Rheology of Polymer/Layered Silicate Nanocomposites

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Rheology of Polymer/Layered Silicate Nanocomposites

Suprakas Sinha Ray^†

Macromolecular Nanotechnology Research Group, Nanoscience Research Centre, CSIR Materials Science and Manufacturing, PO Box 395, Pretoria 0001, Republic of South Africa

Received November 14, 2006

Abstract: Over the last few years, polymer/layered silicate (PLS) nanocomposites have been an area of intense academic and industrial research. No matter the measure-journal articles, patents, and industry research and development (R&D) funding-efforts in PLS-nanocomposite have been exponentially growing worldwide for the last ten years. Aside from the understanding in fabrication, characterization, and improved mechanical and other properties, the measurement of rheological properties of PLS-nanocomposites under molten state is crucial to gain fundamental understanding of the processability of these materials. In the case of PLS- nanocomposites, the measurements of rheological properties are also helpful to find out the strength of polymer-layered silicate interactions and the structure-property relationship in nanocomposites. This is because rheological behaviors are strongly influenced by their nanoscale structure and interfacial characteristics. In this article, recent advances in melt-state rheological properties of PLS-nanocomposites are highlighted. To begin this a very brief description of structure and dynamics of PLS-nanocomposites are also described.

Keywords: polymer/layered silicate nanocomposite, structure, thermodynamics, melt rheology

Introduction

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The word composite is generally used to define any material made of more than one component. Depending on the matrix nature, composite materials can be broadly divided into three categories such as polymeric, metallic, and ceramic. In most of the commercially available composite materials the structural unit is on the micrometer (10^‐6m) length scale and is used mostly to improve the mechanical properties of matrix material. Now one can change the dimension of the structural building block to the nanometer (10^‐9m) scale, the resulting composite ma- terial is called a Nanocomposite. The nanometer is one billionth of meter and ∼10,000 times finer than a human hair. Figure 1 compares the morphology of classical- and nano-filler filled composite.

Here we are interested on polymer nanocomposites that means nanocomposite made with polymer and nanoscale filler. Over the last few years, various types of nano-fillers are in development such as layered silicates

†To whom all correspondence should be addressed.

(e-mail: RSuprakas@csir.co.za)

[1], metal oxide nanoparticles [2,3], inorganic nanotubes [4], expandable graphite [5-7], layered titanate [8], cellu- lose nanowhiskers [9], polyhedral oligomeric silses- quioxanes (POSS) [10], carbon nanotubes [11-17], etc.

Polymer nanocomposites based on layered silicates have attracted great interest in today’s materials research, as it is possible to achieve impressive enhancements of properties when compared with virgin polymer [1,18-20].

These improvements can include high moduli [1,18], increased strength and heat resistance [1,18], decreased gas permeability [21-23] and flammability [24-27], and increased degradability of biodegradable polymers [28-30].

On the other hand, there has been considerable interest in theory and simulations addressing the preparation and properties of these materials [31], and they are also considered to be unique model systems to study the structure and dynamics of polymers in confined environments [32].

The main reason for these improved properties in nanocomposites is the interfacial interaction between the matrix and layered silicate, as opposed to conventional composites [1]. Layered silicates have layer thickness on

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Figure 1. Schematic comparison between microcomposite (classical filler filled) and “Nanocomposite. Reproduced with permission from Elsevier Science Ltd., UK.

Figure 2. Schematic illustration of three different types of ther- modynamically achievable polymer/layered silicate composites.

the order of 1 nm, and very high aspect ratios (e.g. 10∼

1,000). A few weight percent of layered silicate particles that are properly dispersed throughout the matrix can thus create a much larger surface area for polymer filler interactions than do conventional composites. On the basis of the strength of the polymer/layered silicate interfacial interaction, three structurally different types of composites are achievable (see Figure 2): (1) phase- separated composite, when polymer matrix has no interaction with layered silicate, (2) intercalated nanoco- posites, where insertion of polymer chains into the silicate structure occurs in a crystallographically regular fashion, regardless of the polymer-to-layered silicate

ratio, and a repeat distance of few nanometers, and (3) exfoliated nanocomposites, in which the individual silicate layers are separated in the polymer matrix by average distances that totally depend on the layered silicate loading.

This article highlights recent advances in melt-state rheological properties of PLS nanocomposites. To begin this a brief description on structure and properties of layered silicates, nanocomposites preparation techniques and thermodynamics are described. Finally, attention is drawn to how rheological behavior of PLS nanocomposites can be understood directly from their structure.

Structure and Properties of Layered Silicates

The commonly used layered silicates for the preparation of PLS nanocomposites belong to the same general family of 2 : 1 layered- or phyllosilicates [33]. Their crystal structure consists of layers made up of two tetrahedrally coordinated silicon atoms fused to an edge-shared octahedral sheet of either aluminum or magnesium hydroxide. The layer thickness is around 1 nm, and the lateral dimensions of these layers may vary from 30 nm to several microns or larger, depending on the particular layered silicate. Stacking of the layers leads to a regular van der Waal’s gap between the layers called the interlayer or gallery. Isomorphic substitution within the layers (for example, Al⁺³ replaced by Mg⁺² or Fe⁺², or Mg⁺² replaced by Li⁺¹) generates negative charges that are counterbalanced by alkali and alkaline earth cations situated inside the galleries. This type of layered silicate

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Figure 3. Structure of 2:1 phyllosilicates. Reproduced from ref.

[1] by permission of Elsevier Science Ltd., UK.

is characterized by a moderate surface charge known as the cation exchange capacity (CEC), and generally expressed as mequiv/100 gm. This charge is not locally constant, but varies from layer to layer, and must be considered as an average value over the whole crystal.

Layered silicates have two types of structure: tetrahedrally substituted and octahedrally substituted. In the case of tetrahedrally substituted layered silicates, the negative charge is located on the surface of silicate layers and, hence, the polymer matrices can interact more readily with these than with octahedrally substituted material.

Generally, layered silicate minerals are divided into three major groups: (a) the kaolinite group, (b) the semc- tite group and (c) the illite or the mica-clay group.

Among the three major groups, smectite types or more precisely montmorillonite, saponite and hectorite are the most commonly used layered silicates in the filed of polymer nanocomposite technology. Again, among montmorillonite, saponite, hectorite, montmorillonite (MMT) is the most commonly used layered silicates for the fabrication of PLS nanocomposites, because it is highly abundant and inexpensive. MMT is the name given to the layered silicate found near montmorillonite in France, where MMT was first identified by Knight in 1896. The chemical formula of MMT is Mx(Al4‐xMgx)Si8O20(OH)4.

The CEC (generally 90∼110 meq/100 g) and particle length of MMT (100∼150 nm) depends on the source.

The specific surface area of MMT is equal to 750∼800 m²/g and the modulus of each MMT sheet is around 170 GPa [34]. The interlayer thickness of hydrated MMT is equal to 1.45 nm and the average density ρ= 2.385 g/mL. Drying MMT at 150 °C reduces the gallery height to 0.28 nm which corresponds to a water monolayer and hence the interlayer spacing decreases to 0.94 nm and the average density increases to 3.138 g/mL [35]. The unit structure of MMT is presented in Figure 3.

Two particular characteristics of layered silicates are generally considered for PLS nanocomposites fabrication. The first is the ability of the silicate particles to disperse into individual layers inside the polymer matrix.

The second characteristic is the ability to fine-tune their surface chemistry through ion exchange reactions with organic and inorganic cations. These two characteristics are, of course, interrelated, since the degree of dispersion of layered silicate in a particular polymer matrix depends on the interlayer cation. Pristine layered silicates usually contain hydrated Na⁺ or K⁺ ions [1]. Obviously, in this pristine state, layered silicates are only miscible with hydrophilic polymers such as poly(ethylene oxide) (PEO) [36] or poly(vinyl alcohol) (PVA) [37]. To render layered silicates miscible with various polymer matrices, one must convert the normally hydrophilic silicate surface to an organophilic one, making the intercalation of many polymers possible. Generally, this can be done by ion-exchange reactions with cationic surfactants including primary, secondary, tertiary, and quaternary alkylammonium or alkylphosphonium cations. Alkylam- monium or alkylphosphonium cations in the organo- silicates lower the surface energy of the inorganic host and improve the wetting characteristics of the polymer matrix, and result in a larger interlayer spacing.

Additionally, the alkylammonium or alkylphosphonium cations can provide functional groups that can react with the polymer matrix, or in some cases initiate the polymerization of monomers to improve the strength of the interface between the inorganic and the polymer matrix [1].

Preparative Methods

Intercalation of polymer chains inside layered silicate galleries has proven to be a successful approach to fab- ricate PLS-nanocomposites. Thermoplastic polymers based PLS-nanocomposites may be prepared by one of the three following routes: (i) in-situ intercalative polymerization of monomer, (ii) intercalation of polymer or pre-polymer from solution, and (iii) melt-intercalation.

The second method may be processed using emulsion, whereas when the good solvent is not available, researchers have to choose either in-situ polymerization method (i) or direct melt blending (iii).

The direct melt-intercalation process involves annealing a mixture of the polymer and layered silicate (organically modified or not) above the softening point of the polymer, statically or under shear. While annealing, the polymer chains diffuse from the bulk polymer melt into the galleries between the silicate layers. A range of nanocomposites with structures from intercalated to exfoliate can be obtained, depending on the degree of penetration

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of the polymer chains into the silicate galleries. This process was found to be completely reversible in the polystyrene (PS) based nanocomposite because PS has almost no interaction with the organically modified layered silicate (OMLS) surface, but when polymer chains have some favorable interaction with silicate surface, this process found to be irreversible. However, for a large va- riety of nanocomposites, this method requires often the use of extrusion process in order to get a homogeneous dispersion of intercalated silicate layers in the polymer matrix.

Thermodynamics and Molecular Modelling of PLS Nanocomposite Formation

The solution intercalation method is based on a solvent system in which the polymer or pre-polymer is soluble and the silicate layers are swellable. The layered silicate is first swollen in a solvent, such as water, chloroform, or toluene. When the polymer and layered silicate solutions are mixed, the polymer chains intercalate and displace the solvent within the interlayer of the silicate. Upon solvent removal, the intercalated structure remains, resulting in polymer/layered silicate nanocomposite.

For the overall process, in which polymer is exchanged with the previously intercalated solvent in the gallery, a negative variation in the Gibbs free energy is required.

The driving force for the polymer intercalation into layered silicate from solution is the entropy gained by desorption of solvent molecules, which compensates for the decreased entropy of the confined, intercalated chains.

Using this method, intercalation only occurs for certain polymer/solvent pairs. This method is good for the intercalation of polymers with little or no polarity into layered structures, and facilitates production of thin films with polymer-oriented clay intercalated layers.

In recent years, the melt intercalation technique has be- come the standard for the preparation of PLS-nanocomposites. During polymer intercalation from solution, a relatively large number of solvent molecules have to be desorbed from the host to accommodate the incoming polymer chains. The desorbed solvent molecules gain one translational degree of freedom, and the resulting entropic gain compensates for the decrease in conforma- tional entropy of the confined polymer chains. Therefore, there are many advantages to direct melt intercalation over solution intercalation. For example, direct melt intercalation is highly specific for the polymer, leading to new hybrids that were previously inaccessible.

So far, experimental results indicate that the outcome of polymer chains intercalation into two dimensional silicate galleries depends critically on layered silicate surface functionalization and constituent interactions. Pre-

sent author observes that (a) an optimal interlayer structure on the OMLS, with respect to the number per unit area and size of surfactant chains, is most favorable for nanocomposite formation, and (b) polymer intercalation depends on the existence of polar interactions between the OMLS and the polymer matrix.

In order to understand the thermodynamic issue associated with nanocomposite formation during melt-intercalation, Vaia and coworkers [38,39] applied a mean‐

field statistical lattice model, reporting that calculations based on the mean field theory agree well with experimental results. Although there is entropy losses associated with the confinement of a polymer melt during nanocomposite formation, this process is allowed because there is an entropy gain associated with the layer separation, resulting in a net entropy change near to zero.

Thus, from the theoretical model, the outcome of nano- composite formation via polymer melt intercalation de- pends primarily on energetic factors, which may be determined from the surface energies of the polymer and OMLS.

Based on the Vaia and coworkers [38,39] study and the construction of product maps, general guidelines may be established for selecting potentially compatible polymer/OMLS systems. Initially, the interlayer structure of the OMLS should be optimized in order to maximize the configurational freedom of the functionalizing chains upon layer separation, and to maximize potential interaction sites at the interlayer surface. For these systems, the optimal structures exhibit a slightly more extensive chain arrangement than with a pseudo-bilayer. Polymers containing polar groups capable of associative interactions, such as Lewis-acid/base interactions or hydrogen bonding, lead to intercalation. The greater the polar- izability or hydrophilicity of the polymer, the shorter the functional groups in the OMLS should be in order to minimize unfavorable interactions between the aliphatic chains and the polymer.

One of the great advantages of the Vaia mean-field statistical lattice model is the ability to determine analyti- cally the effect of various aspects of the nanocomposite formation. According to this model, the variation of the free energy of mixing and subsequent dependence on enthalpic and entropic factor, suggest the formation of three possible structures-phase separated, intercalated and exfoliated. Although, Vaia model able to address some of the fundamental and qualitative thermodynamic issues associated with the nanocomposite formation, however, some of the assumptions such as the separation of configurational terms and intermolecular interaction and the further separation of the entropic behavior of the con- stituents, somewhat limit the usefulness of the model.

Not only that, this model is based on nanocomposites where polymer chains are completely tethered with the

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silicate surface, which is not the case for most of the polymer nanocomposites.

To overcome the limitation of Vaia model, Balazs and coworkers [40] proposed a model based on a self-consistent field (SCF) calculation, such as the Fleer and Scheutjens theory [41]. In the Fleer and Scheutjens theory, the phase behavior of polymer systems is modeled by combining Markov chain statistics with a mean field approximation. These calculations involve a planar lattice where lattice spacing represents the length of a statistical segment within a polymer chain. Details regarding this theory can be found in ref. [41]. Using this method Balazs and coworkers tried to calculate the interactions between two surfactant-coated surfaces and a polymer melt. They considered two planar surfaces that lie paral- lel to each other in the xy-plane and investigated the ef- fect of increasing the separation between the surfaces in the z-direction. The two surfaces are effectively im- mersed within a polymer melt. As the separation between the surfaces is increased, polymer from the surrounding bath penetrates the gap between these walls. Each surface is covered with monodispersed end-grafted chains, i.e.

surfactants. Now if χsurf represents the Flory-Huggins interaction parameter between the polymers and the under- lying solid substrate and χssurf represents the Flory‐

Huggins interaction between the surfactant and surface.

Therefore, χsurf -χssurf = 0. It should be noted here; in their calculation Balazs and coworkers did not consider electrostatic interaction.

Their calculations show that on increasing the attraction between the polymers and the modified surfaces are qualitatively similar to observation as made by Vaia and coworkers [38,39]. However, Balazs and coworkers found that the actual phase behavior and morphology of the mixer can be affected by the kinetics of the polymers penetrate the gap between the plates. At the beginning, the polymer chains have to penetrate the space between the silicate layers from an outer edge and then diffuse to- ward the centre of the gallery. Now if we consider the case where χsurf < 0 and thus, the polymer and surface experience an attractive interaction. In this case as the polymer diffuses through the energetically favorable gallery, it maximizes contact with the two confining layers.

As a results, the polymer ‘glues’ the two surfaces together as it moves through the interlayer. This ‘fused’ con- dition could represent a kinetically trapped state and con- sequently, increasing the attraction between the polymer and layered silicate sheets would only lead to intercalated, rather than exfoliated structures. On the other hand, in the case where χsurf > 0, the polymer can separate the sheets, as the chain tries to retain its coil-like con- formation and gain entropy. However, recent literatures revealed that the melt mixing of organically modified layered silicates and almost no attractive polymer ma-

trices always lead to the formation of phase separated structure. The SCF calculations and phase diagrams lead to the same conclusion.

To overcome this problem, Balazs and coworkers then proposed the scheme of using a mixture of functionalized and non-functionalized polymers for the melts [42].

While the stickers at the chain ends are highly attracted to the surface, the remainder of polymer does not react with the substrate. Thus, as the polymer chain penetrate the sheets, the majority of the chain is not likely to glue the surfaces together.

Balazs and coworkers [43,44] also proposed a simple model that describes the nematic ordering in the polymer-layered silicate systems. Because of the very high degree of anisotropy (a typical layered silicate platelet is

∼1 nm in thickness and 100∼200 nm in diameter), layered silicate particles experience strong orientational ordering at low volume fractions and can form liquid crystalline phases such as nematic, smectic, or columnar, in addition to traditional liquid and solid phases (see Figure 4). Starting from the Onsager free energy functional for the nematic ordering of rigid rods, they developed a modified expression to combine the disk orientational and positional entropy, steric excluded-volume effects, translational entropy of the polymer, and finally the Flory-Huggins enthalpic interaction. The resulting iso- thropic-nematic phase diagram correctly represents many important features, such as the role of shape anisotropy in depressing the ordering transition and the increase in the size of the immiscibility region with increases in the polymer chain length. Unlike most of the phenomeno- logical theories of polymer-liquid crystal systems [45- 48], in the Onsager-type model the features of the phase diagram are directly derived from the geometric characteristics of the anisotropic component.

Balazs and coworkers first modified and expanded Onsager theory by including nematic, smectic and columnar crystalline phases. To do calculation, they also adopted Somoza-Tarazona formalism of density functional theory (DFT) and then incorporated expressions that describe the entropy of mixing between the different components and the enthalphic interaction between the platelets [34,49]. The resulting free energy function can be minimized with respect to both the orientational and positional single-particle distribution function of the platelets, and thus, potentially, all phases and coexistence regions can be determined. The resulting phase diagram was shown to exhibit a strong dependence on the shape anisotropy of the layered silicate particles, the polymer chain length, and the strength of the interparticle interaction. In particular, an increase in the shape anisotropy for oblate ellipsoidal filler particles leads to the broadening of the nematic phase at the expense of the isotropic region. The increase in the polymer chain

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Figure 4. Possible mesophases of oblate uniaxial particles dispersed in a polymer: (a) isotropic, (b) nematic, (c) smectic A, (d) col- umnar, and (e) crystal. The nematic director n in ordered phases is aligned along Z axis; the disks lie in the XY plane. Dashed lines show smectic layers (c) and columns (d). Reprinted from ref. [44] with permission from American Chemical Society, USA.

length leads to the formation of the crystalline and/or liquid crystalline mesophases and promotes segregation between polymerrich regions and filler particles. Finally, an increase in the strength of the interparticle potential leads to the complete elimination of the nematic phase and to the direct coexistence between isotropic and crystal or columnar phases. The only limitation of this model is that this model cannot determine the topology of the phase diagram and the nature of the ordered phases for intermediate and high volume fractions of colloidal particles.

The huge interfacial area and the nanoscopic dimensions between nanoelements differentiate PLS-nanocomposites from conventional composites and filled

plastics. The dominance of interfacial regions resulting from the nanoscopic phase dimensions implies that the behavior of PLS nanocomposites cannot be understood by simple scaling arguments that begin with the behavior of conventional polymer composites. Since an interface limits the number of conformations that polymer molecules can adopt, the free energy of the polymer at the interfacial region is fundamentally different from that of polymer molecules far away from the interface i.e. bulk.

The influence of interface always depends on the fundamental length scale of the adjacent matrix. In the case of polymer molecules, this is of the order of the radius of gyration of a polymer chain, Rg and this is equal to 5∼10 nm. For this reason, in nanocomposites with a very few

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volume percent of dispersed nanofillers in polymer matrix, the entire matrix may be considered to be a nano- scopically confined interfacial polymer. The restrictions in chain conformations will alter molecular mobility, relaxation behavior, free volume, and thermal transition such as the glass transition temperature.

Recent molecular dynamics simulation studies have shown that the dynamics of the polymer chains undergo radical changes at the interfacial region. For example, both the chain mobility and the chain relaxation times can be slowed by three orders of magnitude near phys- isorbing surfaces. Not only that, extensive surface forces apparatus experiments report how these novel dynamics of nanoconfined polymers are manifested through viscosity increases, two values orders of magnitude higher than the bulk values solid-like responses to imposed shear, and confinement induced ‘sluggish’ dynamics that suggest the existence of a ‘pinned’, ‘immobilized’ layer adjacent to the confining mica surface.

The kinetics of intercalation of polymer chains into the silicate galleries to form layered nanocomposites has been studied by Vaia and coworkers [50]. They have investigated the kinetics of the intercalation of PS above its entanglement molecular weight into octadecylammonium exchanged fluorohectorite. XRD reveals the average layer spacing in the unintercalated silicate to be 2.13 nm.

During intercalation, this spacing increases to 3.13 nm.

Measurement of this layer spacing as a function of time during intercalation yields a time-dependent fraction of intercalated silicate that is directly comparable to the time dependent number of beads in the slit, χ(t) and it is well fit by the prediction of a continuum diffusion model.

In their model Vaia and coworkers considered the diffusion of polymer chains into an empty cylinder with a permeable wall and impermeable caps. The diffusion coefficient determined by fitting χ(t) to this model, effec- tive diffusion coefficient, Deff, is large compared to the equilibrium self-diffusion coefficient in bulk polystyrene of the appropriate molecular weight. These results suggested that the process of intercalation of polymer chains into the two dimensional silicate galleries is limited by the transport of the polymer chains into the primary particles of the silicate and not specifically by transport of polymer chains inside the silicate galleries.

Now to understand the detail formation kinetics and physical properties of PLS-nanocomposites we need a clear molecular picture of the structure and dynamics of confined polymers. Researchers generally used the surfaces force apparatus [51-55] and computation studies [56- 67] to understand the behavior of confined fluids.

Confinement of a fluid on length scales comparable to the molecular size has been demonstrated to dramatically alter its structure and dynamical pictures [51-67]. For example, confined fluids have been shown to solidify or

vertify at temperatures well above the bulk transition temperature [65]. Because of the strong interactions between the confined molecules and the atoms or molecules of the confining medium, the mobility of the molecules in the confined environment is greatly reduced compared to the bulk [51,68-71]. Molecules in films of nanometer thickness organize in layers parallel to the surface. However, the confining medium induces two dimensional orders in these layers [67]. On the other hand, in certain circumstances, confinement phenomenon may also have opposite effect of enhancing molecular mobi- lities in a supercooled thin film, relative to the bulk.

Lee and coworkers [72] have presented an investigation of the molecular mechanism of polymer melt intercalation using molecular dynamics simulations. They tried to find out the motion of polymer chains from a bulk melt into a confined volume. In their model they represented macromolecules by bead-spring chains, leave a reservoir of bulk melt to enter a slit of rectangular cross section and fixed dimension. They adopted a coarse- grained description of polymers, because such a picture has been demonstrated to provide a useful description of melt dynamics over longer time scales than would be ac- cessible with an atomistic model [73,74]. They also considered a slit of fixed dimension to understand the transport of a polymer melt from the bulk into a confined volume of fixed dimension. However, they did not consider the presence of surfactant molecules in the slit and the swelling of the slit during intercalation.

They found that the intercalation process can be approx- imately characterized by an effective diffusion coefficient that is twice as large as the equilibrium self-con- sistence in the bulk melt. Increasing the polymer-silicate interactions is found to induce spontaneous intercalation, but for a high-polymer silicate affinity, the amount of intercalated material at a given time is reduced compared to the case of a weaker polymer-silicate attraction. The crossover from polymer-silicate miscibility to intercalated structures with increasing polymer-surface affinity has already been mentioned by Vaia and coworkers [50].

However, their study suggested that an important role may be played by the relaxations of polymer bridges that connect the two silicate surfaces. The number of these bridges, as well as their dynamical properties, will be controlled by chain length [75].

Manias and coworkers [76] extended Vaia and coworkers studies and presented a systematic study of the kinetics of polymers entering 2 nm wide galleries of mica- type silicates as a function of polymer molecular weight and polymer-surface interactions. They varied the polymer-surface interactions in two ways: either through changing the surface modification, i.e. by varying the organic coverage or through attaching strongly interacting sites-sticky groups-along the polymer chain. The poly-

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mer-surface affinity is the one of the most crucial param- eters during nanocomposite preparation, since it controls the polymer-surface monomeric friction coefficient and thus determines the motion of the polymer next to a solid surface. The interaction between the polymer and silicate surface can be controlled in two novel and well-controlled ways: (i) keeping the polymer the same and modifying the silicate surface, via controlled surface covering by surfactants of varying length at the same grafting density, and (ii) keeping the organically modified surface the same and modifying the polymer friction coefficient, by attaching along the polymer chain a controlled amount of groups that interact strongly with the silicate surface.

They [76] used the PS as polymer and the same octadecylammonium modified fluorohectorite to study the intercalation process. While Vaia and coworkers [50] made all XRD measurements on similar hybrids during in-situ annealing; Manias and coworkers [76] used ex-situ sample for XRD and other measurements. The ex-situ method has several advantages over the in-situ method: First, one can able to anneal the sample under vacuum, thus re- ducing any polymer chain degradation. Second, one can heat and subsequently investigate the same side/surface of the pellet sample by XRD, a procedure that provides a much more accurate control of the annealing temperature than in-situ case, where pellet is heated in one side and is studied by XRD on the opposite side. They made several comparative studies by using concurrent in-situ small-angle neutron scattering (SANS) and intermediate-angle neutron scattering (IANS) to monitor the changes in di- mensions of the polymer, i.e. Rg, and also to follow the changes in the single chain scattering function during intercalation.

Results of concurrent SANS and IANS studies shown that during intercalation the silicate gallery expansion directly reflects the motion inside the 2 nm slit pore of the polymer chains. For the same polymer and the same annealing temperature, they found the experimentally measured effective diffusion coefficient depends strongly on the surfactant used for the modification of fluorohectorite. For several different polymer molecular weights and annealing temperatures they observed that Deff increases markedly with longer surfactant lengths and much more than is expected just from the enhancement of polymer mobility resulting from the polymer dilution by small hydrocarbon oligomers. Longer surfactants result in less silicate surface area exposed to the polymer, thus effectively reduced the density attraction sites. On the other hand, introducing controlled amount of groups along the polymer chain that interact strongly with the silicate surface, resulting in a strong decrease of Deff. Therefore, increasing either the density or the strength of these attractive sites leads to much slower intercalation kinetics. However, an increase in site density or strength

must also increase the driving force for intercalation; evi- dently such increases depress the friction coefficient ζ much more strongly.

To understand the atomistic details of the structure of the confined-intercalated-polystyrene chains inside the two dimensional silicate gallery, Manias and coworkers also carried out molecular dynamic simulation. Details regarding simulation can be found in ref. [76]. They used the rotational-isomeric-state (RIS) model to create initial polymer conformations of PS oligomers. Conformations that fit in the interlayer gallery were chosen, and the PS chains were equilibrated by an off-lattice Monte Carlo scheme that employed small random displacements of the backbone atoms and orientational biased Monte Carlo rotations of the phenyl rings; at the same time, the surfactants were equilibrated by a configurational biased scheme in coexistence with the polymer chains. After equilibrium, MD simulations were used to obtain the structure and density profiles of the intercalated polymer/surfactant films. The numbers of polymer chains and alkylammonium surfactants were chosen so as to match the densities found in the experimental studies. The results suggested that the confined film adopts a layered structure normal to the solid surfaces, with the polar phe- nyls dominating the organic materials adsorbed on the walls, and the aliphatic groups predominantly in the center of the pore. ¹H‐²⁹Si cross-polarized nuclear magnetic resonance (NMR) measurements revealed a coexistence of ultra-fast and solid-like slow segmental dynamics throughout a wide temperature range, below and above Tg, for both the styrene phenyl and the backbone groups.

The mobile moieties concentrate at the center of the slit pore, especially for the higher temperatures. This leads to a strong density inhomogeneity in the direction normal to the surface. Fast dynamics occur in the lower density regions, whereas slower dynamics occur in high-segment‐

density regions close to the surface. This heterogeneous mobility combined with an observed persistence of mobility below the bulk glass transition temperature has im- plications to nanocomposite properties.

Rheology

Let us begin with‐what is rheology? Rheology is the study of the flow of matter. The name derives from the greek word “rheo” which means “to flow”. A solid will usually respond to a stress force by deforming and storing energy elastically. A liquid, however, will flow and dissipate energy continuously in viscous losses. A Newtonian liquid has a linear relationship between the shear rate and stress. Complex fluids are interesting because they exhibit both an elastic and viscous responses and exhibit a non‐linear relationship between

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Figure 5. Schematic diagram describing the end‐tethered nanocomposites. The layered silicates are highly anisotropic with a thickness of 1nm and lateral dimensions (length and width) ranging from ∼100 nm to a few microns. The polymer chains are tethered to the surface via ionic interactions between the silicate layer and the polymer-end. Reproduced from ref [78] by permission of American Chemical Society.

the shear rate and strain.

One of the properties often dealt within rheology is viscosity which measures how thick a fluid is. Some materials are intermediate between solids and fluids and the viscosity is not enough to characterize them. A solid material can be described by its elasticity or resilience:

when it is deformed it will store the energy and fight back. Imagine a spring that regains its original shape after being deformed. The other extreme is a fluid which stores no energy while deformed and just flows. A viscoelastic material is intermediate and stores some energy and flows a little when deformed. Consider a system where a micron-sized probe particle is doped inside a rheologically complex fluid such as polymer network and Brownian fluctuations are recorded. Now question is how to predict the mechanical properties of the medium from the observed Brownian motion of the particle. If the fluid is Newtonian and the probe molecule is spherical, one can easily determine the viscosity of the medium from the measured diffusion coefficient by using the standard Stokes-Einstein (SE) relation. How- ever, a polymer network is generally viscoelastic, with a complex shear modulus having both elastic i.e. solid-like and viscos i.e. liquid-like components of similar magni- tudes over a large range of frequencies. Now question is- is it possible to extend the SE relation to determine shear complex modulus for such a system? An affirmative answer would allow one to measure the local rheological

properties not only in inhomogeneous materials but also at high frequencies that are inaccessible using mechanical rheometer. This subject is known as ‘microrheology’

[77].

Now consider a similar system where nanoscopic-sized particles are dispersed in polymer network, enjoying rap- id development in recent years, known as ‘nanorheol- ogy’. Here we are going to discuss about the melt-rheological behavior of polymer materials where nanometer- size highly anisotropic particles are dispersed, known as polymer nanocomposites. In this section author survey some of the current development in the melt-state rheological behavior of polymer nanocomposites based on layered silicates.

The measurement of rheological behavior of any polymeric material under molten state is crucial to gain fundamental understanding of its processability. In the case of PLS-nanocomposites, the melt-rheological properties are also helpful to find out the degree of polymer-filler interactions and the structure-property relationship in nanocomposites. This is because rheological material functions are strongly influenced by the structure and the interfacial properties.

Linear Dynamic Rheological Properties

Small deformation amplitude oscillatory shear measurements of polymeric materials are generally performed by applying a time dependent strain

γ(t) = γ0 sin(ωt) (1)

and measuring the resultant shear stress

σ(t) = γ0 [G'sin(ωt) + G"cos(ωt) ] (2) where G' and G" are the storage and loss moduli, re- spectively and ω is the frequency.

Krishnamoorti and coworkers [78] first described the flow behavior of delaminated nanocomposites based on poly(ε-caprolactone) (PCL) and nylon 6 (N6). Both nanocomposites were prepared by in-situ polymerization method in the presence of organically modified MMT, which engendered direct grafting of macromolecular chains on the MMT surface-a hairy layered silicate platelets structure (see Figure 5). The authors coded these nanocomposites as ‘end-tethered polymer layered silicate nanocomposites’. Figure 6 represents the rheological master curves based on linear viscoelastic measurements for a series of end-tethered PCL-MMT nanocomposites.

The flow behavior of PCL nanocomposites differed significantly from that of the corresponding neat matrices, whereas the thermorheological properties of the nanocomposites were entirely determined by the behavior of the matrices. The slopes of log G'(ω) and

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Figure 6. Storage modulus (G') for PCL‐based silicate nan- ocomposites. Silicate loadings are indicated in the figure.

Master-curves were obtained by application of time-temper- ature superposition and shifted to T0 = 55 ^oC. Reproduced from ref. [78] by permission of American Chemical Society.

G"(ω) versus log aTω are much smaller than 2 and 1, respectively. Values of 2 and 1 are expected for non- crosslinked polymer melts, but large deviations occur, especially in the presence of a very large amount of layered silicate loading, which may be due to the formation of a network structure in the molten state.

However, nanocomposites based on the in-situ polymerization technique exhibit a fairly broad molar mass distribution in the polymer matrix, which hides the relevant structural information and impedes the inter- pretation of the results. The similar behavior was also observed in the case of end-tethered N6-MMT nanocomposites.

Ren and Krishnamoorti [79] then showed significant modification of flow behavior of nanocomposites when polymer chains are not tethered on the silicate surface.

The nanocmposites of polystyrene-1, 4-polyisoprene (7 mol% 3, 4 and 93 mol% 1, 4) diblock copolymer (PSPI 18) were prepared with (0.7, 2.1, 3.5, 6.7, and 9.5 wt%) dimethyl dioctadecyl ammonium modified MMT by solution mixing in toluene. The homogeneous solution was dried extensively at room temperature and subsequently annealed at 100 ^oC in a vacuum oven for ∼ 12 h to remove any remaining solvent and to facilitate complete polymer intercalation between the silicate layers. The XRD patterns of PSPI18/2C18-MMT and PS/2C18-MMT clearly indicates the formation of intercalated structures, whereas the composite prepared with 1, 4-polyisoprene (PI) showed no change in gallery height. The intercalation of 2C18-MMT by PS and not by PI is consistent with the results of previous experimental work and theories of Vaia and coworkers [37,38] and of Balazs and coworkers [39,40]. The

Figure 7. Time-temperature superposed linear storage modulus (a, top) and loss modulus (b, bottom) for the series of 2C18M-based PSPI18 intercalated hybrids. As expected, the moduli increase with increasing silicate loading at all frequencies. At high frequencies, the qualitative behavior of the storage and loss moduli are essentially unaffected. However, at low frequencies the frequency dependence of the moduli gradually changes from liquidlike to solidlike for nanocomposites with 6.7 and 9.5 wt% silicate. Reproduced from ref. [79] by permission of American Chemical Society.

intercalation of PS into the silicate layers may be due to the slight Lewis base character imparted by the phenyl ring in PS, leading to the favorable interactions with the 2C18-MMT layers.

All rheological measurements were conducted in the temperature range of 80∼105 °C. The time-temperature master curves for neat PSPI and various nanocomposites are presented in Figure 7. It is clear from the figures that viscoelastic behavior of PSPI significantly modified after nanocomposite formation. At high frequency region (where frequency > 10 rad.s), viscoelastic properties of PSPI were unaffected by the presence of silicate particle, while at the low frequency region (region where fre- quency < 10 rad.s), both G' and G" for the nano- composites exhibited a significant diminished frequency

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dependence and this behavior was more prominent with higher loading of silicate layers, similar behavior as observed in the case of end-tethered nanocomposite.

Similar behavior has also been noted for carbon black [80], calcium carbonate [81], and talc [82] filled polymers. However, in those cases the particle volume fractions were substantially higher, 5∼10 wt% for which strong frequency independent behavior is seen for PLS‐

nanocomposites.

Time-temperature superposition required simultaneous horizontal and vertical shifting of the shear moduli: bT G' and bT G" versus aTω, where aT is the horizontal shift factor, bT is the vertical shift factor and ω is the frequency. While aT for all nanocomposites followed the same WLF (Williama-Landel-Ferry) [83] dependence but bT did not. The authors attributed this could be the block architecture of the polymer examined, and the relative proximity of the order-disorder temperature to the temperature of the experimental measurements or may be due to formation of different types of structural orientation in the presence of organocaly.

To understand the viscoelastic behavior observed at low frequencies region in the dynamic measurements for the nanocomposites, they conducted stress relaxation measurements. The stress relaxation data of these nanocomposites in the terminal region showed a solid-like behavior. Like G', the effect was particularly pronounced at ≥ 6.7 wt% silicate loading. The authors attributed this behavior is due to the presence of stacks of intercalated silicate layers and those are randomly oriented in the polymer matrix, forming three dimension network structure.

These stacked intercalated silicate layers have only translational motion. A large-amplitude oscillatory shear is able to orient these structures and reduce the solid like behavior. Another explanation may be the formation of PI micro-domains. Results of XRD revealed that only PS block diffused into the interlayer galleries, leaving the PI block out of the silicate galleries. These PI blocks can form phase-separated micro-domain structures, which ac- tually hinder the orientation of stack intercalated silicate layers under small-amplitude shear.

Another explanation is the physical jamming of the dispersed stack intercalated silicate layers owing to their highly anisotropic nature. Transmission electron micro- scopy image revealed the formation of intercalated structure for the PSPI-based nanocomposites. On the basis of this mesoscopic structure and at low silicate loadings, authors suggested that, beyond a critical volume fraction, the tactoids and the individual layers are incapable of freely rotating and when subjected to shear are prevented from relaxing completely. This incomplete relaxation due to the physical jamming or percolation of the nanoscopic fillers leads to the presence of the pseudo-solid- like behavior observed in both the intercalated and ex-

foliated hybrids.

To verify above percolation theory, Krishnamoorti and coworkers estimated the percolation threshold for these hybrids on the basis of stack size of the intercalated silicate layers. They assumed a hypothetical hydrodynamic sphere surrounding each tactoid and the percolation of these hydrodynamic spheres to signify the onset of incomplete relaxation and the presence of pseudo-solid- like behavior in the intercalated nanocomposites. To find out the relation between the number of silicate layers per tactoid, nper and the layered silicate weight fraction at percolation, wsil, per, as

 _

 





_{ }_{ }  _{ }_

_{ }_{ } 





_

_

_ (3) where Rh is the radius of the hydrodynamic volume (in this case equivalent to the radius of the disk-like layered silicates), øperis the percolation volume fraction, hsil is the thickness of the silicate layers (generally 1 nm), ρorg

and ρsil are the densities of the organic component and layered silicate respectively. By considering the uniform disks diameter (2Rh) of 0.5 µm and the wsil, per to be 0.067, they calculated the average tactoid size to be around 30 layers. This calculated value was well matched with the value calculated from XRD pattern and TEM measurement. Furthermore, this tactoid size implies that the effective anisotropy associated with the filler is 5∼

10. It is this anisotropy, along with the random relative arrangement of the tactoids that leads to the observation of the percolation phenomenon at extremely low loading of the silicate.

Krishnamoorti and coworkers also used above equation to estimate the percolation threshold for the exfoliated nanocomposites by assuming the same geometrical pa- rameters as considered previously. They suggested that in the case of exfoliated nanocomposites, there exist considerable local orientational order which has enormous consequences for the physical and mechanical properties of the hybrids. This orientational ordering along with the end-tethering of the polymer chains on the layered silicate surface in fact responsible for the non-terminal intermediate frequency response in the PCL and N6-based nanocomposites.

Over the last few years, the linear dynamic and steady rheological properties of several nanocomposites have been examined for a wide range of polymer matrices including N6 with various matrix molecular weights [84], polyamides [85-91], polystyrene [92-96], poly- propylene [97-102], polyethylene [103], ethylene vinyl copolymers [104-106], polycarbonate [107-109], poly (ethylene oxide) [110-112], poly(ε-caprolactone) [113, 114], poly(ethylene terephthalate) [115], polylactide

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Figure 8. Complex viscosity versus frequency from a dynamic parallel plate rheometer (solid points) and steady shear viscosity ver- sus shear rate from a capillary rheometer (open points) at 240 °C for (a) pure HMW and its (HE)2M1R1 organoclay nanocomposite, (b) pure MMW and its (HE)2M1R1 organoclay nanocomposite. The nanocomposites contain ∼3.0 wt% MMT. Reproduced from ref.

[84] by permission of Elsevier Science Ltd., UK.

[116-118], poly(butylene succinate) [119-121], poly [(butylene succinate)-co-adipate] [122,123], synthetic biodegradable aliphatic polyester (BAP) [124,125], liquid crystal polymer [126,127], PLS-nanocomposites based on immiscible and miscible polymers blend [128, 129], etc.

Choi and coworers described the melt-state rheological behavior of a series of PEO/organoclay nanocomposites [111,112]. To understand both the effect of degree of dispersion of silicate layers and the interaction between the organoclay surface and polymer matrix, they used three types of surfactants modified montmorillonite. First, they selected two different organoclays, which have different modifier concentration but the same alkylamonium salt.

Second, they differentiated the types of alkylamonium salts having the same organic modifier. All nanocompo- sites were prepared via solvent casting method using chloroform as a co-solvent. XRD patterns and TEM ob- servations revealed the formation of intercalated nanocomposites. Rheological properties measurements

under molten state revealed that an increase in shear viscosity and storage and loss moduli of nanocomposites with clay content. Like other properties, they also fond that rheological properties of nanocomposites directly related to the degree of dispersion of silicate layers in the polymer matrix and also the level of interfacial interactions between the layered silicate surface and polymer chains. They concluded that linear rheological properties are strongly correlated with the mesoscopic structure and it was postulated that the molecular weight and interaction strength would affect the mesoscopic structure and the rheological properties of these nanocomposites.

Fornes and coworkers [84] have conducted dynamic and steady shear capillary experiments over a large range of frequencies and shear rates of pure N6 with different molecular weights, and their nanocomposites with OMLS. Figure 8 shows bilogarithmic plots of the complex viscosity, |η*| vs. ω at 240 ^oC for pure N6 and (HE)2M1R1 nanocomposites based on (a) high molecular weight (HMW), (b) medium molecular weight (MMW),

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Figure 9. Influence of frequency on shear storage modulus for pure LMW, and for LMW, MMW and HMW based nanocomposites. Reproduced from ref. [84] by permission of Elsevier Science Ltd., UK.

and (c) low molecular weight (LMW), obtained using the parallel plate oscillating rheometer. Figure 8 also shows a bilogarithmic plot of the steady-state shear viscosity η versus shear rate  , obtained using a capillary rheometer.

Inspection of these figures reveals a significant difference between the nanocomposites, particularly at low frequencies. The HMW-based nanocomposites show very strong non-Newtonian behavior, and this is more pronounced at low frequencies. On the other hand, the magnitude of the non-Newtonian behavior gradually decreases with decreasing molecular weight of the matrix, and with LMW it behaves like pure polymer. This trend is more clearly observed in the plot of G' vs. ω, due to the extreme sensitivity of G' towards dispersed morphol- ogy in the molten state (see Figure 9). The difference in the terminal zone behavior may be due to different ex- tents of exfoliation of the clay particles in the three types of matrices.

At the other extreme, the steady shear capillary rheology shows a trend with respect to the matrix molecular weight. The HMW and MMW-based nanocomposites exhibit lower viscosities compared to that of their corresponding matrices, whereas the viscosities of LMW- based nanocomposites are higher than the pure matrix.

This behavior is also due the higher degree of exfoliation in the case of HMW and MMW-based nanocomposites compared to the LMW-based nanocomposite.

Finally they considered the differences in the melt viscosity among the three systems. Over the range of frequencies and shear rates tested, the melt viscosity of the

three systems follows the order HMW > MMW > LMW, and hence the resulting shear stresses exerted by the pure polymers also follow the same order. Therefore, during melt mixing the level of stress exerted on the OMLS by the LMW polyamide is significantly lower than those developed in the presence of HMW or MMW polyamides.

As a result, the break-up of layered silicate particles is much easier in the case of HMW polyamides, and layered silicate particle dispersion is ultimately improved.

The role of polymer molecular weight is believed to stem from an increase in the melt viscosity, facilitating the degradation of the taller stacks into shorter ones. The fi- nal step in exfoliation involves peeling away the platelets of the stacks one by one, and this takes time and requires the strong matrix-OMLS interaction to cause spontaneous wetting.

Aubry and coworkers were studied the dynamaic and steady flow properties of a polyamide-12 (PA12) melt layered silicate nanocomposites as a function of the silicate volume fraction, ø [88]. They found above a vol- ume fraction threshold øp is equal to 1.5 %, and below a critical strain γc, the storage (G') and loss moduli (G") were shown to exhibit a low-frequency plateau and the flow curve was shown to exhibit a stress, τy. The study of G', γc, and τy as a function of ø showed the energy needed for removing connectivity on a mesoscale did not depend on the silicate loading. These original properties were attributed to the existence, in the quiescent state, of mesoscopic domains composed of correlated silicate layers. Moreover, the steady shear response of all sam- ples at solid volume fractions above øp showed the existence of a critical shear rate is equal to 1/s, separating a behavior governed by the networked domains from a behavior dominated by the polymer matrix.

Gelfer and coworkers have investigated the particular relationship among the molecular structure of polymer chains, morphology, and rheology of nanocomposites prepared by melt-blending of Cloisite organoclays and ethylene-co-vinyl acetate (EVA) and ethylene-co-methyl acrylate (EMA) random copolymers [104,106]. TEM observation confirmed the mixed intercalated-exfoliated morphology in these nanocomposite systems. The melt- state rheological properties were very similar in EVA- and EMA-based nanocomposites. These materials exhibited a pseudo-solid like rheological behavior, which became more pronounced upon heating, especially at temperature above 180∼200 °C. Authors attributed this behavior is due to the physical gelation due to the formation of 3D tactoid network in the polymer matrix.

They found that this gelation behavior is directly related to the overall content of organoclay and the extent of miscibility between the organoclay and polymer matrix.

The SAXS data indicated that the silicate gallery spacing (d001), intercalated by EVA and EMA chains,

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Figure 10. G' vs. reduced frequency for PS100 matrix (black line), CPS (A), and C-PEA (W). Silicate loading: 5 wt%.

Reproduced from ref. [89] by permission of Wiley-VCH, Germany.

decreases with increasing temperature. At temperature above 200 °C, the desorption of surfactant in organoclay decreases the compability between the clay surface and the polymer matrix, which finally exhibited a LCST (lower critical solution temperature) type behavior between organoclays and polymer, enhancing the state of gelation. This behavior manifested itself by a reverse temperature dependence of viscoelastic properties and strong deviation of rheological behavior from the time‐

temperature superposition principle.

Hoffmann and coworkers [89] recently conducted rheological measurements of nanocomposites of polystyrene (PS100) with synthetic mica (ME100) modified with amine-terminated polystyrene (AT-PS8) and 2-phenyl- ethylamine (PEA) respectively, in order to verify the presence of a particle network formation via inter-particle interaction and self-assembly. Figure 10 represents G'(ω) vs. aTω for PS100 and its corresponding nanocomposites prepared with AT-PS8 modified ME100 (C- PS) and PEA modified ME100 (C-PEA). The rheological responses of PS100 and C-PEA are the same, whereas the rheological response of C-PS is completely different.

At the lowest frequencies, which correspond to the marked region III, G'(ω) strongly increases, and the slope approaches to zero. Such behavior is an indication of network formation involving the assembly of individual plates composed of silicate layers. In the regime of intermediate frequencies (region II, see Figure 10) the G'(ω) value is lower in comparison to that of PS100.

This might be due to dilution of the amine-terminated PS being below the entanglement molecular weight. On the other hand, in high frequency region (region I, Figure 10) the rheological behavior is the same for all systems.

Sinha Ray and coworkers [116-118] have conducted the

Figure 11. Reduced frequency dependence of storage modulus, G' and loss modulus, G" of PLA and various PLACNs.

Reprinted from ref. [118] by permission of Wiley-VCH, Germany.

dynamic oscillatory shear measurements of PLA nanocomposites with intercalated structure. Melt rheological measurements were conducted on Rheometric Dynamic Analyzer (RDAII) instrument using 25 mm diameter parallel plates with a sample thickness of ca. 1.5 mm and in the temperature range of 175∼205 °C. To avoid nonlinear response, the strain amplitude was fixed to 5 % to obtain reasonable signal intensities even at elevated temperature or low w. For each nanacomposite investigated, the limits of linear viscoelasticity were determined by performing strain sweeps at a series of fixed frequencies. The master curves were generated using the principle of time-temperature superposition and shifted to a reference temperature (Tref) of 175 °C, which was chosen as the most representative of a typical processing temperature of PLA.

The master curves for G' and G" of pure PLA and various nanocomposites with different weight percenta- ges of C18 MMT loading are presented in Figure 11. At high frequencies (aT.ω > 10), the viscoelastic behavior of all nanocomposites was the same. On the other hand, at low frequencies (aTω < 10), both moduli exhibited weak frequency dependence with increasing C18MMT content, with a gradual change of behavior from liquid- like [G' ∝ ω² and G" ∝ ω] to solid-like ( G' ≥ G") with

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Figure 12. Flow activation energy as a function of C18- MMTcontent. Reprinted from ref. [118] by permission of Wiley-VCH, Germany.

increasing C18MMT content.

The slope of G' and G" in the terminal region of the master curves of PLA matrix was 1.85 and 1 respectively. On the other hand, the slopes of G' and G"

were considerably lower for all PLACNs compared to those of pure PLA. In fact, for PLACNs with high C18MMT content, G' becomes nearly independent at low aTω and exceeds G", characteristic of materials exhibit- ing a pseudo-solid-like behavior.

Figure 12 represents the C18MMT content dependent (wt%) flow activation energy (Ea) of pure PLA and various nanocomposites obtained from an Arrhenius fit of master curves [118]. There was a significant increase of Ea for PLA/C18 MMT3 compared to that of pure PLA followed by a much slower increase with C18MMT content. This behavior may be due to the dispersion of intercalated and stacked C18MMT silicate layers in the PLA matrix.

The dynamic complex viscosity (｜η*｜) master curves for the pure PLA and nanocomposites, based on linear dynamic oscillatory shear measurements, are presented in Figure 13 [118]. At low aTω region ( < 10 rad.s^‐1), pure PLA exhibited almost Newtonian behavior while all nanocomposites showed very strong shear-thinning tendency. On the other hand, Mw and PDI of pure PLA and various nanocomposites were almost the same, thus the high viscosity of PLACNs were explained by the flow restrictions of polymer chains in the molten state due to the presence of MMT particles.

Figure 14 represents the master curves for G' and G" of neat PBS and various PBSCNs prepared with two different types of OMLS [120]. At all frequencies, both G' and G" of nanocomposites increased monotonically with increasing OMLS loading with the exception of PBS/C18MMT1 and PBS/qC16SAP1 for which viscoe-

Figure 13. Reduced frequency dependence of complex vis- cosity (｜η*｜) of PLA and various PLACNs. Reprinted from ref. [118] by permission of Wiley-VCH, Germany.

lastic behavior was almost identical to that obtained for neat PBS. At high frequencies (aTω< 5), both moduli exhibited weak frequency dependence with increasing clay content, which means that there are gradual changes ofbehavior from liquid-like to solid-like with increasing clay content.

Lepoittevin and coworkers [113] reported the melt rheological properties of PCL-based nanocomposites pre- pared by melt intercalation method. G' and G" of unfilled PCL and PCL filled with 3 wt% of MMA-Na⁺, MMT‐

Alk or MMT-(OH)2 were measured at 80 ºC in the frequency range from 16 to 10^‐2 Hz. The rheological behavior of the PCL filled with 3 wt% of MMT-Alk and MMT-(OH)2 was significantly different compared to unfilled PCL and PCL/MMT-Na composites. The fre- quency dependence of G' and G" was, however, per- turbed by organically modified MMT. The effect was dramatic in the slope of G' which drops from 2 to 0.14 and 0.24 for MMT-(OH)2 and MMT‐Alk, respectively.

The dependence of G' and G" on frequency with the filler content is presented in Figure 15 for the MMT- (OH)2 clay. When the clay content exceeded 1 wt%, not only the classical power laws for the frequency depend- ence of G' and G" were deeply modified, particularly in the case of G' but the moduli increased dramatically at low frequency. This is a characteristic of pseudo-solid- like behavior due to the formation of a network percolating clay lamellae. The same behavior was also observed in the case of PLA or PBS-based nanocomposites [118, 120,121].

Sinha Ray and his coworker recently described the melt-state rheological properties of poly[(butylene succinate)-co-adipate] (PBSA) nanocomposites [122].

Cloisite^® 30B (C30B) organoclay was used for the preparation of nanocomposites. Nanocomposites were

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Figure 14. (a) Reduced frequency dependence of the storage modulus,G'(ω), and the loss modulus, G"(ω), of PBS and various PBSCNs prepared with C18-mmt clay. Tref ) 120 °C. (b) Reduced frequency dependence of the storage modulus, G'(ω), and the loss modulus, G"(ω), of PBS and various PBSCNs prepared with qC16‐sap. Tref ) 120 °C. Reproduced from ref. [120] with permis- sion of American Chemical Society, USA.

prepared at 135 ^oC in a batch mixer. Nanocomposites (PBSACNs) prepared with three different amount of C30B of 3, 6, and 9 wt% were correspondingly abbreviated as PBSACN3, PBSACN6, and PBSACN9.

The structure of the nanocomposites was studied using XRD and TEM that revealed a coexistence of exfoliated and intercalated silicate layers dispersed in the PBSA matrix, regardless of the silicate loading.

The angular ω dependence of storage, G'(ω) and loss modulus, G"(ω) of neat PBSA and PBSACNs with three different wt% of C30B loading are presented in Figure 16, parts a and b, respectively. At all frequencies, both G'(ω) and G"(ω) for PBSACN6 and PBSACN9 in- crease with increasing C30B content with the exception of PBSACN3 for which the viscoelastic response is almost identical to that observed for neat PBSA. At high- ω region, the viscoelastic behavior of all PBSACNs is quite the same, with only a small systematic increase in G'(ω) with the C30B loading, indicating that the ob- served chain relaxation modes are almost unaffected by the presence of the layered silicate particles (Figure 16a and b). However, at the low-ω region, both dynamic moduli exhibit weak frequency dependence with C30B loading. For the 9 wt% of C30B loading, G'(ω) exceeds G"(ω) and becomes nearly independent of ω.

To verify the interesting viscoelastic behavior observed at low frequency region in the dynamic oscillatory shear measurements for the nanocomposites with high wt%

C30B loading, the linear stress relaxation measurements

were conducted. Results of linear relaxation behaviors of neat PBSA and PBSACNs are presented in Figure 17.

For any fixed time after imposition of strain, the relaxation modulus, G(t) increases with increasing C30B loading, similar to that observed in the dynamic oscillatory shear measurements.

Similar to the behavior observed at high frequencies in dynamic mode (Figure 16a and b), the stress relaxation modulus of all PBSACNs is quite of the same magnitude as that of the neat PBSA at short times, with only a slight increase with C30B content (Figure 17). However, at long time periods, the pure PBSA and PBSACN3 relaxes like a viscoelastic liquid, while the PBSACNs with 6 and 9 wt% of C30B loading behave like a pseudo-solid-like material.

Therefore, on the basis of both the dynamic oscillatory shear (Figure 16a and b) and linear stress relaxation measurements (see Figure 17), it is clear that C30B layered silicate has a profound effect on the long relaxation time of the nanocomposites. For PBSACNs with C30B loading in excess of 3 wt% , the liquid-like behavior observed in the case of pure PBSA gradually changes to pseudo-solid-like behavior.

A reasonable explanation for this low frequency viscoelastic behavior of PBSACNs is that the moderate interactions of the PBSA backbone with C30B surface lead to the high degree of confinement of polymer chains inside the silicate layers. As a result anisotropic silicate layers are fully dispersed in the PBSA matrix, as

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Figure 15. Storage modulus (G0) and loss modulus (G00) for unfilled PCL and PCL modified by 3 wt% of MMT-Na, MMT-Alk, MMT-(OH)2 at 80 C [160]. Reproduced from ref.

[113] by permission of Elsevier Ltd., UK.

observed in XRD patterns and TEM images [122].

Because of their highly anisotropic nature, the dispersed layered silicate particles would exhibit local correlations that finally cause the formation of micro-domains or mesoscopic structure [116,124,125]. Like in liquid crystalline and ordered block copolymer systems [130- 133] the presence of these micro-domains or mesoscopic structure causes only a slight increase of the low-ω elastic modulus that varies in low power-law fashion.

This is the case for low C30B concentration (∼3 %).

However, beyond this concentration, exfoliated and or disordered intercalated silicate layers form a network- type structure rendering the system highly elastic as revealed by the low frequency plateau. Such a plateau is similar to the one observed in rubber-toughened polymers [134,135] where rubber particles form a

Figure 16. Frequency dependence of (a) dynamic storage, G' (ω) and (b) loss, G"(ω) moduli of neat PBSA and various PBSACNs. Reprinted from ref. [122] by permission of Wiley‐

VCH, Germany.

percolating network that imparts the blend with solid-like behavior. Such a behavior can also be seen on the dynamic complex viscosity, η*(ω), shown in Figure 18. Little effect of C30B addition is observed at high frequencies, where the relaxation mechanism is mainly dominated by that of the PBSA matrix, whereas at low frequencies, the relaxation is that of particle-particle interactions inside the percolating network of the silicate layers.

In a recent literature Huang and Han reported the melt- state rheological behavior of PLS nanocomposites based on thermotropic liquid-crystalline polymers (TLCP) [126]. In their study they used two different types of TLCPs such as TLCP having pendent pyridyl group (PyHQ12) and TLCP having pendent phenylsulphonyl group (PSHQ12). They also used two different types of