Understanding the Consequences of Intercalation Using Model

Oct 12, 2015 - Department of Plastics Engineering, University of Massachusetts Lowell, 1 University Avenue, Lowell, Massachusetts 01854, United States...
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Understanding the Consequences of Intercalation Using Model Polymer Nanolaminates Erik Dunkerley and Daniel F. Schmidt* Department of Plastics Engineering, University of Massachusetts Lowell, 1 University Avenue, Lowell, Massachusetts 01854, United States S Supporting Information *

ABSTRACT: In spite of significant interest in 2D nanofillers, the effect of polymer intercalation on properties remains poorly understood. We use automated spray deposition to form two families of well-aligned polymer/clay nanolaminates that differ by one methyl group in the clay modifier; this difference results in intercalation in one instance and no intercalation in the other. Dynamic mechanical properties and gas permeability are reported alongside micromechanical and barrier models and a detailed description of the structures and phases present. Observed variations in modulus cannot be explained by micromechanical models often applied to polymer nanocomposites. In contrast with “effective particle” arguments, barrier properties are dominated by single layer aspect ratios. These findings have relevance for all manner of “bricks and mortar” hybrids. Additionally, these systems represent the macroscopic equivalent of the multilayer stacks found in practically all nanocomposites, with wide-ranging implications for nanocomposite design and properties more generally.

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INTRODUCTION In filled polymers in general and nanocomposites in particular, the importance of dispersion is often emphasized.1−3 Consequently, ample research has concentrated on achieving high levels of dispersion during processing,4−11 with intercalation a necessary prerequisite to exfoliation in clay nanocomposites. This is achieved by screening the attractive interactions between adjacent silicate layers, thus allowing them to separate. In the case of sodium montmorillonite (Na+MMT), for instance, the introduction of sufficient water results in what is typically referred to as an exfoliated state.12 Likewise, exchanging the Na+ cations for certain alkylammonium cations13 enables intercalation and sometimes exfoliation in the presence of organic species such as solvents and polymers that would otherwise be unable to penetrate the interlayer spaces.14,15 The ideal modifier should shield the hydrophilic silicate surfaces from one another and interact favorably with the organic species in question. Given some level of compatibility between the organic species and the silicate layer surface, the ideal modifier will induce rapid intercalation of organic species into the gallery spaces via a diffusive mechanism, regardless of system viscosity.16 In contrast, an ineffective modifier will create conditions that disfavor intercalation and subsequent dispersion regardless of the amount of mixing energy applied. Even when high levels of dispersion are possible, exfoliation of the clay platelets is usually incomplete, resulting in some concentration of well-dispersed stacks or tactoids of varying sizes. Whether these stacks contain intercalated polymer or not, when describing materials properties they are often treated as “effective particles” of a © XXXX American Chemical Society

larger size and with a lower aspect ratio. While this treatment is commonly accepted in the field of polymer/clay nanocomposites, the effect of intercalation alone, in the absence of significant changes in dispersion state or composition, remains unknown, as do the properties of the aforementioned multilayer stacks. Gaining such knowledge is critical if these scientifically and commercially important materials are to be better understood, modeled, and controlled. We describe and analyze the behavior of two families of clay nanocomposites spanning a wide range of well-defined compositions and structures, materials effectively identical but for a subtle change in chemistry that controls intercalation state. Previous work has demonstrated that the occurrence of intercalation is strongly affected by the solubility parameter of the modifier tails,17 the length of the modifier tails,18,19 and the density of the modifiers on the platelet surfaces.20,21 Unique to this work, the parameter controlling intercalation is not the number or length of long hydrocarbon tails present in the modifier or the grafting density of modifiers on the layer surface, all of which are constant. Instead, the determining factor is the presence or absence of a single methyl substituent on the ammonium headgroup of the clay modifier. When a dimethylditallowammonium (DMDT) modifier is used, polystyrene (PS) intercalation occurs, and a well-studied nanocomposite system emerges.1,16,22−30 When the exact same montmorillonite is modified by methylditallowammonium (MDT) cations, however, we observe no intercalation of PS. Received: August 13, 2015

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orientation parameters were calculated in Microsoft Excel based on azimuthal scans of the second-order diffraction peak.34 To ensure consistency, all radial integration data were normalized to the minimum observed intensity prior to the first peak before peak area and full width at half-maximum values were calculated. Scanning Electron Microscopy (SEM). Cross-sectional scanning electron micrographs were taken to ascertain film thickness and composite morphology. Films were fractured to expose their cross sections and subsequently mounted vertically, fractured cross-section up. The film cross sections were observed in a high-resolution field emission scanning electron microscope (JEOL JSM-7401F). All samples were coated with gold−palladium using a desktop gold sputter coater (Denton Vacuum Desk IV) and imaged with an accelerating voltage of 2.0−5.0 kV. Density. Film density was measured via an analytical balance (Denver Instrument SI-124) equipped with a density determination kit. Because of the limited amount of material available, one measurement was performed per composite sample. The submerged medium was distilled water, and the procedure was in accordance with ASTM D792. Dynamic Mechanical Analysis (DMA). DMA was performed on each composition using a TA Instruments Q800 dynamic mechanical analyzer. All tests used a standard tension fixture for films. All samples were cut to a width of ∼4 mm and clamped at a length of ∼8 mm. Sample widths were measured using digital calipers with a resolution of 0.01 mm. Sample thickness was measured using a digital dial micrometer with a resolution of 0.001 mm. A preload stress of 250 kPa was used in all cases, and all samples were conditioned at −115 °C for 12 min prior to the start of the test to ensure that the desired start temperature was allowed to stabilize. The samples were then heated from −115 to 120 °C at a rate of 3 °C/min while testing in multifrequency, controlled strain mode using an amplitude of 10 μm and a frequency of 1.0 Hz. All peak deconvolution was performed using the peak analyzer tool in OriginLab Origin Pro 8.x. The peak analyzer was run in Fit Peaks (Pro) mode, with a baseline manually defined by selecting a minimum of 3−5 points along the curve on either side of the transition. Rough peak positions were then inputted, using the minimum number of peaks needed to achieve a high coefficient of determination of the final fit. The thermal transition temperatures of the polymer and modifier were determined in this fashion via deconvolution of the strongest loss modulus and storage modulus derivative peaks (derivative taken with respect to temperature) using a Gaussian model. In the case of the loss modulus peaks in particular, Gaussian and asymmetric double sigmoidal fits were used for samples containing less than 40 vol % modified clay due to the asymmetric shape of the loss modulus peak associated with the neat polymer. Along the same lines, an asymmetric double sigmodial fit was also used for analysis of the storage modulus derivative peak observed for the pure polymer film, consistent with prior work supporting this choice.35 Finally, a combination of constant, linear, logistic, and asymmetric double sigmoidal functions was used to fit trends in transition temperature, full width at half-maximum values, and peak area data as a function of composition. In the case of the perturbed layer thickness calculations in particular, the equations derived from logistic fits of the bulk PS transition peak areas were used to calculate the compositions at which the relative strength of these transitions was reduced by 95%. The Levenberg−Marquardt algorithm was used to optimize the fit parameters in all cases, and all fits involving data with error bars were error-weighted. Oxygen Permeation Analysis (OPA). Oxygen permeation analysis was performed using an Illinois instruments Model 8001 oxygen permeation analyzer. Two samples were simultaneously analyzed for each composition. Ultra high purity nitrogen and air were used for all tests and instrument purging. Each sample was affixed to a brass masking plate (d ∼ 100 mm) using Apiezon T grease so as to cover the 25 mm diameter hole in the center of the plate. The instrument was purged with nitrogen to an initial oxygen concentration of ∼0.14 ppm in the nitrogen carrier gas stream (equivalent to an oxygen transmission rate of 4 cm3/(m2 day), as set in the instrument software) prior to the introduction of oxygen on the

In addition to the significance of this unique observation and its implications as far as the factors governing intercalation are concerned, this result enables an ideal comparison, especially given the lack of experimental data on such highly analogous systems. Differences in intercalation behavior aside, both modified clays possess identically high aspect ratio layers and disperse well in toluene.14,15,17 Given polystyrene’s excellent solubility in toluene, solvent-based processing allows for the formation of films with excellent homogeneity.31 Families of hybrids containing each of the alkylammonium montmorillonites with modified clay contents ranging from 0 to 100 vol % in 10 vol % increments were prepared via automated spray deposition and characterized in terms of morphology, composition, and mechanical and barrier properties. The results were then compared to current analytical models for stiffness and permeability. The conclusions drawn from these results enable us to describe the effects of intercalation alone on properties. In addition to its implications for “bricks and mortar” hybrids more generally, this work provides a rich data set describing in greater detail the properties and behavior of the closest available macroscopic equivalent to the multilayer stacks ubiquitous in the field of polymer/clay nanocomposites.

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EXPERIMENTAL SECTION

Sample Preparation. All samples were prepared from mixtures of modified clay and polymer in solvent. Cloisite 20A and Closite 93A (Southern Clay Products) were chosen for their similarity and ability to disperse in toluene. Closite 20A is a dimethylditallowammonium montmorillonite (DMDT-MMT) while 93A is a methylditallowammonium montmorillonite (MDT-MMT), identical to DMDTMMT except that a hydrogen atom is substituted for one of the two methyl substituents attached to the ammonium headgroup of the alkylammonium modifier. General purpose polystyrene (Scientific Polymer #845) with a molecular weight of 190 000 g/mol and a broad molecular weight distribution, typical of commercial grades, was chosen due to its ability to dissolve in toluene and to intercalate into DMDT-MMT but not into MDT-MMT. All films were formed via a novel spray deposition technique32 and then dried at 125 °C for 4 h in a convection oven to remove residual solvent and relieve any residual stresses present due to the film formation process. Samples were then conditioned at 23 °C and 50% RH for at least 48 h prior to all testing unless otherwise indicated. Thermogravimetric Analysis (TGA). TGA was conducted on each batch of material using a TA Instruments Q50 thermogravimetric analyzer. Prior to testing, samples were dried at 120 °C for a minimum of 2 h to drive off any moisture present. For all tests a platinum pan was used and samples were heated at 20 °C/min to a maximum temperature of 850 °C. In all cases, this was high enough to give a plateau indicative of complete removal of all organic/volatile matter. Samples were analyzed under flowing ultra zero air (60 mL/min) while the balance was purged with ultra high purity nitrogen (40 mL/ min). The nonvolatile content of the nanocomposites was compared to the nonvolatile content of films consisting entirely of the two modified clays studied to confirm the compositions of the hybrids. Wide-Angle X-ray Diffraction (WAXD). WAXD measurements were performed using a Statton box camera and a Rigaku ultraX system with a Cu Kα source (λ = 1.54 Å). The sample-to-detector distance was set to 193 mm for all samples unless otherwise noted. The detector consisted of a reusable X-ray sensitive image plate that was scanned to create a 2-dimensional digital image of the X-ray diffraction pattern. All samples were analyzed under vacuum to minimize scattering due to air. Scanned 2D diffraction patterns were analyzed using FIT2D v12.077 for Windows to produce 2θ and azimuthal integration plots in order to quantify interlayer spacing and degree of orientation utilizing Bragg’s law and the Hermans orientation parameter,33 respectively. Interlayer spacings were calculated from the first-order diffraction peaks observed. Hermans B

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Macromolecules opposite side of the sample. The sampling interval used was 60 min, with the analysis terminated once three OTR values within 1% of each another were measured in succession. OriginLab Origin Pro 8.x was used to perform error-weighted nonlinear regression analysis of the polystyrene nanocomposite data using the various permeation models, with the inorganic content as the independent variable and the oxygen permeation rate or apparent geometric tortuosity factor as the dependent variable. In the former case, the oxygen permeability of the pure polymer and the orientation parameter of each family of materials (calculated as an average over all compositions based on experimentally measured values) must be fixed; in both cases, the aspect ratio of the platelets is allowed to vary.

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structures hypothesized to exist in these two families of materials are shown in Figure 2, with the intercalated structure

RESULTS AND DISCUSSION

Based on comparisons of nanocomposite inorganic content with the inorganic content measured for films consisting solely of the alkylammonium montmorillonite clays studied here, thermogravimetric analysis indicates that the desired compositions (from 0 to 100 vol % modified clay in 10 vol % increments) were achieved in all cases. Actual compositions of the DMDT-MMT hybrids deviated from the desired compositions by less than 1.5 wt % while those of the MDTMMT hybrids deviated by less than 2 wt %. While subsequent discussions will reference the modified clay content in vol % as a means of identification, volume fraction of inorganic material is the relevant composition variable for modeling. The morphology of both families of materials was investigated via 2D wide-angle X-ray diffraction (WAXD, sample scans included in Supporting Information), and the modified clay d-spacings and levels of layer orientation were quantified. For PS/DMDT-MMT hybrids, interlayer spacing increases from an unintercalated value of ∼24 to ∼34 Å in the 0−20 vol % PS range above which point a plateau is reached,1,36 consistent with complete intercalation of all modified clay layers at ∼30 vol % PS or greater. No such trends are observed in the MDT-MMT system, with a constant interlayer spacing of ∼25 Å measured over the entire composition range (Figure 1). The level of orientation is similar regardless of intercalation state (Hermans’ orientation parameter32,37,38 S ∼ 0.65−0.8 in all cases). The overall

Figure 2. Diagrams of proposed morphologies within the intercalated (PS/DMDT-MMT) and nonintercalated (PS/MDT-MMT) hybrids. (A) Organic phases present as a function of composition: blue is “bulk” polystyrene, red is “bulk” polystyrene proximal to modified clay tactoids (perturbed bulk), gray is interphase material (intercalated polstyrene and clay modifier), and green is the unintercalated modified clay phase (clay modifier alone). Inorganic clay layers appear in black. The CVC is the critical volume content of modified clay at which the interphase is the sole organic phase present. (B) Details of the (1) bulk polymer, (2) perturbed bulk polymer, (3) interphase material, and (4) nonintercalated modified clay phases are shown.

proposed in previous work.32 Scanning electron microscopy (SEM, example images included in Supporting Information) of representative materials supports this conclusion; independent of intercalation state, well-oriented structures are observed at high modified clay contents, with a small reduction in orientation at lower modified clay contents. As seen in Figure 3, the densities of the intercalated and nonintercalated systems were found to be similar over the entire composition range. A linear rule of mixtures was sufficient to describe the majority of the macroscopic density variations in these materials. This data rules out the possibility of substantial changes in void content as a function of degree of intercalation. Also of note, the densities of films consisting solely of modified clay were ∼10−20% lower than the values reported by the manufacturer (1.88 g/cm3 for MDT-MMT, 1.77 g/cm3 for DMDT-MMT). TGA rules out the presence of volatile impurities (water, solvent), while the clarity of the films coupled with unsuccessful attempts to densify them in a hot press further confirms that macroscopic voids are not the likely cause. Finally, unit cell based density calculations indicate that densities of 1.55 and 1.63 g/cm3 are to be expected for MDTMMT and DMDT-MMT, respectively. These values closely match what is measured for films consisting solely of modified clay. We have previously shown32 that intercalated PS/DMDTMMT hybrids exhibit a strong loss modulus peak that appears to be a combination of the DMDT modifier transition and the (depressed) polystyrene alpha transition as well as synergistic

Figure 1. WAXD-derived interlayer spacing and Hermans orientation parameter measurements from intercalated (DMDT-MMT) and nonintercalated (MDT-MMT) hybrids with polystyrene. Horizontal dotted lines represent constant fits over the polymer-rich concentration range (for DMDT-MMT) and the entire concentration range (for MDT-MMT). C

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Figure 3. Densities of intercalated (DMDT-MMT) and nonintercalated (MDT-MMT) hybrids with polystyrene. The density of the polystyrene used here was reported by the manufacturer as 1.05 g/ cm3, while densities for the pure MDT-MMT and DMDT-MMT were estimated from to be 1.55 and 1.63 g/cm3, respectively, based on composition and WAXD data.

enhancements in both the storage and loss moduli. To elucidate the effects of intercalation, here the aforementioned PS/DMDT-MMT system has been compared with an analogous nonintercalating system (PS/MDT-MMT). Dynamic mechanical analysis of the latter system (Figure 4) shows similar underlying trends to our previous report on the intercalated system.32 As the presence of clear X-ray peaks in these samples implies the existence of multilayer stacks, such similarities are not surprising, especially in light of the multiscale micromechanical models developed by Sheng et al., who report that strain shielding and the resulting overlap of strain fields lead to increases in effective particle size.39 These authors further note that the morphological transition from intercalated to exfoliated leads to only minor effects on nanocomposite modulus.39 More recent micromechanical modeling efforts by Anoukou et al. have attempted to account for the interphase and the presence of multilayer stacks and indicate that the presence of a stiff interphase represents one means of explaining experimental modulus data for a series of polyamide-6/clay nanocomposites.40−42 The data presented here, on the other hand, support the conclusion that the interphase contributes little to overall differences in the mechanical behavior of these polystyrene-based hybrids, emphasizing the importance of interfacial strength and interphase properties as well as the degree to which such properties may vary from system to system. Given these results, it is important to examine the assumptions of the aforementioned modeling efforts so as to provide additional context. Both groups focus on the behavior of the polymer within intercalated stacks, as opposed to perturbed material surrounding the stacks. Sheng et al. assume a perfect interface and neglect specific contributions from the interphase and/or modifier, while Anoukou et al. assume that the interphase is 8 times stiffer than the matrix while neglecting changes in crystalline morphology or suppression of crystallinity in confined polyamide-6. Likewise, while they note major effects on the crystalline microstructure as a result of nanocomposite formation, they nevertheless assume that the

Figure 4. DMA results showing (A) storage and (B) loss moduli for nonintercalated PS/MDT-MMT hybrids vs composition and temperature. Curves are vertically offset by arbitrary increments for clarity. Equivalent figures for the intercalated hybrids are found in the Supporting Information.

properties of the unfilled matrix polymer are identical to the properties of the bulk polymer in the nanocomposite. The behavior and properties of the polystyrene hybrids reported here highlight clear limitations as far as the general applicability of these assumptions are concerned. In contrast with the elastic response of these hybrids, features in the loss modulus curves for the nonintercalated system are depressed overall compared to the intercalated system.32 This is consistent with the idea that polymer/clay interfaces encourage energy dissipation; the nonintercalated system is expected to contain fewer interfaces of this type and larger regions of bulk polymer, leading to a decrease in mechanical loss. Various thermal transitions are observed as well, including some that appear quite weak. First derivatives of both storage and loss modulus data have been used in the literature to access and quantify subtle transitions;43,44 we apply a similar methodology here. In particular, thermal transitions are assessed through deconvolution of peaks observed in firstderivative plots of storage moduli (Figure 5). Based on these analyses, the temperatures and proposed assignments of the various transitions detected are given in Figure 6. Assignments for the intercalated system were reported D

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in previous work,32 where it was shown that these materials may be described as a combination of four possible phases (shown in Figure 2B), each with its own associated thermal transition: (1) bulk polymer, (2) perturbed bulk polymer, i.e., polymer external to but in close contact with modified clay tactoids, (3) interphase material (polymer and modifier), and (4) nonintercalated modified clay. Of the four phases listed, three (bulk polymer, interphase material, and nonintercalated modified clay) are composition-invariant, with transition temperatures observed to be independent of clay content. The transition temperature of the perturbed bulk polymer, on the other hand, was observed to decrease from levels near the transition temperature of the bulk polymer to levels approaching that of the interphase material as clay content increased. This was ascribed to the fact that the perturbed bulk polymer experienced more confinement and greater interactions with the clay and its modifiers at higher clay concentrations. In contrast with the aforementioned study of intercalated DMDT-MMT hybrids with polystyrene,32 the nonintercalated nature of the MDT-MMT hybrids with polystyrene reported here implies only three phases: bulk polymer, perturbed bulk polymer, and nonintercalated modified clay. The resultant data for both transition temperatures and transition strength (as measured by fractional peak area) are reported in Figure 6. Before an in-depth discussion of the thermal transitions derived from the DMA data, context must be presented. In particular, it is important to note the degree of variation in the literature concerning the effect of clay nanocomposite formation on the glass transition temperature (Tg) of the matrix polymer. Results range from little or no change45 to increases of up to 10 °C46 or decreases of up to 10 °C.47 Work by Rittigstein and Torkelson on solvent-processed nanocomposites based on spherical nanoparticles shows that the solvent choice and processing route can strongly affect the Tg of the resultant nanocomposite systems as well.48 However, details as to how the Tg changes as a function of distance from the nanoparticle surface are generally unknown, and estimates of the thickness of the perturbed polymer layer often tend toward the speculative. Likewise, studies on polystyrene/clay nanocomposites are limited to properties evaluations in all but a few instances. One report describes the melt processing and characterization a polystyrene nanocomposite containing 40 vol % DMDT-MMT and reports a reduction in the strength and breadth of the glass transition.29 In a nonintercalated system, even higher clay concentrations will be required for the polymer chains to experience significant confinement, at which point the volume of perturbed bulk polymer will be minimal and its transition will be quite weak as a result. This is shown graphically in Figure 7, where interparticle distance is represented as a function of modified clay content for a few limiting cases. The interparticle distance is estimated by calculating the surface-to-volume ratio of a stack of a given thickness and then uniformly distributing all polymer external to the stack over the stack surface; the interparticle distance is simply taken as twice the thickness of this coating. Since the surface-to-volume ratio of the stack is the same regardless of layer cross section (disk, square, etc.), no assumption of layer shape is required. As intercalated polyethylene/polyethylene-graf t-maleic anhydride nancomposites containing 5 wt % of dimethylditallowammonium montmorillonite nanoclay have been reported to possess a weight-average stack thickness of 32 nm based on

Figure 5. First-derivative DMA storage modulus curves of nonintercalated PS/MDT-MMT nanocomposites vs composition and temperature. Curves are vertically offset by arbitrary increments for clarity. Equivalent figures for the intercalated hybrids are found in the Supporting Information.

Figure 6. DMA-derived transition data for nonintercalated MDTMMT hybrids with polystyrene. (A) Peak maxima from first-derivative storage modulus data vs volume fraction clay. (B) Peak areas from first-derivative storage modulus data vs volume fraction clay. Lines are presented to guide the eye. Equivalent figures for the intercalated hybrids are found in the Supporting Information.

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nature of the interaction zone around modified clay stacks may play a role in differentiating the intercalated and nonintercalated systems. This argument may be directly evaluated via the analysis of interparticle distances presented in Figure 7. As stack thickness increases, geometry dictates that the surfaceto-volume ratio of the stack rapidly approaches the “infinite” stack limit. Since the bulk polymer transition disappears only at relatively high modified clay concentrations (36.8 vol % DMDT-MMT in the intercalated case; 44.2 vol % MDTMMT in the nonintercalated case), the assumption of a surfaceto-volume ratio equal to the “infinite” stack limit is deemed reasonable. This gives an average distance between tactoids at these critical concentrations of ∼32 nm in the intercalated case and ∼24 nm in the nonintercalated case. As these distances are calculated for systems where only perturbed bulk polymer is present between the tactoids, these values may be taken as estimates of twice the thickness of the perturbed bulk polymer layer surrounding the tactoids. The existence of a thicker perturbed bulk polymer layer in the intercalated case is consistent with a greater degree of mixing and/or interaction. While defining the exact nature of these interaction zones is beyond the scope of this work, the direct estimation of different interaction zone thicknesses from DMA data is a significant result representative of a promising new means of characterizing the polymer nanocomposites more generally. It is also noteworthy that the length scales calculated via this approach compare favorably with recent reports concerning both estimates of the characteristic length scale of the polystyrene glass transition52 and variations in the Tg of polystyrene as a function of free-standing film thickness53 as well as earlier work concerning the range over which the glass transition is perturbed in polymer blends.54 Overall, dynamic mechanical analysis reveals subtle differences between the intercalated and nonintercalated materials. This does not mean that the dynamic mechanical properties of intercalated and immiscible states are effectively the same in low modified clay content nanocomposites, however. Rather, it shows that in systems with similar levels of dispersion, distribution, orientation, and composition that the effects of intercalation on dynamic mechanical behavior are relatively small. This, in turn, emphasizes that differences in dispersion, orientation, and polymer microstructure play a more important role than intercalation. However, it is also true that high levels of dispersion are not characteristic of stable nanostructures found in systems where intercalation is thermodynamically unfavorable. Regardless, the observed trends in the data are readily explained by the coexistence of a limited number of phases in these materials (Figure 2): (1) bulk polymer and (2) perturbed bulk polymer in close proximity to modified clay tactoids (i.e., the polymer-rich phases) as well as (3) the polymer-intercalated interphase and (4) the nonintercalated modified clay (i.e., the clay-rich phases). The derivative storage modulus and loss modulus data produce complementary trends, while offering additional insights into stress transfer and damping behavior as expressed in the loss modulus data. In this context, these results hold clear lessons relevant to broad classes of nanocomposite materials. In particular, for weakly interacting systems, polymer/clay interfaces appear to be lossier than clay/clay interfaces, leading to greater loss modulus values in intercalated systems vs their nonintercalated counterparts. Given a constant area of polymer/clay interface through which to transfer stress, reductions in interfacial strength result in greater mechanical loss (E″) at the cost of

Figure 7. Interparticle distance vs composition for intercalated (DMDT-MMT) and nonintercalated (MDT-MMT) polystyrene hybrids; layers are taken to be 100 nm across, 1 nm thick, and perfectly oriented in-plane; intercalated and nonintercalated interlayer spacings are taken from WAXD data (Figure 1).

TEM image analysis and an intercalated d-spacing of 3.3 nm based on X-ray diffraction,49 10 layers is taken as a reasonable lower limit on stack thickness in our materials. With this background in mind, and a clearer view of the range of interparticle distances expected, we return to the data in Figure 6. Regardless of intercalation state, an invariant transition is observed at ∼100 °C, consistent with accepted values for neat polystyrene50 and assigned to the bulk polymer (1 in Figure 2B). The modified clay content at which this bulk polystyrene glass transition effectively disappears (as judged by a loss of 95% of its original peak area according to a logistic fit of the data) is higher in the nonintercalated system (44.2 vol % MDT-MMT = 17.6 vol % inorganic) than in the intercalated system (36.8 vol % DMDT-MMT = 14.1 vol % inorganic). This follows directly from the fact that the interparticle distance drops more gradually with increasing modified clay content in the nonintercalated case (Figure 7), producing significant confinement only when the polymer content is low. Consistent with this scenario, the perturbed bulk transition does not appear below a modified clay content of 40 vol % and becomes too weak to detect as the modified clay content continues to increase. In the intercalated case, in contrast, the perturbed bulk transition is consistently observed in all nanocomposites except those so starved of polymer that complete intercalation is impossible. Common to both systems, the perturbed bulk transition appears at lower temperatures than the bulk polymer. A number of possible causes for such a decrease have been described in the literature. Plasticization of polymer by the clay modifier is one explanation that is often invoked in the context of polymer/clay nanocomposites.51 Indeed, we observe via DSC and DMA that the addition of low volume fractions of octadecane to polystyrene can reduce Tg to ∼80 °C (data not shown). In the nanolaminates described here, however, any plasticization must be highly localized, given the coexistence of a transition for bulk PS and perturbed PS, both of which are present up to high volume fractions in the intercalated series but less prominent in the nonintercalated case. In addition to the aforementioned arguments concerning interparticle distance and polymer content, differences in the F

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Macromolecules stiffness (E′). In the intercalated case, the addition of more modified clay increases the number of mechanically lossy polymer/clay interfaces while enhancing stiffnessbut this effect is limited to compositions where sufficient polymer exists to give complete intercalation of all modified clay layers. Too great a reduction in interfacial strength precludes intercalation in the first place, significantly reducing the number of lossy polymer/clay interfaces and producing less damping overall. Major increases in interfacial strength, on the other hand, can result in materials with exceptional stiffness, as a recent review of the literature clearly shows55but with little or no damping capacity. It is further concluded that polystyrene/modified clay nanocomposites represent an example of borderline compatibility. As such, the lessons learned here may prove relevant to other commercially relevant borderline-compatible clay nanocomposites based on compatibilized polyolefins as well as all manner of low to moderate polarity commodity plastics (polyesters, vinyls, acrylics, etc.). In addition, these findings highlight clear opportunities to better understanding nanocomposite damping behavior, i.e., through studies of nanocomposites based on amorphous copolymers whose comonomers exhibit similar transition temperatures but significantly different interaction strengths as far as clays are concerned (styrene vs 2-vinylpyridine, acrylates or methacrylates vs acrylamides or methacrylamides, etc.). In addition to dynamic mechanical behavior, it is also worth considering the elastic properties of these materials in a more classical sense. With this in mind, DMA storage modulus data (1 Hz) taken at a temperature below the rich thermal transition behavior described previously (25 °C) and also above the Tg of bulk PS (120 °C) are used to assess the effects of intercalation on the reliability of the modulus predictions of common micromechanical models (Figure 8). In particular, the Halpin− Tsai56 (HS) and Tandon−Weng57 (TW) models have been selected, as these consider the effects of filler and matrix properties, filler aspect ratio, and volume fraction, and are commonly applied in the field of polymer/clay nanocomposites. As these models produce results for platelets either parallel or perpendicular to the direction of the applied stress, a weighted average has been calculated here based on the average Hermans orientation parameter32 obtained from wideangle X-ray diffraction data for each family of materials. Results from the Voigt and Reuss models58 are provided as well, representing upper and lower bounds on modulus, respectively. The latter are based on rule-of-mixtures arguments and account for composition and matrix and filler properties but neglect the existence of discrete particles with defined shapes and aspect ratios. Similar to our previous report concerning the intercalated PS/DMDT-MMT system,32 the nonintercalated system shows a gradual increase in storage modulus that appears to plateau beyond ∼20 vol % inorganic content. Interestingly, the nonintercalated system displays levels of stiffness nearly as high as the intercalated system regardless of temperature, most likely due to the well-ordered and highly oriented nanostructure. This clearly demonstrates that the lack of intercalation has only a minor effect on stiffness as a function of temperature and composition. The plateau in the modulus data observed in both families of material may be related to further weakening in interfacial strength between adjacent clay layers in the absence of any intervening polymer (Figure 2). As inorganic content increases, clay−clay interfaces become more common than

Figure 8. DMA storage moduli for intercalated (DMDT-MMT) and nonintercalated (MDT-MMT) polystyrene hybrids (taken at 1 Hz) at (A) 25 °C and (B) 120 °C (data points). In all cases, continuous curves represent the Reuss (lower solid green) and Voigt (upper solid purple) bounds and the orientation-weighted Halpin−Tsai (dashed blue) and Tandon−Weng (dotted red) results. For modeling purposes, the modulus of the pure polymer is used for the organic phase while a modulus of 160 GPa and a constant layer aspect ratio of 100 are used for the inorganic phase. Analogous plots of storage modulus data taken at −25 and 75 °C appear in the Supporting Information.

polymer−clay interfaces, and the expected increases in stiffness are moderated as a result of the very limited ability of the modifier tails alone to transfer stress. This also explains why the storage modulus of the nonintercalated system tends to fall below that of the intercalated system, given that polymer/clay interfaces are more common in the latter at any given inorganic content. Qualitatively, the modulus predictions of the HS and TW models are similar. At 25 °C, all of the experimental data fall well below the orientation-weighted HS and TW predictions, providing the sort of result that has led many to accept the “effective particle” model of the multilayer stack. Observations taken at 120 °C emphasize the limitations of this approach. At this temperature, the experimental modulus values exceed the HS and TW predictions at all but the lowest inorganic concentrations, in spite of the fact that an aspect ratio of 100 G

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to differences in free volume due to variations in molecular weight distribution and process history.61 Both systems exhibit an exponential decay in oxygen permeability as inorganic content increases. The similarities in the trends for the intercalated (DMDT-MMT) and nonintercalated (MDTMMT) PS based systems indicate that the effect of intercalation on barrier properties is small and that the majority of the improvements appear to arise from a combination of composition and orientation. While a number of models have been used to describe the barrier properties of polymer/clay nanocomposites, little has been done to date to examine the consequences of intercalation on model accuracy. Most models make predictions based on the geometry of the platelets, the permeability of the polymer matrix, and the ratio of polymer to inorganic. More recent work38,62,63 has demonstrated that platelet orientation can be accounted for in permeation modeling as well. Specific to the data shown here, the application of common permeation models63−67 allows comparisons to be made as a function of intercalation state over the majority of the composition range. Correcting for orientation effects improves the fit of the Cussler63,68 and Gusev and Lusti66 models to the oxygen permeability data from the nonintercalated system just as it does in our previous report concerning the intercalated system,38 with the random array Cussler model (shown in Figure 9) producing a slightly higher adjusted r2 value. The random array Cussler model was developed assuming a high filler content composite containing perfectly oriented but randomly distributed high aspect ratio filler particles, while the Gusev and Lusti model was based on an empirical fit of the results of a finite element simulation involving perfectly oriented but randomly distributed disks. Two critical observations may be made here. First, both models give very high adjusted r2 values in spite of the fact that they assume a constant aspect ratio for the layers regardless of hybrid composition. Second, the value of the aspect ratios derived from these fits is close to that of a single clay layer. This is observed in spite of the fact that the materials themselves contain only multilayer stacks whose thickness is expected to vary with composition. We emphasize that this result is in significant contrast to the mechanical properties results reported here, where no single aspect ratio or aspect ratio dependence on composition can provide an accurate fit of storage modulus trends, and variations in interfacial strength and the properties of the various organic phases present dominate. Although in the case of the permeation data the fitted value of the aspect ratio (α) was found to be somewhat lower in the nonintercalated state than the intercalated state, these differences are readily explained by the greater inherent oxygen permeability observed in the pure MDT-MMT vs the pure DMDT-MMT, consistent with the lower theoretical density (and thus greater free volume) of the former as estimated previously in this report. While these conclusions have significant consequences given frequent use of the aspect ratio of multilayer stacks when predicting the performance of nonexfoliated systems, two features are lacking in this rather traditional analysis of the permeation data. First, it is unsatisfying to be unable to account for variations in orientation parameter from one composition to the next. Second, it is equally unsatisfying to be unable to offer more of an explanation for the lack of fit at the highest clay concentrations. These issues lead us to propose a new means of representing and analyzing permeation data in the context of hybrids of this sort, which we term the tau plot. In the sort of

remains in use. These results combined with similar comparisons at other temperatures (shown in the Supporting Information) confirm that it is impossible to define a single aspect ratio, or even a set of composition-dependent aspect ratios, that will enable the aforementioned micromechanical models to predict the observed modulus data at more than one temperature. As effective aspect ratio cannot be a function of temperature, the data must be explained by other means. As has been discussed in previous work for the DMDT-MMT composite series,32 these observations may be attributed to several factors. To summarize, existing models assume that the filler and matrix are isotropic, perfectly bonded, and linearly elastic and often that the filler particles are perfectly aligned, anisotropic, and uniform in size and shape, with particle− particle interactions neglected.59 In addition, many models assume the organic fraction of the material to be homogeneous, in contrast with the results of dynamic mechanical analysis presented here confirming the presence of multiple distinct organic phases with different mechanical characteristics. As is evident from the plots shown in Figure 8, simple orientation corrections do not address the discrepancies observed. Likewise, variations in particle size and shape may be ruled out as an explanation based on prior discussions of aspect ratio. Rather, in addition to the coexistence of multiple organic phases with distinct properties, variations in interfacial stress transfer are posited to be especially important, as demonstrated by the mechanical loss characteristics of these materials. Having examined the morphology and mechanical properties of these materials, we turn our attention to their permeation behavior. Figure 9 presents oxygen permeation data for both

Figure 9. Oxygen permeability data for intercalated PS/DMDT-MMT and nonintercalated PS/MDT-MMT nanocomposites (points) as well as the associated orientation-corrected error-weighted random array Cussler fits (lines). Fitting was performed using the average orientation parameter calculated across all compositions for each system. In the nonintercalated case, as no model allows for an increase in P with increasing ϕ, the last data point was excluded from the fit.

the intercalated and nonintercalated systems in the traditional fashion, as permeability vs a composition variable (in this case, inorganic content in vol %). The measured permeability for the neat PS used in this work was found to be 196.3 ± 0.8 cm3 mm/(m2 day atm), similar to the literature value of 166 cm3 mm/(m2 day atm),50,60 with the discrepancy observed ascribed H

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only in the case of the Gusev and Lusti model (shown), with all other model fits experiencing a reduction in their adjusted r2 value as a consequence of the inclusion of the full set of experimentally derived orientation parameter data. In addition, it is clear that we must exclude not just the last data point but the last two data points from any attempt to model the data for the nonintercalated system, since a reduction in τ with increasing ϕ is not allowed by any of the models referenced here. Likewise, the aspect ratios derived from these best fits of the data have increased to the point where they are nearly identical to what is reported by the manufacturer for these materials (75−150), further emphasizing that the single layer aspect ratio drives the barrier performance of these systems even in the absence of delamination. Finally, the overlay plot of estimated average organic phase density enables us to discuss the origins of this reduction in τapp in the nonintercalated system at the highest clay loading levels in greater detail. In particular, we clearly observe that the aforementioned deviation occurs over exactly the same range of compositions where the average density of the organic phase drops much more rapidly in the nonintercalated system than the intercalated system. This result supports previous arguments concerning density and free volume in a more quantitative fashion and highlights the critical importance of considering not just the concentration of impermeable layers but also the behavior of the organic phase between those layers in determining nanocomposite barrier properties, even (especially) at the highest nanofiller loadings. It is posited that below some critical organic phase density a change occurs in the quantity, distribution, and connectivity of free volume in the system that enables substantial enhancements in P0, the permeability of the organic phase, in our materials, resulting in a significant reduction in the τapp values of the nonintercalated system at the highest clay concentrations. Emphasizing the power of this approach, it may be possible to assess changes in P0 independently (via time lag experiments, for instance) and then incorporate these results into calculations of the apparent tortuosity factor to account for such effects as well.

analysis shown in Figure 9 and reported throughout the literature, the approach is to represent permeability (P) as a function of the volume fraction of platelets (ϕ), as follows: P0(1 − φ)

P=

( 23 )(S + 12 )

1+τ

(1)

Here, the orientation parameter (S) is treated as a constant (usually assumed to be 1), and the aspect ratio (α) is embedded within the expression for the geometric tortuosity factor (τ) in a manner that depends on the model chosen for use. In contrast, the construction of a tau plot involves rearranging this expression to calculate the apparent tortuosity factor (τapp) at each composition, as follows: τapp =

P0 (1 P 2 3

− φ) − 1

(S + 12 )

(2)

This approach is superior to a typical analysis of P vs ϕ in several respects. First, it enables the incorporation of distinct experimental values of the orientation parameter (S) at each individual composition, substantially enhancing the accuracy of any further analyses by removing deviations from a fixed and composition-invariant average value of S as a source of error. Second, it enables the rapid visual selection of the model that best fits the data, given clear expressions of the proportionalities between τ and ϕ in each of the major barrier models mentioned here. In the Nielsen model, τ depends linearly on ϕ;64 in the Cussler model, the dependence is quadratic;63 in the Fredrickson and Bicerano model, the dependence is intermediate between linear and quadratic;65 and in the Gusev and Lusti model, the dependence is exponential.66 Third, this approach more clearly highlights those points that represent deviations from the expected trends. All of these advantages are on display in Figure 10. First, it is observed that the ability to fit the data is retained or improved

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CONCLUSIONS The results presented here demonstrate that while the right combination of polymer and modifier is required for intercalation, even in its absence, well-ordered but nonintercalated hybrids are capable of giving thermomechanical and barrier properties similar to those of intercalated systems. In particular, both the intercalated and nonintercalated systems display high storage moduli, with the nonintercalated system giving values only 10−25% lower than to the intercalated system at modified clay concentrations between 40 and 80 vol %. In contrast, more significant changes in the loss modulus data are observed, with greater mechanical loss observed as a consequence of intercalation. Conventional micromechanical models do a poor job of capturing the storage modulus trends, and two significant conclusions of the mechanical analyses presented here are that the discrepancies observed between model and experiment are not a function of intercalation state and that there is neither an average aspect ratio nor any variation of aspect ratio as a function of inorganic content that enables these models to describe the experimental data as a function of both composition and temperature. Analyses of the derivative storage modulus data provide further insight into temperature-dependent mechanical response of these materials through the identification of a number

Figure 10. Apparent tortuosity factor data for intercalated PS/DMDTMMT and nonintercalated PS/MDT-MMT nanocomposites (points) as well as orientation-corrected fits (solid lines) based on the Gusev and Lusti model, as derived from the oxygen permeability data presented in Figure 9 and the orientation parameter data presented in Figure 1. In the nonintercalated case, as no model allows for a decrease in τ as ϕ increases, the last two data points were excluded from the fit. Also shown are estimates of the average densities of the organic phase present in both systems (dashed lines). I

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concerning the structure, dynamics, and performance of these materials that are directly relevant to numerous reports concerning the preparation, mechanical, and barrier properties of polymer nanolaminates (including layer-by-layer deposited nanocomposites, bricks-and-mortar hybrids, synthetic nacre analogues, nano brick walls, etc.) from a variety of 2D nanofillers.55

of thermal transitions present in these materials. For instance, even though no interphase exists, the nonintercalated composites continue to exhibit transitions corresponding to perturbed polymer and clay modifier, though to a much lesser degree than the intercalated system. This indicates that even in the absence of intercalation, clay modifiers on the outer surfaces of the tactoids still interact with nearby polymer, with consequences for the dynamics of the system. More generally, the observed behavior matches well with the proposed existence of up to four distinct phases in these systems: bulk polymer and bulk polymer perturbed by its close proximity to modified clay tactoids (i.e., the polymer-rich phases) and the interphase and the nonintercalated modified clay (i.e., the clayrich phases). A geometric argument allows for the estimation of the thickness of the perturbed polymer layer and confirms that it is significantly thicker in the intercalated case (∼16 nm) than in the nonintercalated case (∼12 nm) and consistent with literature reports concerning critical length scales related to chain dynamics and confinement effects. The analytical approaches developed in the course of this work hold significant promise for further evalulation of this phenomenon in appropriate materials systems. In particular, much could be learned through the study of analogous nanocomposites based on amorphous copolymers of comonomers with similar transition temperatures but significantly different interaction chemistries (styrene vs 2-vinylpyridine, acrylates or methacrylates vs acrylamides or methacrylamides, etc.). In contrast with the elastic properties, oxygen permeation data demonstrate that the majority of the observed permeation behavior may be captured through the combination of an orientation-corrected tortuosity-based permeation model with a composition-invariant aspect ratio similar to that of a single clay layer, without significant changes as a function of intercalation state. In addition, a new means of analyzing such data via a plot of apparent tortuosity factor vs inorganic content is shown to be superior to this more conventional analysis, with of the incorporation of experimental variations in orientation parameter with composition enabling more effective model selection and improving the fit of the optimal model while more clearly identifying its range of applicability. The overall decrease in permeability and increase in apparent tortuosity remain significant even in the absence of intercalation, at least until the system is highly polymer-starved, at which point the clay modifier makes up the majority of the organic phase and begins to dominate the permeation behavior. In contrast with the micromechanical modeling efforts, permeation modeling is able to describe the majority of the behavior of both the intercalated and nonintercalated systems, with deviations from these models explained via changes in the density of the organic phase occurring over the same composition range. On the whole, these results offer significant insights into the changes that may be expected in macroscopic (thermomechanical and barrier) properties as a result of polymer intercalation and clearly resolve the phases present in various nanocomposite systems (Figure 2). As perfect exfoliation is never achieved, intercalated stacks are commonplace in most polymer/clay nanocomposites, making information on the effects of intercalation critical to the understanding and modeling of real nanocomposite systems. In addition, the performance of weakly interacting nanocomposite systems is of significant commercial relevance given the importance of polyolefins and other low to moderate polarity commodity plastics in a wide range of applications. Finally, this report offers critical details

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b01788. Example 2D wide-angle X-ray diffraction patterns and representative SEM images for intercalated and nonintercalated systems; unit cell based density estimates for (modified) clays, clay modifier phases alone, the intercalated organic interphase alone, and average organic phase density; DMA storage modulus, loss modulus, and first-derivative storage modulus curves for intercalated (DMDT-MMT) polystyrene hybrids; peak maxima and peak area plots derived from first-derivative DMA storage modulus data for intercalated (DMDTMMT) polystyrene hybrids; DMA storage modulus/ micromechanical model plots at −25 and 75 °C (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; Fax ++1.978.934.3089; Ph ++1.978.934.3451 (D.F.S.). Present Address

E.D.: E Ink Corporation 1000 Technology Park Drive Billerica, MA 01821. Funding

This work was supported by the United States Air Force Research Laboratory, via subcontracts with the Universal Technology Corporation. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Scott Fillery at the United States Air Force Research Laboratory (WPAFB, OH) for assistance with SEM analysis, as well as Rich Vaia, Hilmar Koerner, and AFRL/ RXBN for general support, assistance, facilities access, and feedback concerning this report. We thank John Hunt and the Cornell Center for Materials Research for assistance in acquiring the SEM-EDS data referenced in the Supporting Information. We thank the NSF-sponsored Center for HighRate Nanomanufacturing at UML for access to the thermal analysis equipment, the UML Baseball Research Laboratory for sample conditioning facilities, and the Department of Plastics Engineering for general support. Finally, we thank Dr. Mat Celina of Sandia National Laboratories, Prof. William J. Koros of the Georgia Institute of Technology, and Prof. Emmanuelle Reynaud of the University of Massachusetts Lowell for their input and suggestions.

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ABBREVIATIONS DMA, dynamic mechanical analysis; DMDT-MMT, dimethylditallowammonium, montmorillonite; DSC, differential scanJ

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ning calorimetry; HS, Halpin−Tsai; MDT-MMT, methylditallowammonium montmorillonite; MMT, montmorillonite; OPA, oxygen permeation analysis; PS, polystyrene; SEM, scanning electron microscopy; TGA, thermogravimetric analysis; TW, Tandon−Weng; WAXD, wide-angle X-ray diffraction.

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