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Effect of Structural Inhomogeneity on Mechanical Behavior of Injection Molded Polypropylene Investigated with Microbeam X-ray Scattering Yuya Shinohara,*,† Kosuke Yamazoe,† Takashi Sakurai,‡ Shuichi Kimata,‡ Toshiya Maruyama,‡ and Yoshiyuki Amemiya† †

Graduate School of Frontier Sciences, The University of Tokyo, 5-1-5 Kashiwanoha, Kashiwa, Chiba 277-8561, Japan Petrochemicals Research Laboratory, Sumitomo Chemicals Co., Ltd., 2-1 Kitasode Sodegaura, Chiba 299-0295, Japan

‡

ABSTRACT: Relationship between the structure of injectionmolded isotactic polypropylene and its tensile mechanical properties, necking and fracture behaviors in particular, was investigated in terms of micrometer-scale structural inhomogeneity of nanometer- and subnanometer-scale structures. To clarify the micrometer-scale inhomogeneity, we employed scanning microbeam wide-angle X-ray diffraction and smallangle X-ray scattering technique. Four isotactic polypropylene samples were studied, produced using different injectioncondition and thermal treatments. The results of scanning microbeam X-ray scattering measurements showed the presence of two types of micrometer-scale structural inhomogeneity in addition to the orientation of molecules: the distribution of polymorphs and of crystalline ordering. The results of scanning microbeam X-ray scattering of deformed sample showed the disappearance of the β-form isotactic polypropylene crystals at the outer regions accompanied by the plastic deformation. It is indicated that the inhomogeneous distribution of crystalline ordering and the existence of different polymorphs are highly related to the tensile mechanical behavior.

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mesophase.10,11 It has been reported that injection-molded iPP mainly composed of the α-form of crystallites while the β-form and the γ-form exist as minority components.12 Significant differences in the mechanical properties between the different forms of iPP have been reported;13−16 the β-form crystal17−19 is thermodynamically and mechanically less stable than the αform.10 The effect of molding conditions on the final morphology has been widely investigated by using POM, transmission electron microscope, infrared spectroscopy, and X-ray scattering.20−25 It is of practical importance to investigate the mechanical properties of products with relation to the layered structures.3,4,8,25−39 The effects of both the morphology and the nucleating agents on the mechanical properties have been studied from the industrial point of view.8,21,35−38 The tensile mechanical behaviors depend on the morphology of injectionmoldings such as the volume ratio of the skin−core layer, the degree of crystallinity, and the orientation of crystallites.7,23,27,31,39 The thermal treatments also affect the structure and tensile mechanical properties of products;12,26,32,37−40 for example, an increase in crystallinity and crystal size by annealing process has been reported.40,41 Recent solid-state NMR techniques reveal that annealing treatments induce the

INTRODUCTION Isotactic polypropylene (iPP) is widely used in industrial products. Most of the products are manufactured by polymer processing such as injection molding, fiber spinning, and extrusion and by additional processing such as annealing. Among these processing procedures, injection molding is widely employed for shaping the products. In the injection molding process, a flow field changes the chain morphology, which accelerates a nucleation rate in the subsequent crystallization process. The structure of injection molding is greatly affected by molding conditions such as injection speed, injection pressure, processing temperature, holding pressure, cooling time, and cooling temperature. Injection-molded iPP generally consists of a skin−core structure in which several layers are observed with polarized optical microscope (POM).1−8 The skin−core structure is typically classified into three layers; a highly oriented skin layer, a shear layer with molecular chains oriented parallel to the flow direction, and a spherulitic core without chain orientation. In the surface region of the skin layer, so-called shish-kebab structure consisting of folded chain lamellae and extended chain crystals can be formed.6,9 A lower cooling rate and a shear strain history in the core area result in the relaxation of chain molecules leading to the growth of spherulites. The molding conditions affect the formation of these layered structures and also induces the appearance of different crystallographic forms: monoclinic αform, hexagonal β-form, orthorhombic γ-form, and smectic © 2012 American Chemical Society

Received: September 27, 2011 Revised: January 4, 2012 Published: January 24, 2012 1398

dx.doi.org/10.1021/ma202178r | Macromolecules 2012, 45, 1398−1407

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min. Uniaxial tensile tests were performed at room temperature and the strain rate was 50 mm/min. All specimens were cut by using microtome with a thickness of approximately 30 μm along the transverse direction as schematically illustrated in Figure 1. Prior to the X-ray measurements, POM images

rearrangement process of the chain packing in melt and isothermal-crystallized iPP.42−44 Despite a large number of studies, the relationship between the structure and tensile mechanical behavior has not been clarified. One of the reasons is the lack of the detailed structural information in terms of micrometer-scale structural inhomogeneity. It has been noticed that there are cases such that the tensile strength and stiffness does not rise as the crystallinity increases.37 Injection moldings are by nature spatially inhomogeneous; thus spatially resolved structural study is required. Microbeam X-ray scattering is a powerful tool to provide the structural information on local spatial inhomogeneity of nanostructure in a micrometer scale.45,46,68 A conventional Xray beam size for scattering experiments is approximately some hundreds micrometers; this size is too large to obtain the micrometer-scale inhomogeneity and only spatially averaged structural information is obtained. Scanning microbeam X-ray scattering measurements have been successfully applied to the structural analysis of injection molded iPP.8,35,36,47−49 However, the inhomogeneity of the crystalline ordering as well as the crystallinity and orientation of the crystallites has not been investigated, despite its importance for discussing the relationship between tensile mechanical behaviors and morphology induced by injection molding. In this paper, we focused on the relationship between micrometer-scale structural inhomogeneity, that of crystalline morphology and crystalline ordering in particular, and the tensile mechanical behavior of injection-molded iPP by examining the effects of difference in the skin−core morphology and annealing on the structural inhomogeneity and mechanical properties. To reveal the micrometer-scale structural inhomogeneity of nanostructure, we employed scanning microbeam wide-angle X-ray diffraction (WAXD) and small-angle X-ray scattering (SAXS). Furthermore, the structure of deformed specimens was investigated to examine the structural change accompanied by the deformation. It was found out that the inhomogeneous distributions of polymorph and that of crystalline ordering are highly related to tensile mechanical behaviors such as necking and fracture. In situ observation of structural change during a deformation process employing X-ray microbeam will be reported in another paper.

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Figure 1. Schematic illustration of the specimen. The directions of flow and scanning X-rays are shown in the figure: MD, machine direction; ND, normal direction; and TD, transverse direction. were recorded to observe the structures of injection-molded iPP in micrometer-scale. Scanning Microbeam Wide-Angle X-ray Diffraction. Scanning microbeam WAXD measurements were performed at BL-4A, Photon Factory (Japan). The X-ray beam was focused to 5 μm × 5 μm (fullwidth at half of maximum: fwhm) at a sample position with a Kirkpatrick-Baez mirror and the X-ray wavelength was chosen to be 0.113 nm with a multilayer monochromator. An X-ray CCD detector (C4880−50, Hamamatsu Photonics Co. Ltd.) coupled with an X-ray Image Intensifier (Hamamatsu Photonics Co. Ltd.)50 was used to record WAXD images. The distance between the sample and the detector was approximately 170 mm. Two sets of beamstop were placed at the downstream of sample to prevent the damage of the Xray detector and to reduce the background air-scattering. We scanned a sample with the X-ray microbeam from an edge to the other edge of the sample with a step of 10 μm as shown in Figure 1. To estimate the X-ray radiation damage, we continuously irradiated the samples with the X-ray microbeam; the sample started to be damaged after an exposure of 20 s. Thus, we set the exposure time for each WAXD measurement to be 4−6 s to avoid the radiation damage during scanning of the X-ray microbeam. The position of X-ray irradiation was monitored with a POM. For deformed samples, we selected several scanning positions to elucidate the dependence of the structural distribution on the position from a fracture front area. Each scanning was performed more than four times at different positions of the samples to confirm reproducibility. Scanning Microbeam Small-Angle X-ray Scattering. Scanning microbeam SAXS measurements were performed at BL40XU, SPring8. The quasi-monochromatic X-ray beam from a helical undulator was used.51 An X-ray beam of 5 μm was produced by inserting a pinhole of 5 μm diameter. Another pinhole of 100 μm diameter was inserted to remove parasitic scattering. The exposure time was 100−3000 ms, and the X-ray wavelength was 0.10 nm. The distance between the sample and the detector was approximately 2.5 m, which was later calibrated with the diffraction peaks of silver behenate. The position of X-ray irradiation was monitored with a POM. The scanning condition was the same with that for the WAXD measurements.

MATERIALS AND METHODS

Materials. A commercial-grade isotactic polypropylene (Sumitomo Noblen) was used as the sample. The weight-averaged molecular weight Mw was 804 000 and Mw/Mn was 3.5, where Mn is the numberaveraged molecular weight. The samples were molded as tensile specimens with a thickness of 3 mm according to ASTM D638 by using an injection-molding machine (Toshiba Machine IS100EN) at a melt temperature of 260 °C and a mold temperature of 50 °C. The details of injection-molding conditions and the sample codes are listed in Table 1. An injection speed was controlled to prepare the specimens

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Table 1. Sample Fabrication Conditions and Sample Codes sample code

injection speed [mm/s]

holding pressure [MPa]

annealing

PP1 PP1A PP2 PP2A

1000 1000 1400 1400

20 20 65 65

no yes no yes

RESULTS Deformation Behavior. The tensile behaviors of PP1, PP1A, PP2, and PP2A are presented in Figure 2 and are summarized in Table 2. The photograph of deformed specimens is shown in Figure 3. Nonannealed specimens, PP1 and PP2, showed a well-defined necking behavior both in the tensile testing and the photograph. They broke at a draw ratio of approximately 2.7 and showed no strain hardening. There was no significant difference between the tensile

with different skin−core morphology. An injection pressure was automatically controlled to keep the injection speed constant. For the annealed samples, we kept the specimens in an oven at 155 °C for 60 1399


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amorphous regions. The degree of whitening seems to be higher for the annealed specimens. To summarize, the necking behavior highly depends on the annealing. POM Observation. Figure 4 shows the POM images of the cross-section of PP1, PP1A, PP2 and PP2A, which shows a

Figure 2. Macroscopic tensile behavior of (A) PP1 (solid line) and PP1A (dashed line), and (B) PP2 (solid line) and PP2A (dashed line).

Figure 4. POM images of PP1, PP1A, PP2, and PP2A.

skin−core structure; the flow direction and the scanning direction of X-ray microbeam are vertical and horizontal in the image, respectively. The layers of a skin−core structure can be characterized by several regions from its surface to its center.1,49 At least five types of layers were recognized by combining the POM figures with the X-ray results described later; we classified them as follows: (A) a skin layer, (B) a shear layer, (C) a finegrained layer, (D) a core layer, and (E) a β-form rich layer. Before describing the characteristics of these layers, let us recall that we do not intend to discuss the mechanism of crystallization under flow conditions9,20 nor the relationship between the molding conditions and the resultant structure; our focus in this study is to clarify the mechanical properties and the micrometer-scale structural inhomogeneity. There are two types of remarkable morphology in PP1 and PP2: (1) PP2 contains a layer of small spherulite between the outer layers (A and B) and the core (D). In the present study, we call this region f ine-grained layer (C). (2) There is another type of layer at the boundary of the skin layer and the core in PP1. This layer has peculiar crystalline polymorph, containing only the β-form crystals of iPP, the detail of which is described in the next section; thus we call this layer β-form rich layer (E). The thickness of outer layers including the skin layer (A) and the shear layer (B) increased with the increase of the injection speed and the holding pressure (PP2 and PP2A). The β-form rich layer (E) in PP1 disappeared by the annealing treatment, suggesting the transition of β-form crystals into α-form crystals of iPP upon annealing. Distribution of Polymorph in Injection-Molded iPP. Figure 5 shows typical WAXD images and their corresponding one-dimensional WAXD intensity profiles that were obtained from different positions on PP1. At the point 3, four different reflections (110, 040, 130, and 041) from the α-form crystals are distinguished, while two reflections (111 and 13−1) overlap each other. No reflection from the γ-form crystals was observed in the present study. In addition to the reflections from the α-

Table 2. Tensile Testing Results of Injection-Molded iPP yield stress [MPa] yield draw ratio fracture stress [MPa] fracture draw ratio

PP1

PP1A

PP2

PP2A

34.0 1.15 20.2 2.6

34.2 1.1 29.4 2.2

35.0 1.15 21.2 2.8

33.6 1.1 38.0 5.6

Figure 3. Photograph of the fractured specimen. The top specimen is a sample before deformation.

behaviors of PP1 and PP2. In the case of the annealed specimens, PP1A and PP2A, no distinctive necking behavior was detectable in the tensile mechanical behavior, while they show mild plastic deformation in the photograph (Figure 3). Note that PP2A specimen did not break up to a draw ratio of 5.6. Similar stress−strain behavior was reported in the previous studies.37 All the specimens show the occurrence of whitening as shown in Figure 3; this indicates that the cavitation in the 1400


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Figure 5. The WAXD images and their one-dimensional intensity profiles of PP1 at (1) the outermost region of skin layer, (2) the central part of the skin layer, and (3) the core layers. The definition of azimuthal angle is shown in the WAXD image of part 3. The positions of X-ray irradiation are shown by the white arrows in the POM images.

form crystal, those from the β-form crystals (300 and 301) were observed at the point 2. In the following, we focus on the distribution of the 110 (α-form) and the 300 (β-form) reflections. Figure 6 shows the distribution of integrated intensity of the 110 reflection and that of the 300 reflection. The integrated

intensity of diffraction can be used as the estimate of the amount of crystals. The 300 reflection was not observed in the annealed samples, PP1A and PP2A, while it was observed in PP1 and PP2 at the outer regions. The β-form crystals distributed in a wider region in PP2 than in PP1; in PP1, the 300 reflection from the β-form crystals was observed in the skin and the β-form rich layers in both of which no spherulite was observed in the POM images. In these layers, the crystals are highly oriented as shown in Figure 7, in which the azimuthal distribution of 110 reflection intensity is presented. From these results, it is concluded that the β-form crystals exist only in the layers composed of highly oriented structure in PP1. The situation is utterly different in PP2: the 300 reflection from the β-form crystals was observed in the skin, sheared, and finegrained layers. The fine-grained layers were composed of spherulites and there was no orientation characteristic to the injection molding found in 110 refleciton. Another feature of the polymorph’s distribution is that the β-form crystals are rich at around 300 and 2550 μm in PP1. In this region, the 110 reflection is hardly observed and only the 300 reflection is observed; thus, this layer is composed of only the β-form crystals. This peculiar layer is not observed in PP2. Distribution of Crystalline Orientation. Figures 7 and 8 show the distribution of azimuthal orientation of 110 and 040 reflections, respectively. At the skin layer (A) and the shear layer (B) of PP1 and PP1A, both the 110 and 040 reflections

Figure 6. Distribution of WAXD integrated intensity of 110 reflection of the α-form (open circles) and 300 reflection of the β-form (open triangles). (A) PP1, (B) PP1A, (C) PP2, and (D) PP2A. 1401


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Figure 7. Azimuthal distribution of the 110 reflection of the α-form crystals.

Figure 8. Azimuthal distribution of the 040 reflection of the α-form crystals.

are highly oriented, while their orientation at the core layer (D) is randomly distributed. The random distribution of the crystalline orientation originates from the fact that the core is composed of spherulites as shown in the POM images; microbeam X-ray scattering method observes the local structure of cross-hatched crystals of iPP, and the scattering attributing to the parent- and daughter-lamella is observed.45 At the β-phase rich layer (E) of PP1, there is no reflection from the α-form crystals. PP2 and PP2A show a similar behavior: at the skin and sheared layer, both the 110 and 040 reflections are highly oriented. At the fine-grained layer (C), 110 reflection is less oriented, or randomly distributed, whereas 040 reflection is oriented. At the core region (D), both the reflection shows random orientation. Distribution of Crystalline Ordering. The degree of crystalline ordering, which includes the packing order and the crystal size, is estimated from the fwhm of a diffraction peak by using the Scherrer’s formula. Each WAXD diffraction peak was fitted with a Gaussian function and the fwhm of each diffraction peak along the radial direction in the scattering image was calculated. The results are shown in Figure 9, in which the distribution of the crystalline ordering along ⟨110⟩ is clearly observed. When annealed, the crystalline ordering along ⟨110⟩ increases in all of the layers. The distribution of the crystalline ordering along ⟨110⟩ is qualitatively preserved after annealing; this indicates that no drastic structural change occurs during annealing.

Figure 9. (Upper) Distribution of the crystalline ordering along ⟨110⟩ and (lower) its expanded views: (left) PP1 (triangles) and PP1A (circles); (right) PP2 (triangles) and PP2A (circles).

At the outermost region of the skin layers other than that at 2800 μm of PP1 and PP1A, the crystalline ordering along ⟨110⟩ is small compared to those at the neighboring parts. This indicates that the amorphous structure formed at the outermost region of the surface. This idea is supported from the fact that the scattering from amorphous structure was observed in the 1402


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Figure 10. Typical SAXS images of PP1, PP1A, PP2, and PP2A. (A−E) Layers at which the X-ray microbeam irradiates.

Microbeam WAXD Results of Deformed Sample. Next we discuss the structural distribution after tensile testing. Hereafter we focus on the results of PP1 and PP2A because these two samples show utterly different mechanical behaviors. Figure 11 shows an example of image plots representing the

WAXD intensity profile at the outermost surface region in Figure 5. The discrepancy between the values of crystalline ordering along ⟨110⟩ at the both ends of PP1 and PP1A comes from the difficulty in accurately define the position of the surface. The size of microbeam (approximately 5 μm) equaled the width of the outermost region; it was thus difficult to elucidate the detailed structure of the outermost region. In relation to the mechanical properties, the existence of amorphous structure at the outermost region does not have a large influence on the difference of tensile mechanical properties observed in Figure 2, because all the samples show a similar tendency at the outermost region. Next to the outermost region, there are regions at which the crystalline ordering along ⟨110⟩ is large. Then, the crystalline ordering along ⟨110⟩ decreases and shows minimum values at approximately 200 and 2700 μm when we scanned the inner part of the skin region. We emphasize that the crystalline ordering along ⟨110⟩ at these regions is obviously minima for PP1 and PP1A, while it is in the same order with those in the inner regions for PP2 and PP2A. In other words, PP1 and PP1A have a region where the crystalline ordering along ⟨110⟩ is particularly small; they have inhomogeneity in the crystalline ordering along ⟨110⟩ compared to PP2 and PP2A. The crystalline ordering along ⟨110⟩ of PP1 and PP1A increased discontinuously from the innermost region of the shear layer to the outermost region of the core layer. Then, it kept a constant value in the core layer, although the values were sparsely dispersed. By contrast, the crystalline ordering along ⟨110⟩ of PP2 and PP2A gradually increased toward the central part of the core layer from the skin layer; in other words, the inhomogeneity in the crystalline ordering along ⟨110⟩ over the sample width of PP2 and PP2A is qualitatively less than PP1 and PP1A. Results of Scanning SAXS. Figure 10 shows typical SAXS images of PP1, PP1A, PP2, and PP2A at each layer. At the skin layer (A), the shear layer (B), the fine-grained layer (C), and the β-phase rich layer (E), a spot-like pattern along the flow direction was observed. In the core layer (D), a four-point pattern characteristic to the cross-hatched crystals of iPP45 was observed for PP1 and PP2. In the SAXS patterns at the skin, shear, and fine-grained layers of the annealed sample (PP1A and PP2A), the peaks originating from the lamellar long period existed in a lower angle; this indicates that the annealing increased a lamellar long period. At the core of the annealed sample, no distinctive pattern originating from the existence of lamellae is observed in the present SAXS resolution.

Figure 11. Distribution of WAXD intensity of PP1 at (A) 5 mm apart from the fracture point (Figure 12A), (B) the necking area (Figure 12B), (C) the fracture area (Figure 12C), and (D) for the sample without deformation.

intensity distribution of WAXD of deformed PP1 at different scanning positions. The POM images of scanning positions are shown in Figure 12. At the position where the necking did not develop (Figure 11A), the distribution of morphology was similar to that of virgin specimen (Figure 11D); the diffraction from the β-form crystals was observed only at the outer regions. At the necking point (Figure 11B), the diffraction from the βform crystals was observed only at the upper region, which agrees with the characteristics observed in the POM image (Figure 12B); along the scanning position of X-rays, the sample was highly stretched and yielded at the lower region, while not 1403


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Figure 12. POM images of drawing specimen of PP1 (A) 5 mm apart from the fracture point, (B) necking region, and (C) the fracture point. The white arrows in the POM images indicate the positions of irradiation with an X-ray microbeam. The crossed marks indicated (a− k) show the positions of irradiation of microbeam SAXS experiments (Figure 15).

at the upper region. At the fracture area (Figure 11C), the fwhm of diffraction from the α-form crystals clearly increased, and the diffraction from the β-form crystals was no longer observed. It can be concluded that the disappearance of the diffraction from the β-form crystals and the disordering of the α-form crystals relate to the necking and breaking-up behaviors, respectively. The situation is, to the contrary, different for PP2A as shown in Figures 13 and 14. In PP2A, it is hard to distinguish the differences between the image plots of the WAXD intensity distribution before (Figure 13C) and after (Figure 13A,B) deformation. Also, a clear skin−core structure can be observed in the POM image of the deformed sample (Figure 14). Microbeam SAXS Results of Deformed Sample. Typical SAXS images of deformed specimen of PP1 and PP2A are shown in Figure 15 and 16, respectively. At the positions where the necking did not develop (Figure 15i−k), the SAXS patterns were similar to those without deformation; a two-spot pattern at the outer region and a cross-hatched pattern at the inner region. At the necking point (Figure 15e− h), remarkable changes in the SAXS pattern were observed. First, the peaks originating from the lamellar structure disappeared or weakened. Secondary, an isotropic central scattering around the beamstop drastically increased for the upper part of specimen (Figure 15e,f), while it was not observed at the lower part of specimen (Figure 15g,h). This central scattering should correspond to the formation of voids, which agrees with the whitening observed in the photograph of deformed specimen (Figure 3). At the breakup point (Figure 15a−d), further changes were observed; at each position, scattering from the lamellar structure is observed while the orientated central scattering was also observed.

Figure 13. Distribution of WAXD intensity of PP2A at (A) 2 mm apart from the fracture point (Figure 14A), (B) the fracture point (Figure 14B), and (C) the sample without deformation.

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DISCUSSION The structural characteristics obtained in the present study can be summarized as follows: (1) The annealing increases the size of α-form crystals irrespective of molding condition and layers. (2) The β-form crystals exist in the skin (PP1 and PP2), shear (PP1 and PP2), and fine-grained layers (PP2), while they disappear when annealed. (3) The distribution of crystalline ordering along ⟨110⟩ in PP2 and PP2A is rather homogeneous compared to those in PP1 and PP1A; PP1 and PP1A have a region where the crystalline ordering is particularly small. (4) The β-form crystals disappear at necking region when deformed. (5) All the deformed samples show whitening. The most striking differences in the tensile mechanical behavior are found in the necking behaviors. From the tensile testing, PP2A exhibits no distinctive necking and shows tough and strong properties compared to the other samples, PP1, PP1A, and PP2. In this section, the relationship between the tensile mechanical properties and the inhomogeneity of both the crystalline ordering and the polymorph in micrometer scale is discussed. The annealed specimen, PP1A and PP2A, show homogeneous deformation rather than strain localization (necking) observed for PP1 and PP2. Generally, there can be two different origins for the transition from necking to a more homogeneous mode of deformation. First, the deformation will become homogeneous under the condition that the deforma1404


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Figure 14. POM images of drawn specimen of PP2A. (A) 2 mm apart from the fracture point, and (B) the fracture point. The white arrows in the POM images indicate the positions of irradiation with an X-ray microbeam. The crossed marks indicated by parts a−d show the positions of irradiation of microbeam SAXS experiments (Figure 16).

Figure 15. Typical SAXS images of drawn specimen of PP1. The positions of irradiation are shown in Figure 12. The horizontal direction along qx is the drawing direction.

tion proceeds under a constant volume.52−54 The second possibility is the occurrence of stress whitening. In this case, deformation does not proceed under a constant volume since the amorphous regions cavitate. Because of this formation of voids, microscopic strain localization will take place and can proceed on the micrometer-scale, and macroscopically the deformation is fairly homogeneous.55−58 The crystallites work as the point of action for tensile stress, and the necking stress depends on the crystallinity.59 Usually, the deformation under a constant volume is observed in highly oriented samples,47 which is not the case for our samples in which nonoriented spherulites form the core part. Furthermore, strong whitening was observed for the annealed specimen, PP1A and PP2A; thus it can be concluded that the occurrence of stress whitening is the origin of transform from the necking behavior to rather homogeneous deformation behavior. Unfortunately, the SAXS data of deformed specimen of PP2A shows little information about the void formation because of the present SAXS resolution, which is limited by the existence of Fraunhofer diffraction from the first pinhole producing the intense microbeam X-rays. With regard to the structural characteristics leading to the void formation, there are two features: the disappearnace of the β-form crystals in the annealed sample and the increase in the crystalline ordering by the annealing.70,72 As shown in Figure 9, the crystalline ordering along ⟨110⟩ increased when annealed, and at the same time, the annealing induces the disappearance of the β-form crystals; thus, it is difficult to discuss that either of them is a principal factor determining the stress whitening phenomenon. In this regard, it is interesting to see combined

Figure 16. Typical SAXS images of drawn specimen of PP2A. The positions of irradiation are shown in Figure 14. The horizontal direction along qx is the drawing direction.

results of SAXS and WAXD of deformed specimen of PP1. At the necking area (Figure 12B), WAXD results show that the βform crystals disappeared at the lower region where the sample was highly deformed while they existed at the upper region (Figure 11B). On the other hand, an isotropic central scattering around beamstop in SAXS was observed at the upper region (Figure 15e,f) while it was hardly observed at the lower region (Figure 15h). This indicates the possibility that the existence of β-form crystals prevents the voids formation and growth; the βform crystals yield at lower stress71 and thus the amorphous regions cannot cavitate. This discussion will be confirmed by 1405


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crystals (for example, 2 3̅ 1 and 1 6̅ 1 at q ∼ 2.24 Å−1) was observed63,64 in both PP2 and PP2A; thus, the increase in the mechanical properties can be attributed to the absence of structural inhomogeneity. In this study, the relationship is indicated between the inhomogeneous distribution of polymorph/crystalline ordering and the mechanical properties of injection-molded iPP. To elucidate the relationships further, a real-time monitoring of the structural change during sample deformation45,65,66 is a powerful method. Although it is difficult to map the structural distribution during sample deformation, in situ observation of structural change at a single layer will be possible with a recent progress in the X-ray microbeam scattering technique. Also, in situ observation of structural change at each layers during the annealing of injection-molded specimen will give us a clue for understanding the nature of structural change during the annealing.74 Furthermore, the combination of microbeam X-ray scattering technique with a WAXS pole-figure technique67 will be important to quantify the amount of each crystalline phases and amorphous phases at each layer. These studies are now in progress and will be reported in the near future.

further studies with in situ observation of structural change using microbeam X-ray diffraction. The SAXS images of the fractured area (Figure 15a−d) show the anisotropic oriented central scattering pattern with peaks originating from lamellar structure. This central scattering would originate from either of the voids or highly oriented fibril structure. Although there are previous reports in which the voids can be oriented in the deformed sample,73 it is reasonable to suppose that fibril structure like shish-kebab structure is formed at this region,69 because the orientation of the central scattering corresponds to the orientation of lamellar scattering. This also supports the above discussion that the formation of voids in the annealed sample leads to the transformation from necking to rather homogeneous deformation modes; macroscopically homogeneous deformation cannot proceed in PP1 because of the absence of voids as the source of microscopic strain localization. The disappearance of the β-form crystals at the fracture areas and the lower parts of necking areas (Figure 11B,C) indicates a mechanically induced β- to α-form crystals transition during deformation.14,60−62 The fact that the β-form crystals are absent only at the necked region supports the idea that the β-form crystals are stable to the yield point and transform to the αform crystals only during necking.62 It should be noted that the draw ratio of fracture of PP2A is considerably long. Compared to PP1A, the crystalline ordering of PP2A is in the same order but has less inhomogeneity; the crystalline ordering along ⟨110⟩ of PP1A obviously shows minima at 200 and 2700 μm while that of PP2A shows no clear minimum. It is fairly reasonable to assume that the region where the crystalline ordering along ⟨110⟩ shows minima relates to the breakdown. One may attribute the increase in the mechanical properties to the existence of highly ordered crystallites such as α2-form of iPP. To clarify this point, the WAXD intensity profiles of PP2 and PP2A are shown in Figure 17. No diffraction from the α2-

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CONCLUSION We aimed to reveal the relationship between micrometer-scale inhomogeneous distribution of subnm- and nm-scale structure of injection-molded iPP and its tensile mechanical properties. We successfully employed scanning microbeam X-ray scattering techniques to reveal the micrometer-scale inhomogeneous structural distribution. The relationship between structural inhomogeneity and tensile mechanical behavior was revealed for four types of injection-molded specimen. The annealing process leads to the increase in the crystalline ordering and the disappearance of β-form crystals, which would lead to void formation during tensile deformation. The results indicate that the inhomogeneous distribution of polymorph and crystalline ordering is related to the appearance of necking and the breakdown at relatively low strain. Further studies with in situ observation of structural evolution process during sample stretching process45,65,66 will provide us fruitful information regarding the detailed mechanism of structural change upon elongation.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected].

ACKNOWLEDGMENTS Microbeam WAXD experiments were performed under the approval of Photon Factory Proposal Advisory Committee (2008G700, 2010G540). Microbeam SAXS experiments were performed under the approval of SPring-8 Proposal Advisory Committee (2009B1260). We thank Drs. N. Ohta and N. Yagi (JASRI/SPring-8) for their kind help on the microbeam SAXS experiments and Prof. A. Iida (KEK) for his assistance on the microbeam WAXD experiments.

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REFERENCES

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Figure 17. One-dimensional WAXD intensity profiles of PP2 (upper) and PP2A (lower). 1406


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