Fabrication and Properties of Ethylene Vinyl Acetate-Carbon Nanofiber Nanocomposites
© to the authors 2008
Received: 29 June 2008
Accepted: 3 October 2008
Published: 25 October 2008
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© to the authors 2008
Received: 29 June 2008
Accepted: 3 October 2008
Published: 25 October 2008
Carbon nanofiber (CNF) is one of the stiffest materials produced commercially, having excellent mechanical, electrical, and thermal properties. The reinforcement of rubbery matrices by CNFs was studied in the case of ethylene vinyl acetate (EVA). The tensile strength was greatly (61%) increased, even for very low fiber content (i.e., 1.0 wt.%). The surface modification of the fiber by high energy electron beam and gamma irradiation led to better dispersion in the rubber matrix. This in turn gave rise to further improvements in mechanical and dynamic mechanical properties of EVA. The thermal conductivity also exhibited improvements from that of the neat elastomer, although thermal stability of the nanocomposites was not significantly altered by the functionalization of CNFs. Various results were well supported by the morphological analysis of the nanocomposites.
Carbon nanofibers (CNFs) that are much smaller than conventional carbon fibers but significantly larger than carbon nanotubes (CNTs) can be used to produce nanocomposites with excellent properties, which may open up many new applications. They are available in diameters ranging from 70 and 200 nm and length estimated to be 50–100 μm [1, 2]. CNFs generally exhibit a bamboo-like conical structure as observed in transmission electron micrographs (TEM) (J.J. George and A.K. Bhowmick, personal communication).
Use of CNFs as reinforcement to improve properties of various polymer matrix systems like polycarbonate, epoxy, polyethylene, polypropylene, polymethyl methacrylate, polyether ether ketone, and polystyrene has already been demonstrated [3–20]. The results show enhancement in mechanical [6–8, 10, 19], thermal [6–8], and dielectric/electrical [11–15, 17] properties. The key technical challenges which remain for such nanofiber-reinforced polymers are the achievement of a homogeneous dispersion and good interfacial bonding. The smaller diameter and greater surface area of the nanofibers also imply stronger interactions among the nanofibers; hence, it is often difficult to disperse them into a polymer matrix. Thus, if the dispersion of the nanofibers is less than ideal, it impairs the resultant nanocomposite properties. Making composites with optimal properties requires adequate fiber-matrix adhesion, which is governed by the chemical and physical interactions occurring at the interface. If the fiber to matrix adhesion is poor, a composite may fail at the interface, reducing in particular the tensile strength.
In this article, elastomer grade ethylene vinyl acetate (EVA) having 50% vinyl acetate (VA) content has been chosen as the base matrix. The properties of the composites formed with as-received CNFs and various treated nanofibers are compared, with the idea that the modified CNFs contain surface defect sites and surface polar groups, which can form intermolecular interactions with the polar molecules in the matrix polymer. The presence of such defect sites and surface groups on the pristine CNFs is limited. The quality of the nanofiber dispersion in the polymer matrix is observed by TEM and is then correlated with the mechanical, dynamic mechanical, and thermal properties to provide insight into the role of the CNF surface modification and interfacial interactions on the ultimate properties of the resultant nanocomposites. In our earlier communications, we have shown that appropriate modifications of graphite and multiwalled CNTs can enhance the physico mechanical properties of the nanocomposites [21, 22].
CNFs (as-grown grade PR-24 AG Pyrograf—IIITM) were obtained from Applied Sciences Inc., USA. The CNF consists of a mixture of two distinctive structures present in the sample, relatively straight cylindrical tubes and the so-called bamboo tube-like structures, arranged into loose aggregates. The diameter of CNFs varied between 70 and 200 nm and length between 50 and 100 μm. CNFs used had an aspect ratio (length-to-diameter) of over 500 in the as-received state and is free of carbonaceous contamination. High-resolution TEM micrograph of the surface of CNF shows stacking of graphene layers, distance between graphitic planes being 0.334 nm.
EVA elastomer with 50% VA content was supplied by Bayer (now Lanxess), Germany. The cross-linker for the rubber phase, dicumyl peroxide (DCP, 99% pure), was obtained from Hercules India. Triallyl cyanurate (TAC), the co-agent, was procured from Fluka A G, Germany. Tetrahydrofuran (THF) of LR grade, used as the solvent for EVA, was obtained from MERCK (India) Ltd., Mumbai, India.
The CNFs were irradiated by electron beam (EB) accelerator (model ILU-6) at BARC, Mumbai, India. Irradiation doses used were 50, 200, and 800 kGy (dose per pass was fixed at 10 kGy) at room temperature. A FWT-60 dosimeter based on calibration obtained from gamma-radiation was used for the EB dosimetry. The accelerator voltage frequency was 100–120 MHz and the energy range was 0.5–2.0 MeV.
The CNF samples were irradiated with gamma rays at three different doses −1, 5, and 10 kGy, using GC 5000 (Source: Co-60) at a dose rate of 3.2 kGy/h. This was carried out at the gamma irradiation facility in BARC, Mumbai, India.
The amination of CNFs was done by treating 200 mg of the sample with excess of hexamethylene diamine within a 50-mL thick-walled test tube at 130 ± 10 °C in an oil bath for 24 h. The treated sample was then washed with alcohol to remove the excess amine followed by washing with distilled water to remove the alcohol present. The nanofibers were then filtered using nylon membrane filter paper of 0.45-μm pore size. It was then dried in vacuum oven at 80 °C for 4 h.
CNFs (200 mg) were sonicated with H2SO4/HNO3mixture (3:1) for 3 h at 40 °C. The treated samples were washed with distilled water repeatedly until the pH of the mixture came to 6. The nanofibers were then filtered using nylon membrane filter paper of 0.45-μm pore size. These were then dried in vacuum oven at 80 °C for 4 h.
CNFs were functionalized via refluxing with 1 g of vinyl-silane in 25-mL of THF at 80 °C for 8 h. The free radical reaction was initiated by benzoyl peroxide (0.1 g) added to the mixture. Modified nanofibers were washed several times with anhydrous THF.
All the treated and untreated CNFs were analyzed using different morphological, elemental, structural, and thermal characterization techniques. The detailed results of various characterizations were provided elsewhere (J.J. George and A.K. Bhowmick, personal communication).
Untreated carbon nanofiber
CNF-treated with 800 kGy EB
CNF γ-1 kGy
CNF-treated with 1 kGy Gamma irradiation
Amine-treated carbon nanofiber
Acid-treated carbon nanofiber
Silane-treated carbon nanofiber
EVA filled with 1 wt.% of untreated CNF
EVA filled with 4 wt.% of untreated CNF
EVA filled with 8 wt.% of untreated CNF
EVA filled with 1 wt.% of 800 kGy EB irradiated CNF
EVA filled with 1 wt.% of 1 kGy Gamma irradiated CNF
EVA filled with 1 wt.% of amine treated CNF
EVA filled with 1 wt.% of acid treated CNF
EVA filled with 1 wt.% of silane treated CNF
The nanocomposites were synthesized by using a solution-mixing technique. EVA (5 g per batch) was dissolved in 50-mL of THF to make 10% solution of the rubber using a mechanical stirrer. 0.05 g of DCP as the curing agent and 0.05 g of TAC as the co-agent were added to the rubber solution. The solution was thoroughly stirred using a mechanical stirrer. CNFs dispersed in THF were first sonicated for 15 min and subsequently added to the rubber solution while stirring at room temperature (27 °C). The final solution was cast over Teflon trays and kept for air drying followed by vacuum drying at 50 °C till there was practically no weight variation. The dried films were molded in a hot press at a pressure of 5 MPa at 150 °C for an optimum cure time of 25 min, determined from a Monsanto oscillating disc rheometer (100S).
The microscopy was performed using a JEOL JEM-2010 (Japan), TEM operating at an accelerating voltage of 200 kV. The composite samples were cut by ultra-cryomicrotomy using a Leica Ultracut UCT. Freshly sharpened glass knives with cutting edge of 45° were used to get the cryosections of 50–70 nm thickness. Since these samples were elastomeric in nature, the temperature during ultra-cryomicrotomy was kept at −50 °C (which was well below the glass transition temperature of EVA). The cryosections were collected individually on sucrose solution and directly supported on a copper grid of 300-mesh size.
Morphological comparison of untreated and acid-treated CNFs was performed using an SEM model JSM800 manufactured by JEOL at 20 kV acceleration voltage at room temperature.
The mechanical properties of the nanocomposites were evaluated by a universal testing machine (Zwick 1445) on dumbbell specimens, punched out from the cast films using an ASTM Die C. All the tests were carried out as per ASTM D 412-99 method at 25 ± 2 °C at a crosshead speed of 500 mm/min. The average values of three tests for tensile strength, tensile modulus, and elongation at break are reported for each sample.
Dynamic mechanical thermal characteristics of the composite films (0.4–0.6-mm thick) were evaluated by using a DMTA IV (Rheometric Scientific) under tension mode. All the data were analyzed using RSI Orchestrator application software on an ACER computer attached to the machine. The temperature sweep measurements were made from −35 to 20 °C. The experiments were carried out at a frequency of 1 Hz at a heating rate of 2 °C/min. The strain amplitude used in the DMA test was 0.01% and the soak time at −35 °C was 3 min. The storage modulus (E′) and the loss tangent (tan δ) data were recorded for all the samples under identical conditions.
where, V r is the volume fraction of rubber in the swollen gel, D the de-swollen weight of the composites, F the fraction insoluble, T the initial weight of the sample, and A 0 the amount of solvent imbibed. ρ r is the density of the rubber, while ρ s is density of the swelling solvent.
whereW is the power in Watts (here 4 W),K the thermal conductivity,t the thickness of sample,A the area of the sample, and dT the temperature difference between the two plates.
Thermal stability of the composites was investigated by thermo gravimetric analysis (TGA) by using a Perkin–Elmer TGA instrument (Model: Pyris Diamond TG/DTA) from ambient to 800 °C at a programed heating rate of 20 °C/min in nitrogen. A sample weight of approximately 10 mg was taken for all the measurements. The weight loss against temperature was recorded. Differential thermo gravimetric analysis (DTG) of the composites was represented in terms of the first derivative plots of the TGA curves. The data points denote the weight loss/time against temperature at the specified heating rate.
Tensile properties of various nanocomposites
Tensile strength (MPa)
Elongation at break (%)
Modulus at 100% elongation (MPa)
5.05 ± 0.15
490 ± 30
0.74 ± 0.17
8.14 ± 0.20
410 ± 20
1.04 ± 0.21
8.53 ± 0.11
465 ± 20
1.25 ± 0.12
8.60 ± 0.15
440 ± 15
1.38 ± 0.15
8.25 ± 0.12
432 ± 25
1.30 ± 0.10
6.86 ± 0.18
323 ± 30
1.02 ± 0.16
8.50 ± 0.10
436 ± 15
1.36 ± 0.14
Dynamic mechanical analysis data of various nanocomposites
Storage modulus (MPa)
The glass transition temperature (T g) for EVA occurs at −30.8 °C and there is a marginal shift of 2.4 °C inT gtoward higher temperature (Table 3) for EVA-1F. A significant shift of 6.6 °C is observed for the nanocomposite with 4 wt.% untreated CNF. For the composite with 8 wt.% CNF loading, theT gis shifted toward lower temperature region (~2 °C from that of 4 wt.% CNF-filled sample), showing that at higher loadings filler–filler interactions start dominating the filler–polymer interactions. The tan δ peak heights of all the nanocomposites are lower than that of the neat EVA film, which confirms the increase in filler–polymer interaction. When molecular mobility is restricted due to the presence of reinforcing fibers, it results in not only enhanced glass transition temperature, but also in decreased tan δ peak magnitude. The loss modulus versus temperature plots of the nanocomposites with varying filler loadings provide a similar trend (data not shown here). The glass transition temperature appears at still lower temperature. There is only singleT gof the nanocomposites and there is no separate transition for confined EVA chains as effected by CNF.
Thus, the optimum enhancement in the glass transition for EVA-1AF indicates that there exists a significant restriction on the segmental mobility of these polymer chains, which in turn suggests that the polymer chains are in close proximity and interact significantly with the nanofibers and this is further confirmed from the solvent swelling studies.
Volume fraction of rubber in the swollen polymer mass (V r) of various samples
Volume fraction of rubber (V r)
The thermal conductivity shows an increment of 70 and 188% with incorporation of 1 and 4 wt.% of CNFs, respectively. All the modified CNFs except acid treated one give rise to improved thermal conductivity for the nanocomposites, as compared to untreated CNF-filled sample. The samples EVA-1FEBand EVA-1F γ exhibit increments of 2 and 7%, respectively, over that of EVA-1F, whereas EVA-1AF and EVA-1SF show improvements of 10 and 29%, respectively. These enhancements are attributed to the better dispersion and interaction of the nanofibers within the rubber matrix after modification.
Thermal degradation data of various nanocomposites
Temperature at which maximum degradation occurs (°C)
Maximum rate of degradation (°C/min)
Introduction of a small amount of CNFs can lead to improved performance of EVA. At 4 wt.% CNF loading, the modulus and the tensile strength of the nanocomposite increased substantially. However, similar improvements were not observed at higher (8 wt.%) nanofiller loading due to filler agglomeration. Surface treatment of CNFs with high-energy irradiations of EB and gamma irradiation resulted in improved fiber to matrix interaction, which was supported by swelling resistance studies. This in turn enhanced the dispersion and wetting properties of the nanofibers leading to further improvements in mechanical, dynamic mechanical, and thermal properties of the nanocomposites. At similar loading of filler (1 wt.%), EVA-1SF and EVA-1F γ exhibited the best overall properties. The morphological studies revealed that the modified CNFs were better dispersed and distributed in the elastomer matrix at low loadings.
The authors acknowledge the financial assistance provided by DRDO, New Delhi, India.