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Large-Scale Orientation in a Vulcanized Stretched Natural Rubber Network: Proved by In Situ Synchrotron X-ray Diffraction Characterization Gengsheng Weng, Guangsu Huang,* Liangliang Qu, Yijing Nie, and Jinrong Wu State Key Laboratory of Polymer Material Engineering, College of Polymer Science and Engineering, Sichuan UniVersity, 610065, Chengdu, P. R. China ReceiVed: January 31, 2010; ReVised Manuscript ReceiVed: April 10, 2010
In situ studies of strain-induced crystallization in unfilled and multiwalled carbon nanotube (MWCNT)-filled natural rubber (NR) were carried out by using synchrotron wide-angle X-ray diffraction (WAXD). Synchrotron WAXD results indicate that more nuclei appear in the MWCNT-filled NR sample, leading to higher crystallinity, lower onset strain of crystallization, and remarkable enhancement in tensile strength. During deformation, despite the amorphous chains remaining in isotropic orientation, the domains of larger scale (10-100 nm) with high network chain density in the NR matrix are oriented. The MWCNTs induce significant variation of this orientational process, and it is monitored by the stearic acid (SA) crystallites, which are effective nanoprobes of the amorphous phase. The results indicate that a small amount of MWCNTs and SA crystallites can be used as new tools to analyze the microstructural orientation of NR during deformation. The results also yield new insight into the strain-induced crystallization mechanism. 1. Introduction In the community of rubber technologists, natural rubber (NR) has been recognized as a fascinating industrial material that has high elasticity and tensile strength. It is indispensable for pneumatic tires and rubber bearings in a seismic isolation system.1,2 Thin NR films obtained from NR latex are also very important in the biomedical and health care fields. The wide applications of NR can be attributed to its ability to crystallize upon deformation,3–9 which stops crack propagation at ultimate strain.8,9 Strain-induced crystallization origins from the microstructural evolution of NR upon stretching. Toki et al. suggested that the strain-induced crystallization is ascribable to the nonuniformity of NR network structure,7,10–12 which was proved by Ikeda’s group based on small angle neutron scattering analysis.13,14 According to their results, the existence of a denser network domain with a characteristic length scale (Ξ) of 10-100 nm was confirmed. These investigations revealed the complexity of the NR network structure. It is known that NR crystallites always appear in the form of oriented crystals and grow in the direction perpendicular to the molecular chain axis.7,10,15–17 So, the orientational network chain should be the precursor for the formation of NR crystallites. However, Katz’s early work in 1925 has described the existence of the isotropic amorphous halo of stretched NR using conventional wide-angle X-ray diffraction (WAXD).18 Mitchell studied the molecular orientation in cross-linked NR, and the results showed a low orientation degree of molecular chains even at very high strain.19 Although the isotropic halo from the stretched NR may be the result of relaxation during long-time exposure when measured by using a conventional X-ray source, the subsequent synchrotron WAXD measurements still showed a large fraction of isotropic scattering halo during stretching.10,20–22 Besides, it was estimated that only 5-10% of the oriented amorphous phase exists in the stretched NR at a strain of 5.10,23,24 * Corresponding author. E-mail:
[email protected]. Tel.: +86-028-85463433. Fax: +86-028-85463433.
Other approaches such as birefringence and infrared dichroism techniques also showed very a low orientation degree of molecular chains.25,26 On the contrary, Rault et al. studied the chain orientation in NR by means of 2HNMR, and they claimed that no chain segment of NR displays isotropic orientation, which is in complete contradiction with the conclusions of Toki and Tosaka et al.27,28 Thus, the description of amorphous orientation of NR upon deformation is still a controversial topic. Despite the contrary results, Rault et al. did not rule out that different domains (with different strain and different network chain density) exist. Particularly, they employed a small amount of deuterated oligomer chains as probes in the NR network to characterize the orientation of the network chains by deuterium NMR measurements. The orientational order of the probes and of the network chains is identical or proportional, depending on the chemical structure of the probes.29 Therefore, this approach is applicable for the understanding of the microstructure motions of polymers.30,31 Stearic acid (SA), which is an intrinsic ingredient for most of the NR vulcanizates, has also been used as a nanoprobe of the amorphous chains in NR.27 Compared with the “foreign” deuterated oligomers, SA crystallites are more appropriate to reflect the orientation of amorphous chains, because the deuterated oligomers may perturb the microstructure motions of NR. Although the orientation of SA crystallites has been widely observed, the physical evolution process is still unknown.8,11,27,32 Nanofillers are of remarkable influence on the strain-induced crystallization of NR because of their strong interfacial interaction with polymer matrix. By the inclusion of nanofillers, the network structural evolution of a polymer matrix under deformation is quite different. These have been investigated by our previous work and the research of others,33–35 whereas they were limited in the nanoclay-filled systems. In this work, we studied the strain-induced crystallization of multiwalled carbon nanotube (MWCNT)-filled NR with strong interfacial interaction. For comparison, pure NR was also prepared using the same method. The strain-induced crystallization behavior of these samples was investigated by syn-
10.1021/jp100920g 2010 American Chemical Society Published on Web 05/10/2010
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TABLE 1: Recipes of the CNT-Filled NR Composites loading level (phra)
ingredients NR ZnO SA accerlerater accerlerater accerlerater sulfur MWCNTs curing time
100 5 2 0.5-1.5 0.5-1.5 0.5-1.5 0.5-1 2 15
TTb DTDMc DMd (min)
a Parts by weight per hundred parts rubber (phr). b Tetramethylthiuram disulfide. c 4,4′-Dithiodimorpholine. d N-Cyclohexyl-2-benzothiazolylsulfenamide.
chrotron WAXD measurements. The MWCNTs are very important in this research, because they enable us to build contrastive systems of the orientation process. Contrastive observation and analysis revealed that a small amount of MWCNTs and SA crystallites can be used as new tools to analyze the microstructural orientation of NR during deformation. Finally, a comprehensive understanding of the straininduced crystallization mechanism of NR, especially the network structure of NR and its orientational evolution, was proposed. 2. Experimental Section 2.1. Materials. The NR used in this study was ribbed smoked sheet (RSS) No.1 from Indonesia. NR latex (NR content 60 wt %) was purchased from the China Academy of Tropical Sciences. MWCNTs investigated in this study were supplied by Chengdu Organic Chemicals Co., Ltd., Chinese Academy of Sciences. The specific area of the MWCNTs is 400 m2 · g-1, and the diameter is ∼20 nm. 2.2. Preparation of MWCNT/NR Composites. MWCNTs were stirred in a mixture of sulfuric and nitric acids (3:1, 98% and 70%, respectively) at 40 °C for 48 h, and washed several times with deionized water until the filtrate showed neutral pH to remove the amorphous carbon and metallic nanoparticles (used as a catalyst during synthesis). The purified MWCNTs were dispersed in deionized water and sonicated for 30 min. Then an appropriate quantity of NR latex was mixed with MWCNTs in water. After the demulsification and drying process, mater batches were obtained. Two weight percent MWCNT-filled NR (CNR) was prepared by mechanical mixing at room temperature and vulcanized at 143 °C. The formulations of various components of CNR are shown in Table 1. 2.3. Measurements and Characterization. The morphologies of the composites were observed and investigated by means of transmission electron microscopy (TEM, JEOL JEM 2010). It was operated under an acceleration voltage of 200 kV. Rapid solvent swelling measurements (toluene, 72 h at 30 °C) was used to determine the network chain density (n) by the application of the Flory-Rhener equation:36
1 -ln(1 - φr) - φr - χrφ2r ) nV0 φ1/3 r - φr , 2
[
]
Mc ) F/n
(1) where φr is the polymer volume fraction in the swollen network, V0 is the molar volume of the solvent (106.2 mL/mol for toluene), χ is the Flory-Huggins polymer-solvent interaction term (0.393 for NR/toluene), Mc is the average mass of network chains, and F is the density of the rubber (0.92 g/mL for NR).
This value (0.92) is also used for the CNR, considering the low content of MWCNTs. Tensile performances were tested by an Instron 5567 material testing machine. The specimen was a dumbbell-shaped thin strip (25 × 6 × 2 mm) and experiments were performed with a tensile speed of 500 mm · min-1. Synchrotron WAXD experiments were carried out under room temperature at the U7B beamline at the National Synchrotron Radiation Laboratory (NSRL), University of Science and Technology of China, Heifei, China. The wavelength used in U7B was 0.154 nm. The two-dimensional (2D) WAXD patterns were recorded in every 180 s by Mar CCD 165 X-ray detector system. The drawing measurements were performed at room temperature (25 °C). The samples (12 × 1.5 × 1 mm) were mounted between two clamps of a homemade miniature mechanical tester and the drawing speed was 4.2 mm · min-1. 3. Results Figure 1 shows the normalized WAXD patterns with respect to the same thickness change, sample absorption, and beam fluctuation for pure NR and CNR. Both samples show highly oriented crystalline diffraction peaks in the images when strain (R) is greater than ∼3. The intensities of these peaks increase with strain during stretching, and interestingly they are stronger in CNR than in pure NR at the same strain. It is noteworthy that, on one hand, all images exhibit a preferred orientation of crystallographic (001), (002), and (003) planes from SA, indicating that SA crystallites become highly oriented along the stretching direction (see the arrows). On the other hand, there is no strong evidence for the orientation of the amorphous phase based on the WAXD analysis (see the amorphous halo in the images), suggesting that, even if there is some oriented amorphous phase, a large part of amorphous phase is still isotropic during stretching. Air scattering was subtracted from all the WAXD patterns for further analysis accordingly by using the program FIT2D. Figure 2a,b represents the typical three-dimensional (3D) diffraction patterns before and after subtraction. These patterns show the intense anisotropy of both NR and SA crystallites. The equatorial intensity profile of Figure 2b has been shown in Figure 2c. Three diffraction peaks of SA can be clearly observed at 2θ ) 2.1 (001), 2θ ) 4.3 (002), and 2θ ) 6.5 (003). Two diffraction peaks of NR can also be recognized near 2θ ) 14 (200) and 2θ ) 21 (120). In order to monitor the evolution of strain-induced crystallization, the linear diffraction profiles taken along the equator from the 2D WAXD patterns collected at different strains during stretching are shown in Figure 3a (CNR) and b (pure NR). Apparently, the peak intensities of CNR are stronger than those of pure NR. The distinct crystal diffraction peaks of NR first appear at R ) 2.8 in CNR and R ) 3 in pure NR, indicating that CNR crystallizes at a smaller strain than pure NR even at such low MWCNTs content. According to the linear diffraction profiles obtained above, the variation of lateral crystallite size during deformation was analyzed. In this research, the widths of the 200 and 120 planes were estimated to evaluate the values of lateral crystallite size. Each peak on the equatorial diffraction profiles was fitted with a Gaussian function:
I(x) ) h exp[-(x - xc)2 /(2w2)]
(2)
where I(x) is the intensity at position x, and xc refers to the position at the scattering maximum. The parameter h and w
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Figure 1. Sequential variation of WAXD patterns from pure NR (a) and CNR (b). Corresponding strain values are indicated at the left top in each pattern of panels a and b. The enlarged image of the pattern taken at the strain of 4.18 from panel b is shown in panel c, in which indices of crystallographic planes of NR and SA crystallites are indicated.
indicate the peak height and the peak width, respectively. The half-width β was obtained following the conversion procedure described by Tosaka.11 The Scherrer equation was used to calculate the lateral crystallite size:
Lhkl )
Kλ β cos θ
(3)
where Lhkl is the lateral crystallite size in the direction perpendicular to the (hkl) plane, K is the Scherrer factor, and a value of 0.89 was used in this study,37 β is the half-width of the hkl plane in the radial direction, θ is the Bragg angle (half of the scattering angle), and λ is the wavelength. Figure 4 illustrates the variation of L200 and L120. It represents that both L200 and L120 values decrease during stretching. Additionally, the pure NR has larger L200 and L120 values than the CNR. A similar observation has been reported for NR filled with carbon black and calcium carbonate.12 The decrease of lateral crystallite size is attributed to the dense crystal formation on the stretched chains and limited availability of coiled network chains.12 So, it can be concluded that the MWCNTs densify the nuclei in the NR matrix upon deformation.
WAXD patterns were normalized and integrated along the azimuthal direction from 0° to 360°. The normalized air scattering patterns were also integrated and then used for subtracting. The resultant profiles were deconvoluted into individual indexed peaks and amorphous haloes using the Levenberg-Marquardt method.38,39 Gaussian peaks were chosen for this process. The crystallinity is estimated by the following formula:40
Xc )
Ac × 100% Ac + Aa
(4)
where Ac and Aa are the areas of crystalline and amorphous regions, respectively. Figure 5 shows the evolution of crystallinity with strain for pure NR and CNR. The results indicate that the onset strain of crystallization (R°), determined by the interception of the linear regression lines in Figure 5, is 3 and 2.8 for pure NR and CNR. The lower R° of CNR reveals that an earlier orientation of polymer chains occurs, which results from the high level of interaction between polymer chains and MWCNTs due to the high specific area of 400 m2 · g-1.
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Figure 2. The 3D WAXD pattern of CNR at the strain 4.18 before (a) and after subtracting the air scattering (b). The corresponding equatorial diffraction profile of panel b at the 2θ angle range of 1-30° is plotted in panel c. The indices of crystallographic planes of NR and SA crystallites are marked in panel c.
Large-Scale Orientation in a Vulcanized NR Network
Figure 3. Equatorial diffraction profiles taken from 2D WAXD patterns at selected strain values of CNR (a) and pure NR (b) as a function of strain under stretching at the 2θ angle range of 10-28.3°.
This behavior can also be seen in the nanoclay-filled systems.33–35 Here, the most interesting part is that the pure NR exhibits a single crystallization step (the slope s ) 0.09), while CNR has a slope value similar to that of pure NR in the strain range of 2.8-3.8 (s ) 0.08), and a much larger slope value when R is greater than 3.8 (s ) 0.42). This reveals that MWCNTs accelerate the crystallization rate at large deformation. To explain the interesting phenomenon observed above, we investigated the microstructural evolution of MWCNTs in the NR matrix, which may have a significant impact on the mechanism of strain-induced crystallization. MWCNTs are formed from 2D sheets of graphite (known as graphene layers) rolled up to make a defect-free tubular structure.41–43 Previous studies by X-ray diffraction technique show that the diffraction pattern of MWCNTs is dominated by a strong Bragg peak centered at 3.4 Å, which corresponds to the intershell spacing [(002) plane].44,45 So, it is believed that this distinct crystallographic plane can also be seen when proper the content of MWCNTs is filled in the NR matrix. Here, we employed an NR composite filled with 15 wt % MWCNTs for X-ray diffraction analysis. MWCNTs are randomly oriented in the NR matrix without drawing. This random distribution is confirmed by the TEM image in Figure 6. Because of the higher content, we can also see that some MWCNTs contact and entangle each other. Figure 7 shows the 2D diffraction patterns obtained after being stretched by 122% (a) and 365% (b). The diffraction
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Figure 4. Variation of lateral crystallite size: (a) L200 during stretching, (b) L120 during stretching.
Figure 5. Crystallinity as a function of strain for pure NR and CNR.
ring in Figure 7a and the arc in Figure 7b marked by the arrows indicate the (002) plane of the MWCNTs. The rectangular panels attached to each pattern show the intensity with 2θ and azimuth (φ) plots obtained by means of unwrapping the circular diffraction reflections. It is clear that a preferred orientation of MWCNTs parallel to the stretching direction occurs during elongation, since the Bragg intensity is concentrated at two spots centered at φ ≈ 90° and 270°. The (002) Bragg peaks are found to be centered at 26.25° (d spacing ) 3.39 Å). The d spacing is slightly larger than that
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Figure 6. Typical TEM image for 15 wt % MWCNT-filled NR.
of ideal graphite,46–48 which is related to the stacking disorder in the MWCNT nanostructure. Further, the azimuthal intensity distributions of the diffraction patterns corresponding to the (002) plane were determined by the integration along the 2θ axis after subtracting out the air diffraction intensity. A typical azimuthal intensity distribution at the strain of 2.65 is shown in Figure 8, which also suggests that the MWCNTs are aligned along the stretching direction. It is well-known that the distribution of molecular orientation can j 2〉,49–51 which be measured by the orientational order parameter 〈P j 2〉 value ranges is estimated from the azimuthal profile. The 〈P between 0 and 1, where the former corresponds to an isotropic structure, and the latter represents perfect orientation along the director. To calculate the orientational order parameter, the azimuthal profiles can be first fitted to a Maier-Saupe distribution function:52
2
I ) I0 + Aeβcos (φ-φ0)
Figure 8. Typical azimuthal intensity distribution of the (002) plane of MWCNTs obtained at a strain of 2.65 for 15 wt % MWCNT-filled NR. The red curve is the fitting result by using eq 5.
(5)
where I0 denotes the free baseline intensity, φ0 is the azimuth at the position of the maximal intensity, φ is the azimuth, and β is a parameter that determines the width of the distribution. The solid line in Figure 8 indicates the fitted result for azimuthal intensity distribution at the strain of 2.65. The fitted full width at half-maximum of the two peaks is about 45.9°, which means the aligned MWCNTs have a mosaic angle of (23° around the stretching direction at the strain of 2.65. After the fitting procedure, β parameters were obtained, and the orientational j 2〉 can be determined using the following order parameters 〈P 49,53,54 formula:
j 2〉 of Figure 9. Variation of the orientational order parameter 〈P MWCNTs in the 15 wt % MWCNT-filled NR matrix as a function of j 2〉 of MWCNTs after being stretched strain. The arrows indicate the 〈P by 122% and 365% with the corresponding WAXD patterns shown in Figure 7a,b.
∫-11 P2(cos φ)eβcos φd cos φ ∫-11 eβcos φd cos φ 2
j 2〉 ) 〈P
2
(6)
where the function P2(cos φ) is the second-order Legendre polynomial of cos φ, and it often refers to the Hermans Orientation function:19
Figure 7. WAXD patterns obtained from 15 wt % MWCNT-filled NR after being stretched by 122% (a) and 365% (b). The rectangular panels attached to each pattern show the intensity of the (002) plane of MWCNTs with 2θ and azimuth (φ) plots.
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J. Phys. Chem. B, Vol. 114, No. 21, 2010 7185 TABLE 2: Network Chain Density, N (×10-4 mol/cm3), Average Mass of Network Chains, Mc (g/cm3), and the Number of Statistical Segments of a Network Chain, N material
n
Mc
Na
NR CNR
0.98 1.15
9400 8000
138 117
a The length of the statistical segment is approximate to the dimension of the isoprene monomer (molecular weight ) 68 g/ mol),19 so the number of repeat units was used to substitute for n.
Figure 10. Variation of theoretical P2(cos φ) values for NR and CNR as a function of λ.
1 P2(cos φ) ) (3 cos2 φ - 1) 2
(7)
In this work, eq 6 is solved by numerical integration. In terms of the azimuthal plots in Figure 8, β ) 4.25, and a relative j 2〉 ) 0.58 is found for the 15 wt % MWCNThigher value of 〈P filled NR composite at a strain of R ) 2.65, which confirms
the high orientational degree of MWCNTs in the NR matrix during stretching. The evolution of the orientational degree of the MWCNTs with strain is shown in Figure 9. The arrows in Figure 9 indicate the values at the typical strain, whose 2D patterns are shown in Figure 7 accordingly. On the basis of this result, it can be concluded that the MWCNTs in the NR matrix favor gradual orientation with their longitudinal parallel to the drawing direction. In the plot, there is an obvious induction period in the strain range of 0-0.7. Further stretching j 2〉 value, which reaches leads to a remarkable increase of the 〈P a value of 0.72 at a strain of 3.3. Although, the MWCNT content of 15 wt % is much higher than that in CNR, the MWCNTs in j 2〉 values than in the 15 wt % CNR should have larger 〈P MWCNT-filled NR at the same strain. This is because the entanglement of MWCNTs at higher content hampers the orientation process of MWCNTs along the stretched direction
Figure 11. Enlarged images of the center of WAXD patterns from Figure 1a,b represent the sequential change of the orientation of SA crystallites j 2〉 values of pure NR and CNR are shown in panel b as a function of strain. The red solid in pure NR (a (1)) and CNR (a (2)). The variation of 〈P line is only a guide for the eye.
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Figure 12. Schematic representation of the network structural evolution of NR and CNR: (a,b) network structural evolution of NR from point A (before stretching) in panel h, which shows the stress-strain curves of NR and CNR; (e,f) the network structural evolution of CNR from B and C in panel h; (c,g) detail drawing of the large-scale structures (domain A) from panels b and f. Domain B is the matrix with low network chain density (indicated by arrows in panels a and d). (d) One example of the network structure of higher reality.
(see Figure 6). Additionally, considering the excluded volume effect from MWCNTs at higher content,55 the chain distribution of the NR is remarkably perturbed. Thus, our discussion is focused on the CNR. 4. Discussion To achieve strain-induced crystallization, oriented amorphous chains should be the precursors to the induced crystallites. However, our WAXD results described above revealed that the amorphous halo is largely isotropic during deformation, indicating a very low P2(cos φ) value. Previous WAXD research suggested that only a small amount of amorphous phase is oriented (the fraction of oriented amorphous phase is less than 10%).10,11,23,24,56 Infrared dichroism and birefringence results also showed low P2(cos φ) values.10,25,57,58 Such an unexpected phenomenon can be attributed to the heterogeneity of the crosslink distribution. Classical theory predicts the Hermans orientation function as follows:59
P2(cos φ) )
1 2 (λ - λ-1) 5N
(8)
where N indicates the number of statistical segments of a network chain, and λ is the elongation ratio. P2(cos φ) calculated by eq 8 reflects the molecular orientation of rubbers. In order to obtain the P2(cos φ) plots of both samples, n, Mc, and N were estimated on the basis of equilibrium swelling experiments. Their values are shown in Table 2. Considering the nonGaussian behavior at large deformation, the P2(cos φ) values were calculated only in the λ range of 1-3, as shown in Figure 10. The results in Figure 10 show very low P2(cos φ) values,
which are in accord with our amorphous halo analysis and the previous literature. So, all the investigations (by means of WAXD, birefringence, infrared dichroism techniques, etc.) in the literature were limited at the molecular scale. It means the orientating unit is a short section of the polymer chains corresponding to a size the order of a monomer unit (
