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The intensities and lifetimes of positrons annihilating in low and high density poly(ethylene) (LDPE and HDPE) were measured. As expected the orthopositronium (oPs) pick-off lifetime and intensity increases with temperature; the relative values of the intensity are consistent with differential scanning calorimetry (DSC) results. Two other components with shorter lifetimes were deconvoluted from the complex exponential decay profiles. Detailed studies were carried out on quenched and annealed samples of LDPE and HDPE in an attempt to correlate the two shorter components with crystallization; both CsF and plastic scintillators were used. DSC analyses were also carried out. We conclude that the shortest lifetime component is caused by a combination of parapositronium (pPs) decay and the annihilation of free positrons in the crystalline regions and that the second longest lived component can also be associated with crystalline regions especially the crystalline/amorphous boundary.

In this report, we present the results of our investigation into the crystallization behaviour of polyproplene. The crystallization process was followed by hot-stage optical microscopy, differential scanning calorimetry, and dynamic mechanical methods. As well, nuclear magnetic resonance spin-spin relaxation techniques were employed to probe the morphology of the polypropylene. Nucleation in this isotactic polypropylene was athermal and appeared to be heterogeneous. The dependence of the onset temperature of crystallization on the annealing temperature and on the annealing time indicated that PP nuclei were able to survive the melting process. Nuclear magnetic resonance spin-spin relaxation experiments indicated that the PP melt contained a significant proportion of regions of high segment density. It was postulated that these ‘ordered’ regions acted as nucleation sites for PP crystallization, and that the number and size of these regions was determined by the annealing time and temperature. Addition of HDPE to PP resulted in melting-point depression and plasticization of the PP phase at lower HDPE contents. HDPE was able to penetrate the PP phase sufficiently at lower HDPE contents to reduce the number and size of regions of high segment density, thereby delaying the nucleation and subsequent crystallization of the PP phase.

Studies on the crystallization of polymers are usually carried out by dilatometry, differential scanning calorimetry (d.s.c.) and optical microscopy. These methods have been widely used in the analysis of crystallization kinetics using the Avrami equation. Direct determination of nucleation rate and density is only possible by the optical microscope technique. We have employed a new technique for studying the crystallization behaviour of polymers by using a dynamic mechanical method. This method is capable of detecting the onset of nucleation. Semiquantitative data on the nucleation density and the initial crystallization rate may also be deduced by this method. Results from the dynamic mechanical method are cross-compared with d.s.c. and optical microscope data for polypropylene, polyethylene and their blends, and are shown to be mutually consistent.

In size exclusion chromatography (s.e.c.) analysis both the universal calibration curve (UCC) and the secondary calibration curve (molecular weight calibration) can be constructed using the Mark-Houwink-Sakurada (MHS) relation, [η] = KMa. Over a very wide range of molecular weights neither K nor a is constant; therefore, the interpretation of results using one set of K and a values may not be correct (i.e. in the case of commercial resins). An alternative model which relates molecular weight (MW) to [η] over a very wide range of molecular weights is [η] = K π M r 1 2 + K′ M w were Kθ and K′ do not change. The intrinsic viscosities of polystyrene standards with MWs of 103 to 2 × 106 were measured and the results were used to establish a UCC for a multi-detector s.e.c. system. From the UCC two calibration curves were established using the MHS relation and the alternative model. Commercial polystyrene and polyethylene resins were analysed using the same analytical system. At the higher molecular weight end one observes a deviation between the results from absolute MW detectors and the calibration curve based on a particular set of K and a values. On the other hand, MWs obtained from the calibration curve using the alternative model were consistent with results from absolute molecular weight detectors. This indicates that the use of the alternative model will provide a more accurate calibration for s.e.c. analysis when the molecular weight distribution is relatively wide.

Molecular-weight distributions (MWDs) of polyolefins are currently measured by high-temperature size exclusion chromatography. The raw data are converted to MWDs using universal calibration with a differential refractive index (d.r.i.) concentration detector or by direct molecular-weight measurements with a continuous viscometer (c.v.) or a low-angle laser light scattering (l.a.l.l.s.) detector. None of these techniques provides a picture of the true distribution. Simple d.r.i.-universal calibration may fail to sense the high-molecular-weight tail of the distribution, which is important in some applications. The l.a.l.l.s. technique is most sensitive to high-molecular-weight polymers but is also least precise. The c.v. method is presently the most convenient single technique. The most information is obtained, however, when all these detectors are employed. Both the l.a.l.l.s. and c.v., for example, should be used to verify that the polymer sample is actually dissolved, since dissolution of semicrystalline polyolefins is not a trivial exercise.

The general theory of Stockmayer and Benoit and their coworkers yields the weight average molecular weight of a binary copolymer and parameters which reflect the heterogeneity of chemical composition of the copolymer, using light scattering data from solutions in at least three solvents with different refractive indices. Although this theory has been applied successfully to block and graft copolymers and to mixtures of homopolymers, it has not seemed to be valid for random copolymers. In this report, molecular weight and heterogeneity parameters of a number of gel-free styrene-butadiene emulsion copolymers were estimated from results of light scattering measurements in toluene, cyclohexane and 1,2-dichloroethane. The calculated heterogeneity parameters contradict expectations from copolymerization theory and appear to be in error. It is suggested that the particular difficulties with statistical copolymers result from a dependence of specific refractive index increment on polymer molecular weight. This dependence varies with different solvents. The effect is illustrated by experimental results for polystyrenes in toluene, cyclohexane and 1,2-dichloroethane. Accurate study of copolymer characteristics requires selection of solvents such that the differences between specific refractive indices of the appropriate homopolymer solutions do not depend strongly on polymer molecular weight in the molecular weight range which corresponds to the sum of homopolymer sequences in the particular copolymers. The apparent inapplicability of the theory to random copolymers appears to be an artifact of the particular systems which were studied.

This article reports the synthesis and free radical polymerization of ortho-vinylbenzophenone. The glass transition temperature Tg of the homopolymer is 136°C. The products synthesized appeared to be atactic and amorphous. The Mark-Houwink constants for poly (o-vinylbenzophenone) in tetrahydrofuran are K = 4.2 × 10−2 cm3 g−1 and a = 0.765. The pre-exponential constant under theta conditions, Kθ, is estimated to be 5.93 × 10−2 cm3 g−1. The ratio of unperturbed dimensions of the actual polymer and free rotating analogue chain is 3.93, which is almost double that of polystyrene. The Flory-Huggins interaction parameter for poly (o- vinylbenzophenone ) tetrahydrofuran is 0.48 at room temperature. The k p k 1 2 t ratio at 60°C is 1.1 × 10 −2 l 1 2 mol −1 2 s −1 2 . In free radical copolymerizations with styrene at 70°C, r1 (o-vinylbenzophenone) = 1.216, r2 = 0.751. This copolymerizations is virtually random.

Specific volumes of anionic polystyrenes prepared by butyl lithium initiation have been measured at temperatures between 110° and 237°C. Molecular weights of the fractions studied ranged from 2500 to 700 000. Plots of isothermal specific volume against reciprocal molecular weight ( M −1 ) are fitted by two straight lines intersecting at a molecular weight near 10 000. This parallels the behaviour of thermally initiated polystyrene fractions, which have different end-groups. Volumes attributable to end-groups are greater for the anionic polymers. The specific volume-molecular weight-temperature relations in the M ≤ 10 000 region are explained quite well by the equation of Francois and coworkers, in terms of end-group volumes and a uniform volume contribution from internal repeating units. Restraints on the motion of the butyl end-groups in these polymers by the attached polystyrene chain appear to be temperature dependent in terms of this model. Specific volumes of polymers with M ≥ 10 000 are affected by chain ends and by variation of the volumes of non-end repeating units with molecular weight. The present results parallel those which have been reported for partial specific volumes of polystyrenes in dilute solution and suggest that the same factors operate in both cases.

Reactivity ratios can be determined from n.m.r. analyses of binary copolymers. Measurement of the diad or triad sequence distribution provides an estimate of the number average sequence length of each copolymer. Each average sequence length is directly related to the corresponding reactivity ratio in the terminal copolymerization model. Extension to more complicated reaction models is straightforward. A single copolymerization experiment yields values for both reactivity ratios with this procedure. When pentads or triads, centred on one monomer only, are measured the two reactivity ratios cannot normally be estimated unless the copolymer is highly alternating. It is possible, however, to combine such data with an n.m.r. or other analysis of the chemical composition of the copolymer to obtain r 1 and r 2 . When the sequence distribution of runs, centred on only one of the monomers, can be analysed and the chemical composition of the copolymer is not available the procedure can yield a good value for the reactivity ratio of that monomer and a poor estimate for that of the comonomer.

Commercial poly(vinyl alcohol—acetate) copolymers were characterized by 13 C n.m.r. spectroscopy in D 2 O at 100.6 MHz and by differential scanning calorimetric measurements. N.m.r. spectra can be used for quantitative measurements of tacticity, vinyl alcohol—vinyl acetate composition, non-hydrolyzable branch content and blockiness of vinyl acetate residues. Conventional quality control tests for poly(vinyl alcohols) are recognized to be inadequate to account for all the important properties of these polymers. They have been augmented by thermal analyses for estimations of blockiness of alcohol and acetate residues. The crystal melting points of commercial poly(vinyl alcohols) vary significantly with thermal history of the samples. A standardized thermal procedure was used with observed melting temperatures to estimate either the degree of hydrolysis or the overall blockiness of the polymer. The equations used are based on theory but the procedures used industrially should best be regarded as semi-empirical because equilibrium melting points are not normally measured.