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Orientation, mechanical, and optical properties of poly (lactic acid) nanoclay composite films.


It is well established that the effective dispersion of anisotropic particles with high aspect ratios such as short fibers, plates, and whiskers within a continuous polymer matrix, in combination with adequate interfacial adhesion between the filler and polymer, can account for substantially improved reinforcement of the polymer matrix [l]. Layered-silica-based polymer nanocomposites have attracted considerable technological and scientific interest in recent years (2-5), because they have shown dramatic enhancements in the physical, thermal, mechanical, and barrier properties of polymers even with relatively low loading of silicate (5-7). These properties are attributed to the layered silicate and confinement of the polymeric matrix at the nanoscale (8). If the silicate tactoids delaminated completely, an exfoliated nanocomposite is obtained. By contrast, if only an increase in the interlayer distance is obtained, the nanocomposite is called intercalated (9). An increased exfoliation of the tactoids results in the formation of a large interface between the soft polymer phase and hard nanoclay platelets. Consequently, the fraction of these rigid particles that is needed to achieve the same effect in properties decreases as exfoliation increases (10-12). Pioneering advances at Toyota during the early 1990s stimulated the development of various polymer/orga-nosilicate nanocomposites with improved property profiles (13). Many approaches have been reported to produce orga-noclay/polyolefin nanocomposites including in-situ polymerization of monomer in the presence of silicate layers, solvent mixing, and melt compounding (14). Although, in-situ polymerization proved to be more effective to form exfoliated nanocomposites than melt compounding, however, the exfoliated clay platelets to create stacked structures on subsequent melt processing (14), (15). Therefore, compared with in-situ polymerization and solvent mixing, melt compounding is the most convenient and cost effective technique to prepare polyolefin nanocomposites (14). The use of organoclays has been extended into different polymer systems including epoxy, polyurethane, polyimides, polyesters, polypropylene, polystyrene, etc (16).

According to Kaynak and Tasan (17), various parameters affect the development of nanocomposites; the method for production of nanocomposites, the type of resin, the type and content of nanoclay, and the chemical modifications used in producing the nanoclay. It has been already elucidated that exfoliated structures could be prefrably obtained in polar matrices (e.g., PA and PEO) because of the favorable interactions between the polymer chains and the organo-modified clays (18), (19). Zhong and De Kee (14) studied structure of blown films of ethylene vinyl acetate (EVA), low-density polyethylene (LDPE), and high-density polyethylene (HDPE) melt compounded with an organically modified montmorillonite. The morphology analysis showed that all three types of films involve intercalated clay particles while the intercalation extent depended on the matrix type as well as on the molecular weight of the compatibilizers. Wang et al. (20) investigated the dispersion and orientation of compatibilized injection molded isotactic polypropylene (iPP)/organoclay nanocomposites. A much higher orientation of PP was found in the composites compared with the pristine PP. This was interpreted as due to an enhancement in the local stress that occurred in the interparticles region of two-layered tactoids (platelets) with different velocities. Pereira de Abreu et al. (8) showed that exfoliation of Cloisite 15A in the polyolefin films can be achieved by optimizing the processing parameters including temperature, processing time, and feed position. The results of their study revealed that the feed position of the nanoparticles in a twin screw extruder is of vital importance in obtaining an exfoliated film.

Poly (lactic acid) (PLA) is a semicrystalline and biodegradable polymer that can be produced from renewable sources such as corn and whey (21) and is considered as a potential replacement for petrochemical-based materials. Because of its ideal combination of physical properties (high modulus, good film and fiber forming properties, good heat seal characteristics, and barrier to flavor and aroma) and competitive costs, it is being investigated in a number of commodity applications including flexible packaging (films), fibers, and containers (22). However, PLA is brittle and exhibits low thermal stability, medium gas barrier, and low solvent resistance against water (23). As pointed out earlier, the intercalation of nanoplatelets with high aspect ratio such as layered silicate montmoril-lonite clay can lead to an enhancement in the polymer matrix properties in terms of mechanical properties and barrier (23). Ray et al. (24), (25) studied the properties of PLA nanoclay composite films. They observed an intercalation of the layered silicates that resulted in an improvement in modulus of the nanocomposites and a noticeable reduction in oxygen permeability.

Although few authors have investigated the structure of PLA nanocomposites in various processes, no study has been conducted on properties of blown films of PLA nanoclay composites and their multilayer with PE. In this work, trilayer films of LLDPE/PLA-nanocomposite/LLDPE were produced to investigate in details their structure and orientation in relation to final properties. The impact of the nanoclay addition on the orientation, mechanical properties, tear, haze, and clarity were examined and are discussed in this study.



An extrusion grade commercial poly (lactic acid) (PLA 2002D) supplied by NatureWorks and having a melt flow rate (MFR) of 6 g/10 min (under ASTM D1238 conditions of 230 [degrees] C and 2.16 kg) was selected. An LLDPE with MFI of 1 g/10 min (190 [degrees] C, 2.16 kg) from Nova Chemicals was also used. This resin is a butene based Ziegler-Natta (Z/N) catalyzed copolymer (5.4 mol% octene) traded under name PF-0018-F. The PLA and PE had densities of 1.24 and 0.96 g/ [cm.sup.3], respectively. The nanoclay used was organically modified montmorillonite clay (Cloisite 30B) supplied by Southern Clay Products.

Film Preparation

A master batch of PLA containing 10 wt% of nanoclay Cloisite 30B was prepared using a twin screw extruder and was diluted to the concentrations of 1, 2.5, 5, and 7.5 wt%. The processing temperature was 200 [degrees] C, with the extruder profile between 180 and 200 [degrees] C and the flow rate was about 3 kg/hr. A detail description of the screw configuration can be found in our recent publication on PET nanoclay composite films (26). In summary, after the conveying and pressure elements, the first mixing zones with 45 [degrees] right and left hand (LH) kneading block (KB) are located. The second mixing zone includes a short 45 [degrees] kneading block and the third mixing zone consists of 45 [degrees] and 90 [degrees] kneading blocks. The films were produced using a five-layer Brampton Engineering extrusion blowing line. For films preparation, the materials were dry blended with the master batch and extruded in a single screw extruder with typical screw configuration in the multilayer film line. The flow rate was about 3 kg/hr. All the films were produced with the same conditions of draw down ratio (DDR) of 25 and blow up ratio (BUR) of 2. Low frost-line height (FLH) was used in this work to minimize polymer relaxation. The die gap was set at 1.1 mm and the die temperature was maintained at 250 [degrees] C. All the films were prepared with the same thickness of almost 75 [micro] m, whereas all the layers had an equal thickness of around 25 [micro]m.

Film Characterization

X-ray Diffraction. X-ray diffraction (XRD) measurements were carried out using a Bruker AXS X-ray goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper CuK [alpha] radiation ([lambda] = 1.542 [angstrom]) was selected using a graphite crystal monochromator. The sample to detector distance was fixed at 9.2 cm. To get the maximum diffraction intensity, several film layers were stacked together to obtain the total thickness of about 2 mm.

Wide angle X-ray diffraction (WAXD) is based on the diffraction of a monochromatic X-ray beam by the crys-tallographic planes (hkl) of the polymer crystalline phase. Using a pole figure accessory, the intensity of the diffracted radiation for a given hkl plane is measured as the sample is rotated through all possible spherical angles with respect to the beam. This gives the probability distribution of the orientation of the normal to hkl plane with respect to the directions of the sample.

The Herman orientation function Fijof a crystalline axis i with respect to a reference axis j is given by (27):

[F.sub.ij] = (3 [cos.sup.2] ([[empty set].sub.ij) - 1)/2 (1)

where [[empty set].sub.ij] is the angle between the unit cell axes i (a, b, or c) and the reference axis j.

The Herman orientation functions were derived from the 110 and 200 pole figures for the LLDPE. For LLDPE, since the tf-axis of the unit ceils is perpendicular to the 200 plane, its orientation (Fa) relative to the machine direction (MD) can be measured directly as follow:

[F.sub.a] = [F.sub.200] = (3 [cos.sup.2] ([[empty set].sub.200]) - 1)/2 (2)

On the other hand, Fc (orientation of the c-axis) with respect to MD is determined by the combination of data of two planes for LLDPE, which are 110 and 200 (27), (28):

[cos.sup.2] ([[empty set].sub.c]) = 1 - 1.435 [cos.sup.2] ([[empty set].sub.110]) - 0.565 [cos.sup.2] ([[empty set].sub.200]) (3)

where [empty set] c is the angle between the unit cell c-axis and MD. The orientation parameter for the 6-axis can be calculated from the orthogonality relation:

[cos.sup.2] ([[empty set].sub.b]) = 1 - [cos.sup.2] ([V.sub.a] - [cos.sup.2] ([[empty set].sub.c] (4)

The orientation factors from WAXD are mainly due to the crystalline part, therefore no information about the orientation of the amorphous phase can be obtained.

Mechanical and Tear Analysis. Tensile tests were performed using an Instron 5500R machine. The procedure used was based on the D638-02a ASTM standard. A standard test method for the tear resistance of plastic films based on ASTM D1922 was used to obtain the MD and transverse direction (TD) tear resistances. According to this standard, the work required in tearing is measured by the energy loss of the encoder, which measures the angular position of the pendulum during the tearing operation.

Haze and Clarity. Haze measurement was performed in accordance with the procedure specified in the ASTM D 1003-97. The measurements were carried out on a Haze Guard Plus (TM) instrument (Model 4725) made by the BYK-Gardner (Columbia, MD).


Since the properties of semicrystalline polymers depend on their crystalline structure, in this study, a detailed investigation of the crystalline orientation of nanoclay composites was performed using WAXD. Figure 1 shows the surface (MD-TD plane) as well as the cross-section (MD-ND plane) WAXD patterns for the PE/PLA/PE and PE/PLA-nanocomposite/PE films with 2.5 and 7.5 wt% nanoclay. The first and second diffractions at 2 [theta] = 22 and 24 [degrees] represent the patterns for the 110 and 200 LLDPE crystalline planes, respectively. Obviously, no diffraction peaks associated with the PLA at 2 [theta] = 16.6 and 18.9 [degrees] (29) are observed, indicating that the crystalline fraction of the PLA is quite low. From the literature (30), in PE, two major types of crystallization can occur depending on the magnitude of stress in flow; low stress produces kebabs in the form of twisted ribbons resulting in off-axis 110 and meridian 200 diffractions. By contrast, high stress produces fiat kebabs (planar crystal structure) leading to the appearance of equatorial 110 and 200 diffractions. When the magnitude of flow is in-between, an intermediate arrangement is formed, resulting in off-axis 200 and 110 diffractions (30). Four off-axes reflections for the 110 plane of the LLDPE are observed in Fig. 1 that is a typical behavior of the twisted lamellar structure of PE where the tfaxis rotates around the fr-axis, resulting in the rotation of the reciprocal vector of the 110 plane. The diffraction intensity in the equator is somewhat higher in the cross-section than in the surface. In addition, the diffraction patterns for the surface and cross-section of the films are almost similar. It is clear that by the addition of 7.5 wt% nanoclay, the arcs become little sharper and more concentrated in center, implying a slight improvement in LLDPE crystal orientation in the presence of nanoclay that will be discussed later. Figure 1 also elucidates the appearance of the equatorial spots in the cross-section patterns of the nanofilled films, indicating that the clay platelets have been well aligned in parallel to the surface of the nanocomposite films.

Figure 2 presents the cross-section diffraction intensity profiles of the LLDPE crystal unit cell as well as the clay for the neat and nanocomposite multilayer films. In this figure, a major peak at 2 [theta] = 2.7 [degrees] corresponding to a d-spacing of 001 clay plane of 32.7 A ([d.sub.001] = [lambda]/2 sin [theta] where [lambda] is the X-ray wavelength and [theta] is the diffraction angle) is evident. From the XRD analysis, the original Cloisite 30B clay had a d-spacing of 18.5 A. According to Wang and Wilkie (31), in an immiscible clay-matrix polymeric mixture, [d.sub.001] must be identical to that of the pure clay, but if an intercalated or exfoliated nanocomposite is formed, [d.sub.001] will be higher than that of pure clay. Our results show that an intercalation and/or partially exfoliation of the clay particles has been formed in the PLA nanocomposite films, possibly because of a good affinity between the PLA matrix and Cloisite 30B. On the other hand, several studies (32) reported that PLA and its nanoclay composites do not crystallize easily when quenched from the melt, resulting in the production of films with a low level of crystallinity. It is believed that the nanoparticles can be dispersed much better in amorphous polymers than in crystalline ones from the fact that the crystals prevent the diffusion of the particles during their growth (32). This could be one of the reasons explaining why the nanoparticles disperse better in the polymer matrices with a low level of crystallinity (e.g., PLA) than in the matrices with a high crystallinity (e.g., PP).

The crystalline orientation can be analyzed quantitatively from the WAXD pole figures of the 110 and 200 crystallographic planes for the LLDPE, as depicted in Fig. 3. The normal to the 110 plane is the bisector of the a and b axes and 200 is along the a-axis of the crystal unit cells (27). In Fig. 3a, for the pristine PE/PLA/PE film, a significant orientation of the 110 plane along TD and ND is observed. By contrast, for this film, the 200 plane is significantly aligned in MD whereas no orientation of this axis in TD and ND is visible, confirming the formation of a twisted lamellar structure for the LLDPE layer. However, it is clear that the incorporation of the nanofillers reduces the orientation of the 110 plane along ND and improves it in TD slightly. Moreover, obviously, the addition of the nanoclay slightly enhances the orientation of the 200 plane in ND.

Our previous study (33) on the LLDPE/PP-nanoclay/ LLDPE showed some transcrystallization zone around the interface, at which the PE lamellae nucleated on the PP. In other words, crystallization of LLDPE overgrow at the interface. In fact, a transcrystalline layer is formed when a large number of nuclei are formed on an interface such that the crystallites are forced to grow normal to the interface and when a large difference in crystallization temperature is present. In this work, it is believed that the nanoclay particles in the PLA layer act as nuclei sites, leading to the transcrystallization of the LLDPE lamellae in the interface of the PLA nanocomposite layer and LLDPE layer. This could explain the changes observed in the orientation of the LLDPE unit cell axes with incorporating the nanolillers.



The Herman orientation functions of the LLDPE crystalline axes (i.e., a, h, c) along MD, TD, and ND for the neat and nanocomposite multilayer films are presented in Fig. 4. The c-axis orientation along MD is much higher for the PE/PLA/PE film than the PE/PLA-nanoclay/PE films. However, the c-axis orientation in MD remains almost unchanged with introduction of the nanoclay in the range of 1-7 wt%. In Fig. 4b, it is obvious that the tf-axis orientation along MD and TD increases, supporting the pole figures shown in Fig. 3. Moreover, the data of Fig. 4c reveal that the b-axis alignment in TD reduces significantly on incorporating the clay.


The Herman orientation functions (Foot) of the 001 plane of the nanoclay in the multilayer films are elucidated in Fig. 5. It is obvious that the c-axis of the silicate layers is highly aligned in ND, which is also observed from the pole figure of this plane for the films with 2.5 and 7.5 wt% loading clay platelets. The orientation of the 001 nanoclay plane increases by increasing nanoclay content, which is consistent with those reported in our previous study (33) on the monolayer PP and multilayer LLDPE/PP-nanoclay/LLDPE nanocomposite blown films.

It is well established that the structure of the crystalline and amorphous phases as well as the addition of nanoparticles strongly influence the mechanical and tear properties of films. In other words, the mechanical and tear behaviors are closely related to the structure changes. To clearly understand the effects of the nanoclay addition on the mechanical properties of the manufactured films, the Young modulus and tensile strength for all the films were determined, as depicted in Fig. 6. The Young modulus increased by about 30% by introducing 7.5 wt% of nanoclay. This can be explained by the contribution of the much larger modulus of the clay in the PLA nano-composites. From the literature (34), the tensile modulus of polymer nanoclay composites depends on modulus of matrix, modulus of clay platelets, dispersion of clay, clay loading, orientation of clay particles, and orientation of polymer crystallites. It should be pointed out that incorporating 1 wt% of nanoclay slightly reduces the modulus, whereas further increases of the clay drastically enhance the modulus. Furthermore, the Young modulus is somewhat identical in MD and TD, possibly because of the slight orientation differences along MD and TD.

Similar to the modulus, the tensile strength of the films shown in Fig. 6 decreased slightly with introducing 1 wt% of nanoclay, whereas further increases of the clay tactoids improved the strength.

The tear resistance of the multilayer films along MD and TD were measured and are reported in Fig. 7. Tear resistance improves dramatically by the incorporation of the clay particles. This can be explained by considering that the silicate layers are able to inhibit or at least to slow down crack propagation by deviating their tear path, as addressed in Ref. (9). The higher tear resistance of the films observed in TD than in MD is attributed to the alignment of the clay platelets in the film plane.

The haze value of the nanocomposite films decreases by the presence of the clay particles up to 5 wt% and then increases by further addition of the clay, as illustrated in Fig. 8. The former feature is possibly because of the change in the crystalline structure of PLA by the clay incorporation. In fact, by adding a small amount of the clay (< 5 wt%), the platelets surface will act as nuclei sites, leading to the formation of the lamellae with much smaller size than the spherulites and as a consequence the haze value reduces at the low clay content. However, since the clays have a much lower clarity in comparison with the PLA, further addition of the clay (e.g., 7.5 wt%) promotes haze in the films. These results confirm that our previous study on the PP nanocomposite films (29) and findings of Wan et al. (35) who showed that the nanocomposite materials were transparent when the clay content was below a critical loading level.



In this work, we have investigated the orientation, mechanical as well as the optical properties of multilayer films of nanoclay filled PLA and PE obtained from film blowing process. The measurement of d-spacing of the 001 plane of the clay tactoids indicated the intercalation of the silicate layer in the PLA middle layer. In addition, the 001 plane of the nanoclay platelets lied into the normal direction (ND) significantly. The changes in the crystal unit cell axes of the LLDPE layer with adding nanoclay to the PLA layer were speculated to be because of the transcrystallization of the LLDPE lamellae at the interface of the LLDPE and nanofilled PLA. With the incorporation of the clay tactoids, the Young modulus enhanced and the tensile strength remained almost unchanged. These behaviors were discussed in relation to the modulus of nanofiller and clay orientation. In addition, the tear resistance along MD and TD increased with adding clay particles, which was attributed to silicate layers that are able to inhibit or at least to slow down crack propagation by deviating the tear path. Finally, the haze of the nanocomposite films was increased by the presence of clay particles, except in the case of low clay contents of 1 and 2.5 wt%.






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Seyed H. Tabatabaei, Abdellah Ajji

CREPEC, Chemical Engineering Department, Ecole Polytechnique, C.P. 6079, Succ. Centre ville, Montreal, QC H3C 3A7, Canada

Correspondence to: Dr. Abdellah Ajji; e-mail; NSERC (Natural Science and Engineering Research Council of Canada) and from FQRNT (Ponds Quebecois de Recherche en Nature et Technologies).

Published online in Wiley Online Library (

[c] 2011 Society of Plastics Engineers

DOI 10.1002/pen.21968
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Author:Tabatabaei, Seyed H.; Ajji, Abdellah
Publication:Polymer Engineering and Science
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Date:Nov 1, 2011
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