Composites of Polypropylene with Layered Mg-Silsesquioxanes Show

May 9, 2008 - We report the synthesis of vinyl modified magnesium silsesqiuoxanes (“vinyl clay”), and the formation of their composites with isota...
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Ind. Eng. Chem. Res. 2008, 47, 3891–3899

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Composites of Polypropylene with Layered Mg-Silsesquioxanes Show an Unusual Combination of Properties Guruswamy Kumaraswamy,* Yogesh S. Deshmukh, Vikrant V. Agrawal, and Anuya A. Nisal National Chemical Laboratory, Pune, India

We report the synthesis of vinyl modified magnesium silsesqiuoxanes (“vinyl clay”), and the formation of their composites with isotactic polypropylene (iPP) by melt compounding. Vinyl clay is a layered compound with a layer thickness of approximately 1 nm. Vinyl clay does not exfoliate in iPP; rather, it disperses to form a network that exhibits a characteristic low frequency solid-like plateau in the elastic modulus in dynamic melt rheological measurements. Strangely, vinyl clay also plasticizes iPPsthere is a decrease in the high frequency complex viscosity. The decrease in the complex viscosity is higher at higher frequencies, suggesting the influence of slip at the iPP-vinyl clay interface. The combination of the low frequency elastic plateau and plasticization makes the vinyl clay composite significantly more shear thinning than the matrix iPP. In the solid state, vinyl clay-iPP composites exhibit increased tensile modulus (showing ≈50% increase for a 5% loading), but surprisingly, no corresponding decrease in the elongation at break. Thus, while microstructural characterization indicates that only a small fraction, if any, of the vinyl clay is exfoliated, the enhancement in mechanical properties is similar to that observed for iPP-exfoliated montmorillonite nanocomposites. Our compounding protocol is unable to effectively disperse the clay in the iPP at clay loadings greater than about 7.5%. Therefore, the low frequency plateau in the melt elastic modulus and the solid tensile modulus increase with clay loading until 7.5% but exhibit a nonmonotonic decrease at higher clay loadings. 1. Introduction The steady growth in recent years in the volume of synthetic polymers produced globally has come primarily by extending the performance envelope of existing polymers by blending with other polymers1 and/or additives and fillers.2 In this context, nanocomposites, viz. composites of polymers with fillers dispersed at a nanometer scale (nanofillers), are an exciting new frontier. Recent developments in two areas of silicon-incorporating polymer nanocomposites are noteworthyspolyhedral oligomeric silsesquioxane (POSS) polymers and copolymers (wherein the POSS units are ladder- or cage-like) and polymerclay nanocomposites (wherein nanometer-thick sheets of clay are dispersed in the polymer matrix) represent examples of hybrids that aim to advantageously combine the properties of their inorganic (silicon-based) and organic (polymer) components. POSS nanocomposites have been investigated, for example, as dental adhesives, for their fire-retardant properties,3 and for their use in outerspace applications.4 In the early 90s, Toyota rekindled interest in the area of polymer-clay nanocomposites by demonstrating that polymerizing nylon in the presence of surfactant-modified hydrophobic smectite clays led to exfoliation of the clay in the polymer matrix.5,6 The resulting nylon nanocomposite had nanometer-thick platelets of clay dispersed in the polymer, and showed unexpected enhancement in mechanical properties5,6sthus opening the possibility of using these nanocomposites in under-the-hood automotive applications. Shortly after the reports from Toyota, Giannelis’ group demonstrated that polymer-clay nanocomposites could be prepared by direct melt compounding.7,8 Since that time, there has been an explosive growth in the literature on polymer-clay nanocomposites. This literature has been comprehensively summarized in recent reviews.9,10 * To whom correspondence should be addressed. Mailing address: Polymer Science and Engineering, National Chemical Laboratory, Dr. Homi Bhabha Road, Pune-411008, India. Tel.: 91-20-25902182. Fax: 91-20-25902618. E-mail: [email protected].

As polymer-clay nanocomposites typically incorporate only a few weight percent of the inorganic component, the increase in the density of the nanocomposite over the matrix polymer is small and there is a tremendous increase in the modulus to weight ratio. Therefore, the use of polymer nanocomposites in automotive applications has been an active area of research11 due to the possibility of increased fuel efficiency resulting from a decrease in automobile weight. Recently, polypropylene-clay composites have been developed for several interior and exterior autoparts.12 The incorporation of clay into polypropylene13 enhances the strength of this high-volume, relatively inexpensive commodity plastic making it suitable for the cost sensitive yet highly demanding autocomponent industry. Compatibilizing polar clays to render them dispersable in apolar polymers is typically done by surfactant-modification of swellable smectite clays. However, degradation of the surfactant modifiers14 limits the temperature to which the nanocomposite can be used, as surfactant degradation can lead to deintercalation of the polymer and aggregation of the clay and can lead to discoloration of the polymer matrix.15 While the typical threelayer smectites that are used to prepare nanocomposites are not amenable to covalent modification of the basal clay surface, layered silicates such as magadiite can be covalently modified to prepare thermally stable hybrids.16,17 Recently, layered, claylike inorganic-organic metallo-silsesquioxanes have been synthesized that incorporate covalently bound organic groups on a claylike framework.18–22 Mann et al. have complexed proteins with amino modified Mg-silsesquioxanes and have shown that these protein composites stabilize the protein to higher temperatures relative to the native form.23,24 Carrado et al. have synthesized organically modified metallo-silsesquioxanes in the presence of water-soluble polymers such as polyvinyl pyrrolidone, polacrylonitrile, and polyaniline.25–27 The polymers are incorporated into the organically modified metallo-silsesquioxanes to a maximum of 86% by weight of the polymer. However, in these syntheses, the polymer has to be water-soluble and

10.1021/ie071658p CCC: $40.75  2008 American Chemical Society Published on Web 05/09/2008

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compatible with the reaction conditions, and it is difficult to increase the polymer loading (compared with typical polymer clay nanocomposites where property enhancements happen at a clay loading of less than 5%). The research described in this paper explores melt compounding of a vinyl functionalized layered magnesium silsesquioxane with isotactic polypropylene (iPP), to investigate the possibility of nanocomposite formation using these novel materials. We prepare the hybrid silsesquioxane under mild conditions and disperse the silsesquioxane in the polymer by melt compounding. Transmission electron microscopy (TEM) and X-ray diffraction (XRD) suggest that the silsesquioxane does not completely exfoliate in the iPPsthus, the composite is not a “nanocomposite” with exfoliated clay layers dispersed in the iPP. However, we observe that our composite shows mechanical behavior similar to those reported for montmorillonite-iPP nanocomposites. In particular, there is a significant increase in tensile modulus relative to the matrix iPP, and interestingly, there is no corresponding deterioration of the ductility. Further, our composite has an interesting balance of properties: the low shear viscosity increases due to the formation of a network of the silsesquioxane in the polymer, but there is enhanced shear thinning (the high shear viscosity is lower than that of the matrix polymer). Thus, vinyl clay-iPP composites offer the advantage of ease of preparation, enhanced processability due to the apparent increase in shear thinning, and a desirable balance of mechanical strength and ductility. 2. Experimental Details 2.1. Materials. Triethoxy vinyl silane [(CH3CH2O)3SiCHdCH2, 99%] was obtained from Sigma-Aldrich. Sodium hydroxide and magnesium chloride hexahydrate were obtained from Merck. Distilled, deionized water (resistivity ) 18.2 MΩ · cm) from a Millipore Milli-Q system was used for all experiments. Isotatic polypropylene (iPP) with a weight average molecular weight of 160 000 g · mol-1 and polydispersity index of 2.8 (MFI ) 2.7° per 10 min) was obtained from Fina. All chemicals were used as obtained, without further purification. 2.2. Synthesis and Preparation of Composite with iPP. We prepare the layered hybrid by co-condensation of vinylsubstituted trialkoxy silane with magnesium hydroxide under basic conditions as described in the literature.26–28 Triethoxy vinyl silane (2.565 g, 0.27 mol · dm-3) and magnesium chloride hexahydrate (2.03 g, 0.20 mol · dm-3) solutions were prepared in ethanol and then mixed together and stirred for 10 min. To this mixture, we added aqueous sodium hydroxide (0.50 mol · dm-3) dropwise until the pH became 11.5. A white precipitate was obtained that was then aged for a day at room temperature. After aging, the precipitate was repeatedly washed with distilled-deionized water until the filtrate had a neutral pH and was free of chloride ions (confirmed by the addition of silver nitrate to the fresh filtrate solution). Finally, the white precipitate was air-dried, followed by vacuum drying at 80 °C for 12 h, and then crushed in a mortar to get a fine white powder. We refer to this as vinyl clay (talc-like structure with vinyl moieties bonded to framework silicon atoms). Composites of the vinyl clay with iPP were compounded for 10 min at 200 °C at 50 rpm screw speed using a DSM microcompounder (corotating twin-screw microcompounder, 5 mL version). We prepared vinyl clay-iPP composites containing 1%, 2.5%, 5%, 7.5%, and 10% by weight of vinyl clay (weight percentages represent weight of the entire silsesquioxane and not just the rigid inorganic part). In this work, iPP/x refers to a composite of iPP loaded with x% of vinyl clay. A DSM

microinjection molder was used to prepare samples for tensile and Izod impact measurements. We used a nonstandard mold provided by DSM to prepare samples for tensile tests (dimensions: 1.28 mm × 4.6 mm × 52 mm gauge length). Izod impact tests were done according to ASTM D256. As an aside, we note that compounding the vinyl clay with iPP in the presence of a free radical initiator, dicumyl peroxide (DCP) does not lead to covalent coupling of the matrix with the claysprobably due to the low reactivity of the isolated vinyl group on the surface of the silsesquioxane. Rather, it leads to a decrease in the molecular weight of the iPP (data in the Supporting Information). 2.3. Characterization. X-ray diffraction (XRD) was performed in reflection mode on a Rigaku Dmax 2500 using a rotating anode source and Ni-filtered Cu KR radiation. Fourier transform infrared (FTIR) analysis was performed at a resolution of 4 cm-1 on a Perkin-Elmer 16C. Thermogravimetric analysis was carried out using a Perkin-Elmer TGA-7 under nitrogen atmosphere. The samples were initially held at 150 °C in the TGA for 5 min to dry them and were subsequently heated to 900 °C at 10 °C · min-1. CHN analysis was performed using Vario EL from Elemental Analyzer GmbH. Scanning electron microscopy (SEM) was performed using a Leica-S440 at an accelerating voltage of 20 kV. Samples for SEM were prepared by dispersing 50 mg of the sample in 5 mL of chloroform using a sonicating bath and by drying a drop of this dispersion on an SEM stub. Vinyl clay-iPP composites were cryoultramicrotomed for TEM using a Leica Ultracut-UCT equipped with a diamond knife. Samples with a thickness of ≈50 nm were cut and floated onto a grid for imaging on a JEOL1200-EX at an accelerating voltage of 80 kV. Solid properties of the vinyl clay-iPP composite were tested using the microinjection molded specimens on a Universal Testing machine (Instron, Model 4202). At each clay concentration, four samples were tested and the tensile properties were calculated as an average of these measurements. Izod impact strength was measured using an impact tester from Ceast on a notched sample (notch cut to a depth of 2.5 mm). Impact values were obtained as an average over measurements made on four samples. Melt rheology of the composites was investigated using a controlled strain rheometer (ARES, TA Instruments) equipped with a force rebalance transducer, a high-torque motor and a forced convection oven. All experiments were performed under a nitrogen blanket using a rotary parallel plate geometry with 25 mm diameter plates. Disks of the composites (25 mm diameter, 1 mm thickness) for rheological measurements were prepared using compression molding. Dynamic mechanical properties were measured at 180, 190, 200, and 220 °C. A single sample disk was used for the frequency sweep tests at all temperatures: we have verified that there is no molecular weight degradation of the iPP during the duration of our tests. At each temperature, we initially performed strain sweeps at 1 and 100 rad · s-1 to determine the extent of the linear viscoelastic region at these frequencies. We then performed frequency sweeps from 0.1 to 1 rad · s-1 and from 1 to 100 rad · s-1 and selected the strain for each experiment such that (a) the torque measured by the instrument is at least an order of magnitude above the minimum limit claimed by the manufacturer and (b) we are always well in the linear viscoelastic region. We also investigated the effect of annealing the composite at 200 °C for 2 h on its rheological properties by repeating frequency sweep measurements on the annealed sample every 30 min.

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Figure 1. X-ray diffraction of vinyl clay powder. Peaks are indexed by comparison with talc (as discussed in the text) and are indicated in the figure. The d-spacing corresponding to the (001) peak is 1.3 nm.

Figure 2. FTIR of the vinyl clay powder. We show the silicate band in the region of 1000-1100 cm-1 and the characteristic absorptions corresponding to the vinyl group (1642 cm-1) and the MgO-H (3704 cm-1).

3. Results and Discussion 3.1. Organoclay Characterization. We index the X-ray diffraction peaks from the vinyl clay by comparison with natural 2:1 phyllosilicates such as montmorillonite or talc (Figure 1). Peaks are observed near 2θ values of 10, 25, 35, and 60° for copper KR radiation, confirming the structural similarity of natural and vinyl clays.29 However, the vinyl clay shows broader peaks compared to natural clays indicating structural disorder in the synthesized material. The 001 peak indicates a plate-plate spacing of 1.3 nm, which is larger than typically observed for natural clays ≈1.0 nm, since the vinyl clay structure has to accommodate the organic group bonded to every silicon atom in the framework. The upturn in the scattered intensity at low 2θ values has been observed previously for synthetic organoclays and has been interpreted as coming from the colloidal clay particles or from voids between these particles.28 FTIR of the vinyl clay structure shows absorbances in the region of 1000-1100 cm-1 (corresponding to Si-O-Si bond stretching) and at 3704 cm-1 (corresponding to an MgO-H stretching band, Figure 2). We can also clearly observe an absorbance at 1642 cm-1 that is characteristic of vinyl stretching. Our XRD and FTIR data establish that the vinyl clay synthesized has a structure analogous to a 2:1 phyllosilicate with an increased interplate spacing due to the presence of a vinyl moiety bonded to each silicon atom in the clay, but with considerable disruption in structural regularity.

Figure 3. Weight loss from thermogravimetric analysis of the vinyl clay powder is indicated on the left axis. The right axis indicates the corresponding derivative plot.

Figure 4. SEM of the vinyl clay indicating that it is platelike. The plates are stacked to form larger particles. We estimate the characteristic sizescale of the plates from the SEM to be on the order of 500 nm.

The ideal structural formula for the vinyl clay [R8Si8Mg6O16(OH)4 where R ) vinyl], yields a vinyl group weight percentage of around 23.7%. Thermogravimetric analysis of the vinyl clay shows a two stage decrease in weight on heating in nitrogen (Figure 3). There is an initial steep decrease in weight starting at about 330 °C (showing a peak in the derivative curve at about 412 °C) and a more gradual decrease starting at about 450 °C. There is a ≈20% decrease in weight from 330 to 450 °C, that approximately corresponds to the weight fraction of the vinyl group from the ideal structural formula (23.7%). The gradual decrease in weight beyond 450 °C might be due to the dehydroxylation of the phyllosilicatelike framework. Elemental (CHN) analysis of the vinyl clay reveals a carbon content of 20.65% that compares well with the 21.1% expected from the ideal structural formula. SEM of the vinyl clay reveals agglomerates comprised of platelike particles that are approximately 500 nm in size (Figure 4). The vinyl clay does not disperse in water or in organic solventsstherefore, it was not possible to prepare samples suitable for TEM. That the vinyl clay does not disperse in organic solvents suggests that the layered silsesquioxane synthesized might have some interlayer crosslinks that prevent separation of the layers in compatible solvents. 3.2. Composites of iPP and Vinyl Clay. Vinyl clay does not influence the crystallization of the iPP. iPP in the composite crystallizes in the alpha modification regardless of the fraction of vinyl clay present (Figure 5). Further, the crystallization temperatures on cooling from the melt are, within experimental

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Figure 5. X-ray diffraction from the vinyl clay-iPP composites at different loadings of the vinyl clay (from bottom, 0%, viz., near iPP; 1%, 2.5%, 7.5%, and 10% vinyl clay). The curves are shifted vertically for clarity and are plotted on an arbitrary intensity scale. The (001) clay peak is visible at low 2θ ( ≈1 rad · s-1, addition of the vinyl clay to the iPP leads to a small (≈ 10-20%) decrease in the moduli, G′ and G′′ (and therefore in the complex viscosity, η*). Thus, addition of the vinyl clay appears to plasticize the iPP matrix and decreases the complex viscosity at ω >≈ 1 rad · s-1. This behavior is unusualsnanocomposites based on surfactantmodified montmorillonites typically show an increase in viscosity relative to the matrix polymer.32 However, a decrease in the dynamic moduli was reported when polystyrene (PS) was compounded with a clay modified with oligomeric polystyrene.36 In that case, the decrease was attributed to the mixing of the oligomeric PS with the matrix polymersthis mechanism is unlikely to be important in the case of the vinyl clay-iPP composites. A reduction of the composite viscosity relative to the matrix has been reported for a variety of filled systems. For example, calcium carbonate nanoparticles that have not been surfacetreated to compatibilize them with the matrix polymer decrease the viscosity of polystyrene38 and of poly(n-butylmethacrylate)39 when compounded at modest levels ( ≈1 rad · s-1, both elastic and loss moduli are dominated by the iPPshowever, there is a reduction in G′ and G′′ of the composites relative to the iPP (viz. the iPP is plasticized), probably due to slip at the

weak vinyl clay-iPP interface. This plasticization is more prominent with increased clay loading and at higher frequencies. Thus, the rheology of the composite is essentially governed by the polymeric matrix and even exhibits time-temperature superposition44 (TTS fails only for G′ at low frequencies, where we observe the elastic plateau from the vinyl clay network; see data in the Supporting Information). The vinyl clay influences the rheology both at low frequencies (where the well-dispersed fraction of the vinyl clay forms an elastic network, indicating compatibility with the iPP) and at high frequencies (where the matrix iPP is plasticized, probably due to slip at the vinyl clay-iPP interface). This combination of an enhanced low frequency modulus and plasticization at high frequencies leads to an enhanced apparent shear thinning that might be advantageous for processing. We estimate the magnitude of the elastic plateau, KC, for the clay network in the composite as the difference between G′ for the composite and the iPP at the lowest frequency measured, ω ) 0.1 rad · s-1. Thus KC ) G′(iPP ⁄ X) - G′(iPP) at ω ) 0.1 rad · s-1

(1)

KC shows a nonmonotonic dependence on the vinyl clay loading (Figure 6a; a plot of KC as a function of clay loading is shown in the Supporting Information), first increasing until a clay loading of 7.5% and the decreasing when the clay loading is increased to 10%. TEM images suggest a reason for the dependence of KC on vinyl clay concentration. TEMs of our composites show a spatially inhomogeneous distribution of vinyl clay particles in the iPP. The vinyl clay appears to preserve its layered structure after compounding with iPPsthis is in agreement with the XRD data that shows that the 001 clay peak is retained in the nanocomposite (Figure 4). However, for iPP with low clay loadings, for example 2.5%, TEM shows a moderate degree of dispersion of the vinyl clay into smaller stacks (Figure 7a). For the 10% vinyl clay composite, we observe only large particles with a size ≈ O(10 µm) in the TEM (Figure 7b) and we were unable to observe any region where these large particles had broken up into smaller stacks. Thus, our data suggests that, for concentrations of vinyl clay higher than 7.5%, our compounding protocol appears to be unable to disperse the vinyl clay effectively in the iPP and only large particles are observed in the matrix. TEM at other loadings of vinyl clay are presented in the Supporting Information. Our TEM data does not clearly show that the vinyl clay in the composite forms two populations, one of dispersed platelets and the other of aggregates, as suggested by rheology. We believe that we are unable to observe the dispersed vinyl clay platelets because of limitations in TEM resolution, possibly due to the low contrast between the polymeric matrix and the structurally imperfect vinyl clays. We are unable to obtain good TEM images from our composites at higher resolution than is presented here. The elastic plateau, KC, decreases with temperature at a fixed vinyl clay concentration. This is apparent, for example, in the low frequency G′ in the time-temperature superposition plot (see data in the Supporting Information). When we plot ln[KC] against inverse temperature (Figure 8) in an Arrhenius plot, we obtain an “activation energy”, E/R of 4154 ( 1780 K (averaged over all clay concentrations). The large uncertainty in the value of E/R is due to the relatively few data points to which the Arrhenius relation is fitted. The physical meaning of this activation energy is not apparent. However, it is interesting to note that the E/R obtained here is comparable to the flow activation energy of iPP (Eflow/R ) 4652 K).45

3896 Ind. Eng. Chem. Res., Vol. 47, No. 11, 2008 Table 1. Solid Mechanical Properties (from Tensile and Impact Tests) of the iPP and the Vinyl Clay-iPP Composites

Figure 7. TEMs of cryo-ultra-microtomed sections of vinyl clay-iPP composites containing (a) 2.5% and (b) 10% vinyl clay. The scale bars are as indicated.

Figure 8. Temperature dependence of the elastic modulus of the vinyl clay network in the iPP. We present data as an Arrhenius plot for vinyl clay loadings of 1%, 2.5%, 7.5%, and 10%.

Annealing the vinyl clay-iPP composites in nitrogen at 200 °C for up to 2 h does not have any impact on the rheological properties; dynamic mechanical moduli obtained from the composite at the beginning, during, and at the end of annealing are identical within the resolution of our instrument (see data in the Supporting Information). This data indicates that annealing does not lead to further interaction of the vinyl clay with the iPP and suggests that the microstructure structure obtained after compounding is thermodynamically stable.

clay conc (%)

tensile modulus (MPa)

0 5 10

1131 ( 84 1654 ( 410 1226 ( 161

tensile strength elongation at break impact strength (J · m-1) at break (MPa) (mm · mm-1) 28 ( 4.8 31 ( 2.5 21.5 ( 5.8

8.4 ( 1.7 8.9 ( 0.35 5.0 ( 2.9

35.4 ( 3.6 39.8 ( 1.5 33.6 ( 1.5

Finally, we present preliminary measurements of the solid properties of the vinyl clay-iPP composites. The stressextension plots from tensile tests of the composites are qualitatively similar to that of the neat iPP matrix. The tensile stress is linear with extension to roughly the same degree of extension as for the iPPshowever, the composites show a higher tensile modulus relative to iPP (Table 1). The tensile modulus increases by about 50% for the composite containing 5% vinyl clay (1.65 GPa) relative to iPP (1.13 GPa) but then decreases to less than a 10% increase for the 10% vinyl clay composite (1.23 GPa). This agrees well with our melt rheological data that suggests that our compounding is unable to effectively disperse the vinyl clay in iPP above a clay content of 7.5%. The magnitude of the increase in composite modulus over the matrix iPP is significantly higher than for traditional microcomposites46 at the same filler loadingshowever, it compares well with that reported in the literature for iPP-montmorillonite nanocomposites.13 Considering that XRD and TEM indicate that vinyl clay does not exfoliate in the iPP, the substantial increase in modulus obtained at low vinyl clay loadings for our composites is surprising. For conventional polymer-montmorillonite nanocomposites, the increase in modulus typically comes at the expense of decreased toughness (viz. decreased impact strength) and decreased ductility.47–50 However, our vinyl clay-iPP composites show a marginal increase in elongation at break and impact strength at a loading of 5% relative to the iPP, before decreasing at a loading of 10% (Table 1). A simultaneous increase in the modulus, and in the ductility and toughness, is unusual51 and is certainly not anticipated for polymer clay microor nanocomposites. We speculate that the relative flexibility of our structurally imperfect vinyl clays relative to natural clays and the weak vinyl clay-iPP interface might lead to the observed balance of mechanical properties. When the vinyl clay disperses in the iPP, the organoclay layers form networks that result in a higher solid modulusshowever, the failure properties are not adversely affected probably due to dissipation of energy in bending the vinyl clay or due to debonding at the weak iPP-vinyl clay interface. 3.3. Discussion. The rheology and solid properties of the composite are a function of its structure, viz. the dispersion of vinyl clay in the iPP. Here, we discuss the relation of the observed behavior to the structure in conventional clay-based polymer nanocomposites. For conventional surfactant-modified polymer-montmorillonite nanocomposites, enhanced mechanical moduli are observed when the clay platelets exfoliate and are well dispersed in the polymer. Typically, exfoliation of the clay in the polymer is manifested in a disappearance in the (001) XRD peak from the clay. The dispersion of clay in a polymer is determined by the interactions between the clay, the surfactant modifier, and the matrix polymer. The thermodynamic model of Balazs et al.52 suggests that an increase in the length of the surfactant tail favors intercalation of the polymer into the clay galleries as the polymer suffers a lower entropic penalty. This prediction is validated by the results of Reichert et al.50 who show that a minimum of an eight carbon tail surfactant is required for iPP

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to intercalate polar clays. They show that enhancements in the mechanical properties of iPP nanocomposites are observed only for long tail surfactant modified clays and that nanoscopic dispersions (viz. intercalation/exfoliation) are necessary for increase in stiffness and strength. Intercalation or exfoliation of the clay is also influenced by the interaction between the matrix and the polar silicate surface of the clay. Fornes et al.49 found that nylon intercalates clay modified with a single-tail surfactant, but not with a two tailed surfactant. They explain that the second surfactant tail reduces the interaction between the polar nylon and the surface of the clay and hence renders the intercalation of the clay unfavorable. Ogawa et al.17 showed that magadiite covalently modified by a C8 silane does not intercalate n-decane, but swells when exposed to n-octyl alcohol. Their results suggest that the interaction of the polar clay surface with the intercalant cannot be neglected, even after organic modification of the clay surface. Thus, the formation of intercalated/exfoliated nanocomposites of clay with highly apolar polymers such as iPP requires the addition of polar compatibilizers such as maleic anhydride modified iPP that can interact both with the polar silicate tetrahedra at the clay surface and with the apolar matrix polymer.32,50,53,54 Finally, the intercalation of polymer into clay is a function of the degree of modification of the clay surface. Galvin et al. have shown that moderate coverage of the clay surface by modifier leads to intercalation while for dense covererage, the polymer experiences “autophobic dewetting” and does not intercalate the clay galleries.55 They rationalize their results on the basis of Leibler’s criterion56 that states that dewetting happens when σ√N > (N/P)2

(2)

wherein σ is the surface coverage by the modifier, N is the molecular size of the modifier, and P is the polymerization index of the matrix polymer. In our work, the structure of the vinyl clay is talclike, with each silicon atom in the framework covalently attached to a vinyl group. Therefore, the vinyl clay surface is densely covered (σ ≈ 10 nm-2) with a very short (C2) organic group. The modifier size is significantly shorter for the vinyl clay than the critical surfactant length of eight carbon units observed by Reichert et al.50 Leibler’s criterion is also clearly satisfied in our case (N ) 1 for a vinyl group and P ≈ 1350; σ N ) 10 . 5.5 × 10-7 ) (N/P)2) and suggests that iPP will dewet the vinyl clay surface. Our results suggest that the vinyl clay synthesized is comprised of two fractions, one minor fraction that is able to disperse in the iPP matrix and form a network structure and the major fraction that is unable to disperse in iPP (or even in organic solvents), possibly due to interlayer crosslinks formed during the vinyl clay synthesis. Our data present a puzzling (but advantageous for applications) combination of properties: the rheology data show vinyl clay loading dependent decrease in melt viscosity that is more prominent at higher frequencies. This strongly suggests that the iPP-vinyl clay interface is weak and that iPP does not strongly interact with the vinyl clay surface. This agrees well with the results of Reichert et al.50 and with Leibler.56 However, we also report data that suggest faVorable vinyl clay-iPP interactions. For example, we observe the signature of a networklike structure of the vinyl clay in the low frequency elastic modulus of the composite. Further, we see a significant increase in the tensile modulus of our iPP composite, comparable to the enhancements reported for compatibilized iPP-montmorillonite nanocomposites.13,50 TEM also indicates

that at least a few vinyl clay stacks disperse in iPP at low loadings, less than 10%. What then, is the structure of the vinyl clay-iPP composite and what are the interactions that lead to the formation of this structure? We do not have an unambiguous answer to this. It is possible that a fraction of the vinyl clay has a favorable interaction with the iPP matrix and, therefore, disperses to form a network in the melt. The structure of the other fraction appears to be dominated by interlayer crosslinks that preclude dispersion in organic solvents or in the iPP. Further, this fraction appears to be dewet by the iPP leading to slip at the interface. Therefore, the vinyl clay-iPP interactions appear to be tuned such that we observe a weak plateau modulus at low frequencies combined with enhanced shear thinning at higher frequencies and a combination of enhanced modulus in the solid state without compromising on the elongation at break or impact strength. In addition, dispersion of the vinyl clay into the iPP matrix is also a function of the loading of the vinyl clay. For the compounding protocol adopted, we see an optimal balance of properties at loadings lower than 10%. The plateau modulus in the melt state decreases when the loading is increased to 10%, as does the solid tensile modulus. We believe that this is due to a complicated interplay between the complex microstructure of the vinyl clay fractions and iPP in the composite and, the processing conditions. It is curious to note that the optimum loadings that we observe in our systems are similar to those observed for conventional montmorillonite-iPP nanocomposites. For slow relaxing matrices like polymers, the structure in the polymer-clay nanocomposites are often kinetically formed during processing and are not always the result of thermodynamic interactions between the clay, modifier, and polymer. However, in our case, we do not observe any change in the dispersion of the clay or on melt rheological properties even after annealing the composite at 200 °C for 2 h. This suggests that the structure observed in our experiments is not determined by kinetic considerations. 4. Conclusions We report the preparation of composites of isotactic polypropylene with a vinyl-modified Mg silsesquioxane (vinyl clay). Vinyl clay was synthesized under mild reaction conditions, in aqueous solution at room temperature, and was dispersed in the iPP by melt compounding. XRD suggests that the vinyl clay does not completely exfoliate in the iPP. TEM of the composites reveals that vinyl clay disperses into thinner platelets for vinyl concentrations lesser than about 10%. It appears that our compounding protocol is not effective in dispersing the vinyl clay at concentrations above 7.5%. The vinyl clay does not influence the crystallization of the iPP. In the melt state, iPP-vinyl clay composites show a plateau in the elastic modulus at low frequenciesshowever, the loss modulus exceeds the elastic modulus even at low frequencies. Vinyl clay plasticizes the iPPsthis effect is more pronounced at higher frequencies in our dynamic experiments suggesting the possibility of slip at the iPP-vinyl clay interface. The low frequency plateau in the elastic modulus increases with clay loading until a loading of 7.5% (indicating that the clay does not disperse effectively for higher loadings). The elastic plateau decreases with increase in temperature. The tensile modulus of the iPP-vinyl clay composites increase relative to iPP. At a vinyl clay loading of 5%, we see approximately 50% increase in tensile moduluss comparable to the enhancements observed in exfoliated clay-iPP nanocomposites. The increase in modulus is related to the dispersion of the vinyl clay in the iPPstherefore, at a loading

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of 10%, the tensile modulus is lower than at a 5% loading. Surprisingly, the elongation at break for our composites is comparable to that for iPP (in fact, increasing marginally for the composite at a 5% loading) and the impact strength is unaffected. We believe that our work represents the first example of melt-compounded organically modified layered metallosilsesquioxane polymer composites. Acknowledgment We acknowledge NCL for supporting this work (Project: MLP004926). We acknowledge Dr. C. Ramesh for use of the Rigaku-XRD and Deepa Dhobale for use of the FTIR. We gratefully acknowledge Dr. Atanu Basu of the National Institute for Virology, Pune, for TEMs of the 7.5% vinyl clay composite reported in the Supporting Information. Some of the vinyl clay synthesis and compounding with iPP was performed by Mr. Aditya Patwardhan, NCL. We acknowledge Dr. Sujata S. Biswas (NCL) for performing the elemental analysis. Supporting Information Available: Rheological data, polarized optical micrographs of composites, and differential scanning calorimetry (DSC), TEM, and XRD data are presented. This material is available free of charge via the Internet at http:// pubs.acs.org. Literature Cited (1) Utracki, L. A. Polymer Alloys and Blends; Thermodynamics and Rheology; Carl Hanser: Munich, 1989. (2) Pfaendner, R. How will additives shape the future of plastics. Polym. Degrad. Stab. 2006, 91, 2249. (3) Fina, A.; Abbenhuis, H. C. L.; Tabuani, D.; Camino, G. Metal functionalized POSS as fire retardarnts in Polypropylene. Polym. Degrad. Stab. 2006, 91, 2275. (4) Phillips, S. H.; Haddad, T. S.; Tomczak, S. J. Developments in nanoscience: polyhedral oligomeric silsequioxane (POSS)-polymers. Curr. Opin. Solid State Mater. Sci. 2004, 8, 21. (5) Kojima, Y.; Usuki, A.; Kawasumi, M.; Okada, A.; Fukushima, Y.; Kurauchi, T. T. Kamigaito. O. Mechanical properties of nylon 6-clay hybrid. J. Mater. Res. 1993, 8, 1185. (6) Kojima, Y.; Usuki, A.; Kawasumi, M.; Okada, A.; Kurauchi, T.; Kamigaito, O. Synthesis of nylon 6-clay hybrid by montmorillonite intercalated with -caprolactam. J. Polym. Sci. Part A: Polym. Chem. 1993, 31, 983. (7) Vaia, R. A.; Ishii, H.; Giannelis, E. P. Synthesis and properties of two-dimensional nanostructures by direct intercalation of polymer melts in layered silicates. Chem. Mater. 1993, 5, 1694. (8) Giannelis, E. P. Polymer layered silicate nanocomposites. AdV. Mater. 1996, 8, 29. (9) Lebaron, P. C.; Wang, Z.; Pinnavaia, T. Polymer-layered silicate nanocomposites: an overview. J. Appl. Clay Sci. 1999, 15, 11. (10) Ray, S. S.; Okamoto, M. Polymer/layered silicate nanocomposites: a review from preparation to processing. Prog. Polym. Sci. 2003, 28, 1539. (11) Garces, J. M.; Moll, D. J.; Bicerano, J.; Fibiger, R.; McLeod, D. G. Polymer Nanocomposites for Automotive Applications. AdV. Mater. 2000, 12, 1835. (12) Edser, C. Auto applications drive commercialization of nanocomposites. Plast., Addit. Compounding 2002, 4, 30. (13) Manias, E.; Touny, A.; Wu, L.; Strawhecker, K.; Lu, B.; Chung, T. C. Polypropylene/Montmorillonite Nanocomposites. Review of the Synthetic Routes and Materials Properties. Chem. Mater. 2001, 13, 3516. (14) Xie, W.; Gao, Z.; Pan, W.-P.; Hunter, D.; Singh, A.; Vaia, R. Thermal degradation chemistry of alkyl quaternary ammonium Montmorillonite. Chem. Mater. 2001, 13, 2979. (15) Yoon, P. J.; Hunter, D. L.; Paul, D. R. Polycarbonate nanocomposites: Part 2. Degradation and color formation. Polymer 2003, 44, 5341. (16) Ruiz-Hitzky, E.; Rojo, M. Intracrystalline grafting on layer silicic acids. Nature 1980, 287, 28. (17) Ogawa, M.; Okumoto, S.; Kuroda, K. Control of interlayer microstructures of a layered silicate by surface modification with organochlorosilanes. J. Am. Chem. Soc. 1998, 120, 7361.

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ReceiVed for reView December 5, 2007 ReVised manuscript receiVed March 5, 2008 Accepted March 6, 2008 IE071658P