The orientation behaviour of high-molecular polyacrylonitrile

THE ORIENTATION BEHAVIOUR OF HIGH-MOLECULAR POLYACRYLONITRILE * S. L. DOBRETSOV, ~T. V. LOMO~OSOVA, V. P . STELMAKH al~d S. YA. FREI~'KEL' L. Ya. K a r p o v Physicochemical Research Institute (Received 9 October 1970)

I-IIGH-TEMPEILkTUREdrawing is the most commonly used method of increasing the strength of polymeric films and fibres. The strengthening effect is due to the fact t h a t thermal drawing is accompanied b y a reorganization of the polymer structure resulting in preferential orientation of the maeromoleeular chains in the direction of the stretching axis. The orientation processes and the development of fibrillar structure are influenced b y the length of the macromolecules (or the molecular weight, M) of the original polymer. A rise in M up to a certain limit, according to the authors of [1], greatly increases the orientation capacity of the polymer undergoing drawing, b u t with the ordinary methods of drawing a saturation point is reached in regard to the orientation capacity if M is increased to values of the order of 10,000. A number of authors have dealt with the problem of the effect of further increase in M on the orientation capacities of polymers, including t h a t of polyacrylonitrilc (PAN) [2, 3]. However, these authors generally investigated P A N with values of M below l0 s [4], in view of the difficulty of obtaining higher values of M for P A N synthesized b y the usual methods. I t was therefore of interest to investigate P A N with M exceeding 10 e using samples synthesized b y the method described in [5]. The present paper relates to a comparative investigation of changes in the o r i e n t a t i o n a n d p a c k i n g d e n s i t y o f t h e m a c r o m o l e c u l e s in t h e c o u r s e o f t h e r m a d r a w i n g ; t h e s t r e n g t h e n i n g effect o f t h e l a t t e r is c o m p a r e d for P A N w i t h M - - 3 x X 10 s a n d f o r m o r e l o w m o l e c u l a r P A N w i t h M----3 X 105. EXPERIMENTAL The problem of the relationship between strength and orientat ion factors is complicated b y the ever-present scale effect. However a relationship between strength and the d~aw ratio ). and the other orientation parameters investigated in thes e experiments is preserved despite changes in the size of the samples and irrespective of whether flat film samples or monofibre samples were investigated. I n view of the well-known experimental difficulties t h a t are encountered in work with fine ( ~ 10 ~) and non-circular s a m p l e s of monofilaments we had recourse to film samples for investigating the optical, a n d structural orientat3on parameters. The relationship between the draw ratio and the thermomeehanical properties and strength was investigated with fibres prepared by the wet method [6] and with films east * Vysokomol. soyed. A t 4 : No. 5, 1143-1148, 1972. 1278

Orientation behaviour of high-molecular polyacrylonitrile

1279

from 1"5-3 ~/~ solutions of the polymers in dimethylformamide (DMFA). The thermal drawing was carried out under nitrogen. The mechanical and thermomeehanical behaviour (see the isometric heating diagrams (IHM)) of the fibres and films was investigated on a UMIV-4 apparatus using the m e t h o d described in [7]. The I R spectra were recorded on a UR-10 spectrometer using a first class (DS-403-G) Japanese grid spectrometer for the reference measurements. The spectra were recorded in polarized light [8] with two polarizer positions so that the direction of the electrical vector was parallel and perpendicular to the stretching axis. The three bands selected for measuremerit of the dichroic ratio R~DI!/D ± were at 2245 cm -1 (valency vibrations of the C - N group) and 2870 and 2940 cm -1 (symmetrical and asymmetrical valency vibrations of the CH2 group) (a-diehroism). The assignment of the bands and the mode of polarization were taken from paper [9]. The orientation indices were the values of cos ~ O-~(2--R)/(2-FR) (0 is the angle of disorientation of the axes of the molecular chains relative to the stretching axis) determined from the diehroic ratio of the bands at 2870 and 2940 cm -1, in view of the insensitivity of these bands to structural transitions in the molecular chains and the fact (experimentally verified) t h a t the optical density of the bands in question is unaffected by changes in the drawing conditions. The dichroism of the band at 2245 em -~ was investigated because it is only for this band t h a t quantitative measurements are given in [10] in connection with the orientation of PAN. The X- r ay measurements were carried out on a URS-55A apparatus (X-ray diagrams on photofilm) and URS-50IM apparatus with a special attachment for recording the azimuthal intensity distribution of reflections [l 1]. The degree of orientation was estimated from these measurements based on the half-width of the azimuthal intensity distribution A ~ of the reflection observed at an angle of 2 0 ~ 17°12 '.

DISCUSSION OF RESULTS P A N is c l a s s e d w i t h p o l y m e r s o f p a r a c r y s t a l l i n e s t r u c t u r e [12]. T h e c h a r a c t e r i s t i c f e a t u r e o f t h i s s t r u c t u r e is t h e h e x a g o n a l p a c k i n g o f t h e m o l e c u l a r c h a i n s i n t h e p l a n e p e r p e n d i c u l a r t o t h e s t r e t c h i n g axis; a t t h e s a m e t i m e t h e r e is a n a b s e n c e o f o r d e r a l o n g t h e a x e s o f t h e c h a i n s (a t w o - d i m e n s i o n a l c r y s t a l ) [13]. T h e el ee-

~, kg/mm2 a 3o

I00 80

/0 J

2

6

10

/4

!

i

~

5

J

iO

FIG. 1. Plots of ab vs. ~ for P A N fibres with M = 3 × 106 at 150 ° (a) and for PA N films with M = 3 × 108 at 150 ° (1) and with M = 3 × l0 s at 118 ° (2) (b).

1280

S . L . DOBI~ETSOV et al.

iron microscope reveals fibrils ~ 100 A in diameter in oriented PAN, as well as the absence of any alternation in density along the axes of the fibrils [14]. However, even with this type of structure the strength of the oriented polymer is determined b y the orientation of the macromolecules and b y the number of loadbearing chains [15], if we disregard losses of strength due to macrodefects which m a y be present in the actual fibre [16]. TABLE

1. M E C H A N I C A L

AND THERMOMECHANICAL FIBRES

M × 10 -s

~

Td, °C

2 10 14

150 150 150

1"25

ll

118 118 118 118

1 "25 2"5 6 10

150 150 150 150

PROPERTIES

OF THERMALLY

DRAWN

PAN

AND FILMS

oh, k g / m m 2

E, kg/mm 2

(Tlma,X

eb, %

kg/mm2

220 1072 3106

10"5 8"3 7"4

0-5 2.4 3.3

4.4 14.5 21.0 24.8

221 640 701 854

6.4 9.0 6"0 6"1

0.5 1.3 3.4 3"8

4.7 12.2 25.8 35.0

246 526 886 1280

14"5 11"4 6"9 5"0

0.3 0-7 2.1 8.5

Fibres 19 81 119 Films 0.3

5 9

~Vot¢. The values of the tensile strength (ab), the breaking elongation (sb) and the elastic modulus (E) obtained from the tensile plots were based on the average of three to five measurements.

It is seen from Fig. 1 and Table 1 that with increasing 2 the strength and the elastic modulus of the PAN fibres and films axe increased for both values of M, b u t in the case of the films the strength increase is more marked for the higher M, which agrees with the results of an investigation [3] of the effect of M (up to M values of the order of 5 × 10~) on orientation processes during thermal drawing. The higher strength values found for the fibre samples (120 kg/mm 2) compared with the films are apparently due to two factors, the first being the difference in the original state of the polymer in the films and fibres owing to differences in the methods of sample preparation, and the second being the scale factor which influences the extent to which the strength of the polymer is increased b y thermal drawing. For instance, the results of a special experiment showed that the strength values for samples in the form of strips of the same P A N film, differing in width {strips 10 and 2 mm wide) and drawn at 110 ° to ~----7, were 16 and 41 kg/mm ~ respectively. Consequently comparisons of the mechanical properties of the P A N films of different M were in every case related to samples of approximately the same dimensions whether in the original or in the drawn state.

Orientation behaviour of high-molecular polyacrylonitrile

1281

The strength increase with increasing 2 is usually attributed [17] to a higher degree of orientation of the maeromolecules. This conclusion is certainly supported b y the results of the thermomechanical investigations presented in Fig. 2 and Table 1. The I H D for the P A N films with M = 3 × 106 and 3 × 10 ~ with different degrees of thermal drawing show the marked changes brought about in the polymer structure b y drawing, and in the case of the polymer of higher M the value of the first maximum (~1 max) on which estimates of the overall degree of orientation are usually based [2, 3, 18, 19] is considerably higher. G*~kg/mm 2

5 #;

7

5

j 200

~oc

/100

f 200

r,*c

~00

Fro. 2. Isothermal heating diagrams ( I H D ) of P A N films with M = 3 × 10 6 (a) and 3 × 10" (b) at 2=1.25 (1), 2.5 (2), 5 (3), 6 (4), 9 (5), 10 (6) and 11 (7).

However, the results of direct investigations of the orientation using the I R dichroism and the X-ray method of analysis were as follows. As is seen from Fig. 3, showing plots of cos 2 0 vs. 2 for the three absorption bands, the orientation factor increases rapidly with low degrees of drawing (up to )~= 3-4), b u t with further increase in 2, cos ~ 0 remains practically constant, although the internal structural changes continue to accumulate within the range of extensions in question. This conclusion is borne out b y the splitting of the band at 2940 cm -1 for the samples with a high degree of drawing (Fig. 4) and b y the increased intensity of the band at 2245 cm -1. We would point out that there is no splitting of the bands for the more low-molecular PAN. The results of I R analysis therefore reveal peculiarities of the orientation behaviour of P A N resulting from the existence of two processes. The first is the orientation of the macromolecules in the direction of drawing, a rapid process which is already completed at ~-----3-4. We believe the second process is one of refinement of the intermolecular packing of the P A N macro-

1282

S.L. DOBRETSOV~ at.' TA~L~ 2.

C H A N G E S IN" '£~:Lm O R I E N T A T I O N

OF P A N

FILM8 DUE

TO

~.'~LWRMAT, DRAWI~TG $

I reflection 20= 17012' hag-width

ll~'X I0 -6

radial

azimuthal

B

0.3

29000 ,

lO15 •

13030 •

i

1°07 ,

10012 ,

1°00 '

8024 ,

1°09 ,

14042 ,

1°06 ,

10048 ,

l°00 ,

10o33 ,

0054 •

10024,

2+R

0.356 0.548 0-585 0.593 0.623 0.523 0.580 0.557 0.590

100 °

I°12 '

2--R

(2940 cm -1)

Jq

1°24 ,

2.5 4.0 6.0 8.5 6.0 8-5 13.0 19.5

OOSa0--_~

* The drawing temperature, ~Td=135°.

molecules. A c o m p a r a t i v e s t u d y o f t h e changes in t h e o r i e n t a t i o n i n d e x for t h e b a n d s a t 2870 a n d 2940 c m -1 for t h e p o l y m e r s w i t h t h e t w o m o l e c u l a r weights s h o w e d t h a t t h e degree of o r i e n t a t i o n o f t h e m a c r o m o l e c u l e is higher for t h e p o l y m e r w i t h higher M (see Fig. 5). T h e X - r a y s t r u c t u r a l investigations s i m i l a r l y s h o w e d t h a t t h e t h e r m a l d r a w i n g of P A N b o t h increases t h e o r i e n t a t i o n o f t h e o r d e r e d regions of t h e p o l y m e r , 0.9

cosZe

o

oo~_

0 O

vu

0.7 a 0.5

O'~

J = i i I I h i

I

5

0-7 c ~ e

i

I

[

I

1

I

I

I

I

5

fO

I

i

fO

I

I

~

L

15

i

i

i

z

2O

k

c

d

0-5 ~

O~

5

I0 I

~ J d •s

5

=

• u

lO

• -•

W

/5

20

FzG. 3. P10 ts of cos = 0 vs. k for PAN films with M-----3 x 106 (a, c) and 3 X 10~ (b, d) at v = 2245 • (a, b); 2870 (1) and 2940 cm -1 (2) (c, d)i '

1283

Orientation behaviour of high-molecular polyaerylonitrile

based on t h e r e d u c e d h a l f - w i d t h of the a z i m u t h a l intensity distribution of t h e reflections (Fig. 6), a n d refines t h e s t r u c t u r e of the sample as a whole. The l a t t e r m a y be d e d u c e d f r o m the n a r r o w i n g of the reflection a t 2 8 = 17012 ' in t h e r a d i a l direction. I n t h e light of these results it a p p e a r s t h a t there is a difference in t h e o r i e n t a t i o n b e h a v i o u r of t h e P A N of different molecular weights: in t h e case of t h e high-molecular p o l y m e r zt ~ a n d the h a l f - w i d t h B are r e d u c e d with increase in 2 to a greater e x t e n t t h a n occurs with the low-molecular p o l y m e r (see Table 2).

cosZO 0~6F"

f%.

",,

I

i

1,

2 ,

O'#~/~ /Z/~,, ,

,l I

i

....m ----'1u.. I I L I- Ill| 5

15

Fro. 5 I00

"4~°

I:1

5 I

32

U

I

]

28

2A

P~lO-z,cm -1 FiG. 4

, I

I0 A

I

o

z , A ~ . J

,

1

I

I

:

i

5

I,,

J,

I

;,

152; FIG. 6

Fio. 4. I R spectra of PAN films with M = 3 x l0 ° (1) and 3× 105 (2) and )~=9 at 118 °. FIC. 5. Plots of cos 2 0 vs. Afor PAN films with M = 3 × l0 s (I, 3) and 3 × 105 (2, 4) at v = 294(~ (1, 2) and 2870 cm -1 (3, 4) at 150°. FIo. 6. Degree of orientation vs. A for PAN films with ~]I= 3 × 106 (1) and 3 × 105 (2) at 118 and 135° respectively.

The results of the t h e r m o m e c h a n i c a l investigations of the b e h a v i o u r of t h e oriented P A N provide q u a l i t a t i v e confirmation of our a s s u m p t i o n t h a t i m p r o v e d p a c k i n g of the macromolecules occurs with high d r a w ratios. W i t h increasing 2 - the I H D shows a m a r k e d rise in the second m a x i m u m related to intermolecular erosslinking [20] which is facilitated b y the more ordered a r r a n g e m e n t of t h e maeromolecules. Two a l t e r n a t i v e s a p p e a r possible as regards the m e c h a n i s m of the side o r d e r refinement: either a n increase in the n u m b e r of sequences of highly uncoiled chains (reduction in the n u m b e r of defects such as folds a n d loops, the K u v s h i n s k i i -

1284

s. L. DOBRETSOVe$ a$.

Laius mechanism [21]), or transition f r o m u n c o r r e l a t e d to correlated distortions o f t h e paracrystalline lattice [22]. At present it is difficult to say which of t h e two m e c h a n i s m s is actually involved. I n the light of t h e above considerations t h e increased s t r e n g t h observed as a result of t h e r m a l drawing with high d r a w ratios, w h e n t h e r e is no longer a n y increase in t h e orientation, m a y be a t t r i b u t e d to an increase in the n u m b e r of loadbearing chains, a n d t o side order refinement. F u r t h e r investigations will be n e e d e d t o verii~y this e x p l a n a t i o n o f the problem. CONCLUSIONS

(1) A c o m p a r a t i v e s t u d y has been m a d e of t h e orientation b e h a v i o u r o f p o l y acrylonitrile (PAN) of two molecular weights, and it is shown t h a t higher degrees o f o r i e n t a t i o n are obtainable for the p o l y m e r o f higher molecular weight. (2) The results of I R and X - r a y s t r u c t u r a l investigations reveal two areas of change in the o r i e n t a t i o n b e h a v i o u r of P A N in the course o f t h e r m a l drawing. A t low d r a w ratios (4) t h e orientation of the macromolecules is in the direction o f drawing, b u t f u r t h e r increase in ), leads to side order refinement. Two possible mechanisms o f side order r e f i n e m e n t are discussed. (3) T h e s t r e n g t h of t h e p o l y m e r m a y be increased (up to 120 k g / m m ~) b y the o r i e n t a t i o n a l drawing of high-molecular P A N . Translated by R. J. A. I-IENDttY

REFERENCES

1. D. V. S. HURL and R. H. PETERS (Eds.), The Structure of Fibres, 1969 2. L. A. LAIUS and Ye. V. KUVSHINSKII, Mekhanika polimerov, 579, 1967; Vysokomol. soyed. 3: 215, 1961 (Not translated in Polymer Sei. U.S.S.R.); S. KAMALOV, Thesis, 1967 3. S. KAMALOV, A. A. KOROTKOV, V. N. KRASULINA and S. Y&. FRENKEL', Khim. volokna, No. 6, 9, 1966 4. M. P. ZVEREV, A. N. BARASH and K. A. KOSTROVA, Khim. volokna, No. 6, 48, 1968; A. B. PAKSHVER and V. E. GELLER, Khimiya i tekhnologiya volokna nitron, 1960 5. V. I. LUKHOVITSKII, V. V. POLIKARPOV, A. M. LEBEDEVA, R. M. LAGUCHEVA and V. L. KARPOV, Vysokomol. soyed. Bg: 252, 1967 (Not translated in Polymer Sei. U.S.S.R.); AI0: 835, 1968 (Translated in Polymer Sci. U.S.S.R. 10: 4, 969, 1968) 6. J. R. KNUDSEN, Text. Ros. J. 33: 13, 1963 7. S. L. DOBRETSOV, A. I. KURILENKO and V. A. TEM:NIKOVSKII, Mekhardka polimerov, 944, 1966 8. I. I. VETTEGREN, Thesis, 1970 9. R. HOUWINK and A. STAVERMAN, Polymer Chemistry and Technology, Vol. 1, 525, 1965 10. I. B. KLIMENKO mad L. V. SMIRNOV, Vysokomol. soyed. 5: 1520, 1963 (Translated in Polymer Sci. U.S.S.R. 5: 4, 622, 1963); R. ZBINDEN, Infrared Spectroscopy of Highpolymers, 269, 1966 11. B. M. GINZBURG, Thesis, 1966 12. A. N. KITAIGORODSKII, Dokl. AN SSS~ 124: 861, 1959

Some problems of brittle and non-brittle break of polymeric matemals

1285

13. W. O. STATTON, Azm. N.Y. Acad. Sci. 83: 27, 1959; C. R. BOHN, J. R. SCHAEFGEN and W. O. STATTON, J. Polymer Sci. 55: 531, 1961; B. K. WEINSTEIN, X - r a y Diffraction b y Chain Molecules, 304, 1963 14. R. G. SCOTT and A. W. FERGUSON, Text. Res. J. 26: 284, 1956 15. S. N. ZHURKOV, Vestnik AN SSSR, No. 3, 46, 1968 16. M. S. MEZHIROVA and E. A. PAKSHVER, Khimieh. volokna, Iqo. 3, 13, 1968 17. I. I. NOVAK and V. I. VETTEGREN, Vysokomol. soyed. 6: 706, 1964; S. N. ZHURKOV, I. I. NOVAK, B. Ya. LEVIN, A. V. SAVITSKII a n d V. I. BETTEGREN, Vysokomol. soyed. 7: 1203, 1965 (Translated in Polymer Sei. U.S.S.R. 7: 7, 1331,. 1965) 18. L. A. LAIUS and Ye. V. KUVSHINSKII, Vysokomol. soyed. 6: 52, 1964 (Translated in Polymer Sci. U.S.S.R. 6: 1, 60, 1964); G. N. AFANASEVA, IVL I. BESSONOV, L. A. VOLF, A. I. ME()S and S. Ya. FRENKEL', Zh. prikl, kb_imii 37: 1349, 1964 19. S. L. DOBRETSOV, N. V. LOMONOSOVA and V. P. STELMflKH~ Vysokomol. soyed. B l l : 782, 1969 (Not translated in Polymer Sei. U.S.S.R.) 20. L. S. GERASIMOVA, R. A. PALATOVSKAYA, A. B. PAKSHVER and V. A. PANTAYEV, Mekhanika polimerov, 943, 1968; N. V. KOSHELEVA, I. S. OKHRIMENKO and A. D. YAKOVLEV, Vysokomol. soyed. B9: 257, 1967 (Not translated in Polymer Sci. U.S.S.R.) 21. L. A. LAIUS and Ye. V. KUVSHINSKII, Fizika tverdogo tela 5: 3113, 1963 22. R. HOSEMANN, J. Appl. Phys. 34: 25, 1963

SOME PROBLEMS OF BRITTLE AND NON-BRITTLE BREAK OF POLYMERIC MATERIALS* G. L. SLONIMSKII, J~.. A. _A_SKADSKIIa n d V. V. KAZANTSEV• Heteroorganic Compounds Institute, U.S.R.R. Academy of Sciences (Received 11 October 1970) I:N THIS i n v e s t i g a t i o n o u r a i m Was t o s t u d y p r o c e s s e s o f b r i t t l e a n d n o n - b r i t t l e break occurring with solid polymeric materials. Brittle break. According to Alexandrov and Lazurkin [1, 2] polymers become brittle whenever a point is reached, during the lowering of the temperature down to the brittle point Tb, where the forced high-elasticity limit ah.e. exceeds the m a x i m u m brittle strength, a,~ax. I n this case when T ~ T b the material will break rather than soften, as the deterioration in durability proceeds more rapidly than the onset of forced high-elasticity. Let us now consider Tb in relation to the degree of crystallinity of the polymer, a. I n papers [3-5] it was shown t h a t the amorphous regions in crystalline polymers are responsible for their strength (or to be more precise, their "weakness"). A t temperatures above the glass transition temperature (Tg) the polymer is in the high-elastic state. Therefore when T > T g a change in temperature will not lead to any marked difference in the strength of completely amorphous or partially crystalline polymers with identical chemical structures. * Vysokomol. soyed. AI4: No. 5, 1149-1155, 1972.