The mechanical properties of polyethylene terephthalate glass after deformation in adsorbing media

Mechanical properties of polyethylene terephtlmlate glass

2831

REFERENCES 1. Spravoehnik po plasticheskim massam (Plastics Handbook). Izd. " K h i m i y a " , vol. 2, 155, 1969 2. D. SAUNDERS and K. FRISH, K h i m i y a poliuretanov (The Chemistry of Polyurethanes). Izd. " K h i m i y a " , 1968 3. G. A. BULATOV, Penopoliuretany i ikh primenenie na letatelnykh apparatakh. (Polyurethane F o a m s and Their Uses in Aircraft,). "Mashinostroenie", 1970 4. G. A. KUZNETSOV, M. Ye. BLINNIKOV, N. I. NIKIFIROV and V. A. VASIL'EV, Plast. Massy, No. 5, 70, 1970 5. G. A. KUZNETSOV, V. D. GERASIMOV and L. B. SOKOLOV, Plast. Massy, No. 4, 64, 1972 6. I. I. PEREPECHKO, Akustieheskie m e t o d y issledovaniya polimerov (Acoustic S t u d y Methods of Polymers). Izd. " K h i m i y a " , 1973 7. Sb. metodov fiziko-mekhanicheskikh ispytanii penomaterialov (Collection of Physico. Mechanical Testing Methods of Foams). Vladimir, 1967 8. S. H. DENG and L. {3. STUART, Polymer Egng. Sei. 11: 369, 1971 9. J. FERGUSSON, D. J~ HOURSTON, R. MEREDITH and D. PARSOVONDIS, Europ. Polymer J. 8: 369, 1972 10. Yu. Yu. KERCHA, Yu. S. LIPATOV, A. P. GERK0V and R. L. SHAPOVAL, Vysokomol. soyed. A14: 1187, 1972 (Translated in Polymer Sci. U.S.S.R. 14: 5, 1330, 1972) 11. N. I. KORZHUK, V. F. BABICH and Yu. S. LIPATOV, Vysokomol. soyed. B15: 323, 1973 (Not translated in Polymer Sei. U.S.S,R.) 12. G. A. KUZNETSOV, V. D. GERASIMOV and L. B. SOKOLOV, Vysokomol. soyed. 6: I261, 1964 (Translated in Polymer Sei. U.S.S.R. 6: 7, 1391, 1964) 13. D. L. FOLDS, J. Acoust. Soc. America 52: 426, 1972

THE MECHANICAL PROPERTIES OF POLYETHYI,ENE TEREPHTHALATE GLASS AFTER DEFORMATION IN ADSORBING MEDIA* A. L. VOL~SKIL V. I. G~ASIMOV and N. F. BXKEY~V M. V. Lomonosov State University, Moscow (Received 3 November 1974) The mechanical behaviour of amorphous P E T P has been investigated after deformation in n-propanol. Such deformation has resulted in the formation of specific micro-fissures mad the fissured P E T P has a new complex of physieo-mechanical properties when tested. The stretching tests lead to the assumption t h a t the miero-fissurer formation causes a polymer transition to a new structural state. * Vysokomol. soyed. A17; No. 11, 2461-2468, 1975.

2832

A. L. VOLYNSKII 65 0~.

STUDr o f the s t r e t c h i n g d e f o r m a t i o n o f a m o r p h o u s polymers in an adsorbing m e d i u m disclosed a n u m b e r o f interesting features [1, 2]. Such a d e f o r m a t i o n a t t e m p e r a t u r e s below t h e glass t e m p e r a t u r e (Tg) is k n o w n to give rise t o specific micro-fissuring [3] a n d t h e edges of t h e microfissures interlink the 100-200 A long p o l y m e r macromolecules o r i e n t a t e d in the direction of s t r e t c h i n g [4]. This t y p e of p o l y m e r i c material w i t h a super-developed surface will be unstable a f t e r the r e m o v a l o f t h e a d s o r b e n t and will result in its rapid coagulation. A m e c h a n i s m o f f o r m a t i o n o f the d e t e c t e d larger, reversible d e f o r m a t i o n s which t a k e place in p o l y m e r glasses has been suggested on the basis o f chiefly microscopic observations [2]. I t was n a t u r a l to assume t h a t the d e f o r m a t i o n changes t a k i n g place in a n a d s o r b a n t m u s t affect the mechanical properties of the polymer, which c a n n o t recover its original s t r u c t u r e a f t e r t h e s u r f a c t a n t is removed. I t is a c t u a l l y necessary to he~t the p o l y m e r to a t e m p e r a t u r e ~bove t h e Tg in order to r e m o v e t h e m a c r o m o l e c u l a r o r i e n t a t i o n existing inside t h e fibrils. W e describe here t h e results of the structural-mechanical investigation of industrial films based on a m o r p h o u s P E T P a f t e r t h e i r d e f o r m a t i o n in n-propanol. EXPERIMENTAL

Industrial film samples of amorphous PETP 30-350 ~ n thick were used in this investigation. The mechanical properties were tested on paddle-shaped 5 × 20 mm working surface specimen samples which were stretched in n-propanol to variour % elongations at a 45ram/ /rain rate of elongation. The samples were then relased from the clamps and dried. The efficiency of removal of propanol was checked by weighing the sample before deformation and after drying; the weight was found to remain the same. The consideration of the specific weight of the samples and of the analytical balance sensitivity led to the assumption that the amount of n-propanol remaining in the sample did not exceed 0.1%. Some of the tests were made on samples dried after stretching while still clamped because of the findings that shrh~kage takes place during drying [1]. All the tests were made at room temperature, using a Polyani-type dynamometer at 0-5-90 mm/min elongation rates. The small angle X-ray pictures were recorded on the apparatus described by Buts]or and co-workers [5], and the diffraction picture was obtained by means of an electron-optical transducer. The X-ray scattering pictures were examined visually after an about 10 sec exposure on photo-film. The distance of the sample from the transducer was 250 ram, the radiation source was CuK,. A URS-50 instrument was used to produce the wide angle X-ray diagrams under standard conditions. The MIN-8 polarization microscope was used for the microscopic observation. Shrinkage tests were made on samples stretched to various ~o elongations, afterwards releasing them from the clamps and checking the length of the working surfaces over different periods of time. The °//oshrinkage was found after the sample length remained unchanged. RESULTS

The elongation curves for P E T P samples which h a d been first s t r e t c h e d in n - p r o p a n o l to various elongations a n d t h e n dried in the free state are r e p r o d u c e d in Fig. 1. T h e curve for the original P E T P (untreated) is given for comparison. T h e fact t h a t the polymer, having n o w a v e r y large n u m b e r o f cracks, is still

Mechanical properties of polyethylene terephthalato glass

2833

~trong enough to w i t h s ~ n d large deformations, stands out immediately, which ~ance more underlines the principal difference between the micro-fissures created in polymer glasses and the true cracks existing in solids. One can see in Fig. 1 the difference in shape of the elongation curve for fissured samples from t h a t o f the original P E T P . The first part of the elongation curve is sigmoid, which is typical for rubber-like substances. Although the elasticity modulus of such samples drops slightly in comparison with that of the original P E T P , its absolute values remains within the limits typical for polymer glasses. Interesting is Chat the start of neck formation depends on the preliminary elongation in the ~ l s o r b e n t . The elongation at which the neck starts to form is indicated b y the lines of dashes in Fig. 1. The start of neck formation coincides in each ease with the curve going over into the linear part (except in the case of the original P E T P w h e r e this moment coincides with the peak of stress), which agrees well as one ,can see with the % preliminary elongation in the medium. The sample seems to have a " m e m o r y " of how much it had been stretched in n-propanol. A macro,scopic neck forms after the curve reached the linear part.

a', i~glmrnZ

5

I

I

1

!

I

I

3N¢' If/~"

t i

I I

I i

~|i ~I

I i

t I

i t

I

II 50

11

I IOO

5

1

II sq, 150

I 250 c~%

~IG. 1. The stretching deformation curves for PETP samples after deformation in n-propanol ~ad drying in the free state. ~/o Preliminary stretching: 1 -- 25; 2-- 50; 3-- 85; 4-- 100; 5-- 150; fi--original PETP. 10 mm/min rate of stretching. The extent of deformation of the fissured samples, before the curve becomes linear, will obviously differ from that on the linear part. P E T P samples deforme d in an adsorbent medium and dried in the free state are subject to consider~ble shrinkage during deformation in air [2]. The reversible deformation as a function of ~/o elongation shown in Fig. 2 is that for a sample which had been ,stretched 25~/o in n-propanol and then dried in the free state. It is quite clear that considerable reversible deformation exists up to the point at which the neck formation limit is reached, i.e. when the preliminary °/o stretching in the medium is not exceeded. Shrinkage drops markedly after that and the reversible

~834

A . L . VO:LYI~SKII e~ al,

deformation will not be larger (at larger ~ deformation) than one normally encountered when a polymer glass is stretched to neck formation. The relaxation of deformation is thus associated with that part of the sample which had not changed into the neck during the particular ~o elongation. Please note also that shrinkage was very gradual (the sample dimensions taking hours to change), which is evidence for the relaxing nature of the phenomenon. The relaxation processes can be easily detected in the study of the mechanical characteristics of the polymer as a function of time. The maintenance for various times of stretched samples with fixed dimensions will clearly show that shrinkage drops noticeably (Fig. 3). In other words, there is some relaxation taking place of the amorphous polymer below its Tg. The absolute value of deformation is much larger for such a process than that normally found on glass-like polymers. e°/o 50

30

iOo--

I

I Z00

I

I

200

FIG. 2, The reversible deformation c as a function of the ~o elongation for PETP samples stretched in n-propanol, k~25~o, when dried in the free state. The shrinkage phenomena found after removal of the adsorbent had been earlier explained by a peculiar coagulation of the polymer fibrils; but the reverse process is also possible; i.e. a peculiar peptization of the highly disperse polymeric material. The soaking in propanol (wetting) of the samples used in the above experiments (i.e. those showing an about 50~o shrinkage, after which their dimensions remain constant), followed by the removal of the solvent, will produce an almost complete return to the original dimensions. As to the samples dried after elongation in the solvent, having fixed dimensions, these showed noticeable shrinkage during drying after prolonged immersion in the n-propanol (for 1 week). This shrinkage was as large as 50%. The incomplete reversibility of the deformation could be due to development of quite considerable stresses (larger than 3 kgf/mm ~) when samples of fixed dimensions are dried. In many cases, e. g. in polystyrene and polymethyl methacrylate, these stresses cause a macroscopic fracture of the sample during drying. Some local fracture occurred also in P E T P of the fibrillar material inside the separate micro-fissures, or local cold flow

Mechanical properties of polyethylene terephthalate glass

2835

(neck formation), as found under the electron microscope. This will naturally reduce the a m o u n t of reversible deformation. The above raises the question of the mechanism of the observed phenomena, because the described physico-mechanical properties of the P E T P deformed in g~°/o ~g

20

0"5

l.g

2q

Time, hp

FIG. 3. Value e for fissured samples as a function of storage time in the stretched state (~t~ 25~o). The % preliminary elongation in n-propanol was 75%. n-propanol greatly differ from those of the usual polymer glasses. An a t t e m p t to answer this question was made by X-ray structural and microscopic studies of the deformations of fissured samples. The elongation characteristics of a

FIG. 4. The deformation of a PETP sample stretched 50% in n-propanol after drying in the free state: °/odeformation: a-- 15, b --40, c-- 75.

A. L. VOL~SKr~ et al.

2836

P E T P sample which had been deformed in n-propanol and then dried in t h e free state is shown in Fig. 4 up to the moment of neck formation. One can see the micro-fissures becoming much wider with the ~ elongation deformation, although there is no substantial change of the unfissured part of the polymer. One can also see in Fig. 4 that there is no disturbance of the planarity of t h e sample. A small-angle X-ray (SAX) picture of a P E T P sample stretched 125°/o in n-propanol and then dried in the free state is shown in Fig. 5a. The X - r a y

~i~i~ ¸

gZ

:FIG. 5. X-ray pictures of a PETP samples after 125% deformation in n-propano] and drying in the free state: a, b--SAX; c, d--wide angle X-ray. a, c--prior to deformation; b, d--after 125~o deformation in air. diffraction picture indicates the presence of microheterogeneous asymmetrical shapes with an orientation perpendicular to the stretching direction in the sample. This could be connected with the micro-fissures present in the polymer. The more such a sample is stretched~ the clearer will be the equatorial reflexion, which will have the shape shown in Fig. 5b for that part of the sample which ha~l not changed to a neck during stretching. The appearance of this reflexion seems to be due to the uncoiling and orientation of the 150-200 • wide fibrils which coagulate during the solvent (adsorbent) removal. Wide angle X-ray (LAX) pictures were also produced from these samples (Fig. 5c, d). Attention is drawn to the fact that considerable orientation can be found of the macromolecules in the polymer glass before neck formation (Fig. 5d). This result is rather unusual because of the knowledge that orientation will take place only after the change of the material into a neck when an amorphous polymer is stretched at a temperature below the Tg. However, this is easily explained on the basis of the theories [2] developed on the mechanism of deformation. The sample which is deformed in a medium and then dried in the free state will contain chaotically disorientated fibrils so that there is no macromolecular orientation present (Fig. 5c). Its stretching will cause the uncoiling and mutual orientation of separate fibrils. The maeromolecular orientation inside the fibrils is produced during polymer stretching in the adsorbent and will remain after coagulation, so that the fibrillar orientation leads to one of the macromolecules (Fig. 5d). The above is also indicated b y the mechanical behaviour of the fissured P E T P samples dried at fixed dimensions. The strength-deformation curve (Fig. 6) will not be sigmoid in this case because the conditions of sample preparation

Mechanical properties of polyethylene terephthalate glass

283~

were such t h a t the fibrils were unable to change their shape to any extent [2]. T he neck formation will start in such samples at small deformations and will be independent of the % prior elongation in the medium.


2

a I

kg/mmZ

L

I

8

7

b

2 5

t

3 ! 2

1 I 25

1 ,58 Fie. 6

I 7,5 E,%

0 50

100

8~ %

Fro. 7

Fro. 6. The elongation curves of samples deformed 85% in n-propanol, and/--dried in the free state, 2--having fixed dimensions. Elongation rate 10 mm/mln. FIG. 7. The elongation curves of PETP samples stretched in n-propanol by (%): a-- 25; b-- 100; c--150 at rates of (ram/rain): 1--10; 2--30; 3--90. These experimental results permit the conclusion t h a t an uncoiling and mutual orientation takes place of the macromolecular chains during the stretching of fissured P E T P samples when t h e y had been deformed in n-propanol first. All the chains will be naturally orientated when the samples reach the preliminary stretching percentage; this will result in an orientation~l strengthening of the micro-fissured material (as evident in some modulus increase shown on curves 1-5 of Fig. 1 before the start of neck formation), and will be followed by a transition of the material into the macroscopic neck. The mechanical energy seems to be utilized in such a deformation for overcoming the forces of surface tension created by the disengagement of the polymer fibrils. The found relaxation phenomena are due to the mutual rearrangements of the fibrils during sample deformation which makes possible a reduction of the free surface created by fibrillar disengagements under the influence of a mechanical stress. We investigated the influence of the deformation rate on the mechanical properties of P E T P containing various amounts of fibrillar material. Figure 7 shows the results of elongating P E T P samples first stretched in n-propanol t o various ~o elongations, and then dried in the free state. I t shows t h a t a small

~838

A. L. VOLYNSKII et a~.

content of fibrillar material and large rates of deformation make possible the realization of a superstress peak on the deformation-strength curves. It becomes clear that the mechanical properties of the material strongly depend on the deformation rate within the studied range. As with the majority of polymer glasses, the amplitude of the superstress peak noticeably increases together with

FIO. 8. Photo-micrographs of P E T P samples stretched 20 % in n-propanol at rates of (ram/rain): a - - 45; b -- 3. The photographs were produced in crossed polaroids.

the initial modulus. A larger initial % elongation in n-propanol alters the nature of the rate dependence of the mechanical properties~ The super-stress peak will not be obtainable over the whole range of deformation rates and there will be a o',kglmrn z L_

I

25

Ir

50

i

~%

FIG. 9. The elongation curves of PETP samples deformed first in n-propanol at rates of (ram/rain): 1--0.5, 2--80 and then dried in the free state. Final deformation rate = 10 mm/min. marked reduction in the effect the deformation rate has on the mechanical properties. Finally, a 150% preliminary deformation of the sample in the medium will give rise to identical stretching curves. The insensitivity of the material to the deformation rate in the investigated range indicates the quite large rates of the relaxation processes which accompany the deformation; these are uncharacteristic of polymer glasses. An interesting fact arising out of Fig. 7 is that the shape of the curves differs quite considerably at the start for samples deformed

Mechanical properties of polyethylene terephthalate glass

2889

first in n-propanol to various % elongations, although uncoiling and mutual orientation of the fibrils takes place before neck formation. The mechanical properties of the fissured P E T P samples must therefore depend on the length of

FIQ. 10. The external appearance of PETP samples stretched in air: a--original PETP b--sample deformed first in n-propauol and dried. the microchains as well as on the total amount of fibrillar material present, because the former increases with the width increase of the microfissures, i.e. with the °/o elongation in the medium. The latter finding was verified in P E T P

Fie. 11. Photomicrographs of PETP samples stretched in air up to neck formation: a--original PETP; b--sample first stretched in n-propanol and then dried; c--SAX of the neck par~ of PETP stretched 100% in n-propanol. studies after preliminary deformation to identicM length in n-propanol but at various deformation rates. The mechanical properties of such samples, having short fibrils, differed greatly from those having long fibrils (Fig. 9). Although neck formation took place after the same period in time (indicated by the line of dashes in Fig. 9), as could be expected, the curve for the sample with short

2840

A. L. VOLYNSKII ~ a~.

fibrils has a super-stress peak at a larger stress corresponding with the start of uncoiling of fibrils and also a steeper ascent, due to orientation before neck formation (Fig. 8a). The shape of this curve is similar to that of the curves in Fig. 7 which corresponds with small % elongations in n-propanol. These findings permit the conclusion that while the % preliminary elongation in medium primarily determines the treshold of neck formation the rate of preliminary deformation strongly affects the mechanical behaviour of the produced material.

FIG. 12. Electron-scanning photo-micrograph of p a r t of the neck in a P E T P sample deformed first 100% in n-propanol.

Another important characteristic of fissured samples is their neck formation capacity and the presence in the neck of a specific structure. A photograph of P E T P samples stretched in air without prior deformation in n-propanol (a) and after such deformation (b) is shown in Fig. 10. One can see that the fissured sample is capable of producing a distinct neck, although this part, as well as the rest of the affected part remains opaque, which is in contrast with what happens in the case of an unfissured sample. There is also no distinct border line to the neck in the former case (Fig. llb) and the surface is covered with dark lines which are at right angles to the stretching direction. These lines are traces of the micro-fissures which existed in the sample before stretching and are the cause of the neck being opaque, as shown in Fig. 10. According to the SAX scattering diagram such a neck contains a fairly large amount of micro-heterogeneities with a cross section of the order of 100-300 A which are pointing in the stretching direction (Fig. 1lc). The recording method of the SAX beam scattering does not Mlow the observation of the scatter by the neck produced in the original, unfissured PETP. The electron-microscopic study of the neck part produced on stretching the fissured P E T P sample showed the existence of some positions, which had contained microfissures before stretching (Fig. 12), with a fairly large amount of fibrillar material orientated in stretching direction; the latter is probably responsible for the detected SAX scattering. Tranala~d by K. A. ALLE~

Mechanism of polyethylene terephthalate transition

2841

REFERENCES 1. A. L. VOLYNSKII, T. I. KHETSURIANI and N. F. BAKEYEV, Vysokomol. soyed. B16: 564, 1974 (Iqot translated in Polymer Sei. U.S.S.R.) 2. A. L. VOLYNSKII and N. F. BAKEYEV, Vysokomol. soyed. A17: 1610, 1975 (Translated in Polymer Sei. U.S.S.R. 11: 7, 1975) 3. R. P. KAMBOUR, Macromol. Rev. 7: 1, 1973 4. D. G. Le (]RAND, R. P. KOMBOUR and W. R. HAAF, J. Polymer Sci. 10, A-2: 1565, 1972 5. M. M. BUTSLOV, Ya. V. GENIN, V. I. GERASIMOV, A. M. DORFMAN and D. Ya. TSVANKIN, Pribory i Tekhn. Eksperimenta, No. 1, 199, 1972

THE MECHANISM OF POLYETHYLENE TEREPHTHALATE TRANSITION FROM THE AMORPHOUS TO THE ORIENTATED STATE* S. A. G~rBA~ov, Yu. I. MITCHE~KO,A. N. D'YACHKOV and E. 1~I. MZ~,~SHTEn~ All-Union Synthetic Fibres Research Institute (Received 4 November 1974)

Calorimetry, dilatometry, IR, 17MR and ESR spectroscopy, and optical polarization microscopy have been used in the study of the mechR,n[sm of P E T P transition from the amorphous to the orientated state during stretching. A small deformation stress has been found to cause irregular deformation. On the basis of the domain model for amorphous P E T P structure it is concluded that the deformation includes a relative shift of the domains or their parts wihout any significant structural change up to a deformation of 150-200% and a non-uniform straightening of the macrochains at various points of the sample, which, at a specific stage (on reaching a transconformation), creates stable and orientated parts which determine the crystallization characteristics of the deformed PETP. STUDIES of t h e s t r u c t u r e a n d p r o p e r t i e s of a m o r p h o u s p o l y m e r s r e s u l t e d in t h e

c r e a t i o n o f m o d e l s in w h i c h t h e s e p o l y m e r s h a v e a n o r i e n t a t e d s t r u c t u r e [1-3]. A n a t u r a l c o n t i n u a t i o n of s u c h i n v e s t i g a t i o n s is t h e s t u d y of t h e r e a r r a n g e m e n t of amorphous micro-orientated and micro-heterogeneous structures during deform a t i o n , especially d u r i n g t h e t r a n s i t i o n to t h e o r i e n t a t e d state. T h e s e p r o b l e m s were p a r t l y t a c k l e d b y o t h e r a u t h o r s [2-5]. H e r e we a t t e m p t t o find t h e m e c h a n i s m for t h e d e f o r m a t i o n s occurring d u r i n g m o n o a x i a l s t r e t c h i n g in a m o r p h o u s P E T P a n d use as t h e basis t h e p r e s e n c e in t h e p o l y m e r of t h e d o m a i n c o n c e p t s described before [6]. * Vysokomol. soyed. AI7: No. 11, 2469-2474, 1975.