The effect of molecular interactions on the macromolecular structure of oriented polyamide-6 films and changes during elastic stretching

THE EFFECT OF MOLECULAR I~TERACTIONS ON THE MACROMOLECULAR STRUCTURE OF ORIENTED POLYAMIDE-6 FILMS AND CHANGES DURING ELASTIC STRETCHING* B. M. GINZBURG,K. B. KURBANOV and B. A. ASHEROV High Polymers Institute, U.S.S.R. Academy of Science

(Received 26 June 1972) Low and wide angle X-ray scattering studies on oriented polyamide-6 films have shown that molecular interaction forces can greatly influence the macromolecular structure and its changes during deformation a n d heat treatment. When the samples were annealed in a fixed state there is little distortion of the crystal layers or bending ofcrystallites. When annealing is carried out in the free (unsuspended) state, H-bonds prevent "sagging" of the layers a~d distortion becomes much greater at right angles to the H-bonds. The latter is reversible when elastic stretching is applied to the samples at room temperature and the original state is almost restored to t h a t existing in direction of the H-bonds.

L o w angle X-ray (LAX) study of oriented crystalline polymers which are under a stress can give useful information about the correlation between their macromolecular structure (MMS) and mechanical properties [1-11]. Much attention was given in recent studies to the intensity distribution in meridional direction when stress was applied to an oriented sample in the direction of orientation [1-5, 11] (or in direction of coordinate Z in inverse space). Another s t u d y had shown [12] that re-orientation of oriented polyethylene films at various angles (25-75 °) to the primary orientation axis produced bending in the L A X pictures, which was interpreted in terms of crystallite bending [13]. As dispersion of the crystallites exists in truly oriented polymers with respect to orientation one can assume that many of the crystallites will change their shape when an elongation stress is applied in the direction of texture; the result of this will be a change of L A X in direction of the layer lines (or in that of coordinate X of the inversed space). Such changes were actually observed when stress was applied to polyethylene samples [6-10]. With polyamide-6 (PA-6), strong anisotropy of the molecular interaction forces was created b y the presence of the planarity of H-bond positions. This enabled us to examine the marked effect of the molecular interactions on the L A X on PA-6 in this s t u d y (and thus of MMS), as well as the changes which occur when stress is applied in the direction of orientation. * Vysokomol. soyed. A16: :No. 3, 558-565, 1974. 644

M a c r o m o l e c u l a r s t r u c t u r e of oriented p o l y a m i d e - 6 films

645

EXPERIMENTAL P A - 6 films used were 70 p m t h i c k a n d were p r o d u c e d b y e x t r u s i o n of a g r a d e P K - 4 industrial sheeting which was s t r e t c h e d a t r o o m t e m p e r a t u r e a t r i g h t angles to t h e e x t r u s i o n d i r e c t i o n to a p r e - b r e a k i n g state, t h e n t e m p e r e d 2 hr a t 200°C in c l a m p e d (samples I) a n d free (samples 2) states. There was some shrinkage of samples 2 during t e m p e r i n g (about 10 ~ ) . T h e sample dimensions were a p p r o x i m a t e l y 20 × 2 × 0.7 m m 3 w h e n s u b j e c t e d to e x a m i n a tion.

¸

L

D

¢

J F I e . 1. a, b - - T h e LAX~ v, d - - the W A X of samples 1 p h o t o g r a p h e d face-on; e~, ~ : a, c - - 0 ; b, d - - 1 0 . The o r i e n t a t i o n here a n d in all t h e o t h e r figures is vertical. FIG. 5. a, b - - T h e L A X ; c, d - - t h e W A X of samples 2 t a k e n side-on; a, b - - e ~ = 0 ; d - - 1 0 ~ . T h e X - r a y pictures were t a k e n in two positions: t h e fiat film was (1) n o r m a l to t h e p r i m a r y beam; (2) parallel to t h e beam, while the axis of o r i e n t a t i o n was p e r p e n d i c u l a r to t h e prim a r y b e a m (side view). A u t o m a t i c cameras and t h e s a m e p r i m a r y b e a m were used to t a k e t h e L A X a n d wide angle X - r a y ( W A X ) p h o t o g r a p h s . T h e d i s t a n c e f r o m t h e c a m e r a to t h e film was 156 m m in t h e case of L A X , 46 turn in t h e case of W A X . T h e b e a m angle was 20', th(~' source a s h a r p l y

0 34-1 74-3 0.04-0.1

Sc,

I(,)/I(0) 1 1.64-0.3 2.04-0.3 1.04-0.1

A ~z/~t 0.644-0-03 0-544-0.03 0.604-0.03 0.644-0.03

A~z, rain of angle

434-1 354-1 384-1 434-1

c,A 80+1 83+1 864-1 804-1

~M, min of anglo

664-1 644-1 624-1 664-1

704-1 954-1 70+ 1

734-1

544-1

1.74-0.1 1.44-0.1 1.74-0.1

734-1

D,k

A ~x, min of anglo

0 104-0.05 204-0.03 304-0"02 0.04-0.1 (stress)

0 74-3 17J=4 42±6 0.0+0.1

~c, 614-1 574-1 52J=1 434-1 614-1

~,~, rain of a.ng]o 88±2 94+2 1024-2 1244-3 884-2 25.5±1 27.5~1

25.54-1

27.54-1 25.54-1

A ~z, min of anglo 0.45-4-0"03 0.45-{-0.03 0.504-0'03 0.604-0.04 0.454-0"03

2.704-0.05 2.004-0"06 1.70+0-03 1.404-0.03 2.80±0.05

564-1 414-1 354-1 754-1

754-1

A~x, m i n o f anglo

3.04-0.1 1.04-0.1

2.o4-o.1

1 1.44-0.1

I (,)1I (0)

D,A 684-1 914-2 1254-3 1474-3 684-1

TABLE 2. C ] ~ G E S OF THE DISTRIBUTIOI~ INTENSITY PA_RA~ETE:R~ I1~ THE L A X oF SAM~I,ES 2 DURIlqG DEFOR~ff.ATIOI~I (Face-on photographs, pa/pk=0"2-0"5

JP,~'tnarks. ~ M--Angle position of the L A X rcflexion intensity peak; e--larger period; J , , , Z l , x - h a f f - w i d t h of the L A X reflexion intensity in direction Z a n d X; 8 , , 8z-- respective dimension of reflexions (from the black edges on the X - r a y pictures); I(e)-- peak intensity of the reflexion during sample deformation, I(0) -- without deformation.

0 5±0.1 104-0.05 0.04-0.1 (stress)

~,%

TABLE 1. CB~klffGESOF THE ~ I f f S I T Y DISTI~IBUTIOI~PARAMETERS I1~ THE L A X HEI*LEXlOI~S DURING THE DEFOB~f~kTIOlqOF SAMPLES 1 (Face-on photographs, Pa/Pk about 0.55)

Q

647

Macromolecular structure of oriented polyamide-6 films focusing B S V - 1 0

tube. The intensity measurements were carried out on the standard

KRM-I apparatus and on URS-50 IYI. The slits used in the KRM-1 were 0.05, 0.1 and 0.1 m m and the primary beam angle was 2.9', while it was 14' at the receiving slit. Using URS-50 IM, the slits used were 0'5, 0.25 and 0-25 ram, the primary beam angle was 1.5 °. The source was always CuK, radiation filtered through Ni.

I, l~a/s ec a 15

fO

5

I

,

I

20

GO

~, ,wln

FIG. 2

i

1_ fog

I

I

f8

i

I

22

p

I

L ,

Z5 Z8 °

Fro. 3

FIG. 2. The distribution of LAX scattering intensity in meridiona] direction using samples I. e~, ~ : 1--0; 2--5; 3--10. FIG. 3. The distribution of WAX scattering intensity in equatorial direction of highly oriented a-modification PA-6 crystallites with an axis] texture (fibres).

Samples 1. T h e L A X reAiexions (Fig. l a , b) were similar to radial t y p e ones [12]. T h e reflexion i n t e n s i t y increased as a result of s t r e t c h i n g (Fig. 2). T h e d e f o r m a "tion of t h e larger periods So, d e t e r m i n e d f r o m t h e reflcxion p e a k d i s p l a c e m e n t s (Fig. 2), was slightly smaller t h a n t h e m a c r o - d e f o r m a t i o n o f t h e s a m p l e e~ (Table 1). Changes in s h a p e of t h e reflexions were difficult to detect; t h e r e l a t i v e reflexion widths in direction Z r e m a i n e d p r a c t i c a l l y t h e s a m e (Table 1). T h e W A X (Fig. iv, d) were t y p i c a l of t h e a-modification of P A - 6 c o n t a i n i n g a small f r a c t i o n of small erystallites h a v i n g t h e y-modification [14]. T h e " i n t e r n a l " reflexion i n t e n s i t y I200 was slightly larger t h a n t h a t of I00~ in t h e W A X o f t h e original s a m p l e s w h e n t h e flat film was n o r m a l to t h e b e a m . As to t h e a x i a l t e x t u r e t h e o p p o s i t e was true, i.e. 1200 was smaller t h a n Ioo ~ (Fig. 3), b u t t h e difference was small in t h e ease of No. 1 samples, so t h a t one can r e g a r d t h e t e x t u r e to be here close to a n a x i a l one. T h e I2oo/Ioo ~ r a t i o r e m a i n e d a h n o s t c o n s t a n t w h e n stress was a p p l i e d (Fig. ]d).

648

B. Mo GINZBURG e$ ~ .

Samples 2. L A X for the original samples (in face-on photographs) showed fairly diffuse (unresolved) "small boat" type reflexions in direction of the layer lines (Fig. 4a). Such reflexions are typical of fibrils made up of slightly bent crystal-

(

Z~

(

c

r

e FIG. 4. The diffraction picture changes using LAX (upper rowi and WAX (lower row) as a function of sample 2 deformation when taken face-on, e~, ~/o: a--0; b--10; c--20; d--25; e--0 (after stress removal following 25 ~ deformation). lites. Parameter bt/a does not exceed 0.8 according to Gerasimov and Tsvankin [13] in an axial texture (b, a--transverse and longitudinal dimensions of crystallites, t = t a n ~ , ( 9 0 i ~ ) - - a n g l e s of the parallelogram representing the longitudinal section through the bent crystallite). However, our samples do not contain

Macromolecular structure of oriented polyamide-6 films

649

an axial texture. The I,o o in a face-on photograph is much smaller than the I00 , (Fig. 4a, lower series) while the inequality sign changes in the side-on photograph (Fig. 5c). The L A X also differ in shape when photographing in differing directions (Figs. 4a and 5a); the L A X in side-view are more like radial-type reflexions. The lack of axiality means that the bt/a can be even smaller than 0.8 to get the L A X shown in Fig. 4a.

p/ISBC 70 i

4

/0

50

~, rain

20

F I e . 6. T h e dis%ribution of L A X s c a t t e r i n g i n t e n s i t y i n m e r i d i o n a l d i r e c t i o n for s a m p l e s 2; eM, Yo: 1 - 0; 2 - - 1 0 ; 3 - - 2 0 ; 4 - - 3 0 . T h e b l a c k circles are for eM = 0 a f t e r s t r e s s r e m o v a l .

The most interesting result was the change in L A X in samples 2 when subjected to stress; the face-on photographs (Fig. 4b-d) showed in addition to the broadening of the reflexions also their shift towards smaller angles [3] and a narrowing of the reflexions in direction X, so that they take on the appearance of radial t y p e reflexions. The knowledge of the intensity distribution in direction Z (Fig. 6) makes it possible to estimate their half-width in direction X and also (Table 2) the lateral dimension of the coherent scattering region D [15] (disregarding the curvature of the crystallites) from the ratio of the "dimensions" of the reflexions in the L A X picture. The enlargement of relative width in direction Z of the reflexions (Table 2) and the appearance of much diffusive scattering (Fig. 6) when stress is applied should be noted, this could be due to either pores appearing [16], or to a deterioration in the regularity of alternation of crystalline and amorphous zones [17]. The face-on photographs show ec~e M in the initial phases of stress application, after which ec becomes obviously larger than e~ (Table 2). The W A X of samples under stress shows an increase of the I~oo/Ioo ~ ratio (compare Fig. 4a with 4d), the respective reflexions becoming comparable in intensity. The changes in shape of the W A X are hardly notice~,ble in side-on L A X (Fig. 5a, b) when stress is applied. Instead of the original I~o0
650

B.M. GINZBURGet al. RESULTS

Kuksenko and co-workers had shown [18] that the elastic stretching of oriented PA-6 films across their orientation axis causes the crystallites to turn around their b-axis (that of the macromolecules); the I2o0 becomes larger than the lo03 when this happens, although the inequality sign w~s the opposite in the original state. This turning of the erystallites happens in such a manner that the crystallographic planes containing the intermolecular H-bonds lines will dominantly lie in the plane of the polymer film while the H-bond lines themselves will be pointing in the direction of the applied elongation force. TABLE

3. T H E

REFLEXION

INTENSITY

SAMPLES

State of sample

Original Understress Stress removed

RATIO

I~oo/Ioo2

FOR

2

Direction of X-ray beam face-on side-on

I2oo
I2oo>Ioo~ Iloo ~ Ioo, I~oo >Ioo2

The I2oo/Ioo~ ratio changes in samples 2 in the same direction as during lateral re-orientation (Table 3) will be observable in the following cases: 1) when changing from the face-on to the side-on photography in the case of samples free from stress; 2) when stress is applied and the photo is taken face-on, or after stress removal when the photo is taken side-on. B y analogy with [18], one can therefore conclude that the H-bond lines are approximately normal to the plane of the film where samples are not under stress. The result of stress application is that the system of crystallites attempts to form an axial c-texture, b u t returns to the previous planar texture, with the H-bond orientation dominantly at right angles to the film plane, when stress is removed. There is no obvious reason for H-bond orientation in the plane of the oriented polymer film when it is subjected to lateral re-orientation. In addition to the present theories [18] one can also assume that the crystallites will orientate themselves in direction of the longer side during stretching under otherwise identical conditions. It was natural to say that the remote order will extend over a larger distance in the direction of the H-bond than in lateral direction. The measurements of the 002 and 200 reflexion widths m~,de it clear that/~he erystallite dimensions (calculated b y me~ns of the Scherrer formula after correction for the primary beam width) in direction [100] and [001] are about 190 and about 65 A respectively. The L A X scattering dat~ also show the crystalline zones to have larger lateral dimensions in direction of the H-bonds (value D and Table 2). The reason for the H-bond rotation to normality relative to the

Macromolecular structure of oriented polyamide-6 films

651

film plane after the re-orientation of the samples and tempering in the free state is still obscure. B y comparing the L A X of sample 1 with sample 2 (Fig. 1, 4) one comes to the conclusion that the intensified intermolecular reactions maintain the shape of the crystalline zones during the thermal shrinkage of the samples, i.e. that a large change in shape can be detected only when the samples 2 are photographed face-on when the X-ray beam is parallel with the H-bonds. Similarly, the H-bonds produce enhanced resistance to mechanical stresses, i.e. a considerable change in the shape of the L A X reflexions takes place during the lengthwise stretching of the samples 2, and of D, when photographed face-on, b u t hardly at all when side-on (Fig. 4 and Table 2). The reversibility of deformation and MMS changes up to 10re-breaking conditions ( s ~ 3 0 % ) can also be the consequence of strong intermolecular reactions.

FIG. 7. a--Scheme of the longitudinal section through a fibril model with convoluted layers; b--copy reduced in size which is fiat to produce the optical diffraction picture. The sudden increase in intensity of diffusive scattering and of relative width of the low-angle peak in direction Z is a typical feature of the L A X during the deformation of samples tempered in the free state [ll, 19]. This appears to be linked with a formation of zones having sharply differing deformation properties [2, 4, 11, 19]. Those with greatly reduced density of amorphous zones will stretch most under stress and create first a large scattering intensity at the peak, thus contributing most to this intensity; secondly, t h e y produce a larger shift of the reflexion peak into the L A X scattering range, and thirdly, at very non-uniform

652

B . M. GINZBURG et al.

deformations, these zones can lose t h e i r c h a r a c t e r of a m o n o p l a n a r network, so t h a t t h e s c a t t e r p r o d u c e d will c o n t r i b u t e to t h e diffusive c o m p o n e n t of scattering. T y p i c a l also is t h e " s t r e t c h i n g " of t h e low angle p a r t of the reflexion towards t h e c e n t r e of the X - r a y p i c t u r e (Fig. 4b, c), e v e n w h e n t h e r e is n o t m u c h diffusive s c a t t e r (Fig. 6, c u r v e 2). As t h e s y m m e t r y of t h e curves in Fig. 6, p r o d u c e d b y slit collimations, remains m u c h t h e same (compare c u r v e 1 w i t h 2), one gets t h e impression of a r e l a t i v e l y large L A X i n t e n s i t y of the reflexion " t a i l " . This is in a g r e e m e n t with t h e t h e o r y of zones w i t h larger d e f o r m a t i o n c a p a c i t y causing a s u d d e n increase in scattering i n t e n s i t y during d e f o r m a t i o n . V e r y small pa/pk values (Tables 1 a n d 2) resulted (amorphous and crystalline zone intensities respectively) w h e n this ratio was assessed f r o m t h e changes in i n t e n s i t y of low angle reflexion during elastic d e f o r m a t i o n of t h e samples [3]. T h e values of t h e r a t i o were closer to t h e m i n i m u m t h a n to t h e average; such a conclusion appears to be justified n o t o n l y for t h e t e m p e r e d samples. The L A X ohanges in direotion X. T h e scheme of the M ~ S of oriented crystalline p o l y m e r s is o f t e n illustrated as an a l t e r n a t i n g sequence of c o n v o l u t e d crystalline a n d a m o r p h o u s zones in the direction of t h e axis of o r i e n t a t i o n [20]; such a n illustration was confirmed b y e l e c t r o n m i c r o s c o p y (see Figs. 6, 7 [21]). T h e shape o f c u r v e illustrating t h e u p p e r boundaries of the layers directly affects t h e i n t e n s i t y distribution in direction X. This position h a d been generally described b y B o n a r t [22]. One can try to approximate this curve by means efvarious functions. It should be remembered that the function must not be a periodic one, ner must the number of periods be so limited that diffraction effects will be created at the equator of the LAX. The fibril illustrated in Fig. 7a can serve as a diffraction model of alternating convoluted layers in first approximation. The slope of the upper and lower boundaries of a single crystalline zone intersects the macromoleeule and changes the sign in front relative to horizontal directien. The calculation of the diffraction intensity created by an ideally regular fibril of this shape gave the following function for the q-I-th layer line:

in which l is tho length of amorphous zone in direction Z. Equation (1) is easily obtained [13] when the zones with differing slopes ef upper (lower) boundaries are considered non-coherent. The first two terms in the second squared bracket of eqn. (1) thus belongs to the diffraction of fibrils having crystals bent in opposite directions [13], while the third term is the

M~cromoleeular structure of oriented polyamide-6 films

653

result of fibrillar coherence. For some compositional parameters of the fibril illustrated in Fig. 7a (bt/c:n, in which n is integral), the scattering intensity on the meridian of the X-ray picture (X=O)will be zero and the image will be a four point image. The optical diffraction picture (Fig. 7b) of the flat model shown in Fig. 7a con_6rms the correctness of the described assumptions. The changes in layer shape (which is factually the same as the crystal curvature) thus result in LAX changes in direction X. Where the L A X reflexions are dash-like in shape despite the curvature of the crystalline zones, it will mean t h a t the curvature is either slight, or t h a t there is a broad distribution of curvature and the reflexion will therefore consist of numerous two or four point overlapping images.

C

C

FIG. 8. The LAX pictures of: a--high-pressure polyethylene (ttPPE) stretched at room temperature; b--poly-vinylidene-2-fluoride stretched at room temperature and slightly tempered (in the free state at 150°C/30 rain); c--as a but subjected to 20~ elongation; d - - H P P E stretched at 85°C and cooled in the stretched state; e--HI)PE treated as in c then tempered in the free state at 85°C for 2 hr. The distribution of curvature like the curving of the crystalline zones of fibrils themselves, is normally the feature of true systems and is the consequence of irregularities of local i n t e r n a l stresses present in the amorphous zones. I t is well known for example t h a t the LAX of m a n y polymers oriented at room temperature give a four point image (Fig. 8a, b). I f such a sample is subjected to stress at the same te.mperature (or if it is Photographed after orientation without removal from the clamps) the LAX ~vill have the normal shape of a " d a s h " (Fig. 8v), The removal of the external stress will cause the unequal internal stresses existing in various parts to bend the crystallites (increase the convolution intensity of the crystalline layers), and the L A X will be a four point image. The orientation of a polyethylene sample at high temperature, followed by cooling without removal from the clamps, will show in the photograph of the free sample the LAX in the shape of a dash at room temperature (Fig. 8d). One needs only to unclamp the sample at the same high temperature and the internal forces created in the amorphous zones, which increase as a function of temperature, will bend the crystallites, and the LAX will again be a four point image (Fig. 8e). The strong intermolecular reactions taking place in the studied PA-6 samples prevent the formation of a four point L A X picture even when tempering it (200°C). The H-bonds somehow protect the crystal layers from "breaking" and

654

B. M. GI~ZBURG e$ a~.

produce more intense curvature. The planes in which the H-bonds are situated show far less convolutions or curvature of crystalline layers t h a n in direction normal to them, which is schematically illustrated in Fig. 9. The intermoleeular reactions thus can have a substantial effect on the MMS v f polymers and on its deformation changes. !

l/ /

°///

-/

)

I~IG. 9. Scheme of the macromolecular organization in an oriented PA-6 film tempered in the free state:/--sample; 2--H-bonds; 3--side-on; 4--face.on. Finally it should be noted t h a t the narrowing of the L A X reflexions in direction X can be somewhat unexpected during elastic deformations of PA-6 samples. I t appears on t he basis of earlier results [12] t h a t in crystallites the molecular axis deviates from t he t ext ur e axis and thus produces an angle with respect to t h e direction in which the stress is applied; these should show a larger curvature and the reflexion should become more diffuse in direction X. Our tests showed however t h a t t he directional change of reflexion width in direction X depends on th e pre-history (polyethylene samples), i.e. on the ratio of p r i m a r y orientat io n t e m p e r a t u r e to t h a t of deformation, on the cooling conditions after orientation, etc. The results shown in Fig. 8a, c are identical with those got b y others [8, 10], which demonstrates the fundamental possibility of a narrowing of t h e T,A~ reflexions. Translated by K. A. ALr,~.

REFERENCES 1. J. RUNDA.LL, Nature 169: 1029, 1952 2. H. ZAHN and U. WINTER, Kolloid-Z. 128: 142, 1952 3. V. S. KUKSENKO and A. I. SLUTSKER, Fiz. tverd, tela 10: 838, 1968; S.N. ZHURKOV, A. L SLUTSKER and A. A. YASTREBINSKII, Dokl. AN SSSR 153: 303, 1963 4. D. R. BERESFORD and H, BEVAN, Polymer 5: 247, 1964 5. B. M. GINZBURG, lq. SULTANOV, K. B. KURBANOV and Sh. TUICHIEV, Vysokomol. soyed. A13: 1993, 1971 (Translated in Polymer Sci. U.S.S.R. 13: 9, 2242, 1971)

Sedimentatiou, diffusion and viscosity of P B I solutions

655

6. K. ISHII~&WA, K. MXIrASAKA, M. MAYEDA and M. YAMADA, J. Polymer Sci. 7, A-2: 1259, 1969 7. T. SETO and Y. TAJIMA, Japan. J. Appl. Phys. 8: 166, 1969 8. A. COWKING, J. G. RIDER, L L. HAY and A. KELLER, J. Mater. Sci. 3: 646, 1968 9. A. COWKING and J. G. RIDER, J. Mater. Sci. 4: 1051, 1969 10. A. I~EIJ.ER and D. P. POPE, J. Mater. Sci. 6: 453, 1971 11. M. A. GEZALOV, V. S. KUKSENKO and A. I. SLUTSKER, Mekhanlka polimerov, 51, 1972 12. B. M. GINZBURG, 1~. 8ULTANOV and S. Ya. FRENKEL, Vysokomol. soyed. A13: 2692, 1971 (Not translated in Polymer Sci. U.S.S.R.) 13. V. I. GERASIMOV and D. Ya. TSVANKIN, Vysokomol. soyed. A l l : 2652, 1969 (Translated in Polymer Sei. U.S.S.R. 11: 12, 3013, 1969) 14. A. Sh. G O ~ and T. P. TANTSYURA, Vysokomol. soyed. AI0: 724, 1968 (Translated in Polymer Sci. U.S.S.R. 10: 4, 839, 1968) 15. Yu. V. BRESTKIN, B. M. GINZBURG, P. A. IL'CHENKO, K. B. KURBANOV, M. A. MARTYNOV, Sh. TUICH1EY and S. Ya. FREN~EL, Vysokomol. soyed. A15: 621, 1973 (Translated in Polymer Sei. U.S.S.R. 15: 3, 549, 1973) 16. V. 8. KUKSENKO, A. I. 8LUTSKER and A. A. YASTREBII~SKII, Fiz. tverd, tela 9: 2390, 1967; S. N. ZHURKOV, V. S. KUKSENKO and A. I. SLUTSKER, Fiz. tverd. tela 11: 296, 1969; V. S. KUKSENKO and A. 1. SLUTSKER, Fiz. tverd, tela 11: 405, 1969; V. S. KUKSENKO and A. I. SLUTSKER, Mekhanika polimerov, No. 1, 43, 1970 17. A. GU]NIF~ and G. FOURNET, Small-Angle Scattering of X-Rays, J. Wiley Inc., I~.Y.; Chapman and Hall Ltd., London, 1955 18. V. S. KUKSENKO, S. NIZAMIDINOV and A. I. SLUTSKER, Fiz. tverd, tela 9: 1966, 1967 19. A. A. YASTREBINSKII, Dissertation, 1966 20. R. HOSEMANN and 8. N. BAGCHI, Direct Analysis of Diffraction by Matter, North Holland Publ. Co., Amsterdam, 1962; R. HOSEMANN, J. Appl. Phys. 34: 25, 1963 21. Ye. L. GAL'PERII~, V. F. MXNDRUL and V. K. SM][RNOV, Vysokomol. soyed. AI2: 1949, 1970 (Translated in Polymer Sei. U.S.S.R. 12: 9, 2207, 1970) 22. R. BONART, Kolloid-Z. und Z. f'dr Polymere 194: 97, 1964

THE SEDIMENTATION, DIFFUSION AND VISCOSITY OF POLYBUTYL ISOCYANATE SOLUTIONS* V. N. TSVETKOV, I. N. SHTENNIKOVA, ~/[. G. VITOVSXArA,YE. I. RYUMTSEV, T. V. PEKKER, YU. P. GETMANCHUK, P. N. LAVRENKO a n d S. V. BUSHn~ High Polymer Institute, U.S.S.R. Academy of Sciences

(Received 26 June 1972) The progressive diffusion D, sedimentation coefficient So, intrinsic viscosity [t/] and the molecular weights M of a number of polybutyl isoeyanate (PBI) fractions were measured over a wide range of M. The experimental results were used to deter* Vysokomol. soyed. A16: No. 3, 566-574, 1974.