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Morphology of EVA based nanocomposites under shear and extensional flow.

INTRODUCTION

Polymer-layered silicate nanocomposites (PLSN) have gained tremendous interest in both the academic and research fields because of their unique structure as well as physical and chemical properties (1). The mere addition of a small quantity of inorganic material (typically in the range of 5 wt%) offers well-enhanced strength and other properties compared to the pure polymer or conventional composites. These widely reported improvements include better resistance to gas and liquid permeability, higher heat distortion temperatures and better tensile strengths and moduli. These properties make the nanocomposites ideal for industries such as packaging, automotive, wire and cable, and electronics (2).

Structurally, polymer-clay complexes can be classified as nanocomposites or "conventional composites." This depends on the nature of components and the preparation technique. The nature of components refers to the type of silicate material, the organic cation used to render the hydrophilic silicates organophilic, and the nature of the polymer matrix. The preparation technique pertains to mechanical factors that facilitate the penetration or intercalation of polymer chains into the layers of silicate. This may lead to exfoliation or delamination of silicates into individual layers. These mechanical factors include the mechanical shear or extension employed, residence time and type of mixer. Depending on these factors, three morphologies are possible: phase-separated, intercalated, or exfoliated (3). "Phase-separated" refers to composites that maintain immiscibility between the polymer and the inorganic filler. In this morphology, the polymer chains do not penetrate into the clay layers. There is minimal reinforcement by the fillers in this structure. "Intercalated" structures are obtained when polymer chains have penetrated deep between the layers of silicates, resulting in ordered structures. The intercalation of the polymer chains into the layer galleries results in its expansion of the distance between the silicate layers. Because of mechanical shearing forces and interactions between the silicates and polymer chains, the stacks of layered silicates may have dispersed and distributed within the matrix, thus increasing surface area of contact with the polymer. Intercalated structures have been reported to have regions of high and low reinforcements (4). "Exfoliated" morphologies result when individual layers (~1 nm) are well dispersed and distributed throughout the polymer matrix, the average distance between them being dependent on the filler concentration. This structure facilitates maximum reinforcement because of the huge surface area of contact with the matrix.

Layered materials are suitable for the design of nanocomposites. This is due to their lamellar elements that have high in-plane strength and stiffness and a high aspect ratio (> 50). The clay material has a very high surface area of about 750 [m.sup.2]/g (e.g. montmorillonite). Almost all groups of lamellar solids, especially smectite clays, are the main choice for nanocomposite materials for two reasons (5, 6): Their rich intercalation chemistry allows them to be chemically modified and made compatible with organic polymers for dispersal on a nanometer scale; and they can be easily acquired at low costs.

To gain a thorough knowledge of the processing and applications of this relatively new class of material, it is important to understand its rheological and mechanical properties. Rheology of polymer nanocomposites is useful for two reasons (7): It describes the behavior of material while undergoing processes like injection molding in the melt state; and it is sensitive to the microstructure and morphology of the material.

Melt rheological properties are dictated by a combination of mesoscopic structure and the strength of interaction between the polymer and the layered silicates. Further, the mesoscopic structure would be crucially dependent not only on the strength of the polymer/layered silicate interactions, but also on the inherent viscoelastic properties of the matrix in which the layers or collection of layers are dispersed (8). Much of the rheological studies that have been conducted on polymerlayered silicate nanocomposites have concentrated mainly on shear rheological characteristics. Melt extensional properties (e.g. uniaxial and biaxial extension) are one of the most important deformations in polymer processes such as melt spinning, injection molding and film blowing. Yet very limited research has been carried out to determine the extensional flow behavior of layered silicate nanocomposites (9, 10). Owing to the lack of attention given to this area of rheology for polymer nanocomposites, it is unknown as to how the layered silicates behave when subjected to such deformations as well as to their contribution to the overall material property during extension.

It is unclear how the layered silicate structures orient themselves when subjected to an extensional type of deformation. An attempt has been made in this work to examine the structural change of these materials due to extension. An online laser light scattering (LLS) unit has been used to study the morphological change in an extensional deformation field of polymer nanocomposites. This technique has been incorporated into a melt extensional setup to obtain a "pictorial" view of the deformation process.

In this paper, linear viscoelastic, steady shear and melt extensional properties of ethylene-vinyl acetate copolymer (EVA)/layered silicate nanocomposites have been measured. Online measurements of the deformability of the nanocomposites have also been made. The extensional properties have been used to explain the structure of the molten material under extensional deformation.

EXPERIMENTAL

Materials

The polymer used to prepare the nanocomposites was ethylene-vinyl acetate copolymer (EVA) containing 9 wt% vinyl acetate obtained from Elf AtoChem (Australia). The weight-average molecular weight ([M.sub.w]) was 67,320 and its polydispersity ([M.sub.w]/[M.sub.n]) was reported to be 4.6. The density of the polymer was 929 kg/[m.sup.3] and its melt flow index (MFI) under test conditions of 190[degrees]C and 2.16 kg load was 3 (g/10 min). The filler material was Cloisite[R] 15A (C15A) obtained from Southern Clay Products. C15A comes under the group of layered silicates called montmorillonite (MMT). Pristine, hydrophilic [Na.sup.+]-MMT was modified to render it organophilic by cation exchange reaction with dimethyl dihydrogenated tallow quaternary ammonium chloride (11). The cation exchange reaction increases the affinity of polymer chains toward the layered silicates, thus facilitating intercalation process. The cation exchange capacity (CEC) of the clay was 125 meq/100 g and it had a specific gravity of 1.66.

Preparation of Nanocomposite

The EVA/C15A nanocomposites were initially premixed by simply stirring the organoclay and EVA pellets. The premix was then melt-blended by introducing it into a Brabender twin-screw extruder. The extruder was operated at 100[degrees]C and at 70 rpm. Samples with clay loadings of 2.5 wt%, 5 wt% and 7.5 wt% were produced. The extruded materials were then pelletized and compression molded at 120[degrees]C to yield samples of about 2 mm thickness.

Experimental Methods

Wide Angle X-ray Scattering (WAXS)

X-ray diffraction (XRD) was used to analyze the extent of EVA intercalation into the silicate layers by measurement of the degree of layer swelling. Wide angle X-ray scattering (WAXS) was used for this purpose. XRD data was obtained using Philips X-ray generator using 30 kV accelerating voltage and 30 mA current. Intensities from 2[theta] = 1.2[degrees] to 30[degrees] were recorded using Ni filtered Cu-K[alpha] ([lambda] = 0.154 nm). XRD was conducted in transmission mode with 2-mm-thick samples placed on a rotating sample holder and the sample was rotated during the scattering test. This type of sample holder can eliminate the effect of any possible orientations of the structures. Details of this method have been discussed elsewhere (12). Background radiation has been removed from the scattering curves to be able to show scattering intensities up to 2[theta] of about 1.3[degrees].

Environmental Scanning Electron Microscopy (ESEM)

To obtain an image of the state of dispersion and distribution of silicate layers in the polymer matrix, a FEI Quanta 200 environmental scanning electron microscope (ESEM) was used. The instrument was operated using 20 kV electron accelerating voltage and 0.5 Torr pressure. The sample preparation was simple because no conductive coating was required as the gas in the sample chamber prevented any charge buildup on the sample. Samples were freeze-fractured in liquid nitrogen to obtain their cross-sectional images.

Shear Rheology

The 2-mm-thick compression-molded samples were cut into 25-mm-diameter discs. Shear rheological properties were studied using the Advanced Rheometrics Expansion System (ARES) in both dynamic and steady shear modes with parallel plate geometry. The tests were conducted using 25-mm-diameter plates at 130[degrees]C. All measurements were performed with a force transducer with a range of 0.2 to 200 g-cm torque. Dynamic tests include the determination of linear viscoelastic regions (dynamic strain sweep) and dynamic frequency sweeps. Dynamic frequency sweeps were conducted to determine the microstructure and dynamics of the materials concerned. The tests were conducted at frequencies ranging from 100 rad/s to 0.001 rad/s. Nonlinear rheological properties were probed using steady shear mode for shear rates of 0.007 [s.sup.-1] to 10 [s.sup.-1].

Extensional Rheology

To obtain data for uniaxial extensional flows, two different techniques were employed. Transient extensional data were obtained using RME (Rheometrics Melt Extensional) Rheometer. Details of the experimental setup are described elsewhere (13). Melt strength tests were conducted using draw-down experiments. For this purpose, a Gottfert Rheotens melt strength tester (14) was employed. While the melt strength can be used as a measure of melt quality, it is not a well-defined rheological property. This is because of nonuniform strain and temperature along the drawn filament (14).

RME: Samples to be tested were cut into small rectangles of dimensions 1.2 mm (H) X 7.0 mm (W) X 60 mm (L). The samples were held on a table supported by a cushion of air at 130[degrees]C. Constant rotational clamps were used to deform the samples with the application of constant strain rates (0.1 [s.sup.-1], 0.5 [s.sup.-1] and 1 [s.sup.-1]). A displacement transducer was used to measure the extensional force produced by the sample.

Rheotens: Polymer pellets were fed to a Haake single-screw extruder fitted with capillary die of 2-mm diameter and 40-mm length. The extrusion pressure was read by means of a pressure transducer measured upstream of the die. The extrusion flow rate of 6.67 X [10.sup.-5] kg/s was maintained by means of a gear pump. The Rheotens tester was placed at a height of 210 mm below the die. The extruded melts were drawn by four counter-rotating wheels of which the top two are mounted on a beam balance. The bottom rollers serve to guide the drawn filament away from the top rollers. The tensile force in the strand was measured and plotted as a function of roller velocity. The velocity of the rollers increased at a constant acceleration until the extruded strand broke, and the force at which this breakage occurred is termed the "melt strength" (14). Because of very high draw rates, inertial and gravitational forces have been neglected (15). The melt temperature of the drawn strand is dependent on the draw ratio and rate of drawing. The temperature at the die was set at 130[degrees]C. At high draw rates, the extent of ambient cooling of the filament was minimized and the stretching could be considered to be nearly isothermal.

Laser Light Scattering

Scattering techniques have gained prominence in polymer science and are non-invasive. The principle behind the operation of LLS (or for that matter any light scattering) is similar to that of other scattering techniques. An extensive account of light scattering techniques can be found in literature (16-19).

An optical diffraction unit was designed and built for use as an on-line monitoring device. A solid-state laser with an output power of 1 mW was used as a monochromatic light source. The laser beam was focused on the sample to be tested by a small plastic lens and non-coherent radiation was filtered out using two pinholes arranged 50 mm from each other. The scattered beam was captured by a translucent film adhered to a glass plate with dimensions of 100 X 80 mm. The scattering patterns were recorded by a digital camera positioned 350 mm from the screen. Concentric rings in 20-mm-diameter increments were drawn around the point where the primary beam hit the screen. The primary beam was excluded by a beam stop, which consisted of a small piece of black paper adhered to the screen. To avoid the interfering effect of the environment, the screen was covered by black paper forming a tube between the screen and the video camera. All recordings were done in normal laboratory light, which was the source of some background intensities. The camera recorded the scattering pattern as a function of time with 25 patterns recorded per second. Data processing was carried out on individual pictures obtained from the digital video recordings.

Particle scattering caused varying intensities around the incident beam that was estimated to follow a Gaussian distribution as a function of the scattering vector (Eq 1). The particle scattering is due to inhomogeneities in optical densities of the polymer melt. The Guinier approach of scattering intensities relates the width of a Gaussian peak to the radius of gyration of the so-called scattering particle (i.e., inhomogeneities) present in the polymer melt. The Guinier analysis allows for the determination of the radius of gyration. [R.sub.g], of an arbitrarily shaped particle. [R.sub.g] is defined as the root-mean-square of all the distances from the particle to its center of mass. This method is based on a second-order expansion of the structure factor of the particle (20). The Guinier formula (21) is as shown in Eq 2. A detailed review of this analysis can be found elsewhere (17, 20, 21).

S = [4[pi] sin[theta]]/[lambda] (1)

I(s) = [I.sub.0] [e.sup.(-[S.sup.2][R.sub.g.sup.2]/3)] (2)

Guinier's law proceeds by plotting I(s) vs. S to obtain a Gaussian distribution and a plot of In [I(s)] vs. [S.sup.2] produces a linear plot with intercept In [([I.sub.0])] and slope of [R.sub.g.sup.2]/3. A ratio of the radii of gyration parallel and perpendicular to the extruded axis would be used rather than absolute value of [R.sub.g]. This ratio aims to reveal the extent the material can be deformed and whether incorporation of layered silicates improves its deformability. In this study an attempt is made to relate melt strength tests to the light scattering data.

RESULTS AND DISCUSSION

XRD Results

Figure 1 shows XRD scattering curves of EVA/C15A nanocomposites. C15A is observed to have three broad peaks as shown by the [d.sub.001], [d.sub.002] and [d.sub.003] diffraction lines. The polymeric nanocomposites generally have diffraction lines, [d.sub.001] and [d.sub.002]. The [d.sub.001] spacing of C15A is found to be 3.52 nm. The greatest layer swelling is given by the 2.5 wt% sample in which its d-spacing is 4.64 nm. The d-spacing of the 5 wt% and 7.5 wt% samples have been found to be 4.41 nm and 3.92 nm respectively. The increase in d-spacing for the polymer nanocomposites relative to C15A organoclay shows that the silicate layers have expanded because of intercalation of polymer chains into the gallery spaces. The d-spacing calculation was determined using Bragg's Law equation ([lambda] = 2d sin [theta]).

The diffraction shoulders with very wide and less intense peaks for the 5 wt% and 7.5 wt% samples suggest the possibility of layer incoherence or a form of disordered intercalated structure (22). This means that the layered silicates would have become more disordered, while maintaining an average periodic distance. The 2.5 wt% sample, on the other hand, gave a clear peak, indicating ordered intercalated structure. Exfoliation or delamination of layered silicates can be observed only when XRD is beyond the resolution of Bragg-Brentano geometry. This is the case when there is no observable basal reflection (22).

[FIGURE 1 OMITTED]

ESEM

ESEM images of 0 wt% and 5 wt% nanocomposite samples are shown in Figs. 2a and 2b. Figure 2a shows the appearance of the unfilled matrix while Fig. 2b shows the distribution and dispersion of silicate fillers within the polymer matrix. The lighter shades correspond to the silicate material. The images indicate various size distributions of silicates, ranging from tens of nanometers to hundreds of nanometers. This is a general appearance of silicate distribution in the polymer matrix for all the nanocomposite samples prepared. The large particles seen in Fig. 2b are agglomerates of silicate layers and reinforce the XRD finding (Fig. 1) that the nanocomposites formed were not exfoliated.

Shear Rheology

The effects of fillers on the rheological properties of polymer-filler conventional composites and nanocomposites have been studied extensively. The consensus is that the polymer-filler and filler-filler interactions strongly affect the rheological behavior (hence processing) of these composites (2, 3, 7, 8).

Determination of linear viscoelastic region is essential before commencing tests for frequency sweep to ensure that the microstructure of the material would not be affected by shear alignment. The conditions that satisfy linear viscoelasticity are that the stress is linearly proportional to the imposed strain and the torque response involves only the first harmonic (23). In the linear viscoelastic region, both the storage and loss moduli are expected to be independent of strain amplitude, thus satisfying the first condition. The absence of higher harmonics for the stress response ensures that it remains sinusoidal, thus obeying the second condition. The test, conducted at 130[degrees]C and at a constant frequency of 10 rad/s, shows that it is possible to conduct dynamic tests using strains of between 0.1% and 5% for the nanocomposites before nonlinearity starts. However, for the unfilled structure, it is possible to extend the strain to 14% without exceeding the limit of linearity.

The dynamic frequency sweeps were conducted at 4%, 3% and 2% strains for the 2.5 wt%, 5 wt% and 7.5 wt% nanocomposites respectively. A strain of 5% was used for the unfilled system. Figure 3 shows the storage moduli (G') and loss moduli (G") for the unfilled and filled systems. At high frequencies, the G' and G" values coincided with each other, but at lower frequencies, there was an increase in both moduli as filler loading increased. The increase in G' could be attributed to the ability of the filled systems to store elastic energy (25). Unlike the pure polymer, the filled systems did not exhibit any clear terminal behavior, which is characterized by a slope of two in the low-frequency region (24). The slope refers to the slope of G' vs. [omega] curve in the low-frequency region and is called alpha ([alpha]). However, they did exhibit G' dependence on [omega], with the dependence decreasing with increasing silicate loading (Fig. 4). Independence of G' with respect to [omega] would mean pseudosolid-like characteristic owing to a three-dimensional network structure whereby silicate layers act as physical crosslinkers. The decreasing [alpha] would suggest the possibility of a mesostructure with enhanced silicatesilicate interactions as silicate content increases. This is especially so for the 7.5 wt% sample.

[FIGURE 2 OMITTED]

The frequency dependence of the filled nanocomposite system is different from many other reported polymer nanocomposite systems with intercalated or delaminated fillers (2, 3, 7, 8). In the latter systems, there were monotonic increases in moduli at all frequencies. In delaminated structures, the clay particles are well dispersed in the polymer matrix as single layers. This would facilitate an increase in contact area between the polymer chains and silicate surfaces (24). The increase in polymer-silicate and silicate-silicate interactions leads to the formation of a percolated network structure. With this in mind, the EVA/C15A system did not exhibit a percolated network structure, as no change in slope for [alpha] against silicate loading was observed.

[FIGURE 3 OMITTED]

Steady shear viscosity results are shown in Fig. 5. At high shear rates, the viscosity is independent of silicate loadings, and there is rapid shear thinning due to preferential orientation and alignment of silicate layers and polymer chains in the flow direction. At low shear rates, the shear viscosity increases monotonically with silicate loadings. This increase in viscosity is a result of increased interactions between the silicate layers and polymer chains (26). Viscosity of many filled systems is generally higher than that of the unfilled system at all concentrations. The viscosity of a suspension is generally increased because of enhanced energy dissipation in the presence of solids, as was initially shown by Einstein for a dilute suspension of spheres (27). In filled polymers, this is especially true in low concentration regimes where there are little interparticle interactions, as demonstrated by the 2.5 wt% and 5 wt% nanocomposites. However, at higher loadings, as in the 7.5 wt% nanocomposite sample, no zero-shear viscosity was observed and the deviation from dilute suspension behavior was considerable.

[FIGURE 4 OMITTED]

The nanocomposites tested in this work did not demonstrate any yield stress as would be exhibited by a viscosity versus shear rate slope of (-) 1 in the low shear rate region. The presence of yield stress indicates a network structure, as described earlier. The slope for the 7.5 wt% nanocomposite at the low shear rate region is (-) 0.5. Although a significant network structure is absent, the deviation as shown by the slope does demonstrate that there is some significant interparticle interaction at the highest silicate loading tested. Increasing the loading beyond the test range may in fact produce samples that have yield stress, and hence a complex network structure.

[FIGURE 5 OMITTED]

The correlation of first normal stress difference (N1) and shear stress shows the effect of the silicate fillers on the elasticity of the system. From Fig. 6, it is seen that N1 is independent of silicate loadings for the given shear stress range tested. This is a characteristic behavior of anisotropic platelets or disc-like fillers. Similar results have been reported by other researchers studying intercalated systems (28). These observations have been attributed to the ability of the two-dimensional silicates to be preferentially oriented in the direction of shear. This preferential orientation has been experimentally detected using XRD and ESEM, as shown in Figs. 1 and 2. Because of this orientation, effective filler-filler interactions and their subsequent contribution to the elasticity of the system are reduced. The three-dimensional interactions that are present in the quiescent conditions have to give way to interactions in two dimensions (28). However, such independence of N1 with different filler contents was not reported with exfoliated EVA (18% and 28% VA) nanocomposites whereby there was a decrease in N1 with filler concentration (9, 10). This decrease has been attributed to the strong interactions between the clay layers and polymer chains. Interestingly, the independence of N1 with silicate loadings for this system was not repeated for the storage moduli, as already described in Fig. 3.

Extensional Rheology

Limited studies have been conducted on extensional rheological behavior of polymeric nanocomposites (9, 10). This is despite the fact that extensional properties can be very influential in determining the flow behavior. In this section, results of the studies on effect of layered silicates on the melt strength and extensional behavior of EVA nanocomposites will be discussed.

[FIGURE 6 OMITTED]

RME Results

Figure 7 shows logarithmic plots of transient extensional viscosities as a function of time. They were conducted at extensional strain rates of 0.1 [s.sup.-1], 0.5 [s.sup.-1] and 1.0 [s.sup.-1]. For clarity of presentation, the viscosity data for 0.5 [s.sup.-1] and 1.0 [s.sup.-1] were multiplied by factors of 20 and 50 respectively. These measurements reveal that the extensional viscosities gradually increase with time. This region of the curves is termed the linear extensional region (29). However, at longer time scales, the extensional viscosities increase rapidly, until it reaches a plateau. The region of rapid increase in extensional viscosities in the so-called nonlinear extensional region is termed strain hardening. Such a behavior of nanocomposites has been recently reported in the literature (10, 29, 30).

At all strain rates measured, extensional viscosity appeared to be fairly independent of silicate loadings for the unfilled 2.5 wt% and 5 wt% samples. This suggests that the silicate layers have little or no effect on extensional characteristics of the material at low loadings. The 7.5 wt% sample, on the other hand, exhibited a slightly higher extensional viscosity in the linear extensional region at all strain rates measured. The extensional viscosity eventually became independent of silicate loading in the nonlinear extensional region. A possible explanation for the slight difference in extensional behavior for the 7.5 wt% sample is the increased silicate-silicate interactions due to its higher loading. Similar behavior was observed in the viscoelastic characteristics mentioned earlier (Fig. 3a). As the Hencky strain increases (Fig. 8), the silicate particles align orthogonally (10, 29) to the axis of extension, resulting in an increase in three-dimensional interactions. These three-dimensional interactions, often termed "house of cards" arrangements (10, 29, 31), increase the extensional viscosity.

[FIGURE 7 OMITTED]

Previous studies (10) on exfoliated EVA (18 wt% and 28 wt% VA) have shown that nanocomposites of loading 2.5 wt% and 5 wt% exhibited a reduced strain-hardening effect with increase in filler loading. These systems are analogous to microcomposite structures studied by these authors (32-34). Although the "house of cards" structure increases the extensional viscosity, the silicate layers are also reported to aggregate to form microflocs. This results in a decrease in the strain-hardening effect compared to the unfilled material (10, 29). These authors with the help of ESEM and TEM techniques have observed the formation of aggregates, within the nanocomposites when stretched.

Rheotens Results

The Rheotens tests for unfilled and filled systems were conducted at mass flow rates of 6.67 X [10.sup.-5] kg/s and nip roller acceleration of 12 mm/[s.sup.2]. Figure 9 shows the Force-Extensibility curve obtained for the unfilled and filled systems. The force at which the sample breaks, called the Melt Strength, is a relative measure of the extensional strength of the material (35). The maximum draw ratio is regarded as the extensibility or drawability of the melt under the test conditions studied. Figure 9 also indicates that at higher draw ratios it is observed that both drawability and draw-down force of the filled system are marginally higher than that of the unfilled EVA. The melt strength of the samples increases with silicate loadings. The increase in melt strength and drawability, albeit marginal, can be attributed to the interactions between polymer chains and silicate layers. These results support the shear rheological and RME results to a certain extent in that both showed a small rise in melt properties with increase in silicate loading. Previous work (35) has shown that it is not only important to have strong interactions between the silicate layers and polymer chains, but it is also imperative for the system to have well-dispersed layers, in other words, a high degree of exfoliation in which the surface area of contact between the polymer chains and silicate layers has been greatly enhanced.

[FIGURE 8 OMITTED]

[FIGURE 9 OMITTED]

From Fig. 9, it can also be observed that the stability of extruded material at high draw ratios decreases markedly. This instability is characterized by fluctuations in the force-extensibility diagram and is known as draw resonance. This behavior has been observed in many filled systems (36) undergoing uniaxial stretching and melt spinning processes due to the yielding mechanism. The transition from stable to unstable stretching has been shown to depend on the molecular weight of the polymer, structural framework formed from interactions between filler particles, and the changes induced during stretching.

Figure 10 shows the apparent extensional viscosities as a function of apparent extensional strain rates. Both the apparent extensional viscosities and strain rates are calculated from the point of the extruded strand that contacts the nip rollers. Details of these are explained elsewhere (37). The figure indicates that increasing silicate loading increases the apparent extensional viscosities, amounting to strain hardening. The filled system is able to provide greater resistance to extension compared to the unfilled system. This is due to the strong orientation of fillers in direction of stretch and higher degree silicate-silicate interaction.

Laser Light Scattering

The laser light scattering technique was used to analyze structural evolution of the drawn molten material following its exit from the die of the single-screw extruder. Scattering patterns were obtained at different positions along the roller-drawn strands. The positions varied from 25 mm to 165 mm from the die exit at the nip roller acceleration of 12 mm/[s.sup.2]. The video recordings were transferred to individual patterns using EZY[R] software. The scattering patterns were captured as 24 bit/pixel JPG files having 720 X 540 pixels in two perpendicular directions. Red components of the three-color bytes of the pictures were extracted and transferred to an ASCII file.

[FIGURE 10 OMITTED]

A typical two-dimensional scattering image obtained for a drawn 5 wt%-filled EVA is shown in Fig. 11. The particle scattering component was processed using the Guinier concept, i.e., Gaussian curves (Figs. 12a and 12b) were fitted to the central part of the scattering pattern according to the two directions with respect to the stretch. The vertical axis of Fig. 11 corresponds to the direction of stretch. The radius of gyration, [R.sub.g], of the scattering particle was then calculated as outlined earlier in the paper according to the Guinier concept (Eq 2). A detailed data processing description can be found elsewhere (38).

Light scattering was conducted for the unfilled EVA as well as for the 2.5 wt% and 5 wt% nanocomposites. The data processing yielded about a hundred data points for each nanocomposite tested. An average [R.sub.g], both in the direction of extension and orthogonal to it, is plotted in Fig. 13 for the 5 wt% nanocomposite. A linear least square fit was drawn through the points to establish an average for all the positions studied. This fit is equivalent to a master curve of the deformation experienced by the drawn material. The [R.sub.g] was plotted as a function of total extensional strain as experienced by each material element, which is as defined in Eq 3 (39, 40). The ratio in the parentheses is simply the draw or stretch ratio.

[FIGURE 11 OMITTED]

[epsilon] = In ([v.sub.w]/[v.sub.0]) (3)

Figure 13 shows the processed data for [R.sub.g] in two directions. It is clear from the figure that the deformation experienced by the drawn filament at any point is uniaxial since [R.sub.g] corresponding to the direction perpendicular to extensional axis remains almost unchanged with extensional strain. The greatest amount of deformation is experienced in the direction of extension or drawing.

Figure 14 is a plot of deformation ratio as a function of extensional strain, and it describes the extent of deformation experienced by the drawn material. It is interesting to note from this figure that there is an increase in deformation ratio for the unfilled polymer (EVA) with increase in extensional strain. It is presumed that this behavior pattern is due to some inhomogeneity generated within the polymer as a result of the extrusion process. Similar patterns were observed for other unfilled polymer melts, like low-density polyethylene and polypropylene, not shown in this paper. This pattern was absent when the light scattering tests were carried out for polymer melts held between two glass plates. It can be seen that the deformation ratios of filled EVA is much higher than that of the unfilled polymer as the experiment proceeds toward rupture of the drawn filament. The initial deformation ratios of the nanocomposites tested were nearly identical to that of the unfilled material as at this stage the drawing process was just starting. The diffraction patterns obtained here were nearly circular for all materials studied, suggesting that the only form of deformation here originated in the die. As the experiment proceeded, the deformabilities of the two filled systems were almost identical to each other, but higher than that of the unfilled material. The higher extent of deformability is possibly due to increased particle orientation, leading to increase in the degree of filler-filler interactions, enabling the fillers to withstand higher tensile stresses. The scattering patterns exhibited nearly circular shapes at the start of extension to extremely elongated elliptical shapes at higher drawing velocities (Figs. 15a-15d).

[FIGURE 12 OMITTED]

CONCLUSION

Melt-blended EVA (9% VA) and Cloisite[R] 15A analyzed using XRD and ESEM were found to be predominantly intercalated. Shear and extensional rheological tests indicated that polymer chains intercalated into the silicate layers and the intercalated particles served as highly compatible fillers. Laser light scattering studies demonstrated that the filled EVA had higher deformability than the unfilled EVA under extensional deformation.

[FIGURE 13 OMITTED]

[FIGURE 14 OMITTED]

[FIGURE 15 OMITTED]

NOMENCLATURE
G' Storage Modulus
G" Loss Modulus
I(s) Scattering Intensity
I(O) Intensity of Primary Beam
[M.sub.w] Average Molecular Weight
[M.sub.n] Number-Average Molecular Weight
N1 First Normal Stress Difference
[R.sub.g] Radius of Gyration
[v.sub.0] Velocity at Die Exit
[v.sub.w] Velocity at Nip Rollers

Greek Symbols
[alpha] Slope of G' vs. Frequency curve in the Low-Frequency Region
[gamma] Wavelength
[theta] Scattering Angle
[omega] Frequency


[c] 2004 Society of Plastics Engineers

Published online in Wiley InterScience (www.interscience.wiley.com).

DOI: 10.1002/pen.20117

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RANJIT PRASAD, VANJA PASANOVIC-ZUJO, RAHUL K. GUPTA, FERENC CSER, and SATI N. BHATTACHARYA*

Rheology and Materials Processing Centre

School of Civil and Chemical Engineering

RMIT University

Melbourne, Australia

* To whom correspondence should be addressed.

E-mail: Satinath.Bhattacharya@rmit.edu.au
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Author:Prasad, Ranjit; Pasanovic-Zujo, Vanja; Gupta, Rahul K.; Cser, Ferenc; Bhattacharya, Sati N.
Publication:Polymer Engineering and Science
Date:Jul 1, 2004
Words:6270
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