Clay-induced Preferred Orientation in Polyethylene/Compatibilized Clay Nanocomposites

Clay-induced Preferred Orientation in Polyethylene/Compatibilized Clay Nanocomposites

A. FAMULARI,1 P. AROSIO,1 S. FILIPPI,2 C. MARAZZATO,2

P. MAGAGNINI,2 L. MINKOVA,3 AND S.V. MEILLE1

1Dipartimento di Chimica, Materiali ed Ingegneria Chimica del Politecnico di

Milano, Milano, Italy2Dipartimento di Ingegneria Chimica, Chimica Industriale e Scienza dei

Materiali dell’Universita degli Studi di Pisa, Pisa, Italy3Institute of Polymers, Bulgarian Academy of Sciences, Sofia, Bulgaria

Nanocomposites of the organically modified clay Cloisitew 15A (CL15A) dispersed inHDPE-g-MA were prepared by melt-compounding. Microcomposites of the same claywith HDPE were also obtained with similar procedures. The spherulitic morphology ofthe polymer matrix was evidenced by optical microscopy in thin films, whereas thestructure of the up to 2-mm–thick, compression-molded samples was investigated byWAXD and SAXS. Preferred orientation of both the clay and the HDPE crystalliteswere evidenced in the microcomposites and, to a greater extent, in nanocomposites,whereas in HDPE and HDPE-g-MA control specimens hardly any anisotropy wasdetected. The degree of orientation of PE crystals increases with CL15A concentration,but also with clay exfoliation, with lower cooling rates and decreasing sample thickness.The orientation of the clay platelets parallel to the compression-molded surface appearsto be determined by the platelets anisotropy and by shear in the mixing and the com-pression-molding procedures. In turn, it determines the preferred uniaxial orientationof HDPE crystals, which have their crystallographic a axis orthogonal, while b and care coplanar, to the sample surface, as already reported in the literature for melt-crystal-lized HDPE films with thickness below 0.3 mm. It is proposed that the HDPE orientationresults from confined crystallization between parallel clay platelets which are on averageless than 0.1 mm apart. Simple models, qualitatively accounting for the observed orien-tation of HDPE, are discussed. Organized architectures resulting from confined crystal-lization of the polymer matrix in nanocomposites with appropriate anisotropic fillers maybe a general feature, important in determining key properties of these systems.

Keywords nanocomposites, polyethylene, X-ray diffraction, preferred orientation,confined crystallization

Introduction

The combination of organically modified smectites with a number of different organic

polymers has become, in recent years, a major field of interest in materials research.[1 – 8]

A large variety of new microcomposites and nanocomposites has been prepared, the

Received 10 July 2006; Accepted 17 October 2006.Address correspondence to S.V. Meille, Dipartimento di Chimica, Materiali ed Ingegneria

Chimica del Politecnico di Milano, via Mancinelli 7, I-20131 Milano, Italy. E-mail: [email protected]

Journal of Macromolecular Sciencew, Part B: Physics, 46:355–371, 2007

Copyright # Taylor & Francis Group, LLC

ISSN 0022-2348 print/1525-609X online

DOI: 10.1080/00222340601158225

355

latter class resulting when exfoliation of the layered inorganic component occurs. Small

percentages of highly anisotropic clays are able to dramatically modify key properties in

polymer nanocomposites. There is a wide consensus that these effects arise because the

inorganic particles are highly anisotropic, while, at least in one direction, dimensions

are in the nanometer range. In polymer-clay nanocomposites dramatic increments in inter-

facial area as compared to classical composites are achieved, which in turn require appro-

priate compatibility of the inorganic particles with the polymer matrix. Specific bonding

interactions may be important and recent literature shows that nanocomposites form

readily in the case of macromolecules which are polar or can hydrogen bond, like

nylons, whereas with polyolefins appropriate chemical modification is needed.[9 – 12]

In the present paper, a series of nanocomposites has been prepared by mixing in the

melt a high-density polyethylene grafted with maleic anhydride groups (HDPE-g-MA)

and the organically modified clay Cloisitew 15A (CL15A). Structural and morphological

features of these nanocomposites were investigated by optical microscopy and by x-ray dif-

fraction (WAXD and SAXS). Comparison with unmodified HDPE composites obtained

with identical percentages of CL15A, with pure HDPE and HDPE-g-MA, are drawn.

In a number of instances, as for increased barrier properties to gases, the orientation of

the anisotropic inorganic platelets in nanocomposite membranes or films is highly

relevant. When the macromolecular matrix crystallizes, orientation effects may also

result.[9 – 13] Such effects can arise because of a variety of reasons, among which the

most obvious and frequently proposed are: (i) the mechanical processing conditions

(e.g., extrusion); (ii) specific interactions with the oriented anisotropic inorganic

component, (iii) effects relating to the nucleation and growth of polymer crystals.

The focus of the present paper is the investigation of the mutual orientation features in

the mentioned, carefully prepared HDPE-g-MA nanocomposites and HDPE composites

with clay. We anticipate at this point that the preferred orientation of the clay is in line

with reasonable expectation. On the other hand, the orientation features of the HDPE

matrix are somewhat surprisingly the same as those found in PE films of thickness

below 0.3 mm[13] and they appear to occur also in the case of quiescent recrystallization

of our millimeter-thick HDPE microcomposite and HDPE-g-MA nanocomposite

samples. The results indicate that the key factor giving rise to HDPE preferred orientation

is, in the present case, the confined nature of the spherulitic crystallization process among

the oriented clay platelets rather than specific interactions or mechanical effects.

Experimental Section

The CL15A organoclay used in this work was Cloisitew 15A by Southern Clay Products,

Inc. It is a montmorillonite modified with 1.25 meq/g of dimethyldi(hydrogenated-tallow)

ammonium ions and characterized by a 3.2 nm d001-spacing.

HDPE (Eltex B5920, Solvay; MFI ¼ 0.39 dg/min; r ¼ 950 kg/m3), was used as a

reference; HDPE-g-MA (Polybond 3009, Crompton Corp. MFI ¼ 5 dg/min; 1% in

mass maleic anhydride; r ¼ 950 kg/m3) was used as the matrix of the nanocomposites.

Polymer-organoclay microcomposites-HDPE/CL15A (wt.%) and nanocomposites

HDPE-g-MA/CL15A (wt.%) were obtained at 1608C in a Brabender Plasticorder by

10 min, 60 rpm mechanical mixing; weight percent values of organoclay in the

specimens are given in parenthesis. Samples were then compression-molded at 1908C,

to obtain 20 mm disks, 2-mm-thick or else 45 or 90-mm-thick films, and subsequently

either slowly cooled (rate �0.58C min21) in the press or quenched in liquid N2.

A. Famulari et al.356

Optical microscopic observations were performed with an Ortholux Leitz polarizing

microscope equipped with a Linkam TMS 93 hot stage. Small amounts of the samples were

melted at 1908C between glass slides in order to obtain thin films about 60–70 mm thick.

The films were cooled to room temperature at a rate of 18/min. Digitized images of the

polymer spherulites were captured. The Hv diffraction patterns of low-angle light scattering

(SALS) were obtained by applying a Bertrand lens and using the microscope in diffraction

mode. Digitized images of the Hv diffraction patterns were captured, too. At least four Hv

patterns for each sample were obtained using different magnifications. The average dimensions

of the spherulites were determined from the Hv patterns using the Stein equation.[14] An image

processing system (Scion Image) was used to analyze the pattern in order to determine the angle

between the incident and the scattered beams corresponding to the maximum intensity.

Wide angle X-ray diffraction patterns (Cu-Ka) were recorded in reflection mode using a

Siemens D-500 Kristalloflex 810 diffractometer. In some cases, for samples requiring

additional thermal treatments, an Italstructures u/u diffractometer equipped with a high temp-

erature attachment was used. Patterns in transmission geometry were collected with a Bruker

P4 diffractometer equipped with a Hi-Star 2D detector. SAXS patterns (Cu-Ka) were measured

using a Bruker Nanostar (with crossed coupled Gobel mirrors and a Hi-star 2D detector).

Results and Discussion

Optical Microscopy

Films of HDPE, HDPE-g-MA, and also of composites and nanocomposites of the modified

CL15A clay with the two polymers, when crystallized upon cooling from the melt at

18C/min, show a clear spherulitic structure typical for HDPE and functionalized polyethy-

lenes (Fig. 1a–c).[15] Because the dispersed clay particles cannot be observed in the dark

polymer matrix, we are unable to establish whether the clay layers are inside or outside the

spherulites. Micrographs of the neat functionalized polyethylene and its nanocomposite

with 10 wt.% clay are presented in Fig. 1a and 1b, respectively, while Fig. 1c shows

the spherulites obtained for HDPE/CL15A (5 wt.%) composites. The corresponding Hv

small angle light scattering patterns, also included in Fig. 1, show well defined four leaf

features with central reflections, which are typical for the diffraction from non-oriented

nets of spherulites. The average dimensions of the spherulites were determined from

the Hv patterns by means of the equation[14]

Rsph ¼ c l=ðnp sin umÞ

where Rsph is the spherulite radius, c a constant (c ¼ 2.05 for spheres and 1.95 for discs),

l/n the wavelength of the light in a medium of refractive index n, um the angle between the

incident and the scattered beams corresponding to the maximum intensity. The values for

the average radius of the spherulites of representative samples are collected in Table 1. As

expected, radii are only slightly smaller for neat HDPE-g-MA than for our HDPE sample

(7.1 to be compared to 7.6 mm), while the average radius decreases to 4–5 mm in the

presence of the clay. It is well known that the spherulite size is predominantly determined

by the primary nuclei density. The smaller polymer spherulite dimensions in the presence

of the clay are, therefore, plausibly due to an increased nucleation rate implying that the

clay acts as a nucleation agent for the polymer matrix crystallization, albeit not a very

efficient one. This is confirmed by the increased temperature of the onset of polymer crys-

tallization in the presence of the clay observed during cooling at 18C/min (Table 1).

Clay-induced Orientation in Polyethylene Nanocomposites 357

HDPE Orientation

The comparison of diffraction patterns in reflection geometry (Fig. 2) of representative

compression-molded samples of HDPE, HDPE-g-MA, and of various clay composites

of the two polymers allows to establish, to a first approximation, cases in which clay

exfoliation has been achieved. As compared to similarly treated HDPE (Fig. 2a),

HDPE/CL15A (5 wt.%) composites (Fig. 2b) show low angle reflections because of the

interlayer spacing of CL15A, indicating that stacks of clay layers persist. At variance,

comparing compression-molded HDPE-g-MA patterns (Fig. 2c) with those of composites

Figure 1. Optical micrographs and Hv patterns of: (a) HDPE-g-MA, (b) HDPE-g-MA/CL15A

(10 wt.%), and (c) HDPE/CL15A (5 wt.%).

Table 1

Average spherulite radius and crystallization temperature for HDPE and

HDPE-g-MA samples with different CL15A content, crystallized from

the melt at a cooling rate of 18/min

Sample Rsph, mm Tc, 8C

HDPE 7.6 122.7

HDPE/CL15A (5 wt.%) 4.5 123.3

HDPE-g-MA 7.1 119.0

HDPE-g-MA/CL15A (5 wt.%) 4.4 123.0

HDPE-g-MA/CL15A (10 wt.%) 4.3 121.0


of this modified polymer with CL15A (Fig. 2d–f) the only apparent difference at low

angle is some continuous diffraction, suggesting large degrees of clay exfoliation, i.e.,

formation of nanocomposites.

The diffraction patterns in Fig. 2 also suggest substantial preferred orientation effects

of the PE matrix in the nanocomposites under investigation. While for HDPE and HDPE-

g-MA (Fig. 2a, c) the ratio I110/I200 between the uncorrected intensities of reflections 110

and 200 (respectively at 2u ¼ 21.5 and 23.88) is ca. 3.2 and nearly conforms to the value of

about 4.0 generally found in unoriented PE specimens†, nanocomposite patterns evidence

a different situation. Specifically, for the slow-cooled HDPE-g-MA/CL15A samples with

a 5 and 10 wt.% clay content, respectively, the I110/I200 ratio in patterns decreases dra-

matically to about 1.0 and 0.2 (Fig. 2d, e) and suggests that the concentration of exfoliated

clay critically influences the orientation of the polymer matrix. Milder but similar effects

(I110/I200 ratio � 1.0) are observed for quenched nanocomposite samples (Fig. 2f) where

crystallization of HDPE is likely to have occurred at lower temperature than for the slow-

cooled samples. A value of about 2 for the I110/I200 ratio in the case of the HDPE/CL15A

(5 wt.%) microcomposite sample (Fig. 2b) suggests that some orientation of the HDPE

matrix occurs also in the composites, even if to a lesser degree than in nanocomposites.

The preferred orientation mode of the HDPE matrix that accounts for the described

perturbations of the relative intensities of the 110 and the 200 reflections observed in our

nano- and microcomposites is very similar to the one reported in the literature for melt-

crystallized thin films of PE (typically with thickness below 0.3 mm).[13] It can be

characterized straightforwardly by 2D transmission patterns: while spectra obtained with

the beam normal to the sample surface (not shown) are isotropic, those in edge-on

geometry are not, as apparent in Fig. 3 for slow-cooled HDPE-g-MA/CL15A (10 wt.%)

samples respectively 2 mm and 90 mm thick. Although the degree of orientation is

somewhat better for the thinner film, in both cases the 200 PE reflection shows a

maximum in the equatorial direction which corresponds to the normal to the surface for

both samples. This is consistent with preferred orientation of the a lattice vector of PE

Figure 2. XRD patterns (reflection geometry) of compression-molded, 2-mm–thick samples,

cooled at ca. 0.58C min21: (a) HDPE, (b) HDPE/CL15A (5 wt.%), (c) HDPE-g-MA, (d) HDPE-

g-MA/CL15A (5 wt.%), and (e) HDPE-g-MA/CL15A (10 wt.%). For comparison also the pattern

of (a) quenched HDPE-g-MA/CL15A (10 wt.%) sample is shown (f).

†Similar I110/I200 ratios of about 4.0 are obtained in powder patterns calculations in which the110 and the 200 reflections display comparable widths (FWHM ca. 0.78 2u).


crystallites along the same direction, i.e., perpendicular to the compression-molded surface

of the samples. Moreover, the azimuthal orientation of the 110 reflection, although clearly

shown in Fig. 3 for the thinner 90 mm film, is similar for the two samples and the apparent

110 maxima are located at about 208 from the patterns’ meridian, which in turn is coplanar

with the film surface‡.

Azimuthal scans of the 200 reflection are shown in Fig. 4 for some nanocomposites and

microcomposites samples. The profiles pertaining to HDPE-g-MA/CL15A(10 wt.%) nano-

composites show a clearly higher degree of orientation with respect to the corresponding

HDPE/CL15A (10 wt.%) microcomposites as revealed by the width at half-maximum

measurements (FWHM is about 428 and 558 in profiles (b) and (c), respectively). It is

also clear that for similar exfoliation, the degree of orientation increases with clay concen-

tration (see profiles (b) and (d) for HDPE-g-Ma/CL15A nanocomposites with 10 and

5 wt.% clay, respectively). Moreover, thin films show somewhat smaller FWHM with

respect to thick samples (see profiles (a) and (b) in Fig. 4). It is interesting that samples

of different kind, which show similar values of the I110/I200 ratio, also have closely

similar values of azimuthal FWHM for the 200 reflection: for values of the I110/I200 in

reflection geometry patterns around 1, about 0.4 and smaller than 0.2 respectively, values

of ca. 588, 508, and 428 are found for the azimuthal FWHM of the 200 reflection. These

observations indicate that there is a good qualitative correlation between the ratio of the

200 to 110 intensities and the azimuthal FWHM of the 200 reflection.

Integration along the equator at constant 2u of the 2D patterns in Fig. 3a, b gives traces

1 and 2 of Fig. 3c which are expected to be closely comparable to the diffraction patterns

from the same film samples in reflection geometry, i.e., to trace e in Fig. 2. This is

qualitatively the case, and rationalizes the alterations of the I110/I200 ratio in the one-

dimensional patterns in Fig. 2. Somewhat surprisingly, however, it is trace 2 from the

90 mm film, rather than trace 1 from the 2-mm-thick sample of Fig. 3c, that is closer to

trace e of Fig. 2 and which pertains to the 2-mm-thick specimen. Because of the

Figure 3. Two-dimensional transmission WAXS patterns of (a) 2 mm and (b) 90-mm-thick com-

pression-molded, slow-cooled HDPE-g-MA/CL15A (10 wt.%) nanocomposite film samples. The

samples are edge on with respect to the incident X-ray beam i.e., the normal to the film surface is

horizontal and corresponds to the equatorial line in the 2D patterns. Traces 1 and 2 of (c) are the

equatorial scans of the 2D diffraction patterns (a) and (b) and are closely comparable to the diffrac-

tion patterns from the same films in reflection geometry (Fig. 2, trace e).

‡The expected value of the angle deviation from the meridian of the 110 reflection for perfectuniaxial orientation along the a axis would be 34.78. Considering the azimuthal half-width (�308) ofthe reflections it is, however, reasonable that a lower value is observed because of overlapping of thebroad symmetrical reflections on the two sides of the meridian.


dominant contribution of the surface layers to the diffraction in the reflection mode, this

observation implies that the surface layer in our nanocomposite samples is more

oriented than the bulk, especially in the case of thick films, and indicates that the orien-

tation is to some degree influenced by processes occurring at the surface of the sample.

The results of an additional experiment are particularly significant: both nanocompo-

sites and microcomposite films were placed unconstrained in a sample holder and brought

in separate experiments to 140 and 1608C in a u/u diffractometer, where spectra were

recorded to confirm complete melting of the polymer matrix. Diffraction patterns

subsequently collected at room temperature, after quiescent recrystallization by cooling

at a rate of ca. 108C/min, gave diffraction patterns showing that the orientation of the

recrystallized polymer was somewhat lower than before melting (Fig. 5) but in essence

preserved, irrespective of the temperature reached by the melt. Such modest loss of orien-

tation, somewhat more significant for the composite than for the nanocomposite, appear to

Figure 4. Azimuthal scans of the 200 reflection in the 2D XRD patterns of (a) HDPE-g-Ma/CL15A

(10 wt.%) 45 mm film, (b) HDPE-g-Ma/CL15A (10 wt.%), (c) HDPE/CL15A (10 wt.%), and (d)

HDPE-g-Ma/CL15A (5 wt.%) slow-cooled samples. The full width at half-maximum (FWHM) is

about 408 in profile (a), 458 in profile (b), and 55–608 in profiles (c) and (d).

Figure 5. Reflection XRD patterns of compression-molded, 2-mm–thick nanocomposites samples:

HDPE-g-MA/CL15A (10 wt.%) (a) after quiescent melting at 140 or at 1608C and cooling at ca.

108C/min, (b) before the melting-recrystalliztion procedure. The other two pattern were similarly

obtained from HDPE/CL15A (10 wt.%) samples, (c) after quiescent melting at 140 or 1608C,

and (d) before the melting-recrystalliztion step.


depend on the faster cooling rate: further experiments with cooling rates of ca. 18C/min

show negligible loss of orientation. In this respect the comparison of traces as shown in

Fig. 5a, b with those of Fig. 2e, f is quite relevant. The observed behavior indicates that

in the present case shear phenomena have only a marginal impact on the orientation of

the HDPE matrix which corresponds to the one occurring when it is melt-crystallized in

films about three orders of magnitude thinner[13,16] than we are looking at. On the other

hand, it appears relevant to note that the orientation of the clay platelets can reasonably

be expected to be invariant upon melting and recrystallization of the polymer matrix.

Clay Orientation

The experimental data discussed up to this point imply that the presence of 5–10 wt.%

clay has a key role in the orientation of the HDPE matrix, a role which is more pronounced

in nanocomposites, but qualitatively similar and quite apparent also in the case of compo-

sites. It is therefore important to discuss the organization of CL15A in our samples. The

most straightforward evidence concerns HDPE/CL15A (10 wt.%) microcomposites: in

the wide angle diffraction pattern in Fig. 6a (film sample vertical and edge-on with

respect to the x-ray beam) it is very clear that the 00I reflections of the modified clay

are oriented along the equator of the image. This corresponds to an orientation of the

clay stacks with the platelets parallel to the surface of the 2-mm–thick, compression-

molded sample. This interpretation is confirmed by the SAXS pattern (Fig. 6b) obtained

from a sample where the 001 Bragg reflection of CL15A concentrates on the equator,

but where also continuous scattering at lower angle is apparent. The orientation of

HDPE in Fig. 6a is not very apparent, but upon close examination we can see that

reflection 110 is somewhat polarized on the meridian while 200 is more intense in the

equatorial area.

Although in composites the degree of orientation of HDPE is lower, it is not surprising

that its mode is qualitatively similar to the one found in nanocomposites (Fig. 3) in which

CL15A is exfoliated. As the clay platelets are exfoliated, no Bragg maxima due to CL15A

are apparent in diffraction patterns from nanocomposites (Fig. 3a, b). The strong streak

running symmetrically from the center along the equator, i.e., the normal to the film

surface, indicates however that the clay platelets are again preferentially oriented

parallel to the compression-molded sample surface. These scattering effects are clearer

Figure 6. Two-dimensional transmission WAXS (a) and SAXS (b) patterns of 2-mm–thick,

compression-molded, HDPE/CL15A (10 wt.%) composite samples. The samples are placed edge-

on with respect to the incident X-ray beam with the normal to the compression surface coinciding

with the equator in the 2D pattern.


in SAXS patterns (Fig. 7) of 90-mm-thick, slow-cooled HDPE-g-MA/CL15A (10 wt.%)

nanocomposite films, in edge- on geometry. Small angle anisotropic scattering confined

in essence on the equator and decreasing in intensity monotonously with increasing 2u

(Fig. 7a) can be related to the silicate monolayers preferentially oriented orthogonally to

the equator, i.e., in our case parallel to the compression-molded film plane.[12] For the

same nanocomposite sample mounted with the film surface perpendicular to the

incoming x-ray beam (Fig. 7b) the scattering streak at low angle is absent.

The above interpretation is consistent with the observation that nanocomposites show

scattering curves with the intensity decreasing to the background level at values of 2u of

about 9–108 independent of the clay concentration. This behavior fits a model of oriented

inorganic platelets whose scattered intensity, assuming a one-dimensional electronic

density modulation orthogonal to the film surface, can be calculated by means of the

shape factor 2/(qd) sin(qd/2), where d is the layer thickness and q ¼ 4p sinu /l.[17]

The shape factor goes to zero and there is a minimum of the scattered intensity when

qd/2 ¼ p, i.e., when d is about 1 nm (2u � 98), which roughly corresponds to the

thickness of a clay layer in clay-silicates.

To better understand the orientation process of the clay we also performed SAXS

experiments, prior to the compression molding step, on HDPE/CL15A (10 wt.%) micro-

composite samples, obtained by melt-mixing the polymer and the clay in a Brabender

rotor. Samples were obtained detaching residual material from the Brabender’s rotors

and cutting ca. 0.1 mm thick slices parallel to the surface that was in contact with the

rotor, after the preparation of the composite according to the usual procedure. Diffraction

patterns, obtained from such specimens using a 0.1 mm pinhole collimated beam all

showed, in the case of HDPE-CL15A 10% composites, significant orientation of the

CL15A 001 reflection (see Fig. 8). So far we were unable to investigate the clay orien-

tation over smaller domains: it is, however, clear from the diffraction patterns that, even

at the melt-mixing stage, the orientation of stacks of clay platelets strongly correlate

locally over length scales of tens of mm. We consider it very plausible that in the com-

pression molding step, above a critical clay concentration, the degree of correlation

among the preexisting local orientations of the stacks of platelets further increases. As

a result even in our 2-mm-thick samples the platelets are largely iso-oriented, in

essence over the whole sample, with the platelet surfaces parallel to the compression-

molded sample surface.

Figure 7. Two-dimensional transmission SAXS patterns of compression-molded, slow-cooled, 90-

mm-thick HDPE-g-MA/CL15A (10 wt.%) nanocomposite film: (a) the film is vertical and edge-on

with respect to the incident X-ray beam, (b) the film surface is orthogonal with respect to the incident

X-ray beam.


Origin of HDPE Orientation

The last issue we want to address is the origin of the high HDPE orientation in our nano-

composites. As already mentioned, the orientation features of HDPE in the compression-

molded microcomposites or nanocomposite films is characterized by the a lattice vector

orthogonal to the film surface and the b and c vectors randomly oriented, but roughly

in the film plane. This mode of orientation has been described in the literature for films

of HDPE up to a few mm thick. The occurrence of the same orientation of MA-grafted

and unmodified HDPE that we find in clay composites, upon PE crystallization on crystal-

line and amorphous substrates (calcite and rubber),[13] is inconsistent with hepitaxy and

indicates that specific bonding interactions are unlikely to play a major role. A similar

edge-on orientation of PE lamellae on the surface plane of the fillers, was reported in a

TEM study of PE composites with talc and kaolin.[18]

We already evidenced that, for a given clay content, the orientation of HDPE in com-

posites increases with the degree of clay exfoliation and it also grows with clay content for

a given degree of exfoliation. Quite remarkably the orientation in millimeter thick HDPE-

g-MA/CL15A (10 wt.%) nanocomposite samples is comparable in mode and in degree to

the one found in melt-crystallized PE films three orders of magnitude thinner. How can

this be explained?

It is well established that the clay platelets in the nanocoposites are highly anisotropic

with a preferred orientation of the silicate layers parallel to the film surface. Assuming

complete exfoliation, with a 10 wt.% clay nanocomposite and a homogenous dispersion

of the inorganic component into the polymer matrix, an average distance of about

20–100 nm can be estimated between neighboring silicate layers, taking a clay layer

thickness of about 1 nm. If some aggregation of the clay platelets is preserved, the

polymer layer thickness will increase while remaining on an average below 0.1 mm.

The quiescent crystallization of the HDPE melt in the nanocomposites can therefore

be thought to occur confined in 20–100 nm thick polymer layers, giving rise to a two-

dimensional spherulitic growth (see Fig. 9). This is a clear result, even if we assume

that the spherulites observed by optical microscopy, with a typical radius of ca. 5 mm,

i.e., nearly two orders of magnitude (100 times) larger than the polymer layer

Figure 8. Two-dimensional transmission SAXS patterns of a HDPE/CL15A (10 wt.%) composite

sample obtained by mechanical mixing the molten polymer and the clay (without the compression

molding step).


thickness, are not confined by the platelets. In fact, the average distance of ca. 50 nm

between exfoliated platelets imply that the platelets should be largely intraspherulitic.

Therefore, the radial polymer lamellae are in essence confined to grow in a highly aniso-

tropic interplatelet region. Since in PE spherulites[14,19 – 21] it is well established that the

crystallographic b axis is radial, b will be largely constrained to lie parallel to the plane

of the clay platelets, i.e., parallel to the surface plane of the compression-molded

samples. This conclusion is rather straightforward, even without any constraints about

the interfacial energy between PE crystals and substrate, at variance to what was

suggested in some previous work.[13] It is simply based on a faster growth rate, on

average, of PE crystals along the b axis direction.

The in-plane orientation of the PE b axis is indeed compatible with the observed

preferred orientation of the a axis perpendicular to the sample surface. In order to fully

account for this feature it is, however, necessary to constrain not only the b axis but the

also the c axis, i.e., the chain axis of PE, to lie preferentially in the plane of the sample.

Crystalline PE lamellae therefore tend to lie edge-on, with the bc face as the contact

plane, on clay platelets (see inset to Fig. 9) or on other substrates. Although this orientation

appears to be a quite general occurrence with PE thin films, we are not aware of any

detailed discussion of the issue in the literature. We will attempt to briefly address it

with some speculative considerations in the following paragraph.

The first possibility is simply a more effective primary nucleation of PE spherulites

and crystals with chains lying parallel to the substrate surface plane. This sensible hypoth-

esis could account for our main observations. In fact it is the common assumption in

standard treatments of primary heterogeneous nucleation.[19] Interestingly is can be

found also with other polymers, as, for example, in melt-crystallized or solvent cast

Figure 9. Structural model of melt-compounded, compression-molded PE/clay microcomposites

and nanocomposites: spherulites of PE (plausibly anisotropic) and oriented clay platelets are

shown together with the lattice vectors of the orthorhombic PE unit cell. The inset shows an enlarged

view with edge-on PE lamellae growing in the confined space between clay platelets.


films of isotactic polypropylene,[22] poly(3-alkylthiophenes),[23] syndiotactic poly-

styrene[24] at least under specific conditions. Since nucleated lamellae at the surface of

the heterogeneous layer are assumed to have this orientation, a bias favoring similar orien-

tation of other lamellae will be established.

There is also the possibility that it is not the primary nucleation event that determines

the orientation of the lamellae but successive growth, (i.e., secondary nucleation, see

Appendix). In other words the original nuclei may have a wide variety of different orien-

tations with respect to the clay surface. However the polymer lamellae with the “correct”

orientation with respect to the clay surface (i.e., both the b and the c axes in the plane of the

film) experience better secondary nucleation and grow faster. As a consequence preferred

orientation with a perpendicular to the clay layers arises.

Clearly both effects tend in general to dissipate as we go further away from the het-

erogeneous surface and the orientation effects will depend on the film thickness, unless

there are strong homohepitaxial effects: i.e., unless the orientation of lamellae not in

immediate contact with the substrate, is determined by the orientation of lamellae lying

on the substrate.

It seems hard with the present data to discriminate between the two “mechanisms”

and in fact both may play a role. In any case, the confined growth mode of polymer

lamellae implies that in our nanocomposites spherulites are likely to be highly anisotropic

(see Fig. 9) if not bidimensional. The relative dimensions of spherulites and CL15A

particles also appear to imply that clay platelets are at least, in part, intra-spherulitic.

More detailed investigation on these issues is in progress.

Conclusion

Clay lamellae of CL15A in organized stacks and in their exfoliated states, at concentration

values around 10 (wt.%), adopt roughly parallel organization modes in HDPE (and HDPE-

g-MA) composites obtained by melt-compounding, over domains with dimensions

estimated to be of the order of 10 mm or more (Fig. 8). Upon compression molding

these micro and nanocomposites into films, clay platelets and their stacks tend to orient

with the clay platelets parallel to the surface of the compression-molded film. For com-

pression-molded films with thickness 10–100 mm the degree of orientation of clay

lamellae is high and relatively uniform. For films of greater thickness more complex situ-

ations may result, i.e., nonnegligible skin-core effects, with somewhat higher degrees of

orientation of the clay lamellae at the surface.

In compression-molded HDPE/CL15A microcomposites or nanocomposite films the

HDPE matrix crystallizes with a uniaxial preferred orientation mode characterized by the

a lattice vector orthogonal to the film surface and the b and c vectors randomly oriented,

but roughly in the film plane. This mode of orientation, previously described in the litera-

ture for very thin films of HDPE crystallized on various crystalline and amorphous

substrates,[13] points to nonspecific interactions with the substrate. The degree of orien-

tation of HDPE increases with the degree of clay exfoliation for a given clay percentage,

and with clay content, for a given degree of exfoliation. The degree of orientation of

HDPE-g-MA in nanocomposites with 10 wt.% CL15A is comparable to the one in PE

films three orders of magnitude thinner. HDPE orientation is favored also by high crystal-

lization temperatures (or low cooling rates) suggesting that the nucleation step of PE crys-

tallization, plausibly occurring at the polymer-clay interface, plays a significant role for

the orientation of PE crystals.


Upon quiescent recrystallization from the melt of the mentioned composites and

nanocomposites, the mode and the degree of orientation of the clay and of crystalline

HDPE is invariant. The macroscopic preferred orientation of the clay platelets is

plausibly determined simply by effective space-filling requirements and is also

enhanced by shear. Once established, it is largely preserved in the polymer matrix. On

the other hand the orientation of PE is probably determined by the nucleation mode and

by confinement (confined lamellar growth) in the development of spherulites and thus

small shear effects are largely irrelevant. As a consequence, it is reasonable to suggest

that the resulting HDPE spherulites will be significantly anisotropic.

The macroscopic orientation of PE crystals growing in the anisotropic confined inter-

platelet polymer melt, with thickness on average less than 100 nm may in principle be

caused by: (i) the nucleation of PE crystals, plausibly more favored with PE chains

lying flat on the substrate; (ii) kinetic factors, possibly related to the surface energy of

different lateral faces of PE crystals, involving the interaction with the PE melt and/or

the substrate. Although (i) seems more straightforward, (ii) could also play a role as

discussed in the preceding section and in the Appendix.

We finally suggest that, given the nature of the interaction between the polymer and

the surfaces determining the orientation of crystalline HDPE, the interpretation of the

phenomena that we have described may extend well beyond the case we examined: for

example it might apply, complementing shear, to the orientation described for nylon in

nanocomposites in the seminal paper by Kojima et al.[12] Indeed it appears to be quite

general, and in the absence of specific polymer-particle interactions, orientation effects

in very thin (below 0.1 mm) polymer films, should be closely related to those in nano-

composites of the same polymers.

Appendix

The morphology of PE thin films has been analyzed in some detail with respect to the

influence of the substrate on the secondary nucleation process controlling the crystal

growth by Bartczak et al.[13] Assuming, conventionally, that preferential crystal growth

occurs along the b axis direction, when PE lamellar crystals grow in contact with a

substrate surface in very thin films, preferred orientations for the c chain axis in the

plane of the substrate surface (see Fig. 10, case a) is found. The in-plane main chain orien-

tation has been shown to occur also in solvent cast or melt-crystallized films, with some

other polymers[22 – 24] in the absence of shear. In other cases or, under different conditions,

with the same polymers (e.g., syndiotactic polystyrene)[25] orientation of the chain axis

roughly orthogonal to the substrate can also be found (see Fig. 10b, case b). These are

the two chain orientations that we will consider in our treatment.

Starting from classical nucleation concepts,[19,21,26] the change of the free enthalpy

related to a secondary nucleus formation not in contact with the substrate (Figs. 10a

and b, positions 1) can be expressed as:

DG ¼ �abolDgf þ 2bolgcm þ 2aboge ð1Þ

where a, bo, and l are the width, thickness (equal to that of a monolayer for a given super-

cooling) and length of the nucleus, respectively, Dgf is the free enthalpy of melting, gcm is

the interfacial free energy between the bol lateral crystal plane and the polymer melt and ge

is the interfacial free energy between the abo fold surface of the crystal and the polymer


melt. The energy barrier for critical nucleus formation can be calculated imposing the

stationary conditions:

@DG�

@a¼ 0 and

@DG�

@l¼ 0

The resulting critical values

a� ¼ 2gcm=Dgf and l� ¼ 2ge=Dgf

determine the free enthalpy maximum corresponding to the barrier for the critical nucleus

formation

DG� ¼ 4bogcmge=Dgf ð2Þ

It can be shown that the energy barrier for the formation of a critical size nucleus is lower

close to the substrate than in bulk crystallization.[13,19,21,26] In the case of the formation of

a secondary nucleus in the reentrant corner between the substrate and the growing crystal

face (Fig. 10a, position 2: polymer chain axis lying in the surface plane) the free enthalpy

change is

DG ¼ �abolDgf þ bolðgcm þ gcs � gmsÞ þ 2aboge ð3Þ

with critical values

a� ¼ ðgcm þ gcs � gmsÞ=Dgf and l� ¼ 2ge=Dgf

corresponding to a free enthalpy change maximum

DG� ¼ 2boðgcm þ gcs � gmsÞge=Dgf ð4Þ

where gcs and gms represent the interfacial free energy between the lateral plane of the

crystal and the substrate, and between the polymer melt and the substrate, respectively.

The comparison of Equations 2 and 4 shows that the energy barrier related to the

formation of a critical size nucleus in close contact with the substrate is lower than the

barrier for the formation of such a nucleus away from the substrate, provided that

gcm . gcs 2 gms.[13]

Figure 10. Two limiting models for secondary nucleation of chain-folded polymer crystal on a sub-

strate surface: (a) chain axis in the substrate plane, (b) chain axis orthogonal to the substrate plane.

Secondary nuclei growing away from the surface (position 1) and in contact with the surface plane

(position 2) are compared in the two cases.


Considering the case of a secondary nucleus in close contact with the substrate

but with the polymer chain axis normal to the plane of the surface (see Fig. 10b,

position 2), the change of free enthalpy can be written as

DG ¼ �abolDgf þ 2bolgcm þ aboðge þ gcs1 � gmsÞ ð5Þ

where, in the last term, gcs1 represents the interfacial free energy between the fold surface

and the substrate. The critical a and l values are easily calculated as above

a� ¼ 2gcm=Dgf and l� ¼ ðge þ gcs1 � gmsÞ=Dgf

while the corresponding barrier for the critical secondary nucleus formation is

DG� ¼ 2bogcmðge þ gcs1 � gmsÞ=Dgf ð6Þ

Comparing expressions (4) and (6) we can see that the free energy barrier for the

formation of a critical secondary nucleus with the chain axis normal to the substrate

(Fig. 10b, position 2) is larger than that of the corresponding nucleus with the main

chain parallel to the substrate (Fig. 10a, position 2) provided that gcm . ge and that we

can assume gcs1 ffi gcs.

While both of the above hypotheses may need further support, it is quite clear that

energy barriers relative to positions 1 of cases a and b (Fig. 10a and b) coincide. With

respect to position 2 of case b, considering the plausible features of the fold surface of

the secondary nucleus, we could then adopt a different view. We may approximate the

interaction of the secondary nucleus with the substrate as a crystal-amorphous-substrate

interaction rather than a simple crystal-substrate interaction, as we have done so far. If

this description applies, the situation represented in case b at position 2 will differ from

position 1 only marginally. The conditions that make secondary nucleation described at

position 1 disfavored over position 2, case a, would therefore apply also to case b, both

for position 1 and 2, i.e., position 2, case a, is preferred.

Acknowledgment

The authors acknowledge the financial support of MIUR PRIN2004 (A.F., S.F., C.M.,

P.M., P.A., S.V.M.), NATO-grant CBP.EAP.CLG 981257 (P.M., L.M.)- and CARIPLO

Foundation (A.F., P.A., S.V.M.) S.V.M. also acknowledges stimulating discussions with

Dr. Annamaria Buono.

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Clay-induced Preferred Orientation in Polyethylene/Compatibilized Clay Nanocomposites

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