Microstructure and Deformation Behavior of Polyethylene

Sep 29, 2009 - Graduate School of the Chinese Academy of Sciences. ... PE matrix and MMTs, which originated from the presence of a network-like struct...
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Microstructure and Deformation Behavior of Polyethylene/Montmorillonite Nanocomposites with Strong Interfacial Interaction Changyi Ren,†,‡ Zhiyong Jiang,†,‡ Xiaohua Du,†,‡ Yongfeng Men,† and Tao Tang*,† State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Changchun 130022, China, and Graduate School of the Chinese Academy of Sciences, Beijing 100039, China ReceiVed: July 4, 2009; ReVised Manuscript ReceiVed: August 13, 2009

Deformation behavior of polyethylene/modified montmorillonites with polymerizable surfactant (PE/P-MMT) nanocomposite with strong interfacial interaction was studied by means of morphology observation and X-ray scattering measurements. The orientation of PE chains was accompanied by the orientation of well-dispersed MMT platelets due to the presence of strong interfacial interaction, and both of the orientations were parallel to the deformation direction. The high degree of orientation of MMT platelets and PE chains resulted from the synergistic movement of PE matrix and MMTs, which originated from the presence of a network-like structure. Meanwhile, the existence of MMT platelets with good mobility during deformation and strong interfacial interaction with PE matrix could further improve the break energy of material by restraining the initiation and growth of cavities during deformation. In contrast, PE/MMT nanocomposite with no strong interfacial interaction and poor dispersed state of MMT sheets showed the weaker orientation of both PE chains and MMT platelets, and a strong cavitation during deformation. 1. Introduction Polymer/clay nanocomposites (PCNs) have widespread applications in diverse areas including practical life and scientific research, due to their enhanced physical, mechanical, barrier, and flammability properties.1,2 However, persistent challenges such as poor miscibility, dispersion, and interfacial strength have prevented nanocomposites from realizing their full potential. For example, despite an increase in stiffness, most reported PCNs exhibit lower toughness than the matrix polymers.3-6 Actually, there are few cases that the toughness is improved by addition of clay, and the cases about simultaneous improvements in stiffness and toughness are even rare. From the application points of view, brittleness of the nanocomposites severely limits their use in load-bearing applications.7,8 It is, therefore, important to understand the interrelationship between miscibility, interfacial strength, dispersion, and mechanical properties of the nanocomposites in order to design composites with improvements in strength and stiffness accompanied by improvements in toughness. Unfortunately, many investigations had focused on the preparation of PCNs, and little efforts were made in the study of deformation mechanism of PCNs. In most cases, the improvements in toughness are simply attributed to homogeneous dispersion of filler.9-11 However, Gersappe has suggested, by means of molecular dynamics simulation, that the mobility of nanofillers in a polymer, rather than their surface area, controls their ability to dissipate energy, which would increase toughness of the polymer nanocomposites in the matrix with proper thermodynamic state.12 Very recently, Zhou et al. have confirmed this theory in silica nanoparticles toughening polymers.13 From the technical viewpoint, the authors suggested that * To whom correspondence should be addressed. Tel: +86-43185262004. Fax: +86-431-85262827. E-mail: [email protected]. † Chinese Academy of Sciences. ‡ Graduate School of the Chinese Academy of Sciences.

the key issues to provide the nonlayered nanoparticles with sufficient mobility lay in the high mobility of the matrix, low particle-particle attraction and strong filler-matrix interaction. Giannelis and co-workers have reported a unique class of semicrystalline and amorphous polymer/clay nanocomposites having toughness values an order of magnitude higher than that of the unfilled polymer matrix.6,14 In these cases, the nanoparticle mobility during deformation is also a key factor in improving the toughness. The above reports have an agreement that the movement and orientation of nanoparticles help to dissipate energy and increase the toughness of the composites during the deformation. However, the mobility of matrix has also been concerned, which is taken as the precondition for this mechanism to be effective.6,12-14 Haraguchi and co-workers have reported a class of PCNs exhibiting abnormal necking phenomena accompanied by extremely large reversible elongations.15-19 This phenomenon results from the formation of a unique claynetwork morphology, in which polymer chains are linked to the surface of clay platelets. It has been proved that this socalled “plane cross-linking” could lead to better mechanical properties compared with those of conventional chemical “point cross-linking”.19 However, the opinions of nanoparticle mobility and interaction on the interface are still not sufficient to paint a complete picture of the reinforcing and deformation mechanism taking place in many systems, especially in semicrystalline polymer/ clay nanocomposites. Yee and co-workers have thought that the existence of a high concentration of crazes and microcracks in PA6/clay nanocomposite is profitable to the toughness.20 The latest report from Giannelis and co-workers21 has shown that the change from brittle to ductile behavior is originated from physical cross-links introduced by the clay nanoparticles. Effective physical cross-linking depends on high mobility of nanoparticles and strong interactions between the nanoparticles and the polymer. As a result, they can move together with the polymer chains during deformation. This implies that the

10.1021/jp9063164 CCC: $40.75  2009 American Chemical Society Published on Web 09/29/2009

Polyethylene/Montmorillonite Nanocomposites mobility of nanoparticles is not the only factor influencing the toughness of PCNs. They suggested a toughening mechanism via multiple crazing and shear yielding. Unfortunately, crazing is a kind of fracture to the material, which will cause a decrease in stiffness and strength. Consequently, the toughening mechanism of crazing does not work for the systems showing improvements in both strength and toughness simultaneously. Polyethylene (PE) is a typical semicrystalline polymer, which is widely used in many applications due to its low cost and good complex performances. Accordingly, PE/clay nanocomposites with high performance are highly attractive.22-24 It is well-known that the deformation mechanism of PE is rather complicated, because PE is not a simple homogeneous system. Generally, PE is a binary composite composed of periodically stacked lamellar crystals and amorphous phase, and the morphology of crystal phase changes with the strain.25,26 It is believed that the initial tensile deformation includes isolated inter- and intralamellar slip processes after the initial purely Hookean elastic range. Then it changes into a collective activity of slip motions at the yield point. After necking process, fragmentation of lamellae proceeds with deformation, and the fibrils are formed for large strains.27-30 Cavitation is a common phenomenon observed during stretching deformation of semicrystalline polymers with high degree of crystallinity and perfection of crystals, such as high-density polyethylene.31,32 With increasing deformation the cavities change their size, number, and even orientation from perpendicular to parallel to the deformation direction.33-35 Furthermore, the formation of cavities in amorphous phase also depends on stretching rate.34 In this work, we studied the deformation mechanisms of PE/ montmorillonite (PE/MMT) nanocomposites with strong interfacial interaction. The sample was prepared by in situ polymerization of ethylene and modified MMTs with polymerizable surfactant (P-MMTs). For comparison, pure PE (HOMO) and PE/MMT nanocomposites without strong interfacial interaction were also prepared using the same method in the absence or the presence of modified MMTs with unpolymerizable surfactant (12-MMTs). The P-MMTs were well dispersed in the PE matrix than 12-MMTs, as polymerizable organophilic MMTs are more easily exfoliated than unpolymerizable organophilic MMTs in the in situ polymerization.36-41 The deforming behavior of HOMO and these two kinds of PE/clay nanocomposites with different interfacial interaction and microstructure was investigated using tensile test, X-ray scattering, rheology, and various microscopy techniques. It is expected that these will help to give a comprehensive understanding of the mechanism about simultaneous improvements of polymer/clay nanocomposites in strength and toughness. 2. Experimental Section 2.1. Materials. Two kinds of polyethylene/montmorillonite (PE/MMT) nanocomposites were synthesized by in situ ethylene polymerization using metallocene catalysts. The polymerizable surfactant, undec-10-enylammonium chloride, was synthesized to modifier Na+-MMT with a cation exchange capacity (CEC) of 119 mequiv/100 g (MMTs, from Kunimine Co.) via ion exchange in an aqueous medium. Then the resultant modified MMTs with polymerizable surfactant (P-MMTs) were introduced into the reaction system to prepare PE/MMT nanocomposites with strong interaction between the matrix and the MMT platelets, and the obtained nanocomposite was marked as P-PE. As a control, dodecylammonium chloride modified MMTs (12MMTs) were also introduced in the in situ ethylene polymerization, and the resultant sample was marked as 12-PE. Neat

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Figure 1. Sketch drawing of 2-D small-angle X-ray scattering (SAXS) measurement, the tension direction was horizontal in the measurement; signal obtained in the horizontal direction was named by “H”, signal obtained in the vertical direction was named by “V”.

PE was also synthesized for comparison, marked as HOMO. The samples used in the work, 12-PE and P-PE, had the same content of MMT (1.0 wt %). Another special PE/P-MMT nanocomposite sample with a higher content of P-MMTs (3.7 wt %) was used for 1H NMR and 13C NMR measurements to overcome the insensitiveness of the method. The results in Figure S1 in the Supporting Information justified the success of linking P-MMTs to PE backbone. 2.2. Characterizations. Wide-angle X-ray scattering (WAXS) was carried out using a Rigaku model Dmax 2500 with a Cu KR radiation and the interlayer distances (d001) of MMTs were estimated from the position of (001) peaks in the WAXS patterns according to Bragg equation. Synchrotron small-angle X-ray scattering (SAXS) was used to study the cavity formation found in the samples after tensile deformation and the state of dispersion of the clays in the filled ones. Synchrotron SAXS measurements were performed at the beamline BW4 at HASYLAB, DESY, Hamburg, Germany. The energy of the X-ray radiation was 8.979 keV, resulting in a wavelength of 0.13808 nm. The size of the primary X-ray beam at the sample position was 0.4 × 0.4 mm2. The sample to detector distance was 6432 mm. At this distance the effective scattering vector q (q ) 4π sin θ/λ, where 2θ is the scattering angle and λ the wavelength) range is 0.04-0.54 nm-1. The primary X-ray beam was positioned at the middle of the horizontally placed sample bar. The SAXS data were calibrated for background scattering and normalized with respect to the primary beam intensity. Changes in scattering intensities due to varying sample thickness have been corrected for by measuring sample absorption using ionization chambers before and after the sample and performing the respective data correction. Figure 1 shows a sketch drawing of the 2-D SAXS measurement, and rectangular slice cuts passing through the beam center were taken to obtain the onedimensional scattering intensity distributions parallel (marked by “H”) and perpendicular (marked by “V”) to the stretching direction. Azimuthal intensity distribution analysis has been made for both WAXS and SAXS to characterize the orientation degree of PE matrix and clay platelets. The orientation degree can be quantified by following the strain dependence of Hermans’ orientation parameter, S ) {(3〈cos2 Θ〉 -1)/2}, where Θ denotes the angle enclosed by the axis of interest and the unique axis. Molecular weight and molecular weight distribution of the products were measured with gel permeation chromatography (GPC) operated at 135 °C and equipped with four Waters Styragel columns (HMW2, 2 × HMW6W, HMW7) and a refractive index (RI) detector. 1,2,4-Trichlorobenzene was applied as solvent at a flow rate of 1.0 cm3/min. The columns were calibrated with polystyrene standards of narrow molar mass distribution using a universal calibration method. Before examination the samples were dissolved in 1,2,4-trichlorobenzene and then filtrated. Because of low MMTs loading in the samples, only few clay platelets would pass though the columns and the effect on GPC results can be neglected. Field emission scanning electron microscopy (FESEM) and energy-dispersive

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TABLE 1: GPC Results and Crystalline Properties of All the Samples samples

MMTs contenta (wt %)

Mwb ( × 10-5)

PDI

Tm (°C)

∆Hm (J/g)

Xc (%)

dac (nm)

d (nm)

HOMO 12-PE P-PE

null 1.0 1.0

2.27 2.03 1.26

2.7 2.4 2.5

131.6 132.9 133.9

197.3 181.3 171.3

67 62 59

24.5 24.0 26.6

16.5 14.8 15.6

a

Measured by TGA. b Before GPC measurements the P-PE samples were pretreated with concentrated HF to remove MMT platelets.

Figure 2. Polarized light microscope images of HOMO (a, b, c), 12-PE (d, e, f) and P-PE (g, h, i) measured in the melting process; the magnification of all images was the same; the heating rate was 10 °C/min.

X-ray (EDX) analysis were carried out on a XL 30 ESEM FEG scanning electron microscope (Micro Fei Philips, Holland). Samples were coated with an ultrathin film of gold to make them conductive before analysis. A JEOL JEM-2010(HR) 200 kV transmission electron microscope (TEM) was used to investigate the morphology of the clays and the phase distribution in PE/MMTs nanocomposites. Samples were prepared by ultramicrotome at -100 °C. Rheological measurements were performed on a PHYSICA MCR 300 at 190 °C, and the complex viscosities were measured as a function of angular frequency (ranging from 0.01 to 100 rad/s). The static mechanical properties were measured with Instron 1121 tensile testing machine, and the crosshead speed was set at 20 mm/min. The samples were compression molded at 170 °C for 10 min, and the mechanical measurements were performed after 24 h. The data obtained showed good repeatability. Measurements both at room temperature and 120 °C were carried out for different needs. Thermal gravimetric analysis (TGA) was carried out on a Perkin-Elmer TGA-7 series thermal analysis system at a heating rate of 10 °C/min from 50 to 700 °C under nitrogen gas. A DSC 2920 (TA Instruments) was used during the experiments with a heating rate of 10 °C/min. Before DSC measurements, the samples were maintained at 170 °C for 10 min then down to 20 °C at a rate of 10 °C/min, the melting data were taken from the following heating run. The melting point (Tm) denotes the maximum of the thermograms during the second heating. Polarized light microscope was performed on a Carl ZeissA1 microscope equipped with an Infinity 4-11 digital camera from Lumenera Co., Canada. The heating rate was 10 °C/min. Before measurement, the samples were hold at

170 °C for 10 min, and then cooled down to room temperature at a rate of 10 °C/min to perform measurements. 3. Results and Discussion 3.1. Microstructure and Crystalline Properties of PE/ MMTs Nanocomposites. Table 1 shows molecular parameters and crystalline properties of three samples. The content of MMT is 1 wt % in both 12-PE and P-PE. The molecular weight (MW) of PE decreases in the order of HOMO, 12-PE, and P-PE. The decrease of MW of P-PE is mainly caused by the copolymerization of ethylene with other copolymerizable monomers.42,43 The little decrease of 12-PE may be caused by the presence of 12-MMTs. The melting point (Tm) observed in DSC measurements of 12-PE and P-PE is higher than that of HOMO, especially in the case of P-PE. Figure 2 presents polarized light microscope photographs showing the morphological evolution of the samples as a function of temperature. The melting ranges of HOMO, 12-PE, and P-PE are located at 130.0-134.0, 135.2-137.3, and 136.4-138.5 °C, respectively. These results are slightly different from those evaluated from DSC measurements, which might originate from the experimental error between these two techniques, but the sequence of Tm of these samples is the same in the two measurements. There are tiny crystallite structures instead of spherulite structures in all the samples. Previous research has justified that the dispersed platelets of clay affect the microstructural features of semicrystalline polymer matrix, such as crystal structure, crystallinity, lamellae orientation, and spherulite formation.44 The change trend of the crystallinity of PE is contrary to that of the Tm; that is, HOMO > 12-PE > P-PE, due to the confined crystal-

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lization of PE in the PE/MMTs composites.45,46 The crystallinity of nanocomposites usually decreases with the increase of MMT content, the MMT sheets dispersed in the PE matrix confined the PE chains and hinder the crystallization of PE chains. SAXS was used to further measure the long spacing dac, which is defined as the average thickness of lamella together with one interlamellar amorphous layer, measured along the lamella normal and calculated using the Bragg eq 1. The average thickness of lamella d is calculated using eq 2.

dac )

2π qmax

d ) dac × Xc

(1)

(2)

Here qmax represents the position of the intensity maximum in a SAXS pattern; Xc represents the crystallinity of the samples, which is derived from a ratio of melting enthalpy of the sample to that of hypothetical 100% crystalline PE (293 J/g).47 The average thickness of crystalline lamella does not change significantly (Table 1). Detailed morphologies of P-PE and 12PE were investigated by TEM. Before TEM observation, the samples were stained in the vapor of ruthenium tetroxide for 2 h at 30 °C.46 It has been reported that crystalline lamellae are preferential in the proximity of the clay platelets, which is the consequence of “nucleation”.44,49-52 The TEM images in Figure 3 show a good agreement with the conclusion above. For both P-PE (mainly exfoliated structure in Figure 3a) and 12-PE (intercalated structure in Figure 3b), MMT platelets are directly surrounded by PE lamellae, and the PE lamellae are always approximately perpendicular to MMT surface. Combining this result, the higher melting temperatures of 12-PE and P-PE than HOMO should be attributed to restricting effect of the MMTs to the crystallites of PE as there is no significant difference in the thickness of crystalline lamellae among the samples. Furthermore the melting temperature of P-PE is higher than that of 12-PE, which may result from the well dispersed state of P-MMTs and the strong interaction between PE chains and MMT platelets due to the formation of chemical linking in the P-PE. This phenomenon also confirms that P-MMTs have strong interaction with PE crystallites. Rheological analysis was used to measure the interaction between PE and MMTs in the PE/MMT nanocomposites. Generally the linear viscoelastic response of a nanocomposite melt is strongly influenced by the factors, such as dispersion state of particle, interparticle and particle-polymer interactions. Figure 4a shows the plots of storage modulus G′, loss modulus G′′, and complex viscosity |η*| versus frequency ω for the PE/ MMT nanocomposites. 12-PE shows classical viscoelastic behavior characterized by transition from low-frequency Newtonian flow behavior to high-frequency shear thinning nature. In the case of P-PE, although the MW of PE in the P-PE is lower than that of 12-PE, the presence of P-MMTs significantly changes this behavior, especially in the low-frequency region. It shows a strong tendency of storage modulus (G′) to plateau and a significant viscosity (|η*|) upturn at the low-frequency regime, which is widely accepted as the evidence of a networklike structure.53-58 The network-like structure results from ethylene copolymerization with P-MMTs, in which the welldispersed P-MMT platelets serve as cross-linking sites through chemical bonding. This solidlike response has also been proven to be an evidence of improved compatibility and a better interfacial interaction between filler and matrix.59 A recent report

Figure 3. TEM images of P-PE (a) and 12-PE (b); the insets were the selected-area electron diffraction patterns from the marked areas.

has shown that the principal mechanism responsible for the changes in viscoelastic behavior in polymer nanocomposites is through a polymer-mediated filler network and not necessarily good dispersion of the nanoparticles.60 The 12-PE sample does not show the same behavior due to the absence of the similar microstructure. Figure 4b gives the plots of log G′ versus log G′′. The curve for 12-PE has a slope of approximately 2 in the terminal region, as is expected for all ordinary linear polymers.61,62 This also means that there is no strong interfacial interaction between polymer matrix and MMT platelets.59 In contrast, the plot for P-PE has a very small slope, exhibiting solidlike behavior. This phenomenon is proven to be evidence for the existence of strong interaction on the interface. 3.2. Stress-Strain Behavior of PE/MMTs Nanocomposites. Figure 5 shows the stress-strain curves of neat PE (HOMO), 12-PE, and P-PE respectively. Generally, the MW has great impact on the mechanical properties of materials;63 crystallinity and thickness of lamellae can influence the cavitation of samples.32 It is universally accepted that the higher the

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Figure 4. (a) Plots of log G′ (storage modulus), G′′ (loss modulus), and |η*| (complex viscosity) vs log ω (frequency) for 12-PE (9, b, 2) and P-PE (0, O, 4) at 190 °C; (b) plots of log G′ vs log G′′ for 12-PE ([) and P-PE (]) at 190 °C.

Figure 5. Stress-strain curves for pure PE (HOMO), PE/P-MMT nanocomposite (P-PE), and PE/12-MMT composite (12-PE).

MW and crystallinity of polymer are, the stronger the material is. However, in spite of a decrease in the MW and crystallinity compared with that of HOMO (Table 1), P-PE shows improved tensile modulus, yield strength, and tensile strength simultaneously, indicating that such a difference in the MW or crystallinity is not a dominating factor in the deformation of the P-PE compared to HOMO. According to the inset of Figure 5, 12-PE has the highest yield strength compared with those of HOMO and P-PE. Usually filled polymers show higher tensile modulus and higher yield strength, especially for MMTs filled

Ren et al. polymers. The tensile modulus and yield strength of 12-PE are higher than those of P-PE, which might be due to a higher value of crystallinity for 12-PE as compared to P-PE. A significant increase in the elongation at break is observed for P-PE (∼1800%) sample compared with those of 12-PE (∼1000%) and HOMO (∼1100%). In one word, compared to HOMO sample, the P-PE sample shows improvements in the strength and the toughness simultaneously, and its load after yield is still higher than those of 12-PE and HOMO at any strain. Interestingly, there is a stronger strain hardening in the stress-strain curve of P-PE sample compared to HOMO sample. In contrast, the 12-PE sample does not show such a behavior. It is believed that the strong interaction between PE matrix and P-MMT makes stress transferred efficiently from the polymer to the filler and reduces chain slippage. Similar phenomena are observed in other PE/P-MMT samples containing 0.4 - 3.7 wt % P-MMTs. But the extent of improvements in stress and elongation at break varies with P-MMT content, e.g., the elongation at break of the sample containing P-MMT 0.4 wt % approaches 2000%. In this work, two typical samples P-PE (1.0 wt %) and 12-PE (1.0 wt %) with different interfacial interaction are chosen to demonstrate the evolution of the microstructure and the deformation behavior of PE/MMT nanocomposites. 3.3. Orientation Behavior of PE/MMTs Nanocomposites during Deformation. As a typical of semicrystalline polymer, the stress-strain behavior of PE can be influenced by several factors, such as orientation of PE chains and cavitation during stretching. In the case of PE/MMTs nanocomposites, the presence of MMT platelets will introduce other factors into deformation process, such as possible orientation of MMT platelets. When analyzing the deformation process of PE/MMT nanocomposites, it is an important issue how the MMTs affect the PE matrix besides how the MMTs move in the matrix. Joly et al. had found a higher orientation of amorphous chains induced by the addition of nanoclays in organoclay/natural rubber nanocomposites.64 Owing to low content of MMTs in PE/MMT nanocomposites in this work (1.0 wt %), the elongation of the nanocomposites should be the deformation of PE matrix indeed which serves as the continuous phase of the material bearing the loading during drawing. The toughness of the nanocomposites is tightly related to chain mobility of polymer matrix. There are some reports relating the improvement of toughness to higher degree of the matrix orientation, which is accepted as an index of the mobility of the matrix.6,13 Figure 6 presents the azimuthal intensity distributions of (110) reflections obtained by means of wide-angle X-ray scattering (WAXS) experiments for the necked parts of HOMO, 12-PE, and P-PE samples at the elongation of 300%. By measuring the azimuthal intensity distribution of the amorphous halo during deformation, one can determine the amorphous chain orientation induced in the stretching network. Furthermore, it has been found that the orientation of amorphous segments of PE under uniaxial drawing is comparable to the degree of orientation of the crystallites.26 Obviously, the degree of orientation of 12-PE is smaller than those of P-PE and HOMO, and the P-PE sample shows the highest degree of orientation during deformation (Figure 6), which could be attributed to the presence of strong interfacial interaction between PE and MMT platelets in the case of P-PE. A higher degree of orientation means a higher stress during deformation. This is one of the reasons for a higher strength of P-PE than HOMO and 12-PE at the same strain after yield. During the elongation, especially at the necking stage, the MMT platelets change their random alignment into oriented

Polyethylene/Montmorillonite Nanocomposites

Figure 6. Azimuthal intensity distributions of the (110) reflections of deformed HOMO, 12-PE, and P-PE and their corresponding WAXS patterns; deformation direction was horizontal, tensile elongation was 300%; the measurements were performed on the necked parts of the samples.

alignment parallel to the stretching direction with the movement of matrix.14,65 Meanwhile, the polymer chains are elongated parallel to the stretching direction too.66-68 Since the movement of MMTs is driven by the movement of PE matrix, the orientation of MMTs should be accompanied by the orientation of PE chains (including PE lamellae) simultaneously. It is impossible for MMTs to move and rearrange by themselves in PE matrix. Figure 7 shows the WAXS patterns and their corresponding TEM images performed on different positions of deformed 12-PE and P-PE at the necking stage. Here the spots on necked or unnecked part of the samples, where WAXS measurements were carried out, are randomly selected. WAXS patterns at the un-necked regions of both 12-PE and P-PE show almost no orientation of the polymer chains (Figure 7, a and e). TEM images at these regions show no orientation of MMT platelets too (Figure 7, b and f). These images also confirm the dispersion states of both 12-PE (intercalated structure) and P-PE (mainly exfoliated structure). At the necked region, the orientations of both polymer chains and MMT platelets are observed (Figure 7, a, d, g, and h). This indicates that the orientation takes place during the necking process. The degree of orientation of MMT platelets can be tested by means of SAXS. Generally, the formation of cavities within the samples produces much stronger scattering intensity than that of clay platelets, making the latter nearly invisible in the SAXS experiments under current condition. To avoid cavitation within deformed samples, here the samples drawn at 120 °C are used. The SAXS results and the corresponding TEM images of PE/MMT nanocomposites are summarized in Figure 8. It can be seen from the TEM images (inset in Figure 8b), there are no cavities when the stretch is carried out at 120 °C. According to the SAXS patterns in Figure 8a, two kinds of orientation can be observed. One is the orientation of PE lamellae, corresponding to the horizontal direction of the anisotropic SAXS pattern; the other is the orientation of MMT platelets, corresponding to the SAXS pattern along the vertical direction. The intensity where q ranges from 0.1 to 0.4 nm-1 is due to the scattering intensity of oriented PE lamellae. The degree of orientation for MMT platelets can be obtained by measuring the azimuthal intensity distribution of the SAXS intensity where q ranges from 0.04 to 0.1 nm-1 so as to avoid the disturbance of scattering intensity of PE lamellae by the greatest extent. In Figure 8b, the plot of HOMO shows an orientation along the horizontal direction due to the orientation

J. Phys. Chem. B, Vol. 113, No. 43, 2009 14123 of crystalline lamellae. Comparing the plot of P-PE with that of 12-PE, it is obvious that the degree of orientation for P-MMTs in P-PE is higher than that of 12-MMTs in 12-PE. In other words, the P-MMTs in P-PE would experience a higher degree of movements than MMTs in 12-MMTs, which would dissipate more energy. As PE matrix in the nanocomposites is the continuous phase bearing all the loading of the material, the additional energy dissipation from the movements of MMTs would improve the strength of the matrix. This is another reason for higher strength of P-PE than that of 12-PE during deformation.69 It should be noted that the extent of this improvement in strength caused by the orientation of MMTs would be limited, because the content of MMTs in the nanocomposites is very low. 3.4. Cavitation Behavior of PE/MMTs Nanocomposites during Deformation. As cavities can weaken the strength of the material, cavitation behavior is possibly another factor for determining the tensile strength during deformation. Figure 9 shows the TEM images of 12-PE and P-PE at the necked regions after stretching at a strain of 300% at room temperature. Significant difference can be observed from these images. 12PE shows a higher concentration of cavities than P-PE. Clearly, the cavities are initiated in the interfacial region between MMTs and PE matrix. FESEM analysis on the elongated samples provides a similar result (Figure 10). Fractures and cavities are more severe for 12-PE than those of HOMO and P-PE at the same elongation (300%). As MMTs are not exfoliated in 12PE, some tactoids of MMT platelets can be seen easily in the cavities of 12-PE (Figure 10b). This also indicates that the initiation of cavities is preferred at the interface between MMTs and PE matrix. By contrast, the necked region of elongated P-PE is smoother than that of HOMO (Figure 10, part c vs part a). SAXS is usually employed to quantitatively analyze the cavity formation, as revealed by a rapid increase of scattering intensity. Because of the relationship between cavity concentration and scattering intensity, the integral of scattering intensity can be accepted as an index of cavity concentration when other contribution to the intensity remains the same.34 Generally, the cavities show their orientation during deformation. Before necking, the cavities are extended perpendicular to the tensile direction; after necking, they are elongated parallel to the tensile direction.33,34 It has been shown that lamella fragmentation and cavitation during tensile stretching are often detected simultaneously.27,70,71 As the orientation of polymer chains after necking is parallel to the tensile direction, the orientation of crystalline lamellae is perpendicular to the tensile direction. Figure 11 presents the SAXS patterns and SAXS intensity distributions of HOMO, P-PE, and 12-PE. The SAXS patterns clearly show an equatorial streak (Figure 11a), which is originated from the cavities elongated parallel to the stretching direction. The intensity of 12-PE at all measured range is the strongest among three samples, followed by HOMO and P-PE (Figure 11b). This result agrees with the morphological observations (Figure 10). This phenomenon indicates that there are more cavities occurring during the deformation of 12-PE. In contrary, the cavity concentration of P-PE is the lowest; as a result, this sample shows the highest strength during deformation. Therefore, it is confirmed that the difference in cavity concentration is one of the reasons for the change in strength during deformation. Compared to 12-PE, MMT platelets are well dispersed in P-PE due to strong interfacial interaction between PE matrix and MMT surface. The unique microstructure has brought special deformation behavior to the PE/P-MMT nanocomposites. For semicrystalline polymers, it is well accepted that the deformation after the initial purely Hookean elastic range is

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Figure 7. WAXS patterns (a, c, e, g) and their corresponding TEM images (b, d, f, h) performed on different part of deformed 12-PE and P-PE at necking stage.

Figure 8. 2-D SAXS patterns (a) and their azimuthal intensity distributions of the scattering intensity (q ranges from 0.04 to 0.1 nm-1) with their corresponding TEM images (b) of deformed HOMO, P-PE, and 12-PE at 120 °C.

generally a macroscopically homogeneous deformation involving lamellae tilt, slip, and some breakup.72 And it was found that the MMT platelets were directly surrounded by PE lamellae

(Figure 3). As a result, it is vital to know how the clay interacts with the crystalline lamellae of PE. The different interfacial interaction between filler and matrix can cause the difference in the mobility of matrix during deformation. An augment in the orientational order parameter has been associated with the increase in the cross-linking or entanglement density, and the strong interfacial interaction makes the introduced P-MMTs become additional cross-linking and entanglements.62,73,74 In PE/ P-MMT nanocomposites, the role of P-MMT is just like crosslinking sites in cured rubber. The well-dispersed state of P-MMTs and the strong interaction between matrix and PMMTs cause the synergistic movement of PE matrix and P-MMTs, which provides an improved mobility of the whole material during deformation.75 The synergistic movement of two components makes PE chains drawn more effectively compared to HOMO sample and 12-PE sample. The higher degree of orientation of both PE matrix and MMTs for P-PE than that of 12-PE is an indication of such a synergistic movement. In the case of aligned clays there are instead shear stresses as opposed to tensile stresses between the polymer and clays, allowing for more effective stress transfer from the polymer to the filler.76 The elongation of the nanocomposites is the deformation of the matrix indeed, and a good mobility of the matrix during deformation would bring high toughness to the material. This is also the reason that the P-PE sample has a much larger elongation at break than those of HOMO and 12-PE samples. In addition, according to the TEM, FESEM, and SAXS results shown in Figures 8-10, there is significant difference in the cavitations of HOMO, 12-PE, and P-PE. Figure 12 gives the sketch drawing of the morphological evolution during elongation of P-PE and 12-PE. For HOMO, cavities emerge shortly before yielding in the crystalline region, more accurately, the interlamellar amorphous phase in the crystalline region.32,34 For the case of 12-PE, there is no strong interaction between MMT

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Figure 9. TEM images of 12-PE (a) and P-PE (b); the plan view was parallel to the tensile direction (TD).

Figure 11. 2-D SAXS patterns (a) and SAXS intensity distributions taken along the vertical direction (b) of deformed HOMO, 12-PE, and P-PE at tensile elongation of 300%; the elongation was carried out at room temperature; the detected position was at the neck region of the sample; the thickness of the samples was the same, 1.5 mm.

Figure 10. FESEM images of deformed HOMO (a), 12-PE (b), and P-PE (c) on the necked part at 300% elongation.

platelets and crystalline lamellae. Cavities emerge at the interface between the clay platelets and PE matrix, meaning that the weakest part of the composite is at the interface between the clay platelets and PE matrix. In addition, the mobility of MMT platelets is another factor that should be taken into account.6,12-14 As seen in TEM images of 12-PE (Figure 7b), the MMTs of 12-PE are not well exfoliated. Therefore, the aggregates of MMTs that are not easily able to align in the stretching direction would cause concentrated tensile stresses between PE matrix

and MMTs, which would also induce the formation of cavities. These also provide an explanation to the phenomenon observed in Figures 9a and 10b that cavities usually initiate at the interface between clay and matrix. Here, MMTs serve as the initiator of cavitation to form additional cavities, which is the reason that 12-PE has a higher level of cavity concentration than HOMO. For P-PE, the presence of P-MMTs can also restrain the cavitation of PE/P-MMT nanocomposites. The inhibiting effect can apply in the cavity initiation and the cavity growth. The PE chains are directly linked to the surface of MMT platelets through copolymerization with polymerizable modifiers in PE/ P-MMT nanocomposites. The level of this interaction is stronger than the finite entanglements in the interlamellar amorphous

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Figure 12. Schematic representation of the morphology evolution for P-PE (chemical linking between MMT and PE backbone) and 12-PE (without chemical linking between MMT and PE backbone) induced by tensile deformation at room temperature.

phase. So the interaction between lamellae is enhanced through P-MMTs, and some of the potential cavitations are inhibited by the interaction between crystalline lamellae and clay platelets. When a growing cavity meets P-MMT, the growth of such cavity is terminated. In this case, it is difficult to exfoliate the lamellae away from the P-MMT platelets as P-MMT platelets tie lamellae around them together by strong interaction. Cavities cause the reduction in stress during elongation, and the growth of cavities finally leads to rupture of the material. So the inhibiting effect of cavities helps to improve the strength of P-PE compared with other samples. 4. Conclusion The PE/MMT nanocomposites (P-PE) with strong interfacial interaction between PE matrix and MMT platelets have shown completely different deformation behavior during stretching compared with neat PE and PE/MMT nanocomposites without strong interfacial interaction. The existence of the strong interfacial interaction enabled the PE matrix and P-MMTs to form a network-like structure. As a result, the mobility of MMT platelets is improved during the deformation of PE matrix. In this case, the synergistic movement of PE matrix and P-MMTs can be realized easily, which results in increased degree of orientations of PE chains and MMT platelets during tensile deformation. Therefore, more energy would be dissipated during elongation compared to neat PE and PE/MMT nanocomposite with no strong interfacial interaction, which means improved fracture energy. Strong interfacial interaction between two components is also the reason of restraining the cavitation. It is observed that the formation of cavities is obviously inhibited in the case of P-PE, which results in a further increase of the stress during deformation. Therefore, the orientation of polymer chains as well as MMTs and the decrease of intensity of cavitation have contributed to the improvement in fracture energy. Further research should be done to investigate whether the same mechanism could work for other kinds of polymer nanocomposites, especially nanocomposites with brittle polymer matrix; such investigations would help us to understand the truth of deformation behavior of polymer/clay nanocomposite and

design new nanocomposites having stress and toughness improved synchronously. Acknowledgment. This work is supported by the National Natural Science Foundation of China for the Outstanding Youth Fund (50525311), the Fund for Creative Research Groups (50621302) and (20734006, 50873099) Programs, and HASYLAB project II-20052011. The authors thank Stephan V. Roth, Rainer Gehrke, R. Do¨hrmann, and M. Dommach for assistance in SAXS experiments at HASYLAB. Supporting Information Available: 1H NMR and 13C NMR spectra of treated P-PE sample; DSC melting curves (the second run) of undeformed samples; and SAXS patterns of undeformed samples. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Ray, S. S.; Okamoto, M. Prog. Polym. Sci. 2003, 28, 1539. (2) Usuki, A.; Hasegawa, N.; Kato, M. AdV. Polym. Sci. 2005, 179, 135. (3) Liu, L. M.; Qia, F.; Zhu, X. G.; Qi, Z. G.; Chen, G. Q. Acta Polym. Sin. 1998, 3, 304. (4) Kojima, Y.; Usuki, A.; Kawasumi, M.; Okada, A.; Fukushima, Y.; Kurauchi, T.; Kamigaito, O. J. Mater. Res. 1993, 8, 1185. (5) Ranade, A.; D’Souza, N. A.; Gnade, B.; Dharia, A. J. Plast. Film Sheeting 2003, 19, 271. (6) Shah, D.; Maiti, P.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. AdV. Mater. 2005, 17, 525. (7) Dasari, A.; Yu, Z. Z.; Mai, Y. W.; Hu, G. H.; Varlet, J. L. Compos. Sci. Technol. 2005, 65, 2314. (8) Garcia, M.; de Rooij, M.; Winnubst, L.; van Zyl, W. E.; Verweij, H. J. Appl. Polym. Sci. 2004, 92, 1855. (9) Lim, S. K.; Kim, J. W.; Chin, I.; Kwon, Y. K.; Choi, H. J. Chem. Mater. 2002, 14, 1989. (10) Liao, H. T.; Wu, C. S. J. Appl. Polym. Sci. 2005, 97, 397. (11) Gai, J.; Li, H. J. Appl. Polym. Sci. 2007, 106, 3032. (12) Gersappe, D. Phys. ReV. Lett. 2002, 89, 058301. (13) Zhou, T. H.; Ruan, W. H.; Rong, M. Z.; Zhang, M. Q.; Mai, Y. L. AdV. Mater. 2007, 19, 2667. (14) Shah, D.; Maiti, P.; Gunn, E.; Schmidt, D. F.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. AdV. Mater. 2004, 16, 1173. (15) Haraguchi, K.; Ebato, M.; Takehisa, T. AdV. Mater. 2006, 18, 2250. (16) Haraguchi, K.; Takehisa, T.; Fan, S. Macromolecles 2002, 35, 10162.

Polyethylene/Montmorillonite Nanocomposites (17) Haraguchi, K.; Farnworth, R.; Ohbayashi, A.; Takehisa, T. Macromolecules 2003, 36, 5732. (18) Haraguchi, K.; Li, H. J. Macromolecules 2006, 39, 1898. (19) Shibayama, M.; Karino, T.; Miyazaki, S.; Okabe, S.; Takehisa, T.; Haraguchi, K. Macromoleclues 2005, 38, 10772. (20) He, C.; Liu, T.; Tjiu, W. C.; Sue, H.-J.; Yee, A. Macromolecules 2008, 41, 193. (21) Xu, W.; Raychowdhury, S.; Jiang, D. D.; Retsos, H.; Giannelis, E. P. Small 2008, 4, 662. (22) Brintzinger, H. H.; Fischer, D.; Mu¨lhaupt, R.; Rieger, B.; Waymouth, R. M. Angew. Chem. 1995, 107, 1255. (23) Alexandre, M.; Martin, E.; Dubois, P.; Garcia-Marti, M. Macromol. Rapid Commun. 2000, 21, 931. (24) He, A. H.; Hu, H. Q.; Huang, Y. J.; Dong, J. Y.; Han, C. C. Macromol. Rapid Commun. 2004, 25, 2008. (25) Hiss, R.; Hobeika, S.; Lynn, C.; Strobl, G. Macromolecules 1999, 32, 4390. (26) Hong, K.; Strobl, G. Macromolecules 2006, 39, 268. (27) Men, Y. F.; Strobl, G. Chin. J. Polym. Sci. 2002, 20, 161. (28) Men, Y. F.; Rieger, J.; Strobl, G. Phys. ReV. Lett. 2003, 91, 095502. (29) Butler, M. F.; Donald, A. M.; Ryan, A. J. Polymer 1998, 39, 39. (30) Wang, K. H.; Chung, I. J.; Jang, M. C.; Keum, J. K.; Song, H. H. Macromolecules 2002, 35, 5529. (31) Pawlak, A.; Galeski, A. Macromolecules 2005, 38, 9688. (32) Pawlak, A. Polymer 2007, 48, 1397. (33) Men, Y. F.; Rieger, J.; Homeyer, J. Macromolecules 2004, 37, 9481. (34) Pawlak, A.; Galeski, A. Macromolecules 2008, 41, 2839. (35) Zhang, X. C.; Butler, M. F.; Cameron, R. E. Polymer 2000, 41, 3797. (36) Owusu-Adom, K.; Guymon, C. A. Macromolecules 2009, 42, 180. (37) Zeng, C.; Lee, L. J. Macromolecules 2001, 34, 4098. (38) Bottino, F. A.; Fabbri, E.; Fragala`, I. L.; Malandrino, G.; Orestano, A.; Pilati, F.; Pollicino, A. Macromol. Rapid Commun. 2003, 24, 1079. (39) Fu, X.; Qutubuddin, S. Polymer 2001, 42, 807. (40) Zhong, Y.; Zhu, Z. Y.; Wang, S. Q. Polymer 2005, 46, 3006. (41) Samakande, A.; Hartmann, P. C.; Cloete, V.; Sanderson, R. D. Polymer 2007, 48, 1499. (42) Grieken, R.; Martı´n, C.; Moreno, J.; Prieto, O.; Bravo, J. M. Macromol. Symp. 2007, 259, 174. (43) Hammawa, H.; Wanke, S. E. Polym. Int. 2006, 55, 426. (44) Kim, G. M.; Lee, D. H.; Hoffmann, B.; Kressler, J.; Sto¨ppelmann, G. Polymer 2001, 42, 1095. (45) Ma, J. S.; Zhang, S. M.; Qi, Z. N.; Li, G.; Hu, Y. L. J. Appl. Polym. Sci. 2002, 83, 1978. (46) Xu, J. T.; Zhao, Y. Q.; Wang, Q.; Fan, Z. Q. Polymer 2005, 46, 11978. (47) Wunderlich, B. Macromolecular Physics; Academic Press: New York, 1973; Vol. 1. (48) Brown, G. M.; Butler, J. H. Polymer 1997, 38, 3937.

J. Phys. Chem. B, Vol. 113, No. 43, 2009 14127 (49) Lincoln, D. M.; Vaia, R. A.; Wang, Z. G.; Hsiao, B. S.; Krishnamoorti, R. Polymer 2001, 42, 9975. (50) Maiti, P.; Okamoto, M. Macromol. Mater. Eng. 2003, 288, 440. (51) Weon, L.; Sue, H. J. J. Polym. Sci., Polym. Phys. Ed. 2005, 43, 3555. (52) Jimenez, G.; Ogata, N.; Kawai, H.; Ogihara, T. J. Appl. Polym. Sci. 1997, 64, 2211. (53) Goharpey, F.; Nazockdast, H.; Katbab, A. A. Polym. Eng. Sci. 2005, 45, 84. (54) Ye, Z.; AlObaidi, F.; Zhu, S. Ind. Eng. Chem. Res. 2004, 43, 2860. (55) Tang, C. Y.; Xiang, L. X.; Su, J. X.; Wang, K.; Yang, C. Y.; Zhang, Q.; Fu, Q. J. Phys. Chem. B 2008, 112, 3876. (56) Kashiwagi, T.; Du, F. M.; Douglas, J. F.; Winey, K. I.; Harris, R. H.; Shields, J. R. Nat. Mater. 2005, 4, 928. (57) Kharchenko, S. B.; Douglas, J. F.; Obrzut, J.; Grulke, E. A.; Milger, K. B. Nat. Mater. 2004, 3, 564. (58) Du, F.; Scogna, R. C.; Zhou, W.; Brand, S.; Fischer, J. E.; Winey, K. I. Macromolecules 2004, 37, 9048. (59) Lee, K. M.; Han, C. D. Macromolecules 2003, 36, 7165. (60) Xu, L.; Nakajima, H.; Manias, E.; Krishnamoorti, R. Macromolecules 2009, 42, 3795. (61) Han, C. D.; Jhon, M. S. J. Appl. Polym. Sci. 1986, 32, 3809. (62) Han, C. D.; Kim, J. K. Macromolecules 1989, 22, 4292. (63) Fu, Q.; Men, Y.; Strobl, G. Polymer 2003, 44, 1941. (64) Joly, S.; Garnaud, G.; Ollitrault, R.; Bokobza, L.; Mark, J. E. Chem. Mater. 2002, 14, 4202. (65) Carretero-Gonzalez, J.; Verdejo, R.; Toki, S.; Hsiao, B. S.; Giannelis, E. P.; Lo´pez-Manchado, M. A. Macromolecules 2008, 41, 2295. (66) Peterlin, A. J. Mater. Sci. 1971, 6, 490. (67) Peterlin, A. Colloid Polym. Sci. 1975, 253, 809. (68) Jiang, Z.; Tang, Y.; Men, Y.; Enderle, H. F.; Lilge, D.; Roth, S. V.; Gehrke, R.; Rieger, J. Macromolecules 2007, 40, 7263. (69) Zhang, Y.; Yang, J. H.; Ellis, T. S.; Shi, J. J. Appl. Polym. Sci. 2006, 100, 4782. (70) Plummer, C. J. G.; Scaramuzzino, P.; Kausch, H. H.; Phillipoz, J. M. Polym. Eng. Sci. 2000, 40, 1306. (71) Plummer, C. J. G.; Beguelin, P.; Grein, C.; Gensler, R.; Dupuits, L.; Gaillard, C.; Stadelmann, P.; Kausch, H. H.; Manson, J. A. E. Macromol. Symp. 2004, 214, 97. (72) Friedrich, K. In AdVances in Polymer Sciences; Kausch H. H., Ed.; Springer: Berlin, 1983; Vol. 52/53; pp 225-274. (73) Bokobza, L. Polymer 2001, 42, 5415. (74) Bokobza, L. Macromol. Symp. 2001, 169, 243. (75) Wang, K. H.; Chung, I. J.; Jang, M. C.; Keum, J. K.; Song, H. H. Macromolecules 2002, 35, 5529. (76) Fornes, T. D.; Paul, D. R. Polymer 2003, 44, 4993.

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