DTIC ADP011831: Polymer-Layered Silicate Nanocomposites: Emerging Scientific Commercial Opportunities

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TITLE: Polymer-Layered Silicate Nanocomposites: Emerging Scientific
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POLYMER-LAYERED SILICATE NANOCOMPOSITES:
EMERGING SCIENTIFIC AND COMMERCIAL OPPORTUNITIES

Emmanuel P. Giannelis

Department of Materials Science and Engineering
Cornell University, Ithaca, NY 14853, USA

ABSTRACT: Polymer nanocomposites represent a radical alternative to conventionally
(macroscopically) filled polymers. Because of their nanometer-size dispersion the
nanocomposites exhibit markedly improved properties when compared to the pure polymers
or conventional composites. These include increased modulus and strength, outstanding
barrier properties, increased solvent and heat resistance and decreased flammability. In this
paper the physical and mechanical properties of nanocomposites are reviewed and discussed
in terms of their static and dynamic properties.

1. Introduction

Polymer nanocomposites represent a radical alternative to conventionally filled polymers.
Because of their nanometer-size dispersion the nanocomposites exhibit markedly improved
properties when compared to the pure polymers or conventional composites [1], These
include increased modulus and strength, decreased gas permeability, increased solvent and
heat resistance and decreased flammability [2-11], For example, a doubling of the tensile
modulus and strength without sacrificing impact resistance is achieved for nylon-layered
silicate nanocomposites containing as little as 2 vol.% inorganics. In addition, the heat
distortion temperature of the nanocomposites increases by up to 100 °C extending the use of
the composite to higher temperature environments, such as automotive under-the-hood parts.
Furthermore, the heat release rate in the nanocomposites is reduced by up to 63 % at heat
fluxes of 50 kW/m 2 without an increase in the CO and soot produced during combustion.
Applications include low-cost alternatives of high performance composites, food packaging,
microelectronics and biotechnology.

2. Synthesis of Nanocomposites

Melt intercalation of high polymers is a powerful new approach to synthesize polymer-
layered silicate nanocomposites [12], This method is quite general and is broadly applicable
to a range of commodity polymers from essentially non-polar polystyrene, to weakly polar
polyethylene terephthalate) to strongly polar nylon. The nanocomposites are, thus,
processable using current technologies and easily scaled to manufacturing quantities.

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G.M. Chow et al. (eds.), Nanostructured Films and Coalings, 367 - 372 .

© 2000 Kluwer Academic Publishers. Printed in the Netherlands.

368

In general, two types of hybrids are possible: intercalated, in which a single,
extended polymer chain is intercalated between the silicate layers resulting in well ordered
polymer/inorganic multilayers, and dispersed or disordered, in which the silicate layers (1
nm thick) are exfoliated and dispersed in a continuous polymer matrix (Fig 1).

The silicates used belong to the general family of so-called 2:1 layered silicates.
Their crystal structure consists of layers made up of two silica tetrahedra fused to an edge-
shared octahedral sheet of either alumina or magnesia. Stacking of the layers leads to a
regular van der Waals gap between the layers called the interlayer or gallery. Isomorphic
substitution within the layers generates negative charges that are normally counterbalanced
by cations residing in the interlayer.

Pristine layered silicates usually
contain hydrated Na + or K + ions. Ion
exchange reactions with cationic surfactants
including primary, tertiary and quaternary
ammonium ions render the normally
hydrophilic silicate surface organophilic,
which makes intercalation of many
engineering polymers possible. The role of
the alkyl ammonium cations in the
organosilicates is to lower the surface
energy of the inorganic and improve the
wetting characteristics with the polymer.
Additionally, the alkyl ammonium cations
can provide functional groups that can react
with the polymer or initiate a
polymerization of monomers to improve the
strength of the interface between the
inorganic and the polymer.

Similarly to polymer blends any mixture of polymer and layered silicate does not
necessarily lead to a nanocomposite [13]. In most cases the incompatibility of the
hydrophobic polymer and the hydrophilic silicate leads to phase separation resulting in
macroscopically filled systems. In contrast, by using surface modified silicates as noted
earlier one can fine tune their surface energy and render them miscible (or compatible) with
different polymers. The approach is based on a chemical (rather than a mechanical) driving
force, which leads to nanoscopic dispersion.

< a > Intercalated MEmed

Figure 1 Schematic of composite structures obtained
using layered silicates. The rectangular bars represent
the silicate layers, (a) single polymer layers intercalated
in the silicate galleries (intercalated); (b) composites
obtained by delamination of the silicate particles and

3. Structure and Dynamics of Polymer Nanocomposites

The combination of enhanced modulus, strength and toughness is a unique feature of the
nanocomposites. In conventionally-filled polymer systems increases in modulus typically
compromise toughness. Additionally, the decrease in barrier properties of the

369

nanocomposites cannot be explained only on the high aspect ratio afforded by the exfoliation
of the inorganic nanolayers. Alternatively we suggest that the polymer chains at the interface
adopt a different structure and exhibit very different dynamics compared to the chains in the
bulk. Due to the nanodispersion a very large fraction of the polymer is at the interface (close
to 60%) even for a few percent inorganic. As a result these nanoscopically “confined”
polymer chains contribute significantly and to a large extend control the properties of the
hybrid.

Even simple notions regarding the conformations of polymers confined in two
dimensions are not yet fully understood. In three dimensions, it is well known that the
individual molecules in long chain polymers overlap significantly. In two dimensions, it has
been suggested that the different chains overlap only slightly. Therefore, the local and global
conformations of polymers in the nanocomposites are expected to be dramatically different
from those observed in the bulk, not only due to the confinement of the polymer chains but
also due to specific polymer-inorganic surface interactions not normally present in the bulk.

From our current theoretical and experimental studies on nanocomposites a new and
quite unexpected picture is emerging [14]. Despite the presence of the “confining” inorganic
layers, intercalated polymer chains exhibit substantial segmental motion even at temperatures
where the polymer is normally in the glassy state. Thus, in contrast to the bulk polymers
where chain mobility slows precipitously around T g , in the nanocomposites chain mobility
persists well below the bulk Tg. This behavior is counterintuitive as “confinement” of the
polymer chains within ~2 nm is expected to increase their solid-like character and decrease
their mobility.

We start with non-equilibrium dynamics present during polymer intercalation from
the melt. The observation that polymer chains can undergo center of mass transport in
essentially two dimensions is rather surprising because the unperturbed radius of gyration of
the polymer is roughly an order of magnitude greater than the interlayer distance between the
silicate layers. The ability of the polymer chains to undergo center of mass transport during
intercalation is further evidence that the silicate layers do not completely restrict segmental
motions, otherwise large-scale chain motion would not be possible.

Using X-ray diffraction (which monitors the angular shift and integrated intensity
of the silicate reflections) we have studied the intercalation kinetics of polystyrene into
organically modified silicates (Fig. 2) [15]. The effective diffusion coefficient, D e ff, is much
faster than the tracer diffusion coefficient of the bulk polymer or the diffusion coefficient of
the polymer in a thin film. This is because during intercalation polymer chains are moving
down a concentration gradient, whereas in the other two cases polymer motion is entropic in
origin. Furthermore, die diffusion coefficient exhibits an inverse dependence on molecular
weight. Although the diffusion coefficient of polymers near surfaces has been predicted to
have an inverse molecular weight dependence (and not scaled as 1/N 2 , N is the chain length,
characteristic of repetition) this represents the first experimental measurement of the
diffusion of high molecular weight polymer melts in two dimensions.

As the length of the surfactant molecules increases from twelve to eighteen carbon
atoms, C12 to C18, respectively, the effective diffusion coefficient increases. This is because
increasing the length of the surfactant chains effectively reduces the interaction with the
silicate surfaces and thus decreases the stickiness to the surface.

370

Figure 2. Diffusion coefficient as a function of
molecular weight of the polymer [15].

300

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A fi A

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0.0024 0.0026 0.0028 0.0030 0.0032 0.0034

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Figure 3. Spin-echo NMR of PS in bulk and
nanocomposite [16].

We now turn our attention to
equilibrium dynamics for the polymer
chains after they have been intercalated.
Local dynamics of chains “confined”
between the silicate layers were probed by
spin-echo NMR. In a spin-echo experiment
complete refocusing of the signal is
expected as long as there is no change in
resonance frequency before and after the
second pulse. Large intensity losses
therefore take place when large amplitude
dynamics commence as for example those
associated with the liquid state (i.e. above
the glass transition temperature, Tg).

Figure 3 shows the results of the
spin-echo experiment for polystyrene and
polystyrene nanocomposites [16]. To
follow the respective dynamics, polystyrene
deuterated at the backbone, d3, and the ring,
d5, was used. When d3 polystyrene is used,
the intensity of the NMR signal (multiplied
by temperature) remains constant in the
glassy regime followed by a large decrease
above Tg. This is expected as backbone
dynamics are absent below Tg and
commence at Tg. There is some mobility
for the d5 polystyrene below Tg, since the
rings can independently flip 180 ° but a
substantial drop in intensity is found only
above the Tg. In contrast, the
nanocomposites show significant amount of
mobility at least for part of the polymer
even at temperatures well below the Tg.
Additionally, there is no distinct change
from solid-like to liquid-like behavior as in
the bulk polymer.

371

all Carbon*
PS backbone

number density

Figure 4. Computer simulation of PS nanocomposites [15].

Computer simulations offer an explanation for this behavior [15,16], When
confined between the inorganic surfaces the polymer chains order into discrete subnanometer
layers (Fig. 4). This layering, clearly seen in the density profiles, imparts strong density
inhomogeneity in the direction normal to the surface. The fast dynamics arise from areas of
low-density or high free volume, which compensates for the confinement between the
inorganic layers. Neutron scattering measurements support the above structure. The polymer
chains adopt a 2D random-walk structure. Additionally, in contrast to the bulk polymer the
intercalated chains do not show a single characteristic length.

4. Conclusions

Mass transport of polymer chains into the silicate layers is faster than the corresponding self-
diffusion. Thus hybrid formation requires no additional processing time than currently
required for conventional polymer processing techniques such as extrusion.

Despite the presence of the “confining” inorganic layers, intercalated polymer
chains exhibit substantial segmental motion even at temperatures where the polymer is
normally in the glassy state. Thus, in contrast to the bulk polymers where chain mobility
slows precipitously around Tg, in the nanocomposites chain mobility persists well below the
bulk Tg. This behavior is rationalized in terms of the new structure the polymer chains adopt
at the interface. When confined between the inorganic surfaces the polymer chains order into
discrete subnanometer layers. The fast dynamics arise from areas of low-density or high free
volume, which compensates for the confinement between the inorganic layers. Neutron
scattering measurements support the above structure.

Acknowledgements

This work was supported in part by the Cornell Center for Materials Research, AFOSR and
ONR. I would like to thank my coworkers and collaborators S.D. Burnside, H. Chen, J.D.

372

Gilman, J. Genzer, T. Kashiwagi, E. Manias, P.B. Messersmith, E.J. Kramer, R.
Krishnamoorti, R.A. Vaia and D.B. Zax.

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