Relationship between the mechanical behaviour and molecular weight distribution of high density polyethylene

Relationship between mechanical behaviour and molecular weight distribution of P E

1275

8. A. WE1NBERG and E. PROSKAUER, Organic Solvents, 199, 1958 9. Ye. V. KOCHETOV, A. A. BERLIN a n d N. S. YENIKOLOPYAN, Vysokomol. soyed. 8: 1022, 1966 (Translated in Polymer Sei. U.S.S.R. 8: 6, 1122, 1966) I0. Ye. V. KOCHETOV, M. A. MARKEVICH and N. S. YENIKOLOPYAN, Dokl. A N SSSR 180: 143, 1968

RELATIONSHIP BETWEEN THE MECHANICAL BEHAVIOUR AND MOLECULAR WEIGHT DISTRIBUTION OF HIGH DENSITY POLYETHYLENE* A. N. KARASEV, I. l~. AIVDREYEVA, N. M. DOI~IAREVA, K. I. KOS~IATYK~, M. G. KARASEVA and N. A. DOM~ICHEVA Plastics Research I n s t i t u t e

(Received 21 April 1969) THE molecular weight distribution (MWD) of polymers is one of the main factors determining their properties. In the case of a narrow MWD it is sufficient to know the average molecular weight, b u t with polymers having a wide MWD (including polyethylene, particularly when prepared with heterogeneous catalysts such as A1 (C2H5)2C1÷TIC14) it is important to study the behaviour of the polymers in relation to the width of the MWD and the ratio of the individual fractions. Several authors have reported on the molecular weight and width of MWD of polyethylene (PE) in relation to the cracking resistance and other properties of high- and low-density P E [1, 2]. In references [1, 3] a study is made of the properties of high-density polyethylene (HDPE) fractions, b u t the authors' conclusions are based on qualitative determination of changes in MWD (narrowing, broadening, etc.); as the MWD curves are generally not included in the papers, it is difficult to compare the results obtained b y different authors. In this paper we are reporting on MWD analysis of ten samples of H D P E together with a study of its effect on the mechanical behaviour of H D P E in the temperature interval --40-100 ° . EXPERIMENTAL

Study of MWD. The m e t h o d used was fractionation followed b y determination of the molecular weight (MW) of the fractions. Fractional precipitation was carried out at a t e m p e r a t u r e above the melting point of P E in solution [4] to obviate the effect of crystallinity. F r a c t i o n s with M W above 20,000 were obtained b y precipitation from a 1.5% solution * Vysokomol. soyed. A I 2 : No. 5, 1127-1137, 1970.

~276

A . N . KA~ASEV et al.

of the initial sample in tetralin-benzyl alcohol mixture (3 : 2) by lowering the temperature from 165" to 105 °. The low molecular weight residue t h a t remained was fractionated b y precipitation w i t h triethylene glycol from a 0"7-1.0~/o solution in toluene at 107 °. The fractionation was performed in the presence of stabilizer 22-46 in an argon atmosphere. E a c h sample was separated into 15-25 fractions which were vacuum-dried at 55 ° . The fractionation efficiency was estimated from the successive changes in the MW of the fractions over a wide range of values. I n determining the MW of P E difficulties arise both on account of the special nature of the polymer and also because of the high temperature at which the work must be carried out [5, 6]. The viscometric m e t h o d of MW determination was thought to be the most reliable, and we therefore used the light-scattering method to find the dependence of the intrinsic viscosity [t/] on MW, and to determine whether the samples under review were identical in respect to their degree of branching. Measurement of [g] in decalin at 135 ° was carried out with a viseometer fitted with a pendant level, with a shear stress of 9.7 g/cm.sec 2 in the presence of stabilizer 22-46 (not less than 0.1% on wt. of polymer). The Huggius constant k' in the concentration dependence of the specific viscosity was 0.3-0.45 for fractions of all t h e samples apart from the high molecular weight P E , and this shows t h a t no degradation of the polymer occurs during the viscosity measurements. Figure 1 shows plots of tls~/C vs c (~sp is the specific viscosity of a solution of concentration c) for fractions of one of the samples. The light-scattering m e t h o d at an angle of 90 ° was used to determine the weight-average molecular weight Mw of 5-10 fractions of each sample in the MW interval of 8000-470,000. The measurements were c a r r i e d out in monochromatic light with a wavelength of 546 rap, using a laboratory visual nephelometer, the recording portion of which is a PhM-58 phot o m et er graduated for benzene. ~-Chloronaphthalene was used as solvent at 135 °. The refractive index increment for the solutions was taken as 0.19 cm3/g [7]. Figure 2 shows plots of [ g ] = K M ~' on a logarithmic scale for the fractions under review. K and ~ are equal to 3-8 × 10 -4 and 0-74 dl/g respectively, and this is close to the data in reference [8] for linear P E . This dependence was used to find the viscosity-average molecular weights M , of the remaining fractions and of the initial samples. Using the fractionation data and the molecular weights of the fractions we p l o t t e d MWD curves and calculated the weight-average tl~w and the number-averagell//~ molecular weights of the initial samples and the N/w [M~ ratio characterizing their polydispersity. The fractionation of the high molecular weight polyethylene (HMPE, sample 10) and the process of determining the MW of its high molecular weight fractions gave rise to degradation due to the protracted dissolution of the polymer. The section of the MWD curves in the region of M > 103 was therefore obtained b y extrapolation based on the type of MWD and the requirement t h a t (]~wM~)l/~ ~ M ~ of the initial sample (w~, and M~ being the weight fraction and MW of the i t h fraction, respectively). Taking into account t h a t the experimental value of M r for this sample must be slightly underestimated both on account of the degradation and also on account of the shear stress in the measurement of [0] we cannot regard the MW data for this sample as sufficiently reliable. Study of behaviour of H D P E samples. The tensile testing was carried out in accordance with State Standard (GOST) 11262-65 using double-dumbells in a temperature interval of --40-100 °. The speed of the lower clamp of the apparatus was 50 ram/rain. The measuring error for the load and temperature did not exceed ± 3 ~o and q-3 ° respectively. The relative error in measuring the elongation was ± 10% for elongations exceeding 100%. The elongation was measured with marks at a distance of 25 rmn. The ultimate strength was calculated * I n the case of the high molecular weight P E (sample 10), from 180 °.

Relationship between mechanical behaviour and molecular weight distribution of PE

1277

both for the initial cross section (abr) and for the true cross section (atrue) US~rlg the formula /

~b,, "~

where ebr is the breaking elongation, ~o. The samples were cut from sheets 1 ram and 3 m m thick (at 100° the samples with a thickness of 3 m m were used) prepared by press moulding at 170-180 ° (samples 1, 3, 4, 5-9) and 200-210 ° (samples 2, 5, 10) under pressures of 100 and 200 kg/cm 2 respectively. The time kept under pressure was 5 rain, and the samples were cooled down to 60 at the rate of 25-35 deg/min. The heat resistance was determined with a head 6 rn~n in diameter penetrating to a depth of 1 m m under a load of 1 kg with the temperature rising at the rate of 3 deg/5 min [9]. The sample was in the form of a cylinder 7 nun high and 15 nun in diameter. The crack resistance in surfactant substances (v~0) was determined at 50 ° by flexing a strip of the material in a 20~o solution of OP-7 in accordance with State Standard (GOST) 13518-68. The flow melt index (FMI) was determined in accordance with State Standard 11645 at 190 °, load 5 kg. DISCUSSION OF RESULTS

T h e s u b s t a n c e s i n v e s t i g a t e d were t e n s a m p l e s o f H D P E w i t h different M W D p r e p a r e d a t low pressures, s a m p l e s 1-5 b e i n g e x p e r i m e n t a l grade, a n d s a m p l e s 6-10 commercial grade HDPE. In the preparation of samples 6-10 products of t h e i n t e r a c t i o n o f AI(C~Hs)~CI-~TiC14 were u s e d as c a t a l y s t a n d s a m p l e s 1 *, 3 a n d 4 t were p r e p a r e d w i t h t h e s a m e c a t a l y s t m o d i f i e d b y different a d d i t i v e s [10]. S a m p l e s 2 a n d 5 $ were p r e p a r e d in t h e soluble c a t a l y s t AI(C2Hs)~CI~-(CsHs)2TiC19 in a l k y l halide m e d i a [11]. F i g u r e 3 shows t h e i n t e g r a l M W D curves d r a w n on a s e m i l o g a r i t h m i c scale. S a m p l e s 1, 3, 4, 6 - 9 f o r m a g r o u p w i t h fairly similar M W D (within limits o f t h e s c a t t e r o f e x p e r i m e n t a l points) in t h e region o f M W o f 5000-60,000, a n d t h e y differ in t h e high m o l e c u l a r w e i g h t region. F o r e x a m p l e , in a region o f t h e c u r v e c o r r e s p o n d i n g to 50°,/o b y wt. t h e m o l e c u l a r w e i g h t s differ b y less t h a n a f a c t o r o f 2, while in t h e region o f t h e h i g h m o l e c u l a r w e i g h t p a r t o f t h e M W D , w h e r e t h e curves differ m o s t m a r k e d l y (90°/o b y wt.) t h e m o l e c u l a r w e i g h t s differ b y a f a c t o r o f 11.2. T h i s g r o u p o f s a m p l e s m a y t h e r e f o r e be u s e d to d e t e r m i n e t h e effect o f t h e h i g h m o l e c u l a r w e i g h t " t a i l " o f t h e M W I ) on m e c h a n i c a l b e h a v i o u r . T h e M W D c u r v e o f s a m p l e 10 is displaced t o w a r d s higher MW, b u t t h e s h a p e o f t h e c u r v e is similar t o t h o s e described above. S a m p l e s 2 a n d 5 h a v e a r e l a t i v e l y n a r r o w M W D , a n d h a v e v i r t u a l l y no molecules w i t h M W b e l o w 15,000 or a b o v e 300,000 a n d 600,000 r e s p e c t i v e l y . * The samples were synthesized by Ye. V. Veselovskaya (NIIPP). t V. K. Badayev (I~TIPP). In the laboratory of N. M. Chirkov (Inst. of Chem. and Physics).

1278

A.N.

~SEV

T Z T , E 1, I ~ O L E C U L A . R W E I G H T C H A R A C T E R I S T I C S O F

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HI)PE AND TH~ ~U~BER OF SIDE

GROUPS

AND DOUBLE BONDS

M × 10-8

Con~ent of O 0

0

C) C)

r-~

1 2 3 4 5 6 7 8 9 10

1.2 2.2 1.85 1.95 3-5 2.8 2.85 3.151 3.6 ~15

50 60 114 113 90 97 97' 110 2101 218 153 154 160 190 185 255 220 260 1500 1500

16 61 18 19 96 13 19 16 220

3"7 1"8 5"4 5"8 2"3 12 10 16 7

100 225 240 280 440 560 560 760 1120 3700

13"8 23"8 16"1 21"1 30"8 62 84 95 133

0 0 0 0 0 3 6 7 11 45

0-3 0.05 0.3 0.3 0.05 0.35 0.15 0.2 0.2 <0.1

0-7 0-1 0-6 0.6 0.1 0.7 0.6 0.7 0.65 0.05

* Data on the degreeofbranchingand unsaturatlonwereobtainedbyA. L. Gol'denbergand L. I. Zyuzina. t Mxand w~are the molecularweightand weightfractionof the first high molecularweightfractionrespectively. T h e characteristics o f t h e samples h a v e been t a b u l a t e d (see Table 1). Samples 2, 5 a n d 10 differ slightly f r o m t h e others in r e g a r d t o t h e i r c o n t e n t o f m e t h y l groups a n d double bonds. I t m a y be assumed, however, t h a t t h e differences in m e c h a n i c a l b e h a v i o u r t o be described below are caused n o t b y t h e v e r y slight s t r u c t u r a l differences in the P E molecules, b u t b y t h e u n d o u b t e d l y m o r e i m p o r t a n t difference in MWD. Table 1 also gives t h r e e q u a n t i t i e s characterizing t h e high molecular weight p a r t o f t h e MWD: t h e n u m b e r o f fractions h a v i n g M ~ 1 0 e, M10, which is t h e molecular weight value corresponding t o a n o r d i n a t e o f 9 0 ~ on the integral M W D curves (for 10% o f t h e f r a c t i o n o f highest mol.wt, o f t h e P E t h e M W is e q u a l t o M10 or higher); t h e p r o d u c t o f t h e molecular weight o f t h e first (high molecular weight) f r a c t i o n times t h e a m o u n t o f t h e latter. T a b l e 2 gives some o f t h e mechanical properties o f t h e samples. A comparison o f Tables 1 a n d 2 a n d Fig. 3 shows t h a t the d e n s i t y o f t h e samples a n d t h e deformat i o n corresponding t o the yield p o i n t a t r o o m t e m p e r a t u r e are o n l y w e a k l y d e p e n d e n t on MWD. T h e e x c e p t i o n is t h e H M P E (sample 10), the d e n s i t y o f which is m u c h lower t h a n t h a t o f the o t h e r samples. T h e h e a t resistance is slightly higher for samples 2 a n d 5 w i t h t h e n a r r o w M W D (138-140 °) a n d for the H M P E (141°), a n d it is lowest for t h e low molecular weight p o l y e t h y l e n e ( L M P E ) (sample 1,126°).

,

Relationship between mechanical behaviour and molecular weight distribution of P E

1279

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1280

the MWD for samples 1, 3, 4, 6-9, the higher crack resistance is probably due to the presence of 7-10~/o or more of fractions with MW of about l0 s (or higher) appearing in samples 7-9 (see Table 1, and Fig. 4). The absence (or low content) of these samples results in P E with relatively low crack resistance even in the case of samples with lower contents of low molecular weight fractions. For instance, samples 2 and 5 have practically no

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R e l a t i o n s h i p b e t w e e n m e c h a n i c a l b e h a v i o u r a n d molecu]ar w e i g h t d i s t r i b u t i o n of P E

1281

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samples a t - - 4 0 (a), 0 (b) 60 (c) a n d 100 ° (d).

!282

A.N. KAP~SEVet al,

is identical a rise in the content of low molecular weight fractions probably impairs the crack resistance, though no samples of this kind were available for testing. The marked rise in crack resistance observed in papers [2, 13] for samples with FMI in the region of 0.3-1.0 m a y therefore be due to the number of fractions with MW of about l0 s or more in the PE. Let us now determine the relationship between the MWD and mechanical behaviour in tensile tests at different temperatures. Figure 5 shows stress/strain curves for different samples at --40, 0, 60 and 100 °. The shape of the curves depends largely on the MWD of PE, and also on the test temperature. I n the case of samples with narrow MWD we find t h a t at low temperatures a second m a x i m u m appears on the stress/strain curves at elongations of around 1500~, after which the neck t h a t has appeared is stretched to the moment of failure, and in this process the stress is reduced (Fig. 5c, d). This additional stretching of the neck also occurs with the other samples with a long high molecular weight portion on the MWD curve (samples 7-9, with M10 ~500,000), but without the same smooth reduction in stress appearing on the curves. T A B L E 2 . S O M E CI~AI~ACTERISTICS OF H D P E /

Sample No.

%

Density, g/em S

0.956 0.954 0.951

0.951 0.950 0"950 0"950 0.955 0.950 10

0"937

FMI, g/10 rain 50 0.08 6.6 4.2 0.01 3-2 1.6

1.5 0.6 No flow

Heat

resistance, °C

SAMPLES

Deformation corresponding to yield point,

%

126 138 130 131 140 130 130 131 130

9.0 8.5 11.0 11.0 11-0 9.0 10.5 9"0 10.0

141

18-5

Crack resistance, hr Brittle 9-15 10 35 20-30 30-45 320-480 300-800 850-1000 and higher Over 10,000

* The limits of variation for several experiments are indicated. One reason for the scatter of results in respect to the crack resistance m a y well be insufficient care in maintaining the correct cooling rate in the moulding of the samples [12].

The tensile strength and the yield point for different temperatures are given in Table 3. Table 3 and Fig. 6 shows tensile strength as a function of temperature. At temperatures of --40-40 ° there is little difference in the yield points for P E with different MWD (apart from the HMPE, where the yield point is considerably lower t h a n t h a t for the other samples), and the values are in accordance with the density values given in Table 2 [3, 14]. At elevated temperatures the yield point is slightly higher for the samples with narrow MWD.

Relationship between mechanical behaviour and raolecular weight disCribution of PE 1283

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1284

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At temperatures below zero the highest tensile strengths and breaking elongations are observed with the H M P E (sample 10) and the samples with narrow M W D (2 and 5). On comparing samples 2 and 5 with 3, 4 and 6, we conclude t h a t at temperatures below zero the mechanical properties of H D P E are greatly impaired when the latter has up to 40% content of low molecular weight fractions with MW from 3000 to 30,000. e% 35O6 300~ 250t; 20C x

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FIG. 6. Elongation at break vs. temperature for HI)PE samples. The fractured curve is for the elongation calculated for an initial length of 25 rnm. At these temperatures the failure of samples 3, 4 and 6 takes place on the falling branch of the stress/strain curve, and the tensile strength is always lower than the yield point, and cannot always be determined. On increasing the content of high molecular weight fractions to 11% (samples 7, 9) and maintaining the number of low molecular weight fractions there is b u t a small improvement in mechanical behaviour. In the temperature region of 20-40 ° we similarly find that the H M P E and the P E with narrow MWD have the highest strength properties. Where the P E has a

Relationship between mechanical behaviour and molecular weight distribution of PE

1285

considerable content of high molecular weight fractions, as in the case of samples 8 and 9 with M10 >500,000, the tensile strength is higher than that of samples 1 a n d 3 which have shorter high molecular weight portions on the MWD curves. A t 60 ° increase in the length of the high molecular weight "tail" (increase in the value of M10) is accompanied b y reduction in the breaking elongation and improvement in tensile strength. The exceptions are the samples with narrow MWD retaining high values for strength and elongation with low values of M10. Gtr~,k,q/cm z 25O0 2000 1500 foo0 500 0

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:FIG. 7. Effect of the magnitude of Mlo on the true tensile strength at 100° for different HDPE samples. At 80 ° the strength is similarly markedly increased with rise in the value of Mlo, and in this case the P E with narrow MWD is no longer the exception. The change in elongation accompanying change in MWD is slightly smaller than at 60 °. Note that the elongation of the L M P E is high, and close to that of P E with narrow I~IWD, while H M P E has a relatively low elongation. At 100 ° the tensile strength is greatly increased with increase in the value of M10 for all the samples, while the elongation, in contrast to that at 60 °, also rises monotonically and reaches 2170% for sample 9, b u t is again reduced for ItMPE. In the temperature interval 60-100 ° a maximum is seen on curves of breaking elongation versus temperature. The position of this maximum on the temperature axis depends on the content of high molecular weight fractions. A rise in the value of M10 causes displacement of this maximum towards high temperatures. In particular, the maximum for samples 9 and 10 is probably found at a temperature above 100 °. The highest peak value is observed for the elongation of P E with narrow ~IWD (about 3400% ), while H M P E has the lowest elongation in the temperature interval 60-80 ° . It appears therefore that the simultaneous absence of low molecular weight (M~15,000-20,000) and high molecular weight fractions (with M >300,000) in t t D P E ensures the highest value for elongation at the peak on the temperature curve. A marked reduction in elongation accompanying a rise in temperature is characteristic for P E samples with Mlo values below 500,000.

1286

A.N. KARASEVet al.

At temperatures of 60-100 ° the failure of the samples always takes place in the region of the stretched portion of the sample (the neck), and this enables us to find the true tensile strength, atrue. At 60 ° the P E with narrow MWD has the highest true strength, and LMPE has the lowest. As the temperature rises from 60 to 100 ° the true strength is reduced, with the exception of HMPE, for which the true strength remains constant. As the value of M10 rises, the reduction in true strength with rise in temperature becomes smaller. For instance, there was reduction by a factor of 5 in the true strength of LMPE, of 2.2 for sample 9, and no change at all in the strength of HMPE. In the interval 80-160 ° the true strength rises with increase in the value of M~0: at 100 ° there is an 8-fold increase (Table 3). As is seen from Fig. 7, the change in the true strength (in kg/cm 2) at 100 ° with rise in M10 (from 100,000 to 3.7 × l0 s) is expressed by the simple relationship: O'true~----1"65 M~/d-- 165 There are apparently at least two factors underlying the regular rise in true tensile strength at high temperatures accompanying increase in the MW o f high molecular weight fractions. During the crystallization of the polymer from the melt all the supermolecular structural elements, i.e. spherulites, rays and crystallites* are more or less interconnected by "passage" molecules. Moreover, there are always a certain number of physical crosslinks (enmeshed molecules) in the polymer. The former and the latter are obviously formed through long molecules in the polymer. In the stretching of a crystalline polymer of spherulitie structure the boundaries between the spherulites are preserved, as well as the mutual ordering of the small supermolecular structural elements forming the spherulite, though the spherulite itself is transformed into a fibrous formation [15, 16]. The passage molecules present in the original crystalline polymer, and the physical bonds are preserved even in the fibrils of the stretched PE, and these connect the individual crystalline fibrillar segments and determine their strength. As the molecular weight of the high molecular weight fractions rises, t h e number of these passage molecules and the number of supermolecular structural elements connected by them also increases, as well as the number of physical crosslinks. This is the factor underlying the higher tensile strength at high temperatures, and the improvement in the crack resistance of P E observed with rise in the value of Mlo. At lower temperatures the mobility of the molecules is reduced, and the role o f the intermolecular forces as a strength-determining factor becomes more important. Moreover, stable physical crosslinks m a y now be ~ormed by the shorter molecules also. This is reflected in the rise in true tensile strength with reduction * Using the terminology suggested in reference [15].

Relationship between mechanical behaviour and molecular weight distribution of PE

1287

in temperature (the rise is particularly marked in the case of LMPE) and in a lessening of the difference in the strength of samples with different MWD: From these considerations it is obvious that the difference in MWD in P E samples with identical viscosity-average or weight-average MW's (samples 2 and 4; 5 and 7-8) will cause differences in mechanical behaviour. The effect of the individual fractions in P E m a y vary according to the temperature at which the mechanical properties are investigated. When the characteristics selected for studying P E are limited to either the intrinsic viscosity or the average MW, errors m a y arise in elucidating the interrelation between the properties of samples prepared b y different methods and differing in respect to their MWD. In the case of P E these characteristics are based on averaging of the MW of fractions differing b y a factor of several hundreds or even thousands. Moreover, as we have shown, the presence and the amounts of particular fractions in the composition of P E determine its mechanical behaviour, and this m a y not be reflected sufficiently well in the experimentally determined average molecular weights. This applies also to the MFI, which is frequently used to characterize PE. Similarly the good tensile strength values obtained at room temperature cannot always be taken as a guarantee that the polymer will have good tensile strength and high crack resistance at elevated temperatures. CONCLUSIONS

(1) I t has been found that the good strength properties of polyethylene (PE) in tensile tests at high (80-100 °) temperatures depend to a considerable extent on the high molecular weight part of the molecular weight distribution (MWD), and to a lesser extent on the amount of low molecular weight fractions. At sub-zero temperatures not far from room temperature the mechanical behaviour of P E is largely determined b y its content of low molecular weight fractions; a reduction in the latter is reflected in marked improvement in the mechanical properties of the polymer. High molecular weight P E with M w = 1-1.5 × l0 s has the best mechanical properties over a wide temperature interval. (2) P E with narrow MWD (with simultaneous absence of low- and high molecular weight fractions) has the best mechanical properties at sub-zero temperatures and higher extensions at elevated temperatures (maximal elongation at break 3400%) compared with P E having wide MWD b u t similar weight-average MW, though the viscosity of the melt for the latter is much lower. (3) It has been shown that the crack resistance of P E increases markedly with increase in the amount of high molecular weight fractions with an MW of about l0 s or more. The absence of these fractions is reflected in relatively poor crack resistance whether the P E has a wide or narrow MWD. Trans/at~ by R. J. A. I-I~ND~Y

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H. S. KOLES~r~OV and A. YA. VAINER REFERENCES

1. A. S. KENYON, I. O. SALYER, J. E. K U R Z and D. R. BROWN, J. P o l y m e r Sei. C8: 205, 1965 2. Konstruktsionnye svoistva plastmass (Engineering Properties of Plastics) (Ed. b y E. Beer), Ch. 6, Izd. " K b l m l y a " , 1967 3. G. R. WILTJAMSON, B. W R I G H T a n d R. W. HOWARD, J. Appl. Chem. 14: 131, 1964, 4. L. MANDEL'KERN, Kristal. polim. (Crystallization of Polymers). p. 56, Izd. " K h i m i y a " 1968 5. V. KOKLE, F. W. BILLMEYER, Jr. L. T. MUUS and E. J. NEWITT, J. Polymer Sei. 62: 251, 1962 6. H. P. SCHREIBER and M. A. WALDI~_N, J. Polymer Sci. A2: 1655, 1964 7. M. B. HUGLIN, J . Appl. Polymer Sci. 9: 3963, 1965 8. M. O. De La CUESTA and F. W. BILLMEYER Jr., J. Polymer Sci. A I : 1721, 1963 9. I. A. MAIGEL'DINOV, Trans. V I I All-Union Conf. on H i g h Melee. Weight Compounds, Izd. A N SSSR, 1952 10. V. K. BADAYEV, V. P. MARDYKIN and Z. V. ARI~HIPOVA, Plast. massy, No. 12, 6, 1965 11. G. P. BELOV, A. P. LISITSKAYA, N. M. CHIRKOV and V. L TSVETKOVA, Vysokomol. soyed. A9: 1269, 1967 (Translated in P o l y m e r Sci. U.S.S.R. 9: 6, 1417, 1967) 12. J. B. H O W A R D and W. M. MARTIN, SPE Journal 16: 407, 1960 13. L. LANDER a n d R. CAREY, Materie Plast. ed Elastomerie 29: 496, 1963 14. A. R E N F R E W a n d Ph. MORGAN, Polythene. Technology a n d Uses of Ethylene Polymers, L o n d o n - N e w York, 1960 15. F. H. JEIL, Polim. monokrist. (Polymeric Monoerystals). Izd. " K h i m i y a " , 1968 16. V. A. KARGIN and G. L. SLONIMSKII, K r a t . ocherki po fiziko-khim, polim. (Short. Notes on the Physicochemistry of Polymers). p. 170, Izd. " K h i m i y a " , 1967

KINETICS AND MECHANISM OF THE THERMAL DEHYDROBROMINATION OF BROMINATED POLYPHENYLENEETHYL IN SOLUTION* H. S. KOLES~TIKOV(dee.) and A. YA. V~r~ER D. I. Mendeleyev Chemicotechnological Institute, Moscow

(Received 21 April 1969) I ~ an earlier communication [1] we reported on the preparation of brominated polyphenyleneethyl (BPPE) as well as the results of a study of the solvolysis of B P P E in dimethylformamide. I t was found t h a t even at 70 ° an almost completely dehydrobrominated polymer was obtained by heating B P P E in DMFA in the * Vysokomol. soyed. A12: No. 5, 1138-1150, 1970.