NANO LETTERS
Properties of Polyvinylidene Difluoride−Carbon Nanotube Blends
2004 Vol. 4, No. 7 1267-1271
Nicole Levi, Richard Czerw, Shuya Xing, Preethi Iyer, and David L. Carroll* Center for Nanotechnology, Department of Physics, Wake Forest UniVersity, Winston-Salem, North Carolina 27109 Received April 19, 2004; Revised Manuscript Received May 22, 2004
ABSTRACT Single-walled and multiwalled carbon nanotube blends with polyvinylidene difluoride (and its copolymers) have been characterized. The nanotubes are observed to form a well-dispersed, structurally random nanophase within the fluoropolymer matrix. X-ray analysis coupled with differential scanning calorimetry suggests that the nanophase alters crystal formation within the polymer. For most loadings and nanotube types, the piezoelectric β-polymorph is significantly enhanced over other crystal phases. Solution-cast composite thin films exhibit enhancements in both the pyroelectric response and mechanical transduction over pure polymer. This is interpreted as resulting from the change in crystallinity.
The physical properties of piezoelectric and pyroelectric electroactive polymers (EAPs) are sufficiently different from those of conventional ceramic materials that a host of novel applications can be envisaged for these materials. These applications stem from the EAPs availability in large area, thin, flexible films, as well as the possibility of selected area poling of the film through the use of appropriate electrode patterns. Small values of relative permittivity, Ghz response frequencies, wide dynamic range for actuation and transduction, higher dielectric strength than piezoceramic material (30 V/µm versus 1.5 V/µm), excellent elastic compliance, and an acoustic impedance close to that of water (for hydrophonics) are some of the useful parameters of EAPs that promise even further applications. Of all the polymeric materials investigated, polyvinylidene fluoride (PVDF) and its copolymers exhibit the largest piezoelectric and pyroelectric coefficients and have been intensively studied by several research groups.1-4 In fact, such materials have already been commercialized in transducer, pyroelectric, and some actuator applications.5,6 However, as with all EAPs currently, their application potential is still severely limited by their low electromechanical coupling and relatively small force generation capabilities. An interesting approach recently introduced for enhancing the properties of EAPs necessary for transducer and actuator applications is the creation of matrix composites. Typically, these are inorganic-organic blends utilizing ferroelectric ceramic particles as a dispersed phase into a ferroelectric polymer matrix.7-10 They are flexible and light, combining the processibility of polymeric films with the high piezoelectric and pyroelectric responses of ceramics.1,11 Though this approach is promising, it suffers from several important * Corresponding author. E-mail:
[email protected]. 10.1021/nl0494203 CCC: $27.50 Published on Web 06/16/2004
© 2004 American Chemical Society
drawbacks. First, dispersed particles larger than a few hundred nanometers can induce “stress concentrators” in the matrix, leading to fatigue and early failure under cycling. Second, the host materials and ceramic phases tend to pole under quite different conditions, leading to inhomogeneities of the composite overall. Finally, there is a question of miscibility: dispersions of large particles are difficult to process. Alternatively, single-walled carbon nanotube (SWNT) mats have demonstrated amazing actuation properties.12 These organic actuators can generate surprisingly large forces and extension due to a coupling of high tensile strength of the SWNTs and the very high aspect ratios. Further, the heatresistant quality of the carbon means these devices may be used at temperatures up to 1000 °C, far exceeding the capabilities of existing polymer transducers. In fact, macroscopic sheets of carbon nanotubes, working under physiologic conditions and low voltage, show behavior comparable or superior to that of natural muscle.13 Of course, applications over large areas using such actuator technology would be extremely expensive and difficult (they have only limited elastic compliance). However, these developments suggest that SWNTs might be ideal reinforcements as a nanophase dispersed within an EAP matrix. Ideally this would allow a coupling of their tensile strength and actuation properties with that of the polymer at length scales that are very different from those referenced above. In this work, we describe the properties of a set of PVDFCNT matrix composites that exhibit remarkably enhanced pyroelectric and piezoelectric behavior over the pristine polymer. Further, we show that the modifications of the properties can be partially explained by alterations in the polymer microstructure due to the presence of the nanotubes.
These results clearly indicate that such organic-organic matrix composites at the nanoscale hold the promise of creating superior EAP-based actuator and pyroelectric systems. Several sources of carbon nanotubes were used in this work, including purified HiPCO single-walled carbon nanotubes (HiPCo SWNTs), made by Carbon Nanotechnologies Incorporated, as well as our own arc grown and purified SWNT (ARC SWNTs) and multiwalled nanotubes (MWNTs). In the case of SWNTs, cleaning procedures outlined in the literature were employed to provide the cleanest nanotube product possible.14 Generally, SWNT materials have a catalyst content (Fe or Ni/Y) below 1 wt % as determined by X-ray analysis and further characterized by thermal gravimetric analysis (TGA). High resolution electron microscopy indicates that little amorphous carbon survives the cleaning procedures as well. It should be noted that in such a cleaned state, the tubes are relatively difficult to disperse in most solvents. Commercially obtained host polymers used in these studies were the homopolymer poly(vinylidene fluoride) (PVDF) and two PVDF random copolymers.5 The PVDF was obtained in pellet form and a solution was made using HPLC grade N,N-dimethylacetamide (HPLC-DMAc). The random copolymer, poly(vinylidene fluoride-tetrafluoroethylene) (PVDFTFE), was obtained in solution form in HPLC-DMAc, with ∼80% VDF. Finally, the copolymer poly(vinylidene fluoridetrifluoroethylene) (PVDF-TrFE) was also obtained in solution form with VDF content at 65 wt %, again in HPLC-DMAc. We note that in the case of P(VDFx-TFE1-x) and P(VDFxTrFE1-x) copolymers, the co-monomer content, x, is offered from the manufacturer such that maximal piezo response is achieved. Using HPLC-DMAc, the concentration of each polymer solution to be used in film casting was controlled (between 10% and 20 wt %, depending on the thickness of the film needed). Blending of the nanotubes into the polymers was carried out by first dispersing polymer and nanotubes in solvent separately. Nanotube-DMAc solutions (sonicated for 10 min with a high energy ultrasonic probe) were then ultrasonically blended into the polymer-solvent solutions using a low power, water bath sonicator for 60 min at room temperature. Thin films were then cast and drawn using an adjustable “micron film applicator” draw blade on glass substrates. The thickness of the films was controlled by the concentration of the polymer solution (as mentioned above) and the height of the blade of the film applicator. After drawing, the films were placed with the glass substrate, on a hotplate for about 20 min to evaporate the solvent. The surface temperature of the hot plate was in the range of 75-80 °C. Next the samples were annealed in an oven at 70 °C for 38 h. The thickness of the films was typically between 20 and 40 µm as measured with Micromaster Digital Micrometer by Brown & Sharpe. Large scale aggregation characteristics of the matrix were determined using simple light scattering through the thin films. With a Perkin-Elmer 900 UV-vis spectrometer, the films were mounted with the surfaces normal to the incident 1268
Figure 1. (a) UV-vis absorption of PVDF and PVDF-TrFE hosts as a function of HiPCo SWNT content (by weight). (b) Scattering intensity of 500 nm wavelength light by PVDF and PVDF-TrFE (co-polymer) hosts with increasing HiPCo SWNT content.
light and referenced to air. Typical absorption curves are seen in Figure 1a. In this figure two different PVDF polymers (PVDF and PVDF-TrFE) are plotted for four different loadings of HiPCo SWNTs. The absorption spectra of the blends are dominated by Rayleigh scattering from the nanophase. Clearly, the polymer becomes darker with increasing nanotube concentration. This is reflected by the rise in the overall absorption tail of the spectra. Interestingly, however, this increase is rather small from the pure materials up to materials loaded ∼1 wt. %. We note that the increase of the Rayleigh tail as a function of nanotube loading is quite different for PVDF and PVDF-TrFE. In fact, plotted in Figure 1b, we show the scattering intensity at an arbitrarily chosen wavelength (500 nm) to demonstrate the different trends. The absorption/scattering increases more rapidly for PVDF as nanotubes are added than it does for the PVDF copolymer. This suggests that large aggregates begin forming at lower nanotube concentrations in the PVDF than in the PVDF-TrFE host. In fact, the high fluorine content copolymers (TrFE and TFE) both show significantly suppressed Nano Lett., Vol. 4, No. 7, 2004
Rayleigh scattering for the same loading levels. This observation is further substantiated in electron microscopy (not shown). Finally, electrical percolation thresholds also occur at around 0.1 wt % generally and further support the view that the dispersion in these halogen containing polymers is rather good. PVDF and its copolymers are semicrystalline with four primary polymorphs. The crystallinity and volume fraction of each of the polymorphs depends on film preparation generally and the percentage of co-monomer (in the case of the co-polymers). It has been shown, however, that the four primary polymorphs do occur over a wide range of composition.3 The polymorphs are generally referred to as alpha (R) phase or Form II, beta phase (β) or Form I, gamma phase, and delta phase (or polar alpha).16 The R and β phases are the most common of the crystalline forms of PVDF and its copolymers with the preferred polymorph are usually R-phase. β-phase is the most important, however, since it exhibits piezoelectric and pyroelectric properties.17 There are extensive studies on mechanisms for the creation or enhancement of β-phase in PVDF and its co-polymers.18-20 In these studies however, a simple drawing and poling of the PVDF films is used to create the phase structure. The phase structure of our films was examined using X-ray diffraction techniques. Using the 0.154 nm, Cu KR, line on an XDS 2000 from Scintag, 2θ scans were carried out from 2° to 60°. The thin films were unpoled and mounted flat on a stainless steel stage. Shown in Figure 2 are typical 2θ scans of pure PVDF-TrFE and the PVDF-TrFE/HiPCo-SWNT1 at 0.05 wt %. Assigning the features of 2θ at 20.8° with the (110) and (200) reflections of the β phase3,22 and R phase with features around 18 ° and 35°,23 we see a change in the β-to-R phase balance. Peak areas were deconvoluted and integrated using commercially available software. From analysis of peak areas, significantly more β phase occurs in the blends than in the pure polymer. In fact, this is true for PVDF as well as PVDF-TFE and PVDF-TRFE. While HiPCo tubes are used in Figure 2, it is instructive to examine the loading at which the ratio of β to R alpha phase is maximized for different nanotube types. Specifically, we examine PVDF-TrFE loaded with HiPCo SWNT, ARC SWNT, fluorinated HiPCo SWNT, and MWNTs. Each of these nanotube types results in enhanced β-phase, however, again using peak areas for analysis, for HiPCo SWNTs the maximum β:R ratio occurs at a weight loading of 0.01 wt %, for ARC SWNTs the maximum occurs at a loading of 0.05 wt %. Fluorinated HiPCo SWNTs (available from CNI) have a maximum just above 0.001 wt % and MWNTs show a maximum β-to-R alpha phase ration above 0.8 wt %. This unusual result suggests that the nucleation of β-phase over R-phase is related to the overall morphology of the nanoinclusion. This is expected since the nanotubes have different diameter distributions, different bundling characteristics, and perhaps different lengths and persistence lengths within the matrix. Thus, they all represent slightly different surface energies. We further note that in all cases, the total crystallinity due to R and β phases seems to increase slightly. Nano Lett., Vol. 4, No. 7, 2004
Figure 2. (a) X-ray diffraction spectra of PVDF-TrFE without HiPCo SWNTs. (b) X-ray diffraction spectra for PVDF-TrFE/ HiPCo SWNT at 0.05 wt %. The extra features in the blend (∼30°) are related to the occurrence of other polymorphs. There does seem to be some dependence of these crystal types on nanotube concentration. However, this is beyond the scope of the present study.
Likewise, these results are reflected in the differential scanning calorimetry (DSC) of the films as shown in Figure 3. The DSC was performed on a TA Instruments MDSC 2920 calibrated with indium. Aluminum sample pans were used in a nitrogen atmosphere. All measurements were done at a heating rate of 10 °C/s from -50 to 300 °C. Sample film masses used were between 8.73 and 9.66 mg. For each of the polymer hosts used, a shift toward lower temperatures and narrowing of the melting point features is observed, suggesting that the crystallization of films is becoming more monophase, in agreement with XRD results. The same trend is found in the Curie point. The integrated areas under the curves were examined to show that the blends exhibit larger overall heat flow for the same mass. This suggests an increase in volume of material with β-phase. A specific example is the enthalpy change from the pure polymer 18.84 J/g to 20.63 J/g for the PVDF-TrFE blend at 0.0108 wt % of HiPCo with a shift in melting temperature of 150.3 °C to 149.8 °C. Similar results were seen in the cases of PVDF and PVDFTFE. 1269
Figure 4. Pyroelectric coefficient for PVDF-TFE and PVDF-TrFE with a loading of HiPCo SWNTs of around 0.01 wt %. Figure 3. DSC scans of the 150 °C feature associated with crystal melting in PVDF and its copolymers. Notice the shift to lower temperatures and narrowing of the feature as the nanotube concentration is increased. The results presented are for PVDFTrFE loaded with HIPCO SWNTs for the concentrations 0%, 0.088%, and 0.598 wt %. PVDF and PVDF-TFE show the same trends.
In this work, the pyroelectric coefficient was determined directly using the relation Py ) I/(βA), where I is the current, A is the area of the electrode, and β is the rate of the temperature change. The samples used for pyroelectric determinations began with the thin films as described above. Thin aluminum electrodes (1 µm thick) were then evaporated slowly onto each side of the films in a 10 × 28 mm rectangular area. The samples were first poled by applying a DC field (2100 V) at room temperature. Then the electrodes were shorted for at least 24 h before the pyroelectric currents were measured. Thermal cycling was carried out with the sample mounted in a stainless steel cell, heated in a hot water bath. The heating rate was 7 °C/min and current was measued using a Keithley 485 picoammeter. The first thermal run to measure pyroelectric current is called the “depolarization or irreversible” cycle and is a measure of the absorbed charge during the poling process. All subsequent thermal cycles, i.e., heating the short-circuited samples at a constant rate to a desired temperature and subsequent cooling to room temperature, were observed to produce a “reversible” or “true” pyroelectric current. Figure 4 compares the pyroelectric coefficient dependence on temperature for the pure PVDF-TrFE co-polymer, PVDFTFE/HiPCo-SWNT at 0.006 wt % composite, and PVDFTrFE/HiPCo-SWNTs at 0.0072 wt % composite as determined from the second, or subsequent, thermal cycles. The pyroelectric coefficients increase with the increasing temperature in each case due to the corresponding increases in both the thermal expansion coefficient (a secondary effect) and temperature coefficient of the order parameter for dipoles (a primary effect).16 We note that the shapes of these curves are quite different, suggesting that the origins of the temperature dependence may be modified by the presence of nanotubes. Indeed, we would expect significant changes in any property of the materials that depended strongly on the thermal expansion, as the nanotubes will clearly alter 1270
such phenomena. Comparing the pyroelectric coefficients of the three curves, it is found that at room temperature the pyroelectric coefficients of the composites are higher than that of pure PVDF-TrFE copolymer. However, as a function of temperature, the increase in the pyroelectric coefficients of the composites is more gradual than that of the copolymer. We speculate that at higher temperatures, the SWNTs sterically hinder the dipoles, thereby limiting the overall pyroelectricity. Even though the amount of SWNTs is small, the interface between the polymer and SWNTs is large. Because these are thin films, unambiguous determination of the piezoelectric coefficients is difficult at best; however, rough estimates of the voltage transduction can be made. To measure the piezovoltage for a specific force loading, the electroded films (described above) were first poled at a variety of poling voltages and were then clamped at one end so they could be loaded with a sinusoidially varying force at the other, in a pendulum configuration. The loading was driven at ∼2 Hz. The applied load was in the direction of the 28 mm dimension, parallel to the film. The voltage during load cycling was read from the electrodes across the film thickness using a capture oscilloscope. Thus, the voltage was perpendicular to the loading direction and corresponds to the d31 component of the piezotensor. The maximum voltage corresponded to the maximum load in the cycle. Figure 5 compares the maximum voltage output of the PVDF-TrFE/ HiPCo SWNT at 0.1 wt % loading with the pure PVDFTrFE co-polymer. The two polymers are plotted as a function of poling voltage. The dielectric constant of each film was measured, and the d31 coefficient was then determined using the maximum voltage for each sample and the known capacitance of the film. For poling voltages above the saturation value shown in Figure 4, the d31 values were ∼20 pC/N for the pure film24,25 and ∼25 pC/N for the composite. The enhanced pyroelectric and piezoelectric behavior of the nanocomposite materials over the pure materials is most easily explained by the observed increase in β-phase crystallinity. This would imply that this crystal phase would be somehow associated with the nanoinclusion (the nanotube in this case). Specifically, one would expect that the nanotube has nucleated crystal structure. This can be seen in two other suggestive pieces of evidence outside of the XRD and DSC Nano Lett., Vol. 4, No. 7, 2004
suggest that the dispersed nanophase is responsible for morphological changes in the polymer crystallinity. The “selectivity” of polymorph formation has been further confirmed using XRD and DSC analysis. This route to modifying piezopolymer behavior for enhanced actuation and transduction is likely to provide access to novel applications and is an exciting demonstration of properties engineering at the nanoscale. Acknowledgment. The authors gratefully acknowledge funding through AFOSR grant number F49620-99-1-0173. Further, discussions with Dr. J. Coleman and Prof. W. Blau were instrumental in interpretations. References Figure 5. Piezoelectric response of PVDF-TrFE/HiPCo SWNT at 0.0072 wt % as compared to pure PVDF-TrFE.
Figure 6. A transmission electron micrograph of nanotube pullout. The carbon nanotubes shown are multiwalled. The matrix is PVDFTrFE.
studies. First electron micrographs clearly show that fractured surfaces of the composites leave the nanotubes with a polymer coating. As shown in Figure 6, the coating can be quite thick and covers the length of the nanotube. Thus, during pullout, this polymer sheath remains associated with the nanotube, implying some energetic relationship with its surface. Second, the area of the β-phase reflection peaks grow linearly toward a maximum value with loading as would be expected with volumetric exclusion. Beyond this value, the growth in β-phase peaks rolls off and begins decreasing due to reaggregation within the matrix. This relationship was first pointed out by Coleman26 and suggests that one phase grows at the expense of the others. Thus, it seems safe to suggest that some modification of the phase structure of the polymer is responsible for changes in pyroelectric and piezoelectric behavior. This work has examined the pyroelectric and piezoelectric properties of PVDF/nanotube blends. Significant alterations in both pyro- and piezo- properties have been observed. We
Nano Lett., Vol. 4, No. 7, 2004
(1) Furukawa, T. IEEE Trans. Elect. Insulation 1989, 24, 375. (2) Baise, A. I.; Lee, H.; Oh, B.; Salomon, R. E.; Labes, M. M. Appl. Phys. Lett. 1975, 26, 428. (3) Koga, K.; Ohigashi, H. J. Appl. Phys. 1986, 59, 2142. (4) Kepler, R. G.; Anderson, R. A. J. Appl. Phys. 1978, 49, 1232. (5) Kawai, H. Jpn. J. Appl. Phys. 1969, 8, 975. (6) Sessler, G. M. J. Acoustical Soc. Am. 1981, 70, 1596. (7) Janas, V. F.; Safari, A. J. Am. Ceram. Soc. 1995, 78, 2945. (8) Ploss, B.; Ploss, B.; Shin, F. G.; Chan, H. L. W.; Choy, C. L. IEEE Trans. Dielectrics Electrical Insul. 2000, 7, 517. (9) Venkatragavaraj, E.; Satish, B.; Vinod, P. R.; Vijaya, M. S. J. Phys. D: Appl. Phys. 2001, 34, 487. (10) Cui, C.; Baughman, R. H.; Iqbal, Z.; Kazmar, T. R.; Dahlstrom, D. K. Synth. Met. 1997, 85, 1391. (11) DiasIAS, C. J.; Dasgupta D. K. ASGUPTA, Ferroelectric Polymers and Ceramic- Polymer Composites Key Engineering Materials 1994, 92-9, 217. (12) Baughman, R. H.; Cui, C. X.; Zakhidov, A. A.; Iqbal, Z.; Barisci, J. N.; Spinks, G. M.; Wallace, G. G.; Mazzoldi, A.; De Rossi, D.; Rinzler, A. G.; Jaschinski, O.; Roth, S.; Kertesz, M. Science 1999, 284, 1340. (13) Fraysse, J.; Minett, A.; Jaschinski, O.; Journet, C.; Roth, S. Vide Sci. Tech. Appl. 2001, 56, 229. (14) Strong, K. L.; Anderson, D. P.; Lafdi, K.; Kuhn, J. N. Carbon 2003, 41, 1477. (15) 5 Venders were Aldrich Inc. (St. Louis, MO) and Solvay Inc. (Tavaux, France). Polymer properties were generally: molecular weight average (MW) ∼500 000, number average (MN) ∼100 000, and polydispersive index: MW/MN ∼4 for each. We note that polymerization route and specific molecular weights vary between copolymers and homopolymers used in this study. While detailed information regarding physical properties on the polymers used are available from the manufactures websites, the comparisons of physical properties made in this study will be between pure and blend of the same polymer only. (16) The Applications of Ferroelectric Polymers, Wang, T. T., Herbert, J. M., Glass, A. M., Eds; Kluwer Academic: Amsterdam, 1987. (17) Kepler, R.; Anderson, R. AdV. Phys. 1992, 41, 1. (18) Elling, B.; Danz, R.; Weigel, P. Ferroelectrics 1984, 56, 179. (19) Kaura, T.; Nath, R.; Perlman, M. M. J. Phys. D: Appl. Phys. 1991, 24, 1848. (20) Kulek, J.; Hilczer, B.; Kamba, S.; Petzelt, J. Acta Polym. 1995, 46, 152. (21) We use polymer/nanotube at wt % to denote the polymer typenanotube type, and nanotube loading in the blend. (22) Davis, G. T.; McKinney, J. E.; Broadhurst, M. G.; Roth, S. C. J. Appl. Phys. 1978, 49, 4998. (23) Newman, B. A.; Yoon, C. H.; Pae, K. D.; Scheinbeim, J. I. J. Appl. Phys. 1979, 50, 6095. (24) Vuillermoz, B.; Nolf, M.; Toudic, Y. Ferroelectrics 1981, 32, 157. (25) Kunstler, W.; Wegener, M.; Seib, M.; Gerhard-Multhaupt, R. Appl. Phys. A 2001, 73, 641. (26) Private communication with Dr. Coleman, Trinity College Dublin.
NL0494203
1271
