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BY-NC-ND 3.0 license Open Access Published by De Gruyter January 16, 2014

Evaluation of flow-induced nanoclay orientation and microstructural stability in polyethylene/clay nanocomposites via melt rheological and thermal analysis

Nima Mostofi Sarkari, Ali Asghar Katbab and Hossein Nazockdast
From the journal e-Polymers

Abstract

The effects of shear rate upon the flow-induced nanoclay orientation and morphological stability in film blown polyethylene clay nanocomposites were studied by means of linear and nonlinear rheological characterization parallel with differential scanning calorimetry and X-ray diffraction (XRD) analyses. Nanocomposite samples were prepared using a modular twin screw extruder followed by film processing technique. XRD analysis performed on film samples showed that the samples exhibited intercalated/exfoliated microstructure. The 3D physical networks were formed by the clay nanolayers in the structures of undrawn samples. However, the breakdown of the clay physical networks during film processing as a result of the imposed shear field within the die area and also an elongational flow field was evidenced. Time sweep test performed at various shear rates and shearing times using a rheometric mechanical spectrometer showed that, in all samples, the time required for the restructurization of the clay nanolayers during relaxation of the melt was found to be higher than 3600 s.

1 Introduction

Polymer-layered silicate nanocomposites have received substantial attention from both the industrial and the academic fields over the last two decades. This has stemmed from dramatic enhancement in thermal and mechanical properties along with reduced gas permeability and decreased flammability that can be achieved with considerably lower nanoclay loadings compared to conventional polymer composites (1–5). It has been shown that the final microstructure of polymer/clay nanocomposites plays a significant role in determining efficiency in enhanced properties and rheological properties of these composites (3, 6–8).

In general, two types of microstructures, namely, intercalated and exfoliated, have been recognized for polymer/nanoclay nanocomposites. In the intercalated microstructure, the extended polymer matrix chains reside between the silicate layers, resulting in increased interlayer d-spacing. In exfoliated microstructures, the separated silicate layers are randomly dispersed in the polymer matrix (9). Many researchers have reported that polyolefins including polyethylene (PE) are not capable of intercalating and increasing the interlayer d-spacing of silicate layers. Therefore, in order to produce polyolefin/nanoclay nanocomposites with desirable microstructure, it is essential to use functional oligomers as compatibilizer. The most common compound used for producing polyolefin nanocomposites is PE-grafted maleic anhydride (PE-g-MA). PE is used in a number of applications including flexible films, packaging films and agriculture films (10). The main objective of using nanoclay PE nanocomposites is to improve the gas barrier properties of the packaging films. Thus, the effective aspect ratio of nanolayers, which is controlled by the extent of intercalation as well as the nanoclay orientation, can play a significant role.

Many factors have been recognized to influence the microstructure and properties of the final nanocomposite, including molecular weight, melt flowability of the matrix, interfacial compatibilization, clay treatment, weight fraction of the added organoclay (OC), process condition and so on (11–15). The main purpose of optimizing these parameters is to achieve a nanocomposite with higher degrees of exfoliation and intercalation, thereby increasing the average gallery distance.

Nanoclay-filled PE nanocomposite systems have gained attraction both in the academic and the industrial society because conventional equipment (extruder and injection machines) can be used to produce barrier films. The mechanical and barrier properties of the produced films depend strongly on the process itself, on the rheological properties of the filled polymer and on the orientation of nanoclay during processing.

It is well known that flow-induced molecular orientation and its influence on the degree of crystallinity has a significant effect on the optical and mechanical properties of PE films. Several works have been published focusing on the flow-induced orientation hierarchy of nanoparticles and its impact on molecular orientation in polymer-layered silicate nanocomposites (16–24).

The standard fiber orientation model is based on Jeffery’s equation for the motion of a single fiber in an infinite Newtonian matrix (25) with an isotropic rotary diffusion term to account for fiber-fiber interactions (26).

Rheological data on polymer nanocomposites are often interpreted as two distinct structural characteristics: the orientation distribution of individual clay particles and particle network structures, mediated by interparticle interactions; indeed, attempts have been made to model both types of structures to provide insight into the interpretation of the transient rheological phenomena (27). The use of stress growth at the startup of steady shear experiments, also called shear experiments (28), is well known because of the change in orientation and orientation distribution of fibers in a polymer matrix due to shear flow (29). Generally, the linear viscoelastic behavior of nanocomposites shows a significant change from liquid-like flow behavior to pseudo-solid-like behavior, depending on the concentration and the state of dispersion. The solid-like behavior of nanocomposites is often attributed to the existence of a percolation network formed by clay particles (28).

Nawani et al. (16) demonstrated that the use of multidirectional 2D small-angle X-ray scattering can quantitatively describe the orientation distribution of OC in melt-pressed nanocomposite films. They showed that in nanocomposites with mixed intercalation/exfoliation morphology, the dispersed particles possessed partial orientation parallel to the in-plane direction of the film, and the higher content of the clay loading showed a higher clay orientation. Moreover, Dykes et al. (30) studied the shear-induced orientation in polymer/clay dispersions using in situ X-ray scattering and rheological measurements with two different nanoclays. They showed a finite orientation angle, indicating a systematic misalignment of the particle along axes relative to the flow direction. Khalkhal and Carreau (31) also used rheological measurements to study the extent of structure buildup at rest in carbon nanotubes (CNTs) dispersed in an epoxy by a set of transient flow measurements. They showed that in the absence of flow, Brownian motion plays an important role in the structure buildup of the CNT suspensions.

Lele and Mackley (32) studied the flow-induced orientation of intercalated polypropylene (PP) nanocomposites for both compatibilized and uncompatibilized systems using in situ X-ray scattering. They observed for the compatibilized system, following the cessation of shear, a rapid decrease of the particle orientation, with time up to 1000 s, which then remained constant at longer times. They suggested that the rapidly decreasing orientation was due to the coupling of polymer chains and clay platelets. The stress relaxation of the polymer matrix might accelerate the orientation relaxation of the clay particles. Lee and Han (33) observed, for polycarbonate/OC nanocomposites, stress overshoot in startup shear flow experiments. These overshoots were attributed to the presence of strong attractive interactions (hydrogen bonds between the polycarbonate and the surfactant at the surface of the clay). Moreover, most of the studies on flow-induced crystallization have been carried out on relatively pure polymers, without the deliberate addition of particles. Although the combined effect of shear flow and particles on crystallization of polymers has already received some attention in the literature, not many systematic studies have been done so far. Recently, D’Haese et al. (34) studied the combined effect of the presence of particles and flow on the crystallization of filled PP using rheo-optical approach. They found that at high enough flow rates, flow will dominate the nucleation process independent of the presence of particles, whatever their size, shape and concentration.

Many researchers’ (35–37) works have been carried out regarding the orientation of the fiber particles in suspension systems, and their rheological behavior has been modeled in shear flow field. However, few works have been reported on the orientation of nanoclay during film processing of the OC-reinforced nanocomposites. The importance of this aspect is that the mechanical as well as the barrier properties of the prepared film samples are significantly controlled by the clay nanolayers’ orientation and their extent of dispersion within the polymer matrix. Hence, in the present work, the effect of melt processing history, including flow field (shear and elongation), deformation rate and drawdown ratio, on the flow-induced orientation as well as reorientation has been investigated by means of melt linear and nonlinear rheology with thermal analysis.

2 Materials and methods

2.1 Materials

A commercial film-grade linear low-density PE (LLDPE), with melt flow index (MFI 190°C/2.16 kg) of 2.2 g/10 min and commercial name of LL0220 AA supplied by Tabriz Petrochemical Co. (Iran), was used as matrix. PE-g-MA with melting temperature of 134°C and trade name of Fusabond E-100 was purchased from DuPont Co (DE, USA) and utilized as interfacial compatibilizer. Montmorillonite (MMT) organically modified with a quaternary ammonium salt (dimethyl, dehydrogenated tallow, quaternary ammonium) and gallery distance of 24.2 Å with commercial trade name of Cloisite 20A was purchased from Southern Clay Company (TX, USA), and used to prepare nanocomposites.

2.2 Nanocomposite preparation

Nanocomposites were prepared by means of direct melt mixing process using a modular Brabender twin screw extruder model DSE25 (Germany) with screw speed of 150 rpm. For this purpose, all materials were dried at 80°C in a vacuum oven prior to the mixing process. The compositions of the prepared samples and their corresponding designations are shown in Table 1.

Table 1

Compositions of the prepared nanocomposite samples.

Sample designationLLDPE (wt%)Organoclay20A (wt%)Compatibilizer E100 (wt%)
N38839
N580515
N772721
Reference84016

2.3 Film preparation process

In order to study the effects of flow field on nanoclay orientation, film samples were prepared by using flat film method. In this process, the film samples were prepared at different drawdown ratios using a Brabender single screw extruder (Germany) equipped with a slit die (dimensions: width=100 mm, thickness=0.5 mm), and the produced films were cooled at room temperature after being drawn at different speeds. For all the produced films, the cooling air was set at T=25°C. The film sample designations are illustrated in Table 2.

Table 2

Sample designation for films produced from N5 sample.

Sample designationDrawdown speed (m/min)
F-F-N5-11
F-F-N5-22
F-F-N5-44
F-F-N5-66

2.4 WAXD experiments

To investigate the degree of intercalation of the clay particles, wide-angle X-ray diffraction (XRD) (WAXD) analysis was performed on the prepared composite samples using a Philips X’pert MPD diffractometer (Netherlands) with an integral germanium detector and a Cu Kα radiation source ray with a wavelength of 1.54 Å. The system consisted of a rotating anode generator and a wide-angle powder goniometer fitted with a high-temperature attachment. The generator was operated at 40 kV and 40 mA, and samples were scanned within the diffraction angle of 2θ=1.01°–11° at a scan rate of 1.2°/min.

2.5 Transmission electron microscopy

To further study the extent of intercalation and exfoliation of the clay particles, transmission electron microscopy (TEM) was used. The samples were sliced into ultrathin sections (30–50 nm thick) with a microtome in liquid nitrogen.

TEM experiments were conducted on thin sections of the nanocomposites. For this purpose, bright-field TEM was performed using a Zeiss LEO-912 microscope (Germany) and acceleration voltage of 80 kV.

2.6 Rheological measurements

In order to study the rheological behavior of nanocomposite samples and, more importantly, to explore the parameters affecting flow-induced nanoclay orientation and kinetics of nanoclay reorientation, we needed to design a combination of linear and nonlinear viscoelastic experiments.

The melt linear and nonlinear rheological analyses were carried out using a Paar Physica UDS200 rheometric mechanical spectrometer (Austria). Measurements were done utilizing a parallel plate fixture with diameter equal to 25 mm and gap distance of 1 mm. Experiments were carried out in linear and nonlinear viscoelastic regions. In order to find out the linear viscoelastic region, the samples were subjected to strain sweep test at 175°C, and variation of melt complex viscosity vs. strain was recorded at a fixed frequency of 1 s-1. To examine the nonlinear characteristics and degree of network breakdown, the nanocomposite samples were subjected to melt dynamic test at three different shear rates (0.1, 0.5 and 1 s-1) and different times of shearing. The nonlinear melt rheological behavior of prepared composites was investigated by means of the startup of steady shear flow and flow reversal experiments in rotational deformations as well as frequency sweep experiments. In the startup of steady shear flow experiments, a constant shear rate was imposed on the melt samples and the transient stress was monitored with time. In flow reversal experiments, the sample was first sheared for a certain period of time, and then, immediately after cessation of startup flow, the sample was again subjected to the same shear rate as in the first step but in the reverse direction. The difference between the height of the stress overshoot in the shear flow and the reverse flow can be attributed to the orientation of the nanoclay in shear flow. In order to probe the reformation of network-like structure, the storage modulus recovery experiment was performed on nanocomposites samples. The amount of increase in storage modulus with time can be a measure of network restructuring. In this experiment, the samples were first subjected to a constant shear rate for different time periods, and then the flow was ceased and the samples were immediately subjected to a small strain amplitude (0.01) oscillatory deformation at a constant angular frequency (1 s-1) in order to prevent further structural breakdown from occurring during measurement of the storage modulus.

The melt linear viscoelastic measurements were carried out at a constant amplitude of 1% within the frequency range of 625–0.01 s-1 at different temperatures under dry nitrogen atmosphere.

2.7 Differential scanning calorimetry analyses

To study the effect of flow field-induced nanoclay orientation on the crystallinity of the samples, differential scanning calorimetry (DSC) was performed on both drawn film samples and undrawn nanocomposite samples. For this purpose, a Mettler Toledo thermal analyzer (OH, USA) was employed, and tests were carried out under nitrogen atmosphere with a heating rate of 5°C/min.

3 Results and discussion

Figure 1A shows the WAXD pattern of organically modified MMT (Cloisite 20A). As can be seen, the OC exhibits two characteristic peaks at 2θ=3.48° and 2θ=7.08° corresponding to the basal spacing of 24 Å and 12.5 Å, respectively. The peak observed at 2θ=3.48° corresponds to the interlayer spacing of the modified layers of MMT, while the peak at 2θ=7.08° is related to the unmodified layers. Figure 1B shows the WAXD patterns of the nanocomposite sample N5 containing 5% OC. It should be noted that the WAXD peaks due to the PE crystals do not appear in this angular range. It is observed that in the XRD pattern of the N5 sample, the OC characteristic peak has been moved to the lower angle corresponding to the gallery spacing of 29 Å. This indicates the development of an intercalated type of structure for the N5 nanocomposite sample.

Variation of the melt viscosity and melt storage modulus vs. frequency within the linear region for both unfilled PE/PE-g-MA reference sample and intercalated nanocomposite sample containing 3%, 5% and 7% OC are illustrated in Figure 2. As can be observed, all nanocomposite samples exhibit pseudo-solid-like behavior along with viscosity upturn within the low-frequency region. However, the nonterminal characteristic is more pronounced for the N5 and N7 nanocomposites, indicating more potential for the formation of 3D physical networks of the clay nanoplatelets. These results, along with a typical TEM image of the N5 nanocomposite sample shown in Figure 3, reveal a reasonable extent of intercalation and good state of dispersion of nanoclay throughout the PE matrix, which are consistent with the XRD analyses.

Figure 1 XRD patterns for (A) pristine OC and (B) nanocomposite samples containing 5% OC (N5).

Figure 1

XRD patterns for (A) pristine OC and (B) nanocomposite samples containing 5% OC (N5).

Figure 2 Variation of the melt dynamic (A) storage modulus (G′) and (B) complex viscosity (η*) vs. angular frequency for both reference and nanocomposite samples.

Figure 2

Variation of the melt dynamic (A) storage modulus (G′) and (B) complex viscosity (η*) vs. angular frequency for both reference and nanocomposite samples.

Figure 3 TEM image of the interfacially compatibilized nanocomposite sample containing 5 wt% OC.

Figure 3

TEM image of the interfacially compatibilized nanocomposite sample containing 5 wt% OC.

Results of the frequency sweep experiments performed on the N5 nanocomposite sample after being presheared at a shear rate of 0.5 s-1 for different time periods are demonstrated in Figures 4 and 5.

Figure 4 Variation of (A) storage modulus (G′) and (B) melt complex viscosity (η*) vs. frequency for the N5 nanocomposite sample after being presheared at  for different time periods (900, 1800 and 3600 s).

Figure 4

Variation of (A) storage modulus (G′) and (B) melt complex viscosity (η*) vs. frequency for the N5 nanocomposite sample after being presheared at

for different time periods (900, 1800 and 3600 s).

Figure 5 Variation of (A) storage modulus (G′) and (B) melt complex viscosity (η*) vs. frequency for the N5 nanocomposite sample after being presheared at  for different time periods (900, 1800 and 3600 s).

Figure 5

Variation of (A) storage modulus (G′) and (B) melt complex viscosity (η*) vs. frequency for the N5 nanocomposite sample after being presheared at

for different time periods (900, 1800 and 3600 s).

It is clearly observed in Figure 5 that the presheared N5 nanocomposites exhibit lower elastic modulus as well as melt complex viscosity compared to the unsheared N5 nanocomposite sample, implying the breakdown of the physical networks of the clay nanolayers by the imposed preshearing process. However, the extent of viscosity reduction was more significant for the samples that were presheared at higher shear rates (0.5 and 1.0 s-1) and with longer times of shearing (Figure 6).

Figure 6 (A) Melt dynamic storage modulus and (B) complex viscosity vs. angular frequency for the N5 nanocomposite sample after being sheared for 1800 s with different shear rates compared to the unsheared counterpart sample.

Figure 6

(A) Melt dynamic storage modulus and (B) complex viscosity vs. angular frequency for the N5 nanocomposite sample after being sheared for 1800 s with different shear rates compared to the unsheared counterpart sample.

At low shear rates

the storage modulus and complex viscosity remain almost unchanged. This shows that the nanoclay platelets have been oriented only marginally by the shear flow at low shear rates and the network itself has not been disturbed by the deformation. At higher shear rates
and 1.0 s-1), the OC network changes considerably. The oscillatory measurements before and after steady shear at
s-1 show a large decrease in both storage modulus and complex viscosity. This suggests that the network formed between OC platelets completely breaks down at higher shear rates. It should also be noted that the data shown in Figure 5 were measured from high to low frequencies, i.e., the low-frequency data were measured after about 10-min delay times, which explains the increase in low storage modulus in the low-frequency regions.

Similar results have also been observed by Dijkstra et al. (35) for polycarbonate/multiwalled CNT (PC/MWCNT) nanocomposites, which showed a similar behavior for their nanocomposite system after being sheared.

These results led to the conclusion that there exists a minimum shear rate for the breakdown of the physical networks in the structure of the nanocomposite. Breakdown of the physical networks inside the extruder during flat film process can help to increase the extent of platelet orientation within the die area, which would lead to the enhanced barrier properties. In other words, the presence of the physical networks within the internal structure of the LLDPE/OC nanocomposite is beneficial provided that they are broken before the drawn up region during the film process.

However, the degree of orientation of the clay nanolayers within the die area in the film process and the tendency for reorientation with time would determine the degree of microstructural stability and hence the performance of the corresponding film. Therefore, study of the correlation between the extent of shear-induced orientation of the clay nanoplatelets and their restructuring with time was attempted in the present work. For this purpose, samples of the prepared N5 nanocomposite were sheared at various shear rates for different time periods and then immediately subjected to oscillation mode at the angular frequency of 1 s-1 and strain amplitude of 1% (linear region). The initial storage modulus (G′) and its rate of increase with time were studied. The results are demonstrated in Figure 7. As can be observed clearly in this figure, the initial dynamic storage modulus of the samples decreased with shearing time, indicating the destructuring of the clay physical networks by the applied shear fields, which would lead to the partial orientation of the clay nanolayers. The tendency of the nanocomposite melts to restructure the broken physical networks by reorientation of the clay platelets is illustrated in Figure 8. As can be observed in this figure, the sheared sample exhibited lower dynamic storage modulus compared to the unsheared counterpart sample. However, the rate of increase in storage modulus by oscillating time was shown to be higher for the sheared example. This would be attributed to the reorientation tendency for the clay nanolayers to restore their aggregated type of structures, and hence more thermodynamically stable microstructure. From these results, one would conclude that the higher the rate of restructuring by the melted nanocomposite, the lower would be the stability of the microstructure for its corresponding film product.

Figure 7 Initial dynamic storage modulus at 175°C for the N5 nanocomposite sample sheared at various shear rates and different shearing times.

Figure 7

Initial dynamic storage modulus at 175°C for the N5 nanocomposite sample sheared at various shear rates and different shearing times.

Figure 8 Dynamic storage modulus (G′) as a function of oscillating time at 175°C for the samples sheared at  for different shearing times.

Figure 8

Dynamic storage modulus (G′) as a function of oscillating time at 175°C for the samples sheared at

for different shearing times.

Moreover, the nanocomposite has viscoelastic characteristics and hence is not able to recover its original morphology after being sheared. The slight increase in storage modulus with time measured for the unsheared sample indicates the potential for more structurization by the clay nanolayers during melt oscillating test within the linear viscoelastic region and, hence, greater extent of networks. Nevertheless, none of the sheared samples could gain storage modulus comparable to that of the unsheared counterpart within the tested time period, implying that the oriented nanolayers needed a very long time for their reorientation. However, as can be noticed, increase in G′ with oscillating time due to the nanoclay restructuring is very slow and does not seem to reach the initial storage modulus of the unsheared sample. It is clearly seen that low-frequency storage modulus of the nanocomposite sample after being sheared is much lower than values measured before shearing.

The rate of melt dynamic modulus recovery by the N5 nanocomposite samples sheared at various shear rates and different shearing time periods was evaluated by measuring the percentage of modulus recovery at equal oscillating time (1500 s) using the following equation:

In Eq. [1],

and
are the melt storage moduli measured at the start and 1500 s after performing oscillation, respectively. The results are illustrated in Figure 9. As can be observed in the figure, the higher the imposed shear rate
the more would be the extent of modulus recovery, indicating higher instability for the nanocomposite morphology because the degree of network destruction is increased.

Figure 9 Percentage of network recovery after preshearing at T=175°C for shearing times 900, 1800 and 3600 s;  0.5 and 1.0 s-1.

Figure 9

Percentage of network recovery after preshearing at T=175°C for shearing times 900, 1800 and 3600 s;

0.5 and 1.0 s-1.

To distinguish between network breakup and orientation of the clay platelets in shear flow field, flow reversal tests were performed on the nanocomposite samples. The difference between the overshoot peak of stress in the first and second stage (reverse flow) can be a measure of stress needed to orient the nanoclay particles in the flow field. The results obtained for N5 nanocomposite sample subjected to flow reversal tests at various shear rates for different time periods are demonstrated in Figure 10A–C.

Figure 10 Result of flow reversal test with (A)  (B)  and (C)  for N5 nanocomposite at 175°C for different time periods. The curves were shifted upward for clarity.

Figure 10

Result of flow reversal test with (A)

(B)
and (C)
for N5 nanocomposite at 175°C for different time periods. The curves were shifted upward for clarity.

Comparing the results presented in Figure 10A–C implies that at shear rate of 0.1 s-1, the morphology of the nanocomposite remains unchanged by increasing the shearing time, which evidences that network breakdown can hardly occur at this shear rate. However, at shear rates of 0.5 and 1.0 s-1, both breakdown of the networks and orientation of the clay nanolayers could occur, and the extent of network recovery via reorientation of the clay nanolayers decreases.

In Figure 11, the height of the reverse peak in flow reversal tests for different shear rates and shearing times.

0.5 and 1.0 s-1 with shearing time of 900, 1800 and 3600 s) has been compared. For better understanding of the orientation process, these values were normalized to their corresponding shear rate and are presented in the same figure. As can be seen, at low shearing rates (i.e.,
the amount of nanoclay orientation increases linearly with increasing shearing time, but at high shear rates (i.e.,
increasing the time of shearing does not affect the degree of clay nanolayer orientation significantly. By comparing the results in Figure 11B, one would conclude that the slope of the decrease in shear stress peak is almost the same for samples sheared with
and 0.5 s-1, whereas at higher shear rates
the trend is different. Moreover, the extent of nanoclay orientation in the samples sheared at
is much less than those sheared at
and 0.5 s-1.These results suggest that there exists a minimum shear rate at which the orientation of clay nanolayers can occur in the flow field.

Figure 11 (A) Stress overshoot height in reverse flow and (B) normalized stress overshoot height for different shear rates in start-up flow.

Figure 11

(A) Stress overshoot height in reverse flow and (B) normalized stress overshoot height for different shear rates in start-up flow.

To get an insight into the effect of elongational flow field on the network breakup and nanoclay orientation, linear viscoelastic test was performed on the N5 nanocomposite film samples prepared at different elongation rates during flat film processing, and the results are illustrated in Figure 12A and B. As can be observed, the elongational flow field resulted in the decrease of the low-frequency storage modulus and complex viscosity. It should be mentioned that similar experiments were performed on the unfilled reference sample (LLDPE/compatibilizer), and no difference was observed between the undrawn and drawn samples. These results lead to the conclusion that the main causes for the decrease of the storage modulus and complex viscosity of the drawn nanocomposite film are breakdown of the clay physical networks and nanolayer orientation.

Figure 12 Storage modulus (A) and complex viscosity (B) of the drawn nanocomposite film samples compared to the undrawn counterpart sample.

Figure 12

Storage modulus (A) and complex viscosity (B) of the drawn nanocomposite film samples compared to the undrawn counterpart sample.

To further study the extent of clay nanolayer orientation within the die area during flat film processing, nonlinear startup experiment was also performed on the undrawn nanocomposite sample (N5), and corresponding films were prepared at various elongation rates. For this purpose, the subjected shear rate was

during rotational test to prevent further breakdown of networks in order to evaluate the effect of the extension upon the nanolayer orientation within the die zone. The height of the startup test peak can be a measure of the 3D network formed between nanoclay platelets. The results of the startup tests performed on the film samples are shown in Figure 13. As can be seen, by imposing the elongation flow fields the height of the stress overshoot decreases, indicating more network breakup and orientation of nanoclay.

Figure 13 Results of startup tests with  performed on the flat film nanocomposite samples containing 5% nanoclay compared to undrawn nanocomposite sample at T=175°C.

Figure 13

Results of startup tests with

performed on the flat film nanocomposite samples containing 5% nanoclay compared to undrawn nanocomposite sample at T=175°C.

The results of XRD experiments performed on the prepared flat film samples are presented in Figure 14 and summarized in Table 3. As can be seen, the d-spacing of the clay particles within the microstructure of the nanocomposite samples increased from 25 to 29 Å by increasing the draw ratio. These results along with the rheological observations suggest that the elongation flow field imposed on the nanocomposite samples would enhance the extent of intercalation of the polymer chains onto the gallery spaces between the nanoclayers via applied elongational forces.

Table 3

Results of XRD experiments performed on flat films.

Sampled Spacing (Å)
F-F-N5-125.1
F-F-N5-226.0
F-F-N5-428.0
F-F-N5-629.0

The results of DSC experiments performed on the flat film samples are presented in Figure 15A and B and illustrated in Table 4. As can be seen in the table, the melting temperature of nanocomposites and their corresponding films increased compared to that of the reference samples. This can be attributed to the nucleation effect of the nanoclay in the polymer matrix. Similar results for the PP/graphene system were also seen by Xu et al. (38). They showed that for their system, the graphene nanosheets (GNSs) act as nucleating agent, and the crystallization rate of the system is directly proportional to the GNS content. Moreover the presence of GNSs enhanced the effects of shear-induced nucleation as well as orientation of isotactic polypropylene (iPP) crystals.

Figure 14 XRD patterns for nanocomposite cast film samples.

Figure 14

XRD patterns for nanocomposite cast film samples.

Figure 15 DSC graphs for (A) reference film samples and (B) nanocomposite film samples.

Figure 15

DSC graphs for (A) reference film samples and (B) nanocomposite film samples.

Table 4

Results of DSC experiments performed on the flat film samples.

SampleCrystallinity (%)Tm (°C)
F-F-R-139.9125.5
F-F-R-239.2125.7
F-F-R-635.2125.1
N542.8128.1
F-F-N5-139.8126.0
F-F-N5-238.3126.4
F-F-N5-437.7126.0
F-F-N5-635.6127.1

Moreover, it can be observed that the crystallization characteristics of the unfilled reference samples have not been affected by the flow field in the low strain rates, but at higher strain rates, the amount of crystallization decreases. However, the percentage of crystallization decreased for the nanocomposite films by increasing the extent of elongation rate during flat film processing. Moreover, it can be seen that the melting temperature of the films increased by increasing the elongation rate. Derakhshandeh and Hatzikiriakos (39) studied the effects of shear and uniaxial extension on the flow-induced crystallization of PE and showed that strain and strain rate generally enhanced crystallization in both shear and elongation. They also showed that extensional flow is a much stronger stimulus for polymer crystallization compared to shear. These results are in agreement with results obtained from melting temperature. Similar results for isotactic PP have also been seen by Elmoumni and Winter (40). However, decrease in the amount of crystals in cast films by increasing the strain rate can be attributed to the decrease of the film thickness, and hence increase in the rate of cooling, leading to the crystallization.

4 Conclusion

Nanocomposites based on LLDPE/organically modified nanoclay and PE-g-MA as compatibilizer were prepared. The extent of nanoclay orientation and network breakup in simple shear flow field and elongational flow field was studied using rheological and XRD techniques. The melt linear and nonlinear rheological experiments showed that the extent of the clay network breakup and orientation increases by increasing the shear rate and time period imposed on the system. These results suggested that the extent of network breakup and orientation depends on the strain imposed on the nanocomposite. However, there exists a minimum shear rate below which the breakdown of the clay networks could hardly occur; hence, the nanoclay platelets do not orient even at long times of shearing. Flow reversal tests performed on nanocomposite samples at high shear rates revealed that the extent of nanoclay orientation increases linearly by increasing time and rate of shearing. However, at high shear rates the extent of nanoclay orientation reaches almost a constant value and does not change significantly. Moreover, linear and nonlinear rheological tests performed on the flat film samples evidenced that the elongational flow field resulted in higher degree of orientation for the clay nanolayers. Also, the XRD patterns obtained for the drawn nanocomposite films exhibited higher degree of intercalation by increasing the elongation rate.

Time sweep experiments conducted on the nanocomposite samples after being sheared at different shear rates and time periods revealed that the extent of the clay structural recovery increased by increasing the shear rate and time of shearing, resulting in less morphology stability. However, the time needed for reorientation of the nanoclay platelets was found to be longer than 3600 s. These results led to the conclusion that the nanocomposite film produced by this process has enough morphological stability to maintain the orientation before freeze line.


Corresponding author: Ali Asghar Katbab, Polymer Engineering and Color Technology Department, Amirkabir University of Technology, No. 424, Hafez Ave., Tehran, Iran, e-mail:

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Received: 2013-9-12
Accepted: 2013-11-14
Published Online: 2014-01-16
Published in Print: 2014-01-01

©2014 by Walter de Gruyter Berlin Boston

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