An investigation of the effect of sample preparation on the thermal behavior and the small-angle...

An Investigation of the Effect of Sample Preparation on the Thermal Behavior and the

Small-Angle X-Ray Scattering of a Polyurethane Block Copolymer

JOHN A. MILLER and STUART L. COOPER, Department of Chemical Engineering, University of Wisconsin, Madison, Wisconsin 53706

SYIlOpSb

The effect of varying sample preparation parameters on the thermal behavior and on the small-angle x-ray scattering (SAXS) properties of a polyether polyurethane were investigated. The polyurethane studied was a methylene bis(p-phenyl isocyanate) (MDI)/butanediol/ polfltetramethylene oxide) (PTMO) system synthesized in a 6/5/1 mole ratio by a twestep solution polymerization. The F'TMO had a nominal molecular weight of 2000. The samples were compression molded under different conditions for the SAXS experiments. The preparation parameters studied included molding time and temperature, sample thickness, and quenching rate from the molding temperature. The molding temperature has the greatest effect on the SAXS data. In this case the domain size was observed to increase as the molding temperature increased from 130 to 2oo'C. The thermal properties were also found to be strongly dependent on the molding temperature, as measured by differential scanning calorimetry (DSC). An endotherm related to the annealing that occurs during the molding process appears in each sample near the molding temperature. The other preparation parameters have little or no effect on the SAXS and thermal properties of this sample.

INTRODUCTION Much attention has been given by polymer scientists to the relationships

between chemical structure, morphology, and physical properties of mul- tiphase polymer materials. This is especially true for polyurethane elastomers, a block copolymer of the (AB), type.14 These materials consist of alternating hard, rigid blocks and soft, flexible units. The hard segments are typically composed of aromatic or aliphatic diisocyanates linked by short diols or diamines. The soft segment is usually a polyester or polyether diol, although recently hydrocarbon polyols have been used. In the solid state, these components separate into microdomains consisting primarily of either soft segments or hard segments, with some degree of intersegmental mixing. The unique mechanical properties of these thermoplastic rubbers are due to the crosslinking and filler effect that the hard domains impart to the rubbery phase which serves as the matrix in the more elastomeric com- positions. One of the more recently developed tools used to investigate polyurethane morphology is small-angle x-ray scattering (SAXS). This pow- erful technique probes the microstructure with a high resolution; very small features can be detected by SAXS. However, the technique does not present a direct image of the morphology as does electron microscopy. General parametric analysis of the data is normally used, although some success

Journal of Polymer Science: Polymer Physics Edition, Vol. 23, 1065-1077 (1985) @ 1985 John Wiley & Sons, Inc. CCC 0021-8995/85/05106513$04.00

1066 MILLER AND COOPER

has been achieved through modeling. Nonetheless, SAXS has found broad application in structure-property studies of urethane systems.

Bonart5 and Clough and Schneidel.6 were the first to use x-ray scattering in the study of polyurethane systems. Bonart reported an "accumulation of the hard segments" in these materials which was attributed to an ex- tensive hydrogen bonding network. At about the same time, Clough and Schneider found that the SAXS intensity in a series of polyurethanes changed in an orderly fashion as the composition of the materials changed. This was attributed to the degree of aggregation of the hard segments, referred to now as the degree of phase separation. These two studies pro- vided the earliest x-ray evidence for microdomain formation in polyurethane systems. About a year earlier, Cooper and Tobolsky7 inferred microdomain formation in polyurethane systems based on an analysis of the viscoelastic behavior of the materials.

Subsequently Bonart et aL8 studied the effects of annealing and elongation on polyurethane crystallinity using SAXS. Bonart reported an increase in the hard-segment crystallinity with annealing temperature and time. Van Bogart et al.9 also studied annealing effects in polyurethanes using differential scanning calorimetry (DSC). The results suggest a rearrangement of the hard domains to produce a higher degree of local order. Van Bogart et al.1° also investigated a series of polycaprolactone polyurethanes. The samples exhibited a strong dependence of hard-segment crystallinity on thermal history. Clearly some ordering process is occurring in polyurethanes sub- jected to thermal treatment. It appears that the morphological changes that occur during the thermal treatment, including those that accompany the sample preparation, deserve further study.

A number of other workers, notably K~bersteinll- '~ and Wilkes,14" have used SAXS to characterize the morphology of polyurethane systems.1E22 In general, it has been found that the microstructure of polyurethane materials consists of a lamellar or pseudolamellar composite of the hard and soft phases, with an average size on the order of 10 nm. Between the domains is an interfacial region that can be quite sharp or range up to a few na- nometers. Also, the phases in these materials are not pure. A significant fraction of the hard segments are dissolved in or reside within the soft phase. Some mixing of soft segments into the hard domains occurs as well, although probably to a lesser extent.

From the work of Bonart8J9 and Van B ~ g a r t ~ J ~ . ~ it is obvious that thermal treatment and sample preparation can alter the polyurethane microstructure quite markedly. The objective of this study was to investigate the effects of thermal history on the morphology of a polyurethane material. In particular, the thermal history associated with the sample preparation by compression molding was studied, as opposed to longer-term annealing or heat treatment. In order to properly study these effects, a wellcharacterized material was needed. The material chosen was a solution-polymerized polyether polyurethane, described in detail below. The solution polymerization method was chosen to provide a homogeneous medium for the reaction. Recent work by Thomas and co-workers".25 suggests that premature separation of the reactants commonly occurs during the middle stages of bulk polymerization of polyurethanes, thus leading to undesirable chemical het-

THERMAL BEHAVIOR AND SAXS OF POLYURETHANE 1067

erogeneity and subsequent complex polymer morphology. In the present study, the sample preparation parameters studied included molding temperature and time, cooling rate, and sample thickness. The effect of sample thickness will give information on the surface morphology of the polymer relative to the internal bulk morphology. An additional sample was prepared by solution casting to compare this preparation method with compression molding. The morphological changes were measured by small-angle x- ray scattering. Differential scanning calorimetry was also used to monitor changes in structure that gave rise to features in the DSC thermograms.

EXPERIMENTAL

The starting material for this study was prepared via a two-step solution polymerization in dimethyl acetamide (DMAc). The polymer consisted of a 6/51 1 mole ratio of methylene bis(p-phenyl isocyanate) (MDI), butanediol, and poly(tetramethy1ene oxide) (PTMO). The MDI (Eastman Kodak) was purified by melt filtration followed by recrystallization in hexane. The butanediol (Aldrich) and dimethyl acetamide (Aldrich) were vacuum distilled prior to use. The PTMO (DuPont Teracol2000) was dried at 60°C in a vacuum oven for 48 h. A stannous octoate catalyst (M&T Chemicals catalyst T-9) was used as received to promote the reaction.

The MDI was reacted with PTMO for one hour a t 60°C under nitrogen in a 30-w % DMAc solution using 0.15% catalyst. Subsequently the butanediol chain extender was added and the system was allowed to react for an additional three hours. The polymer was then precipitated with water and dried in a vacuum oven at 60°C for one week. The starting material was stored in a desiccator prior to the final sample molding.

The samples used for the SAXS tests were prepared from the starting material described above. Table I lists the samples by group, showing the preparation parameter that was varied. The samples were compression molded into 3-cm diam disks by heating to the molding temperature, a process that required about ten minutes. After being held at the molding temperature for a specified length of time at a pressure of 30 MPa (ca. 4000 psi), the samples were cooled at various rates to room temperature. Natural convection required roughly twenty minutes for the sample and mold to reach room temperature. Forced air convection cooled the sample and mold in about 5 min. Approximately one minute was required when quenching the sample and mold in water at 30"C, while the quench in the ice bath and the quench in the dry ice/acetone bath took even less time. The SAXS experiments were performed within 48 h of the final cooling step, with the samples being kept under desiccation at room temperature until tested.

The thermal properties of the samples were analyzed by differential scanning calorimetry (DSC). The thermal data were obtained shortly after the collection of the SAXS data in order to minimize any room-temperature storage effects. The thermal data were obtained with a Perkin-Elmer DSC I1 interfaced to a TADS data station. Baseline correction and normalization are automatically carried out by this system.

The SAXS data were obtained at the University of Wisconsin Materials Science SAXS facility. The Cu Ka x-rays were generated by an Elliot GX-


TABLE I Sample Preparation Conditions

Molding Molding Sample Weight Temperature time number (s) (T) (mid Cooling method

9 10 11 12

13 14 15 16 17

18 19 20 21

1.5 1.5 1.5 1.5 1.5 1.5 1.5 1.5

2.0 2.0 2.0 2.0

1.0 1 .o 1 .o 1 .o 1 .o

0.5 1.0 1.5 2.0

Variation in molding tempemturn 130 5 140 5 150 5 160 5 170 5 180 5 190 5 200 5

Variation in molding time 170 0 170 5 170 10 170 20

Variation in cooling mte 170 5 170 5 170 5 170 5 170 5

Variation in sample thickness 170 5 170 5 170 5 170 5

Natural convection Natural convection Natural convection Natural convection Natural convection Natural convection Natural convection Natural convection

Forced convection Forced convection Forced convection Forced convection

Natural convection Forced convection Quench in 30°C H20 Quench in ice/water Quench in C02/acetone

Quench in 30°C H20 Quench in 30% H20 Quench in 3072 H20 Quench in 30% H20

22 Sample spin cast at 70°C from 30% DMA solution

21 rotating-anode system operating at a 40-kW accelerating voltage and 50- mA filament current. The x-rays were filtered by a nickel foil and further processed by means of electronic energy discrimination. The latter is done automatically with the use of the TEC model 211 linear position-sensitive x-ray detector and the associated detector electronics. The energy distri- bution of the x-rays is thus quite narrow, consisting almost entirely of the 1.54-8, Cu Ka wavelength. The camera used was a modified Kratky compact camera. The data were desmeared using the iterative technique of Lakem after correction for detector sensitivity and dark current. The data were converted into absolute units by reference to a calibrated Lupolen polyeth- ylene standard.

SAXS DATA ANALYSIS

As the analysis of the SAXS data is crucial in this investigation, the methods used to interpret the data will be described in some detail. Three morphological parameters are of interest. First is the long spacing, which is a rough measure of the distance between like domains. Second is the inhomogeneity length, which is a measure of the average size of all domains. Finally, the degree of phase separation was calculated. This is a relative


parameter based on the electron density differences of the two phases. The interfacial thickness, commonly used as another measure of polyurethane microstructure, was found to be fairly small for the commonly used interfacial profiles, and hence will not be discussed in detail.

The absolute intensity of small-angle scattering from a sample can be written asn

where q is the magnitude of the scattering vector and is equal to 41r/ A sin(28), where A is the incident wavelength and 8 is the scattering angle; I,(q) represents the background scattering due to thermal density fluctua- tions within the phases and &(q) is the scattering due to the morphology and structure of the material. Several expressions for the background term have appeared, and these have recently been reviewed.% The model of VonkB was used to fit the background in this study:

The value of n was fixed at 2 for the data analysis. An accurate determi- nation of the background scattering is necessary to determine the width of the interfacial zone between the phases. The problems associated with this and the errors involved are the subject of a detailed report by Roe.30 In the present study, the model to which the background was fitted and the range over which the fit was performed had little effect on the measured value of the interface; the interface was found to be quite sharp for all profile shapes. One would not expect this to be the case for materials with a broad interfacial zone. For the final analysis the background was fitted using eq. (2) in the high-angle region of the scattering curve, q > 4 nm-I.

Porod3I derived a n expression for scattering from a two-phase system with sharp interfaces in the asymptotic limit as q becomes large:

K is Porod’s constant, and eq. (3) is generally referred to as Porod’s law. R ~ l a n d ~ ~ discusses the limits of validity of Porod’s law in lamellar and pseudolamellar systems. Applying Ruland’s criteria to polyurethane systems, KobersteinI3 states that the region over which PorOd’s law is valid in polyurethanes extends upward in q from twice the value of q at the peak value of the Lorentzcorrected intensity. The Lorentz correction is made by multiplying I (q ) by q2. Application of Bragg’s law to the q value at the peak in q21(q) yields the long spacing Is. Equation (3) can be modified to account for a diffuse interface. This expression is then

H ( q ) is the Fourier transform of the electron density profile across the


interface. A number of different profiles have been proposed and discussed for the interfa~e,~ ' all of which yielded a small value for the interfacial thickness in these samples. Equation (4) was then fitted by using a nonlinear regression routine to the scattering data from twice q p k to the limit of the available high-angle data.

The electron density variance of a two-phase system with sharp boundaries and pure, homogeneous phases is given by

- Ap2 = - 41,) (pH - psY (5)

where is the volume fraction of one of the phases, and pH and ps are the electron densities of each phase. The electron density variance of the actual sample can be determined from the SAXS data by

This expression takes into account the effects of diffuse boundaries and phase mixing. An additional variance parameter that can be used to characterize two-phase systems such as polyurethanes eliminates the effect of a diffuse boundary, but still reflects the presence of phase mixing. This expression is given by

For these materials there is a sharp boundary, hence bp2' = Gy. In calculating the integrals in eqs. (6) and (71, an extrapolation must be

made to q = 0 and q -, 00. The extrapolation to q = 0 was done by assuming q21,(q) approaches zero as q approaches zero. Extrapolation at high values of q can be done using P o d ' s law directly since the samples exhibited a sharp interface. The high-angle extrapolation contributes about 5% to the total integral, while the low-angle extrapolation contributes around 1%. Thus any errors associated with the extrapolation are minimized because of their small contribution.

The parameter for degree of overall phase separation (DPS) was first introduced for polyurethane systems by Bonart and Mulle93-34 as the ratio of the measured to ideal electron density variance:

The value of this parameter is unity for a material that possesses pure phases and sharp boundaries. It decreases as interphase mixing increases and as the interface broadens. To calculate Ap&,,, knowledge of the density of each phase is required. The values used were 1.37 g/cm3 for the hard phase and 0.977 g/cm3 for the soft phase.23 The value of DPS is highly dependent on the accuracy of these numbers. However, errors in these values will only affect the absolute value of DPS; the trends should remain the same.


The inhomogeneity length is a parameter describing the average domain size. It can be visualized as the average length of all possible chords that reside within a particular domain.= PorOd’s inhomogeneity length I, is a harmonic average of the inhomogeneity length of each phase:

where 1, is the inhomogeneity length of phase i and 4, is the volume fraction of that phase. It can be shown that I , is accessible from SAXS data using the relations

The inhomogeneity length can be used as a rough measure of the domain size. The long spacing can be used as a measure of the spacing between domains. We thus have two size parameters with which we can characterize the polymer microstructure.

RESULTS AND DISCUSSION DSC thermograms for samples 1-8 are shown in Figure 1. The glass

transition temperature of the soft segment is about -68°C for each of the samples. An endotherm is seen in each of the traces at about 70°C. This endotherm is the result of annealing at 60°C in the vacuum oven while the starting material was being dried. Annealing endotherms have been observed by Van B0gart~JO9~ and otherss9 in a wide variety of polyurethanes. The annealing endotherm is typically seen at a temperature between 10- 50°C above the annealing temperature.

A prominent feature of the thermograms in Figure 1 is a high-temperature endotherm that appears close to the temperature at which the samples were molded. In effect, the molding is a short-time annealing of the sample. KweP studied the effect of annealing time on the position of the annealing

-100 -50 0 50 100 150 200 TEMPERATURE, Celsius

Fig. 1. DSC thermograms for samples 1-8.


peak relative to the annealing temperature. Kwei found that the annealing endotherm got closer to the annealing temperature as the annealing time was reduced. This is because short annealing times do not allow for the full development of the ordering that accompanies the annealing process. The results obtained in this study are consistent with Kwei's data and analysis. Thermograms were obtained for the other samples as well. No significant differences were seen among the other samples. This implies that over the range of temperatures, times, and quenching rates studied only the molding temperature affects the thermal characteristics of the polymer.

In the DSC traces of samples 6-8, an endotherm is observed at about 10°C. This corresponds to the melting endotherm of the polyether soft- segment crystallites. The endotherm is seen to increase in magnitude as the molding temperature increases. This is due to the increased mobility of the hard segments in the chain above molding temperature of 170"C, which allows the entire chain to relax to a lowerenergy conformation than that found in the sample as it is first precipitated from the DMAc solution with water. This relaxation encourages the formation of soft-segment crystallites as the sample is being cooled in the DSC cell, which in turn is

TABLE I1 SAXS Data Analysis Summary

Pord Degree Long inhomogeneity of phase

Sample spacing length separation number (nm) (nm) (%)

Variation in molding tempemturn 1 12.8 3.69 35 2 12.9 3.66 36 3 12.6 3.77 36 4 13.6 3.85 35 5 13.7 3.96 36 6 14.3 4.01 35 7 15.0 3.99 35 8 16.0 4.16 34

Variation in molding time 9 13.6 3.91 34 10 14.0 3.96 34 11 13.7 3.87 35 12 14.0 3.98 35

13 13.4 3.99 31 14 13.3 3.87 32 15 13.9 3.83 37 16 13.3 3.69 32 17 13.3 3.63 31

18 13.1 3.77 36 19 13.5 3.68 36 20 13.8 3.57 34 21 14.0 3.83 37

22 13.4 3.57 30

Variation in cooling mte

Variation in sample thickness

Solutioncclst sample


partially responsible for the increase in the long spacing seen in the SAXS data, which are summarized in Table 11.

The formation of the soft-segment crystallites may be competitive with the ordering process that occurs during high-temperature annealing. The largest high-temperature endotherm occurs in sample 5, which also fails to show an endotherm around 10°C corresponding to the soft-segment crystallite melting. As the molding temperature increases and the chains be- come more mobile, the magnitude of the soft-segment melting endotherm increases while the magnitude of the high-temperature annealing peak decreases. If the system energetics favor the relaxation of the polymer chains that encourages the formation of the soft-segment crystallites as the sample is cooled over the ordering process responsible for the annealing endotherm, then the results seen in the DSC study are as expected. This competitive ordering would imply that the high-temperature annealing ordering involves both the hard and soft segments, as Van Bogartg has p r e posed. This does not imply that both hard and soft segments are becoming ordered during annealing, but that both segments must undergo some rearrangement for the annealing ordering to occur whatever the mechanism of the annealing ordering.

A typical SAXS curve for these samples is shown in Figure 2. This particular curve represents the scattering data for sample 5. The main feature of this scattering curve is the peak that occurs at around 0.3 nm-’. The results of the SAXS data analysis are summarized in Table 11. The molding temperature has the greatest effect on the SAXS data, which is consistent with the DSC data. The cooling rate and sample thickness have a smaller effect on the SAXS results and essentially none on the DSC results. The molding time was found to have the least effect on either SAXS or DSC

10’

102 - 0 I 2 3 4 5 6

q nm-’ A typical SAXS curve (Sample 5). Fig. 2.


results. From this it can be deduced that the sample morphology is most dependent on the molding temperature.

In the first group of materials, the molding temperature was varied from 130 to 200°C. As the molding temperature increased, both the long spacing and the Porod inhomogeneity length increased, by 25 and 15%, respectively. The degree of phase separation does not change with the molding temperature. This suggests that the domain size grows by coalescence of smaller domains, rather than by accumulation of new material due to improved phase separation. The ease with which the coalescence of domains occurs increases with temperature owing to the increase in chain mobility; hence larger domains are expected at higher molding temperatures, up to some limit determined by the chain dimensions. This coalescence is not complete, as in the case of water droplets coalescing. Instead the smaller packets of hard-segment material in the soft phase migrate to the larger domains, leaving a better-ordered system as a result. The overall degree of phase separation has not changed in this scenario. This result is consistent with the annealing work of Van Bogart,lo who showed that the endotherm associated with annealing increased at higher temperatures, indicative of more ordering of the hard domains.

It is interesting to note that the measured degree of phase separation was unaffected by molding temperature or, for that matter, by any of the preparation treatments. Speckhard et al.,40 in an investigation of polyise butylene polyurethanes, found that the DPS was primarily dependent on the average length of the hard segments. The dependence on the soft-segment length was found to be very small, within the range of experimental error. Miller et al.41 studied two different sets of polyether polyurethanes synthesized in different ways so as to produce materials with different hard- segment sequence length distributions in polymers with the same overall stoichiometry. They report that the materials with a higher fraction of hard segments containing one MDI unit have a greater degree of phase mixing than do their counterparts with the same average hard-segment length. All of these results suggest that phase mixing in polyurethanes depends much more on the segment type, average hard-segment length, and length dis- tribution than on the sample preparation methods.

The time that a sample is held a t the molding temperature has little effect on the microstructure of these materials as measured by DSC and SAXS. This suggests that over the time scale of the experiment, which is fairly limited, the materials have reached a short time equilibrium or quas- iequilibrium state. The true equilibrium state may well be quite different from the measured state, as relaxation mechanisms with large time con- stants would not show up in this time interval. The study of longer time periods will be the subject of future SAXS investigations of annealing effects in polyurethanes.

The rate of sample cooling was found to affect the Porod inhomogeneity length. As the rate of cooling rose, the number of small domains trapped by the rapid loss of chain mobility increased, hence we see a smaller het- erogeneity length. Conversely, at slower cooling rates the small domains have more time to coalesce, and the microstructure becomes more coarse. The results of this study are not entirely conclusive, however. The long


spacing does not change regularly, and although the Porod inhomogeneity length does change monotonically and systematically, the overall change is fairly small. DSC results show no real difference between the samples cooled at different rates.

The effect of sample thickness on the SAXS results is also quite small, and has an inconsequential effect on the DSC data. If the surface of the polyurethane possessed a substantially different morphology than the bulk and if this surface morphology extended to a significant depth, then chang- ing the sample thickness would change the SAXS pattern. This does not appear to be the case. Although there is a change noted in the long spacing, this is more likely due to the different rate of cooling of the sample core. The interior will cool at a slower rate, and hence have a slightly larger domain structure than the surface. As the thicker samples contain a greater percentage of core material, they are expected to show a larger average domain size.

The solutioncast sample does not exhibit substantially different SAXS parameters compared to the compression-molded samples under the conditions tested. It does show the lowest degree of phase separation of all of the samples. This is likely due to a trapping of the hard segments in the soft phase as the solvent is evaporated. DMA is a good solvent for this system, so the chains in solution are fairly extended. If the rate of solvent removal is more rapid than the rearrangement of the chains into a state more near equilibrium, the chains will tend to be more randomly arranged, and hence a greater degree of phase mixing will be observed. It is well known that the choice of solvent affects the final polymer morphology. A solvent that is preferential to one phase will yield a different microstructure than a solvent that preferentially dissolves the other phase. The rate of solvent removal will also affect the final phase structure and purity. A slow removal of solvent will allow the system to more nearly equilibrate. Hoff- mann and B0na1-t~~ discuss this in detail in a study of solventcast polyurethane films. The casting conditions chosen for this study, however, have produced a morphology not unlike that of the compression-molded samples.

The DSC thermogram of the solutioncast sample does differ from that of the molded samples. Greater soft-segment crystallinity is seen on the basis of the magnitude of the soft-segment crystallite melting endotherm. An endotherm is also seen near 80°C due to the casting process at 70°C. There is no high-temperature endotherm. Apparently during the solvent evaporation there is sufficient chain mobility to allow the chains to relax enough that when the sample is cooled in the DSC cell, the soft segments can crystallize to some extent, but the mobility is still low enough to prevent the phase separation from reaching the degree found in the molded samples where the degree of phase mixing is determined by the amount of phase separation that occurs when the solution-polymerized polymer is first precipitated from DMAc solution into water.

CONCLUSION

The effects of sample preparation on the domain morphology of a polyether polyurethane block copolymer prepared by compression molding were


examined. The molding temperature was found to have a significant effect on the small-angle x-ray data and the morphological information derived from it. The molding temperature also significantly affected the thermal behavior of the samples, as evidenced by the DSC traces. The cooling rate and the sample thickness were found to have a small effect on the SAXS results, but this seems inconclusive at this time. The time that the sample was held at the molding temperature had no discernible effect on the morphology, at least over the time scale of this experiment. The cooling rate, sample thickness, and molding time were found to have no substantial effect on the DSC thermograms. The results for the solutioncast film were consistent with those for the compression-molded samples. The study also shows that careful control of sample preparation is necessary for accurate, re- producible SAXS characterization of polyurethane morphology.

The authors would like to acknowledge the support of this research by the Polymer Section of the Division of Materials Research of the NSF through Grant No. DMR 81-06888.

References 1. P. E. Gibson, M. A. Vallance, and S. L. Cooper, in Developments in Block Copolymers,

2. J. W. C. Van Bogart, A. Lilaonitkul, and S. L. Cooper, A m Chem. Soc. Adv. Chem. Ser.,

3. S . L. Aggarwal, Ed., Block Polymers, Plenum, New York, 1970. 4. C. Hepburn, Polyurethane Elastomers, Applied Science, London, 1982. 5. R. Bonart, J. Macromol. Sci. Phys., B2, 115 (1968). 6. S . B. Clough and N. S. Schneider, J. Macromol. Sci. Phys., B2, 553 (1968). 7. S . L. Cooper and A. V. Tobolsky, J. Appl. Polym. Sci., 11, 1361 (1967). 8. R. Bonart, L. Morbitzer, and G. Hentze, J. Macmmol. Sci Phys., B3, 337 (1969). 9. J. W. C. Van Bogart, D. A. Bluemke, and S . L. Cooper, Polymer, 22, 1428 (1981).

Applied Science Series, Elsevier. London, 1982.

176, 3 (1979).

10. J. W. C. Van Bogart, P. E. Gihson, and S. L. Cooper, J. Polym. Sci. Polym Phys. Ed., 21,

11. J. T. Koberstein and R. S. Stein, J. Polym. Sci. Polym. Phys. Ed., 21, 2181 (1983). 12. J. T. Koberstein, Ph.D. Thesis, University of Massachusetts, 1979. 13. J. T. Koberstein and R. S. Stein, J. Polym. Sci Polym. Phys. Ed., 21, 1439 (1983). 14. Z. Ophir and G. L. Wilkes, Polym Prep. Am. Chem Soc. Div. Polym. Chem, 20, 503

15. Y. J. P. Chang and G. L. Wilkes, J. Polym. Sci Polym Phys. Ed., 13, 455 (1975). 16. G. L. Wilkes and R. Wildnauer, J. Appl. Phys., 46,4148 (1975). 17. G. L. Wilkes and J. A. Emerson, J. Appl. Phys., 47, 4261 (1976). 18. C. E. Wilkes and C. S. Yusek, J. Macromol. Sci. Phys., B7, 157 (1973). 19. R. Bonart, L. Morbitzer, and E. H. Muller, J. MacromoL Sci. Phys., B9, 447 (1974). 20. N. S. Schneider, C. R. Desper, J. L. Illinger, A. 0. King, and D. Barr, J. Macr~moL Sci

21. C. S. P. Sung, C. B. Hu. and C. S. Wu, Macromolecules, 13, 111 (1980). 22. C. R. Desper and N. S. Schneider, Am. Chem. Soc. Div. Org. Cwt. Plast. Pmpr., 45, 263

23. J. W. C. Van Bogart, Ph.D. Thesis, University of Wisconsin, 1981. 24. I. D. Fridman and E. L. Thomas, Polymer, 21, 388 (1980). 25. A. L. Chang, R. M. Briber, E. L. Thomas, R. J. Zdrahala, and F. E. Critchtield, Polymer,

26. J. A. Lake, A c h Crystallogr., 23, 191 (1967). 27. W. Ruland, J. Appl. Crystallogr., 4, 70 (1971). 28. J. T. Koberstein, B. Morra, and R. S. Stein, J. Appl. Ctystallogr., 13, 34 (1980). 29. C. G. Vonk, J. Appl. Crystallogr., 6, 81 (1973).

65 (1983).

(1979).

Phys., B11, 527 (1975).

(1981).

23, 1060 (1982).


30. R. J. Roe, J. Appl. Crystalbgr.., 15, 282 (1982). 31. G. Porcd, Kolbid Z., 124, 83 (1951). 32. W. Ruland, Kolbid 2. 2. Polym., 255, 417 (1977). 33. R. Bonart and E. H. Muller, J. Macromol. Sci. Phys., B10, 177 (1974). 34. R. Bonart and E. H. Muller, J. Macmmol. Sci. Phys., B10, 345 (1974). 35. A. Guinier and G. Fournet, Small-Angle Scattering of X Rays, Wiley, New York, 1955. 36. T. K. Kwei, J. Appl. Polym. Sci, 27, 2891 (1982). 37. S. L. Samuels and G. L. Wilkes, J. Polym. Sci. Polym. Phys. Ed., 11, 807 (1973). 38. J. Blackwell and C. D. Lee, J. Polym. Sci. Polym. Phys. Ed., 22, 579 (1984). 39. C. M. Brunette and W. J. MacKnight, Rubber Chern. Technol., 55, 1413 (1982). 40. T. A. Speckhard, P. E. Gibson, S. L. Cooper, V. S. C. Chang, and J. P. Kennedy, Polymer,

41. J. A. Miller, S. B. Lin, K. K. S. Hwang, K. S. Wu, P. E. Gibson, and S. L. Cooper,

42. K. Hoffmann and R. Bonart, Makroml. Chem., 184, 1529 (1983).

to appear.

Macmmokuh, to appear.

Received September 4, 1984 Accepted November 6, 1984

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