Printer Friendly

Morphology, barrier, and mechanical properties of biaxially deformed poly(ethylene terephthalate)-mica Nanocomposites.

INTRODUCTION

Poly(ethylene terephthalate) (PET) is an engineering thermoplastic which finds many applications in industries such as food and beverage packaging, and textiles. It is generally desirable to enhance the mechanical and barrier properties of PET so that lighter products can be manufactured and the shelf life lengthened when food and beverage packaging are concerned. One promising method to achieve these enhancements is by adding nanoclays to the PET matrix. Nanoclays are readily expandable and dispersed in a polymer matrix, provided it is compatible with the matrix and varying exfoliation levels can be achieved depending on the processing conditions. Since, unlike conventional composite materials, the filler has dimensions on a nanometer scale reinforcement can be achieved without significant loss of transparency which is a major advantage in packaging applications. The morphology of polymer-clay nanocomposites is usually characterized as either intercalated or exfoliated. An intercalated structure has polymer chains diffused into the clay layer galleries, expanding them whereas an exfoliated structure has its clay layers completely delaminated and the layered structure is destroyed. Complete exfoliation is generally desirable due to the high aspect ratio of dispersed nanoparticles which is believed to give superior mechanical and barrier properties. However, complete exfoliation is notoriously difficult to achieve and normally nanocomposites reported in literature are only partially exfoliated (1).

Recently, numerous studies on PET-clay nanocompo-sites have been reported in the literature. Kim and Kim processed PET nanocomposites by in situ polymerization using self-prepared montmorillonile (MMT) based organo-clay. It was pointed out that the nanocomposites synthesized had better exfoliation compared to nanocomposites based on unmodified MMT. The tensile strength was improved by 78% for a 0.5 wt% clay loading and the barrier to oxygen improved by 52% at 1 wt% clay loading (2). Hwang et al. attempted to process PET nanocomposites by in situ polymerization and arrived at a similar conclusion in that organically modified clay achieved better exfoliation than the unmodified counterpart. The improved exfoliation with organically was ascribed to the increased clay basal spacing and improved compatibility with the PET matrix. The modulus of PET was increased by 33% at 5 wt% clay loading, however, the tensile strength and elongation at break reached a maximum at 0.5 wt% clay loading and decreased with a further increase in clay loading, which the authors attributed to an increased agglomeration level (3). Their studies showed the significance of the organic surfactant in the nanoclay dispersion. Also, they pointed out that agglomerates of nanoclay particles and delaminated clay sheets can coexist, hence it is necessary to analyze the nanocomposite morphology al both micron and submicron scales. Also, the incorporation of nanoclay could cause degradation of the matrix through side reaction between the surfactant and the polymer chains (4). This degradation can be observed as a reduction in the matrix molecular weight or intrinsic viscosity. Nevertheless, modulus enhancement may still be obtained even if there is a reduction in the molecular weight (5). Bizarria et al. processed PET nanocomposites by direct melt extrusion. The modulus was increased by 25% independent of clay loading even though there was a significant lowering in the intrinsic viscosity of the PET used in the study (5). The modulus enhancement is expected to be greater if molecular weight reduction is avoided. Gurmendi et al. obtained a 41% increase in modulus for a 6 wt% clay loading of melt extruded PET nanocomposites. The nanocomposiles obtained had a partially exfoliated morphology (6). However, micron scale characterizations of the agglomeration were not reported. Sanchez-Gracia et al. processed PET nanocomposiles by melt extrusion and the resulting nanocomposite was claimed to have good exfoliation and no agglomeration. An oxygen barrier enhancement of 55% for a 5 wt% clay loading was reported (7). Kracalik et al. modified the nanoclay using a silanizalion reaction prior to melt exlrusion. The effect of silanization on clay dispersion was dependent on the surfactant type. An increase in modulus of 29% at 5 wt% silicate loading was obtained while elongation lo break was not significantly decreased. This was ascribed to the improved interfacial adhesion due to silanization (8). Apart from directly blending the commercially available organically modified clay with the PET matrix, researchers had attempted to predisperse the clay in a solvent or water before the melt blending process, in order to facilitate exfoliation (9), (10). Chung et al. processed PET nanocomposites by first dispersing the clay in chloroform and then removing the organic modifier with acid before melt extrusion. The authors claimed the nanoclay was fully delaminated and the resultant nanocomposites had a 38% improvement in modulus compared with the unfilled PET. An additional benefit of this novel processing technique is that the degradation of the clay organic surfactant is avoided (9). Ammala et al. processed PET nanocomposites by melt extrusion but before extrusion the clay was dispersed in water and then mixed with a PET ionomer. An exfoliated morphology was obtained due to improved compatibility between clay and PET ionomer (10). However, the mechanical and barrier properties have not yet been reported. An issue for processing PET nanocomposites is the degradation of the organic surfactant of clay due to the high processing temperature of PET (270-300[degrees] C). Stoeffler et al. studied the effect of surfactant degradation on exfoliation level of melt mixed PET nanocomposites. It was found that surfactant degradation could affect the clay dispersion and better exfoliation was achieved by using clay of relatively high thermal stability and large initial basal spacing (11). Since thermally stable clay is not readily commercially available, the processing method reported by Chung et al. (9) is by far more practical for industrial application. Overall, both the in situ polymerization and melt mixing methods have been shown to achieve good exfoliation of PET nanocomposites and the mechanical and barrier property enhancements were comparable, with modulus increases in the range from 25% to 40% and oxygen barrier enhancement of about 50%.

Most of the studies of PET-clay nanocomposites reported in the literature focus on the initial formulation and characterization of the nanocomposites. Little has been reported on secondary processing to form finished products of polymer nanocomposites. Secondary processing such as biaxial preform stretching which occurs in stretch blow moulding or thermoforming plays an important role in determining the final properties of the end product. Characterization of biaxially stretched PET-silica nanocomposites just above the glass transition temperature, [T.sub.g] has been reported and it has been shown that biaxial stretching significantly affects the silica dispersion (12). However, the role of the nanosilica on the process-ability and the effect of stretch-induced morphology on the final properties of the end product has not yet been reported. Another study involving secondary processing where spun fibers were drawn (uniaxial) at just above their [T.sub.g], showed that enhancement in the mechanical properties can be achieved by a secondary processing step. The enhancement was affected by filler agglomeration which alters polymer molecular orientation by making it more difficult for polymer chains to arrange in the fiber axis (13). As revealed in a study of blown film nanocomposites, the filler orientation also plays a role in the reinforcement effect, apart from the exfoliation level (14), as the draw down and inflation of the extrudate film can induce alignment of the fillers. A similar observation was reported in a study into the effect of filler orientation on the mechanical properties of extruded polymer nanocomposites of controlled filler orientation and aspect ratio (14), (15). Shah et al. and Bharadwaj have also pointed out that the same argument can be applied to the barrier properties of polymer nanocomposites (14), (16).

Recently we reported that large-scale biaxial stretching showed an effect on nanoclay (mica) exfoliation in PET nanocomposites and the mica loading affected the processability in a process which imitated stretch blow moulding (17), (18). Subsequently, this article reports the mechanical and barrier properties of PET-mica nanocomposite sheets as a function of mica loading before and after biaxial stretching. The focus of this study is to examine how changes in nanocomposite morphology, induced by biaxial stretching, further affect the mechanical and barrier properties.

MATERIALS AND EXPERIMENTAL

In order to study the effects of biaxial stretching on the structure and properties of PET nanocomposites, the morphology of the nanocomposites was examined in detail using transmission electron microscopy (TEM) and optical microscopy (OM) including quantitative measurement of the nanoparticles dimensions. The quantitative data also served as an input to composite models for analyzing the effects of structural parameters such as the filler aspect ratio, on mechanical and barrier properties. Crystal Unity was measured in order to determine the actual cause of any enhancements in properties since it plays an important part in the mechanical and barrier properties. Tensile and barrier tests were conducted to examine the nanocomposite performances. Dynamic mechanical thermal analysis (DMTA) was performed to study the storage modulus in a range of temperatures including above the glass transition, [T.sub.g]. The effect of incomplete exfoliation on the mechanical and barrier properties was examined using composite models.

Materials

The PET and the mica (Somasif MAE) have been described in previous publications and were also used in this study (17), (18). PET grade T74F9IV080 was supplied by Tergal Fibre. It has a density of 1.4 g c[m.sup.-3] and an intrinsic viscosity of 0.8 dl [g.sup.-1] in 50/50 mixture of phe-nol/ortho di-chlorobenzene al 25[degrees] C and with a polymer concentration of 5 [g.sup.-1]. Synthetic micas (fluorohectorites), Somasif MAE was purchased from Uni-Coop Japan (now CBC Co. Ltd.). The Somasif MAE contains dimethyl di-(hydrogenated tallow) ammonium chloride.

Compounding of PET Nanocomposites

The mica and polymer were dry-blended and dried overnight at 120[degrees] C before extrusion. PET-mica mixtures of 1%, 2%, and 5% loading by weight were extruded using a Colin ZK25 twin screw extruder of 30:1 length-to-diameter (L/D) ratio. Unfilled PET was also extruded for comparison purposes. The temperature profile of the extruder was 230, 275, 270, 270, 265, 260[degrees] C from the feeder to the extruder end and the die temperature was 260 and 255[degrees] C al the inlet and exit, respectively. The die was 100 mm width and the slit was approximately 1.2 mm. The extruded sheet was cooled on a pair of chili rolls at 75[degrees] C rotating at 1 m min ' and the extruder screw speed was set to 150 mm. The extruded sheets had a thickness of approximately 1 mm and width of approximately 9 mm. The output rate of the extrusion was approximately 5.4 kg/h, equivalent to 60 c[m.sup.3] [min.sup.-1]. Under this condition, the extrudate was subjected to draw down. The extruded sheet samples are denoted as PET + 1% MAE, PET + 2% MAE, and PET + 5% MAE according to the mica loading.

Biaxial Stretching of PET Nanocomposites

Simultaneous equal biaxial tests were conducted using the flexible biaxial stretcher at the Queen's University Belfast (19). Square specimens (76 x 76 x 1 mm) were cut from PET-mica nanocomposite sheet and clamped into the biaxial test machine. The specimen was then heated for 3 min by a pair of thermocouple-controlled hot air blowers positioned above and below the specimen so that the specimen achieved a uniform test temperature of 100[degrees] C. Heating was then stopped and the specimen was deformed biaxially by grippers clamping along the specimen edges, at a set speed. The samples were stretched lo stretch ratios of 2, 2.5, and 3 at a nominal strain rate of 8 [s.sup.-1]. The samples were allowed to cool rapidly by natural convection at the end of each test.

Gel Permeation Chromatography (GPC)

The molecular weights of the extruded nanoeomposites and the virgin PET as well as the raw PET granules were measured by gel permeation chromatography (GPC) to examine if nanoclays loading caused degradation to the PET. The measurements were conducted on a Polymer Laboratories GPC 120 with PLgel guard plus two mixed bed-B column at 115[degrees] C; 1,3-cresol with antioxidant was used as the solvent and the flow rate was 0.8 ml [min.sup.-1]. The resulting molecular weight data are shown in Table 1. It is shown that the molecular weight of PET was reduced with increasing mica loading and this was discussed in previous publication (18). The degradation observed is considered modest since the nanoeomposites still showed superior properties than the virgin PET, which are shown later.

TABLE 1. Molecular weight of test materials.

Sample              Clay loading (wt%)  [M.sub.w] (g [mol.sup.-1] )

PET granules                         0                      105,500
Extruded PET sheet                   0                       90,100
PET + 1% MAE                         1                       86.100
PET + 2% MAE                         2                       80,300
PET + 5% MAE                         5                       78,500


Morphology Characterization

Wide angle X-ray diffraction (WAXD) studies were completed using a Philips PW3040 machine using Cu-K[alpha] radiation with a wavelength of 1.54 A at a step size of 0.01[degrees] and a scan rate of 0.3[degrees] [min.sup.-1]. The 20 range was from 2[degrees] to 10[degrees].

The morphology of the nanoeomposites were studied thoroughly using transmission electron microscopy (TEM) and quantitative characterization of the TEM images were performed subsequently to acquire the distribution of the mica tactoids and properties, such as thickness, length, and aspect ratio. Microtoming was performed along the direction of extrusion at the middle of the sheets using a Reichert Jung Ultracut-E microtome using a diamond knife to obtain specimens 70 nm thick. The microtomed specimens were collected on 200 mesh copper grids. TEM analyses was completed using a Philips CM 100 TEM operating at 100 kV accelerating voltage. The thickness, length, and aspect ratio of the mica particles were measured using JMicroVison software (version 1.25) on a set of TEM images taken at 28.5k magnification. The number of layers per tactoid (n) was calculated as follows (20): n = [t.sub.particle] + d - [t.sub.platelet]/d (1)

Here [t.sub.particle] is the thickness of the particle measured on the TEM images, d is the basal spacing measured by WAXD, and [t.sub.platelet] is the thickness of a single mica platelet which is taken as 0.94 nm. Around 600 tacloids were measured for each sample.

Tactoid orientation was analyzed both qualitatively by comparing the TEM images and characterized quantitatively by measuring the distribution tactoid angle relative to a reference (horizontal) axis of a set of TEM images (X8.9k). Around 200 tactoids were measured for each sample.

In our previous publication, we had shown that large scale biaxial stretching on nanocomposites sheet is able to further improve the exfoliation level where more thinner tactoids can be observed on the stretched sheet [171. Previously, only the samples loaded with 2 wt% mica were presented, where the effects of stretching at a constant mica loading were investigated (17). Here, the effects of different mica loadings on the mechanical and barrier properties were examined and comparisons are made to the unstretched extruded sheets.

Micron-sized mica agglomerates were observed using a Nikon Eclipse ME600 optical microscope incorporating a hot-stage. Twenty micrometer-thick sections were microtomed along the machine direction of each sample using a Leica RM2165 microtome machine with a steel blade. Images were taken using a 20X magnification lens when the sections, sandwiched between a pair of glass slides were melted in the hot-stage. Image processing was performed to obtain high contrast black and white images using Adobe Photoshop Elements 2.0 before quantitative analysis using the method described by Chavarria and Paul (21). The processed high contrast black and white images were then exported to the image analysis software, Scion Image to quantify the agglomeration level. Taking the volume fraction as area fraction, the percentage of agglomeration may be defined as:

% Agglomeration = [A.sub.Ob]/[A.sub.Th] x 100%. (2)

The theoretical area (volume) fraction of particles is:

[MATHEMATICAL EXPRESSION NOT REPRODUCIBLE IN ASCII] (3)

Here, [A.sub.Ob] and [A.sub.Th] is the observed and theoretical area fraction of mica, respectively. Mc and Mm is the mass fraction of mica and matrix, respectively. [[rho].sub.c] and [[rhp].sub.m] is the mica and matrix density, respectively.

Areas of at least 0.5 [mm.sup.2] from a total of three to four optical microscopy (OM) images were taken into consideration.

Differential Scanning Calorimetry (DSC)

Samples of around 10 mg were scanned using a Perkin Elmer DSC 6 from 35 [degrees] C to 280 [degrees] C to obtain the melting and cooling thermograms. A temperature ramp rate of 10 [degrees] C [min.sup.-1] was maintained throughout. The initial crystallinity of the samples was calculated as:

[X.sub.m] = ([[delta]H.sub.m] - [[delta]H.sub.cc])/(1-x)/[[delta]H.sub.m.sup.[degrees] (4)

where [[delta]H.sub.m] is enthalpy of melting, [[delta]H.sub.cc] the enthalpy of cold crystallization, x is the organoclay weight fraction, and [[delta]H.sub.m.sup.[degrees]], is the enthalpy of melting for a 100% crystalline PET, which is taken as 140 J/g (5). [[delta]H.sub.m] is determined by plotting a tangent from the upswing during melting to the end of melting in the temperature scan trace. The crystal-Unity values obtained should be treated as comparative indexes rather than absolute values.

Crystallinity Analysis Using Wide Angle X-Ray Diffraction Analysis (WAXD)

Crystallinity ratio of semicrystalline PET nanocompo-siles was determined by fitting the diffraction data with multiple Gaussian functions corresponding to the amorphous and crystalline peaks. The curve fitting was performed using OriginProS software. The analysis assumes the crystalline peak and the clay peaks can be separated. The clay peaks are excluded from the crystallinity analysis. The crystallinity is calculated as:

[X.sub.c] = [A.sub.crystal]/[A.sub.crystal + [A.sub.amorp] x100% (5)

where [A.sub.crystal] and [A.sub.amorp] are the fitted area of crystalline and amorphous part respectively (22).

Mechanical Properties Characterization

The dynamic properties (storage modulus, E' and tan delta) of the stretched samples were measured using a Tritec 2000 Dynamic Mechanical Analyser (DMTA). The experiments were conducted in a temperature range from 30 [degrees] C to 200 [degrees] C at a frequency of 1 Hz. Tests were conducted in tension mode at 0.01 mm controlled displacement which is equivalent to 0.2% strain. The free length of the sample was 5 mm. The glass transition temperature, [T.sub.g] is defined as the temperature at the maximum in the tan delta peak. Tensile tests were conducted at room temperature on an Instron 5564 tensile tester. For modulus measurement, nominal strain was determined using an extensometer attached on the narrow portion of the dumb-bell test specimens (EN ISO 524-2:1996 type IB A) at a test speed of 1 mm [min.sup.-1] and the gauge length was set as 25 mm. Modulus was determined from the slope of the regression of the stress-strain data between 0.05% and 0.25% strain. The yield and fracture behavior were tested at 50 mm [min.sup.-1] on separate samples. Samples were cut using a hand operated cutting press.

Barrier Properties Characterization

The barrier performance of the nanocomposites was compared for their oxygen permeability coefficient. The oxygen barrier property determines the shelf life of products. Oxygen barrier tests were conducted using a Mocon Oxtran model 2/21 at room temperature and 0% relative humidity. The test area for the samples was 50 c[m.sup.2]. Obtaining the permeability coefficient for the thick (1 mm) extruded samples presented a challenge in that the permeability coefficient of the polymer should ideally be derived from the oxygen transmission rate at equilibrium state. However, the thick extruded samples took too long (more than 200 h) to reach the equilibrium state. Therefore, comparisons amongst the extruded samples were made on the oxygen transmission rate at a known time before the equilibrium state was reached. Note, the barrier tester used in this study has two test cells, enabling testing of two samples simultaneously, where each test cell records oxygen transmission rate of the sample sitting within. For this case, while an unfilled sheet specimen was tested in one cell, a nanocomposite sheet specimen was tested in the other cell. The permeability coefficient at equilibrium state was used to compare the thinner stretched samples.

Analysis Using Conventional Composite Models

The Halpin-Tsai model was used to study the effect of fillers aspect ratio, matrix modulus, and filler modulus on the reinforcement effect of nanocomposites. Different levels of exfoliation were incorporated to the model, which will be shown later, to analyze its effect. According to the Halpin-Tsai model, the relative modulus of a nanocomposite is expressed as follows:

E/[E.sub.m] = 1 + 2[alpha][eta][[phi].sub.f]/1[eta][[phi].sub.f] (6)

Here [alpha] is the filler aspect ratio and [[phi].sub.f] is the filler volume fraction. [E.sub.m] is the matrix modulus.

[eta] is given by:

[MATHEMATICAL EXPRESSION NOT REPRODUCIBLE IN ASCII] (7)

Here [E.sub.f] is the filler modulus.

The filler modulus was taken as a 178 GPa which is also the value derived from the muscovite modulus as determined by the ultrasonic pulse method and used by Fornes and Paul in their modelling of mechanical properties of nanocomposites (23). The modulus of virgin PET at stretch ratios 2, 2.5, and 3 were used as the matrix modulus in Eqs. and 7 to examine the reinforcement effect after different stretching history. In order to obtain the lateral dimension of a mica platelet. TEM was employed. Nanopartieles were first dispersed in acetone and then diluted 100 limes. The solution was subjected to ultrasonic vibration before a small quantity was collected on a carbon coated copper grid. TEM micrographs at 15k magnification were taken using a Tecnal F20 TEM (FEI Company) at 200 kV accelerating voltage. Energy dispersive X-ray (EDX) composition analysis was employed to confirm that the particles observed were mica platelets. The chemical formulation of Somasif mica is [Na.sub.2x][Mg.sub.3][Si.sub.4][O.sub.10][F.sub.y]O[[H.sub.1-y].sup.2], where 0.15 < x < 0.50 and 0.8 < y < 1.0. The width of Somasif platelet was determined to be larger than 1000 nm based on TEM observation as shown in Fig. 1. Here, a platelet width of 1000 urn is assumed. The aspect ratio of exfoliated platelet was taken as 1064 which was derived by dividing the Somasif platelet width ([W.sub.platelet] = 1000 nm) by the thickness of single platelet (0.94 nm). Incomplete exfoliation was considered for the Halpin-Tsai prediction by simply replacing the modulus and aspect ratio of the exfoliated clay platelet in Ec/s. 6 and with the modulus and aspect ratio of a stack of platelets.

The modulus of a stack of platelets ([E.sub.stack]) was calculated using the following equations:

[E.sub.stack] = [[PHI].sub.silieate][E.sub.silicate] + [[PHI].sub.gallery][E.sub.gallery], (8)

[[PHI.sub.silieate] = n[t.sub.platelet]/d(n-1)+[t.sub.platelet] (9)

Here [[PHI.sub.silieate] is the volume fraction of the inorganic silicate whereas [[PHI].sub.gallery] is volume fraction of the clay gallery. [E.sub.silicate] is the silicate modulus whereas [E.sub.gallery] is the modulus of substance within silicale gallery, n is the number of platelet layers in a tactoid stack, d is the basal spacing (3 nm as determined by XRD) of the silicate and[t.sub.platelet] is the thickness of a single platelet. [E.sub.silicate] is taken as 178 MPa. The contribution from [E.sub.gallery] assumed to be negligible. For the comparison between the experimental data and the prediction, the average tactoid thickness as obtained from quantitative measurement of the TEM images were used to derive the term [alpha] in Eqs. 6 and 7. The aspect ratios for different levels of exfoliation are expressed as:

[alpha] = [W.sub.platelet]/[t.sub.stack], (10)

[t.sub.stack] = d(n-1) + [t.sub.platelet]. (11)

Here [t.sub.stack] is the thickness of a stack of platelets. The Hal-pin-Tsai model makes a few assumptions such as perfect alignment and bonding to the matrix of the fillers, the fillers are of uniform shape, the fillers do not have any effect on the matrix properties, and the fillers do not agglomerate. More detailed explanation of the model was reported by Forties and Paul (23).

Nielsen's model was used to study the effect of the fillers aspect ratio and the fillers concentration on the barrier behavior of the nanocomposites, and is expressed as follows:

P/[P.sub.m] = 1 - [[phi].sub.f]/1 + 1/2[alpha] [[phi].sub.f] (12)

Here [[phi].sub.f] is the volume fraction of the inorganic silicate and [alpha] is the aspect ratio of the filler. The relative permeability as predicted by Nielsen's model is a function of fillers volume fraction and aspect ratio. Like the Halpin-Tsai model, Nielsen's model makes assumptions about the fillers such as perfect alignment, uniform shape, and no agglomeration and they do not have any effect on the matrix properties. For all the computations using Halpin-Tsai and Neilsen's model, the inorganic surfactant concentration was taken as 40% and it was excluded from the weight fraction of the clay [18]. By excluding the surfactant weight and converting the weight fraction of the inorganic silicate to volume fraction (by taking the density of the silicate and PET as 2.6 g [ml.sup.1] and 1.4 g [ml.sup.1 respectively), the organoclay loading of 1 wt%, 2 wt%, and 5 wt% are converted to 0.32 vol%, 0.65 vol%, and 1.64 vol% silicate loading, respectively.

RESULTS AND DISCUSSION

Morphology and Dispersion of Nanoclay in PET Matrix

A brief review of the stretching effect on the clay dispersion has already been reported by our group earlier (17). In this article, we will focus on the effect of different mica loading and the subsequent stretching on the morphology and the macroscopic properties.

Table 2 shows the effect of stretching and mica loading on the WAXD basal spacing (d-spacing) of the PET-mica nanocomposites. The d-spacing of the Somasif MAE powder and nanocomposites sheets were determined by means of Bragg's equation ([lambda]sin [[theta].sub.001] ). By analyzing the d-spacing data alone, it may be concluded that the mica stacks are not well intercalated by polymer chains because, apart from the PET + 5%MAE, the d-spacing is not increasing for the filled systems. However, literature reports that WAXD should not be used as a stand alone characterization technique for polymer nanocomposites because it may give misleading interpretations if there is a distribution of interlayer distances, a large amount of clays, an uneven distribution of surfactant or a random orientation of the clay (20), (24). The d-spacing may also be reduced if there has been surfactant degradation or compression of the clay stacks upon processing. It is shown later that analysis of TEM images provides a different conclusion.

TABLE 2. Quantitative image analysis dala and d-spacing of PET-mica
nanocomposites.

                  Unstrelched         Stireiched-stretch ratio 3

           PET +     PET + 2%  PET +     PET + 1 %   PET + 2%   PET + 5%
           1 % MAE    MAE (a)  5% MAE     MAE         MAE (a)       MAE

d-Spacing    3.0          3.0    3.1                 3.1     3.1    3.1
(nm)
Average      4.2          8.3    8.7                 4.0     5.7    5.6
tactoid
thickness
(nm)
Stdev        4,3          7.5   10.2                 5.1     7.2    8.9
Average      2.1          3.5    3.6                 2.0     2.5    2.5
number of
platelets
per
tactoids
Stdev        1.4          2.5    3.4                 1.6     2.3    2.8
Average     55.0         66.3   48.8                70.6    93.5   95.8
tactoid
aspect
ratio
Stdev       50.2         69.5   46.4                69.1    81.2   99.2
Average    199.9        213.8  401.1               223.7   352.1  512.3
tactoid
length
(rim)
Stdev      105.0        128.6  269.4               197.7   213.1  357.3

(a) Data reproduced from Ref. 17 with permission from Elsevier Science.
(b) The average tactoid aspect ratio is calculated as: [summation over
(term)] (length/thickness)/Number of tactoids The neat Somasif MAE
has a d-spacing of 3 nm.


Figure 2 shows the morphology of the nanocomposites before and after biaxial stretching to a stretch ratio of 3 whereas Figs. 3 and 4 show the histograms of the distribution of tactoid thickness and aspect ratio obtained from TEM image analysis. The average tactoid thickness, aspect ratio, length, and number of layers per tactoid are shown in Table 2. The data for PET + 2%MAE at stretch ratio 2 and 2.5 are reproduced here in Table 3 for comparison (17). The tactoid thickness and number of layers per tactoid data were inputted to the conventional composite models which are shown later. Qualitatively, from the TEM images, it can be observed that the tactoids are more aligned after biaxial stretching. TEM image analysis revealed that the tactoid thickness was reduced after stretching. There are more tactoids having thickness in the range of 1 to 3 nm in the stretched samples than the unstretched nanocomposites. Taking the PET nanocompo-sites at 1% Somasif MAE loading for example, the percentage of tactoids having thickness of 1 to 2 nm is increased from 12% to 37% after biaxial stretching. For the tactoid aspect ratio, tactoids having an aspect ratio in the range of 100 to 200 is much lower (less than 5% at 2 wt% and 5 wt% mica loading) before stretching. In comparison, the percentage of tactoids having an aspect ratio in that range increased to more than 20% after stretching. It is worth noting that the aspect ratio data, from which is derived the tactoid length is likely to be affected by microtomed section thickness as indicated by Vermogen et al. (20). However, as shown in Table 2, the magnitude of the change in average tactoid length (which is more than 100 nm al 2 wt% and 5 wt% mica loading) is much larger than the ambiguity (about 14 nm) resulting from the microtomed section thickness as indicated by Vermogen et al. Therefore, the increase in tactoid aspect ratio after stretching is unlikely to be purely an effect of sample preparation. It is speculated that stretching has pulled and orientated the initially thicker and randomly orientated tactoids until the mica sheets of each tactoid can slip across in an action that is analogous to the dragging of a deck of cards as shown in Fig. 5. The apparent tactoid lengths and thicknesses are therefore lengthened and thinned, respectively (17).

TABLE 3. Quantitative image analysis data for PET 4-2% MAE at different
stretching conditions.

                   Unstretched  Stretch   Stretch   Stretch
                                  ratio 2  ratio 2.5  ratio 3

Average tactoid            8.3      6.9        5.4      5.7
thickness (nm)
Stdev                      7.5      7.2        5.4      7.2
Average number of          3.5      2.9        2.4      2.5
plaleleis per
tactoicis
Stdev                      2.5      2.3        1.7      2.3
Average tacioid          213.8    328.3      382.5    352.1
length (nm)
Stdev                    128.6    194.8      269.2    213.1
Average tactoid           66.3     65.8       70.0     93.5
aspect ratio
Stdev                     69.5     62.6       54.0     81.2

Data reproduced from Ref. 17 with permission from Elsevier Science.


Comparing the exfoliation level found with different clay loading, an increase in the mica loading percentage caused an increase in the tactoid thickness before and after stretching. When mica loading was increased from 1 wt% to 5 wt%, the average tactoid thickness increased from 4.2 nm to 8.7 nm for the unstretched sample whereas for the stretched sample, the average tactoid thickness increased from 4 nm to 5.6 nm. The tactoid aspect ratio on the other hand was independent of mica loading before stretching. However, the tactoid aspect ratio increased with mica loading after stretching which contrasted with the opposite trend with the tactoid thickness. It is worth noting that PET at 1 wt% mica loading did not exhibit much change in the average tactoid thickness and aspect ratio by stretching. PET loaded with 2 wt% mica had similar average tactoid thickness and aspect ratio to the 5 wt% loaded PET. It is speculated that at higher mica loading, the mica particles have a greater chance of forming agglomerates with neighboring par-tides and therefore thicker tactoids are observed. As a result of the thicker tactoids at higher mica loading, the layers of mica platelets can slip to a greater extent during stretching before eventually being torn apart. Considering the tactoid orientation, stretching does have a pronounced effect on tactoid orientation as shown in the TEM micrographs in Fig. 2 and the histograms in Fig. 6. However, tactoid orientation appears to be independent of different mica concentration.

Optical microscopy was carried out to observe the agglomeration of mica particles. Fig. 7 shows the micrographs of PET loaded with 2 wt% and 5 wt% of mica. Images were not taken for PET loaded with 1 wt% of mica as observations from TEM had shown these materials to have better exfoliation and to be less affected by stretching. As depicted on the micrographs, higher mica loading resulted in a greater level of agglomeration and this was reduced by stretching. These observations are in agreement with the TEM results. The agglomeration level as observed on the micrographs was measured quantitatively and the results are shown in Fig. 8. When unstretched and at low stretch ratio (stretch ratio 2), PET loaded with 5 wt% mica has approximately twice the level of agglomeration of its 2 wt% loading counterpart. For the 5 wt% loaded PET, increasing the magnitude of stretching reduced the agglomeration level down to 5% or less; 2 wt% loaded PET on the other hand shows a reduction in agglomeration upon stretching, and further stretching does not have any significant effect probably because most of the large micron scale agglomerations had already been broken down to smaller sized particles. The effects of the observed morphological properties are related to their mechanical and barrier properties in later section.

Crystallinity

From Table 4, it can be seen that mica loading does not have any obvious effect on the crystallinity of PET for the same stretching conditions. Crystallinity is increased upon stretching due to strain induced crystallization which is behavior characteristic of PET. Mica loading has a nucleating effect and this was indicated by the increased crystallization temperature, [T.sub.C]. There was an increase of around 4[degrees] C in [T.sub.C] for the stretched samples. Although a nucleating effect of nanoclay was expected, the nanocomposiles did not show higher crystallinity and the melting temperature is not affected by mica loading. As the crystallinity of PET may be affected by the heating scan in DSC, WAXD analysis of the crystallinity was carried out to check for the validity. It is shown in Table 5 that the trend of both methods are comparable. Since there was no significant difference in crystallinity between the unfilled and filled PET, this factor was ruled out when examining the mechanical and barrier properties later. In studies of PET nanocomposiles, Calcagno et al. and Sloeffler el al. reported that pristine montmorillonite (Na-MMT) filled PET had higher crystallinity than virgin PET (25), (11). The different observations in the present study are probably due to a different type of filler. The Somasif clay platelet used in this work has larger dimensions than the MMT platelet. The nucleating effect of nanoclay was also reported in other studies of PET nanocomposites (3), (9), (25).

TABLE 4. Crystallinily, cold crystallization temperature ([T.sub.CC]).
crystallization temperature ([T.sub.C]), and melting temperature
([T.sub.M]) of PET nanocomposites at various stretch ratios.

                            Crystallinity  [T.sub.CC.sup.b]
                            (a) (%)        ([degrees]C)

Unstretched  Virgin PET     13 [+ or -] 3  132
             PET + 1% MAE   14 [+ or -] 3  130
             PET + 2% MAE   11             129
             PET + 5% MAE   10             127
Stretch      Virgin PET     36             114
ratio 2
             PET + 1 % MAE  31             112
             PET + 2% MAE   34             112
             PET + 5% MAE   34             112
Stretch      Virgin PET     40 [+ or -] 6  112 [+ or -] 2
ratio 2.5
             PET + 1% MAE   39             112
             PET + 2% MAE   37             110
             PET + 5% MAE   38             109 [+ or -] 2
Stretch      Virgin PET     41 [+ or -] 5  108
ratio 3
             PET + 1% MAE   37             107 [+ or -] 4
             PET + 2% MAE   36             106 [+ or -] 3
             PET + 5% MAE   40             109

                            [T.sub.M.sup.b]  [T.sub.C.sup.b]
                            ([degrees]C)     ([degrees]C)

Unstretched  Virgin PET     251              207
             PET + 1% MAE   251              207
             PET + 2% MAE   251              202
             PET + 5% MAE   25 i             202
Stretch      Virgin PET     251              196
ratio 2
             PET + 1 % MAE  252              196 [+ or -] 5
             PET + 2% MAE   251              195
             PET + 5% MAE   251              200
Stretch      Virgin PET     251              197
ratio 2.5
             PET + 1% MAE   252              196 [+ or -] 4
             PET + 2% MAE   252              201
             PET + 5% MAE   252              202
Stretch      Virgin PET     251              198
ratio 3
             PET + 1% MAE   251              196 [+ or -] 4
             PET + 2% MAE   251              202 [+ or -] 3
             PET + 5% MAE   251              202

(a) Less than 2[degrees] siandard deviation unless indicated.
(b) Less than 1[degrees] standard devialion unless indicated.

TABLE 5. Comparison of the crystallinity ratio obtained from
WAXD and DSC measurement.

                                          WAXD  DSC

Stretched PET (stretch ratio 3)             50   41
Stretched PET + 1% MAE (stretch ratio 3)    48   37
Stretched PET + 2% MAE (stretch ratio 3)    47   36
Stretched PET + 5% MAE (stretch ratio 3)    46   40


Mechanical Properties

Figure 9 shows the tensile mechanical properties of PET nanocomposites at different mica loading and stretch ratio. The modulus of PET was enhanced by increasing mica concentration. It is worth noting that the modulus of virgin PET increases with stretch ratio as a result of strain induced crystallization and orientation of the molecule chains. As reported in the literature, orientation and aspect ratio of tactoids have a significant effect on the mechanical properties of nanocomposites (15). Furthermore, conventional composites based model such as the Haipin-Tsai theory suggests that tactoids orientation, aspect ratio, incomplete exfoliation, and matrix modulus has an effect on the mechanical properties (23). Agglomeration is believed to play an important role in determining the properties of the nanocomposites as it reduces the surface of the filler exposed to the PET matrix. If a significant proportion of the tillers are in agglomerated form, then the effective amount of the reinforcement is less than the nominal filler quantity added into the matrix. The observed modulus enhancement is therefore expected to be less than predicted by theoretical models. All these factors were considered in examining the mechanical properties of the nanocomposite samples here.

For the unslrelched state, modest improvement in modulus was observed with increased mica loading. There was a 10%, 20%, and 30% modulus enhancement at 1 wt%, 2 wt%, and 5 wt% mica loading, respectively. When the unstretched samples were stretched to stretch ratio 2, the greater reinforcing effect of the mica was observed as in Fig. 9. This is due to the better alignment of the tactoids and the reduced agglomeration of the stretched nanocomposites. The virgin PET also shows an increase in modulus with stretch ratio. This behavior is due to the increase in crystallinity caused by strain induced crystallization and chain orientation. At a larger stretch ratio of 2.5, the reinforcing effect decreased for the PET loaded with 1 wt% and 2 wt% mica whereas PET loaded with 5 wt% mica maintained values comparable with those of stretch ratio 2. This behavior corresponds well with their orientation and agglomeration behavior. PET at 2 wt% mica loading showed reduced agglomeration levels at stretch ratio 2. Increasing the stretch ratio did not result in any significant reduction to the agglomeration level. PET loaded with 5 wt% mica on the other hand still showed significant reduction in the agglomeration level at a stretch ratio of 2.5 as the unstretched material presented many more aggregates and therefore larger stretching is needed to break them. The difference in tactoid orientation was only observable before and after stretching, and it is similar for all the stretching conditions and was therefore not enhancing the reinforcement effect further. The level of exfoliation, which is related to tactoid thickness and the aspect ratio, was less dominant when compared with the tactoid orientation and agglomeration level. PET loaded with 2 wt% and 5 wt% mica which had a better exfoliation level after stretching to a stretch ratio of 3 showed 19% and 27% modulus enhancement, respectively, similar to the enhancement level in the unstretched condition, i.e. 20% and 30% respectively. However, this could be due to the embedded filler especially when there are agglomerates of fillers that could make it more difficult for polymer chains to be orientated in one direction since the stretched materials are expected to have a high degree of chain orientation (13). Also, the stretched unfilled PET which has increased crystallinity and chain orientation was much stiffer compared with their unstretched counterparts. An attempt to reinforce a stiffer PET is more difficult than to reinforce a more flexible matrix at unstretched conditions. It is concluded that stretching increases the reinforcing effect of nanoclay as it induces alignment to the filler and reduces the agglomeration level.

The fracture and yield behavior of the materials are shown in Table 6. The fracture stress and strain of the unstretched samples could not be obtained as there were large disturbances to the stress-strain data. Mica concentration did not have a significant effect on the yield stress at the unstretched sheet and at a stretch ratio of 2. At stretch ratios of 2.5 and 3, increasing the mica concentration only had a slight effect. Fracture strain, apart from stretch ratio 2 samples, was reduced by mica loading due to brittlcness of the materials at high mica concentration. At higher mica loading percentages, mica tactoids act as stress concentration points and hence reduced the elongation to break. Increasing the mica concentration for the samples of stretch ratios 2.5 and 3, enhanced the fracture stress slightly. Samples of stretch ratio 2 exhibited dissimilar behavior than the rest. Their fracture stresses were lower than the virgin PET, but their elongations to break were higher.

TABLE 6. Crystallinity and mechanical properties of the nanocompoistes.

                            Crystallinity  Modulus  Yield   Fructure
                              (DSC) (%)               stress   stress
                                                      (MPa)     (MPa)

Unstretched  Virgin PET                13     2937      59       N/A
             PET + 1% MAE              14     3219      60       N/A
             PET + 2% MAE              11     3519      61       N/A
             PET + 5% MAE              10     3824      59       N/A
Stretch      Virgin PET                36     3528      78       121
ratio 2
             PET + 1 % MAE             31     4206      74        98
             PET + 2% MAE              34     4554      75        96
             PET + 5% MAE              34     4937      X7        97
Stretch      Virgin PET                40     3933      79       124
ratio 2.5
             PET - 1 % MAE             39     4621      86       120
             PET - 2% MAE              37     4780      87       119
             PET - 5% MAE              38     5894     100       133
Stretch      Virgin PET                41     4738      90       128
ratio 3
             PET + 1 % MAE             37     4988      97       133
             PET + 2% MAE              36     5651      98       127
             PET + 5% MAE              40     6023     111       135

                            Fracture
                             strain
                              (%)

Unstretched  Virgin PET          N/A
             PET + 1% MAE        N/A
             PET + 2% MAE        N/A
             PET + 5% MAE        N/A
Stretch      Virgin PET          193
ratio 2
             PET + 1 % MAE       246
             PET + 2% MAE        225
             PET + 5% MAE        188
Stretch      Virgin PET          146
ratio 2.5
             PET - 1 % MAE       150
             PET - 2% MAE        127
             PET - 5% MAE        127
Stretch      Virgin PET          100
ratio 3
             PET + 1 % MAE        94
             PET + 2% MAE         92
             PET + 5% MAE         74


The mechanical properties of PET nanocomposites were analyzed by the conventional composite, Halpin-Tsai model. Since the nanoclay, Somasif MAE, is a layered silicate, incomplete exfolialion reduces the effective filler modulus and aspect ratio by increased layer thickness. The Halpin-Tsai predictions are shown in Fig. 10. The moduli of virgin PET of stretch ratios 2, 2.5, and 3 were used as the matrix modulus in Eqs. 6 and 7.

The organic surfactant is subtracted from the clay weight fraction and converted to the volume fraction as shown earlier in the Experimental section. The 1 wt%, 2 wt%, and 5 wt% loading of Somasif MAE gives 0.32 vol%, 0.65 vol%, and 1.64 vol% silicate loading, respectively. However, for the reason of consistency with the nomenclature used in previous sections, the original wt% loading of clay (1 wt%, 2 wt%, and 5 wt%) is used as nominal clay loading in the Halpin-Tsai and Nielsen plots instead of the actual silicate vol% loading.

Based on the Halpin-Tsai predictions for completely exfoliated nanocomposites the relative modulus was predicted to be 1.79 at 5 wt% organoclay loading (corresponding to 1.64 vol% silicate), at stretch ratio 2. At stretch ratio 3, the relative modulus was predicted to be 1.59 at 5 wt% organoclay loading. The lower relative modulus predicted at stretch ratio 3 is due to the higher matrix modulus as the relative modulus is the ratio of the composites modulus to the matrix modulus. Tncomplete exfoliation has a major effect on the predicted modulus as shown in Fig. 10. The largest reduction in the relative modulus happens when the filler deviates from a completely exfoliated platelet to a stack of silicates consisting of two layers. Silicate stacks containing two silicate layers was predicted to have a relative modulus of 1.36 at 5 wt% organoclay loading under a stretch ratio of 2 which is far below the value of 1.79 obtained under complete exfoliation conditions. Similar behavior was also reported by Forties and Paul (23).

By comparing the experimental data with the theoretical predictions, as shown in Fig. 10, the predicted number of layers of silicate stacks acting as fillers is in the range between one and two layers. The prediction of the relative modulus based on the average tactoid thickness as measured quantitatively from TEM images are also included in Fig. 10. Here, the Halpin-Tsai theory underpredicled the experimental data. The Halpin-Tsai theory assumes that the filler particles are perfectly aligned and are of uniform shape, there is perfect bonding to the matrix and the filler does not change the properties of the polymer matrix. The mica dispersed in the PET matrix is highly orientated after undergoing large scale biaxial stretching as shown in the earlier publication (17), and it therefore meets the perfect alignment condition. However, in the real case the fillers are unlikely to be of uniform shape as the slacking structure could have been distorted (see Fig. 5) as a result of stretching. Quantitative measurements from the TEM images indicated that the stacking structure of the mica tactoids was distorted. The distorted stacking structure of the layered silicate would give a higher aspect ratio with the same number of layers per stack resulting in the Halpin-Tsai theory underpredicting the relative modulus. The underprediction by the Halpin-Tsai theory could also be a sign of the existence of a nanofillers-matrix interaction. It is noteworthy that the Halpin-Tsai prediction was also affected by the filler modulus. The filler modulus used here was assumed to be 178 GPa. However, it was shown by Chen et al. that the clay modulus can reasonably range from 178 to 275 GPa (26). The predicted modulus is shifted upward by increasing the filler modulus as shown in Fig. 11.

Figure 12(a) shows the effect of mica loading on the storage modulus, E' for PET nanocomposites at a stretch ratio of 3. The mica loading has a significant effect on the storage modulus of PET in the temperature range 70 to 100CC. A rapid decrease in E' was observed in the virgin PET due to the transition from glass to rubber-like behavior. The PET nanocomposites exhibited the same behavior, but the loss in E' due to glass transition was not as significant as observed on virgin PET. Further heating above 120 C declined the E' of virgin PET rapidly. The E' of the nanocomposites on the other hand declined at a slower rate. The 0 was enhanced not only at room temperature, but also at high temperature. In fact, enhancement at high temperature, which is above the glass transition ([approximately equal to]93[degrees] C), was even more significant. There was a 110% increment in E' at 5 wt% mica loading at 120[degrees] C. Since it was revealed earlier that there was no significant change in the crystallinity, the enhancement observed here was mainly due to the reinforcing effect of the fillers with increasing mica loading. The higher reinforcing effect above the glass transition temperature, [T.sub.g] is due to the softening of the polymer matrix when it transforms from the glassy to the rubbery state, but the filler remains rigid throughout this temperature range. The reinforcing effect is therefore more pronounced on a softer matrix. A simitar observation was reported for nylon 6-clay nanocomposites (23). Figure 12(b) shows the effect of mica loading on the tan delta peak. Tan delta which is conventionally viewed as a viscoelastic index, shows a suppression of its peak value by mica loading in the glass transition region, indicating the nanocomposites are more elastic in that temperature range. The [T.sub.g], as depicted by tan delta peak, was shifted towards higher temperatures with increasing mica loading. The improved rigidity of the nanocomposites at high temperature implies that the formulation of nanocomposites could extend the applications of PET in containers for hot filling where high temperature rigidity is critical.

Barrier Properties

Figure 13 shows the oxygen transmission rate of unstretched PET nanocomposites of 2 wt% and 5 wt% mica loading. Data for virgin PET is included for comparison. Since the transmission rate curves shown have not yet reached their equilibrium state, the permeability coefficients for the unstretched samples are not determined from them. However, the effects of mica on PET barrier properties can still be compared qualitatively from the curves. The nanocomposites showed much lower oxygen transmission rates than virgin PET. One of the reasons of the barrier enhancement is the increase in the tortuosity of the gas diffusion path by mica filler and hence increases the barrier property of the nanocomposites. Other factor such as the flipping of phenyl ring on the PET chains may affect the barrier properties (27). However, it is unclear if the mica loading could affect the flipping of the phenyl ring. The stretched samples, on the other hand, can reach the equilibrium stale faster due to their much lower thickness. Their permeability coefficient can be derived from the transmission rate data. The oxygen permeability coefficient of biaxially stretched PET nanocomposites, of stretch ratio 3, is shown in Table 7. The nanocomposites have better barrier properties than the virgin PET. There was a 22% reduction in the oxygen permeability coefficient at 5 wt% mica loading.

TABLE 7. Oxygen permeability of biaxialiy sirelched (stretch ratio 3)
PET nanocompositcs.

            [O.sub.2] permeability coefficient [[cm.sup.3]
                        mm]/[[m.sup.2] day bar]

Virgin PET                                            2.35
(a)
1% MAE                                                2.15
2% MAE (a)                                            1.82
5% MAE                                                1.84

(a) Data reproduced from Ref. 17 with permission from Elsevier Science.


Figure 14 shows Nielsen's model prediction of the barrier property of PET nanocomposites. As in the Halpin-Tsai plot, the nominal wt% of mica loading is plotted instead of the actual vol% of mica loading. Incomplete exfoliation is accounted in the model by reducing the filler aspect ratio which is equivalent to increasing the number of platelet layers per tactoid. The relalive permeability prediction of Nielsen's model was reduced by increasing the exfoliation level (increasing aspect ratio). Based on Nielsen's model, complete exfoliation gives a relative permeability of 0.1 at 5 wt% organoclay loading (corresponding to a 1.64 vol% silicate loading). A large increment to the relative permeability was observed when the exfoliation level decreased from one layer to two layers of silicates. It is shown in Fig. 14 that the experimental data falls in the range of 5 to 20 layers. However, quantitative analysis of the TEM images shows that the average layer is in the range of 2 to 3 layers. Nielsen's model is based on several assumptions. One of the assumptions is that the tillers do not affect the properties of matrix. In fact, the crystallization of the PET matrix is strongly affected by nanofiller loading. Although we report earlier that the percentage of crystallinity was not affected by mica loading, it is likely that the structure and type of the crystal may be different, for example the nanoclay might promote the formation of smaller and incomplete crystals. This phenomenon has been reported in the literature (3), (25). Hwang et al. observed that the PET nanocomposites had smaller crystals compared with virgin PET (3). A similar observation was also reported by Calcagno et al. (25). In a study of the barrier properties of virgin PET, Natu el al. showed that for thermally crystallized PET of the same level of crystallinity but different crystals size, larger crystals were formed leading to better barrier properties as there was more tortuosity compared with their smaller crystals counterparts (28). This might have caused the Nielsen's model to overpredict the barrier property of the nanocomposites.

CONCLUSIONS

The effects of biaxial stretching on the morphology of PET-mica and the resulting mechanical and barrier properties were examined. Increasing the mica loading had a prominent effect on the mechanical properties even though there is a tendency to affect the level of agglomeration and exfoliation. Stretching tends to align the mica tactoids and reduces the agglomeration level, hence increasing the reinforcing effect. Mica loading enhanced the modulus of PET especially at stretch ratios of 2 and 2.5 as a result of more aligned tactoids and reduced agglomeration. The reinforcing effect lessened at a stretch ratio of 3 when the PET matrix was more crystallized and had increased rigid orientated amorphous phase compared with the stretch ratio of 2 and 2.5 and further improvements in tactoids alignment and the reduction in the agglomeration level was negligible. DMTA analysis revealed that mica loading improved the mechanical properties of PET at elevated temperatures which possibly extend its usefulness in applications such as hot filling of containers. The storage modulus, E', was enhanced at room temperature and the enhancement was even greater above the glass transition temperature, [T.sub.g]. Also, the decrease in E' of the nanocomposites was at a slower rate above 120[degrees] C. The barrier properties were enhanced by the mica loading in both the unstretched and stretched states (stretch ratio 3). Attempts were made to compare the results from both the Halpin-Tsai model and the Nielsen model with the mechanical and barrier data. It was found that the data falls in the incomplete exfoliation region. The Halpin-Tsai theory undeipredicted the relative modulus of the PET nanocomposites whereas the Nielsen's model overpredicts the relative permeability.

ACKNOWLEDGMENTS

The authors thank Dr. Mahendrasingam for his advice on crystanillinily analysis using WAXD.

REFERENCES

(1.) D.R. Paul and L.M. Robeson, Polymer, 49, 3187 (2008).

(2.) S.H. Kim and S.C. Kim, J. Appl. Polym. Sci., 103, 1262 (2007).

(3.) S.Y. Hwang, W.D. Lee, J.S. Urn, K.H. Park, and S.S. Im, J. Polym. Sci. Pol. Phys., 46, 1022 (2008).

(4.) T.D. Fornes, P. J. Yoon, and D.R. Paul, Polymer, 44, 7545 (2003).

(5.) M.T.M. Bizarria, A.L.F. De, M. Giraldi, CM. de Carvalho, J.I. Velasco, M.A. d'Avila, and L.H.I. Mei, J. Appl. Polym. Sci., 104, 1839 (2007).

(6.) U. Gunnendi, J.I. Eguiazabal, and J. Nazabal, Macromol. Mater. Eng., 292, 169 (2007),

(7.) M.D. Sanchez-Garcia, E. Gimenez, and J.M. Lagaron, J., Plast. Film. Sheet., 23, 133 (2007).

(8.) M. Kracalik, M. Studenovsky, J. Mikesova, A. Sikora, R. Thomann, I. Friedrich, and J. Simonik, J. Appl. Polym. Sci., 106, 926 (2007).

(9.) J.W. Chung, S.B. Son, S.W. Chun, T.J. Kang, and S.Y. Kwak, Polym. Degrad. Stab., 93, 252 (2008). (13.)

(10.) A. Ammala, B. Bell, and D. Katherine, Compos. Sci. Technol., 68, 1328 (2007).

(11.) K. StoeHler, P.G. Larleur, and J. Denault, Polym. Degrad. Stab., 93, 1332 (2008).

(12.) S. Jeol, F. Fenouillot, A. Rousseau, K. Masenelli-Varlot, C. Gaulhier, and J.F. Briois, Macromolecules, 40, 3229 (2007).

(13.) D.W. LilchJield and D.G. Baird, Polymer, 49. 5027 (2008).

(14.) R.K. Shah, R.K. Krishnaswamy, S. Takahashi, and D.R. Paul, Polymer, 47, 6187 (2006).

(15.) J.I. Weon, and H.J. Sue, Polymer, 46, 6325 (2005).

(16.) R.K. Bharadwaj, Macromolecules, 34, 9189 (2001).

(17.) R.S. Rajcev, E. Harkin-Jones, K. Soon, T. McNally, G. Menary, C.G. Armstrong, and P.J. Martin, Eur. Polym. J., 45, 332 (2009).

(18.) K. Soon, E. Harkin-Jones, R. Rajeev, G. Menary, T. McNally, P.J. Martin, and C.G. Armstrong, Polym. Int., 58, 1134(2009).

(19.) P.J. Martin. C.W. Tan, K.Y. Tshai, R. McCool, G. Menary, C.G. Armstrong, and E. Harkin-Jones, Plast. Rubber Compos., 34, 276 (2005).

(20.) A. Vermogen, K. Masenelli-Varlot. and R. Seguela, Macro-molecules, 38, 9661 (2005).

(21.) F, Chavarria. and D.R. Paul, Polymer. 45, 8501 (2004).

(22.) P.W. Zhu, J. Tong, A. Philips, and G. Edward, Macromole-cules. 39, 1821 (2006).

(23.) T.D. Forncs, and D.R. Paul, Polymer. 44, 4993 (2003).

(24.) K. Masenelli-Varlot, G. Vigier, A. Vcnnogen, C. Gaulhier, and J.Y. Cavaillc, J. Poiym. Sci. PartB: Polym. Rhys., 45.1243 (2007).

(25.) C.I.W. Calcagno, CM. Mariani. S.R. Teixeira. and R.S. Mauler, Polymer, 48, 966 (2007).

(26.) B. Chen, and J.R.G. Evans, Scripta Mater., 54, 1581 (2006).

(27.) R.M, Gohil, J. Appl. Polymer. 48, 966 (2007).

(28.) A.A. Natu. E.A. Lolgren. and S.A. Jabarin, Polym. Eng. Sci., 45, 400 (2005).

Correspondence to: K. Soon; e-mail: Ksoon01@qub.ac.uk

Contract grant sponsor: Engineering and Physical Sciences Research Council (EPSRC), UK.

DOI 10.1002/pen.22114

Published online in Wiley Online Library (wileyonlinelibrary.com).

[c] 2011 Society of Plastics Engineers

Kok Soon, Eileen Harkin-Jones, Rajvihar S. Rajeev, Gary Menary, Peter J. Martin, Cecil G. Armstrong School of Mechanical and Aerospace Engineering, Queen's University Belfast, BT9 5AH, UK
COPYRIGHT 2012 Society of Plastics Engineers, Inc.
No portion of this article can be reproduced without the express written permission from the copyright holder.
Copyright 2012 Gale, Cengage Learning. All rights reserved.

Article Details
Printer friendly Cite/link Email Feedback
Author:Soon, Kok; Harkin-Jones, Eileen; Rajeev, Rajvihar S.; Menary, Gary; Martin, Peter J.; Armstrong, Cec
Publication:Polymer Engineering and Science
Article Type:Report
Geographic Code:4EUUK
Date:Mar 1, 2012
Words:9699
Previous Article:Effect of phenol functionalized carbon nanotube on mechanical, dynamic mechanical, and thermal properties of isotactic polypropylene nanocomposites.
Next Article:Effect of carbon black structure on low-strain conductivity of polypropylene and low-density polyethylene composites.
Topics:

Terms of use | Privacy policy | Copyright © 2019 Farlex, Inc. | Feedback | For webmasters