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Exfoliation Kinetics of a Phyllosilicate-Organoclay Structure in a Polymer Matrix during Melt-Extrusion Process.

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Dev. Chem. Eng. Mineral Process. I2(1/2), pp. 149- 158, 2004.
Exfoliation Kinetics of a PhyllosilicateOrganoclay Structure in a Polymer Matrix
during Melt-Extrusion Process
C.Y. Lew, W.R. Murphy and G.M. McNally
Medical Polymers Research Institute, Ashby Building,
Stranmillis Road, Belfast BT9 5AH, Northern Ireland, UK
Nanostructured polymer-composite materials were prepared by melt-exfoliating the
smectite structure of a synthetic tetrasilisicfluoromica in a nylon-12 matrix, by means
of a melt-extrusion process using a single-screw extruder. The structure and
morphology of these materials were evaluated using the X-ray diffraction (XRLl) and
transmission electron microscopy (TEM) techniques. Uniform dispersion of exfoliated
silicate layers were observed throughout the polymer matrix. Rheological studies
using a dual capillary rheometer showed that the melt viscosity of the nanocomposites
was dependent on the degree of clay exfoliation, and in addition the orientation of the
layered-silicates in the polymer matrix. The exfoliation and orientation of layeredsilicates were in turn dependent on the processing conditions adopted. Typically a
lower processing shear would improve the intercalation of the layered-silicates while
increasing shear enhanced their exfoliation. Improved exfoliation was shown to
decrease the polymer melt viscosity, which may be attributed to the plasticising effect
of the clay platelets.
The rapid expansion of indusnial and economic activities results in a continuous
demand for new and low-cost materials to meet increasingly rigorous conditions.
Traditionally polymers are commonly admixed with a variety of natural or synthetic
fillers to improve their performance. Polymers filled with small amounts of layeredsilicate dispersed at the nanoscale level are currently the most promising materials.
They are characterised by a well balanced combination of improved properties, thus
extending the technological limits of conventional macro- or microscopically
dispersed filled-compositematerials.
C.Y. Lew, El?.Murphy and G.M. McNa&
Polymer-phyllosilicate nanocomposites consisting of highly anisotropic nanoscale clay platelets dispersed in a polymeric matrix have been the subject of intensive
research in recent years. This activity is motivated by the realisation that these
materials show a combination of low weight, ease of recycling, improved mechanical,
chemical, permeability, flame-retardancy, and optical properties, with a low level of
clay incorporation, typically between 1 to 5 wt% [l-51. Development of polymerphyllosilicate nanocomposites is one of the recent advances in polymer technology.
This material offers attractive diversification and application of conventional
polymeric-composite materials. Since the possibility of direct melt-intercalation was
first demonstrated by Giannelis [7], the melt-intercalation technique has effectively
become the mainstream method for the preparation of polymer-phyllosilicate
The extent of property enhancement depends on various factors such as the aspect
ratio of the silicate layers, the degree of silicates dispersion and orientation in the
polymer matrix, and the clay-matrix interface homogeneity [6-71. Although for most
applications an exfoliated system is preferable [8-91, intercalated systems have shown
greater property improvement in niche applications. For example, Gilman et al. have
shown that an intercalated structures exhibited better flame-retardancy properties over
the exfoliated equivalent [ 10-111. Conventionally, the structure of the nanocomposite
systems may be controlled by altering the length, density, and type of surfactants used
in the modification of the layered-silicates [ 121.
Previous studies have reported that the arrangement and orientation of the silicatelayers in the polymer matrices have a significant effect on the structure and properties
of the nanocomposites [13]. For instance, the melt viscosity of polymer
nanocomposites has been reported to vary appreciably from a disordered-exfoliated
system to an ordered-exfoliated configuration. In addition, the mechanical and
thermal properties of the nanocomposites are also dependent on the structure and
orientation of the silicate-layers in the polymer matrices.
T h s work investigated the intercalation and exfoliation mechanisms of the
layered-silicates during a melt-extrusion process. Results had showed that the
precedence of either intercalation or exfoliation of the clay structure is dependent on
the processing conditions adopted. The dry-mixture homogeneity of the polymer/clay
particles prior to extrusion also appeared to have a significant effect on the clay
exfoliation during the extrusion process. The dry-mixture uniformity was related to
the polymer particle size and its particle size distribution. These findings give a better
understanding of the exfoliation mechanisms associated with the delamination of
layered-silicates in a thermodynamical shear environment. This enables the
optimisation of clay exfoliation in a polymer nanocomposite system using
conventional extrusion equipment, and hence expands the commercial opportunities
for this technology.
Experimental Details
(a) Materials. The nylon-12 used in this work was a commercially available
polyamide from EMS-Chemie with an average M, of 13,100 (from gel permeation
chromatography). The organoclay used was a synthetic tetrasilisic fluoromica
Exfoliation Kinetics of n Phyllosilicate-Organoclay in a Polymer Matrix
supplied by Uni-COOP Japan with an average particle size of 1 to 5pm, cation
exchange capacity (CEC) of 2 meqtg, and a peak interlayer spacing of 3.4 nm
(calculated from XRD). The fluoromica was produced by heating of talc and Na2SiF6
at lugh temperature for several hours in an electric finace, followed by subsequent
ion-exchange with quaternary alkylammonium cations. Three different types of
polymer-organoclay mixtures with 5 wt% of organoclay were prepared and their
formulation is shown in Table 1, where P, MC, W, and o represent the polymer
pellets, micropellets, powder, and organoclay respectively.
Nylon 12 Pellets
Nylon 12 Micropellets
Nylon 12 Powder
+ Organoclay
+ Organoclay
+ Organoclay
(b) Preparation of Samples. The polymer powder and micropellets were
cryogenically ground from the original nylon-12 pellets, giving an average particle
size of about 200 prn (ASTM E l l ) and 605 prn (optical microscopy) respectively.
The mixtures were dried at 80°C for I6 hours prior to compounding.
Compounding of the various mixtures was carried out using a Killion-KN150
single-screw extruder fitted with a 38 mm diameter, 25:l (L:D) barrier-design screw.
The extrusion temperature profile was 185°C at the feed to 225°C at the die. The three
blends in Table 1 were compounded using a screw speed of 37.5 rpm. An additional
four nylon-12 nanocomposite blends each containing 5 wt% of organoclay were
compounded using the extrusion conditions shown in Table 2.
Table 2. Extrusion processing conditions.
The notations Lo, lo, Mo, and Ho represent blends of nylon-12 and organoclay
compounded at low, intermediate, medium, and high shear rates in the extruder. The
unfilled base nylon-12 pellets were passed through the extruder for comparison
purposes. Standard ASTM test samples were produced by injection moulding using
an Arburg 320s Allrounder. Barrel to die temperatures were ramped from 180 to
235°C at 1100 bar injection pressure. The mould temperature was maintained at 43°C.
C. Y. Lew,
W.R.Murphy and G.M. McNally
(c) Analysis. Wide angle X-ray diffraction (WAXD) patterns of the organoclay in
polymer matrix were analysed using a Siemens D5000 diffiactometer (40 kV,40 mA)
with CuKa radiation ( h = 1.5406A), scanned at 0.3"/min. The morphology of the
nanocomposites was studied using a Philips CMlOO TEM. Ultrathin (50-60 nm thick)
sections were prepared using an Ultracut-E Reichert-Jung ultramicrotome. Sectioning
was performed perpendicular to the flow direction of the injection-moulded samples.
Rheological behaviou of the nanocomposites was studied at 200°C using a Rosand
RH7 dual capillary rheometer.
Results and Discussion
(0 The effect of polymer-organoclay mixtureformulation on the structure-proper@
of nanocomposites
The WAXD diffractograms in Figure 1 show typical spectra for the organoclay and
the nylon-l2/organoclay nanocomposites. The apparent shift in the characteristic
peaks of the organoclay would indicate that intercalation of the organoclay periodic
framework had occurred in the nanocomposites as a result of melt compounding. The
disappearance of the charateristic peaks between 4.5 and 5.5" would illustrate that
delamination of the layered-silicates had also occurred.
Theta (degre.1
Figure 1. WAXD pattern of pristine organoclay and polymer-organoclay
The diffractograms in Figure 2 show the effects of the original particle size of the
polymer, prior to compounding, on the final structures of the compounded
nanocomposites. The spectra shows a progressive left-shift of the characteristic peaks
with progressive decrease in polymer particle size, i.e. from pellet to micropellet to
Exfoliation Kinetics of a Phyllosilicate-Organoclayin a Polymer Matrix
powder. This shift of the characteristic peaks would indicate an increase in interlayer
spacing of the layered-silicate in the nanocomposites matrices, with decrease in
original polymer particle size in the original dry blend prior to compounding. This
may be due to the faster heat transfer rate for polymer of smaller particle size, which
requires a shorter time to reach the melt stage during an extrusion, thus increasing the
time for diffusion of the polymer chain into the layered-silicate galleries. In addition,
smaller particle size would give a more hamogeneous polymerklay dry-mix prior to
compounding, which could lead to improved dispersion of the clay.~particles in the
polymer matrix during the extrusion process.
2 Theta
Figure 2. Eflects of original polymer particle size on the intercalation and exfoliation
level of organoclay.
The rheological behaviour of the nanocomposites was investigated using a dual
capillary rheometer and the results in Figure 5(a) show the changes in viscosities of
the nanocomposites over a shear rate range of 30 s-' to 1000 s-' at 200°C. The results
show that the viscosities of these nanocomposites is significantly lower than that of
the nylon-12, particularly for the mixture blended from polymer powder ( Wo).The
nanocomposite with the highest viscosity was the one compounded using polymer
pellets. These results would tend to indicate that the particle size and particle size
distribution of the polymer resin used in the compounding process had a significant
effect on the rheological properties of the nanocomposites, which in turn were found
to be related to the WAXD spectra trends observed in Figure 2. For instance, a
progressive decrease in the polymer particle size of the dry polymerklay blend
resulted in a simultaneous increase in the melt viscosity of the nanocomposites (see
Figure 5a), and the degree of intercalation or delamination of the layered-silicate as
shown by the WAXD spectra in Figure 2.
C.Y. Lew, W.R. Murphy and G.M. McNallj
Figure 3. WAXD patterns of nanocomposites produced by different shear level
in an extrusion environment.
Microstructure examination of these nanocomposites using TEM is shown in
Figure 4(a) to (c). The observed change in size of the layered-silicates agreed well
with the WAXD results, which show a change in the degree of intercalation and
delamination with particle size and particle size distribution of the polymerklay dry
blend. For example, the nanocomposites compounded from the blend using polymer
pellets show some thicker layered-silicate stacks, randomly dispersed in the polymer
matrices (see Figure 4a), whereas very fine and well-orientated silicate layers were
observed in the blend produced from polymer powder (see Figure 4c).
This would indicate that the lower viscosity recorded for the Wo was attributed to
a greater degree of delamination of layered-silicates in the nanocomposites, thus
generating thinner layered-silicates that could cause a plasticising effect in the
polymer melt. This would result in increased slippage of adjacent polymer chains
during extrusion. The TEM microstructures in Figure 4 also show significant
differences in the orientation of the clay platelets in the polymer matrices of the
nanocomposites. These changes in alignment of the silicate layers could be attributed
to the change in the polymer melt rheology during the extrusion process [13]. The
orientation of these silicate layers would also influence changes in melt viscosity of
these nanocomposites as previously explained by Lew et al. [13]. However, other
factors such as polymer degradation and chain-scission effect by the clay platelets
may also have an effect on viscosity [14].
Exfoliation Kinetics of a Phyllosilicate-Organoclay in a Polymer Matrix
Figure 4. Transmission electron microscopy photomicrographs of organoclay silicate
layers in the nylon-I2 polymer matrix.
(ii) The effects of thermodynamic shear conditions on the structure and properties
of the nanocomposites
The effect of extrusion shear level on the structure and properties of the
nanocomposites were investigated by subjecting the polymer-organoclay mixture to
different levels of shear rate during the compounding process. The WAXD
diffiactograms in Figure 3 shows the structures of various nanocomposite samples
processed under different levels of shear rate in an extruder.
Results from the diffiactograms show that an increase in the extruder screw speed
(i.e. shear rate) would result in an improvement in the degree of delamination of the
layered-silicates structure, with concurrent decrease in the level of intercalation.
Typically, the delamination of a layered-silicate structure would be a result of
expansion of the clay galleries by polymer chains during the intercalation process.
When this interlayer swelling increases beyond the binding force of the silicate
interlayers, the periodic-layered structure of the clay would collapse with the
assistance of shear stress. However, in the absence of either sufficient shear intensity
or residence time during an extrusion process, there may be a divergence from the
conventional delamination process.
For example, the structures of layered-silicate of Ho in Figure 3 appeared to be
delaminated without further and significant improvement in its interlayer spacing.
This would suggest that the thicker silicate tactoid was broken into some thinner
stacks under the experience of very high shear stress, imparted from the extruder
screw, rather than due to the finite intercalation of its galleries by polymer matrices.
On the contrary, a marked expansion in interlayer spacing of the layered-silicate
galleries (however without further or significant improvement in the degree of
delamination) was seen for the Lo sample, which was processed under a very long
extrusion residence time at a low shear rate. This would suggest that the enhanced
intercalation of the clay galleries was due to the penetration, and hence continuous
diffusion, of the polymer chains into the silicate interlayers. However, it may appear
that when an equilibrium state for the inter-polymer/galleriesdiffusion was achieved,
further delamination of the layered-silicate would not be realised due to insufficient
force required to shear-break the layered-silicate structure.
C.Y. Lew, W.R.Murphy and G.M. McNaIly
Figure 5. Rheological behaviour of nanocomposites preparedji-om polymer of
diferent original particle size.
The rheological studies (see Figure 5b) correlate well with the WAXD results in
Figure 3, which show that the delamination level of the layered-silicates was
improved concomitantly with the compounding screw speed (shear intensity), while
the interlayer galleries expansion improved progressively with improvement in the
extrusion residence time. For instance, a lower melt viscosity measured for
nanocomposite obtained with highest extrusion shear rate recorded weaker peaks
intensities in the WAXD analysis (see Figure 3), i.e. greater degree of delamination of
the layered-silicates. On the other hand, the melt viscosity.of the nanocomposites was
lugher with an increase in the expansion of interlayer clay galleries.
Under a higher thermodynamic shear condition, the polymer matrix network
would experience an increasing ease of flow and improved chain alignment in the
extruder, and hence resulted in a decrease in melt viscosity. For a conventional
delamination process of layered-silicate to occur, it must be preceded by satisfactory
diffision of the polymer chain into the clay galleries (i.e. intercalation). Theoretically,
those polymer chain with greater degree of flow alignment and lower viscosity
(corresponding to a smaller molecular size in a conventional diffusion process) should
diffuse more easily into the clay galleries, and hence also improve its interlayer
spacing. However a more crucial mechanism, before the polymer/galleries interface
diffision process could occur, is the initial polymer/galleries interface penetration.
Due to the inherent radius of gyration for polymer (approximately 10 nm for most
polymeric materials), which is much greater than the interlayer gap of the clay
galleries, further diffusion of the polymer chain may not occur unless the edge gap of
the layered-silicate was expanded. However at a very low polymer melt viscosity (e.g.
Ho in Figure 5b) and a very high extrusion shear rate, these polymer chains would
tend to slip over the edge surface of the layered-silicate, and thus reduce the
acluevement of the intended initial penetration stage necessary for an interlayer gap
Exfiliation Kinetics of a Phyllosilicate-Organoclayin a Polymer Matrix
opening. This would explain the WAXD trend in Figure 3, where improvement in
delamination (without further improvement in the interlayer spacing) was observed
with increasing exmsion shear rate. Therefore, the observed delamination appeared
to be as a result of high shear stress imparted by the extruder screw.
On the other hand, when processed under a longer residence time and lower shear
rate, the polymer would experience a much higher viscosity. However, due to a low
shear rate, slippage of polymer chain over the edge surface of the layered-silicate
would be reduced considerably and hence allow for the intended polymer/galleries
interface penetration process to be initiated. This primary edge-gap opening would
enable or ease the diffusion of the polymer chain into the layered-silicate framework.
However due to inadequate shear, when an equilibrium diffusion state was reached, a
delamination process would not occur or occur only slightly. This explains the
WAXD trend in Figure 3, where improvement in intercalation was not followed by
The rheological analysis in Figure 5b shows the variation in the melt viscosity of
the nanocomposites, processed using different thermodynamic shear conditions.
Results show that lowest melt viscosity was recorded for Ho (processed at greatest
shear) and highest viscosity was measured for Lo (processed at longest extrusion
residence time). The higher viscosity observed for Lo was due to the greater degree of
delamination. The lower viscosity for Ho was attributed to greater degree of interlayer
expansion (see Figure 3). An expansion of the clay galleries to finite spacing without
delaminating them could result in a loss of the intended nanolayer reinforcement
mechanism, and thus the clay would behave similar to conventional filler particles,
increasing the viscosity of the polymer system.
This study describes the intercalation and delamination mechanisms of a
phyllosilicate organoclay in polymer matrices under different thermodynamic shear
(extrusion) conditions. All the nanocomposites studied were successfully
delaminated, as indicated in the TEM photomicrographs. In addition, significant
structure-property relationshps were found between the relative intercalation/
delamination levels of the clay and the dynamic melt viscosity of the polymer, based
on the rheological analysis, TEM, and WAXD within the scanning range of 2" to 11'.
A substantial decrease in melt viscosity recorded for all the nanocomposites further
indicates some great extent of delamination of the layered-silicate into thinner
periodic layers, which appeared to cause a plasticising (slippage) effect on the
polymer chains.
This work has shown the significance of polymer particle sizes in the
polymerklay mixture for optimisation of the delamination of layered-silicate
framework during the extrusion process, especially in the absence of adequate shear
or residence time. A smaller polymer particle size in the pre-compounded
polymer/clay blend appeared to give a greater degree of delamination. In contrast to
the conventional delamination mechanism, usually preceded with an intercalation
process, this typical convention also appeared to diverge depending upon the
extrusion conditions adopted. For example, a very h g h level of extrusion shear rate
C. Y. Lew, W.R. Murphy and G.M. McNally
which resulted in shorter residence time and lower polymer melt viscosity appeared to
decrease the intercalation of the clay galleries due to various factors explained
previously. As a result of understanding the various factors affecting the delamination
mechanism of the layered-silicate in the polymer matrix during the extrusion process
including the orientation of the clay platelets (not discussed here), it would be
possible to fine-tune the extrusion parameters in order to optimise the delamination of
layered-silicates in the nanocomposites.
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polymer, process, matrix, structure, melt, phyllosilicate, kinetics, organoclay, extrusion, exfoliation
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