Multiwalled Carbon Nanotube/Polymer Nanocomposites:Processing and Properties

F. DALMAS,1,2 L. CHAZEAU,1 C. GAUTHIER,1 K. MASENELLI-VARLOT,1 R. DENDIEVEL,2

J. Y. CAVAILLE,1 L. FORRO3

1GEMPPM, INSA de Lyon, 20, Avenue A. Einstein, 69621 Villeurbanne Cedex, France

2GPM2, INP Grenoble, ENSPG, BP46, 38402 St. Martin d’Heres, France

3IPCM, Ecole Polytechnique Federale de Lausanne, CH-1015 Lausanne, Switzerland

Received 9 July 2004; revised 5 January 2005; accepted 5 January 2005DOI: 10.1002/polb.20409Published online in Wiley InterScience (www.interscience.wiley.com).

ABSTRACT: Nanocomposite materials were prepared with an amorphous poly(styrene-co-butyl acrylate) latex as a matrix with multiwalled carbon nanotubes (MWNTs) asfillers. The microstructure of the related films was observed by transmission electronmicroscopy, which showed that a good dispersion of MWNTs within the matrix wasobtained. The linear and nonlinear mechanical behavior and the electrical propertieswere analyzed. Mechanical characterization showed a mechanical reinforcementeffect of the MWNTs with a relatively small decrease of the elongation at break. Thecomposite materials exhibited an elastic behavior with increasing temperature,although the matrix alone became viscous under the same conditions. The electricalconductivity of the composite filled with 3 vol % MWNTs was studied during a tensiletest, which highlighted the late damage of the material. VVC 2005 Wiley Periodicals, Inc.

J Polym Sci Part B: Polym Phys 43: 1186–1197, 2005

Keywords: carbon nanotubes; electrical properties; mechanical properties; nano-composites; reinforcement

INTRODUCTION

Since their discovery in 1991 by Ijima,1 in addi-tion to their amazing electronic properties,2 car-bon nanotubes (CNTs) have exhibited interestingintrinsic mechanical properties. Indeed, experi-mental studies3–5 and theoretical modeling6–8

have demonstrated high Young’s modulus (ca. 1 TPa),stiffness, and flexibility of CNTs. These outstand-ing properties, combined with CNTs’ low density1

and high aspect ratio (up to 1000), offer scope forthe development of CNT-reinforced compositematerials, with polymer matrices in particular.

Some articles have already reported resultsabout CNT/polymer composites with severaltypes of matrices: poly(methyl methacrylate),9–11

epoxy resin,12–14 and polystyrene.15,16 Reportedresults on the mechanical behavior of such mate-rials have shown that the macroscopic behaviorof these composites depends on many factors,such as the types of nanotubes used (single-walled or multiwalled) and their geometric char-acteristics, the nature of the polymer used (CNTscan have a nucleating effect in semicrystallinematrices or can induce a modification of the epoxymatrix crosslinking), and the elaboration process(in situ polymerization,9 solvent method,15 ormelt mixing17).

In a glassy or semicrystalline matrix (whoserigidity is relatively high), the applicability ofsuch flexible nanofibers as a means of improving

Correspondence to: F. Dalmas (E-mail: florent.dalmas@gpm2.inpg.fr)

Journal of Polymer Science: Part B: Polymer Physics, Vol. 43, 1186–1197 (2005)VVC 2005 Wiley Periodicals, Inc.

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mechanical properties is still an open question,as suggested by the contradictory results foundin the literature.11–16,18 In the case of rubberymatrices, CNTs (which have high flexibilityproperties because of their elevated aspect ratio)could form an entangled network within thematrix, whose effect should be visible at largedeformation when this network is stretched.The presence of such a connective network ofCNTs (which are conductive) within an insulat-ing matrix can also be characterized by electri-cal conductivity measurements. An electricalpercolation threshold can be determined butdepends strongly on many parameters.19–23

The goals of this work were to study the effi-ciency of multiwalled carbon nanotubes (MWNTs)as fillers in a soft thermoplastic matrix and tovalidate the interest in in situ alternating cur-rent (ac) electrical measurements under large-strain testing. To do this, model systems wereelaborated and tested. MWNTs were dispersedin water with an anionic surfactant. This stablesuspension was mixed with an aqueous suspen-sion of polymer, that is, a latex, to produce com-posite films. We report here on the mechanicaltesting and electrical measurements of thesematerials.

EXPERIMENTAL

Characterization

A drop of a diluted MWNT suspension in etha-nol was deposited on amorphous carbon-coatedtransmission electron microscopy (TEM) gridsand observed at room temperature with a JEOL200CX transmission electron microscope (for low-magnification observations) and a JEOL 2010FEGmicroscope (for high-resolution imaging).

To check nanocomposite homogeneity, conven-tional TEM observations were also performed oncomposite materials. The samples were cut bycryoultramicrotomy with an Ultracut ReichertS. The temperature of the samples was set to�20 8C, and the speed of the glass knife was4 mm/s. Thin sections (ca. 60 nm thick) weredry-collected and placed on a 400-mesh coppergrid. X-ray diffraction patterns were also col-lected from both pure MWNTs and compositematerials at room temperature. Data were col-lected with a Cu Ka source (k ¼ 1.54 A) and atwo-dimensional imaging plate detector in thetransmission mode.

For mechanical characterization in the linearrange, dynamic mechanical analysis (DMA)measurements were performed in torsion modeat a fixed frequency (0.1 Hz) from 170 to 450 Kat a heating rate of 1 K/min. The sample dimen-sions were around 10 � 3 � 0.6 mm3.

Nonlinear mechanical behavior was analyzedvia tensile tests. Measurements were performedat room temperature, that is, 30 K above theglass-transition temperature (Tg) of the matrixon an MTS device (MTS 1/ME). A constantcrosshead speed was maintained with an initialstrain rate of 2.10�2 s�1.

In situ electrical conductivity measurementswere carried out during a tensile test, as alreadydone for other conductive nanocomposites.24

Parallelepipedic samples (ca. 5 � 15 � 0.7 mm3)were coated at their ends with a silver paint toensure good electrical contact. Electrodes andsamples were carefully isolated from the tensilemachine. The AC complex electrical conductivitymeasurements were performed at the ambienttemperature for several frequencies rangingfrom 10 mHz to 1 MHz with a Solartron 1226bridge with a low applied field of about 1 V/cm.The complex admittance Y* was recorded versustime. From this admittance, the conductivity r*ecan be deduced with the following equation:

r�e ¼ Y� L

WT¼ Y� L

2

V0ð1Þ

where W, L, and T are the width, length, andthickness of the sample during the test, respec-tively, and V0 ¼ L0 � T0 � W0 is the sample ini-tial volume. In Figures 7 and 8, we have plottedthe conductivity in place of the admittancebecause it is an intrinsic parameter that shouldremain constant without any changes in thearrangement of the conductive component,unlike the admittance, which depends on thegeometry variation of the samples. However,eq 1 uses the assumption of a constant volumeof the sample during stretching. This will be dis-cussed further on in this article.

Materials

MWNTs

These were synthesized at Ecole PolytechniqueFederale de Lausanne from the catalytic decompo-sition of acetylene at 720 8C on a supported cobalt/iron catalyst as already described.25 For

CARBON NANOTUBE/POLYMER NANOCOMPOSITES 1187

purification, the raw MWNT sample was sonicatedin 30%HNO3 for 3 h, then filtered and washed withdistilled water, and finally dried at 120 8C. Asshown in Figure 1(a), the purified MWNTs used inthis study were pure, very long (ca. 10 lm), flexible,and entangled. Couteau et al.25 previously showedby energy-dispersive X-ray analysis that, with thispurification process, the peaks for Fe (Ka 6.4 eVandKb 7.06 eV) and Co (Ka 6.93 eVand Kb 7.65 eV) aredrastically reduced. This means that the metallic

catalysts are mainly removed by the purificationstep. A rapid image analysis made on several TEMpictures gives a diameter distribution between30 and 100 nm and centered around 40 nm.Figure 1(b) allows us to appreciate the crystallinestructure of these MWNTs: we can easily see ananotube with 34 walls, that is, 34 graphitic planesparallel to the tube axis with very few defects.

The Latex Matrix

The latex matrix was prepared by the emulsioncopolymerization of styrene (35 wt %) and butylacrylate (65 wt %) at the Laboratoire de Chimie etProcedes de Polymerisation (Ecole Superieure deChimie Physique Electronique Lyon). The polymer-ization occurred in a micellar aqueous solutionof a mixture of an anionic [C12H25O(CH2CH2O)4SO3Na] surfactant and a nonionic [C12H25O(CH2CH2O)19H] surfactant. A surfactant-stabilizedaqueous suspension of poly(styrene-co-butyl acryl-ate) [P(S-BuA)] was obtained containing 43 wt %spherical polymer particles with an average diame-ter of 145 6 10 nm (determined by light scatter-ing). The Tg of P(S-BuA) was determined by differ-ential scanning calorimetry and found to be 265 K.

Composite Processing

The purified nanotubes were dispersed in a dis-tilled water solution of 0.9 g L�1 sodium dode-cylbenzene sulfonate (SDBS; i.e., C18H29SO3Na)surfactant in a sonication step for 5 min with a20-mL suspension volume; this step was per-formed with a Branson Sonifier with a 13-mmprobe tip at 20 kHz and a power source of 25 W.Several weight ratios of the nanotubes to thesurfactant were investigated for solution stabil-ity, and an optimum ratio of 5:1 by weight wasfound; this allowed us to obtain an aqueous sus-pension stable for at least 2 weeks. This weightratio corresponded to 7 surfactant molecules/nm2 of available nanotube surface (the calcula-tion was made with an average tube diameter of40 nm) or a surface available for each moleculeof 14 A2, which seems reasonable.

The MWNT dispersion was stirred with thislatex suspension to obtain composite films withMWNT volume fractions of 1.5 and 3 vol %.After being mixed, the preparations were cast inan aluminum mold with a Teflon coating andput in a drying oven at 30 8C in vacuo for 5 daysto allow water evaporation and film formation(i.e., particle coalescence).

Figure 1. (a) Low-magnification and (b) high-resolu-tion TEM micrographs of CNTs after purification innitric acid.

1188 DALMAS ET AL.

RESULTS AND DISCUSSION

Both composite materials have a good level ofdispersion of the CNTs within the P(S-BuA)matrix; this is demonstrated in Figure 2 for 3vol % MWNTs. Indeed, no aggregate can beobserved, and the nanotubes seem to be isolatedand homogeneously distributed. On a compositemicrosection, MWNTs seem to be shorter andless entangled than they are in the correspond-ing macroscopic sample.

MWNTs have been studied previously byX-ray diffraction techniques.26–28 The diffractionspectrum is dominated by a strong Bragg peakcentered at 3.4 A, which corresponds to theintershell spacing within the nanotubes [re-ferred to as the (002) peak]. This slightlylarger d-spacing (compared with that of graph-ite, 3.35 A) is related to the turbostratic struc-ture of CNTs.27 Figure 3(a) displays the powderdiffraction pattern of pristine nanotubes. As dis-cussed earlier, the (002) Bragg peak was found

to be centered at 26.18 in 2h (which correspondsto a d-spacing of 3.41 A), with a constant inten-sity with respect to the azimuth angle. This dia-gram is a classical randomly oriented MWNTdiagram. Figure 3(c) shows the two-dimensional(2D) X-ray diffraction pattern of a compositefilled with 3 vol % MWNTs in the transmissionmode when the X-ray beam is perpendicular tothe cast film plane. This diffraction pattern isthe superposition of the nanotube diffractionring (with the same intershell spacing) with thatof the amorphous matrix [Fig. 3(b)]. Jin et al.28

showed that after nanotube alignment within acomposite (by mechanical stretching), the (002)Bragg peak intensity is no longer constant withthe azimuth angle, /; rather, it is concentratedat two spots centered at / & 90 and 2708. It isthus obvious that there is no tube alignment inthe composites studied here and that the nano-tube dispersion is isotropic in the plane. Whenthe X-ray beam is parallel to the cast film plane,the same diffraction pattern shown in Figure 3(c)

Figure 2. Low-magnification TEM micrographs of a microsection from a P(S-BuA)film filled with 3 vol % CNTs.

CARBON NANOTUBE/POLYMER NANOCOMPOSITES 1189

is obtained. One could have expected a decreasein the Bragg peak intensity at the equator.Indeed, it is well known that composites made bythe evaporation process present an orientation ofthe fillers in the cast film plane. In the case ofrigid stick fillers, this anisotropy can be high-lighted by 2D X-ray diffraction, with a decreasein the peak intensity at the equator. In our case,it is difficult to discuss this phenomenon becauseof the nanotubes’ flexibility properties, which canexplain the fact that the diffraction still occurs inall directions even if the material presents a pla-nar anisotropy.

The storage modulus G0 (which is indicative ofthe elastic behavior of the polymer) and the lossfactor tan d, determined by DMA measurementsfor the pure P(S-BuA) matrix and compositematerials filled with 1.5 and 3 vol % MWNTs,are plotted versus the normalized temperaturein Figures 4 and 5, respectively. DMA testinghighlighted the fact that the SDBS surfactanthas a plasticizing effect on the P(S-BuA) matrix.Indeed, if we look at the main relaxation temper-ature Ta (defined as the temperature at whichthe tan d peak occurs), this mechanical relaxa-tion occurs at 282 K for pure P(S-BuA) and isshifted to 280 and 278 K with 1.5 and 3 vol %MWNTs, respectively. It has been confirmed bydifferential scanning calorimetry with a PerkinElmer Pyris Diamond apparatus at a heatingrate of 10 K/min that Tg of P(S-BuA) is decreasedby the addition of SDBS by 4 K with 3 vol %nanotubes. For this reason, the DMA results areplotted versus the normalized temperature, T/Ta.The unfilled P(S-BuA) matrix displays classicalbehavior of an amorphous thermoplastic poly-mer; below Ta, the polymer is in the glassystate with a roughly constant modulus (ca. 1GPa). When Ta is reached, a strong decrease inG0 (by more than 3 decades) occurs, correspond-ing to a peak in tan d. Above Ta the material isin the rubbery state until the material starts toflow (here at approximately 350 K); then itbehaves as a viscous fluid and the experimentalapparatus fails. As shown in both Figures 4and 5, when the polymer is filled with CNTs,this temperature is strongly increased up to425 K for both nanocomposites. This phenom-enon is certainly due to the formation of a

Figure 3. 2D X-ray diffraction patterns obtainedfrom (a) the purified MWNT powder, (b) the P(S-BuA)matrix, and (c) the 3 vol % MWNT composite.

1190 DALMAS ET AL.

connective network of CNTs within the polymermatrix. The entanglement level of this networkand the interactions between the polymer chainsand the nanotube surface can explain the delayof the material flowing and the existence ofthis wide modulus plateau. In addition to thisphenomenon, the value of G0 above Ta isstrongly increased for MWNT-filled composites.For instance, at the ambient temperature(300 K), G0 of the film containing 3 vol %MWNTs is six times higher than that of theunfilled P(S-BuA).

Various models exist for the mechanical prop-erties of heterogeneous materials. The Halpin–Kardos29 model is a mean-field approach for iso-tropic short-fiber composites. This approach usesthe Halpin–Tsai30 semiempirical equation, whichdescribes the modulus of an aligned short-fibercomposite, Mc, by

Mc

Mm¼ 1þ ngVf

1� gVfð2Þ

where g is given by

g ¼ ðMf=MmÞ � 1

ðMf=MmÞ þ nð3Þ

where Vf is the volume fraction of fibers, Mf isthe fiber modulus, Mm is the matrix modulus,and n is a factor that depends on the shape ofthe filler particle and on the type of modulus Mto be calculated. The ply shear modulus Gc canbe predicted with n equal to 1, the longitudinalmodulus E11 can be predicted with n equal to2 L/d (L and d are the fiber length and diameter,respectively), and the in-plane moduli E22 andE33 can be predicted with n equal to 2.31 Halpinand Kardos29 considered short-fiber isotropiccomposites as laminated composites composed ofunidirectional plies. The modulus of each ply isgiven by eq 2. In this work, the calculation usedfour plies; each ply is rotated 458 from the pre-vious ply. The continuous curve in Figure 4shows the result of this model for a compositefilled with 3 vol % fibers. To fit the experimentaldata [Fig. 4(c)] in the glassy state, we used forthe fibers a longitudinal modulus Ef11 equal to300 GPa and transversal moduli Ef22 and Ef33

Figure 4. G0 as a function of normalized temperature for (a) the unfilled P(S-BuA)matrix, (b) the related composite filled with 1.5 vol % CNTs, and (c) the related com-posite filled with 3 vol % CNTs. (Results from the Halpin–Kardos model with 3 vol %fibers are shown as a solid line.)

CARBON NANOTUBE/POLYMER NANOCOMPOSITES 1191

equal to 15 GPa. This value of 300 GPa for theMWNT longitudinal Young’s modulus is lowerthan what is found in the literature [Demczyket al.5 found (with bending tests) a longitudinalmodulus of 910 GPa for MWNTs]. As shown inFigure 2, our explanation is that the CNTsremain curved and entangled within the compo-site. Indeed, Fisher et al.32 showed by finite ele-ment calculation that the effective modulus ofthe tubes in composite materials is lower thantheir real longitudinal modulus and depends ontheir aspect ratio and their waviness ratio. Inthe rubbery state, the fiber aspect ratio is themain parameter influencing the reinforcementlevel. To describe the modulus of the 3 vol %composite above Tg, an aspect ratio (ratio of thelength to the diameter of the fiber) of 200 isused. Indeed, this value is reasonable if we lookat Figure 1(a), which shows that all the tubesseem to have a length higher than 5 lm with anaverage diameter around 40 nm.

However, the mean-field approach alone can-not explain the fact that the nanotubes delaythe flow of the material. Such a model takes intoaccount the modulus of each phase, as well as

the geometry of the particles, but the particlesare assumed to have no interaction with oneanother. Thus, neither the effect of entangle-ments between the tubes nor the possible inter-actions between the nanotubes via the matrix istaken into account. A new modeling approachbased on the discretization of CNTs has to bedeveloped to characterize the influence of thesetube entanglements and of the strength of thenanotube–matrix–nanotube contacts on themechanical behavior of composites.

Tensile tests were performed on these compo-sites. The corresponding results are plotted inFigure 6. The unfilled P(S-BuA) exhibits thetypical behavior of an amorphous thermoplasticabove Tg. First, the stress increases with thestrain; this is characteristic of the entropic elas-ticity due to the entanglements between poly-mer chains. Then, it reaches a maximum anddecreases with increasing strain because ofchain disentanglement. A film of the surfactant(without nanotubes) combined with the polymermatrix was processed and tested in the sameway. This test showed that the only effect of sur-factant on the matrix is a plasticizing effect

Figure 5. Tan d versus the normalized temperature for (}) the unfilled P(S-BuA)matrix, (�) the related composite filled with 1.5 vol % CNTs, and (*) the relatedcomposite filled with 3 vol % CNTs.

1192 DALMAS ET AL.

already shown by DMA results. P(S-BuA) sam-ples, both pure and with surfactant, display thesame tensile behavior at the same T/Ta ratio.With the addition of MWNTs, the material keepsthe same type of behavior. However, as deter-mined from the initial slope of the curves inFigure 6, the Young’s modulus of the compositesis increased with the presence of nanotubes (ina way consistent with the DMA results). Themost interesting point highlighted by these ten-sile tests is that, in addition to this mechanicalreinforcement, even the composite with thehighest content keeps a high ultimate strain (>2in true strain, i.e., >600% stretching).

CNTs, which are good conductive fillers, allowthe composite to be conductive beyond a certainfiller content, which is called the electrical per-colation threshold.19–23 Figure 7 shows the evo-lution of the real part of the electrical conductiv-ity with the frequency for the materials studiedhere. The P(S-BuA) matrix displays classicaldielectric material behavior, with the electricalconductivity increasing with the frequency forhigh frequency values (>1 Hz). For the compo-site filled with 1.5 vol % CNTs, the same type of

behavior can be observed with a conductivitylevel higher than that for the matrix; that is,the composite displays a capacitive behavior.This result corresponds to a conductive filler/polymer composite below the percolation thresh-old. The same type of result was obtained in thesame range of conductivity values by Sandleret al.19 for epoxy/MWNT composites. The electri-cal behavior of the 3 vol % MWNT material iscompletely different; the conductivity is muchhigher (28 S/m) and remains constant with thefrequency: the material exhibits purely realresistivity. In this case, the material is abovethe percolation threshold; there is a connectivenetwork of nanotubes within the matrix. Theelectrical percolation threshold is thereforebetween 1.5 and 3 vol % CNTs.

One can try to model this composite conduc-tivity, rcomp, with a simple mixing law:

rcomp ¼ Vmatrmat þ Vnetrnet ð4Þ

where Vnet is the volume fraction of the conduc-tive network, Vmat is the volume fraction of thematrix embedding it, and rnet is the conductivity

Figure 6. True stress–strain curves for tensile tests performed on (a) the unfilledP(S-BuA) matrix, (b) the related composite filled with 1.5 vol % CNTs, and (c) therelated composite filled with 3 vol % CNTs.

CARBON NANOTUBE/POLYMER NANOCOMPOSITES 1193

of the network. A first scheme considers the com-posite as a pure nanotube network embedded bythe polymer matrix with the surfactant. Then, inthe 3 vol % MWNT composite, Vnet is 3% and ameasurement of the conductivity of a pure nano-tube film gives a rnet value equal to 1285 S/m.From eq 4, the theoretical rcomp value is found tobe 45 S/m, which is higher than the experimentalvalue. A second scheme considers the compositeas an SDBS–nanotube network embedded inthe composite. Then, Vnet is taken to be 4.7 vol %,and rnet, deduced from the measurement on afilm prepared by evaporation of the SDBS–nano-tube stable aqueous suspension, is taken to be330 S/m. The deduced value for the composite isthen 16 S/m, which is below the experimentalvalue.

Thus, the experimental value is between thevalues calculated from the two schemes. Thissuggests that the conductivity of the compositeresults from the conductivity of a nanotube net-work whose nanotube contacts are influenced bythe presence of the surfactant. This fact is con-firmed by the conductivity values measured forthe pure nanotube film and the film preparedfrom the evaporation of the SDBS–nanotube

suspension. Indeed, the conductivity of the lat-ter, 330 S/m, is lower than the conductivity cal-culated from a mixing law based on the assump-tion that this film is a nanotube network in anSDBS matrix (850 S/m).

To understand the microstructural evolutionof these materials during the tensile test, ACelectrical conductivity for different frequencieswas followed in situ for the 3 vol % MWNT-filledcomposite; related results of the real and imagi-nary parts of the conductivity are shown inFigure 8. Note that the evolution of the electri-cal conductivity is calculated with the assump-tion that the sample volume remains constantduring the deformation. For a high strain level,this assumption is questionable. However, thecalculation with a constant sample cross section(i.e., with the assumption of the maximumincrease of the sample volume during stretch-ing) gives the same evolution of conductivity,only shifted toward lower deformation levels.

First, it can be observed that the materialremains conductive until it reaches a high strainlevel, e ¼ 1.7 (which corresponds to the yield pointof the material, as shown in Figure 6). Indeed, inthis region, the real part of the conductivity is

Figure 7. Real part of the conductivity versus the frequency of the applied field forunstrained samples of (}) the pure P(S-BuA) matrix, (�) the composites filled with1.5 vol % MWNTs, (*) the composite filled with 3 vol % MWNTs, (*) the SDBS sur-factant, (n) the SDBS/MWNT film, and (u) the pure MWNTs.

1194 DALMAS ET AL.

very high, on an order of magnitude of 20 S/mand almost independent of the frequency. Thisindicates a resistive behavior. As previously dis-

cussed for frequency-dependence measurementson the unstrained samples, this high level of elec-trical conductivity signifies the presence of a CNT

Figure 8. (a) Real and (b) imaginary parts of the electrical conductivity versus thetrue strain under tensile testing for the composite containing 3 vol % CNTs.

CARBON NANOTUBE/POLYMER NANOCOMPOSITES 1195

percolating network within the matrix. One canobserve a small increase in the real conductivityat the beginning of the deformation that can bedue to a rearrangement of the nanotubes, whichin turn leads to an increase of contact nodesbetween neighboring nanotubes. After havingreached a maximum, the real conductivity slowlydecreases with increasing strain, whereas theimaginary part increases, up to the yield point ofthe material. At this deformation level, the realpart of the conductivity begins to decreasestrongly and becomes dependent on the fre-quency, whereas the imaginary part passesthrough a peak indicative of material damage. Athigher deformation levels, the material is dam-aged, that is, the connective network of nano-tubes responsible for electrical percolation is bro-ken. The nanotubes are then no longer in resis-tive contact but form capacitances betweenthemselves, as shown by the capacitive behaviorof the material.

CONCLUSIONS

MWNT/P(S-BuA) matrix composites were proc-essed with two different concentrations ofMWNTs. The use of an anionic surfactantallowed stabilization of the nanotubes in water,which led to a good level of dispersion (i.e.,random) of MWNTs in the composites. Twoeffects of the CNTs on the mechanical behaviorof the material were highlighted: a strongimprovement in thermomechanical stability andimproved mechanical reinforcement. The in-crease in modulus is quite low in comparisonwith other high aspect-ratio fillers such as cel-lulose whiskers, but MWNTs allow the materialto keep a high level of strain at break. Theyallow the material to be electrically conductingwith a filler content above 1.5 vol %. With 3 vol %MWNTs, the material remains conductive up tohigh strain because of the late damage of the con-nective network of entangled flexible fillers.

The authors thank C. Graillat (LCPP–CPE Lyon) forhis help and fruitful advice on latex synthesis, as wellas M. Kneveler for his contribution to some experi-mental data. The work in Lausanne was supported bythe Swiss National Science Foundation and its NCCR‘‘Nanoscale Science’’. This work is performed in theframe of the European CNT-network and the GDREno 2756 ‘‘Science and applications of the nanotubes –NANO-E’’.

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