Clay

Article pubs.acs.org/Macromolecules

Shear-Induced Orientation in Well-Exfoliated Polystyrene/Clay Nanocomposites Laura M. C. Dykes,† John M. Torkelson,†,‡ and Wesley R. Burghardt*,†,‡ †

Department of Chemical and Biological Engineering and ‡Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois 60208, United States ABSTRACT: We report measurements of shear-induced particle orientation in highly exfoliated polystyrene/clay nanocomposites. Samples were prepared using an in situ polymerization technique, in which native clay is organically modified with a cationic surfactant that incoporporates a polymerizable vinylbenzyl moiety. Controlled radical polymerization was used during the synthesis to limit the molecular weight and polydispersity of the nanocomposite polymer matrix. Flowinduced orientation was measured in the flow-gradient plane of shear flow using synchrotron-based small-angle X-ray scattering. Despite the small rotational diffusivity expected for the clay particles, significant particle orientation was only observed at relatively high rates in steady shear or at high frequencies in large-amplitude oscillatory shear. Measurements of orientation upon flow cessation provided direct evidence of a structural relaxation process that was orders of magnitude faster than estimates of rotational Brownian motion. It is suggested that this fast relaxation arises from either relaxation of shear-induced distortion of partially flexible exfoliated clay sheets in a highly entangled nanoparticle network or coupling of particle and polymer dynamics.

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state.11−15 To better anticipate how processing will impact particle orientation in nanocomposites, it is desirable to develop detailed understanding of the impact of flow on nanocomposite structure through in situ experiments in well-defined flows. We have recently reported detailed studies of shear flow-induced particle orientation in two different dispersions of intercalated clay particles in low-molecular-weight polymer melts that are viscous liquids at room temperature, using synchrotron-based X-ray scattering methods.16,17 In both cases, many aspects of the observed orientation dynamics could be rationalized by appealing to concepts from the hydrodynamics of anisotropic (disk-like) Brownian particles in shear, although the propensity of clay particles to form particulate network structures also had a large impact on both particle orientation and the macroscopic rheology of these samples. Given the strong technological interest in exfoliated polymer/clay nanocomposites, it is valuable to extend the scope of such in situ structural studies beyond the case of intercalated clay dispersions. Here we present studies of shearinduced orientation in highly exfoliated polystyrene (PS)/clay nanocomposites. We again employ X-ray scattering as a tool to probe particular orientation under flow. Although anisotropic wide-angle X-ray diffraction from the periodic layering within intercalated clay particles provides a direct representation of the particle orientation distribution function,13,16,17 this option does not exist for exfoliated samples. However, the highly anisotropic shape of disk-like clay particles also induces

INTRODUCTION Polymer nanocomposites offer opportunities to engineer polymer materials with enhanced properties through introduction of modest quantities of nanometer-scale fillers. Improvements in heat distortion temperature, stiffness, flame resistance, and barrier properties have all been demonstrated in polymer/ clay nanocomposites.1−3 Use of other types of nanoparticles such as carbon nanotubes4 or graphene sheets5 can further expand the palette of property enhancement by imparting high electrical and/or thermal conductivity to polymers. In preparing nanocomposites, it is frequently a challenge to develop processing methods that facilitate thorough dispersion of the nanoparticles in the matrix polymer. In the case of polymer/ clay nanocomposites, complete exfoliation of individual clay sheets is usually the most desired outcome, since high exfoliation maximizes the nanoparticle/polymer interfacial area. Frequently, however, conventional processing techniques like melt mixing lead to intercalated samples in which polymer penetrates into the intergallery spaces of organically modified clay, expanding but not destroying their characteristic layered structure. Many classes of polymer nanocomposites employ highly anisotropic (rod-like or disk-like) particles. Particle orientation induced during processing of nanocomposites will have a strong impact on mechanical,6 barrier,7 and transport properties.4 Flow-induced particle orientation is also frequently invoked as a possible explanation of rheological phenomena in nanocomposites.8−10 While there is ample evidence of flow- and processing-induced orientation in polymer/clay nanocomposites, most studies have relied on ex situ structural studies of nanocomposites that had been quenched into the solid © 2012 American Chemical Society

Received: June 6, 2011 Revised: December 21, 2011 Published: January 18, 2012 1622

dx.doi.org/10.1021/ma2012738 | Macromolecules 2012, 45, 1622−1630

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were purchased from VWR. Natural montmorillonite, trade name Cloisite NA+, was obtained from Southern Clay Products. Synthesis. Synthesis of surfactant dimethyhexadecylstyrylammonium chloride (VB16) and its subsequent use in clay surface modification follow procedures described by Zhu et al.28 N,NDimethyl-n-hexadecylamine (13.64 g) and 4-vinylbenzyl chloride (7.82 g) were mixed in equal mole ratio and stirred overnight in 50 mL of ethyl acetate. The resulting white precipitate was recrystallized using ethyl acetate, yielding 11.84 g of VB16, as confirmed by 1H NMR. Ion exchange was used to modify the clay surface with the VB16 surfactant. 25 g of unmodified montmorillonite was stirred overnight in distilled water. 100 mL of a 27 mM solution of VB16 in distilled water was added dropwise to the clay suspension. The mixture was stirred for over 3 h and then filtered and vacuum-dried. Nanocomposites of PS and VB16-modifed montmorillonite clay were synthesized via in situ polymerization. While Zhu et al.28 used conventional free-radical polymerization, here we employed nitroxidemediated controlled free-radical polymerization in order to produce samples with narrower polymer molecular weight distribution and to control the molecular weight to lower levels in order to avoid excessive viscoelasticity in the polymer melt. Polymerizations were initiated via thermal decomposition of benzoyl peroxide (BPO) in the bulk, while the concentration of propagating radicals was maintained at a low level through introduction of a stable nitroxide radical (here TEMPO).30 110 mL of deinhibited styrene was added to a 150 mL round-bottom flask, together with TEMPO (0.4 mol % in styrene) and BPO (TEMPO:BPO mole ratio = 3:1). VB16-modified montmorillonite was added at varying weight fractions. The flask was sealed under N2 and stirred for at least 1 h. After three N2/vacuum cycles, the flask was placed in an oil bath at 125 °C. The mixture was allowed to react for 72 h, resulting in a solid product that was recovered by breaking the flask. After breaking the solid polymer product into pieces, it was placed in a vacuum oven at 100 °C overnight to remove unreacted monomer. Characterization. The clay content of the final products was determined using thermogravimetric analysis on a Mettler Toledo TGA. Samples were heated from 25 to 650 at 10 °C/min. The majority of weight loss occurred at temperatures between 350 and 450 °C. The clay content was determined from the residual mass remaining at 650 °C. The molecular weight of the polystyrene in the nanocomposite samples was determined using gel permeation chromatography (GPC). To prevent damage to the GPC columns, the clay was first removed from the samples prior to testing using a clay stripping procedure described by Xu et al.29 Each nanocomposite sample was refluxed for 3 h in a 1:3 volume ratio mixture of methanol and THF with 2.5 wt % LiBr. After centrifugation, the supernatant liquid was filtered with a 0.2 μm filter and the PS was precipitated in methanol. TGA was used to confirm removal of the clay. GPC was performed in tetrahydrofuran (THF) using a Waters 810 calibrated with polystyrene standards. Polystyrene/THF solutions were filtered a second time with a 0.2 μm filter prior to analysis. Onset glass transition temperatures were determined using a Mettler Toledo Star DSC 822. Measurements were taken upon a second heating cycle, during a temperature ramp from 25 to 160 °C at a rate of 10 °C/min. Table 1 lists composition and PS number-average molecular weight (Mn), polydispersity (Mw/Mn), and glass transition temperature (Tg) for the samples studied here. The molecular weight of the PS matrix is fairly consistent across the samples, although the 0.25 wt % clay sample has a markedly lower Mn and, correspondingly, a lower Tg. Transmission electron microscopy (TEM) was performed using a JEOL 1230 microscope operating at 80 kV. Samples of 70 nm thickness were sectioned using a Leica Ultracut S ultramicrotome. TEM images show a well-exfoliated structure consisting mostly of individually dispersed clay layers, although layered “tactoid” structures are occasionally observed (Figure 1). Rheological characterization was performed using a Rheometrics Scientific ARES controlled-strain rheometer. Dynamic shear testing was performed using 25 mm parallel plate fixtures, after first

anisotropy in the small-angle scattering arising from the particle as a whole, which provides an alternative route to quantify flowinduced orientation. Our earlier work included comparison between wide- and small-angle scattering as probes of particle orientation, and the methods gave generally consistent results, albeit with some quantitative differences ascribed to the difference in length scale probed in the different scattering regimes.17 In the present study, we rely on anisotropic smallangle scattering as a probe of shear-induced orientation in exfoliated nanocomposites. Most methods for producing well-exfoliated polystyrene/clay nanocomposites have employed in situ polymerization. Two general approaches have been used. In the first, an initiator is tethered to the clay surface, such that all growing polymer chains are tethered to the clay surface. This approach was first employed by Weimer et al.,18 who used a nitroxide initiator such that the reaction proceeded via a controlled radical polymerization mechanism. Fan et al.19,20 used thermal free radical initiators based on azobis(isobutyronitrile), to which quaternized amine groups were added to facilitate attachment to clay surfaces via ion exchange. Uthirakumar et al.21−23 similarly employed azo-based cationic initiators in their studies of tethered radical initiators. In an alternative in situ polymerization approach, polymerizable surfactants are attached to the clay surface using ion exchange, providing a different route to allow direct coupling of growing polystyrene chains to the clay surface. In this strategy, initiation occurs in the bulk. Many such studies have used polymerizable surfactants employing a vinylbenzyl moiety. Tseng et al.24 explored modifying clay surfaces with combinations of the surfactant vinylbenzyldimethylethanolammonium chloride (VBDEAC) and a long chain alkyl surfactant, cetylpyridinium chloride (CPC). Use of VBDEAC alone did not lead to wellexfoliated nanocomposites. Tseng et al. concluded that adding CPC with its longer chain helped facilitate intercalation of styrene monomer into the organically modified clay and reported optimal results at a VBDEAC:CPC ratio of 1:1.24 A similar two-surfactant strategy was employed by Zhong et al.,25 in which the polymerizable surfactant was based on an acrylic group. Alternatively, Fu et al.26,27 and Zhu et al.28 employed a single surfactant incorporating both a vinylbenzyl moiety and a long chain alkyl group; both groups had success in producing well-exfoliated nanocomposites using these polymerizable surfactants in conjunction with bulk-initiated free radical polymerization. The surfactant used by Zhu et al., dimethylhexadecylstyrylammonium chloride (denoted VB16), was subsequently used in preparation of exfoliated nanocomposites for rheological studies.29 In this work we have adopted the synthesis procedures of Zhu et al.28 and Xu et al.,29 employing the polymerizable surfactant VB16 to facilitate preparation of well-exfoliated polystyrene/clay nanocomposites. Since our ultimate goal is to study orientation dynamics in shear flow, we have deviated from their procedures by using controlled radical polymerization to limit (and better define) the degree of polymerization achieved, in order to avoid excessive “background” elasticity from the polymer melt matrix.

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EXPERIMENTAL SECTION

Materials. Reactants N,N-dimethyl-n-hexadecylamine, 4-vinylbenzyl chloride, benzoyl peroxide, 2,2,6,6-tetramethylpeperidine 1-oxyl (or TEMPO), and styrene were purchased from Sigma-Aldrich. Solvents 1623


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elasticity from the nanocomposite matrix during shear experiments. The zero shear viscosity of the pure PS sample is 700 Pa·s at 160 °C. Using a clay platelet diameter d = 200 nm estimated from TEM images, the rotational diffusivity of an isolated particle is estimated as Dr = 3kBT/4ηd3 = 8 × 10−4 s−1. Consistent with the higher viscosity of this PS melt matrix, this estimated rotational diffusivity is substantially smaller than was the case in recent room-temperature investigations of orientation in intercalated clay dispersions in oligomeric matrices.16,17 Addition of clay leads to enhanced elastic character in the nanocomposite samples (Figure 3). Since each sample resulted from

Table 1. Characterization of PS/Clay Nanocomposite Samples

a

clay content (wt %)a

Mn (×103)b

Mw/Mnb

Tg (°C)c

0 0.25 1.2 2.6 4.4

24 16 26 20 23

1.28 1.55 1.16 1.49 1.39

93 83 92 91 87

Determined via TGA. bDetermined via GPC. cDetermined via DSC.

Figure 1. Transmission electron micrograph of 2.6 wt % PS/clay nanocomposite.

Figure 3. (a) Storage modulus and (b) loss modulus measured in PS/ clay nanocomposites at a temperature of 160 °C: pure PS (●); 1.4 (○), 2.6 (▲), and 4.4 wt % clay (△) nanocomposites. (Data on 0.25 wt % clay nanocomposite not presented here, since the lower polymer molecular weight in this sample substantially reduces the matrix viscosity.)

conducting oscillatory strain sweep experiments to determine the extent of the linear viscoelastic regime. The pure PS sample exhibits typical melt viscoelasticity for a marginally entangled polymer of relatively narrow molecular weight distribution (Figure 2). The well-

its own polymerization, data in Figure 3 reflect variability in polymer molecular weight distribution as well as clay loading. However, the progressive increase in storage modulus with clay content, leading to a low-frequency solid-like plateau, is typical of behavior in polymer/clay nanocomposites32−35 and is generally attributed to the development of a percolated nanoparticle network. (Data for the 0.25 wt % sample are not included here, since the combined effects of low Mn and Tg in this sample lead to a significantly lower viscosity than in the pure PS sample.36) The 2.6 wt % sample was subjected to more extensive rheological testing using a Paar Physica UDS rheometer equipped with an electrically heated plate temperature controller. Experiments were performed using cone and plate fixtures with a cone angle of 4°. Results of these tests are presented below. X-ray Scattering. Flow-induced changes in PS/clay nanocomposite structure were studied using an annular cone and plate shear cell designed to facilitate small-angle X-ray scattering studies of fluid structure in the flow-gradient (1−2) plane of shear flow.37 Experiments were conducted at 160 °C under an inert helium atmosphere. Measurements were performed at beamline 5ID-D (DND-CAT) of the Advanced Photon Source at Argonne National Laboratory, using a sample−detector distance of 0.9 m and X-ray energy of 15 keV. Two-dimensional X-ray scattering patterns were collected using a Gemstar CCD detector which allowed data acquisition at rates up to 7 frames/s; as will be seen below, this relatively high rate of data acquisition was necessary to resolve fast transient processes in the orientation dynamics of these exfoliated PS/

Figure 2. Linear viscoelasticity of PS homopolymer synthesized via controlled radical polymerization. Storage modulus (unfilled symbols) and loss modulus (filled symbols) data were measured at temperatures of 140 (◇, ◆), 150 (△, ▲), 160 (□, ■) and 170 °C (○,●) and shifted to a reference temperature of 160 °C. defined terminal regime in this polymer allows extraction of an average relaxation time from the low-frequency limit of G′/ωG″.31 At a temperature of 160 °C (where melt X-ray scattering measurements were performed), the relaxation time is ∼0.004 s. The fast relaxation of the polymer reduces complications associated with background 1624


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clay nanocomposite melts. Samples were first compression-molded at 160 °C into donut-shaped pieces, sized appropriately to fill the gap between the annular shear cell fixtures. Care was taken to minimize sample deformation during loading of the shear cell, but some deformation is inevitable as the fixtures are squeezed together. Given the possibility of incomplete structural relaxation upon cessation of shear flow, testing was performed in order of increasing shear rate (unidirectional shear) or frequency/strain (large-amplitude oscillatory shear). Two-dimensional (2D) X-ray scattering patterns demonstrate a lack of a diffraction peak associated with clay layering (Figure 4).

anisotropy factor =

4⟨qxqy⟩2 + ⟨qxqx − qyqy⟩2

(2)

Defined in this way, and given the normalization used in eq 1, the anisotropy factor ranges from 0 for an isotropic scattering pattern to 1 for the hypothetical case in which all scattering is concentrated along a single direction within the 2D scattering pattern. The average particle orientation angle in the x−y plane, χ, is extracted from the principal directions of ⟨qq⟩:

χ=

⎛ 2⟨q q ⟩ ⎞ x y 1 ⎟ tan−1⎜ ⎜ ⎟ 2 q q q q ⟨ − ⟩ y y ⎠ ⎝ x x

(3)

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RESULTS Enhanced low-frequency elasticity is polymer/clay nanocomposites is generally attributed to a percolated nanoparticle network, which may be disrupted by application of shear flow. Large-amplitude oscillatory shear (LAOS) is seen to strongly influence the elastic properties of the 2.6 wt % sample (Figure 5a). During a brief interval of LAOS with 100% strain, the storage modulus is rapidly and dramatically suppressed. Much of this reflects the intrinsically nonlinear response at large strains; data collected during this interval must be understood as “effective” storage modulus. However, upon returning to small strains in the linear viscoelastic regime, G′ is much smaller than it was prior to the LAOS interval. The storage modulus gradually recovers over an extended period of time (Figure 5a) but remains well below its value prior to application of LAOS. Such gradual recovery of elasticity is frequently observed in polymer/clay systems.9,16,39 Interestingly, a period of steady shear at a low rate (>30 strain units at 0.02 s−1) leads to “rejuvenation” of the network structure (Figure 5b); the lowfrequency storage modulus following steady shear at 0.02 s−1 is an order of magnitude larger than after LAOS. Shear stress overshoots upon flow inception are often used as another probe of elastic networks in nanocomposites.10,16,34,39 The relative strength of elastic networks present in this 2.6 wt % sample following LAOS and steady shearing is also manifested in the transient shear stress response observed upon subsequent flow inception (Figure 5c). Following LAOS (a condition in which elasticity is considerably suppressed), the stress growth upon inception of shear flow is monotonic. Conversely, following steady shear flow at 0.02 s−1 (a condition with more pronounced elasticity), there is a pronounced stress overshoot upon inception of shear. Linear viscoelastic data presented earlier in Figure 3 were all collected after first performing strain sweep experiments up to strain amplitudes of 100%. These data, then, are affected by elastic network disruption associated with LAOS. As an example, the results for the 2.6 wt % sample in Figure 3 agree closely with the post-LAOS data in Figure 5b. As a result, data in Figure 3 under-represent the true degree of elasticity present in these well-exfoliated nanocomposite samples. However, they still convey the qualitative changes associated with clay loading, since an identical shear history was used for all samples. While shear flow at low rates leaves a robust elastic network intact, increasing prior steady shear rate does lead to a progressive suppression of elasticity measured using oscillatory shear after flow cessation (Figure 6a). Such systematic decreases in elasticity with increasing prior shear rate are commonly observed in nanocomposites17,39,40 and attributed to

Figure 4. Two-dimensional X-ray scattering patterns measured in 2.6 wt % PS/clay nanocomposite during steady shear flow in (a) forward and (b) reverse direction at a shear rate of 6.0 s−1 at 160 °C. Axes used for analysis of shear-induced anisotropy are defined in part (a). The patterns cover a scattering vector range |qx|, |qy| ≤ 0.25 Å−1. Conventional wide-angle X-ray diffraction measurements with Cu Kα radiation showed diffraction peaks at scattering angle 2θ = 8.6° and 4° for respectively the native Cloisite Na+ and VB16-modified montmorillonite clays.36 The latter feature would fall within the qrange probed in these 2D patterns; the absence of a discernible peak provides evidence of a well-exfoliated structure, consistent with previous reports using in situ polymerization of PS in the presence of VB16-modified clay28,29 and with the TEM evidence presented in Figure 1. For intercalated nanocomposites, orientation may be characterized using anisotropy in either the wide-angle diffraction peak associated with clay layering or the small-angle scattering arising from the clay particle as a whole.17 In well-exfoliated samples, the former option is removed; however, two-dimensional small-angle scattering patterns in the PS/clay samples are clearly rendered anisotropic by application of shear flow (Figure 4). To characterize shear-induced orientation, we compute a second moment tensor of the scattered intensity distribution:

⎡⟨q q ⟩ ⟨q q ⟩⎤ ∫ qqI(q) dq dq x y ⎥ x y ⎢ x x ⟨qq⟩ = ⎢ = ⎥ 2 ⟨q q ⟩ ⟨qyqy⟩ ∫ q I(q) dqx dqy ⎣ x y ⎦

(1)

where “x” denotes the flow direction and “y” denotes the velocity gradient direction. The denominator on the right-hand side of eq 1 is essentially the scattering invariant.38 Normalization of ⟨qq⟩ in this way accounts for frame-to-frame variability in X-ray beam intensity and also allows direct comparisons across samples with varying scattering power. For shear flow as depicted in Figure 4a, the “shear” component, ⟨qxqy⟩, is negative due to the concentration of scattered intensity in the second and fourth quadrants. In presenting results below, we plot the positive quantity −⟨qxqy⟩. Similarly, anisotropy in the normal components of ⟨qq⟩ is characterized by the quantity −⟨qxqx − qyqy⟩; the higher intensity observed along the gradient (y) axis is consistent with preferred alignment of thin disks along the flow direction. These two quantities are analogous to the shear stress and first normal stress difference in shear rheometry. An alternative representation of shear-induced orientation may be obtained from the difference in the principal values of ⟨qq⟩: 1625


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Figure 6. Effect of steady shear rate on 2.6 wt % sample rheology. (a) Effect of prior steady shear rate on storage modulus measured 150 s following flow cessation, for shear rates of 0.02 (▲), 0.1 (◇), 0.4 (■), and 2.0 s−1 (○). (b) Steady shear viscosity η(γ̇) (●). Also plotted is the complex viscosity |η*(ω)| measured 150 s following cessation of shear flow at 0.02 (△) and 2 s−1 (○).

Figure 5. Impact of flow history on elastic network structure of 2.6 wt % PS/clay nanocomposite. (a) Storage modulus measured as a function of time at ω = 5 rad/s and strains of 1.0 (●) and 100% (○). (b) Storage (filled) and loss (open) moduli measured 4000 s after application of LAOS (5 rad/s, 100% strain; ○, ●) and 150 s after application of steady shear (0.02 s−1; △, ▲). (c) Shear stress measured as a function of strain after inception of shear flow at 0.02 and 0.04 s−1, with shear history indicated in parentheses.

progressive degradation of nanoparticle networks. Steady viscosity shows shear thinning at low rates (Figure 6b), another common manifestation of particle networks in nanocomposites39−41 (steady data were obtained after preshearing for at least 30 strain units and using data collected in both clockwise and counterclockwise directions to guard against autozero errors associated with yield stresses). In materials where fluid structure will be a complex function of shear history, there is little reason to expect adherence to the empirical Cox−Merz rule (equivalence between steady shear viscosity η(γ̇) and complex viscosity |η*(ω)| when evaluated at γ̇ = ω31). As is frequently observed in nanocomposites,42,43 we find violation of the Cox−Merz rule here as well (Figure 6b). A final manifestation of the impact of steady shear rate on fluid structure is found in measurements of shear stress relaxation following flow cessation (Figure 7). At low shear rates, a large fraction of the steady state stress is observed to relax slowly; over 20% of the steady stress remains unrelaxed after 100 s. With increasing shear rate, both the fraction of

Figure 7. Normalized shear stress relaxation measured in 2.6 wt % sample upon cessation of steady shear flow at indicated rate.

“slowly relaxing” stress and the relaxation time decrease dramatically. These results closely resemble observations of Mobuchon et al.,39 who used data of this sort to divide stress in steady shear flow into “elastic” and “viscous” contributions. The decrease in the slowly relaxing fraction with increasing shear rate is, again, attributed to progressive degradation of nanoparticle network structure in steady shear. We turn now to measurements of clay particle orientation during shear flow. Shear rate has a significant impact on the degree of particle orientation measured in the 2.6 wt % nanocomposite sample at steady state (Figure 8). (Steady state orientation data were obtained from time-dependent results in transient shear reversal testing protocols, after at least 30 strain units had been applied.) The shear and normal components of anisotropy in the second moment tensor ⟨qq⟩ are both 1626


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Figure 9. Anisotropy factor measured as a function of shear rate at 160 °C in PS/clay nanocomposites with clay loadings of 0.25 (△), 1.4 (◇, ◆), 2.6 (○), and 4.4 wt % (■). Multiple symbols for 1.4 wt % sample reflect duplicate X-ray scattering measurements using different sample loadings.

loadings of the 1.4 wt % sample showed good consistency. In light of the qualitative similarities in steady state orientation behavior among the samples, in the remainder of this paper we concentrate attention on the 2.6 wt % sample. More extensive orientation data on the other samples are available elsewhere.36 Reversal of shear flow direction leads to reorientation of the clay particles (Figure 10). Here we adopt a convention that the

Figure 8. Shear-induced orientation measured in steady shear flow in 2.6 wt % PS/clay nanocomposite sample at 160 °C. (a) Shear (◇) and normal (◆) components of second moment tensor, defined in eq 1. (b) Anisotropy factor (◇) and orientation angle (◆), defined in eqs 2 and 3.

negligible at low rates (Figure 8a). Both −⟨qxqy⟩ and −⟨qxqx − qyqy⟩ initially increase at rates higher than ∼1 s−1. The normal anisotropy grows rapidly with increasing shear rate, while the shear component appears to pass through a maximum at high rates. These data may be recast as anisotropy factor and orientation angle using eqs 2 and 3 (Figure 8b). At low rates where the anisotropy factor is small, the orientation angle reaches an asymptote of ∼35°, indicating substantial misalignment of clay platelets away from the flow direction, although the orientation angle does not reach the low shear rate asymptote of 45° expected for systems that relax to isotropy in the absence of flow. As anisotropy increases at higher shear rates, χ steadily decreases, indicating that particles increasingly orient toward the flow direction. However, at the highest shear rates studied, particles are still substantially misaligned relative to the flow direction, with an average orientation angle of ∼12°. Experiments on samples at various clay loadings show similar orientation behavior (Figure 9), with monotonic growth in steady state anisotropy factor with increasing shear rate. (The weak decrease in anisotropy seen in some samples at low rates is attributed to unrelaxed orientation induced by squeezing flow during sample loadings. It appears that shear flow at sufficiently low rates was actually able to progressively degrade some of this residual orientation, perhaps related to the ability of shear flow at low rates to rejuvenate elastic network structures seen in Figure 5b. Relaxation data presented below provide further insights into the incomplete relaxation of particle orientation in these samples.) In general, higher anisotropy is achieved in samples with higher clay content. Experiments on the most concentrated (4.4 wt %) sample could not be extended to high rates due to flow instabilities,36 presumably associated with its significantly higher elasticity (Figure 3b). While constraints on available synchrotron beam time generally preclude extensive studies of reproducibility, repeated experiments using multiple

Figure 10. Transient orientation response in (a) shear and (b) normal components of second moment tensor measured in 2.6 wt % PS/clay nanocomposite following reversal of shear flow at rates of 1.0 (■), 2.5 (△), 4.0 (●), 6.0 (◇), and 10.0 s−1 (▼).

shear flow prior to reversal is in the negative direction. The “shear” component of the second moment tensor −⟨qxqy⟩ changes signs, reflecting the primary reorientation process apparent in the two-dimensional scattering patterns for forward and reversed shear flow (Figure 4). During this reorientation, the normal anisotropy exhibits an undershoot, indicating that there is a transient state of lower anisotropy prior to reestablishing the steady orientation state in the reversed direction. The reorientation process is quite rapid, essentially complete within ∼10 shear strain units. 1627


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Upon flow cessation there is a rapid but partial relaxation of shear-induced anisotropy (Figure 11). At higher shear rates, a

Figure 11. Transient orientation response in (a) shear and (b) normal components of second moment tensor measured in 2.6 wt % PS/clay nanocomposite following cessation of shear flow at rates of 1.0 (■), 2.5 (△), 4.0 (●), 6.0 (◇), and 10.0 s−1 (▼).

significant degree of relaxation is observed within the first several data points at times below 1 s, while a large fraction of the ultimate relaxation occurs within ∼5 s. Particularly for the normal components, −⟨qxqx − qyqy⟩, the relaxation of anisotropy is incomplete. As experiments were performed in order of increasing shear rate, there was a progressive buildup of residual orientation in the sample. Orientation dynamics in the PS/clay nanocomposites were also studied in large-amplitude oscillatory shear (LAOS), performed with different specimens than used in unidirectional shearing experiments in order to minimize the impact of residual orientation. With the fast data acquisition afforded by the bright synchrotron source, it was possible to track timeresolved orientation dynamics during the application of oscillatory shear. At fixed strain amplitude of 200%, frequency is found to have a significant effect on the qualitative character of the orientation response (Figure 12). At low frequency, the shear component of anisotropy (−⟨qxqy⟩) shows weak oscillations while the normal anisotropy (−⟨qxqx − qyqy⟩) is essentially unaffected (Figure 12a). Increasing the frequency by an order of magnitude leads to an increase in the amplitude of the oscillatory response of −⟨qxqy⟩ and also gives rise to perceptible fluctuations in −⟨qxqx − qyqy⟩ occurring at twice the imposed frequency (Figure 12b). An even higher frequency further increased the magnitude of the −⟨qxqy⟩ response and led to a clear dc offset in the normal anisotropy signal in addition to high-frequency fluctuations. (At these higher frequencies, these fluctuations could no longer be well-resolved with the available data acquisition rate.) The behavior of −⟨qxqx − qyqy⟩ at high frequencies is analogous to the response of the first normal stress difference in viscoelastic liquids subjected to

Figure 12. Time-dependent response in shear (●) and normal components (○) of second moment tensor measured in 2.6 wt % PS/ clay nanocomposite during large-amplitude oscillatory shear flow at a strain amplitude of 200% and frequencies of (a) 0.01, (b) 0.1 and (c) 1 Hz. Arrows in (c) illustrate the method used to extract amplitudes of shear and normal anisotropy components.

LAOS, which oscillates at twice the input frequency, and remains positive, oscillating relative to a dc offset value.44 To enable a more quantitative representation of the impact oscillatory shear parameters have on orientation state, we characterize LAOS-induced orientation in terms of the peak-topeak amplitude of oscillations in the shear component, −⟨qxqy⟩, and the time-averaged (dc) amplitude of the normal anisotropy, −⟨qxqx − qyqy⟩, as indicated by the arrows in Figure 12c. Mirroring the behavior seen in steady shear flow, only weak shear-induced orientation is found at low frequencies, while the amplitudes of both shear and normal components of ⟨qq⟩ increase markedly at higher frequencies (Figure 13). While LAOS-induced effects in −⟨qxqy⟩ are observable at the lowest frequencies (Figure 12a), it is likely that the low frequency asymptote in −⟨qxqx − qyqy⟩ is associated with residual anisotropy induced during sample loading. At fixed frequency, higher strain amplitude leads to higher anisotropy; the effect is larger in the normal component.

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DISCUSSION

Several features of these shear-induced orientation data in these exfoliated PS/clay nanocomposites resemble our prior observations for intercalated clay dispersions in oligomeric, room-temperature melts.16,17 First, much of the impact of shear 1628


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While the resistance of these exfoliated PS/clay nanocomposites to shear-induced particle orientation is puzzling, it is important to note that there is strong internal consistency among the results presented here. The lack of significant flowinduced anisotropy at low rates and/or frequencies implies that there must be some fast structural relaxation process, and, indeed, the relaxation data show evidence of some relaxation mechanism that is orders of magnitude faster than that expected for rotational Brownian motion. Under these conditions, the structural state of the nanocomposite melt at low shear rates will not differ appreciably from that found at equilibrium. If Brownian motion fails to explain the fast relaxation process that underlies many of the present observations, what is its origin? In considering this question, we first note that envisioning exfoliated clays to be flat, rigid platelets is an oversimplification. Electron micrographs of exfoliated nanocomposites show ample evidence of considerable flexibility of the clay sheets.23,26−28 In a well-exfoliated sample, the high aspect ratio of the clay sheets further means that they will be extremely crowded, a concept that is commonly invoked to relate low-frequency elasticity in polymer/clay nanocomposites to their degree of exfoliation.33 We hypothesize that the orientation dynamics observed here may be dominated by shear-induced distortions within a highly crowded network of semiflexible clay sheets. That is, the dominant balance governing the development of orientation is not hydrodynamic vs Brownian forces, but rather hydrodynamic forces vs elastic deformation within the nanoparticle network. This line of reasoning mirrors arguments we have advanced as a possible explanation of unexpected relaxation observed in non-Brownian multiwalled carbon nanotube dispersions.40,46 While consistent with the available evidence, this hypothesized mechanism is not the only possible origin of fast dynamics. The synthesis procedure used here should result in grafting of polymer chains to the clay particles; it is possible that this could couple the particle orientation dynamics more directly to the dynamics in the polymer matrix, providing an alternative mechanism for a fast relaxation process. To investigate this further, it would be beneficial to study shear-induced orientation dynamics in wellexfoliated polymer/clay nanocomposites that do not involve chemical tethering of polymer to the nanoparticles. While the unexpected rapid relaxation process discovered here plays a major role in determining the nature of flowinduced particle orientation in these samples, other factors also come into play. The fact that only partial relaxation of orientation is observed upon flow cessation mirrors observations in model intercalated clay dispersions16,17 and probably reflects the complexities of orientation dynamics within a highly crowded environment. This “hindered” relaxation may also explain the failure of the orientation angle to reach the low rate asymptote of 45° that is expected in the weak flow limit of fluids with relaxing microstructure in shear flow.

Figure 13. Frequency dependence of (a) shear and (b) normal anisotropy measured in 2.6 wt % PS/clay nanocomposite during largeamplitude oscillatory shear flow at strains of 100% (△) and 200% (●).

flow on elastic properties seems to be largely independent of particle orientation per se. For instance, Figures 6 and 7 show large shear-rate-dependent changes in rheological properties at rates below 2 s−1, a regime in which shear-induced particle orientation is only very weak. There are also gross similarities in many aspects of the particle orientation response: (i) increase in anisotropy and decrease in orientation angle with steady shear rate; (ii) rapid reorientation upon shear flow reversal; (iii) partial relaxation of flow-induced orientation upon flow cessation; and (iv) strain dependence of shear-induced orientation in large-amplitude oscillatory shear. Despite these similarities to earlier work, there is also a remarkable discrepancy. In the model oligomeric dispersions, many aspects of the orientation response could be qualitatively interpreted using concepts of orientation dynamics anticipated in weakly Brownian particulate dispersions. In those systems, orientation relaxed slowly, on time scales consistent with estimates of rotational diffusivities.16,17 Similarly, the progressive increase in orientation with increasing shear rate could be rationalized in terms of increasing Peclet number (Pe = γ̇/ Dr), as hydrodynamic forces increasingly overwhelm the randomizing influence of Brownian motion. Translating these concepts to the present study, we first consider the hydrodynamic effects responsible for promoting particle orientation. The common expectation is that exfoliated clay platelets, with their higher aspect ratio, should be more susceptible to orient in shear than lower aspect ratio clay tactoids, owing to the characteristics of the tumbling “Jeffrey orbits” exhibited by disklike particles in shear flow.45 Similarly, the higher matrix viscosity in the current studyand correspondingly smaller rotational diffusivityshould make it easier for the applied flow to promote high orientation. For example, even at a shear rate of 0.1 s−1 (the lowest used in the X-ray orientations studies), the estimated Peclet number is 125; under all conditions studied here, Brownian effects should be minimal. Why, then, is it so difficult to obtain significant flow-induced orientation?

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CONCLUSIONS The exfoliated PS/clay nanocomposites studied here are unexpectedly resistant to shear flow-induced orientation. Despite the small magnitude of the estimated rotational diffusivity of a single clay platelet, significant orientation in steady shear flow only develops at shear rates greater than 1 s−1. Similarly, flow-induced orientation in large-amplitude oscillatory shear only develops at frequencies greater than 0.1 Hz. The inability of shear at low rates and frequencies to produce 1629


Macromolecules

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significant alignment suggests the existence of a structural relaxation process that is much faster than that of rotational Brownian motion of the clay particles. Direct evidence of such a fast process is found in measurements of orientation relaxation following cessation of shear flow, although the samples do not revert to a randomly oriented state as shearing at progressively higher rates leads to higher degrees of residual orientation. It is hypothesized that this fast relaxation arises from either relaxation of shear-induced distortion of partially flexible exfoliated clay sheets in a highly entangled nanoparticle network or a direct coupling of particle and polymer dynamics through the chemical tethering of particle chains that occurs during the in situ polymerization synthesis of these materials.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected].

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ACKNOWLEDGMENTS This work was funded by the NSF-MRSEC program (DMR0520513) at the Materials Research Center of Northwestern University. L. Dykes acknowledges support from a GEM Fellowship. We acknowledge Dr. Charlene Wilke of Northwestern’s Biological Imaging Facility for performing TEM imaging and thank Professor David Venerus for providing access to the UDS rheometer. Finally, we thank the staff of the DuPont−Northwestern−Dow Collaborative Access Team (DND-CAT) at Sector 5 of the Advanced Photon Source for their help with setup and execution of synchrotron experiments. DND-CAT is supported by the E.I. DuPont de Nemours & Co., the Dow Chemical Company, and the National Science Foundation through Grant DMR-9304725 and the State of Illinois through the Department of Commerce and the Board of Higher Education Grant IBHE HECA NWU 96. Use of the Advanced Photon Source was supported by the U.S. Department of Energy, Basic Energy Sciences, Office of Energy Research, under Contract W-31-102-Eng-38.

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