The structure of highly oriented polypropylene and its effect on physico-mechanical properties

Polymer Science U.S.S.R. Vol. 211,No. 10, pp. 2380-2386, 1986 Printed in Poland

0032-3950/86 $10.00+ .00 @) 1987 Pergamon Journals Ltd.

THE STRUCTURE OF HIGHLY ORIENTED POLYPROPYLENE AND ITS EFFECT ON PHYSICO-MECHANICAL PROPERTIES* A. A. TURETSKII,A. O. BARANOV,S. N. CHVALUN,N. A. YERINA,

Yu. A. ZUBOV,E. V. PRUT, N. F. BAKEYEVand N. S. YENIKOLOPYAN Institute of Chemical Physics, U.S.S.R. Academy of Sciences L. Ya. Karpov Physico-Chemical Research Institute

(Received 25 February 1985) The structure of highly oriented polypropylene was studied by X-ray, IR-spectroscopic and thermophysical methods. At elongations 2> 15, the crystallites and the polymer chains in amorphous regions were shown to be fully oriented in the strain direction, while the long period, and the longitudinal and transversal crystallite dimensions were independent of 2. With increasing 2, the crystallinity degree as determined by X-ray exhibits but insignificant growth, whereas that determined by DSC measurements increases by 30~. The effect of t h e structure of highly oriented polypropylene on the elastic modulus was analyzed. A model of the structure of highly oriented polypropylene is proposed.

TIlE production of materials with high modulus and high strength presently appears as one o f the m o s t i m p o r t a n t p r o b l e m s o f p o l y m e r physics a n d technology. The p r o blems c o n n e c t e d w i t h the c h a r a c t e r i z a t i o n o f the structure o f highly o r i e n t e d p o l y m e r s a n d with the effects o f s t r u c t u r e o n the p h y s i c o - m e c h a n i c a l p r o p e r t i e s have n o t yet been fully a n d clearly solved [1, 2], a n d the p r e s e n t w o r k is d e v o t e d to t h e i r study. P r e v i o u s l y [3] it h a s b e e n shown t h a t it is p o s s i b l e to o r i e n t p o l y p r o p y l e n e (PP) to elongations 2 ~ 3 0 . T h e elastic m o d u l u s d e t e r m i n e d in this case f r o m the initial slope o f the s t r e s s s t r a i n d i a g r a m s at 20 ° is equal to --~30 G P a , a n d f r o m a c o u s t i c m e a s u r e m e n t s to ,-~34" 8 G P a [4]. It has been suggested that such high values o f the elastic m o d u l u s c a n be exp l a i n e d by c h a r a c t e r i s t i c features o f the structure o f crystalline a n d a m o r p h o u s d o m a i n s [4, 5]. T h e r e f o r e we have d e c i d e d to study the structure o f highly o r i e n t e d PP by s t r u c t u r a l a n d t h e r m o p h y s i c a l m e t h o d s , a n d to d e t e r m i n e its c o r r e l a t i o n with m e c h a n i c a l c h a r acteristics. The studied PP samples have been characterized in [3]. The initial PP film, ~400 am thick, was obtained by extrusion with inflation [6]. From it, dumbbells were cut in the extrusion direction, to be isothermally stretched at 145 ° with the testing machine "Instron°l122", in two steps. In the first step, dumbbells of working length 35 mm and width 5 mm were elongated to 2=8-9; in the second step, strips 25 or 50 mm long were used, cut from the central part of the stretched dumbbells. The rate of clamp motion was constant and equal to 20 mm/min. The elongation 2 was determined by the measurement of the displacement of markers applied to the initial sample. After stretching the samples were cooled in the free state. * Vysokomol. soyed. A28: No. 10, 2141-2146, 1986. 2380

Structure of highly oriented polypropylone

2381

The macroscopic density of the samples was determined in gradient columns with water and ethanol, at 30 °, with an error of +0-001 g/cm 3. The degree of orientation was measured by the IR-spectroscopic method with the spectrometer "Digilab-FTS-15C", according to [7]. Crystallite orientation was determined by means of the band at 998 cm T M , the mean orientation of crystallites and of the chains in amorphous domains by means of the band at 975 cm- 1. The heat of melting was determined by the DSC method using the apparatus DSM-2M at a heating rate of 16 deg/'min. Sample weight was 5-6 mg. Both free and fixed samples were measured. In both these cases, the heat of melting was identical and only dependent on ,L In evaluating the calorimetric crystallinity value Wan of the studied samples, the heat of melting of crystalline PP, d H c , was taken as equal to 146"65 J/g [8]. For small-angle X-ray measurements, the apparatus KRM-I (CuK~-radiation, Ni-filter) was used. The long periods were determined by means of the Bragg formula from the maximum position on the meridional scattering curve. Intensity was normalized for sample thickness, absorption and primary beam intensity. The density difference of the crystalline and amorphous phases zip was determined by the absolute intensity method described in [9]. The photoroentgenograms of the samples were obtained using the apparatus URS-2 with CuK,-radiation. The sample-film distance (fiat camera) was 32.3 mm. The effective transversal crystallite dimension 11lo was determined from the integral half-width of the corresponding reflection using the apparatus DRON-2, according to [10]. The error in the determination of the angular position of the maximum was + 5 x 10- a deg, of the corresponding interplanar spacing d~ lo + 1 × x 10-2 A~,and of the crystallite dimension ltlo +_5 A. The longitudinal effective crystallite dimension was determined from the integral half-width of the 113 reflection after separation of the experimental line shape corresponding to the 113 and 123 reflections. The error of the 111a measurement was + 8 A.. The shape of the reflection was recorded by means of the apparatus DRON-1. The true integral reflection half-width was calculated by the method of [l l] using a standard. Besides calorimetric crystallinity, also the X-ray crystallinity of the samples, Wr, was determined, from the integral intensity of the 110 reflection [10]. The error of W, measurement was + 5 %. For highly oriented samples with 2 > 20, the determination of the integral intensity of the 110 reflection is difficult because of the large number of pores. This was indicated by the clouding of samples and by the high intensity of the small-angle diffusion scattering. The mechanical properties of the highly oriented PP films were studied with the testing machine 'Instron-1122" [3]. The elastic modul us E was evaluated for several strain values e from a stressstrain in diagram, from the initial slope at e=0-45%, and also at e = 1 and 2%.

It is well k n o w n that the stretching o f semi-crystall ine polymers leads to the format i o n o f fibrillar material, with high anisotropy of ph ysico-mechanical properties [1, 2]. W i t h PP, already at elongations 2 > 10, the crystallites appear almost fully o r i e n t e d in the d e f o r m a t i o n d i r e c t i o n (Fig. 1, curve 1). U n f o r t u n a t e l y , at smaller elongations 2, the degree of o r i e n t a t i o n could n o t be m e a s u r e d with sufficient accuracy by the I R m e t h o d because of large sample thickness. A t further elongation, the degree o f crystallite o r i e n t a t i o n cos 2 0c changes very little, while in the a m o r p h o u s d o m a i n s the degree of o r i e n t a t i o n increases. The change o f t h e m e a n d e g r e e o f o r i e n t a t i o n of the sample, cos 2 0m, (0 is the angle b e t w e e n the stretching direction a n d the chain axis c) in d~pendence on 2 is shown in Fig. I (curves 3 a n d 4). I n this case, for each value of 2, two values o f cos 2 0m wzre calculated, b : cause t w o values o f the angle between the t r a n s i t i o n m o m e n t a n d molecular axis are givert

:2382

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in the literature [7] for the respective IR band: atv=0 and 18°. The figure shows that for 2 i> 15, cos ~ 0m--*I, irrespective of the value of atv. At lower 2 values, the mean degree of orientation is higher in the former case than in the latter. Thus it may be stated that a t 2/> 15, both the crystallites and the polymer chains in the amorphous domains are almost fully oriented along the stretching direction.

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FIQ. 1. Effectof 2 on cosz 0c (1), Ag~/z (2) and cos2 Om(3, 4). ~to= 18 (1, 4) and 0° (3). FIG. 2. Plot of L (1), 11~0(2) and l~3 (3) vs. 2. The analysis of photoroentgenograms of oriented samples with various 2 values has shown that the presently studied PP has the at-modification structure [12]. The photoroentgenograms of the initial films exhibit a concentration of diffraction rings, characteristic of a PP texture with an axis parallel to the crystalline direction 201 [13]; this may be explained by partial polymer orientation in the process of film preparation (extrusion with inflation). The photoroentgenograms of the oriented samples exhibit a well-pronounced c-axis texture, and their orientation direction coincides with the texture c-axis. An increase in the elongation 2 is accompanied by improved crystallite orientation, as shown by the decrease in the azimuthal half-width dV~ of reflection 110 (Fig. 1, curve 2), although its change is quite small for the increase in 2 from 9 to 32"5. These results are in agreement with spectroscopic data. It was found that the identity period as determined from the inter-layer spacing dl t3 is independent of the strain ratio and is equal to 2.160+0.002 A. A comparison of the small-angle photoroentgenograms of samples with 2= 1 and 9 indicates that the reflection form changes from a ring pattern with concentration on the meridian to a line pattern, due to the transition from the initial spherulitic structure with some lamellae orientation, to a fibrillar structure characterized by the value of the long period L. The dependence of L, and also of the longitudinal (lt lo) and transversal (111~) crystallite dimensions on 2 is shown in Fig. 2. It appears that the long period L

2383

Structure of highly oriented polypropylene

increases on transition from lamellar to fibrillar structure, and does not change at further elongation. At 2 > 20, L can no longer be evaluated, due to the disappearance o f the smalhangle reflection on the meridional scattering curve. The value of the transverse crystallite dimension decreases at the transition from 2 = 1 to 2 = 9 , and becomes inde-

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FiG. 3. Effect of 2 on Ap, as determined by the methods of absolute (l) and relative (2) intensities (a), and the curves of small-angle meridional scattering of PP samples with 2= 1 (1), 15 (2) and 20 (3) (b). pendent of 2 in highly oriented samples (Fig. 2, curve 2). The value of the longitudinal crystallite dimension, ll 1~-, increases at first, and becomes practically constant at 2 > 12. It should be noted that with PP at all strain ratios, 1ll3 does not exceed L, in contrast to highly oriented H D P E [10]. One of the important characteristics of highly oriented semicrystalline polymers is the density difference of the crystalline and amorphous phases zip. In this work, the dependence of Ap on 2 was determined by two methods: by the absolute and by the relative intensities (Fig. 3a). The absolute intensity method (AI) is based on the mean square fluctuation of electron density along the fibril, (At/) 2, which is connected with the scattering intensity detected at the scattering angle 20 by a slit parallel to the fibril axis; from this the density difference of the crystalline and amorphous phases zip can be calculated [14]. The obtained dependence is shown in Fig. 3a. The intensity distribution of small-angle meridional scattering in samples with various strain ratios is shown in Fig. 3b. It is apparent that the intensity of the small-angle reflection, which primarily depends on zip [15], decreases with increasing 2. It is interesting to compare the values of Ap as determined by the AI method and by the measurement of the relative intensity of the small-angle reflection. Figure 3a shows that these values of zip practically coincide for 2 = 9 and 20. The dependences on 2 of the crystallinity degrees as determined by the X-ray (W,) and calorimetric methods (Wan) are shown in Fig. 4. The crystallinities of the initial films as determined by the IR method and the macroscopic density values are also shown in the same figure. The sample density, p=0.907 g/cm 3, remained unchanged with in-

etaL

2384

A.A. TURETSKll

creasing elongation 2. Evidently such p(2) dependence is connected with inner pore formation or stretching. The calorimetric crystallinity degree increases monotonously from 60 to 90 % with increasing 2, while the corresponding increase in Wr is insignificant (Fig. 4). The difference in the course of the 2 dependences of Wr and may be connected with the

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FI6. 4. Plot of crystallinity degreee W, as determined by X-ray (1) and calorimetric (2) methods vs. 2. A -IR-spectroscopic method, B - W as determined from polymer density. peculiar structure of the amorphous domains at high 2. Thus with increasing 2, in the amorphous domains a large number of extended macromolecules are formed, connecting the neighbouring crystallites. Bundles of such chains will substantially contribute to without affecting Wr. Contrary to the planar chains of HDPE, the helical PP molecules are less probable to form extended coherent linear systems, including both chains in the crystallites and in tie segments in amorphous domains, of dimensions exceedingL. The small probability of the formation of such systems is caused, first [16], by the possible change of the sense of rotation at the phase boundary in the fibril; second, by the extension of the helical PP molecule in amorphous domains, with a change of the identity period, due to the relatively low elastic modulus of the PP chain as compared to HDPE. In the process of o.rientation stretching, some chains may be drawn out of the crystallite, causing defects in the crystal lattice and disturbing the close packing of the chain monomer units. Actually, at increasing elongation 2 the reflection 110 was found to be displaced to smaller angles, corresponding to an increase in the interlayer spacing dido.

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! 14.053 6.300

8'9 14'042 6"310

14"0 14.040 6-310

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32"5 13.997 6.330

The existence of a large number of extended chains in the amorphous domains of highly oriented PP samples is confirmed by the appearance of one-dimensional diffraction on the photoroentg.~nog:ams, i.e. of a relatively uniform intensity distribution along the layer line. The increased number of extended tie macromolecules in the amorphous domains and their high orientation lead to the growth of the elastic modulus E and of the ultimate

2385

Structure of highly oriented polypropylene

tensile strength trB with increasing elongation 2 (Fig. 5). It appears that up to the extension 2,-, 25 the values of E, trB and strain at break, eB undergo considerable changes, and at 2 > 2 5 they are practically constant. It should b,~ noted that in [3] the E0.) dependence has a somewhat different appearance, because the samples at high strain ratios were cooled in the fixed state. At high 2 values the samples begin to cloud, they become opaque due to pore formation caused by "cold flow" of the polymer. Also the considerable scatter of the E values (Fig. 5, curve I) is evidently connected with the "cold flow".

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F I G . 5. Plot of E (1), trB (2) and eB (3) vs. 2. FIG. 6. Plot of E calculated for various deformation values e, vs. 2. e.=0.45 (/), 1.00 (2) and 2.00% (3). The points correspond to E values given in [17].

Th': plots of E vs. 2 calculated for various e fi'om a-e diagrams are shown in Fig. 6. The Figu,:e shows a considerable decrease of E with increasing e. A comparison of the :esults shown in Fig. 6 (curve 1) with literature data indicates reasonably good agreement up to 2 ~ 2 6 30. The value of % at high ,;. also is in good agreement with the results of [17]. Thus the data of the present work indicate that the number of extended tie chains in amorphous domains of oriented PP increases with the strain ratio, resulting in an increase o f the elastic modulus E, and of the ultimate tensile strength ot~, and in a decrease of the strain at break, eB.

Transhttedby D. DOSKOCILOV,~

2386

A.A. TURETSKIIet al. REFERENCES

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