Dynamic Transitions and Creep in Carbon Nanofiber

Article pubs.acs.org/IECR

Dynamic Transitions and Creep in Carbon Nanofiber/ Polydimethylsiloxane Nanocomposites with Meticulously Architectured Polymer−Filler Interfaces Nabarun Roy† and Anil K. Bhowmick*,†,‡ †

Indian Institute of Technology, Kharagpur-721302, India Department of Chemistry, Indian Institute of Technology, Patna- 800013, India

‡

ABSTRACT: This work intends to produce smart materials for use in advanced applications such as artificial load resisting tissues and articular implants by exploiting the good strength of carbon nanofibers (CNFs) and the excellent biocompatibility of polydimethylsiloxane (PDMS). In this work, a series of CNF/PDMS nanocomposites was fabricated by varying the method of preparation and nanofiber surface modification. Dynamic mechanical properties and creep behavior of these hybrid materials were studied in order to comprehend the reinforcing effect of the nanofibers on the polymer matrix. The improvement was overwhelming with 74% enhancement in storage modulus and 1100% reduction in creep compliance for just 4 phr of amine modified CNF. Property improvement in the nanocomposites is a function of the extent of nanofiber dispersion in the polymer matrix which was examined by High Resolution Transmission Electron Microscopy (HRTEM) and Field Emission Scanning Electron Microscopy (FESEM).

1. INTRODUCTION With the transformation from conventional fillers to nanofillers, the field of composite science and technology has undergone a mammoth change thereby extending its reach to a wider gamut of applications.1−4 These unconventional nanoparticles have immense potential to bestow their virtues and hence renovate a conventional material suitable for unconventional applications. Myriads of variables owe their responsibility in improving these properties, most important being extent of dispersion, mean aspect ratio of the particles, nature of the interface, and nanofiller surface modification. Hence, in order to prepare nanocomposites with improved properties, one should optimize the variables and figure out the contribution of each one of them in property improvement. An elaborate study of nature of interface of nanocomposites reveals that polymer chains near the nanofiller surface execute an altogether different chain dynamics.5−7 The surface atoms of the nanofiller being at a higher energy state interact with the surrounding polymer molecules and hence restrict their mobility.8 This phenomenon is entirely reliant to the compatibility of polymer and filler and hence to polymer− filler adhesion.9Apart from interfacial interactions, proper dispersion of nanofillers in a polymer matrix is an important criterion which influences the bulk properties of the polymer. Some important factors in improving nanofiller dispersion in a polymer matrix are method of nanocomposite preparation and nanofiller surface modification. Our previous studies elucidate the dependence of various properties of nanocomposites on these factors.10,11 PDMS shows a unique behavior at low temperature viz., low temperature treatment of PDMS gives materials with high crystallinity.12,13 However, the molecular architecture of crystalline PDMS is still a matter of tussle between scientists working with this interesting macromolecule.14,15 In the beginning of the past decade, Albouy16 confirmed the helical © 2012 American Chemical Society

structure of PDMS in the crystalline state through XRD studies. Crystallization of polymers is a phenomenon known to be contributed by well balanced thermodynamic and kinetic parameters. The intricate macromolecular chain dynamics accomplish themselves in spherulite formation with welldefined crystalline and amorphous zones. Literature acknowledges several investigations on kinetics of cold crystallization and dynamics of amorphous segments of PDMS at low temperature.17−19 Very recently, Lund et al. pursued a detailed study of the dynamics of PDMS molecules at low temperature in the light of cold crystallization phenomenon.20 In another contemporary study, Fragiadakis and Pissis justified through various analytical techniques that chain dynamics in PDMS is affected dramatically by nanofiller incorporation.21 The same group in another study explained an independent slower αrelaxation besides the one associated with glass transition and related this to the steady enhancement of the relaxation times of the molecules near the nanofiller surface.22 Though a good amount of work has been pursued on silica/PDMS nanocomposites, detailed dynamic mechanical studies of PDMS nanocomposites with nanocarbons, particularly carbon nanotubes and nanofibers, is rather very scanty.4,23 From the application point of view, another important property of polymers is creep. Although amorphous polymers undergo almost complete dimensional recovery, crystalline polymers exhibit momentous discrepancies in this behavior due to their structural complicacy.24−26 In addition, it is a well established fact that nanofiller incorporation into a polymer matrix significantly reduces creep deformation. Since PDMS Received: Revised: Accepted: Published: 9571

April 20, 2012 June 18, 2012 June 25, 2012 June 25, 2012 dx.doi.org/10.1021/ie301027s | Ind. Eng. Chem. Res. 2012, 51, 9571−9580

Industrial & Engineering Chemistry Research

Article

Figure 1. Synthesis of in situ and ex situ nanocomposites with sample designation. Here, PD and C stand respectively for hydroxyl end-capped PDMS and carbon nanofibers. A stands for amine modified CNF, E stands for ex situ prepared nanocomposites, and 1,2,4 etc. stand for parts of CNF per hundred grams of rubber (phr).

2. EXPERIMENTAL SECTION

possesses microcrystalline domains, it is also expected to show some anomaly in creep behavior. Combining all the points mentioned earlier, it can be figured out that although the literature speaks out of some works on CNF/PDMS nanocomposites, detailed correlation of dynamic mechanical properties of these nanocomposites with the extent of dispersion has not yet been executed. Hence, tuning the degree of dispersion by manipulating the method of nanocomposite preparation and nanofiller surface modification and investigating its effect on dynamic properties of PDMS nanocomposites is completely new and unexplored. In addition, investigation of creep behavior of PDMS in the presence of carbon nanofiber is promising since there is no mention of such work in the literature. In this paper, we have synthesized CNF/PDMS nanocomposites by tuning various factors: method of nanocomposite preparation (in situ polymerization and ex situ solution casting technique) and chemical functionalization of the nanofiber. The nanocomposites prepared were subjected to dynamic mechanical analysis by manipulating the test conditions: temperature sweep at a particular frequency, frequency sweep through the same range of temperature, and creep studies at a constant applied stress. In addition, we compared the storage modulus of the nanocomposites with the help of the Halpin−Tsai model. The results obtained from these studies were correlated with the morphological analysis by Wide Angle X-ray Diffraction (WAXD), High Resolution Transmission Electron Microscopy (HRTEM) and Field Emission Scanning Electron Microscopy (FESEM). This type of study has a thrust when practical application is a point of concern. PDMS has an undisputed acceptance as a biocompatible material and has been used in biomedical application for over four decades. Hence, incorporation of CNF into this biocompatible polymer may end up in the development of some advanced materials likely to be used in smart applications such as artificial connective tissues and articular implants.

2.1. Materials. Octamethylcyclotetrasiloxane, [(CH3)2SiO]4 (D4) purity >99% (GC), was supplied by Momentive Performance Materials, Bangalore, India. CNFs and tetraethoxysilane (TEOS) were obtained from Applied Sciences Inc., USA and Acros Organics, New Jersey, USA, respectively. Dibutyltin dilaurate (DBTDL) was procured respectively from Sigma Aldrich, USA and Aldrich Chemicals, Bangalore, India. Hexamethylenediamine (HMDA) was procured from Merck, Merck Schuchardt OHG, Germany. Potassium hydroxide was purchased from Merck, Mumbai, India. 2.2. Synthesis of Pristine PDMS and PDMS Based Composites. Functionalization of the nanofiller and synthesis of the in situ and ex situ nanocomposites were pursued in accordance with our previous works as shown in Figure 1.10 2.3. Characterization of Synthesized CNF/PDMS Nanocomposites. 2.3.1. High Resolution Transmission Electron Microscopy (HRTEM). The samples for HRTEM analysis were prepared by ultracryomicrotomy with a Leica Ultracut UCT (Leica Microsystems GmdH, Vienna, Austria). Freshly sharpened glass knives with cutting edges of 45° were used to obtain cryosections of about 100−150 nm thickness at −150 °C. Microscopy was performed with JEOL 2100, Japan. Transmission electron microscope was operated at an accelerating voltage of 200 kV. 2.3.2. Field Emission Scanning Electron Microscopy (FESEM). Field Emission Scanning Electron Microscopy (FESEM) of the fractured surface of the nanocomposite was carried out in order to study the fiber reinforcement phenomenon along with the extent of dispersion. This was pursued with FESEM S4800 Hitachi microscope. The optimized conditions for the analysis were an acceleration voltage 10.0 kV and a working distance of 10.1 mm. 9572

dx.doi.org/10.1021/ie301027s | Ind. Eng. Chem. Res. 2012, 51, 9571−9580


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2.3.3. X-ray Diffraction (XRD) Studies in Wide Angle Range. The wide angle X-ray diffraction profiles of the samples were documented in a Philips X-ray diffractometer (model PW1710) utilizing crystal monochromated Cu Kα radiation in the 2θ range 10−70°. In all the experiments, an operating voltage of 40 kV and a current of 20 mA were kept constant. The crystallite size was calculated according to the Scherrer equation27 Lhkl = Kλ /β cos θ

(1)

where Lhkl is the size of the crystallites at reflection of hkl, θ is half the Bragg’s angle, K is the Scherrer factor, and β is the full width at half-maximum. λ is the wavelength of the X-ray used. 2.3.4. Dynamic Mechanical Analysis (DMA). 2.3.4.1. Temperature Sweep Study. The dynamic mechanical data of the nanocomposites (12.59 mm × 6.65 mm × 1.2 mm) were acquired using a DMA of TA Instruments (model Q800). The runs were performed at a constant frequency of 1 Hz, a strain of 0.05%, and a temperature range from −125 to 50 °C at a heating rate of 3 °C/min. The data were scrutinized by TA Universal analysis software. Storage modulus (E′) and loss tangent (tanδ) were recorded as a function of temperature for all the samples under the same conditions. 2.3.4.2. Frequency Sweep Study. The study was carried out in the range of −125 to 25 °C with Q800 where data were recorded at the interval of 10 °C. The sample was subjected to isothermal condition at each temperature for 3 min followed by frequency sweep throughout the entire frequency range of 0.05 to 15 Hz. The strain was set at 0.05%. Storage modulus was recorded and analyzed by TA data analysis software. 2.3.4.3. Creep Study. Creep compliance of the nanocomposite samples was measured by a dynamic mechanical analyzer (Q800 of TA Instruments) in the tensile mode. In the experiment, the sample with a dimension of 6.25 mm wide × 30 mm long × 1.5 mm thick was subjected to a constant stress of 0.025 MPa, and the resulting strain was recorded at 30 °C. Compliance D(t) was calculated from the stress and strain data using eq 228

D(t ) =

ε(t ) σ0

Figure 2. Representative HRTEM images of (a) PD C4E (ex situ prepared nanocomposite with unmodified CNF), (b) PD C4 (in situ prepared nanocomposite with unmodified CNF), and (c) PD C4A (in situ prepared nanocomposite with amine modified CNF) (CNF loading in each case is 4 phr).

nanocomposites prepared with amine modified CNF, the amine functional groups interact with the growing polymer chains, thereby further improving the extent of dispersion. This is reflected in the HRTEM micrograph shown in Figure 2(c). In our previous publication, we quantified the dispersion of the nanofibers in the nanocomposites prepared by various methods with the unmodified and the amine modified nanofibers.10,11 In situ prepared nanocomposites with amine modified CNF gave the best and most uniform dispersion (defined by the parameter D0.1, where a higher magnitude of the parameter signifies more homogeneous and better dispersion) among all the nanocomposites. The D0.1 was found to be the highest for the amine modified CNF filled nanocomposites (28.31%), while ex situ prepared nanocomposites exhibited the poorest dispersion with a very low value of 11.26%.10 The effect of difference in dispersion is well reflected in the results of dynamic mechanical analysis and creep studies mainly discussed in the subsequent sections. 3.2. Fracture Morphology Studies through FESEM. The FESEM micrographs offer a handful of evidence of strong polymer−filler interface formation in the case of amine modified CNF/PDMS nanocomposites. Figure 3(a)-(b) shows the fractured surface of the nanocomposite. A very high quality dispersion of the nanofibers at the fractured surface can be unanimously established from Figure 3(b) for amine modified samples. In the nanocomposite, cracks get generated and hence propagate through the polymer and polymer−filler interface due to which fractured surface is unusually rough (Figure 3(b)).

(2)

where ε(t) and σ0 are respectively strain and stress on the sample.

3. RESULTS AND DISCUSSION 3.1. HRTEM Studies. Morphology studies have been reported here to explain the results of dynamic transitions and properties. Figure 2(a)-(c) shows respectively the representative HRTEM images of conventional solution cast ex situ nanocomposites with unmodified CNF and in situ prepared nanocomposites with unmodified CNF and amine modified CNF respectively at 4 phr loading of the filler. These low resolution micrographs display the effect of various preparative methods in influencing the extent of dispersion of the nanofillers in the matrix. The ex situ prepared nanocomposite shows prominent agglomeration (Figure 2(a)). However, in the in situ prepared nanocomposite, monomer molecules entrapped in between the CNFs tend to deagglomerate the latter during polymerization. This is evident from the HRTEM image shown in Figure 2(b) where dispersion is better compared with the one shown for the ex situ prepared nanocomposites. In the case of in situ 9573



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Figure 3. (a) shows fractured irregular surface for PD C8A and (b) shows good nanofiller dispersion at the interface with strong interface formation.

Figure 4. (a) WAXD profiles for CNF/PDMS nanocomposites and (b) scheme for disruption of microcrystalline domains by CNFs during nanocomposite formation (circles show breakage of crystalline domains by CNF).

respectively with filler incorporation. The peak around 12.5° in the XRD profile of PDMS corresponds to centered tetragonal unit cell.29 However, careful observation reveals a shift in peak position for the nanocomposites toward lower 2θ value with simultaneous broadening. With a gradual increase in filler amount, shift in peak position toward lower Bragg angle and peak broadening become more prominent. This may be due to the fact that the process of crystalline domain formation in the presence of the filler is restricted during nanocomposite preparation. A schematic representation of this phenomenon is depicted in Figure 4(b). This is the consequence of decrease in orderness of the crystalline phase, dislocations, point defects, etc. among the crystalline domains.29 In the presence of amine modified CNF, the peak around 12° shows a much more dramatic shift in position and a greater extent of peak broadening compared with the nanocomposites prepared with the unmodified CNF. This is probably due to the fact that unlike pristine CNF, amine functional groups on the filler surface interact with the growing polymer chain through H-bonding interaction. This phenomenon restrains the process of microcrystalline domain formation to a greater extent compared with unmodified CNF based nanocomposites by restricting the polymer chains to attain their native orientation.

A scrupulous examination of the micrographs reveals that strong polymer−filler interfaces are formed not only at the junction between the polymer and the nanofibers but also along the surface of the nanoparticles (these are being shown with circles in Figure 3(a) and 3(b)). The junctions of the nanofillers protruding out from the broken polymer surface seem to have fused with the polymer matrix. This is an indication of the strong polymer−filler interface formation. Once the crack approaches a nanofiber, it propagates through another direction thereby accounting for the unevenness of the fractured surface. The nanocomposite on its own part absorbs energy and thus requires a greater load to break. This is evident from the nanofiber “pull-out” and broken ends. These are not reported before for these systems. The effect of nanofiber reinforcement is reflected in dynamic mechanical property improvement of the nanocomposites. 3.3. WAXD Analysis. WAXD studies offer a concise idea about the effect of filler on the matrix crystalline morphology in the nanocomposites. Figure 4(a) shows the WAXD pattern of the unfilled and the CNF filled PDMS nanocomposites in the 2θ range of 10−15°. This portion of the XRD profiles is of importance, since the change in the peak position and full width at half-maximum give a clear idea of the shift in d-spacing and change in microcrystallite size distribution of the matrix 9574



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ogy of the matrix. It is observed from WAXD studies that with an increase in filler concentration the size of the crystalline domains decreases. A good amount of storage modulus is owed to crystalline nanodomains which, in fact, act as cross-links imparting strength to the matrix.31 Thus, a decrease in the crystalline domain size in the matrix leads to a simultaneous decrease in storage modulus. On the other hand, modulus values for the nanocomposites increase with an increase in polymer−filler interaction due to strong interface formation. As a consequence, the two opposing factors lead to a marginal increase in modulus in the low temperature zone for the nanocomposites as observed in Figure 5. Thus, the magnitude of the low temperature modulus is a product of the conflicting effects played by a decrease in domain size and an increase in polymer−filler interaction. However, at room temperature and above, a reinforcing effect of the nanofillers takes an upper hand, which is reflected in the modulus values in the higher temperature span. Ex situ nanocomposites show a poor dispersion state of nanofibers which is reflected in the significantly low modulus value at room temperature and above compared with the in situ prepared nanocomposites over the entire temperature range. This is observed from the results in Table 1. Prominent

Quantification of the size of the microcrystalline domains in PDMS in the virgin state as well as in the presence of a filler has been done by using eq 1 (Scherrer equation). While for PD C0 crystalline domain size is 2.5 nm, it decreases to 2.3 nm for PD C8 and 2.0 nm for PD C8A. The nanofibers, being comparatively much larger in dimension, affect the size and number of the crystalline domains in the polymer. However, the ex situ prepared nanocomposite PD C8E does not show much decrease in crystalline domain size compared with PD C0. Thus, in situ preparation of nanocomposites and chemical modification of nanofiller contribute toward a decrease in the size of crystalline domains of PDMS matrix, the effect of which may be reflected in the dynamic mechanical properties and creep properties of the nanocomposites in the subsequent sections. 3.4. DMA Temperature Sweep Studies. Figure 5 shows the plot of storage modulus versus temperature for neat PDMS

Table 1. Storage Modulus Values for Unfilled and Filled PDMS Vulcanizates at Different Temperatures E′ at different temperatures sample PD PD PD PD

Figure 5. Plot of storage modulus versus temperature for various CNF/PDMS nanocomposites and comparison with the Halpin−Tsai model.

C0 C8 C8A C8E

−100 °C (MPa)

−70 °C (MPa)

25 °C (MPa)

930 1260 1550 1445

400 600 810 645

0.34 0.83 4.36 0.60

agglomeration of the nanofiller generates voids in the system thereby weakening it. For instance, while PDC8 shows a modulus value of 0.83 MPa, it is just 0.60 MPa for PDC8E. 3.4.2. Effect of Filler Functionalization on Storage Modulus of Nanocomposites. Filler functionalization is carried out in order to verify its effect on its dispersion state in the matrix. HRTEM analysis shows that a greater extent of filler dispersion is achieved with amine modified CNF filled nanocomposites. As a consequence, the reinforcing effect of the filler is at its maximum as observed from the storage modulus values for the nanocomposite over the entire range of temperature. Moreover, additional interaction in the form of Hbonding is imposed upon the matrix by the amine functionalities present on the filler surface due to modification. This is reflected in the room temperature modulus of PD C8A (4.36 MPa) which is much higher compared with that of PD C8 (0.83 MPa) (the modulus of the unfilled PDMS being just 0.34 MPa). However, increases in low temperature storage modulus (above Tg) was also prominent for nanocomposite with higher filler loading when compared with the unfilled vulcanizate. While for the unfilled PDMS vulcanizates, it is 2187 MPa at −125 °C, an increase of around 70% is observed for PD C8A. This is probably due to better dispersion of the nanofiller facilitated by the additional interaction between the amine groups on the nanofiber surface and the macromolecular backbone. The results are compiled in Table 1. 3.4.3. Modeling of Tensile Modulus of the Nanocomposites. The experimental results are compared with predictions from the Halpin−Tsai model32 in order to

and its in situ and ex situ prepared nanocomposites with unmodified and amine modified CNF as the reinforcing filler. Magnitude of storage modulus at room temperature shows a decent increase with an increase in the amount of nanofiller. This is due to the reinforcing effect of the latter on the matrix which, in a way, is the outcome of its good dispersion in the polymer matrix. While the unfilled PDMS shows three prominent transitions, these are altogether smeared in the nanocomposites, particularly for the in situ prepared ones. This is probably due to reduction in size and number of crystalline domains of the PDMS matrix in the presence of CNF. 3.4.1. Effect of Dispersion on the Elastic Properties of the Nanocomposites. Two dispersion techniques have been adopted in preparing the nanocomposites, viz., in situ and ex situ techniques. HRTEM analysis shows better nanofiller dispersion for in situ prepared nanocomposites, and this is reflected in the results of dynamic mechanical studies. It is worth mentioning that the reinforcing effect of filler is more vivid from the high temperature storage modulus for the nanocomposites.30 While the unfilled PDMS shows prominent transitions for crystalline phase development (around −75 °C) and crystalline melting (around −35 °C), blurring of both these transitions in the case of nanocomposites speaks of improper microcrystalline domain formation. Thus, low temperature storage modulus is governed by two conflicting factors viz., polymer−filler interaction and change in crystalline morphol9575



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−75 °C is converted into a hump covering a region of 50 °C. This is probably due to the loss in crystalline domains in the matrix in the case of the nanocomposites. In fact, with the increase in filler loading the size of the hump gradually decreases and hence for PD C8A the hump disappears. 3.5. Frequency Sweep Studies. Literature reveals indeed very less amount of work on frequency dependent modulus determination of silicone rubber. For instance, Hadjistamov studied the viscoelastic behavior of silicone oils of varying molecular weights through oscillatory measurements and framed the storage and loss modulus master curves.35 In a different study, Small and Wilson carried out the frequency sweep studies of silicone rubber and generated the shear storage and loss modulus TTS master curves.36 Zhang et al. investigated the frequency dependence of silicone resin applicable in high power LED packaging.37 However, there is no report on the frequency sweep studies of CNF/PDMS nanocomposites in the literature which warrants a detailed investigation. Figure 7(a) shows a compilation of the storage modulus (E′) versus temperature plot over a range of frequencies (0.05 to 15

investigate the effect of nanofiller aspect ratio and extent of dispersion on storage modulus values. The Halpin−Tsai model is expressed as (1 + ξηϕ) E = Em (1 − ηϕ)

(3)

where ξ is dependent on geometry and orientation of the filler, and η is given as Ef

η=

Em Ef Em

−1 +ξ

(4)

E, Em, and Ef are Young’s modulus of nanocomposite, unfilled elastomer, and nanofibers, respectively. ξ is equal to twice the aspect ratio of the carbon nanofibers. Figure 5 shows comparison of the nanocomposites prepared with the unmodified and the amine functionalized nanofibers at 4 phr loading with the Halpin−Tsai model. It is found that storage modulus values for PD C4A are in good coherence with the model in the high temperature region (above −30 °C). The modulus value at 25 °C for PD C4A matches exactly with the model, whereas for PD C4 it is slightly lower compared with the model. This difference is due to extent of dispersion. The low temperature modulus for the nanocomposites, however, shows a huge deviation from the model. This is ascribed to decrease in modulus which is consistent with the reduction in size and number of crystalline domains of the polymer in the presence of CNF as discussed earlier and is evident from XRD studies. 3.4.4. Effect of Polymer−Filler Interaction on Tg. Figure 6 shows the plot of tan δ versus temperature for the unfilled and

Figure 6. Plot of tan δ versus temperature for various CNF/PDMS nanocomposites.

the nanofiber filled PDMS vulcanizates. Polymer−filler interaction is evidenced from the decreasing height of the peak at Tg (−116 °C) for the nanocomposites, which is in accordance with the observation of Fornes and Paul.33 Careful observation of the tan δ plot for unfilled PDMS vulcanizate reveals three prominent peaks: the one around −116 °C corresponds to Tg, the second one around −75 °C is due to crystalline domains in PDMS,12 and the third one around −35 °C is for melting of crystallites.34 The peaks correspond to the transitions in the storage modulus plot. It is found that in the nanocomposites the crystalline peak around

Figure 7. (a) Storage modulus comparison of the unfilled and 8 phr amine modified CNF filled PDMS vulcanizates throughout a frequency range of 0.05 to 15 Hz and (b) storage modulus master curves for unfilled and amine modified CNF filled PDMS vulcanizates.

Hz). At a particular temperature, E′ increases with an increase in frequency for both the systems. In the plot, the different storage modulus values were obtained for different frequencies at a particular temperature, for instance the one shown with a dotted arrow. However, a frequency dependent increase is 9576



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where J(t) and J0 are the creep compliance after time t and instantaneous creep compliance, and Ji and τi are the constants characteristic of the system. The strain-time plot for PDMS vulcanizates in Figure 8 shows good coherence with the characteristic strain-time plot

more prominent in the case of the unfilled vulcanizates as observed from the inset in Figure 7(a). The data at low frequency are found to be uneven, particularly for the unfilled one. This trend is consistent with the temperature sweep data at high temperature. This happens since at lower frequency the sample is subjected to deformation on an extended time scale. In order to elaborate the time−temperature character of the various transitions, temperature independent behavior of the filled and unfilled polymer systems is investigated by shifting the storage modulus data along the frequency axis.38,39 A reference temperature (−65 °C) is chosen, and the modulus values are shifted at the reference temperature. Through time− temperature superposition principle a storage modulus master curve is framed by using a horizontal shift factor aT which is defined in the Williams−Landel−Ferry (WLF) eq 5 as40 log aT =

−C1(T − Tg) C2 + (T − Tg)

(5)

Here Tg is the glass transition temperature, and C1 and C2 are constants related to the molecular structure. The master curves for unfilled PDMS vulcanizate, and the systems with the best dynamic mechanical properties PD C8A are compared and shown in Figure 7(b). The master curves offer a clear insight into the nanofiller reinforcement of the polymer matrix through increased magnitude of storage modulus for the nanocomposite particularly in the low frequency zone. The constants C1 and C2 are determined, and the values are reported in Table 2. Values of C1 and C2 are found to increase

Figure 8. Comparison of strain versus time curves for various CNF loaded PDMS nanocomposites (under a constant stress of 0.025 MPa and at 30 °C temperature).

for viscoelastic materials. The plot clearly shows the various stages of creep viz., primary creep with relatively high strain rate, the secondary creep due to attainment of steady state. However, the tertiary creep leading to yielding of the material is not observed in any of the samples. Figure 8 shows the representative strain response to creep and recovery for the unfilled and CNF filled PDMS nanocomposites. The temperature chosen is 30 °C, and a constant stress of 0.025 MPa is applied. Because of the elastic nature of PDMS, instant deformation is the consequence of application of the constant stress. This is followed by a continuous and steady increase in strain as a function of time with a subsequent attainment of steady-state creep. On removal of the stress, there is a sudden fall in strain which is followed by a smooth decrease due to elastic recovery of the filled polymeric material. Strain recovery is generally below 100% because of the viscous nature of the material. The creep profiles of the CNF filled nanocomposites are of no exception. However, prominent reduction in the extent of deformation is perceivable with increasing filler content. 3.6.2. Creep Compliance. Figure 9 shows a comparative analysis of creep compliance for various unmodified and amine modified CNF filled nanocomposites respectively as a function of filler loading. It is quite vivid that with an increase in filler loading, in both cases there is a drastic fall in the magnitude of creep compliance. For instance, for the unmodified and the amine modified CNF based nanocomposites at 4 phr filler loading there is 225% and 1100% decrease in this parameter. This indicates improved creep resistance for the nanocomposites which is the consequence of enhanced polymer− filler interaction. Improvement in creep resistance is more pronounced for amine modified filler based nanocomposites as compared with those prepared with unmodified nanofibers. This is due to better nanofiller distribution and hence strong polymer−filler interface formation as evident from the HRTEM, FESEM, and DMA temperature sweep studies. This observation has coherence with those of Gojny et al.47

Table 2. WLF Constants for Unfilled and Amine Modified CNF Filled PDMS Vulcanizates sample

C1

C2

PD C0 PD C8A

10.4 39.0

152.3 289.3

with the addition of CNF which is due to reduction in fractional free volume since free volume f 0 is related to C1 as follows38

f0 =

B 2.303C1

(6)

where B is a constant and is assumed to be unity. 3.6. Creep Studies. 3.6.1. Creep Behavior. Literature shows that studies on creep behavior of silicone rubber and its composites are very scanty. While Wang et al.41,42 studied the creep behavior of silicone rubber filled with carbon black, glass beads were used as the reinforcing material by Chen et al.43,44 in studying the same behavior. However, there is still a wider scope of research in this field in order to explore the effect of nanofillers on the creep behavior of this elastomer. In addition, this behavior of silicone rubber is important, while application in implants is taken under consideration. Polymers being viscoelastic materials show “creep” behavior when subjected to a constant stress. The stress−strain curves are modeled by Kelvin and Voigt as45,46 J(t ) = J0 +

n

⎡

i=2

⎣

⎛ t ⎞⎤ ⎟⎥ ⎝ τi ⎠⎥⎦

∑ Ji ⎢⎢1 − exp⎜−

(7) 9577



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Figure 10. Comparison of the strain% for the PD C4 nanocomposite at different equilibrating temperatures.

Figure 9. Comparison of the creep compliance of unmodified and amine modified CNF filled PDMS vulcanizates with an unfilled one as a function of filler loading (constant stress of 0.025 MPa and a temperature of 30 °C were imposed).

4. CONCLUSIONS Integration of carbon nanofiber significantly changes the molecular dynamics of the PDMS matrix. Polymer−filler interaction which results in strong interface formation is the sole reason for this phenomenon. This was confirmed by FESEM analysis of the fractured surface. The nature of the interface was intricately designed by factors such as method of nanocomposite preparation and nanofiber surface functionalization, the consequence of which was reflected in the various dynamic mechanical studies. Storage modulus of the nanocomposites improved significantly upon nanofiber incorporation. While room temperature modulus of unfilled PDMS was found to be 0.34 MPa only, it increased to 4.36 MPa for 8 phr amine modified CNF loaded sample. In addition, creep compliance also showed prominent reduction upon in situ preparation of nanocomposites and nanofiller functionalization. For instance, reduction of 225 and 1100% in creep compliance was observed for the in situ prepared nanocomposites with 4 phr unmodified and the amine modified CNF. The factors which determined property improvement were strong polymer−filler interface formation and extent of nanofiber dispersion in the polymer matrix. These were examined respectively by FESEM and HRTEM analyses. However, nanocomposite formation also controls certain subsidiary factors such as microcrystalline domain formation (determined by WAXD studies), which received prominence in this study.

According to this group, nanofiller surface functionalization facilitates deagglomeration of the nanofillers and hence aids better dispersion and stronger interface formation. It is worth mentioning that unfilled PDMS shows a higher magnitude of creep compliance compared to CNF filled nanocomposites irrespective of the loading and surface modification. The method of nanocomposite preparation has an important role to play in creep behavior determination. Compared to in situ prepared nanocomposites, the conventional ex situ nanocomposites exhibit higher values of creep compliance, the key factor being the extent of nanofiber dispersion in the polymer matrix. This is explained by the HRTEM analysis of the nanocomposites. In the presence of the nanofiller, size of the crystallites decreases as evident from WAXD studies. Contradicting to this phenomenon is the improved polymer−filler interaction with strengthened interface regions. This facilitates generation of a close-packed structure which significantly improves the creep behavior.48 During recovery of the sample, the defects originated in the deformed sample can either relapse back to their original state or can recombine with other defect at the boundary through energy release and get converted into defect at similar energy level.49 The recovery compliance also follows the same trend as that of creep compliance. 3.6.3. Effect of Temperature. Creep deformation increases with an increase in temperature in the nanocomposite owing to viscous flow of the material at high temperature.50 The samples are subjected to creep studies at −50 °C, 0 °C, 30 °C, and 100 °C as evident from Figure 10. Unlike the other samples, the one at 100 °C shows tertiary creep (flow of the sample) irrespective of nanofiller loading. Hence, in this case thermal softening rules over the nanofiller strengthening. PDMS, due to its inherent weakness, undergoes rapid softening once the sample reaches the temperature. The extent of deformation goes on increasing until the time the sample is able to withstand the applied stress. This is followed by mechanical failure. With an increase in temperature the strain recovery gradually diminishes which ultimately ends up in failure of the sample. Thus, temperature has an important role in determining creep in filled and unfilled samples.

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AUTHOR INFORMATION

Corresponding Author

*Phone: 91-3222-283180/91-612-2277380. Fax: 91-3222220312/91-612-2277384. E-mail: [email protected]/ [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge the financial grant for this work by Council of Scientific and Industrial Research (CSIR), New Delhi.

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