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NMR Spin-Spin Relaxation Studies of Silicate-Filled Low Molecular Weight Poly(dimethylsiloxane)s Terence Cosgrove,* Claire Roberts, and Tania Garasanin School of Chemistry, University of Bristol, Cantock’s Close, Bristol, BS8 1TS United Kingdom
Randall G. Schmidt and Glenn V. Gordon Dow Corning Corporation, Midland, Michigan 48686-0994 Received April 30, 2002. In Final Form: September 13, 2002 Nuclear magnetic resonance (NMR) spin-spin relaxation measurements were used to investigate the mobility of poly(dimethylsiloxane) (PDMS) polymers, below the molecular entanglement point, when mixed with trimethylsilyl-treated polysilicate nanoparticles. The results showed that a high molecular weight polysilicate caused a dramatic reduction in the overall PDMS chain mobility at all concentrations. The relaxation decays were deconvoluted into multiple exponential decays using nonlinear least squares and the DISCRETE algorithm. The components of these decays were associated qualitatively with adsorbed and nonadsorbed polymer segments. When compared with differential scanning calorimetry measurements, the reduction in the mobility of the PDMS chains as seen in the NMR experiments corresponded to a shift in the glass transition to higher temperatures, a decrease in the specific heat increment at the glass transition, and a loss in the ability of the polymer to crystallize at high concentrations of polysilicate.
Introduction It is of both fundamental and practical interest to understand the influence of fillers on the physical, primarily flow and mechanical, properties of polymer melts and networks. To this end two poly(dimethylsiloxane) (PDMS) low molecular weight samples were blended with silicate particles of varying size on the nanometer scale. In this study, the silicate particles were trimethylsilylated to avoid the complications of interfacial interactions and agglomeration. It is envisaged that the understanding gained from this approach can be extended to active fillers such as fumed silica where interactions play an important role. This paper focuses primarily on NMR techniques to investigate the role of the nanosized silicate particles on the segmental mobility of a 5 and 12 kg mol-1 polymer. The mobility of polymer melts has been widely discussed in the literature, from both theoretical1 and experimental2 points of view. In particular, nuclear magnetic resonance (NMR) relaxation data are useful in determining the segmental mobility of chains, especially in different physical environments.3,4 A two-state model proposed by Brereton and co-workers1 was shown to successfully describe the segmental mobility of bulk PDMS melts, giving a semilogarithmic dependence of the inverse of the spin-spin relaxation time (T2) on the weight-average molecular weight (M h w).2 In the Brereton study, the polymers were narrow fractions with higher molecular weights; however, a slight nonexponential character was found with respect to the decay of the transverse nuclear * To whom correspondence should be addressed. E-mail:
[email protected]. (1) Brereton, M. G.; Ward, I. M.; Boden, N.; Wright, P. Macromolecules 1991, 24, 2068-2074. (2) Cosgrove, T.; Griffiths, P. C.; Hollingshurst, J.; Richards, R. D. C.; Semlyen, J. A. Macromolecules 1992, 25, 6761-6764. (3) Litvinov, V. M.; Spiess, H. W. Makromol. Chem., Macromol. Symp. 1991, 44, 33-36. (4) Cosgrove, T.; Griffiths, P. C. Adv. Colloid Interface Sci. 1992, 42, 175-204.
magnetization signal that was partially ascribed to a finite distribution of molecular weights between entanglements.5 Upon adsorption at an interface, polymer segments can no longer orient isotropically; consequently, the spinspin relaxation times will be reduced. The transverse relaxation time, T2, decreases with decreasing segmental mobility in the extreme narrowing regime corresponding to
ω2τc2 < 1
(1)
where ω is the Larmor frequency and τc is a correlation time describing a particular motion (e.g., rotation or translation). However, beyond a certain mobility limit, T2 becomes independent of τc. For polymers adsorbed as trains, in direct contact with the surface, very short values of T2 can be found.6 For segments that retain some mobility, T2 will depend on the local monomer concentration and the segmental mobility. In a previous study of a terminally attached polymer,7 it was found that the relaxation decay could be fitted by means of a distribution of relaxation times dependent on the volume fraction profile of the attached chains. In a more recent paper,8 the measured T2 values for silica-filled PDMS mixtures showed that the presence of the filler particles dramatically changed the mobility of the polymer chains. For silica concentrations greater than 25 wt %, most of the polymer segments in the dispersion experienced some degree of restricted mobility. This effect was molecular weight dependent with the high molecular weight chains being more influenced at a given silica concentration. The incorporation of silica in PDMS also (5) McCall, D. W.; Douglas, D. C.; Anderson, E. W. J. Polym. Sci. 1962, 59, 301-316. (6) Cosgrove, T.; Vincent, B.; Sissons, D. S.; Cohen Stuart, M. Macromolecules 1981, 14, 1018-1020. (7) Cosgrove, T.; Ryan, K. Langmuir 1990, 6, 136-142. (8) Cosgrove, T.; Turner, M. J.; Thomas, D. R. Polymer 1997, 38, 3885-3892.
10.1021/la025884+ CCC: $22.00 © 2002 American Chemical Society Published on Web 11/22/2002
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Table 1. Molecular and Physical Properties of Poly(dimethylsiloxane)sa
Table 2. Molecular and Physical Properties of Silicate Materialsa
sample name M h w/kg mol-1 M h w/M h n Rg/nm Rh/nm F/g cm-3 η0/Pa s 5K 12K
5.2 12.2
1.07 1.03
2.0 3.1
1.5 2.4
0.960 0.965
0.047 0.121
M h w is the weight-average molecular weight, M h w/M h n is the molecular weight distribution, Rg is the radius of gyration, Rh is the hydrodynamic radius, F is the density at 25 °C, and η0 is the zero-shear-rate viscosity at 25 °C. a
affects the ability of the melt to crystallize,9 and this is another indication of the interaction of the filler surface with the polymer chain in that system. These particular polysilicate nanoparticles dispersed in low molecular weight melts have also been studied by pulsed field gradient NMR.10 In that paper, it was discerned that a monolayer of PDMS coated the particles and that the layer thickness increased with melt molecular weight. In the present study, NMR spin-spin relaxation measurements were used to investigate the mobility of PDMS mixed with trimethylsilyl-treated silicates. These samples were used to minimize hydrogen-bonding effects with the polymer. Differential scanning calorimetry (DSC) measurements were obtained to corroborate the NMR data. This paper will focus on two low molecular weight PDMS polymers, both of which are below the critical molecular weight that characterizes the onset of entanglement effects in PDMS (Mc ∼ 30 kg mol-1).11 The effect of the silicates on higher molecular weight polymers (M h w g Mc) is reported in a companion paper.12 Experimental Section Materials. The linear PDMS and the silicate-based materials were supplied by Dow Corning Corp. Polymer and silicate were blended together at different weight ratios, and measurements were conducted approximately 1 month after preparation. An anionic polymerization process was used to obtain relatively monodisperse PDMS from hexamethylcyclotrisiloxane. The molecular and physical properties are given in Table 1. The information on molar and hydrodynamic size was obtained from size exclusion chromatography (SEC) which used Polymer Laboratories Mixed-D columns with a Waters 2690 HPLC+2410 refractive index detector coupled to a Viscotek T60a right-angle laser light scattering/viscometer dual detector and toluene as an eluent at 35 °C. The density was measured at 25 °C using an Anton PAAR DMA48 density meter. The zero-shear-rate viscosity at 25 °C was obtained from a Rheometric Scientific RDAII using 50-mm-diameter cone-and-plate fixtures with a cone angle of 0.04 rad. The molecular and physical properties of the silicate-based materials are given in Table 2. The lower molecular weight silicate, designated R1, was made from the acid-catalyzed hydrolysis and condensation of tetraethoxysilane, endcapped with hexamethyldisiloxane. Gas chromatography revealed that the R1 silicate was essentially a mixture of 68 wt % tetrakis(trimethylsiloxy)silane and 21 wt % hexabis(trimethylsiloxy)silane plus residual amounts of larger oligomeric species. The trimethylsilylated silicate designated R3 was obtained from acidcatalyzed polymerization of sodium silicate in a process described elsewhere13 and was glassy at room temperature. The molar mass was determined by SEC using Polymer Laboratories (9) Ebengou, R. H.; Cohen-Addad, J. P. Polymer 1994, 35, 29622969. (10) Roberts, C.; Cosgrove, T.; Schmidt, R. G.; Gordon, G. V. Macromolecules 2001, 34, 538-543. (11) Lee, C. L.; Polmanteer, K. E.; King, E. G. J. Polym. Sci. A-2 1970, 8, 1909-1916. (12) Cosgrove, T.; Roberts, C.; Choi, Y.; Schmidt, R. G.; Gordon, G. V.; Goodwin, A. J.; Kretschmer, A. Langmuir 2002, 18, 10075. (13) Daudt, W.; Tyler, L. U.S. Patent 2,676,182, 1954 (Dow Corning Corp.).
sample
M h w/kg mol-1
M h w/M hn
Rg/nm
F/g cm-3
η0/Pa s
R1 R3
0.5 14.1
1.15 2.99
0.35 2.2
0.88 1.17
0.0037 >1026 b
a M h w is the weight-average molecular weight, M h w/M h n is the molecular weight distribution, Rg is the radius of gyration, F is the density at 25 °C, and η0 is the zero-shear-rate viscosity at 25 °C. b R3 is a glass at 25 °C; the value was extrapolated from data of blends of R3 with PDMS.
Mixed-D columns, with a Miran 1A-CVF HPLC infrared detector, calibrated with silicate molecular weight fractions, and using chloroform as an eluent at 35 °C. The molecular size of the R1 silicate was calculated from molecular modeling [Cerius2 release 3.8, Molecular Simulations, San Diego, CA], whereas the other silicate was measured from hydrodynamic considerations using the universal calibration SEC technique. NMR Experimental Studies and Data Analysis. The NMR spin-spin relaxation results were obtained using a JEOL FX200 modified with a Surrey Medical Image Systems vector processor and a digital radio frequency console. All the results were obtained using the Carr-Purcell-Meiboom-Gill (CPMG)14,15 pulse sequence with 180° pulse spacings between 100 and 500 µs. A total of up to 8192 data points were collected by sampling the echo maxima. The resultant data were fitted to multiple exponential decays using a nonlinear least-squares analysis program and the DISCRETE16 algorithm. The “best” fit was obtained by allowing the program to fit the data to up to n distinct exponentials in the form n
y(t) )
∑y
0 i
exp(-t/T2i) + b
(2)
i)1
where y0i is the protonic proportion of component i with relaxation time T2i and b is the baseline, which should be zero except for any DC offsets and nonrandom noise. The fitted baseline has been subtracted from the data in all the figures shown. The choice of the best fit16 was based on a nonlinear correction to the standard deviation of the fit from the data for different integers of n. The 1H solid-state spectrum of the glassy R3 silicate was obtained from a solid-state Bruker MSL300 spectrometer using a single 90° pulse with a length of 2.7 µs.12 Differential Scanning Calorimetry. The thermal properties of the polymer/silicate blends were analyzed by DSC using a TA Instruments DSC 2920 with a specified sensitivity of 0.2 µW. The calorimeter was calibrated using indium as a reference. Each experiment was conducted at a heating rate of 10 K min-1 using 7-8 mg of sample sealed in an aluminum pan. The experimental uncertainty was found to be (1 K based on duplicate measurements.
Results and Discussion Pure Components. Figure 1 shows the experimental NMR T2 relaxation decays using the CPMG sequence obtained for the oligomeric silicate, R1, compared to those obtained for the 5K and 12K polymers. The relaxation behavior of the polysilicate R3, which is a glassy solid at room temperature, is also shown both as a CPMG decay signal and from a solid-state experiment using a single 90° pulse with a length of 2.7 µs. The values obtained by fitting the decays to single- or double-exponential decays and the calculated relaxation times are given in Table 3. The oligomeric silicate, R1, was calculated to have a radius of 0.35 nm and was a free-flowing liquid with a viscosity of 3.7 mPa s at 25 °C. The observed relaxation time (1.51 s) of R1 was consistent with both of these properties, and the decay obtained was a very good single (14) Carr, H. Y.; Purcell, E. M. Phys. Rev. 1954, 94, 630. (15) Meiboom, S.; Gill, D. Rev. Sci. Instrum. 1958, 29, 688-691. (16) Provencher, S. J. Chem. Phys. 1976, 64, 2772-2777.
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Figure 1. NMR spin-spin relaxation decays for the 5K and 12K linear PDMS (Table 1) and for silicates R1 and R3 (Table 2). For the R3 silicate material, results from both the CPMG (b) and FT solid-state (O) experiments are shown.
Figure 2. Relaxation decays for blends of the R1 silicate with the 12K PDMS (all normalized to 1): 100 (O), 52 (4), 27 (0), and 0 (3) vol % R1. Shown inset is the effect of R1 content on the viscosity of the 5K and 12K polymers. The data have been normalized to unit intensity.
Table 3. NMR Relaxation Data for Pure Components material 5K PDMS 12K PDMS R1 silicate R3 silicateb
T2,A/s 10-1
4.72 × 9.20 × 10-2 1.51 × 100 1.07 × 10-4
T2,B/s 10-1
y0A/y0B a
9.14 × 2.43 × 10-1
6.1 8.8
2.90 × 10-4
0.96
a
Ratio of intensities for a double-exponential fit when appropriate. b Fitted to a Gaussian for the short T2 and a Lorentzian for the longer T2.
exponential. This result was indicative of the sensitivity of the NMR relaxation times to both rotational and translational diffusion. A comparison of the silicate materials was quite revealing, showing the transition from molecular behavior (R1) to particulate behavior (R3) as perceived by NMR. The two pure polymer melts displayed T2 values intermediate between those observed for the silicates R1 and R3. The 12K polymer had a relatively narrow molecular distribution (M h w/M h n ) 1.03), but the T2 relaxation was still not well described by a single exponential. Instead, the decays from both PDMS polymers were fitted rather well by a double exponential or by a distribution of exponentials. The ratios of the components using the double-exponential fit are shown in Table 3. The more mobile component (longer T2) made only a small contribution to the intensity, and this is overemphasized by plotting the data logarithmically. There are also errors in obtaining the baseline due to poor signal-to-noise at very long times which can also adversely affect the fit and hence the visualization. The nonexponential behavior was due to a combination of finite dispersity of chain lengths and nonaveraged dipolar interactions. The extra free volume of the chain ends could lead to enhanced mobility, but it is unlikely that this effect would not be averaged out in the melt. These nonaveraged interactions have been discussed in detail by CohenAddad17,18 and become progressively more important above the chain entanglement molecular weight though these two samples are well below this value. In a recent paper where we measured the attenuation function in a pulsedfield gradient experiment10 which is sensitive to polydispersity, there was no clear evidence of nonlinear behavior. Although this issue is not central to the current paper, it does warrant further investigation. Blends. Figure 2 shows the relaxation decays found by mixing 12K PDMS with R1. Essentially, the relaxation (17) Cohen-Addad, J. P.; Viallat, A. Polymer 1986, 27, 1855-1863. (18) Cohen-Addad, J. P.; Domard, M.; Boileau, S. J. Phys. Chem. 1981, 75, 4107-4114.
Figure 3. T2 relaxation decays for blends of 12K PDMS and polysilicate R3 with a pulse interval of 1 ms. The decays are shown in decreasing R3 concentration in going from left to right and correspond to 66, 55, 45, 31, 22, 8, and 0 vol % R3. The data have been normalized to a relative intensity of 1 for the pure polymer.
decays become longer as R1 is simply acting as a solvation agent. This is consistent with the viscosity data shown inset for the blends based on the 12K and the 5K PDMS. As both the polymer and R1 decays are completely visible in the CPMG experiment, the data have all been normalized to unit intensity. In sharp contrast, Figure 3 shows the relaxation decays found for a series of blends with the R3 polysilicate. A pronounced effect of the R3 polysilicate on reducing the polymer mobility can be seen immediately from the raw data. This clearly showed that the polysilicate functioned as a reinforcing filler despite the fact that it has a radius of only ∼3 nm. A second very clear observation seen in Figure 3, where the data have been normalized to a relative intensity of 1 for the pure polymer, is the drop in signal intensity on complexation with R3, and this is discussed more fully below. For these systems, the relaxation times were essentially determined by a combination of segment mobility and diffusion. However, beyond a certain filler concentration, spatial constraints become a factor and the number of polymer chains that are free to diffuse becomes very small as a result of which only segmental motion is effective as a relaxation path. As discussed previously, the pure 12K polymer required at least a double exponential to adequately describe the relaxation decays, although the faster component only comprised a maximum of ∼9% of
Relaxation Studies of Low Molecular Weight PDMS
Figure 4. The variation in T2 (O) and CPMG intensity (9) using a single exponential fit as a function of the volume fraction of R3 in 12K PDMS. The solid line represents the theoretical intensity if only a polymer signal is measured.
the intensity. It might therefore be expected that at least two components would be evident in the blended samples. Although the pure R3 component exhibited solid-state NMR behavior, it was possible that it could have been solvated by the matrix polymer and thus appeared as a part of the CPMG decays. However, a transitional environment exists in a polymer melt, and the proton signal from the R3 component in the polymer melt may become progressively less visible as the R3 concentration increases and both rotational and translational diffusion are suppressed. This complicating factor must be considered in the analysis. To investigate this possibility, Figure 4 shows the CPMG T2 and absolute intensity results of fitting the raw data to a single-exponential decay function. The T2 values decreased rapidly with increasing R3 concentration and reflected the overall increase in viscosity of the blend due to the reinforcing effect of the filler (not shown). The solid line represents the intensity that would be expected if all of the polymer signal were visible, and clearly the measured intensity shows that virtually no signal was apparent from R3 under these experimental conditions. Since the first data point was recorded at 1 ms, it was possible that some intensity had already been lost due to fast relaxation processes by this point, especially at the higher R3 concentrations. However, by substituting the T2 values and their relative proportions obtained from the DISCRETE fitting process into eq 2, it was possible to extrapolate back to time zero. Figure 5 shows the extrapolated intensity at time zero (0) and the calculated polymer intensity assuming that all the PDMS protons were visible (straight line). The results clearly showed that no signal was apparent from R3 and that the entire polymer signal was visible at low R3 inclusions. However, as the R3 concentration increased less polymer signal was observed than could be predicted. Using the equation
Γ ) (WadsPDMS/WR3)RF/3 where R is the particle radius, F is the particle density, WR3 is the weight of R3, and WadsPDMS is the weight of adsorbed PDMS, it was possible to estimate the adsorbed amount, Γ (mg m-2), from where the normalized PDMS signal would hit the baseline (dashed line in Figure 5). This procedure resulted in Γ ) 0.34 ( 0.04 mg m-2, which can be compared with that obtained from diffusion measurements10 (0.59 ( 0.02 mg m-2). The difference between these two values may be attributable to probing slightly different regions of the polymer layer (trains vs loops and tails).
Langmuir, Vol. 18, No. 26, 2002 10083
Figure 5. Variation in CPMG intensity as a function of the volume fraction of R3 silicate in 12K PDMS. The solid line represents the theoretical intensity if only a polymer signal is measured. Measured CPMG intensity from extrapolated singleexponential fits (9), the difference between the theoretical and measured intensities (b), and the calculated signal from R3 using the relaxation data given in Table 3 (O); the dashed line shows the extrapolated signal.
Figure 6. Multiple relaxation decay fits to a 31/69 v/v blend of R3 silicate filler and 12K PDMS. The solid line through the experimental data is the quadruple-exponential fit. The inset shows the residual plots for 1 (O), 2 (4), 3 (0), and 4 (b) exponentials.
To further interpret the data, the relaxation decays were fitted to a sum of exponentials using the DISCRETE algorithm as a first approximation. Figure 6 shows a typical fit of eq 2 to a blend containing 69 vol % 12K PDMS. The inset shows a plot of the residuals: a singleexponential fit (O) was clearly inadequate in describing the measured relaxation decay, whereas the quality of the fit improved up to four exponential terms as would be expected. However, the populations of the shorter T2 components were quite small and it is not our intent to overinterpret the data. The algorithm chose the best fit as described above, Figure 7 shows the respective T2 values based on the optimum fitting, and Figure 8 shows the intensity of these contributions as a function of R3 concentration. All the relaxation times essentially decreased monotonically with increasing R3 content. For the lowest R3 concentration, the data showed only two apparent relaxation times, but the ratio of these times was quite different from that of the pure polymer. This was an indication of a change in the dynamic environments for some of the polymer segments. Above 20 vol % R3, four exponentials were evident. However, based on the data from the pure polymer, we combined the contributions from the two longest decays as being representative of the free polymer.
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Figure 7. The variation in T2 components for blends of 12K PDMS with the R3 silicate.
Cosgrove et al.
Figure 9. T2 relaxation decays for blends of 5K PDMS and polysilicate R3 with a pulse interval of 1 ms. The decays are shown in increasing R3 concentration from 0 to 50 wt % in going from right to left at intervals of 10 wt %. The data were normalized to a relative intensity of 1 for the pure polymer. Available data at 5 and 15 wt % R3 in the 5K PDMS were excluded for clarity.
Figure 8. The variation in the proportions of the T2 components for blends of the 12K PDMS with the R3 silicate: free polymer (O), perturbed polymer (9), and bound polymer (0).
The first and rather striking observation in Figure 7 was the one component (0) with a T2 that is approximately constant across the composition range with a value of 1 ( 0.5 ms. Figure 8 shows that the increase in the intensity of this component was monotonic and was consistent with a bound polymer layer. The intermediate relaxation time (9) corresponded to polymer segments that were perturbed by the proximity of R3 molecules and which could be loops or tails that retain some local mobility but lose translational mobility. The amount of unperturbed polymer segments dropped dramatically as R3 content increased beyond 45 vol %. The series of experiments were repeated for the 5K polymer with the R3 silicate. Figure 9 shows a selected set of CPMG decays at R3 concentration intervals of 10 wt %. An initial analysis of the results showed that the R3/5K PDMS system was very similar to that observed with the 12K PDMS: the polymer was reinforced by the inclusion of the polysilicate filler. Upon fitting the data to a single exponential (not shown), the trend of decreasing T2 with increasing R3 concentration was the same. However, the lower molecular weight polymer exhibited higher T2 values, which indicated a greater mobility that could be expected for a smaller polymer and manifested by the measured viscosity (Table 1). As the concentration of R3 increased, the difference between the T2 values decreased indicating that the particles were having a significant effect upon the segmental mobility of the polymer. The series of R3/5K PDMS data were also fitted using the DISCRETE algorithm and deconvoluted using the same procedure as described above for the 12K polymer. Figure 10 shows the individual T2 components as a
Figure 10. The variation in T2 components for blends of the 5K PDMS with the R3 silicate.
function of R3 content. The results obtained from fitting the data to a multiexponential decay were remarkably similar to the 12K polymer results up a volume fraction of 0.5. Above this value, the fraction of free polymer decreased more rapidly for the 12K polymer and one component assigned to free polymer could no longer be identified. In comparing the relative populations of polymer in different environments, the blends with 5K PDMS appeared to contain less unbound polymer and more trains (strongly perturbed polymer segments) than the 12K samples; however, it should be emphasized that this was only a small effect. Also, four exponentials appeared to be necessary at a lower R3 concentration (10 vol %), where the contributions from the two longest decays were representative of the free polymer. An interesting comparison can be made with DSC data. Figure 11 shows the DSC results in the range of 120-320 K as a function of the volume fraction of R3 silicate in the 5K polymer. For the pure polymer, the four thermal transitions typically associated with linear PDMS were detected: the glass transition, an exothermic crystalline formation (cold crystallization), and the endothermic melting of two crystalline forms. The glass temperature, Tg, was 145.5 K at the half-height of the endothermic shift. The extrapolated melting temperature Tm was 220.1 K, whereas the heat of fusion ∆Hf was 2.81 kJ mol-1 of the dimethylsiloxane repeat unit. These results were in reasonable agreement with those found in the advanced thermal analysis system databank.19
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Figure 12. The effect of the R3 silicate filler on the thermal properties of the 5K PDMS: Tg (O), ∆cP(Tg) (0), Tm (2), and ∆Hf (1).
suppressed as quantified by ∆Hf. These results provided corroborating evidence that there were no “free” PDMS chains remaining above a critical concentration of the R3 silicate and that adsorption has changed both the chain dynamics and topology significantly. Conclusions
Figure 11. DSC results obtained at a heating rate of 10 K min-1 as a function of the volume fraction of R3 silicate in the 5K PDMS.
Figure 12 summarizes the consequences of incorporating the R3 silicate as a filler on the thermal properties of the polymer matrix. First, the glass transition, which defines the onset of cooperative segmental motion including those from neighboring molecules, increased proportionally with the volume fraction of R3. Second, the glass transition is accompanied by an incremental change in heat capacity ∆cP(Tg). The contributions to this discontinuity in heat capacity include conformational factors, which can be related to the translational and rotational degrees of freedom of the molecule. There was an apparent increase in ∆cP(Tg) at the lowest loading of R3 (0.04 volume fraction) followed by a decreasing trend with R3 content. Third, as with Tg, the extrapolated melting temperature increased, but at a faster rate, up to a R3 volume fraction of 0.17, above which the crystallization of PDMS was completely (19) http://funnelweb.utcc.utk.edu/∼athas/databank/other/other/ siloxane/pdms/pdms.html.
The NMR and calorimetric data presented showed the changes in mobility that occurred in PDMS melts filled with silicate-type reinforcing fillers. A multiple-exponential spin-spin relaxation behavior was observed, and three different regimes of polymer segmental mobility can be identified: (1) free polymer, which required two exponentials to adequately describe its relaxation; (2) a perturbed region ascribed to loops and tails; and (3) a highly constrained population of trains. As the polysilicate concentration increased, the bound population increased as free polymer was lost. The polysilicate had a more pronounced effect on the higher molecular weight polymer. The DSC results also showed a loss of free polymer and significant changes in chain dynamics and topology with increasing polysilicate concentration. Acknowledgment. The authors acknowledge Dow Corning Limited, Barry, for supporting studentships and in particular Mr. A. J. Goodwin and Mr. J. P. Hannington [C.R., T.G.], EPSRC [C.R.], and the University of Bristol [T.G.] for support in carrying out this work. T. L. Sanders, Jr., is thanked for collecting the DSC data. LA025884+
