylene brushes emerging from the graphene promote the mixing with the ethylenic phase of the elastomer leading to a homogeneous distribution of the GPE in this phase, in agreement with TEM observations (see Figure S2 in Supporting Information). Furthermore, the electrical conductivity of this kind of nanocomposites is known to be dependent on the temperature, and different behaviors have been observed depending on the type of matrix or filler. We have found that the variation of conductivity with temperature is nonlinear and that the behavior changes with the filler content.48,49 Some authors have demonstrated that the conductivity vs temperature curve is composed of different segments, each governed by different transport mechanisms,48 principally variable range hopping (VRH)50−52 and thermal fluctuation induced tunneling (TFIT).50,53,54 Among the nanocomposites investigated in the present work, those with filler content near the percolation threshold concentration logically show a stronger dependence on the temperature, displaying a perfect exponential growth of conductivity with temperature (Figure 7B). This curve can be adjusted using the VRH model, where the conductivity is expressed as follows:48

the derivatives curves. For the 5.2 vol % nanocomposite it becomes a clear shoulder. A most important challenge for polymer nanocomposites based on carbon nanomaterials, and elastomers should not be an exception, is obtaining electrical conductivity, which impacts directly on a wide range of applications, including actuators, bodily motion sensors, or structural damage or fatigue sensors, among others.24 The variation of the electrical conductivity of the SEBS/GPE composites as a function of GPE volume fraction at room temperature, obtained by the four-probe method, is shown in Figure 7A. A typical percolation

Figure 7. (A) Electrical conductivity of the SEBS/GPE composites as a function of GPE volume fraction. The inset corresponds to the variation of log σ with log(ϕ − ϕc), ϕc being the percolation threshold (eq 3). (B) Dependence of the dc conductivity of SEBS−GPE3 with the reciprocal of temperature (eq 4).

phenomenon, i.e., a rapid increase of the electrical conductivity when the filler forms an infinite network of connected paths across the insulating matrix, is evident. From a sigmoidal fit of the percolation curve, a percolation threshold of ϕc = 0.76 wt % was obtained, slightly higher than that reported for similar polymers like poly(styrene-co-butadiene-co-styrene) (SBS)42 and polystyrene (PS).43 The conductivity of a composite, above the percolation threshold, obeys the following power law:44 ⎡ ϕ − ϕ ⎤t c ⎥ σ = σf ⎢ ⎢⎣ 1 − ϕc ⎥⎦

σ = 2.7 × 10−10 exp( −0.021/kT α)

(4)

where k is the Boltzmann constant and α is the temperature exponent. The linear plot of log σ against the inverse of temperature yields a value of α = 0.24, matching perfectly with the typical values for classical semiconductors (α = 0.25). This suggests that for low filler amounts the mechanism that governs the charge transport is the hopping of carriers between neighboring conductive graphene laminates near ethylenic domains in SEBS.

(3) F


Article

Macromolecules As the filler content increases, the variation of conductivity with temperature cannot be adjusted considering a single mechanism, especially at low temperature. In the case of the 2.5 vol % nanocomposite, the VRH model applies above 363 K, a temperature in the vicinity of the glass transition of styrene domains (∼373 K), while at lower temperatures the conductivity adjusts more favorably with the TFIT model to T−1/7.49 The nanocomposite with the highest volume of graphene (5.2 vol %) displays the same behavior, with a change of transport mechanism around 352 K. The different conductivity mechanism for low and high filler contents can be correlated to the distinct morphologies found by AFM, TEM, and SAXSwell-developed in the former case and significantly more disordered in the latter. Figure 8 illustrates the fwhm of

temperature where a rearrangement of domains is facilitated and the average size is increased. This could explain the change in conductivity mode from tunneling at low temperatures to hopping at higher ones. The changes in the mechanical properties of SEBS/GPE with various GPE contents are shown in Table 3. It can be clearly Table 3. Mechanical Parameters Determined from Tensile Testsa

a

sample

GPE (vol %)

E (MPa)

SEBS SEBS−GPE2 SEBS−GPE3 SEBS−GPE4 SEBS−GPE5 SEBS−G

0 0.7 1 2.5 5.2 b

75.9 22.4 35.6 76.8 153.6 39.7

modulus 100% (MPa) 3.5 5.2 7.2 10.09

± ± ± ±

0.3 0.8 0.6 0.05

7.06 ± 0.03

εb (%) 641 647 511 314 45 527

Young’s modulus (E), elongation at break (εb). bUnmodified G.

seen that the elongation at break decreases by 20% with the incorporation of 1.0 vol % of GPE, to a value that is still relatively high. Therefore, this sample preserves to a significant extent the mechanical properties of SEBS but incorporates reasonable electrical conductivity which makes it very attractive for use in applications such as actuators or sensors. Regarding the Young’s modulus, while the nanocomposites with low graphene concentration (up to 1.0 vol %) display values slightly lower than pure SEBS, at the highest composition a value 2 times higher was observed. The case of the sample with 2.5 vol % of graphene is especially interesting as it still behaves as an elastomer with an elongation at break higher than 300%, it displays a Young’s modulus similar to SEBS, and it is electrically conductive reaching 0.024 S cm−1. As a final point, Table 3 also includes the tensile properties in terms of the elastic modulus at 100% strain. A progressive enhancement in this value is observed as the amount of filler increases. It must be noted that the modulus at 100% strain could not be measured for the nanocomposite with the highest filler content (5.2 vol %), as it breaks at lower deformations as a consequence of the aggregation of the filler throughout the matrix that creates points of rupture in the sample. The evaluation of the modulus at 100% strain represents additional information on the nanocomposites performance and highlights the relevance of an appropriate dispersion of the filler.

Figure 8. Full width at half-maximum (fwhm) of the hexagonal (100) main diffraction peak as a function of graphene content at three different temperatures.

the main diffraction peak (100) for the three conductive nanocomposites (GPE ≥ 1 vol %). Values for SEBS and the nanocomposites with graphene contents below the percolation threshold have also been included for comparison. For the sake of simplicity, a SEBS pattern exhibiting limited eccentricity, and a pure hexagonal arrangement at room temperature was selected. The background-subtracted I−q profiles of all materials employed to calculate fwhm were obtained by integration of the scattered intensity over a small range of azimuthal angles (±5°) along the direction of preferential orientation. Figure 8 shows that the nanocomposites with GPE concentration ≤1 vol % exhibit a significantly lower fwhm than those with higher filler content and that all fwhm values decrease with increasing temperature. This observation has two important consequences. First, the size of the morphological domains for 2.5 and 5.2 vol % GPE is significantly lower than that for 1 vol % of filler. This, together with the fact that nanocomposites with high filler content exhibit a significant isotropy of domain orientation (see Figure 4E,F) and a poorly developed morphology (see Figure 5D), could favor the proximity between fillers leading to a preference for tunnelingtype conductivity in contrast to hopping type. The latter behavior seems to be promoted in the case of the lowest graphene content above the percolation threshold where the local orientation and significant size of the morphological domains could promote the separation of graphene sheets among neighboring domains. It is also noteworthy that the fwhm for high graphene contents decreases significantly with

■

CONCLUSIONS A new family of nanocomposites composed of one of the most employed thermoplastic elastomers, SEBS, and graphene has been prepared. The appropriate chemical functionalization of graphene provides the means for a selective interaction of the filler with specific matrix domains. In this study, the polyethylene brushes emerging from graphene are found to mix preferentially with the ethylenic phase of the elastomer, locating the conductive filler near electron-poor domains. The nanocomposites present a relative low percolation threshold and reasonable values of electrical conductivity, the first time that such values are reported for SEBS. Moreover, the conductive nanocomposites exhibit the elastomeric character of SEBS up to significant filler loadings. The compendium of conductivity and elasticity in the same material makes it a good candidate to be employed as actuator or electromechanical sensor. G


Article

Macromolecules

■

Transfer and RAFT Agents for Producing Graphene Polymer Brushes and Polystyrene Nanocomposites. Macromolecules 2012, 45 (17), 7083−7090. (10) Ye, Y.-S.; Chen, Y.-N.; Wang, J.-S.; Rick, J.; Huang, Y.-J.; Chang, F.-C.; Hwang, B.-J. Versatile Grafting Approaches to Functionalizing Individually Dispersed Graphene Nanosheets Using RAFT Polymerization and Click Chemistry. Chem. Mater. 2012, 24 (15), 2987−2997. (11) He, H.; Gao, C. General Approach to Individually Dispersed, Highly Soluble, and Conductive Graphene Nanosheets Functionalized by Nitrene Chemistry. Chem. Mater. 2010, 22 (17), 5054−5064. (12) Kan, L.; Xu, Z.; Gao, C. General Avenue to Individually Dispersed Graphene Oxide-Based Two-Dimensional Molecular Brushes by Free Radical Polymerization. Macromolecules 2011, 44 (3), 444−452. (13) Wu, S.; Li, H.; Huang, G.; Wu, J. Nucleating effect of multiwalled carbon nanotubes and graphene on the crystallization kinetics and melting behavior of olefin block copolymers. RSC Adv. 2014, 4 (36), 19024−19033. (14) Trung, T.; Ramasundaram, S.; Hong, S.; Lee, N.-E.; Lee; Trung, S.; Hong, N.-E. Flexible and Transparent Nanocomposite of Reduced Graphene Oxide and P(VDF-TrFE) Copolymer for High Thermal Responsivity in a Field-Effect Transistor. Adv. Funct. Mater. 2014, 24 (22), 3438−3445. (15) Popescu, M.-T.; Tasis, D.; Tsitsilianis, C. Ionizable Star Copolymer-Assisted Graphene Phase Transfer between Immiscible Liquids: Organic Solvent/Water/Ionic Liquid. ACS Macro Lett. 2014, 3 (10), 981−984. (16) Yang, H.; Paek, K.; Kim, B. Efficient Temperature Sensing Platform Based on Fluorescent Block Copolymer-Functionalized Graphene Oxide. Nanoscale 2013, 5 (13), 5720−4. (17) Tan, Y.; Fang, L.; Xiao, J.; Song, Y.; Zheng, Q. Grafting of copolymers onto graphene by miniemulsion polymerization for conductive polymer composites: improved electrical conductivity and compatibility induced by interfacial distribution of graphene. Polym. Chem. 2013, 4 (10), 2939−2944. (18) Beckert, F.; Rostas, A.; Thomann, R.; Weber, S.; Schleicher, E.; Friedrich, C.; Mülhaupt, R. Self-Initiated Free Radical Grafting of Styrene Homo- and Copolymers onto Functionalized Graphene. Macromolecules 2013, 46 (14), 5488−5496. (19) Yu, A.-D.; Liu, C.-L.; Chen, W.-C. Supramolecular block copolymers: graphene oxide composites for memory device applications. Chem. Commun. (Cambridge, U. K.) 2012, 48 (3), 383−385. (20) Jiang, K.; Ye, C.; Zhang, P.; Wang, X.; Zhao, Y. One-Pot Controlled Synthesis of Homopolymers and Diblock Copolymers Grafted Graphene Oxide Using Couplable RAFT Agents. Macromolecules 2012, 45 (3), 1346−1355. (21) Cao, Y.; Lai, Z.; Feng, J.; Wu, P. Graphene oxide sheets covalently functionalized with block copolymers via click chemistry as reinforcing fillers. J. Mater. Chem. 2011, 21 (25), 9271−9278. (22) Peponi, L.; Tercjak, A.; Verdejo, R.; Lopez Manchado, M.; Mondragon, I.; Kenny, J. M. Confinement of Functionalized Graphene Sheets by Triblock Copolymers. J. Phys. Chem. C 2009, 113 (42), 17973−17978. (23) Papageorgiou, D. G.; Kinloch, I. A.; Young, R. J. Graphene/ elastomer nanocomposites. Carbon 2015, 95, 460−484. (24) Sadasivuni, K. K.; Ponnamma, D.; Thomas, S.; Grohens, Y. Evolution from graphite to graphene elastomer composites. Prog. Polym. Sci. 2014, 39 (4), 749−780. (25) Liang, J.; Xu, Y.; Huang, Y.; Zhang, L.; Wang, Y.; Ma, Y.; Li, F.; Guo, T.; Chen, Y. Infrared-Triggered Actuators from Graphene-Based Nanocomposites. J. Phys. Chem. C 2009, 113 (22), 9921−9927. (26) Boland, C. S.; Khan, U.; Backes, C.; O’Neill, A.; McCauley, J.; Duane, S.; Shanker, R.; Liu, Y.; Jurewicz, I.; Dalton, A. B.; Coleman, J. N. Sensitive, High-Strain, High-Rate Bodily Motion Sensors Based on Graphene−Rubber Composites. ACS Nano 2014, 8 (9), 8819−8830. (27) Roh, E.; Hwang, B.-U.; Kim, D.; Kim, B.-Y.; Lee, N.-E. Stretchable, Transparent, Ultrasensitive, and Patchable Strain Sensor for Human−Machine Interfaces Comprising a Nanohybrid of Carbon

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b00490. Additional results on the characterization of samples by TEM and SAXS (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*Phone +34-912587432; Fax +34-915644853; e-mail horacio@ ictp.csic.es (H.J.S.). Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS Financial support from MINECO (Ministerio de Economiá y Competitividad), Spain, for funding the research reported under grants MAT2013-47898-C2-1-R and MAT2013-47898C2-2-R is gratefully acknowledged. H.J.S. acknowledges the MINECO for a “Ramón y Cajal” Senior Research Fellowship, and P.E.-J. and S.Q.-D. acknowledge a FPI Fellowship. The authors are indebted to Mr. J. González-Casablanca and R. Castro from Universidad Rey Juan Carlos (URJC) of Madrid for their assistance in the TEM measurements and to Mr. David Gómez and Dr. Pilar Posadas of the Characterization Service of the Institute of Polymer Science & Technology, for the assistance in the SEM and AFM measurements, respectively. The SAXS experiments were performed at NCD-BL11 beamline at ALBA Synchrotron with the collaboration of ALBA staff, and we are especially endeavored to Dr. Eva Crosas for her dedication and availability.

■

REFERENCES

(1) Hu, K.; Kulkarni, D.; Choi, I.; Tsukruk, V. Graphene-polymer nanocomposites for structural and functional applications. Prog. Polym. Sci. 2014, 39 (11), 1934−1972. (2) Roy, N.; Sengupta, R.; Bhowmick, A. K. Modifications of carbon for polymer composites and nanocomposites. Prog. Polym. Sci. 2012, 37 (6), 781−819. (3) Salavagione, H. J.; Diez-Pascual, A. M.; Lazaro, E.; Vera, S.; Gomez-Fatou, M. A. Chemical sensors based on polymer composites with carbon nanotubes and graphene: the role of the polymer. J. Mater. Chem. A 2014, 2 (35), 14289−14328. (4) Salavagione, H. J.; Martínez, G.; Ellis, G. Recent Advances in the Covalent Modification of Graphene With Polymers. Macromol. Rapid Commun. 2011, 32 (22), 1771−1789. (5) Potts, J. R.; Dreyer, D. R.; Bielawski, C. W.; Ruoff, R. S. Graphene-based polymer nanocomposites. Polymer 2011, 52 (1), 5− 25. (6) Salavagione, H. J.; Gómez, M. A.; Martínez, G. Polymeric Modification of Graphene through Esterification of Graphite Oxide and Poly(vinyl alcohol). Macromolecules 2009, 42 (17), 6331−6334. (7) Salavagione, H. J., Innovative Strategies to Incorporate Graphene in Polymer Matrices: Advantages and Drawbacks from an Applications Viewpoint in Innovative Graphene Technologies: Developments & Characterisation. In Innovative Graphene Technologies:Developments and Characterisation; Tiwari, A., Ed.; Smithers-Rapra: Shawbury, 2013; Vol. 1, pp 177, 221. (8) Salavagione, H. J.; Martínez, G. Importance of Covalent Linkages in the Preparation of Effective Reduced Graphene Oxide−Poly(vinyl chloride) Nanocomposites. Macromolecules 2011, 44 (8), 2685−2692. (9) Beckert, F.; Friedrich, C.; Thomann, R.; Muelhaupt, R.; Mülhaupt, R. Sulfur-Functionalized Graphenes as Macro-ChainH


Article

Macromolecules Nanotubes and Conductive Elastomers. ACS Nano 2015, 9, 6252− 6261. (28) Li, X.; Zhang, R.; Yu, W.; Wang, K.; Wei, J.; Wu, D.; Cao, A.; Li, Z.; Cheng, Y.; Zheng, Q.; Ruoff, R. S.; Zhu, H. Stretchable and highly sensitive graphene-on-polymer strain sensors. Sci. Rep. 2012, 2, 870. (29) You, F.; Wang, D.; Cao, J.; Li, X.; Dang, Z.-M.; Hu, G.-H. In situ thermal reduction of graphene oxide in a styrene−ethylene/butylene− styrene triblock copolymer via melt blending. Polym. Int. 2014, 63 (1), 93−99. (30) Shi, H.; Shi, D.; Li, C.; Luan, S.; Yin, J.; Li, R. K. Y. Preparation of functionalized graphene/SEBS-g-MAH nanocomposites and improvement of its electrical, mechanical properties. Mater. Lett. 2014, 133 (0), 200−203. (31) Mao, C.; Huang, J.; Zhu, Y.; Jiang, W.; Tang, Q.; Ma, X. Tailored Parallel Graphene Stripes in Plastic Film with Conductive Anisotropy by Shear-Induced Self-Assembly. J. Phys. Chem. Lett. 2013, 4 (1), 43−47. (32) Shen, B.; Zhai, W.; Chen, C.; Lu, D.; Wang, J.; Zheng, W. Melt Blending In situ Enhances the Interaction between Polystyrene and Graphene through π−π Stacking. ACS Appl. Mater. Interfaces 2011, 3 (8), 3103−3109. (33) Périé, T.; Brosse, A.-C.; Tencé-Girault, S.; Leibler, L. Nanostructured films and composites from carbon nanotubes dispersed by ABC block terpolymers in selective solvent. Polymer 2011, 52 (14), 3065−3073. (34) Périé, T.; Brosse, A.-C.; Tencé-Girault, S.; Leibler, L. Mechanical and electrical properties of multi walled carbon nanotube/ABC block terpolymer composites. Carbon 2012, 50 (8), 2918− 2928. (35) Behling, R. E.; Wolf, L. M.; Cochran, E. W. Hierarchically Ordered Montmorillonite Block Copolymer Brushes. Macromolecules 2010, 43 (5), 2111−2114. (36) Peponi, L.; Tercjak, A.; Gutierrez, J.; Stadler, H.; Torre, L.; Kenny, J. M.; Mondragon, I. Self-Assembling of SBS Block Copolymers as Templates for Conductive Silver Nanocomposites. Macromol. Mater. Eng. 2008, 293 (7), 568−573. (37) Castelaín, M.; Martínez, G.; Marco, C.; Ellis, G.; Salavagione, H. J. Effect of Click-Chemistry Approaches for Graphene Modification on the Electrical, Thermal, and Mechanical Properties of Polyethylene/ Graphene Nanocomposites. Macromolecules 2013, 46 (22), 8980− 8987. (38) Castelaín, M.; Martínez, G.; Ellis, G.; Salavagione, H. J. Versatile chemical tool for the preparation of conductive graphene-based polymer nanocomposites. Chem. Commun. 2013, 49 (79), 8967−8968. (39) Hammersley, A. P.; Svensson, S. O.; Thompson, A. Calibration and correction of spatial distortions in 2D detector systems. Nucl. Instrum. Methods Phys. Res., Sect. A 1994, 346 (1−2), 312−321. (40) Bockstaller, M. R.; Mickiewicz, R. A.; Thomas, E. L. Block Copolymer Nanocomposites: Perspectives for Tailored Functional Materials. Adv. Mater. 2005, 17 (11), 1331−1349. (41) Mamodia, M.; Indukuri, K.; Atkins, E. T.; De Jeu, W. H.; Lesser, A. J. Hierarchical description of deformation in block copolymer TPEs. J. Mater. Sci. 2008, 43 (22), 7035−7046. (42) Liu, Y.-T.; Xie, X.-M.; Ye, X.-Y. High-concentration organic solutions of poly(styrene-co-butadiene-co-styrene)-modified graphene sheets exfoliated from graphite. Carbon 2011, 49 (11), 3529−3537. (43) Stankovich, S.; Dikin, D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen, S. T.; Ruoff, R. S. Graphene-based composite materials. Nature 2006, 442 (7100), 282− 286. (44) McLachlan, D. S.; Chiteme, C.; Park, C.; Wise, K. E.; Lowther, S. E.; Lillehei, P. T.; Siochi, E. J.; Harrison, J. S. AC and DC percolative conductivity of single wall carbon nanotube polymer composites. J. Polym. Sci., Part B: Polym. Phys. 2005, 43 (22), 3273−3287. (45) Bauhofer, W.; Kovacs, J. Z. A review and analysis of electrical percolation in carbon nanotube polymer composites. Compos. Sci. Technol. 2009, 69 (10), 1486−1498.

(46) Chiteme, C.; McLachlan, D. S. ac and dc conductivity, magnetoresistance, and scaling in cellular percolation systems. Phys. Rev. B: Condens. Matter Mater. Phys. 2003, 67 (2), 024206. (47) Balberg, I.; Balberg. Limits on the continuum-percolation transport exponents. Phys. Rev. B: Condens. Matter Mater. Phys. 1998, 57 (21), 13351−13354. (48) Syurik, J.; Ageev, O. A.; Cherednichenko, D. I.; Konoplev, B. G.; Alexeev, A. Non-linear conductivity dependence on temperature in graphene-based polymer nanocomposite. Carbon 2013, 63 (0), 317− 323. (49) Alexander, M. G. Anomalous temperature dependence of the electrical conductivity of carbon-poly(methyl methacrylate) composites. Mater. Res. Bull. 1999, 34 (4), 603−611. (50) Jeon, K.; Lumata, L.; Tokumoto, T.; Steven, E.; Brooks, J.; Alamo, R. G. Low electrical conductivity threshold and crystalline morphology of single-walled carbon nanotubes − high density polyethylene nanocomposites characterized by SEM, Raman spectroscopy and AFM. Polymer 2007, 48 (16), 4751−4764. (51) Psarras, G. C. Hopping conductivity in polymer matrix−metal particles composites. Composites, Part A 2006, 37 (10), 1545−1553. (52) Peng, H. Aligned Carbon Nanotube/Polymer Composite Films with Robust Flexibility, High Transparency, and Excellent Conductivity. J. Am. Chem. Soc. 2008, 130 (1), 42−43. (53) Jović, N.; Dudić, D.; Montone, A.; Antisari, M. V.; Mitrić, M.; Djoković, V. Temperature dependence of the electrical conductivity of epoxy/expanded graphite nanosheet composites. Scr. Mater. 2008, 58 (10), 846−849. (54) Nogales, A.; Broza, G.; Roslaniec, Z.; Schulte, K.; Šics, I.; Hsiao, B. S.; Sanz, A.; García-Gutiérrez, M. C.; Rueda, D. R.; Domingo, C.; Ezquerra, T. A. Low Percolation Threshold in Nanocomposites Based on Oxidized Single Wall Carbon Nanotubes and Poly(butylene terephthalate). Macromolecules 2004, 37 (20), 7669−7672.

I


Development of Advanced Elastomeric Conductive Nanocomposites

Recommend Documents