Synthesis and Characterization of Nanostructured Copolymer-Grafted

Article pubs.acs.org/IECR

Synthesis and Characterization of Nanostructured CopolymerGrafted Multiwalled Carbon Nanotube Composite Thermoplastic Elastomers toward Unique Morphology and Strongly Enhanced Mechanical Properties Feng Jiang,† Yaqiong Zhang,† Zhongkai Wang,† Huagao Fang,‡ Yunsheng Ding,‡ Hangxun Xu,† and Zhigang Wang*,† †

CAS Key Laboratory of Soft Matter Chemistry, Department of Polymer Science and Engineering, Hefei National Laboratory for Physical Sciences at the Microscale, University of Science and Technology of China, Hefei, Anhui Province 230026, P. R. China ‡ Provincial Key Laboratory of Advanced Functional Materials and Devices, Institute of Polymer Materials and Chemical Engineering, School of Chemistry and Chemical Engineering, Hefei University of Technology, Hefei, Anhui Province 230009, P. R. China S Supporting Information *

ABSTRACT: Considering that multiwalled carbon nanotubes (MWCNTs) can be used as anisotropic and stiff nano-objects acting as minority physical cross-linking points dispersed in soft polymer grafting matrixes, a series of copolymer-grafted multiwalled carbon nanotube composite thermoplastic elastomers (CTPEs), MWCNT-graf t-poly(n-butyl acrylate-co-methyl methacrylate) [MWCNT-g-P(BA-co-MMA)], with minor MWCNT contents of 1.2−3.8 wt % was synthesized by the surfaceinitiated activators regenerated by electron transfer for atom-transfer radical polymerization (ARGET ATRP) method. Excellent dispersion of the MWCNTs in the CTPEs was demonstrated by SEM and TEM, and the thermal stability properties and glass transition temperatures of the CTPEs were characterized by thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC), respectively. Mechanical property test results demonstrated that the CTPEs exhibit obviously enhanced mechanical properties, such as higher tensile strength and elastic recovery, as compared with their linear P(BA-co-MMA) copolymer counterparts. The microstructural evolutions in the CTPEs during tensile deformation as investigated by in situ smallangle X-ray scattering (SAXS) revealed the role of the MWCNTs, which can provide additional cross-linking points and transform soft elastomers into strong ones.

1. INTRODUCTION Thermoplastic elastomers (TPEs), which combine elasticity with thermoplastic properties, have a myriad of important applications in daily life.1−3 Most TPEs are linear ABA triblock copolymers, such as polystyrene-b-polybutadiene-b-polystyrene (SBS) and polystyrene-b-polyisoprene-b-polystyrene (SIS).4 The phase-separated hard domains provide physical crosslinks and endow sufficient tensile strength during deformation. However, synthesis of these triblock copolymers is challenging, requiring high chain-extension efficiency among blocks and stringent polymerization conditions. Recently, we proposed a novel architecture (architecture III) toward third-generation TPEs by using semirigid cellulose chains as the minority physical cross-linkers and grafted random copolymer chains from the semirigid cellulose backbone chains as the soft polymer matrix. Such an architecture design circumvents the conventionally challenging synthesis of TPEs and has great potential for being put into practice for large-scale production.5 In principle, the semirigid polymer backbone chains can be replaced by other stiff nano-objects, for example, single- or multiwalled carbon nanotubes (SWCNTs or MWCNTs, respectively), which is the main topic of this work. Carbon nanotubes (CNTs) have attracted tremendous attention since their discovery6 owing to their inherent excellent electrical and mechanical properties. From a structural © 2014 American Chemical Society

standpoint, CNTs can be considered as very stiff macromolecules with persistence lengths of 26−174 μm or even longer.7 CNTs exhibit extremely high tensile strength and tensile modulus values of 11−150 GPa and 0.27−0.95 TPa, respectively.8,9 Therefore, CNTs are regarded as ideal fillers for improving the mechanical and electrical properties of polymer composites. Composites consisting of CNTs and polymer matrixes have great potential for important applications, such as in reinforced materials and electronic devices.10−12 Minute amounts of CNTs can have significant influences on the properties of CNT/polymer composites. According to our previous work, the aspect ratio of CNTs is a key factor affecting CNT network formation and the corresponding percolation threshold.13−17 Unfortunately, because of the high aspect ratios and strong π−π interactions of CNTs, it is difficult to obtain homogeneous dispersions of CNTs in polymer matrixes.18,19 Filler dispersion and interfacial interactions are crucial for enhancing the mechanical properties of polymer composites.20 Various methods have been employed to optimize the dispersion of CNTs, such as oxidation,21,22 sonication,23 melt Received: Revised: Accepted: Published: 20154

October 10, 2014 November 20, 2014 December 9, 2014 December 17, 2014 dx.doi.org/10.1021/ie504005f | Ind. Eng. Chem. Res. 2014, 53, 20154−20167

Industrial & Engineering Chemistry Research

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Figure 1. Illustration of the synthesis of copolymer-grafted MWCNT CTPEs, MWCNT-g-P(BA-co-MMA) copolymers.

compounding,24,25 solution blending,26 coagulation,27,28 use of surfactants,29 and in situ polymerization.30 Among these methods, in situ polymerization is versatile and robust for the preparation of polymer-grafted nanotubes and the subsequent processing of the corresponding polymer composites. A “graftfrom” approach, namely, surface-initiated atom-transfer radical polymerization (ATRP), has been utilized to graft various polymers from CNTs,31−34 through which CNT/polymer composites can be fabricated with high grafting densities and improved mechanical properties due to the efficient homogeneous dispersion of the CNTs in the polymer matrix and interfacial stress transfer between the CNTs and the polymer.35 Hwang et al. used poly(methyl methacrylate)- (PMMA-) grafted MWCNTs to reinforce commercial PMMA and found that the functionalized MWCNTs were dispersed homogeneously in the PMMA composites and that a tensile load could be transferred to PMMA-grafted MWCNTs.36 The mechanical properties of CNT/polymer composites were found to increase dramatically at low CNT contents.7 Addition of CNTs to various elastomers such as polyurethanes,37,38 silicon rubber,39,40 liquid-crystalline elastomers,41 and acrylic elastomers42 has also been found to impart attractive properties. However, little research has been performed on copolymer-grafted CNT composite thermoplastic elastomers (CTPEs) synthesized by directly grafting soft polymers from the surface of CNTs. We consider that MWCNTs can be homogeneously dispersed in this type of CTPE and that the MWCNT/polymer interface becomes strong through chemical bonding by graft polymerization. The key factor that governs the enhanced mechanical properties of MWCNT-based CTPEs is efficient load transfer between the walls of the MWCNTs and the polymer matrix.36,43 In this work, in an attempt to thoroughly understand the effects of stiff nano-objects on the enhanced mechanical properties of CTPEs, a series of copolymer-grafted MWCNT CTPEs was synthesized, in which the random copolymers were grafted from the stiff nano-objects (MWCNTs). Specifically, we applied the activators regenerated by electron transfer for atomtransfer radical polymerization (ARGET ATRP)44 method to synthesize the CTPEs, namely, MWCNT-graf t-poly(n-butyl acrylate-co-methyl methacrylate) [MWCNT-g-P(BA-coMMA)] copolymers (Figure 1). The thermal properties and glass transition temperatures of the CTPEs were characterized by thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC), respectively. The phase morphologies were observed by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The mechanical properties of the CTPEs were measured in both monotonic tensile test mode and step cyclic tensile test mode. The microstructural evolutions during the tensile deformation and elastic recovery processes were examined by the in situ small-

angle X-ray scattering (SAXS) technique. For comparison purposes, the linear poly(n-butyl acrylate-co-methyl methacrylate) [P(BA-co-MMA)] copolymer counterparts were synthesized and subjected to the same characterizations.

2. MATERIALS AND METHODS 2.1. Materials. Multiwalled carbon nanotubes (MWCNTs) with high aspect ratios were supplied by ChengDu Organic Chemistry Co. Ltd., Chinese Academy of Sciences (average diameter, 8−15 nm; average length, ∼50 μm; specific surface area, >233 m2/g; purity, >95%; ash content, 45 wt %) in the MWCNT-g-P(BA-co-MMA) copolymers. To clarify, we applied the Fox equation to describe the correlation between Tg and the monomer composition of the statistical copolymers48,49

Figure 5. SEM micrographs of fractured surfaces (at different magnifications) of samples (a,a′) BA6500-MWCNT1.2 and (b,b′) BA6500-MWCNT3.8.

matrix is basically featureless because of the random copolymerization of MMA and BA components, and the bright spots in the micrographs represent MWCNTs because of their high electrical conductivity. During fracture of the film samples, some MWCNTs were pulled out of the film surfaces and exposed to SEM observation. It can be seen that the MWCNTs were dispersed homogeneously in the copolymer matrix with no obvious large aggregates, if compared with the results observed for conventional polymer composites.13,21,22 For MWCNT-g-P(BA-co-MMA) copolymers, the P(BA-co-MMA) 20158

dx.doi.org/10.1021/ie504005f | Ind. Eng. Chem. Res. 2014, 53, 20154−20167


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chains cannot be separated from the MWCNTs because they are tethered onto the MWCNTs. Therefore, the dispersion of MWCNTs is sufficiently satisfied for MWCNT-g-P(BA-coMMA) copolymers. TEM observations can provide clearer morphological visualizations of the dispersion of the MWCNTs in MWCNT-g-P(BA-co-MMA) copolymers. Figure 6 shows

Table 2. Mechanical Properties of MWCNT-g-P(BA-coMMA) Copolymers average ERa (%) sample ID

first series

second series

stress at break (MPa)b

strain at break (%)b

BA7000-MWCNT1.3 BA6500-MWCNT1.2 BA6500-MWCNT1.7 BA6500-MWCNT3.8 BA6000-MWCNT1.4 BA5500-MWCNT1.5

93 93 93 92 96 45

91 92 91 91 96 46

2.1 4.7 7.8 9.9 9.5 16.8

720 621 520 293 420 254

Calculated from ER=100% × (εmax-ε (0,εmax))/εmax, where εmax is the maximum strain and ε(0,εmax) is the strain at zero stress in the cycle after the maximum strain εmax. bDetermined by monotonic stress− strain curves. a

and then necking and strain hardening occur until fracture at an approximate strain of 250% (Figure 7a). The ultimate tensile strength of MWCNT-g-P(BA-co-MMA) CTPEs increased with increasing MMA content, whereas the elongation at break decreased as the MMA content increased, indicating that the MMA content is crucial for the tensile properties of the CTPEs. This result is in accordance with the change in Tg for MWCNT-g-P(BA-co-MMA) copolymers as discussed above. The mechanical properties of the CTPEs are also obviously related to the MWCNT content. Figure 7b shows the nominal stress−strain curves of CTPEs with varying MWCNT contents. It was found that, at the same strain, the tensile stress of the CTPEs increased with increasing MWCNT content, indicating that the tensile modulus of the CTPEs can obviously be enhanced through the introduction of MWCNTs in the copolymer matrix. On the other hand, the strain at break tended to decrease with increasing MWCNT content. The monotonic nominal stress−strain curves of CTPEs with similar low MWCNT contents and their linear counterparts are displayed in Figure 8. The curves of the CTPEs are all above the curves of their corresponding linear counterparts. Strain hardening can be clearly seen for the CTPEs, whereas strain hardening was relatively weak for the linear P(BA-co-MMA) copolymers. The tensile stress values at break for the CTPEs were much higher than those for their linear counterparts (Figure 8e). The changing trends in the stress at break shown in Figure 8e obviously demonstrate that the chemically bonded MWCNTs in the CTPEs play a vital role as well for strong enhancements in mechanical properties, apparently through the

Figure 6. TEM micrographs of samples (a) BA6500-MWCNT1.2 and (b) BA6500-MWCNT3.8.

typical TEM micrographs of samples BA6500-MWCNT1.2 and BA6500-MWCNT3.8 (with a high MWCNT density in Figure 6b), which reveal the well-distributed nanotubes within the copolymer matrix phase. Note that the fact that the MWCNTs dispersed as individual nanotubes in the matrix without aggregation is in agreement with the SEM results, demonstrating that the graft copolymerization method is efficient for synthesizing MWCNT-g-P(BA-co-MMA) copolymers. The enhancement of the mechanical properties of CTPEs depends strongly on the extent of MWCNT dispersion and the strength of interfacial adhesion between the nanotubes and the polymers because the stress can be easily transferred between stiff nanotubes and the grafted copolymer matrix, which will result in significant influences on the mechanical properties of the CTPEs even though the MWCNT contents are less than 4.0%, as discussed in a later section of this article. 3.3. Strongly Enhanced Mechanical Properties of MWCNT-g-P(BA-co-MMA) CTPEs. The monotonic nominal stress−strain curves of the CTPEs are shown in Figure 7 and clearly indicate the elastomeric behavior of the CTPEs. The mechanical properties of the CTPEs and their linear copolymer counterparts are summarized in Tables 2 and S2 (Supporting Information), respectively. For sample BA5500-MWCNT1.5, the stress−strain curve is linear up to a yield point of 4% strain,

Figure 7. Monotonic nominal stress−strain curves for MWCNT-g-P(BA-co-MMA) copolymers synthesized with (a) varying monomer feed ratios and similar MWCNT contents and (b) the same monomer feed ratio and varying MWCNT contents. 20159



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Figure 8. (a−d) Monotonic nominal stress−strain curves for samples (a) BA7000 and BA7000-MWCNT1.3, (b) BA6500 and BA6500MWCNT1.2, (c) BA6000 and BA6000-MWCNT1.4, and (d) BA5500 and BA5500-MWCNT1.5. (e,f) Changes in (e) stress at break and (f) strain at break with MMA content for MWCNT-g-P(BA-co-MMA) CTPEs and linear P(BA-co-MMA) copolymers.

progressively larger because of the plastic deformation due to the Mullins effect, which is known as stress softening.50 Figure 9b shows the nominal stress−strain curves during the second step cyclic tensile deformation series for the same sample, BA6500-MWCNT1.2, that had experienced the first step cyclic tensile deformation series shown in Figure 9a. Note that, for the second series, the strain values (and sample area) were based on the dimensions of the sample at the beginning of this series; in other words, the sample length and area were measured again after the first series. It can be noticed from Figure 9a,b that the shape of the ascending curves is concaveupward, the hysteresis in each cycle decreases, and this trend becomes more obvious during the second step cyclic tensile deformation series. Similar phenomena were also observed for the other CTPEs, as shown in Figure S4 (Supporting Information). Note that the areas of these samples were nearly unchanged after the first step cyclic deformation series. However, in the case of sample BA5500-MWCNT1.5, the

introduction of strain-hardening effects for the CTPEs. In general, strain hardening has been observed for polymer elastomers with fillers such as carbon blacks or carbon nanotubes added by blending.37,40 However, much higher contents of added fillers are required for this purpose. The results in Figure 8f further indicate that the strain at break decreased when MWCNTs were introduced into the copolymer matrix for the CTPEs. Therefore, it was worth investigating the elastic recovery properties of the CTPEs by applying step cyclic tensile tests, as discussed next. Elasticity is crucial for applications of thermoplastic elastomers, and step cyclic tensile tests were applied to examine the elastic recovery properties of the CTPEs. Figure 9a shows the nominal stress−strain curves for sample BA6500MWCNT1.2 during the first step cyclic tensile deformation series with the maximum strain values increasing from 25% to 50%, 75%, 100%, 125%, 150%, and so on up to 300%. Note that, in each step cycle, the residual strain at zero stress was 20160



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Figure 9. (a−d) Nominal stress−strain curves for samples (a,b) BA6500-MWCNT1.2 and (c,d) BA5500-MWCNT1.5 during the (a,c) first and (b,d) second cyclic tensile deformation series. (e,f) Changes in elastic recovery as a function of maximum strain during the (e) first and (f) second cyclic tensile deformation series for the CTPEs. The short dashed lines in panels e and f indicate an elastic recovery value of 90%. (g) Changes in average ER with MMA content for the CTPEs and P(BA-co-MMA) copolymers.

residual strain at zero stress rose quickly during the first series, as shown in Figure 9c, because of its large plastic deformation. After the first cyclic deformation series, the sample area changed significantly as a result of plastic deformation, leading to a significant increase in stress during the second cyclic

deformation series, as shown in Figure 9d, where the asterisk indicates fracture of the sample. The changes in elastic recovery (ER) during the first and second cyclic tensile deformation series with maximum strain, εmax, are shown in Figure 9e,f. ER is defined as ER = {[εmax −ε(0,εmax)]/εmax} × 100% , where εmax and ε(0,εmax) are the 20161



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Figure 10. Changes in true plastic (εH,p) and elastic (εH,e) strains as functions of the true maximal strain (εHmax) for samples (a,b) BA6500MWCNT1.2 and (c,d) BA5500-MWCNT1.5 during the (a,c) first and (b,d) second cyclic deformation series.

enhancements in both the tensile strength and elastic recovery in this type of CTPE. The true plastic strain, εH,p (strain remaining when unloaded to zero stress), and the true elastic strain, εH,e (strain recovered after unloading), can be extracted from the true stress versus true strain curves according to the literature, and the maximum true strain of cycle N is denoted as εHmax(N).51,52 Figures 10 and S6 (Supporting Information) show the evolutions of the true plastic and elastic strains as functions of the maximum true strain in each cycle for the CTPEs. The true plastic strain was low at small deformation and increased as the maximum true strain increased. It can be more clearly seen that the true elastic strain values were all higher than the true plastic strain values (except for sample BA5500-MWCNT1.5), indicating sufficient elasticity for the CTPEs, which is in accordance with the high elastic recovery values. For sample BA5500-MWCNT1.5, the true plastic strains were higher than the true elastic strains, indicating that the tensile deformation was mainly dominated by plastic deformation. The evolutions of the true plastic and elastic strains as functions of the maximum true strain for linear P(BA-co-MMA) copolymers are shown in Figure S7 (Supporting Information). For sample BA7000, the true elastic strain values were all higher than the true plastic strain values, consistent with its high elastic recovery values (see Figure S5, Supporting Information). However, for samples BA6000 and BA5500, the true plastic strain values were all much higher than the true elastic strain values, consistent with their high plastic deformation behaviors. The data points for sample BA6500 show a crossover, indicating an intermediate elastic recovery performance for this sample. Overall, the linear P(BA-co-MMA) copolymers showed much less elastic deformation than their corresponding CTPEs.

maximum strain and the strain at zero stress in the cycle after the maximum strain, respectively. For the CTPEs, a significant increase in elastic recovery was observed after several initial loading and unloading cycles during both the first and second cyclic deformation series, and the ER values could be raised well above 90% (Figure 9g), which implies that the copolymer chains can return to their initial configurations with minor plastic deformation with the assistance of the chemically bonded MWCNTs in the CTPEs. The average ER values during the second cyclic deformation series were similar to those during the first series, as listed in Table 2. Again, it is noticed that the ER values for sample BA5500-MWCNT1.5 were much lower than those for the other samples (Figure 9g). This result can be explained by the much higher MMA content in this sample with a higher Tg value because the chains in the amorphous state with higher Tg values are less easy to return to their random coil configuration after stress is released. The nominal stress−strain curves and changes in elastic recovery with maximum strain in each step cycle during the first and second cyclic tensile deformation series for linear P(BA-coMMA) copolymers are shown in Figure S5 (Supporting Information), demonstrating that the linear P(BA-co-MMA) copolymers have significantly larger plastic deformation than the MWCNT-g-P(BA-co-MMA) CTPEs during both the first and second step cyclic tensile deformation series. Therefore, the ER values of the linear P(BA-co-MMA) copolymers are much lower than those of the MWCNT-g-P(BA-co-MMA) CTPEs (Figure 9g; Tables 2 and S2, Supporting Information). Although BA7000 seems to have a reasonably high elastic recovery value, its nominal stress of less than 0.4 MPa is weakest among all of the samples. This comparison demonstrates the significant role that stiff MWCNTs chemically bonded to the copolymer matrix play in making strong 20162



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Figure 11. (a) Nominal stress−strain curve for sample BA6500-MWCNT1.2 during tensile deformation (loading) and recovery (unloading). (b) Photographs of sample BA6500-MWCNT1.2 taken at different strains during loading (A, strain = 0%; B, strain = 400%) and unloading (C, strain = 60%). (c,d) Selected SAXS patterns of sample BA6500-MWCNT1.2 taken at different strains during (c) loading and (d) unloading. The stretching direction is horizontal.

Figure 12. (a,b) Typical azimuthal scattering intensity profiles at different strains for sample BA6500-MWCNT1.2 during (a) loading and (b) unloading. (c) Variation in the orientational order parameter, ⟨P̅2⟩ for MWCNTs in sample BA6500-MWCNT1.2 during loading and unloading. The arrows indicate the loading and unloading processes. The solid lines are guides to the eyes. (A,B) Schematic illustrations of physical crosslinking networks in MWCNT-g-P(BA-co-MMA) CTPEs prior to loading (cartoon A) and during loading (cartoon B). The stretching direction is horizontal. The gray rods represent MWCNTs, and phase domains in red circles represent the physical cross-linking points due to aggregation of MMA chain segments. MWCNTs provide additional multiple cross-linking points through grafting.

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orientation along the director. Prior to calculating ⟨P̅2⟩, the Maier−Saupe distribution function given by

3.4. Evolution of Microstructure during Tensile Deformation of MWCNT-g-P(BA-co-MMA) CTPEs by SAXS Measurements. Investigating the microstructural evolution of CTPEs during tensile deformation is of great interest from both technological and scientific viewpoints. Thus, in situ SAXS measurements were performed during tensile deformation (loading) and subsequent recovery (unloading) processes. As shown in Figure 11a, sample BA6500-MWCNT1.2 was subjected to one cycle of consecutive loading and unloading, resulting in a residual strain of 60%. As displayed in Figure 11b, sample BA6500-MWCNT1.2 was black due to graft copolymerization of MMA and BA monomers on MWCNTs, and the sample could be stretched up to several times the initial length and then retract back, indicating its sufficient elastic recovery properties. The nominal stress−strain curves under the loading and recovery paths look quite different, which is termed hysteresis.53 The lack of a complete recovery can be explained by slippage of the trapped entanglements in the soft copolymer domains during tensile deformation. This type of slippage results in a delay of failure of the glassy MMA domains (high Tg) and stiff MWCNT domains in the sample, leading to the observed elongation. Because the slippage of chains is not instantaneously reversible, the overall recovery of the sample is affected. Selected SAXS patterns of sample BA6500-MWCNT1.2 taken at various strains during the loading and unloading processes are shown in panels c and d, respectively, of Figure 11. An initial isotropic SAXS pattern (0% strain in Figure 11c) was observed before deformation, indicating that the MWCNTs were randomly distributed in the copolymer matrix. As the strain increased during loading, the SAXS pattern changed into an ellipsoidal form with its short axis along the stretching direction (horizontal direction in Figure 11c), indicating that the MWCNTs were forced to become parallel to the stretching direction. It must be emphasized that, during loading, the linear P(BA-co-MMA) copolymers did not show any detectable SAXS patterns because of the absence of electron density contrast in these samples. During subsequent unloading, the SAXS pattern gradually recovered almost to the initial state as strain decreased (Figure 11d), indicating a sufficient elastic recovery for the CTPEs. The changes in the orientational order parameter during loading and unloading for the CTPEs are analyzed in detail next. The azimuthal scattering intensity profiles at different strains during the loading and unloading processes, shown in Figure 12a,b, were extracted from the SAXS patterns (Figures 11c,d) by integration at a scattering vector of q = 0.043 nm−1 after subtraction of the air scattering intensity. Before deformation, the azimuthal scattering intensity profile did not show any peak, indicating a random orientation of MWCNTs in the sample (Figure 12, cartoon A), whereas peaks in the azimuthal scattering intensity appeared at ϕ ≈ 90° and 270° (not shown in Figure 12) when the sample was stretched, indicating that the MWCNTs were forced to become parallel to the stretching direction during loading (Figure 12, cartoon B). The azimuthal scattering intensity peaks became narrower with increasing strain, indicating that the orientation of the MWCNTs became stronger at higher strain. The orientational order parameter, ⟨P̅2⟩, can be estimated from the azimuthal scattering intensity profile, which represents the distribution of molecular orientations around the director. The values of ⟨P̅2⟩ range between 0 and 1, with the former corresponding to an isotropic structure and the latter to perfect

I = I0 + A exp[β cos2(ϕ − ϕ0)]

(3)

where I0 is the baseline intensity, β is a parameter that determines the width of distribution, ϕ is the azimuthal angle, and ϕ0 is the azimuthal angle at the position with the maximum intensity, was used to fit the azimuthal scattering intensity profile.54 The solid lines in Figure 12a,b are the fitted curves, from which β can be obtained, and then ⟨P̅2⟩ can be obtained as54,55 1

⟨P2̅ ⟩ =

∫−1 P2(cos ϕ) exp(β cos2 ϕ) d cos ϕ 1

∫−1 exp(β cos2 ϕ) d cos ϕ

(4)

where the function P2(cos ϕ) is the second-order Legendre polynomial of cos ϕ, given by P2(cos ϕ) =

1 (3 cos2 ϕ − 1) 2

(5)

Equation 5 can be solved by numerical integration. A relatively high value of ⟨P̅2⟩ = 0.65 was obtained for sample BA6500-MWCNT1.2 at a strain of 400%, indicating a high orientation degree of MWCNTs in the copolymer matrix due to stretching. The changes in the orientational order parameter of the MWCNTs with strain during loading and unloading are shown in Figure 12c. Obviously, the MWCNTs in the copolymer matrix oriented gradually with their longitudinal axis parallel to the stretching direction as the strain increased during loading, indicated by ⟨P̅2⟩ increasing from 0 to 0.65. However, during unloading, ⟨P̅2⟩ dropped more rapidly as strain decreased. An induction period for orientation of the MWCNTs cannot be seen in Figure 12c, because the dispersion of MWCNTs in the copolymer matrix was homogeneous, as demonstrated by TEM and SEM observations, and the individual isolated MWCNTs at such a low content (1.2 wt %) with grafted P(BA-co-MMA) chains facilitated their orientation during loading and disorientation during unloading. The ⟨P̅2⟩ values were much lower at the same strains during unloading than during loading, indicating that the orientation of MWCNTs depends on deformation. Much higher stress is needed to orientate MWCNTs to a higher orientation degree, as strain increases during loading, bringing out the strainhardening behavior. It is well-known that, in the case of elastomers, the major effect of loading is a stretching of the network chains, which substantially reduces entropy. Thus, the retroactive force arises primarily from the tendency for the system to increase its entropy toward the maximum value it stored in the initial state. 56 Therefore, the stress decreases rapidly during unloading, as shown in Figure 11a. Figure 12c shows that, after stress release (with 60% plastic strain), the MWCNTs did not fully recover to the initial state (⟨P̅2⟩ = 0.03), indicating that the MWCNTs still remained slightly oriented after the loading and unloading processes. For most thermoplastic elastomers, the well-ordered microphase-separated morphology provides acceptable mechanical properties, whereas, in the case of densely grafted copolymers or multigraft copolymers, a large number of branched points is even more important.57 For MWCNT-g-P(BA-co-MMA) CTPEs, MWCNTs are homogeneously dispersed in the grafted copolymer matrix through chemical bonding, and these nicely 20164



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clarify the relationship between the microstructures and mechanical properties. The microstructural evolutions during loading and unloading of the CTPEs as investigated by in situ SAXS indicate that the homogeneously dispersed MWCNTs in the copolymer matrix are oriented gradually with their longitudinal axis parallel to the stretching direction as strain increases. Because of successful stress transfer from the grafted copolymer chains to the chemically bonded MWCNTs, which are orientated during tensile deformation, an apparent strainhardening behavior can be observed for the CTPEs. The CTPEs further demonstrate obviously enhanced elastic recovery properties as compared with their linear P(BA-coMMA) copolymer counterparts because of the existence of physical cross-linking networks from MMA chain segment aggregation and additional multiple physical cross-linking points of MWCNTs that are connected by grafted P(BA-coMMA) copolymer chains. The CTPEs also show applicable thermal stability. The results in this work confirm that the graft copolymerization of comonomers on stiff nano-objects with low amounts serving as multiple physical cross-linking points can achieve high-performance composite thermoplastic elastomers. This efficient and robust synthetic approach can be further applied to the synthesis of other similar types of highperformance composite thermoplastic elastomers.

dispersed stiff nanotubes can also act as multiple physical crosslinking points. Because of the high stiffness and high aspect ratio of the nanotubes, the synthesized MWCNT-g-P(BA-coMMA) CTPEs in this work exhibited strongly enhanced mechanical tensile properties, such as much higher tensile strength and higher elastic recovery, as compared with their linear P(BA-co-MMA) copolymer counterparts. We emphasize here that the MWCNT contents in the CTPEs were not necessarily very high for the purpose of enhancing the mechanical properties. The MWCNT contents were less than 4.0% in this work. Noticeably, both the MMA content and homogeneously dispersed MWCNTs showed significant effects on the mechanical performances of MWCNT-g-P(BA-co-MMA) CTPEs as thermoplastic elastomers. The tensile strength and elongation at break of the CTPEs are related to Tg, which can be determined from the MMA content in the grafted copolymer chains. During tensile deformation, the initial tensile strength mainly depends on the MMA content, because the Young’s modulus (influenced by Tg) increases with increasing MMA content. If the MMA content is low, sparse glassy MMA domains exist in the bulk material, and the soft chains have the ability to dissipate energy, leading to a high elongation during tensile deformation. Because the mechanical properties of the CTPEs are actually governed by MMA content (MMA chain segment aggregation as physical cross-linking points as schematically illustrated in Figure 12, cartoon A) and introduction of MWCNTs into the CTPEs as multiple physical cross-linking points (Figure 12, cartoon A), it is necessary to illustrate the roles that MMA and MWCNTs play at different tensile deformation stages. At the initial stage (small strain), the stress needed is low because the deformation is attributed to the random copolymer chains. At subsequent high strain, the physical cross-linking points by MMA chain segment domains and stiff MWCNTs begin to play roles together, causing an obvious increase of stress. Because of the strong chemical bonding between the MWCNTs and the copolymer chains, further tensile deformation leads to strong orientation of the MWCNTs along the stretching direction, and the stress can be transferred to the stiff MWCNTs because of the chemically bonded interfaces between the MWCNTs and P(BA-co-MMA) copolymer chains, providing strong strain hardening (Figure 12, cartoon B). Ultimately, further deformation leads to failure of the CTPE materials. In the CTPE materials, the cross-linking networks from both MWCNTs and MMA chain segment aggregation can give high elastic recovery properties to the CTPE materials that are absent for their linear P(BA-co-MMA) copolymer counterparts.

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ASSOCIATED CONTENT

S Supporting Information *

Additional experimental details and supplementary data. This material is available free of charge via the Internet at http:// pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +86 0551-63607703. Fax: +86 0551-63607703. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Z.W. acknowledges financial support from the National Science Foundation of China (Grant 51473155) and the National Basic Research Program of China (Grant 2012CB025901). Professor Yongfeng Men at Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, is acknowledged for providing use of the SAXS facility.

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REFERENCES

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4. CONCLUSIONS In summary, we report enhanced mechanical properties of copolymer-grafted MWCNT composite thermoplastic elastomers (CTPEs) synthesized by the surface-initiated ARGET ATRP grafting strategy. In the CTPEs, the stiff carbon nanotubes act as minority stiff nano-object domains (only 1.2−3.8 wt %), and the grafted P(BA-co-MMA) random copolymer chains serve as the polymer matrix. A wide spectrum of CTPEs with enhanced mechanical properties, such as higher tensile strength and higher elastic recovery properties, can be prepared by simply adjusting the feed molar ratio of the monomers for copolymerization and the content of introduced MWCNTs. The role of MWCNTs as additional cross-linking points in the CTPEs was comprehensively investigated to 20165



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Synthesis and Characterization of Nanostructured Copolymer-Grafted

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