Polymer International Polym Int 57:805–836 (2008)

ReviewNanocomposites based on polyolefins andfunctional thermoplastic materialsFrancesco Ciardelli,1,2∗ Serena Coiai,1,3 Elisa Passaglia,4 Andrea Pucci1 andGiacomo Ruggeri1,2

1Department of Chemistry and Industrial Chemistry, University of Pisa, Via Risorgimento 35, 56126 Pisa, Italy2PolyLab, INFM-CNR, Pisa, Italy3Italian Center for Packaging (CIP), Via delle Industrie 25/8, 30175 Venezia, Italy4CNR-ICCOM Pisa section, c/o Department of Chemistry and Industrial Chemistry, Via Risorgimento 35, 56126 Pisa, Italy

Abstract: Polyolefins are today the most used thermoplastic materials thanks to the high technology andsustainability of the polymerization process, their excellent thermomechanical properties and their goodenvironmental compatibility, including easy recycling. In the last few decades much effort has been devotedworldwide to extend the applications of polyolefins by conferring on them new properties through mixingand blending with different materials. In this latter context, nanocomposites have recently offered new excitingpossibilities. This has been made possible on the basis of the improvement of polyolefin functionalization processeswith the availability of several olefin homo- and copolymers bearing a small (generally less than 1 mol%) amountof backbone grafted polar groups. These are indeed adequate to endow favourable interface interactions with polarmacromolecules and inorganic compounds, leading first to compatible blends and then to microcomposites. Thesuccessful use of nanostructured dispersed materials has opened, on a similar basis, the way to nanocompositesas described in this review. This review provides a broad and updated description of the synthetic routes tonanostructured biphase materials having the typical structural properties of polyolefins (continuous matrix) butshowing enhanced thermomechanical properties, thermostability, lower flammability, lower gas permeability andelectrical and optical properties, thanks to the presence of an extended interphase interaction with very differentnanodispersed species. 2008 Society of Chemical Industry

Keywords: nanocomposites; polyolefins; functionalized polyolefins; layered silicates; carbon nanotubes; noblemetal nanoparticles; thermomechanical properties; flammability; electrooptical properties

INTRODUCTIONSince the word nanotechnology was proposed for thefirst time to describe an approach to new technologiesbased on molecular-scale science, nanocompositesappeared the most direct way to bring the new scienceinto practical use. Now, after only a few (no morethan three) decades, the number of scientific papers,technical achievements and applications is so largethat even simply listing them would take much morespace than that available for this review.

Well aware of the situation, we decided to focusthis review on the most relevant aspects concerningthe development of nanocomposites for structural,functional and optoelectronic applications. To providethe review with better value and adequately criticizedaffordable information, we selected examples wherethe polymer matrix and the dispersed nanophase bothfall in the area of expertise where the authors ofthe present review paper were involved during theirresearch activity.

Accordingly, the polymer matrices in the examplesdiscussed here are based on olefin homopolymers andcopolymers, which are well-defined macromolecularsystems where macromolecular interactions andinterfaces are mainly hydrophobic. The effect ofpolar interactions is, however, also considered inspecific cases where the simple hydrocarbon structurecontains some more or less polar groups in differentamounts deriving from copolymerization or backbonefunctionalization.

The role of these functional groups can then bediscussed according to their molecular structure, con-centration and distribution along the macromoleculesof the polymer matrix to contribute to the understand-ing of the design parameters granting nanocompositesformation and stability. Indeed when referring to theselast two aspects one should keep in mind the two maindistinct routes available to nanocomposite prepara-tion. The most sustainable process is certainly thatbased on the reactive formation of a nanophase duringblending with a polymer. The most popular system

∗ Correspondence to: Francesco Ciardelli, Department of Chemistry and Industrial Chemistry, University of Pisa, Via Risorgimento 35, 56126 Pisa, ItalyE-mail: fciard@dcci.unipi.it(Received 28 September 2007; revised version received 21 November 2007; accepted 26 November 2007)Published online 12 March 2008; DOI: 10.1002/pi.2415

2008 Society of Chemical Industry. Polym Int 0959–8103/2008/$30.00

F Ciardelli et al.

falling in this group is the exfoliation of layered sil-icates, or minerals in general, during mixing withthe polymer, where both polymer/inorganic materialinterfacial interactions and thermomechanical stressinduced by the machine work together to separate thefiller into nanoparticles which are then distributed togive the nanostructured system.

In a similar way, polymerization reactor blend-ing can give exfoliation because of the improvedmonomer/clay interaction and polymerization energyevolution. The formation of nanoparticles through achemical reaction performed in the presence of a poly-mer or the blending of preformed stable nanoparticleswith a polymer will be examined in this paper andrelevant examples described in detail. Thus we thinkthat the examples selected can provide a good cover-age as far as the synthetic procedure and the role ofnanophase/polymer interaction are concerned.

There is another aspect of general interest as regardsboth the structure and the properties of nanocompos-ites: this is the nanoshape of the nanoparticles inthe final system. Indeed it is accepted to consider asnanocomposites all those systems where the dispersedphase exhibits a nanosize in at least one dimension.This issue is also discussed here with good coverage,as lamellar silicates give after exfoliation monodimen-sional nanoparticles, while nanotubes are consideredas a bidimensional nanophase (nano section and microlength) and substantially spherical metal particles arenanosized in three dimensions.

Following these general lines this review will reporta substantial number of examples we arbitrarilyconsidered suitable to provide the reader withilluminating information about the main topic,disregarding a full coverage which would be impossiblefor space limitations and merely informative ratherthan formative and impressive.

Before moving further ahead, we wish to apologizeto those whose work is not included in ourreport. Selection was not based on quality butonly arbitrarily made with the objective of providingeffective information in a limited space. Also, catalyticapplications are not considered.

ORGANIC–INORGANIC HYBRIDNANOCOMPOSITESThe inorganic componentOrganic–inorganic nanometric hybrids from layeredinorganic crystals refer generally to a particle-filledmaterial where nanoparticles have only one dimensionin the nanometre range. These fillers usually existin the form of sheets of one to a few (1–3)nanometres thick, and hundreds and even severalthousand nanometres long, thus with very high aspectratios. Layered inorganic crystals can be considered asactive fillers able to swell the interlayer spaces (underspecific conditions) and to modify their structure frommicrosized particles to nanodispersed particles owingto the intercalation/exfoliation processes by polymer

chains (supported by using an intercalating agent).The characteristic properties of nanocompositesthus formed are generally reported as remarkablyimproved when compared to those of pristinepolymers or traditionally microfilled composites.Improvements include higher modulus, increasedstrength and heat resistance and decreased gaspermeability and flammability, all properties reallydepending on the degree of nanodispersion oflayered inorganic substrates modulated by tuning theinterfacial interaction with the polymer matrix.

The use of these particular fillers overcomesthe limited commercial availability of anisotropicnanofillers and the potential health hazards relatedto their handling and inhalation. A wide range ofnatural and synthetic inorganic fillers can be used andare susceptible to intercalation by macromolecules;1

among them layered silicates are really the mostemployed (particularly with polyolefin matrices) dueprobably to their availability (as natural clay) andease of surface treatment, and they have beenstudied for a long time2 and recently reviewed.3,4

Other interesting inorganic layered materials arelayered double hydroxides,5 often synthetic substratescharacterized by similar shape and structure to naturalclay that have more recently appeared in the field ofpolymer nanocomposites.

Layered silicates have a crystal structure consistingof layers made of two tetrahedrally coordinated siliconatoms fused to an edge-shared octahedral sheet ofeither aluminium or magnesium hydroxide. Stackingof the layers leads to a regular van der Waals gapbetween the layers, which is precisely the interlayeror gallery. Isomorphic substitutions within the layers(most commonly Al3+ replaced by Mg2+ or Fe2+,or Mg2+ replaced by Li+) generate negative chargesthat are counterbalanced by alkali and alkaline earthcations situated inside the galleries (Fig. 1). Thistype of layered silicate is thus characterized by amoderate surface charge, directly located on thesurface of layered silicate, and even by a highercharge if tetrahedrally substituted with respect to theoctahedrally substituted compound.6

The charge is not locally constant, but varies fromlayer to layer and its average value over the wholecrystal provides the cation exchange capacity (CEC),that is the capacity to substitute the cations betweenthe layer with other inorganic or organic salts, generallyexpressed as meq (100 g)−1 (Table 1).

Layered double hydroxides (LDHs) are a classof anionic clays whose structure is based onbrucite (Mg(OH)2)-like layers in which some ofthe divalent cations have been substituted bytrivalent ions to form positively charged sheets.LDHs can be represented by the general formula[MII

1−xMIIIx(OH)2]x+ · [(An−)x/n · mH2O], where MII

and MIII are divalent and trivalent metal cations,respectively, and A is the interlayer anion (Fig. 2).8–12

LDHs consist of stacks of positively charged mixedmetal hydroxide layers with hydrated anions between

806 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 1. Comparison of talcum (left) with water-swellable hectorite (right) (Reproduced from Zilg C, Dietsche F, Hoffmann B, Dietrich C andMulhaupt R, Macromol Symp 169:65 (2001) by permission of John Wiley & Sons, Ltd.).7

Table 1. Structure and chemistry of layered silicate most commonly used6

SilicateLocation of

isomorphous substitution FormulaCEC

(meq (100 g)−1)Particle

length (nm)

Montmorillonite (MMT) Octahedral Mx(Al4−xMgx)Si8O20(OH)4 110 100–150Hectorite Octahedral Mx(Mg6−xLix)Si8O20(OH)4 120 200–300Saponite Tetrahedral MxMg6(Si8−xAlx)Si8O20(OH)4 86.6 50–60

M, monovalent cation; x, degree of isomorphous substitution (between 0.5 and 1.3).

the sheets to maintain overall charge neutrality.They are available both as naturally occurring andsynthetic minerals resembling natural hydrotalcite,which has the formula Mg6Al2(OH)16CO3 · 4H2O,and for this reason they are also known as hydrotalcite-like materials.

The most common method for LDH synthesisis coprecipitation of MII and MIII salts (chloride,nitrate, sulfate, carbonate, etc.) from homogeneoussolution.8–13 LDHs grow as hexagonal crystals,where each cation is octahedrally surrounded byhydroxide groups, the octahedral sharing edges11

forming two-dimensional sheets. Therefore, LDHsheets are constituted of one polyhedra-made layer,often corrugated, and for this reason they areconsidered more flexible than other bidimensionalframeworks such as 2:1 layered silicate.11 Most LDHsare binary systems, i.e. with two kinds of metal cationswithin the hydroxide layers, but ternary LDHs have

Figure 2. Generic LDH structure (Reproduced from Leroux F andBesse JP, Chem Mater 13:3507 (2001) by permission of the AmericanChemical Society).11

Figure 3. AEC versus layer charge for some LDH compositions(Reproduced from Leroux F and Besse JP, Chem Mater 13:3507(2001) by permission of the American Chemical Society).11

also been synthesized. Moreover, the identities of MII

and MIII are highly tunable and different type ofcations have been accommodated in the layers of theLDH materials, such as most of the transition metalsof the first row as divalent cations, combined with Fe,Al and Ga as trivalent cations.14,15 Lately it has beenshown that tetravalent cations such as Zr4+ and Sn4+

can also be incorporated into the brucite-like LDHlayers.16 This is of importance because the presence ofinterlayer species is directly related to the net chargevalues and anion exchange capacity (AEC) that forthese systems can vary in a large range between 450and 200 meq (100 g)−1 (Fig. 3).11

Polym Int 57:805–836 (2008) 807DOI: 10.1002/pi

F Ciardelli et al.

The ability of these microparticles to disperse intoindividual layers (nanoparticles) in polyolefin matricesis correlated with the possibility of modulating thesurface chemistry through ion exchange reactionswith well-tuned surfactant agents able to favour theintercalation process, and therefore can be at firstdirectly defined on the basis of CEC and AEC.

Pristine layered inorganic materials, due to theirhydrophilic characteristics, are miscible only withhydrophilic polymers, as in the case of poly(ethyleneoxide) (PEO) and poly(vinyl alcohol) (PVA).17

Pristine LDHs are not suitable for intercalation ofpolymer chains because the intergallery spacing istoo short (originally about 7.6 A corresponding tothe basal peak (003) of the XRD pattern) to allowlarge molecules to penetrate. In order to render thehydrophilic inorganic substrates more organophilic,ion exchange modification with cationic or anionicsurfactants bearing long alkyl chains is generallyperformed. These voluminous organic ions increasethe interlayer distance, lower the surface energy byimproving the wetting characteristic of the layers andmake intercalation with apolar polymers possible.

Exfoliation of layered host crystal tonanostructured polymeric materialsGeneral aspectsThree main composite morphologies1 can be obtainedby mixing polymer chains and layered inorganic hostcrystal: (a) microcomposites where polymer moleculesare not able to penetrate within the galleries, andthe interlayer distance remains unchanged as wellas the order of the stacked layers; (b) intercalatednanocomposites in which a single or more polymerchains are intercalated between the host layersand a repeat distance is expanded, but only to a

limited extent (usually of the order of 20–80 A); and(c) exfoliated or delaminated nanocomposites where theorder of the stacked layers is completely destroyed andthe single layers are uniformly dispersed in the polymermatrix. The separation distance between the layers istypically 10 nm or more apart, depending also on thecontent of the inorganic materials and the aspect ratio.

In the scheme successively proposed by Vaia18

and later partially re-proposed by Schadler19 theintercalated and the exfoliated structures are dividedinto ordered and disordered substructures dependingon the capacity of the layers to swell in the adoptedconditions, and the maximum separation is dependenton the volume fraction of separate layers. Also there isan intermediate morphology termed partially exfoliatedassociated with dispersed exfoliated layers togetherwith small stacks of intercalated layers. Other similarclassification schemes refer directly to the distancebetween the layers to define the ordered exfoliatednanocomposites and so to the gallery volume occupiedby the polymer chains (Fig. 4).20

Again Okamoto and co-workers21 define floccu-lated nanocomposites as materials where intercalatedstacked silicate layers can aggregate due to the hydrox-ylated edge–edge interactions between adjacent layers(Fig. 5).

In all the nanostructures (or sub-nanostructures)the polymer/inorganic interface surface is very largeand can be considered the major component of thepolymer/nanofiller mixture.

In addition to their potential applications, nanocom-posites are unique model systems to study the staticsand dynamics of polymers in confined environments.The local and global conformations of the polymerswithin the host galleries dramatically differ from thoseobserved in the bulk not only due to the confinement

Figure 4. Morphological structures and substructures of nanocomposites from layered silicate (Reproduced from Schadler LS. Polymer-based andpolymer-filled nanocomposites. In Nanocomposite Science and Technology, ed. by Ajayan PM, Schadler LS and Braun PV, Wiley-VCH, Weinheim,pp 77 (2003) by permission of John Wiley & Sons, Ltd.).19

808 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 5. Morphological structures of nanocomposites from layered host crystal (Reproduced from Ray SS, Okamoto K, Okamoto M,Macromolecules 36:2355 (2003) by permission of the American Chemical Society).21

of polymer chains but also due to the specific poly-mer–surface interactions, normally not observed inthe bulk6,22,23 and that directly refer to the interfaceregion.

Modification of layered inorganic crystalFor the very broad class of polyolefin nanocomposites,modification with ammonium cations characterizedby long alkyl chains (C14–C18) (to obtain theso-called organophilic layered silicate, OLS) is themost common way of obtaining the hydrophobicinteractions at the interface necessary to favour theintercalation of the aliphatic macromolecules. Thismodification produces first a reduction of surface freeenergy, thus leading to weaker forces between thelayers, as recently confirmed.24

It is generally accepted that the extent of swellingdepends on the length of the alkyl chain andthe CEC of the clay. For example the interlayerdistance of fluoromica modified with protonated n-alkylammonium salts increases with increasing lengthof the n-alkyl chain.7 Early studies rationalized theseresults by considering that the alkyl chains pack in thelayered structures forming mono- or bilayers where thechain axes are parallel to the clay platelets or radiateaway from the surface forming a precise angle.25

Such structures are based on the assumption ofa nearly all-trans conformation for the long alkylchains in the organoclay (OLS). Recent Fouriertransform infrared (FTIR) and NMR studies showedthat both trans and gauche conformations exist, whichsuggests more random packing arrangements.26–28

The extent of swelling (d-spacing increasing) dependson the mass of organic material in the gallery permass of clay (montmorillonite in this case)3 and theinterlayer distance increases linearly with the massratio of intercalated surfactant (thus by increasingboth chain length and CEC values). The packingof surfactant alkyls is disordered, but the density ofthe organic materials in the galleries is somewhathigher than would be expected for a correspondingbulk liquid. The head groups are essentially tetheredto the clay surface and the long alkyl chains tendto arrange in bilayer and trilayer structures adoptingpredominantly trans conformations although there areextensive gauche conformations. In the presence of OHgroups the surfactants are packed in an even densermolecular structure as compared with those withoutthese functional groups. This effect is attributed to

hydrogen bonding of –OH groups with the oxygen ofthe clay surface.

New in-depth studies of the orientation of surfactantalkyl chains between the layers evidence the presenceof ammonium salt not truly bonded to the nanoplateletsurface (an excess with respect to the CEC) which isnot located in the galleries but among the particles,24

and this negatively affects the mechanical and flameretardant properties of the nanocomposites.29

The exchange capacity of LDHs is quite elevated(which corresponds to higher layer charge density)rendering more difficult the formation of the nanocom-posite: a high AEC corresponds to the presence oflayers tightly stacked via the attractive forces withthe interlayer anions filling the gallery, and this isunfavourable for either ion exchange or exfoliation.This may explain the relatively small number of LDHnanocomposites reported in the literature.

In order to make such layered materials suit-able for intercalation and ultimately exfoliation bypolymer molecules, water-soluble sulfonate, sulfate,carboxylate10,30 and phosphonate11 organic dyes havebeen intercalated in LDHs by following four generalapproaches (Table 2):

(i) anion exchange of a LDH precursor;(ii) direct synthesis by coprecipitation;(iii) rehydration of a calcinated LDH precursor; and(iv) thermal reaction.

A common problem with all these methods is thatwhen using anions other than carbonate it is importantto avoid contamination from CO2 since the carbonateanion is readily incorporated and tenaciously heldin the interlayer. Consequently, decarbonated anddeionized water is often used and exposure of thereacting material to the atmosphere is kept to aminimum.

The interlayer spacing and structure of organic–inorganic hybrid LDHs can be controlled by the choiceof the appropriate anion. In order to improve thecompatibility of LDHs towards polyolefin matricesas well as to increase the interlayer spacing thuspromoting polymer chain intercalation during meltmixing, LDH modification with organic moietieshaving a long aliphatic chain is sought. In thisperspective sulfate anions (CnH2n+1SO4

−) with shortand long chains were incorporated by anion exchangein LDH containing nitrate or chloride anions.41 The

Polym Int 57:805–836 (2008) 809DOI: 10.1002/pi

F Ciardelli et al.

Table 2. Preparation of organo-modified LDHs

Method Characteristics Organo-LDH preparation Reference

Anion exchange Monovalent anions, such as OH−, NO3−,

Cl−, can be easily displaced from LDHinterlayers and a huge variety of organicanions can be intercalated

The modification is carried out bydispersing the LDH in aqueous solutioncontaining an excess of the organicanion

14, 30–32

Coprecipitation The anion that is to be introduced musthave a high affinity for the hydroxidelayers, otherwise the counter anions ofthe metal salts may be incorporated,thus contaminating the LDH. Metalnitrate or chloride salts are commonlyutilized because of the low selectivity ofLDHs towards these anions

MII and MIII hydroxide layers are nucleatedand grown from aqueous solutionscontaining the anion that is to beincorporated into the LDH. The methodis based on the addition of metal saltssolution to a base solution containingthe organic anions at a constant pH

11, 14, 33

Rehydration of calcinatedLDH precursor

This method is known as memory effect.The hydrotalcite is firstly transformedinto a mixed magnesium aluminiumoxide by moderate thermal treatmentdue to a partial dehydroxylation ofhydroxide layers. The LDH structure isthen reconstructed either on cooling inair (uptake of carbonate anions) or bysoaking in an aqueous solution

The rehydration is performed in a nitrogenatmosphere to avoid carbonation. Theuse of a swelling agent such as glycerolcan assist the incorporation of theorganic guests

8–13, 34–38

Thermal reaction This procedure was firstly reported in 1994 An intimate mixture of hydrotalcite andorganic acid is heated at a temperature20–30 ◦C above the melting point of theacid. In addition to the correspondingorgano-LDH phase, the thermal reactionproduct often contains an unreactedphase

39, 40

products of the exchange were found to have basalspacing in the range 21.1 (n = 8) to 32.6 A (n = 18)and underwent additional swelling in the presence ofn-alkyl alcohols or n-alkylamines. It was proposed thatthe alkyl sulfate anions adopt a preferred monolayerarrangement between the hydroxide layers and anear perpendicular orientation of the alkyl chains.42

However following washing and drying, a reductionin the interlayer spacing revealed that the chainsadopt a slant angle of 56◦ to the surface of thehydroxide layers. In successive studies it was possibleto obtain interlayer spacings distributed among threemean values: 26, 36 and 47 A10,43 due to differentinterlayer arrangements of the anion. The shortestinterlayer spacing was attributed to a perpendicularmonolayer arrangement, the interlayer spacing of 36 Awas associated with a bilayer arrangement in whichthe anion is tilted at an angle of approximately 40◦to the surface of the hydroxide layers, whereas the47 A spacing was consistent with a vertical end-to-endbilayer arrangement. A further anion widely studied isthe dodecylbenzene sulfonate ion (RSO3

−) which wasdemonstrated to have a high affinity in Mg2Al-LDH,exceeding that of sulfate and approaching thatof carbonate, thus suggesting that non-electrostaticinteractions, such as hydrophobic interactions amongidrocarbon chains, make an important contribution.Even if the binding of a particular anion in a LDHis usually attributed to electrostatic interaction withthe hydroxide layers in the case of organic anions,

especially those with a long straight hydrocarbonchain, hydrophobic interactions among amphiphilicanions may make an essential contribution. In thiscase the most probable anion configuration within theinterlayer spacings is made of two interpenetratinganti-parallel half monolayers, which is consistent withthe observed d value44 (Fig. 6).

The coexistence of alkyl anions in the interlamellardomain would occur not only due to the availablearea, but also due to the stepwise mechanism inwhich the intercalation may occur, with the initialstep consisting of the intercalation of micelles. Theinteraction between the hydrophobic chains of thesurfactant is determinant for the intercalation inLDHs, due to the aggregation into micelles andthe tendency of self-organizing in the interlayerdomain. This feature is probably responsible for thewell-organized structure of the obtained LDHs andthe increase of structural order can be correlatedwith the increase of the chain length of differentsurfactants. Moreover the molecular packing withinthe interlayer gap may allow the molecules to connectto each other, thus forming a polymer as illustratedin the case of styrene sulfonate incorporated intoLDH.45 Another important parameter of comparisonbetween long alkyl chain sulfates and alkylbenzenesulfonates is related to the thermal stability of modifiedLDHs: the breaking of the surfactant head–tail bondwhere the loss of layered structure happens at lowertemperatures (100–200 ◦C) for the sulfate anions

810 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 6. Proposed packing of dodecylbenzene sulfonate in the LDHinterlayer (Reproduced from Xu, ZP, Braterman PS, J Mater Chem13:268-273 (2003) by permission of the Royal Society of Chemistry).44

compared to the sulfonate which maintains the layeredstructure up to higher temperature (ca 400 ◦C). Inaddition the modification of LDHs with long alkylchains was carried out by using monocarboxylic anddicarboxylic aliphatic acids and aromatic acids.10,46

The intercalation reaction does not seem to dependon the alkyl chain length, and it has also beendemonstrated46 that fatty acids such as stearateanions can be intercalated. In the case of aliphaticmonocarboxylic acids CnH2nCO2

− a different carbonnumber leads to different gallery height and anionorientation.39,47,48 The gallery height of the sebacate-LDH, obtained by subtracting the 4.8 A brucitelayer thickness from the interlayer separation, wasclose to the physical length of the sebacate anion.Dimensional analysis suggested, therefore, that thesebacate anion was oriented with its long axisapproximately perpendicular to the hydroxide layersin a monolayer arrangement.

Extruder blendingPolyolefin-OLS nanocomposites. The confinement ofmacromolecules inside galleries results in a decrease in

the overall entropy of the polymer chains that may becompensated by the increased conformational freedomof the surfactant molecules, in less confined statusdue to the separation of the layers and thanks to theinteraction with polymer chains.6,49

During the intercalation of polymer chains threedifferent types of interactions are established orchange their degree of extent due to the presence ofpolymer inside the galleries: the polymer–surfactantinteraction, the surfactant–surface interaction and thepolymer–surface interaction (Fig. 7).

The formation of a stabilized hybrid dependson the interlayer structure of the modified layeredsilicate;50 at the same time the polar interactionbetween the inorganic surface and the polymer playsa very important role. If this interaction is notvery effective, a good dispersion of the particlesmay be reached through the help of strong shearforces during the preparation and the processing ofthe nanocomposite materials; the system, however,remains thermodynamically unstable. This can beobserved, e.g. in the case of isotactic polypropylene(iPP) or linear low-density polyethylene (LLDPE)nanocomposite materials prepared by melt mixingof the polymers with surface-modified natural clay(ion exchanged with dimethydioctadecylammoinumions) using high shear forces. If such a mixture isheated (e.g. during processing) to temperatures abovethe melting temperature, immediately a (partial) re-agglomeration of the particles takes place.51 In othercases, microcomposites (immiscible system)52 or atleast intercalated structures have been obtained bymixing either polyethylene or polypropylene withOLS53 using both static and dynamic (extruder ormixer) melt compounding processes.

To overcome the problem and to establish moreeffective interactions between the surface and thepolymer, synthetic efforts are focused on designingintercalants/dispersants effective at very low levels, forexample as in the case of fluorinated surfactants,54

or to introduce functional groups onto polyolefinchains that can be used as polymers or suitablecompatibilizers like in maleic anhydride (MA)-graftedfunctionalized polyolefins, ammonium-terminated (orOH-terminated) polyolefins and suitable block copoly-mers. Complete reviews, published between 2000and 2003, including more than polyolefin layered

MMT modified by alkyl ammonium saltConfined-intercalated polymer chains

Polymer chains

hh

S

S

Polymer-surfactantinteraction Surfactant surface

interaction

Polymer-surfaceinteraction

Figure 7. Possible situation of intercalated polymer chains. The entropy loss due to the confinement (the macromolecules pass fromthree-dimensional coil to two-dimensional structure) is compensated by the entropy gain of the surfactant chain and by the balance of differentinteractions.

Polym Int 57:805–836 (2008) 811DOI: 10.1002/pi

F Ciardelli et al.

Figure 8. Effect of organophilic and polar interactions on the morphologies of composites (Reproduced from Passaglia E, Sulcis R, Ciardelli F,Malvaldi M, Narducci P, Polym Int 54:1549 (2005) by permission of John Wiley & Sons, Ltd.).56

silicate materials, address all these aspects.1,54,55 Inparticular, it has been shown that the presenceof MA derivatives grafted onto backbone ensuresthe occurrence of effective hydrophilic interactionsbetween the layered silicate surface and the poly-olefin, providing good mechanical and gas barrierproperties.55,56 The possible interactions investigatedfor poly(ethylene-co-propylene) involve the formationof hydrogen bonds between –OH present on the layersurface and the carbonyl groups of the functionalizedpolyolefin (Fig. 8).56 The importance of the presenceof maleate groups is brought out by the fact thatthe same unmodified polymer provides the forma-tion of intercalated nanostructure only if the silicateis swollen in the presence of a surfactant (alkylammo-nium onium).

The dispersion level is strictly dependent on theamount of silicate, and the evolution of different mor-phologies from disordered exfoliated to predominantlyintercalated has been studied in the range 6 to 36 wt%OLS.57

For polyethylene/clay nanocomposite formation,58

an exfoliated morphology may be obtained fora grafting level higher than the critical value ofMA (established as 0.1 wt%) and a number ofmethylene groups in alkylamine chains higher than16. Similar results had been previously reported59–63

showing again a minimal functionalization value ofpolypropylene (PP) chains (oligomer and polymer) toobtain at least an intercalated morphology.

Figure 9. Molecular structure of chain-end-functionalized polyolefin(right) and side-chain-functionalized polyolefin between the layersand related morphologies of nanocomposites (Reproduced fromWang ZM, Nakajimia H, Manias E, Chung TC, Macromolecules36:8919 (2003) by permission of the American Chemical Society).64

The use of ammonium-terminated PP (PP-t-NH3)in comparison with random PP copolymers with–MA, –OH or methylstyrene has been reported:the experimental results demonstrate the advantageof chain-end-functionalized PP with respect to morefunctionalized copolymers. The terminal hydrophilicNH3

+ functional group anchors the PP chains(by cation exchange) on the inorganic surfaces,and the hydrophobic high-molecular-weight andsemicrystalline PP ‘tail’ gives an exfoliated structurewhich is maintained even after further mixing with

812 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

neat PP. In contrast, side-chain-functionalized or PPblock copolymer form multiple contacts with each ofthe clay surfaces, as illustrated in Fig. 9, which notonly results in aligning the polymer chains parallel tothe clay surfaces but also can bridge consecutive clayplatelets promoting intercalated structures.64

Simulation studies have also confirmed thisbehaviour by considering the addition to polymer/claymixtures of end-functionalized chains in compari-son with more functionalized polymer systems. Theincreased polymer–filler attractive interactions maycreate bridges between adjacent silicate layers, leadingto poor intercalation.65,66

Polyethylene (PE) bearing random-size or end-chain dimethylammonium chloride and the diblockcopolymer poly[ethylene-block-(methacrylic acid)](PE-b-PMAA)67 showed that one functional quater-nary ammonium end group is not apparently enough toalter the structure even of organophilic clay (in com-parison with Na montmorillonite), in contrast withresults obtained for PP. Better results were obtainedwith PE bearing more than one ammonium group asside chain as well as by increasing the concentration ofPE-b-PMAA.68,69 Indeed with the random copolymerpoly[ethylene-co-(vinyl alcohol)] 31/69 by weight, astrictly intercalated structure with the largest interlayerdistance was obtained; the exfoliation was hindered bythe ‘bridging effect’ between the layers caused by theestablished strong hydrogen bonds. Poly[ethylene-co-(methacrylic acid)] 89/11 by weight gave a moreeffective dispersion, suggesting no ‘bridging’ of thelayers, due not only to a lower polar comonomercontent, but also to the acidic nature. The blockcopolymer poly[ethylene-block-(ethylene glycol)] (PE-b-PEG) characterized by a very low molecular weightand containing 2.6 ethylene oxide units per molecule,a sort of end-functionalized PE oligomer, generatedexfoliated nanocomposites with high dispersion level.Similar results were also obtained for a maleated PEwith a very low content of grafted MA.68,69

PEO is clayophilic and known to readily intercalateunmodified clays.17,70,71 Thus non-ionic surfactantsbased on PEO blocks, like PE-b-PEO, were used ascompatibilizers and clay modifiers in PP51 and PE72

composites (Fig. 10). Linear PE-b-PEO copolymerswith a short PEO block and an alkyl chain with about30 carbon units (C30) have been successfully usedfor PP nanocomposites starting from both unmodifiedand modified clays.73,74

Regarding the influence of the molecular weight ofcompatibilizers or polymer matrix on the dispersionlevel, general opinion indicates a better exfoliationcapability for low-molecular-weight polymers, particu-larly in the case of PP nanocomposites59,60,75 until alimit value due to a negative effect on processingparameters/conditions.76 Stress shear factors play animportant role in the success of nanocompositeformation; thus the viscosity and the rheologicalproperties of the polymer compatibilizer have to matchmachine requirements/parameters. In fact the most

Figure 10. Schematic of the action of block copolymer for exfoliationof clay sheets within a polymer matrix (Reproduced from Fisher H,Mat Sci Eng C23:763 (2003) by permission of Elsevier Ltd.).51

common extrusion procedure provides the formationof masterbatches68 between the compatibilizers andthe layered silicate (with higher content, from 30 to60 wt%) and successive dispersion in the polyolefinmatrix. Recently increasingly sophisticated extrudertechnology has allowed one to optimize the processingconditions on the basis of characteristic parameterslike barrel temperature and screw profile.77–79

Polyolefin-LDH nanocomposites. There are much fewerreports in the literature about the preparationof polymer-LDH nanocomposites12 compared withlayered silicate materials based on (a) intercalation ofthe monomer followed by polymerization in situ; and(b) direct intercalation of polymer chains from eithersolution or the melt.

The former methodology refers in particular to thepreparation of polyacrylates, polymethacrylates andpolystyrene (PS) LDH-based nanocomposites.80–85

Exfoliated PS-LDH nanocomposites86 were also pre-pared by the solution intercalation method withenhanced thermal stability and applied to prepareLLDPE-LDH nanocomposites by using dodecyl sul-fate (DS)-modified MgAl and ZnAl hydrotalcites87,88

in xylene, yielding mixed intercalated–exfoliatedstructures, even if thermodynamically a large degreeof LDH exfoliation is very difficult to achieve becauseof its high surface charge density.89,90

The preparation of LDH nanocomposites by themelt blending process is certainly more versatileand environmentally friendly from the technologicalpoint of view than the solution intercalation proce-dure. Specifically, organo-modified LDH layers weredispersed in various polymers such as poly(methylmethacrylate),91 polyimide,92 epoxy,93 poly(ethyleneterephthalate),94 nylon,95 PS96 and polyolefins. Afew studies were carried out on the preparation ofPE-LDH nanocomposites principally by using LDHsmodified with long alkyl chains.97 It was in fact shownthat LDH in carbonate form does not exfoliate, andis not able to intercalate PE chains when dispersedinto the molten polymer.98 Sulfonate and stearateanions were preferentially selected for the modifica-tion because of their easy intercalation and formation

Polym Int 57:805–836 (2008) 813DOI: 10.1002/pi

F Ciardelli et al.

Largecrystallites/platelets

Expandedplatelets

Fragmented or distorted plateletsand individual layers

LDH SDBS-LDH Nanocomposites

Figure 11. Nature of dispersion of LDH particles in nanocomposites (platelets containing a multiple number of hydroxide layers are shown indotted circles) (Reproduced from Costa FR, Wagenknecht U, Jehnichen D, Goad MA, Heinrich G, Polymer 47:1649 (2006) by permission of ElsevierLtd.).100

of organic–inorganic hybrids with a large interlayerdistance.99 The data reported in the literature concern-ing the preparation of PE/organo(alkylchain)-LDHnanocomposites are not totally coherent. For instance,a melt-compounding process100–102 in a Brabendermixer of LDPE with LDH nanocomposites usingan organo-LDH modified with sodiumdodecylben-zene sulfonate (SDBS) and MA-grafted high-densitypolyethylene (HDPE-g-MA) as compatibilizer did notshow any significant change in the position of the DBS-LDH (003) XRD basal peak for any LDH loading andno single LDH layers were evidenced by transmissionelectron microscopy (TEM) analysis, which insteadshowed the presence of localized domains formed bythe dispersed particle fragments.

Further investigations showed that the method gavemostly intercalated or flocculated composites withLDH particles dispersed in the form of thin platelets(Fig. 11).

The use of the compatibilizer HDPE-g-MA asmatrix caused the disappearance (for low DBS-LDHconcentrations) of the (003) basal peak and TEManalysis showed for these specific cases the highlyexfoliated nature of the dispersed particles (Fig. 12).However with increasing LDH concentration, thedispersion state changed and the XRD patterns werenewly characterized by the presence of the basalpeak and by weak higher order peaks. Nevertheless,the resulting platelets were more intimately coatedby the polymer phase compared with unmodifiedLLDPE, but showed also a tendency to formstructural association or clusters with increasing LDHloading.

In contrast, complete exfoliation of HDPE/LDHnanocomposites was observed by melt intercala-tion using DS-LDH and controlling its amount.103

Similarly the XRD patterns of medium-densityPE/stearate-ZnAl LDH composites, containingincreasing amounts of filler (5, 10, 15 wt%), obtainedby melt mixing,99 showed the absence of the typi-cal X-ray reflection peaks of the LDH, as a strongindication of the occurrence of exfoliation.104

Figure 12. XRD spectra of HDPE-g-MA/LDH (PBXLDH)nanocomposites (Reproduced from Costa FR, Wagenknecht U,Jehnichen D, Goad MA, Heinrich G, Polymer 47:1649 (2006) bypermission of Elsevier Ltd.).100

A paper concerning the preparation of PP/DBS-ZnAl LDH nanocomposites synthesized via meltintercalation has also been published.105

Reactor blendingIt is essential that the polymerization proceeds, at leastin part, between the layers leading to the formation,or in the limiting case, of an exfoliated system. Forpolyolefin nanocomposites, apparently, this is onlypossible when the interlayer gaps contain active centrescatalyzing olefin polymerization.

A synthetic hectorite with intercalated zirconocenecatalyst,106 after treatment of the silicate withmethylaluminoxane (MAO), gave (by addition ofan excess of MAO and propylene to the inor-ganic–catalyst system) PP nanocomposites withreasonably high activity where PP chains havelow molecular weight (oligomers). Nanocompos-ites of HDPE and LLDPE107 were preparedthrough the polymerization of ethylene and long-chain α-olefin by various catalytic systems inthe presence of previously dispersed organo-modified clay (OLS). Brookart’s single-component

814 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

palladium-based complex and OLS fluorohectorite108

allowed researchers to obtain polyethylene nanocom-posites with exfoliated morphology. Titanium-based Ziegler–Natta catalyst109 intercalated inOLS-montmorillonite (MMT), with an ammo-nium salt bearing –OH functionalities that offersfacile reactive sites for anchoring the catalystin between silicate layers allowed the formationof a nanocomposite system in subsequent ethy-lene polymerization. A very similar approachwas reported by using alkyltriphenylphosphonium-modified MMT/TiCL4/MgCl2 catalytic system forthe polymerization of propylene.110,111

More recently, PE-based nanocomposites have beenprepared using the so-called polymerization fillingtechnique (PFT).112,113 Pristine MMT and hectoriteare first treated with MAO before being contacted by

a titanium-based constrained geometry catalyst. Thismethodology consists of attaching the polymerizationcatalyst to the surface and into the interlayers of thesilicate and polymerizing ethylene in situ. Interestingly,the authors report that the exfoliated morphology is,however, thermodynamically unstable: for sufficientlyfluid PE matrices the exfoliated structure collapsesin the melt (during processing) into a non-regularstructure as ideally depicted in Fig. 13, presumablydue to the lack of active surface–polymer chainsinteractions.

From the examples discussed above, we can notethat the validity of the in situ polymerization approachor PFT has been demonstrated using metallocene,Ziegler–Natta, Brookart type and nickel/palladium-based catalysts.

1) MAO2) Catalyst

3) EthyleneLayeredsilicate

Exfoliatednanocomposites

within PE particles

processing

Collapsed structureof elongated stacks

of silicate

Figure 13. Preparation of PE nanocomposites using the PFT and collapse of the structure owing to processing.

CatMe

Si(CH3)2

Me

ClCl Zr

Cat*P

Ph Ph

Pd

PPh Ph

OTs

O2HOTs

c

e

Na+ Na+Na+

Na+

RNH3Cl

-NaCl +H3N

+H3N

NH3+

NH3+

NH3+

Cat

Cat

Cat

C2H4

Cat*

C2H4/CO

C3H6/C2H4/COa

b

+H3N

+H3NNH3

+

MeMe

NNNi

Br Br

d

+

-

Figure 14. In situ polymerization of polyolefin nanocomposites.114

Polym Int 57:805–836 (2008) 815DOI: 10.1002/pi

F Ciardelli et al.

However, naturally occurring silicate materials arecompletely incompatible with these catalyst systems,and extensive pre-treatment is required before use.Intercalation of large organic molecules also bearingfunctionalities able to react with catalyst or pre-catalyst(activators) serves to widen the gap between thelayers, so as to favour the insertion of bulky tran-sition metal catalysts (Fig. 14(a), catalysts c and d).Since this treatment is generally insufficient to renderthe material catalyst friendly, a large excess of co-catalyst, MAO for instance, is commonly requiredto neutralize traces of water or of acidic groups onthe silicate surface, particularly for the PFT whereunmodified layered silicates are used. On the otherhand, the use of only unmodified clays could beone of the key advantages of the in situ polymeriza-tion methodology. This has indeed been reached114

in the case of polyketone nanocomposites obtainedusing the precursor [(dppp)Pd(OTs)(H2O)]OTs(dppp = 1,3-bis(diphenylphosphino)propane; OTs =p-toluensulfonate; Fig. 14(b), catalyst e) which isknown to generate an effective catalyst for thehomogeneous copolymerization of CO and alkenes(Fig. 14(b)).

Polyketone nanocomposites (copolymers and ter-polymers) with different morphologies (from inter-calated to exfoliated, depending on the amount ofstarting silicate) have been obtained without the needof a previous treatment of Na-MMT to increase theinterlayer distance and to render the environmentmore friendly to catalysts, as shown by the good activ-ity values obtained.

Properties of layered host crystalnanocompositesRecent literature about PE,67–69,74,115–132 PP,64,73,75,

77–79,133–156 syndiotactic polypropylene,157–160

poly(ethylene-co-propylene),56,161,162 polybutadieneand natural rubber163–165 and poly(ethylene-co-vinylacetate)166–169 layered silicate nanocompositesrefers in detail to preparation,116–118,121–123,127–130,132,

136,161,163,164,169 morphology development64,67,69,78,79,

120,133–135,139,140,144,149,150,154,157,162 and thermal,75,

119,126 mechanical,75,76,131,142,143,151,153,160 rheo-logical, 155,156 barrier124,137 and flame resistance115,

126–138,141 properties.

Thermal propertiesEven moderate changes of melting and crystallizationtemperatures as well as degree of crystallinity may havea strong influence on the structure of composites andthereby on their mechanical properties. A detailedstudy concerning the crystallization behaviour ofneat PP-g-MA and its nanocomposites170 indicatedhow clay particles act as nucleating agents forthe crystallization of the matrix, but the lineargrowth rate and overall crystallization rate are notinfluenced significantly by the presence of clay. Insome intercalated PP nanocomposite171 systems afaster crystallization could be noted by calculating

the crystallization rate coefficient (CRC) defined asthe variation in cooling rate required to changethe undercooling of the polymer melt by 1 ◦C, butduring isothermal crystallization kinetics experimentsgenerally no significant interferences were observedbecause of the presence of clay. In other papers172 anincrease of the melting temperature as the content ofmodified clay increases was reported and explainedconsidering an increase in PP lamellar thickness;at the same time the very small decrease of theassociated enthalpy was attributed to the loweringof the total amount of thicker lamellae present in thematerial, suggesting for this nanocomposite a largeramorphous part and more space for chain movement,in agreement with the results of barrier properties.No data or description of morphologies are reportedas backing or support for these results as stemmingfrom nanocomposite formation. In contrast, Modestiet al.173 reported a remarkable increase of crystallinityfor PP nanocomposites explained by consideringthat the nanoclay platelets dispersed in the matrixpromote heterogeneous nucleation, thus increasingthe crystallization rate and causing the crystallinity toincrease.

Typical DSC traces for PE nanocomposites pre-pared from PE and PE-g-MA as compatibilizer andOLS are shown in Fig. 15.121

The increase of crystallization onset and peaks(reported also in similar works174,175) confirms that theclay can provide heterogeneous nuclei with restrictionof the conformational transitions and mobility ofpolymer chains in the PE matrix, thus hindering theformation of large crystalline domains in the restrictedand confined space. Although the clay can reduce thesize of spherulites, it has no significant effect on thetotal non-isothermal crystallization rate in spite of theincrease of its activation energy.175

The simultaneous presence of mixed interca-lated/exfoliated structures is generally obtained bymixing a polyolefin (in the presence of compati-bilizer) and a clay (modified) in the melt or insolution. Exfoliated PE/MMT nanocomposite176 withlower MMT content (sample A) and intercalatedPE/MMT nanocomposite with higher MMT content(sample B) were obtained by in situ polymerizationand modulation of the amount of polymerized ethy-lene. The non-isothermal crystallization behaviour ofthese two nanocomposites was compared (Fig. 16(a))and melting and crystallization temperatures reportedfor different cooling rates (Fig. 16(b)).

The intercalated sample B always exhibits lowercrystallization temperature (Tc) and melting tempera-ture (Tm) than the exfoliated sample A, suggesting thatconfinement of PE in the MMT layer suppresses crys-tallization of PE. Moreover, sample A has higher Tc

than neat PE, showing the nucleation effect of MMT.In contrast, the intercalated sample B has lower crys-tallization temperature than neat PE, which can beattributed to the confinement of MMT layers to PEcrystals. Nevertheless, both the exfoliated sample A

816 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 15. Non-isothermal DSC traces of PE/OLS and PE-g-MA/OLSnanocomposites (Reproduced from Zhai H, Xu W, Guo H, Zhou Z,Shen S, Song Q, Eur Polym J 40:2539 (2004) by permission of JohnWiley & Sons, Ltd.).121

and the intercalated sample B start to crystallize athigher temperature than neat PE. As a result, MMTstill has a nucleation effect on the crystallization ofPE in the intercalated sample B and both nucleationand suppression effects coexist as also reflected by theinduction period and faster overall crystallization rateof the intercalated sample B. The Avrami exponentsfor sample A are close to 3.0, indicating a three-dimensional growth of the PE crystals. In contrast, theintercalated sample B has Avrami exponents close to2.0 at various cooling rates, implying that the growth ofPE crystals in sample B is two-dimensional accordingto their confinement between the clay layers. Thesmaller d-spacing between the MMT layers and themelting temperature just slightly lower than that ofnormal PE indicated that the PE chains in sample Bare parallel to the MMT layers, but not perpendicularto the clay layers (Fig. 17).176

The glass transition temperature (Tg) is stronglyaffected by the addition of nanoparticles and,particularly when there is a good filler–particle

Figure 16. (a) Non-isothermal crystallization DSC curves of PE/MMTat a cooling rate of 10 ◦C min−1 and (b) melting and crystallizationtemperatures at various cooling rates (Reproduced from Xu JT, WangQ, Fan ZQ, Eur Polym J 41:3011 (2005) by permission of ElsevierLtd.).176

Figure 17. Schematic of PE crystals intercalated between MMTlayers (Reproduced from Xu JT, Wang Q, Fan ZQ, Eur Polym J41:3011 (2005) by permission of Elsevier Ltd.).176

interaction, Tg of the amorphous polymer tendsto increase by decreasing the size of the particlesor by increasing the filler content; this behaviouris generally associated to the confinement effectgenerating a reduction in chain mobility untilsuppression of cooperative segmental motions of theconfined macromolecules (Tg). These effects, wellknown for polar polymers,6 were also evidenced for

Polym Int 57:805–836 (2008) 817DOI: 10.1002/pi

F Ciardelli et al.

amorphous hydrocarbon polymers such as naturalrubber (NR).177,178 Tg of NR increases in the presenceof organoclay (Fig. 18) due to the confinement of theelastomer segment into the organoclay nanolayers.

Accordingly poly(ethylene-co-propylene)/organo-clay nanocomposites prepared by intercalation/exfoliation processes in the melt showed larger Tg withincreased silicate content (Fig. 19); in this case dif-ferent Tg values of bulk and confined polymer chainswere determined after solvent extraction separation.179

The glass transition temperature (position ofmaximum of the tan δ peak) of the PP matrixof LDH/PP composites160 was shifted to highertemperatures compared to PP according to the

Figure 18. Storage modulus versus temperature for pure NR andNR/organoclay nanocomposite (10 wt%) (Reproduced from Wang Y,Zhang H, Wu Y, Yang J, Zhang L, Eur Polym J 41:2776 (2005) bypermission of Elsevier Ltd.).178

0.0 0.2 0.4 0.6 0.8 1.0

226

227

228

229

230

231

Gla

ss tr

ansi

tion

tem

pera

ture

(K

)

Mixture composition

Experimental dataB=1B=0.25B=2.5

Figure 19. Tg of poly(ethylene-co-propylene)/organoclaynanocomposites versus the OLS content. (The fitting is carried out byconsidering the Gordon–Taylor equation as a model to interpolate thedata for B values reported in the legend).179

restriction of chain mobility. However, Tg of LDH/PPwas lower than that of PP/MMT.

Thermal stability and flammabilityThe thermal stability of polymeric materials is usuallystudied by TGA using nitrogen or air (oxygen).Generally, the incorporation of clay into a polyolefinmatrix enhances its thermal stability by acting assuperior insulator and mass transport barrier to thevolatile products generated during decomposition. Abroad rage of literature (even by referring only to thelast three years) reports this behaviour according to thetype of clay modifier, compatibilizer, composition andprocessing conditions56,59,67,82,85,90,173,180 (Fig. 20).

The onset temperature can be affected by the ther-mal stability of the modifier: the use of oligomericallymodified clay (using a functional terpolymer with lowmolecular weight)126,181 provides PE nanocompos-ites with earlier onset temperature than the virginpolymer, but the mid-point is higher. The deacyla-tion of ethylene–(vinyl acetate) (EVA) copolymer innanocomposites was accelerated and occurred at tem-peratures lower than pure polymer due to the catalysisof strongly acidic sites created by thermal decom-position of the silicate modifier.182,183 Indeed thealkylammonium derivative decomposes with Hofmannelimination or with an SN2 nucleophilic substitutionreaction at a temperature as low as 155 ◦C, creatingan acid-activated clay,183,184 thus affecting particu-larly the flammability properties of the nanocomposite(Fig. 21).

EVA-based nanocomposites showed under non-oxidative decomposition a negligible reduction oftheir thermal stability as compared to the matrix.In contrast, when decomposed in air the same

Figure 20. Thermal stability of PP nanocomposites undernon-oxidative atmosphere (Reproduced from Zhang J, Wilkie CA,Polymer 47:5736 (2006) by permission of Elsevier Ltd.).180

Figure 21. Degradation of an ammonium salt to yield an olefin, an amine and a proton as the counterion of the clay (Reproduced from Zanetti M,Polymer Nanocomposites 253 (2006) by permission of Woodhead Publishing Ltd.).183

818 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

nanocomposite exhibited a rather large increase inits thermal stability due to char formation actingas a physical barrier between the polymer and thesuperficial zone where the combustion of the polymeroccurs.185

The cone calorimeter allows the determinationof fire-relevant properties such as heat release rate(HRR), peak of HRR (PHRR), smoke productionand CO2 and CO yields. Generally the formationof a nanocomposite leads to a reduction in peakHRR, change in char structure and decrease in massloss rate during combustion. Typical HRR evolutionswith time for different micro- and nanocomposites incomparison with the pure PP matrix were recentlydiscussed186 (Figs 22 and 23) by taking into accountthe effects of different modifiers of the clays and thedifferent morphologies obtained due the presence ofcompatibilizer.

It was suggested186 that the improved flammabilityproperties (Fig. 22) of these materials are dueto a difference in condensed-phase decompositionprocesses and not to a gas-phase effect as generallyreported. The authors suggested also that the flameretardant mechanism of PP/clay nanocomposite isassociated with thermooxidative degradation of thematrix, diffusion of volatile decomposition productsand heat transfer in the condensed phase. PP/PP-g-MA blend behaved very similarly to pure PP.

The PHRR of PP/Na-MMT microcomposite is 33%lower than that of pure PP (Fig. 23), while that ofPP/H-MMT microcomposite is 44% lower. Althoughthe dispersion state of clay particles in PP/H-MMT isthe same to that in PP/Na-MMT, the PHRR of PP/H-MMT is much lower. The results indicate that theaddition of clay can decrease the PHRR of PP matrix,and that the acidic sites created on clay layers (due tothe Hofmann reaction occurring during heating) candecrease the PHRR.

Figure 22. HRR plots for pure PP, PP/PP-g-MA, PP/OLS andPP/PP-g-MA/OMMT, where PP-g-MA is a functionalized PP used ascompatibilizer (15 wt%) and OMMT is an organophilic clay obtainedby cationic exchange with dioctadecyldimethylammonium chloride(Reproduced from Qin H, Zhang S, Zhao C, Hu G, Yang M, Polymer46:8386 (2005) by permission of Elsevier Ltd.).186

Figure 23. HRR plots for pure PP, PP/C18, PP/Na-MMT, PP/H-MMTand PP-OMMT, where C18 is made of OMMT, which is anorganophilic clay obtained by cationic exchange withoctadecyltrimethylammonium chloride, Na-MMT is a sodiummontmorillonite and H-MMT is a protonic clay containing H+ acidicsites within the galleries (Reproduced from Qin H, Zhang S, Zhao C,Hu G, Yang M, Polymer 46:8386 (2005) by permission of ElsevierLtd.).186

The PHRR of PP/H-MMT is close to that ofPP/OMMT and PP/PP-g-MA/OMMT, though theclay dispersion state of the first composite is only at themicroscale level compared to the other two compositesdefined as intercalated–exfoliated nanocomposites.This suggests that the acidic sites on clay layers playan important role in the reduction of HRR. Comparedto pure PP, the ignition time is shorter and the initialHRR is higher in PP/clay composites, but the initialmass loss of PP/Na-MMT is more serious than that ofPP nanocomposites suggesting that the shorter ignitiontime and the higher initial HRR is not only due tothe volatiles from the decomposition of the organicmodification of the clay, but also to the decompositionof the polymer matrix catalysed by the clay itself.Thus, the barrier effect of exfoliated layered silicatesfor volatiles gives a minor contribution to the delayof thermooxidative degradation and the decrease ofHRR in the nanocomposite.

The thermal stability of polyolefin/LDH nanocom-posites was recently tested and these kinds of layeredinorganic materials were considered as promising can-didates as fire retardants. The thermal decompositiontemperature measured at 50% mass loss for PE-g-MA/LDH nanocomposites, prepared by adding 2 and5 wt% of dodecyl sulfate-LDH140 was found to bearound 50–60 ◦C higher than that of PE-g-MA. More-over these nanocomposites showed a fast charringprocess at low temperature with formation of a charredlayer, which enhances the material thermal stability athigher temperatures. Similar results were obtainedfor PP,155 LLDPE139,149 and HDPE155/organo-LDHnanocomposites synthesized both by melt and solutionintercalation.

In this case too, nanocomposites showed muchslower degradation rate below 400 ◦C than pure poly-mers (Fig. 24), which was attributed both to the

Polym Int 57:805–836 (2008) 819DOI: 10.1002/pi

F Ciardelli et al.

Figure 24. Thermogravimetric curves of LLDPE andLLDPE/hydrotalcite-exfoliated nanocomposites (2, 5, 10 and 20 wt%)under air flow (Reproduced from Chen W, Qu B, J Mater Chem14:1705 (2004) by permission of the Royal Society of Chemistry).187

molecular dispersion of LDH hydroxide nanolayersin the polyolefin matrix and to the formation ofcharred layers in the first decomposition step. TheTGA data for LLDPE/hydrotalcite nanocompositesevidenced for the exfoliated structures a more effectiveenhancement of the thermal stability than interca-lated structures. However, the thermal properties ofthese nanocomposites resulted not only from the mor-phological structure but also were determined by thechemical components of the clays. In fact, this effectwas efficient up to 5 wt% of organo-LDH; succes-sively the decomposition temperature decreased mostlikely because the increased amount of LDH modifiercontent in the composites gave less stable charred lay-ers or led to formation of some aggregates that couldact as heat source domains during the degradationstep.187 LLDPE/LDH nanocomposites149 have muchhigher thermal stability than MMT nanocompositesfor the same filler content and similar structures. Also,dynamic FTIR spectroscopy, morphological evolutionand isoconversional kinetic analysis evidenced twodifferent mechanisms of enhanced thermal stabilityfor LLDPE/MMT and LLDPE/LDH nanocompos-ites. The former was mainly based on the protectivecharred layer formation by the MMT catalytic dehy-drogenation of PE molecules, whereas the latter on thebarrier effect of LDH layers with very high activationenergy, which prevents the diffusion of oxygen fromgas phase into the polymer nanocomposites, and thusnot only protects the C–C main chain from thermaldegradation but also hinders the dehydrogenation pro-cess of PE molecules. Costantino et al.98 studied thecombustion behaviour of PE sterarate anion-modifiedhydrotalcite nanocomposites evidencing that the pres-ence of hydrotalcite lamellae shields PE from thermaloxidation, thus producing a barrier effect to oxygendiffusion. The formation of this protective layer on thespecimen surface, which retains an intumescent quitecompact surface up to the end of the test, was evi-denced during the combustion in the cone calorimeter

test. In fact, due to the presence of 5 wt% of hydrotal-cite there was a reduction of 55% in the maximum ofHRR and a delay of about 50 s. The analysis showedthat the shape of the mass and mass loss rate (MLR)and of heat release and HRR curves was similar,suggesting that the lower HRR of nanocomposites wascaused by the reduction of MLR and that the enhancedflame retardancy of PE/hydrotalcite nanocompositeswas due to modifications taking place in the con-densed phase during polymer combustion. Finally,the reduction of fire performance index evidenced aconsistent reduction of fire risk.

Mechanical propertiesMechanical properties of polymer–clay nanocompos-ites are related to their micro-/nanostructure, whichin turn is directly related to the exfoliation and dis-persion of clay platelets in the polymer matrix. Gooddispersion of the clay platelets in the polymer matrixgenerally yields enhanced Young’s modulus, storagemodulus and tensile strength, but significantly reducedtensile ductility and impact strength compared to neatpolymer.73,86,87,90,97,99,103,111,188

In the case of PP nanocomposites, most studiesreport the tensile properties as a function of claycontent (Fig. 25).54

In PP/layered silicate nanocomposites, there is asharp increase in tensile modulus for very smallclay loading (about 3 wt%), followed by a muchslower increase beyond a clay loading of 4 wt%. Thisbehaviour is characteristic of polymer layered silicatenanocomposites. With an increase in clay content,strength does not change markedly compared to neatPP, and there is only a small decrease in the maximumstrain at break.

Conventional composites of PP with the samefillers do not exhibit as much of an improvementin their tensile modulus. In contrast, as the PP/layeredsilicate interaction is improved, for example whenMA functional groups are incorporated (Fig. 25(B))into the polymer, the stress is much more efficientlytransferred from the polymer matrix to the inorganicfiller, resulting in a greater improvement in tensileproperties.

Very similar behaviour is reported for PE nanocom-posites: typical stress–strain curves for different LDPEnanocomposites are shown in Fig. 26.189

The yield stress is similar for LDPE and allcompatibilized compositions, at about 9 MPa, andslightly larger, at about 9.5 MPa, for the non-compatibilized LDPE/MMT 94:6 wt% system. Othermechanical parameters, like elongation and stress atbreak, depend strongly on the clay content and itsdispersion. Non-compatibilized composition of LDPEwith modified MMT exhibits poor drawability incontrast to the compatibilized nanocomposite with thesame clay content of 6 wt%. The elongation achievedfor the nanocomposite with 6 wt% of modified MMT,larger than that for the nanocomposite with 3 wt%of modified MMT, results from the exfoliation of

820 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 25. (A) Tensile characterization (Young’s modulus, yield stress, strain at break) of PP/MMT nanocomposite by Instron testing (filledsquares). For comparison, conventionally filled macrocomposites are also shown (open circles).54 (B) Relative Young’s moduli of various PPnanocomposites, each normalized by the Young’s modulus of the respective PP. The comparison is made between PP and PP-MA used as matrixand different types of modified MMT (Reproduced from Manias E, Touny A, Wu L, Strawhecker K, Lu B, Chung TC, Chem Mater 13:3516 (2001) bypermission of the American Chemical Society).54

clay particles, and is also caused by a relativelyhigh content of compatibilizer. LDPE/PE-g-MA blendexhibits lower yield stress and larger elongation tobreak than neat LDPE. The presence of MMT-ODAclay in LDPE causes an increase, while the presenceof relatively soft compatibilizer causes a decrease inthe modulus of elasticity.

The effect of clay content on tensile propertiesbehaviour of MA-compatibilized PE-OLS nanocom-posites shows that the increase in the tensile strength ishigher at low clay content (Fig. 27(a)), indicating thatthe clay layers are better exfoliated. The strain at breakdecreases with increasing clay content as expected190

(Fig. 27(b)).The relative elastic modulus and stress at break

of melt-compounded PE-organo-MMT increase as afunction of the d-spacing of the organoclay,191 that is,exfoliation. The relative yield stress and yield strainversus d-spacing of PE–2.8 vol.% MMT are shown inFig. 28. The modest increase in yield stress suggeststhe existence of weak attraction forces between theMMT and PE, while the decrease in yield strain isattributed to local strain amplification in the polymerconfined between the particles.

The properties of LDH/polyolefin composites havenot been discussed exhaustively; however, a fewpapers mention mechanical characterization. Forexample, tensile properties such as Young’s modulus,stress at yield and strength and elongation at breakwere investigated for dodecyl sulfate-LDH/LLDPE

Figure 26. Stress–strain dependencies during tensile drawing forvarious LDPE nanocomposites. Ratios are in terms of wt% andMMT-ODA is a MMT modified with octadecylamine (Reproduced fromMorawiec J, Pawlak A, Slouf M, Gelaski A, Piorkowska E, KrasnikowaN, Eur Polym J 41:1115 (2005) by permission of Elsevier Ltd.).189

nanocomposites187 prepared by adding 2, 5, 10 and20 wt% of organo-LDH to the polymer matrix.A significant increase of Young’s modulus overpure LLDPE (59% higher) was observed in thecase of 20 wt% organo-LDH loading. However,both the strength and elongation at break decreasedseverely compared to LLDPE. From XRD a decreaseof the crystallinity was evidenced for all LDHnanocomposites compared with the LLDPE matrixbut the critical parameter for these systems appeared

Polym Int 57:805–836 (2008) 821DOI: 10.1002/pi

F Ciardelli et al.

Figure 27. Tensile properties of (a) modulus and strength and(b) strain at break190 (clay content expressed in wt%) (Reproducedfrom Lee JH, Jung D, Hong CE, Rhee KY, Advani SG, Comp Sci Tech65:1996 (2005) by permission of Elsevier Ltd.).

to be the degree of dispersion of LDH layers. In thisspecific case, a certain degree of aggregation of dodecylsulfate-LDH nanolayers was observed, which couldbe another possible reason for the decrease of strengthand elongation at break. Similar results were obtainedfor EVA samples filled with 50 wt% of unmodifiedhydrotalcite.192

Gas permeabilityThe gas barrier properties of polymer–clay nanocom-posites, with far less inorganic content of layered-silicate fillers, are remarkably superior to those of neatpolymers or their conventional counterparts. This isgenerally attributed to the fact that the dispersed clayat the nanometre scale improves the barrier proper-ties by creating a maze of tortuous paths that retardsthe progress of the gas through the matrix.55 Modelsable to describe permeability in nanofilled polymersare developed by taking into account the generaltortuosity arguments elaborated by considering corre-lation between the sheet length, concentration, relativeorientation and state of aggregation, all parametersproviding evidence of better barrier materials usingthe nanocomposites approach.193–197

The reduction of permeability arises from the longerdiffusive path that the pentrants must travel in thepresence of the filler (layered silicate in the present

2.0 3.0 4.0

d-spacing [nm]

Figure 28. Relative yield stress and strain of 2.8 vol.% HDPE-OLSnanocomposites plotted as a function of their d-spacing (Reproducedfrom Osman MA, Rupp JEP, Suter UW, Polymer 46:1653 (2005) bypermission of Elsevier Ltd.).191

case). A sheet-like morphology is particularly efficientat maximizing the path length. The tortuosity factor(f or τ depending on the symbolism) is defined as theratio of the actual distance (d′) that a penetrant musttravel to the shortest distance (d) that it would havetravelled in the absence of the layered silicate and isexpressed in terms of the length (L), width (W ) andvolume fraction (�) of the sheets (Fig. 29):

τ = d′

d= 1 +

(L

2W

)�

The effect of tortuosity on the permeability isexpressed as Ps/Pp = (1 − �)/τ , where Ps and Pp rep-resent the permeabilities of the permeant in polymer-filled composite and in the pure matrix, respectively.As an example, the dependence of the relative perme-ability on the sheet length at several different volumefractions of silicate is shown in Fig. 30.193

The curves are gradually displaced to progressivelylower relative permeabilities as a function of increasingconcentration and sheet length. This is to be expectedgiven the extremely large length-to-width ratio of sil-icate sheets (30 to ca 2000 nm in length, 1 nm thick)that serves to significantly increase the tortuosity. Also,there is no significant reduction in the relative perme-ability regardless of sheet length beyond � = 0.05.

There are different theories providing an expres-sion for τ elaborated on the basis of geometricalfactors which characterize the plate-like shape of

822 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 29. Schematic representation of the gas path in a nanocomposite (path of the penetrant = d′) and in a pure polymer (path of thepenetrant = d).

Figure 30. Dependence of the relative permeability on the sheetlength at several different volume fractions of silicate (Reproducedfrom Bharadwaj RK, Macromolecules 34:9189 (2001) by permission ofthe American Chemical Society).193

nanoparticles. The first deals with what will be called‘ribbons’ here. These are rectangular platelets withfinite width, w, and thickness, t, but infinite length.Their aspect ratio is α = w/t. The second category ofplate-like theories employs circular discs of diameter dand thickness t, with aspect ratio α = d/t. These theo-ries allow for random placement of discs in space butassume they are oriented within the plane of the filmor membrane. A clear explanation of these theories isgiven by Takahashi et al.195

Also, the sheet orientation considered normal tothe direction of diffusion sheets results in the highesttortuosity. A range of relative orientations of thesheets with respect to each other and the plane of thefilm can be obtained by defining an order parameterS representing the angle between the direction ofpreferred orientation (n) and the sheet normal (p)unit vectors. The two extremes, planar and orthogonalalignment of the sheets, may be treated by simplyinterchanging L and W (Fig. 31).

The tortousity factor is modified to include theorientational order, and the relative permeability isshown to decrease by reaching a completely planararrangement of clay platelets. Similar predictions havebeen made more recently.196 Also, the effect of theexfoliation on the relative permeability (Fig. 32) was

considered to be the critical factor in determining themaximum performance of polymer nanocompositesfor barrier applications, and the relative permeabilitywas observed to be extremely sensitive to incompletedelamination particularly for short sheet lengths.

CARBON NANOTUBES AS DISPERSED PHASEStructure of carbon nanotubesCarbon nanotubes (CNTs) are an extremely impor-tant class of nanostructured materials due to theirunique mechanical, electrical and thermal proper-ties. CNTs are the third allotropic form of carbonand were synthesized for the first time by Iijimain 1991.198 Their exceptional properties depend onthe structural perfection and high aspect ratio (typ-ically ca 300–1000). Single-walled CNTs (SWNTs)consist of a single graphene sheet (monolayer of sp2-bonded carbon atoms) wrapped into cylindrical tubeswith diameters ranging from 0.7 to 2 nm, and havetypical lengths of micrometres; multi-walled CNTs(MWNTs) consist of sets of concentric SWNTs andare therefore characterized by larger diameters.199–201

In particular, depending on the rolling direction ofthe graphene layer, different CNT structures may begenerated showing either metallic or semiconductingcharacteristics (Fig. 33).202,203

Only armchair SWNTs are metallic while allthe others are semiconductors with a band gapinversely dependent on the nanotube diameter.In addition, individual CNTs show exceptionalmechanical properties with tensile modulus204 higherthan 640 GPa and tensile strength205 up to 180 GPa.Electrons as well as phonons propagate well along theconjugated structure, conferring on CNTs thermalconductivity199,206 higher than 3000 W m−1 K−1.

Preparation of nanostructured bidimensionalsystemsThe excellent characteristics shown by both SWNTsand MWNTs suggested their incorporation into sev-eral polymeric matrices to form a new class of nanos-tructured materials (nanocomposites) characterized byremarkable ultimate properties for a variety of applica-tions. CNT polymeric composites were prepared forthe first time by Ajayan et al.207 in 1994. Since then,

Polym Int 57:805–836 (2008) 823DOI: 10.1002/pi

F Ciardelli et al.

Figure 31. Effect of sheet orientation on the relative permeability in exfoliated nanocomposites (Reproduced from Bharadwaj RK, Macromolecules34:9189 (2001) by permission of the American Chemical Society).193

Figure 32. Effect of incomplete exfoliation on relative permeability(Reproduced from Bharadwaj RK, Macromolecules 34:9189 (2001) bypermission of the American Chemical Society).193

much effort has been devoted to the production of thisnew class of extremely strong and lightweight materialsfor electrochemical, field emission and nanometer-sizeelectronic devices, sensors and functional materials forair and space technologies.199,200

However, while CNTs potentially represent one ofthe most important fillers for polymers, the majordrawback for their utilization is the existence ofstrong van der Waals interactions between individualnanotubes that make them difficult to separate fromeach other (strong tendency to form microsized

aggregates consisting of bundles of 100–500 CNTspacked together) and to disperse them uniformly in acomposite at the nanoscale.

Today, there are three fundamental methods forthe nanostructured incorporation of CNTs into athermoplastic polymer matrix: solution blending, meltblending and in situ polymerization.200,201,208 However,an efficient nanodispersion is not always achieved.

Solution blendingThe solution process is at present the most effectivemethodology for producing these nanocomposites ata small sample level. A solvent is used to dissolvefirstly the CNTs, the dissolution being in generalattained by ultrasonication and/or opportune amountsof surfactants in order to produce a metastablesuspension of nanotubes. The polymer, swollenseparately in the same solvent, is then added tothe mixture. The composite is then obtained aftersolvent evaporation at reduced time by spin-coatingthe nanotube/polymer suspension; this last procedureis often employed to prevent nanotube re-aggregation.

Another interesting approach developed to reducethe nanotube agglomeration is the hot-coagulationmethod in which the nanotube/polymer suspension ispoured into an excess of non-solvent. During polymerprecipitation (coagulation) the macromolecular chainsare able to entrap the nanotubes preventing themfrom re-aggregation. The coagulation method wasrecently adapted209 also to polyolefins (The polymermatrices discussed in this review) and based on the

824 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

Figure 33. Structures of three atomically perfect SWNT structures obtained by graphene wrapping.

difference in polymer solubility (swellability) withtemperature. In particular, CNTs were suspended in1,2-dichlorobenzene (DCB) by ultrasonication and ahot PE/DCB solution was added at about 100 ◦C.The temperature was then decreased to 70 ◦C andthe polymer allowed to precipitate, thus incorporatingthe nanotubes. The composite was then recovered byfiltration and dried under vacuum at high temperature.

Alternative solution methods have been developedtaking into account the increased solubility of nan-otubes after acid functionalization. Acid-treated CNTsbecome suspendable in ethanol by ultrasonication.The addition of the suspension to a decaline solutioncontaining polyethylene gave, after further stirring andsolvent evaporation, the expected composite.210

Another interesting approach regarding solutionblending of functionalized CNTs was proposed byBlake et al.211,212 MWNTs were alkylated with n-butyllithium by nucleophilic addition of the alkylgroups on the surface. The produced (n-Bu-MWNT−)Li+ derivative was soluble in tetrahydrofuran (THF)and reactive against halogenated polymeric speciessuch as chlorinated PP. This last reaction, performedin THF solution, promoted the covalent attachmentof PP macromolecules on the nanotube surface viaelimination of lithium chloride. The chemical linkbetween CNTs and the PP matrix provided MWNT-grafted PP (MWNT-g-PP) composite materials withstrongly enhanced interfacial interactions between thetwo components, thus potentially solving the problemsencountered of the nanotube dispersion in the polymermatrix. In addition, MWNT-g-PP can act as aneffective compatibilizer in CNTs/polyolefin blends oras a useful masterbatch for formulations based onvarious polymers.

Although improving nanotube solubility in solvents,CNT acid functionalization is not a simple methodsince it involves complex and time-consuming purifi-cation steps.202,213–216 In addition, the acid treatmentmay shorten the nanotube length, thus decreasingits aspect ratio, which is fundamental for compositeproperties.

An innovative dispersion method was recently pro-posed by Rastogi and co-workers for the prepa-ration of SWNT ultrahigh molecular weight PE(SWNT/UHMWPE) and SWNT/HDPE

composites.217,218 SWNTs were first suspended inwater with the help of a surfactant under ultrasonica-tion. The suspension was then nebulized and sprayedover the PE powder, favouring the adsorption ofSWNTs on the polymer surface. For the UHMWPEmatrix, the composite was then dissolved in hot xyleneand highly homogeneous nanocomposite films wereobtained after solvent evaporation.

Melt blendingThe melt blending process involves the dispersionof CNTs into a polymer melt using well-known melt-processing techniques (e.g. extrusion and compressionmoulding) of polymers. High temperature and shearforces in the polymer fluid are able to break thenanotube bundles, and the high viscosity of the meltprevents their formation during cooling. The meltblending process allows the preparation of large-scaleCNTs/polymer mixtures, but this is less effectivethan solution blending. In addition, due to the highviscosity of the melt at high nanotube loadings, onlysmall CNT concentrations effectively provide highlyhomogeneous nanostructured polymer composites.

Examples for the preparation by processing in themelt of CNTs/polyolefin nanocomposites are pref-erentially devoted to PE217,219 and PP220–225 matri-ces. Tang et al.226 efficiently prepared MWNT/HDPEnanocomposites by melt blending following two dif-ferent steps. The first step allowed the preparationof composite pellets by mechanically stirring the meltpolymeric mixture containing MWNTs. In the secondstep the composite was extruded at 170 ◦C and com-pression moulded. According to the authors the initialstep was necessary for feeding the nanotubes into theextruder.

In situ polymerizationThe preparation process of CNTs/polymer nanocom-posites by in situ polymerization is generally used inthe case of easily polymerizable monomers such asepoxy resins or styrene.200,201 In order to facilitatethe dissolution process in the polymerization solventor directly in the monomer, the nanotubes are eitheracid-functionalized or exfoliated under ultrasonicationin the presence of an appropriate surfactant.

Polym Int 57:805–836 (2008) 825DOI: 10.1002/pi

F Ciardelli et al.

Epoxy resin precursors containing dispersed CNTsare successively polymerized by curing in the presenceof a hardener. CNTs/polystyrene nanocomposites aregenerally obtained by the mini-emulsion polymeriza-tion technology.

Dubois and co-workers227 recently applied thein situ polymerization technique to olefin monomersby using a method derived from the polymerization-filling technique (PFT). In particular, MAO wasanchored on a MWNT surface by a non-covalentfunctionalization. A zirconocene catalyst (Cp2ZrCl2;Cp = cyclopentadienyl) was then added to the func-tionalized nanotube suspension in heptane and thereaction with MAO promoted the formation of themethylated cationic Cp2ZrMe+ species electrostat-ically linked to the nanotube surface and able toinitiate ethylene polymerization. Following this strat-egy, PE grew near the nanotube surface until reachinga molecular weight high enough to promote poly-mer deposition onto nanotubes. MWNT/PE compos-ites were obtained with a variable polymer content(from 51 to 83% by weight) and crystallinity (from33 to 59%, respectively) and were characterized bya highly homogeneous nanostructured dispersion ofthe carbon-based fillers. The as-prepared MWNT/PEnanocomposite may be additionally used as master-batch in many polymer formulations obtained bymelt blending with polyolefins such as HDPE orpoly(ethylene-co-vinylacetate) (EVAc).

Mechanical propertiesSince individual CNTs show extraordinary mechan-ical properties because of their low density andremarkable aspect ratio, their application as nanos-tructured fillers in composite materials looks par-ticularly attractive.203,228–232 It is generally acceptedthat the incorporation of well-dispersed CNTs intopolymer matrices provides composite materials withincreased tensile modulus and strength. In particular,this enhancement is directly linked with the nan-otube concentration, the degree of dispersion and theanisotropic distribution in the host polymer matrix.

For example, Tang et al.226 demonstrated that theincorporation of 1 to 5 wt% of MWNTs into HDPEprovides nanocomposites with increased stiffness(from 3.62 to 7.88%, respectively) and greater abilityto absorb energy before breaking (from 1.64 to 4.95%,respectively). Also, Ruan et al.210 reported an increaseof 25% in Young’s modulus and of 47.6% in yieldingstress for MWNT/UHMWPE composites containing1 wt% of CNTs.

However, the enhancement in tensile modulus andstrength is generally accompanied by a decrease inpolymer ductility as evidenced by the reduction ofthe strain at break.200 For example, 5 wt% car-bon nanofibre/PP composites showed a 50% increasein PP fibre modulus and a 30% decrease in theelongation at break.220 Different behaviour can beobserved in the case of well-dispersed CNTs/polymernanocomposites where the nanofiller is anisotropically

distributed within the polymer matrix after tensiledrawing. These nanocomposites show increased ten-sile strength coupled with increased ductility. Theexample reported by Ruan et al.210 showed thatanisotropic 1 wt% MWNT/UHMWPE nanocompos-ites (draw ratio = 25) were characterized by simulta-neous enhancement of tensile strength (25%) andcomposite ductility (140%). This phenomenon isattributed to the increased mobility of the PE macro-molecular chains, which increase the nucleating effectproduced by the nanofillers during polymer drawing at120 ◦C. The secondary crystals generated after draw-ing are actually supposed to be much more mobilethan the PE primary crystalline phase, thus inducinghigher ductility of the derived composite. Therefore,the incorporation of well-dispersed CNTs may addto the already high-performance UHMWPE films andfibres new innovative characteristics, thus reducing,for example, the problem related to the low creepresistance, and the low frictional resistance of CNTsmay widen the composite application to biomaterials.

Electrical propertiesThe application of CNTs as electrically conduct-ing fillers in polymers has been reviewed by severalauthors.200,205,219,233–235 In particular, high electri-cal conductivity (higher than 10−2 –10−1 S cm−1) andlow percolation threshold (less than 0.1 wt%) havebeen reached for SWNTs or MWNTs efficientlydispersed into epoxy,207,236–240 poly(methyl methacry-late) (PMMA),241–245 polycarbonate (PC),246 PS244

and poly(vinyl acetate) (PVAc).247 As for mechani-cal properties, the electrical conductivity of CNTs/polymer nanocomposites is strictly dependent on sev-eral factors such as nanotube aspect ratio, degreeof dispersion and anisotropic distribution within thepolymer matrix.

In contrast, very few examples are reported forCNTs/polyolefin electroconductive composites, prob-ably due to the difficulty of obtaining homogeneouscomposites with low percolation threshold. For exam-ple, Lemstra and co-workers217 obtained uniformlydispersed SWNT/HDPE composites with electri-cal conductivity of 10−10 S cm−1 with 2.6 wt% ofnanofiller, which was increased to 10−3 S cm−1 ondoubling the CNT concentration. According to thoseauthors, a concentration higher than 4 wt% (per-colation threshold) of SWNTs dispersed in the PEcomposite allows the formation of an interconnectedstructure of nanotubes, which is responsible for thecomposite electrical conductivity. Analogously, thesame research group reached similar values of con-ductivity by dispersing only 0.6 wt% of SWNTs in anUHMWPE matrix.218 The very low percolation valueobtained was attributed to the very effective prepara-tion technique of the composite; however, even lowerthreshold values may be further obtained by improvingthe CNT dispersion in solvents.

826 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

NOBLE METAL NANOPARTICLES ASDISPERSED PHASEThe development of methods to control size, mor-phology and aggregation of inorganic nanoparticlesis a subject of particular interest, since these vari-ables dramatically influence the optical properties, andtherefore offer ideal means for their modulation.248–252

Different from smooth metal surfaces or metal pow-ders, clusters of noble metals, such as gold, silver orcopper, assume a real and natural colour due to theabsorption of visible light at the surface plasmon reso-nance (SPR) frequency. Under the irradiation of light,the conduction electrons of the metal nanoparticlescollectively oscillate at the resonance frequency provid-ing the absorption of photons. Some of these photonsare successively converted in phonons or in latticevibrations, an effect associated with absorption. Asdescribed by the Drude–Lorentz–Sommerfeld248,249

theory and shown by a huge number of experimentaldata,249,253 a decrease in metal particle size leads to abroadening of the absorption band, a decrease of themaximum intensity and often a hypsochromic (blue)shift: these effects depend also on the cluster topologyand packing.

In addition, colloidal inorganic nanocrystals madeof a few hundred up to a few thousand atoms (groupsII–VI) are receiving considerable attention due to theirappealing properties (optical, electrical and catalytic)derived from the zero-dimensional quantum confinedcharacteristics.254–258 The colour tuneability of bothabsorption and emission of these semiconductingnanocrystals as a function of size is actually one ofthe most attractive properties.259

Dispersion of metal nanoparticles in polyolefinsThe dispersion of metal particles at the nanometrescale (MNPs) into a polymer matrix has attracted addi-tional attention due to the need of functional materialscharacterized by innovative optical properties.257,260

An even more expanding area of nanoscience isdevoted to the study and preparation of innovativenanocomposite structures in order to create opticallyfunctional materials.261,262 Noble metal nanoparticlesor semiconducting nanocrystals efficiently incorpo-rated into several polymeric matrices enhance theiroptical properties (absorption, luminescence and non-linearity) thanks to the size and growth stabilizationprovided by the macromolecular support.261,263–272

Mixing of preformed nanoparticlesThe most common procedure to obtain a dispersionof MNPs in a polymer matrix is to prepare a colloidalsolution of stabilized MNPs and then to mix it withthe desired polymer in a mutual solvent and cast a filmby evaporation from the solution.263,273 In contrast,few examples are reported showing the dispersion ofpreformed MNPs in a polymer matrix by melt mixingat high temperature.274,275

Usually a water-soluble metal salt is dispersedinto an organic solvent using a tetraalkylammonium

bromide as phase transfer agent and successivelyreduced with sodium borohydride in the presence of analkylthiol as surface stabilizer to prevent coalescenceof growing nanoparticles.276,277

In addition to thiols, different surface stabilizer havebeen used such as amines,278 poly(vinyl pyrrolidone)(PVP)279 and poly(sodium acrylate).280

By using the colloid chemistry technique describedabove, MNPs have been dispersed inUHMWPE,281,282 HDPE,285 PVA,263,268,283 polydi-methylsiloxane284,285 and poly(styrene-block-ethylene/propylene).286

The optical response of metal nanoparticles can befinely modulated and enhanced through the introduc-tion of photoactive organic molecules. Gold nanopar-ticles (DT) were prepared by the reduction of anAu(III) salt in the presence of dodecyl mercaptane,and gold nanoparticles (TT) functionalized on thesurface with a chromophore were prepared by thesame procedure using a mixture of dodecyl mercap-tane and 5′′-thiooctadecyl-[2,2′ : 5′,2′′]terthiophene-5-thiol (C18S-TT-SH), a thiol derivative of a stronglydichroic terthiophene-based chromophore.282

Bright-field TEM images of DT and TT nanoparti-cles from colloidal dispersions show for both systemsan approximately spherical shape with dimensions ofabout 3–4 and 2 nm, respectively. The difference insize between DT and TT is ascribed to the higherelectron-donating capability of the heteroaromaticstructure of the terthiophene derivative.

From these nanoparticles, 4 wt% polymeric goldnanocomposites (UH DT and UH TT) were pre-pared by casting a p-xylene solution of gold nanopar-ticles and UHMWPE and recovering the film aftersolvent evaporation at room temperature. The aggre-gation behaviour of DT nanoparticles is stronglyinfluenced by the mixing process with UHMWPEthat leads to the size increasing from 3–4 to 4–5 nm.This phenomenon is not observed for TT nanoparti-cles embedded in the PE matrix, probably due to theirhigher thermal stability induced by the terthiophene-based thiol (Fig. 34).

Formation of nanoparticles in the presence ofmacromolecular matrixAnother approach for the preparation of nanocom-posites containing metal nanoparticles involves thein situ formation of the nanoparticles directly withinthe polymer matrix.266,268,270,287–299 This process issimple and just requires the reduction of the metal ionprecursors by a photochemical or a thermal-inducedprocess. However, in contrast to nanoparticles pre-pared by the most common colloid chemistry method,the control of the size distribution of the particles pre-pared in situ is often more difficult to realize due to theinfluence of several factors such as the polymer matrixcomposition and the time and the energy density ofthe photo- or thermo-irradiation process.

Field-responsive gold nanoparticles were produceddirectly inside a vinyl alcohol-containing polymer

Polym Int 57:805–836 (2008) 827DOI: 10.1002/pi

F Ciardelli et al.

1 9 10 11 12 13 14 150

2

4

6

8

10

12

14F

requ

ency

(%

)

particle diameter (nm)

4 wt.% UH_DT

2 3 4 5 6 7 8

1 2 5 70

5

10

15

20

25

Fre

quen

cy (

%)

particle diameter (nm)

4 wt.% UH_TT

3 4 6

Figure 34. Bright-field transmission electron micrographs andparticle size distributions of 4 wt% UH DT (top) and UH TT (bottom)films (Reproduced from Pucci A, Tirelli N, Willneff EA, SchroederSLM, Galembeck F, Ruggeri G, J Mater Chem 14:3495 (2004) bypermission of the Royal Society of Chemistry).282

matrix by a photo-reduction process.300 Dispersionsof HAuCl4 and ethylene glycol in polymer films, suchas poly[ethylene-co-(vinyl alcohol)] (EVAl copolymerswith 0.27 and 0.44 ethylene molar fraction) andPVA, were irradiated with a strong UV sourceproviding gold nanoparticles even after a short periodof time (ca 5 min).290,297,298 The resulting goldparticles were efficiently stabilized by the presenceof electron-donor hydroxyl groups composing thepolymer matrices, which prevented agglomerationand formation of microsized phase-separated metalaggregates. Interestingly, the formation of silvernanoparticles in PVA thin films has been recentlyanalysed in real time using spectroscopic ellipsometryby Oates and Christalle.301

Optical properties of MNP/polyolefin dispersionsNoble metal nanoparticles incorporated in polymericmatrices, depending on particle size, shape and aggre-gation, may confer tunable absorption and scatter-ing characteristics on the derived thin films.253,302

When dispersed into polymers in non-aggregated

form, nanoparticles with very small diameters (a fewnanometres) allow the design of materials with muchreduced light scattering properties, overcoming thewidely encountered problem of opacity of heteroge-neous composites for optical applications.260,263

Even more interesting is the fact that nanopar-ticle dispersions in a polymer matrix can berendered macroscopically anisotropic, a featurethat has allowed their use in nonlinear opticaldevices and linear absorbing polarizers, e.g. fordisplay applications.263,271,274,275,281,282,303–305 Highlydichroic noble metal nanoparticles are efficientlyobtained after mechanical drawing of the polymermatrix. The uniaxial orientation of the macromolecu-lar fibres promotes the anisotropic distribution of boththe crystalline and the amorphous phases, which even-tually determines the alignment of the metal particlesalong the drawing direction.263,273

For example, PVA263 and HDPE film compositeswith alkylthiol-coated gold and silver particles, onceuniaxially oriented by stretching, present angulardependencies of absorption intensity and the colourof transmitted light. The absorption of photons isdominated by the excitation of surface plasmons in themetal particles and their aggregates.248,249 In addition,the dichroic response of gold nanocomposites mayeven be reversible if polydimethylsiloxane (PDMS) isused as polymer matrix.284 Owing to the elastomericproperties of crosslinked PDMS dichroic absorptionstates can be obtained by stretching and the opticallyisotropic conditions generated again after polymerrelaxation and by using a proper swelling agent.

The optical response of metal nanoparticles can befinely modulated and enhanced through the introduc-tion of photoactive organic molecules, possibly com-bined with control of the nanoparticle dimensions.306

The presence of direct electronic interactions betweenmetal and metal-bound chromophores is of particularinterest, because it could allow for a fine modulationof the optical properties by inducing an energy trans-fer from the excited state of the chromophore to thesurface plasmon resonance of the metal.307,308

Dispersion of gold nanoparticles and gold-bindingchromophores in a stretched polymer matrix of PEgives nanocomposites with unusual and anisotropicoptical properties. Two types of gold nanoparticleswere studied (DT and TT), and were preparedfollowing literature procedures according to thewell-known colloid chemistry technique.309 Uniaxialorientation of gold nanoparticles was achieved byfilm stretching in the solid state providing anisotropicdistribution along the drawing direction (drawing ratiodefined as the ratio between the final and the initiallength of the film). As seen from the TEM image (insetof Fig. 35), nanoparticle size and size distributionwere substantially unchanged after thermomechanicalelongation, thus suggesting aggregation phenomena tobe negligible.

UV-visible absorption spectroscopy in polarizedlight performed on oriented nanocomposites revealed

828 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

300 400 500 600 700 800 900

0.00

0.06

0.12

0.18

0.24

0.30

0°30°45°60°90°

Abs

orba

nce

(a.u

.)

wavelength (nm)

Figure 35. UV-visible spectra of 4 wt% UH TT oriented film (drawingratio = 30) as a function of the angle between the polarization of lightand the drawing direction of the film. The inset shows a TEM image ofthe same film (Reproduced from Pucci A, Tirelli N, Willneff EA,Schroeder SLM, Galembeck F, Ruggeri G, J Mater Chem 14:3495(2004) by permission of the Royal Society of Chemistry).282

that gold plasmon absorption (550 nm) is character-ized by poor dichroism, in line with literature reportsfor other small-sized gold nanoparticles.263,274,275,304

In contrast, good dichroism (dichroic ratios between14 and 30 recorded at 400 nm, depending on thedrawing ratio) was observed for the TT nanoparticleterthiophene band; this phenomenon indicates a sen-sitivity of the adsorbed chromophores to mechanicalorientation (Fig. 35). It is important to note that theseterthiophene chromophores linked to gold nanoparti-cles show a higher dichroic ratio for absorption thanthe micro-aggregates of very similar dyes dispersed inoriented PE.

The anisotropic response of these gold-boundchromophores appears to be very efficient at themolecular level thanks to the particles’ nanostructureprovided by the preparation procedure. Anotherimportant feature of UHMWPE gold nanocompositesis the occurrence of luminescence even in the absenceof the aromatic chromophore, showing a well-definedemission centred at 435 nm upon excitation at the goldabsorption band (ca 288 nm). This phenomenon hasbeen reported in the literature,310 being discoveredfor the first time by Mooradian,311 who observedvisible photoluminescence from copper and goldfilms. A possible model of this emission process312

involves excitation of electrons from occupied dbands into unoccupied sp states above the Fermilevel. In addition, the photoluminescence of UH TTnanocomposites is strongly enhanced compared tothat of DT nanoparticles dispersed in UHMWPE.This enhancement was ascribed to the interactionbetween terthiophene and gold which may modifythe electronic levels of the metal and the emissionprocess involving the sp states.313 Moreover, since TTnanoparticles in PE were smaller (about 2 nm) thanDT nanoparticles (ca 4–5 nm), the reduced diameterprovides high emission intensities, due to a more

325 350 375 400 425 450 475 500 525 550 575

0

50

100

150

200

20 30 405

10

15

20

25

0°90°

Em

issi

on (

a.u.

)

wavelength (nm)

λexc = 288 nm

UH_DTUH_TT

drawing ratio

Dic

hroi

c ra

tio

Figure 36. Fluorescence emission of oriented 4 wt% UH TT (dodecylmercaptane/C18S-TT-SH/gold) nanocomposites (drawing ratio = 40),using the gold excitation wavelength (288 nm) and (inset) emissiondichroic behaviour of all 4 wt% gold nanoparticle dispersions inUHMWPE as a function of the drawing ratio. Fluorescencemeasurements: λexc = 288 nm, λem = 435 nm.

efficient coupling of the incident radiation to theirsurface plasmon.314

When exciting TT nanocomposites at the chro-mophore excitation wavelength (390 nm), the typicalterthiophene fluorescence centred at 470 nm (andobserved in PE dispersions without gold) completelydisappeared, likely due to a quenching process pro-moted by the noble metal,307,308 suggesting terthio-phene derivatives are present in a gold-bound form.

Even more interestingly, the emission stimulatedby irradiation in the gold band region (288 nm) wasdichroic for both UH DT and UH TT nanocom-posites. This dichroic response of gold nanoparticlesanisotropically distributed along the PE fibres mightbe due to electromagnetic energy transport betweenplasmon–polariton modes of closely spaced orientedparticles.315 This phenomenon has previously beenreported, e.g. in metal nanoparticle plasmon waveg-uides, and is generally observed in systems that containclosely spaced optically excited atoms, molecules ornanocrystalline semiconductors.316,317 UH TT sam-ples, however, exhibited a much higher dichroism(at a drawing ratio of 30, the dichroic ratios are,respectively, 23 and 9 for UH TT and UH DT) andsensitivity to the drawing extent of the PE matrix(Fig. 36).

This extent of the dichroic enhancement in theUH TT systems is unlikely to be explained by theeffect of particle size alone, even if we take intoaccount that small particles should be more sensitiveto mechanical orientation in a highly viscous matrix.It appears that the dichroic enhancement may ratherbe due to an energy transfer mechanism between thewell-oriented chromophores and gold that preservesthe polarization of the radiation. This would be inline with the previously reported observation thatterthiophene nanoparticles have a less intense but stilldichroic318 absorption band (ε ∼ ε370/3) in the 260to 300 nm spectral region. This band is likely to be

Polym Int 57:805–836 (2008) 829DOI: 10.1002/pi

F Ciardelli et al.

responsible for the chromophore excitation observedin the present experiments.

Also, in the case of gold nanoparticles prepareddirectly inside a vinyl alcohol-containing polymermatrix by UV irradiation, the ultimate opticalproperties of the AuNPs/polymer dispersion filmdepend strictly on the experimental conditions.The solid-state photoinduced reaction (Au3+ → Au0)was monitored by analysing the film with differentspectroscopic techniques such as FTIR (evolution ofthe typical carbonyl band at about 1740–1720 cm−1

showing oxidation of the ethylene glycol duringthe process), UV-visible (disappearance of the Au3+

absorption band at 300 nm; Fig. 37), evolution ofthe surface plasmon absorption band attributed tonanostructured Au0 and XRD (evolution of diffractionpeaks assigned to the face-centred cubic unit cell ofgold).278

Depending on polymer composition (PVA orEVAl) and irradiation time (5 min to 2 h), goldnanoparticles showed a broad size distribution,ranging from 3–4 to 20 nm (Fig. 38), conferring on thepolymeric nanocomposite variable optical propertiesas a function of the position of the surface plasmonabsorption band.

As reported before for thiol-bound gold nanoparti-cles in oriented PE composites, the uniaxial stretchingof PVA or EVAl films promotes the anisotropic particledistribution along the drawing direction. Atomic forcemicroscopy (AFM) performed on PVA Au nanocom-posite indicates the effective alignment of the vesicle-like metal structures on the polymer surface (Fig. 39).

The induced macroscopic dichroic behaviour wasdetected using UV-visible absorption spectroscopy inpolarized light (Fig. 39).

The anisotropically distributed and interacting goldparticles are known to display a bathochromicallyshifted (i.e. red-shifted) absorption band when thepolarization vector of the photons is aligned with thestretching direction of the film, and a hypsochromic(blue) shift for the cross-polarized absorption.249,263

Such dichroic behaviour consisting of a maximumabsorption shift of about 30 nm is even betterevidenced by optical microscopy in polarized light(Fig. 40). The surface plasmon resonance shift fromφ = 0◦ (parallel to the drawing direction) to φ = 90◦(perpendicular) is associated with a clear changein colour from blue to purple of the orientednanocomposite film.

These kinds of nanostructured materials can beapplied in different fields that range from sensors(e.g. sensitive devices for UV-degradable substances,molecular strain sensors for polymer matrices) tophotonics (e.g. linear absorbing polarizers, displays,nonlinear optical devices). The results presented canbe viewed as a demonstration of the general validityof this approach aimed at conferring stimuli-sensitiveproperties by the efficient dispersion of nanostructuredmetal particles in thermoplastic polymer systems.

Also, template-synthesized gold nanoparticles ori-ented in PE demonstrate remarkable visible regionpolarization spectra. The plasmon resonance extinc-tion bands observed with the incident electric fieldpolarized parallel, perpendicular and at intermedi-ate angles to the direction of friction orientation areconsistent with the long axis of the particles beingaligned with the gross orientation axis. The experi-mental spectra agree qualitatively with the predictionsof the Rayleigh, Maxwell Garnett and dynamicalMaxwell Garnett theories in the case of larger radiusnanoparticles.319

CONCLUDING SUMMARYThe previous sections provide an impressive indicationof the possibilities offered by the nanodispersionof three rather different materials for conferring onpolyolefins markedly improved and new properties.Such results are obtained thanks to the recentlyavailable routes to overcome the severe problemsarising from the highly hydrophobic properties ofthe polymeric materials, which compete with thepolar nature of the most useful nanoparticles.Indeed polyolefins are characterized by a verylow surface tension, as typical for all lipid-typematerials, while layered clay, acid-treated carbonnanotubes and metal nanoparticles are polar species.As the interface interaction grows from less than10% in microcomposites to more than 90% innanocomposites, putting together the nanosized polarmaterials with polyethylene or polypropylene in ananostructured composite is much more complicatedthan in a classic microcomposite. Accordingly theexamples described show how the establishment ofa thermodynamically favourable surface interactionbetween the nanodispersed and continuous phase isdeterminant during preparation to reach the desiredmorphology, and also plays an important role inmaintaining the nanostructure of the final materialduring ageing and performance.

A compatibilizer is therefore necessary, capable oflocating itself at the interface. Examples of compati-bilizers are the well-known functionalized polyolefinswith grafted polar side chains, such as maleated ethy-lene and propylene polymers, as well as long-chainparaffins bearing a terminal polar group, such asalkylthiols for gold particles and alkylammonium saltsfor layered clays. While the surface modification of thedispersed phase can be made on the starting materialand its efficiency evaluated in advance, the role of thefunctionalized polyolefins is more difficult to controlas functionalization occurs during the preparation inboth melt mixing and reactor blending processes.

The treatment of the nanophase is often aimed notonly at improving its compatibility with the matrix butalso to assist the separation of the nanodimensionedpart from the starting micro-aggregate. This isthe case for layered clay where the use of long-chain alkylammonium salts is necessary to expand

830 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

300 400 500 600 700 800

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

Abs

orba

nce

(a.u

.)

wavelength (nm)

PVA_Au 0PVA_Au 10

320 nm

550 nm

Au3+

Au0

Figure 37. UV–visible absorption spectra and digital images ofPVA Au film before and after UV irradiation (10 min) and bright-fieldtransmission electron micrograph of the irradiated sample.

0 10 20 30 40 50 60

0

5

10

15

20

25

30

35

40

Fre

quen

cy (

%)

particle diameter (nm)

EVAl27_Au 5 min.

PVA_Au 5 min.

Figure 38. Particle size distributions of EVAl27 Au and PVA Aunanocomposite films irradiated for 5 min.

Figure 40. Optical microscopy images of oriented EVAl44 Au film(drawing ratio = 5) with polarization direction of the incident light(white arrows) parallel (φ = 0◦) and perpendicular (φ = 90◦) to thedrawing direction (yellow arrows) (Reproduced from Pucci A, BernaboM, Elvati P, Meza LI, Galembeck F, de Paula Leite CA, et al., J MaterChem 16:1058 (2006) by permission of the Royal Society ofChemistry).300

the interlayer distance, thus facilitating the polymerpenetration which is particularly important in the caseof nanocomposite preparation according to the moreaccepted melt mixing process. The resulting modifiedclay is called ‘organophilic clay’ as it is assumed thatthe polar silicate surface is covered by an organic layer.However, even in this case, the use of a maleated orfunctionalized polyolefin makes the obtaining of anexfoliated nanostructured composite easier. It is likelythen that the treatment with a long-chain ammoniumsalt, which gives an increase of the interlayer distance,does not provide a complete organophilic coating ofthe silicate layer surface. Rather some of these polargroups remain exposed and need to be balanced by

Figure 39. Surface plot of AFM deflection image of oriented PVA Au film irradiated for 30 min and UV-visible spectra in polarized light of the samefilm with polarization direction (grey arrows) respectively parallel and perpendicular to the drawing direction of the film (black arrow). Drawingratio = 5.

Polym Int 57:805–836 (2008) 831DOI: 10.1002/pi

F Ciardelli et al.

the presence of carboxylated side chains grafted to thepolyolefin backbone.

In the case of nanotubes, the first difficulty inobtaining a nanodispersion is provided by the strongtendency to form aggregates in the form of bundles.These bundles are quite often disaggregated intosingle nanotubes by chemical treatment formingpolar functional groups on the nanotube surface.As a consequence the separated nanotubes presenta rather polar surface and their compatibility with thepolyolefin matrix can only be obtained by a furthertreatment covering the tube with a paraffin coating.A well-established approach is based on the in situpolymerization of ethylene (reactor blending), afterthe catalyst has properly impregnated the nanotubebundle. Indeed the in situ polymerization of alkenes,a process that is scientifically very significant, offersalso an additional possibility for preparing polyolefin-based nanocomposites. This approach has certainlygreat scientific interest but does not appear at presentto be as sustainable as melt mixing. Reactor blendingin fact requires one to perform the polymerizationprocess with quite sensitive organometallic catalystsunder appropriate inert operation conditions and thenhighly purified reagents, including the fillers expectedto provide the nanophase.

In conclusion we can now consider the obtaining ofnanocomposites with polyolefin matrices as a maturetechnology as shown by the large amount of successfulexamples reported in the relevant literature. The use ofnanophase surface modification and/or functionalizedpolyolefins allows the nanodispersion of clays, carbonnanotubes and various metal nanoparticles into widelyused polymers such as thermoplastic or elastomericethylene and propylene polymers. Melt mixing ismuch less used as a dispersion process; indeed reactorblending still has some technological problems to besolved, while remaining a very interesting approachfrom the fundamental viewpoint.

Nanodispersed clays are certainly the most conve-nient for a partial improvement of thermal stabilityand elastic modulus increase of the starting polyolefin.Also gas permeability of thin flexible films of polyethy-lene or polypropylene can be reduced by about 50%.A necessary improvement in this case would be thereplacement of the long-chain ammonium salts withmore environmentally friendly and less expensive addi-tives to be used as interlayer expanders.

Carbon nanotubes are clearly more expensive anddifficult to handle in this context. However, theycan provide really unique mechanical and electricalproperties and therefore deserve more attention inthe future. The results reported up to now arecertainly significant, but a real exploitation is stillfar away; evaluation of the real nanostructure of thefinal composite is not fully available and preparationmethods are still under development.

Metal nanoparticles have been used for severaldecades for catalytic processes, often embedded inpolymer matrices. The birth of nanotechnology has

brought the scientific community to regard thesesystems from a more general viewpoint, thus extendingthe interest towards optoelectronic properties of theresulting nanocomposites, the preparation of whichis favoured by the possible synthesis of nanoparticleseither in advance or in situ. Moreover the well-definedgeometry and surface nature of these nanoparticlesfacilitate in general the compatibilization with anddispersion in the hydrophobic matrix. Thus thesehybrid materials can be regarded as very important forfuture advanced applications including the possibilityof inducing nanophase organization through externalstimuli.

ACKNOWLEDGEMENTThis work was performed with the financial supportof the FIRB Project RBNE03R78E 005 of the ItalianMinistry of University and Research (MIUR).

REFERENCES1 Alexandre M and Dubois P, Mater Sci Eng 28:1 (2000).2 Ogawa M and Kuroda K, Bull Chem Soc Jpn 70:2593 (1997).3 Paul DR, Zeng QH, Yu AB and Lu GQ, J Colloid Interface Sci

292:462 (2005).4 Picard E, Gauthier H, Gerard JF and Espuche E, J Colloid

Interface Sci 307:364 (2007).5 Constantino VRL and Pinnavaia TJ, Inorg Chem 34:883

(1995).6 Giannelis EP, Krishnamoorti R and Manias E, Adv Polym Sci

138:107 (1999).7 Zilg C, Dietsche F, Hoffmann B, Dietrich C and Mulhaupt R,

Macromol Symp 169:65 (2001).8 Leroux L and Taviot-Gueho C, J Mater Chem 15:3628 (2005).9 Evans DG and Duan X, Chem Commun 485 (2006).

10 Newman SP and Jones W, New J Chem 105 (1998).11 Leroux F and Besse JP, Chem Mater 13:3507 (2001).12 Ding P, Chen W and Baojun QU, Prog Nat Sci 16:573 (2006).13 Miyata S, Clays Clay Miner 28:50 (1980).14 Miyata S, Clays Clay Miner 31:305 (1983).15 Lopez-Salinas E, Garcia-Sanchez M, Montoya JA, Acosta DR,

Abrasolo JA and Schifter I, Langmuir 13:4748 (1997).16 Velu S, Suzuki K, Okazaki M, Osaki T, Tomura S and

Ohashi F, Chem Mater 11:2163 (1999).17 Strawhecker KE and Manias E, Chem Mater 15:844 (2003).18 Vaia RA, Structural characterization of polymer-layered silicate

nanocomposites, in Polymer–Clay Nanocomposites, ed. byPinnavaia TJ and Beall GW. John Wiley & Sons, Chichester,pp. 229 (2000).

19 Schadler LS, Polymer-based and polymer-filled nanocom-posites, in Nanocomposite Science and Technology, ed. byAjayan PM, Schadler LS and Braun PV. Wiley-VCH, Wein-heim, pp. 77 (2003).

20 Ginzburg VV, Singh C and Balazs AC, Macromolecules 33:1089(2000).

21 Ray SS, Okamoto K and Okamoto M, Macromolecules 36:2355(2003).

22 Bansal A, Yang H, Li C, Cho K, Benicewicz BC, Kumar SK,et al., Nat Mater 4:693 (2005).

23 Alcoutlabi M and McKenna GB, J Phys: Condens Matter17:R461 (2005).

24 Kadar F, Szazdi L, Fekete E and Pukansky B, Langmuir22:7848 (2006).

25 Lebaron PC, Wang Z and Pinnavaia TJ, Appl Clay Sci 15:11(1999).

26 Vaia RA, Teukolsky RK and Giannelis EP, Chem Mater 6:1017(1994).

832 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

27 Suresh R, Vasudevan S and Ramanathan KV, Chem Phys Lett371:118 (2003).

28 Zeng QH, Yu AB, Lu GQ and Standish RK, J Phys Chem B108:10025 (2004).

29 Morgan AB and Harris JD, Polymer 44:2313 (2003).30 Carlino S, Solid State Ionics 98:73 (1997).31 Brindley GW and Kikkawa S, Clays Clay Miner 28:87 (1980).32 Khan A and O’Hare D, J Mater Chem 12:3191 (2002).33 Van Oosterwyck-Gastuche MC, Brown G and Mortland MM,

Clay Miner 7:177 (1967).34 Sato T, Fujita H, Endo T, Shimada M and Tsunashima A,

React Solids 5:219 (1988).35 Carlino S, Hudson MJ, Husain SW and Knowles JA, Solid

State Ionics 84:117 (1996).36 Dimotakis ED and Pinnavaia TJ, Inorg Chem 29:2393 (1990).37 Leroux F, Adachi-Pagano M, Intissar M, Chauviere S,

Forano C and Besse JP, J Mater Chem 11:105 (2001).38 Singh M, Ogden MI, Parkinson GM, Buckley CE and Con-

nolly J, J Mater Chem 14:871 (2004).39 Carlino S and Hudson MJ, J Mater Chem 4:99 (1994).40 Carlino S and Hudson MJ, J Mater Chem 5:1433 (1995).41 Kopka H, Beneke K and Lagaly G, J Colloid Interface Sci

123:427 (1988).42 Crepaldi EL, Pavan PC, Tronto J and Valim JB, J Colloid

Interface Sci 248:429 (2002).43 Chibwe K and Jones W, Chem Mater 1:489 (1989).44 Xu, ZP and Braterman PS, J Mater Chem 13:268 (2003).45 Moujahid EM, Besse JP and Leroux F, J Mater Chem 12:3324

(2002).46 Meyn M, Beneke K and Lagaly G, Inorg Chem 29:5201 (1990).47 Miyata S and Kumura T, Chem Lett 843 (1973).48 Chibwe K and Jones W, J Chem Soc, Chem Commun 926

(1989).49 Vaia RA and Giannelis EP, Macromolecules 30:7990 (1997).50 Vaia RA and Giannelis EP, Macromolecules 30:8000 (1997).51 Fisher H, Mater Sci Eng C23:763 (2003).52 Hotta S and Paul DR, Polymer 45:7639 (2004).53 Kawasumi M, Hasegawa N, Kato M, Usuki A and Okada A,

Macromolecules 30:6333 (1997).54 Manias E, Touny A, Wu L, Strawhecker K, Lu B and

Chung TC, Chem Mater 13:3516 (2001).55 Ray SS and Okamoto M, Prog Polym Sci 28:1539 (2003).56 Passaglia E, Sulcis R, Ciardelli F, Malvaldi M and Narducci P,

Polym Int 54:1549 (2005).57 Koo CM, Ham HT, Kim SO, Wang KH and Chung IJ,

Macromolecules 35:5116 (2002).58 Wang KH, Choi MH, Koo CM, Choi YS and Chung IJ,

Polymer 42:9819 (2001).59 Kato M, Usuki A and Okada A, J Appl Polym Sci 66:1781

(1997).60 Kawasumu M, Hasegawa N, Kato N, Usuki A and Okada A,

Macromolecules 30:6333 (1997).61 Hasegawa N, Kawasumi M, Kato M, Usuki A and Okada A, J

Appl Polym Sci 67:87 (1998).62 Marchant D and Jayaraman K, Ind Eng Chem Res 41:6402

(2002).63 Rogers K, Takacs E and Thompson MR, Polym Test 24:423

(2005).64 Wang ZM, Nakajimia H, Manias E and Chung TC, Macro-

molecules 36:8919 (2003).65 Ginzburg VV and Balazs AC, Adv Mater 23:1805 (2000).66 Sinsawat A, Anderson KL, Vaia RA and Farmer BL, J Polym

Sci B: Polym Phys 41:3272 (2003).67 Chrissopoulou K, Altintzi I, Anastasiadis SH, Giannelis EP,

Pitsikalis M, Hadjichristidis N, et al., Polymer 46:12440(2005).

68 Mainil M, Alexandre M, Monteverde F and Dubois P, JNanosci Nanotechnol 6:337 (2006).

69 Osman MA, Rupp JEP and Suter UW, Polymer 46:8202(2005).

70 Fisher HR, Gielgens LH and Koster TPM, Acta Polym 50:122(1999).

71 Lee JW, Lim YT and Park OO, Polym Bull 45:191 (2000).72 Liao B, Song M, Liang H and Pang Y, Polymer 42:10007

(2001).73 Moad G, Dean K, Edmond L, Kukaleva N, Li G, Mayadunne

RTA, et al., Macromol Mater Eng 291:37 (2006).74 Shah RK and Paul DR, Polymer 47:4075 (2006).75 Gianelli W, Ferrara G, Camino G, Pellegatti G, Rosenthal J

and Trombini RC, Polymer 46:7037 (2005).76 Wang Y, Chen FB, Li YA and Wu KC, Composites B 35:111

(2004).77 Vermogen A, Masenelli-Varlot K, Seguela R, Duchet-Rumeau

J, Boucart S and Prele P, Macromolecules 38:9661 (2005).78 Modesti M, Lorenzetti A, Bon D and Besco S, Polymer

46:10237 (2005).79 Wang K, Liang S, Zhang Q, Du R and Fu Q, J Polym Sci B:

Polym Phys 43:2005 (2005).80 O’Leary S, O’Hare D and Seeley G, Chem Commun 14:1506

(2002).81 Wang GA, Wang CC and Chen CY, J Inorg Organomet Polym

Mater 15:239 (2005).82 Chen W and Qu B, Polym Degrad Stab 90:162 (2005).83 Costache MC, Wang D, Heidecken MJ, Manias E and

Wilkie CA, Polym Adv Technol 17:272 (2006).84 Qiu, L, Chen W and Qu, B, Colloid Polym Sci 283:1241 (2005).85 Moujahid EM, Besse JP and Leroux F, J Mater Chem 13:258

(2003).86 Qiu L, Chen W and Qu B, Polym Degrad Stab 87:433 (2005).87 Chen W, Feng L and Qu B, Chem Mater 16:368 (2004).88 Chen W and Qu B, Chem Mater 15:3208 (2003).89 Lan T, Kadiratna PD and Pinnavaia TJ, Chem Mater 7:2144

(1995).90 Hibino T and Jones W, J Mater Chem 11:1321 (2001).91 Li B, Hu Y, Liu J, Chen Z and Fan W, Colloid Polym Sci

281:998 (2003).92 Hsueh HB and Chen CY, Polymer 44:1151 (2003).93 Hsueh HB and Chen CY, Polymer 44:5275 (2003).94 Lee WD, Im SS, Lim HM and Kim KJ, Polymer 47:1364

(2006).95 Zammarano M, Bellayer S, Gilman JW, Franceschi M, Beyer

FL, Harris RH, et al., Polymer 47:652 (2006).96 He FA, Zhang LM, Yang F, Chen LS and Wu Q, J Polym Res

13:483 (2006).97 Qiu L, Chen W and Qu B, Polymer 47:922 (2006).98 Costantino U, Gallipoli A, Nocchetti M, Camino G, Bel-

lucci F and Frache A, Polym Degrad Stab 90:586 (2005).99 Itoh T, Ohta N, Shichi T, Yui T and Takagi K, Langmuir

19:9120 (2003).100 Costa FR, Wagenknecht U, Jehnichen D, Goad MA and

Heinrich G, Polymer 47:1649 (2006).101 Costa FR, Goad MA, Wagenknecht U and Heinrich G, Poly-

mer 46:4447 (2005).102 Costa FR, Satapathy BK, Wagenknecht U, Weidisch R and

Heinrich G, Eur Polym J 42:2140 (2006).103 Ding P and Qu B, J Polym Sci B: Polym Phys 44:3165 (2006).104 Liao CS and Ye WB, Electrochim Acta 49:4993 (2004).105 Ding P and Qu B, Polym Eng Sci 46:1153 (2006).106 Tudor J, Willington L, O’Hare D and Royan B, Chem Commun

2031 (1996).107 Heinemann J, Reichert P, Thomann R and Mulhaupt R,

Macromol Rapid Commun 20:423 (1999).108 Bergman JS, Chen H, Giannelis EP, Thomann MG and

Coates GW, Chem Commun 2179 (1999).109 Jin YH, Park HJ, Im SS, Kwak SY and Kwak S, Macromol

Rapid Commun 23:135 (2002).110 Shin AYA, Simon LC, Soares JBP and Scholz G, Polymer

44:5317 (2003).111 He A, Wang L, Li J, Dong J and Han CA, Polymer 47:1767

(2006).112 Alexandre M, Martin E, Dubois P, Garcia-Marti M and

Jerome R, Macromol Rapid Commun 21:931 (2000).113 Alexandre M, Dubois P, Sun T, Garces JM and Jerome R,

Polymer 43:2123 (2002).

Polym Int 57:805–836 (2008) 833DOI: 10.1002/pi

F Ciardelli et al.

114 Miller HA, Moneti S, Vizza F, Passaglia E, Bianchini C,Bronco S, et al., e-Polymers. Available: http://www.e-polymers.org/journal/papers/epassaglia 130 406.pdf.

115 Zhang J and Wilkie CA, Polym Degrad Stab 80:163 (2003).116 Zheng X and Wilkie CA, Polym Degrad Stab 80:441 (2003).117 Kuo SW, Huang WJ, Huang SB, Kao HC and Chang FC,

Polymer 44:7709 (2003).118 Su S, Jiang DD and Wilkie CA, Polym Degrad Stab 83:321

(2004).119 Su S and Wilkie CA, Polym Degrad Stab 83:347 (2004).120 Liang G, Xu J, Bao S and Xu W, J Appl Polym Sci 91:3974

(2004).121 Zhai H, Xu W, Guo H, Zhou Z, Shen S and Song Q, Eur

Polym J 40:2539 (2004).122 Motha K, Hippi U, Hakala K, Peltonen M, Ojanperae V,

Loefgen B, et al., J Appl Polym Sci 94:1094 (2004).123 Ray S, Galgali G, Lele A and Sivaram S, J Polym Sci A: Polym

Chem 43:304 (2005).124 Osman MA, Rupp JEP and Suter UW, J Mater Chem 15:1298

(2005).125 Zhang J, Jiang DD and Wilkie CA, Thermochim Acta 430:107

(2005).126 Zhang J, Jiang DD and Wilkie CA, Polym Degrad Stab 91:298

(2006).127 Ivanyuk AV, Gerasin VA, Rebrov AV, Pavelko RG and

Antipov EM, J Eng Phys Thermophys 78:926 (2005).128 Zhang J, Gupta RK and Wilkie CA, Polymer 47:4537 (2006).129 Kaminsky W, Funck A and Wieman K, Macromol Symp 239:1

(2006).130 Chaiko DJ, e-Polymers. Available: http://www.e-polymers.org/

papers/mchaiko 180 506.pdf [15 February 2008].131 Araujo EM, de Oliveira AD, Barbosa R and Alves TJ, Mater

Sci Forum 530–531:709 (2006).132 He FA, Zhang LM, Yang F, Chen LS and Wu Q, J Macromol

Sci A: Pure Appl Chem 44:11 (2007).133 Morgan AB and Gilman JW, J Appl Polym Sci 87:1329 (2003).134 Koo CM, Kim MJ, Coi MH, Kim SO and Chung IJ, J Appl

Polym Sci 88:1526 (2003).135 Koo CM, Kim MJ and Chung IJ, Macromolecules 36:2748

(2003).136 Pavlikova S, Thoman R, Reichert P, Mulhaupt R, Marcincin A

and Borsig E, J Appl Polym Sci 89:604 (2003).137 Gorrasi G, Tammaro L, Tortora M, Vittoria V, Kaempfer D,

Reichert P, et al., J Polym Sci B: Polym Phys 41:1798 (2003).138 Tang Y, Hu Y, Wang S, Gui Z, Chen Z and Fan W, Polym Int

52:1396 (2003).139 Tidjani A, Wald O, Pohl MM, Hentschel MP and Schartel B,

Polym Degrad Stab 82:133 (2003).140 Hasegawa N and Usuki A, J Appl Polym Sci 93:464 (2004).141 Bartholmai M and Schartel B, Polym Adv Technol 15:355

(2004).142 Parija S, Nayak SK, Verma SK and Tripathy SS, Polym Comp

25:646 (2004).143 Koo CM, Kim MJ, Wang KH and Chung IJ, J Polym Sci B:

Polym Phys 43:158 (2004).144 Leuteritz A, Pospiech D, Kretzschmar B, Willeke M,

Jehnichen D, Jentzsch U, et al., Macromol Symp 221:53(2005).

145 Kotec J, Kelnar I, Studenovsky M and Baldrian J, Polymer46:4876 (2005).

146 Szazdi L, Pukanszky B and Foeldes E, Polymer 46:8001(2005).

147 Benett EM, Causin V, Marega C, Marigo A, Ferrara G, Fer-raro A, et al., Polymer 46:8275 (2005).

148 Dudkina MM, Tenkovtsev AV, Pospiech D, Jehnichen D,Hauler L and Leuteritz A, J Polym Sci B: Polym Phys 43:2493(2005).

149 Causin V, Marega C, Marigo A and Ferrara G, Polymer46:9533 (2005).

150 Abranyl A, Szazdi L, Pukanszky Jr B, Vancso GJ and Punan-szky B, Macromol Rapid Commun 27:132 (2006).

151 Lim JW, Hassan A, Rahmat AR and Wahit MU, J Appl PolymSci 99:3441 (2006).

152 Jayaraman K and Kuman S, Polym Nanocomposite 130 (2006).153 Szazdi L, Pukanszky B, Vancso GJ and Pukanszky B, Polymer

47:4638 (2006).154 Szazdi L, Abranyl A, Pukanszky Jr B, Vancso GJ and Pukan-

szky B, Macromol Mater Eng 291:858 (2006).155 Lee SH, Cho ENR and Youn JR, J Appl Polym Sci 103:3506

(2007).156 Rohlmann CO, Failla MD and Quinzani LM, Polymer 47:7795

(2006).157 Lele A, Mackley M, Galgali G and Ramesh C, J Rheol 46:1091

(2002).158 Mlynarcikova Z, Kaempfer D, Thomann R, Mulhaupt R, Bor-

sig E and Marcincin A, Polym Adv Technol 16:362 (2005).159 Gregoriou VG, Kandilioti G and Bollas ST, Polymer 46:11340

(2005).160 Pucciariello R, Villani V, Guadagno L and Vittoria V, Polym

Eng Sci 46:1433 (2006).161 Hasegawa N, Tsukigase A and Usuki A, J Appl Polym Sci

98:1554 (2005).162 Acharya H, Srivastava SK and Bhowmick AK, Polym Eng Sci

46:837 (2006).163 Varghese S and Karger-Kocsis J, Polymer 44:4921 (2003).164 Madhusoodanan KN and Varghese S, J Appl Polym Sci

102:2537 (2006).165 Kim JT, Lee DY, Oh TS and Lee DH, J Appl Polym Sci

89:2633 (2003).166 Pasanovic-Zujo V, Gupta RK and Bhattacharya SN, Rheol

Acta 43:99 (2004).167 Pastore EO, Frache A, Boccaleri E, Marchese L and

Camino G, Macromol Mater Eng 289:783 (2004).168 Lee SS, Hur MH, Yang H, Lim S and Kim J, J Appl Polym Sci

101:2749 (2006).169 Tillekeratne M, Jollands M, Cser F and Bhattacharya SN, J

Appl Polym Sci 100:2656 (2006).170 Maiti P, Nam PH, Okamoto M, Kotaka T, Hasegawa N and

Usuki A, Polym Eng Sci 42:1864 (2002).171 Bertini F, Canetti M, Audisio G, Costa G and Falqui L, Polym

Degrad Stab 91:600 (2006).172 Krump H, Luyt AS and Hudec I, Mater Lett 60:2877 (2006).173 Modesti M, Lorenzetti A, Bon D and Besco S, Polym Degrad

Stab 91:672 (2006).174 Xie Y, Yu D, Kong J, Fan X and Qiao W, J Appl Polym Sci

100:4004 (2006).175 Yuan Q, Awate S and Misra RDK, J Appl Polym Sci 102:3809

(2006).176 Xu JT, Wang Q and Fan ZQ, Eur Polym J 41:3011 (2005).177 Liu L, Luo Y, Jia D, Fu W and Guo B, J Elast Plast 38:147

(2006).178 Wang Y, Zhang H, Wu Y, Yang J and Zhang L, Eur Polym J

41:2776 (2005).179 Prevosto D, Di Cuia F, Lucchesi M, Passaglia E and

Bertoldo M, European Polymer Conference EPF 2007, Por-toroz, 2–6 July (2007). book of abstracts P3.4.80CDRom.

180 Zhang J and Wilkie CA, Polymer 47:5736 (2006).181 Zhang J, Jiang DD and Wilkie CA, Polym Degrad Stab 91:641

(2006).182 Zanetti M and Costa L, Polymer 45:4367 (2004).183 Zanetti M, Polym Nanocomposites 253 (2006).184 Lewin M, Pearce EM, Levon K, Mey-Marom A, Zam-

marano M, Wilkie CA, et al., Polym Adv Technol 17:226(2006).

185 Beyer G, J Fire Sci 23:75 (2005).186 Qin H, Zhang S, Zhao C, Hu G and Yang M, Polymer 46:8386

(2005).187 Chen W and Qu B, J Mater Chem 14:1705 (2004).188 Pegoretti A, Dorigato A and Penati A, eXPRESS Polym Lett

1:123 (2007).189 Morawiec J, Pawlak A, Slouf M, Gelaski A, Piorkowska E and

Krasnikowa N, Eur Polym J 41:1115 (2005).190 Lee JH, Jung D, Hong CE, Rhee KY and Advani SG, Compos

Sci Technol 65:1996 (2005).191 Osman MA, Rupp JEP and Suter UW, Polymer 46:1653

(2005).

834 Polym Int 57:805–836 (2008)DOI: 10.1002/pi

Nanocomposites based on polyolefins and thermoplastics

192 Camino G, Maffezzoli A, Braglia M, De Lazzaro M andZammarano M, Polym Degrad Stab 74:457 (2001).

193 Bharadwaj RK, Macromolecules 34:9189 (2001).194 Gusev AA and Lusti HR, Adv Mater 13:1641 (2001).195 Takahashi S, Goldberg HA, Feeney CA, Karim DP, Far-

rell M, O’Leary K, et al., Polymer 47:3083 (2006).196 Lu C and Mai YW, Phys Rev Lett 95:088303:1 (2005).197 Sorrentino A, Tortora M and Vittoria V, J Appl Polym Sci B:

Polym Phys 44:265 (2006).198 Iijima S, Nature 354:56 (1991).199 Baughman RH, Zakhidov AA and de Heer WA, Science

297:787 (2002).200 Moniruzzaman M and Winey KI, Macromolecules 39:5194

(2006).201 Grossiord N, Loos J, Regev O and Koning CE, Chem Mater

18:1089 (2006).202 Dyke CA and Tour JM, J Phys Chem A 108:11151 (2004).203 Thostenson ET, Ren Z and Chou TW, Compos Sci Technol

61:1899 (2001).204 Uchida T and Kumar S, J Appl Polym Sci 98:985 (2005).205 de Heer WA, MRS Bull 29:281 (2004).206 Berber S, Kwon Y-K and Tomanek D, Phys Rev Lett 84:4613

(2000).207 Ajayan PM, Stephan O, Colliex C and Trauth D, Science

(Washington, DC) 265:1212 (1994).208 Grossiord N, Loos J and Koning CE, J Mater Chem 15:2349

(2005).209 Haggenmueller R, Fischer JE and Winey KI, Macromolecules

39:2964 (2006).210 Ruan S, Gao P and Yu TX, Polymer 47:1604 (2006).211 Blake R, Coleman JN, Byrne MT, McCarthy JE, Perova TS,

Blau WJ, et al., J Mater Chem 16:4206 (2006).212 Blake R, Gun’ko YK, Coleman J, Cadek M, Fonseca A,

Nagy JB, et al., J Am Chem Soc 126:10226 (2004).213 Bahr JL and Tour JM, J Mater Chem 12:1952 (2002).214 Georgakilas V, Kordatos K, Prato M, Guldi DM, Holzinger M

and Hirsch A, J Am Chem Soc 124:760 (2002).215 Sun Y-P, Fu K, Lin Y and Huang W, Acc Chem Res 35:1096

(2002).216 Xie X-L, Mai Y-W and Zhou X-P, Mater Sci Eng R49:89

(2005).217 Zhang Q, Rastogi S, Chen D, Lippits D and Lemstra PJ,

Carbon 44:778 (2006).218 Zhang Q, Lippits DR and Rastogi S, Macromolecules 39:658

(2006).219 Haggenmueller R, Zhou W, Fischer JE and Winey KI, J

Nanosci Nanotechnol 3:105 (2003).220 Kumar S, Doshi H, Srinivasarao M, Park JO and Schi-

raldi DA, Polymer 43:1701 (2002).221 Bhattacharyya AR, Sreekumar TV, Liu T, Kumar S, Eric-

son LM, Hauge RH, et al., Polymer 44:2373 (2003).222 Coleman JN, Cadek M, Blake R, Nicolosi V, Ryan KP, Bel-

ton C, et al., Adv Funct Mater 14:791 (2004).223 Krishnamoorti R and Yurekli K, Curr Opin Colloid Interface Sci

6:464 (2001).224 Xia H, Wang Q, Li K and Hu GH, J Appl Polym Sci 93:378

(2004).225 Assouline E, Lustiger A, Barber AH, Cooper CA, Klein E,

Wachtel E, et al., J Polym Sci B: Polym Phys 41:520 (2003).226 Tang W, Santare MH and Advani SG, Carbon 41:2779 (2003).227 Bonduel D, Mainil M, Alexandre M, Monteverde F and

Dubois P, Chem Commun 781 (2005).228 Coleman JN, Khan U and Gun’ko YK, Adv Mater 18:689

(2006).229 Ramanathan T, Liu H and Brinson LC, J Polym Sci B: Polym

Phys 43:269 (2005).230 Liu T, Phang IY, Shen L, Chow SY and Zhang WD, Macro-

molecules 37:7214 (2004).231 Jin L, Bower C and Zhou O, Appl Phys Lett 73:1197 (1998).232 Gao G, Cagin T and Goddard III WA, Nanotechnology 9:184

(1998).233 Celzard A, McRae E, Deleuze C, Dufort M, Furdin G and

Mareche JF, Phys Rev B 53:6209 (1996).

234 Du F, Fischer JE and Winey KI, Phys Rev B 72:121404/121401(2005).

235 McEuen PL, Bockrath M, Cobden DH and Lu JG, Microelec-tron Eng 47:417 (1999).

236 Bai JB and Allaoui A, Composites A 34:689 (2003).237 Barrau S, Demont P, Perez E, Peigney A, Laurent C and

Lacabanne C, Macromolecules 36:9678 (2003).238 Bryning MB, Islam MF, Kikkawa JM and Yodh AG, Adv

Mater 17:1186 (2005).239 Sandler JKW, Kirk JE, Kinloch IA, Shaffer MSP and Win-

dle AH, Polymer 44:5893 (2003).240 Sandler J, Shaffer MSP, Prasse T, Bauhofer W, Schulte K and

Windle AH, Polymer 40:5967 (1999).241 Du F, Fischer JE and Winey KI, J Polym Sci B: Polym Phys

41:3333 (2003).242 Benoit JM, Corraze B, Lefrant S, Blau WJ, Bernier P and

Chauvet O, Synth Met 121:1215 (2001).243 Park SJ, Cho MS, Lim ST, Choi HJ and Jhon MS, Macromol

Rapid Commun 24:1070 (2003).244 Regev O, ElKati PNB, Loos J and Koning CE, Adv Mater

16:248 (2004).245 Sung JH, Kim HS, Jin H-J, Choi HJ and Chin I-J, Macro-

molecules 37:9899 (2004).246 Clingerman ML, Weber EH, King JA and Schulz KH, J Appl

Polym Sci 88:2280 (2003).247 Wang G, Tan Z, Liu X, Chawda S, Koo JS, Samuilov V, et al.,

Nanotechnology 17:5829 (2006).248 Klabunde KJ, Nanoscale Materials in Chemistry. Wiley Inter-

science, New York (2001).249 Kreibig U and Vollmer M, Optical Properties of Metal Clusters,

Springer Series in Materials Science 25. Springer (1995).250 Kelly KL, Coronado E, Zhao LL and Schatz GC, J Phys Chem

B 107:668 (2003).251 Hu M, Chen J, Li ZY, Au L, Hartland GV, Li X, et al., Chem

Soc Rev 35:1084 (2006).252 Wilcoxon JP and Abrams BL, Chem Soc Rev 35:1162 (2006).253 Daniel MC and Astruc D, Chem Rev (Washington, DC)

104:293 (2004).254 Weller H, Adv Mater 5:88 (1993).255 Murphy CJ, Anal Chem 74:520A (2002).256 Sakurai H, Organomet News 100 (2004).257 Willner I and Willner B, Pure Appl Chem 74:1773 (2002).258 Murakata T, Higashi Y, Yasui N, Higuchi T and Sato S, J

Chem Eng Jpn 35:1270 (2002).259 Cozzoli PD, Pellegrino T and Manna L, Chem Soc Rev 35:1195

(2006).260 Beecroft LL and Ober CK, Chem Mater 9:1302 (1997).261 Gehr RJ and Boyd RW, Chem Mater 8:1807 (1996).262 Godovski DY, Volkov AV and Moskvina MA, Adv Sci Technol

(Faenza, Italy) 4:163 (1995).263 Caseri W, Macromol Rapid Commun 21:705 (2000).264 Godovsky DY, Adv Polym Sci 153:163 (2000).265 Sih BC and Wolf MO, Chem Commun 3375 (2005).266 Porel S, Singh S and Radhakrishnan TP, Chem Commun 2387

(2005).267 Nicolais L and Carotenuto G, Metal–Polymer Nanocomposites.

John Wiley, New York (2005).268 Perez-Juste J, Rodriguez-Gonzalez B, Mulvaney P and Liz-

Marzan LM, Adv Funct Mater 15:1065 (2005).269 Lee MH, Ju JJ, Chung WJ, Park SK, Kim MS, Seo HS, et al.,

Proc SPIE 5724:277 (2005).270 Stepanov AL, Technical Phys (Translation of Zhurnal Tekhnich-

eskoi Fiziki) 49:143 (2004).271 Hao E, Schatz GC and Hupp JT, J Fluor 14:331 (2004).272 Tamborra M, Striccoli M, Comparelli R, Curri ML, Petrella A

and Agostiano A, Nanotechnology 15:S240 (2004).273 Caseri W, Chem Nanostruc Mater 359 (2003).274 Dirix Y, Bastiaansen C, Caseri W and Smith P, J Mater Sci

34:3859 (1999).275 Dirix Y, Bastiaansen C, Caseri W and Smith P, Adv Mater

11:223 (1999).276 Brust M, Walker M, Bethell D, Schiffrin DJ and Whyman R,

J Chem Soc, Chem Commun: 801 (1994).

Polym Int 57:805–836 (2008) 835DOI: 10.1002/pi

F Ciardelli et al.

277 Lu H, Lu GH, Kessinger AM and Foss CA, J Phys Chem B101:9139 (1997).

278 Leff DV, Brandt L and Health JR, Langmuir 12:4723 (1996).279 Carotenuto G, Appl Organomet Chem 15:344 (2001).280 Hussain I, Brust M, Papworth AJ and Cooper AI, Langmuir

19:4831 (2003).281 Pucci A, Elvati P, Ruggeri G, Liuzzo V, Tirelli N, Isola M,

et al., Macromol Symp 204:59 (2003).282 Pucci A, Tirelli N, Willneff EA, Schroeder SLM, Galem-

beck F and Ruggeri G, J Mater Chem 14:3495 (2004).283 Mbhele ZH, Salemane MG, van Sittert CGCE, Nedeljkovic

JM, Djokovic V and Luyt AS, Chem Mater 15:5019 (2003).284 Uhlenhaut DI, Smith P and Caseri W, Adv Mater 18:1653

(2006).285 Pastoriza-Santos I, Perez-Juste J, Kickelbick G and Liz-

Marzan LM, J Nanosci Nanotechnol 6:453 (2006).286 Bockstaller MR, J Phys Chem B 107:10017 (2003).287 Mallick K, Witcomb MJ and Scurrell MS, Eur Phys J: Appl

Phys 29:45 (2005).288 Salvati R, Longo A, Carotenuto G, De Nicola S, Pepe GP,

Nicolais L, et al., Appl Surf Sci 248:28 (2005).289 Mallick K, Witcomb MJ and Scurrell MS, J Mater Sci 39:4459

(2004).290 Dong S, Tang C, Zhou H and Zhao H, Gold Bull (London,

UK) 37:187 (2004).291 Kaneko K, Sun HB, Duan XM and Kawata S, Appl Phys Lett

83:1426 (2003).292 Henneke DE, Malyavanatham G, Kovar D, O’Brien DT,

Becker MF, Nichols WT, et al., J Chem Phys 119:6802(2003).

293 Callegari A, Tonti D and Chergui M, Nano Lett 3:1565(2003).

294 Seibel M, in PCT Int Appl. Plasco Ehrich Plasma-CoatingGmbH, Germany, p. 12 (2001).

295 Mallick K, Wang ZL and Pal T, J Photochem Photobiol A 140:75(2001).

296 Gaddy GA, McLain JL, Steigerwalt ES, Broughton R, SlatenBL and Mills G, J Cluster Sci 12:457 (2001).

297 Tanahashi I and Kanno H, Appl Phys Lett 77:3358 (2000).298 Zhou Y, Wang CY, Zhu YR and Chen ZY, Chem Mater

11:2310 (1999).299 Itakura T, Torigoe K and Esumi K, Langmuir 11:4129 (1995).300 Pucci A, Bernabo M, Elvati P, Meza LI, Galembeck F, de

Paula Leite CA, et al., J Mater Chem 16:1058 (2006).301 Oates TWH and Christalle E, J Phys Chem C 111:182 (2007).302 Heilmann A, Polymer Films with Embedded Metal Nanoparticles.

Springer, Berlin (2003).303 McConnell WP, Novak JP, Brousseau III LC, Fuierer RR,

Tenent RC and Feldheim DL, J Phys Chem B 104:8925(2000).

304 Dirix Y, Darribere C, Heffels W, Bastiaansen C, Caseri W andSmith P, Appl Opt 38:6581 (1999).

305 Crespo P, Litran R, Rojas TC, Multigner M, de la Fuente JM,Sanchez-Lopez JC, et al., Phys Rev Lett 93:087204/087201(2004).

306 Chandrasekharan N, Kamat PV, Hu J and Jones II G, J PhysChem B 104:11103 (2000).

307 Thomas KG and Kamat PV, Acc Chem Res 36:888 (2003).308 Thomas KG and Kamat PV, J Am Chem Soc 122:2655 (2000).309 Brust M and Kiely CJ, Colloids Surf A 202:175 (2002).310 Zheng J and Dickson RM, J Am Chem Soc 124:13982 (2002).311 Mooradian A, Phys Rev Lett 22:185 (1969).312 Apell P, Monreal R and Lundqvist S, Phys Scr 38:174 (1988).313 Huang HH, Ni XP, Loy GL, Chew CH, Tan KL, Loh FC,

et al., Langmuir 12:909 (1996).314 Wilcoxon JP, Martin JE, Parsapour B, Wiedenman B and

Kelley DF, J Chem Phys 108:9137 (1998).315 Maier SA, Brongersma ML, Kik PG, Meltzer S, Requicha

AAG and Atwater HA, Adv Mater 13:1501 (2001).316 Citrin DS, Opt Lett 20:901 (1995).317 Dexter DL, J Chem Phys 21:836 (1953).318 Pucci A, Tirelli N, Ruggeri G and Ciardelli F, Macromol Chem

Phys 206:102 (2005).319 Al-Rawashdeh NAF, Sandrock ML, Seugling CJ and Foss

Jr CA, J Phys Chem B 102:361 (1998).

836 Polym Int 57:805–836 (2008)DOI: 10.1002/pi