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Molecular orientation behavior of mesomorphic isotactic polypropylene under uniaxial and biaxial deformation.


Isotactic polypropylene (iPP) is a typical crystalline polymer that is widely used for packaging and industrial laminate films because it has the advantages of strength at higher temperatures, gas barrier properties, and resistance to abrasion and oil and grease. It is well known that iPP is a typical polymorphic material with three main crystal forms, such as monoclinic [alpha]-modification, hexagonal [beta]-modification, and triclinic [gamma]-modification, of which the [alpha]-modification is the most stable crystalline form (1). In addition, iPP with intermediate order between crystalline and amorphous phases is obtained on rapid cooling from the melt. Natta et al. (2) termed this structure the "smectic" state and suggested that the smectic structure is composed of parallel 3/1 helices similar to the [alpha]-modification but disorder exists in the packing of the chains perpendicular to their axes. It is referred to as a "mesomorphic phase" or "mesophase". The mesophase can be considered as a frozen intermediate ordering state during crystallization, which is caused by a quenching solidification process that hinders the molecular motions necessary for crystallization (3). As a consequence, mesomorphic iPP films are transparent and their tensile stress at room temperature is considerably lower than that of [alpha]-crystalline iPP films. The mechanical behavior of mesomorphic iPP seems to be nearly that of [alpha]-crystalline iPP at higher temperatures (4).

In many industrial processes, iPP is biaxially extended at high temperatures near the melting point and at extremely high elongation speeds. There are typically two types of stretching methods: inflation molding in which films are blown in tubular shape; and tenter molding in which films are elongated in the machine direction and/or the transverse direction by tenter clips (5-9). From a visco-elastic view point, in such conditions the iPP is considered to be in the solid phase rather than the liquid phase, resulting in that the industrial elongation processes involve inhomogeneous cold-drawing or necking deformation reflecting on the heterogeneous structure, such as crystalline and amorphous phases. Consequently, the mechanical characteristics of mesomorphic iPP at room temperature under slower elongation speeds may give a physical insight into the industrial elongation processes. In addition, in tubular and tenter stretching methods, the films before stretching are considered to include the mesomorphic phase because they were quenched rapidly by cooled air or chill roll immediately on discharge from the die in the molten state.

Most studies of the deformation mechanism of iPP films have been conducted on biaxially drawn films (10), (11), but very few investigations have been done on their deformation behavior during biaxial stretching. In this article, we present an experimental method for in situ measurements of birefringence under various stretching modes including biaxial as well as uniaxial stretching. The method is applied to a mesomorphic iPP film and we compare the molecular orientation and the local strain states under biaxial and uniaxial stretchings.



The material used in this work was a commercial iPP with high tacticity (98%), weight-average molecular weight [M.sub.w] of 380 k, and polydispersity [M.sub.w]/[M.sub.n] of 4.9. iPP pellets were compression molded in a laboratory hot press at 230[degrees]C. The pellets were completely melted for 10 min between two aluminum sheets under 15 MPa pressure to prepare films with about 80 [micro]m thick. On removal from the hot press, the samples were immediately plunged directly into an ice water bath maintained at 0[degrees]C to prepare the mesomorphic iPP films.


Wide angle X-ray diffraction (WAXD) was carried out at room temperature with a Rigaku Ultima IV diffractometer and Cu-K[alpha] radiation at a 2[degrees] [min.sup.-l] scanning rate over the diffraction angle (2 [theta]) range 5-40[degrees]. The diffraction pattern (see Fig. 1) showed a broad amorphous background superimposed on two broad peaks of the 14.5 and 21.5[degrees], reflections of the mesomorphic modification (12). The formation of a supermolecular structure, such as spherulites can be revealed by transmission optical microscopy between crossed polarizers. It was confirmed that there are no supermolecular structure in the present sample films, suggesting that the sample film is isotropic and shows a high transparency because of no light-scattering because of the spherulitic structure.


Density of the film was determined by a conventionally flotation method at room temperature. Ethyl alcohol was used as a medium. The density value of the sheet was estimated to be 884 kg/[m.sup.3], which corresponds to the volume fraction of mesophase of 48% (4).

Mechanical Experiments

Square specimens (50 mm X 50 mm) cut out from the compression molded films were stretched in tension by a biaxial tensile tester manufactured by Kato Tech in air at room temperature, which is an appropriate temperature for mesomorphic iPP materials. The sample specimens were clamped by four chucks on each side. The initial gauge area was 30 mm X 30 mm. The maximum detectable load was about 490 N, and the maximum stretching displacement was 90 mm. The biaxial tensile tester was specially designed for both pairs of four chucks on two stretch directions to symmetrically move from the central point of the square film so that the central point remained at the initial position during stretching. The stretching speeds in each direction could be controlled independently. The deformation fields of stretched films were estimated by means of black marks drawn on the specimen surfaces. The modes of biaxial stretching experiments performed in this work were simultaneous equibiaxial stretching and planar extension or stretching at constant width. For comparison, we performed uniaxial stretching tests on rectangular specimens with width of 20 mm and length of 50 mm. All tests were performed at a fixed elongation speed of 3 mm/min.

Optical Experiments

The biaxial tensile tester was equipped with a video-camera system that was used to perform in situ rheo-optical measurements in transmission mode. The experimental setup is illustrated schematically in Fig. 2. The polarized visible light passed through the deforming film and was recorded by the video camera mounted on the rheo-optical system. The isotropic film was dark but the stretched film was colored. The degree of the molecular orientation state of the whole film specimen was qualitatively visualized by observing the degree of coloring of the deformed film.


The rheo-optical birefringence investigations were performed using the technique of crossed polarizers (see Fig. 3). The biaxial tensile tester was designed for each pair of clamps to symmetrically move from the central point of the film so that the beam spot remains at the initial position during a whole stretching. The sample was placed between the polarizer and the quarter-wave plate with their axes at 45[degrees] to the strain axis in the specimen (13). An He-Ne laser (632.8 nm, 5 mW) was used as incident beam source. The transmitted beam was received by a photodetecter after passing through a rotating polarizer at about 0.1 Hz (14).


The photo detector was fixed to measure light intensity transmitted from the lower polarizing plate. When the sample was isotropic, the rotation period of the lower polarizing plate was in agreement with the period of the light intensity measured by the photo detector, whereas differences between both those periods appeared for oriented anisotropic samples. The birefringence values were determined from the time lag:

[DELTA]n = [[[omega]([t.sub.light] - [t.sub.plate])[lambda]]/2[pi]d] (1)

where [omega] is the rotation frequency (around 0.1 Hz), [t.sub.light] is the period of light intensity, [t.sub.plate] is the rotation period of the lower polarizing plate, [lambda] is the wavelength (632.8 nm) of the incident beam, and d is the sample thickness. The extension axis of the positive birefringence is taken to be the X-axis (machine direction) and that of negative birefringence was the Y-axis (transverse direction).

Off-line measurements of the orientation states were carried out using the drawn specimens after testing. Deformation, thickness, and retardation of films after stretching were directly determined from the deformed shapes of black marks on the drawn specimens Black circles with diameter of 2 mm were marked on the virgin specimen and the thicknesses around each marked area were measured in advance. Deformations were measured from the deformed shape of the marks printed on the specimen surface; the thickness of each marked area was measured with a Mitsutoyo IDC-112B micrometer. The birefringence values were evaluated from the measured retardation and the thickness of the specimen. The retardation was measured by a Berek compensator, which was fitted to an Olympus BX-51 polarizing microscope.


In situ observations of a mesomorphic iPP specimen uniaxially deformed under crossed polarizers are shown in Fig. 4. The stretching direction was fixed on the horizontal direction with an angle of 45[degrees] under crossed polarizers. The mesomorphic iPP was transparent even in the necked region so that the local strained state or molecular orientation can be visualized by variation of the interference colors caused by the optical retardation, according to the Michel-Levy interference chart (15). The mesomorphic iPP film prepared by compression-molding was isotropic, showing a black color in the original state (Fig. 4a). Extension was applied to the film, inducing some molecular orientation (see Fig. 4b), and the neck was initiated somewhere in the specimen interior. Once the neck had stabilized, the color of the necking portion remained almost unchanged, indicating that the thickness and birefringence of the necked portion were unchanged and the neck propagated in a steady-state fashion extending around the gauge area. According to the stability of the purple color of the necked portion, inside the neck the film was highly oriented in the direction of the stretching X-axis, but at the neck boundary the orientation was found to change rapidly toward that direction.


Figure 5 shows the images of a mesomorphic iPP specimen deformed by the biaxial tester with constant width under crossed polarizers. The specimen necked and extended around the grip area, leaving the interior almost undeformed. In some cases, multiple localized necks appeared around the grips and extended around the grip area to coalesce into one large neck.


Figure 6 shows images of the orientation state together with simultaneous equibiaxial stretch. Initially, the specimen necked around the grips, leaving the specimen interior undeformed. The necks propagated around the grip area from four edges, resulting in a mosaic structure consisting of necked blocks appealing over the whole of the biaxially deformed specimen. The chain molecules within each block preferentially orientates either the X or the Y-axis and not in the intermediate direction. Some dark blocks remained until the final stage of the experimental range (Fig. 6f). This is due to the fact that there is no propagation of necking or that chain molecules orientate to almost the same degree in both X- and Y-axes. It should be noted here that some dark areas appeared after they were previously colored. In these regions the films were stretched equibiaxially, but the propagation mode of necking was sequential equibiaxial not simultaneous equibiaxial. Thus, after necking propagated in the direction of either the X- or the Y-axis, necking propagated in the opposite direction at the same area.


Figure 7a-c show the rheo-optical results of in situ birefringence measurements during uniaxial, stripbiaxial (constant width), and equibiaxial stretchings. Here, the tensile stress was determined by dividing the tensile load by the initial cross sectional area and the tensile strain was calculated from the ratio of the increment of the length between grips in X-direction to the initial gauge length. In these figures, the birefringence values measured simultaneously with tensile loads are plotted against nominal strain. In biaxial tests, the stress values in the X- and Y-axes are plotted on the same figures.


The stress-strain curve of the uniaxial stretching mode is typical of that of a ductile polymer with a maximum corresponding to the yield point. Beyond the yield maximum, the molecular chains are aligned in the stretching direction through yield flow, leading to necking. In the uniaxial test the birefringence values were almost zero in the vicinity of the sampling area, whereas it was located outside the neck. During passage of the neck through the sampling area, the birefringence increases in a stepwise manner toward a certain value between 0.03 and 0.04, and the birefringence was almost unchanged when the sampling area completely entered the neck entity. This result reflects rapidly changing orientation toward the stretching direction at the neck boundary.

In the case of the stripbiaxial stretching mode, there was a large difference in stress between along and normal to the stretch direction and the axial stress was about three times larger than the uniaxial stress. The birefringence behavior was similar to that in the uniaxial case and almost the same value of birefringence was attained.

By contrast, there was no difference in the stresses in both directions in the equibiaxial test mode, indicating that axial deformation proceed simultaneously in both stretching directions. As shown by snap shots of the biaxial test specimen during deformation (see Fig. 6), the specimen showed necking that extended around the grip edge, leaving the interior of the specimen undeformed. As the sampling area was at the center of the specimen, the birefringence was close to zero until the necking portion passed through the sampling point. When the sampling area touched the neck shoulder, dramatic molecular orientation, and longitudinal extension started. After the birefringence increased positively, it oppositely decreased to the negative one. This is consistent with the results of in situ observation of deformed specimens (see Fig. 6); once necking propagated to the direction of cither the X- or the Y-axis, the necking propagated to the opposite direction at the same area.

The local strain state of drawn specimens removed from the tensile tester was examined by means of black marks printed onto the specimen surface before testing. In Fig. 8, the natural logarithm of A/[A.sub.0], where A is the area of deformed marks after stretching and, [A.sub.0] is the area of the original marks before stretching, is plotted against the natural logarithm of t/[t.sub.0], where t is the thickness after stretching and [t.sub.0] is the thickness before stretching for uniaxial, stripbiaxial, and equibiaxial stretching. It is interesting to note that all data fall on the same line. This curve was found to be divided into two regions with different slopes below and above In(A/[A.sub.0]) is 0.5. According to clastic mechanics, the relationship between ln(A/[A.sub.0]) and ln(t/[t.sub.0]) is described by the equation:


ln([t/[t.sub.0]]) = - [v/1 - v]ln([A/[A.sub.0]]) (2)

where v is the Poisson ratio. Consequently, the value of v can be calculated from the slope of the straight lines in each region. In the region corresponding to In(A/[A.sub.0]) [less than or equal to] 0.5, the slope is about 0.49 hence, v is 0.33, which is a typical value for thermoplastic materials. In the region corresponding to ln(A/[A.sub.0])[greater than or equal to] 0.5, the slope becomes about 1.00, giving v = 0.5. This means that the volume of the film remains constant in this region. The value of 0.5 for ln(A/[A.sub.0]) corresponds to the onset of necking. In the prenecking region, there is little intermolecular slippage under tensile deformation whereas in the necking region, there is intermolecular slippage, thus the volume remains unchanged. The values of d in Eq 1 were estimated from the data in Fig. 8.

Figure 9 shows the distribution of local strains in each stretching mode where the strain in the direction of the Y-axis is plotted against the strain in the direction of the X-axis. In the case of uniaxial stretch, when the film was elongated in the direction of X-axis, the strain in the Y-axis direction was negative because of Poisson shrinkage; that is, the neck-in. In stripbiaxial stretch, it was confirmed that there was no strain in the direction of the Y-axis because deformation in the transverse direction was restrained. In the case of equibiaxial stretch, it is interesting to note that the film specimen had undergone prior deformation in the direction of the X- or the Y-axis in the initial strain region of equibiaxialstretching, and there were a few points where the strains were the same in the both directions or in the equibiaxialstretching region. In higher strains or necking regions, the strain points go to the equibiaxialstretching region. This is consistent with the result of in situ measurements of birefringence.


The compression ratio t/[t.sub.0] in the thickness direction plotted against the stretching strains in both X- and Y-axis directions is shown in Fig. 10. In the initial stretching region a neck was stabilized and propagated in the direction of either the X- or the Y-axis, and the thickness decreased drastically in the neck. Above the transverse strains of 2, the compression ratio t/[t.sub.0] at the position where necking propagated in the direction of either the X- or the Y- axis was found to have leveled off, and it turned in the another direction above the strain of about 4. The deformations under uniaxial and biaxial stretchings can be expressed by the draw ratio A/[A.sub.0], irrespective of the stretching mode. Figure 11 shows the relationship between the draw ratio and birefringence in each stretching mode. In the uniaxial and stripbiaxial modes, the birefringence increased in proportion to the draw ratio. On the other hand, in equibiaxial stretch, the birefringence values fell on the same line as the uniaxial and stripbiaxial birefringences at lower strains, but deviated from the line above strains of 2, 3, and 5 and gradually decreased to zero. The stretching ratio of 5 was the maximum strain in the stripbiaxial mode. The maximum points of birefringence at strains of 2 and 3 correspond to the turning points of neck propagation where the direction of necking changed to the perpendicular direction as shown in Fig. 9.



In Fig. 12, birefringence is plotted against the strains in both X-and Y-axis directions in simultaneous equibiaxial stretching. At the positions where the sample was preferentially strained in the direction of the X- or the Y-axis, the birefringence monotonically increased with increasing strain. At positions where the sample was strained in both X- and Y-axis directions, the birefringence values were located on the intermediate region over the saddle point and decreased to nearly zero in the higher strain region, that is, the equibiaxial strained region.



Biaxial and uniaxial deformation behaviors of mesomorphic isotactic polypropylene film prepared by quenching at 0[degrees]C was investigated at room temperature. We constructed an instrument for in situ measurement of local strains and birefringence under uniaxial, stripbiaxial and equibiaxial stretching and we compared the deformation behaviors under uniaxial and biaxial stretching.

Deformation under uniaxial, stripbiaxial, and equibiaxial stretchings mainly proceeds through generation of necking beyond the yielding point, and subsequently propagation of necking, whereas the stress is kept constant in each stretching mode. In the case of the stripbiaxial stretching mode, the stress level along the stretching direction is about three times larger than the stress normal to the stretching direction. The necking behavior is similar to the case of uniaxial stretching, and the sample reaches almost the same level of orientation. In equibiaxial stretching, propagation of necking does not proceed simultaneously in two transverse directions. Necking propagates to one of the two stretching directions, leading to mosaic structures consisting of necked blocks. Subsequently, the necking caused by another stretching process occurs and propagates in the opposite direction in the blocks.


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Correspondence to: Koh-hei Nitta; e-mail:

Published online in Wiley Online Library (

[C] 2010 Society of Plastics Engineers

Shingo Yoshida, (1), (2) Kazuomi Ishii, (2) Takanobu Kawamura, (2) Koh-hei Nitta (2)

(1) Kobe Fundamental Research Laboratory, Sumitomo Bakelite Co., Ltd. 1-5, Murotani 1-Chome, Nishi-ku, Kobe, 651-2241, Japan

(2) Graduate School of Natural Science and Technology, Kanazawa University, Kakuma Campus, Kanazawa, 920-1192, Japan

DOI 10.1002/pen.21787
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Author:Yoshida, Shingo; Ishii, Kazuomi; Kawamura, Takanobu; Nitta, Koh-Hei
Publication:Polymer Engineering and Science
Article Type:Report
Geographic Code:9JAPA
Date:Feb 1, 2011
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