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Mechanical properties of polymer compositions based on polyethylene and prepared by polymerization in situ
Mechanical properties of polymer compositions based on PE
933
REFERENCES i. E. BUNCEL and A. DAVIES, J. Chem. Soc., 7, 1550, 1958 2. N. A. IZMAILOV, L. L. SPIVAK and V. I. LEVINKOVA, Uchenye zapiski K h G U (Academic Notes of Univ. of Kharkov) p. 30, 1950 3. N. P. SLUCHEVSKAYA, V. A. YABLOKOV, N. V. YABLOKOVA and Yu. A. ALEKSANDROV, Z k obshch, khimii 46: 1540, 1976 4. Yu. A. ALEXANDROV, N. V. YABLOKOVA and V. V. GORBATOV, Basic Research in Homogeneous Catalysis, vol. 3, p. 359, Univ. Coll. Station Texas, N. Y.-London, 1979 5. R. FOSTER and C. A. FYCE, Chem. Commun. 3: 642, 1965 6. V. P. ZUBOV and V. A. KABANOV, Khimiya i tekhnol, vysokomol, soyed., VINITI, Moscow, 9: 56, 1977 7. ~ . C. A P P L E T O N and J. TYRREL, J. Phys. Chem. 81: 1201, 1977 8. H. HIRAI and M. KAMIYAMA, J. Polymer Sci. Polymer Chem. Ed. 12: 2701, 1974
Polyrve ScienceU.S.S.R. Vol. 27, No. 4, pp. 933-941, 1985 Priqt.z~' in Poland
0032 3950/85 $10.00+.00 ~(; 1986 Pergamon Press Ltd.
MECHANICAL PROPERTIES OF POLYMER COMPOSITIONS BASED ON POLYETHYLENE AND PREPARED BY POLYMERIZATION IN SITU* A. L. VOLYNSKII, A . SH. SHTANCHAYEV, V. D. ZANEGIN, V. I. GERASIMOV a n d N. F. BAKEYEV Lomonosov State University, Moscow
(Received 21 August 1983)
A study has been made of polymer compositions prepared via polymerization of a runlber of liquid monomers in a PE matrix (polymerization in situ). The compositions have high strength and elasticity values if the polymeric component incorporated in PE is in the g'assy state. Structural changes accompanying nonelastic deformation of the compositions e~ e r a wide range of temperature have been investigated by X-ray analysis. A deformation ~'L~chanism is proposed on the basis of experimental results.
l-r is k n o w n [1,2] t h a t u n i a x i a l d r a w i n g o f high d e n s i t y P E ( H D P E ) films in a c o m p a t i b l e l i q u i d causes m a r k e d l y i n c r e a s e d swelling o f t h e p o l y m e r . P o s t p o l y m e r i z a t i o n o f a low molecular weight component
in t h e P E m a t r i x results is n o v e l p o l y m e r i c m a t e r i a l s
b a s e d on P E [3]. O u r a i m in t h e p r e s e n t w o r k was to i n v e s t i g a t e s o m e m e c h a n i c a l p r o p * Vysokomol. soyed. A27: No. 4, 831-837, 1985.
934
A . L . VOLYNSKIIet al.
erties o f m a t e r i a l s o f t h i s type b a s e d on H D P E a n d P M M A , PS o r p o l y - n - b u t y l m e t h acrylate (PBMA). The study objects were commercial HDPE films of thickness 50 ltm. The films were drawn in liquid monomers containing an initiator (benzoyl peroxide) and, in some cases, a crosslinking agent (ethylene glycol dimetharcylate (EGD)). Drawing to 100, 200, 300 and 400 ~ extensions was with the aid of specially designed equipment. After stretching, the lengths of samples were determined in stainless-steel frames, and the samples were transferred to a heat-controlled vessel containing an inert liquid heated to 80 °. Polymerization was conducted in this vessel for 8 hr. After polymerization the samples were annealed under isothermal conditions at 100 ° for 1 hr. Details of the experimental procedure and data on the composition of the resulting materials, along with results of the X-ray analysis were published in [3]. Mechanical testing was carried out with the aid of an Instron tensile machine at a constant rate of stretch (20 ram/rain). A s was n o t e d above, m o n o m e r is i n c o r p o r a t e d in the H D P E d u r i n g the s t r e t c h i n g o f the polymer. A c c o r d i n g l y P E phase in the f o r m i n g m a t e r i a l will be m o r e o r less oriented. T o shed light on the extent to which m e c h a n i c a l p r o p e r t i e s o f the m a t e r i a l are influenced by the o r i e n t e d p o l y m e r m a t r i x let us first o f all e x a m i n e the m e c h a n i c a l b e h a v i o u r o f H D P E stretched in a plasticizing liquid (heptane).* A f t e r stretching in the liquid m e d i u m lengths o f the H D P E samples were d e t e r m i n e d a n d the l a t t e r were d r i e d in air a n d a n n e a l e d u n d e r i s o m e t r i c c o n d i t i o n s for 1 h r ar 100 °, i.e. u n d e r w e n t h e a t t r e a t m e n t similar to t h a t to which the o b t a i n e d (composite) m a t e r i a l s were subjected.
MPa
q 0 80
120 ~.3 80
~//~2
qO
qo 0
1 I
I
200
[
I
400 2,%
I
0
I
100
I
I
200
f
~%
FIG. 1. Tensile curves of HDPE (a) preoriented in n-heptane (degrees of extension 100 (1), 200 (2), 300 (3) and 400~ (4)) and for composites based on HDPE-PBMA (b) prepared by polymerization in situ by stretching HDPE in n-butyl methacrylate (degrees of extension 100 (1), 200 (2), 300 (3) and 400 ~ (4)). F i g u r e l a shows the tensile curves p l o t t e d at r o o m t e m p e r a t u r e for H D P E films t h a t h a d first been stretched in n - h e p t a n e to v a r i o u s degrees o f extension, a n d were then subjected to the a b o v e described h e a t t r e a t m e n t . I t can be seen t h a t the cold d r a w ing o f H D P E in the liquid m e d i u m leads to m a j o r changes in its m e c h a n i c a l p r o p e r t i e s . * The nature of the plasticizing liquid in which stretching takes place hardly affects the mechanical properties of the HDPE. The only matter of essential importance is that the degree of compatibility of the components should be adequate.
Mechanical properties of polymer compositions based on PE
935
An increase in the draw ratio is accompanied by well known changes due to an increase in the degree of orientation of the polymer: the initial modulus and the yield point are both higher, while the breaking elongations are reduced. For the samples with extensions of 300 and 4 0 0 ~ there is no yield point, and rupture occurs at small elongations ( ~ 50 ~o)The incorporation of PBMA in H D P E does not result in any marked change in mechanical properties of the composite material (Fig. lb). It is readily apparent that features of mechanical behaviour in the presence of PBMA are the same as for the pure H D P E . The main difference lies in the fact that tensile curves for the compositions are displaced towards low stress, and their respective mechanical properties are lower (,-,2 times) than for the pure HDPE. In this connection it should be noted that the pure H D P E samples and the corresponding composite materials differ in thickness by a factor of ~ 2 . The point is, that after drawing in heptane contraction of the H D P E takes place in the drying and annealing processes owing to removal of the liquid that was incorporated in the polymer structure in the drawing process [1, 2]. It appears that no such contraction takes place during polymerization, since considerable amounts of the P B M A remain in the H D P E structure, though the amounts of H D P E in cross sections of the samples are identical in both cases. This means that mechanical behaviour of the compositions based on P B M A - H D P E is mainly related to properties of the PE phase, while the P B M A behaves like an inert filler. It can be seen that the incorporation of a rubber-like polymer, PBMA (T~= 20 °) into the intercrystallite regions of H D P E , at the testing temperature, impairs mechanical properties of the resuiting composition, as has often been observed in the blending of PE with rubbers [4]. Let us now turn to changes in the mechanical properties of the compositions i f the polymers that are incorporated in the PE are in the glassy state at the temperature of the experiments. In Fig. 2 we have tensile curves for the compositions based on H D P E and PS or P M M A . Here also the obtained composite materials " r e m e m b e r " that the PE phase has been oriented, and so as the extent of prior drawing of H D P E in one or other monomer increases, we find that the modulus and the yield point for the materials are increased accordingly. It is seen from Fig. 2a and b that the second component, that is in the glassy state, likewise has a marked effect on properties of the materials, which is primarily reflected in higher initial moduli for the latter. At the same time properties are acquired that are not typical for either one of the components. It can be seen that the materials based on H D P E and PS, as well as on H D P E - P M M A are capable of considerable plastic deformations. Certainly, in contradistinction to the case (discussed above) of the H D P E - P B M A system, the incorporation of PS or P M M A in H D P E means that considerable breaking elongations can be obtained even for the compositions that are based on H D P E which has undergone 300 and 400 o~ extensions and which in the pure form is incapable of any such marked elongations (Fig. ta). At room temperature neither PS nor P M M A are themseh'es capable of any significant nonelastic deformations, and rupture occurs at 3-5 ~ elongations. Thus there is some sort of synergism in the mechanical behaviour when two polymers which, individually,. are incapable of marked deformations, jointly become capable of the latter.
936
A.L. VOLYNSKIIet al.
On comparing the systems based on PBMA, PS and P M M A we surmise that, in order to obtain the effect described above, the second component incorporated in H D P E by polymerization in situ has to be in the glassy state. Verification of this .assumption is obtainable by adding to n-butyl methacrylate (used as a medium in which H D P E is stretched) a sufficient amount of a crosslinking agent, D M E G . Crosslinking leads to a higher glass-transition for the PBMA, such that it rises above room temperature. In this case PBMA, like PS and PMMA, will be in the glassy state testing temperature. It is seen from Fig. 2c that the effect in question, viz. higher plasticity and increased strength, is quite marked for the system HDPE-crosslinked PBMA. This means that the chemical nature of the second component introduced into H D P E has no determining influence of mechanical properties of the resulting composite: the main factor is, apparently, its high modulus.
c, MPo
b
El 120 2
s
/
80
1
qO I
100
I
I
2_.00
I
100
I
I
I
200
,t,%
~, MPa
~
120
q
C
80 40 I
I
t
I
I00
200
;t , %
FIG. 2. Tensile curves of com.posites based on HDPE-PS (a) and HDPE-PMMA (b, c) prepared by polymerization in situ in styrene (a) or in methyl methacrylate (b, c) (degrees of extension 100 (1), 200 (2), 300 (3) and 400~o (4)) without a crosslinking agent (b) or in presence of 20 wt. ~ DMEG (c). Another major factor determining the mechanical behaviour of the composites is the high degree of interdispersion of incompatible components. It is known [3] that a very highly dispersed type of structure is formed in the case of in situ polymerization processes. Annealing at a temperature above the PE melting point leads to complete disorientation of the polymer and to phase distribution leading to a considerable enlargement of phase domains of the components [3]. It was found that the films suffer a considerable loss of strength and plasticity as a result of annealing, and that quite low levels of stress ( ~ 5 MPa) and slight elongations (2-5 ~o) are sufficient to cause rupture of the materials. Actually the composite structure formed in this case is of the type that is normally obtained through the blending of polymers from the melt, with
Mechanical properties of polymer compositions based on PE
937
the result that the mechanical properties are generally quite low. This is observed, for instance in the mixing of PE and PS melts in an extruder [5] leading of the formation of composites whose properties occupy a position that is intermediate between the properties of the pure components over a wide range of compositions, so that no synergism is involved in changes in the properties. For instance, the breaking elongation for blends of PE and PS each containing 50 ~o of the components, i.e. similar in composition to those discussed above, was found to be from 2 to 8 ~ , while the tensile strength was less than 10 MPa [6]. It is seen that the method of polymerization in situ results is composites possessing incomparably better mechanical properties. We now come to the matter of reasons underlying the observed effect. One would naturally assume that the deformation mechanism for components forming separate phases in the structure of a composite is bound to be different. Neither PS nor P M M A are themselves capable of any significant nonelastic deformations at room temperature, so it would appear that their deformation in the composite structure is due to a change in the shape of the structural elements. According to the electron microscopy data the PS or the P M M A phase in a composite forms a highly dispersed skeletal structure that is capable of deformation owing to a shape effect in the absence of any appreciable orientation of the macromolecules in the same way as may occur in the case of steel springs or, for instance, in the case of the so-called cryptoheterogeneous polymer systems [7]. The latter finding has been verified by infrared spectroscopy. We investigated the I R dichroism of absorption bands relating to PS in a sample of the P S - H D P E composition after tensile testing. Although this material underwent an extension of practically 2 0 0 ~ (Fig. 2a, curve 3) no significant dichroism in the bands for PS was detectable in the structure of the composite. This shows that no molecular orientation has occurred. It is known [3] that both components in the composite structure form a continuous phase and, accordingly, will undergo considerable deformation at a macroscopic level. The absence of diehroism bears out an assumption that deformation of the glassy component in the composite structure is mainly due to a change in the shape of structural elements. At the same time the rigid glassy component is very firmly attached to PE phase owing to chemical grafting or to molecular catenane interpermeation, as was demonstrated in [3]. Such a strong interconnection means that there is a particular mode of deformation lbr the PE phase, since the deformation mechanism for a material formed from two continuous phases will be determined by the more rigid component with a higher elastic modulus. This is fairly obvious, since it is only when deformation of the more rigid component has started that deformation of the phase having a lower modulus becomes possible. In view of this it is clear that the PE phase that likewise has a skeletal fine porous structure and is firmly linked to the glassy matrix will "follow" the latter in the deformation process, i.e. will likewise undergo deformation mainly in accordance with a mechanism of change in the shape of structural elements. Deformation of this type cannot occur in the case of pure PE because of its monolithic character. In this case as soon as the yield point has been attained we have the onset of cold flow due to molecular orientation of the polymer at low levels of prior deformation, or else there
938
A.L. VOLYNSKIIet al.
is a process of breakdown in cases where the plastic deformation capacity has been exhausted as a result of the initial drawing where samples have undergone preliminary 300-400 ~ extensions. In cases where the polymer incorporated in PE by means of polymerization in situ is above or in the region of the glass-transition temperature its rigidity will be insufficient to dictate or compel realization of a mechanism of deformation of the PE structure via a change in the shape of structural elements. In this case it is once again the more rigid phase that begins to "dictate" its own deformation mechanism, the more rigid phase being now a system of interconnected H D P E crystallites. The deformation mechanism of crystalline polymers, particularly PE, has been quite adequately investigated, and there is no point in describing it here. We would only point out that in the presence of a rubbery component H D P E deformation takes place by a normal mechanism as is evident, in particular, from the undoubted similarity of the sets of tensile curves in Fig. la and b. In view of these considerations one can account for the influence of crosslinking o f PBMA affecting mechanical properties of the composites. As a result of crosslinking the glass transition temperature of PBMA is higher than the (room) temperature as which mechanical testing was carried out. It is quite natural that in this case the rigidity and the modulus of PBMA phase should be markedly increased, leading (for the reasons outlined above) to an increase in the plasticity of the resultant composites.
4
MPa
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I
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lO0
100
I
I
300 A,%
~, t'l P a
rio
t l ~
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c ..----2 ~
1 I
300 ,~..,';~
Fro. 3. Tensile curves of HDPE (a) and for composites based on HDPE.-PS (b) and HDPE-PMMA (c) prepared at 100°. Degrees of prior stretching: 100 (1), 200 (2), 300 (3) and 400~ (4). If the proposed mechanism of plasticity of the composites holds, it must follow that the plasticity effect will disappear in the vicinity of Ts for polymers incorporated in the PE. This is fully substantiated by a study of mechanical properties of the composites at 100 °, i.e. at a temperature close to the glass transition temperature of PS in PMMA.
Mechanical properties of polymer compositions based on PE
939
Figure 3 shows the tensile curves for the pure PE, and for the P E - P M M A . and PE-PS ~ystems. It is natural that at this temperature a slight increase in-plasticity should appear for the pure HDPE; however, there is no increase in plasticity for any of the composites. Although the presence of a second component does slightly alter the shape of the tensile curves, which is reflected in an inconsiderable increase in the initial elastic modulus and in the appearance of clear-cut overstrain peaks, the breaking elongations remain the same as for the pure HDPE, i.e. there is no longer any increased plasticity effect. These findings fully bear out assumptions that were made above regarding the deformation mechanism for the composites above and below Tg for the incorporated polymeric filler. It is obvious that the proposed variation in the deformation mechanism of the composite varies according to whether a second polymer component added to the PE is ,above or below the glass transition temperature will be accompanied by different modes of structural transformation in the deformed material. Features of one or other type o f deformation were detected in an X-ray study of samples of composities under deformation at various temperatures. Figure 4 shows the wide and low angle X-r adiograms for samples of the compositebased on H D P E - P S subjected to extension to the point of rupture at room temperalure and at 100 °. It was found that the molecular orientation of H D P E in the composite stretched at 100 ° is significantly higher than in that stretched at room temperature (irrespective of the degree of prior drawing of H D P E in the monomer), although the draw ratios were broadly speaking identical in both cases. This finding was unexpected, since it is well known that the higher the drawing temperature, the lower is the attainable degree of orientation of the polymer, other conditions being identical. The reason why this is so is directly related to the difference (discussed above) in the deformation mechanisms of composites containing a polymeric filler in the glasslike or the rubbery state. Certainly, deformation of the polymeric porous carcass via a change in the shape of structural elements occurring at low temperatures need not result in any significant degree of molecular orientation. However stretching of the composite containing a rubbery filler, i.e. stretching above its glass transition t e m p e r a t u r e - a t 100°-takes place in a way such that the PE carcass is capable of deformation via orientation of the polymer in much the same way as in the case of pure monolithic PE. It is clear that the degree of molecular orientation attained in the former case is significantly lower than in the latter, though stretching takes place at a lower temperature in the former case. This conclusion regarding the deformation mechanism of the composites is supported by low angle X-ray scattering. It is seen from Fig. 4c, d, f, h that low temperature stretching of the composites is accompanied by the appearance of a large number of microvacancies, whereas orientation of the polymer at 100 ° takes place fairly uniformly without any significant loss of soundness of the material. This is because deformation of two interconnected rigid carcasses, which takes place at low temperatures through a change in the shape of the structural elements is naturally accompanied by partial separation of the latter, leading to the emergence of a great many microvacancies. At high temperature the polymeric filler is at a temperature above Tg, i.e. it has a high
A . L . VOLYNSKII et al.
940
IX',,
-7
,~t
!
.
N
FIG. 4. Wide-angle (a, b, e,f) and low-angle (c, d, g, h) X radiograms for composites based on H D P E PS samples stretched at room temperature (a, e, e, g) and at 100 ° (b, d, f , h). Degrees of prior stretching of H D P E in styrene: 100 (a-d) and 400 ~ (e-h).
Mechanical properties of polymer compositions based on PE degree o f pliability a n d does n o t i m p e d e d e f o r m a t i o n o f the P E carcass which, u n d e r these conditions, d e t e r m i n e s the d e f o r m a t i o n m e c h a n i s m o f the c o m p o s i t e as a whole. The soft filler follows u p the P E carcass a n d no s e p a r a t i o n o f the c o m p o n e n t s t a k e s place, so the m o n o l i t h i c n a t u r e o f the m a t e r i a l is preserved. T h u s it a p p e a r s t h a t the i n c o r p o r a t i o n o f an a m o r p h o u s c o m p o n e n t in H D P E b y p o l y m e r i z a t i o n in situ results in c o m p o s i t e s possessing high strength a n d plasticity. T h e high degree o f plasticity o f these m a t e r i a l s is due to a change in the d e f o r m a t i o n m e c h a n i s m o f the m a t e r i a l c o n t a i n i n g the glassy filler. I n this case elongation o f the m a t e r i a l is o b t a i n e d m a i n l y on a c c o u n t o f a change in the shape o f elements of the highly dispersed skeletal s t r u c t u r e o f the c o m p o s i t e materials. Translated by R. J. A. 1-|E,',D~aY REFERENCES
1. A. V. YEFIMOV, V. V. BONDAREV, P. V. KOZLOV and N. F. BAKEYEV, Vysckomol. soyed. A24: 1690, 1982 (Translated in Polymer Sci. U.S.S.R. 24: 8, 1927, 1982) 2. A. L. VOLYNSKII, A. Sh. SHTANCHAYEV and N. F. BAKEYEV, Vysokomol. sc:.cd. 26: 2374, 1984 (Translated in Polymer Sci U.S.S.R. 26: 11, 2654, 1984) 3. A. L. VOLYNSKII, A. Sh. SHTANCHAYEV and N. F. BAKEYEV, V vsek~mol, soyed. A26: 2445, 1984 (Translated in Polymer Sci. U.S.S.R. 26: 11, 2741, 1984) 4. N. ASANDEI, G. ADAMESCU, M. RUSU, S. PETROVAN ar, d M. NICU, B~li! Inst. Polyt. Din. Tasi 26: 53, 1980 5. D. HEIKENS, Kemisk. Industr. 31: 165, 1982 6. W. M. BARENTSEN and D. HEIKENS, Polymer 14: l 1, 579, 1973 7. G. M. SIN1TSYNA, In: Uspekhi kolloidnoi khimii (Advances Colloid Chemistr~ p, 331, Nauka: Moscow, 1973
933
REFERENCES i. E. BUNCEL and A. DAVIES, J. Chem. Soc., 7, 1550, 1958 2. N. A. IZMAILOV, L. L. SPIVAK and V. I. LEVINKOVA, Uchenye zapiski K h G U (Academic Notes of Univ. of Kharkov) p. 30, 1950 3. N. P. SLUCHEVSKAYA, V. A. YABLOKOV, N. V. YABLOKOVA and Yu. A. ALEKSANDROV, Z k obshch, khimii 46: 1540, 1976 4. Yu. A. ALEXANDROV, N. V. YABLOKOVA and V. V. GORBATOV, Basic Research in Homogeneous Catalysis, vol. 3, p. 359, Univ. Coll. Station Texas, N. Y.-London, 1979 5. R. FOSTER and C. A. FYCE, Chem. Commun. 3: 642, 1965 6. V. P. ZUBOV and V. A. KABANOV, Khimiya i tekhnol, vysokomol, soyed., VINITI, Moscow, 9: 56, 1977 7. ~ . C. A P P L E T O N and J. TYRREL, J. Phys. Chem. 81: 1201, 1977 8. H. HIRAI and M. KAMIYAMA, J. Polymer Sci. Polymer Chem. Ed. 12: 2701, 1974
Polyrve ScienceU.S.S.R. Vol. 27, No. 4, pp. 933-941, 1985 Priqt.z~' in Poland
0032 3950/85 $10.00+.00 ~(; 1986 Pergamon Press Ltd.
MECHANICAL PROPERTIES OF POLYMER COMPOSITIONS BASED ON POLYETHYLENE AND PREPARED BY POLYMERIZATION IN SITU* A. L. VOLYNSKII, A . SH. SHTANCHAYEV, V. D. ZANEGIN, V. I. GERASIMOV a n d N. F. BAKEYEV Lomonosov State University, Moscow
(Received 21 August 1983)
A study has been made of polymer compositions prepared via polymerization of a runlber of liquid monomers in a PE matrix (polymerization in situ). The compositions have high strength and elasticity values if the polymeric component incorporated in PE is in the g'assy state. Structural changes accompanying nonelastic deformation of the compositions e~ e r a wide range of temperature have been investigated by X-ray analysis. A deformation ~'L~chanism is proposed on the basis of experimental results.
l-r is k n o w n [1,2] t h a t u n i a x i a l d r a w i n g o f high d e n s i t y P E ( H D P E ) films in a c o m p a t i b l e l i q u i d causes m a r k e d l y i n c r e a s e d swelling o f t h e p o l y m e r . P o s t p o l y m e r i z a t i o n o f a low molecular weight component
in t h e P E m a t r i x results is n o v e l p o l y m e r i c m a t e r i a l s
b a s e d on P E [3]. O u r a i m in t h e p r e s e n t w o r k was to i n v e s t i g a t e s o m e m e c h a n i c a l p r o p * Vysokomol. soyed. A27: No. 4, 831-837, 1985.
934
A . L . VOLYNSKIIet al.
erties o f m a t e r i a l s o f t h i s type b a s e d on H D P E a n d P M M A , PS o r p o l y - n - b u t y l m e t h acrylate (PBMA). The study objects were commercial HDPE films of thickness 50 ltm. The films were drawn in liquid monomers containing an initiator (benzoyl peroxide) and, in some cases, a crosslinking agent (ethylene glycol dimetharcylate (EGD)). Drawing to 100, 200, 300 and 400 ~ extensions was with the aid of specially designed equipment. After stretching, the lengths of samples were determined in stainless-steel frames, and the samples were transferred to a heat-controlled vessel containing an inert liquid heated to 80 °. Polymerization was conducted in this vessel for 8 hr. After polymerization the samples were annealed under isothermal conditions at 100 ° for 1 hr. Details of the experimental procedure and data on the composition of the resulting materials, along with results of the X-ray analysis were published in [3]. Mechanical testing was carried out with the aid of an Instron tensile machine at a constant rate of stretch (20 ram/rain). A s was n o t e d above, m o n o m e r is i n c o r p o r a t e d in the H D P E d u r i n g the s t r e t c h i n g o f the polymer. A c c o r d i n g l y P E phase in the f o r m i n g m a t e r i a l will be m o r e o r less oriented. T o shed light on the extent to which m e c h a n i c a l p r o p e r t i e s o f the m a t e r i a l are influenced by the o r i e n t e d p o l y m e r m a t r i x let us first o f all e x a m i n e the m e c h a n i c a l b e h a v i o u r o f H D P E stretched in a plasticizing liquid (heptane).* A f t e r stretching in the liquid m e d i u m lengths o f the H D P E samples were d e t e r m i n e d a n d the l a t t e r were d r i e d in air a n d a n n e a l e d u n d e r i s o m e t r i c c o n d i t i o n s for 1 h r ar 100 °, i.e. u n d e r w e n t h e a t t r e a t m e n t similar to t h a t to which the o b t a i n e d (composite) m a t e r i a l s were subjected.
MPa
q 0 80
120 ~.3 80
~//~2
qO
qo 0
1 I
I
200
[
I
400 2,%
I
0
I
100
I
I
200
f
~%
FIG. 1. Tensile curves of HDPE (a) preoriented in n-heptane (degrees of extension 100 (1), 200 (2), 300 (3) and 400~ (4)) and for composites based on HDPE-PBMA (b) prepared by polymerization in situ by stretching HDPE in n-butyl methacrylate (degrees of extension 100 (1), 200 (2), 300 (3) and 400 ~ (4)). F i g u r e l a shows the tensile curves p l o t t e d at r o o m t e m p e r a t u r e for H D P E films t h a t h a d first been stretched in n - h e p t a n e to v a r i o u s degrees o f extension, a n d were then subjected to the a b o v e described h e a t t r e a t m e n t . I t can be seen t h a t the cold d r a w ing o f H D P E in the liquid m e d i u m leads to m a j o r changes in its m e c h a n i c a l p r o p e r t i e s . * The nature of the plasticizing liquid in which stretching takes place hardly affects the mechanical properties of the HDPE. The only matter of essential importance is that the degree of compatibility of the components should be adequate.
Mechanical properties of polymer compositions based on PE
935
An increase in the draw ratio is accompanied by well known changes due to an increase in the degree of orientation of the polymer: the initial modulus and the yield point are both higher, while the breaking elongations are reduced. For the samples with extensions of 300 and 4 0 0 ~ there is no yield point, and rupture occurs at small elongations ( ~ 50 ~o)The incorporation of PBMA in H D P E does not result in any marked change in mechanical properties of the composite material (Fig. lb). It is readily apparent that features of mechanical behaviour in the presence of PBMA are the same as for the pure H D P E . The main difference lies in the fact that tensile curves for the compositions are displaced towards low stress, and their respective mechanical properties are lower (,-,2 times) than for the pure HDPE. In this connection it should be noted that the pure H D P E samples and the corresponding composite materials differ in thickness by a factor of ~ 2 . The point is, that after drawing in heptane contraction of the H D P E takes place in the drying and annealing processes owing to removal of the liquid that was incorporated in the polymer structure in the drawing process [1, 2]. It appears that no such contraction takes place during polymerization, since considerable amounts of the P B M A remain in the H D P E structure, though the amounts of H D P E in cross sections of the samples are identical in both cases. This means that mechanical behaviour of the compositions based on P B M A - H D P E is mainly related to properties of the PE phase, while the P B M A behaves like an inert filler. It can be seen that the incorporation of a rubber-like polymer, PBMA (T~= 20 °) into the intercrystallite regions of H D P E , at the testing temperature, impairs mechanical properties of the resuiting composition, as has often been observed in the blending of PE with rubbers [4]. Let us now turn to changes in the mechanical properties of the compositions i f the polymers that are incorporated in the PE are in the glassy state at the temperature of the experiments. In Fig. 2 we have tensile curves for the compositions based on H D P E and PS or P M M A . Here also the obtained composite materials " r e m e m b e r " that the PE phase has been oriented, and so as the extent of prior drawing of H D P E in one or other monomer increases, we find that the modulus and the yield point for the materials are increased accordingly. It is seen from Fig. 2a and b that the second component, that is in the glassy state, likewise has a marked effect on properties of the materials, which is primarily reflected in higher initial moduli for the latter. At the same time properties are acquired that are not typical for either one of the components. It can be seen that the materials based on H D P E and PS, as well as on H D P E - P M M A are capable of considerable plastic deformations. Certainly, in contradistinction to the case (discussed above) of the H D P E - P B M A system, the incorporation of PS or P M M A in H D P E means that considerable breaking elongations can be obtained even for the compositions that are based on H D P E which has undergone 300 and 400 o~ extensions and which in the pure form is incapable of any such marked elongations (Fig. ta). At room temperature neither PS nor P M M A are themseh'es capable of any significant nonelastic deformations, and rupture occurs at 3-5 ~ elongations. Thus there is some sort of synergism in the mechanical behaviour when two polymers which, individually,. are incapable of marked deformations, jointly become capable of the latter.
936
A.L. VOLYNSKIIet al.
On comparing the systems based on PBMA, PS and P M M A we surmise that, in order to obtain the effect described above, the second component incorporated in H D P E by polymerization in situ has to be in the glassy state. Verification of this .assumption is obtainable by adding to n-butyl methacrylate (used as a medium in which H D P E is stretched) a sufficient amount of a crosslinking agent, D M E G . Crosslinking leads to a higher glass-transition for the PBMA, such that it rises above room temperature. In this case PBMA, like PS and PMMA, will be in the glassy state testing temperature. It is seen from Fig. 2c that the effect in question, viz. higher plasticity and increased strength, is quite marked for the system HDPE-crosslinked PBMA. This means that the chemical nature of the second component introduced into H D P E has no determining influence of mechanical properties of the resulting composite: the main factor is, apparently, its high modulus.
c, MPo
b
El 120 2
s
/
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1
qO I
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I
I
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I
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I
I
I
200
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~
120
q
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80 40 I
I
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I
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FIG. 2. Tensile curves of com.posites based on HDPE-PS (a) and HDPE-PMMA (b, c) prepared by polymerization in situ in styrene (a) or in methyl methacrylate (b, c) (degrees of extension 100 (1), 200 (2), 300 (3) and 400~o (4)) without a crosslinking agent (b) or in presence of 20 wt. ~ DMEG (c). Another major factor determining the mechanical behaviour of the composites is the high degree of interdispersion of incompatible components. It is known [3] that a very highly dispersed type of structure is formed in the case of in situ polymerization processes. Annealing at a temperature above the PE melting point leads to complete disorientation of the polymer and to phase distribution leading to a considerable enlargement of phase domains of the components [3]. It was found that the films suffer a considerable loss of strength and plasticity as a result of annealing, and that quite low levels of stress ( ~ 5 MPa) and slight elongations (2-5 ~o) are sufficient to cause rupture of the materials. Actually the composite structure formed in this case is of the type that is normally obtained through the blending of polymers from the melt, with
Mechanical properties of polymer compositions based on PE
937
the result that the mechanical properties are generally quite low. This is observed, for instance in the mixing of PE and PS melts in an extruder [5] leading of the formation of composites whose properties occupy a position that is intermediate between the properties of the pure components over a wide range of compositions, so that no synergism is involved in changes in the properties. For instance, the breaking elongation for blends of PE and PS each containing 50 ~o of the components, i.e. similar in composition to those discussed above, was found to be from 2 to 8 ~ , while the tensile strength was less than 10 MPa [6]. It is seen that the method of polymerization in situ results is composites possessing incomparably better mechanical properties. We now come to the matter of reasons underlying the observed effect. One would naturally assume that the deformation mechanism for components forming separate phases in the structure of a composite is bound to be different. Neither PS nor P M M A are themselves capable of any significant nonelastic deformations at room temperature, so it would appear that their deformation in the composite structure is due to a change in the shape of the structural elements. According to the electron microscopy data the PS or the P M M A phase in a composite forms a highly dispersed skeletal structure that is capable of deformation owing to a shape effect in the absence of any appreciable orientation of the macromolecules in the same way as may occur in the case of steel springs or, for instance, in the case of the so-called cryptoheterogeneous polymer systems [7]. The latter finding has been verified by infrared spectroscopy. We investigated the I R dichroism of absorption bands relating to PS in a sample of the P S - H D P E composition after tensile testing. Although this material underwent an extension of practically 2 0 0 ~ (Fig. 2a, curve 3) no significant dichroism in the bands for PS was detectable in the structure of the composite. This shows that no molecular orientation has occurred. It is known [3] that both components in the composite structure form a continuous phase and, accordingly, will undergo considerable deformation at a macroscopic level. The absence of diehroism bears out an assumption that deformation of the glassy component in the composite structure is mainly due to a change in the shape of structural elements. At the same time the rigid glassy component is very firmly attached to PE phase owing to chemical grafting or to molecular catenane interpermeation, as was demonstrated in [3]. Such a strong interconnection means that there is a particular mode of deformation lbr the PE phase, since the deformation mechanism for a material formed from two continuous phases will be determined by the more rigid component with a higher elastic modulus. This is fairly obvious, since it is only when deformation of the more rigid component has started that deformation of the phase having a lower modulus becomes possible. In view of this it is clear that the PE phase that likewise has a skeletal fine porous structure and is firmly linked to the glassy matrix will "follow" the latter in the deformation process, i.e. will likewise undergo deformation mainly in accordance with a mechanism of change in the shape of structural elements. Deformation of this type cannot occur in the case of pure PE because of its monolithic character. In this case as soon as the yield point has been attained we have the onset of cold flow due to molecular orientation of the polymer at low levels of prior deformation, or else there
938
A.L. VOLYNSKIIet al.
is a process of breakdown in cases where the plastic deformation capacity has been exhausted as a result of the initial drawing where samples have undergone preliminary 300-400 ~ extensions. In cases where the polymer incorporated in PE by means of polymerization in situ is above or in the region of the glass-transition temperature its rigidity will be insufficient to dictate or compel realization of a mechanism of deformation of the PE structure via a change in the shape of structural elements. In this case it is once again the more rigid phase that begins to "dictate" its own deformation mechanism, the more rigid phase being now a system of interconnected H D P E crystallites. The deformation mechanism of crystalline polymers, particularly PE, has been quite adequately investigated, and there is no point in describing it here. We would only point out that in the presence of a rubbery component H D P E deformation takes place by a normal mechanism as is evident, in particular, from the undoubted similarity of the sets of tensile curves in Fig. la and b. In view of these considerations one can account for the influence of crosslinking o f PBMA affecting mechanical properties of the composites. As a result of crosslinking the glass transition temperature of PBMA is higher than the (room) temperature as which mechanical testing was carried out. It is quite natural that in this case the rigidity and the modulus of PBMA phase should be markedly increased, leading (for the reasons outlined above) to an increase in the plasticity of the resultant composites.
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Fro. 3. Tensile curves of HDPE (a) and for composites based on HDPE.-PS (b) and HDPE-PMMA (c) prepared at 100°. Degrees of prior stretching: 100 (1), 200 (2), 300 (3) and 400~ (4). If the proposed mechanism of plasticity of the composites holds, it must follow that the plasticity effect will disappear in the vicinity of Ts for polymers incorporated in the PE. This is fully substantiated by a study of mechanical properties of the composites at 100 °, i.e. at a temperature close to the glass transition temperature of PS in PMMA.
Mechanical properties of polymer compositions based on PE
939
Figure 3 shows the tensile curves for the pure PE, and for the P E - P M M A . and PE-PS ~ystems. It is natural that at this temperature a slight increase in-plasticity should appear for the pure HDPE; however, there is no increase in plasticity for any of the composites. Although the presence of a second component does slightly alter the shape of the tensile curves, which is reflected in an inconsiderable increase in the initial elastic modulus and in the appearance of clear-cut overstrain peaks, the breaking elongations remain the same as for the pure HDPE, i.e. there is no longer any increased plasticity effect. These findings fully bear out assumptions that were made above regarding the deformation mechanism for the composites above and below Tg for the incorporated polymeric filler. It is obvious that the proposed variation in the deformation mechanism of the composite varies according to whether a second polymer component added to the PE is ,above or below the glass transition temperature will be accompanied by different modes of structural transformation in the deformed material. Features of one or other type o f deformation were detected in an X-ray study of samples of composities under deformation at various temperatures. Figure 4 shows the wide and low angle X-r adiograms for samples of the compositebased on H D P E - P S subjected to extension to the point of rupture at room temperalure and at 100 °. It was found that the molecular orientation of H D P E in the composite stretched at 100 ° is significantly higher than in that stretched at room temperature (irrespective of the degree of prior drawing of H D P E in the monomer), although the draw ratios were broadly speaking identical in both cases. This finding was unexpected, since it is well known that the higher the drawing temperature, the lower is the attainable degree of orientation of the polymer, other conditions being identical. The reason why this is so is directly related to the difference (discussed above) in the deformation mechanisms of composites containing a polymeric filler in the glasslike or the rubbery state. Certainly, deformation of the polymeric porous carcass via a change in the shape of structural elements occurring at low temperatures need not result in any significant degree of molecular orientation. However stretching of the composite containing a rubbery filler, i.e. stretching above its glass transition t e m p e r a t u r e - a t 100°-takes place in a way such that the PE carcass is capable of deformation via orientation of the polymer in much the same way as in the case of pure monolithic PE. It is clear that the degree of molecular orientation attained in the former case is significantly lower than in the latter, though stretching takes place at a lower temperature in the former case. This conclusion regarding the deformation mechanism of the composites is supported by low angle X-ray scattering. It is seen from Fig. 4c, d, f, h that low temperature stretching of the composites is accompanied by the appearance of a large number of microvacancies, whereas orientation of the polymer at 100 ° takes place fairly uniformly without any significant loss of soundness of the material. This is because deformation of two interconnected rigid carcasses, which takes place at low temperatures through a change in the shape of the structural elements is naturally accompanied by partial separation of the latter, leading to the emergence of a great many microvacancies. At high temperature the polymeric filler is at a temperature above Tg, i.e. it has a high
A . L . VOLYNSKII et al.
940
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.
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FIG. 4. Wide-angle (a, b, e,f) and low-angle (c, d, g, h) X radiograms for composites based on H D P E PS samples stretched at room temperature (a, e, e, g) and at 100 ° (b, d, f , h). Degrees of prior stretching of H D P E in styrene: 100 (a-d) and 400 ~ (e-h).
Mechanical properties of polymer compositions based on PE degree o f pliability a n d does n o t i m p e d e d e f o r m a t i o n o f the P E carcass which, u n d e r these conditions, d e t e r m i n e s the d e f o r m a t i o n m e c h a n i s m o f the c o m p o s i t e as a whole. The soft filler follows u p the P E carcass a n d no s e p a r a t i o n o f the c o m p o n e n t s t a k e s place, so the m o n o l i t h i c n a t u r e o f the m a t e r i a l is preserved. T h u s it a p p e a r s t h a t the i n c o r p o r a t i o n o f an a m o r p h o u s c o m p o n e n t in H D P E b y p o l y m e r i z a t i o n in situ results in c o m p o s i t e s possessing high strength a n d plasticity. T h e high degree o f plasticity o f these m a t e r i a l s is due to a change in the d e f o r m a t i o n m e c h a n i s m o f the m a t e r i a l c o n t a i n i n g the glassy filler. I n this case elongation o f the m a t e r i a l is o b t a i n e d m a i n l y on a c c o u n t o f a change in the shape o f elements of the highly dispersed skeletal s t r u c t u r e o f the c o m p o s i t e materials. Translated by R. J. A. 1-|E,',D~aY REFERENCES
1. A. V. YEFIMOV, V. V. BONDAREV, P. V. KOZLOV and N. F. BAKEYEV, Vysckomol. soyed. A24: 1690, 1982 (Translated in Polymer Sci. U.S.S.R. 24: 8, 1927, 1982) 2. A. L. VOLYNSKII, A. Sh. SHTANCHAYEV and N. F. BAKEYEV, Vysokomol. sc:.cd. 26: 2374, 1984 (Translated in Polymer Sci U.S.S.R. 26: 11, 2654, 1984) 3. A. L. VOLYNSKII, A. Sh. SHTANCHAYEV and N. F. BAKEYEV, V vsek~mol, soyed. A26: 2445, 1984 (Translated in Polymer Sci. U.S.S.R. 26: 11, 2741, 1984) 4. N. ASANDEI, G. ADAMESCU, M. RUSU, S. PETROVAN ar, d M. NICU, B~li! Inst. Polyt. Din. Tasi 26: 53, 1980 5. D. HEIKENS, Kemisk. Industr. 31: 165, 1982 6. W. M. BARENTSEN and D. HEIKENS, Polymer 14: l 1, 579, 1973 7. G. M. SIN1TSYNA, In: Uspekhi kolloidnoi khimii (Advances Colloid Chemistr~ p, 331, Nauka: Moscow, 1973