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poly (ethylene terephthalate) blends
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Composites: Part A 39 (2008) 164–175 www.elsevier.com/locate/compositesa
Morphology, static and dynamic mechanical properties of in situ microfibrillar composites based on polypropylene/poly (ethylene terephthalate) blends K. Jayanarayanan a, Sabu Thomas b, Kuruvilla Joseph c,* a
Department of Chemical Engineering and Materials Science, Amrita Vishwa Vidyapeetham, Coimbatore 641 105, Tamil Nadu, India b School of Chemical Sciences, Mahatma Gandhi University, Priyadarshini Hills P.O, Kottayam 686 560, Kerala, India c Department of Chemistry, Indian Institute of Space Science and Technology, ISRO P.O, Thiruvananthapuram 695 022, Kerala, India Received 23 June 2007; received in revised form 5 September 2007; accepted 18 November 2007
Abstract In situ composites based on blends of polypropylene (PP) and poly (ethylene terephthalate) (PET), were prepared by melt extrusion, followed by continuous drawing and isotropization. The blending of the mixture was carried out in a single screw extruder and the isotropization of the stretched blend was achieved by injection moulding. Scanning electron microscopy (SEM) studies showed that the extruded blends were isotropic, but both phases became highly oriented after drawing (stretching). The stretched blends were converted into in situ composites after injection moulding at temperatures below the melting point of PET. The size of the PET fibrils generated decreased with increasing stretch ratio. The tensile and impact properties were found to be higher for the samples drawn at stretch ratios 5 and 8. Dynamic mechanical properties such as the storage modulus (E0 ), loss modulus (E00 ) damping behavior (tan d) and static mechanical properties were investigated as a function of stretch ratio. The E0 values were found to be increasing up to a stretch ratio 8. The tan d and E00 modulus spectra showed a strong influence of the microfibrils on the a and b relaxation of PP. Finally, the tensile properties obtained experimentally were compared with those determined using theoretical equations. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Polypropylene; A. Poly (ethylene terephthalate); B. Morphology; C. Dynamic mechanical analysis
1. Introduction Blending of polymers is an effective method to generate newer materials offering better properties than the individual ones. The blends of polyolefins and poly (ethylene terephthalate) (PET) are of interest since the latter can improve the mechanical properties of the relatively weak partner. The usage of polypropylene as one of the phases in polymer blends is widely favored due to its excellent processability and low cost. Several studies were carried out on blends of PP and PET [1–5] to analyze the morphology
*
Corresponding author. Tel.: +91 471 2564808; fax: +91 471 2564806. E-mail address: [email protected] (K. Joseph).
1359-835X/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesa.2007.11.008
and mechanical behavior. The incorporation of PET fibres into PP to convert them into thermoplastic composites was tried out by Lopez and Arroyo [6]. The incorporation of PET fibers is of interest as a nucleating agent for the crystallization of PP [7,8]. The presence of PET fibres encourages transcrystallization of PP as revealed by these studies. In situ composites were made from the blends of liquid crystalline polymer (LCP) and polypropylene [9–11]. After melt-blending LCP attain fibrillar morphology to provide excellent properties to the resultant blend. Evstatiev and Fakirov [12] introduced a novel concept of developing microfibrillar reinforced composites (MFC) from polymer blends. Only immiscible, fibre forming polymers with considerable difference in their melting points (Tm) are capable of getting converted into MFCs. There are essentially three
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steps involved in the preparation of polymer/polymer MFCs, viz: (a) blending of the component polymers, usually in an extruder, (b) drawing or orientation and (c) annealing or the isotropization step. In the first step, melt-blending of two or more immiscible thermoplastic polymers take place with or without the presence of a compatibilizer. In the second step, the extruded blend is either cold drawn or hot drawn to achieve fibrillar morphology for the constituent polymers. In the third step, the drawn blends are annealed or isotropized which destroys the fibrillar morphology of the polymer with lower Tm. The fibrils of the polymer with higher Tm are preserved and are randomly distributed in an isotropic matrix of the polymer with lower Tm. Since the matrix and the reinforcement are thermoplastic polymers and the reinforcements (fibrils of the polymer with the higher Tm) are generated in situ during the process they are rightly called as in situ reinforced polymer/polymer composites. The sizes (diameter) of the fibrils generated are in the order of few microns and hence the name microfibrillar composites (MFCs). In the early days of the development of MFCs, the fibrillization step was achieved in a separate step after the meltblending was completed [12–22]. Recently, a new method of continuous drawing of the blend online, downstream the extrusion die was developed for the fibrillization step. The polyolefinic materials were blended with polyamides and polyesters and were converted into MFCs [23,24]. Evstatiev et al. [25] suggested a scheme for production of MFCs, which could be scaled up to industrial standards, using LDPE/PET blends. They made use of recycled PET obtained from bottles as reinforcement for LDPE, and thereby suggested the scheme as an efficient method for the recycling of PET. Another method was employed by Li et al. [26,27] to produce PE/PET MFCs, where the blended extrudate in the form of a strip was stretched while they were still in the hot condition. The tensile strength and modulus increased with an increase in hot stretching ratio and beyond a critical value it reduced. Recently, Friedrich et al. [28] developed in situ composites at a constant draw ratio from PP/PET (60/40) blend in the presence and absence of a compatibilizer. The composites were developed by compression moulding as well as injection moulding, and it was observed that composites from compression moulding were found to have better mechanical properties since the uniaxial orientation of the PET fibrils are maintained. Shorter fibrils were obtained in the blends with higher concentration of the compatibilizer. This indicates that the compatibilizer although improves interfacial adhesion, reduces the aspect ratio of the fibrils. The dynamic mechanical properties of the in situ composites are not reported widely in the literature except may be for those using liquid crystalline polymers. The main objective of this work is to study the morphology development of the in situ composites prepared from PP and PET (85/15) blends obtained at different draw ratios and the evaluation of their static and dynamic mechanical properties. The dynamic mechanical properties of compos-
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ites are of significance since the relaxations of the polymer under the various conditions of temperature and stresses can be studied. Dynamic mechanical analysis (DMA) over a wide range of temperatures permits the determination of the viscoelastic behavior of polymeric materials. The stiffness and damping characteristics of the composites can be analyzed using DMA which could be co-related to the morphology. The dynamic mechanical properties analysis is a convenient method for estimating the polymer transition temperatures which can influence the fatigue and impact properties of the composite. 2. Experimental work The polymers used were PP (Repol-H110MA, Reliance, India, MFI – 11.0 g/10 min) and PET (940400-B, Futura Polymers, India, I.V.0.814 dl/g). After drying PET for 12 h at 100 °C it was tumble mixed with PP at a constant weight ratio of 15/85. The mixture was then melt blended in a single screw extruder (Screw diameter – 20 mm, L/D ratio – 30) using a set temperature profile of 225, 235, 250, 255 and 260 °C and the screw rotating at 30 rpm. The scheme of the experimental setup is given in Fig. 1. The extrudate was in the form of circular strands of diameter 2 mm and were taken to a cooling water bath for solidification. Then these strands were passed through a hot air oven maintained at 100 °C. Subsequently these strands were taken to a take up device for continuous drawing. The take up device consists of a pair of nip rolls whose peripheral velocity (V1) is maintained same as the velocity of the extrudate. Beyond the nip rolls there was a pair of stretch rolls of the same diameter as that of the nip rolls, but whose speed (V2) can be varied to attain different draw ratios and thereby reduction in the cross-sectional dimensions of the strands. The ratio of the speeds of the stretch rolls to the nip rolls (V2/V1) was taken as the draw (stretch) ratio. The experiment was carried out for draw ratios 1, 2, 5, 8, 10 and the specimens were named as neat blend (NB), H2, H5, H8 and H10, respectively. The term H indicates that the stretching was carried out in the hot condition of the blend. Downstream the stretching unit, the strands were granulated to approximately 3 mm length. The granules thus produced were injection moulded into tensile test and impact test samples in a Ferromatic Milacron-Sigma 50T machine. The extruded neat blend was injection moulded at a temperature profile of 185, 225, 245 and 270 °C from the feed zone to the nozzle and it was designated as NBI. All the drawn blends were injection moulded at a set temperature profile of 160, 170, 190 and 205 °C, from the feed zone to the nozzle. The samples taken were designated as H2I, H5I, H8I and H10I corresponding to the stretch ratios. The term I indicates the samples were isotropized by injection moulding technique. For example, H5I indicates that PP/PET blend in the ratio 85/15 was hot stretched at a V2/V1 of 5 followed by the isotropization by injection moulding technique.
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V2
V1
Extruder Cooling water bath
Nip rolls
Oven
Stretch rolls
Grinder
Drawn Blend
Injection moulding machine
Fig. 1. Scheme of the experimental setup for microfibrillar in situ composites preparation.
3. Characterization methods 3.1. Morphological analysis A JEOL JSM 840 SEM with an acceleration voltage of 20 kV was used for studying the morphology of the specimens. To extract the PET phase from the specimens, a mixture of phenol/1, 1, 2, 2, tetra chloroethane in 60/40 wt.% was used as the solvent. Similarly to remove PP and retain the PET fibrils the specimens were treated with hot xylene. All the specimens were coated with a thin gold layer prior to SEM analysis. The diameter of about 150 fibrils was measured from the micrographs of stretched samples using image analysis software. 3.2. Mechanical properties The tensile properties of the in situ composites at different stretch ratios were measured at room temperature according to ASTM D-638 using the injection moulded dumbbell specimens. All the tests were conducted at a constant crosshead speed of 50 mm/min. Impact Strength was measured using an Izod impact tester using notched specimens according to ASTM D-256. The flexural properties of the samples were measured according to ASTM D-790 using injection moulded rectangular specimens. Five specimens were tested in each case and the average values were reported. 3.3. Dynamic mechanical properties Rectangular specimens having size 60 mm 13 mm 3.3 mm were used for the dynamic mechanical experiments. Dynamic mechanical thermal analyzer NETZSCH DMA 242 was used for the evaluation of storage modulus (E0 ), loss modulus (E00 ) and mechanical damping factor (tan d). Three point-bending modes were used. The temperature
range over which properties were measured was 20 to 150 °C at a heating rate of 5 °C/min. The tests were carried out at a frequency of 1 Hz. 4. Results and discussion 4.1. Morphology development of in situ composites The morphology of the blend after extrusion and drawing are shown in Fig. 2. As evident from Fig. 2a, after extrusion, the neat blend (stretch ratio 1) indicates typical incompatible blend morphology with discrete domains of minor component dispersed in a continuous phase of the major component. The PET phase in the form of spherical and elliptical particles with sectional dimensions varying from 1 to 6 lm is found to be distributed in the major phase PP. The samples obtained after drawing with ratios 2, 5, 8 and 10 show orientation and fibrillation of the PET phase. Large bunches of continuous fibres of PET are observed in the micrographs (Fig. 2b–e) for the samples H2, H5, H8 and H10, respectively. In the micrographs of the stretched blends, any apparent loss of orientation is due to the floating effect of the fibrils in the solvent during extraction of PP phase and drying. Also some residual PP phase is seen as attached to the fibrils. The fibril diameter after stretching was estimated, however, the aspect ratio could not be analyzed due to the twisted and curly nature of the fibrils. The PET fibrils has a mean diameter of 8.6 lm in the case of H2. When the stretch ratio is increased to 5, the mean diameter reduces to 4.9 lm and at stretch ratio 8 it is 4.1 lm. At a stretch ratio of 10, the mean diameter increases to 6.9 lm. A careful examination of the SEM micrograph (Fig. 2e) showed the breakage of the fibrils which is indicative of the large extensional deformation the blend being subjected to at draw ratio of 10. The increase in mean diameter at draw ratio 10 is due to the presence of some large fibrils.
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Fig. 2. SEM images of samples after extrusion and drawing (a) neat blend after extrusion with PET phase extracted, (b–d) drawn blends at stretch ratios 2, 5 and 8, respectively, with PP phase extracted and (e), (f) drawn blend at stretch ratio 10 with PP phase extracted.
SEM analysis and particle size estimation of another specimen of H10 (Fig. 2f) confirmed the presence of large and small diameter fibrils (ranging from 1 to 10lm). It was also observed that many of the smaller diameter fibrils were discontinuous due to the breakage. It has been reported in the case of in situ composites based on liquid crystalline polymers [29] and PE/PET MFCs [27], at very high draw ratios the elongated molten droplets cannot withstand the large extensional deformation and break up into smaller entities leading to the reduction in aspect ratio. At a stretch ratio of 10, the PET phase does not get sufficient time to complete the orientation and get transformed into continuous fibres. At the same time, large extensional deformation forces at this stretch ratio results in breakage of many fibrils. It should be reported here that
there was physical breakage of the extruded strands at draw ratios of 12 and experiments could not be carried out at higher draw ratios. Injection moulding at temperatures above the melting point of PP but below that of PET leads to the melting of PP phase, but the fibrillar morphology of PET is preserved. After injection moulding, PET fibrils loose their orientation and is randomly distributed in the PP matrix. Since these fibrils are exposed to temperatures above the melting point of PP during injection moulding coupled with the high shear rate, there is a reduction in the diameter of the PET microfibrils. This phenomenon is due to the ‘break up behavior’ [30] of the fibrils during relaxation at elevated temperatures which is schematically represented in Fig. 3. First the long fibril breaks up into a number of
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A
B
C
A
Long Fibril
B
Short Fibrils
C
Spheres
Fig. 3. Break up behavior of fibrils during injection moulding.
shorter ones, then the short rod is pulled back to spheres by interfacial tension. During isotropization, the break up of the long fibrils should be limited to the formation of short fibrils to complete the production of in situ composites, rather than going back to the spherical shape. The SEM micrograph (Fig. 4a) of injection moulded samples from neat blend (NBI) does not reveal any presence of fibrils as expected. The SEM micrograph of H2I (Fig. 4b) shows short fibrils with cross-sectional dimensions varying from 7 to 9 lm. The number of fibrils obtained after isotropization are less which can be attributed to (a) the lower conversion rate of the PET phase into fibrils during stretching since the draw ratio is less and (b) the break up behavior of the fibrils during isotropization leading to the formation of spherical entities rather than short rods. For H5I, microfibrils of about 5 lm diameter are observed. The aspect ratio of the fibres is relatively higher as evident from the micrograph (Fig. 4c). In the case of H8I (Fig. 4d), fine fibrils of about 3 lm, but of a lower aspect ratio in comparison with H5I, which are strongly attached to the matrix, are observed. For H10I, very short fibrils of approximately 4 lm diameter are seen (Fig. 4e). It should be noted that for H10, breakage of the fibrils was observed and that could be an additional reason along with the break up behavior for the reduction in the length of the PET fibrils in H10I. The variation of PET microfibril diameter with stretch ratio after extrusion and drawing and after isotropization (injection moulding) is delineated in Fig. 5. Based on the above studies a scheme for the morphology development of the blend during drawing and injection moulding is presented in Fig. 6. After melt extrusion at 270 °C, the extrudate is cooled and then raised to the orientation temperature of 100 °C for continuous drawing. The fibrils of PP and PET are highly oriented at this stage as represented in Fig. 6b. In the next step, during injection moulding at 205 °C, which is higher than the Tm of PP but lower than that of PET, PP phase melts and forms the isotropic matrix for the composite. The PET microfi-
brils which are highly oriented prior to injection moulding do not melt at this temperature. However, they loose their orientation to get randomly distributed in the PP matrix as represented in Fig. 6c. This scheme has been widely reported [22,25] as the route to develop microfibrillar composites from two individual polymers. 4.2. Static mechanical properties The tensile, flexural and impact properties of PP, NBI and all the in situ composites are given in Table 1. The low tensile strength of NBI in comparison with PP can be attributed to the spherical domains of the PET phase in the blend which remain incompatibile with the PP phase. The tensile strength of H2I is slightly higher than NBI but still lower than PP. The tensile strength of H5I is the maximum amongst all the samples analyzed. H8I shows a slight reduction in the tensile strength in comparison with H5I. Further increase in stretch ratio leads to reduction in the tensile properties. The tensile modulus is found to be maximum for H5I. There is an increase of 41% in the tensile modulus of H5I when compared to that of NBI. H8I also exhibits good tensile properties but slightly lower than H5I. The influence of stretch ratio on the tensile strength and tensile modulus of the blend is presented in Fig. 7. It can be inferred that the tensile strength and tensile modulus of the in situ composites increases with stretch ratio, reaches a maximum value between stretch ratios 5 and 8 and then reduces. The stress–strain curves of NBI, H2I, H5I, H8I and H10I are plotted in Fig. 8. It could be seen from the plot that the neat blend (NBI) yields and fails at a low stress indicating that PET phase in the form of spherical particles are not able to act as reinforcement for the PP matrix. H2I yields and fails at a higher stress in comparison with H10I. It should be remembered that the microfibrils obtained in the case of H2I were longer than those obtained in the case of H10I. H5I and H8I exhibit a relatively ductile failure in comparison with NBI, H2I and H10I. From Fig. 8, we
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Fig. 4. SEM images of injection moulded (isotropized) blends a) injection moulded neat blend with PP phase extracted, (b–e) isotropized drawn blends at stretch ratios 2, 5, 8 and 10, respectively with pp phase extracted.
could see that necking is appreciable in the case of H5I and H8I as evident from the strain these samples are able to take up beyond the yield stress before failure. High yield stress followed by a pronounced necking region makes H5I and H8I samples stiffer and tougher. H10I exhibits poor tensile properties, which may be attributed to the low aspect ratio of the fibrils as evidenced from the micrographs. The tensile modulus of all the in situ composites was found to be higher than PP and NBI. It was reported recently [27] that the tensile properties of HDPE/PET microfibrillar composites increased as the stretch ratio increases and then tends to decrease at very high stretch
ratios. High tensile modulus coupled with high elongation at break as exhibited by H5I and H8I indicate the improvement in the toughness of the composites in comparison with NBI. The length of the PET microfibrils were maximum in the case of H5I and H8I as evident from the corresponding SEM images. They are acting as excellent stress transfer agents, which contribute to the enhancement in the toughness of H5I and H8I. Another reason for the improvement of the tensile and impact properties could be the development of a trancrystalline layer [8] of PP around the PET fibrils especially in the case of H5I and H8I which can encourage better adhesion between the phases.
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a
pared to NBI. There is an improvement in the flexural properties of the in situ composites with stretch ratio. However, beyond stretch ratio 5, flexural strength and modulus decrease. H10I exhibits poor flexural properties, which indicates the ineffectiveness of short PET microfibrils as reinforcement for PP. The impact strength of H8I is the maximum amongst all the samples followed by H5I. For H8I, microfibrils with the least diameter were obtained as observed from the micrographs. Further, the relatively continuous nature of fibrils allow better stress transfer from the matrix polymer (PP) during impact. H10I exhibits poor impact properties since the short fibrils cannot act as effective stress transfer agents. It should also be mentioned that the impact strength is the lowest for NBI, which is indicative of the immiscibility of the two polymers.
Fibril dimater (microns)
10
8
6
4
2
0 0
Fibril diameter (microns)
b
3
6 Stretch ratio
9
12
10
4.3. Dynamic mechanical analysis
8
The dynamic mechanical analysis of the injection moulded samples was carried out from 20 to 165 °C at 1 Hz.
6
4.3.1. Storage modulus The storage modulus E0 obtained from the DMA is closely related to the load bearing capacity of the material and when the experiment is carried out in a three point-bending mode, it is similar to the flexural modulus measured as per the ASTM D-790 [31]. In Fig. 9, variation of E0 with temperature for PP, NBI, H2I, H5I, H8I and H10I at 1 Hz is shown. At low temperatures, E0 values of PP, H2I, H5I and H8I are very close to each other. It is an indication, that, at low temperatures the fibrils do not contribute much in imparting stiffness to the composite. In the case of PP, there is a sharp fall in E0 on passing through the glass transition temperature (Tg). It is due to the increased molecular mobility of the polymer chains above Tg. There is a large fall in modulus with increasing temperature for PP in comparison with other samples, the stiffness at high temperature being determined by the amorphous regions, which are very submissive above glass transition temperature.
4
2
0 0
3
6
9
12
Stretch ratio
Fig. 5. Variation of PET microfibril diameter with stretch ratio (a) after extrusion and drawing and (b) after isotropization (injection moulding).
There is considerable improvement in the flexural properties of the in situ composites in comparison with PP and NBI. The flexural strength of H5I was found to be increased by 20% and flexural modulus by 50% when com-
a
b
c PP isotropized
PP
Extrusion / Drawing
Injection Moulding
270°C/100°C
205°C
PET PETMicrofibril
Fig. 6. Morphology development scheme of the blend during extrusion and injection moulding.
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Table 1 Tensile, flexural and impact properties of PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios Sample
Tensile strength (MPa)
Tensile modulus (MPa)
Elongation at break (%)
Flexural strength (MPa)
Flexural modulus (MPa)
Impact strength (J/m)
PP NBI H2I H5I H8I H10I
29.2 23.3 27.9 32.4 31.5 27.1
1470 1514 1662 2137 1927 1642
33.9 24.2 28.6 34.1 32.1 27.8
36.4 39.1 41.9 47.1 44.4 38.3
742 673 787 1015 837 735
26.1 21.7 25.6 30.6 36.7 25.2
The drop in the modulus on passing through the Tg is much less for H5I and H8I in comparison with PP. Even for NBI, H2I and H10I the difference between the modulii of the glassy state and rubbery state is smaller when compared to PP. The variation of storage modulus with stretch ratio at different temperatures is delineated in Fig. 10. At temperatures above the Tg of PP, the storage modulus increases with draw ratio, reaches a maximum value between the stretch ratios 5 and 8 and then declines. This substantiates the SEM observations, where long microfi-
a Tensile strength ( M Pa)
32
28
24
brils of PET were obtained for H5I and H8I. The longer microfibrils contribute in imparting stiffness to the composite. The influence of draw ratio could further be explained by studying the normalized storage modulus with temperature at different draw ratios as depicted in Fig. 11. Normalized storage modulus can be defined as the ratio of the storage modulus of the composite to the storage modulus of the matrix at the same temperature. It is evident from the plot that as the draw ratio increases the normalized storage modulus increases, reaches a maximum value and then decreases. Also at a particular draw ratio the normalized storage modulus increases with increase in temperature. This indicates the restriction of the mobility of the PP molecules due to the presence of the PET micro fibrils at elevated temperatures. The shorter microfibrils obtained in the case of H10I are not as efficient as those from H5I and H8I in restricting the segmental mobility of PP chains. The effect of the microfibrils as a reinforcing agent is appreciable at 50 °C, beyond which PET becomes flexible on approaching its Tg.
35
20 0
3
6
9 30
Stretch ratio
b
Stress (MPa)
Tensile modulus (M Pa)
2500
2100
25
NBI
20
H2I
H5I
15
H8I
10
1700
H10I
5
0
1300
0
0
3
6
10
20
30
40
50
60
9
Stretch ratio
Fig. 7. The influence of stretch ratio on (a) tensile strength and (b) tensile modulus for PP/PET (85/15) blend.
Strain (%)
Fig. 8. Stress–strain curves for injection moulded neat blend and microfibrillar in situ composites prepared at different stretch ratios.
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1.8 PP NBI
E' (MPa)
H2I H5I
6000
H8I H10I
4000
1.6
Normalised E'
8000
1.4
1.2
1 2000
10ºC
25ºC
50ºC
80ºC
0.8 1
3
0 -20
0
20
40
60
80
100
120
5
7
9
11
Stretch Ratio
140
Fig. 11. Variation of normalized storage modulus of microfibrillar in situ composites with stretch ratio at temperatures 10, 25, 50 and 80 °C.
Temperature (ºC)
Fig. 9. Variation of storage modulus with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
4.3.2. Loss modulus The loss modulus E00 is the viscous response of viscoelastic materials. It is a measure of the energy dissipated or lost as heat per cycle of sinusoidal deformation, when different systems are compared at the same strain amplitude. Fig. 12 shows the variation of loss modulus with temperature for PP, net blend and in situ composites at 1 Hz. In all the cases the value of the loss modulus increases, reaches a maximum value and then decreases. Polypropylene exhibits three relaxations localized in the range of 80 °C (c), 10 °C (b) and 100 °C (a). The loss modulus value is maximum for PP at Tg (b transition) when compared to the blend and in situ composites. There is a steady fall of the modulii
beyond the Tg for PP and NBI. The loss modulus of H5I and H8I beyond 35 °C is found to be the highest which is an indication of the improved energy dissipation, which could contribute to their better impact properties. In the case of H10I, viscous dissipation is poor as evident from low loss modulus values in comparison with H5I and H8I. The higher loss modulus on approaching the a transition is an evidence for the restriction on the segmental mobility of the polymer molecules [32]. Generally, it is considered that the lowering and the broadening of the loss modulus peak is an indication of the greater constraints on the relaxation of the amorphous phase [33,34] and thereby improvement in the mechanical properties. But it has also been reported that higher E00 values at Tg is an indication of the good interfacial adhesion
8000 900 PP
750
6000
NBI H2I H5I
E" (MPa)
E' (MPa)
600
4000
H8I
450 H10I
300
2000 150
25ºC
10ºC
50ºC
80ºC
0
0
1
3
5
7
9
11
Stretch ratio Fig. 10. Variation of storage modulus of microfibrillar in situ composites with stretch ratio at temperatures 10, 25, 50 and 80 °C.
-20
10
40
70
100
130
Temperature (°C)
Fig. 12. Variation of loss modulus with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
K. Jayanarayanan et al. / Composites: Part A 39 (2008) 164–175
between matrix and the fibre [35]. In this particular case, the lowest peak at Tg was obtained for H10I which exhibited poor impact properties. H2I has the highest E00 values at Tg amongst the in situ composites, but it should also be noted that values beyond 30 °C is lower than that of H5I and H8I. Hence it has to be inferred that the increased loss modulus on approaching the a transition contributes much more than the lowering of the E00 peak towards the mechanical properties.
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of the chains of the matrix polymer. Elevation of Tg is taken as a measure of the interfacial interaction. The intensity of the tan d peak is maximum for neat PP and progressively it is found reducing with the increase in draw ratio. The damping peak in the in situ composites showed a decreased magnitude of tan d in comparison to PP and NBI. The damping peak is associated with the partial loosening of the polymer structure so that groups and small chain segments can move. The broadening of the tan d peak is appreciable for H5I and H8I which is due to the restricted relaxation. It is interesting to study the behavior of the samples at the a relaxation. The peak heights of the in situ composites are more than that of PP. It can be attributed to the higher loss modulus values of these samples at this temperature. There is a positive shift in the a relaxation temperature (Ta) of PP phase for the in situ composites in comparison with neat polymer. Ta for neat PP is 93.5 °C and for NBI it is 89 °C. For the in situ composites H2I, H5I, H8I and H10I Ta is 101.9, 102.8, 103 and 95.25 °C, respectively. The shift in Ta in the case of H10I is marginal (1.7 °C) when compared to all other in situ composites which indicates interfacial interaction is poor at this draw ratio. This fact is consistent with the interpretation of the loss modulus curve and it is further substantiated with the experimental results.
4.3.3. Tan d The ratio of loss modulus to storage modulus is measured as the mechanical loss factor or tan d. The damping properties of the material give the balance between the elastic phase and viscous phase in a polymeric structure. The damping behavior of the composites in the transition region is governed by (a) mechanical relaxation of the matrix and the fibre, (b) the interface between the fibre and the matrix and (c) fibre loading and fibre length. In the present investigation, the variation of tan d as a function of temperature at 1 Hz is represented in the Fig. 13. Since the Tg of PET is higher than 60 °C, the effect of PET fibrils on the relaxation of PP phase can be easily analyzed. Neat PP has a Tg of 10.5 °C as obtained from the tan d curve. For NBI, there is only a marginal increase to 11.4 °C. The Tg value is maximum for H8I which is 21 °C and it is 10.5 °C more than the neat PP. This result is in contrast with one of the earlier studies [6] where PP was mixed with chemically treated PET fibres. They reported that Tg of PP in the composites produced with treated PET fibres which when introduced separately was found to be less than that of neat PP. But, it is widely agreed [33–35] that an increase in the temperature of the tan d peak is indicative of the restriction of the mobility
4.4. Theoretical prediction of tensile properties Halpin–Tsai equation has been proved useful in predicting the longitudinal modulus of aligned short-fiber composites, e.g., in situ composite with liquid crystalline polymer [36,37]. Li et al. [26] and Fuchs et al. [38] tried to apply this equation in predicting the tensile properties
0.16
PP
NBI
H2I
H5I
H8I
H10I
0.14
Tan delta
0.12 0.1 0.08 0.06 0.04 0.02 0 -20
0
20
40
60
80
100
120
140
160
Temperature (ºC) Fig. 13. Variation of mechanical damping factor with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
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of MFCs. Here, both the fiber modulus (Ef) and the fiber aspect ratio (L/D) are dependent on the stretching ratio and the dispersed phase content of the extrudate. According to Halpin–Tsai model [39], the Young’s modulus, Mc, and the tensile strength, Tc, of a composite are given by the following equations: 1 þ Ag1 V f Mc ¼ Mm ð1Þ 1 g1 V f 1 þ Ag2 V f ð2Þ Tc ¼ Tm 1 g2 V f where g1 and g2 are calculated using the following equations: g1 ¼ g2 ¼
Mf 1 Mm Mf þA Mm Tf 1 Tm Tf þA Tm
ð3Þ ð4Þ
The terms Mc, Mm, Mf and Tc, Tm, Tf are the tensile modulii and tensile strengths of the composite, matrix and fibre, respectively, and Vf is the fibre volume fraction. The value of A is twice the aspect ratio of the fibres. Halpin and Kardos [39] modified Eqs. (1) and (2) by including the term u which in turn is dependent on Umax, maximum packing fraction, of the reinforcement in the following manner: 1 þ Ag1 V f Mc ¼ Mm ð5Þ 1 ug1 V f 1 þ Ag2 V f Tc ¼ Tm ð6Þ 1 ug2 V f 1 Umax ð7Þ u¼1þ U2max The value Umax is 0.78 for square arrangement of fibres, 0.91 for hexagonal array of fibres and 0.82 for random packing of fibres. The value of Umax in this case was taken as 0.82 since the PET microfibrils are distributed randomly in the PP matrix. The value of A estimated using SEM images were 12.5, 64, 60 and 10, respectively, for H2I, H5I, H8I and H10I.The tensile strength and tensile modulus of PET fibrils were taken as 62 and 3250 MPa, respectively, for the calculations. The tensile strength and tensile modulus of PP (matrix) was experimentally obtained as 29.2 and 1470 MPa.
The tensile strength and modulus values obtained using Halpin–Tsai and modified Halpin–Tsai equations along with corresponding experimental results are given in Table 2.The theoretical values obtained by using both the equations are almost same for the in situ composites at a particular stretch ratio. The tensile strength value obtained by theoretical prediction for H8I shows a positive deviation of 8% from the experimental value, whereas in the case of H5I, the deviation is reduced to 5%.The tensile strength obtained experimentally for H2I and H10I are much lower than the values estimated theoretically. The tensile modulus values obtained theoretically are in close agreement with the experimental values obtained for H2I and H10I. In the case of H8I, the theoretical value for the tensile modulus shows a negative deviation of 12%, whereas in the case of H5I the negative deviation increases to 17%. The theoretical equations were found to be useful to model the tensile strength of H5I and H8I and tensile modulii of H2I and H10I. 5. Conclusion In situ composites were prepared from the blends of polypropylene and poly (ethylene terephthalate) by continuous drawing followed by injection moulding. The SEM observations show high level of orientation for the drawn blends. The PET fibrils with the lowest mean diameter of 4.1 lm were obtained at draw ratio of 8. After injection moulding at a temperature below the Tm of PET, fibrils with high aspect ratio and good orientation were obtained for samples drawn at stretch ratio 5. Beyond stretch ratio 8, the breakage of the fibrils were observed during stretching which produced very short randomly distributed fibrils after injection moulding.The tensile, flexural and impact properties were found increasing with stretch ratio. However, too high a stretch ratio is not desirable for the reinforcement of PP. The effect of temperature on dynamic mechanical properties like storage modulus, loss modulus and mechanical loss factor of PP, neat blend and in situ composites prepared at different draw ratios were studied. The E0 values were found decreasing drastically for neat PP beyond Tg. The presence of microfibrils showed a positive effect on the modulus at temperatures above Tg especially for the samples drawn at stretch ratio 5 and 8. The tan delta peak height at b transition (Tg) reduced and shifted to higher values with draw ratio. The maximum shift was obtained for samples drawn at a stretch ratio of
Table 2 Experimental and theoretical tensile strength, tensile modulus values obtained for in situ composites Sample
H2I H5I H8I H10I
Experimental
Halpin–Tsai
Modified Halpin–Tsai
Tensile strength (MPa)
Tensile modulus (MPa)
Tensile strength (MPa)
Tensile modulus (MPa)
Tensile strength (MPa)
Tensile modulus (MPa)
27.9 32.4 31.5 27.1
1662 2137 1927 1642
33.7 34.0 34.0 33.7
1718 1732 1731 1714
33.8 34.0 34.0 33.7
1718 1733 1732 1715
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8. The E00 curves also showed positive shift in the peak values the highest being for samples drawn at stretch ratio 5. These samples exhibited high loss modulus values beyond 25 °C which is indicative of the better energy dissipation. The experimental values for tensile strength and tensile modulus were compared with those obtained from the theoretical equations. Based on the above studies it can be concluded that microfibrils of PET obtained at draw ratios between 5 and 8 can effectively reinforce PP. References [1] Papadopoulou CP, Kalfoglou NK. Comparison of compatibilizer effectiveness for PET/PP blends: their mechanical, thermal and morphology characterization. Polymer 2000;41:2543–55. [2] Lepers JC, Favis BE, Tabar RJ. The relative role of coalescence and interfacial tension in controlling dispersed phase size reduction during the compatibilization of poly(ethylene terephthalate)/polypropylene blends. J Polym Sci Polym Phys 1997;35:2271–80. [3] Pang YXD, Jia M, Hu HJ, Hourston DJ, Song M. Effects of a compatibilizing agent on the morphology, interface and mechanical behaviour of polypropylene/poly (ethylene terephthalate) blends. Polymer 2000;41:357–65. [4] Bae TY, Park KY, Kim DH, Suh KD. Poly (ethylene terephthalate)/ polypropylene reactive blends through isocyanate functional group. J Appl Polym Sci 2001;81:1056–62. [5] Song M, Pang YX. Correlation among morphology, interface, and mechanical properties: experimental study. J Macromol Sci Phys B 2001;40:1153–67. [6] Lopez MA, Arroyo M. Thermal and dynamic mechanical properties of polypropylene and short organic fiber composites. Polymer 2000;41:7761–7. [7] Wang C, Liu CR. Transcrystallization of polypropylene composites: nucleating ability of fibres. Polymer 1999;40:289–98. [8] Saujanya C, Radhakrishnan S. Structure development and properties of PET fibre filled PP composites. Polymer 2001;42:4537–48. [9] Postema AR, Fennis PJ. Preparation and properties of self reinforced polypropylene/liquid crystalline polymer blends. Polymer 1997;38:5557–64. [10] Chiou YP, Chiou KC, Chang FC. In situ compatibilized polypropylene/liquid crystalline polymer blends. Polymer 1996;37:4099–106. [11] Saengsuwan S, Sauvarop BL, Mitchell GR, Olley RH. Thermotropic liquid crystalline polymer (Rodrun LC5000)/polypropylene in situ composite films: rheology, morphology, molecular orientation and tensile properties. Polymer 2003;44:3407–15. [12] Evstatiev M, Fakirov S. Microfibrillar reinforcement of polymer blends. Polymer 1992;33:877–80. [13] Fakirov S, Evstatiev M, Schultz JM. Microfibrillar reinforced composite from drawn poly (ethylene terephthalate)/nylon-6 blend. Polymer 1992;34:4669–79. [14] Fakirov S, Evstatiev M, Petrovich S. Microfibrillar composites from binary and ternary blends of polyesters and nylon 6. Macromolecules 1993;26:5219–23. [15] Evstatiev M, Fakirov S, Friedrich K. Effect of blend composition on the morphology and mechanical properties of microfibrillar composites. Appl Compos Mater 1995;2:93–106. [16] Serhatkulu T, Erman B, Bahar I, Fakirov S, Evstatiev M, Sapundjieva D. Dynamic mechanical study of amorphous phases in poly (ethylene terephthalate)/nylon-6 blends. Polymer 1995;36:2371–7. [17] Evstatiev M, Nicolov N, Fakirov S. Morphology of microfibrillar reinforced composites PET/PA 6 blend. Polymer 1996;37:4455–63. [18] Evstatiev M, Schultz JM, Petrovich S, Georgiev G, Fakirov S, Friedrich K. In situ polymer/polymer composites from poly (ethylene terephthalate), polyamide-6 and polyamide-66 blends. J Appl Polym Sci 1998;67:723–37.
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[19] Krumova M, Fakirov S, Balta Calleja FJ, Evstatiev M. Structure development in PET/PA6 microfibrillar – reinforced composites as revealed by microhardness. J Mater Sci 1998;33:2857–68. [20] Sapoundjieva D, Denchev Z, Evstatiev M, Fakirov S, Stribeck N, Stamm M. Transcrystallization with reorientation in drawn PET/PA 12 blend as revealed by WAXS from synchrotron radiation. J Mater Sci 1999;34:3063–6. [21] Evstatiev M, Schultz JM, Fakirov S, Friedrich K. In situ fibrillar reinforced PET/PA-6/PA-66 blend. Polym Eng Sci 2001;41:192–204. [22] Sarkissova M, Harrats C, Groeninckx G, Thomas S. Design and characterization of microfibrillar reinforced composite materials based on PET/PA 12 blends. Composites: Part A 2004;35:489–99. [23] Monticciolo A, Cassagnau P, Michel A. Fibrillar morphology development of PE/PBT blends: rheology and solvent permeability. Polym Eng Sci 1998;38:1882–9. [24] Pesneau I, Kadi AA, Bousmina, Cassagnau PH, Michel A. From polymer blends to in situ polymer/polymer composites: morphology control and mechanical properties. Polym Eng Sci 2002;42:1990–2004. [25] Evstatiev M, Fakirov S, Krasteva B, Friedrich K, Covas JA, Cunha AM. Recycling of poly(ethylene terephthalate) as polymer–polymer composites. Polym Eng Sci 2002;2:826–35. [26] Li ZM, Yang MB, Lu A, Feng JM, Huang R. Tensile properties of poly(ethylene terephthalate) and polyethylene in situ microfiber reinforced composite formed via slit die extrusion and hot-stretching. Mater Lett 2002;56:756–62. [27] Li ZM, Yang MB, Lu A, Feng JM, Huang R, Feng JM, et al. In situ microfiber reinforced composite based on PET and PE via slit die extrusion and hot stretching: influences of hot stretching ratio on morphology and tensile properties at a fixed composition. Polym Eng Sci 2003;43:615–28. [28] Friedrich K, Evstatiev M, Fakirov S, Evstatiev O, Ishii M, Harrass M. Microfibrillar reinforced composites from PET/PP blends: processing, morphology and mechanical properties. Compos Sci Tech 2005;65:107–16. [29] Lin QH, Jho J, Yee AF. Effect of drawing on structure and properties of a liquid crystalline polymer and polycarbonate in situ composite. Polym Eng Sci 1993;33:789–98. [30] Li X, Chen M, Huang Y, Cong G. In situ generation of polyamide-6 fibrils in polypropylene processed with a single screw extruder. Polym Eng Sci 1999;39:881–6. [31] Saha AK, Das S, Bhatta D, Mitra BC. Study of jute fibre reinforced polyester composites by dynamic mechanical analysis. J Appl Polym Sci 1999;71:1505–13. [32] Lopez MA, Biagotti J, Kenny JM. Comparative study of the effects of different fibres on the processing and properties of ternary composites based on PP-EPDM blends. Polym Compos 2002;23:779–89. [33] Mohanty S, Verma SK, Nayak SK. Dynamic mechanical and thermal properties of MAPE treated jute/HDPE composites. Compos Sci Tech 2006;66:538–47. [34] Pothan LA, Oommen Z, Thomas S. Dynamic mechanical analysis of banana fiber reinforced polyester composites. Compos Sci Tech 2003;63:283–93. [35] Joseph PV, Mathew G, Joseph K, Groeninckx G, Thomas S. Dynamic mechanical properties of short sisal fiber reinforced polypropylene composites. Composites: Part A 2003;34:275–90. [36] Crevecoeur G, Groeninckx G. Morphology and mechanical properties of thermoplastic composites containing a thermotropic liquid crystalline polymer. Polym Eng Sci 1990;30:532–42. [37] Machiels AGC, Denys KFJ, Van Dam J, Posthuma De Boer A. Effect of processing history on the morphology and properties of polypropylene/thermotropic liquid crystalline polymer blends. Polym Eng Sci 1997;37:59–72. [38] Fuchs C, Bhattacharyya D, Friedrich K, Fakirov S. Application of Halpin–Tsai equation to microfibril reinforced polypropylene/poly (ethylene terephthalate) composites. Compos Interf 2006;13:331–44. [39] Halpin JC, Kardos JL. The Halpin–Tsai equations: a review. Polym Eng Sci 1976;16:344–52.
Composites: Part A 39 (2008) 164–175 www.elsevier.com/locate/compositesa
Morphology, static and dynamic mechanical properties of in situ microfibrillar composites based on polypropylene/poly (ethylene terephthalate) blends K. Jayanarayanan a, Sabu Thomas b, Kuruvilla Joseph c,* a
Department of Chemical Engineering and Materials Science, Amrita Vishwa Vidyapeetham, Coimbatore 641 105, Tamil Nadu, India b School of Chemical Sciences, Mahatma Gandhi University, Priyadarshini Hills P.O, Kottayam 686 560, Kerala, India c Department of Chemistry, Indian Institute of Space Science and Technology, ISRO P.O, Thiruvananthapuram 695 022, Kerala, India Received 23 June 2007; received in revised form 5 September 2007; accepted 18 November 2007
Abstract In situ composites based on blends of polypropylene (PP) and poly (ethylene terephthalate) (PET), were prepared by melt extrusion, followed by continuous drawing and isotropization. The blending of the mixture was carried out in a single screw extruder and the isotropization of the stretched blend was achieved by injection moulding. Scanning electron microscopy (SEM) studies showed that the extruded blends were isotropic, but both phases became highly oriented after drawing (stretching). The stretched blends were converted into in situ composites after injection moulding at temperatures below the melting point of PET. The size of the PET fibrils generated decreased with increasing stretch ratio. The tensile and impact properties were found to be higher for the samples drawn at stretch ratios 5 and 8. Dynamic mechanical properties such as the storage modulus (E0 ), loss modulus (E00 ) damping behavior (tan d) and static mechanical properties were investigated as a function of stretch ratio. The E0 values were found to be increasing up to a stretch ratio 8. The tan d and E00 modulus spectra showed a strong influence of the microfibrils on the a and b relaxation of PP. Finally, the tensile properties obtained experimentally were compared with those determined using theoretical equations. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Polypropylene; A. Poly (ethylene terephthalate); B. Morphology; C. Dynamic mechanical analysis
1. Introduction Blending of polymers is an effective method to generate newer materials offering better properties than the individual ones. The blends of polyolefins and poly (ethylene terephthalate) (PET) are of interest since the latter can improve the mechanical properties of the relatively weak partner. The usage of polypropylene as one of the phases in polymer blends is widely favored due to its excellent processability and low cost. Several studies were carried out on blends of PP and PET [1–5] to analyze the morphology
*
Corresponding author. Tel.: +91 471 2564808; fax: +91 471 2564806. E-mail address: [email protected] (K. Joseph).
1359-835X/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.compositesa.2007.11.008
and mechanical behavior. The incorporation of PET fibres into PP to convert them into thermoplastic composites was tried out by Lopez and Arroyo [6]. The incorporation of PET fibers is of interest as a nucleating agent for the crystallization of PP [7,8]. The presence of PET fibres encourages transcrystallization of PP as revealed by these studies. In situ composites were made from the blends of liquid crystalline polymer (LCP) and polypropylene [9–11]. After melt-blending LCP attain fibrillar morphology to provide excellent properties to the resultant blend. Evstatiev and Fakirov [12] introduced a novel concept of developing microfibrillar reinforced composites (MFC) from polymer blends. Only immiscible, fibre forming polymers with considerable difference in their melting points (Tm) are capable of getting converted into MFCs. There are essentially three
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steps involved in the preparation of polymer/polymer MFCs, viz: (a) blending of the component polymers, usually in an extruder, (b) drawing or orientation and (c) annealing or the isotropization step. In the first step, melt-blending of two or more immiscible thermoplastic polymers take place with or without the presence of a compatibilizer. In the second step, the extruded blend is either cold drawn or hot drawn to achieve fibrillar morphology for the constituent polymers. In the third step, the drawn blends are annealed or isotropized which destroys the fibrillar morphology of the polymer with lower Tm. The fibrils of the polymer with higher Tm are preserved and are randomly distributed in an isotropic matrix of the polymer with lower Tm. Since the matrix and the reinforcement are thermoplastic polymers and the reinforcements (fibrils of the polymer with the higher Tm) are generated in situ during the process they are rightly called as in situ reinforced polymer/polymer composites. The sizes (diameter) of the fibrils generated are in the order of few microns and hence the name microfibrillar composites (MFCs). In the early days of the development of MFCs, the fibrillization step was achieved in a separate step after the meltblending was completed [12–22]. Recently, a new method of continuous drawing of the blend online, downstream the extrusion die was developed for the fibrillization step. The polyolefinic materials were blended with polyamides and polyesters and were converted into MFCs [23,24]. Evstatiev et al. [25] suggested a scheme for production of MFCs, which could be scaled up to industrial standards, using LDPE/PET blends. They made use of recycled PET obtained from bottles as reinforcement for LDPE, and thereby suggested the scheme as an efficient method for the recycling of PET. Another method was employed by Li et al. [26,27] to produce PE/PET MFCs, where the blended extrudate in the form of a strip was stretched while they were still in the hot condition. The tensile strength and modulus increased with an increase in hot stretching ratio and beyond a critical value it reduced. Recently, Friedrich et al. [28] developed in situ composites at a constant draw ratio from PP/PET (60/40) blend in the presence and absence of a compatibilizer. The composites were developed by compression moulding as well as injection moulding, and it was observed that composites from compression moulding were found to have better mechanical properties since the uniaxial orientation of the PET fibrils are maintained. Shorter fibrils were obtained in the blends with higher concentration of the compatibilizer. This indicates that the compatibilizer although improves interfacial adhesion, reduces the aspect ratio of the fibrils. The dynamic mechanical properties of the in situ composites are not reported widely in the literature except may be for those using liquid crystalline polymers. The main objective of this work is to study the morphology development of the in situ composites prepared from PP and PET (85/15) blends obtained at different draw ratios and the evaluation of their static and dynamic mechanical properties. The dynamic mechanical properties of compos-
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ites are of significance since the relaxations of the polymer under the various conditions of temperature and stresses can be studied. Dynamic mechanical analysis (DMA) over a wide range of temperatures permits the determination of the viscoelastic behavior of polymeric materials. The stiffness and damping characteristics of the composites can be analyzed using DMA which could be co-related to the morphology. The dynamic mechanical properties analysis is a convenient method for estimating the polymer transition temperatures which can influence the fatigue and impact properties of the composite. 2. Experimental work The polymers used were PP (Repol-H110MA, Reliance, India, MFI – 11.0 g/10 min) and PET (940400-B, Futura Polymers, India, I.V.0.814 dl/g). After drying PET for 12 h at 100 °C it was tumble mixed with PP at a constant weight ratio of 15/85. The mixture was then melt blended in a single screw extruder (Screw diameter – 20 mm, L/D ratio – 30) using a set temperature profile of 225, 235, 250, 255 and 260 °C and the screw rotating at 30 rpm. The scheme of the experimental setup is given in Fig. 1. The extrudate was in the form of circular strands of diameter 2 mm and were taken to a cooling water bath for solidification. Then these strands were passed through a hot air oven maintained at 100 °C. Subsequently these strands were taken to a take up device for continuous drawing. The take up device consists of a pair of nip rolls whose peripheral velocity (V1) is maintained same as the velocity of the extrudate. Beyond the nip rolls there was a pair of stretch rolls of the same diameter as that of the nip rolls, but whose speed (V2) can be varied to attain different draw ratios and thereby reduction in the cross-sectional dimensions of the strands. The ratio of the speeds of the stretch rolls to the nip rolls (V2/V1) was taken as the draw (stretch) ratio. The experiment was carried out for draw ratios 1, 2, 5, 8, 10 and the specimens were named as neat blend (NB), H2, H5, H8 and H10, respectively. The term H indicates that the stretching was carried out in the hot condition of the blend. Downstream the stretching unit, the strands were granulated to approximately 3 mm length. The granules thus produced were injection moulded into tensile test and impact test samples in a Ferromatic Milacron-Sigma 50T machine. The extruded neat blend was injection moulded at a temperature profile of 185, 225, 245 and 270 °C from the feed zone to the nozzle and it was designated as NBI. All the drawn blends were injection moulded at a set temperature profile of 160, 170, 190 and 205 °C, from the feed zone to the nozzle. The samples taken were designated as H2I, H5I, H8I and H10I corresponding to the stretch ratios. The term I indicates the samples were isotropized by injection moulding technique. For example, H5I indicates that PP/PET blend in the ratio 85/15 was hot stretched at a V2/V1 of 5 followed by the isotropization by injection moulding technique.
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V2
V1
Extruder Cooling water bath
Nip rolls
Oven
Stretch rolls
Grinder
Drawn Blend
Injection moulding machine
Fig. 1. Scheme of the experimental setup for microfibrillar in situ composites preparation.
3. Characterization methods 3.1. Morphological analysis A JEOL JSM 840 SEM with an acceleration voltage of 20 kV was used for studying the morphology of the specimens. To extract the PET phase from the specimens, a mixture of phenol/1, 1, 2, 2, tetra chloroethane in 60/40 wt.% was used as the solvent. Similarly to remove PP and retain the PET fibrils the specimens were treated with hot xylene. All the specimens were coated with a thin gold layer prior to SEM analysis. The diameter of about 150 fibrils was measured from the micrographs of stretched samples using image analysis software. 3.2. Mechanical properties The tensile properties of the in situ composites at different stretch ratios were measured at room temperature according to ASTM D-638 using the injection moulded dumbbell specimens. All the tests were conducted at a constant crosshead speed of 50 mm/min. Impact Strength was measured using an Izod impact tester using notched specimens according to ASTM D-256. The flexural properties of the samples were measured according to ASTM D-790 using injection moulded rectangular specimens. Five specimens were tested in each case and the average values were reported. 3.3. Dynamic mechanical properties Rectangular specimens having size 60 mm 13 mm 3.3 mm were used for the dynamic mechanical experiments. Dynamic mechanical thermal analyzer NETZSCH DMA 242 was used for the evaluation of storage modulus (E0 ), loss modulus (E00 ) and mechanical damping factor (tan d). Three point-bending modes were used. The temperature
range over which properties were measured was 20 to 150 °C at a heating rate of 5 °C/min. The tests were carried out at a frequency of 1 Hz. 4. Results and discussion 4.1. Morphology development of in situ composites The morphology of the blend after extrusion and drawing are shown in Fig. 2. As evident from Fig. 2a, after extrusion, the neat blend (stretch ratio 1) indicates typical incompatible blend morphology with discrete domains of minor component dispersed in a continuous phase of the major component. The PET phase in the form of spherical and elliptical particles with sectional dimensions varying from 1 to 6 lm is found to be distributed in the major phase PP. The samples obtained after drawing with ratios 2, 5, 8 and 10 show orientation and fibrillation of the PET phase. Large bunches of continuous fibres of PET are observed in the micrographs (Fig. 2b–e) for the samples H2, H5, H8 and H10, respectively. In the micrographs of the stretched blends, any apparent loss of orientation is due to the floating effect of the fibrils in the solvent during extraction of PP phase and drying. Also some residual PP phase is seen as attached to the fibrils. The fibril diameter after stretching was estimated, however, the aspect ratio could not be analyzed due to the twisted and curly nature of the fibrils. The PET fibrils has a mean diameter of 8.6 lm in the case of H2. When the stretch ratio is increased to 5, the mean diameter reduces to 4.9 lm and at stretch ratio 8 it is 4.1 lm. At a stretch ratio of 10, the mean diameter increases to 6.9 lm. A careful examination of the SEM micrograph (Fig. 2e) showed the breakage of the fibrils which is indicative of the large extensional deformation the blend being subjected to at draw ratio of 10. The increase in mean diameter at draw ratio 10 is due to the presence of some large fibrils.
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Fig. 2. SEM images of samples after extrusion and drawing (a) neat blend after extrusion with PET phase extracted, (b–d) drawn blends at stretch ratios 2, 5 and 8, respectively, with PP phase extracted and (e), (f) drawn blend at stretch ratio 10 with PP phase extracted.
SEM analysis and particle size estimation of another specimen of H10 (Fig. 2f) confirmed the presence of large and small diameter fibrils (ranging from 1 to 10lm). It was also observed that many of the smaller diameter fibrils were discontinuous due to the breakage. It has been reported in the case of in situ composites based on liquid crystalline polymers [29] and PE/PET MFCs [27], at very high draw ratios the elongated molten droplets cannot withstand the large extensional deformation and break up into smaller entities leading to the reduction in aspect ratio. At a stretch ratio of 10, the PET phase does not get sufficient time to complete the orientation and get transformed into continuous fibres. At the same time, large extensional deformation forces at this stretch ratio results in breakage of many fibrils. It should be reported here that
there was physical breakage of the extruded strands at draw ratios of 12 and experiments could not be carried out at higher draw ratios. Injection moulding at temperatures above the melting point of PP but below that of PET leads to the melting of PP phase, but the fibrillar morphology of PET is preserved. After injection moulding, PET fibrils loose their orientation and is randomly distributed in the PP matrix. Since these fibrils are exposed to temperatures above the melting point of PP during injection moulding coupled with the high shear rate, there is a reduction in the diameter of the PET microfibrils. This phenomenon is due to the ‘break up behavior’ [30] of the fibrils during relaxation at elevated temperatures which is schematically represented in Fig. 3. First the long fibril breaks up into a number of
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A
B
C
A
Long Fibril
B
Short Fibrils
C
Spheres
Fig. 3. Break up behavior of fibrils during injection moulding.
shorter ones, then the short rod is pulled back to spheres by interfacial tension. During isotropization, the break up of the long fibrils should be limited to the formation of short fibrils to complete the production of in situ composites, rather than going back to the spherical shape. The SEM micrograph (Fig. 4a) of injection moulded samples from neat blend (NBI) does not reveal any presence of fibrils as expected. The SEM micrograph of H2I (Fig. 4b) shows short fibrils with cross-sectional dimensions varying from 7 to 9 lm. The number of fibrils obtained after isotropization are less which can be attributed to (a) the lower conversion rate of the PET phase into fibrils during stretching since the draw ratio is less and (b) the break up behavior of the fibrils during isotropization leading to the formation of spherical entities rather than short rods. For H5I, microfibrils of about 5 lm diameter are observed. The aspect ratio of the fibres is relatively higher as evident from the micrograph (Fig. 4c). In the case of H8I (Fig. 4d), fine fibrils of about 3 lm, but of a lower aspect ratio in comparison with H5I, which are strongly attached to the matrix, are observed. For H10I, very short fibrils of approximately 4 lm diameter are seen (Fig. 4e). It should be noted that for H10, breakage of the fibrils was observed and that could be an additional reason along with the break up behavior for the reduction in the length of the PET fibrils in H10I. The variation of PET microfibril diameter with stretch ratio after extrusion and drawing and after isotropization (injection moulding) is delineated in Fig. 5. Based on the above studies a scheme for the morphology development of the blend during drawing and injection moulding is presented in Fig. 6. After melt extrusion at 270 °C, the extrudate is cooled and then raised to the orientation temperature of 100 °C for continuous drawing. The fibrils of PP and PET are highly oriented at this stage as represented in Fig. 6b. In the next step, during injection moulding at 205 °C, which is higher than the Tm of PP but lower than that of PET, PP phase melts and forms the isotropic matrix for the composite. The PET microfi-
brils which are highly oriented prior to injection moulding do not melt at this temperature. However, they loose their orientation to get randomly distributed in the PP matrix as represented in Fig. 6c. This scheme has been widely reported [22,25] as the route to develop microfibrillar composites from two individual polymers. 4.2. Static mechanical properties The tensile, flexural and impact properties of PP, NBI and all the in situ composites are given in Table 1. The low tensile strength of NBI in comparison with PP can be attributed to the spherical domains of the PET phase in the blend which remain incompatibile with the PP phase. The tensile strength of H2I is slightly higher than NBI but still lower than PP. The tensile strength of H5I is the maximum amongst all the samples analyzed. H8I shows a slight reduction in the tensile strength in comparison with H5I. Further increase in stretch ratio leads to reduction in the tensile properties. The tensile modulus is found to be maximum for H5I. There is an increase of 41% in the tensile modulus of H5I when compared to that of NBI. H8I also exhibits good tensile properties but slightly lower than H5I. The influence of stretch ratio on the tensile strength and tensile modulus of the blend is presented in Fig. 7. It can be inferred that the tensile strength and tensile modulus of the in situ composites increases with stretch ratio, reaches a maximum value between stretch ratios 5 and 8 and then reduces. The stress–strain curves of NBI, H2I, H5I, H8I and H10I are plotted in Fig. 8. It could be seen from the plot that the neat blend (NBI) yields and fails at a low stress indicating that PET phase in the form of spherical particles are not able to act as reinforcement for the PP matrix. H2I yields and fails at a higher stress in comparison with H10I. It should be remembered that the microfibrils obtained in the case of H2I were longer than those obtained in the case of H10I. H5I and H8I exhibit a relatively ductile failure in comparison with NBI, H2I and H10I. From Fig. 8, we
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Fig. 4. SEM images of injection moulded (isotropized) blends a) injection moulded neat blend with PP phase extracted, (b–e) isotropized drawn blends at stretch ratios 2, 5, 8 and 10, respectively with pp phase extracted.
could see that necking is appreciable in the case of H5I and H8I as evident from the strain these samples are able to take up beyond the yield stress before failure. High yield stress followed by a pronounced necking region makes H5I and H8I samples stiffer and tougher. H10I exhibits poor tensile properties, which may be attributed to the low aspect ratio of the fibrils as evidenced from the micrographs. The tensile modulus of all the in situ composites was found to be higher than PP and NBI. It was reported recently [27] that the tensile properties of HDPE/PET microfibrillar composites increased as the stretch ratio increases and then tends to decrease at very high stretch
ratios. High tensile modulus coupled with high elongation at break as exhibited by H5I and H8I indicate the improvement in the toughness of the composites in comparison with NBI. The length of the PET microfibrils were maximum in the case of H5I and H8I as evident from the corresponding SEM images. They are acting as excellent stress transfer agents, which contribute to the enhancement in the toughness of H5I and H8I. Another reason for the improvement of the tensile and impact properties could be the development of a trancrystalline layer [8] of PP around the PET fibrils especially in the case of H5I and H8I which can encourage better adhesion between the phases.
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a
pared to NBI. There is an improvement in the flexural properties of the in situ composites with stretch ratio. However, beyond stretch ratio 5, flexural strength and modulus decrease. H10I exhibits poor flexural properties, which indicates the ineffectiveness of short PET microfibrils as reinforcement for PP. The impact strength of H8I is the maximum amongst all the samples followed by H5I. For H8I, microfibrils with the least diameter were obtained as observed from the micrographs. Further, the relatively continuous nature of fibrils allow better stress transfer from the matrix polymer (PP) during impact. H10I exhibits poor impact properties since the short fibrils cannot act as effective stress transfer agents. It should also be mentioned that the impact strength is the lowest for NBI, which is indicative of the immiscibility of the two polymers.
Fibril dimater (microns)
10
8
6
4
2
0 0
Fibril diameter (microns)
b
3
6 Stretch ratio
9
12
10
4.3. Dynamic mechanical analysis
8
The dynamic mechanical analysis of the injection moulded samples was carried out from 20 to 165 °C at 1 Hz.
6
4.3.1. Storage modulus The storage modulus E0 obtained from the DMA is closely related to the load bearing capacity of the material and when the experiment is carried out in a three point-bending mode, it is similar to the flexural modulus measured as per the ASTM D-790 [31]. In Fig. 9, variation of E0 with temperature for PP, NBI, H2I, H5I, H8I and H10I at 1 Hz is shown. At low temperatures, E0 values of PP, H2I, H5I and H8I are very close to each other. It is an indication, that, at low temperatures the fibrils do not contribute much in imparting stiffness to the composite. In the case of PP, there is a sharp fall in E0 on passing through the glass transition temperature (Tg). It is due to the increased molecular mobility of the polymer chains above Tg. There is a large fall in modulus with increasing temperature for PP in comparison with other samples, the stiffness at high temperature being determined by the amorphous regions, which are very submissive above glass transition temperature.
4
2
0 0
3
6
9
12
Stretch ratio
Fig. 5. Variation of PET microfibril diameter with stretch ratio (a) after extrusion and drawing and (b) after isotropization (injection moulding).
There is considerable improvement in the flexural properties of the in situ composites in comparison with PP and NBI. The flexural strength of H5I was found to be increased by 20% and flexural modulus by 50% when com-
a
b
c PP isotropized
PP
Extrusion / Drawing
Injection Moulding
270°C/100°C
205°C
PET PETMicrofibril
Fig. 6. Morphology development scheme of the blend during extrusion and injection moulding.
K. Jayanarayanan et al. / Composites: Part A 39 (2008) 164–175
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Table 1 Tensile, flexural and impact properties of PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios Sample
Tensile strength (MPa)
Tensile modulus (MPa)
Elongation at break (%)
Flexural strength (MPa)
Flexural modulus (MPa)
Impact strength (J/m)
PP NBI H2I H5I H8I H10I
29.2 23.3 27.9 32.4 31.5 27.1
1470 1514 1662 2137 1927 1642
33.9 24.2 28.6 34.1 32.1 27.8
36.4 39.1 41.9 47.1 44.4 38.3
742 673 787 1015 837 735
26.1 21.7 25.6 30.6 36.7 25.2
The drop in the modulus on passing through the Tg is much less for H5I and H8I in comparison with PP. Even for NBI, H2I and H10I the difference between the modulii of the glassy state and rubbery state is smaller when compared to PP. The variation of storage modulus with stretch ratio at different temperatures is delineated in Fig. 10. At temperatures above the Tg of PP, the storage modulus increases with draw ratio, reaches a maximum value between the stretch ratios 5 and 8 and then declines. This substantiates the SEM observations, where long microfi-
a Tensile strength ( M Pa)
32
28
24
brils of PET were obtained for H5I and H8I. The longer microfibrils contribute in imparting stiffness to the composite. The influence of draw ratio could further be explained by studying the normalized storage modulus with temperature at different draw ratios as depicted in Fig. 11. Normalized storage modulus can be defined as the ratio of the storage modulus of the composite to the storage modulus of the matrix at the same temperature. It is evident from the plot that as the draw ratio increases the normalized storage modulus increases, reaches a maximum value and then decreases. Also at a particular draw ratio the normalized storage modulus increases with increase in temperature. This indicates the restriction of the mobility of the PP molecules due to the presence of the PET micro fibrils at elevated temperatures. The shorter microfibrils obtained in the case of H10I are not as efficient as those from H5I and H8I in restricting the segmental mobility of PP chains. The effect of the microfibrils as a reinforcing agent is appreciable at 50 °C, beyond which PET becomes flexible on approaching its Tg.
35
20 0
3
6
9 30
Stretch ratio
b
Stress (MPa)
Tensile modulus (M Pa)
2500
2100
25
NBI
20
H2I
H5I
15
H8I
10
1700
H10I
5
0
1300
0
0
3
6
10
20
30
40
50
60
9
Stretch ratio
Fig. 7. The influence of stretch ratio on (a) tensile strength and (b) tensile modulus for PP/PET (85/15) blend.
Strain (%)
Fig. 8. Stress–strain curves for injection moulded neat blend and microfibrillar in situ composites prepared at different stretch ratios.
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1.8 PP NBI
E' (MPa)
H2I H5I
6000
H8I H10I
4000
1.6
Normalised E'
8000
1.4
1.2
1 2000
10ºC
25ºC
50ºC
80ºC
0.8 1
3
0 -20
0
20
40
60
80
100
120
5
7
9
11
Stretch Ratio
140
Fig. 11. Variation of normalized storage modulus of microfibrillar in situ composites with stretch ratio at temperatures 10, 25, 50 and 80 °C.
Temperature (ºC)
Fig. 9. Variation of storage modulus with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
4.3.2. Loss modulus The loss modulus E00 is the viscous response of viscoelastic materials. It is a measure of the energy dissipated or lost as heat per cycle of sinusoidal deformation, when different systems are compared at the same strain amplitude. Fig. 12 shows the variation of loss modulus with temperature for PP, net blend and in situ composites at 1 Hz. In all the cases the value of the loss modulus increases, reaches a maximum value and then decreases. Polypropylene exhibits three relaxations localized in the range of 80 °C (c), 10 °C (b) and 100 °C (a). The loss modulus value is maximum for PP at Tg (b transition) when compared to the blend and in situ composites. There is a steady fall of the modulii
beyond the Tg for PP and NBI. The loss modulus of H5I and H8I beyond 35 °C is found to be the highest which is an indication of the improved energy dissipation, which could contribute to their better impact properties. In the case of H10I, viscous dissipation is poor as evident from low loss modulus values in comparison with H5I and H8I. The higher loss modulus on approaching the a transition is an evidence for the restriction on the segmental mobility of the polymer molecules [32]. Generally, it is considered that the lowering and the broadening of the loss modulus peak is an indication of the greater constraints on the relaxation of the amorphous phase [33,34] and thereby improvement in the mechanical properties. But it has also been reported that higher E00 values at Tg is an indication of the good interfacial adhesion
8000 900 PP
750
6000
NBI H2I H5I
E" (MPa)
E' (MPa)
600
4000
H8I
450 H10I
300
2000 150
25ºC
10ºC
50ºC
80ºC
0
0
1
3
5
7
9
11
Stretch ratio Fig. 10. Variation of storage modulus of microfibrillar in situ composites with stretch ratio at temperatures 10, 25, 50 and 80 °C.
-20
10
40
70
100
130
Temperature (°C)
Fig. 12. Variation of loss modulus with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
K. Jayanarayanan et al. / Composites: Part A 39 (2008) 164–175
between matrix and the fibre [35]. In this particular case, the lowest peak at Tg was obtained for H10I which exhibited poor impact properties. H2I has the highest E00 values at Tg amongst the in situ composites, but it should also be noted that values beyond 30 °C is lower than that of H5I and H8I. Hence it has to be inferred that the increased loss modulus on approaching the a transition contributes much more than the lowering of the E00 peak towards the mechanical properties.
173
of the chains of the matrix polymer. Elevation of Tg is taken as a measure of the interfacial interaction. The intensity of the tan d peak is maximum for neat PP and progressively it is found reducing with the increase in draw ratio. The damping peak in the in situ composites showed a decreased magnitude of tan d in comparison to PP and NBI. The damping peak is associated with the partial loosening of the polymer structure so that groups and small chain segments can move. The broadening of the tan d peak is appreciable for H5I and H8I which is due to the restricted relaxation. It is interesting to study the behavior of the samples at the a relaxation. The peak heights of the in situ composites are more than that of PP. It can be attributed to the higher loss modulus values of these samples at this temperature. There is a positive shift in the a relaxation temperature (Ta) of PP phase for the in situ composites in comparison with neat polymer. Ta for neat PP is 93.5 °C and for NBI it is 89 °C. For the in situ composites H2I, H5I, H8I and H10I Ta is 101.9, 102.8, 103 and 95.25 °C, respectively. The shift in Ta in the case of H10I is marginal (1.7 °C) when compared to all other in situ composites which indicates interfacial interaction is poor at this draw ratio. This fact is consistent with the interpretation of the loss modulus curve and it is further substantiated with the experimental results.
4.3.3. Tan d The ratio of loss modulus to storage modulus is measured as the mechanical loss factor or tan d. The damping properties of the material give the balance between the elastic phase and viscous phase in a polymeric structure. The damping behavior of the composites in the transition region is governed by (a) mechanical relaxation of the matrix and the fibre, (b) the interface between the fibre and the matrix and (c) fibre loading and fibre length. In the present investigation, the variation of tan d as a function of temperature at 1 Hz is represented in the Fig. 13. Since the Tg of PET is higher than 60 °C, the effect of PET fibrils on the relaxation of PP phase can be easily analyzed. Neat PP has a Tg of 10.5 °C as obtained from the tan d curve. For NBI, there is only a marginal increase to 11.4 °C. The Tg value is maximum for H8I which is 21 °C and it is 10.5 °C more than the neat PP. This result is in contrast with one of the earlier studies [6] where PP was mixed with chemically treated PET fibres. They reported that Tg of PP in the composites produced with treated PET fibres which when introduced separately was found to be less than that of neat PP. But, it is widely agreed [33–35] that an increase in the temperature of the tan d peak is indicative of the restriction of the mobility
4.4. Theoretical prediction of tensile properties Halpin–Tsai equation has been proved useful in predicting the longitudinal modulus of aligned short-fiber composites, e.g., in situ composite with liquid crystalline polymer [36,37]. Li et al. [26] and Fuchs et al. [38] tried to apply this equation in predicting the tensile properties
0.16
PP
NBI
H2I
H5I
H8I
H10I
0.14
Tan delta
0.12 0.1 0.08 0.06 0.04 0.02 0 -20
0
20
40
60
80
100
120
140
160
Temperature (ºC) Fig. 13. Variation of mechanical damping factor with temperature for PP, neat blend and microfibrillar in situ composites prepared at different stretch ratios.
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of MFCs. Here, both the fiber modulus (Ef) and the fiber aspect ratio (L/D) are dependent on the stretching ratio and the dispersed phase content of the extrudate. According to Halpin–Tsai model [39], the Young’s modulus, Mc, and the tensile strength, Tc, of a composite are given by the following equations: 1 þ Ag1 V f Mc ¼ Mm ð1Þ 1 g1 V f 1 þ Ag2 V f ð2Þ Tc ¼ Tm 1 g2 V f where g1 and g2 are calculated using the following equations: g1 ¼ g2 ¼
Mf 1 Mm Mf þA Mm Tf 1 Tm Tf þA Tm
ð3Þ ð4Þ
The terms Mc, Mm, Mf and Tc, Tm, Tf are the tensile modulii and tensile strengths of the composite, matrix and fibre, respectively, and Vf is the fibre volume fraction. The value of A is twice the aspect ratio of the fibres. Halpin and Kardos [39] modified Eqs. (1) and (2) by including the term u which in turn is dependent on Umax, maximum packing fraction, of the reinforcement in the following manner: 1 þ Ag1 V f Mc ¼ Mm ð5Þ 1 ug1 V f 1 þ Ag2 V f Tc ¼ Tm ð6Þ 1 ug2 V f 1 Umax ð7Þ u¼1þ U2max The value Umax is 0.78 for square arrangement of fibres, 0.91 for hexagonal array of fibres and 0.82 for random packing of fibres. The value of Umax in this case was taken as 0.82 since the PET microfibrils are distributed randomly in the PP matrix. The value of A estimated using SEM images were 12.5, 64, 60 and 10, respectively, for H2I, H5I, H8I and H10I.The tensile strength and tensile modulus of PET fibrils were taken as 62 and 3250 MPa, respectively, for the calculations. The tensile strength and tensile modulus of PP (matrix) was experimentally obtained as 29.2 and 1470 MPa.
The tensile strength and modulus values obtained using Halpin–Tsai and modified Halpin–Tsai equations along with corresponding experimental results are given in Table 2.The theoretical values obtained by using both the equations are almost same for the in situ composites at a particular stretch ratio. The tensile strength value obtained by theoretical prediction for H8I shows a positive deviation of 8% from the experimental value, whereas in the case of H5I, the deviation is reduced to 5%.The tensile strength obtained experimentally for H2I and H10I are much lower than the values estimated theoretically. The tensile modulus values obtained theoretically are in close agreement with the experimental values obtained for H2I and H10I. In the case of H8I, the theoretical value for the tensile modulus shows a negative deviation of 12%, whereas in the case of H5I the negative deviation increases to 17%. The theoretical equations were found to be useful to model the tensile strength of H5I and H8I and tensile modulii of H2I and H10I. 5. Conclusion In situ composites were prepared from the blends of polypropylene and poly (ethylene terephthalate) by continuous drawing followed by injection moulding. The SEM observations show high level of orientation for the drawn blends. The PET fibrils with the lowest mean diameter of 4.1 lm were obtained at draw ratio of 8. After injection moulding at a temperature below the Tm of PET, fibrils with high aspect ratio and good orientation were obtained for samples drawn at stretch ratio 5. Beyond stretch ratio 8, the breakage of the fibrils were observed during stretching which produced very short randomly distributed fibrils after injection moulding.The tensile, flexural and impact properties were found increasing with stretch ratio. However, too high a stretch ratio is not desirable for the reinforcement of PP. The effect of temperature on dynamic mechanical properties like storage modulus, loss modulus and mechanical loss factor of PP, neat blend and in situ composites prepared at different draw ratios were studied. The E0 values were found decreasing drastically for neat PP beyond Tg. The presence of microfibrils showed a positive effect on the modulus at temperatures above Tg especially for the samples drawn at stretch ratio 5 and 8. The tan delta peak height at b transition (Tg) reduced and shifted to higher values with draw ratio. The maximum shift was obtained for samples drawn at a stretch ratio of
Table 2 Experimental and theoretical tensile strength, tensile modulus values obtained for in situ composites Sample
H2I H5I H8I H10I
Experimental
Halpin–Tsai
Modified Halpin–Tsai
Tensile strength (MPa)
Tensile modulus (MPa)
Tensile strength (MPa)
Tensile modulus (MPa)
Tensile strength (MPa)
Tensile modulus (MPa)
27.9 32.4 31.5 27.1
1662 2137 1927 1642
33.7 34.0 34.0 33.7
1718 1732 1731 1714
33.8 34.0 34.0 33.7
1718 1733 1732 1715
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