Noncovalent Functionalization of Multiwalled Carbon Nanotube by a

Article pubs.acs.org/JPCC

Noncovalent Functionalization of Multiwalled Carbon Nanotube by a Polythiophene-Based Compatibilizer: Reinforcement and Conductivity Improvement in Poly(vinylidene fluoride) Films Amit Mandal and Arun K. Nandi* Polymer Science Unit, Indian Association for the Cultivation of Science, Jadavpur, Kolkata 700 032, India S Supporting Information *

ABSTRACT: A new compatibilizer (P2) containing thiophene moiety and poly(dimethylamino ethyl methacrylate) (PDMAEMA) group is prepared by atom transfer radical polymerization (ATRP). The dispersion of multiwalled carbon nanotube (MWNT) with P2 in N,N-dimethylformamide (DMF) is stable for >3 months. The UV−vis spectrum of the MWNT/P2 dispersion shows a blue shift, and the photoluminescence (PL) spectrum also indicates a quenching suggesting a significant interaction between MWNT and P2. SEM and TEM micrographs suggest a good dispersion and coating of P2 on MWNT surface. The Raman D/G band intensity ratio remains unchanged, and a small shift of >CO vibration peak to lower energy suggests an interaction between P2 and MWNT. The presence of CH−π and π−π interactions between P2 and MWNT is evidenced from UV−vis, PL, Raman, FTIR, and NMR spectral results facilitating the wrapping of P2 on MWNT surface. Poly(vinylidene fluoride) (PVDF)/P2/MWNT composites (PCMx, x indicates wt % of MWNT) are prepared by solvent casting method with 0.02% (w/v) P2 and varying concentration of MWNT in DMF. The SEM and TEM micrographs show a homogeneous dispersion of MWNT/P2, and the optical micrographs indicate a loss of spherulitic morphology in the composites. FTIR spectra suggest the formation of piezoelectric β-phase PVDF and the presence of specific interaction between >CO group of P2 and >CF2 group of PVDF. The storage modulus (G′) shows a highest increase (122%) in PCM0.5 among the PVDF/MWNT composites. The increase in Young’s modulus, tensile strength, and toughness in PCM 0.05 is two to three times higher than that without P2. Analysis of Young’s modulus data suggests a random distribution of MWNT in the composite. The PCM1 has the highest conductivity (2 × 10−2 S/cm), and the system shows a very low percolation threshold at 0.06 wt % MWNT. The conductivity data obeys the 3-D percolation model. By comparing the above mechanical property and conductivity data with that of other MWNT/PVDF composites, it is argued that the noncovalently functionalized MWNT with P2 is a better reinforcing agent.

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INTRODUCTION A high aspect ratio, low mass density, and high mechanical strength render the carbon nanotubes (CNTs) as an ideal reinforcing filler producing an ultrastrong lightweight polymer nanocomposite.1 More intriguing is that the excellent thermal, mechanical, and novel electrical properties of CNTs afford great opportunities for fabricating multifunctional materials.2−5 The usual poor improvement in the mechanical properties of CNT/polymer composites can be attributed to both poor nanotube dispersion or exfoliation and to weak nanotubes/ matrix load transfer. This is because of strong van der Waals forces among the nanotubes and the chemical inertness of the pristine CNTs inducing weak interfacial adhesion with the polymer matrix.3−6 These result in intertube slippage with applied stress, making a poor load transfer from the matrix to the nanotubes. To exploit the high mechanical properties of CNTs in the composites, the nanotubes are to be welldispersed into the matrix, and the nanotube/matrix interface interaction is to be strong. The scientists working in this area © 2012 American Chemical Society

are still facing challenges for homogeneous dispersion of CNTs on a single nanotube level and a good alignment in the polymer matrix strengthening the CNT/polymer interfacial interaction.6−10 The surface functionalization of CNTs is an effective means to improve the dispersion of CNTs, and, more importantly, it engineers the nanotube/polymer interface at a molecular level required for strong interfacial interactions.11−16 Both covalent and noncovalent approaches are attempted for the surface modification of CNTs.13−28 The covalent approach is a good method in functionalizing CNTs that can moderately improve the mechanical properties of the composites; it is usually achieved by complicated chemical reactions disrupting the longrange π conjugation of CNTs and also making it of shorter length. This leads to the decreased electrical conductivity and Received: March 1, 2012 Revised: April 6, 2012 Published: April 9, 2012 9360

dx.doi.org/10.1021/jp302027y | J. Phys. Chem. C 2012, 116, 9360−9371

The Journal of Physical Chemistry C

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diminished mechanical properties of the composites.5,13 The high density of functionalization needed for effective stress transfer disrupts the extended π conjugation and is typically not effective for individually dispersing high concentration of CNTs in the polymer matrix.16−20 Noncovalent functionalization using polymer or surfactant has been shown to be effective in dispersing the nanotubes by utilizing the multiple weak interactions such as van der Waals interactions, π−π interactions, electrostatic interactions, and so on. Such noncovalent interactions avoid the destruction of the chemical structure and retain the electric and mechanical properties of CNTs to a full extent.21−29 Recently, two types of materials have attracted a significant amount of attention as dispersants of CNTs. One type of these dispersing materials is block copolymers,24,25 and the solubility of the CNTs can be tuned by the composition of the block copolymers. Other types of these materials are conjugated oligomers and polymers with anchored functional groups dispersing the CNTs by π−π interactions with the conjugated part and specific interaction of the matrix polymer with the anchored functional groups. Therefore, it plays a role of compatibilizer between CNTs and a matrix polymer, and this method has been gaining popularity recently.26−30 Conjugated and aromatic condensation polymers, for example, polythiophene, 27−30 polybenzimidazole, poly(phenylene ethynylenes),31and polyimide,32,33 interact with CNTs via strong π−π interaction. Zhai et al. have extensively functionalized the CNTs using conjugated block copolymers containing nonconjugated blocks with tunable functionality. They have dispersed the CNTs uniformly in various solvents with conjugated block copolymers, for example, poly(3hexylthiophene) (P3HT)-b-polystyrene (PS), P3HT-b-poly(methyl methacrylate) (PMMA), P3HT-b-poly(acrylic acid) (PAA), and so on. In the dispersion process, conjugated P3HT blocks become attached to CNTs through π−π interactions, whereas the nonconjugated blocks located at the outermost surface of CNTs enhance the solubility of CNTs in different solvents.24,25 In another typical work, CNTs were noncovalently functionalized by poly(vinyl benzyl oxyethyl naphthalene)-g-poly(methyl methacrylate) and is dispersed in the matrix of poly(styrene-co-acrylonitrile).28 The resultant composites displayed significantly improved mechanical and electrical properties as compared with pristine CNTs. So, the preparation of multifunctional materials requires an effective compatibilizer, having favorable interaction between the surface of CNTs and the host polymer, simultaneously. Poly(vinylidene fluoride) (PVDF) is a semicrystalline electroactive polymer and is technologically important because of its piezo and pyroelectric properties.34,35 It has a wide range of applications in supercapacitors, actuators, batteries, membranes, and in different optoelectronic devices.36,37 It has five different crystalline polymorphs (α, β, γ, δ, ε,), and βpolymorph is the most important for its piezo and pyroelectric properties.34,38−40 Recently, PVDF composites with different nanomaterials (e.g., clay, metal nanoparticles, CNTs, and graphene) have demonstrated the formation of a piezoelectric β-polymorph with a highly improved mechanical and electrical properties.37−45 So, nanocomposites of PVDF and modified CNTs are gaining importance for the development of new functional materials.43−45 In this Article, we report the design and synthesis of a new reactive noncovalent polymeric compatibilizer for multiwalled carbon nanotube (MWNT) and PVDF. The new compatibil-

izer is composed of a polythiophene backbone grafted with poly(dimethylamino ethyl methacrylate) (PDMAEMA) and is synthesized following Scheme 1. The polythiophene backbone Scheme 1. Synthesis of Compatibilizer (P2) and Schematic Illustration of Noncovalent Functionalization of MWNTs by P2

strongly interacts with the MWNT surface via π−π interaction and PDMAEMA interacts with the PVDF matrix, facilitating a good dispersion, as evidenced from FE-SEM and HR-TEM data. The interaction between MWNTs and compatibilizer is examined by Raman, FT-IR, 1H NMR, UV−vis and PL spectra, and thermal data. The mechanical and conducting properties of the composites are measured as a function of MWNT content in the composites, and attempts are made to explain the results with the existing theories.

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EXPERIMENTAL SECTION Materials. 3-Thiophene ethanol (99%), triethylamine, 2bromoisobutyryl bromide, copper(II) bromide (99%), and 1,1,4,7,10,10-hexamethyltriethylenetetramine (HMTETA) (99%) were purchased from Aldrich Chemical and were used as received. Copper(I) bromide (Aldrich) was purified by washing with 10% HCl, followed by methanol and diethyl ether in a Schlenk tube under a nitrogen atmosphere and stored. The monomer 2-(dimethylamino ethyl methacrylate) (DMAEMA) (Aldrich) was purified by passing through basic alumina column prior to use in the polymerization reaction. Methylene chloride, chloroform, and anisole used as solvent were purchased from Rankem (Kolkata, India) and were dried by refluxing over CaH2, followed by distillation, and stored in a N2 atmosphere. PVDF (Mw = 180 000, PDI = 2.57, (H−H) defect = 4.33 mol % measured from 19F NMR spectra) and MWNT (purity 95%, length 0.5−20 μm, and outer-diameter 20−30 nm) were purchased from Aldrich Chemical. The polymer was recrystallized from a dilute solution (0.3% w/v) in acetophenone. Synthesis of Compatibilizer. Synthesis of 3-[1-Ethyl2(2-bromoisobutyrate)] Thiophene (EBIBT). The overall synthesis route is presented in Scheme 1. In a 250 mL two9361



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necked round-bottomed flask, 3-thiophene-ethanol (3.48 g, 27 mmol) and triethylamine (3.65 g, 36.5 mmol) were dissolved in 75 mL of anhydrous CH2Cl2 under a N2 atmosphere with stirring and cooled in an ice/water bath. Then 2-bromoisobutyryl bromide (8.32 g, 36.8 mmol), dissolved in 15 mL of CH2Cl2 and transferred into a pressure-equalizing funnel was added dropwise to the flask over 30 min. After complete addition, the reaction mixture was stirred for 24 h at 30 °C. The white precipitate was filtered, and the filtrate was diluted to 100 mL with CH2Cl2 and was washed with 1% HCl, saturated NaHCO3, and distilled water. The organic layer was collected and dried with anhydrous Na2SO4. The product was purified by a silica column chromatography in a mixture of hexane/ethyl acetate (95:5, v/v) as an eluting solvent. The final product was dried under vacuum, and the yield was 5.55 g (73.5%). Copolymerization of EBIBT and 3-Hexylthiophene [P(EBIBT-co-HT)] (P1). Anhydrous FeCl3 (11.05 g, 66.4 mmol) was dispersed in 50 mL of anhydrous CHCl3 in a 100 mL twonecked round-bottomed flask under a N2 atmosphere. EBIBT (2.32 g, 8.3 mmol) and 3-hexylthiophene (HT) (1.40 g, 8.3 mmol) were dissolved in 20 mL of CHCl3 in a pressureequalizing funnel and were added dropwise to the flask over 30 min. After addition, the mixture was stirred at 30 °C for 2 days. The reaction mixture was poured in 1 L of methanol and stirred for another 3 h. The solid was collected on a thimble and was washed with methanol four times. The resulting precipitate was extracted with methanol by soxheletion for 3 days and was then dried under vacuum at room temperature overnight. The darkbrown polymer obtained was dissolved and is refluxed in a mixture of 250 mL of CHCl3 and 100 mL of concentrated ammonia for 2 days to remove the unused FeCl3 in the polymer product. The solution was concentrated and was precipitated in 100 mL of methanol. The red precipitate was washed with methanol by three to four times and was dried under vacuum at room temperature overnight (Yield = 1.95 g (26.3%)). The polymer P(EBIBT-co-HT) was then characterized by gel permeation column chromatography (GPC) (Waters) using PS standard and μ-styragel column with THF as eluent. Synthesis of 2,5-Poly[(3-{1-ethyl-2(2-(poly(N,N-dimethylaminoethylmethacrylate))} thiophene)-co-hexylthiophene] [P(EBIBT-co-HT)-g-PDMAEMA] (P2). In a 100 mL schlenck flask (flask A), P(EBIBT-co-HT) macroinitiator (100 mg, 0.22 mmol Br) was dissolved in 15 mL of anisole with stirring. Then, CuBr2 (2.4 mg, 10.2 × 10−3 mmol) was added and was deoxygenated by three freeze−thaw cycles. In another 100 mL round-bottomed flask (flask B), CuBr2 (2.44 mg, 10.3 × 10−3 mmol), CuBr (28.2 mg, 0.21 mmol), and HMTETA (100 μL, 0.43 mmol) are charged, and then DMAEMA (5 mL, 32 mmol) was added to flask B. The mixture was deoxygenated by N2 bubbling for at least 1 h. The mixture in flask B was then transferred into flask A, and flask A was deoxygenated by three cycles of freeze−thaw and was stirred for 1 h in an oil bath at 80 °C. Then, the reaction mixture was exposed to air and diluted with 100 mL of THF. The reaction product was passed through an Al2O3 column to remove the copper catalyst. The resulting polymer solution was concentrated and was precipitated in 100 mL of hexane. The viscous precipitate was redissolved with 5 mL of THF and was precipitated in 100 mL of hexane. The precipitation process was repeated; finally, the polymer P(EBIBT-co-HT)-g-PDMAEMA was dried in vacuum at 60 °C for 3 days (yield = 1.2 g). Preparation of PVDF/P2/MWNT Nanocomposites Films. Nanocomposite films of PVDF/P2/MWNT with

different loadings of MWNTs in the presence of compatibilizer (P2) were prepared by solution blending method. A typical procedure of PVDF/P2/MWNT nanocomposites containing 0.2 wt % MWNT is as follows: A dispersion of 5 mg of P2 and 2 mg of MWNT in 25 mL of N,N-dimethylformamide (DMF) was sonicated for 1 h. Then, the dispersed solution was added to a solution containing 1 g of PVDF in 25 mL of DMF. The mixture was sonicated for an additional 30 min and was precipitated in excess methanol to remove ungrafted compatibilizer. The precipitate was collected by filtration and was washed with methanol. The filter cake was dried at 60 °C under vacuum for 3 days. For the formation of composites of different MWNTs concentration, the compatibilizer concentration is kept fixed (5 mg in 25 mL of DMF), but MWNT concentration is varied, and the composites are designated as PCM0.05, PCM0.2, PCM0.5, and PCM1, where the number indicates weight percentage of MWNTs in the composites. Nanocomposite films were cast from DMF on Teflon dishes at ∼100 °C for solvent evaporation in a controlled air stream for 6−8 h. The resulting films were further dried in vacuum at 60 °C for 3 days to remove DMF completely. Characterization. Microscopy. The dispersion of functionalized MWNTs with compatibilizer into the polymer matrix was studied using a transmission electron microscope (TEM, JEOL 2010EX) operated at an accelerated voltage of 200 kV. A drop of the dilute solution of the homogeneous mixture on the carbon-coated copper grid was initially dried in air at 30 °C and finally in vacuum at 30 °C for 2 days. The surface morphology of the films was observed through a field-emission scanning electron microscope (FE-SEM, JEOL GSM-5800) at an acceleration voltage of 5 kV after platinum coating. The texture of the sample was also studied using a polarized optical microscope (Leitz Biomed) fitted with a digital camera. Spectroscopy. 1H NMR spectra were measured from a Bruker 300 MHz spectrometer using CDCl3 as the solvent. FTIR spectra of the samples were obtained from thin films cast from their hybrid solution of DMF using a Perkin-Elmer FT-IR instrument (spectrum 100). The UV−vis spectra were taken in DMF solutions of the samples from 190 to 1100 nm using a UV−vis spectrophotometer (Hewlett-Packard, model 8453) at 30 °C. The fluorescence spectra were obtained from a Fluoromax-3 instrument (Horiva Jovin Yvon). Raman spectra were recorded with a LabRam HR800 high-resolution Raman spectrometer (Horiba-Jobin Yvon) using a 632.8 nm wavelength laser source. XRD Study. The wide-angle X-ray scattering (WAXS) experiments of the samples were carried out in a Bruker AXS diffractomer (model D8 Advance) using a Lynx Eye detector. The instrument was operated at a 40 kV voltage and at a 40 mA current. The samples were scanned in the range of 2θ = 5−35° at the scan rate 0.5 s/step with a step width of 0.02°. Thermal Study. The thermal analysis of the samples was performed in a Perkin-Elmer differential scanning calorimeter (Diamond DSC-7 with Pyris software) under a N2 atmosphere. The instrument was calibrated with indium standard before each set of experiments. The solvent-cast samples (3 to 4 mg) were taken in aluminum pan and were heated from −50 to 210 °C at a heating rate 40 °C/min under a nitrogen atmosphere. The sample was then kept at 210 °C for 5 min and was cooled at the cooling rate of 5 °C/min to −50 °C, where it was kept for 10 min for crystallization. Then, the samples were again heated to 210 °C at the heating rate 20 °C/min. The melting (Tm) and crystallization (Tc) temperatures were taken as the 9362



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peak temperature and the glass-transition temperature (Tg) was recorded as the inflection point of the heat-capacity jump from the second heating curve. The crystallinity of the samples was calculated from the enthalpy of fusion obtained from the thermogram by dividing with the enthalpy of fusion of perfect PVDF crystal (104 J/g).36 The thermogravimetric analysis (TGA) experiment was performed using a TA Instrument (model SDT Q 600) under a nitrogen atmosphere at a heating rate of 10 °C/min. Mechanical Property. The dynamic mechanical properties of the samples were measured using a dynamic mechanical analyzer (DMA) (TA Instruments, model Q-800). Films of dimension (25 × 5 × 0.25) mm were made from the composite samples by solution casting from DMF solution on a die, and the films were installed on the tension clamp of the calibrated instrument. The samples were heated from −100 to 150 °C at a heating rate of 10 °C/min. The storage modulus (G′), loss modulus (G″), and tan δ values were measured with a static force of 0.02 N at a constant frequency (1 Hz). Tensile tests were carried out on solvent-cast films of uniform thickness using a universal testing machine (Zwick Roell, model Z005) at a strain rate of 1 mm/min at room temperature (30 °C). Each experiment was repeated four times to observe reproducibility. The averages of these results are presented in the manuscript. Conductivity. The conductivity (σ) of the samples was measured with an electrometer (Keithley, model 617) by the two probe technique at 30 °C using the equation σ=

Figure 1. 1H NMR spectra of (a) EBIBT, (b) P(EBIBT-co-HT) (P1), and the compatibilizer (c) P(EBIBT-co-HT)-g-PDMAEMA (P2).

that overlaps with that of ‘p’ protons. So the integrated area of ‘c’ is taken as 3/2 of ‘g’, and taking the ratio of peak area (q/c), the average grafted chain length is found to be 12. This yields the molecular weight (M̅ n) of P2 ≈ 102 000. In Figure 1 of the Supporting Information, the GPC result indicates the molecular weight of P2, 108 000, supporting the NMR data analysis. Dispersion of MWNTs with the Compatibilizer P2 and Spectral Property. In the inset of Figure 2, the photographs of P2 solution in DMF (0.1% w/v) and that of P2 solution containing (0.02% w/v) MWNTs are presented. It is apparent from the Figure that MWNT/P2 mixture produces a homogeneous dispersion that is stable for a long time (>3 months). From the UV−vis spectra (Figure 2a), it is apparent that the P2 solution has a peak at 430 mm due to π−π* transition of conjugated thiophene units, but on addition of MWNTs, the peak intensity decreases abruptly with a small blue shift. The π−π interaction between MWNTs and the PT part of the compatibilizer is probably the cause for such decrease in intensity and blue shift.47 The baseline of the MWNT/P2 spectrum shows a gradual decrease with wavelength and is suggestive of Rayleigh scattering by MWNTs in the solution, and it decreases with increase in wavelength.44 The photoluminescence (PL) spectrum of P2 solution (0.1% w/v) (Figure 2b) shows a green emission at 557 nm on irradiation at 430 nm; also, the MWNT/P2 solution shows a green emission at 554 nm with a much lower PL intensity. This is because MWNTs act as a superquencher and arise from an efficient energy transfer of excitons through π−π interaction.48 In Figure 2 of the Supporting Information, we have compared the NMR spectra of P2 and MWNT/P2 solution in CDCl3. The NMR spectra give evidence of CH−π interaction as there is a shift of ‘l’ proton of P2 from 4.055 to 4.060 ppm, ‘p’ protons from 2.561 to 2.569 ppm, and ‘q’ from 2.277 ppm to a doublet at 2.295 and 2.281 ppm. These downfield shifts of different −CH protons suggest the existence

1 d × R a

where d is the thickness, a is the area measured by a screw gage, and R is the resistance of the sample measured from the electrometer. The dc conductivity of the films was measured by sandwiching the samples between the two gold electrodes.

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RESULTS AND DISCUSSION Characterization of P(EBIBT-co-HT) (P1) and [P(EBIBTco-HT)-g-PDMAEMA] (P2). The P1 copolymer is prepared as presented in Scheme 1 using oxidative coupling technique with FeCl3 of EBIBT and HT in chloroform medium at 0 °C, and the compatibilizer P2 is prepared by atom transfer radical polymerization (ATRP) of P1 with DMAEMA using CuBr/ CuBr2 catalyst and HMTETA ligand in anisole medium at 80 °C. The CuBr2 is used to reduce the irreversible termination reaction of active sites during polymerization process.46 The structure of macroinitiator, P1, and P2 are characterized by 1H NMR spectra presented in Figure 1. The positions of the marked protons of EBIBT are shown by the corresponding peaks of the NMR spectrum. In the NMR spectrum of P1 copolymer, new peaks of HT (‘f’, ‘g’, ‘h’, ‘i’) apart from that of EBIBT appear as evident from the spectrum. The peak ‘a’ of EBIBT is absent due to polymerization, and peaks ‘b’ and ‘f’ become broader, indicating copolymerization of EBIBT and HT. The molecular weight (M̅ n) of P1 copolymer measured by gel permeation chromatography (GPC) is 17 250 and polydispersity index (PDI) = 3.2. From the ratio of integral area of ‘c’ and ‘g’ protons, the composition of the copolymer in terms of n/m ratio is found to be 3/2. In the NMR spectra of P2-grafted polymer, the corresponding peaks ‘b’, ‘d’, ‘f’, and ‘g’ of P(EBIBT-co-HT) are not practically observed because of the increased intensity of grafted PDMEMA protons for its large concentration, and peaks ‘c’ and ‘g’ together yield a broad peak 9363



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Figure 3. FE-SEM and TEM images of pristine MWNT (a,c) and corresponding images of the compatibilizer (P2)-wrapped MWNT (b,d), respectively.

1450 cm−1 peak at P2 corresponds to the stretching vibration of thiophene ring, and it shifts to 1446 cm−1 in MWNT/P2 composites, suggesting the presence of π−π interaction. The 2932 cm−1 peak of P2 arises for C−H stretching vibration of P2, and it shifts to lower energy at 2930 cm−1, suggesting the existence of CH−π interaction.27b In Figure 4b, the Raman spectra of MWNT and MWNT/P2 hybrid are shown. The ratio of disordered mode ‘D’ band (1320 cm−1) and the tangential mode ‘G’ band (1572 cm−1) of pristine MWNT is 0.838, and that of MWNT/P2 hybrid structure is 0.842, indicating that the structure of MWNT remains totally unaffected due to the noncovalent functionalization of P2. The shift of the tangential ‘G’ band to higher frequency in MWNT/P2 hybrid (1576 cm−1) may be attributed to the strong π−π interaction between MWNT and PT moiety of P2. The ‘G’ band of MWNTs is accompanied by a shoulder at 1597 cm−1 (D′ band), which can be attributed to the disordered induced lattice distortion5b and is unaffected due to binding with the compatibilizer. In Figure 4c, the TGA thermograms of MWNT, P2, and MWNT/P2 hybrid are presented; the latter two systems exhibit a two-stage degradation temperature. The first-stage degradation occurs at 286 and 292 °C, and the second-stage degradation temperatures are 401 and 414 °C for P2 and MWNT/P2 hybrid, respectively. Possibly the side chain of P2 (side chain in the grafted PDMAEMA) degrades at first,49 and the main chain part degrades at the higher temperature. Therefore, both of these degradation temperatures increase due to wrapping of P2 on MWNT. The MWNT acts as a barrier to the heat flow to the polymer matrix, increasing the degradation temperature. This further supports the presence of strong interaction of P2 with MWNT. PVDF and MWNT/P2 Composites (PCM). Morphology. In Figure 5a,b and Figure 3 of the Supporting Information, the SEM micrographs of PCM composites are shown and MWNTs (white spots) are found to remain well-dispersed within the PVDF matrix. It is to be noted that PVDF fibrils are splaying from the center of a spherulites in PCM0.05, but this

Figure 2. (a) UV−vis spectra of P2 and MWNT/P2 solution in DMF (inset: visual images of left P2 solution and right MWNT/P2 dispersion in DMF) and (b) emission spectra of P2 and MWNT/P2 solution in DMF on irradiation at 430 nm (inset: visual images on irradiation with 430 nm light of left P2 and right MWNT/P2 dispersion in DMF).

of a strong CH−π interaction of P2 with MWNT.27b Therefore, it is the dimethyl aminoethyl protons that are interacting here with MWNT by CH−π interactions. This noncovalent functionalization of MWNT would help in the increase of mechanical and electrical properties of the composites in PVDF In Figure 3a,b, the SEM morphology of pristine MWNT and MWNT/P2 systems are presented, and the micrographs indicate that the pristine MWNTs are entangled with each other, but in the later system, the MWNTs are well-dispersed and are fatter than the former. To get a direct evidence of coating with P2, we present the TEM micrographs in Figure 3c,d, and it is clearly evident that MWNTs are better dispersed in the presence of P2. The average thickness of pristine MWNT is (24.5 ± 2.7) nm, and that of P2-coated MWNT is (38.4 ± 3.7) nm. From these data, it is evident that the average thickness of the P2 polymer on the MWNT surface is 4.9 nm, which indicates a substantial coating of the P2 polymer on the MWNT surface. This coating is occurring due to the π−π interaction between PT moiety of P2 and graphite rings of MWNT surface as evidenced from the UV−vis and PL spectra. In Figure 4a, the FTIR spectra of the P2, MWNTs and MWNT/P2 systems are compared. The >CO vibration peak of P2 at 1732 cm−1 peak is shifted to 1728 cm−1 in the MWNT/P2 system. A probable reason for this shift to lower energy might arise from the delocalization of carbonyl electrons to the π-clouds of MWNT, causing a decrease in carbonyl bond order. The presence of noncovalent interaction between P2 and MWNT is also further evidenced from the FTIR spectra. The 9364



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Figure 5. Representative FE-SEM and HR-TEM images of PVDF- and P2-functionalized MWNT nanocomposites. (a,c) PCM0.05 and (b,d) PCM0.5 samples, respectively (inset in micrographs c and d indicates a good contrast between MWNTs and polymer showing concentration of polymer chains surrounding MWNTs).

Figure 4. (a) FTIR spectra of P2, MWNTs, and MWNT/P2. (b) Raman spectra of pristine MWNTs, and MWNT/P2, respectively. (c) TGA weight loss curves for pristine MWNTs, P2, and MWNT/P2, respectively.

not true in the case of PCM0.5. The SEM images show that the MWNTs are homogeneously and singly dispersed in the PVDF matrix. In the corresponding TEM micrographs, Figure 5c,d and Figure 4 of the Supporting Information, the MWNTs are found to be well-dispersed within the PVDF matrix, and at the inset of each Figure the polymer layers surrounding the MWNTs are clearly shown. It is apparent from the Figure that the thickness of this absorbed polymer layer is larger than that in Figure 3d, and it is due to the specific interaction between the compatibilizer P2 and the PVDF. It is to be noted here that with the increase in MWNT concentration the MWNTs are yet dispersed as single, suggesting that the compatibilizer P2 is very effective in the homogeneous dispersion of MWNTs. In Figure 6a−c, the optical micrographs of PVDF, PCM0.05, and PCM0.2 are compared. The PVDF shows well-defined spherulitic morphology, but in the optical micrograph of PCM0.05 (Figure 6b) spherulitic boundaries are observed; however, we do not observe any birefringence pattern. This is

Figure 6. Polarized optical micrographs of (a) PVDF, (b) PCM0.05, and (c) PCM0.2 samples produced by isothermal crystallization at 135 °C temperature. (d) Digital photographic image of PCM0.2 composite film.

probably due to the formation of a major fraction of β phase (see later) that has no birefringence property,42 but in the micrograph of PCM0.2 neither the boundary nor the spherulitic morphology is observed, suggesting the complete transformation of β phase (cf. WAXS results). The white PVDF film turns brown in the composites (Figure 6d) due to the optical absorbance of P2. PVDF is known to form a lamellar structure, which supramolecularly organizes into spherulitic structure, but in the PCM composites, the lamellar structure may be hampered to some extent, causing a detoriation of spherulitic morphology. Hence, the nanocomposites formation 9365



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The CF−CH−CF bending vibration of PVDF at 875 cm−1 shows a shift to higher energy (879, 882, 883, and 885 cm−1) in the composites with increasing MWNT concentration due to the increased specific interaction.50 With increase in MWNT concentration, the P2 becomes more uncoiled, yielding more sites (>CO group) available for interaction with the >CF2 groups of PVDF. So, from the FTIR spectra, it may be concluded that P2 acts as a good compatibilizer, the thiophene ring interacts with the aromatic ring of MWNTs, and the >C O groups of P2 react in the >CF2 groups of PVDF.51 The formation of β polymorph of PVDF in the composites is also evident from the WAXS spectra (Figure 7b). PVDF has peaks at 2θ = 17.8, 18.5, 20.0, and 26.7° characterizing the αphase formation,52,53 but in the composite PCM0.2 to PCM1, only the diffraction peaks at 2θ = 20.1 characterizing the fully βphase formation is observed. In PCM0.05, a mixture of α- and β-phase formation is clearly evident from the WAXS patterns. This result suggests that in the immediate vicinity of the MWNTs surface the nucleation of the β-phase crystal occurs and it grows to form β-phase crystals, but at very low MWNT concentration (0.05 wt %) the nucleation cannot be all β phase for the lack of sufficient MWNT surface, and some α-phase crystals nucleate at the bulk, causing a mixture of α and β phases. No doubt, it is a better method to obtain β phase than that previously reported where higher concentration of esterfunctionalized MWNTs (2 wt %) is needed to obtain a fully β phase.44 Thermal Properties. From the DSC thermograms (Figure 8), it is evident that the melting point (Tm) of PVDF (Table-1)

with the MWNT/P2 may be given in a schematic model (Scheme 2), where the MWNT/P2 is present within the amorphous region of PVDF crystal. Scheme 2. Schematic Illustration for the Dispersion of Noncovalently Wrapped MWNTs with P2 in the PVDF Matrix

Structure. From the FTIR spectra (Figure 7a), it is apparent that pure PVDF has transmittance peaks at 975, 794, 764, and

Figure 8. DSC melting thermograms of neat PVDF, P2, and PCM nanocomposites at indicated compositions of MWNTs.

gradually deceases with increase in MWNT concentration, and a 7 °C decrease in melting point is observed in the PCM1 Figure 7. (a) FTIR and (b) WAXS spectra of solvent-cast PCM composite films at indicated MWNTs concentration.

Table 1. Summary of Glass-Transition (Tg) and Melting (Tm) Temperature and Crystallinity of PVDF and PVDF/ P2/MWNT Nanocomposites Measured by DSC

530 cm−1, characterizing the formation of α phase, and the composites have transmittance peaks at 510, 840, 1234, and 1275 cm−1. Although the former three peaks are common to both β and γ phases, the presence of 1275 cm−1 peak characterizes the PVDF to crystallize in β phase in the composites.41 The P2 has >CO stretching peak at 1728 cm−1, and in the composites with PVDF this peak shows a gradual shift to lower energy from 1725 cm−1 in PCM0.05 to 1721 cm−1 in PCM0.5. This is due to the dipolar interaction between >CO groups of P2 and the >CF2 groups of PVDF.

samples P2 PVDF PCM0.05 PCM0.2 PCM0.5 PCM1 9366

melting temperature (Tm)

glass-transition temperature (Tg)

heat of fusion (J/g)

crystallinity of PVDF (%)

48.6 32.2 30.6 29.4 28.9

46.4 30.7 29.2 28.1 27.6

49.2 169.5 166.2 164.4 163.8 162.5

46.4 44.7 44.5 44.2



Article

from the Table it is apparent that the increase in G′ is significant in the viscoelastic region for all composites. The PCM0.5 shows the highest increase (122.5%) over that of pure PVDF at 50 °C, and it is attributed to the high aspect ratio of MWNTs. It is to be noted that the percentage increase at these 0.5 wt % MWNT concentrations is highest among the other functionalized MWNTs (66−77%) at this composition.5b,44 The Tan δ versus temperature plot of PVDF shows three peaks at −41, 8, and 102 °C corresponding to the glass transition (T g ) and transition temperatures at crystal amorphous interface (Tr(I)) and within the crystalline zone (Tr(II)), respectively.55,56 In the PCM composites, the Tr(I) is not at all observed, but the other two transitions with the appearance of a new peak at ∼40 °C are observed. The new peak at ∼40 °C is due to the Tg of PDMAEMA chains of P2 polymer.54 Here we observe a small decrease in Tg of PDMAEMA with increasing MWNT concentration, and it is also supported from the DSC data (Table 1). The small mismatch between the Tg data of P2 polymer may be attributed to the different techniques used in the work. The DSC technique uses the heat-capacity variation, whereas tan δ uses the damping characteristics of the material. The heat capacity change monitors the segmental motion characterizing the Tg, but damping characterizes the resistivity to vibration. However, in both of the methods, the Tg decreases, although small, suggesting that the P2 segments get more relaxed with increasing MWNT concentration as the intramolecular interaction of P2 decreases due to wrapping with MWNT. However, it should be noted that both Tg and Tr(II) of PVDF increase to some extent with increasing MWNT concentration, suggesting that the surface forces of compatibilizer-wrapped MWNTs hinder the segmental motion of PVDF at both amorphous and crystalline zones. However, we did not observe any characteristic peak of Tr(I) in the PCM composites, probably due to more diffuse interphase in the composites. Mechanical Properties. In Figure 10a, the stress−strain plots of the PCM composites are shown. The Young’s modulus, tensile strength, elongation at break, and toughness are computed from the stress−strain plots for four different measurements, and they are plotted in the bar diagrams (Figure 10b−e). From the plots, it is clear that Young’s modulus of the composites gradually increases with increase in MWNT content but the elongation at break gradually decreases with increase in MWNTs content. The cohesive force between the filler particles causes the composite to be stiffer, and the decrease in elongation at break is due to slipping of MWNTs on applied stress with increase in MWNT concentration. Both the tensile strength and toughness show a maximum at ∼0.3 wt % MWNTs content. The reason is not yet known, and probably at 0.3 wt % it reaches the mechanical percolation threshold, and with further increasing MWNT concentration some agglomeration of MWNTs may occur, decreasing the tensile strength and toughness at higher compositions. Figure 10a also compares the stress−strain plot of PM (i.e., composite without the compatibilizer P2) with the PCM composites, and it is much lower than any of the PCM composites due to lack of good dispersion of MWNTs causing a lesser reinforcement. In Table 1 of the Supporting Information, the mechanical properties with their percent increase from that of PVDF are presented. PCM1 shows a maximum increase of Young’s modulus by 64.5%, and PCM0.3 shows a maximum increase of toughness of 178.7% and tensile strength of 354. 8%. In the

sample. The enthalpy of fusion values (Table 1) is also decreasing from 48.6 to 28.9 J/g. The decrease in melting point in the PCM0.05 is very large (3.5 °C) because of the interaction of PVDF with the PDMAEMA part of P2, but with increase in MWNT concentration (P2 concentration remains constant), there is also a 2.5 °C decrease in melting point, suggesting that the P2 chains becomes more uncoiled (due to π−π interaction), yielding more sites to interact with PVDF. The same is also true for the decrease in enthalpy of fusion values, supporting the fact that the compatibilizer become more uncoiled upon addition of MWNTs. The P2 polymer exhibits a glass transition (Tg) at 49.2 °C due to the amorphous PDMAEMA chain,54 and it gradually decreases with increase in MWNT concentration in the composites showing a maximum of 4.7 °C decrease. This decrease is probably due to the interaction of PDMAEMA with the amorphous portion of PVDF crystal. Therefore, the compatibilizer P2 acts as a good interfacial adhesion agent between MWNTs and PVDF matrix, and it plays an important role for dispersing MWNTs in the PVDF matrix. Dynamic Mechanical Properties. In Figure 9a,b, the storage modulus (G′) and tan δ plots of the composites are presented. The G′ of PVDF gradually decreases with increase in temperature, although the decrease is not linear and it also occurs for the composites. It is to be noted that with increase in MWNTs concentration the G′ values gradually increase, showing a maximum rise for PCM0.5 sample. The G′ values of the different compositions are presented in Table 2, and

Figure 9. Dynamic mechanical property versus temperature plots of PVDF and the PCM nanocomposite films. (a) Storage modulus and (b) Tan delta. 9367



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Table 2. Summary of Glass-Transition (Tg), Transition Temperatures (Tr), and Storage Modulus (G′) Values of PVDF and Their Nanocomposites with MWNTs at Different Temperatures Measured by DMA transition temperature (Tr) sample

glass-transition temperature (Tg)

Tr(I)

Tr(II)

G′ (MPa) at −50 °C

% increase

G′ (MPa) at 0 °C

% increase

G′ (MPa) at 50 °C

% increase

PVDF PCM0.05 PCM0.2 PCM0.5

−41.2 −40.5 −41.1 −39.3

8.4 40.6 38.8 37.8

101.8 102.5 103.1 105.4

4965 5591 6491 7384

12.6 30.7 48.7

1994 3080 3558 4257

54.6 78.4 113.5

1338 2085 2549 2980

55.8 90.5 122.5

Conductivity. In Figure 12, the variation of dc conductivity of PCM nanocomposites with MWNT content in PVDF is shown. PVDF is an insulator having conductivity of 4.6 × 10−13 S/cm, but the conductivity increases sharply with increasing MWNT concentration and shows a leveling at 0.5% (w/w) at a conductivity value of 10−2 S/cm. The PCM1 shows the highest conductivity of 2 × 10−2 S/cm. The electrical percolation threshold has been calculated from the conductivity versus composition plot, and it is taken as the composition having dc conductivity value at 10−7 S/cm. Therefore, the percolation threshold has a value equal to 0.06 weight fraction of MWNT. The insulator−conductor transition is usually explained from the percolation theory:5b,60−62

Table, the mechanical properties of PM0.05 are compared with that of PCM0.05, and the results clearly indicate a significantly lower increase in Young’s modulus (4.7%), tensile strength (35%), and toughness (5%) than those in PCM0.05, where the increases from PVDF are 10.6, 120.3, and 16.4% respectively. Also, comparing the increase in mechanical property data of PMMA-functionalized MWNT/PVDF composites (F0.5) containing ∼0.07% neat MWNT (increase in Young’s modulus 4.7%, tensile strength 37.6%, and toughness 7.7%)5b with the present results for 0.05 wt % MWNTs (increase in Young’s modulus 10.6%, tensile strength 120.3%, and toughness 16.4%), it is certain that the compatibilizer P2 is very much efficient in the enhancement of mechanical properties, as the same MWNTs were used in both of the works. The Young’s modulus of the nanocomposites has been analyzed by the Halpin−Tsai equation to understand the orientation of MWNT in the composites. The modified Halpin−Tsai equation57,58 relating the Young’s modulus (Ec) of the composite to that of polymer (Ep) and MWNT (ENT) is given for the random orientation:

σ = σ0(p − pc )t

(1)

where (E NT /E P) − 1 (E NT /E P) + 2(lNT/D NT)

ηT =

(E NT /E P) − 1 (E NT /E P) + 2

and for the parallel orientation of the nanotubes to the surface of the sample films ⎡ 1 + 2(lNT/D NT)η VNT ⎤ L ⎥ Ec″ = ⎢ ⎢⎣ ⎥⎦ 1 − ηLVNT

(3)

where σ is conductivity, σ0 is a constant, p is the weight fraction of f-MWNT, pc is the percolation threshold, and t is the critical exponent. Hence a plot of log σ versus log (p − pc) would yield a straight line, as shown in the inset of the Figure. The slope of the straight line is t = 2.26 ± 0.16 and intercept (σ0) = 2.1 × 10−5 S/cm. Hence the critical exponent (t) value of the PCM composites is close to the universal value of 3-D percolation model (t = 1.94).60 The conductivity of pristine MWNT is calculated by taking the above parameters with P = 100, and it has the value 7.3 × 10−1 S/cm. This value is about three times higher than that of PMMA-functionalized MWNT5b and three orders lower than that of pure MWNT(1.83× 103 S/cm).63 The MWNTs are possibly well-wrapped by the P2 polymer through π−π and CH−π interaction, causing a significant hindrance in charge transport and lowering the conductivity by three orders from that of pure MWNTs. The three times higher conductivity in the present system over that of PMMAfunctionalized MWNT/PVDF composite is due to more intact MWNTs that became distorted due to covalent functionalization using nitrene chemistry.5b The percolation threshold values of HO-MWNT/PVDF is reported to be 0.25%, and unfunctionalized MWNT/PVDF is 0.49%.64 These results suggest the supremacy of noncovalent functionalization of MWNTs over that of covalent functionalization in the PVDF composites in enhancing both the mechanical and conductivity properties.

⎡ 3 1 + 2(lNT/D NT)η VNT 5 1 + 2ηTVNT ⎤ L ⎥E p + Ec = ⎢ ⎢⎣ 8 1 − ηLVNT 8 1 − ηTVNT ⎥⎦

ηL =

when P > Pc

(2)

where Ec″ represents the Young’s modulus of the nanocomposites with MWNTs aligned parallel to the surface of the sample film lNT, DNT refers to the length and diameter of the MWNTs, and VNT is the volume fraction of MWNTs in the nanocomposites. Taking Ep of PVDF = 1.25 GPa, ENT =1280 GPa, lNT =12 μm, DNT = 10 nm, density of MWNT = 2.1 g/mL (Aldrich), and density of PVDF = 1.92 g/mL,9,43,59 the Ec and Ec″ values are calculated and are plotted with different volume fraction of MWNTs in Figure 11. From the Figure, it is apparent that the measured data mostly correspond to the random distribution of MWNTs in the PCM composites, and the deviation above 0.3 wt % might be related to some agglomeration of MWNTs in these higher concentrations.

■

CONCLUSIONS A compatibilizer P2 containing thiophene moiety and a reactive PDMAEMA group is prepared by ATRP and is characterized by NMR, FTIR, UV−vis, and PL techniques. The dispersion of MWNTs with P2 in DMF is stable for long time (>3 months), and the UV−vis spectrum of the MWNT/P2 dispersion shows a blue shift with a significant decrease in intensity from that of P2. The PL spectrum also indicates a quenching, suggesting a 9368



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Figure 10. Mechanical properties of PVDF and PCM nanocomposites at different MWNT concentration: (a) representative stress−strain curves at 30 °C and the bar diagrams of (b) tensile modulus, (c) tensile strength, (d) elongation at break, and (e) toughness, respectively.

significant interaction between MWNTs and P2 by π−π interaction. SEM and TEM micrographs suggest a good dispersion and coating of P2 on MWNT. The intensity ratio of D and G bands remains unchanged on this noncovalent doping process, and small lowering of >CO vibration peak position suggests some interaction between P2 and MWNTs. The presence of CH−π and π−π interaction between P2 and

MWNT is also evidenced from FTIR and NMR spectral results facilitating the wrapping of P2 on MWNT surface. The SEM and TEM micrographs of PCM composites show that the MWNTs are homogeneously dispersed, and the optical micrographs indicate that spherulitic morphology of PVDF is lost on increasing MWNTs concentration. The β-phase PVDF is produced in the composites, and its total for MWNT 9369



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Article

ASSOCIATED CONTENT

S Supporting Information *

GPC traces, SEM, TEM, NMR, mechanical property data, and so on. This material is available free of charge via the Internet at http://pubs.acs.org.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We gratefully acknowledge the DST New Delhi (grant no. SR/ S1/PC/26/2009) and ‘Unit of Nanoscience at IACS’ for financial support. A.M. acknowledges CSIR, New Delhi for the fellowship.

Figure 11. Young’s modulus versus loading of MWNT (vol %) of the PCM nanocomposites films. The theoretical data are derived from the Halpin−Tsai model using random and unidirectional orientation of MWNTs.

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REFERENCES

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Figure 12. dc conductivity of PVDF nanocomposites as a function of the weight percentage of P2-functionalized MWNTs at 30 °C. The inset shows the linear dependence of log (σdc) with log (p − pc) according to the percolation scaling law where the red line corresponds to the best fitted line.

concentration is ≥0.5% (w/w). The FTIR spectra suggest specific interaction between the >CO group of P2 and the >CF2 group of PVDF. The decrease in melting point of PVDF with increasing MWNT concentration indicates significant interaction between P2 with PVDF and uncoiling of P2 on MWNTs surface. The G′ value shows the highest increase (122%) for the addition 0.5% MWNTs reported in the literature. The increase in Young’s modulus, tensile strength, and toughness in PCM0.05 is two to three times higher than that of the composite without compatibilizer. In a comparison of the mechanical property data with that reported for covalent functionalized samples, it is found that the compatibilizer P2 is very much efficient in the enhancement of mechanical properties. The Young’s modulus data are analyzed using the Halpin and Tsai equation, and it suggests a random distribution of MWNT in the composite, particularly at lower concentration (≤0.3 wt %). The PCM1 shows the highest conductivity (2 × 10−2 S/cm) and shows a very low percolation threshold of MWNTs reported in literature for PVDF system. Analysis with percolation theory suggests that the conductivity obeys the 3-D percolation model and the pure MWNT/P2 has three times higher conductivity over that of PMMA-functionalized MWNTs. Therefore, noncovalent functionalization of MWNTs is found to be superior to that of covalent functionalization in enhancing both the mechanical and conductivity properties in the PVDF composites. 9370



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Noncovalent Functionalization of Multiwalled Carbon Nanotube by a

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