Clay Intercalation and its Influence on the Morphology and Transport

Article pubs.acs.org/JPCC

Clay Intercalation and its Influence on the Morphology and Transport Properties of EVA/Clay Nanocomposites Runcy Wilson,† Sajeev Martin George,† Hanna J. Maria,† Tomás S. Plivelic,‡ Anil Kumar S,*,§ and Sabu Thomas*,†,⊥,∥,○ †

School of Chemical Science, Mahatma Gandhi University, Kottayam, Kerala, 686560, India MAX IV Laboratory, Lund University, SE-221 00 Lund, Sweden § Department of Chemistry, N.S.S College, Ottapalam, Kerala, 679103, India ⊥ Center for Nanoscience and Nanotechnology, Mahatma Gandhi University, P.D Hills, Kottayam, Kerala, 686560, India ∥ Universiti Teknologi MARA, Faculty of Applied Sciences, 40450 Shah Alam, Selongor Malaysia ○ Center of Excellence for Polymer Materials and Technologies, Tehnoloski park 24, 1000 Ljubljana, Slovenia ‡

ABSTRACT: Nanocomposite membranes based on poly(ethylene-co-vinyl acetate) copolymer (18% vinyl acetate content) and two different organomodified clays have been prepared by mechanical mixing using two roll mill method. The morphology of the nanocomposites was investigated using small angle X-ray scattering, scanning electron microscopy, and transmission electron microscopy. The mechanical and thermal studies were also performed using universal testing machine and differential scanning calorimeter, respectively. Samples with low filler content showed excellent dispersion of layered silicates resulting in a partially exfoliated structure. The diffusion and transport of organic solvents through the membranes have been investigated in detail as a function of clay content, nature of solvent and clay, and temperature in the temperature range of 28−70 °C. The influence of free volume on the transport properties of the membranes was studied using positron annihilation lifetime spectroscopy. The solvent uptake was minimum for composites with 3 wt % of filler, and it get increased with increasing filler content, which is presumably due to aggregation of clay filler at higher loading. The transport phenomenon was found to follow an anomalous mode. Activation parameters were estimated, and the molar mass between cross-links was calculated. Finally, the experimental transport data were compared with theoretical predictions.

1. INTRODUCTION

exfoliated hybrids, where the silicate layers are fully separated and dispersed homogeneously throughout the polymer matrix.7 The diffusion process is a kinetic parameter depending on the free volume within the material, segmental mobility of polymer chains, and the size of the penetrant molecule.8 The diffusion and transport in polymer matrix nanocomposites depend upon the nature of the fillers, the degree of adhesion, and their compatibility with the polymer matrix. If the filler used is inert and it is compatible with the polymer matrix, it will take up the free volume within the polymer matrix and create a tortuous path for the permeating molecules. The degree of tortuousity is dependent on the volume fraction and the shape and orientation of the particles. On the other hand if the filler is incompatible with the polymer matrix, voids tend to occur at the interface, which tends to increase the free volume of the system and thereby increases the permeability through it.9

Polymer−clay nanocomposites are a new class of composite materials consisting of a polymer matrix with dispersed clay nanoparticles. The interest in such systems (organic−inorganic hybrid material) is due to the fact that, the ultrafine or nano dispersion of filler, as well as the local interactions between the matrix and filler, lead to a higher level of properties than for equivalent micro- and macrocomposites.1 In recent years, more attention has been given to incorporate nanomaterials into the polymer matrices to obtain high performance nanocomposites. Polymer/clay nanocomposites have drawn considerable interest because of their enhanced properties, including flame resistance,2 mechanical properties,3 gas barrier properties,4 thermal stability,5,6 and biodegradability when compared to pristine polymers. The enhancement of the polymer performance depends on the microstructure of the polymer layered silicate nanocomposites which are normally classified into two types: intercalated hybrids, in which a part of the polymer chain intrudes between the silicate layers to form a well-ordered multilayer with alternating polymer and clay layers, and © 2012 American Chemical Society

Received: March 6, 2012 Revised: July 6, 2012 Published: August 15, 2012 20002

dx.doi.org/10.1021/jp302177y | J. Phys. Chem. C 2012, 116, 20002−20014

The Journal of Physical Chemistry C

Article

Drozdov et al.10 conducted moisture diffusion tests on vinyl ester resin−Montmorillonite clay nanocomposites and showed that the diffusivity decreases with higher clay content. Messersmith and Giannelis11 studied the permeability of liquids and gases in nanocomposites, and they observed that permeability in nanocomposites is dramatically reduced compared to the unfilled polymer, and they also reported that the permeability is significantly reduced in the nanocomposites compared to conventionally filled polymers having higher filler content. The incorporation of the nanoclay particles can considerably reduce the diffusion and permeation of molecules through the polymer matrix.12,13 The type of the modifier used and the amount of clay particle and its compatibility with the host polymer matrix greatly influence the permeability of the membrane. The diffusion of molecules and its kinetics, through nanoclay-based nanocomposites, were also examined by many researchers, and a substantial decrease in permeability in comparison with those of pristine polymer matrix was reported.14,15 The reduction in permeability was attributed to the extremely high aspect ratio of clay platelets, which increased the tortuosity of the diffusion path of the molecules into the nanocomposite. Ethylene−vinyl acetate copolymer (EVA) is a random copolymer synthesized from ethylene monomer and vinyl acetate comonomer. The copolymer is often produced by continuous bulk polymerization under high pressure similarly to low density poly ethylene (LDPE). In this study, nanocomposites with poly(ethylene-co-vinyl acetate) (EVA) and two different organically modified clay with varying clay loadings (3, 5, 7 wt %) were prepared. The morphology of nanocomposites was studied using scanning electron microscopy (SEM), transmission electron microscopy (TEM), and small angle X-ray scattering (SAXS). The effect of free volume on the transport properties was investigated by positron annihilation lifetime spectroscopy (PALS). The mechanical and thermal properties of the samples were also discussed. The transport properties of aromatic hydrocarbons through the nanocomposites were analyzed in detail and the experimental data compared with the theoretical predictions.

with a nip gap of 1.3 mm and at a friction ratio of 1:1.4. The nip gap, mill speed ratio, time of mixing, and temperature of the rolls were kept constant for all the mixes. The vulcanization behaviors of samples were studied by using Monsanto Rheometer. The cure times obtained were tabulated in Table 2. The sheeted out samples were compression molded in a Table 2. Cure Time of EVA/Clay Nanocomposites

Cloisite 10A Cloisite 25A

modifier concentration

2MBHT: dimethyl, benzyl, hydrogenatedtallow, quaternary ammonium 2MHTL8: dimethyl, dehydrogenated tallow, 2ethylhexyl quaternary ammonium

125 meq/100 g of clay 95 meq/100 g of clay

19.00 15.00 12.40 12.20 13.30 12.20 12.10

3. CHARACTERIZATION TECHNIQUES OF NANOCOMPOSITES 3.1. Small Angle X-ray Scattering (SAXS). The smallangle X-ray scattering (SAXS) experiments were carried out at I711 beamline of the MAX-lab Synchrotron, Sweden.16 A monochromatic beam of 1.1 Å wavelengths was used, and the sample detector distance was 1245.47 mm for all the samples. Two-dimensional pictures were recorded using a 2D-CCD detector (165 mm active area from MAR research, GmbH) with 10 min of data acquisition. Multilayer samples of 1 mm thickness were prepared by stacking pieces cut from a single film. The original orientation of the samples was preserved in the stack. The samples were placed in a multiple position sample holder and measured in an evacuated chamber. X-ray scattering data were analyzed with the program FIT2D.17 Average radial intensity profiles, I(q), as a function of the scattering vector q (q = 4π/λ sin(θ), where 2θ is the scattering angle and λ is the wavelength) were obtained by integrating the data in the complete image. For the comparison of the scattering curves, the intensities were normalized by the integrated intensity incident on the sample during the exposure, and corrected by sample absorption, parasitic scattering, and background detection. 3.2. Scanning Electron Microscopy (SEM). The SEM micrographs of the polymer/clay nanocomposites were taken in JEOL (Japan), Model JSM 6390 electron microscope with an accelerating voltage of 10 kV. The specimens were prepared by cryogenically fracturing the vulcanized sheets in liquid nitrogen and coating with platinum to avoid the electrostatic charge dissipation. 3.3. Transmission Electron Microscopy (TEM). Transmission electron micrographs of the nanocomposites were taken in JEOL JEM 2100 transmission electron microscope with an accelerating voltage of 200 keV. Ultrathin sections of bulk specimens (∼100 nm thickness) were obtained at −85 °C using an ultra microtome fitted with a diamond knife.

Table 1. Descriptive Properties of the Organo-Modified Nanoclays organic modifier used

cure time (min)

EVA A3 A5 A7 D3 D5 D7

hydraulic press at 160 °C under a load of 24.5 × 104 N .The samples were prepared with different clay content and designated as A3, A5, A7, D3, D5, and D7, where A and D represent the Cloisite 10A and Cloisite 25A organoclays, respectively. The subscript denotes the number of grams of nanoclay used per 100 g of the polymer, and EVA represents the pristine polymer with zero filler content.

2. EXPERIMENTAL SECTION 2.1. Materials. Poly(ethylene-co-vinyl acetate), EVA with vinyl acetate content of 18 mol %, was obtained from Taisox industries Limited, Taiwan. The cross-linking agent used was dicumyl peroxide (DCP). The nanoclay Cloisite 10A and Cloisite 25A were obtained from Southern Clay Products, United States. The descriptive properties of nanoclays are detailed in Table 1. The solvents benzene, toluene, and xylene used in the study were of reagent grade. 2.2. Sample Preparation. Organically modified nanoclay− EVA membranes were prepared by mechanical mixing using dicumyl peroxide as the curing agent. The amount of the DCP used was 2 g. The mixing was done in a two roll mixing mill

nanoclays

sample code

20003



Article

3.4. Differential Scanning Calorimetry (DSC). The nonisothermal crystallization and the melting experiments of samples were carried out with a differential scanning calorimeter instrument (a Perkin-Elmer Diamond DSC). Samples weighing about 10 mg were cut off for the DSC, and the temperature ranges from −30 to 150 °C at a heating rate of 10 °C min−1 under nitrogen atmosphere. The first cooling and the heating cycle were taken for the study. The crystallinity (Xc) of the samples was calculated according to equation Xc =

ΔHf × 100% ΔHf*

Q t (mol %) = [mass of solvent sorbed by the polymer] /[molecular mass of the solvent /initial mass of polymer] × 100

(2)

4. RESULTS AND DISCUSSION 4.1. Small Angle X-ray Scattering Analysis. The SAXS patterns of the EVA/clay nanocomposites were shown in Figure 1. The X-ray scattering methods (SAXS) have been used to characterize the morphology and structure of the polymer− silicate hybrid by monitoring the position, shape, and intensity of the nanocomposite characteristic peaks. When the insertion of polymer chains into the organoclay stacks occurs, an increase of the overall particle clay volume and corresponding interlayer silicate spacing could be obtained, which in turn give rise to the

(1)

where ΔH*f is the enthalpy of fusion of the perfect polyethylene (PE) crystal and ΔHf is the enthalpy of fusion of EVA samples, respectively. The value of ΔH*f for 100% crystalline PE is 277.1 J g−1.18 3.5. Mechanical Tests. The mechanical properties of the samples were studied using universal testing machine in accordance with ASTM D 412-2002. The instrument used was Tinius Oslen H50 kT UTM. The experiments were conducted for dumbbell shaped samples with gauge length of 12 mm at a cross head speed of 500 mm min−1 at room temperature to rupture by an electro mechanical machine equipped with a 50 kg load cell. Five specimens were considered for each experiment. 3.6. Positron Annihilation Life Time Spectroscopy (PALS). The positron lifetime spectrometer used in the present study consists of a fast−fast coincidence system with BaF2 scintillators coupled to photomultiplier tubes of type XP2020/ Q with quartz windows as detectors. A 17 μ Ci 22Na positron source deposited on a pure Kapton foil of 0.0127 mm thickness was placed between two identical pieces of samples under investigation. This sample source sandwich is positioned between the two detectors of PALS to acquire a lifetime spectrum. The spectrometer was operated at 220 ps time resolution. All lifetime measurements were performed at room temperature, and two to three positron lifetime spectra with more than a million counts under each spectrum were recorded. 3.7. Swelling Experiments. Circular samples were punched out from molded samples by means of a sharp edged steel die. The thickness of the sample was measured at several points using a micrometer screw gauge and the average was taken as the initial thickness of the sample. These samples were immersed in different solvents taken in test bottles and kept in a thermostatically controlled oven. At regular intervals of time, samples were removed and the surface was dried between filter paper. They were weighed immediately and then replaced into the diffusion bottles. The weighing was continued until equilibrium swelling was attained. Since each weighing was carried out with in 30 s, the error due to the evaporation of solvents may be neglected. The experiments were duplicate or triplicate in most cases, and the standard deviation was found to range from 0.07 to 0.1. The mol uptake (Qt) for the solvent by 0.1 kg of the polymer for a specific time, t, under solvent effects, is calculated using the expression.19

Figure 1. SAXS patterns of EVA nanocomposites. (a) EVA with Cloisite 10A nano clay. (b) EVA with Cloisite 25A nano clay. The curves were multiplied by an arbitrary scale factor to better visualize the sample behavior. The solid line indicates the peak positions with dspacings d1 and d2 for the lowest clay concentration on the samples. The arrows show the peaks associated with the lamellar structure of the semicrystalline polymeric matrix. 20004



Article

relatively better homogeneous dispersion of Cloisite 25A organoclay in the EVA matrix for all the clay content than Cloisite 10A based samples. From the images, it can be seen that the surface morphology changes according to the clay loading being rougher for larger amounts of clays. The formation of microfiller (agglomeration of individual silicate layers) is obtained due to the difficulties to achieve homogeneous dispersion and can result in deterioration of nanocomposite properties ((for example, poor reinforcing effect at higher loadings as shown later in the text). 4.3. Transmission Electron Micrographs. TEM images shown in Figures 3 and 4 provide visible evidence of the dispersion of clay platelets in the nanocomposites. The higher electron density of the silicates compared to EVA gives them a much darker appearance. The well-defined small and highly oriented clay stacks are responsible for individual silicate layers in the matrix, but these layers lack a very good distribution within the polymer matrix at higher clay loadings, and this result in the agglomeration of clay platelets. The overall morphology can be considered as a coexistence of exfoliated and intercalated silicate layers. Prasad et al. have already reported that besides the ethylene backbone providing a point of contact between the EVA chain and clay layers via the surface modifier, the VA group can help by attaching itself to an unmodified region of the clay layers that is virtually hydrophilic.23 The strong interaction between EVA and clay might occur to such extent that the pure clay of dense layered structure was shown to be broken apart and dispersed with individual platelets in the matrix polymer. These samples exhibit a better dispersed morphology consisting primarily of individually dispersed clay platelets. Peeterbreock et al.22 reported that the exfoliation and distribution of the clay nanoplatelets in the EVA matrix appear to depend on the nature of the clay organomodifier; they observed the best results for Cloisite 30B, which was organomodified by ammonium cations bearing hydroxyl groups. This was due to the interactions of hydroxyl groups of clay and the acetate functions of the EVA matrix. This extent of interaction can be varied by using different modifier. In this study, two different clays were used (Cloisite 10A and Cloisite 25A), and it can be noted the result was in excellent correlation with the SAXS analysis. 4.4. Differential Scanning Calorimetric Measurements. Figure 5a shows the melting behavior of the EVA/ clay nanocomposites. From the TEM analysis, it is seen that the polymer nanocomposites having partially exfoliated structure, the mobility of the chains has been restricted. The variation of the melting temperature can be related to the lower molecular mobility of the polymeric chains. The lower molecular mobility reduces the entropy gain connected to the melting transition, and this determined the increase of the melting temperature. At higher temperatures, multimelting peaks are seen for neat EVA. This may be caused by the nature of random copolymer (EVA), which leads to various degrees of the imperfection of the ethylene-sequence crystallite. It is concluded that thermally stable clay nanoplatelets with high aspect ratio will cause better thermal barrier effects on the polymer matrix. Figure 5b shows that the crystallization peak slightly shifts to higher temperature as the filler content increases, which indicates that the nanoclay particles acted as a heterophase nucleating agent and also the oriented clay platelets can increase the stereoregularity of the polyethylene macromolecular chain. Chaudhary et al.24 found

shifting of scattering peaks to lower angles. The widening of the scattering peaks and the decrease of the scattered intensity for different samples can be attributed to a beginning exfoliation of clay stacks. Scattering peaks cannot be seen in the case of fully exfoliated structures where the silicate layers are completely and uniformly dispersed into a continuous polymer matrix. The SAXS curves for neat EVA and nanocomposites (A and D series) are shown in Figure 1a and b. All patterns, including the pure polymer, present a first peak centered on q = 0.06 Å−1 (indicated by an arrow in both figures). This feature can be associated with the lamellar periodicity produced in semicrystalline polymers and is out of the scope of the present paper. On the other hand, at high q values (q > 0.1 Å−1) two additional peaks are observed in all nanocomposites curves. The d-spacings of the samples obtained from the f itted position of the maximum (qm) and using the relationship d = 2π/qm are summarized in Table 3 (named d1 and d2 for the first and Table 3. d-Spacing Values of Different Nanocomposites sample

d1 (Å)

d2 (Å)

reported dmc values in the literature (Å)

A3 A5 A7 D3 D5 D7

34.18 33.60 34.07 31.46 32.49 31.57

17.55 17.55 17.55 16.72 16.33 16.12

Cloisite 10 A20 19.2 Cloisite 25 A21 20.7

second peak, respectively). In order to fit the experimental data in this q range, the addition of two Lorentzian functions and an exponential baseline was used. Both sets of samples (Cloisite 10A and Cloisite 25A) present at the d1 peak, and in all cases, interlayer distances are larger than the basal distances, dmc, reported in the literature for pure organoclays (see Table 3).20,21 This is a clear indication that the nanocomposites present intercalated polymer/clay morphology. From the analysis of the d2 peaks, further morphological information can be extracted. Despite the fact that they have dvalues closer to satisfy a multiple relationship with d1-spacings (see Table 3), the similar or even smaller full width halfmaximum values (fwhm) of the d2 peak shapes seems to indicate that two periodic structures of silicate layers are present in the nanocomposites. This second periodic structure can be due to poorly intercalated polymer/clay structures or unmodified organoclay periodicity. This is especially clear for Cloisite 10A based nanocomposites where dmc and d2 are within the experimental error. From the intercalation spaces d1, a comparison between the effects of different organoclays and clay loadings in the nanocomposites can be established. A samples series (Cloisite 10A based nanocomposites) presents larger d1 values and larger degree of intercalation (d1−dmc) than D samples series (Cloisite 25A based nanocomposites). This was in agreement with the observations of Peeterbroeck et al.;22 according to them, an increase in interlayer spacings indicates a good affinity of EVA with the clay galleries. The last parameter defers in almost 4 Å between organoclays. The results can be an indication of better affinity between EVA chains and 2MBHT organic modifier than 2MHTL8. 4.2. Scanning Electron Micrographs. The scanning electron micrographs of pure clays Cloisite 10A and Cloisite 25 A are shown in Figure 2a and b, while the nanocomposites images are shown in Figure 2c−i. The micrographs show a 20005



Article

Figure 2. SEM micrographs for EVA/clay nanocomposites. (a) Cloisite 10A. (b) Cloisite 25A. (c) SEM image of A3. (d) SEM image of A5. (e) SEM image of A7. (f) SEM image of D3. (g) SEM image of D5. (h) SEM image of D7.

Figure 3. Transmission electron micrographs of EVA/Cloisite 10A nanocomposites. (a) Nanocomposite with 3 wt % of Cloisite 10A. (b)Nanocomposite with 5 wt % of Cloisite 10A. (c) Nanocomposite with 7 wt % of Cloisite 10A.

4.5. Mechanical Studies. Figures 6−9 show the effect of Cloisite 10A and Cloisite 25A content on tensile strength and elongation at break of EVA/clay nanocomposites. With a low concentration of clay, at 3 wt %, tensile strength is improved marginally in comparison to pure EVA (see Figures 7 and 8). Furthermore, the maximum response is achieved at 5 wt % clay content. Low clay contents with a homogeneous special distribution results in much more interface

that the crystallization process for EVA-28 is significantly affected by the presence of clays. From the Table 4, it can be noted that the crystallization temperatures and the melting temperatures of EVA/clay nanocomposites increased accordingly with the increase of clay content. The ΔHm value and the percentage of crystallinity was found to be increasing up to 5% clay loading and then decreased due to the agglomeration of clay at higher loading. 20006



Article

Figure 4. Transmission electron micrographs of EVA/Cloisite 25A nanocomposites. (a) Nanocomposite with 3 wt % of Cloisite 25A. (b) Nanocomposite with 5 wt % of Cloisite 25A. (c) Nanocomposite with 7 wt % of Cloisite 25A.

Table 4. DSC Analysis for EVA/Clay Nanocomposites

a

sample

Tma (°C)

Tcb (°C)

ΔHmc (J g−1)

Xc (%)

EVA A3 A5 A7

67.59 79.03 79.43 79.63

55.01 59.57 60.72 61.27

45.87 35.68 40.51 21.45

16.55 12.87 14.61 7.74

Melting temperature. melting.

b

Crystallization temperature.

c

Enthalpy at

Figure 6. Stress−strain graph of EVA and D series of nanocomposites.

but there is no significance increment due to the agglomeration of clay in polymer matrix. The homogeneous dispersion of individual layers shows an effective reinforcing effect for polymer matrix.25 According to Liang et al.26 the nano dispersed clay with high aspect ratio possesses a higher stress bearing capability and efficiency. Stronger interactions between nanoclay layers and polymer molecules associated with larger contact surfaces result in more effective constraint of the motion of polymer chains. Above 5 wt % organoclay loading, the formation of homogeneous nanostructure formation is more difficult. The organoclay is found in the form of microscale filler at this higher concentration due to the poor EVA polymer/clay mixture. The stress−strain graphs of the Cloisite 25A clay nanocomposites are shown in the Figure 6. All the curves show a similar pattern with a strong tendency for strain induced

Figure 5. DSC curve of A series of nanocomposites. (a) Crystallization exotherms of A series of nanocomposites. (b) Melting endotherms of A series of nanocomposites.

area and this provides an excellent interaction between clay and the EVA chains. This is also evident from the TEM images. At higher clay contents, 7 wt %, the interface area may be raised 20007



Article

folded chain segments which are probably held together by van der Waals forces. Under tension, the folded chains may get separated and pulled out. This increases the surface free energy of the system by increasing the exposed surface area of the particles leading to an increase in elongation at break. When the tension is relaxed, the secondary (van der Waals) bonds tend to reform leading to a more compact structure. However, for the increased filler loading (>5 wt %) both these parameters show a decreasing trend due to the filler agglomeration. As shown in TEM and SEM investigation above, at high filler loading, the organoclay is found in the form of microscale filler due to preferred stacking of the individual silicate layers in ordered structures (tactoids). The microstructured organoclay reduces the EVA/organoclay nanocomposite properties. So, the sample is in the form of microcomposite structure (conventional composite). The comparison of the tensile strength and elongation at break of the nanocomposites with the percentage of clay loading is shown in Figure 9. In both cases the composite having 5 wt % showed a higher value.

Figure 7. Elongation at break versus percentage of clay loading of D series.

Figure 8. Comparison of tensile strength of A5 and D5 nanocomposites.

Figure 9. Comparison of elongation at break of A 3 and D3 nanocomposites.

crystallization at higher strain. The addition of clay improves the mechanical properties of the pristine polymer. The elongation at break of the polymer increases by the addition of clay up to 5 wt %, and this is evident from Figure 7. Srivastava et al.27 studied the properties of EVA nanocomposites in detail and found that the tensile strength and elongation at break increase upon nanofiller addition up to a certain extent; thereafter both these values show a dip. This behavior can be explained in terms of two aspects. In the initial region, the filler and polymer matrix show good interaction. The strong interfacial interaction between the filler and the matrix forms some shear zones when the composites are under stress and strain. Because of the strong interaction and development of shear zones, tensile strength of the nanocomposites is increased. The good adhesion between the polymer and the filler will lead to a failure process in which the fracture goes from particle to particle rather than following a direct path, and this result in the higher elongation at break of filled samples than the neat one. This increase in elongation could be explained further by the straightening of the randomly oriented chains under tension and this will result in the sliding of the polymer chains. This sliding leads to the propagation of cracks in the neck region and the molecules can diffuse into it causing repair of the crack. There are regions along the nanoclay particles composed of

On the comparison of the tensile strength and the elongation at break for both nanocomposites series (Cloisite 10A and Cloisite 25A base materials), the composites with Cloisite 25A composites having 5 wt % shows better properties (see Figures 6−9). This fact can be attributed to the better interaction of the clay with the polymer for the D5 sample than the A5 sample.

5. TRANSPORT CHARACTERISTICS 5.1. Diffusion Coefficient. The diffusion coefficient (D) was calculated using the equation28 ⎡ θ ⎤2 ⎥ D = π h⎢ ⎢⎣ 4Q ∞ ⎥⎦

(3)

where θ is the slope of the initial portion of the plots of Qt versus √t and h is the thickness of the sample. The calculated values of diffusion coefficients are given in Table 5. The result shows that the liquid barrier properties of EVA nanocomposites are remarkable. Due to the nanometric level dispersion of the organic and inorganic phases, the available free volume decreases and the platelet morphology of the silicates results in an increase in tortuousity of the path which leads to the reduced diffusivity. D series sample showed the least diffusivity because the dispersion of clay particle is 20008



Article

Table 5. Values of Diffusion Coefficient (D × 1012 m2 s−1) at 28 °C sample

benzene

toluene

xylene


3.10 1.49 1.91 2.76 1.41 1.59 1.65

2.62 1.11 1.04 2.01 0.72 0.78 1.59

2.10 0.69 0.85 1.05 0.55 0.72 0.79

structural characteristics of the polymer and gives idea about the nature of the interaction between polymer and the solvent. The k values of layered silicates filled samples are lower than the pristine polymer. This indicates that the presence of filler can reduce the interaction between polymer and the solvent. 5.3. Influence of Nanoparticles on Sorption and Diffusion. Sorption curves were analyzed to obtain conclusions from the diffusion experiments. Figure 10a and b show

maximum for the D sample. The diffusivity also increases as a function of filler concentration. This can be explained in terms of aggregation of fillers at higher filler loading which leads to an increase in the free volume of the samples. The decrease in diffusivity with increase in the size of the penetrant has been reported by many researchers.29,30 It is also clear from the table that as the size of the penetrant molecule increases the solvent uptake decreases. The diffusivity is maximum for benzene and minimum for xylene. 5.2. Sorption Behavior. The sorption behavior for the system is studied by the equation31 log

Qt Q∞

= log k + n log t (4)

where Qt is the mole percent solvent uptake, and k is a constant that depends on the structural characteristics and polymer/ solvent interaction, whereas the constant, n, determines the mode of sorption mechanism. The values of n and k are determined by power regression analysis of the linear portion of plots Qt versus square root of time .To ensure linearity, values up to 50% of the equilibrium uptake were only used. The values of n and k are placed in Table 6. According to the n values Table 6. Analysis of Sorption Data at 28 °C n values solvents

EVA

A3

benzene toluene xylene

0.50 0.66 0.53

0.68 0.79 0.66

solvent

EVA

A3

A5

A7

D3

D5

D7

benzene toluene xylene

1.51 1.65 1.56

1.50 1.59 1.55

1.45 1.36 1.43

1.44 1.45 1.44

1.31 1.25 1.36

1.42 1.33 1.40

1.45 1.38 1.48

A5 0.61 0.64 0.60 k×

D3

D5

D7

0.61 0.65 0.70 0.66 0.61 0.60 10−2 values

A7

0.65 0.71 0.62

0.63 0.67 0.60

Figure 10. (a) Qt versus time (min1/2) of A series of nanocomposites. (b) Qt versus time (min1/2) of A3 and D3 nanocomposites.

the sorption curves of EVA and nanocomposites. The experiment was conducted at 28 °C with benzene as solvent. The sorption curves showed that the benzene uptake is highest in the case of pure EVA and lowest for D sample and A sample takes an intermediate position. The two main factors that influence the swelling are the availability of free volume in the matrix and the chemical compatibility between the polymer chain and the penetrant molecules. Hence a higher material volume and higher flexibility of the network allow an increased solvent uptake for the unfilled sample. The decreased sorption and diffusion values of EVA nanocomposites (both A and D) is explained as follows. The impermeable clay layers dictate a tortuous pathway for the penetrant molecule to pass through the nanocomposites. The swelling of the nanocomposites is influenced by two factors,

obtained from the above equation, three basic modes of transport are distinguished, If n = 0.5, the diffusion mechanism is Fickian, in that case, the rate of diffusion of permeant molecules is much less than the polymer segment mobility. If n = 1, the mechanism is non-Fickian, this may be considered in systems in which permeant diffusion rates are much faster than polymer relaxation process. If n lies between 0.5 < n < 1, the diffusion mechanism is anomalous, and it occurs when the permeant mobility and polymer segment relaxation rates are similar. From the n values, it can be understood that all systems exhibit anomalous behavior. The anomalous behavior is due to the slow viscous molecular relaxation of the polymer structure in the presence of layered silicate. The value of k implies the 20009



Article

namely the geometry of the filler and the molecular level interaction of the matrix and the filler. Because of the plateletlike morphology of the silicates, the nanofilled samples exhibit reduced equilibrium uptake values owing to the increase in the tortousity of the path. This situation is schematically represented in the Figure 11. Among the two filled samples, D

is assumed that there is an electron layer in the region R < r < R0, with R0 = R + δR where δR represents the thickness of the electron layer or the probability of the overlap of Ps wave function and electron wave function. The expression relating the free volume radius R (nm) and the o-Ps pickoff lifetime τ3 (ns) according to Nakanishi et al.32 is ⎛⎛ R ⎞ ⎛ 1 ⎞ ⎛ 2πR ⎞⎞ ⎛1⎞ ⎜ ⎟ = 21 − ⎜⎜⎜ ⎟ + ⎜ ⎟sin⎜ ⎟⎟⎟ ⎝ τ3 ⎠ ⎝⎝ R 0 ⎠ ⎝ 2π ⎠ ⎝ R 0 ⎠⎠

(5)

Here, the value of δR = 0.1656 nm was determined by fitting experimental τ3 values to data from molecular materials with well-known hole size like zeolites.33 Using this value of R, the free volume size (Vf) is calculated as Vf = (4/3)R3. Then, the relative fractional free volume is evaluated as the product of free volume (Vf) and o-Ps intensity, I3 (%). The transport properties and the overall performance of the membranes were strongly influenced by the free volume present in the nanocomposites. The diffusion of permeate through polymeric membranes can be described by two theories, namely molecular and free volume theory. According to the free volume theory, the diffusion is not a thermally activated process as in the molecular model, but it is assumed to be the result of random redistribution of free volume voids in the polymer matrix. Cohen and Turnbull36 developed the free volume models that describe the diffusion process when a molecule moves into a void larger than a critical size, Vc. Voids are formed during the statically redistribution of free volume within the polymer.34 The influence of fractional free volume percentage is shown in Figure 13.

Figure 11. Schematic representation of tortous pathway.

sample exhibited reduced swelling because of high polymer/ filler interaction. This is mainly due to the difference in the dispersion of clay particles in the matrix. The influence of clay loading on sorption is shown in Figure 12. The equilibrium

Figure 12. Q∞ versus filler loading of D series of composites. Figure 13. Fractional free volume versus filler loading of A and D series of nanocomposites.

uptake value increases as a function of filler loading. This can be explained in terms of aggregation of filler with increase in the concentration of filler. Thus a microphase separation was formed between the polymer and clay platelets. 5.4. Influence of Free Volume. In PALS analysis for polymers and polymer based composites, mainly two parameters are of relevance, namely o-Ps lifetime, τ3, and o-Ps intensity, I3 .The o-Ps lifetime, τ3, measures the size of the free volume holes (Vf), and I3 is a relative measure of the number of free volume sites in the polymer matrix. The free volume cavity radius (R) is related to the o-Ps pickoff lifetime (τ3) by a simple relation. The underlying assumption in the formulation of this relation is that the o-Ps atom in a free volume cell can be approximated to a particle in a potential well of radius R0. The potential is infinite if r > R0 and constant for r ≤ R0. Further, it

It is seen from the figure that the relative fractional free volume percentage of the unfilled sample decreases upon addition of layered silicates. The decrease is due to the interaction between the layered silicates and the polymer due to its platelet structure and high aspect ratio of the fillers. In the presence of layered silicates, the mobility of polymer chain segments is reduced.35 This results in reduced free volume concentration. The contact area between the filler and polymer is higher in the D series of nanocomposites samples, which in turn reduces the free volume concentration of the D sample. It is also found that relative fractional free volume percentage increases with the 20010



Article

trend. It is also found that the slope of the linear portion increases with temperature showing that the transport properties are temperature activated. From Figure 16, it can be noted that the D3 composites having lesser diffusion compared to the A3 nanocomposites even at higher temperature.

clay loadings. This is assessed with the agglomeration of the clay at higher loading. Hence the permeability of nanocomposites is mainly influenced by the free volume effects The relation between diffusion and free volume was also studied according to Cohen−Turnbull theory in which36 ⎡ γV ⎤ D = D0 exp⎢ c ⎥ ⎣ Vf ⎦

(6)

where D is the diffusion coefficient, D0 is a constant, γ is an overlap factor [(γ0) ∼ 1, for most polymers], Vc is the critical size of the molecule, and Vf is the free volume. It is clear from Figure 14 that the relationship between the observed diffusion data and free volume follows the Cohen− Turnbull theory within the experimental error limits.

Figure 16. Qt versus time (min1/2) of A3 and D3 nanocomposites at 50 °C.

The energy of activation for the diffusion and permeation process is calculated from the Arrhenius relationship37 X = X 0e−E / RT

(7)

where X is P (permeation coefficient, obtained as a product of diffusion coefficient, obtained as a product of diffusion coefficient or sorption coefficient) and X0 is P0 or D0 which is a constant. The values of activation energy for diffusion, ED, and the activation energy for permeation, Ep, were estimated. From the difference between Ep and ED, the heat of sorption, ΔH, was estimated. The values of Ep, ED, and ΔHs are compiled in Table 7. It can be seen from the table that the activation energy for diffusion and permeation for unfilled EVA is lower than that of polymer nanocomposite films.

Figure 14. Log D versus 1/Vf.

5.5. Temperature Effects and Activation Parameters. The temperature dependence of the diffusion process through EVA/clay nanocomposites was studied at three different values: 28, 50, and 70 °C. In Figure 15, Qt mole percent uptake is plotted as a function of time at various temperatures for A3 sample. The solvent used in this case was toluene. It has been observed in all cases that maximum solvent uptake happens at higher temperatures. All the other systems showed the same

Table 7. Activation Parameters of Diffusion (±0.07) sample

ED (kJ mol−1)

Ep (kJ mol−1)

ΔHs (kJ mol−1)


7.34 26.7 19.2 16.1 38.2 31.7 31.2

6.21 9.8 2.4 3.8 17.4 18.1 19.2

1.13 16.9 16.8 12.3 20.8 13.6 12.0

The large aspect ratio of the clay platelets effectively increased the diffusion path, which was responsible for the increased activation energy. The value of ΔHs gives additional information about the molecular transport through the polymer matrix. ΔHs is a composite parameter involving contribution from Henry’s law and Langmuir type sorption. (i) Henry’s law is needed for the formation of a site and the dissolution of the species into that sitethe formation of the site involves an endothermic contributionand (ii) Langmuir’s (hole filling) type sorption mechanism, in which case the site already exists

Figure 15. Influence of temperature on diffusion of A3 nanocomposites. 20011



Article

in the polymer matrix and sorption by hole filling gives exothermic heat of sorption. All values are positive suggesting that sorption is mainly dominated by Henry’s law, i.e. the formation of sites and the filling of these sites by penetrant molecules. 5.6. Comparison with Theory. The theoretical sorption curves were generated using38 Qt Q∞

8 =1− 2 π

∞

∑ n=0

χ=

(dϕ/dT ){[ϕ/(1 − ϕ] + N ln(1 − ϕ) + Nϕ} 2ϕ(dϕ/dT ) − ϕ2N (dϕ/dT ) − ϕ2 /T (10)

in which N is calculated from the below equation N=

⎡ D(2n + 1)2 π 2t ⎤ 1 exp⎢ − ⎥ 2 ⎣ ⎦ (2n + 1) h2

(ϕ2/3/3) − (2/3) ϕ1/3 − (2ϕ/3)

Experimentally obtained values of diffusion coefficients of A3 composites are substituted in the equation and the resulting curve is shown in Figure 17.

Mc(aff) = Figure 17. Comparison of theoretical curve versus experimental curve of D3 nanocomposite.

ρp Vϕ2c 2/3ϕ2m1/3(1 − μ/υϕ2m1/3) −[ln(1 − ϕ2m) + ϕ2m + χϕ2m 2]

(12)

where μ and υ are the number of effective chains and junctions.45 ϕ2m is the polymer volume fraction of swelling at equilibrium, and ϕ2c, the polymer volume fraction during cross linking. The calculated Mc values are detailed in Table 8. James and Guth46 proposed the phantom network model where the chain may move freely through one another. According to the theory, the molecular weight between the cross links for the phantom limit of model Mc(ph) was calculated by eq 13

It can be seen that the theoretical and experimental results are not in good agreement. Similar results were reported by Muralidharan et al.39 where the theoretical data did not agree with experimental results. In fact the theoretical curves were drawn based on the Fickian diffusion model. The deviation from Fickian mode is due to the fact that addition of clay strongly hinders with the diffusion process. 5.7. Network Structure Analysis. The diffusion is influenced by the polymer morphology, and hence in addition to experimental variables, we have estimated the molecular mass between the cross links The molecular mass between cross links was estimated using the Flory−Rehner equation40

Mc(ph) =

(1 − 2/χ )ρp Vϕ2c 2/3ϕ2m1/3 −[ln(1 − ϕ2m) + ϕ2m + χϕ2m 2]

(13)

where χ is the junction functionality. The values are given in the Table 9. Joseph et al.48 observed that Mc(chem) values are close to Mc(aff) values which suggests that, in the highly swollen state, the chains in the blends and in the component polymers deform affinely, i.e. the chains in the network are freely moving without fluctuating the junction points. It is found that Mc values of EVA/clay nanocomposites are close to Mc(aff). This shows that, in the solvent swollen state, the network deforms affinely. 47

−ρp V Ø1/3 ln(1 − Ø) + Ø + χϕ2

(11)

The calculated Mc values are in Table 8. The highest value of Mc was shown by pure EVA which supported the observation of high solvent uptake by it. As the filler content increases, the agglomeration takes place leading to poor polymer/filler interaction. Among all the filled samples, the Mc is the lowest for A3 series of samples, which shows the maximum degree of filler−polymer interaction. Lower values of Mc indicate that the network is more restrained and this results in lower swelling of these samples. The study of deformation of polymeric networks during the transport of molecules through the polymer, gives us a better understanding of the transport phenomena. Flory and Rhener relations were well developed for a network deforming affinity, that is, the components of each chain vector transform linearly with macroscopic deformation and the junction points are assumed to be embedded in the network without fluctuations. The molecular weight between the cross links for the affine limit of the model [Mc(aff)] was calculated using the equation43,44

(8)

Mc =

42

(9)

where ρp is the density of the polymer, V is is the molar volume of the solvent, Ø is the volume fraction of the polymer in the fully swollen state, and χ is the polymer solvent interaction parameter which is given by the equation41 Table 8. Mc Values (g cm−3) solvent

EVA

A3

A5

A7

D3

D5

D7

benzene xylene toluene

16842 15964 14381

9485 7337 6994

9788 8694 7249

10261 9868 8944

9187 6839 6590

9346 7142 6908

9626 8269 7946

20012



■

Table 9. Mc(aff) and Mc(Ph) (g cm−3) Values A3


9293 7186 6854

solvent

A3

A5


4647 3593 3427

4729 4249 3477

■

A5

A7

D3

D5

D7

8999 6671 6426

9004 6918 6587

9247 8040 7490

A7

D3

D5

D7

4856 4834 4254

4499 3336 3213

4502 3459 3294

4624 4020 3745

9458 9712 8497 9667 6954 8507 Mc(Ph) values

AUTHOR INFORMATION

Corresponding Author

Mc(aff) values solvent

Article

*E-mail: [email protected] (A.K.S.), sabupolymer@ yahoo.com (S.T.). Tel. no.: (1) 09446045664 (2) 0944223452, +91-481-2730003, 2731036. Fax: +91-481-2731002. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are thankful to UGC, New Delhi, and KSCSTE, Trivandrum, Nanomission of DST India. SAXS data were collected at the 711 beamline of the MAX-lab synchrotron facility using beam time granted to T.S.P. The financial support from the Ministry of Higher Education, Science and Technology of the Republic of Slovenia through the contract No. 3211-10-000057 (Center of Excellence Polymer Materials and Technologies) is acknowledged.

CONCLUSIONS Poly(ethylene-co-vinyl)/clay nanocomposites with two different organomodified clays have been prepared and their morphology, mechanical properties, and their liquid transport characteristics were investigated. Samples with varying amount of organic modifier (Cloisite 10A and Cloisite 25A) were also prepared, and results were compared with unfilled one. The morphology of the nanocomposites was evaluated using SAXS, transmission, and scanning electron microscopy. SAXS results showed intercalation of the polymeric chains between the silicate layers in all cases. The interlayer spacing, and the degree of intercalation for the nanocomposite showed dependence with the kind of organoclay employed, being larger for Cloisite 10A based nanocomposites. No significant differences were observed in the average periodicities of the structures (d-values) between different clay loadings TEM images showed that sample with 3 wt % of clay showed a good dispersion of clay particles resulting in a partially exfoliated structure (mixture of intercalated and exfoliated structures). The dispersion of nanoparticles seems to be better for Cloisite 25A samples series compared with Cloisite 10A samples. For both kinds of nanocomposites, the agglomeration of the particles increase at higher clay loading. The mechanical properties of the nanocomposites were analyzed. Samples with 5 wt % Cloisite 25A clay showed superior performance in tensile strength and elongation at break. The influence of nanoclay loading, penetrant size, and temperature on the diffusion process was analyzed. The maximum solvent uptake decreases with the addition of nanoclays. Due to the enhanced polymer/filler interaction, Cloisite 25A sample exhibited lower solvent uptake. However, the solvent uptake tendency increased at higher clay loading. This is ascribed to the poor physical interaction between the matrix and filler, leading to aggregation of fillers. A similar behavior was found for Cloisite 10A sample series. The mode of diffusion deviates from the Fickian behavior. The transport behavior of the systems was correlated with the morphology of the system. The diffusion coefficient values also showed the above trend. The transport coefficients increase with an increase in temperature. The activation energy for diffusion was determined using an Arrhenius relationship, the values of which support the above results. Different diffusion models were applied to analyze the transport data. The molar mass between cross links was calculated using Flory−Rehner theory. The phantom and affine models were used to analyze the deformation of network during swelling. It was found that affine models agree well with the experiment. The experimental sorption curve was found to deviate from the theoretical curve, which is fully a Fickian one.

■

REFERENCES

(1) Alexandre, M.; Dubois, P. Mater. Sci. Eng. 2000, 28, 1−63. (2) Schmidt, D.; Shah, D.; Giannelis, E. P. Curr. Opin. Solid State Mater. Sci. 2002, 6, 205−212. (3) Pinnavaia, T. J.; Beall, G. W. Polymer−Clay Nanocomposites; John Wiley and Sons: New York., 2000. (4) Gilman, J. W. Appl. Clay Sci. 1999, 15, 31−49. (5) Ahn, J.; Chung, W.-J.; Pinnau, I.; Song, J.; Du, N.; Robertson, G. P.; Guiver, M. D. J. Membr. Sci. 2010, 346, 280−287. (6) Chavarria, F.; Paul, D. R. Polymer 2006, 47, 7760−7773. (7) Park, J. U.; Choi, Y. S.; Cho, K. S.; Kim, D. H.; Ahn, K. H.; Lee, S. J. Polymer 2006, 47, 5145−5153. (8) Alentiev, A.; Semenova, S. I.; Sanopoulou, M. J. Appl. Polym. Sci. 2005, 95, 226−230. (9) George, S. C.; Thomas, S. Prog. Polym. Sci. 2001, 26, 985−1017. (10) Drozdov, A. D.; Christiansen, J. C.; Gupta, R. K.; Shah, A. P. J. Polym. Sci. B Polym. Phys. 2003, 41, 476−92. (11) Messersmith, P. B.; Giannelis, E, P. J. Polym. Sci. A Polym. Chem. 1995, 33, 1047−57. (12) Zhong, Y.; Janes, D.; Zheng, Y.; Hetzer, M.; Kee, D. D. Polym. Eng. Sci. 2007, 47, 1101−1107. (13) Bharadwaj, R. K. Macromolecules 2001, 34, 9189−9192. (14) Musto, P.; Mascia, L.; Mensiteri, G.; Ragosta, G. Polymer. 2005, 46, 4492−4503. (15) Vlasveld, D. P. N.; Groenewold, J.; Bersee, H. E. N.; Picken, S. J. Polymer 2005, 46, 12567−12576. (16) Knaapila, M.; Svensson., C.; Barauskas, J.; Zackrisson, M.; Nielsen, S. S.; Toft, K. N.; Vestergaard, B.; Arleth, L.; Olsson, U.; Pedersen, J. S.; Cerenius, Y. J. Synchrotron. Radiat. 2009, 16, 498−504. (17) Hammersley, A. P.; Svensson, S. O.; Thompson, A.; Graafsma, H.; Kvick, E.; Moy, J. P. Rev. Sci. Instrument. 1995, 14, 2729−2734. (18) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer hand book; Wiley-Inter Science: New York, 1976. (19) Mathew, A. P.; Packirisamy, S.; Stephen, R.; Thomas, S. J. Membr. Sci. 2002, 201, 213−227. (20) Mohomane, S. M.; Djokovic, V.; Thomas, S.; Luyt, A. S. Polym.Test. 2011, 30, 585−593. (21) Byoung, C. C.; Tae, K, C.; Mi, H. C.; Yong-Chan, C.; Jihua, C.; David, M.; Robert, C. C. J. Appl. Polym. Sci. 2007, 106, 712−721. (22) Peeterbroeck, S.; Alexandre, M.; Jérôme, R. Ph. Dubois. Polym. Degrad. Stab. 2005, 90, 288−294. (23) Prasad, R.; Gupta, R. K.; Cser, F.; Bhattacharya, S. N. J. Polym. Eng. 2005, 25, 305−330. (24) Chaudhary, D. S.; Prasad, R.; Gupta, R. K.; Bhattacharya, S. N. Thermochim. Acta 2005, 433, 187−95. (25) Grant, B. D.; Bret, C. H.; Robert, M. B. Polymer 2005, 46, 6706−6714. 20013



Article

(26) Liang, Y.; Wang, Y.; Wu, Y.; Lu., Y; Zhang, H.; Zhang, L . Polym. Test. 2005, 24, 12−17. (27) Srivastava, S. K.; Pramanik, M.; Acharya, H. J. Polym. Sci. B Polym. Phys. 2006, 44, 471−480. (28) Joseph, A.; Mathai, A. E.; Thomas., S. J. Membr. Sci. 2003, 220, 13−30. (29) Lucht, L. M.; Peppas, N. A. J. Appl. Polym. Sci. 1987, 33, 1557− 1566. (30) Aminabhavi, T. M.; Khinnavar, R. S. Polymer 1993, 34, 1006− 1018. (31) Southern, E.; Thomas, A. G. Trans. Farad. Soc. 1967, 63, 1913− 1921. (32) Nakanishi, H.; Jean, Y. C.; Schrader, D. M.; Jean, Y. C. Positron and Positronium Chemistry; Elsevier: Amsterdam., 1988. (33) Jean, Y. C. J. Microchem. 1990, 42, 72−102. (34) Stephen, R.; Ranganathaiah, C.; Varghese, S.; Joseph, K.; Thomas, S. Polymer 2006, 47, 858−870. (35) Pramanik, M.; Acharya, H.; Srivastava, S. K. Macromol. Mater. Eng. 2004, 289, 562−567. (36) Turnbull, D.; Cohen, M. H. J. Chem. Phys. 1961, 34, 120−125. (37) Mathai, A. E.; Thomas, S. J. Macromol. Sci. B Phys. 1996, 35, 229−253. (38) Gent, A. N.; Tobias, R. H. J. Polym. Sci. B Polym. Phys. 1982, 20, 217−232. (39) Muralidharan, M. N.; Anil, K. S.; Thomas, S. J. Membr. Sci. 2008, 315, 147−154. (40) Flory, P. J.; Rehner., J., Jr. J. Chem. Phys. 1943, 11, 512−521. (41) Unnikrishnan, G.; Thomas, S. Polymer 1994, 35, 5504−5510. (42) Kumar, S. A.; Kumaran, M. G.; Thomas, S. J. Polym. Sci., Part B: Polym. Phys. 2007, 45, 2470−2480. (43) Flory, P. J. Principles of Polymer Chemistry; Cornell University Press: Ithaca., 1953. (44) Treloar, L. R. G. The Physics of Rubber Elasticity; Clarendon Press: Oxford., 1975. (45) Mark, J. E.; Erman, B. Rubber Like Elasticity-A Molecular Approach; Wiley Inter Science: New York., 1988. (46) James, M.; Guth, E. J. Chem. Phys. 1947, 15, 669−683. (47) Cowie, J. M. G. Polymers: Chemistry and Physics of Modern Materials; Chapman and Hall: New York, 1991. (48) Joseph, A.; Mathai, A. E.; Thomas, S. J. Membr. Sci. 2003, 220, 13−30.

20014


Clay Intercalation and its Influence on the Morphology and Transport

Recommend Documents