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Dielectric loss in polyethyleneterephthalate
Dielectric loss m polyethyleneterephthalate
1401
(2) A method has been developed and is proposed for quantitative determination of alkoxy groups in polyphenylalkoxysiloxanes from infrared spectroscopic data. Translated by E. 0. P H i ~ r e s REFERENCES
1. A. I. SIDNEV, F. N. VISHNEVSK17, A. F. MOISEYEV, I. A. ZUBKOV and A. N. PRAVEDNIKOV, Vysokomol. soyed. A12: 355, 1970 (Translated in Polymer Sei. U.S.S.R. 12: 2, 1970) 2. A . P . TERENT'EVA, Ye. A. BONDAREVSKAYA, T. V. KIRILLOVA, N. A. GRADSKOVA and Ye. D. KROPOTOVA, Zh. analit, khim. 22: 454, 1967 3. A. P. TERENT'EVA, Ye. A. BONDAREVSKAYA and T. V. KIRILLOVA, Zh. analit. khim. 22: 1242, 1967 4. R. ZBINDEN, I K - s p e k t r o s k o p i y a ~%-sokopolimerov (Infrared Spectroscopy of High Polymers). p. 34, Izd. "Mir", 1966 5. A. I. SIDNEV, B. M. KOVARSKAYA and A. N. PRAVEDNIKOV, Vysokomol. soyed. B9: 134, 1967 (Not translated in Polymer Sei. U.S.S.R.) 6. Sh. A. SAMSONIYA, A. I. SIDNEV, O. V. SMIRNOVA, Yu. V. KHVASHCHEVSKAYA, I. A. ZUBKOV and H. S. KOLESNIKOV, Vysokomol. soyed. B10: 344, 1968 (Not translated in Polymer Sci. U.S.S.R.) 7. A. M:IYAKE, J. Chem. Soc., J a p a n , Ind. Chem. Soc. 62: 1449, 1959 8. W. D. D. BRYANT and R. C. VOTER, J. Amer. Chem. Soe. 75: 6113, 1953 9. F. P. REDING and C. M. LOVELL, J. P o l y m e r Sci. 21: 157, 1956 10. K. NAKANISI, Infrakrasnye spektry i stroyenie organicheskikh soyedniemi (Infrared Spectra and the Structure of Organic Compounds). p. 23, Izd. "Mir" 1965 11. C. W. YOUNG, P. C. SERVAIS, C. C. CURRIE and H. I. HUNTER, J. Amer. Chem. Soc. 70: 3758, 1948 12. A. A. V. STUART, C. LAHAN and M. BRUDERVELD, Rec. tray. chim. 74: 747, 1955 13. I. SPIELTER, D. C. PRIEST and C. N. HARRIS, J. Amer. Chem. Soc. 77: 6227, 1955
DIELECTRIC LOSS IN P O L Y E T H Y L E N E T E R E P H T H A L A T E * V. I. BEKICHEV and G. M. BARTENEV V. I. Lenin State Pedagogical Institute, Moscow
(Received
Ii
February 1969)
THERE is now extensive experimental n~aterial indicating the existence of two types of relaxational dielectric loss in linear, amorphous polymers, namely dipole-segmental loss associated with motion of the macromolecules within fairly large molecular volumes, and dipole-group loss associated with motion of the monomer units or of side groups, occurring within smaller volumes [1]. * Vysokomol. soyed. A12: ~-o. 6, 1240-1245, 1970.
1402
V.I.
BEKICHEV a n d O. BI. B.~-~tTE~-EV
Dipole-segmental losses occur in linear, amorphous polymers in the glassy state, even in those one would regard as being weakly polar, such as polystyrene, which has a dipole moment (~) of 0-2-0-3 debye [3]. The presence of a c~-stalline phase imparts some specific effects to the nature of the dependence of dielectric loss on temperature and frequency, but these effects have so far received little study. This is evidently because there are not so m a n y crystalline polymers as there are amorphous polymers. In most instances one is dealing with crystalline polymers that exist almost exclusively in the crystallized state {polyethylene, polytetrafluoroethylene etc.), in which it is impossible to follow changes in the dielectric properties with change in the degree of crystallization. Polyethyleneterephthalate (PETP) differs in that its glass temperature is above room temperature and therefore it can be obtained in both the amorphous and crystalline states. In view of the extensive use of P E T P in eleotrical insulation it is also of much interest to study the dependence of the dielectric properties of this polymer on its prehistory, which has a substantial effect on the structure of articles made from it [4, 5]. Reddish showed that there are two relaxation maxima in the temperature dependence of the dielectric loss factor (tan ~) of PETP. He attributed the lowtemperature maximum to the presence of terminal hydroxyl ~ o u p s in the polymer, and the high-temperature maximum to the presence of some amorphous phase, assuming that in the amorphous phase the P E T P molecules have the cis-configuration, in contrast to the crystalline phase, in which the molecules have the trans-configuration, as a result of which the dipole moment of the ester groups is self-compensating and losses are possible only in the amorphous phase [6]. The incorrectness of Reddish's conclusions has been discussed in papers by Mikhailov and his collaborators [3, 7, 8], who have proved convincingly that in polyesters, to which class of polymer P E T P belongs, dipole-group losses are due to relaxation of the C 0 0 group, and the most probable relaxation time for this process is determined by the mobility of chain segments joined directly to this group. The nature of the high-temperature maximum in P E T P remains obscure, because it is known that complete compensation of dipole moments in molecules with the trans-configuration, for example in ethylene succinate and adipate polymers does not occur, in spite of the fact the dipole moments of the polar groups of these polymers act in opposite directions [9]. In the present paper an attempt is made to examine the problem of the nature of the high-temperature tan ~ maximum in P E T P in relation to its physical state. The specimens studied were samples of Soviet P E T P of molecular weight 22,000, in the form of strips of thickness 150-200/1, obtained by rapid cooling of the molten pol:~uner on a metal drum immediately after its emergence from the slit of the extruder head. Highly ordered and highly crystalline films were obtained from the strip by the following procedure. The strip was heated to 90-110 °
Dielectric loss in polyethyleneterephthalate
1403
and stretched in the transverse direction to 3.5 times its original width, held in the stretched state at 165 ° and then cooled. The X-ray diffraction pattern of the strip was t}-pical of an amorphous isotropic structure and the film gave a pattern of a highly ordered, highly crystalline structure. Similar X-ray studies of P E T P obtained under conditions identical with those used in the present work are described in papers b y Kozlov and :Berestneva [4, 5, 10]. We have previeusly described the methods of obtaining the temperature dependence of tan 5 and the thermomechanical curves [11]. m/1%
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FIG. I. Therrnomeehanieal curves of aanorphous isotropic (1) and orientated (2) P E T P and the ternpersture dependence of tan 6 of amorphous P E T P st 400 llz (3). FIG. 2. Temperature dependence of tan 6 for P E T P at 400 llz: /--original, amorphous polymer; 2--specimen held at 120 ° for I0 rain; 3--30 rnin; 4--60 rnin; 5--film obtained by stretching amorphous specimen at 85 ° and holding for 1 rain at 165 °.
The thermomechanical curves of the amorphous P E T P in the temperature interval of 20-115 ° are typical of the curves of linear, amorphous polymers. The unorientated specimens begin to deform appreciably at 60 ° (Fig. 1, curve 1) while the unorientated specimens begin to contract at this same temperature (Fig. 1, curve 2). This contraction is atributable to the onset of chain mobility, which, as is well known, occurs near to the softening point (glass temperature). After this the shape of the thermomechanical curve 2 is roughly the same as that of the curve of the unorientated specimen. At 115 ° the increase in deformation slows down and the curves form a plateau, extending to 220 °. The formation of the plateau is explained b y the sharp increase in the rate of crystallization in the viscous-flow state, as a result of which flow is retarded, disappearing altogether at 120 °. Deformation increases again at 220 ° as a result of melting, which finally leads to rupture of the specimen at 240 °. ~Vhen amorphous samples are kept at 120 ° or higher the amorphous halo gradually disappears from the X-radio-
1404
V. I. B]:KICHEV and G. M. B,U~TE.','EV
grams and the number of D~bye rings increases. Thus the original specimens (strips) of P E T P are in the glassy state up to 60 °, in the high-elastic state in the temperature interval of 60-95 ° and in the viscous-flow state from 95 to 120 °. At 120 ° amorphous specimens crystallize e,nd then melt at 220-240 °. As for all linear amorphous polymers tan J at low frequencies (400 Hz) passes through a maximum in the softening point region (Fig. 1, curve 3), which characterizes dipole-segmental loss. It was not possible to obtain a sufficient number of points for determination of the energy of activation for this loss from the shift in the tan j maximum because it shifts right up to the crystallization temperature with only a two-fold increase in frequency. Usually immediately after the maximum for dipole-segmental loss, after a minimum has been passed tan j for linear, amorphous polymers increases approximately exponentially as a result of loss of conductivity [2]. This increase in tan J occurs at 110 ° for amorphous P E T P , but the sharp increase in the rate of crystallization at 115-120 ° changes the polymer from the softened, amorphous state to the hard, crystalline state. As a result of this, increase in the loss of conductivity is retarded and tan J again passes through a small maximum. From data presented in Sazhin's monograph [12] the specific volume resistance increases in P E T P at approximately the same temperatures, then after reaching a maximum falls according to an exponential law. This indicates that here the tan J maximum is explained not by any relaxation process but b y change in the electrical conductivity due to crystallization of the polymer. The essential nature of the processes involved can be understood from the data presented in Fig. 2, from which it can be seen that the tan J maximum shifts to higher temperatures as crystallization proceeds. When amorphous P E T P is kept at 120 ° for 10 and 30 min crystallization, as indicated b y decrease in the rate of loss of conductivity at 120 °, is not complete. After 60 min at 120 ° crystallization is fairly complete. Further heating at this temperature does not alter the relaxation picture. Above 120 ° curves 1-4 merge as a result of continued crystallization during the measurement of tan J and normally increase in the loss of conductivity appears in them at 140% The temperature of the tan J maximum for the highly orientated, highly e~-stalline film (curve .5) is higher and the value of tan J at the maximum lower than for the freely crystallized specimens 2-4. This shows that preliminary orientation promotes more extensive crystallization. Thus crystallization of P E T P has a substantial effect on both the nature of the processes involved in dipole-segmental relaxation, and on its electrical conduetivity. As the degree of crystallization increases the most probable relaxation time is increased and the spectrum of relaxation times is broadened, resulting in broadening of the region of the tan J maximum. The value of tan J at the maximum for a highly crystalline sample (Fig. 2, curve -5) is one seventh of the corresponding value for amorphous P E T P (Fig. 2, curve 1). Crystallization causes decrease in the loss in conductivity. Note that good agreement was obtained in the temperature of the tan J maximum for all Soviet and foreign P E T P films
1405
Dielectric loss in polyethy leneterephthalate
studied. The curves obtained (Fig. 2, curve 5) coincide with similar curves in references [6] and [13]. The nature of the change in the dipole-segmental processes as crystallization proceeds indicates that the high-temperature maximum in crystalline P E T P is due to the presence of an amorphous phase. This conclusion is supported b y the energy of activation for this loss, determined for the highly orientated, highly crystalline film, the structure of which does not alter when it is heated to 150 °. The dependence of log fro on reciprocal temperature for this film is given in Fig. 3 (here fm is the frequency and T the temperature (in °K) at which the tan 5 maximum appears). The energy of activation of 64 kcal/mole found from Fig. 3 is characteristic of dipole-segmental relaxation loss. It is welI known that extrapolation of the curve of logf~-----~0 (tIT) to a frequency of 1 Hz. gives a temperature close to the glass temperature [7]. In Fig. 3 this temperature is 84 °, which is in agreement with glass temperatures quoted in the log fro, H z 6
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:FIG. 3. Temperature dependence of the frequency corresponding to the dielectric loss for PETP film obtained by stretching an amorphous sample. FIG. 4. Thermomechanical curve (1) and the temperature dependence of tan 5 of PETP film at 400 Hz (2). literature for P E T P without reference to its structure [14]. As was stated above,. it was not possible to construct a similar curve for the amorphous pol}nner. h'evertheless there are grounds for assuming that here we are dealing with two glass temperatures for P E T P , i.e. the glass temperature of the amorphous polymer (60 °) and the glass temperature of the amorphous phase coexisting with t h e crystalline phase (84°).
1406
•
v. I. BEKICttEV o,nd G..~,I. B.~TE_~'EV
The change in the nature of the dipole-segmental relaxation processes when P E T P crystallizes is similar to the change occurring in the crystallization of neoprene rubber [3]. In both instances crystallization brings about an increase in the relaxation time, broadening of the spectrum of relaxation times and a decrease in tan ~f at the maximum. The molecule of neoprene rubber does not exhibit cis-trans isomerism because each monomer unit contains only one chlorine atom. _~[ikhailov and his collaborators have shown by calculation, however, that when the crystallized rubber is stretched the dipole moment and the distribution parameter ~ decrease, indicating broadening of the spectrum of relaxation times and hence decrease in the value of tan ~ at the maximum. Therefore, the change in the nature of the relaxation processes during crystallization of neoprene is explained only by the assumption that the crystalline formations act as cL'osslinkages in the amorphous regions, altering the local mobility of the polymer molecules and at the same time the orientational mobility of the polar radicals bound chemically to them. The results presented in this paper show that when amorphous P E T P is heated to 120 ° the same relaxation processes occur as in any linear amorphous polymer. Since dipole-group relaxation is due to motion of CO0 groups, as has been shown in references [3], [7] and [8], dipolesegmental relaxation is associated with these same groups. The change in the nature of the latter t~-pe of relaxation as crystallization of P E T P proceeds is similar to the change in crystallizing rubbers. It must therefore obviously be assumed that the cause of these changes is the same, i.e. the "crosslinking" effect of c~-stalline regions. The high degree of crosslinking of polymers of high crystallinity can reduce dipole losses to such an extent that either they can be detected only by very sensitive apparatus or they disappear altogether, as in the case of the transition from amorphous to crystalline sugar, studied by Aleksandrov, Kobeko and Kuvshinskii [2]. This is confirmed by Fig. 4, which shows the temperature dependence of tan ~ (400 Hz) and the thermomechanical curve of low-pressure polyethylene (LPPE), which has a fairly high concentration of polar radicals [3]. It is obvious that if it were possible to obtain amorphous L P P E close to the glass temperature, as in the case of polystyrene there would be a dipole-segmental loss maximum. I~, is seen however from Fig. 4 that the sensitivity of the .~YLE-1 bridge is insufficient to reveal any relaxation maxima at all, which is attributable to the high degTee of crystallinity of LPPE. ~foreover, the increase in loss of conductivity that is usual in linear polymers occurs in L P P E only after melting. This is clearly seen in the thermomechanical curve (:Fig. 4, curve 1). ~Iikhailov and his collaborators were able to find low-, medium- and high-frequency relaxation in L P P E only by the use of sensitive apparatus [3]. I t is evident that decrease in dipole losses as a result of crystallization is a property common to both polymers and simple substances like glucose and sugar [2]. This can be made use of in practice to altel' not only the mechanicM, but also the electrical insulation properties of polymeric insulating films. By crystallizing film in the orientated state it is possible both to reduce dielectric loss and increase
Dielectric loss in polyethyleneterephthalate
1407
the heat resistance in comparison with a film of amorphous polymer. From this point of view polycarbonates, the glass temperature of which can be as high as 160-180 ° [15], are very interesting. In conclusion it should be noted t h a t from the nature of the dielectric relaxation processes in partially crystallized P E T P (Fig. 2) conclusions can be ch'ax~a about its physical states, and the degree of crystallinity can be determined by the use of suitable refined methods. CONCLUSIONS
(1) A study has been mexle of dipole-segmental relaxation in amorphous polyethyleneterephthalate (PETP) as it undergoes crystallization. (2) I t was found t h a t the decrease in the amount of amorphous phase is accompanied by a shift in the position of the tan 5 maximum to higher temperatures. At the same time the value of tan 5 at the maximum decreases and the region of the maximum broadens, indicating broadening of the spectrum of relaxation times. (3) I t is shown t h a t dipole tosse, in crystalline P E T P are due to the presence of an amorphous phase. (4) It has been proved t h a t the glass temperature of the amorphous phase in highly crystalline, highly orientated P E T P is higher than the glass temperature of the amorphous isotropic polymer. (5) Comparison of the change in the relaxation processes in P E T P during crystallization with the change in other crystallizable polymers in which there is no cis-trans isomerism suggests t h a t these changes are clue to the "crosslinking" effect of micro-regions of crystalline phase. Translated by E. O. p~r:rT,treS REFERENCES
1. G. P. ~ ~ O V , Sb. Relaksatsionnye yavleniya v tverdykh telakh (Collectedpapers. Relaxation Effects in Solid Bodies). p. 76, Izd. "Metallurgiya", 1968 2. P. P. KOBEKO, Amorfnye veshchestva (Amorphous Substances). Izd. Akad. _~'auk SSSR, 1952 3. G. P. I~KHAILOV, Uspekhi khimii 24: 875, 1955 4. P. ¥. KOZLOVand @. L. BERESTNEVA, Vysokomol. soyed. 2: 590, 1960 (Not translated in Polymer Sci. U.S.S.R.) 5. G. L. BERESTNEYA and P. V. KOZLOV,Vysokomol. soyed. 2: 601, 1960 (Not translated in Polymer Sci. U.S.S.R.) 6. W. REDDISH, Trans. Faraday Soc. 46: 459, 1950 7. G. P. ~IIKFIAILOVand IV[.P. EIDEL'NANT, Vysokomol. soyed. 2: 287, 1960 (Not transIated in Polymer Sci. U.S.S.R.) 8. G. P. MIKHAILOV and T. I. BORISOV, Uspekhi khimii 30: 895, 1961 9. D. R. PELMORE and E. L. S~.IONS, Proc. Roy. Soc. A175: 468, 1940 10. G. L. BERESTNEVA, D. Ya. TSVANKIN and P. V. KOZLOV, Vysokomol. soyed. 3: 1786, 1961 (Not translated flu Polymer Sci. U.S.S.R.) 11. V. I. BEKICHEV, Zavod. lab, 35: 748, 196g
1408
T . A . AFrOV.~ et al.
12. B. I. SAZIffLN', Elektroprovodnost' polimerov (Electrical Conductivity of Polymers). Izd. " K h i m i y a " , 1964 13. R. BOYER, Perekhody i relaksatsionnye yavleniya v polimerakh (Transitions and Relaxation Phenomena in Polymers). Izd. "Mir", 1968 14. V. V. KORSH~AK, Termostoikie polimery (Heat-Resistant Polymers). Izd. "*-auka", 1969 15. G. SI-~'ELL, K h i m i y a i fizika polikarbonatov (Chemistry and Physics of Polycarbonates). Izd. " K h i m i y a " , 1967
THE HOMOPOLYMERIZATION AND COPOLYMERIZATION WITH STYRENE OF DIVINYLBENZENE ISOMERS* T. A. :-~PTOVA, YIj. Y.-k. BABUSHKIN, YE. A. GUKASOVA, YE. V. YEGOROV, G. V. KOROLEV, S. B. ~a,I.4.KAROVA, B. 1~. S3[IRN0V a n d T. M. CHERB-YAVSKAYA Branch of the I n s t i t u t e of Chemical Physics, U.S.S.R. A c a d e m y of Sciences All-Union Research I n s t i t u t e of Chemical Reagents and Highly Purified Chemicals
(Received 24 February 1969) DESPITE the fact t h a t copol?-mers of styrene (St) and divinylbenzene (DVB) are used very
extensively as ion-exchangers, detailed s t u d y of their copolymerization and of the homopolymerization of DVB has begun only very recently. This is due both to the experimental difficulty of studying three-dimensional copolymerization and the impossibility of determining the composition of these copolymers by ordinary analytical methods and, until recently, to the absence of reliable methods of preparation and isolation of the individual DVB isomers. Published information on the homopollunerization of DVB and its copolymerization with styrene is very limited. The polymerization of DVB in toluene and tert.butylbenzene was studied by the viscometric method in references [1] and [2]. A t 60 °, depending on conversion, the following ratios of the rates were obtained: Wm/I~= 1'50-1"58, Wm/W..m+~ ----1"26-1"29, W2m+a~/Wp=l'19-1"12.t These ratios are a little higher at 80 °. The rates of polymerization of the series of isomers decrease in the following order: m - D V B > ( 2 m + lp)D V B > commercial D V B > the p-isomer. The energy of activation for polymerization increases with increase in conversion and decreases in the following order p - D V B > ( 2 m + lp)D V B > m-DVB, the respective values being 20.8-27.1, 17.8-23.7 and 17.8-23-3 kcal/mole. F o r cormnercial DVB the energy of activation varies with the degree of conversion from 14.9 to 23.8 kcal/mole. The rates of polymerization of all the DVB isomers are higher than the rate of polymerization of styrene. I n contrast to the m- and p-isomers, when o-DVB polymerizes it usually forms soluble polymers containing five-membered rings (intraanolec* Vysokomol. soyed. A12: No. 6, 1246-1253, 1970. t F o u n d from the time of p61~anerization for different DVB isomers to the same time of flow of the solution. Win, Wp and W~+l~ are the rates of polymerization of the rn., pand a 2 : 1 mixture of the m- and p-isomers of DVB respectively.
1401
(2) A method has been developed and is proposed for quantitative determination of alkoxy groups in polyphenylalkoxysiloxanes from infrared spectroscopic data. Translated by E. 0. P H i ~ r e s REFERENCES
1. A. I. SIDNEV, F. N. VISHNEVSK17, A. F. MOISEYEV, I. A. ZUBKOV and A. N. PRAVEDNIKOV, Vysokomol. soyed. A12: 355, 1970 (Translated in Polymer Sei. U.S.S.R. 12: 2, 1970) 2. A . P . TERENT'EVA, Ye. A. BONDAREVSKAYA, T. V. KIRILLOVA, N. A. GRADSKOVA and Ye. D. KROPOTOVA, Zh. analit, khim. 22: 454, 1967 3. A. P. TERENT'EVA, Ye. A. BONDAREVSKAYA and T. V. KIRILLOVA, Zh. analit. khim. 22: 1242, 1967 4. R. ZBINDEN, I K - s p e k t r o s k o p i y a ~%-sokopolimerov (Infrared Spectroscopy of High Polymers). p. 34, Izd. "Mir", 1966 5. A. I. SIDNEV, B. M. KOVARSKAYA and A. N. PRAVEDNIKOV, Vysokomol. soyed. B9: 134, 1967 (Not translated in Polymer Sei. U.S.S.R.) 6. Sh. A. SAMSONIYA, A. I. SIDNEV, O. V. SMIRNOVA, Yu. V. KHVASHCHEVSKAYA, I. A. ZUBKOV and H. S. KOLESNIKOV, Vysokomol. soyed. B10: 344, 1968 (Not translated in Polymer Sci. U.S.S.R.) 7. A. M:IYAKE, J. Chem. Soc., J a p a n , Ind. Chem. Soc. 62: 1449, 1959 8. W. D. D. BRYANT and R. C. VOTER, J. Amer. Chem. Soe. 75: 6113, 1953 9. F. P. REDING and C. M. LOVELL, J. P o l y m e r Sci. 21: 157, 1956 10. K. NAKANISI, Infrakrasnye spektry i stroyenie organicheskikh soyedniemi (Infrared Spectra and the Structure of Organic Compounds). p. 23, Izd. "Mir" 1965 11. C. W. YOUNG, P. C. SERVAIS, C. C. CURRIE and H. I. HUNTER, J. Amer. Chem. Soc. 70: 3758, 1948 12. A. A. V. STUART, C. LAHAN and M. BRUDERVELD, Rec. tray. chim. 74: 747, 1955 13. I. SPIELTER, D. C. PRIEST and C. N. HARRIS, J. Amer. Chem. Soc. 77: 6227, 1955
DIELECTRIC LOSS IN P O L Y E T H Y L E N E T E R E P H T H A L A T E * V. I. BEKICHEV and G. M. BARTENEV V. I. Lenin State Pedagogical Institute, Moscow
(Received
Ii
February 1969)
THERE is now extensive experimental n~aterial indicating the existence of two types of relaxational dielectric loss in linear, amorphous polymers, namely dipole-segmental loss associated with motion of the macromolecules within fairly large molecular volumes, and dipole-group loss associated with motion of the monomer units or of side groups, occurring within smaller volumes [1]. * Vysokomol. soyed. A12: ~-o. 6, 1240-1245, 1970.
1402
V.I.
BEKICHEV a n d O. BI. B.~-~tTE~-EV
Dipole-segmental losses occur in linear, amorphous polymers in the glassy state, even in those one would regard as being weakly polar, such as polystyrene, which has a dipole moment (~) of 0-2-0-3 debye [3]. The presence of a c~-stalline phase imparts some specific effects to the nature of the dependence of dielectric loss on temperature and frequency, but these effects have so far received little study. This is evidently because there are not so m a n y crystalline polymers as there are amorphous polymers. In most instances one is dealing with crystalline polymers that exist almost exclusively in the crystallized state {polyethylene, polytetrafluoroethylene etc.), in which it is impossible to follow changes in the dielectric properties with change in the degree of crystallization. Polyethyleneterephthalate (PETP) differs in that its glass temperature is above room temperature and therefore it can be obtained in both the amorphous and crystalline states. In view of the extensive use of P E T P in eleotrical insulation it is also of much interest to study the dependence of the dielectric properties of this polymer on its prehistory, which has a substantial effect on the structure of articles made from it [4, 5]. Reddish showed that there are two relaxation maxima in the temperature dependence of the dielectric loss factor (tan ~) of PETP. He attributed the lowtemperature maximum to the presence of terminal hydroxyl ~ o u p s in the polymer, and the high-temperature maximum to the presence of some amorphous phase, assuming that in the amorphous phase the P E T P molecules have the cis-configuration, in contrast to the crystalline phase, in which the molecules have the trans-configuration, as a result of which the dipole moment of the ester groups is self-compensating and losses are possible only in the amorphous phase [6]. The incorrectness of Reddish's conclusions has been discussed in papers by Mikhailov and his collaborators [3, 7, 8], who have proved convincingly that in polyesters, to which class of polymer P E T P belongs, dipole-group losses are due to relaxation of the C 0 0 group, and the most probable relaxation time for this process is determined by the mobility of chain segments joined directly to this group. The nature of the high-temperature maximum in P E T P remains obscure, because it is known that complete compensation of dipole moments in molecules with the trans-configuration, for example in ethylene succinate and adipate polymers does not occur, in spite of the fact the dipole moments of the polar groups of these polymers act in opposite directions [9]. In the present paper an attempt is made to examine the problem of the nature of the high-temperature tan ~ maximum in P E T P in relation to its physical state. The specimens studied were samples of Soviet P E T P of molecular weight 22,000, in the form of strips of thickness 150-200/1, obtained by rapid cooling of the molten pol:~uner on a metal drum immediately after its emergence from the slit of the extruder head. Highly ordered and highly crystalline films were obtained from the strip by the following procedure. The strip was heated to 90-110 °
Dielectric loss in polyethyleneterephthalate
1403
and stretched in the transverse direction to 3.5 times its original width, held in the stretched state at 165 ° and then cooled. The X-ray diffraction pattern of the strip was t}-pical of an amorphous isotropic structure and the film gave a pattern of a highly ordered, highly crystalline structure. Similar X-ray studies of P E T P obtained under conditions identical with those used in the present work are described in papers b y Kozlov and :Berestneva [4, 5, 10]. We have previeusly described the methods of obtaining the temperature dependence of tan 5 and the thermomechanical curves [11]. m/1%
fan&lO z ,
/
I20
6
#0'
2
0
80
IZ~0 FIG. i
i
220 T,°C
0
fan~40 z
8
= 3
20
60
I00
Iz~O
180
T,°C
FIG. 2
FIG. I. Therrnomeehanieal curves of aanorphous isotropic (1) and orientated (2) P E T P and the ternpersture dependence of tan 6 of amorphous P E T P st 400 llz (3). FIG. 2. Temperature dependence of tan 6 for P E T P at 400 llz: /--original, amorphous polymer; 2--specimen held at 120 ° for I0 rain; 3--30 rnin; 4--60 rnin; 5--film obtained by stretching amorphous specimen at 85 ° and holding for 1 rain at 165 °.
The thermomechanical curves of the amorphous P E T P in the temperature interval of 20-115 ° are typical of the curves of linear, amorphous polymers. The unorientated specimens begin to deform appreciably at 60 ° (Fig. 1, curve 1) while the unorientated specimens begin to contract at this same temperature (Fig. 1, curve 2). This contraction is atributable to the onset of chain mobility, which, as is well known, occurs near to the softening point (glass temperature). After this the shape of the thermomechanical curve 2 is roughly the same as that of the curve of the unorientated specimen. At 115 ° the increase in deformation slows down and the curves form a plateau, extending to 220 °. The formation of the plateau is explained b y the sharp increase in the rate of crystallization in the viscous-flow state, as a result of which flow is retarded, disappearing altogether at 120 °. Deformation increases again at 220 ° as a result of melting, which finally leads to rupture of the specimen at 240 °. ~Vhen amorphous samples are kept at 120 ° or higher the amorphous halo gradually disappears from the X-radio-
1404
V. I. B]:KICHEV and G. M. B,U~TE.','EV
grams and the number of D~bye rings increases. Thus the original specimens (strips) of P E T P are in the glassy state up to 60 °, in the high-elastic state in the temperature interval of 60-95 ° and in the viscous-flow state from 95 to 120 °. At 120 ° amorphous specimens crystallize e,nd then melt at 220-240 °. As for all linear amorphous polymers tan J at low frequencies (400 Hz) passes through a maximum in the softening point region (Fig. 1, curve 3), which characterizes dipole-segmental loss. It was not possible to obtain a sufficient number of points for determination of the energy of activation for this loss from the shift in the tan j maximum because it shifts right up to the crystallization temperature with only a two-fold increase in frequency. Usually immediately after the maximum for dipole-segmental loss, after a minimum has been passed tan j for linear, amorphous polymers increases approximately exponentially as a result of loss of conductivity [2]. This increase in tan J occurs at 110 ° for amorphous P E T P , but the sharp increase in the rate of crystallization at 115-120 ° changes the polymer from the softened, amorphous state to the hard, crystalline state. As a result of this, increase in the loss of conductivity is retarded and tan J again passes through a small maximum. From data presented in Sazhin's monograph [12] the specific volume resistance increases in P E T P at approximately the same temperatures, then after reaching a maximum falls according to an exponential law. This indicates that here the tan J maximum is explained not by any relaxation process but b y change in the electrical conductivity due to crystallization of the polymer. The essential nature of the processes involved can be understood from the data presented in Fig. 2, from which it can be seen that the tan J maximum shifts to higher temperatures as crystallization proceeds. When amorphous P E T P is kept at 120 ° for 10 and 30 min crystallization, as indicated b y decrease in the rate of loss of conductivity at 120 °, is not complete. After 60 min at 120 ° crystallization is fairly complete. Further heating at this temperature does not alter the relaxation picture. Above 120 ° curves 1-4 merge as a result of continued crystallization during the measurement of tan J and normally increase in the loss of conductivity appears in them at 140% The temperature of the tan J maximum for the highly orientated, highly e~-stalline film (curve .5) is higher and the value of tan J at the maximum lower than for the freely crystallized specimens 2-4. This shows that preliminary orientation promotes more extensive crystallization. Thus crystallization of P E T P has a substantial effect on both the nature of the processes involved in dipole-segmental relaxation, and on its electrical conduetivity. As the degree of crystallization increases the most probable relaxation time is increased and the spectrum of relaxation times is broadened, resulting in broadening of the region of the tan J maximum. The value of tan J at the maximum for a highly crystalline sample (Fig. 2, curve -5) is one seventh of the corresponding value for amorphous P E T P (Fig. 2, curve 1). Crystallization causes decrease in the loss in conductivity. Note that good agreement was obtained in the temperature of the tan J maximum for all Soviet and foreign P E T P films
1405
Dielectric loss in polyethy leneterephthalate
studied. The curves obtained (Fig. 2, curve 5) coincide with similar curves in references [6] and [13]. The nature of the change in the dipole-segmental processes as crystallization proceeds indicates that the high-temperature maximum in crystalline P E T P is due to the presence of an amorphous phase. This conclusion is supported b y the energy of activation for this loss, determined for the highly orientated, highly crystalline film, the structure of which does not alter when it is heated to 150 °. The dependence of log fro on reciprocal temperature for this film is given in Fig. 3 (here fm is the frequency and T the temperature (in °K) at which the tan 5 maximum appears). The energy of activation of 64 kcal/mole found from Fig. 3 is characteristic of dipole-segmental relaxation loss. It is welI known that extrapolation of the curve of logf~-----~0 (tIT) to a frequency of 1 Hz. gives a temperature close to the glass temperature [7]. In Fig. 3 this temperature is 84 °, which is in agreement with glass temperatures quoted in the log fro, H z 6
"4Wz'%
sI
2 -20
o
2,," \ \ 2
\
\
-$0
\
f 0
,
\
I,
I
2.~
2.~ Fro. 3
0
\
l
(8,°c)
--20 FIG 4
:FIG. 3. Temperature dependence of the frequency corresponding to the dielectric loss for PETP film obtained by stretching an amorphous sample. FIG. 4. Thermomechanical curve (1) and the temperature dependence of tan 5 of PETP film at 400 Hz (2). literature for P E T P without reference to its structure [14]. As was stated above,. it was not possible to construct a similar curve for the amorphous pol}nner. h'evertheless there are grounds for assuming that here we are dealing with two glass temperatures for P E T P , i.e. the glass temperature of the amorphous polymer (60 °) and the glass temperature of the amorphous phase coexisting with t h e crystalline phase (84°).
1406
•
v. I. BEKICttEV o,nd G..~,I. B.~TE_~'EV
The change in the nature of the dipole-segmental relaxation processes when P E T P crystallizes is similar to the change occurring in the crystallization of neoprene rubber [3]. In both instances crystallization brings about an increase in the relaxation time, broadening of the spectrum of relaxation times and a decrease in tan ~f at the maximum. The molecule of neoprene rubber does not exhibit cis-trans isomerism because each monomer unit contains only one chlorine atom. _~[ikhailov and his collaborators have shown by calculation, however, that when the crystallized rubber is stretched the dipole moment and the distribution parameter ~ decrease, indicating broadening of the spectrum of relaxation times and hence decrease in the value of tan ~ at the maximum. Therefore, the change in the nature of the relaxation processes during crystallization of neoprene is explained only by the assumption that the crystalline formations act as cL'osslinkages in the amorphous regions, altering the local mobility of the polymer molecules and at the same time the orientational mobility of the polar radicals bound chemically to them. The results presented in this paper show that when amorphous P E T P is heated to 120 ° the same relaxation processes occur as in any linear amorphous polymer. Since dipole-group relaxation is due to motion of CO0 groups, as has been shown in references [3], [7] and [8], dipolesegmental relaxation is associated with these same groups. The change in the nature of the latter t~-pe of relaxation as crystallization of P E T P proceeds is similar to the change in crystallizing rubbers. It must therefore obviously be assumed that the cause of these changes is the same, i.e. the "crosslinking" effect of c~-stalline regions. The high degree of crosslinking of polymers of high crystallinity can reduce dipole losses to such an extent that either they can be detected only by very sensitive apparatus or they disappear altogether, as in the case of the transition from amorphous to crystalline sugar, studied by Aleksandrov, Kobeko and Kuvshinskii [2]. This is confirmed by Fig. 4, which shows the temperature dependence of tan ~ (400 Hz) and the thermomechanical curve of low-pressure polyethylene (LPPE), which has a fairly high concentration of polar radicals [3]. It is obvious that if it were possible to obtain amorphous L P P E close to the glass temperature, as in the case of polystyrene there would be a dipole-segmental loss maximum. I~, is seen however from Fig. 4 that the sensitivity of the .~YLE-1 bridge is insufficient to reveal any relaxation maxima at all, which is attributable to the high degTee of crystallinity of LPPE. ~foreover, the increase in loss of conductivity that is usual in linear polymers occurs in L P P E only after melting. This is clearly seen in the thermomechanical curve (:Fig. 4, curve 1). ~Iikhailov and his collaborators were able to find low-, medium- and high-frequency relaxation in L P P E only by the use of sensitive apparatus [3]. I t is evident that decrease in dipole losses as a result of crystallization is a property common to both polymers and simple substances like glucose and sugar [2]. This can be made use of in practice to altel' not only the mechanicM, but also the electrical insulation properties of polymeric insulating films. By crystallizing film in the orientated state it is possible both to reduce dielectric loss and increase
Dielectric loss in polyethyleneterephthalate
1407
the heat resistance in comparison with a film of amorphous polymer. From this point of view polycarbonates, the glass temperature of which can be as high as 160-180 ° [15], are very interesting. In conclusion it should be noted t h a t from the nature of the dielectric relaxation processes in partially crystallized P E T P (Fig. 2) conclusions can be ch'ax~a about its physical states, and the degree of crystallinity can be determined by the use of suitable refined methods. CONCLUSIONS
(1) A study has been mexle of dipole-segmental relaxation in amorphous polyethyleneterephthalate (PETP) as it undergoes crystallization. (2) I t was found t h a t the decrease in the amount of amorphous phase is accompanied by a shift in the position of the tan 5 maximum to higher temperatures. At the same time the value of tan 5 at the maximum decreases and the region of the maximum broadens, indicating broadening of the spectrum of relaxation times. (3) I t is shown t h a t dipole tosse, in crystalline P E T P are due to the presence of an amorphous phase. (4) It has been proved t h a t the glass temperature of the amorphous phase in highly crystalline, highly orientated P E T P is higher than the glass temperature of the amorphous isotropic polymer. (5) Comparison of the change in the relaxation processes in P E T P during crystallization with the change in other crystallizable polymers in which there is no cis-trans isomerism suggests t h a t these changes are clue to the "crosslinking" effect of micro-regions of crystalline phase. Translated by E. O. p~r:rT,treS REFERENCES
1. G. P. ~ ~ O V , Sb. Relaksatsionnye yavleniya v tverdykh telakh (Collectedpapers. Relaxation Effects in Solid Bodies). p. 76, Izd. "Metallurgiya", 1968 2. P. P. KOBEKO, Amorfnye veshchestva (Amorphous Substances). Izd. Akad. _~'auk SSSR, 1952 3. G. P. I~KHAILOV, Uspekhi khimii 24: 875, 1955 4. P. ¥. KOZLOVand @. L. BERESTNEVA, Vysokomol. soyed. 2: 590, 1960 (Not translated in Polymer Sci. U.S.S.R.) 5. G. L. BERESTNEYA and P. V. KOZLOV,Vysokomol. soyed. 2: 601, 1960 (Not translated in Polymer Sci. U.S.S.R.) 6. W. REDDISH, Trans. Faraday Soc. 46: 459, 1950 7. G. P. ~IIKFIAILOVand IV[.P. EIDEL'NANT, Vysokomol. soyed. 2: 287, 1960 (Not transIated in Polymer Sci. U.S.S.R.) 8. G. P. MIKHAILOV and T. I. BORISOV, Uspekhi khimii 30: 895, 1961 9. D. R. PELMORE and E. L. S~.IONS, Proc. Roy. Soc. A175: 468, 1940 10. G. L. BERESTNEVA, D. Ya. TSVANKIN and P. V. KOZLOV, Vysokomol. soyed. 3: 1786, 1961 (Not translated flu Polymer Sci. U.S.S.R.) 11. V. I. BEKICHEV, Zavod. lab, 35: 748, 196g
1408
T . A . AFrOV.~ et al.
12. B. I. SAZIffLN', Elektroprovodnost' polimerov (Electrical Conductivity of Polymers). Izd. " K h i m i y a " , 1964 13. R. BOYER, Perekhody i relaksatsionnye yavleniya v polimerakh (Transitions and Relaxation Phenomena in Polymers). Izd. "Mir", 1968 14. V. V. KORSH~AK, Termostoikie polimery (Heat-Resistant Polymers). Izd. "*-auka", 1969 15. G. SI-~'ELL, K h i m i y a i fizika polikarbonatov (Chemistry and Physics of Polycarbonates). Izd. " K h i m i y a " , 1967
THE HOMOPOLYMERIZATION AND COPOLYMERIZATION WITH STYRENE OF DIVINYLBENZENE ISOMERS* T. A. :-~PTOVA, YIj. Y.-k. BABUSHKIN, YE. A. GUKASOVA, YE. V. YEGOROV, G. V. KOROLEV, S. B. ~a,I.4.KAROVA, B. 1~. S3[IRN0V a n d T. M. CHERB-YAVSKAYA Branch of the I n s t i t u t e of Chemical Physics, U.S.S.R. A c a d e m y of Sciences All-Union Research I n s t i t u t e of Chemical Reagents and Highly Purified Chemicals
(Received 24 February 1969) DESPITE the fact t h a t copol?-mers of styrene (St) and divinylbenzene (DVB) are used very
extensively as ion-exchangers, detailed s t u d y of their copolymerization and of the homopolymerization of DVB has begun only very recently. This is due both to the experimental difficulty of studying three-dimensional copolymerization and the impossibility of determining the composition of these copolymers by ordinary analytical methods and, until recently, to the absence of reliable methods of preparation and isolation of the individual DVB isomers. Published information on the homopollunerization of DVB and its copolymerization with styrene is very limited. The polymerization of DVB in toluene and tert.butylbenzene was studied by the viscometric method in references [1] and [2]. A t 60 °, depending on conversion, the following ratios of the rates were obtained: Wm/I~= 1'50-1"58, Wm/W..m+~ ----1"26-1"29, W2m+a~/Wp=l'19-1"12.t These ratios are a little higher at 80 °. The rates of polymerization of the series of isomers decrease in the following order: m - D V B > ( 2 m + lp)D V B > commercial D V B > the p-isomer. The energy of activation for polymerization increases with increase in conversion and decreases in the following order p - D V B > ( 2 m + lp)D V B > m-DVB, the respective values being 20.8-27.1, 17.8-23.7 and 17.8-23-3 kcal/mole. F o r cormnercial DVB the energy of activation varies with the degree of conversion from 14.9 to 23.8 kcal/mole. The rates of polymerization of all the DVB isomers are higher than the rate of polymerization of styrene. I n contrast to the m- and p-isomers, when o-DVB polymerizes it usually forms soluble polymers containing five-membered rings (intraanolec* Vysokomol. soyed. A12: No. 6, 1246-1253, 1970. t F o u n d from the time of p61~anerization for different DVB isomers to the same time of flow of the solution. Win, Wp and W~+l~ are the rates of polymerization of the rn., pand a 2 : 1 mixture of the m- and p-isomers of DVB respectively.