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Effect of molding temperature on crystalline and phase morphologies of HDPE composites containing PP nano-fibers
Polymer 45 (2004) 5719–5727 www.elsevier.com/locate/polymer
Effect of molding temperature on crystalline and phase morphologies of HDPE composites containing PP nano-fibers Jianxiong Lia,1, Qiang Wanga, Chi-Ming Chana,*, Jingshen Wub a
Department of Chemical Engineering, The Hong Kong University of Science and Technology, Academic Building, Clear Water Bay, Kowloon, Hong Kong b Department of Mechanical Engineering, The Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong Received 7 January 2004; received in revised form 29 April 2004; accepted 24 May 2004 Available online 19 June 2004
Abstract A high-density polyethylene (HDPE)/isotactic polypropylene (PP) (75/25) blend containing 25 wt% of PP was fibrillated by roller drawing at 138 8C. The fibrillated blend was processed again at temperatures ranging from 155 to 200 8C by compression molding or extrusion. The effects of molding temperature on the morphology and mechanical properties of the blend were investigated. Wide angle X-ray scattering (WAXS) and transmission electron microscopy (TEM) were used to study the morphology of the samples. The roller-drawn blend exhibited a fibrous structure with the chain direction aligned parallel to the drawing direction. After molding at 155 8C, the HDPE formed parallelstacked lamellae retaining the parallel orientation after the melting of the PE crystals. As the molding temperature increased the parallel orientation gradually vanished and some of the parallel-stacked lamellae changed into twisted lamellae. The PP phase existed as fibrils in the PE matrix and the crystals stayed with their molecular chain aligned parallel to the fibrillation direction even when the molding temperature was far above the melting temperature of PP. Nevertheless, the orientation of the crystals did not change as the molding temperature increased from 155 to 165 8C. The internal structure of the PP fibrils changed from a needle structure to a parallel-stacked one. The PP fibrils induced the crystallization of the PE melt, leading to the formation of a trans-crystalline layer at their surface. As the molding temperature increased, more PE lamellae protruded into the PP fibrils and the interface between the PP fibrils and the PE matrix became diffuse. q 2004 Elsevier Ltd. All rights reserved. Keywords: Nanocomposite; Morphology; Roller drawing
1. Introduction It has been widely established that the mechanical properties of crystalline polymeric materials strongly depend on processing conditions and techniques used to process the materials. The same polymeric material can be processed into a soft and flexible product or a strong and stiff product under different conditions. Many scientists and engineers [1,2] have devoted much effort to developing stronger polymeric materials. These practices include fiber reinforcement [3] and orientation reinforcement [2,4,5]. In the latter case the polymer chains are aligned by means of stretching the polymers below and close to their melting points, such means include solid extrusion [4,5] and die/ roller drawing [6,7]. There have been many studies on the reinforcement of * Corresponding author. Tel./fax: þ 852-2358-7125. E-mail address: [email protected] (C.M. Chan). 1 Present address: ASM Assembly Automation Ltd., 4/F, Watson Center, 16 Kung Yip St., Kwai Chung, Hong Kong. 0032-3861/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2004.05.049
crystalline plastics with liquid crystalline polymer (LCP) fibers generated in situ [8]. Under specific conditions, polymer/LCP blends are injected into a mold; the LCP aligns in the flow direction and forms LCP fibers to reinforce the matrix. Also, there are reports on the conversion of melt-spun fibers into composites [9,10]. Highly aligned fibers are partially melted and compression molded. In this way, the fibers are joined together by melting the skin layer, forming a composite. Three basic morphologies of crystals have been observed in such oriented polymers [1]: fringed micelles, parallelstacked lamellae and fine needle crystals. Normally, fringed micelles appear in a sample stretched at a temperature below its glass transition temperature [11]. An example of fringed micelles is strain-induced crystallization of PET at room temperature. Stacked lamellae occur when the sample is drawn from a melt and then crystallizes [12]. Needle crystals are generated during large plastic deformation of a crystalline polymer at a temperature between the polymer’s glass transition temperature and its melting point [13,14]. Of
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course, there have been samples generated that show some combination of the above basic morphologies, such as the so-called shish-kebab structure [14,15], which consists of parallel-stacked lamellae and needle crystals. The crystalline structures in oriented polymers are believed to be metastable. The structure of the crystalline phase may change, depending on the temperature and time. For example, upon annealing, defects in needle crystals could agglomerate. Periodic arrangements of the defect-free and highly defective zones along the needle may appear. With further annealing at an elevated temperature, the needle crystals may be transformed into a shish-kebab structure and finally into a parallel-stacked structure [14,15]. Recently, a new method to prepare thermoplastic fibrilreinforced composites has been investigated in this laboratory [16]. Two crystalline polymers, commercial high-density polyethylene (HDPE) and isotactic polypropylene (PP), were first melt-blended together and then fibrillated through roller drawing. Finally, the fibrillated HDPE/PP blend was melt-processed again at temperatures above the melting point of HDPE. Preliminary results indicated that PP fibrils with a diameter of about 100 nm could survive after molding at temperatures up to 185 8C. The modulus and tensile strength in the fibrillation direction were increased by more than three times. Meanwhile, the modulus and tensile strength in the transverse direction increased by more than 20%. Our results also indicated that the mechanical properties of the sample varied with molding temperature. The finding that the strength in the transverse direction increased with increasing molding temperature and reached a maximum at 185 8C is interesting. In the present work, the effects of molding temperature on the morphology and mechanical properties of the HDPE/ PP nanocomposites will be examined.
2. Experimental 2.1. Preparation of the fibrillated HDPE/PP blend HDPE (Philips HMMPE) was blended with 25 wt% of isotactic polypropylene (Himont 6501) using a Haake TW 100 twin screw extruder. The HDPE/PP (75/25) pellets were extruded into 2.4 mm thick tapes and fibrillated by roller drawing at 138 8C [17]. After fibrillation, the thickness of the tapes was reduced to 0.3 mm and the drawing ratio was calculated to be approximately 10. A thermal analysis of the tapes [17] indicated that the PE component has a melting point at 139 8C and the fusion of the PE crystals is completed below 150 8C, while the fusion of the PP crystals, with a melting point of 167 8C, started at a temperature lower than 155 8C and finished below 175 8C.
laid parallel with respect to the fibrillation direction in a hot mold. The hot mold with the fibrillated sheets were first preheated in a press for 3 min under a pressure of about 3 MPa; then, the pressure was increased to 10 MPa and released immediately to degas; finally they were molded into 1.5 mm plates under a pressure of about 15 MPa. After molding for 25 min, the resulting plates were cooled first with compressed air from the pre-set molding temperature to below 120 8C (about 1 h) and then with tap water to below 50 8C under pressure. In the present work, five molding temperatures, 155, 165, 175, 185 and 200 8C, were selected. In the molding process, all PE crystals were melted while the partial melting of PP was changed to complete melting as the molding temperature increased. 2.3. Identification of the crystalline phase and orientation A powder X-ray diffraction system, Philips PW 1830, was utilized to identify the crystals and their orientations. The Cu anode X-ray generator was operated at 40 kV and 40 mA. The Cu Ka radiation impinged on the specimens with the beam perpendicular to the fibrillation direction. The diffraction experiment was conducted in the transmission mode and the output signal was recorded on a flat film at a diffraction distance of 90 mm. 2.4. Examination of morphology Thin strips with a cross-section of about 0.2 £ 0.3 mm2 were cut from the molded samples perpendicular to the fibrillation (or extrusion) direction. They were embedded in an epoxy resin and cured at 45 8C for 48 h. The embedded specimens were trimmed first with razor blades and then with glass knives on a Reichert-Jung Ultracut R microtome. The trimmed specimens were exposed to the vapor of an aqueous solution of ruthenium tetroxide [18,19] in a sealed test tube for 48 h. After staining with ruthenium tetroxide, the specimens were washed with a 3% aqueous solution of sodium periodate and distilled water. Then the samples were dried in a desiccators for 24 h. On a Reichert-Jung R microtome, ultrathin sections with a thickness about 60 nm were cut from the stained specimens with a diamond knife after a top layer of approximately 400 nm was removed. The sections were mounted on 200-mesh copper grids. After drying in a desiccator for 24 h, they were examined with a JEOL JEM-100 CX II transmission electron microscope (TEM), which was operated at an accelerating voltage of 80 kV.
3. Results 3.1. Crystalline phase and orientation
2.2. Preparation of HDPE/PP nanocomposites The fibrillated tapes were cut into short pieces and were
Fig. 1 shows the X-ray diffraction pattern of the fibrillated HDPE/PP blend. In Fig. 1, the fibrillation
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Fig. 1. The X-ray diffraction pattern of the HDPE/PP (75/25) blend after fibrillation by roller-drawing at 138 8C.
direction is horizontal. One can see that the azimuths of the diffractions are very small and they appear as spots, indicating a high degree of orientation of the crystals in the sample. The diffraction pattern is similar to that of PP fibers except three diffraction spots are from HDPE. Twelve diffraction spots align in the vertical direction and four more diffraction spots appear about 408 away from the meridian (vertical line). Beginning from the inner one, the diffraction spots are named Spots 1 to 6 and the one that is off the meridian is named Spot 7. The diffraction radii of the spots were measured and the diffraction angle, 2u; and the d-space of the reflecting planes were calculated. The results are summarized in Table 1. Also, the values of the d-spaces reported in the literature [20 –23] are listed in the last row of the table. Compared with the data from the literature, Spots 1 to 3 as well as Spot 7 are recognized to come from the monoclinic PP crystals (a) and the corresponding crystallographic planes are a(110), a(040), a(130) and a(111), respectively. Spot 4 is a reflection from (001) of the monoclinic PE crystals (m) while Spots 5 and 6 are the reflections from (110) and (200) of the orthorhombic PE crystals (o). However, it should be pointed out that some reflections of the monoclinic PP crystals—the primary diffraction of a(131) and the secondary diffraction of a(110)) can contribute to the intensity of Spot 5. But due to the low PP content and the weak diffraction intensity of the
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primary diffraction of a(131) and the secondary diffraction of a(110) compared with that of a(130), the crystallographic plane (110) of the orthorhombic PE crystals is considered to be the main source for Spot 5. Fig. 1 indicates that three crystalline phases exist in the fibrillated HDPE/PP blend—the monoclinic PP crystals and both the monoclinic and orthorhombic PE crystals. The monoclinic PP crystals align with their c-axes along the fibrillation direction as do the orthorhombic PE crystals. However, for the monoclinic PE crystals, the c-axis is perpendicular to the fibrillation direction. X-ray diffraction patterns of the HDPE/PP nanocomposites prepared by molding the HDPE/PP blend at temperatures from 155 to 185 8C are shown in Fig. 2. After molding, the radius of the diffraction spots for the PP and PE crystals stays almost the same. However, a new diffraction spot (Spot X) is observed near the reflection of a(110) (Spot 1). Based on the diffraction radius of the new spot, the ˚ . In PP, corresponding d-space was calculated to be 6.8 A there are three possible crystalline phases—the monoclinic (a), hexagonal (b) and triclinic (g) phases [20,21]. Normally, the monoclinic phase predominates in the samples. However, the hexagonal and triclinic phases may exist under special conditions, such as high pressure in the triclinic phase [23] and a special nucleating agent [24] in the hexagonal phase. Of these crystals, the g(100) has the ˚ as reported by Jones [20]. largest d-space of 6.37 A However, it is still much smaller than the value for this new reflection plane. Up to now, no report on co-crystallization of HDPE and PP has been found in the literature. When highly orientated thin PE and PP films were molded at temperatures between their melting points only epitaxially crystallized PE was observed [25]. The new diffraction spot is unlikely to be from a new crystalline phase. It has been reported that the unit cell constants of polymer crystals may change under stress, so the d-space of the reflecting plane may change [26]. The new diffraction spots may result from the stretched a(110) plane. However, if this is true, the diffraction intensity of the stretched a(110) should be stronger before molding. Because heating should result in the relaxation of stretched crystals, the intensity should decrease as the molding temperature increases. The fact is that the diffraction signal around the diffraction angle ð2uÞ; 12.998, is vague in the fibrillated blend (cf. Fig. 1). Only after molding does it become more obvious (cf. Fig. 2).
Table 1 The diffraction angle and corresponding crystallographic plane of the fibrillated HDPE/PP blend Sport Intensity Diameter [mm] 2u ˚) d-space (A Reflecting plane ˚) d-spacea (A a
N Weak 41.5 12.99 6.8 g(100)? 6.37
1 Middle 45.5 14.36 6.30 a(110) 6.26
Corresponding d-space from the literature.
2 Middle 55.5 16.96 5.22 a(040) 5.19
3 Weak 61.5 18.67 4.75 a(130) 4.77
4 Strong 63.5 19.51 4.54 m(001) 4.55
5 Very strong 72.0 21.58 4.11 o(110) 4.10
6 Middle 81.0 23.99 3.70 o(200) 3.69
7 Weak 70.5 21.17 4.19 a(111) 4.19
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Fig. 2. The X-ray diffraction pattern of the fibrillated HDPE/PP blend after molding at different temperatures: (a): 155 8C; (b): 165 8C; (c): 175 8C and (d): 185 8C.
Nevertheless, as the molding temperature was increased to 200 8C, the diffraction signal around 12.998 could not be identified (cf. Fig. 3). At this moment, it seems to be difficult to determine the source of the new diffraction sport. Fig. 2 shows a significant change on the diffraction patterns of m(001), o(110) and o(200) from the PE crystals. After molding at 155 8C, the diffractions of m(001), o(110) and o(200) changed from spots to circle rings (cf. Fig. 2 (a)). However, the diffraction intensities of the rings are not uniform. In the meridian regions, the intensities are much stronger. This indicates that the c-axis of the PE crystals still had the preferred orientation. But, as the molding temperature increased from 155 to 185 8C, the diffraction rings became more and more uniform (cf. diffractions of m(001) and o(200) in Fig. 2). The orientation of the PE crystals gradually became random. In contrast to the PE crystals, the azimuths of the diffractions of the PP crystals did not change much after molding at 155 8C. The PP crystals still aligned with the
Fig. 3. The X-ray diffraction pattern of fibrillated HDPE/PP blend after molding at 200 8C.
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c-axis towards the fibrillation direction. This is easy to understand because at this molding temperature, the PP crystals had not melted. Even when the molding temperature increased from 155 to 185 8C, the azimuthal extension of the diffractions did not increase much. This implies that the PP crystals still aligned with the c-axis along the fibrillation direction even after the sample was heated at 185 8C for 25 min. Only when the molding temperature increased to 200 8C, did the PP crystals lose their preferred orientation and did the sample give five uniform diffraction rings (cf. Fig. 3). 3.2. Morphology and interphase structure Fig. 4 shows the TEM micrographs of the sample prepared by molding a fibrillated HDPE/PP blend at 155 8C. In this sample, the PP component appears as straight fibrils in the PE matrix. The size of the fibrils ranges from 30 to 150 nm. In the PE matrix, the crystals are in a lamellar structure and have a thickness of approximately 15 nm. Instead of developing into spherulites, the lamellae form a highly ordered arrangement between the PP fibrils. They stack together almost parallel to each other and are aligned approximately perpendicular to the long axial of the fibrils, forming a transcrystalline interphase [27]. The coarse dark bands between the lamellae are believed to be the lessordered materials as they absorbed more staining agents. This is very similar to the observations made by Li et al. that a transcrystalline interphase was obtained between the isotactic PP matrix and poly(ethylene terephthalate) microfibers after a slit-extrusion, hot-stretching and quenching of a blend containing these two polymers [28]. The PP fibrils contain finer crystalline structures. At a high magnification (cf. Figs. 4(c) and 5(a)), it can be seen that there are some needle-shaped, fine crystals (marked with N). The needles are about 3 nm in width and tens of nm in length. Normally, they align in the fibril direction. However, a few fine crystals perpendicular to the fibril direction can be observed as well (marked with S in Fig. 5 (a)). Thus, the PP fibrils exhibit some features of a shishkebab structure. The surface of the PP fibrils is rough and the roughness of the fibrils is in the order of ten nm (marked with R in Fig. 4 (c)). Nevertheless, the boundary between the PE matrix and the PP fibrils appears to be rather sharp. Fig. 6 shows the TEM micrographs of a sample molded at 165 8C. The PE crystals similarly remain in their stackedlamellar structure, although the increase in the molding temperature led to a decrease in the thickness of the PE lamellae, from about 15 nm to about 10 nm. Additionally, a few twisted PE lamellae can be found in the PE phase. Furthermore, the PE lamellae are no longer aligned perpendicularly to the PP fibrils (cf. Fig. 6). The internal structures of the PP fibrils changed significantly as the molding temperature increased to 165 8C. No needle crystals are found; instead, some stacked lamellae are observed. The thickness of these PP lamellae varies from 4
Fig. 4. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 155 8C. B: boundary between the PP fibril and PE matrix; N: needle crystals; S: shish-kebab structure.
to 10 nm. The surface of the PP fibrils is still rough and the boundary between the PP fibrils and the PE matrix is still sharp. However, some PE lamellae are found to be occasionally sticking into the PP fibrils (marked with P in Fig. 6 (a)). Fig. 7 shows the morphologies of the sample molded at 185 8C. When the molding temperature increased from 165 to 185 8C, the thickness of the PE lamellae seemed to increase a little. Near the PP fibrils, the PE lamellae show a stacked pattern. However, between the PP fibrils, many twisted PE lamellae can be seen, just as the lamellar
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Fig. 6. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 165 8C. P: parallel stacked PP lamellae. Fig. 5. The fine structure of the PP fibrils of the HDPE/PP composite molded at 155 8C: (a) an enlarged view of an area of Fig. 1 and (b) PE lamellae and interface of between the PE matrix and PP fibrils.
structure observed in the sample crystallized from the PE melts without disturbance [29,30]. In contrast to the marked structural changes that occurred between 155 and 165 8C, there were no large changes in the internal crystalline structure of the PP fibrils when the molding temperature increased from 165 to 185 8C. The crystals are still in a stacked-lamellar structure. However, a striking change is observed at the interface between the fibrils and matrix. Some PE lamellae extended from the PE phase and penetrated into the PP fibrils. The boundary became rather diffuse. At a high magnification (Fig. 7 (b)), the surface of the PP fibrils is seen to have a saw-tooth structure with a period of 10 nm. However, it should be pointed out that there are more defects or disordered materials at the interface between the PP fibrils and the PE matrix because the interface always appears darker than other regions in the examined sections. Fig. 8 (a) shows a TEM micrograph of the sample molded at 200 8C. As the molding temperature was further increased to 200 8C, the PP component no longer existed as
fibrils but rather as particles, dispersing throughout the PE matrix. The size of the PP particles ranges from 0.5 to 2 mm. This morphology is very similar to that of the extruded sample (cf. Fig. 8 (b)). It should be pointed out that the interfacial bonding is weak, as evidenced by the voids in the dark rings around the PP particles. This observation is also supported by the mechanical property data indicating that the strong and ductile composite became a weak and brittle material as the molding temperature increased to 200 8C [11,17]. Fig. 9 shows the TEM micrographs of the sample prepared by extrusion of the fibrillated HDPE/PP at temperatures below 160 8C. On the stained section of this sample, PP fibrils can be observed (ref. Fig. 9 (a)). Beside the long PP fibrils, many circular and elliptical particles are observed (marked with C and E in Fig. 9 (a)). These circular and elliptical particles are believed to be the cross-sections of the PP fibrils due to the fact that the cutting plane was perpendicular or at an inclined angle to the fibril. This implies that the PP fibrils have a random orientation in the extruded sample. The size of the fibrils ranges from 50 to 300 nm, which is a little larger than those in the compression-molded samples. At a higher magnification
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Fig. 7. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 185 8C.
(cf. Fig. 9 (b)), a fine crystalline structure, similar to the structures observed in the sample molded at 155 8C, can be found in the PE matrix and PP fibrils. However, the lamellae appear to be finer. The PE lamellae are about 8 nm thick and the size of the PP crystals is about 2 nm.
4. Discussion The orthorhombic PE phase occurs most frequently in a normal, melt-crystallized sample. However, when a sample is subjected to stress or deformation, monoclinic PE crystals may be generated in the sample [22,23]. In the fibrillation process of the present work, the PE crystals were definitely stretched in the melt because of the high temperature of the rollers and the effect of the drawing [31]. Therefore, in the fibrillated HDPE/PP sample, some monoclinic PE crystals are present. It is evident that after molding at 155 8C, below which the fusion of PE was completed as measured by differential scanning calorimetry (DSC) [17], the diffraction spots of the PE crystals changed to circular rings because of the loss of the orientation induced by the fibrillation. However, the intensity of the reflections from m(001), o(110) and o(200)
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Fig. 8. The morphology of the HDPE/PP (75/25) blend; (a): prepared by compression-molding the fibrillated blend at 200 8C and (b): extruded tape before fibrillation.
appeared to be much stronger in the meridian (cf. Fig. 2). This implies that the PE crystals did not lose their orientation completely after molding at 155 8C for 25 min. The preferred orientation of the PE crystals can be explained by the incomplete randomization of the PE chains (molecular shape memory [32]) in the molding process. During the molding at 155 8C, the PE crystals disintegrated but many PE molecular chains still aligned in the fibrillation direction. Consequently, cooling caused the alignment of the chain axis of the PE molecules along the fibrillation direction. Nevertheless, the intensity of the spots decreased gradually as the molding temperature increased. The diffraction intensity became uniform in all directions at 185 8C, indicating that a complete randomization of PE chain coils had been achieved. This is also supported by the matrix morphology of the sample molded at 185 8C showing many twisted lamellae between the PP fibrils. DSC experiments [17] showed that the melting of the PP crystals in the fibrillated blend began at 155 8C and finished below 175 8C, with a melting peak at 167 8C. The molding temperature of 155 8C is lower than the melting point of the PP component, which should still exist as fibrils. Actually,
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Fig. 9. TEM micrographs of the HDPE/PP (75/25) composite prepared by extrusion of the fibrillated blend at 155 8C; (a): low magnification and (b) high magnification.
the molding at 155 8C can more appropriately be considered as annealing the PP fibrils at this temperature. Therefore, the diffraction azimuths (orientation) of the PP crystals did not change much. However, the defects in the original needle crystals might aggregate, resulting in the formation of needle crystals with periodic arrangements of defect-free and highly defective zones in the PP fibrils [15,16]. As an intermediate stage in the transformation of the needle crystals into the parallel-stacked structure during annealing, a shish-kebab structure was generated. As the molding temperature increased to above 165 8C, most of PP crystals should have disintegrated although the overall orientation of the PP chains did not change greatly within that period of time. Hence, the PP crystals changed from the needle to parallel-stacked structure while the diffraction azimuths (overall orientation) of the PP crystals remained almost the same. Even when the molding temperature was raised above 175 8C and the PP crystals disintegrated completely, the complete randomization of the
stretched polymer chains through the thermal motion of the chain segments needed a period of time and might not have been achieved within 25 min at molding temperatures below 185 8C. The long-range alignment of the molecular chains remained in the PP phase. As a result, the PP molecules were transformed into parallel-stacked lamellae with the chain direction aligned along the fibrillation direction on cooling, similar to the crystallization of a drawn polymer melt [13]. Below 185 8C, the contracting force of the stretched polymer chains was not strong enough to overcome the frictional force at the interface between the PP fibrils and the PE matrix, allowing the fibrils to shrink back. As a result, the PP phase stayed as fibrils. A similar phenomenon has been observed in multi-layer composites of polybutene-1 (PB-1) and PP thin films [25]. The orientation of the PB-1 crystals remained even when the composites had been heated to a temperature above the melting point of PB-1, provided that the PB-1 film was thin enough. When the molding temperature was further elevated to 200 8C, the stretched polymer chains possessed enough thermal energy to overcome the friction at the interface, forming random coils. The PP fibrils shrank into particles; consequently, an isotropic diffraction pattern was obtained. In our previous paper [17], we have shown that the mechanical properties of the sample varied with the molding temperature. The strength in the transverse direction increased with increasing molding temperature and reached a maximum at 185 8C. As the molding temperature increased, the degree of the thermal motion of the chain segments increased. The inter-diffusion of the molecules increased at the interface; consequently, more PE lamellae penetrated into the PP fibrils. The change in the structure of the interface between the PP fibrils and the PE matrix is the explanation for the increase in the tensile strength in the transverse direction of the composites with increasing molding temperature to 185 8C.
5. Conclusion The fibrillated HDPE/PP blend, which exhibits a fibrous structure, consists of at least three crystalline phases: the monoclinic PP phase, and the orthorhombic and monoclinic PE phases. The monoclinic PP crystals and orthorhombic PE crystals align with their molecular chain (c-axis) along the fibrillation direction while the monoclinic PE crystals orient with their c-axis perpendicular to the fibrillation direction. It will take a rather long period of time for a stretched polymer chain to return to a random molecular coil even after disintegration above its melting point. After molding at 155 8C, the PE matrix exhibits a parallel-stacked lamellar structure and still retains this orientation after fibrillation. However, the preferred orientation of the PE gradually vanishes as the molding temperature increases, the parallel-stacked lamellae change into twisted lamellae
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and the diffraction pattern becomes uniform rings as the molding temperature increases to 185 8C. The PP phase exists as fibrils in the PE matrix and the crystals retain their molecular chains in the fibrillation direction when the molding temperature is below 185 8C. Although the orientation of the crystals does not change as the molding temperature increases from 155 to 185 8C, the internal structure of the PP fibrils changes from a needle to parallel-stacked structure. As the molding temperature increases to 200 8C, a complete randomization of the PP chains is achieved and the PP phase contracts into particles, losing all previous orientations. The PP fibrils induce crystallization of the PE melt, leading to a trans-crystalline layer surrounding the PP fibrils. As molding temperature increases, more PE lamellae protrude into PP fibrils, producing a diffused interface between the PP fibrils and the PE matrix.
Acknowledgements This work was supported by the Hong Kong Government Research Grants Council under grant No. HKUST 6043/01P.
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Effect of molding temperature on crystalline and phase morphologies of HDPE composites containing PP nano-fibers Jianxiong Lia,1, Qiang Wanga, Chi-Ming Chana,*, Jingshen Wub a
Department of Chemical Engineering, The Hong Kong University of Science and Technology, Academic Building, Clear Water Bay, Kowloon, Hong Kong b Department of Mechanical Engineering, The Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong Received 7 January 2004; received in revised form 29 April 2004; accepted 24 May 2004 Available online 19 June 2004
Abstract A high-density polyethylene (HDPE)/isotactic polypropylene (PP) (75/25) blend containing 25 wt% of PP was fibrillated by roller drawing at 138 8C. The fibrillated blend was processed again at temperatures ranging from 155 to 200 8C by compression molding or extrusion. The effects of molding temperature on the morphology and mechanical properties of the blend were investigated. Wide angle X-ray scattering (WAXS) and transmission electron microscopy (TEM) were used to study the morphology of the samples. The roller-drawn blend exhibited a fibrous structure with the chain direction aligned parallel to the drawing direction. After molding at 155 8C, the HDPE formed parallelstacked lamellae retaining the parallel orientation after the melting of the PE crystals. As the molding temperature increased the parallel orientation gradually vanished and some of the parallel-stacked lamellae changed into twisted lamellae. The PP phase existed as fibrils in the PE matrix and the crystals stayed with their molecular chain aligned parallel to the fibrillation direction even when the molding temperature was far above the melting temperature of PP. Nevertheless, the orientation of the crystals did not change as the molding temperature increased from 155 to 165 8C. The internal structure of the PP fibrils changed from a needle structure to a parallel-stacked one. The PP fibrils induced the crystallization of the PE melt, leading to the formation of a trans-crystalline layer at their surface. As the molding temperature increased, more PE lamellae protruded into the PP fibrils and the interface between the PP fibrils and the PE matrix became diffuse. q 2004 Elsevier Ltd. All rights reserved. Keywords: Nanocomposite; Morphology; Roller drawing
1. Introduction It has been widely established that the mechanical properties of crystalline polymeric materials strongly depend on processing conditions and techniques used to process the materials. The same polymeric material can be processed into a soft and flexible product or a strong and stiff product under different conditions. Many scientists and engineers [1,2] have devoted much effort to developing stronger polymeric materials. These practices include fiber reinforcement [3] and orientation reinforcement [2,4,5]. In the latter case the polymer chains are aligned by means of stretching the polymers below and close to their melting points, such means include solid extrusion [4,5] and die/ roller drawing [6,7]. There have been many studies on the reinforcement of * Corresponding author. Tel./fax: þ 852-2358-7125. E-mail address: [email protected] (C.M. Chan). 1 Present address: ASM Assembly Automation Ltd., 4/F, Watson Center, 16 Kung Yip St., Kwai Chung, Hong Kong. 0032-3861/$ - see front matter q 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2004.05.049
crystalline plastics with liquid crystalline polymer (LCP) fibers generated in situ [8]. Under specific conditions, polymer/LCP blends are injected into a mold; the LCP aligns in the flow direction and forms LCP fibers to reinforce the matrix. Also, there are reports on the conversion of melt-spun fibers into composites [9,10]. Highly aligned fibers are partially melted and compression molded. In this way, the fibers are joined together by melting the skin layer, forming a composite. Three basic morphologies of crystals have been observed in such oriented polymers [1]: fringed micelles, parallelstacked lamellae and fine needle crystals. Normally, fringed micelles appear in a sample stretched at a temperature below its glass transition temperature [11]. An example of fringed micelles is strain-induced crystallization of PET at room temperature. Stacked lamellae occur when the sample is drawn from a melt and then crystallizes [12]. Needle crystals are generated during large plastic deformation of a crystalline polymer at a temperature between the polymer’s glass transition temperature and its melting point [13,14]. Of
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course, there have been samples generated that show some combination of the above basic morphologies, such as the so-called shish-kebab structure [14,15], which consists of parallel-stacked lamellae and needle crystals. The crystalline structures in oriented polymers are believed to be metastable. The structure of the crystalline phase may change, depending on the temperature and time. For example, upon annealing, defects in needle crystals could agglomerate. Periodic arrangements of the defect-free and highly defective zones along the needle may appear. With further annealing at an elevated temperature, the needle crystals may be transformed into a shish-kebab structure and finally into a parallel-stacked structure [14,15]. Recently, a new method to prepare thermoplastic fibrilreinforced composites has been investigated in this laboratory [16]. Two crystalline polymers, commercial high-density polyethylene (HDPE) and isotactic polypropylene (PP), were first melt-blended together and then fibrillated through roller drawing. Finally, the fibrillated HDPE/PP blend was melt-processed again at temperatures above the melting point of HDPE. Preliminary results indicated that PP fibrils with a diameter of about 100 nm could survive after molding at temperatures up to 185 8C. The modulus and tensile strength in the fibrillation direction were increased by more than three times. Meanwhile, the modulus and tensile strength in the transverse direction increased by more than 20%. Our results also indicated that the mechanical properties of the sample varied with molding temperature. The finding that the strength in the transverse direction increased with increasing molding temperature and reached a maximum at 185 8C is interesting. In the present work, the effects of molding temperature on the morphology and mechanical properties of the HDPE/ PP nanocomposites will be examined.
2. Experimental 2.1. Preparation of the fibrillated HDPE/PP blend HDPE (Philips HMMPE) was blended with 25 wt% of isotactic polypropylene (Himont 6501) using a Haake TW 100 twin screw extruder. The HDPE/PP (75/25) pellets were extruded into 2.4 mm thick tapes and fibrillated by roller drawing at 138 8C [17]. After fibrillation, the thickness of the tapes was reduced to 0.3 mm and the drawing ratio was calculated to be approximately 10. A thermal analysis of the tapes [17] indicated that the PE component has a melting point at 139 8C and the fusion of the PE crystals is completed below 150 8C, while the fusion of the PP crystals, with a melting point of 167 8C, started at a temperature lower than 155 8C and finished below 175 8C.
laid parallel with respect to the fibrillation direction in a hot mold. The hot mold with the fibrillated sheets were first preheated in a press for 3 min under a pressure of about 3 MPa; then, the pressure was increased to 10 MPa and released immediately to degas; finally they were molded into 1.5 mm plates under a pressure of about 15 MPa. After molding for 25 min, the resulting plates were cooled first with compressed air from the pre-set molding temperature to below 120 8C (about 1 h) and then with tap water to below 50 8C under pressure. In the present work, five molding temperatures, 155, 165, 175, 185 and 200 8C, were selected. In the molding process, all PE crystals were melted while the partial melting of PP was changed to complete melting as the molding temperature increased. 2.3. Identification of the crystalline phase and orientation A powder X-ray diffraction system, Philips PW 1830, was utilized to identify the crystals and their orientations. The Cu anode X-ray generator was operated at 40 kV and 40 mA. The Cu Ka radiation impinged on the specimens with the beam perpendicular to the fibrillation direction. The diffraction experiment was conducted in the transmission mode and the output signal was recorded on a flat film at a diffraction distance of 90 mm. 2.4. Examination of morphology Thin strips with a cross-section of about 0.2 £ 0.3 mm2 were cut from the molded samples perpendicular to the fibrillation (or extrusion) direction. They were embedded in an epoxy resin and cured at 45 8C for 48 h. The embedded specimens were trimmed first with razor blades and then with glass knives on a Reichert-Jung Ultracut R microtome. The trimmed specimens were exposed to the vapor of an aqueous solution of ruthenium tetroxide [18,19] in a sealed test tube for 48 h. After staining with ruthenium tetroxide, the specimens were washed with a 3% aqueous solution of sodium periodate and distilled water. Then the samples were dried in a desiccators for 24 h. On a Reichert-Jung R microtome, ultrathin sections with a thickness about 60 nm were cut from the stained specimens with a diamond knife after a top layer of approximately 400 nm was removed. The sections were mounted on 200-mesh copper grids. After drying in a desiccator for 24 h, they were examined with a JEOL JEM-100 CX II transmission electron microscope (TEM), which was operated at an accelerating voltage of 80 kV.
3. Results 3.1. Crystalline phase and orientation
2.2. Preparation of HDPE/PP nanocomposites The fibrillated tapes were cut into short pieces and were
Fig. 1 shows the X-ray diffraction pattern of the fibrillated HDPE/PP blend. In Fig. 1, the fibrillation
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Fig. 1. The X-ray diffraction pattern of the HDPE/PP (75/25) blend after fibrillation by roller-drawing at 138 8C.
direction is horizontal. One can see that the azimuths of the diffractions are very small and they appear as spots, indicating a high degree of orientation of the crystals in the sample. The diffraction pattern is similar to that of PP fibers except three diffraction spots are from HDPE. Twelve diffraction spots align in the vertical direction and four more diffraction spots appear about 408 away from the meridian (vertical line). Beginning from the inner one, the diffraction spots are named Spots 1 to 6 and the one that is off the meridian is named Spot 7. The diffraction radii of the spots were measured and the diffraction angle, 2u; and the d-space of the reflecting planes were calculated. The results are summarized in Table 1. Also, the values of the d-spaces reported in the literature [20 –23] are listed in the last row of the table. Compared with the data from the literature, Spots 1 to 3 as well as Spot 7 are recognized to come from the monoclinic PP crystals (a) and the corresponding crystallographic planes are a(110), a(040), a(130) and a(111), respectively. Spot 4 is a reflection from (001) of the monoclinic PE crystals (m) while Spots 5 and 6 are the reflections from (110) and (200) of the orthorhombic PE crystals (o). However, it should be pointed out that some reflections of the monoclinic PP crystals—the primary diffraction of a(131) and the secondary diffraction of a(110)) can contribute to the intensity of Spot 5. But due to the low PP content and the weak diffraction intensity of the
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primary diffraction of a(131) and the secondary diffraction of a(110) compared with that of a(130), the crystallographic plane (110) of the orthorhombic PE crystals is considered to be the main source for Spot 5. Fig. 1 indicates that three crystalline phases exist in the fibrillated HDPE/PP blend—the monoclinic PP crystals and both the monoclinic and orthorhombic PE crystals. The monoclinic PP crystals align with their c-axes along the fibrillation direction as do the orthorhombic PE crystals. However, for the monoclinic PE crystals, the c-axis is perpendicular to the fibrillation direction. X-ray diffraction patterns of the HDPE/PP nanocomposites prepared by molding the HDPE/PP blend at temperatures from 155 to 185 8C are shown in Fig. 2. After molding, the radius of the diffraction spots for the PP and PE crystals stays almost the same. However, a new diffraction spot (Spot X) is observed near the reflection of a(110) (Spot 1). Based on the diffraction radius of the new spot, the ˚ . In PP, corresponding d-space was calculated to be 6.8 A there are three possible crystalline phases—the monoclinic (a), hexagonal (b) and triclinic (g) phases [20,21]. Normally, the monoclinic phase predominates in the samples. However, the hexagonal and triclinic phases may exist under special conditions, such as high pressure in the triclinic phase [23] and a special nucleating agent [24] in the hexagonal phase. Of these crystals, the g(100) has the ˚ as reported by Jones [20]. largest d-space of 6.37 A However, it is still much smaller than the value for this new reflection plane. Up to now, no report on co-crystallization of HDPE and PP has been found in the literature. When highly orientated thin PE and PP films were molded at temperatures between their melting points only epitaxially crystallized PE was observed [25]. The new diffraction spot is unlikely to be from a new crystalline phase. It has been reported that the unit cell constants of polymer crystals may change under stress, so the d-space of the reflecting plane may change [26]. The new diffraction spots may result from the stretched a(110) plane. However, if this is true, the diffraction intensity of the stretched a(110) should be stronger before molding. Because heating should result in the relaxation of stretched crystals, the intensity should decrease as the molding temperature increases. The fact is that the diffraction signal around the diffraction angle ð2uÞ; 12.998, is vague in the fibrillated blend (cf. Fig. 1). Only after molding does it become more obvious (cf. Fig. 2).
Table 1 The diffraction angle and corresponding crystallographic plane of the fibrillated HDPE/PP blend Sport Intensity Diameter [mm] 2u ˚) d-space (A Reflecting plane ˚) d-spacea (A a
N Weak 41.5 12.99 6.8 g(100)? 6.37
1 Middle 45.5 14.36 6.30 a(110) 6.26
Corresponding d-space from the literature.
2 Middle 55.5 16.96 5.22 a(040) 5.19
3 Weak 61.5 18.67 4.75 a(130) 4.77
4 Strong 63.5 19.51 4.54 m(001) 4.55
5 Very strong 72.0 21.58 4.11 o(110) 4.10
6 Middle 81.0 23.99 3.70 o(200) 3.69
7 Weak 70.5 21.17 4.19 a(111) 4.19
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Fig. 2. The X-ray diffraction pattern of the fibrillated HDPE/PP blend after molding at different temperatures: (a): 155 8C; (b): 165 8C; (c): 175 8C and (d): 185 8C.
Nevertheless, as the molding temperature was increased to 200 8C, the diffraction signal around 12.998 could not be identified (cf. Fig. 3). At this moment, it seems to be difficult to determine the source of the new diffraction sport. Fig. 2 shows a significant change on the diffraction patterns of m(001), o(110) and o(200) from the PE crystals. After molding at 155 8C, the diffractions of m(001), o(110) and o(200) changed from spots to circle rings (cf. Fig. 2 (a)). However, the diffraction intensities of the rings are not uniform. In the meridian regions, the intensities are much stronger. This indicates that the c-axis of the PE crystals still had the preferred orientation. But, as the molding temperature increased from 155 to 185 8C, the diffraction rings became more and more uniform (cf. diffractions of m(001) and o(200) in Fig. 2). The orientation of the PE crystals gradually became random. In contrast to the PE crystals, the azimuths of the diffractions of the PP crystals did not change much after molding at 155 8C. The PP crystals still aligned with the
Fig. 3. The X-ray diffraction pattern of fibrillated HDPE/PP blend after molding at 200 8C.
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c-axis towards the fibrillation direction. This is easy to understand because at this molding temperature, the PP crystals had not melted. Even when the molding temperature increased from 155 to 185 8C, the azimuthal extension of the diffractions did not increase much. This implies that the PP crystals still aligned with the c-axis along the fibrillation direction even after the sample was heated at 185 8C for 25 min. Only when the molding temperature increased to 200 8C, did the PP crystals lose their preferred orientation and did the sample give five uniform diffraction rings (cf. Fig. 3). 3.2. Morphology and interphase structure Fig. 4 shows the TEM micrographs of the sample prepared by molding a fibrillated HDPE/PP blend at 155 8C. In this sample, the PP component appears as straight fibrils in the PE matrix. The size of the fibrils ranges from 30 to 150 nm. In the PE matrix, the crystals are in a lamellar structure and have a thickness of approximately 15 nm. Instead of developing into spherulites, the lamellae form a highly ordered arrangement between the PP fibrils. They stack together almost parallel to each other and are aligned approximately perpendicular to the long axial of the fibrils, forming a transcrystalline interphase [27]. The coarse dark bands between the lamellae are believed to be the lessordered materials as they absorbed more staining agents. This is very similar to the observations made by Li et al. that a transcrystalline interphase was obtained between the isotactic PP matrix and poly(ethylene terephthalate) microfibers after a slit-extrusion, hot-stretching and quenching of a blend containing these two polymers [28]. The PP fibrils contain finer crystalline structures. At a high magnification (cf. Figs. 4(c) and 5(a)), it can be seen that there are some needle-shaped, fine crystals (marked with N). The needles are about 3 nm in width and tens of nm in length. Normally, they align in the fibril direction. However, a few fine crystals perpendicular to the fibril direction can be observed as well (marked with S in Fig. 5 (a)). Thus, the PP fibrils exhibit some features of a shishkebab structure. The surface of the PP fibrils is rough and the roughness of the fibrils is in the order of ten nm (marked with R in Fig. 4 (c)). Nevertheless, the boundary between the PE matrix and the PP fibrils appears to be rather sharp. Fig. 6 shows the TEM micrographs of a sample molded at 165 8C. The PE crystals similarly remain in their stackedlamellar structure, although the increase in the molding temperature led to a decrease in the thickness of the PE lamellae, from about 15 nm to about 10 nm. Additionally, a few twisted PE lamellae can be found in the PE phase. Furthermore, the PE lamellae are no longer aligned perpendicularly to the PP fibrils (cf. Fig. 6). The internal structures of the PP fibrils changed significantly as the molding temperature increased to 165 8C. No needle crystals are found; instead, some stacked lamellae are observed. The thickness of these PP lamellae varies from 4
Fig. 4. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 155 8C. B: boundary between the PP fibril and PE matrix; N: needle crystals; S: shish-kebab structure.
to 10 nm. The surface of the PP fibrils is still rough and the boundary between the PP fibrils and the PE matrix is still sharp. However, some PE lamellae are found to be occasionally sticking into the PP fibrils (marked with P in Fig. 6 (a)). Fig. 7 shows the morphologies of the sample molded at 185 8C. When the molding temperature increased from 165 to 185 8C, the thickness of the PE lamellae seemed to increase a little. Near the PP fibrils, the PE lamellae show a stacked pattern. However, between the PP fibrils, many twisted PE lamellae can be seen, just as the lamellar
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Fig. 6. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 165 8C. P: parallel stacked PP lamellae. Fig. 5. The fine structure of the PP fibrils of the HDPE/PP composite molded at 155 8C: (a) an enlarged view of an area of Fig. 1 and (b) PE lamellae and interface of between the PE matrix and PP fibrils.
structure observed in the sample crystallized from the PE melts without disturbance [29,30]. In contrast to the marked structural changes that occurred between 155 and 165 8C, there were no large changes in the internal crystalline structure of the PP fibrils when the molding temperature increased from 165 to 185 8C. The crystals are still in a stacked-lamellar structure. However, a striking change is observed at the interface between the fibrils and matrix. Some PE lamellae extended from the PE phase and penetrated into the PP fibrils. The boundary became rather diffuse. At a high magnification (Fig. 7 (b)), the surface of the PP fibrils is seen to have a saw-tooth structure with a period of 10 nm. However, it should be pointed out that there are more defects or disordered materials at the interface between the PP fibrils and the PE matrix because the interface always appears darker than other regions in the examined sections. Fig. 8 (a) shows a TEM micrograph of the sample molded at 200 8C. As the molding temperature was further increased to 200 8C, the PP component no longer existed as
fibrils but rather as particles, dispersing throughout the PE matrix. The size of the PP particles ranges from 0.5 to 2 mm. This morphology is very similar to that of the extruded sample (cf. Fig. 8 (b)). It should be pointed out that the interfacial bonding is weak, as evidenced by the voids in the dark rings around the PP particles. This observation is also supported by the mechanical property data indicating that the strong and ductile composite became a weak and brittle material as the molding temperature increased to 200 8C [11,17]. Fig. 9 shows the TEM micrographs of the sample prepared by extrusion of the fibrillated HDPE/PP at temperatures below 160 8C. On the stained section of this sample, PP fibrils can be observed (ref. Fig. 9 (a)). Beside the long PP fibrils, many circular and elliptical particles are observed (marked with C and E in Fig. 9 (a)). These circular and elliptical particles are believed to be the cross-sections of the PP fibrils due to the fact that the cutting plane was perpendicular or at an inclined angle to the fibril. This implies that the PP fibrils have a random orientation in the extruded sample. The size of the fibrils ranges from 50 to 300 nm, which is a little larger than those in the compression-molded samples. At a higher magnification
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Fig. 7. TEM micrographs of the HDPE/PP (75/25) composite prepared by compression-molding the fibrillated blend at 185 8C.
(cf. Fig. 9 (b)), a fine crystalline structure, similar to the structures observed in the sample molded at 155 8C, can be found in the PE matrix and PP fibrils. However, the lamellae appear to be finer. The PE lamellae are about 8 nm thick and the size of the PP crystals is about 2 nm.
4. Discussion The orthorhombic PE phase occurs most frequently in a normal, melt-crystallized sample. However, when a sample is subjected to stress or deformation, monoclinic PE crystals may be generated in the sample [22,23]. In the fibrillation process of the present work, the PE crystals were definitely stretched in the melt because of the high temperature of the rollers and the effect of the drawing [31]. Therefore, in the fibrillated HDPE/PP sample, some monoclinic PE crystals are present. It is evident that after molding at 155 8C, below which the fusion of PE was completed as measured by differential scanning calorimetry (DSC) [17], the diffraction spots of the PE crystals changed to circular rings because of the loss of the orientation induced by the fibrillation. However, the intensity of the reflections from m(001), o(110) and o(200)
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Fig. 8. The morphology of the HDPE/PP (75/25) blend; (a): prepared by compression-molding the fibrillated blend at 200 8C and (b): extruded tape before fibrillation.
appeared to be much stronger in the meridian (cf. Fig. 2). This implies that the PE crystals did not lose their orientation completely after molding at 155 8C for 25 min. The preferred orientation of the PE crystals can be explained by the incomplete randomization of the PE chains (molecular shape memory [32]) in the molding process. During the molding at 155 8C, the PE crystals disintegrated but many PE molecular chains still aligned in the fibrillation direction. Consequently, cooling caused the alignment of the chain axis of the PE molecules along the fibrillation direction. Nevertheless, the intensity of the spots decreased gradually as the molding temperature increased. The diffraction intensity became uniform in all directions at 185 8C, indicating that a complete randomization of PE chain coils had been achieved. This is also supported by the matrix morphology of the sample molded at 185 8C showing many twisted lamellae between the PP fibrils. DSC experiments [17] showed that the melting of the PP crystals in the fibrillated blend began at 155 8C and finished below 175 8C, with a melting peak at 167 8C. The molding temperature of 155 8C is lower than the melting point of the PP component, which should still exist as fibrils. Actually,
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Fig. 9. TEM micrographs of the HDPE/PP (75/25) composite prepared by extrusion of the fibrillated blend at 155 8C; (a): low magnification and (b) high magnification.
the molding at 155 8C can more appropriately be considered as annealing the PP fibrils at this temperature. Therefore, the diffraction azimuths (orientation) of the PP crystals did not change much. However, the defects in the original needle crystals might aggregate, resulting in the formation of needle crystals with periodic arrangements of defect-free and highly defective zones in the PP fibrils [15,16]. As an intermediate stage in the transformation of the needle crystals into the parallel-stacked structure during annealing, a shish-kebab structure was generated. As the molding temperature increased to above 165 8C, most of PP crystals should have disintegrated although the overall orientation of the PP chains did not change greatly within that period of time. Hence, the PP crystals changed from the needle to parallel-stacked structure while the diffraction azimuths (overall orientation) of the PP crystals remained almost the same. Even when the molding temperature was raised above 175 8C and the PP crystals disintegrated completely, the complete randomization of the
stretched polymer chains through the thermal motion of the chain segments needed a period of time and might not have been achieved within 25 min at molding temperatures below 185 8C. The long-range alignment of the molecular chains remained in the PP phase. As a result, the PP molecules were transformed into parallel-stacked lamellae with the chain direction aligned along the fibrillation direction on cooling, similar to the crystallization of a drawn polymer melt [13]. Below 185 8C, the contracting force of the stretched polymer chains was not strong enough to overcome the frictional force at the interface between the PP fibrils and the PE matrix, allowing the fibrils to shrink back. As a result, the PP phase stayed as fibrils. A similar phenomenon has been observed in multi-layer composites of polybutene-1 (PB-1) and PP thin films [25]. The orientation of the PB-1 crystals remained even when the composites had been heated to a temperature above the melting point of PB-1, provided that the PB-1 film was thin enough. When the molding temperature was further elevated to 200 8C, the stretched polymer chains possessed enough thermal energy to overcome the friction at the interface, forming random coils. The PP fibrils shrank into particles; consequently, an isotropic diffraction pattern was obtained. In our previous paper [17], we have shown that the mechanical properties of the sample varied with the molding temperature. The strength in the transverse direction increased with increasing molding temperature and reached a maximum at 185 8C. As the molding temperature increased, the degree of the thermal motion of the chain segments increased. The inter-diffusion of the molecules increased at the interface; consequently, more PE lamellae penetrated into the PP fibrils. The change in the structure of the interface between the PP fibrils and the PE matrix is the explanation for the increase in the tensile strength in the transverse direction of the composites with increasing molding temperature to 185 8C.
5. Conclusion The fibrillated HDPE/PP blend, which exhibits a fibrous structure, consists of at least three crystalline phases: the monoclinic PP phase, and the orthorhombic and monoclinic PE phases. The monoclinic PP crystals and orthorhombic PE crystals align with their molecular chain (c-axis) along the fibrillation direction while the monoclinic PE crystals orient with their c-axis perpendicular to the fibrillation direction. It will take a rather long period of time for a stretched polymer chain to return to a random molecular coil even after disintegration above its melting point. After molding at 155 8C, the PE matrix exhibits a parallel-stacked lamellar structure and still retains this orientation after fibrillation. However, the preferred orientation of the PE gradually vanishes as the molding temperature increases, the parallel-stacked lamellae change into twisted lamellae
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and the diffraction pattern becomes uniform rings as the molding temperature increases to 185 8C. The PP phase exists as fibrils in the PE matrix and the crystals retain their molecular chains in the fibrillation direction when the molding temperature is below 185 8C. Although the orientation of the crystals does not change as the molding temperature increases from 155 to 185 8C, the internal structure of the PP fibrils changes from a needle to parallel-stacked structure. As the molding temperature increases to 200 8C, a complete randomization of the PP chains is achieved and the PP phase contracts into particles, losing all previous orientations. The PP fibrils induce crystallization of the PE melt, leading to a trans-crystalline layer surrounding the PP fibrils. As molding temperature increases, more PE lamellae protrude into PP fibrils, producing a diffused interface between the PP fibrils and the PE matrix.
Acknowledgements This work was supported by the Hong Kong Government Research Grants Council under grant No. HKUST 6043/01P.
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