Structures and Properties of Supramolecular Assembled Fullerenol

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J. Phys. Chem. B 2004, 108, 5937-5943

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Structures and Properties of Supramolecular Assembled Fullerenol/Poly(dimethylsiloxane) Nanocomposites Jianying Ouyang,† Shuiqin Zhou,*,† Feng Wang,† and Suat Hong Goh‡ Department of Chemistry, The College of Staten Island and the Graduate Center, City UniVersity of New York, 2800 Victory BouleVard, Staten Island, New York 10314, and Department of Chemistry, National UniVersity of Singapore, Singapore 117543 ReceiVed: January 29, 2004; In Final Form: March 4, 2004

Supramolecular assembled fullerenol/poly(dimethylsiloxane) nanocomposites were prepared from the solution casting of the mixtures of multihydroxylated [60]fullerene (fullerenol) and bis(3-aminopropyl)-terminated poly(dimethylsiloxane) (PDMS-di-NH2) in tetrahydrofuran at different molar ratios. Dynamic light scattering measurements reveal that the fullerenol-PDMS complexes could be formed due to the strong hydrogen bonding interactions between the hydroxyl groups of fullerenol and the amino groups of PDMS-di-NH2. The size and size distribution of the complex particles depend on the molar ratios of OH/NH2 groups and the solution concentrations. Small-angle X-ray scattering results indicate that nanodomains of fullerenol aggregates are confined homogeneously in PDMS matrix. The higher fullerenol content in the composites leads to the larger size of fullerenol nanodomains. This novel structural feature of the nanocomposites together with the unique chemical structure and dielectric property of fullerenol molecules results in superior thermal and thermal mechanical stability, severely suppressed crystalline phase, elastic mechanical response, and attractive dielectric properties of the fullerenol-PDMS nanocomposites; i.e., the high content of fullerenol in the nanocomposites increases the permittivity but dramatically decreases the loss factor of the materials.

Introduction [60]Fullerene-based polymeric materials have received much attention in recent years in view of the attractive superconducting, optical, catalytic, and medicinal properties of [60]fullerene (C60).1-9 Although various methods have been developed to link C60 onto polymers,10-20 it is simpler and more convenient to embed C60 physically in polymer matrixes. However, the incompatibility between C60 and polymers leads to the formation of large C60 aggregates in the polymer matrixes, compromising the performance of the materials. The use of C60 derivatives reduces the compatibility problem, and in some cases the performance of polymeric materials embedded with C60 derivatives is better than that of materials embedded with pristine C60.21,22 The functionalization of C60 and polymers with suitable pendent or end functional groups can lead to strong interactions between the C60 derivatives and polymers and enable the C60 derivatives to be well dispersed and adhere strongly to the polymer matrixes. The polymers could be either single pendent or cross-linked by C60 species, leading to a significant improvement in the storage moduli of the materials.23,24 A supramolecular network-like structure of C60-containing polymeric materials could be constructed via the covalent reactions between the multifunctional groups appending on the C60 cage and the terminal functional groups of polymers. For example, Chiang25 prepared fullerenol-cross-linked polyurethanes by means of isocyanate-hydroxy condensation reaction. The resultant starburst polymer network with fullerenol as a mo* To whom correspondence should be addressed. E-mail: zhoush@ postbox.csi.cuny.edu. † College of Staten Island, City University of New York. ‡ National University of Singapore.

lecular core exhibits significantly enhanced thermal mechanical properties. Goswami26 reported that the hydroxyl groups of fullerenol selectively undergo nucleophilic addition reaction with carbonyl groups of a cycloaliphatic epoxy resin. The formed starlike polymers show improved thermal stability. It has been pointed out that C60-based materials must possess film-forming ability in many applications.27,28 Poly(dimethylsiloxane) (PDMS) is a good choice as the polymer network chains to host C60 due to its many useful properties, such as flexibility, permeability to gases, low glass transition temperature, very low surface energy, good thermal stability, and biocompatibility.29,30 It is desirable to make novel materials with a combination of the outstanding properties of both C60 and PDMS. Nanodomains of gold, metal oxides, clay, and zeolite have been, respectively, incorporated into polysiloxane matrixes to form various nanocomposite materials;31-35 however, little has been studied dealing with the incorporation of C60 into PDMS. Kraus36 demonstrated that the reactive fullerene bisadduct allows the insert of C60 units into the backbone of PDMS whereas reactive monoadducts may be attached to the side chain of PDMS. The resultant polymers show good film-forming ability. In the present work, a series of supramolecular assembled C60-containing PDMS nanocomposites have been prepared based on the strong hydrogen bonding interactions between the multihydroxylated C60 derivative fullerenol (Fol) and the bis(3-aminopropyl)-terminated poly(dimethylsiloxane) (PDMS-diNH2) in tetrahydrofuran (THF). Free-standing films could be obtained from the solution casting of these Fol/PDMS-di-NH2 complex solutions in THF. Depending on the mixing ratios of Fol to PDMS-di-NH2, the size of the Fol nanodomains homogeneously confined in PDMS matrix are controllable. This novel structural feature is different from traditional highly crosslinked polymer networks or the covalently linked C60-containing

10.1021/jp0495850 CCC: $27.50 © 2004 American Chemical Society Published on Web 04/20/2004

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TABLE 1: Preparation Concentrations and Compositions of Various Samples PDMS-di-NH2 (mg/mL) Fol-(OH)11 (mg/mL) molar ratio of OH/NH2

composite 1

composite 2

composite 3

composite 4

30.0 1.29 0.75:1

29.9 1.73 1:1

30.3 3.47 2:1

29.9 6.87 4:1

copolymers.25,26,36 Benefiting from the unique chemical structure and dielectric property of Fol molecules, the resultant nanocomposites with controllable size of Fol nanodomains embedded in PDMS matrix exhibit superior thermal and thermal mechanical stability, elastic mechanical response, and unique dielectric properties; i.e., the high content of Fol in the nanocomposites increases the permittivity but dramatically decreases the loss factor. Experimental Section Materials and Composites Preparation. [60]Fullerene (C60) (99.9%) was purchased from SES Research Company, USA. Fol in red powders was synthesized and characterized according to the literature.37 The carbon/oxygen ratio of Fol was found to be 60:11 using X-ray photoelectron spectroscopy, which is in good agreement with the literature value of an average 10-12 hydroxyl addents per C60 molecule.37 PDMS-di-NH2 was purchased from Sigma-Aldrich Company, USA. It is a colorless free-flowing liquid, with viscosity of 50 cSt and amine number of 0.6-0.8 mequiv/g stated by the manufacturer, and its numberaverage molecular weight and polydispersity are 3140 and 1.78 as determined by gel permeation chromatography (GPC) with polystyrene standards used for calibration. THF is of AR grade and used as received. To prepare Fol/PDMS-di-NH2 nanocomposites, appropriate amounts of PDMS-di-NH2 and Fol were dissolved in THF and the mixtures were continuously stirred overnight followed by solvent evaporation at room temperature. Four nanocomposites were prepared, where the concentration of PDMS-di-NH2 was kept constant at 30 mg/mL, while the Fol concentration was increased to control the molar ratios of hydroxyl groups of Fol to amino groups of PDMS-di-NH2 to 0.75:1, 1:1, 2:1, and 4:1, respectively (see Table 1). The samples were finally dried in vacuo at 50 °C for 3 days. It was observed that Fol was distributed into the composites, and excess PDMS-di-NH2 was squeezed out of composites 1 and 2 to form small regions of colorless liquid in the corner, which was not apparent in composites 3 and 4. Composite 1 is a reddish gel, and the other composites are red free-standing films as shown below.

Small-Angle X-ray Scattering (SAXS) Measurements. SAXS measurements were performed at the Advanced PolymersPRT Beamline (X27C), National Synchrotron Light Source (NSLS) at Brookhaven National Laboratory (BNL), using a laser-aided prealigned pinhole collimator. The incident beam wavelength (λ) was tuned at 0.1366 nm. A two-dimensional imaging plate was used in conjunction with an image scanner as the detection system.39 The sample-to-detector distance for SAXS was 795 mm. The scattering vector q is expressed as q ) (4π/λ) sin(θ/2), with θ being the scattering angle between the incident and the scattered X-rays. The d spacing of the ordered structures can be calculated as d ) 2π/q. Thermal Analyses. The measurement of glass transition temperatures (Tg’s) was made with a TA Instruments 2920 differential scanning calorimeter (DSC) under a flow of 50 mL min-1 purified nitrogen at a heating rate of 20 °C min-1. The midpoint of the abrupt increase of heat capacity in the DSC curve was taken as Tg. Each sample was scanned several times to erase its thermal history and check the reproducibility of the results. To study their crystallization behavior, samples were annealed in situ at -95 °C for 30 min, followed by heating at 20 °C min-1 under purified nitrogen. Thermal stability was analyzed under a flow of 60 mL min-1 purified nitrogen at a heat rate of 10 °C min-1 using a Hi-Res TGA 2950 thermogravimetric analyzer supplied by TA Instruments. Thermal mechanical analyses (TMA) were performed in compression mode with a flat-tipped standard expansion probe using a TA instruments TMAQ400 under a flow of 50 mL min-1 purified nitrogen. Samples were cooled in situ to -150 °C with liquid nitrogen, and then heated at a rate of 5 °C min-1. Dielectric Analyses. The dielectric analysis was carried out on a TA Instruments Dielectric Analyzer 2970, covering a frequency range from 300 Hz to 100 kHz and a temperature range from -150 to 100 °C. The dielectric spectra were obtained in ceramic parallel plate mode during heating scan at a rate of 3 °C min-1 in purified nitrogen. Rheometry. Viscoelastic properties of various samples were investigated using a shear strain controlled rheometer (ARES RFS rheometer, TA Instruments) with 25 mm parallel plates. The frequency sweep measurements were carried out from -20 to 20 °C at 10 °C intervals at a constant 10% strain with frequency ranging from 0.1 to 100 rad/s. The temperature was controlled by a Peltier assembly with an accuracy of 0.01 °C. Results and Discussion

Dynamic Laser Light Scattering (LLS) Measurements. A standard laser light scattering spectrometer (BI-200SM) equipped with a BI-9000 AT digital time correlator (Brookhaven Instrument Inc.) and a He-Ne laser (35 mW, 633 nm) was used to perform dynamic light scattering studies over a scattering angular range of 30-120°. The temperature was controlled at 25 ( 0.05 °C. The complex solutions were filtered through 0.5-µm Millipore Millex filter to remove dust particles. The Laplace inversion of each measured intensity-intensity time correlated function in a dilute solution can result in a characteristic line width distribution G(Γ). For a purely diffusive relaxation, G(Γ) can be converted to a hydrodynamic radius distribution by using the Stokes-Einstein equation.38

Supramolecular Assembly of Fol and PDMS-di-NH2. The use of hydrogen bonding interactions to construct polymeric structures has received increasing attention in recent years.40-42 In our designs, the PDMS-di-NH2 chain has two terminal propylamino groups and the Fol possesses an average of 11 hydroxyl groups attached on the C60 core. We expect that strong hydrogen bonding interactions between the amino and hydroxyl groups will drive the supramolecular assembly of Fol molecules and PDMS chains. Figure 1 shows the apparent hydrodynamic radius (Rh) distributions of the Fol/PDMS-di-NH2 complex particles formed at 30 mg/mL PDMS-di-NH2 but different molar ratios of the OH/NH2 groups in the Fol/PDMS-di-NH2 mixtures in THF,

Properties of Fullerenol/PDMS Nanocomposites

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Figure 1. Apparent hydrodynamic radius (Rh) distributions of Fol/ PDMS-di-NH2 complex particles in THF formed at 30 mg/mL PDMSdi-NH2 but different ratios of OH/NH2 groups of 0.75:1 (O), 1:1 (0), 2:1 (2), and 4:1 (]), respectively, measured at θ ) 30°.

measured at a scattering angle of θ ) 30°. Clearly, the supramolecular assembled PDMS-Fol complex particles were formed since neither PDMS-di-NH2 nor Fol solutions show selfaggregation in THF. It should be mentioned that PDMS-di-NH2 solutions show very weak scattering intensity due to the very similar refractive index between PDMS and THF. GPC has been used to examine the self-association behavior of PDMS-di-NH2 in THF. PDMS-di-NH2 solution only showed one GPC peak with a number-average molecular weight (Mn) value of 3140 and polydispersity of 1.78, suggesting no self-aggregation of PDMS-di-NH2 in THF. For Fol solution in THF, we did a DLS measurement at a concentration of 5 mg/mL at θ ) 30°, which shows a single species with an approximate 〈Rh〉 of 0.8 nm. Although the DLS experimental error is relatively big for such small particles with low scattering intensity, this result indicates that the Fol molecules in dilute THF solution have nearly no self-association. The DLS result described below for solution of composite 4 with 〈Rh〉 of about 0.56 nm for excess Fol molecules provides further support that the Fol molecules are nearly monomeric in dilute THF solution. In contrast, the mixing of the Fol solution with the PDMS-di-NH2 solution produced large particles in THF. At a low molar ratio of OH/NH2 groups of 0.75:1 in composite 1, the PDMS chains are in excess. Nearly monodisperse small complex particles with an average 〈Rh〉 about 81 nm were formed. The increase in the molar ratio of the OH/NH2 groups to 1:1 in composite 2 increased the 〈Rh〉 of complex particles to 199 nm. This increase in the size of the complexes is reasonable since the additional Fol can complex more excess PDMS chains to form larger particles. Surprisingly, the further increase in the molar ratio of the OH/NH2 groups to 2:1 in composite 3 did not affect the size and size distribution of the complex particles. However, the scattering intensity from the solution of composite 3 was much stronger than that from the solution of composite 2, implying that more Fol molecules aggregated to form larger Fol nanodomains in composite 3 than in composite 2, because the PDMS chains linked between these Fol nanodomains inside the complex particles have no contribution to scattering intensity due to the nearly same refractive index of PDMS and THF. Interestingly, the continuous addition of Fol molecules to the PDMS-di-NH2 solution could eventually break the large complex particles into small particles coexisting with the excess Fol molecules. For example, at the molar ratio of OH/NH2 groups of 4:1 in the solution of composite 4, two species were determined from dynamic LLS. The small species with 〈Rh〉 of 0.56 nm could be attributed to the excess free Fol molecules in THF, while the broadly distributed species with 〈Rh〉 of 8.2 nm should be attributed to the Fol/PDMS-di-NH2 complex particles.

Figure 2. Angular dependence of Rh distributions of Fol/PDMS-diNH2 complex particles in THF (a) formed at 30 mg/mL PDMS-diNH2 and a molar ratio of OH/NH2 groups at 2:1, and (b) 10 times dilution of the mixture solution from (a).

Figure 2a shows the Rh distributions of the complex particles in the solution of composite 3 at a molar ratio of OH/NH2 groups of 2:1 determined at different scattering angles. No apparent angular dependence was observed for the nearly monodisperse Fol/PDMS-di-NH2 complex particles with 〈Rh〉 of about 195 ((5) nm, indicating a globular conformation for the supramolecular assembled complex particles. However, the large complex particles could break into smaller particles upon dilution. As shown in Figure 2b, the 10 times dilution of the solution of composite 3 not only reduced the size of the complex particles, but also broadened the size distributions. Furthermore, the complex particles formed from the dilute precursor solutions showed clear angular dependence. These results indicated that the complex particles formed at dilute solutions were more irregular in size and shape. It is envisaged that, as the mixture solution of Fol and PDMS-di-NH2 becomes more concentrated upon evaporation, the degree of association between the two components will be increased, which eventually leads to flexible, elastic, and free-standing network-like gels or films. Although the fullerene might covalently react with the terminal -NH2 groups of PDMS-di-NH2, we believe that the supramolecular assembly between Fol and PDMS-di-NH2 in solutions is predominantly driven by the hydrogen bonding interactions between the -OH groups of Fol and the -NH2 groups of PDMS, which can be weakened upon dilution resulting in irregular size and shape of complex particles. To further confirm the negligible effect of possible covalent reactions between Fol and PDMS-di-NH2, the homogeneous mixture of fullerene and PDMS-di-NH2 in toluene was stirred for 72 h. Without the design of hydrogen bonding interactions between Fol and PDMS-di-NH2, no composites could be formed due to the severe phase separations of fullerene and PDMS-di-NH2. Figure 3 shows the SAXS profiles of pure PDMS-di-NH2 and dried Fol/PDMS-di-NH2 nanocomposites formed at different molar ratios of OH/NH2 groups. While PDMS-di-NH2 had no scattering peaks in the experimental q range, the complexation of Fol to the PDMS matrix produced a single scattering peak, implying the homogeneous distributions of the Fol domains in the PDMS matrix but no highly ordered close-packing structure.

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Figure 3. SAXS profiles of the dried Fol/PDMS-di-NH2 nanocomposites. Curve 1: PDMS-di-NH2. Curves 2, 3, 4, and 5: composites 1, 2, 3, and 4, respectively.

The increase of Fol content in the nanocomposites gradually increased the scattering intensity of the peak due to the increase of electron density difference between Fol nanodomains and the PDMS matrix. Meanwhile, the increase of Fol content in the composites also shifted the peak position to lower q values. The peak positions centered at 0.93, 0.81, 0.76, and 0.65 nm-1 for composites 1, 2, 3, and 4 correspond to a gradual increase in the interdistances of Fol nanodomains at 6.8, 7.8, 8.3, and 9.7 nm, respectively. In our design for the Fol/PDMS-di-NH2 nanocomposite formation, the Fol molecules are constrained by the end-functionalized PDMS-di-NH2 chains, which means that the distance between the neighbored Fol nanodomain surfaces should be a constant and equivalent to the PDMS chain length on average. Therefore, the increase in the interdistances of Fol nanodomains could only be explained by the increase in the size of individual Fol domains. As shown in dynamic LLS results, the complexation of Fol with PDMS-di-NH2 chains in solutions is concentration dependent. In dilute solutions, the excess Fol molecules can freely diffuse in THF. However, when the mixture solution becomes more concentrated upon dilution, it is possible for the Fol molecules to self-assemble to form larger nanodomains instead of a single Fol molecule to complex with the PDMS-di-NH2 chains, especially when the Fol molecules are in excess to complex all the NH2 end groups of PDMS chains. Thermal Analyses. Figure 4a shows the thermogravimetric curves in terms of weight loss for the Fol/PDMS-di-NH2 nanocomposite films formed at different molar ratios of OH/ NH2 groups. In comparison with the PDMS-di-NH2 precursor, the main degradations of Fol/PDMS-di-NH2 composites move to higher temperatures, which can be seen more clearly in the derivative weight-loss profiles as shown in Figure 4b. PDMSdi-NH2 exhibited two depolymerization peaks, reaching the maximum rate at 424 and 589 °C, respectively. In contrast, composites 1, 2, and 3 moved them to 574 and ca. 640 °C, respectively, indicating that the thermal stability of PDMS-diNH2 is significantly enhanced after complexing with Fol molecules. For composite 4, the two depolymerization peaks were located at 514 and 619 °C, respectively. Clearly, composite 4 still dramatically enhanced the thermal stability of PDMSdi-NH2, but less prominently in comparison with composites 1, 2, and 3. The formation of relatively large Fol domains with excess Fol molecules in composite 4 may be responsible for the less prominent enhancement of the thermal stability, because the uncomplexed Fol molecule attains a maximum degradation rate at 255-350 °C. The DSC measurements show that the glass-rubber transition temperatures (Tg) of composites 2, 3, and 4 are, respectively,

Figure 4. Thermogravimetric analyses for various samples: (a) weightloss profile; (b) derivative weight-loss profile. Curve 1: PDMS-diNH2. Curves 2, 3, 4, and 5: Fol/PDMS-di-NH2 composites 1, 2, 3, and 4, respectively. Curve 6: Fol.

Figure 5. Melting behavior of various samples during DSC heating scans after annealing at -95 °C for 30 min. Curve 1: PDMS-di-NH2. Curves 2, 3, 4, and 5: Fol/PDMS-di-NH2 nanocomposites 1, 2, 3, and 4, respectively. Curves vertically shifted for clarity.

-119.4, -118.1, and -117.0 °C, which are slightly higher than that of PDMS-di-NH2 at -121.1 °C. The restriction of the chain mobility of PDMS-di-NH2 arising from the complexation with Fol molecules is responsible for the increase of Tg. However, the restriction is weak since PDMS-di-NH2 is only end-linked by Fol. The PDMS chain length (Mn ) 3140) between Fol crosslinking points is long enough to maintain the excellent mobility of PDMS chains. In other words, the advantage of PDMS chain flexibility is remained in our nanocomposites due to the very low cross-linking density. Figure 5 shows the melting behavior of various Fol/PDMSdi-NH2 composites during DSC heating scans. After annealing at -95 °C for 30 min, PDMS-di-NH2 sample shows double endothermic peaks commonly observed for linear PDMS at -52.8 and -41 °C, respectively, which is ascribed to the melting-recrystallization of the original crystallites.43 With increasing content of Fol in the nanocomposites, the endothermic peaks gradually diminished, and the high-temperature endothermic peaks gradually shifted to lower temperatures and eventually disappeared. The fusion enthalpy was greatly reduced from 14.2 J/g for PDMS-di-NH2 to 6.0, 1.6, 0.7, and 0.7 J/g (uncorrected by weight percentage of Fol) for composites 1, 2, 3, and 4, respectively. It may be concluded that the mobility of

Properties of Fullerenol/PDMS Nanocomposites

Figure 6. Thermal mechanical stability of various samples: (a) PDMSdi-NH2; (b) curves 1, 2, 3, and 4 for Fol/PDMS-di-NH2 nanocomposites 1, 2, 3, and 4, respectively.

Figure 7. Temperature dependence of permittivity ′ and loss factor ′′ of Fol at 0.3 (O, 0), 3 (×), and 30 kHz (9). The inset in ′′ plot shows the shoulder peak of Fol at low temperatures.

whole PDMS chains is restricted due to the end linkage with Fol, which can greatly suppress the crystallization of PDMS chains and inhibit the melting-recrystallization process. TMA results reveal that Fol/PDMS-di-NH2 nanocomposites have superior thermal mechanical stability in comparison to their polymer precursor. As shown in Figure 6a, PDMS-di-NH2 experienced a glass-rubber transition at -109 °C indicated by a sharp increase of dimension change, and then was penetrated by the probe at -41 °C. For the Fol/PDMS-di-NH2 nanocomposites as shown in Figure 6b, the softening temperatures dramatically increased with increasing content of Fol. For example, the thermal penetration for composite 1 was around 0 °C, which was increased to 41 °C for composite 2. No thermal penetration was observed for composites 3 and 4 within the experimental temperature range. Dielectric Property. Figure 7 shows the dielectric permittivity (′) and loss factor (′′) of Fol as a function of temperature at several frequencies. Both ′ and ′′ values increased dramati-

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Figure 8. Temperature dependence of ′ and ′′ of PDMS-di-NH2 (star symbols), Fol/PDMS-di-NH2 composite 2 (solid symbols), and composite 3 (open symbols). ′: 0.3 (square), 3 (up triangle), and 30 kHz (circle). ′′: at 0.3 kHz.

cally with increasing temperature for the highly polar Fol molecules. For example, the ′ and ′′ values increased from 6.5 to over 100, and from 0.6 to over 100, respectively, when the temperature was increased from -150 to 0 °C at 300 Hz. The two shoulder-like peaks in its ′′ spectrum (see the inset in Figure 7b) located at -100 and -30 °C at 300 Hz might arise from the relaxations of absorbed water44 and the hydroxyl addents of Fol, respectively. The abrupt increases of both ′ and ′′ at high temperatures are due to direct current (dc) conductivity. Figure 8 shows the temperature dependence of ′ and ′′ of Fol/PDMS-di-NH2 nanocomposites at several frequencies. In the ′ spectra, the ′ value of PDMS-di-NH2 increases sharply from -150 to -130 °C and decreases gradually later. The increase of ′ value corresponds to the glass transition of PDMSdi-NH2, and the later decrease may be due to its crystallization.45 Compared to PDMS-di-NH2, the Fol/PDMS-di-NH2 nanocomposites show broader glass transitions (′ increases from -150 to -110 °C). The interactions between Fol and PDMS-di-NH2 can suppress the mobility of dipoles, which may account for the lower ′ value of composite 2 than that of PDMS-di-NH2. Very interestingly, composite 3 shows a higher ′ value than that of PDMS-di-NH2. As discussed in the SAXS results, the higher Fol content in the Fol/PDMS-di-NH2 nanocomposites could form relatively larger Fol domains from the self-assembly of Fol molecules constrained by the end-linked PDMS chains. In these Fol domains, only those OH groups located on the domain surface were linked with the PDMS chains. The part of Fol molecules without interactions with PDMS-di-NH2 could enhance the ′ value, which explains the higher ′ value of composite 3 than that of PDMS-di-NH2. In the ′′ spectra, PDMS-di-NH2 showed a sharp ′′ relaxation peak centered at -142 °C, which could be assigned to the motions of unstrained segments in the purely amorphous state45 and correspond to the glass-rubber transition. The multiple ′′ peaks of PDMS-diNH2 from -60 to 30 °C might be attributed to the melting of its crystalline phase. The Fol/PDMS-di-NH2 nanocomposites showed much smaller values of ′′ compared to PDMS-di-NH2 and Fol, due to the severe restraints of dipole mobility by

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Ouyang et al. in the cases of composites 3 and 4, indicating that the elastic response dominated at high contents of Fol in the nanocomposites. The comparison of Tan δ at 20 °C is shown in Figure 10. All nanocomposites presented lower Tan δ values than that of the PDMS-di-NH2 polymer precursor, suggesting an increase in the stored mechanical energy and a decrease in the dissipated mechanical energy. Similar results were obtained when the measurements were conducted at -20 °C. Conclusions

Figure 9. Frequency dependence of storage moduli G′ (O) and loss moduli G′′ (4) of various samples at 20 °C.

Figure 10. Frequency dependence of Tan δ for various samples at 20 °C: PDMS-di-NH2 (b); Fol/PDMS-di-NH2 nanocomposites 1 (9), 2 (0), 3 (2), and 4 (4), respectively.

interactions. The glass transitions of the nanocomposites were broadened and moved to higher temperatures by ca. 20 °C, and the single melting peaks were moved to lower temperatures, which were centered at -45 and -55 °C for composites 2 and 3, respectively. The melting behavior shown by dielectric analysis is reminiscent of that revealed by DSC. These changes also arise from the restriction effect of Fol on the PDMS crystallization. The ′′ value does not change much with increasing Fol content. It is noteworthy that, while the Fol/ PDMS-di-NH2 nanocomposites show much smaller ′′ values than the PDMS-di-NH2 below 30 °C, their ′ values can be adjusted to be higher than that of PDMS-di-NH2 when suitable Fol content is controlled in the nanocomposites. This unique dielectric property of Fol/PDMS-di-NH2 nanocomposites is a great advantage to make novel materials with high ′ but much lower ′′ values. Viscoelastic Properties. Figure 9 compares storage moduli (G′) and loss moduli (G′′) of Fol/PDMS-di-NH2 nanocomposites with those of the polymeric precursor at 20 °C. The incorporation of Fol into PDMS-di-NH2 matrix dramatically increases both G′ and G′′ by orders of magnitude. Both G′ and G′′ were increased with increasing content of Fol in the composites. A viscous behavior was observed with G′′ > G′ in the cases of PDMS-di-NH2 and composites 1 and 2, while G′ exceeded G′′

Supramolecular assembled fullerene/PDMS nanocomposites could be prepared from the solution casting of mixtures of Fol and PDMS-di-NH2 in THF at different ratios. Dynamic LLS measurements reveal that there are strong associations between the Fol and the PDMS-di-NH2 in solution due to the strong hydrogen bonding interactions between the OH and NH2 groups. Depending on the molar ratios of OH/NH2 groups and the solution concentrations, various complex particles in different sizes and size distributions could be formed. SAXS profiles indicated that the Fol nanodomains are homogeneously distributed in the PDMS matrix. The higher Fol content in the composites could lead to the formation of relatively larger Fol nanodomains. Compared to the PDMS-di-NH2 precursor, the nanocomposites exhibit higher glass transition temperatures, superior thermal and thermal mechanical stability, and greatly suppressed crystalline phase. Moreover, the nanocomposites possess very attractive dielectric properties; i.e., high content of Fol increased permittivity while it severely decreased the loss factor of the nanocomposite materials. The investigation on viscoelastic properties show that the incorporation of Fol into PDMS matrix dramatically increases the storage and loss moduli of the composites; meanwhile, elastic response gradually exceeded viscous response with increasing content of Fol in the composites. The Tan δ values of the composites are lower than that of the polymeric precursor. Acknowledgment. We acknowledge the financial support of this work from the US National Science Foundation (CHM 0316078). S.Z. is grateful to Dr. Igors Scis for his help on the X-ray beamline setup for the SAXS experiments. References and Notes (1) Jensen, A. W.; Wilson, S. R.; Schuster, D. I. Bioorg. Med. Chem. 1996, 4, 767. (2) Withers, J. C.; Loutfy, R. O.; Lowe, T. P. Fullerene Sci. Technol. 1997, 5, 1. (3) Imahori, H.; Sakata, Y. AdV. Mater. 1997, 9, 537. (4) Prato, M. J. Mater. Chem. 1997, 7, 1097. (5) Smalley, R. E.; Yakobson, B. I. Solid State Commun. 1998, 107, 597. (6) Sun, Y. P.; Riggs, J. E. Int. ReV. Phys. Chem. 1999, 18, 43. (7) Diederich, F.; Go´mez-Lo´pez, M. Chem. Soc. ReV. 1999, 28, 263. (8) Wudl, F. J. Mater. Chem. 2002, 12, 1959. (9) Rincon, M. E.; Hu, H.; Campos, J.; Ruiz-Garcia, J. J. Phys. Chem. B 2003, 107, 4111. (10) Chen, Y.; Huang, Z. E.; Cai, R. F.; Yu, B. C. Eur. Polym. J. 1998, 34, 137. (11) Geckeler, K. E.; Samal, S. Polym. Int. 1999, 48, 743. (12) Dai, L.; Mau, A. W. H.; Zhang, X. J. Mater. Chem. 1998, 8, 325. (13) Sun, Y. P.; Lawson, G. E.; Huang, W.; Wright, A. D.; Moton, D. K. Macromolecules 1999, 32, 8747. (14) Huang, X. D.; Goh, S. H. Macromolecules 2000, 33, 8894. (15) Zhou, P.; Chen, G. Q.; Hong, H.; Du, F. S.; Li, Z. C.; Li, F. M. Macromolecules 2000, 33, 1948. (16) He, J. D.; Wang, J.; Li, S. D.; Cheung, M. K. J. Appl. Polym. Sci. 2001, 81, 1286. (17) Ford, W. T.; Lary, A. L.; Mourey, T. H. Macromolecules 2001, 34, 5819. (18) Brabec, C. J.; Dyakonov, V.; Sariciftci, N. S.; Graupner, W.; Leising, G.; Hummelen, J. C. J. Chem. Phys. 1998, 109, 1185.

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