Silicate Dispersion and Mechanical Reinforcement in Polysiloxane

Chem. Mater. 2010, 22, 167–174 167 DOI:10.1021/cm9026978

Silicate Dispersion and Mechanical Reinforcement in Polysiloxane/Layered Silicate Nanocomposites Daniel F. Schmidt*,† and Emmanuel P. Giannelis Department of Materials Science and Engineering, Cornell University, Ithaca, New York 14853. Present address: Department of Plastics Engineering, University of Massachusetts-Lowell, One University Avenue, Lowell, MA 01854.

†

Received August 31, 2009. Revised Manuscript Received October 31, 2009

We report the first in-depth comparison of the mechanical properties and equilibrium solvent uptake of a range of polysiloxane nanocomposites based on treated and untreated montmorillonite and fumed silica nanofillers. We demonstrate the ability of equilibrium solvent uptake data (and, thus, overall physical and chemical cross-link density) to serve as a proxy for modulus (combining rubber elasticity and Flory-Rehner theory), hardness (via the theory of Boussinesq), and elongation at break, despite the nonideal nature of these networks. In contrast, we find that tensile and tear strength are not well-correlated with solvent uptake. Interfacial strength seems to dominate equilibrium solvent uptake and the mechanical properties it predicts. In the montmorillonite systems in particular, this results in the surprising consequence that equilibrium solvent uptake and mechanical properties are independent of dispersion state. We conclude that edge interactions play a more significant role than degree of exfoliation, a result unique in the field of polymer nanocomposites. This demonstrates that even a combination of polymer/nanofiller compatibility and thermodynamically stable nanofiller dispersion levels may not give rise to reinforcement. These findings provide an important caveat when attempting to connect structure and properties in polymer nanocomposites, and useful guidance in the design of optimized polymer/layered silicate nanocomposites in particular. Introduction Recent developments in nanoscience and nanotechnology have spawned a wide range of novel and interesting materials, but the realization of practical utility has come relatively slowly in comparison. With that said, there are success stories, and polymer nanocomposites may be one of them, with numerous companies now producing such materials on an industrial scale. The interest in these materials stems from the ability of polymer nanocomposites (when thermodynamically compatible and effectively processed) to give rise to improvements in a wide range of materials properties (mechanical, thermal, barrier, fire, etc.) with the addition of only a small amount of nanofiller. A prerequisite is nanofiller dispersion; however, this statement alone is misleading, because it has been clearly shown that dispersion in the absence of thermodynamic compatibility does not give rise to property enhancements.1 In compatible systems where dispersion is realized, on the other hand, there are several examples in the literature of systems whose overall ability to absorb mechanical energy, or toughness, has actually increased with the addition of layered silicate nanofillers, *Author to whom correspondence should be addressed. E-mail: [email protected].

(1) Alexandre, M.; Dubois, P.; Sun, T.; Garces, J. M.; Jerome, R. Polymer 2002, 43, 2132. (2) Zilg, C.; M€ ulhaupt, R.; Finter, J. Macromol. Chem. Phys. 1999, 200, 661. r 2009 American Chemical Society

both in impact2-4 and tension.5-8 Such behavior, which is atypical of most filled polymers, indicates that, under the right circumstances (i.e., when homogeneity of dispersion and distribution are high and large agglomerates are rare or absent), dispersed nanoparticles are able to enhance the ability of the polymer matrix to dissipate energy. Such enhancements must originate from one of two factors: changes in polymer microstructure (alterations in crystalline phase, crystallite size and morphology, degree of crystallinity, confinement effects, cross-link/entanglement density, phase behavior in multiphase systems, etc.) or interfacial stress transfer and modulations thereof (due to irreversible bond formation, reversible intermolecular interactions, interfacial slip, etc., with the importance of nanofiller size, shape, and properties implicit here). Beyond these concerns, there is also the matter of the matrix polymer chosen for this work, i.e., poly(dimethylsiloxane) (or PDMS). As a polysiloxane, or silicone, PDMS belongs to a family of unique polymers whose backbones consist entirely of silicon-oxygen bonds and (3) Reichert, P.; Nitz, H.; Klinke, S.; Brandsch, R.; Thomann, R.; M€ ulhaupt, R. Macromol. Mater. Eng. 2000, 275, 8. (4) Lu, H.; Liang, G.-Z.; Ma, X.; Zhang, B.; Chen, X. Polym. Int. 2004, 53, 1545. (5) Wang, Z.; Pinnavaia, T. Chem. Mater. 1998, 10, 3769. (6) Shah, D.; Maiti, P.; Gunn, E.; Schmidt, D. F.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. Adv. Mater. 2004, 16, 1173. (7) Liu, W.; Hoa, S. V.; Pugh, M. Polym. Eng. Sci. 2004, 44, 1178. (8) Shah, D.; Maiti, P.; Gunn, E.; Schmidt, D. F.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. Adv. Mater. 2004, 41, 3264.

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are therefore highly physically, chemically, and thermally stable. The high ionic character of these bonds allows for very high levels of molecular flexibility, and, as a result, PDMS has an exceptionally low glass transition temperature of 146 K.9 As liquids at room temperature, PDMS and several other linear silicones have found application as lubricants and heat-transfer fluids. Similarly, they are routinely cross-linked to produce silicone elastomers with excellent physical, chemical, and thermal stability; good low-temperature flexibility; clarity; and low surface energy. These find use in applications from sealants and caulking compound to nontoxic antifouling coatings for ships to microfluidics cells and medical implants for reconstructive surgery. However, the low surface energy of silicones such as PDMS implies a lack of intermolecular interactions, and this translates into one of the major weaknesses of silicone elastomers: their mechanical stability is so poor that they must be reinforced to be useful. Fumed silica, which consists of agglomerated aggregates of amorphous silica nanoparticles on the order of nanometers to tens of nanometers in size, has been one of the most favored reinforcing (nano)fillers for silicones, studied and used for decades, thanks to the strong hydrogen-bond mediated interactions between hydroxyl groups on the fumed silica surface and siloxane bonds in the PDMS backbone.10 In contrast, while thousands of papers have been published on the subject of layered silicate nanocomposites, only a handful of detailed reports have focused on polysiloxanes,11-20 despite the desirable properties of these materials, their need for reinforcement, and their compatibility with silica-based nanofillers. Previously, we reported that polysiloxanes are fundamentally incapable of dispersing typical alkylammoniummodified layered silicates unless polar pendant and/or end-groups are present in the polysiloxane chain.21 In particular, poly(dimethylsiloxane) (PDMS) was demonstrated to display end-group mediated intercalation in combination with dimethyl dioctadecylammonium montmorillonite (2C18MMT), with nonintercalated, weak intercalated, and apparently blank X-ray diffraction patterns realized with 0.37, 1.3, and 1.5 mmol Si-OH groups per gram of layered silicate, respectively. Here, we (9) Grulke, E. A. In Polymer Handbook, 4th edition; Brandrup, J., Immergut, E. H., Grulke, E. A., Abe, A., Bloch, D. R., Eds.; John Wiley & Sons: New York, 2005; Chapter VII. (10) Cohen Addad, J.-P. Surf. Sci. Ser. 2000, 90, 621. (11) Burnside, S. D.; Giannelis, E. P. Chem. Mater. 1995, 7, 1597. (12) Wang, S.; Li, Q.; Qi, Z. Key Eng. Mater. 1998, 137, 87. (13) Wang, S.; Long, C.; Wang, X.; Li, Q.; Qi, Z. J. Appl. Polym. Sci. 1998, 69, 1557. (14) Takeuchi, H.; Cohen, C. Macromolecules 1999, 32, 6792. (15) Anastasiadis, S. H.; Karatasos, K.; Vlachos, G.; Manias, E.; Giannelis, E. P. Phys. Rev. Lett. 2000, 84, 915. (16) Burnside, S. D.; Giannelis, E. P. J. Polym. Sci. B: Polym. Phys. 2000, 38, 1595. (17) LeBaron, P. C.; Pinnavaia, T. J. Chem. Mater. 2001, 13, 3760. (18) Wang, J.; Chen, Y.; Jin, Q. Macromol. Chem. Phys. 2005, 206, 2512. (19) Kaneko, M. L. Q. A.; Yoshida, I. V. P. J. Appl. Polym. Sci. 2008, 108, 2587. (20) Simon, M. W.; Stafford, K. T.; Ou, D. L. J. Inorg. Organomet. Polym. 2008, 18, 364. (21) Schmidt, D. F.; Clement, F.; Giannelis, E. P. Adv. Funct. Mater. 2006, 16, 417.

Schmidt and Giannelis

present the results of a systematic study of the mechanical properties, equilibrium swelling, and sol fraction measurements of cross-linked silanol-terminated PDMS networks as a function of nanofiller type and content, and base resin formulation. The results reported connect the origins of silicate dispersion with the resultant mechanical properties in the PDMS/layered silicate nanocomposite family, and represent, to the best of our knowledge, the first systematic study on the subject. Experimental Section Nanofillers Used. Unmodified montmorillonite (NaþMMT, ∼0.9 mequiv/g) and Cloisite 20A (2C18MMT) were provided by Southern Clay Products. CabOSil L-90 (untreated, with a Brunauer-Emmett-Teller (BET) specific surface area (SSA) of 100 m2/g), LM-130 (untreated, BET SSA = 130 m2/g), M-5 (untreated, BET SSA=200 m2/g), HS-5 (untreated, BET SSA=325 m2/g), and TS-530 (hexamethyldisilazane-treated, BET SSA = 225 m2/g) fumed silicas were provided by Cabot Corporation. Preparation of Nanocomposites Based on Silanol-Terminated PDMS. Silanol-terminated PDMS resins (PS340, PS343.5, PS345.5, and PS347.5) were obtained from United Chemical Technologies and used as received. The following molecular weight data was provided by United Chemical Technologies, based on gel permeation chromatography (GPC), using a PDMS standard: for PS340, Mn ≈ 2.6K, Mw/Mn ≈ 1.0; for PS343.5, Mn ≈ 49K, Mw/Mn ≈ 1.4; for PS345.5, Mn ≈ 74K, Mw/Mn ≈ 1.6; for PS347.5, Mn ≈ 111K, Mw/Mn ≈ 1.6. 60 g samples, consisting of 6 g of nanofiller, 9 g of PS340 (Mn ≈ 2.6K), and 45 g of a high molecular weight silanolterminated PDMS (PS343.5, PS345.5, PS347.5) were combined using an orbital mixer (FlackTek DAC-150FV Speed Mixer, 3000 rpm, 1 min, with multiple mixing cycles and stepwise mixing necessary with fumed silica nanofillers, because of their low bulk density, coupled with limited mixing cup volume). The cross-linker, poly(diethoxysiloxane) (PSI-023, Gelest; aka Ethyl Silicate 50 or ES-50, DPn ≈ 3-6, assuming a linear chain), was then added, using a reactivity ratio of 11:1 (11 Si-O-CH2CH3 for every Si-OH). The samples were mixed once more using an orbital mixer (3000 rpm, 30 s), after which the catalyst, tin(II) 2-ethylhexanoate (Aldrich), was added, using a 5:1 ratio of catalyst to crosslinker by volume, and the samples mixed a final time using an orbital mixer (3000 rpm, 30 s). Samples were then spread on a smooth polyethylene film between two steel draw-down bars (∼1 cm ∼30 cm ∼2 mm), using a steel trowel, and cured for several days at room temperature. Analysis. Mechanical properties were measured at a Dow Corning facility in Midland, MI. Stress-strain curves were obtained via tensile testing (at a crosshead speed of 500 mm/min), from which values for the stress at break (i.e., tensile strength), elongation at break, toughness, and the various 50% and 100% moduli (where applicable) were derived. Three to five stress-strain

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Table 1. Silanol-Terminated Polydimethylsiloxane Formulations Studied Nanofiller Content sample ID 49K-Na-25, -50, -100 49K-2C18-25, -50, -100 49K-L-100 49K-LM-100 49K-M-25, -50, -100 49K-HS-100 49K-TS-25, -50, -100 74K-Na-100 74K-2C18-100 74K-L-100 74K-LM-100 74K-M-100 74K-HS-100 111K-Na-100 111K-2C18-100 111K-L-100 111K-LM-100 111K-M-100 111K-HS-100 a

base resin Mn (5:1 w/w blends) ∼49K/∼2.6K ∼49K/∼2.6K ∼49K/∼2.6K ∼49K/∼2.6K ∼49K/∼2.6K ∼49K/∼2.6K ∼49K/∼2.6K ∼74K/∼2.6K ∼74K/∼2.6K ∼74K/∼2.6K ∼74K/∼2.6K ∼74K/∼2.6K ∼74K/∼2.6K ∼111K/∼2.6K ∼111K/∼2.6K ∼111K/∼2.6K ∼111K/∼2.6K ∼111K/∼2.6K ∼111K/∼2.6K

nanofiller þ

Na MMT 2C18MMT CabOSil L-90 CabOSil LM-130 CabOSil M-5 CabOSil HS-5 CabOSil TS-530 Naþ MMT 2C18MMT CabOSil L-90 CabOSil LM-130 CabOSil M-5 CabOSil HS-5 Naþ MMT 2C18MMT CabOSil L-90 CabOSil LM-130 CabOSil M-5 CabOSil HS-5

(wt %) 2.5, 5, 10 2.5, 5, 10 10 10 2.5, 5, 10 10 2.5, 5, 10 10 10 10 10 10 10 10 10 10 10 10 10

(vol % inorganic)

SiOH contenta (mmol/g nanofiller)

0.8, 1.7, 3.5 0.6, 1.1, 2.3 4.8 4.8 1.1, 2.3, 4.7 4.7 1.1, 2.3, 4.7 3.5 2.3 4.8 4.8 4.7 4.7 3.5 2.3 4.8 4.8 4.7 4.7

6.3, 3.1, 1.5 6.3, 3.1, 1.5 1.5 1.5 6.3, 3.1, 1.5 1.5 6.3, 3.1, 1.5 1.4 1.4 1.4 1.4 1.4 1.4 1.3 1.3 1.3 1.3 1.3 1.3

Initial SiOH content from PDMS end-groups prior to cross-linking.

curves were collected per material. Durometer hardness (Shore A) values were obtained using a standard durometer hardness tester, with three hardness values collected per system. Finally, tear-strength (tear B) tests were performed (at a crosshead speed of 500 mm/min), again with three measurements per system. For each equilibrium solvent uptake experiment, a specimen with a mass of ∼0.1 g (( 40%) was cut from a larger sheet of material with a razor blade, weighed, then placed in a clean 20-mL glass vial. The vial was then filled with toluene and sealed, and the sample swollen to equilibrium. In practice, no more than three days were required to reach equilibrium swelling.22 Sample mass in the equilibrium swollen state was measured by removing the sample from the toluene, gently patting it with a paper towel to remove liquid solvent present on the sample surface, and weighing it; these actions were performed as rapidly as possible to prevent appreciable solvent loss. Finally, following equilibrium solvent uptake, the sample was air-dried under a fume hood for a minimum of 24 h and weighed once more. The polymer volume fraction in the swollen network is calculated using the relation j=(Vdry - Vfiller)/(Vswollen Vfiller), where Vdry is the volume of the dry network following the equilibrium solvent uptake experiment, Vfiller the volume of the nanofiller (considering the inorganic component only), and Vswollen the volume of the swollen sample at equilibrium. For these calculations, the densities used were 0.8669 g/mL for toluene, 1 g/cm3 for cross-linked PDMS, 2.86 g/cm3 for the inorganic component of all montmorillonite nanofillers (based on the density of Cloisite Naþ, according to Southern Clay Products23), and 2.2 g/cm3 for all fumed silica nanofillers (22) Burnside, S. D. Synthesis and Characterization of Polymer Matrix Nanocomposites and Their Components (Silicates, Solvent Uptake, Thermal Stability), Ph.D. dissertation, Department of Materials Science and Engineering, Cornell University, Ithaca, NY, 1997. (23) See http://www.scprod.com/product_bulletins/PB%20Cloisite% 20NAþ.pdf, viewed March 2009.

(this is the value reported in Cabot Corporation technical datasheets for all of the CabOSil grades studied). Results Based on our previous work21 showing its ability to disperse in these systems, 2C18MMT was chosen as the basis for the bulk of these studies. This was compared with NaþMMT as well as four unmodified and one hexamethyldisilazane-treated fumed silica of varying particle sizes. High molecular weight PDMS base resins were used to ensure highly elastomeric properties following cross-linking. A small amount of very low molecular weight silanol-terminated PDMS was added in all cases to ensure consistent levels of clay dispersion and to overcome clay-induced cure inhibition. The formulations studied are listed in Table 1. Analogous unfilled PDMS samples were also prepared, but, because the unfilled resins do not suffer from cure inhibition, these materials were observed to be translucent and highly overcrosslinked, precluding direct comparisons between unfilled PDMS and PDMS nanocomposites. Because unfilled silicones are exceptionally weak and difficult to test accurately, emphasis is placed on comparisons between different nanofiller types and concentrations and different resin molecular weights. Solvent uptake data are of critical importance, because solvent uptake reflects polymer/nanofiller interactions that give rise to physical crosslinks and mechanical reinforcement. Three series of comparisons are made. In the first, nanofiller content and type are fixed and PDMS molecular weight is varied. The commercial PDMS resins used here have polydispersity indices in the range of ∼1.1-1.5, making it possible to discuss molecular weight effects with some confidence. The second series compares variations in nanofiller type, with a fixed filler content. The third series examines the effects of nanofiller content in selected systems. Consistent with our previous work, weak intercalated peaks were observed for all PDMS

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nanocomposites described here, based on 2C18MMT. Silanol content is dominated by the low molecular weight PDMS additive and has been chosen to produce partial exfoliation, based on our prior work.21 As such, similar dispersion levels are expected for all nanocomposites with the same 2C18MMT content, regardless of the molecular weight of the high molecular weight component. In contrast, the NaþMMT based controls are known to show no nanoscale dispersion, given the total lack of compatibility between PDMS and unmodified clay.

these values is performed using the power series recommended by the same source,9 the following expression for χ, as a function of concentration, results (r2 =0.99522): χ ¼ χ0 þ χ1 j þ χ2 j 2

ð4Þ

where χ0 ¼ 0:450 ( 0:011 χ1 ¼ 0:272 ( 0:052

Discussion Solvent uptake and mechanical properties are often correlated using rubber elasticity and the classic FloryRehner approach,24,25 assuming an ideal, filler-free network and Gaussian chain statistics. In contrast, our materials consist of bimodal network with nonideal crosslinks and containing various types and concentrations of nanofillers, making it important to confirm the validity of this approach. The elastic modulus (E) may be defined according to the theory of rubber elasticity as follows: E ðMPaÞ ¼

3FRT Mc

ð1Þ

where F is the density of the polymer network (in units of g/cm3), R the gas constant (R=8.314472 J mol-1 K-1),26 T the absolute temperature (in Kelvin), and Mc the (average) molecular weight between crosslinks (in units of g/mol). The Flory-Rehner approach provides the following relationship between polymer volume fraction (j) and molecular weight between crosslinks (Mc, in units of g/mol) in a swollen polymer network: Fν 1=3 j lnð1 -jÞ þ j þ χj ¼ j Mc 2 2

ð2Þ

χ2 ¼ 0:094 ( 0:050 Substituting eq 4 into eq 3 gives us the final form of the expression that we will use to calculate the relationship between elastic modulus and solvent uptake in the materials studied here: " # 3RT lnð1 -jÞ þ j þ ðχ0 þ χ1 j þ χ2 j2 Þj2 E ¼ ð5Þ ν j1=3 -ðj=2Þ When j is in the range of 0.3-0.5, as we measure experimentally, the values of E vs j predicted by eq 5 may be fit by a more convenient power law of the form E=a þ bjc. Therefore, we have chosen to fit our experimental data for tensile modulus in this fashion. An overlay of the predictions based on eq 5 and the experimental data appears in Figure 1. While there is a discrepancy between the model predictions and the experimental data, what is striking is just how close they are, given the assumptions being made. In fact, the model data may be made to coincide almost exactly with the data if the value of χ0 in eq 4 is increased from 0.450 to 0.495. The value of χ0, in turn, is significantly influenced by molecular weight through its dependence

Here, χ is the Flory-Huggins polymer-solvent interaction parameter, F is again the density of the polymer network (expressed in units of g/cm3), and ν the molar volume of the swelling solvent (given in units of cm3/mol). An expression for E, as a function of j, may be derived by rearranging eq 2 and substituting it into eq 1: E ¼ -

3RT lnð1 -jÞ þ j þ χj2 ½ ν j1=3 -ðj=2Þ

ð3Þ

For our measurements (in toluene at room temperature), we set values of T = 293 K and ν = 106.8 cm3/mol,9 giving the initial constant term a value of 68.4. For χ, the Polymer Handbook provides χ at 20 °C for PDMS (50 kg/mol)/toluene combinations with j = 0.0, 0.2, 0.4, 0.6, 0.8, and 1.0.9 When a nonlinear regression of (24) Flory, P. J.; Rehner, J. J. Chem. Phys. 1943, 11, 521. (25) Flory, P. J. J. Chem. Phys. 1950, 18, 108. (26) Mohr, P. J.; Taylor, B. N.; Newell, D. B. Rev. Mod. Phys. 2008, 80, 633.

Figure 1. Model elastic and experimental 50% modulus versus polymer volume fraction data for all networks tested. Predictions come from rubber elasticity and the Flory-Rehner swelling model. Experimental data comes from tensile testing and equilibrium solvent uptake in toluene.

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Figure 2. Experimental 100% modulus versus polymer volume fraction data for all networks tested. Experimental data comes from tensile testing and equilibrium solvent uptake in toluene.

on the second virial coefficient,9 making this shift plausible. This result confirms that the behavior of our networks is largely captured by the theoretical considerations expressed in eq 5. Having addressed correlations between solvent uptake and 50% modulus, we examine whether the former tracks with other mechanical properties as well. A similar treatment of 100% modulus data is shown in Figure 2. The proposed power law again describes the majority of the data and produces a value for the scaling exponent (c) statistically identical to those presented in Figure 1. This shows that equilibrium solvent uptake measurements provide a reasonable indication of 100% modulus data as well as 50% modulus data. While not an inherent materials property, because of issues of indenter geometry and contact mechanics, hardness may be correlated to elastic modulus as well. Recent work has demonstrated that the theory of Boussinesq has the potential to relate Shore A durometer hardness (H) and elastic modulus (E), according to the expression shown below.27 ! C1 þ C2 H 1 -ν2 E ¼ ð6Þ 100 -H 2rC3 Here, C1, C2, and C3 are constants, ν is the Poisson’s ratio of the elastomer, and r is the radius of the indenter (r = 0.395 mm for the Shore A test). Rearranging to solve for H as a function of E, we have H ¼

100E -C1 ½ð1 -ν2 Þ=ð2rC3 Þ E þ C2 ½ð1 -ν2 Þ=ð2rC3 Þ

ð7Þ

Experimental durometer hardness (Shore A) data versus the 50% and 100% moduli are shown in Figure 3, along with nonlinear regressions, following the functional form of eq 7. (27) Kunz, J.; Studer, M. Kunststoffe 2006, 96, 92.

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Figure 3. Experimental 50% and 100% moduli versus hardness for all networks tested. The experimental moduli come from tensile testing, while hardness comes from durometer hardness testing. The lines indicate nonlinear regressions with the functional form of the Boussinesq model.

Figure 4. Experimental and predicted hardness versus polymer volume fraction for all networks tested. Experimental data comes from durometer hardness testing and equilibrium solvent uptake in toluene. The model prediction is based on the Boussinesq model27 fit to the 50% modulus data shown in Figure 3, coupled with the expression derived for 50% modulus versus polymer volume fraction in Figure 1.

The correlations observed here imply that 50% modulus in particular (and the 100% modulus, to a lesser extent) may be used to predict hardness. It follows that hardness should correlate with equilibrium solvent uptake, as is observed in Figure 4. The next question is whether the “ultimate” properties of these materials (elongation at break, tensile strength, tear strength) track with equilibrium solvent uptake data. A relationship with elongation at break would seem plausible in that the same factors that limit swelling; physical and chemical crosslinks;should limit the maximum elongation an elastomeric polymer network should be able to accommodate before breaking. As Figure 5 shows, the expected correlation is indeed observed.

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Figure 5. Elongation at break versus polymer volume fraction for all networks tested. Experimental data comes from tensile testing and equilibrium solvent uptake in toluene.


Figure 6. Tensile and tear strength versus polymer volume fraction for all networks tested. Experimental data comes from tensile and tear strength testing and equilibrium solvent uptake in toluene.

A simple linear relationship between elongation at break and polymer volume fraction j is observed, implying that the elongation at break, like j, is governed mainly by the effective chemical and physical cross-link density, not the presence of large inhomogeneities (agglomerates, voids, etc.). Finally, plots of tensile and tear strength versus polymer volume fraction at equilibrium solvent uptake are presented in Figure 6. There is a tendency for samples with low polymer volume fractions at equilibrium solvent uptake to have low tensile strengths, but no obvious correlations with tear strength. This implies that tensile and especially tear strength are likely to be more sensitive to large inhomogeneities (agglomerates, voids, etc.) in these samples than effective chemical and physical cross-link density. To summarize the aforementioned results: (1) Polymer volume fraction at equilibrium solvent uptake predicts moduli and hardness, consistent with rubber elasticity, Flory-Rehner swelling, and Boussinesq contact mechanics. (2) Polymer volume fraction at equilibrium solvent uptake predicts elongation at break, implying that this property is determined mainly by effective chemical and physical cross-link density, not the presence of large inhomogeneities. (3) Polymer volume fraction at equilibrium solvent uptake predicts tensile strength very poorly and fails to predict tear strength, implying that these quantities (tear strength especially) are determined by factors not captured in these models. These conclusions inform comparisons of the materials produced. In particular, variations in polymer volume fraction j at equilibrium solvent uptake with changes in base resin molecular weight and nanofiller content capture trends in moduli, hardness, and elongation at break more succinctly than individual consideration of these properties. A plot of polymer volume fraction j at

Figure 7. Polymer volume fraction (j) versus base resin molecular weight and nanofiller content for all PDMS networks tested. Experimental data comes from equilibrium solvent uptake in toluene.

equilibrium solvent uptake versus base resin molecular weight and nanofiller content appears in Figure 7. Two major trends are observed. First, lower base resin molecular weights lead to higher values of j, as expected, given that lower base resin molecular weights should correlate with increased (chemical) cross-link density. Second, increased nanofiller content leads to higher values of j, as expected given that compatible nanofillers should increase the (physical) cross-link density. The effectiveness of nanofillers in increasing physical cross-link density is dependent on effective interfacial area (which is a function of specific surface area and dispersion state) and interfacial interaction strength;in this case, through interactions between silanols (Si-OH) and siloxane bonds (Si-O-Si). Fully exfoliated montmorillonite possesses a higher specific surface area than fumed silica. On the other hand, montmorillonite is

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Figure 8. (a) Tensile strength and (b) tear strength versus base resin molecular weight and nanofiller content for all networks tested. Experimental data come from tensile testing, tear strength testing (Tear B), and equilibrium solvent uptake in toluene.

completely free of hydroxyls, with the exception of the layer edges, which comprise only a small fraction of the total surface area.28 Thus, strong silanol-siloxane interactions can only occur between the (many) PDMS repeats and the (few) layer edges or between the (few) silanol endgroups of the PDMS and the layer faces. In contrast, untreated fumed silicas studied here (CabOSil L-90, LM130, M-5, and HS-5) possess (many) surface hydroxyls that interact with (many) PDMS repeats, giving a stronger interface versus fumed silica treated with hexamethyldisilazane (CabOSil TS-530) and possessing a lower surface hydroxyl concentration. Considering only the untreated fumed silicas, higher nanofiller content increases the value of j, but increasing the specific surface area (from CabOSil L-90 to LM-130 to M-5 to HS-5) has little effect. A comparison of the treated (CabOSil TS-530) and untreated (CabOSil M-5) fumed silicas of similar specific surface area reveals that the latter gives equal or higher values of j at all nanofiller concentrations, which is consistent with stronger polymer/nanofiller interactions. Wholly different results are observed in the case of the montmorillonite nanofillers. The lack of significant differences in j trends observed between the Naþ MMT and the 2C18MMT systems is surprising, in that it indicates that the dispersion of montmorillonite nanolayers has little or no effect. The only explanation for such a situation is that interactions with the layer faces do not contribute to the overall cross-link density. As similar (28) For a sheet with a thickness of 0.95 nm, a diameter/width of 100 nm, and a density of 2.86 g/cm3, the specific surface area is 750 m2/ g, which is consistent with that reported by Southern Clay Products (see http://www.nanoclay.com/keyproperties.asp, viewed March 2009). Of this total specific surface area, 736 m2/g (98.1%) comes from the layer faces and 14 m2/g (1.9%) comes from the layer edges.

reductions in the value of j are nevertheless observed as montmorillonite concentration increases, regardless of modification/dispersion state, the network and the montmorillonite layers must be interacting, and in a fashion that is not dependent on layer separation. Strong interactions between the silanol-terminated PDMS network and the silanol-functional montmorillonite layer edges represent a logical explanation for this result. Consistent with this explanation, and given the very low specific surface area of the layer edges, the values of j are consistently lower in the montmorillonite-based systems than in the fumed silica-based systems. Having addressed changes in j as a proxy for changes in stiffness, hardness, and extensibility, we now consider tensile and tear strength, which do not correlate strongly with j. Plots of tensile and tear strength versus base resin molecular weight and nanofiller content appear in Figures 8a and 8b, respectively. In Figure 8a, tensile strength seems to be strongly affected by base resin molecular weight, with higher base resin molecular weights producing lower tensile strengths in the case of the montmorillonite-based systems, regardless of clay treatment or dispersion level. This effect may be related to the presence of additional large-scale inhomogeneities as base resin molecular weight increases, or to the inherent properties of the networks based on highermolecular weight base resins. The tensile strengths of the fumed silica-based systems show no sensitivity to base resin molecular weight or nanofiller content, indicating that strong polymer/nanofiller interactions overwhelm molecular weight effects. This also implies the existence of similar defect (agglomerates, voids) concentrations in these materials, which is consistent with the compatibility between fumed silica and PDMS.

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In Figure 8b, on the other hand, tear strength has a tendency to increase with nanofiller content, independent of nanofiller type. In the untreated fumed silica filled systems, tear strength is observed to increase with base resin molecular weight as well. In contrast, the tear strengths observed in the montmorillonite systems are effectively independent of base resin molecular weight. This again points to a fundamental difference in reinforcement behavior. In particular, in both the tensile and tear strength data, it is clear that increasing base resin molecular weight favors higher performance in fumed silica filled systems than in montmorillonite-based systems, which is consistent with the argument for stronger interactions in the former. Similarly, it is observed in both cases that the dispersion state of the montmorillonite does not affect the strength of the nanocomposite, which is consistent with the argument that interactions with the layer edges are more important in these systems than in conventional layered silicate nanocomposites. To summarize this discussion, we conclude the following for the materials reported here: (1) Interfacial interactions seems to dominate the data on polymer volume fraction at equilibrium solvent uptake (which is proportional to modulus and hardness, inversely proportional to elongation at break), with the order of interfacial strength observed to be (untreated fumed silica) > (treated fumed silica) > (montmorillonite); this is consistent with the potential for strong hydrogen-bond-mediated silanol-siloxane interactions. (2) Data on polymer volume fraction at equilibrium solvent uptake for the montmorillonite systems excludes nanoscale dispersion state as a factor in determining mechanical properties and points to a unique situation where interactions at the layer edges are more important. (3) Only the montmorillonite systems show decreases in tensile and tear strength with increasing base resin molecular weight, independent of dispersion state, emphasizing the importance of interfacial interactions in the context of changes in end group concentration. Conclusions Poly(dimethylsiloxane) (PDMS) networks have been produced based on untreated and surface-treated montmorillonites and fumed silicas. Rubber elasticity, coupled with the Flory-Rehner description of network swelling, describes the behavior of these systems well. This has been successfully extended to durometer (Shore A) hardness using an approach based on the theory of Boussinesq,27 and correlations between elongation at break and equilibrium solvent uptake also have been demonstrated. By examining the dependence of polymer


volume fraction at equilibrium solvent uptake on resin molecular weight and nanofiller content, we observe that interfacial strength is the dominant factor in determining the 50% and 100% moduli, hardness, and elongation at break obtained in all systems. In contrast, tensile and tear strength show minimal correlation with equilibrium solvent uptake, and appear to be more dependent on other factors. Specific to the montmorillonite nanofillers, we have previously demonstrated that PDMS is unique among polymer/layered silicate nanocomposites due to the extreme sensitivity of dispersion levels in this family of materials to small concentrations (on the order of 1 mmol per gram of nanofiller) of polar substituents.21 We now show that this translates into unique behavior from the standpoint of mechanical reinforcement as well, both in comparison with fumed silica based PDMS nanocomposites and with the bulk of the polymer/layered silicate nanocomposite literature. While dispersion of montmorillonite layers in PDMS may be achieved given the right PDMS formulation and an appropriately modified montmorillonite, the level of reinforcement is unaffected versus a nondispersed montmorillonite control. Reinforcement in this system seems to be governed by interactions between siloxane bonds in the PDMS chains and silanols present at the edges of the montmorillonite layers, a conclusion supported by the sum of the equilibrium solvent uptake and mechanical properties data presented. Pinnavaia and co-workers seem to have been the first to explicitly note the potential importance of polymer/ layer edge interactions in polymer/layered silicate nanocomposites, as well as the need to further evaluate their role in affecting mechanical properties.29 Subsequent reports linking polymer/layer edge interactions to mechanical properties show that small changes in elastic properties may be realized in polyolefin/layered silicate nanocomposites through silane modifications of the layer edges.30,31 In contrast, in the nanocomposites described here, mechanical reinforcement seem to be dominated by interactions, not exfoliation. Acknowledgment. This work was supported by the DowCorning Corporation and the United States Office of Naval Research (ONR). Specific thanks go to Dr. Deborah Bergstrom and Dr. Timothy Chao for arranging for mechanical properties testing at Dow-Corning’s Midland, MI facility, and to Jason Fisk, Kermit Kwan, and Dr. Carl Fairbank for performing the testing. EPG acknowledges the support of Award No. KUS-C1-018-02, made by King Abdullah University of Science and Technology (KAUST). (29) Shi, H.; Lan, T.; Pinnavaia, T. J. Chem. Mater. 1996, 8, 1584. (30) Zhao, C.; Feng, M.; Gong, F.; Qin, H.; Yang, M. J. Appl. Polym. Sci. 2004, 93, 676. (31) Lee, J. W.; Kim, M. H.; Choi, W. M.; Park, O. O. J. Appl. Polym. Sci. 2006, 99, 1752.

Silicate Dispersion and Mechanical Reinforcement in Polysiloxane

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