Improving Reinforcement of Natural Rubber by Networking of Activated Carbon Nanotubes

CARBON

4 6 ( 20 0 8 ) 1 0 3 7–10 4 5

available at www.sciencedirect.com

journal homepage: www.elsevier.com/locate/carbon

Improving reinforcement of natural rubber by networking of activated carbon nanotubes Sanjib Bhattacharyyaa, Christophe Sinturela, Ouziyine Bahloula, Marie-Louise Saboungia, Sabu Thomasb, Jean-Paul Salvetata,* a

Centre de Recherche sur la Matie`re Divise´e, CNRS Universite´ d’Orle´ans, UMR6619, 1B rue de la Fe´rollerie, 45071 Orle´ans cedex 2, France School of Chemical Sciences, Mahatma Gandhi University, Priyadarshini Hills P.O. Kottayam, Kerala 686 560, India

b

A R T I C L E I N F O

A B S T R A C T

Article history:

Reinforcement of natural rubber was achieved using carboxylated multiwalled carbon

Received 31 January 2008

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

Accepted 14 March 2008

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

Available online 19 March 2008

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 measurement at room temperature revealed a low percolation threshold ( 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 level of reinforcement filler–matrix and filler–filler interactions

* Corresponding author: Fax: +33 2 38255376. E-mail address: [email protected] (J.-P. Salvetat). 0008-6223/$ - see front matter  2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.carbon.2008.03.011

1038

CARBON

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 filler aggregates. Formation of rigid filler network usually provides 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 effect [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-fillers. Also in present days carbon black is the principal reinforcement 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 reinforcing 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 natural fibers is particularly interesting for increasing the modulus 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 significant level of reinforcement has been achieved [33]. By contrast, 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 rubber (NR), Bokobza and Kolodziej have made detailed investigation 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 reinforcement occurs mainly at low-strain values, whereas at highstrain 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 lowstrain 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 synthesized in our laboratory by catalytic decomposition of acetylene at 600 C on a CoxMg(1x)O solid solution [37].

2.1.

Purification and carboxylation of MWCNTs

Prepared MWCNTs by CVD technique contain amorphous carbon and metallic nanoparticles (used as catalyst during synthesis) 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 sonicator, and then poured into a Petri dish and dried at 60 C for 24 h to obtain freestanding composite films. For the composite 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 electron microscopy (TEM) on carbon coated copper grid in Philips CM-20 TEM. Fig. 1 shows the TEM picture of purified and carboxyl 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 investigated 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

CARBON

4 6 ( 20 0 8 ) 1 0 3 7–10 4 5

1039

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 socalled reduced (or Mooney) stress ½r  ¼ r=ða  a12 Þ as a function of the extension ratio a, or a1 (Fig. 3). In the James and Guth approach, [r*] = E, the elastic modulus, which is proportional to the cross-link density and to the temperature [1]. In the Mooney–Rivlin model for large deformation, [r*] varies linearly with a1, according the expression ½r ¼ 2C 1 þ 2C 2 a1  [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 representation, which is often used by rubber scientists to discuss

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.

of a Pico+ AFM (Molecular Imaging), with standard Si cantilevers 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 measured 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 composite 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 (tan d) were measured in the temperature 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 impedance meter at room temperature in the frequency range 100 Hz–1 MHz. Non-metallized samples were sandwiched between 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. 3.1.

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

Results and discussion 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

Fig. 3 – Reduced stress, [r* ] ¼ r/(a  a12 ), vs. inverse extension ratio, a1 (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.

1040

CARBON

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 a1 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 < a1 < 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 effect disappeared after vulcanization, and the authors explained 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 elastomers. The difficulty lies in introducing the effect of fillers in the constitutive equation of rubber, i.e., expressing the reduced 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 fillers. The stress–strain is then described by r ¼ EðK  1=K2 Þ. 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 E0 ð1 þ 0:67f / þ 1:62f 2 /2 Þ 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 volume 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 structural characteristics deduced from TEM observation. However, 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–filler interactions to interpret the tensile behavior of our samples 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

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.

valid at low filler concentrations before network formation, and at high-strain when network destruction occurs. Loading–unloading cycles showed a strong stress-softening (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 fraction (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. 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.

CARBON

4 6 ( 20 0 8 ) 1 0 3 7–10 4 5

1041

centration (plot not shown), while elongation at break decreased (Fig. 2). There is almost 20% reduction in strain to failure 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 phenomenon suggests the homogeneous nature of our composite films. During the first loading it is highly probable that some degree 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 mechanisms 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 temperature and strain amplitude. The goal is to evaluate the degree of filler–rubber or filler–filler interactions, which provide additional levels of static and dynamic reinforcement. Fig. 6 shows the variation of storage modulus (log E) (Fig. 6a) and the loss factor tan d (Fig. 6b) at 10 Hz as a function of temperature (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 corresponding relaxation process where the loss factor tan d passes through a maximum around 205 K. It may be of interest to note that the temperature corresponding to the maximum of tan d decreases when MWCNT concentration increases. A similar 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 nanofillers. 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 value rather than a linear increase such as predicted by the rule-of-mixture. This suggests that nanotube aggregation occurs 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 nanotubes 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 modulus 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-

Fig. 6 – (a) Storage modulus and (b) tan d 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 tan d decreases linearly vs. MWCNTs concentration.

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 dynamicmechanical 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 r0 j/  /c jt 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

1042

CARBON

4 6 ( 2 0 0 8 ) 1 0 3 7 –1 0 4 5

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.

500 % extension. The high-strain induces an irreversible increase of the resistivity of one order of magnitude for the 8.3 wt% filled samples, demonstrating that the nanotube network 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 attribute 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 formation at the ‘‘grain boundaries’’. Grunlan et al. also observed a very strong increase of the storage modulus at low concentration 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 treatment (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, decrease the porosity and improve the adhesion between nanotube 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 technique, although the modulus at higher strain (for instance 200%) is larger in their case. They also found a significant decrease of the tensile strength and elongation at break compared 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 segregation 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, probably 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 staining the TEM grid of our materials with uranyl acetate, a classical 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 lowstrain regime (