Production, properties and impact toughness of die-drawn toughened polypropylenes.
The research described here forms part of an extensive Joint Research Project between IMI. Boucherville, Canada, and the IRC in Polymer Science and Technology, University of Leeds, U.K. The research Is a comparative study of the production and properties of oriented sheets produced either by roller-drawing (IMI) or die-drawing (IRC) and includes homopolymers and specially toughened blends of polypropylene and polyethylene terephthalate. In this paper we discuss the production and properties of the die-drawn toughened polypropylenes.
Because of their high strength to weight ratios, polymers are cost effective, easy to fabricate, and corrosion resistant, and are starting to replace conventional engineering materials like metallic alloys and ceramics for automotive, construction, household appliances etc. There are a host of important properties such as stiffness, strength, and toughness that a polymer must satisfy for a practical application. By orienting the polymers, the long chain molecules can be aligned along the draw direction, creating molecular orientation and thereby increasing the stiffness and toughness along the draw direction (1). The properties, especially toughness, perpendicular to the draw direction reduce with increasing draw ratio. The purpose of this study is to improve the transverse toughness of oriented polypropylene by incorporation of a polyethylene based elastomer.
Polypropylene, though cheap, easily recyclable, and easy to process, has comparatively poor stiffness and impact strength, thus limiting its applications. It exhibits brittle fracture especially under conditions of plane-strain, low temperature and high strain rates. When unnotched or tested under low strain rate conditions, the sample is ductile and cold draws. The impact strength of polypropylene can be enhanced by the incorporation of a dispersed rubber phase such as the ethylene-propylene copolymer (2, 3), ethylene-propylene diene terpolymer (4, 5), ethylene-octene copolymer (6, 7) and styrene/ethylene-butylene/styrene tri-block copolymer (8, 9). The dispersed phase acts as stress concentrators, promoting crazing and/or shear yielding of the polypropylene matrix (10). According to Olf and Peterlin (11) for a pure polypropylene, crazing is the main fracture energy--dissipating process. In the case of rubber-toughened polypropylene, shear yielding and crazing were observed to be the two stages of fracture ( 12). shear yielding being more dominant and making a significant contribution to toughness enhancement (13).
In the present study, a polyethylene-based elastomer was dispersed in the polypropylene matrix and the blend was then evaluated in terms of drawability, mechanical performance, and impact toughness. The results were compared with those of oriented polypropylene homopolymer. The toughness of oriented polypropylene and oriented impact modified polypropylene drawn to a draw ratio of 4 was investigated under impact loading conditions. The effect of draw ratio on the toughness of the oriented polypropylene has been examined in depth in an associated paper (14).
Polymers can exhibit either brittle or ductile fracture depending on the test temperature and the rate at which they are tested. Brittle fracture is accompanied by unstable crack propagation, a mode of failure that should be avoided in practice. On the other hand, during ductile fracture a considerable amount of energy is dissipated around the vicinity of the crack tip and the crack propagates in a stable manner. Many tough polymers that appear to exhibit ductile failure at room temperature may fracture in a brittle manner at relatively low temperatures. Therefore it is crucial to determine the transition temperature at which the failure mode of the material changes from ductile to brittle. We have therefore extended our fracture studies to observe the fracture behavior of isotropic and oriented polypropylene and blends at different temperatures in order to determine the brittle-ductile transition temperature ([T.sub.BD]). There have been studies where the brittle-ductile transition is determined by changing the test speed (15, 16), but this is outside the scope of the present work.
Impact strength or the specific fracture energy, which is defined as the energy absorbed per unit area of the fracture surface is commonly advocated as the measure of material toughness. in spite of its technological importance, this property is strongly dependent on the geometry and the shape of the test specimen and hence cannot represent the inherent material property. By applying the fracture mechanics approach, a toughness parameter independent of the specimen geometry can be determined.
The mechanics of brittle fracture have been described by Linear Elastic Fracture Mechanics (LEFM), where the critical strain energy release rate, [G.sub.c], is used for characterizing the toughness. [G.sub.c] provides a measure of the energy required to extend a crack over a unit area. The LEFM approach for evaluating the toughness ([G.sub.c])
from impact testing of polymers exhibiting brittle fracture was addressed independently by Brown (17) and Marshall et al. (18) and is given by
[G.sub.c] = U/B * 1/C * dC/da
where U is the fracture energy, B is the sample thickness, a is the crack length and C is the compliance of the sample the expression for which can be found in reference (17) or (18). Equation 1 indicates a linear relationship between total energy to rupture the sample (U) in impact and a geometry factor (BC (da/dC). The gradient of this plot gives the toughness, [G.sub.c], The estimated [G.sub.c] value is independent of the crack length and is solely a material parameter. The intercept on the Y-axis gives an estimate of the kinetic energy lost on accelerating the sample from rest (17, 18).
LEFM theory assumes that there is minimal plastic flow at the crack tip and the energy required to fracture the sample is that stored elastically up to the point of fracture. However, for ductile fracture, especially the fracture of ductile and toughened polymers, plasticity precedes stable crack propagation, causing a continuous transfer of energy from the striker to the sample during the test. Because of the excessive yielding during ductile fracture, it is inappropriate to use LEFM parameters for characterizing the ductile fracture. For such cases, the elastic-plastic J-integral concept (19) has to be employed to assess the toughness. The J value can be calculated from the fracture energy from the following equation,
U = J/2 B(W - a) (2)
where U is the fracture energy, B is the sample thickness, Wis the sample width and a is the crack length. From Eq 2, the slope of the plot between (2 U/B) and ligament length (W-a) will yield the toughness J value. The fracture energy, U, is a combination of the energy required to initiate a crack and propagate it through the sample until final rupture. Thus for the elastic case, J will be almost be equal to [G.sub.c], as the propagation energy for a brittle fracture is negligible.
Various J-integral techniques are currently being used to study the initiation toughness, [J.sub.c], under impact loading conditions (20-24). All these techniques involve stopping the test after the crack has extended by a small fraction or by employing high-speed photography to monitor the advancement of the crack during the test. As these were not possible with our experimental arrangement, we measured the total toughness, [J.sub.I], which includes both the initiation and propagation components. The subscript I indicates that the samples were tested under mode I.
The material investigated in this study was a commercial grade polypropylene, Profax 6823 from Montell (melt flow index = 0.5) supplied in the form of pellets. The polyethylene elastomer blended with polypropylene was Engage 8150 elastomer manufactured by Dow DuPont. Engage 8150 elastomer is an ethyleneoctene copolymer containing 25% octene. This was also supplied in the form of pellets. In this study. three materials were characterized in terms of their processing-structure-property relationships. They are A-polypropylene homopolymer. B-90% polypropylene + 10% by weight Engage 8150 elastomer, and C- 75% polypropylene + 25% by weight Engage 8150 elastomer.
The blends were compounded in a co-rotating twin-screw extruder, at 200[degrees]C and 200 rpm. and pelletized. For studying the die-drawing behavior, sheets approximately 5 mm thick and 100 mm wide were extruded at 200[degrees]C. The edges of the extruded sheets were discarded and sheets of 65-mm width were machined. The thickness of the extruded sheet was not altered. In order to aid the start-up procedure for the die-drawing process, a tag 50 mm long and 1 mm thick with a taper of 15[degrees] semi-angle was machined on one end of each sheet. On each billet, lines of 20 mm spacing were drawn along the length of the sheet. The ratio of the spacing after ([L.sub.D]) and before drawing (LB) gave the actual draw ratio, [R.sub.A].
The billet was placed in the heated chamber at 110[degrees]C and the tag gripped by the haul-off unit to which the load cell is attached. The material was allowed to soak at the draw temperature for 1 hour to establish good thermal equilibrium. After the soaking period, the tag was drawn at a slow speed of 50 mm/min until the starting face was fully drawn out and oriented material emerged. The drawing process was stopped and the tag was cut. The product was re-gripped and the drawing was started at the desired rate. The drawing load was monitored throughout the run and the average value of the load was used for the calculation of draw stress. To investigate the drawing behavior, A, B, and C were drawn over a range of drawing speeds between 30 and 540 mm/min. A schematic sketch of the die-drawing process is shown in the Fig. 1.
The tensile moduli of the isotropic and die-drawn sheets of material A, B, and C were determined using a video extensometer in the Instron testing machine. Dumbbell samples of gauge length 40 mm and width 3 mm were stamped from the drawn sheets both parallel and perpendicular to the draw direction. The samples were tested at a crosshead speed of 5 mm/min in a temperature-controlled room maintained at 20[degrees]c. The modulus was measured at selected draw ratios, and in each case at least eight samples were tested and the average values were quoted.
Wide Angle X-Ray Diffraction
The wide-angle X-ray diffraction (WAXS) patterns of the isotropic and the oriented sheets of draw ratio 10 were obtained using a flat camera assembly with CuK[alpha] radiation generated at 40 kV and 30 mA. The specimen to film distance was 50 mm. The draw direction in all the cases was perpendicular to the incident X-ray beam.
The sample dimensions for the impact tests are 60 X 10 X 5 [mm.sup.3]. Isotropic samples for the impact tests were cut from the extruded sheet parallel to the extrusion direction and annealed in an oven for 1 hour at 120[degrees]C to remove the orientation effects during the extrusion process. For characterizing the impact toughness on oriented polypropylene and blends, thick plaques of A, B, and C were compression molded and then welded together to produce thick and long billets, which were annealed at 120[degrees]C to remove the orientation effects due to the compression molding process. Following annealing, the samples were die-drawn to a draw ratio of 4 at 1 110[degrees]C. Impact samples were cut from the drawn samples with the initial notch parallel (PL) and perpendicular (PR) to the draw direction as shown in Fig. 2. The impact tests were performed in the Rosand instrumented impact-testing machine fitted with a temperature cabinet. The tests were undertaken at different temperatures to determine the bri ttle-ductile transition point of each material.
At least 15 samples with different crack lengths were tested at each temperature. The samples were notched on one side of the sample half way through the length. The crack tips were sharpened by gently tapping a fresh razor blade for brittle materials and slicing a serrated razor blade for ductile materials. The ratio of crack length to specimen width was kept between 0.45 and 0.55, following the recommendations suggested in the testing protocol by the European Group on Fracture (EGF) (25). The samples were left to soak at the test temperature for 90 minutes. After the soaking time, the impact tests were performed at 1 m/s. The load-displacement plots were downloaded and analyzed on a computer interfaced with the impact machine.
RESULTS AND DISCUSSION
The variation of draw stresses with draw ratio for materials A, B, and C is shown in Fig. 3. The draw stress was based on drawing load and the cross-sectional area of the drawn sheet. In general, for all the three materials, the draw stress increased with draw ratio. For a given draw ratio, the draw stress required to draw the blends is very much less than the stress required to draw material A to similar draw ratio, implying that the amount of work done to draw the blends through the die is comparatively less than for the polypropylene homopolymer.
Figure 4 shows the effect of draw ratio on the tensile moduli of die-drawn A, B, and C sheets. In all cases, the moduli increased with draw ratio although at different magnitudes. The enhancement of modulus with drawing has been interpreted in terms of the morphology of the drawn sample (26, 27). Taraiya et al. (28] concluded that for oriented polypropylene, both the taut tie molecules and extended chain crystals contributed to the improved properties along the draw direction. Increasing the elastomer content in the polypropylene matrix decreased the modulus of the blend. The difference in modulus between the isotropic and drawn homopolymer and the blends is primarily due to the presence of the soft rubbery phase in the blend.
In the transverse direction, the moduli for all three samples decreased with increasing draw ratio, as shown in Fig. 5. This is due to molecular orientation in the draw direction, the extent of which increases with draw ratio. In sample A for a draw ratio of 6, the enhancement of the modulus along the draw direction is approximately 65% with reference to the isotropic. The mechanical anisotropy of modulus due to molecular orientation causes the modulus to drop by over 30% in the direction perpendicular to the principal draw direction. The modulus in this direction decreased on blending, as shown in Fig. 5, and the fall is proportional to the elastomer content.
Wide Angle X-Ray Diffraction
The WAXS patterns for the isotropic and oriented samples of [R.sub.A] = 10 for A, B, and C showed strong reflections from diffraction planes having Miller indices (110), (040) and (130), as shown in Fig. 6. Drawing to [R.sub.A] = 10 causes sharpening of the diffraction arcs, indicating crystalline alignment. From the WAXS patterns in Fig. 6, it may be observed that for a [R.sub.A] = 10, the crystalline orientation in the blends B and C is comparatively lower than that in homopolymer A. The presence of the crystalline rings along the diffraction arcs in the WAXS patterns for B and C at higher draw ratio suggests the presence of crystalline orientation and isotropic domains in the sample, confirming that the crystalline orientation in the blends is comparatively less than the drawn polypropylene homopolymer, A. The elastomer phase in the blend restrains the crystalline orientation in polypropylene by maintaining some of the isotropic domains. The suppression of the orientation in polypropylene as a result of ble nding is another reason for the decrease in the modulus of the drawn blends.
Impact Tests on Isotropic A, B, and C
For brittle fracture, [G.sub.c] values were obtained from the slope of the plot of total energy required to rupture the sample (U) in impact and the geometry factor, BC [da/dC]. For ductile fracture, the slope of the plot of 2U versus ligament area (B(W-a)) gave the toughness, J. Representative plots for the calculation of [G.sub.c] and J values are shown in Figs. 7 and 8. The toughness values of the isotropic samples of A, B, and C obtained at different temperatures are shown in Fig. 9, and the toughness values at selected temperatures are listed in Table 1.
Sample A exhibited a brittle-ductile transition temperature [T.sub.BD] at around 40[degrees]C, B showed a transition at 10[degrees]C and C had a [T.sub.BD] of 40[degrees]C. This was confirmed from the fracture surfaces of A, B, and C tested above and below the [T.sub.BD] as shown in Figs. 10 to 12. All three materials exhibited ductility above the transition temperature, as evident from the presence of deformation marks on the fracture surface. Below the [T.sub.BD] all three materials fractured in a brittle manner. In this regime, formation and subsequent breakdown of crazes is usually known to be the predominant fracture mechanism (13), whereas above the [T.sub.BD], shear yielding is thought to be the main toughening phenomenon (12, 29).
At 20[degrees]C, polypropylene homopolymer is brittle in notched impact tests with a low toughness. By dispersing an elastomer phase in the polypropylene matrix, the toughness of the blend has been enhanced and the fracture behavior changed from brittle to ductile mode as revealed by the fracture surface (cf. Figs. 10a to 12). It is interesting to note from Fig. 9 that in the brittle regime (below the [T.sub.BD] for the homopolymer and the blends), the blends are still tougher than the homopolymer, which suggests that the elastomer particles maintain the craze integrity that would otherwise lead to uninhibited growth of crazes, resulting in premature failure by crack initiation and growth. Above the [T.sub.BD], shear yielding is the main deformation mechanism for both the homopolymer and the blends. Increasing the elastomer content leads to the initiation and growth of multiple shear bands in the matrix at the site of the elastomer particles, which postpones the failure to higher strain levels than those fou nd for the unmodified polymer. The considerable amount of plastic energy dissipation that occurs during the formation and growth of shear bands results in an increase in the macro-toughness of the blends above the transition temperature.
Impact Tests on Die-Drawn A, B and C With the Notch Parallel to the Draw Direction (A-PL, B-PL & C-PL)
The toughness values of A-PL, B-PL and C-PL drawn to a draw ratio of 4 are shown as a function of temperature in Fig. 13. Because of the anisotropic nature of the oriented polymer (cf. WAXS patterns, Fig. 6), the toughness of all the samples in this direction dropped when compared to that of the isotropic state as shown in Figs. 14a to 14c. Moreover, on drawing, owing to the reduced mobility of the chains, the transition temperature increased when compared to the isotropic state. A-PL exhibited very poor impact toughness and failed in a brittle manner up to 60[degrees]C. Very small plasticity was noticed in the fracture surface above 60[degrees]C (Fig. 15). B-PL showed a transition around 20[degrees]C (Fig. 13). The fracture surface below the [T.sub.BD-PL] appeared smooth with very little evidence of plasticity on the fracture surface (Fig. 16a). Above [T.sub.BD-PL], discrete bands of ductility were present on the fracture surface (Fig. 16b), indicating that some plasticity had occurred prior to and during fr acture. C-FL had a transition at -20[degrees]C. The fracture surface above the [T.sub.BD-PL] was marked by enhanced ductility, as opposed to the smooth fracture surface in the case of brittle fracture (Fig. 17). Even though the toughness of C in this direction (PL) decreased when compared to its isotropic state, it is encouraging to note that C-PL is 85% tougher and has a [T.sub.BD] around 50[degrees]C less than the isotropic polypropylene homopolymer, A (cf. Fig. 9 and Fig. 13).
Impact Tests on Die-Drawn A, B and C With the Notch Perpendicular to the Draw Direction (A-PR, B-PR & C-PR)
The photographs of the fractured samples when A-PR, B-PR, and C-PR were tested at 20[Degrees]C are shown in Fig. 18. In all cases, the fracture process was incomplete with the specimens exhibiting a hinge type break with the specimens still held intact. This is described as "H-type failure" according to ISO standard 179. A plastic zone developed opposite to the crack front, because of localized shear yielding due to compression and prevented the propagation of the main crack through the sample.
When A-PR was tested at 20[degrees]C, the crack bifurcated producing subsidiary cracks (Fig. 18a). The formation of secondary cracks resulted in a fibrillar type of fracture, causing the sample to split perpendicularly to the axis of loading i.e., along the draw direction, deflecting the direction of the primary crack to propagate through the length of the sample. The damage zone in this case is reminiscent of that of a unidirectional fiber composite (30-32) and oriented Perspex (33), when the initial crack was perpendicular to the principal fiber or orientation direction.
The fracture event when B-PR was tested at 20[degrees]C can be divided into two stages. First, at the crack tip, an initial lobe shaped damage zone with a small plastic zone formed at the tip of the main crack. Further loading resulted in the termination of this initial damage zone and initiation of another damage zone at the foot of the pre-existing zone. The depth of propagation of this damage zone was bigger than the former and no plastic zone was visible. The main crack propagated stably in the damage zone (Fig. 18b). The formation of a plastic zone at the tip of the initial crack is probably due to the temperature buildup at the crack tip. Because of the high rate of testing together with the low thermal conductivity of the polymer, the heat transfer to the rest of the sample is highly retarded. Thus, the plastic zone was confined to the crack tip. This plastic zone masked the growth of the secondary crack, and as a result, the depth of propagation of the secondary crack in the first stage is minimal com pared to that seen in the second.
When C-PR was tested at 20[degrees]C, a plastic zone was seen to form at the crack tip, as shown in Fig. 18c. Further loading resulted in the stable propagation of the primary crack In the stress-whitened zone. At this temperature, no secondary cracks were generated at the crack tip. Because of the complicated fracture behavior, the toughness along this direction was not quantified, and the results have been interpreted only qualitatively for comparison purposes. The crack tip blunting mechanism in anisotropic materials with the crack growing perpendicular to the orientation direction was first addressed by Cook and Gordon (34). According to their theory (34), in an anisotropic material, two components of stress act at the vicinity of the crack tip. In addition to the stress acting along the loading direction, secondary stresses are generated perpendicular to the loading direction or parallel to the primary crack propagation direction. In an anisotropic material, because of the reduction in the transverse str ength, secondary cracks would open that would blunt and deflect the main crack to follow the weaker direction. This deflection of the main crack by the generation of secondary cracks has been viewed as increasing material toughness (34-38).
From the size of the damage zone ahead of the main crack front in Fig. 18, it can be inferred that the homopolymer A-PR was tougher in this direction than the blends B-PR and C-PR. The toughness along this direction depends on the degree of orientation of polypropylene. As confirmed from the WAXS patterns (Fig. 6). A had a higher degree of orientation than the blends B and C. Increasing the elastomer content decreased the orientation in polypropylene and hence the toughness. This is clearly exemplified in the fracture behavior of C-PR. where no damage zone was observed (Fig. 18c).
The inclusion of an elastomer phase in the polypropylene matrix also resulted in a decrease in the [T.sub.BD-PL] for the blend. Below the [T.sub.BD-PL], the toughness perpendicular to the draw direction is low. Consequently, secondary cracks would propagate easily along this direction, whereas above the [T.sub.BD-PL], the toughness increased appreciably and prevented the secondary cracks from propagating.
The [T.sub.BD-PL] for oriented A is around 60[degrees]C. When tested at 20[degrees]C. which is 40[degrees]C below the [T.sub.BD-PL], secondary cracks would propagate easily along this direction because of the weaker plane perpendicular to the main crack. Conversely, for C, the [T.sub.BD-PL] is 40[degrees]C below the temperature at which the sample was tested. As a result of this difference, no secondary cracks were visible. [T.sub.BD-PL] for B is around 20[degrees]C, which is close to the temperature at which the sample was tested. Thus a mixed mode fracture behavior with both the secondary and the main crack propagation could be seen in this sample.
On a decrease in the test temperature to -20[degrees]C, the size of the damage zone for sample A-PR was greater than that observed at 20[degrees]C (Fig. 19a) because of the further reduction in the toughness in the weaker direction at low temperatures. The fracture behavior of B-PR (Fig. 19b) resembled that of A-PR, with a large damage zone because the test temperature was less than the [T.sub.BD-PL]. In the case of C-PR, the size of the plastic zone was very small and confined to the tip of the main crack (Fig. 19c). It is interesting to note that no secondary cracks were seen in sample C-PR in the weaker direction, even though the test temperature was below its [T.sub.BD-PL]. This is due to the improvement in toughness in the transverse direction as a result of blending. Moreover, the inclusion of elastomer decreased the orientation in the longitudinal direction when compared to that of the polypropylene homopolymer, as shown earlier.
When tested at 60[degrees]C, the formation of secondary cracks is minimal and the main crack is propagating stably through the sample (Fig. 20). In the case of A, a mixed mode fracture behavior with both the secondary and the main crack propagation could be seen because the test temperature was close to its [T.sub.BD-PL]
The effect of impact modifier on the die-drawing behavior and mechanical properties of polypropylene has been investigated. The addition of modifier particles reduced the draw stress significantly, allowing the blend to be easily oriented at low draw temperatures. Parallel to the draw direction, the tensile modulus of the oriented homopolymer and that of the blends increased with draw ratio, the properties being lower for the blends than the homopolymer at all draw ratios. In the transverse direction, the modulus decreased as a result of anisotropy along the draw direction. The WAXS patterns suggest that for comparable draw ratios, the crystalline orientation of the blends was lower than the homopolymer.
Under impact loading, the isotropic blends showed a pronounced increase in toughness over the polypropylene homopolymer. When drawn to a draw ratio of 4, the toughness of the pure homopolymer and the blends when tested with the initial notch parallel to the draw direction decreased when compared with the isotropic material. When tested with the initial notch perpendicular to the draw direction, the blends and the homopolymer showed considerable improvement in the impact toughness.
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[FIGURE 8 OMITTED]
[FIGURE 9 OMITTED]
[FIGURE 13 OMITTED]
[FIGURE 14 OMITTED]
Table 1 [G.sub.c] and J Values for Isotropic (a) A, (b) B and (c) C. Temperature [degrees]C [G.sub.c] J kJ/[m.sub.2] Temperature KJ/[m.sub.2] [degrees]C 0 3.9 [+ or -] 0.8 -- -40 20 4.3 [+ or -] 0.8 -- -20 50 -- 23.6[+ or -]1.3 20 80 -- 30.3[+ or -]1.4 30 Temperature [degrees]C [G.sub.c] J kJ/[m.sub.2] KJ/[m.sub.2] 0 4.1[+ or -] 0.8 -- 20 5 [+ or -] 0.7 -- 50 -- 69 [+ or -] 7.7 80 -- 96.5 [+ or -] 4.7 Temperature [degrees]C [G.sub.c] kJ/[m.sub.2] J kJ/[m.sub.2] -60 5.4 [+ or -] 0.9 -- -20 -- 36.9 [+ or -] 9.2 -10 -- 85 [+ or -] 19.6 20 -- 169 [+ or -] 5.8
The financial support of this work by the National Research Council, Canada and British Council, UK, is gratefully acknowledged. J. Mohanraj is grateful to the Overseas Research Student Awards scheme, UK for a Postgraduate research scholarship award. The authors would also like to thank Dr. P. J. Hine, Dr. A. K. Taraiya and Dr. M. Bonner for their useful discussions and Mr. D. Stephenson for his assistance in sample preparation.
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J. MOHANRAJ (a)
N. CHAPLEAU (b)
A. AJJI (b)
R. A. DUCKETT (a)
I. M. WARD (a) *
(a.) IRC in Polymer Science & Technology University of Leeds Leeds, U.K.
(b.) IMI-NRC 75, de Mortagne Boucherville, Quebec, Canada
* Corresponding author, Email: firstname.lastname@example.org
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|Author:||Mohanraj, J.; Chapleau, N.; Ajji, A.; Duckett, R. A.; Ward, I. M.|
|Publication:||Polymer Engineering and Science|
|Date:||Jun 1, 2003|
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