Organoclay Nanocomposites from Ethylene–Acrylic Acid Copolymers

Organoclay Nanocomposites from

Ethylene–Acrylic Acid Copolymers

Sara Filippi,1 Cristina Marazzato,1 Pierluigi Magagnini,*1 Liliya Minkova,2 Nadka Tzankova Dintcheva,3

Francesco P. La Mantia3

1 Dipartimento di Ingegneria Chimica, Chimica Industriale e Scienza dei Materiali, Universita degli Studi di Pisa,Via Diotisalvi 2, Pisa 56126, ItalyFax: (þ39) 050 511266; E-mail: [email protected]

2 Institute of Polymers, Bulgarian Academy of Sciences, Acad. G. Bontchev str., Sofia 1113, Bulgaria3 Dipartimento di Ingegneria Chimica dei Processi e dei Materiali, Viale delle Scienze, Universita degli Studi di Palermo,Palermo 90128, Italy

Received: May 26, 2006; Revised: July 17, 2006; Accepted: July 18, 2006; DOI: 10.1002/mame.200600217

Keywords: ethylene–acrylic acid copolymers; nanocomposites; organoclay

Introduction

A vast number of papers have been published in the last few

years on polymer/clay hybrids,[1–8] and an increasing number

of the most recent studies are dealt with the preparation and

characterisation of smectite–clay nanocomposites with poly-

olefinic matrices.[5,9–51] The expected improvement in one or

more of the properties of these commodity polymers, e.g.,

barrier, mechanical, fire resistance, etc., might, in fact, be

foreseen to balance the cost increase due to the nanofiller

(whose addition, if made by melt compounding, should not

require significant changes in the usual processing techni-

ques). The strong hydrophobicity of polyolefins, however,

makes their interaction with the silicate surface quite difficult,

even for clays organically modified by substitution of the

metal cations with e.g., quaternary ammonium ions containing

one or more long alkyl groups. In particular, for the composites

prepared by melt compounding from the different polyethy-

lene (PE) grades,[9–40] most studies have shown that no

significant intercalation is obtained, unless the PEs are blended

with appropriate oligomeric or polymeric compatibilisers.

The literature describing the use of copolymers of ethylene

with polar monomers, either as matrices or as compatibilisers

for PE/clay nanocomposites, is quite rich. In particular,

Summary: A study of the structure–property relationshipsfor nanocomposites prepared by melt compounding fromethylene–acrylic acid copolymers of varied composition andmolecular architecture, and organoclays modified withdifferent ammonium ions has been made by DSC, POM,SEM, TEM, WAXD, and rheological and mechanical tests.Within the series of clays investigated, the best levels ofdispersion were displayed by those organically modified withquaternary ammonium ions containing two long alkyl tails.The relevant nanocomposites were shown to possess mixedexfoliated and intercalated morphology. The spacing of theintercalated clay stacks, most of which comprise few silicatelayers, was found to be independent of clay loading, in therange of 2–50 phr, and to change with the molecular archi-tecture of the matrix polymer. An indication that the excesssurfactant present in some of the clays, and the organicmaterial added in others to expand the interlayer spacing,were expelled from the clay galleries during melt blendingand acted as plasticisers for the matrix polymer, was obtainedfrom WAXD and rheological characterisations. TEM micrograph of the nanocomposite of EAA1 with 11 phr

of 15A.

Macromol. Mater. Eng. 2006, 291, 1208–1225 � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1208 DOI: 10.1002/mame.200600217 Full Paper

several studies have been made on silicate hybrids prepared

from ethylene-vinyl acetate (EVA) copolymers, or their

blends with PE,[4,6] for applications requiring enhanced

barrier properties and flame resistance. The copolymers,

either commercial or synthesised on purpose, of ethylene

with maleic anhydride (MA) have also been investigated

extensively,[10–15,17,19–25,29,33,41–45,48] mainly because of

their effectiveness as compatibilisers for PE-based nano-

composites. Actually, most of the studies on nanocomposites

containing PE-g-MA, either as the matrix or the compatibil-

iser, report high levels of organoclay exfoliation, although the

results are not always consistent with respect to the minimum

concentration of MA in the copolymer and the optimal

compatibiliser-to-clay ratio.

It may be surprising that the copolymers of ethylene with

acrylic acid (AA) or methacrylic acid (MAA), though being

largely employed as adhesion promoters, blends compatibil-

isers etc., have received comparatively little attention inview

of the preparation of PE-based nanocomposites. Preston

et al.[46] described the synthesis of nanocomposites from

three different ethylene copolymers and found (by WAXD)

that the ethylene-co-methyl acrylate-co-acrylic acid terpol-

ymer gave rise to better intercalation, with an organically

modified bentonite, than the two others: EVA and poly-

[(ethylene)-co-(methyl acrylate)]. These authors concluded

that the AA units interact effectively with the clay surfaces.

In a very recent paper,[27] Fang and coworkers described the

synthesis of an HDPE-g-AA copolymer with 8.4 wt.-% AA

and its use for the preparation of nanocomposites with a

bentonite modified with trimethyloctadecylammonium ions.

On the basis of WAXD patterns and TEM micrographs, these

authors concluded that the nanocomposites are characterised

by good ‘exfoliation or intercalation’ and suggested, from

FTIR analyses, that chemical reactions occur between the

carboxyl groups of the HDPE-g-AA matrix and the clay.[27]

Finally, Chrissopoulou et al.[29] synthesised a diblock copo-

lymer of ethylene and MAA (PE-b-PMAA) and used it

as a compatibiliser for nanocomposites of HDPE with a

commercial organoclay (Cloisite1 20A, by Southern Clay

Products). They observed a gradual shift in the basal dif-

fraction peak of the clay from 2y¼ 3.158 to 2y¼ 2.458,accompanied by a significant decrease in intensity, as the

concentration of PE-b-PMAA in the polymer matrix was

increased up to 15 wt.-%. On the basis of diffractometric

data, these authors concluded that most probably, as the

amount of the PE-b-PMAA increases, the system attains a

mixed intercalated and exfoliated morphology, resulting

from intercalation of the copolymer chains into the galleries

of the MMT driven by the polarity of the carboxyl groups.[29]

Other studies[47,48] have dealt with nanocomposites from

ethylene-MAA (EMAA) ionomers. In particular, Hsiao

et al.[47] studied the hybrids prepared from EMAA ionomers,

containing�3.5 mol-% of MAA partly (80%) neutralised with

sodium or magnesium, and a commercial organoclay (Nano-

mer1 I30 E, by Nanocor Inc.). Through small-angle X-ray

scattering (SAXS) and TEM studies, they found that the

nanocomposite from the Na ionomer was predominantly

exfoliated, whereas that prepared from the Mg ionomer was

only slightly exfoliated and showed a distinct scattering

peak corresponding to a d001 spacing of �3 nm. A detailed

paper by Shah et al.[48] dealt with the behaviour of nano-

composites prepared by melt processing from a commercial

EMAA ionomer (Surlyn1 8945, du Pont; MAA content

5.6 mol-%; Na-neutralisation 39%) and a series of organo-

clays. They found that the extent of clay dispersion and the

mechanical properties varied appreciably with the chem-

istry of MMT modification. In particular, they demon-

strated, by TEM, WAXD, stress-strain and Izod-impact

measurements, that improved levels of clay exfoliation and

enhanced mechanical properties could be obtained by the

use of: (i) amine modifiers with more than one alkyl tail,

(ii) longer alkyl tails, (iii) amines containing 2-hydroxy-

ethyl rather than methyl groups and (iv) an excess amount of

the amine surfactant with respect to the cation exchange

capacity (CEC) of the pristine MMT.

In a previous communication from our laboratories,[51] a

preliminary X-ray and TEM investigation of the morphology

of nanocomposites from different ethylene-AA copolymers

(EAA) loaded with fixed amounts (5 phr) of organoclays

modified with different proportions of the same dimethyldi-

hydrogenated tallow quaternary ammonium surfactant has

been reported. All nanocomposites displayed mixed, interca-

lated and exfoliated morphology. Together with a population

of individual platelets, thin stacks of a few irregularly spaced

silicate layers were seen in the TEM images, whose average

periodicity was shown by WAXD to depend on the molecular

architecture of the matrix copolymers. In fact, a d001 spacing

of about 4 nm was measured for the branched EAA matrices,

whereas a considerably lower MMT expansion (d001ffi 3 nm)

was found with a linear HDPE-g-AA (HDAA) matrix. The

other investigated variables such as the polymer molar mass,

the AA concentration (in the range of 6.2–11 wt.-%) and the

excess of surfactant present in two of the investigated clays,

were shown to have negligible effect on morphology.

In this work, the morphological characterisation has been

extended to EAA and HDAA nanocomposites filled with a

series of commercial and experimental organoclays containing

different quaternary ammonium ion modifiers, and their ther-

mal, rheological and mechanical properties have been investi-

gated. Moreover, the effect of increasing the clay loading up to

50 phr has been studied in order to test the possibility of

obtaining concentrated nanoblends which might be promising

for use as compatibilisers for PE-based nanocomposites.

Experimental Part

Materials

The polymers used as matrices are indicated in Table 1,together with their sources and some of their properties. The

Organoclay Nanocomposites from Ethylene–Acrylic Acid Copolymers 1209

Macromol. Mater. Eng. 2006, 291, 1208–1225 www.mme-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Escor1 (EAA) samples are random copolymers with a signi-ficant amount of long branches. Polybond1 (HDAA) is a linearcopolymer produced by grafting AA onto HDPE.

The organoclays used in this study are listed in Table 2 alongwith their sources, the milliequivalent exchange ratio (MER),the loss on ignition determined by thermogravimetric analysis(TGA) as the weight loss in air at 900 8C, and the d001 spacingdetermined by X-ray analysis. Most organoclays were suppliedby Southern Clay Products, Inc.; they were prepared by a cationexchange reaction between an Na-MMT with a CEC of0.926 meq�g�1 (Cloisite1–Naþ) and appropriate amine sur-factants. Additional organoclays were either supplied by SudChemie (Nanofil1 848 and Nanofil1 SE3000) or prepared inour laboratory from Cloisite1–Naþ (M3C18). The nomencla-ture used to describe the organic modifier’s structure is thatproposed by Shah et al.:[48] H for hydrogen, M for methyl, T fortallow (a blend of unsaturated alkyl groups with approximatecomposition: 65% C18; 30% C16; 5% C14); (HT) for hydro-genated tallow, (HE) for 2-hydroxy-ethyl.

Chemicals, e.g. octadecyltrimethylammomium bromide,and solvents were supplied by Aldrich.

Techniques

The M3C18 organoclay was synthesised as follows: 30 g ofCloisite1–Naþ was stirred for 2 h in 400 mL of a 50:50 v/vH2O/EtOH solution at 75 8C; 13.3 g of trimethyloctadecylam-

monium bromide, dissolved in 100 mL of a similar H2O/EtOHmixture, was added to the clay solution and the mixturewas mechanically stirred for 2 h at 75 8C and then sonicated for2 more hours at 60 8C. The precipitate was separated with asintered glass filter, and washed with warm H2O/EtOH solutionup to complete elimination of bromine ions in the filtrate(AgNO3 test). The organoclay was dried in a vacuum oven at40 8C for 24 h, milled in a mortar and sieved (200 mesh).

The composites were prepared by melt compounding in aBrabender Plasticorder static mixer of 50 mL capacity,preheated to 120 8C (150 8C for the HDAA-based composites).The rotor speed was maintained at 30 rpm for about 2.5 min andwas then increased gradually (in 30 s) to 60 rpm. The overallblending time was 10 min, unless otherwise stated. Dependingon the viscosity, the temperature rose by 10–20 8C above theset value, due to stress heating. Prior to charging into the mixer,the polymers were dried in a vacuum oven at 60 8C for at least24 h and stored in a desiccator; the organoclays were used asreceived. The polymer was premixed in a beaker with thechosen quantity of clay powder, which showed good adhesionto the pellets, and the mixture was fed into the mixing bowlwith caution in order to minimise any loss of powder. Theapplied torque and the blend temperature were recorded duringthe whole compounding period. At the end, the moltencomposites were extracted from the mixer and cooled naturallyin the atmosphere. Blank samples to be employed as referenceswere prepared by processing pure polymers with the same

Table 1. Polymers used for the nanocomposites preparation.a)

Sample name Commercial designation (manufacturer) AA content Density MFIb)

mol-% kg�m�3 dg�min�1

EAA1 Escor1 5100 (Exxon-Mobil Chemical) 4.6 940 8.0EAA2 Escor1 5000 (Exxon-Mobil Chemical) 2.5 931 8.0EAA3 Escor1 5001 (Exxon-Mobil Chemical) 2.5 931 2.0HDAA Polybond1 1009 (Crompton Corp.) 2.4 950 5.0

a) Manufacturer’s data.b) Melt flow index (ASTM D1238).

Table 2. Organoclays used in this work.

Clay Commercial designation (Supplier) Surfactant structure MERa) Organic contentb) d001b)

meq�g�1 % nm

6A Cloisite1 6A (South. Clay Prod.) M2(HT)2 1.40 45.2 3.4915A Cloisite1 15A (South. Clay Prod.) M2(HT)2 1.25 42.4 3.1020A Cloisite1 20A (South. Clay Prod.) M2(HT)2 0.95 38.5 2.53SE3000 Nanofil1 SE3000 (Sud Chemie) -c) -c) 54.5 3.5093A Cloisite1 93A (South. Clay Prod.) MH(HT)2 0.90 36.0 2.4730B Cloisite1 30B (South. Clay Prod.) M(HE)2T 0.90 30.0 1.84N848 Nanofil1 848 (Sud Chemie) H3C18 -c) 25.4 1.82M3C18 Experimental M3C18 1.22b) 30.4 1.96

a) Supplier’s data.b) Experimental data.c) Undisclosed.

1210 S. Filippi, C. Marazzato, P. Magagnini, L. Minkova, N. T. Dintcheva, F. P. La Mantia


procedure. The clay concentrations in the composites are givenin phr (grams of organoclay per hundred grams of polymermatrix).

DSC analyses were made under nitrogen, on samples of 5–8mg, with a Pyris Perkin Elmer apparatus calibrated withindium and tin standards. The scanning rate was 10 8C�min�1.The first heating scan was carried out from 20 to 190 8C and thistemperature was maintained for 5 min to exclude thermalhistory effects. The values of the temperatures and enthalpychanges associated with the crystallisation and the fusion of thesamples were recorded from the cooling and second heatingscans. The enthalpies were normalised to the polymer contentand the degrees of crystallinity were calculated from theexperimentalDHm figures with reference to the 293 J�g�1 valueof 100% crystalline PE.[52]

Thermogravimetric analyses were made with a Q500 TAInstruments balance using samples of about 10 mg. Tests werenormally made in air (60 mL�min�1), in a temperature rangebetween 30 and 900 8C, with a scanning rate of 10 8C�min�1.

Optical microscopic observations were made on a LeitzOrtholux polarizing microscope equipped with a Linkam TMS93 hot stage and a digital camera JVC TK-1085E. A smallfragment of the samples was placed on a thoroughly cleanedglass slide, melted at 190 8C on the microscope hot stage, andpressed with a cover slip to obtain a film of 80–100 mmthickness. Micrographs were taken both at 190 8C and at roomtemperature after cooling at a rate of 1 8C�min�1. The tem-peratures TcPOM corresponding to formation of the first bire-fringent crystals during the cooling cycle, as observed throughcrossed polarisers, were recorded. At least four Hv patterns ofsmall-angle light scattering (SALS) were recorded, at differentmagnifications, for each crystallised sample at room temper-ature, using the microscope in the diffraction mode with aBertran lens. The Hv patterns were used to calculate the ave-rage spherulite dimensions with the Stein equation.[53] Animage processing system (Scion Image) was used to analysethe patterns in order to determine the angle of the incident andscattered beams corresponding to maximum intensity.

Scanning electron microscopy (SEM) observations of cryo-fractured composite samples, coated with gold, were madewith a Jeol JSM-5600LV microscope.

Samples for WAXD analysis were prepared by compressionmoulding, using a laboratory Carver press preheated to a tem-perature of 190 8C (unless otherwise stated). The mouldconsisted of a 2-mm thick stainless steel plate, with four holesof 20-mm diameter, which was sandwiched within two stain-less steel plates covered with anti-adherent film. Four sampleswere prepared at a time; after moulding, they were let to coolslowly in the press. WAXD measurements were made in thereflection mode with a Siemens D500 Krystalloflex 810apparatus with an X-ray wavelength of 0.1542 nm at a scanrate of 1.08�min�1. The analyses were normally made on theouter surface (skin) of the compression moulded discs. How-ever, for some of the nanocomposites, the core morphologywas also analysed with the X-ray beam reflected off the surfacecreated by abrading the discs with a fine sandpaper until thethickness was reduced to �1 mm.

TEM observations were made with a ZEISS EM 900 micro-scope, at accelerating voltages of 50 and 80 KeV, in the Genovadivision of the Institute for Macromolecular Studies (ISMac)

of C.N.R. Samples for TEM analysis were taken either frompieces of the composites as received from the Brabender mixeror from the compression moulded discs used for WAXD.Ultrathin sections�50 nm thick were cut with a diamond knifeat a temperature of �130 8C (�145 8C for the HDAA nano-composites) using a Leica Ultracut UCT ultramicrotomeequipped with a Leica EM FCS cryosystem. Sections werecollected on the surface of a water–dimethylsulphoxide (60:40 v/v) bath cooled to �60 8C, taken on TEM grids and driedwith filter paper. The micrographs were digitised using a high-resolution plane scanner and the silicate lamellae were identi-fied on the magnified electronic images with lines drawnmanually. With the use of the Image-Pro1 Plus software, thedistances between adjacent lines were automatically recorded,thereby allowing the determination of the gallery height d2 andof the spacing d1 of the tactoids to be compared with the d001

figures measured by WAXD.Samples for rheological and mechanical analyses were

made by compression moulding with a Carver press preheatedto 190 8C. The rheological tests were carried out in shear flowusing a Rheometrics RDA II apparatus in the plate–plategeometry (the plate’s diameter was 25 mm) at 160 8C in thefrequency range of 0.1–500 rad�s�1, with 5% strain. Themechanical measurements were made with a universal Instronmachine mod. 4443, according to ASTM D882. Averagevalues were calculated for elastic modulus (E), yield strength(YS), tensile strength (TS) and elongation at break (EB) fromthe results of at least five measurements. The reproducibility ofthe results was within �5%.

Results and Discussion

Calorimetry and Polarised Optical Microscopy

The crystallisation temperature Tc, measured by DSC at

10 8C�min�1 (peak values), and TcPOM, measured by POM

at 1 8C�min�1 (first appearance of birefringent crystallites),

and the fusion temperatures Tm, measured by DSC at

10 8C�min�1 on the second heating scan (peak values), are

given in Table 3 for the polymers and some of their nano-

composites, together with the degrees of crystallinity calcu-

lated from the fusion enthalpies DHm and the average

spherulite sizes measured by SALS. The POM micrographs

of two representative copolymers and their nanocomposites

with 5 phr of 15A, taken under crossed polarisers on films

cooled to room temperature from 190 8C at 1 8C�min�1, are

shown in Figure 1, together with the relevant Hv patterns.

The crystallisation temperatures Tc of the branched EAA

copolymers are considerably lower than that (95–100 8C)

typical for LDPE homopolymers with comparable molec-

ular architecture. The Tc lowering, which is stronger for

EAA1, is clearly due to the presence of the AA comonomer

units disturbing the crystalline regularity. The linear AA-

grafted copolymer (HDAA) expectedly shows a much

higher Tc. The strong effects of both molecular architecture

and AA content are also demonstrated by the Tm data and,

even more, by the degrees of crystallinity calculated from

the fusion enthalpies DHm.



On the other hand, only slight differences were observed

between the Tc, Tm and DHm figures of the polymers and

those of the relevant nanocomposites. Since these data were

obtained from asymmetric DSC peaks with long tails on the

low temperature sides, they are probably affected by consi-

derable uncertainty and we believe that the recorded dif-

ferences should not be given much weight. On the basis of

the DSC data, therefore, one might feel inclined to infer

that the dispersed silicate particles have negligible effect

on the crystallisation of these copolymers. However, from

the TcPOM figures given in Table 3, an indication that the

presence of clay particles may lead to an increase in the

crystallisation temperature of the polymers can be drawn.

Moreover, as it is shown by the micrographs in Figure 1 and

the spherulite size data in Table 3, the definite increase in

crystalline nuclei concentration observed in the presence of

clay provides strong indication in favour of a heterogeneous

nucleating effect played by the dispersed inorganic parti-

cles. Non-isothermal and isothermal crystallisation studies

carried out on these materials by one of us have confirmed

that the nucleation rate is increased in the presence of clay

and that the spherulite growth changes from three-dimen-

sional to two-dimensional.[54]

Mesoscale Morphology

Traditionally,[4,6] three main types of composites have been

considered to develop when clays and polymers are blended:

(i) a phase-separated microcomposite, obtained when poly-

mer and clay are incompatible and the latter is dispersed

within the matrix as micron-sized agglomerates; (ii) an inter-

calated nanocomposite, formed when, thanks to appropriate

entropic and energetic factors causing strong interfacial inter-

actions between the polymer and the silicate, one or more

extended polymer chains penetrate the clay galleries and

swell the silicate layer stacks without causing complete des-

truction of periodicity; (iii) an exfoliated nanocomposite,

resulting from complete clay delamination and disordered

and uniform dispersion of individual silicate platelets within

the polymer matrix. However, in recent years, several papers

have appeared in which nanocomposites with mixed morpho-

logy (exfoliated and intercalated) have been described and an

increasing number of the examples of composites comprising

exfoliated and intercalated structures together with multilayer

tactoids or clay agglomerates at the micrometre length scale

have been given.[51,55–58] Thus, POM and SEM, whose utili-

sation was normally confined to investigations of the

crystallisation processes and crystalline morphology or to

studies of the anisotropic phases formed in nanocomposites

with high clay loadings,[21,28,32,41,58–63] may provide, toge-

ther with TEM, very useful information in view of a full-scale

description of nanocomposites morphology. In fact, X-ray

diffraction and TEM, which have been traditionally used to

elucidate the morphology of nanocomposites, may some-

times fail, for different reasons, to see micrometre-sized clay

particles eventually surviving, together with exfoliated layers

and/or intercalated stacks, e.g. after insufficient melt com-

pounding, if their population is rarefied.

Table 3. Thermal properties, degree of crystallinity and average spherulite radii of the polymers and of some of their nanocomposites.

Material Tc Tm DHm Crystallinity TcPOM Spherulitesize

8C 8C J�g�1 % 8C mm

EAA1 80.1 94.3 53 18 85.3 2.06EAA1þ 5 phr 15A 80.7 94.6 51 17 88.0 1.80EAA1þ 11 phr 15A 80.0 95.0 50 17 88.1 1.64EAA1þ 18 phr 15A 79.0 94.6 53 18 - -EAA1þ 50 phr 15A 78.6 93.8 54 18 - -EAA2 86.7 103.8 82 28 - -EAA2þ 5 phr 15A 86.7 102.2 84 29 - -EAA2þ 11 phr 15A 86.5 102.3 81 28 - -EAA3 89.3 104.5 73 25 94.0 2.75EAA3þ 5 phr 15A 89.3 105.1 71 24 97.0 2.69EAA3þ 11 phr 15A 88.7 105.5 74 25 96.5 2.28EAA3þ 5 phr 93A 88.8 104.8 74 25 95.0 2.23EAA3þ 5 phr 30B 89.1 104.5 74 25 94.9 2.11EAA3þ 5 phr N848 88.9 104.9 72 24 95.5 2.19EAA3þ 5 phr SE3000 89.8 103.8 71 24 97.1 1.89EAA3þ 5 phr M3C18 89.0 104.6 72 25 96.2 2.13HDAA 117.0 129.7 175 60 124.0 9.26HDAAþ 5 phr 15A 118.3 129.1 165 56 124.4 5.98HDAAþ 11 phr 15A 116.4 130.9 178 61 - -HDAAþ 18 phr 15A 116.1 129.7 185 63 - -HDAAþ 50 phr 15A 116.2 128.9 183 62 - -



In Figure 2, the micrographs taken under crossed

polarisers on molten films of some of the materials studied

in this work are shown. At least five specimens, taken out of

the bulk of each nanocomposite as obtained from melt

compounding, were analysed by POM and the micrograph

with the average morphological aspect was chosen for the

figure. Despite the relatively low magnification, in fact, the

micrographs taken on samples of the same material appea-

red sometimes not fully alike, thus showing that, at least

with some of the employed organoclays e.g., N848 and

M3C18, the dispersion at micrometre scale was not uniform.

Figure 2(a–c) illustrates the evolution of the mesoscale

morphology of the EAA1/15A nanocomposite with 5 phr

organoclay, during the melt compounding process. Some

clay agglomerates up to 50 mm in size, together with much

smaller tactoids of a few microns, are seen in the composite

prepared with a melt compounding time of only 5 min

[Figure 2(a)]. After 10 min [Figure 2(b)], and even more

after 45 min [Figure 2(c)], the large agglomerates disappear

and just a few small tactoids are still visible. Thus, the 10-

min compounding time, normally employed for preparing

the nanocomposites studied in this work, proved sufficient

to break the large agglomerates. It may be interesting to

observe that, as it will be shown later, the tactoids at the

micrometre scale still visible by POM in the EAA1 nano-

composite prepared with 5 phr of 15A and 10 min blending

[Figure 2(b)] are too few or too small to be revealed by

WAXD and, for statistical reasons, even by high-resolution

TEM.

The other micrographs in Figure 2(d–i) illustrate the

behaviour of some of the other organoclays. A comparison

of Figure 2(b), 2(d) and 2(e) suggests that the higher the

excess of the M2(HT)2 surfactant used for modifying the

MMT, the easier is the organoclay dispersion. As demons-

trated by the micrograph in Figure 2(i), the worst dispersion

behaviour was observed for the N848 organoclay, modified

with octadecylamine. The remaining organoclays (the

micrograph of the M3C18 nanocomposite is not shown in

Figure 2) displayed intermediate behaviour. In summary,

the different organoclays employed in this work for the

preparation of nanocomposites with an EAA1 matrix can

be ranked in the following order of increasing dispersion

effectiveness: N848<M3C18 < 30B < SE3000 < 93A�20A < 15A� 6A.

A few POM observations made on films of nanocom-

posites prepared with EAA2, EAA3 and HDAA as the

matrix polymers (not shown in Figure 2) confirmed this

trend.

Cryo-fractured surfaces of the composites were also

analysed by SEM to gain additional information on meso-

scale morphology (Figure 3). The nanocomposites display-

ing just a few small birefringent spots in their POM

micrographs failed to show the presence of clay agglom-

erates when viewed by SEM [cf. Figure 3(a)]. This was not

surprising because SEM can only see the clay particles on

Figure 1. Polarised optical micrographs of (a) EAA1, (b) EAA1with 5 phr of 15A, (c) HDAA and (d) HDAAwith 5 phr of 15A. Thesamples were cooled from 190 8C to room temperature at1 8C�min�1. The insets show the Hv patterns.



the fractured surface, whereas POM observations were

made in transmission on �100 mm thick films. However,

when the number of clay agglomerates or micron-sized

tactoids is sufficiently high, the probability that they re-

main exposed on the cryo-fractured surface increases. An

example is given in Figure 3(b) and 3(c), where the SEM

micrographs of the cryo-fractured surface of the EAA1

nanocomposite with 5 phr of 30B are shown. According to

the Cloisite1 30B data sheet, this organoclay contains,

together with 90 vol.-% of particles less than 13 mm in

size, about 10 vol.-% of larger agglomerates. Actually, in

Figure 3(b) and 3(c), several clay particles may be clearly

seen having dimensions up to several micrometres.

Morphology of the Nanocomposites with theM2(HT)2-Modified Clays

In a previous paper,[51] a preliminary discussion on the

morphology of some of the nanocomposites studied in this

work has been made. Through X-ray and TEM character-

isations, it was found that the addition (5 phr) of one of the

organoclays modified with the M2(HT)2 surfactant (6A and

15A) leads to nanocomposites containing a large number of

exfoliated silicate layers and stacks of 2–4 platelets with an

average periodicity of about 4 nm, for the branched EAA

matrices, and about 3 nm for the linear HDAA. It was also

found that solvent extraction of the excess of surfactant

from 15A does not lead to appreciable changes in the X-ray

patterns of the composites, suggesting that a mechanism of

polymer intercalation with concurrent counter-diffusion of

excess surfactant takes place during melt compounding.

The morphology studies carried out in this work confirm

these conclusions. However, the question of whether the

exfoliated/intercalated morphology revealed by WAXD

and TEM for the EAA/15A nanocomposites corresponds

to thermodynamic equilibrium or is the result of, e.g. insuf-

ficient kneading time, is still open. The kinetics of nano-

composites formation by melt compounding has been

studied quite extensively[21,57,64–66] and it has been found

that break-up of clay agglomerates may be considerably

expedited by shear forces, whereas exfoliation of silicate

layer stacks does probably occur through a peeling-off

mechanism driven mainly by chemical compatibility bet-

ween clay and polymer. In Figure 4, the X-ray patterns of

three samples of EAA1 nanocomposites with 5 phr of 15A

prepared with 5, 10 and 45 min blending times, and of a

fourth sample obtained by annealing for 2 h at 190 8C that

prepared with 10 min blending are shown. The POM

Figure 2. Polarised optical micrographs of molten composites containing 5 phr of different organoclays,taken at 190 8C; (a) EAA1/15A, prepared by 5 min compounding, (b) EAA1/15A, prepared by 10 mincompounding, (c) EAA1/15A, prepared by 45 min compounding, (d) EAA1/6A, (e) EAA1/20A, (f) EAA1/93A, (g) EAA1/SE3000, (h) EAA1/30B and (i) EAA1/N848.



characterisation described above had demonstrated that the

sample prepared with the shortest kneading time (5 min)

still contains large-sized agglomerates. However, the X-ray

analysis illustrated in Figure 4 fails to provide any clear

indication of the presence of unintercalated clay particles in

this sample. Moreover, the diffractograms in Figure 4 show

that the intercalated stacks revealed by the basal reflection

at 2yffi 2.28 do not disappear when kneading is continued

for 45 min or the sample is annealed for 2 h at 190 8C(the latter treatment apparently causes only a slight shift

of the peak towards wider angles, probably due to partial

decomposition). This suggests that, in the conditions of

melt compounding employed in this work, a thermody-

namic equilibrium between exfoliated and intercalated clay

layers is reached quite rapidly and is practically independ-

ent of the thermo-mechanical history, being solely deter-

mined by the interfacial interactions between the polymer

and the organically modified clay surface.

The effect of higher clay loadings, up to 50 phr, was

studied to gain information on the possibility of preparing

concentrated EAA/clay master blends to be used as com-

patibilisers for PE-based nanocomposites. The dependence

of morphology on clay concentration is illustrated in

Figure 5–7, where the X-ray patterns of the nanocompo-

sites EAA1/20A, EAA1/15A and HDAA/15A are shown. A

comparison of Figure 5 and 6 confirms quite clearly that the

EAA1 nanocomposites prepared from 20A and 15A

possess very similar morphology. In fact, for both materials,

an intercalated structure characterised by a d001 spacing of

about 4 nm, which is independent both of the clay concen-

tration and the original basal spacing of the organoclay

(2.53 and 3.10 nm, respectively, for 20A and 15A), is

revealed by X-ray analysis. Moreover, the patterns display a

dramatic increase in intensity of the basal reflection and of

the higher order peaks as the amount of clay is increased.

This supports the view that the morphology of these

nanocomposites is only dependent on the chemical struc-

ture of the copolymer and the surfactant used for clay

modification, whereas the excess of surfactant initially pre-

sent in the 15A organoclay is probably expelled during melt

compounding. However, a deeper inspection of the low-

angle reflections of these two sets of nanocomposites

Figure 3. SEM micrographs of (a) EAA1 with 5 phr 15A, (b) and(c) EAA1 with 5 phr 30B.

10987654321

2θ (deg)In

tens

ity (

a.u.

)

a)

d 001 = 4.0 nm

d)

c)

b)

Figure 4. X-ray patterns of EAA1 nanocomposites with 5 phr15A prepared by (a) 5 min compounding, (b) 10 min compound-ing, (c) 45 min compounding and (d) 10 min compounding and 2 hannealing at 190 8C.



indicates that for the 15A nanocomposites they are

considerably sharper, suggesting that the structural order

and the coherence length of the intercalated tactoids is

improved if polymer intercalation takes place in the pre-

sence of an excess of surfactant.

The X-ray patterns of the HDAA/15A nanocomposites

shown in Figure 7 confirm that the spacing of the inter-

calated structure is invariant with clay content. As already

pointed out,[51] for this polymer too the average periodicity

was found to be independent of whether 20A, 15A or 6A

were used as nanofillers. The appreciably smaller interlayer

spacing of the HDAA nanocomposites (about 3 nm) can

tentatively be attributed to the lesser lateral bulkiness of the

linear chains of this copolymer which causes smaller

expansion of the silicate layer stacks. Interestingly, the

shape of the low-angle reflections shown in Figure 5 and 7

suggests that the intercalated clay tactoids of the HDAA

nanocomposites are less ordered than those in the corres-

ponding EAA1 nanocomposites.

The crystalline structure of the pure polymer matrices is

revealed, in the wide angle region of the spectra shown in

Figure 5–7, by the typical 110 and 200 reflections of PE at

about 2y¼ 21.3 and 23.58, respectively. These reflections

are considerably more intense for HDAA than for the

branched EAA1, as expected on the basis of the higher

crystallinity of the former copolymer (cf. Table 3). There-

fore, the effect of increasing the organoclay amount, though

being qualitatively similar for the two polymers, is more

evident for the set of HDAA nanocomposites (Figure 7). It

may be observed that the position of the two reflections on

the 2y scale does not change, and this means that the clay

addition does not alter the crystalline structure of the

polymers. However, the dramatic reduction of the ratio I110/

I200 of the uncorrected intensities of the reflections, which is

seen to occur as the clay loading is increased, demonstrates

302520151050

2θ (deg)

Inte

nsity

(a.

u.)

EAA1 + x phr 20A

x = 50

x = 18

x = 11

x = 5

x = 0

d 001 = 4.0 nm

d 001 = 2.5 nm 20A

Figure 5. X-ray patterns of 20A and EAA1/20A nanocompo-sites with different clay contents.

302520151050

2θ (deg)

Inte

nsity

(a.

u.)

EAA1 + x phr 15A

x = 50

x = 18

x = 11

x = 5

x = 2

x = 0

d 001 = 4.0 nm

Figure 6. X-ray patterns of EAA1/15A nanocomposites withdifferent clay contents.

302520151050

2θ (deg)In

tens

ity (

a.u.

)

HDAA + x phr 15A

x = 50

x = 18

x = 11

x = 5

x = 0

d 001 = 3.2 nm

Figure 7. X-ray patterns HDAA/15A nanocomposites withdifferent clay contents.



that the polymer lamellae are strongly oriented parallel to

the faces of the compression moulded discs. This effect was

also observed for nanocomposites comprising different

functionalised PE matrices and was shown to be partic-

ularly evident for highly exfoliated samples cooled slowly

from the melt.[67,68] WAXD and SAXS studies carried out

on quenched and slowly cooled nanocomposites of HDPE-

g-MA with 5 and 11 phr of 15A showed[69,70] that the clay

platelets are oriented preferentially in the compression

moulding plane and cause a confined polymer crystallisa-

tion leading to highly oriented crystallites.

In order to see if there were any significant differences

between the morphologies of the skin and the core of

compression moulded nanocomposite samples, some of the

discs were milled down to a thickness of 1 mm, as described

in the Experimental Part, and the X-ray patterns were taken

on both the core and the outer surfaces. The results are

presented in Figure 8 for some of the EAA1/15A nano-

composites. It is clearly observed that the skin and core

patterns are always very similar and that the basal ref-

lections of the intercalated clay stacks are seen at the same

2y angles. In some instances, however, the intensities of the

clay reflections were found to be slightly lower in the core,

suggesting that, occasionally, the clay platelets may either

gather or become more oriented in the skin.

The TEM characterisation of the nanocomposites with

the M2(HT)2 clays carried out in this work confirmed the

information of the X-ray study and was also in agreement

with the preliminary conclusions drawn in our previous

paper.[51] In particular, the low magnification (7–20 K)

TEM micrographs of these materials, not shown, indicate

that the organoclays are dispersed at the nanometre level

within the matrix polymers, thus confirming the POM

results discussed above. The higher magnification (50–

85 K) images of the EAA1 and HDAA nanocomposites

with 5 and 11 phr of 15A, shown in Figure 9, display

individual clay platelets accompanied by stacks of a few

layers. In most specimens taken from the nanocomposites

as received from the Brabender mixer, almost no preferred

orientation of the platelets and the tactoids could be seen

[Figure 9(c)and 9(d)]. In the compression moulded WAXD

specimens, on the contrary, the clay layers were more or less

oriented in a preferred direction [Figure 9(a)]. The gallery

heights d2 and the spacings d1 of the tactoids were measured

on the digitised micrographs with an Image-Pro1 Plus

software. The average figures obtained (on several tens of

measurements) for the EAA1/15A nanocomposites were in

good agreement with the WAXD data: d1¼ 4.1–4.3 nm and

d2¼ 2.3–2.5 nm. The scattering of the data was fairly

strong, also because of the observed poor parallelism of the

platelets within each sheaf: individual d1 measurements

ranged from about 3 to over 5 nm. The d1 spacing measured

for the HDAA/15A nanocomposites varied between 3.6 and

4.6 nm, with a mean value of about 4.1 nm. Thus, as com-

pared to the d001 figures measured from the X-ray patterns

(about 3.2 nm), higher values were found from the TEM

images of these nanocomposites, and this can probably be

ascribed to the uneven morphology of the matrix which

hampers a clear identification of the edges of individual

platelets. Furthermore, the preparation of thin sections of

these materials required more severe conditions; indeed, it

was necessary to go down to �145 8C to obtain satisfactory

sections, whereas a temperature of �130 8C proved suffi-

cient for the EAA1-based materials.

Although the structure of the organic modifier of SE3000

is undisclosed, we briefly discuss the behaviour of the

relevant EAA1 nanocomposites in this paragraph because

their morphology is very similar to that of the nano-

composites containing the M2(HT)2 clays. The similarity is

clearly demonstrated by the WAXD patterns in Figure 10,

compared with e.g., those in Figure 6, as well as by the

TEM micrographs of the EAA1/SE3000 nanocomposites

(not shown), which are practically identical to those in

Figure 9(a) and 9(b). The observation that the d001 spacings

measured for all these nanocomposites are very similar,

although the organic content of SE3000 is considerably

larger than those of 20A, 15A and 6A, made us specu-

late that SE3000 might contain an unbounded organic

compound (excess surfactant, or else) which is expelled,

during the compounding process, through a counter-flow

10987654321

2θ (deg)

Inte

nsity

(a.

u.)

a)b)

c)d)

e)f)

Figure 8. X-ray patterns of EAA1/15A nanocomposites withdifferent clay contents; (a) and (b) 5 phr, (c) and (d) 11 phr, (e) and(f) 50 phr. The (a), (c) and (e) patterns were taken on the core, andthe (b), (d) and (f) patterns were taken on the skin of compressionmoulded samples.



mechanism similar to that suggested for the M2(HT)2 clays.

The FTIR spectrum of SE3000 is compared in Figure 11

with that of 20A. The main difference is represented by the

two absorbance peaks at about 1 245 and 1 740 cm�1,

visible in the spectrum of SE3000. Different solvents were

used to try to extract the unbounded organic material from

this organoclay. Ethyl alcohol, which proved effective for

the extraction of the excess surfactant from 15A,[51] showed

poor efficiency with SE3000. In contrast, a deep extraction

was obtained with chloroform, but the clay powder re-

covery, accomplished by centrifugation, was incomplete as

�20% of it remained in suspension. The organic content of

the residue, measured by TGA, was of 32.3%. Tetrahy-

drofuran (THF) showed intermediate extraction power and

allowed to isolate the extracted organic material in fairly

pure form by evaporation of the clear solution obtained by

filtration. The inset of Figure 11 contains the region bet-

ween 1 150 and 1 800 cm�1 of the FTIR spectra of 20A

(curve a), of the residue of the extraction with chloroform

(curve b) and of the organic material recovered from the

THF solution (curve c). These spectra demonstrate that the

two absorbance peaks at about 1 245 and 1 740 cm�1 of the

SE3000 are in fact due to an unbounded organic material

and that the latter is different from the surfactant used for

modifying the MMT by cation exchange. The FTIR spect-

rum of the THF extract (curve c) resembles closely that of

an EVA sample with 14 wt.-% vinyl acetate (curve d). The

similarity of the spectra (a) and (b) in the inset of Figure 11,

as well as that of the diffraction patterns of the nanocom-

posites prepared from SE300 and the M2(HT)2 clays, seems

to indicate that all these clays contain the same quaternary

ammonium ion modifier and that they differ from each other

for the organic substance employed to expand their struc-

ture: excess surfactant for the 6A and 15A, and a low molar

mass polymer with a structure similar to EVA, for the

SE3000. To further confirm that the unbounded organic

material present in the SE3000 is practically unessential in

view of the preparation of EAA1-based nanocomposites,

Figure 9. TEM micrographs of the nanocomposites of EAA1, (a) and (b), and HDAA, (c) and (d), with5 phr, (a) and (c), and 11 phr, (b) and (d), of 15A.



we used the residue of the extraction of SE3000 with

chloroform to prepare a new composite with EAA1 and

found that the relevant X-ray pattern (not shown for the sake

of brevity) was in fact perfectly similar to that of the EAA1/

SE3000 nanocomposite (Figure 10).

Morphology of the Nanocomposites Loaded withDifferent Organoclays

The X-ray patterns of the nanocomposites of EAA1 with

5 phr of 93A, 30B, N848 and M3C18 are presented in

Figure 12, together with those of the clays, and the TEM

micrographs are shown in Figure 13. The patterns of all

composites show distinct peaks indicative of the presence of

unexfoliated clay tactoids. However, the position of these

peaks changes markedly with the structure of the clay

modifier and differs from that found for the M2(HT)2

nanocomposites described above. The TEM images in

Figure 13 also point to appreciable differences in the degree

of dispersion of the four clays, and demonstrate that the

extent of exfoliation is, for all of them, considerably lower

than that observed for the M2(HT)2 clays.

The effect of substitution of one of the methyl groups of

the M2(HT)2 surfactant of the 20A with hydrogen is

illustrated by the WAXD pattern and the TEM micrograph

of the EAA1/93A composite shown in Figure 12 and 13(a).

The morphologies of the EAA1 composites with 93A and

20A are qualitatively similar: individual exfoliated clay

layers are accompanied by fairly small tactoids. However,

the level of exfoliation of 93A is considerably lower and the

degree of expansion of the layer stacks (d001ffi 2.8 nm) is

definitely smaller. The different behaviour of the two clays

may be attributed to the lower shielding ability of the

hydrogen atom of the tertiary ammonium ion of 93A, as

compared to the methyl group of the 20A, and to the conse-

quently lower compatibility of the former clay for the

largely aliphatic EAA1 copolymer.

The three other clays were organically modified with

surfactants containing only one long alkyl group. One-

tailed organoclays such as 30B and M3(HT) (the latter is

very similar to the M3C18 used here) had been shown by

others[25,48] to give rise to significant shifts of the basal

reflection to larger 2y upon mixing with LLDPE or with a

PE ionomer and the reason for this shift was not fully

understood. The earlier interpretation[71] in terms of degra-

dation with loss of surfactant caused by the high temper-

atures (250 8C) used to process nylon-based nanocomposites

was subsequently questioned, as the same behaviour was

observed in preparations carried out at 200 8C.[25] The thermal

decomposition of the alkylammonium ions of the 30B clay

taking place at temperatures of 220–250 8C was also studied

by Dharaiya and Jana,[72] who found that the shift of the X-ray

diffraction peak to wider angles is accompanied by a reduction

of the height of the absorbance band due to –OH stretching in

the region 3 000–3 500 cm�1 of the FTIR spectrum, and by a

significant decrease of the surface polarity, as determined by

contact angle measurements. Tidjani et al.[73] found that a shift

of the 30B peak towards higher 2y angles was only observed

when the clay was melt-blended in air with PP-g-MA, whereas

such shift did not occur when compounding was carried out in

a nitrogen flow, and attributed the effect to an oxidation

process.

The diffractograms in Figure 12 demonstrate that the

peak shift behaviour has in fact been observed also in this

work for the 30B and M3C18 clays. In fact, the d001 spacing

of 30B is lowered from 1.84 to 1.55 nm, and that of M3C18

from 1.96 to 1.52 nm, after melt compounding with EAA1.

As described in the Experimental Part, we carried out the

melt compounding in a nitrogen flow and the processing

temperatures were normally rather low (120–150 8C), al-

though higher temperatures (190 8C) were used for the

preparation of X-ray specimens by compression moulding.

In order to see whether the latter high-temperature treat-

ment could be responsible for the observed effect, addi-

tional specimens of both the EAA1/30B and EAA1/M3C18

nanocomposites were prepared by moulding at 120 8Ceither the products prepared with the usual compounding

procedure (120 8C and 60 rpm) and new samples prepared at

100 8C and 30 rpm. The EAA1/30B samples displayed X-

ray patterns identical to that shown in Figure 12, thus

demonstrating that the structural tightening of the 30B clay

takes place even under very mild processing conditions. In

contrast, in the X-ray spectra of the samples of EAA1/

M3C18 the peak was approximately at the same position

(2yffi 4.58) of the virgin organoclay, suggesting that M3C18

302520151050

2θ (deg)

Inte

nsity

(a.

u.)

EAA1 + x phr SE3000

x = 50

x = 18

x = 11

x = 5

x = 0

d 001 = 4.2 nm

d 001 = 3.5 nm SE3000

Figure 10. X-ray patterns of SE3000 and EAA1/SE3000 nano-composites with different clay contents.



is appreciably more stable than 30B and undergoes the

spacing shrinkage only when a high temperature (190 8C) is

used in the preparation of the compression moulded X-ray

specimens.

The thermal stability of the two clays was studied further

by treating virgin clay samples in a vacuum oven at 160 and

190 8C (i.e., the maximum temperatures possibly reached

during melt compounding and compression moulding,

respectively) and analysing them after different treatment

times. The results are shown in Figure 14 for the 30B

samples. The shape of the peak recorded after increasing

times suggests that a new reflection centred at about 2y¼5.78 (i.e., the position observed for the composites) appears

and gradually replaces that of the virgin 30 B clay at

2y¼ 4.88. However, when the 30B is thermally treated in a

vacuum oven the process appears to be much slower than

that taking place during melt-blending with the EAA1. In

fact, after 3 h at 160 8C, the new reflection is seen as a

shoulder on the wide-angle side of the original peak, and it

is only after a 3-h treatment at 190 8C that the shape of the

diffraction pattern actually becomes similar to that of the

EAA1/30B composite. Surprisingly, however, the TGA

weight loss of the 30B sample held 3 h at 190 8C was found

to be only 4% lower than that of the virgin clay, and the

FTIR spectrum revealed just a slight reduction of intensity

of the hydroxyl absorbance at 3 000–3 500 cm�1. As for the

M3C18 clay, on the contrary, no peak shift was observed

even after 3 h at 190 8C under vacuum, confirming that the

thermal stability of M3C18 is considerably higher than that

of 30B. The present results do not allow a definite under-

standing of the mechanism of the structural tightening of

these two clays. However, they show that, at any temper-

ature, the process is amazingly accelerated when the clays

are dispersed within a polymer matrix. To our knowledge,

this had not been recognised before.

The TEM micrographs of the EAA1/30B and EAA1/

M3C18 composites shown in Figure 13(b) and 13(c) demon-

strate that these are actually microcomposites. In agreement

5001000150020002500300035004000

Wavenumber (cm-1)

Tra

nsm

ittan

ce

1150135015501750

a

b

c

d

20A

SE3000

1245

1740

Figure 11. FTIR spectra of 20A and SE3000. The inset shows the 1 150–1 800 cm�1 regionof the FTIR spectra of (a) 20A, (b) residue of the extraction of SE3000 with chloroform,(c) organic material extracted from SE3000 with THF and (d) EVAwith 14 wt.-% vinyl acetate.

10987654321

2θ (deg)

Inte

nsity

(a.

u.)

93A

EAA1/93A

EAA1/30B

30B

EAA1/N848

N848

M3C18

EAA1/M3C18

Figure 12. X-ray patterns of the composites of EAA1 with 5 phrof 93A, 30B, N848 and M3C18, compared with those of the clays.



with the POM and SEM observations described previously, in

fact, the TEM images display a great number of fairly large

clay aggregates. The present results differ from those found

by Shah et al.[48] for the nanocomposites of an EMAA iono-

mer with 30B. In fact, despite the observed shift of the X-ray

peak to lower d-spacings, these authors[48] found a partially

exfoliated morphology for the Surlyn1/30B nanocomposite,

revealed by TEM, as opposed to the unexfoliated structure of

similar composites based on M3T and M3(HT). Apparently,

more favourable chemical interactions between 30B and the

matrix polymer are granted by the sodium carboxylate rather

than by the free carboxyl groups.

The behaviour of the third one-tailed organoclay (N848)

is far from clear. The TEM analysis of the relevant EAA1

composite, in agreement with the POM and SEM obser-

vations, demonstrates the presence of very large clay aggre-

gates, together with smaller tactoids [Figure 13(d)]. This

might lead to the reasonable conclusion that this material is

a conventional microcomposite with unexfoliated and

unintercalated clay particles poorly dispersed within the

polymer matrix. On the other hand, the WAXD pattern

shown in Figure 12 indicates the presence of a reflection at

2yffi 2.48 which, according to common understanding,

suggests the intercalation of polymer within the clay gal-

leries. The only indication of a shift in the diffraction peak

of the clay to wider angles, similar to that observed for the

other one-tailed clays discussed above, is the weak and

broad band centred at 2yffi 5.38, whose position is different

from that expected for a higher order of the 2yffi 2.48reflection. Thus, on the basis of the X-ray pattern, it seems

that just a fraction of this clay undergoes the structure

tightening observed for the other one-tailed organoclays

and gives rise to the aggregates observed by TEM, while the

remaining fraction is intercalated by the EAA1 chains.

Although we cannot offer either a definite evidence or a

reasonable mechanistic explanation for this hypothetical

behaviour, a circumstantial evidence in favour of it is

represented by the fact that nanocomposites prepared from

EAA1 and an organoclay synthesised in our laboratory[74]

from Cloisite1–Naþ and octadecylamine were found to

Figure 13. TEM micrographs of the composites of EAA1 with 5 phr of (a) 93A, (b) 30B, (c) M3C18

and (d) N848.



display X-ray patterns identical to that shown in Figure 12

for the EAA1/N848.

Rheology and Mechanical Testing

The effect of the addition of 5 or 11 phr of 15A into the four

copolymers on their rheological behaviour is illustrated in

Figure 15, where the dimensionless viscosity of the nano-

composites (i.e., the ratio, at any frequency, of their

viscosity to that of the matrix) is plotted versus frequency.

It may be observed that the viscosity of the nanocomposites

is generally higher than that of the matrix, especially at low

frequencies, and that the difference increases as the clay

concentration is increased. The non-Newtonian behaviour

of the materials is markedly enhanced in the presence of

clay, notably for concentrations higher than 5 phr. This is

typical for filled polymers, including nanocomposites. Only

for the EAA3/15A nanocomposite, Figure 15 shows that the

viscosity ratio becomes slightly lower than unity at frequ-

encies exceeding 100 rad�s�1. This might be due to the

excess of surfactant of the 15A clay, which was shown to

migrate into the matrix during melt-blending, thus playing a

plasticising role which is obviously more pronounced for

this higher molar mass copolymer. The curves in Figure 15

also show that the apparent viscosity ratio is higher for the

EAA1 composite, probably because of the higher AA

concentration in the matrix. The increase in viscosity

caused by the addition of 15A, however, is considerably

stronger for the linear HDAA copolymer. This is demon-

strated even more clearly by the curves of the nanocompo-

sites with 11 phr of 15A, whose low-frequency trend

suggests a pronounced solid-like behaviour. The latter be-

haviour was also clearly demonstrated by the curves of the

storage modulus (not shown) displaying strong slope

changes at low frequencies.

The effect of the different clays studied in this work on

the rheological behaviour of the composites is compared

with that of 15A in Figure 16, where the viscosity ratios

calculated for the EAA1 nanocomposites with 5 phr of clay

are shown. Clearly, none of the clays cause a viscosity

increase comparable with that observed for 15A. This was

not unexpected, considering the morphological character-

istics of the different composites. The modest viscosity

increase of the EAA1/SE3000 nanocomposite, whose

morphology was shown to resemble closely to that of the

M2(HT)2 clays, may probably be accounted for considering

that, as suggested before, an appreciable amount of un-

bounded organic material of the SE3000 clay is supposed to

migrate to the matrix polymer during blending, thus acting

as a plasticiser. The viscosity of the EAA1/30B composite

was found to be lower than that of the matrix in the whole

frequency range investigated. The only possible interpre-

tation we can tentatively offer for this finding is again

connected with matrix polymer plasticisation caused by the

organic material produced by the clay decomposition dis-

cussed above.

The mechanical properties of the pure polymers and of

some of their nanocomposites are reported in Table 4. The

TS and the EB of the three unfilled EAA copolymers are

very similar, whereas the E and the YS of EAA1 are lower

than the corresponding values of EAA2 and EAA3. The

reduced rigidity of EAA1 can be attributed to the higher

content of AA monomer units that lowers the degree of

crystallinity, as reported in Table 3. On the other hand, the

high crystallinity of HDAA is responsible for the higher

rigidity and lower ductility of this material.

The addition of clays enhances the rigidity of all the

matrices, as demonstrated by the increase in E and YS,

without significantly reducing the EB; the TS variations are

in general quite small. Among the nanocomposites with

5 phr of 15A, the strongest increase of E (about 60%) was

displayed by that with an EAA3 matrix. For this material,

the TS was almost unchanged while the EB was only about

15% lower than that of the pure matrix. The strongest

reduction of EB was recorded for the highly crystalline

HDAA/15A nanocomposite. The effect of nanofiller con-

tent on the mechanical properties is also illustrated in

Table 4 with reference to the EAA1 and HDAA nano-

composites with 15A. With both these matrices, an increase

10987654321

2θ (deg)

Inte

nsity

(a.

u.)

30B

160°C 1h

160°C 2h

160°C 3h

190°C 3h

190°C 2h

190°C 1h

EAA1/30B

d 001 = 1.55 nm

d 001 = 1.84 nm

Figure 14. X-ray patterns of 30B (lower curve) and of the EAA1composite with 5 phr of 30B (upper curve). The intermediatecurves are for 30B annealed at 160 or 190 8C for the indicatedtimes.



in nanofiller content leads to an appreciable increase in

modulus, whereas the TS and EB are lowered moderately.

The mechanical properties of the EAA1 composites

containing 5 phr of the different organoclays studied in this

work are also shown in Table 4. For all these materials, the YS

is higher than that of the matrix by 10–20% and the TS is

slightly lower. However, the dependence of these two

properties on the structure of the organic clay modifier is

rather irregular. Actually, in view of the multiplicity of para-

meters that may exert an influence on both morphology and

mechanical properties of these materials, such as interfacial

adhesion between clay and polymer, polymer crystallinity,

decomposition of the clay modifier, polymer plasticisation

due to organic materials released by the clay during proces-

sing etc., the available results are not sufficient for drawing

definite conclusions on the relationships between mechanical

properties and clay structure. The only indication provided by

the data in Table 4 is that the E values of the nanocomposites

are about 20% higher than that of the pure polymer (except for

the nanocomposite with the 93A clay, modified with a ternary

ammonium ion, whose E is about 30% higher) and that this

property seems to be scarcely influenced by the level of dis-

persion of the nanofiller. The latter morphological character-

istic, on the other hand, appears to have some effect, though

1

2

3

4

5

6

7

8

9

10001001010,1

Frequency (rad/s)

Vis

cosi

ty r

atio

EAA1/15A

EAA2/15A

EAA3/15A

HDAA/15A

11 phr

5 phr

Figure 15. Dimensionless viscosity (ratio of the apparent viscosity to that of thematrix polymer) of the EAA1, EAA2, EAA3 and HDAA nanocomposites with 5 and11 phr of 15A.

0,5

1,0

1,5

2,0

2,5

10001001010,1

Frequency (rad/s)

Vis

cosi

ty r

atio

15A93ASE3000N848M3C1830B

EAA1with 5 phr of:

Figure 16. Dimensionless viscosity (ratio of the apparent viscosity to that of thematrix polymer) of the EAA1 composites with 5 phr of different organoclays.



slight, on the EB values that are only lowered by about 10%

for the intercalated nanocomposites with good clay dis-

persion, such as those with 15A and SE3000, whereas they

decrease by about 20% for the other composites.

Conclusion

Ethylene–acrylic acid copolymers have been shown to

interact efficiently with organoclays containing quaternary

ammonium ion modifiers with two methyl groups and two

long alkyl tails (M2HT2). Mixed exfoliated and intercalated

morphologies were observed for these systems and were

shown to correspond to thermodynamic equilibrium. In

fact, an increase in the blending time from 5 to 45 min,

though improving the fragmentation of clay agglomerates,

had no effect on the gallery height of the silicate layer

stacks, as observed by X-ray diffraction. TEM observation

of the nanocomposites with M2HT2 clays revealed a popu-

lation of well-dispersed individual platelets and thin stacks

of a few irregularly spaced silicate layers. The average

spacing, measured from X-ray patterns, was found to

depend on the molecular architecture of the matrix polymer,

rather than on the concentration of AA units and the

molecular weight. The nanocomposites from the branched

EAA copolymers had a d-spacing of about 4 nm, indepen-

dent of that of the original clay (i.e., of the excess of

surfactant used for clay modification). The behaviour of the

linear HDAA graft copolymer is perfectly similar, but the

intercalated silicate stacks have smaller periodicity (about

3 nm), probably because the extended chains of this

copolymer have thinner lateral dimensions. The finding that

the clay stack expansion is not reduced as a result of an

increase in clay loading up to 50 phr suggests that these

EAA copolymers may be used as compatibilisers for PE-

based nanocomposites. Studies are in progress on the

characterisation of nanocomposites produced by dilution of

EAA/M2HT2 and HDAA/M2HT2 concentrated blends with

different PE grades.

Acknowledgements: Thanks are due to Exxon-Mobil Chemicaland Crompton Corp. for kindly providing samples of the polymersused in this work. The research was carried out within the PRIN2004 098403 project funded by the Italian Ministry of Education,University and Research (MIUR). The cooperation among theInstitutions was also supported by NATO (grant CBP.EAP.CLG981257). The authors express their gratitude towards ProfessorGiovanna Costa and Professor Barbara Valenti of the Institute forMacromolecular Studies (ISMac) of C.N.R., Genova, for theiraccurate reading of the text and insightful discussions. We are alsothankful to Dr. Alessandro Sola, Dr. Andrea Buti and Dr. ElenaMameli for their help in preparing the samples and carrying onseveral characterisation experiments.

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Table 4. Mechanical properties of some composites.

Material E YS TS EB

MPa MPa MPa %

EAA1 96 9.5 19.5 460EAA1þ 5 phr 15A 115 11.2 17.9 416EAA1þ 11 phr 15A 138 13.2 15.6 386EAA1þ 5 phr 93A 127 11.2 18.2 382EAA1þ 5 phr SE3000 115 10.3 20.1 416EAA1þ 5 phr N848 118 10.4 19.2 396EAA1þ 5 phr M3C18 117 10.2 17.6 373EAA1þ 5 phr 30B 119 11.5 19.2 385EAA2 112 10.7 18.8 452EAA2þ 5 phr 15A 151 12.2 17.8 405EAA3 117 10.9 19.8 468EAA3þ 5 phr 15A 184 12.9 19.3 395HDAA 585 - 30.3 19.8HDAAþ 5 phr 15A 818 - 29.5 6.7HDAAþ 11 phr 15A 863 - 27.8 4.9



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Organoclay Nanocomposites from Ethylene–Acrylic Acid Copolymers

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