polymer nanocomposite foams

Journal ofMaterials Chemistry A

FEATURE ARTICLE

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Polymer nanocom

LdEUaEPcC

aCabot Corporation, Billerica, MA, 01821, U

Tel: +1 978-670-6287bDepartment of Materials Science and Engine

Rensselaer Polytechnic Institute, Troy, NY, 1

Cite this: J. Mater. Chem. A, 2013, 1,3837

Received 27th August 2012Accepted 17th December 2012

DOI: 10.1039/c2ta00086e

www.rsc.org/MaterialsA

This journal is ª The Royal Society of

posite foams

Limeng Chen,*a Deniz Rende,b Linda S. Schadlerb and Rahmi Ozisikb

Polymer nanocomposite foams, polymer foams with nanoparticles, are an intriguing class of materials with

unique structure and properties. The shape, size and surface chemistry of nanoparticles can be tailored to

control the foam structure, and therefore, foam properties. Nanoparticles also add functionality to polymer

foams. In this paper, we briefly review the recent developments in polymer nanocomposites and

nanocomposite foams. This is followed by an extensive discussion regarding the role of nanoparticles in

foam morphology and properties. Finally, the current and future trends of polymer nanocomposite

foams are summarized. Both challenges and opportunities in this field are discussed.

1 Introduction

Polymer nanocomposites used to create micro- and nanoporousfoams are an intriguing class of materials for applicationsranging from thermal insulation to tissue engineering scaf-folds.1 They exhibit increased strength,2,3 surface area, anddamping properties4 over traditional foams, and provide anoutstanding opportunity to control foam structure and prop-erties. They are leading to new applications in hydrogenstorage,5 electromagnetic shielding,6 sensors,7 and thermallyinsulating materials.8 This review provides an overview of hownanollers can be used to modify polymer foam structure andproperties, and some of the fundamental understanding of themechanisms leading to the observed behavior.

A polymer foam (Fig. 1a) can be characterized by its density,cell size, and wall thickness, and comes in two main categories:

imeng Chen received a B.S.egree in Materials Science andngineering from Zhejiangniversity in Hangzhou, Chinand a Ph.D. degree in Materialsngineering from Rensselaerolytechnic Institute. He isurrently a R&D Engineer atabot Corporation.

SA. E-mail: [email protected];

ering, Rensselaer Nanotechnology Center,

2180, USA

Chemistry 2013

open cell foam (cells are connected) and closed cell foam (cellsare isolated). Foams are widely observed in nature in the form ofbone, natural sponge, coral, and natural cork. Inspired by thesematerials, processing of polymer foams has received consider-able attention.9,10 Their lightweight porous structure makesthem an excellent substitute for various functional materialsthat are used as barriers such as thermal and sound barriers,shock absorbers, absorbents, and cushions.11,12 However, poly-meric foams suffer from low mechanical strength, poor surfacequality, and low thermal and dimensional stability.1,13,14

The properties of polymer foams not only depend on theintrinsic properties of the polymer, but also on the foammorphology, such as cell density, cell size and size distribu-tion.14 Therefore, one of the most important aspects of poly-meric foam research is to enhance the strength of the foamwithout compromising its light weight. There are twoapproaches to solve this issue. The rst one is to decrease theaverage cell size without decreasing the foam density. Thesecond approach is to use llers to reinforce the polymermatrix. Of particular interest here is the use of nanollers to act

Deniz Rende is currentlyworking as a post-doctoralresearcher at Rensselaer Poly-technic Institute, USA. Shereceived her Ph.D. degree fromthe Department of ChemicalEngineering at Bogazici Univer-sity, Turkey. Her currentresearch focuses on supercriticaluid assisted processing ofpolymer nanocomposites andfoams.

J. Mater. Chem. A, 2013, 1, 3837–3850 | 3837

http://dx.doi.org/10.1039/c2ta00086e

http://pubs.rsc.org/en/journals/journal/TA

http://pubs.rsc.org/en/journals/journal/TA?issueid=TA001012

Fig. 1 (a) A scanning electron micrograph showing a typical closed cell foamstructure, (b) a graph showing the number of particles vs. particle size for a givenloading of particles, (c) the influence of cell size on the crushing strength of acarbon foam19 (Copyright at Elsevier Ltd.), and (d) the decrease of thermalconductivity of poly(methyl methacrylate) foams with increasing cell density20

(Copyright at John Wiley and Sons).

Journal of Materials Chemistry A Feature Article

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as heterogeneous nucleation sites to decrease cell size and atthe same time to reinforce the polymeric matrix.

Polymer nanocomposites are dened as composites with apolymer matrix, and llers that are less than 100 nm in at leastone dimension. While still in the early stages of development,polymer nanocomposites exhibit unique combinations ofproperties not achievable with other composite systems. Forexample, highly conductive transparent polymers15 or compos-ites with simultaneous improvement in strength and ductility16

have been demonstrated. This multifunctionality is possiblebecause: (1) some nanollers have unique properties (such ascarbon nanotubes); (2) nanollers are small optical, electrical,and mechanical defects that provide an opportunity to addfunctionality, without a decrease in transparency, breakdownstrength, and ductility; and (3) nanollers have a much largersurface area than traditional llers, leading to a larger three

Linda S. Schadler received a B.S.degree from Cornell Universityand a Ph.D. degree from theUniversity of Pennsylvania inMaterials Science and Engi-neering. She is currently theAssociate Dean of AcademicAffairs in the School of Engi-neering and the Russel SageProfessor of Materials Scienceand Engineering at RensselaerPolytechnic Institute.

3838 | J. Mater. Chem. A, 2013, 1, 3837–3850

dimensional interfacial region with properties different fromthose of the bulk polymer.

Polymer nanocomposites also offer an exciting opportunityin the formation of polymer foams for four key reasons:

(1) The nanollers act as heterogeneous nucleation sites.17

While traditional llers also act as heterogeneous nucleationsites, the small size of the nanollers means that at a givenconcentration, there are more of them per unit volume (Fig. 1b),which creates a larger number of potential nucleation sites.While this is tempered by a reduced nucleation efficiency,17 theaddition of nanoparticles tends to lead to smaller cell size.18

This ability to control the size of cells is important because cellsize and cell density can impact the properties of the foamsignicantly, as shown in Fig. 1c (ref. 19) and Fig. 1d (ref. 20).

(2) The nanollers offer the same advantages that they offerin the bulk – they can provide multifunctionality by addingconductivity or mechanical strength to the foam21 while notdecreasing other properties. The added functionality is impor-tant because, for example, we can create conductive foams foruse in electromagnetic interference (shielding)22 or foams foruse in sandwich structures23 with signicantly higher strength.

(3) The nanollers change the rheological properties. Forexample, nanollers can increase the viscosity preventing voidcoalescence, which further decreases the average cell size.24

(4) The nanoller surface chemistry can be tailored tocontrol the dispersion of the nanollers and cell size distribu-tion, and to lower the energy barrier for bubble nucleation inorder to change the structure and density of the foam.18,25

Several types of nanollers can be used to create nano-composite foams and they each have their strengths. Forexample, carbon based nanollers (carbon nanotubes and gra-phene) can add strength and conductivity.2,21 The atter surfaceof graphene provides a better nucleating surface than thecurved surfaces of carbon nanotubes17 but this disadvantage iseliminated because nanotubes provide multiple nucleationsites along their length.26

Clay nanoparticles offer another excellent opportunity tocontrol foam morphology.27,28 The at surface of the clay is anexcellent nucleation site, and if the clay dispersion is controlled,uniform cell size and signicant improvements in mechanicalproperties are achieved.8,29

Rahmi Ozisik received a B.S. inMechanical Engineering fromBogazici University in Istanbul,a M.S. in Polymer Engineeringfrom the University of Akron,and a Ph.D. in Polymer Sciencefrom the University of Akron. Heis currently an AssociateProfessor of Materials Scienceand Engineering at RensselaerPolytechnic Institute.

This journal is ª The Royal Society of Chemistry 2013


Feature Article Journal of Materials Chemistry A

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Spherical nanollers have also been used in nanocompositefoams,18,23 but they do not provide the same level of reinforce-ment as platelet and rod shaped llers. They are more valuablein adding optical or biological functionality to the foam.

With the ability to tailor the cell morphology and the prop-erties of the polymer, there are many exciting applications thatcould be developed including

� electromagnetic interference shielding,6

� hydrogen storage5 – foam structure inside the tank givesstability,

� thermal insulation30 – for buildings, additives have alreadybeen added for IR attenuation. If we can get the diameter ofcells below 70 nm, which is less than the mean free path ofphonons in air, we will further reduce thermal conductivity dueto scattering, and

� composite sandwich structures.31

This paper reviews the use of nanollers as heterogeneousnucleation sites to control cell morphology. The impact ofnanoller geometry and surface chemistry is also addressed.This is followed by a section reviewing the mechanical, elec-trical, and thermal properties of nanocomposite foams with afocus on highlighting the role of the nanollers in altering theproperties.

2 Foam morphology

The properties of foams depend strongly on the morphology ofthe foam: cell size, cell size distribution, number of cells or celldensity, and cell wall thickness. In this section we review theparameters that inuence the foam morphology.

2.1 Structural features

2.1.1 Open and closed cell structures. During the forma-tion of cells in a polymer matrix, the following sequence ofevents takes place: (i) at the initial stages, small sphericalbubbles are formed with a small reduction in the density of thematerial. (ii) As the bubbles grow spherically, the lowest densityis reached, and a closed cell structure is achieved. (iii) Theclosed cell structure is distorted as the cells grow, impinge uponeach other and form polyhedral structures. (iv) As the cellscoalesce, an open cell structure is achieved by further rupture ofcell walls.32

Depending on the processing methodology, two distinctstructures are observed in the nal foamed polymer, as shownin Fig. 2. In the closed cell morphology, the compartments areisolated from each other, which makes the foam more rigid

Fig. 2 (a) Open33 (Copyright at Nature Publishing Group) and (b) closed cellstructure of foams.


compared to open cell foams. Open cell morphology consists ofcompartments (pores, bubbles) that are connected to eachother, making the material soer and more absorbent. There-fore, the interplay between open and closed cell structures playsa key role in the nal microstructure of the foam and deter-mines their nal properties and end use.

2.1.2 Cell size and cell density. Cell density is dened asthe number of cells per unit volume. In general, for foams withthe same bulk density, cell density and cell size are inverselycorrelated. The cell size distribution also has a signicantinuence on the mechanical properties. Cell size, size distri-bution, and cell density depend on the cell nucleation and cellgrowth mechanisms, all of which contribute to the nal bulkdensity of the product.9,34 For example, the cell size is inverselyproportional to the number of nucleation sites, and heteroge-neous nucleation, by, for example nanollers can increase thenucleation rate35 and decrease the cell size distribution.36

2.2 Thermodynamics of nucleation

Classical nucleation theory considers the change in the Gibbsfree energy due to the appearance of one phase within anotherand has two components: energy liberated by the coalescence ofatoms or molecules (bulk free energy) and energy required tomaintain a border between the two phases (surface free energy).

DG ¼ sA + DGVV (1)

where DGV is the bulk free energy per unit volume, V is thevolume, s is the surface free energy per unit area, and A is thesurface area of the new phase. V and A depend on the shape ofthe nucleus. In the case of homogeneous nucleation, additionof atoms or molecules to the nucleus costs energy becauseinitially the surface energy term dominates; however, when acritical radius is reached, further growth of the nucleusbecomes energetically favorable. The critical nucleus size (rc)can be obtained by setting the rst derivative of DG with respectto size (r) to zero and is given by the following equation:

rc ¼ a

�s

DGV

�(2)

where a is a shape factor and for a spherical nucleus it is equalto�2. Any nucleus that has a size greater than rc is energeticallystable and will continue to grow.

In the case of heterogeneous nucleation, where foreignparticles or surfaces exist, the classical equations changeslightly because surfaces promote wetting – atoms and mole-cules will have a tendency to aggregate at the foreign surface. Inthis case, eqn (1) can be rewritten as follows:

DGhet ¼ DGhom

�f

2

�(3)

where f is the energy reduction factor and is a function of thecontact angle (q) and the curvature (r) of the foreign surface. Todate, the most complete form of the energy reduction factorequation has been provided by Fletcher for nucleation over aspherical surface with a single curvature.

J. Mater. Chem. A, 2013, 1, 3837–3850 | 3839



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f ðm; xÞ ¼ 1þ�1�mx

b

�3

þ3mx2

�x�m

b� 1

�

þ x3

"2� 3

�x�m

b

�þ�x�m

b

�3#

(4.1)

where

m ¼ cos q

x ¼ R

rc¼ 1

rrc

b ¼ ð1þ w2 � 2mxÞ1=2(4.2)

The inuence of the surface curvature (r) is actually coupledto the critical radius (eqn (4.2)). This necessitates the inclusionof the shape of the surface directly into the equations in acomplex way.

The dependence of f on the contact angle, which depends onthe surface chemistry (via m in eqn (4)) and R/rc (via x in eqn (4)),is presented in Fig. 3. Two regions are clearly separated by R/rc¼1. It is important to note that the equation for f provided byFletcher assumes a single curvature because it is derived for asphere, and no analytical description of f exists for a generalsurface with two curvatures. For the region R/rc[ 1, the case ofa at surface, f reaches its lowest value at any given q. In situ-ations where R/rc < 1, f is always greater than that for atsurfaces. Therefore heterogeneous nucleation is most favorableon at surfaces. Thus, nanollers with an rc on the order ofnanometers are not obvious nucleating agents. It is importantto note, however, that the number of potential nucleation sitesincreases tremendously for a given loading for nanollers.

The inuence of the contact angle is straightforward; for q ¼180�, the surface has no effect, f ¼ 2 and DGhet ¼ DGhom, hencethe homogeneous nucleation equations are recovered; forq ¼ 0�, f reaches an asymptote for values R/rc > 1 such that nonucleation can take place because the attraction to the surfaceis overwhelmingly stronger than any energetic gain obtained via

Fig. 3 Energy reduction factor (f) as a function of contact angle (q) and R/rc for aspherical particle of size R.

3840 | J. Mater. Chem. A, 2013, 1, 3837–3850

nucleation (the atoms and molecules prefer to adsorb to thesurface rather than interact with each other); for any otherintermediate value of q, f < 2 and DGhet < DGhom and the foreignsurface promotes nucleation.

2.3 Factors that inuence foam morphology

2.3.1 Nucleation mechanism. As described in the previoussection, nucleation can occur via two different mechanisms:homogeneous or heterogeneous nucleation. In homogeneousnucleation, concurrent initiation and growth of cells areobserved leading to a wide cell size distribution in the nalmicrocellular structure. In heterogeneous nucleation, however,the initiation of cell growth is supported by the existence offoreign surfaces, promoting simultaneous growth of cells insidethe polymer matrix, resulting in a narrow cell size distribution,as shown in Fig. 4. Although at surfaces provide the lowestenergy barrier to nucleation, the presence of nano-sized llersstill decreases the energy-barrier for cell nucleation comparedto that required for homogeneous nucleation and forces thenucleation to occur at the ller–polymer interface.37–39 The largenumber of nanollers provides many more nucleation sites,leading to a smaller average cell size. Therefore, the addition ofinorganic particles (llers) induces heterogeneous nucleation,provides a large number of nucleation sites, and provides asmall average cell size and narrow cell size distribution. Thehighest nucleation efficiency is achieved when nucleation onthe ller surface is energetically favorable and the ller isuniformly distributed and dispersed inside the polymer. Thesefactors can be controlled through ller type, ller size, and llersurface chemistry.9 Fig. 5 shows a summary of literature datahighlighting the increase in cell number density and decreasein cell size of various polymer foams with the addition ofnanollers. Ncomposites and Nneat polymer foam represent the cellnumber density of composite foams and neat polymer foamsfoamed under the same conditions. Dcomposites and Dneat polymer

represent the average cell diameter of composite foams andneat polymer foams foamed under the same conditions.

2.3.1.1 Filler aspect ratio. The effect of ller aspect ratio isnot well understood. When multi-walled carbon nanotubes

Fig. 4 Schematics of foammorphology differences between homogeneous andheterogeneous nucleation.



Fig. 5 A summary of literature data showing the increase in cell numberdensity and decrease in cell size of PMMA foams due to the addition ofnanofillers.18,20,27,36,40–47


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were used as llers, Chen et al.26 reported that because of thedifferences in the geometry and defect structure, the ends andsidewalls of the nanotubes acted as two distinct nucleatingcenters. At low saturation pressures, it was shown that nucle-ation occurred mainly at the nanotube ends because the criticalfree energy of nucleation is lower at the nanotube ends than atthe sidewalls; therefore, nanocomposites with shorter nano-tubes showed greater bubble density. On the other hand, athigh saturation pressures, nucleation at the nanotube sidewallsbecomes feasible and longer nanotubes acted as multiplenucleation sites leading to an increased bubble density in thehigher aspect ratio MWNT nanocomposites.26

2.3.1.2 Filler surface chemistry. The ller surface chemistryaffects the contact angle which controls both dispersion andcell nucleation efficiency, which impacts the nal cell size anddensity. For example, in supercritical carbon dioxide assistedprocessing of polymer nanocomposite foams, the interaction ofCO2 with uorinated silica nanoparticles (lower contact angle)was found to yield higher cell density and smaller cell sizescompared to unmodied nanosilica.18 Recent studies by Yehet al.20 highlighted the importance of surface modication ofsilica nanoparticles, which improved the dispersion of silicananoparticles in PMMA, leading to more heterogeneous cellnucleation sites during the foaming process. Chen et al.25

reported that the cell density of PMMA–MWNT nanocompositefoams could be greatly increased if the MWNT was covalentlymodied with glycidyl phenyl ether (GPE). GPE functionalizedMWNT had a higher affinity for CO2 than pristine MWNT,leading to a lower contact angle of CO2 on the MWNT surface.Table 1 summarizes some reported research work on thesurface modication of nanollers for controlling the structureand properties of polymer nanocomposite foams.

2.3.2 Foaming method and foaming agent. Conventionalfoamed products can be produced either by chemical or phys-ical blowing agents. Chemical foaming agents are those thatdecompose at a certain temperature whereas physical blowing


agents undergo an irreversible phase change (change from adense phase to a gas phase). Due to the challenges in theremoval of side-products, chemical foaming is oen replaced byphysical foaming. Physical foaming dates back to the 1990s andinvolves the saturation of a polymer matrix with a gas at highpressure followed by a rapid decrease in pressure. This rapiddecrease in pressure leads to phase change to a gaseous state,and volume expansion of the gas results in the formation ofbubbles inside the polymer matrix. The lack of hazardouschemical solvent use during physical foaming makes thistechnique the preferred method to produce polymeric foams.

Supercritical uids have become an attractive option forpolymer foaming processes as physical blowing agents. Super-critical uids (SCFs), especially supercritical CO2, have a rela-tively high solubility in polymers, plasticizing the polymers andlowering their glass transition temperature. Moreover, CO2 isstable, non-toxic,50,51 non-ammable, low-cost, easy to obtainfrom air and recycle, and has an easily attainable low criticaltemperature and moderate critical pressure compared to othersupercritical uids. It diffuses into the polymer melt, reducesthe viscosity and surface tension of the polymer melt and assistspolymer processing.3,52,53 In addition, carbon dioxide is easilyrecyclable and can be removed from the reaction systemthrough simple depressurization at ambient temperatures,allowing for low-energy, low-cost processing protocols.54

Foaming with supercritical carbon dioxide can be performed ineither batch processing conditions, where the samples are keptin a pressure chamber and saturated with supercritical carbondioxide,18,25,55 or in continuous extrusion foamingprocesses.3,56–58 Aer a certain period of time under supercriticalcarbon dioxide, the rapid decrease in pressure generates athermodynamic instability that causes supersaturation of CO2

dissolved in the polymer matrix and therefore, nucleation ofbubbles occurs. The saturation pressure, the saturationtemperature, and the saturation time are the critical parametersin determining the nal morphology of the polymer foam.59,60 Inrecent years, supercritical carbon dioxide has been used to foamvarious polymers, including poly(methyl methacrylate),2,18

polystyrene,39 poly-D,L-lactic acid,48 polyethylene,38 andpolyurethane.13

2.3.3 Foaming conditions. The foaming conditions are themajor determinants of nal polymer foam morphology. Forsupercritical carbon dioxide assisted foaming there are fourfundamental conditions: (i) pressure, (ii) temperature, (iii)saturation time (soaking time), and (iv) depressurization rate.These variables affect both homogeneous and heterogeneousnucleation in the same way. At high pressures, depending onthe pressure release rate, the driving force to initiate cell growthis higher. At high pressures, more CO2 molecules are absorbedinto the polyurethane yielding high cell density.13 Dai observeda dramatic decline in the cell density at low pressures. At hightemperatures, carbon dioxide acts as a plasticizer and thepolymer absorbs more carbon dioxide, leading to a decrease inthe glass transition temperature of the polymer. When foamingtakes place at temperatures lower than the glass transitiontemperature of the polymer, not all nucleated bubbles can growdue to the high viscosity (resistance) of the polymer bulk. This

J. Mater. Chem. A, 2013, 1, 3837–3850 | 3841


Table 1 A summary of literature research on the surface modification of nanofillers for controlling the structure and properties of polymer nanocomposite foams

Matrix Filler Foaming agent Surface modier Reference

PMMA Nanosilica CO2 Tridecauoro-1,1,2,2-tetrahydrooctyl silane Goren et al.18

PMMA MWNT CO2 Glycidyl phenyl ether Chen et al.25

PMMA Nanosilica N2 3-(Trimethoxysilyl)-propyl methacrylate Yeh et al.20

Poly(D,L lactic acid) Nanoclay CO2 Alkylammonium surfactants Tsimpliaraki et al.48

Polystyrene Nanoclay CO2 2-Methacryloyloxyethylhexadecyldimethylammoniumbromide

Zeng et al.27

PU Nanoclay Pentane Dibutyldimethoxytin Cao et al.49

Poly(vinyl chloride) Wood ber/MWNT Not specied Sodium hypochlorite solution Ghasemi et al.50

Fig. 6 A summary of literature data showing the percentage of increase in theelastic modulus of polymer foams due to the addition of various nanofillers. Foreach polymer matrix reported, neat polymer foams and nanocomposite foamswere foamed under the same conditions with a similar density afterfoaming.2,21,23,62,63,65,66


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leads to lowered cell density at low temperatures. At hightemperatures (above the glass transition temperature),although bubble growth is easier due to low polymer viscosity,growing bubbles can easily coalesce, leading to decreased celldensity. The highest cell density is observed in the vicinity of theglass transition temperature as the viscosity at this state isconsiderably higher than at high temperatures preventingbubble coalescence but low enough that bubbles can grow. At aconstant saturation pressure and saturation time, Dai andcoworkers reported that cell sizes increase with increasingtemperature. In agreement with increased cell sizes, the celldensity and mass density of the polymer foam decreased.13

In batch processing of polymer nanocomposite foams withsupercritical carbon dioxide, the pressure drop rate and vitri-cation temperature are additional parameters that have aninuence on the nal microcellular structure. Antunes andcoworkers reported that higher expansion rate polypropylenefoams could be obtained by increasing the rate of pressure drop.These structures exhibit slightly lower specic storage moduliwhen compared with their solid counterparts.61 Vitrication isgenerally performed aer processing to stop cell growth.

3 Properties of polymer nanocompositefoams

Polymer foams are utilized in various applications due to theiruniquemulti-functionality. With the increase in the demand forlightweight structures, polymer foams with superior propertiesat low density have become popular. However, properties ofneat polymer foams are usually inferior compared to the prop-erties of their solid counterparts due to the inclusion of airbubbles. In order to achieve a high functionality to weight ratiofor polymer foams, nanollers have been widely explored toenhance the properties of polymer foams by enhancing thepolymer matrix and controlling the foam structure.21,23,62–65 Forexample, Fig. 6 shows a summary of literature data, high-lighting the increase in the elastic modulus of polymer foamswith the addition of nanollers. In the following sections, themechanical, acoustic, electrical and thermal properties ofpolymer foams will be introduced, and the inuence of nano-llers on such properties will be discussed.

3.1 Mechanical properties

Traditionally, polymer foams are used in body protectionapparels, packaging and cushioning for delicate devices as

3842 | J. Mater. Chem. A, 2013, 1, 3837–3850

shock absorption materials.67 With increasing concerns aboutglobal warming and the endless rise in oil prices, the need forfuel efficient transportation vehicles continues to increase. As aresult, lightweight materials are increasingly in demand fortransportation vehicles to reduce fuel consumption. Lowdensity metal alloys, ber composites and polymer compositesare commonly used in such applications. Polymer foams canalso be a good candidate for such applications due to their lowdensity. However, the mechanical properties of polymer foamsare almost always reduced due to density reduction,68 whichsignicantly limits their potential as lightweight structuralmaterials. A combination of light weight and high strength isthus a major research objective in the research area of polymerfoams. The mechanical properties of polymer foams are notonly dependent on the intrinsic properties of the polymer, butalso on the density and microstructure of the foam, such as cellsize,69 cell anisotropy70 and cell orientation.71 In general, theelastic modulus and strength of polymer foams will increasewith increasing relative density, where relative density isdened as the ratio of density of the foam and the density of thesolid polymer before foaming. Fig. 7 summarizes data from theliterature showing the change of relative modulus (Er) with



Fig. 8 A typical compressive stress–strain curve of rigid polymer foams.


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relative density (rr) of polymer foams, where relative modulus isdened as the elastic modulus of foams divided by the elasticmodulus of the composing solid polymer or polymer compos-ites.2,72–76 While keeping the relative density constant,enhancing the polymer matrix properties can enhance themechanical properties of polymer foams. Nanollers arecommonly used to enhance the mechanical properties of poly-mer foams through matrix enhancement and microstructureoptimization.23,28,63

A signicant amount of work has been done to dene therelationships between mechanical properties, foam density andfoam microstructure by theoretical modeling, numericalsimulation or experimental exploration.

Fig. 8 shows a typical compressive stress–strain curve of rigidpolymer foams. There are three regions: (I) linear elasticity atlow stresses; (II) a wide plateau due to cell collapse and (III) asteep stress increase aer most cells have collapsed.14

One of the most commonly used models to quantitativelypredict the mechanical properties of polymeric foams wasdeveloped by Gibson and Ashby.78 In this model, the polymerfoams are assumed to be composed of unit cells with cell edges,cell walls (for closed cell foams) and gas inside the cells. Allthree components contribute to the mechanical properties offoams. The contribution from the gas compression is negligiblewhen the gas pressure is atmospheric pressure. The Gibson–Ashby model describes the relationship between the relativeproperties of the polymer foams and the relative density asfollows:

Er ¼ C1f2rr

2 + C2(1 � f)rr (5)

sr ¼ C01f

32r

32r þ C

02ð1� fÞrr (6)

where C1, C2, C01, and C

02 are empirical constants. Gibson and

Ashby proposed that empirically, C1, C2 and C02 are equal to 1,

and C01 is close to 0.3. f is the volume fraction of polymer used

Fig. 7 A summary of literature data showing the change of relative modulus ofpolymer foams with the change of relative density.2,72–74,76,77


for constructing cell edges, and is used as a tting constant. Foropen cell foams, f is equal to 1. rr is the relative density offoams, which is dened as the bulk density of a foam divided bythe bulk density of its solid counterpart (rr¼ rfoam/rsolid). Er andsr are the relative modulus and relative collapse strength ofpolymer foams, which are the modulus and collapse strength ofthe foam divided by the modulus and yield strength of the bulksolid matrix respectively.

At large compressive strains, the cells completely collapsed.The material is defoamed and behaves like a solid polymer. As aresult, the stress–strain curve rises steeply and its slopeapproaches the value of the bulk modulus.

According to Gibson and Ashby's model, the mechanicalproperties of polymer foams depend on the mechanical prop-erties of the bulk solid polymer and the relative density of thepolymer foam. If the relative density is constant, the mechanicalproperties depend only on the mechanical properties of thesolid polymer. This model was found to be applicable in lowdensity foams.75,79 Specically, Kabir et al.71 studied the tensileand mode-I fracture behavior of cross-linked polyvinyl chlorideand rigid polyurethane (PU) foams, and found that the foamdensity and solid polymer intrinsic property had a signicantinuence on the tensile strength, tensile modulus and fracturetoughness of PU foams. In the case of nanocomposite foams,the nanollers can reinforce the solid matrix polymer, resultingin foams with enhanced mechanical properties.14,79,80 Sahaet al.81 found that with the addition of 0.5 weight percent ofcarbon nanobers, the mode-I fracture toughness of rigidpolyurethane foams increased by 28%. Shen et al.82 found thatthe tensile modulus of polystyrene foams increased withincreasing the concentration of carbon nanobers (CNF). Whenthe ber concentration reached 5 wt%, the tensile modulus ofthe CNF–polystyrene nanocomposite foams was 1.07 GPa,which is comparable to that of solid polystyrene (1.26 GPa).

The Gibson–Ashby model focuses on low relative densityfoams (rr < 0.2). These foams are represented as lattice orrod-like structures in which cell walls and struts are modeledas structural shell and beam elements, respectively. The

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formulation of shell and beam elements assumes that thelongitudinal dimension is much higher than the axial dimen-sion. However, when the relative density of the foam is high,such an assumption is no longer valid.83 Therefore, beam andshell elements are not suitable for modeling medium–highrelative density foams. In addition, the f value in the Gibson–Ashby model, e.g. the volume fraction of polymer used forconstructing cell edges, is a constant that does not depend onfoam density. This may not be true for real samples. For foamswith high relative density, it is very hard to determine f. Saint-Michel et al.68 studied the mechanical properties of high densitypolyurethane foams with a relative density range of 0.3–0.8. Itwas found that the Gibson–Ashby approach was not accurate inmodeling the mechanical properties of such high densityfoams. Instead, a 2 + 1 phase model proposed by Christensenand Lo,84 which treats high density foam as a macroscopicallyisotropic composite medium containing spherical void inclu-sions, was found to be more accurate than the Gibson–Ashbymodel.

To more accurately predict the mechanical properties offoams with high relative density, new cubic-cell models weredeveloped by Naguib et al.85 and Chen et al.2 respectively. Inboth models, foams were assumed to be composed of discretecubic cells, as shown in Fig. 9. Naguib's model85 assumed thecomposition unit to be cubic cells with only cell edges, i.e. cellcages, to model the mechanical properties of open cell foams,while cubic cells with both cell edges and cell faces were used inChen's model2 for closed cell foams. Both models dened therelationships between the relative density of foams (rr) and cellshape factors including cell edge thickness (t) and cell edgelength (l). Chen's model further dened the relationshipbetween the f factor, i.e. the volume fraction of polymer used toconstruct cell edges, and the ratio of t over l, by assuming thatcell edges and cell faces had the same thickness t. Both modelst well with experimental data for foams with relative density inthe range of 0.2–0.5. The major difference between these twomodels and the Gibson–Ashby model is the adjoining methodthat serves to connect the cells. Naguib and Chen's modelsassume discrete cubic cells with solid continuum spacing to bethe composition units, while the Gibson–Ashby model assumesinterconnected beam and shell elements. Chen and coworkers2

also reported that the addition of multi-walled carbon nano-tubes (MWNT) could increase the compressive modulus andcollapse strength of poly(methyl methacrylate) (PMMA) foams

Fig. 9 Modeling of a polymer foam: (a) macroscopic arrangement of unit cellsinside the foam. (b) The unit cell of Naguib's model for open cell foams.85 (c) Theunit cell of Chen's model2 for closed cell foams. Copyright at Elsevier.

3844 | J. Mater. Chem. A, 2013, 1, 3837–3850

signicantly not only because of the reinforcement of the solidPMMA matrix by MWNT, but also because of the reduction ofcell size due to heterogeneous nucleation at the MWNT surface.At the same relative density, the relative modulus and relativecollapse strength of the nanocomposite foams were muchhigher than neat PMMA foams, as shown in Fig. 10. Because therelative modulus and relative collapse strength already excludethe inuence of the solid polymer properties, the increase inrelative modulus and relative strength in the nanocompositefoams was due to the decrease in cell size. The authors proposedthat, at constant relative density, increasing cell density couldresult in more cell interactions, which changed the shapes ofthe cells, thus changing the ratio of cell edge thickness to celledge length (t/l). Differences in t/l led to a difference in themechanical properties of foams. It was found that the t/l ratiofor PMMA foams lled with MWNT was much higher than neatPMMA foams with similar relative density, and a high t/l ratioled to high modulus and collapse strength.

None of the models described above dene an explicit rela-tionship between the mechanical properties of polymer foamsand cell size. Generally speaking, the dependence of mechan-ical properties on cell size varies signicantly from one propertyto another property and from one bulk solid polymer to anotherbulk solid polymer. Elastic modulus is less sensitive to thechange of cell size than collapse strength and bending strength.The probability of nding a critical aw in a long cell strut ishigher than the probability of nding a critical aw in a shortcell strut. As a result, for foams that are composed of brittlesolid matrix material, those with large cells, and thus long cellstruts, are easier to crush than those with small cells.19 Forpolymer foams composed of glassy polymers or highly crystal-line polymers, the cell size effect on crushing strength andbending strength is more signicant due to the brittleness ofthe polymer.86 Brezny and Green19 studied the inuence of cellsize on the mechanical behavior of brittle carbon foams andfound that at the same density, the compressive elastic modulusand fracture toughness of carbon foams were insensitive to cellsize change, while the crushing and bending strength scaleinversely with cell size. Alvarez et al.83 studied the structure–property relationships of high density foams by nite elementmodeling. They found a linear relationship between the elasticmodulus and relative density of the foam when the density ishigher than 0.5. They also found that foams with an average cellsize of 10 mm and foams with an average cell size of 100 mmshow a similar modulus at the same density, thus theyconcluded that the cell size has a negligible inuence on theelastic modulus of the foams. Weller and Kumar69 investigatedthe effect of cell size on the tensile properties of polycarbonatefoams. They obtained foams with different cell sizes whilekeeping the relative density 0.5, and found that the tensilemodulus, tensile strength, elongation to break, and toughnessare not signicantly affected when the average cell size is variedfrom 2.8 to 37.1 mm. Doroudiani and Kortschot75 did asystematic investigation on the tensile property–structure rela-tionship of expanded microcellular polystyrenes (EPS) based onstatistical experiment design. They found no signicant inu-ence of cell size on the tensile properties of polystyrene foams



Fig. 10 Comparison of the relative collapse strength (a) andmodulus (b) of nanocomposite foams filled with differentMWNT (F-C100, F-C20) and the relative collapsestrength and modulus of neat PMMA foams. The predicted values by the Gibson–Ashby model and Chen's model are also shown in the plot.2 Copyright at Elsevier Ltd.


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with the same density. However, Saha et al.14 found that whenthe foam density was the same, polyvinylchloride (PVC) foamwith small bubbles had a higher compressive modulus andyield stress than PVC foam with large bubbles. Reglero Ruizet al.87 found that at the same density, the elastic modulus ofrubber particle reinforced PMMA foams increased withdecreasing bubble size. Wang et al.72 utilized nite elementanalysis (FEA) combined with a 3D representative volumeelement (RVE) tetrakaidecahedral model to study the property–structure relationships of polymethacrylimide (PMI) foams.They found that the foam cell size and cell thickness can greatlyinuence the elastic properties of PMI foams: thin and mediumthickness cell foams with large cell size tended to have smallerYoung's and shear moduli, while the thick cell foams with largecell size tended to have larger Young's and shear moduli. Theeffect was especially large when the cell size and cell thicknesswere large.

It can be concluded that nanollers play a vital role in theenhancement of the mechanical properties of foams becausethey enhance the matrix and modify the foam structure. Whilehuge progress has beenmade with respect to understanding themechanical properties of polymer foams, more work is neededto fully reveal the relationships between various mechanicalproperties and foam structure, so that the properties of foamscan be optimized.

3.2 Acoustic properties

Another notable property of polymer foams that can be inu-enced by nanollers is acoustic damping efficiency. Polymerfoams are widely used as sound insulation materials because ofthe dissipation of sound energy as heat due to viscous frictionbetween polymer chains and air friction inside cells.78 Nano-llers can be used to improve acoustic damping properties offoams by increasing polymer chain friction to increase heatdissipation in the solid matrix, by increasing the rigidity of cellwalls to enhance sound wave scattering and reection, or byreducing cell size to increase tortuosity.88 Lee and coworkers89

used nano-silica to enhance the sound absorption properties ofpolyurethane (PU) foams and found that the sound absorptionratio of PU foams increased with increasing loading of nano-silica over a wide frequency range of 150–1800 Hz. At very low


frequencies (100–150 Hz), however, unlled samples had ahigher sound absorption ratio than nano-silica lled samplesdue to a resonance effect. The improvement in sound absorbingproperties was attributed to the increase in energy dissipationas heat through hysteresis with the addition of nano-silica andthe increased scattering or reection at cell walls because thenano-silica increased the stiffness of the cell walls. Verdejoet al.4 found that 0.1 wt% of carbon nanotubes was enough tosignicantly raise the sound absorption ratio of exible PUfoams in the frequency range of 1000–7000 Hz due to theenhanced mechanical energy damping caused by interfacialsliding and stick-slip behavior at the polymer–carbon nanotubeinterface.

3.3 Electrical properties

The discovery of various conductive nanoparticles, such ascarbon nanobers, carbon nanotubes,90 and graphene,91 has ledto the development of electrically conductive polymer foams.66,92

Such materials can be widely utilized in areas such as electro-static discharge (ESD) protection, electromagnetic interference(EMI) shielding,93,94 and lightening-protection panels. Electri-cally conductive llers form an electron transport path in thefoam structure through cell walls and cell struts95 and thusresult in an increase in electrical conductivity of polymer foams.The electrical conductivity of polymer foams is mainly inu-enced by the electrical conductivity of the matrix polymercomposites, foam density and microstructure of the foams.

Xu and coworkers95 reported the synthesis of an ultra-lightelectrically conductive composite foam with a density of 0.05 gcm�3 based on carbon nanotube (CNT) and rigid polyurethane(PU). They found an interesting density-dependent conductor–insulator transition, which revealed the light-weight limit of aconductive polymer composite foam. Xu and coworkersproposed three reasons for the conductor–insulator transition:(1) with the decrease in density, the cell walls of the foam gotthinner and thinner, and there was a transition from threedimensional (3D) percolation to two dimensional (2D) perco-lation. This transition increased the percolation threshold andreduced the conductivity of the cell walls. (2) Foams with lowerdensity possessed more porous structures, which led to less of achance for CNT bridging across cells through cell-strut and thus

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a less conductive path. (3) The authors also proposed that theCNT content in the cell walls was reduced with the decrease infoam density, which also decreased the conductivity of cellwalls. But this argument was not valid. CNT in cell walls couldnot leave cell walls when the foam density changed, so theloading of CNT in cell walls could not change. Another possiblereason for the reduction in conductivity with decreasing foamdensity is that the signicant alignment of CNT due to thestretching of cell walls could reduce the number of contactsbetween CNT96 and the conductivity of the cell walls decreasedwith increasing stretching of cell walls when foam densitydecreased.

Athanasopoulos and coworkers97 studied the process–struc-ture–property relation of CNT–polyurethane nanocompositefoams using percolation laws. Polyurethane foams withdifferent CNT concentrations and different density wereprepared. The change of electrical conductivity of foams withCNT concentrations at constant density was tted using aStatistical Percolation Model. They found that the percolationthreshold of foams increased with decreasing foam density.And the dimensionality parameter, a parameter that describesthe structure of a conductive network, indicated that the CNTsin the PU foams retained their distribution in a 3D percolationnetwork instead of a 2D network even if the foam density waslow. In addition, a Statistical Percolation Model was also foundto be applicable in predicting the change in electrical conduc-tivity of foams with varying relative density at constant CNTconcentration. The statistical percolation threshold of relativedensity was calculated by tting the model with experimentaldata. The statistical percolation threshold of relative densitydecreased with increasing CNT concentration. At the end oftheir report, a material design/selection map for electricallyconductive foams was presented. Such a map provides astraightforward reference for designing lightweight electricallyconductive PU foams. However, it should be noted that the mapwas limited to foams whose structure was not signicantlyinuenced by conductive llers. The foam structure of otherpolymer systems might be very sensitive to ller concentra-tions.41 As a result, the foam structure parameters such as cellsize and cell density might also play a role in determining theelectrical conductivity of nanocomposite foams.

Xiang et al.98 reported conductive polyurethane (PU) foamswith a negative temperature coefficient of resistivity (resistivitydecreased with increasing temperature). The resistivity of solidconductive polymer composites usually increases with temper-ature because of the breakage of conductive pathways when thecomposite expands upon heating. The negative temperaturecoefficient was attributed to the change in foam structure withtemperature. The increase in temperature increased the volumeof gas in the cells, which in turn squeezed the cell walls and thecarbon nanotubes. The higher contact pressure decreasedthe interfacial resistance, which led to decreased resistivity. Thetransition was reversible and the coefficient depended on thefraction of gas (or the composite density) and the modulus ofthe cell walls (or the modulus of the solid polymer). The samegroup reported CNT/highly crosslinked polyurethane foamsexhibiting excellent electrical stability with temperature.21 The

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matrix did not change volume as signicantly with temperatureand thus did not change the CNT–CNT contact pressure.

Electromagnetic interference (EMI) shielding has been oneof the most studied applications of conductive nanocompositefoams.92,93,99 Electromagnetic interference (EMI) may be denedas electromagnetic radiation emitted by electrical circuits undercurrent operation. These EMI signals are undesirable becausethey disturb the normal function of the electronic appliancesand they potentially cause irradiative damage to the humanbody. Conventional EMI shielding materials are metals andtheir composites, which have high shielding effectiveness dueto their high conductivity and high dielectric constant. Whilemetals have superior EMI shielding properties, they have manydisadvantages such as high density and low corrosion resis-tance. Due to the desire for lightweight EMI shielding systems,conducting polymer and polymer based conducting compositesare receiving more attention.6,100 Polymer nanocompositefoams, compared to solid conductive nanocomposites, not onlyhave lower density but also a large air–polymer interface areadue to the large number of cells, which will create multiplereections of electromagnetic waves inside the shieldingmaterial.26,101,102 Thus, nanocomposite foams will absorb andattenuate instead of reecting electromagnetic radiation, whichwill reduce damage to internal electronic circuits caused byreected radiation. Table 2 summarizes the EMI shieldingeffectiveness of various polymer nanocomposite foams reportedin the literature. The shielding effectiveness (SE) is dened asthe ratio between the incoming power (Pi) and outgoing power(Po) aer the electromagnetic wave travels through the testedsample. Generally, shielding effectiveness (SE) is expressed indecibels (dB): SE ¼ 10 log(Pi/Po). Table 2 shows that the SE ofpolymer nanocomposite foams can be at least 19 dB, whichmeans that 80% of the electromagnetic wave is shielded. Forsome systems, such as the MWNT–polycaprolactone nano-composite foams reported by Thomassin and coworkers,99 theSE was higher than 60 dB, resulting in nearly complete shield-ing of electromagnetic waves in the test frequency range.However, the shielding effectiveness of test samples is directlyinuenced by the sample thickness, which was not reported insome of the literature. It is highly suggested that furtherresearch on the EMI shielding properties of materials includesthe information of sample thickness in order to make areasonable comparison between different sets of samples andbetween results from different research groups.

3.4 Thermal insulation properties

One of the major functions of polymer foams is to providethermal insulation. Thermal insulation materials are widelyused in buildings and motor vehicles to reduce energyconsumption. The effective thermal conductivity of polymerfoams is due to heat ow in the solid polymer and the cell gas,thermal radiation, and convection processes. For polymerfoams with small cells, the convection contribution to thermalconductivity can be neglected,104 and heat ow in low densityfoams, which is commonly used for thermal insulation, ismainly determined by gas properties due to the high volume



Table 2 A summary of literature data showing the EMI shielding effectiveness of polymer nanocomposite foams

Matrix Filler Filler loadingEMI frequency(GHz)

Shieldingeffectiveness (dB) Reference

Polycaprolactone MWNT 0.25 vol% 25–40 60–80 Thomassin et al.99

Polystyrene MWNT 7 wt% 8.2–12.4 18.2–19.3 Yang et al.93

Polystyrene CNF 15 wt% 8.2–12.4 >19 Yang et al.92

Fluorocarbon elastomer MWNT 12 wt% 8.2–12.4 50 Fletcher et al.103

PMMA Graphene 1.8 vol% 8–12 13–19 Zhang et al.22


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fraction of gas. As a result, the effective thermal conductivity ofpolymer foams depends primarily on the heat ow in the cellgas and radiative heat transfer. For most foams used as thermalinsulation materials, the heat ow contribution to the thermalconductivity depends on the volume fraction and thermalconductivity of the gas, but is not signicantly inuenced by thestructure of the foams unless the cell size is below themean freepath of air (70 nm), when the air molecules will not collide witheach other and consequently minimize kinetic energy andsubsequent heat transfer (i.e. Knudsen effect).30 When the cellsize is below the mean free path of air, the conduction throughthe gas phase can be dened by the following equation:

lg ¼ lg0fg

1þ bT

PgDcell

(7)

where lg is the thermal conductivity of the gas phase insidefoam cells, lg0 is the thermal conductivity of gas under atmo-spheric conditions, Pg is the gas pressure inside foam cells, fg isthe volume fraction of the gas phase, T is the absolutetemperature and b is the Knudsen coefficient. Polyurethane(PU) aerogels and xerogels64 have been successfully producedwith a cell size smaller than 100 nm possessing extremely lowthermal conductivity but poor mechanical properties due to thesynthesizing process. Future efforts are necessary for thedevelopment of a polymeric nanoporous structure possessing acombination of low thermal conductivity and superiormechanical properties, and nanoparticles might play a positiverole in the process.

With xed foam density, the volume fraction of cell gas isxed, so the heat ow in polymer foams is dependent on gasthermal conductivity. In the foam industry, gases with lowthermal conductivity are commonly used to generate cells inorder to reduce the thermal conductivity of foams.30,105

However, the thermally insulating gases in foam cells willdiffuse into the atmosphere and eventually all cells will be lledwith air and the thermal conductivity will increase accordingly.In order to maintain the thermal insulation properties of foams,llers were commonly used to reduce the gas diffusion rate,8,106

thus keeping the thermal insulating gas inside foams for amuch longer time. Kim and coworkers8 utilized ultrasonicationto achieve a uniform dispersion of organoclays in polyurethane(PU) and generated PU–clay nanocomposite foams using envi-ronmental friendly blowing agents. They had several importantndings: (1) the thermal conductivity of PU foams depended onthe thermal conductivity of the foaming agents. Low thermalconductivity foaming agents led to foams with low thermal


conductivity; (2) clays slowed the diffusion of foaming agentsfrom foam cells to atmosphere, and thus slowed down theamelioration of thermal insulation properties of PU foams; and(3) with the same foaming agent, the addition of clays decreasedthe cell size, which resulted in decreased thermal conductivityof the foam.

The contribution from thermal radiation to the effectivethermal conductivity of polymer foams is not negligible. Forexample, it was reported that as much as 7–34% of thermalconductance in polystyrene was through radiative heat trans-fer.107 For foams with the same density, heat radiation can bereduced by reducing cell size or increasing radiation absorptionthrough the addition of infrared light absorbing llers107 suchas carbon black,108 graphite30 and graphene.109 Loeb104 proposedthat the effective radiative thermal conductivity (k) of a sphericalcell was:

k ¼ 16

3resT3 (8)

where r is the diameter of the cell, e is the emissivity of theradiating surface, s is Stefan's radiation constant and T is theabsolute temperature.

Reducing the cell size of foams while keeping the foamdensity constant can be achieved by changing the foamingconditions such as foaming agents, foaming pressure andtemperature, or by using llers as heterogeneous cell nucleationsites, which can increase cell density and decrease cell size infoams as discussed in Section 2. Yeh et al.20 reported that, whenholding the foaming conditions constant, the thermalconductivity of PMMA foams could be reduced with the addi-tion of vinyl modied silica. The authors attributed the thermalconductivity difference to the inclusion of a larger amount of airin PMMA–silica composite foams compared to neat PMMAfoams. However, the reduction in cell size in the compositefoams could lead to lower radiative heat transfer, which couldalso have reduced the thermal conductivity. Patro et al.110

reported a 10% reduction in thermal conductivity of poly-urethane foams with the addition of 4.46 wt% vermiculite clay.The thermal conductivity reduction was attributed to thereduction in cell size and the increase in closed cell content.Harikrishnan et al.111 reported a 5.4% reduction in thermalconductivity of PU foams with the addition of only 0.5 wt% ofcarbon nanobers in the foams. The thermal conductivityreduction was believed to be due to both reduced cell size andincreased radiation absorption/reection induced by carbonnanobers. Trying to decrease thermal conductivity of a mate-rial by the incorporation of a thermally conductive ller seems

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to be contradictory. The thermally conductive ller can increasethe thermal conductivity of the insulating matrix, thusincreasing the heat ow rate in the matrix. However, Kompanet al.109 showed that the reduction in radiative heat transfer dueto the addition of graphene clusters outweighed the increase inheat ow rate of the solid matrix, as long as the matrix materialwas insulating and the concentration of graphene clusters waslower than the percolation threshold. As a result, low loading ofcarbon llers can be an efficient way to enhance the thermalinsulation properties of polymer composite foams.

3.5 Thermal stability

The thermal stability of polymer composite foams can also beenhanced by llers. Silica nanoparticles were shown to slightlyincrease the thermal decomposition temperature of rigid poly-urethane foams63 and poly(methyl methacrylate) foams.20 Yanget al.46 prepared graphene oxide (GO) lled polystyrene foamsand found that graphene oxide could slightly improve thethermal stability of the PS–GO composites. Yan et al.21 showedthat the incorporation of 0.5 wt% carbon nanotubes (CNTs)could signicantly improve the thermal stability of poly-urethane foams, with the degradation temperature increasingfrom 450 �C to 499 �C at 50% weight loss. The nely dispersedCNTs restricted the thermal motion of the PU chains as well asthe diffusion of volatile decomposition products, improving thethermal stability of the CNT–PU composites. A noticeableimprovement in the thermal stability of silicone nanocompositefoams was achieved by the addition of functionalized graphenesheets (FGS).66,112 Verdejo and coworkers66 reported that 0.25 wt% of FGS could both increase the thermal degradationtemperature of silicone foams by 57.7 �C and slow down thethermal degradation rate. In another paper published by thesame group,62 the authors compared the thermal stability ofCNT/silicone foams and FGS/silicone foams, and the datashowed that 0.25 wt% FGS were as efficient as 1 wt% CNT inenhancing the thermal stability of silicone foams. Threereasons were proposed for the enhancement in thermalstability: (1) dispersed nanollers might hinder the ux ofdegradation products, delaying the onset of degradation; (2)polymer chains near nanollers might degrade more slowly;and (3) higher thermal conductivity of nanocomposites couldease heat dissipation within the matrix.

In this section, ve properties of polymer foams were intro-duced in detail. The role of nanollers in the improvement offoam properties was discussed, and recent studies regardingnew nanocomposite foams with superior properties werereviewed. In general, nanollers can improve the properties ofpolymer foams by improving the properties of the matrix, bychanging the structure of the foams through heterogeneousnucleation and by providing additional functionality to the neatpolymer foams. In many cases, nanocomposite foams possess acombination of different superior properties due to the additionof nanollers, such as high EMI shielding effectivenesscombined with high toughness,22 high elastic moduluscombined with high thermal stability,23,62 and high electricalconductivity combined with high compressive modulus.21

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4 Conclusion and outlook

Advances in nanotechnology have opened up signicantopportunities for innovation in functional polymer foams. Forexample, since nanoparticles act as cell nucleation sites, bycontrolling the size, aspect ratio, and surface chemistry ofnanoparticles, the foammorphology including the number andsize of cells can be modied. Nanoparticles can also inuencethe foam morphology by changing the mechanical/rheologicalproperties of the polymer matrix.111 Finally, nanoparticles canadd additional functionality such as conductivity, thermalstability and gas impermeability. Thus, nanoparticles provide anew mechanism for tailoring the structure and properties ofpolymer foams and make polymer nanocomposite foams apromising class of materials to be further studied and used invarious applications.

The literature is full of novel nanocomposite foams, with allkinds of nanoparticles being utilized. However, few studies werefound to have systematically studied the structure–propertyrelationships of polymer nanocomposite foams. Most existingmodels regarding the properties of polymer foams have theirlimitations. New models based on systematic experimentalstudies and theoretical calculations are highly desired.

With the continuing development of nanotechnology, newnanoparticles will be discovered or synthesized, which may ndtheir applications in polymer foams. The inuence of the size,shape and surface chemistry of existing nanoparticles on foamstructure and properties still needs further study. Particularly,new surfactant or new surface modication techniques could bevery useful in facilitating good dispersion of nanoparticles incomposites, which is critical for the large-scale production ofpolymer nanocomposite foams and uniform foam structurecontrol.

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