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The effect of stretching along the orientation axis on the supermolecular structure of polyethylene films affected by shrinkage
Supermolecular structure of polyethylene films affected by shrinkage
1525
The conformational models of POBH seen in Fig. 2 are proposed in the light of the available data. It is possible to conclude on the basis of these results that the difference in the chemical properties of imidazole groups of the mono- and disubstituted POBH are due to the dissimilarity of the conformational state of these polymers. The authors thank M. F. Vyalykh for assisting with eleetrophoretic separation of the compounds, and are grateful for the help given by N. F. Kazanskaya in determining the pK values. Translated by R. J. A. HE~VD~Y REFERENCES
1. I. N. TOPCHIEVA, A. B. SOLOV'EVA and V. A. KABANOV, Dokl. AN SSSR 199:1084 1971 2. I. N. TOPCHIEVA, A. B. SOLOV'EVA and V. A. KABANOV, Vysokomol. soyed. A14: 825, 1972 (Translated in Polymer Sci. U.S.S.R. 4: 918, 1972) 3. Ye. M. SAMSONOV, Sb. Fiziko-khimich. metody isueheniya, aaaliza i fraktsionirovaniya biopolimerov (Physicoehemical methods of Investigation, Analysis and Fraetionation of Biopolymers). p. 40, Izd. "Nauka", 1966 4. D. GRINSHTEIN and M. VINITS, Khimiya aminokislot i peptidov (Chemistry of Amino Acids and Peptides). p. 784, Izd. "Mir", 1964 5. I. M. HEISS and K. I~a,TSEK, Paper chromatography, p. 797, 1963 6. S. KORMAN and N. T. CLARKE, J. Biol. Chem. 221: 113, 1956 7. A. ALBERT and E. SERGENT, Ionization Constants of Acids and Bases, p. 21, 1964
THE EFFECT OF STRETCHING ALONG THE ORIENTATION AXIS ON THE SUPERMOLECULAR STRUCTURE OF POLYETHYLENE FILMS AFFECTED BY SHRINKAGE* B. M. GmZBU~G, K. B. KV~BA~OV, Y~. A. LEOSKO, M. A. MA~TY~OVand D. I~ASItIDOV High Polymer Institute, U.S.S.R. Academy of Sciences
(Received 30 October 1972) Changes in the supermolecular structure of low-density polyethylene films were investigated by X-ray analysis with low and wide angle diffraction patterns. The films underwent preliminary orientation at temperatures of 100, 85, 60 and 20° and * Vysokomol soyed. A16: No. 6, 1317-1323, 1974.
1526
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GINZBURG et al.
were s u b j e c t e d to s h r i n k a g e a t t h e s a m e t e m p e r a t u r e s . T h e c h a n g e s i n s u p e r m o l e c u l a r s t r u c t u r e (SMS) were t h e r e s u l t of g r a d u a l e x t e n s i o n a t 20 ° a l o n g t h e t e x t u r e axis. A m o d e l of t h e c h a n g e s in SMS h a s b e e n p r o p o s e d .
I~ EARLIER investigations [1] it was shown that the stretching of oriented films of low density polyethylene (PELD) at various angles to the c-texture axis accompanied by the "skewing" of crystallites'. During the stretching of actual oriented systems along the texture axis one finds that for some of the crystallites the direction of stretching makes a certain angle with the c-axis, and one would therefore expect changes in the shape of these crystallites to be reflected in low angte X-ray diffraction patterns. These changes in the low angle diffraction patterns of polyethylene films were considered in papers [2-4], and the results obtained in the latter investigations are partly in agreement with those presented below, but by varying the conditions of sample preparation (mainly the temperature of primary orientation T1) we obtained a number of new results, and the total amount of experimental data available is considered solely from the point of view of the skewing of crystallites then their subsequent destruction after a critical degree of skew has been reached, and the new supermoleeular structure (SMS) which is then formed. Secondary stretching along the texture axis was carried out at room temperature only (T2~ 20°). After undergoing the primary orientation the samples were subjected to shrinkage at temperature T1. This resulted in a slight deviation of the preferential direction of the c-axes of the crystallites from the texture axis, but in this case subsequent secondary stretching causes major and easily recordable changes in the low angle diffraction patterns, the interpretation of which is conducive to better understanding of changes in the patterns appearing in the course of the stretching of samples that had not been subjected to shrinkage. MATERIALS AND METHODS ! P E L D films (the c h a r a c t e r i s t i c s of t h e P E L D are give.n i n [2] w i t h a t h i c k n e s s of 1.7 to 1-8 m m were p r e p a r e d b y c o m p r e s s i o n m o u l d i n g f r o m t h e m e l t (200 °) a t a p r e s s u r e of 50 a t m ; t h e s a m p l e s were t h e n q u e n c h e d i n w a t e r a t 2 2 °. T h e films h a d a fine s p h e r u l i t i e s t r u c t u r e (R ~ 2 - 3 / l m ) . A f t e r t h i s t h e films u n d e r w e n t u n i a x i a l d r a w i n g a t t e m p e r a t u r e s (T1) of 20, 60, 85 a n d 100 ° u p t o a s t r a i n in t h e v i c i n i t y of t h e p r e - b r e a k i n g s t r a i n , a n d were t h e n freed f r o m t h e c l a m p s , a n d h a v i n g b e e n allowed to c o n t r a c t freely a t t e m p e r a t u r e T1 for 1 hr, t h e films were t r a n s f e r r e d t o r o o m c o n d i t i o n s . All t h e films u n d e r w e n t g r a d u a l s t r e t c h i n g u n d e r r o o m c o n d i t i o n s , a n d as t h i s proceeded, low a n d wide a n g l e p h o t o g r a p h s were t a k e n s i m u l t a n e o u s l y , s h o w i n g one a n d t h e s a m e p a r t of a s a m p l e , u s i n g c a m e r a s of t h e t y p e p r o p o s e d i n [5]. A KI~M-1 i n s t r u m e n t (slit c o l l i m a t i o n ) was u s e d t o m e a s u r e t h e i n t e n s i t y of t h e l o w - a n g l e s c a t t e r i n g . T h e t r a n s v e r s e d i m e n s i o n s of c r y s t a l l i t e s (110 a n d 200 reflections) a n d t h e l o n g i t u d i n a l d i m e n s i o n s (002 reflection) were m e a s u r e d w i t h a U R S - 5 0 I M i n s t r u m e n t . Ni-filtered C u K ~ - r a d i a t i o n w a s used.
Supermolecular structure of polyethylene films affected by shrinkage
1527
DISCUSSION OF RESULTS
T1 --~100 °. The low angle diffraction patterns of the original samples (Figs. la, 2a) * are characteristic for a n axial t e x t u r e of crystalline layers inclined towards t h e t e x t u r e axis [6, 7] a n d the ratio of the t r a n s v e r s e dimensions of t h e
I ~h
FIG. 1. Typical small and wide angle X-ray diffraction pictures; T~--100°C; a--original sample; ~0~0; b--40; c--60%. layers to the longitudinal dimensions considerably exceeds unity. J u d g i n g b y t h e wide angle p a t t e r n s (Fig. la) the t e x t u r e of the crystallites is close to a c-texture; t h e m a x i m a of t h e l l 0 and 200 reflections are near the equator. The X - r a y p h o t o g r a p h s o b t a i n e d w h e n t h e p r i m a r y b e a m was parallel to the plane of the film a n d p e r p e n d i c u l a r to the t e x t u r e axis, show t h a t the t e x t u r e of the films is close to a n axial t e x t u r e (Figs. 1, 5, 6). As the stretching proceeds, the following qualitative changes are d e t e c t a b l e in the X - r a y p h o t o g r a p h s .
®
C((
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I( `°
a
I
O
O
d FIG. 2. Schematic representations of some of the low and wide angle X-ray diffraction patterns: a--T~=100 °, original sample; b--ditto, e0=60%; e - - T l ~ 8 5 °, e0-----60°/o; d--T~ ~60 °, original sample; e--ditto, e0= 30%. * In Fig. 1 and in the other Figures the direction of the texture axis and of the secondary drawing is vertical; unless otherwise stated, the primary beam direction is perpendicular to the texture axis and to the plane of the polymer film, and the sample is under a load that has caused deformation.
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GINZBU~G et al.
1. The centres of the low angle reflections move away from the meridian (Fig. lb, c), which, in view of the small changes in the longitudinal dimensions of the crystallites, as well as the smaller transverse dimensions of the latter (see below) can only mean an increase in the angle of slope of the layers or in the angle of skew ~] of the crystallites [1, 6].
I!
i!
0
a
b FIG. 3 (a, b)
2. The arcs of equation reflections in the wide angle pictures grow wider in the azimuthal direction, and in the final stages of deformation (Figs. lc, 2b) the splitting of the arcs is clearly seen. 3. Reflections of the "spot" or radical [6] types in the low angle pictures become more and more anisotropie in their dimensions, and in the final stages of deformation these appear as "striations" (Figs. lc, 2b). The change to reflections of the striation type means that the transverse dimensions of the layers have become smaller. 4. The direction of the "striations" in the low angle pictures approximates more and more to a horizontal direction; a low angle picture comes to resembie a well-resolved diffraction pattern of the "four-petal" type, (Figs. la, 2b) in contrast to the "four-spot" (Fig. la). 5. Up to deformations of around 20% changes in the shape of a low angle pattern are practically reversible, although the intensity of the reflections begins to decrease already with smaller deformations (see below). After deformations
Supermolecular structure of polyethylene films affected by shrinkage
1529
of 40-60% unloading of the samples no longer produces any marked changes in the diffraction patterns of the latter. I n view of all the changes detected in the diffraction patterns, the following model of the changes in SMS is proposed. Before going on to describe the model we would observe that the crystalline
°
C
I/
ii :
c(
• IJ
d
FIG. 3. Schematic representation of changes in SMS corresponding to changes in the low and wide angle X-ray diffraction patterns at ~P1= 100°: a--original structure; b--rotation of crystalline layers accompanied by tilting of the macromolecular axes away from tlie texture axis (or extension of the helices); c--disruption of the crystalline layers; d--deformation of disoriented fibrillar structure (the scattering pattern corresponding to the model of structure c is depicted above by the fractured line in the schematic diagram of the low
angle diffraction pattern.
layers form what appear to be spirals. Experimental grounds favouring this assumption will be discussed below, but to begin the discussion we will disregard the helical nature of the layers, althouth the latter appear in helical form in Fig. 3 which illustrates the scheme of changes in SMS. Now, as the stretching proceeds, the layers rotate as a whole, tending to take up a position so that the long sides of the layers will be along the texture axis (Fig. 3a, b). Behaviour of this sort on the part of the layers is in agreement with para. 1. The rotation of the layers leads to a tilting of the c-axes (macromolecular axes) away from the texture axis (which agrees with para. 2). Rotation of the layers is simultaneously accompanied by a reduction in their
1530
B. ]Yl. GI~znuRG et al.
transverse dimensions (Fig. 3a), which agrees with paras. 3, 4. The latter consideration is also in agreement with the widening of the wide-angle equatorial reflections in the radial direction; the dimensions of crystallites along the v-axes (judging b y the width of the 002 reflections) remain virtually unaltered. I,,pa 'se/sec I20
~0
6O
~"
//
\/
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\ \\',\ \,, \
\',&'.\
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~""~ 80 ~o,~ l ~
FIG. 4. Curves of intensity distribution of low angle X-ray scattering along the meridional direction: Tx=100°; co=0 (1), 20 (2) and 50% (4); unloading after Co=30 (3), 50% (5). Let us now consider how far rotations of the layers are in agreement with rotations of the macromolecules. During stretching the 110 reflections are tilted 12 ° at most away from the equator, and the 200 reflections, b y not more than 8 to 10°. Taking the axial nature of the texture into account, and assuming the polyethylene lattice to be orthorhombie [8], it is possible to find the angle ~ f deviation of the c-axes from the texture axis. This angle ~ is approximately
Supermoleeular structure of polyethylene films affected by shrinkage
1531
12 ° (Fig. 3a). In view of the fact that in the original samples there was already a small deviation of the preferential direction of the c-axes from the texture axis, the magnitude of ~ is the maximum possible value for the angle of rotation of the macromolecules as result of the stretching of the samples. The tilting of the layers is characterized b y the angle between the texture axis and the direction from the centre of the low angle picture towards the maxim u m intensity of the low angle reflection [6]. The estimates that have been m a d e show that the layers rotate through 17-20 ° , i.e. rather more than the macromolecules in the crystallites (Fig. 3c). The difference in the slopes is naturally explained b y the fact that b y way of macromolecules passing through from one side to the other the forces brought to bear during stretching tend to turn the macromolecules of crystallite layers in a direction contrary to the rotation of the layers themselves In this w a y appreciable shear deformations occur within crystallite layers; the degree of shear m a y be as much as 20 °.
FIG. 5. Typical small and wide angle X-ray diffraction pictures at TI= 85°C; e0 (%); a--0; b-- 20; c-- 60. The proposed model of changes in SMS brings to mind structural changes occurring in the a~texture of polyethylene [7] in the early stages of drawing. A major difference, however, is seen in the tilting of the c-axes away from t h e texture axis, which increases as the stretching proceeds, and in the considerable resistance of the original crystallites to deformation, as is shown b y the fact that only a small proportion of them are destroyed, and a new SMS i s formed. As a result of 60% extension the intensity of the low angle reflections due to the original SMS is reduced b y approximately 45~o (see Fig. 4, cf. curves I and 5). This means that at most 45~o of the original SMS is broken down. The new SMS appears in the low angle diffraction pattern as a "connecting arch" of fairly low intensity, joining the low angle reflections due to the original SMS (Fig. lc, 2b). The arch-like (not striated) form of the reflection due to t h e new SMS is a natural consequence of macromolecules of the original SMS tilting away from the texture axis so that the forming fibrils of the new SMS have marked dispersion according to orientation (Fig. 3d). The long period typical for the new SMS is equal to ~ 155 /~, and coincides with the long period of t h e original SMS (Fig. la), if the latter is reckoned in terms of the distance from t h e reflection to the centre of diffraction pattern. The new SMS apparently consist~
1532
B. M. GINZBURO e t a l .
of fibrils formed as a result of separation of the original crystalline layers (Fig. 3c, d) * T1----85°. When T 1falls, resistance of the crystals to mechanical loads is reduced [1]. The impairment of mechanical resistance is apparently due to reduced crystallite dimensions and to impairment of the internal perfection of the crystallites; this is reflected in the widening of the wide angle reflections. Although the original low angle diffraction patterns for T I = 85 ° are similar to those observed for T l = 1 0 0 °, dissimilar changes in the patterns are observed during secondary stretching. The reflections are at first located along the layer lines (Fig. 5b),
FIO. 6. Small and wide angle X-ray diffraction pictures at T1 (°C): a, b--60; c, d--20; e0 ~o: a, c--0; b--30; c--20. but afterwards the original SMS is disrupted (approximately 70 to 80% disruption) and replaced by a new SMS that is reflected in the form of a "connecting arch" in the low angle diffraction pattern, and is now of considerable intensity (Figs. 5c and 2c). This must mean that a considerable amount of new SMS is formed, and that the contribution of the latter to the wide angle scattering should on the whole improve the orientation of the c-axes (of. Fig. 5a-c), i.e. there is a narrowing of the arcs of equatorial reflections in the wide angle diffraction pattern. The long period of the new SMS is practically identical with the periodicty of the layers in the original structure. I f the deformation does not exceed 40~/o, transitions to the new SMS are to a considerable extent reversible: after unloading, the connecting arch in the low angle diffraction pattern practically disappears, and reflections from the original SMS regain their former appearance. After 60% deformations unloading leads only to slight changes in the low angle diffraction pattern. * Jointly the c and d structures make the low angle diffraction pattern schematically portrayed in Fig. 2c.
Supormolecular structure of polyethylene films affected by shrinkage
1533
Let us now consider Fig. 5c where there are striations passing from the weakened reflections of the original SMS to the centre of the low angle diffraction pattern. The appearance of these striations is attributed to the formation of screw structures [7, 9]. If structures of this type are also present in the samples treated at T I = 100 ° (the proximity of the reflections to the strong diffuse scattering may prevent the "crosses" from appearing in the respective low angle diffraction patterns) there will then be a natural explanation for the splitting of the equatorial reflections in the wide angle diffraction patterns for TI--~ ]00 °. Let us say that the crystalline layers form helices (Fig. 3a), and that the macromolecules in the layers are directed approximately along the texture axis. Stretching of the helices will result in tilting of the macromolecules and in the splitting referred to above (Fig. 3b). Further changes in the diffraction patterns may, as before be attributed to the disruption of crystallites (Fig. 3c) and to the formation of a disoriented fibrillar structure (Fig. 3d). T1~---60°. The change in the wide angle diffraction patterns is the same as in the case of T1--~85 °. The quite obvious similarity is in the narrowing of the arcs of the l l 0 and 200 reflections (Fig. 6). As regards changes in the low angle patterns (Figs. 6a, b, 2d, e), these are fairly specific, and a proper understanding of the latter would probably call for investigation of a whole series of samples that had undergone orientation and quenching at different temperatureS. Certainly it is seen that during stretching an arc is formed in the low angle diffraction patterns with the ends of the arc clearly bending towards small scattering angles. (Figs. 6b, 2e). This transformation of the low angle diffraction pattern can be explained if one assumes that the reflections of the old SMS arc removed from the meridian at T~----60° and that the presence of new SMS is reflected in the formation of a meridional reflection in the form of a connecting arch. Owing to the proximity of the long periods of the initial SMS and the new SMS the respective reflections merge and form a common arc. Since the crystallites of the original SMS become skewed and rotate as a whole, reflections due to the latter move away from the meridian; at the same time the transfer of material from the old SMS to the new leads to intensified scattering in the vicinity of the meridian, but following the sizes of the scattering angles along the meridian it will be seen that these are somewhat larger. The effect of all these factors is elongate the "common" low angle. The azimuthal narrowing of the 110 and 200 reflections is due to transition of the SMS from the original to the new form, the latter pradominating in the final stages of stretching. T t = 2 0 °. The secondary stretching is carried out at a temperature practically equal to T~. An ordinary four petalled low angle diffraction pattern (Fig. 6c) is typical for the original samples. As in earlier investigations [2-4], it was found that, stretching resulted in transition to a normal striated (or rather "two-petalled") type of low angle diffraction pattern (Fig. 6d); at the same time careful examination of the latter revealed the formation of reflection "tails" similar to those observed at T ~ 6 0 °, "curling" towards small scattering angles. The
1534
B. M. GINZBURGet al.
orientation of crystallites improves during the stretching process. With deformations of e z 2 0 % or less, all changes in the low and wide angle diffraction patterns are reversible. At ~ 2 7 - 3 0 ~ o changes in the low angle patterns are irreversible, and remain as striations after unloading; at the same time cracks directed along the orientation axis appear in the samples. /
60
40
?0
~~- -' ~c 'r£' ~
I~
j
I
I ,
FIG. 7. e~ vs. to; T i l l 0 0 (1), 85 (2), 60 (3) and 20° (4).
Results of quantitative measurement of the low angle scattering intensity. The position of the reflections along the meridional direction is displaced, with stretch~ ing, towards small scattering angles (Fig. 4). I n the case of relatively small deformations (5-10%) changes in intensity (I) are completely reversible; within the limits of reversibility the value of I is invariably increased with stretching. The higher the value of T1, the more limited is the range of deformation in which changes in I are reversible (~5~o at T1----100°; 10~o at 60°; 20% at 20°). Given deformations exceeding 5-20% I is reduced, and with unloading only a partial recovery in I is observed. The reduction in I may be due, firstly to partial breakdown of the original SMS, and secondly to diminution of the transverse dimensions of crystallites in the latter. For TI~--100 ° deformation of the long periods ~ , determined from the displacement of reflections in the meridional direction is first of all equal, within limits of fairly big errors in measurement, to the macrodeformation of the samples Zo; thereafter, at e0zS0~o, the magnitude of ~ clearly exceeds that of t 0 (Fig. 7). As T1 falls, the tendency of e~ to exceed the macroscopic deformation t0 gradually disappears, and at T1~20 and 60 °, ~'4 clearly falls behind to. Where e~ exceeds
Thermooxidative degradation of polysulphamides
1535
~0 this is attributed [10] to microinhomogeneity of the sample as to deformation behaviour. As T1 rises, this inhomogeneity becomes more marked. The authors thank S. Ya. Frenkel for his support in connection with this investigation. Translated by R. J. A. I-IE~DRY
REFERENCES 1. B. M. G]NZBURG, N. S U L T ~ O V and S. Ya. FRENY~L, Vysokomo]. soyed. AI3: 2691, 1971 (Translated in P o l y m e r Sci. U.S.S.R.: 12: 2397, 1971) 2. A. COWKING, Y. G. RIDER, J. L. HAY and A. KELLER, J. Mater. Sci. 3: 646, 1968 3. A. COWKING a n d Y. G. RIDER, J. Mater. Sci. 4: 1051, 1969 4. A. K E L L E R a n d D. P. POPE, J. Mater. Sci. 6: 453, 1971 5. V. I. GERASIMOV and D. Ya. TSVANKIN, P r i b o r y i tekhnika eksperimenta, No. 2, 204, 1968 6. V. I: GERASIMOV and D. Ya. TSVANKIN, Vysokomol. soyed. B l l : 1652, 1969 (Not translated in P o l y m e r Sci. U.S.S.R.) 7. V. I. GERASIMOV and D. Ya. TSVANKIN, Vysokomol. soyed. A12: 2599, 1970 (Translated in Polymer Sci. U.S.S.R.: 11, 2944, 1970) 8. C. W. BUNN, Trans. F a r a d a y Soc. 35: 482, 1939 9. A. N. J. HEYN, Text. Res. J. 19: 163, 1949; J. Amer. Chem. Soc. 7@: 3138, 1948; 71: 1873, 1949; 72: 2284, 1950 10. M. A. GEZALOV, V. S. KUKSENKO and A. I. SLUTSKER, Mekhanika polimerov, No. 1, 51, 1972
THE THERMAL AND THERMOOXIDATIVE DEGRADATION OF POLYSULPHAMIDES* I . I . LEVA~TOVSKAYA, B. M. KOVARSKAYA, V. F . YAVOROVSK~¥A, V . P . I:)SI-IE~TITSY:bTA,G. V. DRALYUK, T. T. YAKOVEI~KO a n d P . A . GVOZD Plastics Research I n s t i t u t e Polytechnical Institute, L v o v {Received 1 November 1972)
The kinetics of the thermal and thermooxidative degradation of polysulphamides based on hexamethylenediamine and p,p'-diphenyldisulphochloride and p,p'-diphenylhydroxidedisulphochloride were investigated. This was carried out b y using the thermo, gravimetric m e t h o d under isothermal conditions and during dynamic heating, as well as b y investigation of the kinetics of oxygen absorption. The TG curves plotted under cond itions of dynamic heating in air are characterized b y two sections corresponding to * Vysokomol. soyed. A16: No. 6, 1324-1328, 1974.
1525
The conformational models of POBH seen in Fig. 2 are proposed in the light of the available data. It is possible to conclude on the basis of these results that the difference in the chemical properties of imidazole groups of the mono- and disubstituted POBH are due to the dissimilarity of the conformational state of these polymers. The authors thank M. F. Vyalykh for assisting with eleetrophoretic separation of the compounds, and are grateful for the help given by N. F. Kazanskaya in determining the pK values. Translated by R. J. A. HE~VD~Y REFERENCES
1. I. N. TOPCHIEVA, A. B. SOLOV'EVA and V. A. KABANOV, Dokl. AN SSSR 199:1084 1971 2. I. N. TOPCHIEVA, A. B. SOLOV'EVA and V. A. KABANOV, Vysokomol. soyed. A14: 825, 1972 (Translated in Polymer Sci. U.S.S.R. 4: 918, 1972) 3. Ye. M. SAMSONOV, Sb. Fiziko-khimich. metody isueheniya, aaaliza i fraktsionirovaniya biopolimerov (Physicoehemical methods of Investigation, Analysis and Fraetionation of Biopolymers). p. 40, Izd. "Nauka", 1966 4. D. GRINSHTEIN and M. VINITS, Khimiya aminokislot i peptidov (Chemistry of Amino Acids and Peptides). p. 784, Izd. "Mir", 1964 5. I. M. HEISS and K. I~a,TSEK, Paper chromatography, p. 797, 1963 6. S. KORMAN and N. T. CLARKE, J. Biol. Chem. 221: 113, 1956 7. A. ALBERT and E. SERGENT, Ionization Constants of Acids and Bases, p. 21, 1964
THE EFFECT OF STRETCHING ALONG THE ORIENTATION AXIS ON THE SUPERMOLECULAR STRUCTURE OF POLYETHYLENE FILMS AFFECTED BY SHRINKAGE* B. M. GmZBU~G, K. B. KV~BA~OV, Y~. A. LEOSKO, M. A. MA~TY~OVand D. I~ASItIDOV High Polymer Institute, U.S.S.R. Academy of Sciences
(Received 30 October 1972) Changes in the supermolecular structure of low-density polyethylene films were investigated by X-ray analysis with low and wide angle diffraction patterns. The films underwent preliminary orientation at temperatures of 100, 85, 60 and 20° and * Vysokomol soyed. A16: No. 6, 1317-1323, 1974.
1526
B.M.
GINZBURG et al.
were s u b j e c t e d to s h r i n k a g e a t t h e s a m e t e m p e r a t u r e s . T h e c h a n g e s i n s u p e r m o l e c u l a r s t r u c t u r e (SMS) were t h e r e s u l t of g r a d u a l e x t e n s i o n a t 20 ° a l o n g t h e t e x t u r e axis. A m o d e l of t h e c h a n g e s in SMS h a s b e e n p r o p o s e d .
I~ EARLIER investigations [1] it was shown that the stretching of oriented films of low density polyethylene (PELD) at various angles to the c-texture axis accompanied by the "skewing" of crystallites'. During the stretching of actual oriented systems along the texture axis one finds that for some of the crystallites the direction of stretching makes a certain angle with the c-axis, and one would therefore expect changes in the shape of these crystallites to be reflected in low angte X-ray diffraction patterns. These changes in the low angle diffraction patterns of polyethylene films were considered in papers [2-4], and the results obtained in the latter investigations are partly in agreement with those presented below, but by varying the conditions of sample preparation (mainly the temperature of primary orientation T1) we obtained a number of new results, and the total amount of experimental data available is considered solely from the point of view of the skewing of crystallites then their subsequent destruction after a critical degree of skew has been reached, and the new supermoleeular structure (SMS) which is then formed. Secondary stretching along the texture axis was carried out at room temperature only (T2~ 20°). After undergoing the primary orientation the samples were subjected to shrinkage at temperature T1. This resulted in a slight deviation of the preferential direction of the c-axes of the crystallites from the texture axis, but in this case subsequent secondary stretching causes major and easily recordable changes in the low angle diffraction patterns, the interpretation of which is conducive to better understanding of changes in the patterns appearing in the course of the stretching of samples that had not been subjected to shrinkage. MATERIALS AND METHODS ! P E L D films (the c h a r a c t e r i s t i c s of t h e P E L D are give.n i n [2] w i t h a t h i c k n e s s of 1.7 to 1-8 m m were p r e p a r e d b y c o m p r e s s i o n m o u l d i n g f r o m t h e m e l t (200 °) a t a p r e s s u r e of 50 a t m ; t h e s a m p l e s were t h e n q u e n c h e d i n w a t e r a t 2 2 °. T h e films h a d a fine s p h e r u l i t i e s t r u c t u r e (R ~ 2 - 3 / l m ) . A f t e r t h i s t h e films u n d e r w e n t u n i a x i a l d r a w i n g a t t e m p e r a t u r e s (T1) of 20, 60, 85 a n d 100 ° u p t o a s t r a i n in t h e v i c i n i t y of t h e p r e - b r e a k i n g s t r a i n , a n d were t h e n freed f r o m t h e c l a m p s , a n d h a v i n g b e e n allowed to c o n t r a c t freely a t t e m p e r a t u r e T1 for 1 hr, t h e films were t r a n s f e r r e d t o r o o m c o n d i t i o n s . All t h e films u n d e r w e n t g r a d u a l s t r e t c h i n g u n d e r r o o m c o n d i t i o n s , a n d as t h i s proceeded, low a n d wide a n g l e p h o t o g r a p h s were t a k e n s i m u l t a n e o u s l y , s h o w i n g one a n d t h e s a m e p a r t of a s a m p l e , u s i n g c a m e r a s of t h e t y p e p r o p o s e d i n [5]. A KI~M-1 i n s t r u m e n t (slit c o l l i m a t i o n ) was u s e d t o m e a s u r e t h e i n t e n s i t y of t h e l o w - a n g l e s c a t t e r i n g . T h e t r a n s v e r s e d i m e n s i o n s of c r y s t a l l i t e s (110 a n d 200 reflections) a n d t h e l o n g i t u d i n a l d i m e n s i o n s (002 reflection) were m e a s u r e d w i t h a U R S - 5 0 I M i n s t r u m e n t . Ni-filtered C u K ~ - r a d i a t i o n w a s used.
Supermolecular structure of polyethylene films affected by shrinkage
1527
DISCUSSION OF RESULTS
T1 --~100 °. The low angle diffraction patterns of the original samples (Figs. la, 2a) * are characteristic for a n axial t e x t u r e of crystalline layers inclined towards t h e t e x t u r e axis [6, 7] a n d the ratio of the t r a n s v e r s e dimensions of t h e
I ~h
FIG. 1. Typical small and wide angle X-ray diffraction pictures; T~--100°C; a--original sample; ~0~0; b--40; c--60%. layers to the longitudinal dimensions considerably exceeds unity. J u d g i n g b y t h e wide angle p a t t e r n s (Fig. la) the t e x t u r e of the crystallites is close to a c-texture; t h e m a x i m a of t h e l l 0 and 200 reflections are near the equator. The X - r a y p h o t o g r a p h s o b t a i n e d w h e n t h e p r i m a r y b e a m was parallel to the plane of the film a n d p e r p e n d i c u l a r to the t e x t u r e axis, show t h a t the t e x t u r e of the films is close to a n axial t e x t u r e (Figs. 1, 5, 6). As the stretching proceeds, the following qualitative changes are d e t e c t a b l e in the X - r a y p h o t o g r a p h s .
®
C((
J
®
))
I( `°
a
I
O
O
d FIG. 2. Schematic representations of some of the low and wide angle X-ray diffraction patterns: a--T~=100 °, original sample; b--ditto, e0=60%; e - - T l ~ 8 5 °, e0-----60°/o; d--T~ ~60 °, original sample; e--ditto, e0= 30%. * In Fig. 1 and in the other Figures the direction of the texture axis and of the secondary drawing is vertical; unless otherwise stated, the primary beam direction is perpendicular to the texture axis and to the plane of the polymer film, and the sample is under a load that has caused deformation.
1528
B.M.
GINZBU~G et al.
1. The centres of the low angle reflections move away from the meridian (Fig. lb, c), which, in view of the small changes in the longitudinal dimensions of the crystallites, as well as the smaller transverse dimensions of the latter (see below) can only mean an increase in the angle of slope of the layers or in the angle of skew ~] of the crystallites [1, 6].
I!
i!
0
a
b FIG. 3 (a, b)
2. The arcs of equation reflections in the wide angle pictures grow wider in the azimuthal direction, and in the final stages of deformation (Figs. lc, 2b) the splitting of the arcs is clearly seen. 3. Reflections of the "spot" or radical [6] types in the low angle pictures become more and more anisotropie in their dimensions, and in the final stages of deformation these appear as "striations" (Figs. lc, 2b). The change to reflections of the striation type means that the transverse dimensions of the layers have become smaller. 4. The direction of the "striations" in the low angle pictures approximates more and more to a horizontal direction; a low angle picture comes to resembie a well-resolved diffraction pattern of the "four-petal" type, (Figs. la, 2b) in contrast to the "four-spot" (Fig. la). 5. Up to deformations of around 20% changes in the shape of a low angle pattern are practically reversible, although the intensity of the reflections begins to decrease already with smaller deformations (see below). After deformations
Supermolecular structure of polyethylene films affected by shrinkage
1529
of 40-60% unloading of the samples no longer produces any marked changes in the diffraction patterns of the latter. I n view of all the changes detected in the diffraction patterns, the following model of the changes in SMS is proposed. Before going on to describe the model we would observe that the crystalline
°
C
I/
ii :
c(
• IJ
d
FIG. 3. Schematic representation of changes in SMS corresponding to changes in the low and wide angle X-ray diffraction patterns at ~P1= 100°: a--original structure; b--rotation of crystalline layers accompanied by tilting of the macromolecular axes away from tlie texture axis (or extension of the helices); c--disruption of the crystalline layers; d--deformation of disoriented fibrillar structure (the scattering pattern corresponding to the model of structure c is depicted above by the fractured line in the schematic diagram of the low
angle diffraction pattern.
layers form what appear to be spirals. Experimental grounds favouring this assumption will be discussed below, but to begin the discussion we will disregard the helical nature of the layers, althouth the latter appear in helical form in Fig. 3 which illustrates the scheme of changes in SMS. Now, as the stretching proceeds, the layers rotate as a whole, tending to take up a position so that the long sides of the layers will be along the texture axis (Fig. 3a, b). Behaviour of this sort on the part of the layers is in agreement with para. 1. The rotation of the layers leads to a tilting of the c-axes (macromolecular axes) away from the texture axis (which agrees with para. 2). Rotation of the layers is simultaneously accompanied by a reduction in their
1530
B. ]Yl. GI~znuRG et al.
transverse dimensions (Fig. 3a), which agrees with paras. 3, 4. The latter consideration is also in agreement with the widening of the wide-angle equatorial reflections in the radial direction; the dimensions of crystallites along the v-axes (judging b y the width of the 002 reflections) remain virtually unaltered. I,,pa 'se/sec I20
~0
6O
~"
//
\/
\2"\~\,
\'x \ i \
\ \\',\ \,, \
\',&'.\
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\",,\"× \ I
I 20
I
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I
I gO
~""~ 80 ~o,~ l ~
FIG. 4. Curves of intensity distribution of low angle X-ray scattering along the meridional direction: Tx=100°; co=0 (1), 20 (2) and 50% (4); unloading after Co=30 (3), 50% (5). Let us now consider how far rotations of the layers are in agreement with rotations of the macromolecules. During stretching the 110 reflections are tilted 12 ° at most away from the equator, and the 200 reflections, b y not more than 8 to 10°. Taking the axial nature of the texture into account, and assuming the polyethylene lattice to be orthorhombie [8], it is possible to find the angle ~ f deviation of the c-axes from the texture axis. This angle ~ is approximately
Supermoleeular structure of polyethylene films affected by shrinkage
1531
12 ° (Fig. 3a). In view of the fact that in the original samples there was already a small deviation of the preferential direction of the c-axes from the texture axis, the magnitude of ~ is the maximum possible value for the angle of rotation of the macromolecules as result of the stretching of the samples. The tilting of the layers is characterized b y the angle between the texture axis and the direction from the centre of the low angle picture towards the maxim u m intensity of the low angle reflection [6]. The estimates that have been m a d e show that the layers rotate through 17-20 ° , i.e. rather more than the macromolecules in the crystallites (Fig. 3c). The difference in the slopes is naturally explained b y the fact that b y way of macromolecules passing through from one side to the other the forces brought to bear during stretching tend to turn the macromolecules of crystallite layers in a direction contrary to the rotation of the layers themselves In this w a y appreciable shear deformations occur within crystallite layers; the degree of shear m a y be as much as 20 °.
FIG. 5. Typical small and wide angle X-ray diffraction pictures at TI= 85°C; e0 (%); a--0; b-- 20; c-- 60. The proposed model of changes in SMS brings to mind structural changes occurring in the a~texture of polyethylene [7] in the early stages of drawing. A major difference, however, is seen in the tilting of the c-axes away from t h e texture axis, which increases as the stretching proceeds, and in the considerable resistance of the original crystallites to deformation, as is shown b y the fact that only a small proportion of them are destroyed, and a new SMS i s formed. As a result of 60% extension the intensity of the low angle reflections due to the original SMS is reduced b y approximately 45~o (see Fig. 4, cf. curves I and 5). This means that at most 45~o of the original SMS is broken down. The new SMS appears in the low angle diffraction pattern as a "connecting arch" of fairly low intensity, joining the low angle reflections due to the original SMS (Fig. lc, 2b). The arch-like (not striated) form of the reflection due to t h e new SMS is a natural consequence of macromolecules of the original SMS tilting away from the texture axis so that the forming fibrils of the new SMS have marked dispersion according to orientation (Fig. 3d). The long period typical for the new SMS is equal to ~ 155 /~, and coincides with the long period of t h e original SMS (Fig. la), if the latter is reckoned in terms of the distance from t h e reflection to the centre of diffraction pattern. The new SMS apparently consist~
1532
B. M. GINZBURO e t a l .
of fibrils formed as a result of separation of the original crystalline layers (Fig. 3c, d) * T1----85°. When T 1falls, resistance of the crystals to mechanical loads is reduced [1]. The impairment of mechanical resistance is apparently due to reduced crystallite dimensions and to impairment of the internal perfection of the crystallites; this is reflected in the widening of the wide angle reflections. Although the original low angle diffraction patterns for T I = 85 ° are similar to those observed for T l = 1 0 0 °, dissimilar changes in the patterns are observed during secondary stretching. The reflections are at first located along the layer lines (Fig. 5b),
FIO. 6. Small and wide angle X-ray diffraction pictures at T1 (°C): a, b--60; c, d--20; e0 ~o: a, c--0; b--30; c--20. but afterwards the original SMS is disrupted (approximately 70 to 80% disruption) and replaced by a new SMS that is reflected in the form of a "connecting arch" in the low angle diffraction pattern, and is now of considerable intensity (Figs. 5c and 2c). This must mean that a considerable amount of new SMS is formed, and that the contribution of the latter to the wide angle scattering should on the whole improve the orientation of the c-axes (of. Fig. 5a-c), i.e. there is a narrowing of the arcs of equatorial reflections in the wide angle diffraction pattern. The long period of the new SMS is practically identical with the periodicty of the layers in the original structure. I f the deformation does not exceed 40~/o, transitions to the new SMS are to a considerable extent reversible: after unloading, the connecting arch in the low angle diffraction pattern practically disappears, and reflections from the original SMS regain their former appearance. After 60% deformations unloading leads only to slight changes in the low angle diffraction pattern. * Jointly the c and d structures make the low angle diffraction pattern schematically portrayed in Fig. 2c.
Supormolecular structure of polyethylene films affected by shrinkage
1533
Let us now consider Fig. 5c where there are striations passing from the weakened reflections of the original SMS to the centre of the low angle diffraction pattern. The appearance of these striations is attributed to the formation of screw structures [7, 9]. If structures of this type are also present in the samples treated at T I = 100 ° (the proximity of the reflections to the strong diffuse scattering may prevent the "crosses" from appearing in the respective low angle diffraction patterns) there will then be a natural explanation for the splitting of the equatorial reflections in the wide angle diffraction patterns for TI--~ ]00 °. Let us say that the crystalline layers form helices (Fig. 3a), and that the macromolecules in the layers are directed approximately along the texture axis. Stretching of the helices will result in tilting of the macromolecules and in the splitting referred to above (Fig. 3b). Further changes in the diffraction patterns may, as before be attributed to the disruption of crystallites (Fig. 3c) and to the formation of a disoriented fibrillar structure (Fig. 3d). T1~---60°. The change in the wide angle diffraction patterns is the same as in the case of T1--~85 °. The quite obvious similarity is in the narrowing of the arcs of the l l 0 and 200 reflections (Fig. 6). As regards changes in the low angle patterns (Figs. 6a, b, 2d, e), these are fairly specific, and a proper understanding of the latter would probably call for investigation of a whole series of samples that had undergone orientation and quenching at different temperatureS. Certainly it is seen that during stretching an arc is formed in the low angle diffraction patterns with the ends of the arc clearly bending towards small scattering angles. (Figs. 6b, 2e). This transformation of the low angle diffraction pattern can be explained if one assumes that the reflections of the old SMS arc removed from the meridian at T~----60° and that the presence of new SMS is reflected in the formation of a meridional reflection in the form of a connecting arch. Owing to the proximity of the long periods of the initial SMS and the new SMS the respective reflections merge and form a common arc. Since the crystallites of the original SMS become skewed and rotate as a whole, reflections due to the latter move away from the meridian; at the same time the transfer of material from the old SMS to the new leads to intensified scattering in the vicinity of the meridian, but following the sizes of the scattering angles along the meridian it will be seen that these are somewhat larger. The effect of all these factors is elongate the "common" low angle. The azimuthal narrowing of the 110 and 200 reflections is due to transition of the SMS from the original to the new form, the latter pradominating in the final stages of stretching. T t = 2 0 °. The secondary stretching is carried out at a temperature practically equal to T~. An ordinary four petalled low angle diffraction pattern (Fig. 6c) is typical for the original samples. As in earlier investigations [2-4], it was found that, stretching resulted in transition to a normal striated (or rather "two-petalled") type of low angle diffraction pattern (Fig. 6d); at the same time careful examination of the latter revealed the formation of reflection "tails" similar to those observed at T ~ 6 0 °, "curling" towards small scattering angles. The
1534
B. M. GINZBURGet al.
orientation of crystallites improves during the stretching process. With deformations of e z 2 0 % or less, all changes in the low and wide angle diffraction patterns are reversible. At ~ 2 7 - 3 0 ~ o changes in the low angle patterns are irreversible, and remain as striations after unloading; at the same time cracks directed along the orientation axis appear in the samples. /
60
40
?0
~~- -' ~c 'r£' ~
I~
j
I
I ,
FIG. 7. e~ vs. to; T i l l 0 0 (1), 85 (2), 60 (3) and 20° (4).
Results of quantitative measurement of the low angle scattering intensity. The position of the reflections along the meridional direction is displaced, with stretch~ ing, towards small scattering angles (Fig. 4). I n the case of relatively small deformations (5-10%) changes in intensity (I) are completely reversible; within the limits of reversibility the value of I is invariably increased with stretching. The higher the value of T1, the more limited is the range of deformation in which changes in I are reversible (~5~o at T1----100°; 10~o at 60°; 20% at 20°). Given deformations exceeding 5-20% I is reduced, and with unloading only a partial recovery in I is observed. The reduction in I may be due, firstly to partial breakdown of the original SMS, and secondly to diminution of the transverse dimensions of crystallites in the latter. For TI~--100 ° deformation of the long periods ~ , determined from the displacement of reflections in the meridional direction is first of all equal, within limits of fairly big errors in measurement, to the macrodeformation of the samples Zo; thereafter, at e0zS0~o, the magnitude of ~ clearly exceeds that of t 0 (Fig. 7). As T1 falls, the tendency of e~ to exceed the macroscopic deformation t0 gradually disappears, and at T1~20 and 60 °, ~'4 clearly falls behind to. Where e~ exceeds
Thermooxidative degradation of polysulphamides
1535
~0 this is attributed [10] to microinhomogeneity of the sample as to deformation behaviour. As T1 rises, this inhomogeneity becomes more marked. The authors thank S. Ya. Frenkel for his support in connection with this investigation. Translated by R. J. A. I-IE~DRY
REFERENCES 1. B. M. G]NZBURG, N. S U L T ~ O V and S. Ya. FRENY~L, Vysokomo]. soyed. AI3: 2691, 1971 (Translated in P o l y m e r Sci. U.S.S.R.: 12: 2397, 1971) 2. A. COWKING, Y. G. RIDER, J. L. HAY and A. KELLER, J. Mater. Sci. 3: 646, 1968 3. A. COWKING a n d Y. G. RIDER, J. Mater. Sci. 4: 1051, 1969 4. A. K E L L E R a n d D. P. POPE, J. Mater. Sci. 6: 453, 1971 5. V. I. GERASIMOV and D. Ya. TSVANKIN, P r i b o r y i tekhnika eksperimenta, No. 2, 204, 1968 6. V. I: GERASIMOV and D. Ya. TSVANKIN, Vysokomol. soyed. B l l : 1652, 1969 (Not translated in P o l y m e r Sci. U.S.S.R.) 7. V. I. GERASIMOV and D. Ya. TSVANKIN, Vysokomol. soyed. A12: 2599, 1970 (Translated in Polymer Sci. U.S.S.R.: 11, 2944, 1970) 8. C. W. BUNN, Trans. F a r a d a y Soc. 35: 482, 1939 9. A. N. J. HEYN, Text. Res. J. 19: 163, 1949; J. Amer. Chem. Soc. 7@: 3138, 1948; 71: 1873, 1949; 72: 2284, 1950 10. M. A. GEZALOV, V. S. KUKSENKO and A. I. SLUTSKER, Mekhanika polimerov, No. 1, 51, 1972
THE THERMAL AND THERMOOXIDATIVE DEGRADATION OF POLYSULPHAMIDES* I . I . LEVA~TOVSKAYA, B. M. KOVARSKAYA, V. F . YAVOROVSK~¥A, V . P . I:)SI-IE~TITSY:bTA,G. V. DRALYUK, T. T. YAKOVEI~KO a n d P . A . GVOZD Plastics Research I n s t i t u t e Polytechnical Institute, L v o v {Received 1 November 1972)
The kinetics of the thermal and thermooxidative degradation of polysulphamides based on hexamethylenediamine and p,p'-diphenyldisulphochloride and p,p'-diphenylhydroxidedisulphochloride were investigated. This was carried out b y using the thermo, gravimetric m e t h o d under isothermal conditions and during dynamic heating, as well as b y investigation of the kinetics of oxygen absorption. The TG curves plotted under cond itions of dynamic heating in air are characterized b y two sections corresponding to * Vysokomol. soyed. A16: No. 6, 1324-1328, 1974.