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Mechanical properties of carbon nanoparticle-reinforced elastomers
Composites Science and Technology 63 (2003) 1647–1654 www.elsevier.com/locate/compscitech
Mechanical properties of carbon nanoparticle-reinforced elastomers Mark D. Frogley, Diana Ravich, H. Daniel Wagner* Department of Materials and Interfaces, Weizmann Institute of Science, Rehovot 76100, Israel Received 3 September 2002; received in revised form 17 December 2002; accepted 28 December 2002
Abstract Silicone based elastomers have been mixed with single-wall carbon nanotubes or larger carbon nanofibrils. Tensile tests show a dramatic enhancement of the initial modulus of the resulting specimens as a function of filler load, accompanied by a reduction of the ultimate properties. We show that the unique properties of the carbon nanoparticles are important and effective in the reinforcement. The modulus enhancement of the composites initially increases as a function of applied strain, and then at around 10–20% strain the enhancement effect is lost in all of the samples. This ‘‘pseudo-yield’’ in elastomeric (or rubber) composites is generally believed to be due to trapping and release of rubber within filler clusters. However, insitu Raman spectroscopy experiments show a loss of stress transfer to the nanotubes suggesting that instead, the ‘‘pseudo yield’’ is due to break-down of the effective interface between the phases. The reorientation of nanotubes under strain in the samples may be responsible for the initial increase in modulus enhancement under strain and this is quantified in the Raman experiments. # 2003 Elsevier Ltd. All rights reserved. Keywords: B Interface; B. Mechanical properties, Raman spectroscopy; Nanomaterials; Elastomeric composites
1. Introduction The high and reversible deformability of elastomers is of great industrial importance. Typically however, the initial modulus and durability of such materials is low, and an additional reinforcing phase is required for practical use. Carbon black and silica particles have been used extensively for this purpose [1–4]. For the composite to be effective, there must be a strong interaction between the matrix and the stiffer phase and this can be achieved using a filler with a large surface-area-to-volume ratio. Optimally this means using small fillers with a large aspect ratio, and recently, particles with nanoscale dimensions such as flakes [5–7], nanofibers [8–9] or hollow nanotubes [10– 15] have become the subject of extensive research. Unprecedented improvement of the mechanical properties has been observed in these nanocomposites. The key issues for nanoparticles, which ultimately determine their usefulness as fillers, are their effective dispersion in * Corresponding author. Tel.: +972-8-934-2594; fax: +972-8-9344137. E-mail address: [email protected] (H.D. Wagner). 0266-3538/03/$ - see front matter # 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0266-3538(03)00066-6
the matrix (with large surface/volume ratios the particle-particle interactions are strong) and the nature of the interface with the matrix. But for nanoparticles, how do we test these properties? Bulk tests, such as frictional wear or tensile testing, can tell us if a material performs well, but to understand and improve the performance, results must be related to the interaction between individual nanofillers and the matrix. This is directly possible for some properties, and the filler dispersion as well as wetting, buckling and pull-out [5,8,13,16–18] have been observed by high-resolution electron microscopy. However, this technique yields mostly qualitative information since the local stresses and strains on the nanotubes remain unknown. Conventional tests of the fiber–matrix interaction such as fiber pull-out (whilst recording the stress in the fiber), or fragmentation tests can be performed in principle with nanoscale fibers, but as well as the technical difficulty involved, interpretation of the results is difficult. It is not yet clear whether standard micromechanical theories can be extrapolated to the nanoscale, where individual fibers are comparable in size to the polymer chains. In this paper we study the mechanical properties of a silicone rubber reinforced
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with single-wall carbon nanotubes or with larger, carbon nanofibrils. We perform tensile tests on bulk specimens, which show dramatic improvement in the mechanical properties as a function of filler content. For the nanotube samples we use Raman spectroscopy of the nanotubes to relate the macroscopic results to the nanoscale behaviour.
2. Experimental The elastomer was RTV silicone rubber (Epusil 186, Polymer G’vulot ltd., Israel). The fillers were either single-wall carbon nanotubes (SWNTs) produced by the HiPCo technique, (Carbon Nanotechnologies Inc., USA) or vapor-grown carbon nanofibrils (Polygraf III, Applied Sciences Inc., USA). Nanofibrils are highly graphitic single fibers with lengths up to a few hundred microns and diameters around 200 nm. Composites with good fiber dispersion, as determined by electron microscopy, were prepared by first dissolving the RTV in toluene (1 mg/ml) to reduce the viscosity, and separately dispersing the fillers in toluene by ultrasonication (0.1 mg/ ml, 10 min, 6 W). Then the two dispersions were mixed and further ultrasonicated for 10 min. The toluene was removed by evaporation at 50 C over several days with continuous stirring and finally the hardener was mechanically mixed in, and air bubbles were removed by vacuum pumping. Films of thickness 200–300 mm were produced by shearing the uncured mixture across a glass plate with a doktor blade. After curing for 24 h at room temperature, simple beams of width 1.5 mm and length 20 mm were cut from the films. Samples with filler concentrations up to 1 wt.% (SWNT) and 4 wt.% (fibrils) were prepared but samples with higher concentrations were too viscous for bubbles to be effectively removed and for spreading uniform films. For tensile tests an Instron 4502 machine was used and samples of gauge length 10 mm were stretched at 100 mm/min until failure. Ten samples of each composition were tested for statistical accuracy. Throughout this work, the quoted strain is the true strain, i.e. the percentage elongation, and the stress is the nominal or engineering stress (the force divided by the cross sectional area of the unloaded sample). For Raman measurements on the nanotube composites, the samples were stretched using a home made mini-tensile tester [19], which was placed under a Renishaw Raman microscope. Spectra were obtained in the backscattering geometry (shown inset in Fig. 5) using 632.8 nm laser light, which was polarized along the strain axis of the tensile tester. The laser beam was focused to a spot 20 mm in diameter on the sample to avoid heating effects, and the same point on the sample was used for all measurements.
3. Results and discussion 3.1. Mechanical tests Fig. 1 shows representative stress–strain curves for pure RTV and for RTV/nanoparticle composites. For both SWNTs and fibrils, the initial modulus (measured by fitting a straight line to the data below 10% strain) is radically improved as the filler concentration increases as shown in Fig. 2. For RTV/SWNT composites the modulus increase is approximately linear with weight fraction, with a slope of 200%/wt.% or a maximum of 200% increase in stiffness for 1% SWNT. This compares with 120%/wt.% for the nanofibrils measured here and about 100%/wt.% for a nanofibril/rubber composite made by Canales and Cavaille [20]. For carbon black filled rubbers, the rate is as little as 4%/wt.% [1,2] and for silica particles in styrene-butadiene rubber [3], where there is chemical bonding between phases, the modulus increases about 2–3 times for a loading of 24 wt.% (a rate of increase a few times higher than for carbon black). We conclude from these results that the
Fig. 1. Stress–strain curves for RTV/carbon nanofiber composites. The stress is the engineering stress (applied force divided by the unloaded cross sectional area) and the strain is the true strain.
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modulus of the two phases as well as filler packing effects, but for low loadings of rod-like fillers which are much stiffer than the matrix (as in our samples) it reduces to [22]: 1 þ 2fc Epara ¼ E0 1c 1 þ 0:5c Eperp ¼ E0 1c Erand ¼ 0:2Epara þ 0:8Eperp
Fig. 2. (a) Initial modulus of the composites as a function of filler weight fraction. Error bars represent the scatter of ten experiments for each composition. The solid curves are the theory of Guth [Eq. (1)] with filler aspect ratios, f, as shown. The dashed curve is typical of conventional fillers such as carbon black and silica particles. (b) The modified Halpin–Tsai theory [Eq. (2)] is compared with the data. In the key, O refers to fillers perfectly oriented along the strain axis and R refers to fillers randomly oriented in three dimensions.
specific properties of the carbon nanotubes and fibrils are important in the reinforcement. The well known model of Guth [1] for the modulus of rod-like particle reinforced elastomers is based solely on the aspect ratio and volume fraction of the filler, and does not take any other properties of the filler into account, except to assume that the fillers are stiff compared with the matrix: E ¼ E0 1 þ 0:67fc þ 1:62f 2 c2 ð1Þ where E and E0 are the moduli of the composite and the pure matrix respectively, and f and c are the aspect ratio and volume concentration of fillers. Another commonly used model for filled polymers is the modified HalpinTsai theory [21], which in principle takes account of the
ð2Þ
for the Young’s modulus measured parallel (Epara) or perpendicular (Eperp) to perfectly oriented fibers or in any direction (Erand) for fibers randomly oriented in three dimensions. In Fig. 2a, Guth’s theory [Eq. (1)] is compared with the experimental data, and shows good agreement using f 120 for the nanotubes and f 70 for the fibrils. These values are consistent with observations of the brittle fibrils after they were broken up by the ultrasonication procedure and with SWNT bundles of diameter 10–15 nm, and lengths of a few microns. SWNT bundles have a stiffness comparable to single tubes [23,24] and are often observed in the RTV and other matrices by electron microscopy [16,18]. Individual SWNTs are rarely seen. The curve for a fiber of aspect ratio 1000, which is typical for an individual SWNT, does not agree with the data, but is a clear incentive to develop methods to separate the nanotubes. In Fig. 2b, the modified Halpin–Tsai equation is compared with the same data. Assuming a random distribution of fibers, aspect ratios of 700 and 500 are needed for the nanotubes and fibrils respectively to achieve a good fit: much higher than observed. Raman measurements described later suggest that there is a degree of fiber alignment during sample preparation (shearing) but it is far from perfect (Fig. 8, or see e.g. Ref. [25]). Even perfectly oriented fibers require aspect ratios of 140 and 100 for the theory to fit the data. Lower aspect ratios lead to a prediction of a lower modulus and so for both theories, the measured level of reinforcement is either consistent with or better than predicted by theory, which is reasonable because the theories ignore the increase in surface-area-tovolume ratio for smaller fillers, even when the aspect ratio is kept constant. Note that, to use Eqs. (1) and (2), the measured filler weight fractions were converted to volume fractions, c, using densities of 1800, 1340 and 1215 kg m3 for fibrils, SWNTs [26] and the matrix respectively. Others have used nanofillers with high aspect ratios. Kim and Reneker [9] used electrospun fibers of polybenzimidazole (PBI) in a styrene-butadiene rubber. The fibers had an average diameter of 300 nm and f1000 (after processing) and for a 10 wt.% loading, the
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modulus was enhanced by a factor of 10. This result is close (by weight) to that of the carbon fibrils. Osman et al. [5] used plate-like mica particles with various aspect ratios as fillers in PDMS, and showed that the aspect ratio plays a crucial role in the reinforcement. The best reinforcement was seen for the highest aspect ratio (f=120) where the modulus was improved by a factor of 13 for a 14 vol.% filler load [5]. However, the Guth model, which is for rod-like fillers [Eq. (1)] suggests a much stronger increase is possible for SWNTs. If the volume fraction of SWNT bundles could be increased to 14% the expected modulus increase would be almost 400 times (extrapolating the fit of Fig. 2). Several works have appeared regarding PDMS/OMT (organo-montmorillonite) nanocomposites [6-7], which are plate-like in nature, but the enhancement of stiffness is somewhat less impressive than those discussed so far. SWNT composites represent the highest initial level of reinforcement, by weight of filler, of any elastomeric material to date. This reinforcement comes at a price. The ultimate strain of the SWNT and fibril composites is significantly less than the pure RTV (Fig. 1) and the tensile test results are summarized in Fig. 3. All the data fall on one curve, with the reduction in ultimate strain being quite dramatic at first, and then leveling off to around 200% strain (60% of the value for pure RTV) for large concentrations. Very similar behavior was seen for the high aspect ratio (f=120) PDMS/mica nanocomposites [5], for which the value stayed on the low plateau up to the highest concentrations. This behavior is also seen in other fiber-filled rubbers [9,27] whereas for low loads of near-spherical fillers, a less marked reduction, or even an increase [5,9] in the ultimate strain has been
observed. In any case, for most applications requiring a large initial reinforcement, the ultimate strain is not critical. Fig. 4 shows the modulus measured at 80% strain, as a function of filler content. The modulus is constant within the experimental scatter, that is, there is no obvious effect of the filler at high strain, in contrast to the huge reinforcement at low strain. Several studies [3] have revealed such behavior, termed a ‘pseudo-yield’ in filled rubbers and it has been attributed to breakdown of secondary particle structure, (aggregates) and the resulting release of unstrained, trapped rubber. The degree of secondary structure should not be significant for low volume fractions however, especially if the fillers are well separated like the fibrils. For our SWNT composites the tight binding of the SWNT bundles makes it unlikely that rubber will penetrate the bundle, and so any trapped rubber would have to be in some tertiary structure formed from many bundles for SWNTs. We have seen no evidence of microscale stucture in our samples. We propose that, at least in our composites, the pseudo-yield is due to weakening of the frictional interface between the two phases—that is, filler dewetting—and we return to this point later. The mere existence of the pseudo-yield demonstrates that the fillers, and their interface with the matrix, are responsible for the reinforcement of the polymer, and not curing effects such as increased crosslinking [22] induced by the fillers. If this were the case, the stiffness of the composites at high strain (Fig. 4) would be larger for higher crosslinking densities (higher filler loads). The marked increase in viscosity of the nanotube-monomer mixture (before curing), as a function of filler load, supports this argument.
Fig. 3. Strain to failure as a function of filler weight fraction for the composites. Error bars represent the scatter of ten experiments for each composition.
Fig. 4. The modulus of the composites, measured at 80% tensile strain, as a function of filler content.
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Having shown that carbon nanofibers have a dramatic effect on the bulk mechanical properties of RTV rubber we now relate this, for the carbon nanotube composites, to the interaction between the individual nanotube bundles and the matrix. In stiff polymers, like poly (methyl methylacrylate) (PMMA) and polyurethane acrylate (PUA) or epoxy resins, which have elastic moduli about 1000 times greater than rubbers, embedded nanotubes undergo significant mechanical compression during the polymer curing process, and are also strongly affected by external stress applied to the cured composite. This has been observed by looking at changes in the Raman spectrum of the nanotubes, which are related to mechanical deformation of the tubes [10–12,15,25,28–30]. Specifically, the D* Raman mode of SWNTs exhibits a large spectral shift when the composite is strained in uniaxial tension, and the shift is approximately linear with the applied strain until the polymer yields [25,29–31] which happens at about 1% strain for an epoxy matrix as shown inset in Fig. 6 [31]. After the polymer yields, there is no further change in the Raman spectrum, and this indicates that stress is no longer efficiently transferred from the matrix to the nanotubes. Raman spectroscopy is therefore sensitive to the strength of the interface between the individual nanotubes and the matrix, and we can interpret a quench in the Raman strain-shift of the D* mode as an indication of interfacial breakdown. Fig. 5 shows a typical spectrum of the D* peak for SWNTs in RTV. Fig. 6 shows the peak position for nanotubes in the 0.3
wt.% composite, as a function of applied tensile strain. The Raman wavenumber decreases linearly with strain at first, and tends to a constant value by about 50% strain. There are two important features of this data. The first is the shape of the curve, which tells us about the rate of stress transfer to the nanotubes. If the interfacial strength is constant, we expect a linear wavenumber–strain relationship, even for large deformations of the nanotubes (Raman strain shifts of 10–20 cm1) as seen e.g. inset in Fig. 6. Non-linear behavior indicates a change in stress transfer efficiency. Above 40% strain in Fig. 6, the Raman wavenumber has an almost constant value, implying that stress is no longer transferred effectively to the tubes. This quench in strain-shift corresponds to the pseudo-yield phenomenon, which is apparent over the same strain region in the stress–strain curves of Fig. 1. The pseudo-yield is shown more clearly, for the 0.3 wt.% SWNT sample, in a plot of modulus enhancement against strain (Fig. 7). The modulus enhancement factor at a given strain is defined as the modulus of the composite at that strain divided by the equivalent modulus for pure RTV, and was calculated from the stress-strain curves of Fig. 1. This is a relative parameter, so it does not matter if the true or engineering stress is used. The modulus enhancement increases initially with strain, and then falls quickly after about 12% strain. This is the same strain value at which the Raman data departs from a linear function in Fig. 6, suggesting that the ‘‘pseudo-yield’’ is due to a loss of stress transfer between the individual nanotube bundles and the matrix, which may be a result of the straightening of the polymer chains. This is further
Fig. 5. The Raman spectrum of an RTV/0.3 wt.% SWNT composite in the vicinity of the D* band for SWNTs (solid curve). The crosses represent the mixed Gaussian/Lorentzian function used to fit the three peaks near 2420, 2500 and 2600 cm1 (D*). Inset: setup for polarized Raman spectroscopy of the sample film in the backscattering geometry. E is the electric vector of the laser light.
Fig. 6. The Raman wavenumber of the D* band for SWNTs embedded in RTV rubber is plotted against tensile strain applied to the composite. The initial slope is 0.08 cm1/% strain (solid line). Inset: The strain induced shift of the Raman wavenumber of the D* band for SWNTs embedded in an epoxy matrix is plotted against tensile strain applied to the composite. After the polymer yields at 1% strain, there is no significant change in the wavenumber.
3.2. Nanoscale behaviour—Raman spectroscopy
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Fig. 7. The modulus enhancement factor for the RTV/0.3 wt.% SWNT composite as a function of applied strain. The enhancement factor is the modulus of the composite divided by the modulus of pure RTV at each strain. Increased noise below 5% strain is from division of small stress values.
evidence that the yield is not due to a release of trapped rubber from within filler clusters (the suggested mechanism for conventional filled composites) in this material, as discussed earlier. The second interesting feature in the Raman data of Fig. 6 is the magnitude of the spectral shift. The total
Fig. 8. Polarized Raman intensity for SWNTs embedded in RTV as a function of the angle between the sample axis (the tensile strain axis) and the polarizer axis. A constant (normalized) intensity of 1 is expected for totally unoriented tubes, whereas the lower dashed curve is expected for perfectly oriented tubes. The data points are for an RTV/0.3 wt.% SWNT composite at various strains, showing that the nanotubes become oriented as the tensile strain is increased. The solid curve is a theoretical fit to the data at 200% strain using the method of Hwang et al. [14] as explained in the text. All data are normalized to the intensity at =0.
shift is only 2 cm1, over 50% strain, which is much less than 7–18 cm1 shift after 1% strain for SWNTs in stiffer polymers [25,29–31]. In stiffer matrices, the rate of Raman shift with strain depends on the thermal compression of the matrix around the nanotubes, which occurs on cooling after curing at high temperature because of the mismatch in thermal expansion coefficient of the two phases [29]. For tensile tests at lower temperatures, the rate of wavenumber shift with applied tensile strain is higher and this is thought to be due to a more effective interface resulting from enhanced lateral clamping of the nanotubes [32]. The RTV/SWNT composites, however, were cured at room temperature over a long time and so the radial thermal compression or ‘clamping stress’ is low. This results in a poor interface, which must be non-strain-identical in nature: if the nanotubes were being stretched at the same rate as the matrix, the resulting wavenumber strain-shift would be comparable with the epoxy data in Fig. 6. The axial thermal stress may play a more significant role. Traditional carbon fibers embedded in a matrix which undergoes significant thermal contraction in the post-curing cool-down are seen to go into compression via shear stress transfer across the interface [33,34]. If the curing takes place at room temperature over a long time, the residual axial stress in the fiber is much lower, or zero [34]. Similarly, for nanotubes, a large compressive Raman shift (increase in D* wavenumber of order 10 cm1) is observed after cooling the high-temperature-cure polymers [11,28,29]. The compressive stress stores energy in the nanotube, like a coiled spring and so when the polymer is subsequently stretched, the nanotube will return to its original length as quickly as the surrounding polymer allows. Hence the tensile Raman strainshift is large. This has been observed in various SWNT composites and the increase in wavenumber due to thermal compression is always similar to the decrease in wavenumber on subsequent stretching to yield. In silicone rubber however, which is cured at room temperature, there is no thermal stress, and so the nanotube is close to its equilibrium (in the uncured RTV liquid) length. Under tensile strain therefore, extension of the nanotubes relies directly on the weak interface so the wavenumber strain-shift is much smaller, as observed. We rule out direct effects of radial stress on the Raman wavenumber since there is no shear stress transfer mechanism in the radial direction and the small normal stresses in the polymers (of the order of MPa) are orders of magnitude too small to significantly affect the D* wavenumber as seen in high-pressure experiments on SWNT bundles [29]. Nanotube bending effects also must not affect the wavenumber shift since the degree of bending will be much higher in the RTV (over 50% strain) than over 1% strain in epoxy, whereas the wavenumber shift is much lower in the RTV.
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3.3. Orientation of the nanotubes In composites at large strains, originally unoriented fibers, for example, will align strongly in the direction of applied tensile strain, which is a simple consequence of the change in shape of the rubber. The effect is important because oriented fibers provide better reinforcement (along the orientation direction) than unoriented fibers [22]. Nanotube reorientation can be monitored by polarized Raman spectroscopy because the intensity of the polarized Raman spectrum of a carbon nanotube depends strongly on the angle, y, between the nanotube axis and the optical polarization direction [35,36]. For single-wall carbon nanotubes excited with polarized light of wavelength 632.8 nm, the Raman scattering is resonant [36,37] so that for a single nanotube, or for tubes oriented perfectly along one direction, the total intensity of the Raman modes varies as [14,36]: RðÞ / cos4
ð3Þ
For a sample of randomly oriented nanotubes the total Raman intensity as a function of polarization angle is a constant, and any systematic deviation from this indicates a degree of nanotube alignment. Fig. 8 shows the normalized intensity of the D* mode as a function of the angle, , between the sample axis and the optical polarization direction, in the plane of the film. The sample axis is defined as the direction of shear flow (used to produce the RTV films), which is also the direction of applied tensile strain and therefore is the nominal axis of nanotube alignment. With no applied strain, (up triangles) there is already some orientation, and this is due to the shear flow processing method. As tensile strain is applied the degree of orientation increases and by 200% strain the data is close to cos4, the dashed curve in the figure, which is expected for perfect alignment. To quantify the degree of alignment from the Raman data we use the method of Hwang et al. [14]. This assumes that the number of nanotubes N() at an angle to the alignment axis (in the plane of the film) is a Lorentzian function centered on the alignment direction, and that the contribution of each nanotube to the Raman signal varies as cos4(-), where - is equal to and is the angle between that nanotube and the polarizer axis, from Eq. (3). One then obtains a best-fit to the data by varying the width (FWHM) of the Lorentzian, and this width is the measure of the degree of alignment. We obtain a FWHM of 30 for the data at 200% strain (Fig. 8), i.e. half of the nanotubes lie within 15 of the alignment axis. Strictly the procedure is only valid for nanotubes in the plane of the film but our sample approximates this at high strain because during the shearing process used to produce the films, the alignment of nanotubes into the plane of the film is substantially stronger than the alignment of tubes within the plane, towards the shearing direction [25,38]. This
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effect is magnified at high tensile strain. At lower strains, both the two-dimensional approximation and the assumption of a Lorentzian nanotube distribution will not be valid [39] so we do not attempt to quantify the alignment here, but note that using suitable Raman measurements, any distribution of nanotube orientations can in principle be determined, and so full characterization of the nanotube reorientation under strain is possible. After the release of tensile stress from the sample, the Raman intensity data is almost flat, (solid diamonds in Fig. 8) indicating that much of the original orientation of the nanotubes in the sheared film has been lost and that therefore some plastic deformation has occurred in the composite.
4. Conclusions SWNTs provide an unprecedented level of reinforcement (by weight) to an RTV rubber matrix. The experimental evidence for SWNT and other carbon fillers suggests that this is due to the high aspect ratio and low density of the nanotube bundles and that well dispersed single nanotubes should provide even better reinforcement. A ‘pseudo-yield’ occurs in the composites under strain and this has been observed both in macroscopic tensile tests and in situ measurements of the strain dependent Raman spectrum of the nanotubes themselves. The Raman measurements provide extra insight into the nature and efficiency of the interface in the composites all well as quantification of the straininduced reorientation of the nanotubes towards the axis of tensile strain.
Acknowledgements We are grateful to Dr. C.A. Cooper and Miss Q. Zhao for valuable contributions to this work. This project was supported by the (CNT) Thematic European network on ‘‘Carbon Nanotubes for Future Industrial Composites’’ (EU), the G. M. J. Schmidt Minerva Centre of Supramolecular Architectures, and by the Israeli Academy of Science. H.D. Wagner is the recipient of the Livio Norzi Professorial Chair.
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Mechanical properties of carbon nanoparticle-reinforced elastomers Mark D. Frogley, Diana Ravich, H. Daniel Wagner* Department of Materials and Interfaces, Weizmann Institute of Science, Rehovot 76100, Israel Received 3 September 2002; received in revised form 17 December 2002; accepted 28 December 2002
Abstract Silicone based elastomers have been mixed with single-wall carbon nanotubes or larger carbon nanofibrils. Tensile tests show a dramatic enhancement of the initial modulus of the resulting specimens as a function of filler load, accompanied by a reduction of the ultimate properties. We show that the unique properties of the carbon nanoparticles are important and effective in the reinforcement. The modulus enhancement of the composites initially increases as a function of applied strain, and then at around 10–20% strain the enhancement effect is lost in all of the samples. This ‘‘pseudo-yield’’ in elastomeric (or rubber) composites is generally believed to be due to trapping and release of rubber within filler clusters. However, insitu Raman spectroscopy experiments show a loss of stress transfer to the nanotubes suggesting that instead, the ‘‘pseudo yield’’ is due to break-down of the effective interface between the phases. The reorientation of nanotubes under strain in the samples may be responsible for the initial increase in modulus enhancement under strain and this is quantified in the Raman experiments. # 2003 Elsevier Ltd. All rights reserved. Keywords: B Interface; B. Mechanical properties, Raman spectroscopy; Nanomaterials; Elastomeric composites
1. Introduction The high and reversible deformability of elastomers is of great industrial importance. Typically however, the initial modulus and durability of such materials is low, and an additional reinforcing phase is required for practical use. Carbon black and silica particles have been used extensively for this purpose [1–4]. For the composite to be effective, there must be a strong interaction between the matrix and the stiffer phase and this can be achieved using a filler with a large surface-area-to-volume ratio. Optimally this means using small fillers with a large aspect ratio, and recently, particles with nanoscale dimensions such as flakes [5–7], nanofibers [8–9] or hollow nanotubes [10– 15] have become the subject of extensive research. Unprecedented improvement of the mechanical properties has been observed in these nanocomposites. The key issues for nanoparticles, which ultimately determine their usefulness as fillers, are their effective dispersion in * Corresponding author. Tel.: +972-8-934-2594; fax: +972-8-9344137. E-mail address: [email protected] (H.D. Wagner). 0266-3538/03/$ - see front matter # 2003 Elsevier Ltd. All rights reserved. doi:10.1016/S0266-3538(03)00066-6
the matrix (with large surface/volume ratios the particle-particle interactions are strong) and the nature of the interface with the matrix. But for nanoparticles, how do we test these properties? Bulk tests, such as frictional wear or tensile testing, can tell us if a material performs well, but to understand and improve the performance, results must be related to the interaction between individual nanofillers and the matrix. This is directly possible for some properties, and the filler dispersion as well as wetting, buckling and pull-out [5,8,13,16–18] have been observed by high-resolution electron microscopy. However, this technique yields mostly qualitative information since the local stresses and strains on the nanotubes remain unknown. Conventional tests of the fiber–matrix interaction such as fiber pull-out (whilst recording the stress in the fiber), or fragmentation tests can be performed in principle with nanoscale fibers, but as well as the technical difficulty involved, interpretation of the results is difficult. It is not yet clear whether standard micromechanical theories can be extrapolated to the nanoscale, where individual fibers are comparable in size to the polymer chains. In this paper we study the mechanical properties of a silicone rubber reinforced
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with single-wall carbon nanotubes or with larger, carbon nanofibrils. We perform tensile tests on bulk specimens, which show dramatic improvement in the mechanical properties as a function of filler content. For the nanotube samples we use Raman spectroscopy of the nanotubes to relate the macroscopic results to the nanoscale behaviour.
2. Experimental The elastomer was RTV silicone rubber (Epusil 186, Polymer G’vulot ltd., Israel). The fillers were either single-wall carbon nanotubes (SWNTs) produced by the HiPCo technique, (Carbon Nanotechnologies Inc., USA) or vapor-grown carbon nanofibrils (Polygraf III, Applied Sciences Inc., USA). Nanofibrils are highly graphitic single fibers with lengths up to a few hundred microns and diameters around 200 nm. Composites with good fiber dispersion, as determined by electron microscopy, were prepared by first dissolving the RTV in toluene (1 mg/ml) to reduce the viscosity, and separately dispersing the fillers in toluene by ultrasonication (0.1 mg/ ml, 10 min, 6 W). Then the two dispersions were mixed and further ultrasonicated for 10 min. The toluene was removed by evaporation at 50 C over several days with continuous stirring and finally the hardener was mechanically mixed in, and air bubbles were removed by vacuum pumping. Films of thickness 200–300 mm were produced by shearing the uncured mixture across a glass plate with a doktor blade. After curing for 24 h at room temperature, simple beams of width 1.5 mm and length 20 mm were cut from the films. Samples with filler concentrations up to 1 wt.% (SWNT) and 4 wt.% (fibrils) were prepared but samples with higher concentrations were too viscous for bubbles to be effectively removed and for spreading uniform films. For tensile tests an Instron 4502 machine was used and samples of gauge length 10 mm were stretched at 100 mm/min until failure. Ten samples of each composition were tested for statistical accuracy. Throughout this work, the quoted strain is the true strain, i.e. the percentage elongation, and the stress is the nominal or engineering stress (the force divided by the cross sectional area of the unloaded sample). For Raman measurements on the nanotube composites, the samples were stretched using a home made mini-tensile tester [19], which was placed under a Renishaw Raman microscope. Spectra were obtained in the backscattering geometry (shown inset in Fig. 5) using 632.8 nm laser light, which was polarized along the strain axis of the tensile tester. The laser beam was focused to a spot 20 mm in diameter on the sample to avoid heating effects, and the same point on the sample was used for all measurements.
3. Results and discussion 3.1. Mechanical tests Fig. 1 shows representative stress–strain curves for pure RTV and for RTV/nanoparticle composites. For both SWNTs and fibrils, the initial modulus (measured by fitting a straight line to the data below 10% strain) is radically improved as the filler concentration increases as shown in Fig. 2. For RTV/SWNT composites the modulus increase is approximately linear with weight fraction, with a slope of 200%/wt.% or a maximum of 200% increase in stiffness for 1% SWNT. This compares with 120%/wt.% for the nanofibrils measured here and about 100%/wt.% for a nanofibril/rubber composite made by Canales and Cavaille [20]. For carbon black filled rubbers, the rate is as little as 4%/wt.% [1,2] and for silica particles in styrene-butadiene rubber [3], where there is chemical bonding between phases, the modulus increases about 2–3 times for a loading of 24 wt.% (a rate of increase a few times higher than for carbon black). We conclude from these results that the
Fig. 1. Stress–strain curves for RTV/carbon nanofiber composites. The stress is the engineering stress (applied force divided by the unloaded cross sectional area) and the strain is the true strain.
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modulus of the two phases as well as filler packing effects, but for low loadings of rod-like fillers which are much stiffer than the matrix (as in our samples) it reduces to [22]: 1 þ 2fc Epara ¼ E0 1c 1 þ 0:5c Eperp ¼ E0 1c Erand ¼ 0:2Epara þ 0:8Eperp
Fig. 2. (a) Initial modulus of the composites as a function of filler weight fraction. Error bars represent the scatter of ten experiments for each composition. The solid curves are the theory of Guth [Eq. (1)] with filler aspect ratios, f, as shown. The dashed curve is typical of conventional fillers such as carbon black and silica particles. (b) The modified Halpin–Tsai theory [Eq. (2)] is compared with the data. In the key, O refers to fillers perfectly oriented along the strain axis and R refers to fillers randomly oriented in three dimensions.
specific properties of the carbon nanotubes and fibrils are important in the reinforcement. The well known model of Guth [1] for the modulus of rod-like particle reinforced elastomers is based solely on the aspect ratio and volume fraction of the filler, and does not take any other properties of the filler into account, except to assume that the fillers are stiff compared with the matrix: E ¼ E0 1 þ 0:67fc þ 1:62f 2 c2 ð1Þ where E and E0 are the moduli of the composite and the pure matrix respectively, and f and c are the aspect ratio and volume concentration of fillers. Another commonly used model for filled polymers is the modified HalpinTsai theory [21], which in principle takes account of the
ð2Þ
for the Young’s modulus measured parallel (Epara) or perpendicular (Eperp) to perfectly oriented fibers or in any direction (Erand) for fibers randomly oriented in three dimensions. In Fig. 2a, Guth’s theory [Eq. (1)] is compared with the experimental data, and shows good agreement using f 120 for the nanotubes and f 70 for the fibrils. These values are consistent with observations of the brittle fibrils after they were broken up by the ultrasonication procedure and with SWNT bundles of diameter 10–15 nm, and lengths of a few microns. SWNT bundles have a stiffness comparable to single tubes [23,24] and are often observed in the RTV and other matrices by electron microscopy [16,18]. Individual SWNTs are rarely seen. The curve for a fiber of aspect ratio 1000, which is typical for an individual SWNT, does not agree with the data, but is a clear incentive to develop methods to separate the nanotubes. In Fig. 2b, the modified Halpin–Tsai equation is compared with the same data. Assuming a random distribution of fibers, aspect ratios of 700 and 500 are needed for the nanotubes and fibrils respectively to achieve a good fit: much higher than observed. Raman measurements described later suggest that there is a degree of fiber alignment during sample preparation (shearing) but it is far from perfect (Fig. 8, or see e.g. Ref. [25]). Even perfectly oriented fibers require aspect ratios of 140 and 100 for the theory to fit the data. Lower aspect ratios lead to a prediction of a lower modulus and so for both theories, the measured level of reinforcement is either consistent with or better than predicted by theory, which is reasonable because the theories ignore the increase in surface-area-tovolume ratio for smaller fillers, even when the aspect ratio is kept constant. Note that, to use Eqs. (1) and (2), the measured filler weight fractions were converted to volume fractions, c, using densities of 1800, 1340 and 1215 kg m3 for fibrils, SWNTs [26] and the matrix respectively. Others have used nanofillers with high aspect ratios. Kim and Reneker [9] used electrospun fibers of polybenzimidazole (PBI) in a styrene-butadiene rubber. The fibers had an average diameter of 300 nm and f1000 (after processing) and for a 10 wt.% loading, the
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modulus was enhanced by a factor of 10. This result is close (by weight) to that of the carbon fibrils. Osman et al. [5] used plate-like mica particles with various aspect ratios as fillers in PDMS, and showed that the aspect ratio plays a crucial role in the reinforcement. The best reinforcement was seen for the highest aspect ratio (f=120) where the modulus was improved by a factor of 13 for a 14 vol.% filler load [5]. However, the Guth model, which is for rod-like fillers [Eq. (1)] suggests a much stronger increase is possible for SWNTs. If the volume fraction of SWNT bundles could be increased to 14% the expected modulus increase would be almost 400 times (extrapolating the fit of Fig. 2). Several works have appeared regarding PDMS/OMT (organo-montmorillonite) nanocomposites [6-7], which are plate-like in nature, but the enhancement of stiffness is somewhat less impressive than those discussed so far. SWNT composites represent the highest initial level of reinforcement, by weight of filler, of any elastomeric material to date. This reinforcement comes at a price. The ultimate strain of the SWNT and fibril composites is significantly less than the pure RTV (Fig. 1) and the tensile test results are summarized in Fig. 3. All the data fall on one curve, with the reduction in ultimate strain being quite dramatic at first, and then leveling off to around 200% strain (60% of the value for pure RTV) for large concentrations. Very similar behavior was seen for the high aspect ratio (f=120) PDMS/mica nanocomposites [5], for which the value stayed on the low plateau up to the highest concentrations. This behavior is also seen in other fiber-filled rubbers [9,27] whereas for low loads of near-spherical fillers, a less marked reduction, or even an increase [5,9] in the ultimate strain has been
observed. In any case, for most applications requiring a large initial reinforcement, the ultimate strain is not critical. Fig. 4 shows the modulus measured at 80% strain, as a function of filler content. The modulus is constant within the experimental scatter, that is, there is no obvious effect of the filler at high strain, in contrast to the huge reinforcement at low strain. Several studies [3] have revealed such behavior, termed a ‘pseudo-yield’ in filled rubbers and it has been attributed to breakdown of secondary particle structure, (aggregates) and the resulting release of unstrained, trapped rubber. The degree of secondary structure should not be significant for low volume fractions however, especially if the fillers are well separated like the fibrils. For our SWNT composites the tight binding of the SWNT bundles makes it unlikely that rubber will penetrate the bundle, and so any trapped rubber would have to be in some tertiary structure formed from many bundles for SWNTs. We have seen no evidence of microscale stucture in our samples. We propose that, at least in our composites, the pseudo-yield is due to weakening of the frictional interface between the two phases—that is, filler dewetting—and we return to this point later. The mere existence of the pseudo-yield demonstrates that the fillers, and their interface with the matrix, are responsible for the reinforcement of the polymer, and not curing effects such as increased crosslinking [22] induced by the fillers. If this were the case, the stiffness of the composites at high strain (Fig. 4) would be larger for higher crosslinking densities (higher filler loads). The marked increase in viscosity of the nanotube-monomer mixture (before curing), as a function of filler load, supports this argument.
Fig. 3. Strain to failure as a function of filler weight fraction for the composites. Error bars represent the scatter of ten experiments for each composition.
Fig. 4. The modulus of the composites, measured at 80% tensile strain, as a function of filler content.
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Having shown that carbon nanofibers have a dramatic effect on the bulk mechanical properties of RTV rubber we now relate this, for the carbon nanotube composites, to the interaction between the individual nanotube bundles and the matrix. In stiff polymers, like poly (methyl methylacrylate) (PMMA) and polyurethane acrylate (PUA) or epoxy resins, which have elastic moduli about 1000 times greater than rubbers, embedded nanotubes undergo significant mechanical compression during the polymer curing process, and are also strongly affected by external stress applied to the cured composite. This has been observed by looking at changes in the Raman spectrum of the nanotubes, which are related to mechanical deformation of the tubes [10–12,15,25,28–30]. Specifically, the D* Raman mode of SWNTs exhibits a large spectral shift when the composite is strained in uniaxial tension, and the shift is approximately linear with the applied strain until the polymer yields [25,29–31] which happens at about 1% strain for an epoxy matrix as shown inset in Fig. 6 [31]. After the polymer yields, there is no further change in the Raman spectrum, and this indicates that stress is no longer efficiently transferred from the matrix to the nanotubes. Raman spectroscopy is therefore sensitive to the strength of the interface between the individual nanotubes and the matrix, and we can interpret a quench in the Raman strain-shift of the D* mode as an indication of interfacial breakdown. Fig. 5 shows a typical spectrum of the D* peak for SWNTs in RTV. Fig. 6 shows the peak position for nanotubes in the 0.3
wt.% composite, as a function of applied tensile strain. The Raman wavenumber decreases linearly with strain at first, and tends to a constant value by about 50% strain. There are two important features of this data. The first is the shape of the curve, which tells us about the rate of stress transfer to the nanotubes. If the interfacial strength is constant, we expect a linear wavenumber–strain relationship, even for large deformations of the nanotubes (Raman strain shifts of 10–20 cm1) as seen e.g. inset in Fig. 6. Non-linear behavior indicates a change in stress transfer efficiency. Above 40% strain in Fig. 6, the Raman wavenumber has an almost constant value, implying that stress is no longer transferred effectively to the tubes. This quench in strain-shift corresponds to the pseudo-yield phenomenon, which is apparent over the same strain region in the stress–strain curves of Fig. 1. The pseudo-yield is shown more clearly, for the 0.3 wt.% SWNT sample, in a plot of modulus enhancement against strain (Fig. 7). The modulus enhancement factor at a given strain is defined as the modulus of the composite at that strain divided by the equivalent modulus for pure RTV, and was calculated from the stress-strain curves of Fig. 1. This is a relative parameter, so it does not matter if the true or engineering stress is used. The modulus enhancement increases initially with strain, and then falls quickly after about 12% strain. This is the same strain value at which the Raman data departs from a linear function in Fig. 6, suggesting that the ‘‘pseudo-yield’’ is due to a loss of stress transfer between the individual nanotube bundles and the matrix, which may be a result of the straightening of the polymer chains. This is further
Fig. 5. The Raman spectrum of an RTV/0.3 wt.% SWNT composite in the vicinity of the D* band for SWNTs (solid curve). The crosses represent the mixed Gaussian/Lorentzian function used to fit the three peaks near 2420, 2500 and 2600 cm1 (D*). Inset: setup for polarized Raman spectroscopy of the sample film in the backscattering geometry. E is the electric vector of the laser light.
Fig. 6. The Raman wavenumber of the D* band for SWNTs embedded in RTV rubber is plotted against tensile strain applied to the composite. The initial slope is 0.08 cm1/% strain (solid line). Inset: The strain induced shift of the Raman wavenumber of the D* band for SWNTs embedded in an epoxy matrix is plotted against tensile strain applied to the composite. After the polymer yields at 1% strain, there is no significant change in the wavenumber.
3.2. Nanoscale behaviour—Raman spectroscopy
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Fig. 7. The modulus enhancement factor for the RTV/0.3 wt.% SWNT composite as a function of applied strain. The enhancement factor is the modulus of the composite divided by the modulus of pure RTV at each strain. Increased noise below 5% strain is from division of small stress values.
evidence that the yield is not due to a release of trapped rubber from within filler clusters (the suggested mechanism for conventional filled composites) in this material, as discussed earlier. The second interesting feature in the Raman data of Fig. 6 is the magnitude of the spectral shift. The total
Fig. 8. Polarized Raman intensity for SWNTs embedded in RTV as a function of the angle between the sample axis (the tensile strain axis) and the polarizer axis. A constant (normalized) intensity of 1 is expected for totally unoriented tubes, whereas the lower dashed curve is expected for perfectly oriented tubes. The data points are for an RTV/0.3 wt.% SWNT composite at various strains, showing that the nanotubes become oriented as the tensile strain is increased. The solid curve is a theoretical fit to the data at 200% strain using the method of Hwang et al. [14] as explained in the text. All data are normalized to the intensity at =0.
shift is only 2 cm1, over 50% strain, which is much less than 7–18 cm1 shift after 1% strain for SWNTs in stiffer polymers [25,29–31]. In stiffer matrices, the rate of Raman shift with strain depends on the thermal compression of the matrix around the nanotubes, which occurs on cooling after curing at high temperature because of the mismatch in thermal expansion coefficient of the two phases [29]. For tensile tests at lower temperatures, the rate of wavenumber shift with applied tensile strain is higher and this is thought to be due to a more effective interface resulting from enhanced lateral clamping of the nanotubes [32]. The RTV/SWNT composites, however, were cured at room temperature over a long time and so the radial thermal compression or ‘clamping stress’ is low. This results in a poor interface, which must be non-strain-identical in nature: if the nanotubes were being stretched at the same rate as the matrix, the resulting wavenumber strain-shift would be comparable with the epoxy data in Fig. 6. The axial thermal stress may play a more significant role. Traditional carbon fibers embedded in a matrix which undergoes significant thermal contraction in the post-curing cool-down are seen to go into compression via shear stress transfer across the interface [33,34]. If the curing takes place at room temperature over a long time, the residual axial stress in the fiber is much lower, or zero [34]. Similarly, for nanotubes, a large compressive Raman shift (increase in D* wavenumber of order 10 cm1) is observed after cooling the high-temperature-cure polymers [11,28,29]. The compressive stress stores energy in the nanotube, like a coiled spring and so when the polymer is subsequently stretched, the nanotube will return to its original length as quickly as the surrounding polymer allows. Hence the tensile Raman strainshift is large. This has been observed in various SWNT composites and the increase in wavenumber due to thermal compression is always similar to the decrease in wavenumber on subsequent stretching to yield. In silicone rubber however, which is cured at room temperature, there is no thermal stress, and so the nanotube is close to its equilibrium (in the uncured RTV liquid) length. Under tensile strain therefore, extension of the nanotubes relies directly on the weak interface so the wavenumber strain-shift is much smaller, as observed. We rule out direct effects of radial stress on the Raman wavenumber since there is no shear stress transfer mechanism in the radial direction and the small normal stresses in the polymers (of the order of MPa) are orders of magnitude too small to significantly affect the D* wavenumber as seen in high-pressure experiments on SWNT bundles [29]. Nanotube bending effects also must not affect the wavenumber shift since the degree of bending will be much higher in the RTV (over 50% strain) than over 1% strain in epoxy, whereas the wavenumber shift is much lower in the RTV.
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3.3. Orientation of the nanotubes In composites at large strains, originally unoriented fibers, for example, will align strongly in the direction of applied tensile strain, which is a simple consequence of the change in shape of the rubber. The effect is important because oriented fibers provide better reinforcement (along the orientation direction) than unoriented fibers [22]. Nanotube reorientation can be monitored by polarized Raman spectroscopy because the intensity of the polarized Raman spectrum of a carbon nanotube depends strongly on the angle, y, between the nanotube axis and the optical polarization direction [35,36]. For single-wall carbon nanotubes excited with polarized light of wavelength 632.8 nm, the Raman scattering is resonant [36,37] so that for a single nanotube, or for tubes oriented perfectly along one direction, the total intensity of the Raman modes varies as [14,36]: RðÞ / cos4
ð3Þ
For a sample of randomly oriented nanotubes the total Raman intensity as a function of polarization angle is a constant, and any systematic deviation from this indicates a degree of nanotube alignment. Fig. 8 shows the normalized intensity of the D* mode as a function of the angle, , between the sample axis and the optical polarization direction, in the plane of the film. The sample axis is defined as the direction of shear flow (used to produce the RTV films), which is also the direction of applied tensile strain and therefore is the nominal axis of nanotube alignment. With no applied strain, (up triangles) there is already some orientation, and this is due to the shear flow processing method. As tensile strain is applied the degree of orientation increases and by 200% strain the data is close to cos4, the dashed curve in the figure, which is expected for perfect alignment. To quantify the degree of alignment from the Raman data we use the method of Hwang et al. [14]. This assumes that the number of nanotubes N() at an angle to the alignment axis (in the plane of the film) is a Lorentzian function centered on the alignment direction, and that the contribution of each nanotube to the Raman signal varies as cos4(-), where - is equal to and is the angle between that nanotube and the polarizer axis, from Eq. (3). One then obtains a best-fit to the data by varying the width (FWHM) of the Lorentzian, and this width is the measure of the degree of alignment. We obtain a FWHM of 30 for the data at 200% strain (Fig. 8), i.e. half of the nanotubes lie within 15 of the alignment axis. Strictly the procedure is only valid for nanotubes in the plane of the film but our sample approximates this at high strain because during the shearing process used to produce the films, the alignment of nanotubes into the plane of the film is substantially stronger than the alignment of tubes within the plane, towards the shearing direction [25,38]. This
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effect is magnified at high tensile strain. At lower strains, both the two-dimensional approximation and the assumption of a Lorentzian nanotube distribution will not be valid [39] so we do not attempt to quantify the alignment here, but note that using suitable Raman measurements, any distribution of nanotube orientations can in principle be determined, and so full characterization of the nanotube reorientation under strain is possible. After the release of tensile stress from the sample, the Raman intensity data is almost flat, (solid diamonds in Fig. 8) indicating that much of the original orientation of the nanotubes in the sheared film has been lost and that therefore some plastic deformation has occurred in the composite.
4. Conclusions SWNTs provide an unprecedented level of reinforcement (by weight) to an RTV rubber matrix. The experimental evidence for SWNT and other carbon fillers suggests that this is due to the high aspect ratio and low density of the nanotube bundles and that well dispersed single nanotubes should provide even better reinforcement. A ‘pseudo-yield’ occurs in the composites under strain and this has been observed both in macroscopic tensile tests and in situ measurements of the strain dependent Raman spectrum of the nanotubes themselves. The Raman measurements provide extra insight into the nature and efficiency of the interface in the composites all well as quantification of the straininduced reorientation of the nanotubes towards the axis of tensile strain.
Acknowledgements We are grateful to Dr. C.A. Cooper and Miss Q. Zhao for valuable contributions to this work. This project was supported by the (CNT) Thematic European network on ‘‘Carbon Nanotubes for Future Industrial Composites’’ (EU), the G. M. J. Schmidt Minerva Centre of Supramolecular Architectures, and by the Israeli Academy of Science. H.D. Wagner is the recipient of the Livio Norzi Professorial Chair.
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