NANOCOMPOSITES,
POLYMER–CLAY

Introduction

Polymer matrix nanocomposites are a fairly new class of engineered materials
which offer for a broad range of properties, an interesting and even radical al-
ternative to more conventional filled polymers, yet at much lower filler loadings.
They can be defined as polymer matrix systems in which the dispersed inorganic
reinforcing phase has at least one of its dimensions in the nanometer range, which
is quite close to the scale of elementary phenomena at the molecular level. The
resulting unique combination of large interfacial area and small interparticle dis-
tance strongly influences nanocomposite behavior.

Current status of research and industrial development of polymer nanocom-

posites clearly outlines the prominent position of clay nanocomposites and the
present review is mainly devoted to the latter materials.

From a general point of view, filler aspect ratio is a pertinent parameter to

distinguish between various types of nanocomposites. Table 1 summarizes typ-
ical dimensions of particles under concern. Spherical silica particles are an ex-
ample of isotropic nanoparticles which either provide increased composite stiff-
ness while retaining matrix transparency, or exhibit novel optical properties by
forming colloidal crystals. Although it is obviously critical to the optical behav-
ior, it is generally observed that optimum mechanical properties are not achieved
in conjunction with the best state of dispersion (1). When only two dimensions
are in the nanometer range, fiber-like structures, such as whiskers or carbon
nanotubes, with aspect ratios ranging between 50 and 1000 are dealt with. For
instance, cellulose whiskers extracted from tunicate shells have been shown to
dramatically improve composite stiffness in the case of poly(vinyl chloride) or

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Encyclopedia of Polymer Science and Technology. Copyright John Wiley & Sons, Inc. All rights reserved.

Vol. 3

NANOCOMPOSITES, POLYMER–CLAY

337

Table 1. Typical Nanofiller Dimensions

Material

Shape

Typical dimensions

Silica particles

Spheres

Diameter: 30–150 nm

Cellulose whiskers Rigid rods

Diameter: 15 nm; length: 1

µm

Carbon nanotubes Flexible tubes (multiwall) Diameter: 30 nm; length: 10–50

µm

Layered silicates

Flexible discs

Diameter: 50–500 nm; thickness: 1 nm

poly(styrene-co-butyl acrylate) matrix in the rubbery state. In the latter material
a 3 order of magnitude modulus improvement is achieved at only 6 mass% whisker
content. Percolation of a rigid whisker network is evoked to account for such a
property increase (2). Reaching a conductivity percolation threshold at very low
loadings, owing to the nanoscale dispersion of carbon nanotubes with large aspect
ratio, has also been of prime importance for designing conductive polymer blends
with electrostatic paintability, while limiting resin embrittlement (3). Preliminary
experiments with carbon nanotube–polymer composites also underline the poten-
tial of this nanofiller for entering the field of composite materials for structural
applications as well (4). Opportunities for functional and/or structural applica-
tions of carbon nanotubes will mainly depend on the capacity to reach large-scale
production at moderate cost and to monitor composite elaboration while retaining
nanotube integrity.

Among the variety of composites that display unique structure and behavior

at the nanometer level, as compared to classical micrometer scale particulate filled
materials, the use of layered silicates as a reinforcing phase is by far the most suc-
cessful way of designing polymer nanocomposites with a broad range of markedly
modified properties. A report from Toyota Central Research Laboratories of the
development of a polyamide-6 (PA6)-based clay nanocomposite with remarkable
thermomechanical behavior, at low clay loadings relative to conventional filler
additives (below 5 mass% instead of 20–30 mass%), triggered extensive research
efforts worldwide (5). In fact the benefits were shown not only for strength and
stiffness, but also for thermal stability and barrier properties (6). Accordingly, the
following presentation focuses on these materials, starting with a brief description
of the layered silicates commonly used, followed by nanocomposite structural char-
acterization and elaboration routes. Various properties of interest are reviewed
together with the currently emerging structure–property relationship schemes.

Polymer-Layered Silicate Nanocomposites

Layered silicates, the more widely used in polymer nanocomposites, belong to the
same structural group, the 2:1 phyllosilicates, and more specifically to the smectite
group. They comprise natural clay minerals such as montmorillonite, hectorite,
and saponite and also synthetic layered minerals, fluorohectorite, laponite, or
magadiite. An idealized structure for montmorillonite is presented in Figure 1.
Elementary clay platelets consist of a 1-nm-thick layer made of two tetrahedral
sheets of silica fused to an edge-shared octahedral sheet of alumina or magnesia.
Isomorphic cation substitution results in an excess of negative charges within

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NANOCOMPOSITES, POLYMER–CLAY

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Fig. 1.

Structure of montmorillonite.

Al or Mg;

OH;

O;

exchangeable cations.

the layer. Stacking of the layers leads to a regular van der Waals gap called
gallery or interlayer. Cations located in the galleries (eg Ca

2

+

, Na

+

) counterbal-

ance the excess layer charges. They are usually hydrated. This negative surface
charge is quantified as the cation-exchange capacity (CEC), usually in the range
from 80 to 150 meq/100 g for smectites. The interlayer space can be penetrated
by organic cations or polar organic liquids as well. Exchange reactions with or-
ganic cations (aminoacids, alkylammonium ions, etc) enable to render the silicate
surface organophilic. A key parameter of the stacking is the basal spacing or
d-spacing, which ranges from 0.96 nm (ie, the layer thickness) in the fully col-
lapsed state to about 2 nm, depending on the nature of the interlayer cation and
the amount of adsorbed water (7). Individual lamellae have a high aspect ratio,
with a diameter typically in the range 50–500 nm. Primary crystals (also called
tactoids) consist of 8–10 lamellae, with usually disordered stacking. Their aggre-
gation leads to a turbostratic structure. This organization is reflected in the x-ray
diffraction (xrd) pattern where diffuse hk bands rather than sharp hkl reflections
are observed. The basal reflection (001) is of more interest since it is used to de-
rive the d-spacing. It will be dependent on the amount and nature of intercalated
molecules in the galleries.

Characterization of Nanocomposite Microstructure.

Mixing clay

with a polymer does not necessarily lead to a nanocomposite. Elaboration strate-
gies are aimed at monitoring dispersion of the inorganic compound at the nanome-
ter level, that is down to the elementary clay platelet. Figure 2 provides a
schematic illustration of the various microstructures readily achievable, namely
a conventional filled polymer with clay particles in the micrometer range, an

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NANOCOMPOSITES, POLYMER–CLAY

339

(a)

(b)

(c)

Fig. 2.

Schematic representation of clay platelet dispersion as (a) primary crystals, (b)

intercalated, or (c) exfoliated structures.

intercalated nanocomposite in which extended polymer chains are intercalated
in the gap between silicate layers while the stacking order is retained (note that
in this case the host gallery height is much smaller than the radius of gyration of
the polymer chain), and a delaminated (or exfoliated) nanocomposite where clay
layers are individually dispersed in the host polymer matrix.

The two basic tools used to elucidate nanocomposite morphology are x-ray

diffraction (xrd) and transmission electron microscopy (tem). They provide com-
plementary information on clay dispersion in the host matrix.

The basal plane reflection (00l) diffraction peak yields a direct evaluation of

the d-spacing between the clay lamellae, as long as layer registry is preserved to
some extent. It allows therefore to follow the structural evolution from the pristine
clay stacking to any intercalated state. With changes in d values in the range 1–4
nm, such data are accessible at low diffraction angle, ie, for 1

< 2 < 9

◦

. Follow up

of the increase in peak width at half maximum also reflects the increase in the
degree of disorder in layer stacking during the intercalation process [see eg, (8)].

Figure 3 illustrates the change in d-spacing of a stearylammonium-modified

montmorillonite (C

18

Mt) blended with a maleic anhydride-modified polypropylene

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NANOCOMPOSITES, POLYMER–CLAY

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0.5

2

3.5

5

6.5

8

9.5

2

Θ -Cu, deg

0.5

2

3.5

5

6.5

8

9.5

2

Θ -Cu, deg

Intensity, a.u.

6.3 nm

5.7 nm

3.4 nm

2.1 nm

(a)

(b)

(c)

(d)

I

6.0 nm

5.0 nm

3.5 nm

2.1 nm

(a)

(b)

(c)

(d)

II

Intensity, a.u.

Fig. 3.

Evolution of the d-spacing showing montmorillonite intercalation in PP-MA

oligomer (I), and no subsequent change upon further blending with PP (II). Reproduced
from Ref. 9.

oligomer (PP-MA). As the amount of the latter component is increased, the inter-
layer distance increases drastically (spectra Ia–d), indicating the efficiency of the
intercalation process. Spectrum IIa reveals the lack of intercalation of C

18

Mt by

pure PP. Spectra IIb–d do not show any change in the state of intercalation as-
sessed in Spectra I when the PP-MA intercalated organoclay is further dipersed
in a PP matrix. It is concluded that PP penetration in the galleries is not favored
in this particular case (9).

Whenever the xrd pattern becomes silent, meaning that the interlayer spac-

ing is no longer accessible by conventional wide-angle x-ray diffraction, the need
still remains to have access to a mesoscale description of the spatial distribution
of the clay platelets. Small-angle x-ray scattering may provide some answers re-
garding average interparticle distance, and also about platelet orientation with
respect to process geometry (10).

Transmission electron microscopy has the great advantage to give a di-

rect view of the microstructure. Typical micrographs reveal alternating dark and
bright bands refering to the silicate layers and polymer matrix respectively. The
illustrations of Figure 4 show exfoliated and intercalated-exfoliated structures in
the case of 1 and 5 mass% dimethylditallow ammonium-modified montmorillonite

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NANOCOMPOSITES, POLYMER–CLAY

341

500 nm

200 nm

Fig. 4.

Transmission electron microscopy of water-aided melt-dispersed organoclay in

PA6. Courtesy of Dr M. van Es, DSM Reseach, Geleen.

100 nm

Fig. 5.

Transmission electron microscopy of a melt-intercalated organoclay tactoid in a

PP matrix. Courtesy of Dr M. Bacia, UST Lille.

dispersed by melt extrusion in a PA6 matrix. Additionally tem enables imaging
of any mesoscale long-range ordering of the clay. It also reveals platelet flexibil-
ity promoted by the high aspect ratio and nanometer thickness as illustrated in
Figure 5 for an intercalated organoclay tactoid in a PP matrix. An accurate de-
scription of polymer clay nanocomposites is therefore available from a combination
of electron microscopy and xrd techniques.

Elaboration of Polymer–Organoclay Nanocomposites

Various methods have been developed in order to prepare polymer-layered silicate
nanocomposites. These include in situ polymerization, polymer intercalation in

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NANOCOMPOSITES, POLYMER–CLAY

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solution, emulsion polymerization in the presence of layered silicates, and melt
intercalation. A few examples are given below to illustrate the different strategies.
Review articles are available for more detailed information [see eg (11–13), and
references therein].

In situ Polymerization.

Early works in the 1960s (14) demonstrated the

feasability of intercalation polymerization of methyl methacrylate (MMA) after
insertion–adsorption of the polar monomer between the lamellae of a sodium
montmorillonite. Polymerization was initiated by free-radical catalysts or with
γ irradiation. Renewed interest for the method in the last decade came from the
work on PA6–clay hybrids at Toyota Central Research Laboratories (5), follow-
ing a scheme adopted earlier by Unitika Co. (15). Sodium montmorillonite was
first cation exchanged with

ω-amino acids [H

3

N

+

(CH

2

)

n

− 1

COOH]. X-ray diffrac-

tion data show that the basal spacing is highly sensitive to the length of the
alkyl chain. For n

> 11, the ω-amino acid chain lies slanted to the layer and pro-

vides optimum swelling behavior by

ε-caprolactam. 12-Amino-l-lauric acid (n =

12) modified montmorillonite (12-Mt) was used for performing the in situ ring-
opening polymerization of

ε-caprolactam. The carbonyl end groups of 12-Mt ini-

tiate polymerization and an increasing amount of polyamide-6 is bonded to the
clay as the 12-Mt clay content is increased as clearly established by nmr studies
(16). The resulting hybrid is either exfoliated at low 12-Mt content or increas-
ingly intercalated beyond 10 wt% clay. In an alternative approach, the same team
successfully developed a so-called one-pot synthesis of PA6–clay hybrid without
preliminary cation exchange of the montmorillonite. Ring-opening polymerization
of

ε-caprolactam with 6-amino-l-caproic acid as an accelerator was performed in

a water dispersion of montmorillonite in the presence of an acid. Phosphoric acid
seems to be the best candidate to achieve true exfoliation in this particular pro-
cess (17). 12-Amino-l-lauric acid was also successfully used both as a fluorinated
silicate modifier and as a monomer in order to prepare a polyamide-12-based
nanocomposite with exfoliated–intercalated structure (18). Some attempts to pro-
duce polystyrene (PS)-based nanocomposites through the in situ polymerization
route have been reported (19). Intercalated structures are obtained. Intercala-
tive polymerization of

ε-caprolactone has also been achieved in the presence of

α-protonated amino acid-exchanged montmorillonite. Upon heating, the organic
acid groups initiate ring-opening polymerization of the monomer and the resulting
polymer is ionically bound to the silicate platelets with a good level of delamina-
tion as revealed by xrd (20). In situ polymerization has also been extended to
polyolefins, polyesters, or polycarbonate in recent years.

The production of thermoset-based nanocomposites by the same method has

been investigated by many authors (21–23). In the case of epoxy–clay nanocom-
posites, the organophilic clay is first swollen in the mixture of epoxy prepolymer
and curing agent. The gel state and final structure are strongly dependent on the
nature of the onium ion, cross-linking amine, and curing conditions. In particu-
lar, larger chain length of the alkylammonium and ion-exchange with protonated
primary amines should be preferred. Polyurethane networks are equally good can-
didates for clay nanolayer reinforcement (22). The approach consists in solvating
polyol precursors in montmorillonite exchanged with long-chain onium ions and
further adding the diisocyanate curing agent. Intercalated tactoids are obtained
in the final cured nanocomposite.

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NANOCOMPOSITES, POLYMER–CLAY

343

Intercalation in Solution.

In the case of water-soluble polymers, it is pos-

sible to prepare solvent-cast nanocomposites by using water as a cosolvent. Pris-
tine montmorillonite can be easily dispersed in water, owing to its hydrophilic char-
acter, and blending for instance with polymers such as poly(ethylene oxide) (PEO)
or poly(vinyl alcohol) (PVOH) is thus achievable (see eg (24–26)). Intercalated–
exfoliated structures are observed for PVOH-based nanocomposites. On the con-
trary, in the presence of PEO, intercalated silicate layers are organized as large
clay tactoids; this indicates that reaggregation of the initial silicate water sus-
pension occurred during the film casting process. A similar elaboration strategy
has also been developed, starting from an organoclay. In the examples of poly(

ε-

caprolactone) and poly(l-lactide), chloroform was used as a cosolvent. In both situ-
ations no evidence of intercalation could be found but the clay tactoids displayed a
remarkable geometric arrangement, with their surfaces lying parallel to the cast
film surface (27). These few illustrations are indicative of the high sensitivity of
the final materials structure to the nature of the host matrix and to the interplay
of polymer–polymer and polymer–clay interactions.

The above technique can be advantageously adapted to a situation where

the polymer of interest is not soluble in any solvent, as is the case for polyimides.
The polyimide precursor, ie, a poly(amic acid) solution in dimethylacetamide, is
prepared and blended to a dispersion of an organomodified montmorillonite in the
same solvent. Dimethylacetamide is then gradually removed, and upon heating
the poly(amic acid) film at 300

◦

C under nitrogen atmosphere the polyimide–clay

hybrid is obtained. When an ammonium salt of dodecylamine is used in the ion-
exchange process, true exfoliation is achieved and furthermore the clay platelets
align parallel to the film surface (28).

Emulsion Polymerization.

Considering the high hydrophilic character

of sodium montmorillonite, it was anticipated that polymerization in an aque-
ous medium might provide an alternative route for polymer–clay nanocomposite
preparation. The first report dealing with such an approach concerned the emul-
sion polymerization of MMA dispersed in a water phase in the presence of Na

+

-

montmorillonite. The structural analysis confirms intercalation of the PMMA in
the clay galleries. A PS–clay nanocomposite has been elaborated according to the
same procedure, with the same resulting intercalated morphology. Intercalation
by the emulsion technique was also achieved for an epoxy system, again without
requiring any ion-exchange treatment (29).

Melt Intercalation.

Although emulsion polymerization or in situ polymer-

ization may be considered as viable routes for industrial-scale production, as ex-
emplified for PA6 with the latter process, melt compounding remains the most
obvious route for cost-effective development of nanoreinforced polymers. The first
industrial applications refer to the PA6–clay hybrid in situ polymerization process
patented by Toyota (30), but growing interest in achieving clay nanodispersion by
melt compounding is observed worldwide (31). Patent literature is clearly indica-
tive of the current trend with many references to polyamides and other polar
polymers such as polyesters or polyimides [see eg (32)]. Emphasis is put on the
nature of the surfactant and processing conditions. Nonpolar polymers such as
polyolefins are also highly attractive candidates with research efforts primarily
driven by the automotive market. Polypropylene–clay hybrids have been prepared
by melt mixing an organoclay, maleic anhydride-modified PP oligomers (PP-MA),

344

NANOCOMPOSITES, POLYMER–CLAY

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and PP. The polar PP-MA intercalates in the clay galleries and the quality of clay
dispersion in the hybrids are clearly affected by the degree of miscibility of the
polar oligomers in PP (9,33).

Direct melt intercalation of PEO in pristine montmorillonite by static anneal-

ing of cold-pressed powder mixture slightly above the melting temperature of PEO
has been reported (34). The resulting d-spacing of the order of 1.8 nm indicates
that the PEO chains are constrained in a 0.8-nm interlayer space. Differential
scanning calorimetry studies reveal that the polymer is deprived of any thermo-
dynamic transition. Neither the heat capacity jump characteristic of the glass
transition nor the melting endotherm is observed. Chain dynamics appear quite
peculiar in these systems. Thermally stimulated current results point at the essen-
tially noncooperative nature of the motions of the confined chains. Intercalation
kinetics have also been followed by xrd monitoring of the basal spacing reflection
in model systems consisting of monodisperse PS and organically modified fluo-
rohectorite. The most striking result is that the mobility of the polymer chains
in this confined environment seems larger than that in the bulk melt (34,35).
Spin-echo nmr experiments using deuterated PS suggest the coexistence of multi-
ple environments, from solid-like to liquid-like, for intercalated chains. Molecular
dynamics simulations relate this complex dynamic behavior to strong density in-
homogeneity normal to the surface [(36), and Refs. 45 to 49 therein].

An extensive investigation of nanocomposite preparation with the aid of a

swelling agent that is known to intercalate the clay provides some additional
interesting experimental ground on the phase behavior. The results underline
the influence of the polymer/swelling agent miscibility (as assessed by a solubil-
ity parameter approach) on nanocomposite formation. The example of an epoxy
monomer as the swelling agent shows that either complete miscibility or strong
immiscibility are preferable (37).

Structure Development in Polymer–Clay Nanocomposites

Whatever the elaboration route is, understanding phase behavior of the resulting
nanocomposites is of prime importance to achieve reliable material development.

A lattice-based mean field model has been developed in order to address the

problem of nanocomposite formation (38). By deriving the evolution of the free en-
ergy of the system with interlayer spacing, the model provides some basic predic-
tions regarding the equilibrium states (ie, exfoliated, intercalated, or immiscible),
in relation to the enthalpic and entropic factors of the interacting constituents,
namely the silicate, the tethered surfactant chains, and the polymer. Polymer
confinement results in entropy loss but the latter may be compensated in part by
the entropy gain induced through the increase in conformational freedom of the
tethered surfactant chain upon layer separation. As a consequence, melt inter-
calation is predicted to depend primarily on energetic (enthalpic) factors. Polar
polymers or polymers containing groups interacting with the silicate surface will
favor polymer–clay hybrid formation (35,38).

A theoretical investigation of the phase behavior of model analogues

of polymer–clay nanocomposites has been conducted (39). Combining a self-
consistent field model with density functional theory, the investigation underlines

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NANOCOMPOSITES, POLYMER–CLAY

345

some trends regarding phase morphology and stability for these systems. In par-
ticular, the calculations point at the key influence of the surfactant chain length.
Polymer-like values lead to enhanced miscibility of the clay platelets and polymer
matrix (exfoliated structure), even at moderate level of interactions between the
grafted chains and polymer melt (39). The equilibrium behavior of a mixture of
functionalized and nonfunctionalized chains (taken chemically identical) in the
presence of two infinite planar surfaces has alternatively been considered using
scaling theory. Functionalized chains have a telechelic architecture; ie, the active
groups are located at each end of the chain. Qualitative phase diagrams are de-
rived with prime consideration of the respective influences of interaction energy
between the surface and the end group, and volume fraction

φ of functionalized

chains. Provided the interaction energy

ε is high enough, exfoliation occurs more

easily as

φ is increased, whereas low ε values result in an immiscible mixture

(40).

Nanocomposites Behavior

Thermal Stability.

Thermal stability improvement was already rec-

ognized in the pioneering work by Blumstein on PMMA intercalated in
montmorillonite. Intercalated PMMA degraded at a temperature 50

◦

C higher

than that of conventional unfilled PMMA. In recent years, thermogravimetric
analysis of various polymer–clay systems have confirmed this observation even
for low nanofiller loadings [see eg (13,41), and references therein]. A particularly
striking example is that of cross-linked poly(dimethylsiloxane) incorporating 10
mass% exfoliated organomontmorillonite (42) for which thermal stability under
nitrogen flow is enhanced by 140

◦

C. Overall, restricted thermal motion in silicate

interlayers and hindering of the diffusion of decomposition products are certainly
key factors, but polymer structure and nature of the degradation mechanisms
and degradation conditions are equally important to account for the disparities
observed in literature.

Owing to what has been said previously on the role of the organic modi-

fier of the layered silicate in elaboration and phase behavior of the polymer–clay
nanocomposites, understanding of the key structural factors that influence its
thermal degradation is of prime importance. Major concern is of course toward
processing conditions and fire-retardant behavior as well. Recent work (43) on
alkyl quaternary ammonium montmorillonite by a combination of thermogravi-
metric analysis, Fourier transform infrared spectroscopy, and mass spectrometry
points at the complex degradation behavior of the organic surfactant. The initial
degradation temperature is insensitive to chain architecture and length, or ex-
change ratio. Compared to the parent alkyl ammonium, the thermal stability of
a fraction of the surfactant is substantially lowered because of catalytic sites on
the layered silicates. Polymer processing generally implying temperatures higher
than 180

◦

C, chemical degradation of the surfactant is expected to occur. Many

questions remain regarding the role of the decomposition products on the melt
intercalation mechanism and subsequent nanocomposite phase stability. In this
respect the potential of nmr techniques to follow the fate of the organic modifier
seems quite promising (44).

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Crystal Organization.

Semicrystalline polymer nanocomposites present

a unique interplay between nanoscale morphology of crystal lamellae on one hand
and clay platelet organization on the other. Considering the importance of inter-
facial interactions and the confined chain environment, one may expect drastic
changes in crystal organization. The case of PA6–clay hybrids is rather well doc-
umented. The preferred orientation of the silicate layers under melt-flow condi-
tions, together with polymer confinement, will affect crystallization behavior. For
instance, in injection-molded bars, close to the surface both the silicate layers
and polymer chain axes (and hence lamella thickness) are parallel to the surface,
whereas the chain axes rotate by 90

◦

inside the bar and remain perpendicular

to the silicate layers (45). Clay also favors nucleation of the

γ phase (16,44), con-

trary to bulk PA6 which predominantly crystallizes in the more stable

α form. The

elevated temperature (205

◦

C) crystal morphology of PA6–clay hybrids has been ex-

amined by performing simultaneous small- and wide-angle x-ray scattering (46).
These results clearly establish that the nanoscale correlations of the silicate layer
organization (40–60 nm) affect polymer crystallization, resulting in a less-ordered
crystal

γ phase. Evidence is also provided of the impact of polymer-layer interac-

tions (tethered vs nontethered chains), the more defective lamellae pertaining to
the in situ polymerized (tethered) nanocomposites. Spherulitic structure is unable
to develop as in bulk polymers, which ought to influence the nonelastic mechanical
response. Similar crystallization behavior is also observed in other semicrystalline
polymers such as poly(

ε-caprolactone) or poly(ethylene terephtalate) nanocompos-

ites (27,47). The observation of high melting temperature phases, though defec-
tive, might come from a reduced entropy of fusion

S

f

due to the confined crys-

tallization environment.

Fire-Retardant Behavior.

Controlling polymer flammability remains a

key issue in numerous applications of engineering plastics and commodity
polymers as well. The fire-retardant additive approach provides cost-effective
solutions, but generally at the expenses of some physical and mechanical proper-
ties. There is also growing pressure for environmentally safe products and pro-
cesses, including recyclability and use of halogen-free compounds. For these rea-
sons, recognition of improved flammability properties in the case of polymer–clay
nanocomposites has triggered the development of extensive research programs on
a large variety of materials (41,48,49).

Cone calorimetry is used to evaluate the flammability under fire-like condi-

tions. Relevant parameters such as the rate of heat release (HRR) and its peak
value, heat of combustion (Hc), smoke yield (specific extension area, SEA), and
carbon monoxide yield are obtained. Table 2 shows some typical data for lay-
ered silicate nanocomposites based on organically treated montmorillonite, with
polyamide 6, poly(propylene-graft-maleic anhydride), and polystyrene as the host
matrix. Nanocomposites under investigation have either delaminated (PA6) or
intercalated–delaminated structures. In all cases there is a substantial reduc-
tion in peak HRR value (50–75%), whereas Hc and CO formation show little
variation.

Table 2 also compares the PS–clay nanocomposite with a PS–clay mix for

which intercalation does not occur. The kinetics of heat release are displayed
in Figure 6. In the case of the mix, particle dispersion is only achieved at the
primary particle micrometer level. The peak HRR value remains identical to

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NANOCOMPOSITES, POLYMER–CLAY

347

Table 2. Cone Calorimeter Data

a

Total

Mean

Residue

Peak

Mean

Mean

heat

Mean

CO

yield,

HRR,

HRR,

Hc,

released,

SEA,

yield,

%

kW/m

2

kW/m

2

MJ/kg

MJ/m

2

m

2

/kg

kg/kg

Nylon-6

1.0

1011

603

27

413

197

0.01

Nylon-6–clay 2%

3.0

686

390

27

406

271

0.01

delaminated

Nylon-6–clay 5%

5.7

378

304

27

397

296

0.02

delaminated

PS

0

1118

703

29

102

1464

0.09

PS–clay mix 3%

3.2

1080

715

29

96

1836

0.09

immiscible

PS–clay 3%

3.7

567

444

27

89

1727

0.08

intercalated

PP-g-MA

0

2028

861

38

219

756

0.04

PP-g-MA–clay 5%

8.0

922

651

37

179

994

0.05

intercalated

a

After Ref. 48.

0

50

100

150

200

250

300

Time, s

0

200

400

600

800

1000

1200

1400

Heat release rate, kW/m

2

Fig. 6.

Kinetics of heat release rate for PS-based compounds. Reproduced from Ref. 48.

PS pure;

PS immiscible (3% silicate);

PS intercalated (3% silicate).

that of pure PS while the intercalated PS–clay nanocomposite shows a 50%
reduction.

The observed reduced flammability in the nanocomposites may not be at-

tributed to an additional retention of carbon alone since the residue yields are
not markedly increased. Some key insights are provided by radiative gasifica-
tion experiments which enable to follow pyrolysis either in a nitrogen or in a
nitrogen–oxygen atmosphere. The example of PA6 nanocomposites is revealing

348

NANOCOMPOSITES, POLYMER–CLAY

Vol. 3

enough of enhanced char formation and reduced mass loss rate in comparison to
pure PA6. The current interpretation of these results is that the nanocomposite
flame-retardant mechanism occurs through the build up of a reinforced char layer,
which acts as an insulator, and a mass transport barrier so as to slow down the
escape of the decomposition products. Developments according to this concept in-
clude the use of a PA6–clay nanocomposite in intumescent formulations as in the
case of ethylene–vinyl acetate copolymers with ammonium polyphosphate (50).
Enhanced flame-retardant performance is related to the formation of a thermally
stable phosphocarbonaceous structure in the char, and the blend even shows a
slight improvement in mechanical behavior.

Processing conditions also strongly influence the flame-retardant behavior.

For example, in the case of PS-based nanocomposites, extrusion above 180

◦

C under

partially oxidative conditions yields an intercalated nanocomposite but with no
flammability improvement, whereas the melt-extruded system at 170

◦

C under

nitrogen or vacuum exhibits flame-retardant efficiency (41). The way thermal
degradation of the organic modifier alters the flammability reduction mechanism
has yet to be understood.

Barrier Properties.

Early work in the Toyota Research group acknowl-

edged the great potential of polymer–clay nanocomposites to reduce moisture
absorption and decrease water and gas permeability, even at low clay loadings
(6). A further advantage for packaging applications lies in the fact that lowering
of permeability is achieved while preserving transparency, owing to the suitable
dispersion of platelets smaller than the wavelength of visible light.

The example of polyimide clay films illustrates the dramatic decrease of

permeability coefficients. Only 2 mass% montmorillonite loading reduces the per-
meability by more than 50% of the pure polymer value for water vapor, oxygen,
or helium. Notwithstanding possible changes in diffusion and/or solubility, it has
been postulated that the major role of the clay platelets consists in substantially
increasing the path length of the permeant, that is by creating a highly tortuous
path, due to the high aspect ratio of the clay. A simple theory derived by Nielsen
expresses the relative permeability as follows:

P

c

/P

o

= 1/[1+(L/2W)V

f

]

in which V

f

is the volume fraction of plates, L is the plate length, and W is the

thickness. P

c

and P

o

stand for the nanocomposite and polymer permeability re-

spectively.

Using equivalent loadings of clay but varying the aspect ratio yields results in

fairly good agreement with the theoretical prediction (28). In the same way, a sig-
nificant reduction in water vapor permeability was observed in the case of a poly(

ε-

caprolactone)–organomontmorillonite nanocomposite, showing a fivefold reduc-
tion at only 4.8 vol% clay whereas it is only halved at best with a 20 vol% conven-
tionally filled silicate composite (20). The tortuosity model shows discrepancies be-
tween actual aspect ratios measured by tem and values deduced from curve fitting.
Among the reasons for that are the deviations from the ideal dispersion of platelets
parallel to the film surface, and possible aggregation of individual platelets.

Although tortuosity plays a role in barrier enhancement, some other factors

ought to be taken into account (see B

ARRIER

P

OLYMERS

). Recent work on MXD6

Vol. 3

NANOCOMPOSITES, POLYMER–CLAY

349

Table 3. Comparison of PA6 and PA6–Clay Hybrids Mechanical
Properties

a

Property

PA6 clay hybrid

PA6

Tensile strength, MPa

b

At 23

◦

C

97

69

At 120

◦

C

32

27

Tensile modulus, GPa

c

At 23

◦

C

1.9

1.1

At 120

◦

C

0.6

0.2

Flexural strength, MPa

b

At 23

◦

C

143

89

At 120

◦

C

33

12

Flexural modulus, GPa

c

At 23

◦

C

4.3

2.0

At 120

◦

C

1.2

0.3

Izod impact, J/m

d

18

21

Charpy impact, KJ/m

2e

6.1

6.2

HDT (1.82 MPa),

◦

C

152

65

a

After Ref. 5.

b

To convert MPa to psi, multiply by 145.

c

To convert GPa to psi, multiply by 145,000.

d

To convert J/m to ft

·lbf/in., divide by 53.38.

e

To convert kJ/m

2

to ft

·lbf/in.

2

, divide by 2.4.

nanocomposites is indicative of an oxygen permeation reduction in humid envi-
ronment far beyond what is expected from the increase in path length alone (51).
In the light of the findings regarding chain packing and dynamics in such con-
fined environments (36), models for the prediction of barrier properties ought to
take into account the changes induced in terms of solubility and diffusion. A con-
ceptual model has been proposed (52). So far no general predictive approach is
available.

Mechanical Behavior.

Being able to improve strength and stiffness with

limited alteration of toughness is a goal readily achievable with polymer–clay
nanocomposites (see M

ECHANICAL

P

ROPERTIES

; R

EINFORCEMENT

).

Table 3 gathers some key data of the original work by the Toyota group (5),

which show the dramatic influence of organomontmorillonite on mechanical prop-
erties of PA6–clay hybrids at low mineral loading (4.7 mass%). The improvement
in strength is claimed to have little or no influence on impact properties as evalu-
ated from Izod or Charpy tests. Increase in modulus is paralleled by a substantial
rise in heat distortion temperature. A concept of constrained polymer region re-
lated to the ion-bonding strength of clay and PA6 is introduced to account for the
observed behavior.

Linear Elastic and Rubber Elastic Behavior.

Although stiffening is quite

noticeable in the glassy regime of the amorphous phase, the most spectacular ef-
fect is seen in the rubber elastic regime phase, as already evoked in the case of
reinforcement by cellulose whiskers (2). The PA6–clay hybrids example presented
in Table 3 is quite representative of the situation encountered with semicrystalline
thermoplastics, but elastomeric networks benefit as well of clay layer dispersion
with a two- to threefold increase in modulus for polyurethane or epoxy networks

350

NANOCOMPOSITES, POLYMER–CLAY

Vol. 3

(22). In the meantime, improved elongation at break is observed, contrary to what
is seen in classical filled systems, presumably due in part to dangling chain for-
mation in the network (see E

LASTICITY

, R

UBBER

-

LIKE

). The dynamic mechanical

loss peak related to the glass-transition mechanism is equally informative on the
extent of polymer–clay interaction. This shows mainly in the reduction of the loss
peak area and/or in the evolution of the peak temperature with clay content and
elaboration conditions (5,53). Predictive modeling of both storage and loss mod-
ulus faces a complex challenge in order to account for the mechanical coupling
between the phases, the potential existence of an interphase, and/or a certain
degree of connectivity between the fillers. In the latter situation, a percolation
approach should be useful (2,54). Otherwise models derived from Halpin–Tsai
equations seem quite promising for modulus prediction in relation to clay platelet
arrangement (55,56).

Plasticity and Rupture.

The main drawback identified regarding the solid-

state drawing behavior is certainly the limited elongation at break encountered
for most thermoplastic–clay nanocomposites (9,33) in the vicinity or below the
glass-transition temperature. Nanovoiding and subsequent extensive fibrillation
of the polymer matrix is clearly evidenced from volume strain measurements
during drawing (53) and from in situ tem observations (57). Such enhancement of
nanoscale plasticity offers an opportunity for optimizing the stiffness/toughness
balance. However critical microvoids may develop from areas where load transfer
is no more achievable because of splitting of clay platelet aggregates. This points at
the most critical issue in nanocomposite development, ie, monitoring of elaboration
and processing conditions.

Key Role of Processing.

Scarce work has been devoted to the influence of

processing on microstructure and properties of polymer–clay nanocomposites (58,
59). It is shown that twin-screw extrusion enables achieving a significant degree
of dispersion of the clay platelets, provided residence time and degree of shearing
are optimized in conjunction with the nature of the organoclay. Thermal stability
of the organic modifier is again at the heart of the problem.

In the same way as demonstration was made in the last decade of the im-

portance of processing to design polymer blends, taking the full benefits of the
interesting combination of properties displayed by polymer nanocomposites will
mainly rely on key developments in the field of processing. Automotive and pack-
aging markets are undoubtedly the driving force for it.

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J

EAN

-M

ARC

L

EFEBVRE

Universit´e des Sciences et Technologies de Lille

NOVOLAKS.

See P

HENOLIC

R

ESINS

.

NUCLEIC ACIDS.

See P

OLYNUCLEOTIDES

.

NYLON.

See P

OLYAMIDES

.

OLEFIN-SULFUR DIOXIDE POLYMERS.

See P

OLYSULFONES

.

ORGANOMETALLIC POLYMERS.

See M

ETAL

C

ONTAINING

P

OLYMERS

.

ORIENTED FILMS.

See F

ILMS

, O

RIENTATION

.