C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5

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Improving reinforcement of natural rubber by networkingof activated carbon nanotubes

Sanjib Bhattacharyyaa, Christophe Sinturela, Ouziyine Bahloula, Marie-Louise Saboungia,Sabu Thomasb, Jean-Paul Salvetata,*

aCentre de Recherche sur la Matiere Divisee, CNRS Universite d’Orleans, UMR6619, 1B rue de la Ferollerie, 45071 Orleans cedex 2, FrancebSchool of Chemical Sciences, Mahatma Gandhi University, Priyadarshini Hills P.O. Kottayam, Kerala 686 560, India

A R T I C L E I N F O

Article history:

Received 31 January 2008

Accepted 14 March 2008

Available online 19 March 2008

0008-6223/$ - see front matter � 2008 Elsevidoi:10.1016/j.carbon.2008.03.011

* Corresponding author: Fax: +33 2 38255376.E-mail address: salvetat@cnrs-orleans.fr

A B S T R A C T

Reinforcement of natural rubber was achieved using carboxylated multiwalled carbon

nanotubes (c-MWCNT) dispersed with sodium dodecyl sulfate. The structure of the rein-

forced latex films was investigated by TEM and AFM. The tensile and dynamic-mechanical

tests demonstrated a strong enhancement in the Young’s modulus (�10-fold), tensile

strength (�2-fold) and storage modulus (�60-fold) at low-strain in the rubbery state with

up to 8.3 wt% of MWCNTs, with a small reduction in elongation at break. Dielectric mea-

surement at room temperature revealed a low percolation threshold (<1 wt%) associated

with the formation of an interconnected nanotube network. Latex film formation plays a

critical role in the network formation due to the segregation effect at the surface of latex

beads. We observed large Payne and Mullins effects due to the mechanical behavior of

the nanotube network. The disruption of the network during stretching induces both an

increase of electrical resistivity and mechanical stress-softening.

� 2008 Elsevier Ltd. All rights reserved.

1. Introduction

Polymer/nano-filler composites have received intense atten-

tion and research in the past decade, driven by the unique

properties of nanostructures and potential to create new

materials with superior properties. The goal of polymer/

nano-filler composite sciences is not only to fill a matrix with

nanometer sized particles, but also to modify the matrix tex-

ture by interaction with nano-fillers during processing, in or-

der to make new functional materials, especially with

improved mechanical properties. In the case of reinforced

polymer composites, it is important to distinguish two ther-

modynamic regimes (glassy and rubbery regions) related to

the viscoelastic behavior of polymer matrices. This because

of the different reinforcing mechanisms involved in the

glassy and rubbery regimes of polymeric matrices. In the

er Ltd. All rights reserved(J.-P. Salvetat).

glassy state of the polymer matrix (T < Tg the glass transition

temperature), reinforcement is achieved with high aspect ra-

tio fibers having modulus at least one order of magnitude

higher than that of the matrix. The high specific surface area

of the nano-fillers is expected to provide enhanced interphase

effects and tensile strength. Reinforcement depends essen-

tially on the efficiency of load transfer to the fibers and is

influenced by the dispersion state and aspect ratio of the fill-

ers, and as well as on the quality of the interphase. High levels

of reinforcement are achieved when all factors are carefully

controlled. However, in the rubbery state of the matrix

(T > Tg), elasticity is dominated by the entropy effects rather

than elemental bond stiffness [1–3]. In this regime a first level

of reinforcement is ‘‘hydrodynamical’’ in origin and can be

achieved with dispersed nanoparticles [4–8]. In the second le-

vel of reinforcement filler–matrix and filler–filler interactions

.

1038 C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5

play an important role in the mechanical properties of filled

elastomers due to the increase in cross-link density and the

occlusion of a certain amount of the elastomeric matrix by fil-

ler aggregates. Formation of rigid filler network usually pro-

vides additional levels of reinforcement. In the case of silica

particles and carbon black as fillers, it has been shown that

the aggregation and networking of fillers improve the

mechanical properties compared to a dispersed state [9–11].

Filler–matrix and filler–filler interactions are unanimously

associated with the decrease of the dynamical modulus with

the strain amplitude, which is known as the Payne effect [12–

14]. Complementary approaches to understand the Payne ef-

fect [13,15] involve filler network formation–disruption [12]

and adsorption–desorption of polymer chains at the surface

of filler [16]. Filler network is also involved in strain-softening

or Mullins effect [17]. Reinforcement of rubber compounds

dates back to 1904 when carbon black was used as nano-fill-

ers. Also in present days carbon black is the principal rein-

forcement element for natural rubber, along with silica

[11,18–21] and clay nanoparticles [22–25]. A huge amount of

data is available on the topic, especially on the influence of

surface treatment and morphology of the fillers on the rein-

forcing efficiency. Short fibers of various types have also been

studied as reinforcing elements for elastomers [26–29]. The

advantage of elongated, fibrous structures over the spherical

ones has been emphasized, and the networking ability of nat-

ural fibers is particularly interesting for increasing the modu-

lus at low filler content [30–32].

Due to very good intrinsic mechanical properties, high

aspect ratio, low density, and high surface area, carbon

nanotubes (CNTs) or nanofibers seem to be ideal candidate

for the application as fillers in polymer matrices [33]. Most

efforts have dealt with glassy matrices, such as polyvinyl

alcohol, polymethyl methacrylate, epoxy, etc, where signifi-

cant level of reinforcement has been achieved [33]. By con-

trast, there are few reports on using carbon nanotubes as

reinforcing fillers in rubber matrices, although the first work

by Froglet et al. showed that MWCNTs are very promising

for this application [34]. In the specific case of natural rub-

ber (NR), Bokobza and Kolodziej have made detailed investi-

gation of the reinforcing effect by MWCNT and also

concluded that CNTs are efficient reinforcing elements

[35]. They suggest that surface modification can be used to

modify the reinforcing ability of CNTs. Atieh et al. also

achieved significant reinforcement of natural latex with

MWCNT [36].

In this work, we have used ‘‘activated’’ MWCNTs to improve

the mechanical properties of NR especially in the low-strain

dynamical regime, reaching the same level of reinforcement

obtained with the inclusion of cellulosic fibers in NR matrix

[31]. We argue that the large increase of storage modulus is

related to the formation of a network of non-covalently

cross-linked CNTs during latex film formation, as evidenced

by structural and electrical characterizations. In particular,

the most exciting part of the observation is that the reinforce-

ment occurs mainly at low-strain values, whereas at high-

strain levels the behavior remains very similar to that of pure

natural rubber. This implies that the networking effect of CNTs

dominates over the ‘‘hydrodynamical’’ contribution at low-

strain amplitude.

2. Experimental

The two main materials we used in this work are as follows:

(1) Prevulcanised natural rubber ‘HR Revultex’ in ammonia

solution were kindly supplied by the company Revertex

(Malaysia) SDN. Berhad, Malaysia and the total solid

content was 60.5%.

(2) Multiwalled carbon nanotubes (MWCNTs) were synthe-

sized in our laboratory by catalytic decomposition of

acetylene at 600 �C on a CoxMg(1�x)O solid solution [37].

2.1. Purification and carboxylation of MWCNTs

Prepared MWCNTs by CVD technique contain amorphous car-

bon and metallic nanoparticles (used as catalyst during syn-

thesis) as impurities. Metal nanoparticles were removed by

treating the raw nanotubes with 37% hydrochloric acid at

80 �C. For removal of amorphous carbon, MWCNTs were heat

treated at 500 �C in the presence of carbon dioxide for 24 h.

For creation of carboxyl functionality on the side walls of

MWCNTs the purified nanotubes were refluxed in 2.8 M nitric

acid for 38 hrs, and washed several times with deionized

water until the filtrate shows neutral pH. Presence of carboxyl

groups has been confirmed by X-ray photoelectron

spectroscopy.

2.2. Preparation of MWCNT/natural rubber composites

For the preparation of composites containing 8.3 wt% of

MCWNTs, 120 mg of MWCNTs and 120 mg of sodium dodecyl

sulfate (SDS) were dispersed in 20 ml of water by sonicating

the mixture in bath sonicator for 15 min. The dispersion

was then added to the ammonia solution of rubber containing

1.2 g of prevulcanized natural rubber. Mixture was stirred

magnetically for 24 h, sonicated for 15 min in a bath sonica-

tor, and then poured into a Petri dish and dried at 60 �C for

24 h to obtain freestanding composite films. For the compos-

ite films with different wt% of MWCNTs we kept the amount

of natural rubber fixed at 1.2 g and we varied the amount of

MWCNTs and SDS keeping the same weight ratio between

them as indicated earlier. We have prepared composite films

with five different wt% of MWCNTs, i.e., 1, 2, 2.8, 5.4 and

8.3 wt%.

2.3. Characterization

Purified MWCNTs were characterized by transmission elec-

tron microscopy (TEM) on carbon coated copper grid in Philips

CM-20 TEM. Fig. 1 shows the TEM picture of purified and car-

boxyl functionalized MWCNTs. High resolution picture in the

inset shows the ‘‘crystalline’’ nature of the inner shells of

MWCNTs while outer shells are significantly damaged by

the oxidation treatment. TEM image analysis gives a diameter

distribution between 10 and 20 nm.

Surface chemistry of MWCNTs after oxidation was investi-

gated by XPS, and approximately 10 at% of C–O and C@O

bounds were found grafted at the surface.

Atomic force microscopy (AFM) images were obtained in

air at room temperature using the intermittent contact mode

Fig. 1 – TEM images of activated CNTs. The chemical

oxidation process allows activating the surface without

destroying the morphology. In inset, HR image showing the

damaged outer shells of activated CNTs.

Fig. 2 – Stress–strain curves for pure latex films and

composites.

Fig. 3 – Reduced stress, [r*] ¼ r/(a� 1a2), vs. inverse extension

ratio, a�1 (Mooney–Rivlin plot). The strong increase of [r*] at

low-strain for highest filler concentration is remarkable and

cannot be explained by a simple hydrodynamical

reinforcement.

C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5 1039

of a Pico+ AFM (Molecular Imaging), with standard Si cantile-

vers resonating at approximately 300 kHz.

The non-linear mechanical behavior was measured using

a Llyod LF Plus testing machine in tensile mode, with a load

cell of 1 kN capacity. The specimen was a dog bone shaped

thin strip (20 · 4 · 2 mm) and the experiment performed with

a cross-head speed of 30 mm/min. Stress versus strain curves

were plotted and the tensile or Young’s modulus (E) was mea-

sured from the slope of the low-strain region in the vicinity of

r = e = 0 ([dr/de]e!0). Five samples were tested for each com-

posite film and the averages of the values were taken.

The dynamic-mechanical thermal analysis (DMTA) spectra

were taken on rectangular specimens (20 · 10 · 0.1 mm) in

tensile mode at a frequency of 10 Hz using a Metravib DMTA

equipment. DMTA spectra, viz. storage modulus (E 0) and

mechanical loss factor (tand) were measured in the tempera-

ture range 170–290 K at a heating rate of 2 K/min. For Payne

effect, samples were measured at room temperature, at

10 Hz, in the dynamic strain range 5–450 lm.

AC conductivity was measured by Agilent 4284A imped-

ance meter at room temperature in the frequency range

100 Hz–1 MHz. Non-metallized samples were sandwiched be-

tween two copper electrodes in capacitor configuration for

the most resistive samples. For 5.4 and 8.3 wt% filled samples,

longitudinal resistivity could be measured as well with mm or

cm gaps between electrodes. DC resistance was extrapolated

from the measured AC resistance values.

3. Results and discussion

3.1. Tensile behavior and hydrodynamical effects

The true stress vs. strain curves for unfilled natural rubber

matrix and up to 8.3 wt% carbon nanotubes NR composites

at 20 �C are represented in Fig. 2. In all cases the typical

non-linear ‘‘S curve’’ described in the kinetic theory of rubber

elasticity [1] is observed. It is useful in that case to plot a so-

called reduced (or Mooney) stress ½r�� ¼ r=ða� 1a2Þ as a function

of the extension ratio a, or a�1 (Fig. 3).

In the James and Guth approach, [r*] = E, the elastic modu-

lus, which is proportional to the cross-link density and to the

temperature [1]. In the Mooney–Rivlin model for large defor-

mation, [r*] varies linearly with a�1, according the expression

½r� ¼ 2C1 þ 2C2a�1� [1]. Although many refinements have been

developed to describe the functional dependence of r vs. a

[38,39], we limit our analysis within the Mooney–Rivlin repre-

sentation, which is often used by rubber scientists to discuss

Fig. 4 – Initial tensile modulus vs. CNT volume fraction.

Dashed line is a fit using the Guth formula for anisotropic

fillers. Full line is a fit using a percolation approach suited to

describe formation of the nanotube rigid network.

Fig. 5 – Two loading–unloading cycles performed on 8.3 wt%

filled film. The modulus at low-strain strongly decreases

after the first loading, suggesting a disruption of the

nanotube network. A permanent deformation is observed

after first loading.

1040 C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5

stress–strain curves. Fig. 3 shows that the modulus (i.e., [r*]) is

increasing with the increase of filler volume fraction, which is

in qualitative agreement with the Mooney–Rivlin equation, at

least in the moderately large deformation regime. The upturn

observed at low a�1 is related to stress-induced crystallization

of natural rubber [1,40,41]. However, the strong increase of the

reduced stress [r*] in the range 0.4 < a�1 < 1 (see 8.3 wt% CNT

filled NR on Fig. 3), which corresponds to a decrease of the

modulus when strain increases, is rather unusual; a simple

hydrodynamic reinforcement is expected to be weakly strain

dependent. A similar behavior although lower in magnitude

has been observed for unvulcanized NR reinforced with

in situ generated silica nanoparticles [42]. In their case the ef-

fect disappeared after vulcanization, and the authors ex-

plained that in the unvulcanized state silica particles tends

to agglomerate and form a rigid network [42].

Description of filled elastomers is a hard task from the

point of view of theory [43] since one has to deal with an

effective medium composed of two very different materials

in terms of elasticity. Approaches used for composites when

the matrix is in its glassy state are poorly suited to filled elas-

tomers. The difficulty lies in introducing the effect of fillers in

the constitutive equation of rubber, i.e., expressing the re-

duced stress at least as a function of filler fraction. One of

the earliest attempts [44] introduced the notion of amplified

strain, with the amplified stretch written as K = 1 + X(a � 1),

where X is a constant depending of the volume fraction of fill-

ers. The stress–strain is then described by r ¼ EðK� 1=K2Þ.According to the Einstein–Smallwood–Guth hydrodynamical

model (ESG model) [1], the reinforcement of elastomers by

rod like particles only depends on the aspect ratio and volume

fraction of the fillers

E ¼ E0ð1þ 0:67f/þ 1:62f 2/2Þ ¼ E0 � XESG ð1Þ

where E and E0 are the moduli of the composite and the pure

matrix, respectively, and f and / are the aspect ratio and vol-

ume concentration of fillers. XESG was used originally as the

amplification constant by Mullins and Tobin [44]. Several works

have extended the ESG model and all converge to a virial-type

expansion of X in powers of the volume fraction [45].

The variations of Young’s modulus as determined from the

slope of the stress vs. strain curves (Fig. 2) in the vicinity of

r = e = 0, with carbon nanotube volume fractions are shown

in Fig. 4. The dashed line in Fig. 4 represents the best fitting

curve by Eq. (1), giving a value of 32 for the aspect ratio of

the fillers, which is in reasonable agreement with the struc-

tural characteristics deduced from TEM observation. How-

ever, this hydrodynamic approach cannot explain the rapid

decrease of E when strain increases as already mentioned

(Fig. 3). So it is necessary to include filler–matrix and filler–fil-

ler interactions to interpret the tensile behavior of our sam-

ples at low and moderate strain. We have also fitted our

data with a percolation model such as used for networks of

cellulose whiskers [31], but actually with limited success. It

is more realistic in fact to expect both hydrodynamical and

networking effects to be involved in the reinforcement. We

believe that the network approach is more appropriate to take

into account the decrease of E with strain, and a number of

other results as described below. The hydrodynamical one is

valid at low filler concentrations before network formation,

and at high-strain when network destruction occurs.

Loading–unloading cycles showed a strong stress-soften-

ing (Mullins) effect especially at 8.3 wt% CNT filling (Fig. 5),

which is hardly explained by a hydrodynamical model [46].

It has to be noted, however, that the minimum modulus for

8.3 wt% CNT is still two times that of pure NR (0.63 vs.

0.25 MPa, Fig. 3), a value reached only with high volume frac-

tion (approximately 50%) of carbon black or silica particles.

In contrast to the Young modulus, the tensile strength of

the composites is found to increase linearly with CNT con-

Fig. 6 – (a) Storage modulus and (b) tand vs. temperature for

NR (s) and composites with 1 (m), 2.8 (�), 5.4 (d) and 8.3 (h)

wt% MWCNTs. Inset of (b) shows that maximum of tand

decreases linearly vs. MWCNTs concentration.

C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5 1041

centration (plot not shown), while elongation at break de-

creased (Fig. 2). There is almost 20% reduction in strain to fail-

ure in the case of composite with 8.3 wt% filler. During the

experiment it was observed that the strain is homogeneously

distributed in macroscopic scale and uniform along the whole

sample, until it breaks and the absence of any necking phe-

nomenon suggests the homogeneous nature of our composite

films.

During the first loading it is highly probable that some de-

gree of CNT alignment occurs in the strain direction [47]. This

effect if present should increase the effective modulus. Such

an increase is indeed observed at high elongation (Fig. 3) but

usually attributed to recrystallization of rubber. Further study

would be needed to quantify alignment effects properly.

3.2. Dynamical analysis

To go further in the understanding of reinforcement mecha-

nisms by MWCNTs, dynamic-mechanical measurements

were performed on pure natural rubber matrix and up to

8.3 wt% carbon nanotubes composites, as a function of tem-

perature and strain amplitude. The goal is to evaluate the de-

gree of filler–rubber or filler–filler interactions, which provide

additional levels of static and dynamic reinforcement.

Fig. 6 shows the variation of storage modulus (logE) (Fig. 6a)

and the loss factor tand (Fig. 6b) at 10 Hz as a function of tem-

perature (170–290 K) for the various nanotube compositions.

At low temperature the polymer is in the glassy state with a

modulus around 0.6 GPa. With increasing temperature the

elastic tensile modulus suddenly drops down by 3 orders of

magnitude corresponding to the glass–rubber transition. This

modulus drop can be ascribed to an energy dissipation

phenomenon involving cooperative motions of long chain

sequences. This phenomenon is also reflected in the corre-

sponding relaxation process where the loss factor tand passes

through a maximum around 205 K. It may be of interest to

note that the temperature corresponding to the maximum of

tand decreases when MWCNT concentration increases. A sim-

ilar effect was observed with cellulosic whiskers in natural

rubber [31]. The inverse trend is usually observed in reinforced

composites and interpreted as a decrease of chain mobility (or

equivalently an increase of Tg) due to interaction with nano-

fillers. The origin of the apparent decrease of Tg is not clear

at the moment and necessitates more investigation.

For the composite films in the glassy state (T < 200 K), there

is a significant increase in the storage modulus for 1 wt%

(�1.82 GPa) of filling, after that it has an almost constant va-

lue rather than a linear increase such as predicted by the

rule-of-mixture. This suggests that nanotube aggregation oc-

curs even at low volume fraction during film formation,

which decreases the effective load transfer efficiency [33].

Above Tg, the composite films exhibit a huge increase in

the storage modulus with increase in the fillers wt%, for

example in case of composite with 8.3 wt% of carbon nano-

tubes the relaxed modulus at Tg + 70 K is 75 times higher than

that of the pure matrix. Moreover, the initial tensile modulus

determined from tensile tests is lower than the storage mod-

ulus for CNT-NR composites where as they are equal for pure

NR films. This is to be related to the observed Payne effect

when the dynamic strain amplitude is increased. The differ-

ence between initial tensile modulus E(0) and E 0 increases

with the filler content and decreases when strain amplitude

increases (Fig. 7). This strong Payne effect is consistent with

the strong Mullins effect observed in cyclic tensile tests, and

both have probably the same origin.

3.3. CNT networking

All these results suggest that there is a significant contribution

from the networking of CNT on the static and dynamic-

mechanical properties of the composites. Formation of this

network can be followed by electrical measurement. The plot

of the extrapolated DC resistivity versus the MWCNTs loading

expressed in volume fraction is depicted in Fig. 8.

The solid line represents the fitting curve with a simplified

bond percolation model r ¼ r0j/� /cjt for / > /c) which gives

the percolation threshold /c � 0.004 and power law exponent

t � 3. We have also measured the resistivity of the samples

along the tensile direction before and after stretching up to

Fig. 7 – The storage modulus of the MWCNT-filled NR films

decreases when the strain amplitude increases (Payne

effect). Measurements were performed at room temperature.

Fig. 8 – Resistivity vs. CNT volume fraction, showing an

electrical percolation behavior due to CNT network

formation.

1042 C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5

500 % extension. The high-strain induces an irreversible in-

crease of the resistivity of one order of magnitude for the

8.3 wt% filled samples, demonstrating that the nanotube net-

work is partially broken. We therefore correlate this increase

of resistivity to the mechanical softening observed on the

same samples.

Following Grunlan et al. [48,49], we are tempted to attri-

bute the nanotube network formation to segregation effects

due to latex film formation. TEM (Fig. 9a and b) investigations

showed clearly that nanotubes strongly adhere to the latex

particle surface and that the latex spheres act as an exclusion

volume for carbon nanotubes. The available space for

MWCNTs during film formation is drastically reduced, which

lowers the percolation threshold and favors network forma-

tion at the ‘‘grain boundaries’’. Grunlan et al. also observed

a very strong increase of the storage modulus at low concen-

tration of single walled carbon nanotubes [48]. However,

development of some porosity, concomitant with SWCNT

network formation, induced a decrease in the modulus above

2 wt% in their case. We did not encounter such a problem

with our materials. AFM imaging before and after heat treat-

ment (Fig. 10a and b) also confirms, because no aggregates

were observed, that CNTs are properly embedded in the films

and do not impede latex spheres coagulation.

It is worth comparing our results with those obtained by

Bokobza and Kolodziej on similar materials [35]. The main

difference between their materials and ours concerns the film

formation, which involved in their case hot pressing at 170 �C,

which is expected to increase the cross-linking density, de-

crease the porosity and improve the adhesion between nano-

tube and latex particles. The level of reinforcement they

reached at low-strain (in both static and dynamic tests) is less

than what we achieved using the water evaporation tech-

nique, although the modulus at higher strain (for instance

200%) is larger in their case. They also found a significant de-

crease of the tensile strength and elongation at break com-

pared to the pristine NR film. That is hardly attributable to

dispersion effects of the MWCNTs since the DC resistivity

dependence is very similar to our results, suggesting that seg-

regation effects also occurred in their case. It is remarkable

that we obtained the same high value of resistivity for

8.3 wt% CNT, which should be lower for a percolating network

of carbon nanotubes, suggesting the presence of adsorbed

molecules at the inter-tube contacts.

We thus propose that the level of reinforcement in our

case is due to the formation of a MWCNT rigid network, prob-

ably mediated by rubber or protein molecules adsorbed at the

surface of oxidized MWCNTs. We have indeed observed the

presence of those molecules at the MWCNT surface by stain-

ing the TEM grid of our materials with uranyl acetate, a clas-

sical procedure for enhancing the contrast of biological

materials (Fig. 8b). We also suggest that filler–filler interaction

dominates over filler–rubber matrix interaction in the low-

strain regime (<100%). The network is very efficient to in-

crease low-strain modulus, while at high-strain the stress–

strain behavior of NR is recovered with some degree of strain

amplification effect due to exclusion volume by MWCNTs.

This rigid network represents a second level of reinforcement

adding to the hydrodynamic contribution. The carboxyl

groups we introduced by the chemical treatment may interact

strongly with the absorbed molecules, which help to increase

the rigidity of inter-nanotube links. SDS may also play a role

in the molecular interaction between CNT and polyisoprene

chains by reducing CNT-NR matrix adhesion. The correlation

observed between the stress-softening and the irreversible in-

crease of the electrical resistivity after stretching is another

strong argument in favor of the CNT networking effect.

Such a high level of reinforcement of NR (for similar frac-

tion of fillers, i.e., 8.3 wt%) is found with chitin, cellulose

whiskers, and starch nanocrystals that form a rigid network

by hydrogen bonding [31]. Nair and Dufresne have shown that

interaction of chitin whiskers with NR can be modified using

Fig. 9 – (a) TEM of a latex film formed on the grid. (b) TEM of the same latex film after staining with uranyl acetate, which

increases the image contrast and shows that adsorbed molecules are wrapped around CNTs. Staining also allowed

observation that MWCNTs were segregated at the latex spheres periphery.

Fig. 10 – AFM image of the surface of a composite with 2.8 wt% MWCNT. Before (a) and after (b) coagulation of the latex beads

at 60 �C. The derivative of topography signal is shown to highlight the morphology of the film after water evaporation. It was

not possible to resolve individual nanotubes at the film surface. No clusters of MWCNT were observed, either, showing that

macroscopic dispersion was good, in agreement with optical microscopy observations.

C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5 1043

coupling agents, and that networking effects decrease when

increasing filler/matrix interactions [50], in agreement with

the analysis of Bokobza and Chauvin [42]. However, cellulose

or chitin whiskers interaction with NR is strong enough so

that elongation at break is always significantly reduced for

similar volume fraction of fillers.

4. Conclusion

We have been able to demonstrate an efficient and successful

way to prepare reinforced and conducting natural rubber

composite materials through incorporation of up to 8.3 wt%

MWCNT activated by oxidation. The composites showed a

1044 C A R B O N 4 6 ( 2 0 0 8 ) 1 0 3 7 – 1 0 4 5

very high degree of increase both in the tensile and storage

modulus in the rubbery region at room temperature, probably

due to rigid networking effect favored by segregation effects

coming from latex structure and cross-linking via functional

groups at the surface of CNT and organic molecules present

in the natural latex solution. It is thus an interesting fact that

segregation effects are detrimental in the glassy state but

profitable in the rubbery region.

Acknowledgments

Fabienne Warmont (CRMD) is warmly thanked for TEM imag-

ing. JPS is grateful to Georges Desarmot (Onera) for his critical

reading of the manuscript and numerous fruitful discussions

on reinforcement of polymers by nano-fillers.

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