crystallization-induced phase separation in mixtures of model linear and short-chain branched...

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Crystallization-inducedphase separation inmixtures of model linearand short-chain branchedpolyethylenesMasaya Ueda a b & Richard A. Register aa Department of Chemical Engineering ,Princeton University , Princeton, New Jersey,08544b Mitsubishi Engineering Plastics, TechnicalCenter , 5-6-2 Higashi-Yawata, Hiratsuka City,Kanagawa, 254, JapanPublished online: 19 Aug 2006.

To cite this article: Masaya Ueda & Richard A. Register (1996) Crystallization-induced phase separation in mixtures of model linear and short-chain branchedpolyethylenes, Journal of Macromolecular Science, Part B: Physics, 35:1, 23-36

To link to this article: http://dx.doi.org/10.1080/00222349608220374

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J. MACROMOL. X I . - P H Y S . , B35(1), 23-36 (1996)

Crystallization-Induced Phase Separation in Mixtures of Model Linear and Short-Chain Branched Polyethylenes

MASAYA UEDA* and RICHARD A. REGISTER? Department of Chemical Engineering Princeton University Princeton, New Jersey 08544

ABSTRACT

The possibility of cocrystallization is investigated in a model system of linear and short-chain branched polyethylene. Both components have low molecular weight and narrow molecular weight distribution; the branched component, hydrogenated polybutadiene (HBD), contains one ethyl branch per 60-70 backbone carbon atoms. Cocrystallization was not observed even for the most rapid cooling conditions attainable, despite the essentially identical unit cells of the two materials. A combination of microscopy and scattering techniques was used to demonstrate that the HBD and linear polyethylene (LPE) crystallize into separate lamellar stacks. The rapid growth rate of the LPE creates volume-filling spherulites which trap the HBD in micron-size inclusions, within which it sub- sequently crystallizes on further cooling.

*Permanent address: Mitsubishi Engineering Plastics, Technical Center, 5-6-2 Higashi- Yawata, Hiratsuka City, Kanagawa, 254 Japan. ?To whom correspondence should be addressed.

23

Copyright 0 1996 by Marcel Dekker, Inc.

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24 UEDA AND REGISTER

INTRODUCTION

Cocrystallization between unbranched and branched polyethylenes has been investigated by several groups over the last decade, with some groups reporting extensive or complete cocrystallization, and others reporting little or no cocrystallization. The differing conclusions arise principally from the diversity of systems which have been investigated, and particularly the identity of the branched component. Many investigations have employed materials with long-chain branching, but we focus here on molecules whose architectures are essentially linear, and which have either no branches or only short-chain branches. Most such investigations have employed mixtures of commercial high-density polyethylene (HDPE, unbranched) and linear low-density polyethylene (LLDPE, branched). Perhaps the most comprehensive investigation of cocrystallization in the HDPEILLDPE system was performed by Tashiro et al. [l-61, who used differential scanning calorimetry (DSC), Fourier-transform infrared (FTIR) spectroscopy, wide- angle x-ray scattering (WAXS), small-angle x-ray scattering (SAXS), and small-angle light scattering (SALS) to verify that the two components co- crystallized, even on slow cooling. Their LLDPE contained 17 ethyl branches per 1000 backbone carbons. Both components were of commercial molecular weights (M, = 60-80 kg/mol) and were highly polydisperse. Extensive cocrystallization was also found by Gedde and coworkers [7-91 in blends of low molecular weight HDPE (M, = 2500 g/mol) with an LLDPE similar to that studied by Tashiro et al. Extensive or complete cocrystallization has also been reported in other HDPEILLDPE blends by Edward [ 101 and Hu et al. [ 111.

Taken together, these results indicate that commercial LLDPE containing 15-20 ethyl branches per 1000 backbone carbons should easily cocrystallize with HDPE under conditions typically encountered in processing. However, it is known that commercial LLDPE has a large compositional polydispersity [ 121, or short-chain branching distribution. Thus, it is perhaps more correct to consider HDPE/LLDPE blends to contain a contin- uum of components, rather than just two, with branch contents extending from zero (HDPE) upward. The crystallization behavior in such a system may well differ from that in a simple binary system of linear and uniformly branched components.

Model LLDPE systems can be prepared by hydrogenation of polybutadiene with varying levels of 1,2 addition, yielding polyethylenes with vari- able ethyl branch contents [ 131. The anionic polymerization of the precursor polybutadiene ensures that the chains have narrow molecular weight and compositional distributions, and no long-chain branching; the sequenc- ing of the ethyl branches in the hydrogenated butadiene (HBD) product has

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CRYSTALLIZATION-INDUCED PHASE SEPARATION 25

been shown to be essentially random [14]. Blends of HBD polymers with HDPE, particularly narrow-distribution HDPE fractions, are more repre- sentative of true binary blends. In contrast to the numerous studies of HDPE/LLDPE blends, very few studies of HDPE/HBD blends have been reported in the literature. Alamo et al. [ 151 investigated blends of narrow- distribution, high molecular weight (M, = 100 kg/mol) HDPE and hydrogenated polybutadiene (HBD) containing 22 ethyl branches per 1000 backbone carbons. Using a combination of DSC, SALS, transmission electron microscopy (TEM), and Raman longitudinal acoustic mode measurements, they demonstrated that cocrystallization occurs when the blends are rapidly quenched, but that segregation occurs during isothermal crystallization. Crist and Yang [ 161 examined similar HDPE/HBD blends of lower molecular weight (M, = 30 kg/mol) by DSC and concluded that segregation was the rule. Possible partial cocrystallization was noted for samples quenched at 80°C/min in the DSC, as indicated by the emergence of a third melting endotherm on reheating, at a temperature between the melting endotherms of the HDPE and HBD components.

Here, we investigate the crystallization behavior of low molecular weight, narrow-distribution linear polyethylene (LPE) and HBD, using both thermal and morphological characterization techniques. The low molecular weight ensures complete miscibility between the two components in the melt, given the known interaction parameters between LPE and similar HBD [17], as well as extending the range of molecular weights used in previous investigations [ 15,161. The relatively narrow crystallite size distribution allows the presence or absence of cocrystallization to be straightfor- wardly determined.

EXPERIMENTAL

Materials and Sample Preparation

The linear polyethylene (LPE) studied here, POLYWAX@ 3000, was generously provided by Petrolite Corporation (Tulsa, OK). Petrolite literature indicates that this polymer has M, = 3000 g/mol (from vapor phase osmometry), M,/M, = 1.10, and no detectable branching. These results were confirmed by independent measurements reported by Pearson et al. [18], who found the level of branching to be below 0.8 branches per 1000 backbone carbons. The hydrogenated polybutadiene was synthesized and hydrogenated by Dr. L. J. Fetters and Mr. W. G . Funk of Exxon Research and Engineering Company. The anionic synthesis of the precursor polybutadiene was done in cyclohexane with an alkyllithium initiator, using high- vacuum techniques [19]. Hydrogenation [13] was done in a Parr reactor

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using gaseous H2, and a Pdo catalyst supported on CaCO,. FTIR of the product confirmed that the residual content of double bonds was less than 1 per lo00 backbone carbons. Size-exclusion chromatography on the satu- rated polymer yielded M, = 3800 g/mol and M,/M, = 1.06. The ethyl branches in the hydrogenated polymer result from 1,2 butadiene additions during the polymerization, which are easily quantified by 'H nuclear mag- netic resonance on the unsaturated precursor. Examination in this laboratory [ 201 of numerous polybutadienes synthesized identically indicates that the ethyl branch content should be approximately 16 ( + 2) per 1000 backbone carbons.

A 50/50 w/w blend of the two polymers was prepared by codissolution in toluene at 85OC (4.0 wtVo). The hot toluene solution was poured into 10 volumes of methanol cooled with an ice/water bath to rapidly and completely precipitate both constituents. The precipitate was vacuum dried at 7OoC for 3 days to ensure complete removal of solvent. To obtain film specimens of approximately 1 -mm thickness, the pure components and the blend precipitate were melted at 16OOC for 1 min in polytetrafluoroethylene dishes, allowed to flow under their own weight into films, and then cooled in air at 5°-100C/min to room temperature.

Characterization Methods

Differential scanning calorimetry (DSC) employed a Perkin-Elmer DSC-7, calibrated with indium and mercury, and operating under nitrogen. The melting and crystallization temperatures T, and T, were taken as the maxima during heating or cooling at 10°C/min; for the heating runs, unless otherwise indicated, the samples were cooled from the 2OOOC melt to O°C at 10°C/min prior to scanning. Enthalpies were converted to crystalline weight fractions w, by taking the heat of fusion of the perfect polyethylene crystal [21] as 277 J/g. The broad melting tail on the HBD makes the calculation of w, correspondingly uncertain.

SAXS was used to probe the room temperature morphology of the samples, using a compact Kratky camera and a Braun position-sensitive detector. Data were collected and reduced using procedures described else- where [22] and expressed as the background-subtracted, desmeared abso- lute intensity Z/ZeV versus the scattering vector q = (4~/X)sin 8; Z, is the scattering from a single electron, V is the illuminated scattering volume, X is the radiation wavelength, and 28 is the scattering angle. In anticipation of a lamellar structure, the intensities are shown after multiplication by q2, to easily determine the lamellar repeat distance [23].

Bright-field polarized and unpolarized optical microscopy employed a Zeiss Axioskop and a Mettler FP2 hot stage. The sample, contained be-

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CRY STALLIZATION-INDUCED PHASE SEPARATION 27

tween a microscope slide and cover slip, was inserted into the hot stage, heated to 150°C and held for 2 min, then cooled at 10°C/min to either 13OOC (LPE, blend) or llO°C (HBD). Crystallization was monitored on cooling from this temperature at 2OC/min.

SALS in the depolarized (H,) mode employed a home-built apparatus [24] using a 15-mW He-Ne laser (Melles-Griot). The scattering pattern impinges on a back-projection screen and is recorded with a CCD camera (NEC TI-22A) connected to a videocassette recorder. Images were digitized with a frame grabber (PC-EYE, Chorus Data Systems) for quantitative analysis of the scattered intensities. The samples were contained within a temperature-controlled heater block described previously [25]. The material was contained between two 1-in. diameter glass disks (Gray Glass Co.). The thickness of the sample (10-40 pm) was controlled by inserting 10-pm stainless steel shims between the glass disks. Cooling at roughly 2OC/min is accomplished by natural convection across the face of the heater block. To account for temperature gradients across the block which are likely to occur during cooling, the cell was calibrated by observing the crystallization of LPE on cooling and comparing the value indicated by the block temperature controller (CN9000A, Omega Engineering) with the temperature (1 18OC) at which DSC indicated that the LPE had crystallized 20% during cooling at 2 OC/min (1 18 OC).

RESULTS AND DISCUSSION

The melting and crystallization curves of LPE, HBD, and their 50/50 (w/w) blend obtained by DSC at 10°C/min are plotted in Figs. 1 and 2, respectively. While the melting and freezing transitions of LPE are repre- sented by a sharp single peak, the exotherms and endotherms for HBD show long tails to lower temperature due to the random ethyl branching and resulting wide distribution of ethylene sequence lengths. Melting points (endotherm peak temperatures, T,,,), freezing points (exotherm peak temperatures, T,), and degrees of crystallinity (weight fraction crystallinity, w,) are given in Table 1. Note that the LPE is nearly 100% crystalline (mea- sured heat of fusion 265 J/g), which is certainly not typical of commercial HDPE having higher molecular weight and polydispersity. By contrast, the HBD is only about 50% crystalline, due to the presence of the ethyl branches.

Two peaks, corresponding to melting and crystallization of the individual components, are evident in the blend. In Fig. 1, the higher-temperature endotherm (T, = 126.OOC) in the blend is shifted to lower temperature by 1.8OC relative to the melting point of the pure LPE. Similarly, the crystallization temperature T, is shifted downward by l.O°C in the blend

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I , , , 4, ,H:, , 40 60 80 100 120 140 160

Temperature ("C)

FIG. 1. DSC heating scans (lO°C/min); top to bottom: pure LPE, 50/50 (w/w) LPE/HBD blend, and pure HBD. Prior to scanning, samples were heating to 2OO0C, then cooled to O°C at 10°C/min.

relative to pure LPE. However, these small shifts should not be construed as proof of cocrystallization; indeed, the heating thermogram for the blend strongly resembles a simple superposition of the LPE and HBD thermograms, allowing for minor shifts in peak position.

Cocrystallization should be favored by increasing the cooling rate from the melt. Figure 3 shows the effect of quench rate on the thermogram of the blend; the heating rate in each case is 10°C/min, following quenches at 2, 10, or 400°C/min (nominal) as indicated. As the cooling rate is increased, the temperatures of both the lower-temperature (HBD) and higher- temperature (LPE) endotherms are reduced; the difference is roughly 4OC for both peaks when the thermograms for the slow-cooled (2OC/min) and quenched (400°C/min nominal) materials are compared. However, the shape of the thermogram is not noticeably affected by the cooling rate. The overlap between the peaks makes it difficult to unambiguously determine the areas of the individual peaks, but to within the uncertainty of this determination, there is no change in the relative peak areas with cooling rate. This stands in contrast to the results of Alamo et al. [ 151, who found that "DSC-quenched" 50/50 blends of higher molecular weight HDPE/ HBD showed little or no endotherm corresponding to the HBD component.

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WE f

blend

I I I I 60 80 1 do 120 140

Temperature ("C)

FIG. 2. DSC cooling scans ( 10°C/min, cooling from 200°C); top to bottom: pure LPE, pure HBD, and their 50/50 (w/w) blend.

It also contrasts with the results of Crist and Yang [ 161 in an intermediate molecular weight HDPEIHBD system, which showed a peak between the HDPE and HBD endotherms for a DSC-quenched sample.

For polymer systems where the two components have different unit cells [26] , WAXS can easily assess whether cocrystallization occurs or not. However, because HDPE and HBD have essentially identical unit cells [27], interpretation of the WAXS pattern would necessarily rest on quantitative analysis of the line shape, making an unambiguous determination difficult. To assess whether cocrystallization is occurring in these polyethylene blends, the most direct structural technique is SAXS, since the linear and branched components have very different lamellar thicknesses. Figure 4

TABLE 1

Characteristics of Blend Components

Material M, (g/mol) MJM, T,,, ("C) T, ("C) w,

LPE 3300 1.10 127.8 114.3 0.96 HBD 3800 1.06 106.8 97.4 0.49

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30

1 1

UEDA AND REGISTER

FIG. 3. DSC heating scans ( 10°C/min) of the 50150 (w/w) LPE/HBD blend, after cooling from 2OOOC at various rates; top to bottom: 2OC/min, 10°C/min, and 400°C/min (nominal).

shows the q2-corrected SAXS patterns for LPE, HBD, and the 50/50 blend after cooling at 5°-100C/min from the melt. While HDPE commonly shows only one or perhaps two broad SAXS peaks, the LPE shows three sharp peaks at q values of 0.262, 0.50, and 0.74 nm - I . These correspond to the first-, second- and third-order reflections from a lamellar structure with a Bragg spacing d = 27r/q* = 24.0 nm. Since the LPE is nearly 100% crystalline by DSC, the Bragg spacing should correspond closely to the lamellar thickness. From the known polyethylene unit cell [28], a lamellar thickness of 24 nm corresponds to 94 unit cells in the chain axis direction; a string of 94 unit cells corresponds to a molecular weight of 2630 g/mol. Since M, = 3000 g/mol for the LPE, it is apparent that the LPE forms essentially extended-chain crystals. Thus, the low molecular weight and narrow distribution of the LPE lead to a very narrow lamellar thickness distribution, providing the well-ordered SAXS pattern. By contrast, the HBD shows one narrow and one broad reflection at q = 0.426 and 0.86 nm-', indicating a greater range of lamellar thicknesses and that the aver- age chain undergoes several folds. Comparing this HBD with similar materials of higher molecular weight previously studied in this laboratory [20], the lower molecular weight material has a similar q* but a much narrower

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10000

n v)

a d

W

1000

O.$O 0.15 0.50 0.75 1.60 1.15 l ' I I I I I ) , I I

9 bm-') FIG. 4. SAXS patterns of the pure LPE (-), pure HBD (--), and their 501

50 (w/w) blend (- -). Patterns taken at room temperature after cooling from the melt at 5°-100C/min.

peak. Thus, even though the HBD does not form extended-chain crystals, it appears that the low molecular weight still produces a more narrow distribution of Bragg spacings. Since the HBD has less than 50% by volume crystallinity, it is unclear whether this improved order results from a narrower distribution of crystallite thicknesses or from a narrower thickness distribution for the intercrystallite amorphous regions.

The SAXS pattern for the blend is clearly a simple superposition of the SAXS patterns of the two pure components, indicating that the materials do not cocrystallize into the same lamellae. Moreover, it is clear that the LPE and HBD must be segregated on the scale of many lamellar thicknesses, forming separate lamellar stacks; otherwise, the peaks would broaden relative to the pure components, and no substantial broadening is evident in the blend SAXS pattern.

The combination of SAXS and DSC shows that these materials do not cocrystallize, even for the maximum quench rate possible in the DSC. While this indicates that the size scale of segregation must at least correspond to the lamellar stacks (order 0.1 pm or larger), it does not place an upper limit on the segregation length scale. From the DSC cooling scans shown in Fig. 2, it is apparent that the LPE component in the blend crystallizes first. Two

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possibilities suggest themselves, depending on the relative rates of diffusion and crystallization for the two polymers. In one case, the uncrystallized HBD might completely segregate from the growing LPE spherulite; upon further cooling, the HBD might form spherulites of its own. Another possibility is that the HBD is trapped in small pockets within the LPE spherulites, forming separate lamellar stacks but not separate spherulites.

Hot-stage optical microscopy, where the sample was cooled at 2OC/ min, shows that both pure components as well as the blend form rather large spherulites (roughly 50 pm diameter) [24]. In the blend, rapid forma- tion of spherulites begins at l 15 OC; these spherulites apparently grow to fill the entire sample volume by 113OC, well above the crystallization temperature of HBD. This suggests that the HBD is principally contained in small pockets within the LPE spherulites. Additional evidence that the HBD does not crystallize into separate spherulites is provided by the DSC cooling scans shown in Fig. 2. In contrast to the DSC heating scan, the cooling scan of the blend does not resemble a superposition of the scans of the two pure components. The pure LPE and HBD both have sharp leading (high- temperature) edges, as is expected for spherulitic primary crystallization. Both have extended tails to low temperature (particularly the HBD), reflect- ing secondary crystallization into thin lamellae. In the blend, the leading edge of the high-temperature peak is also sharp, but the low-temperature peak shows only a broad leading edge. This suggests that the primary spherulitic crystallization process involves only LPE, and that the HBD crystallizes in a secondary process within the LPE spherulites.

Additional confirmation of this conclusion was obtained by small-angle light scattering. The H, SALS patterns show the typical four-leaf-clover pattern characteristic of spherulites [29], though the large spherulite size makes the pattern difficult to completely resolve from the incident beam [24]. Samples were examined on cooling at approximately 2OC/min. In the blend, once the LPE component has crystallized at 115OC, the scattering pattern does not change with further cooling, but the intensity increases. By capturing a series of scattering patterns during the cooling process, reducing each to a plot of the intensity I versus q, where I is determined at an azimuthal angle of 4 5 O (i.e., along the four lobes of the pattern), and integrating Z with respect to q, Fig. 5 is obtained, which shows the integrated SALS intensity as a function of temperature for the two pure components and the 50/50 blend. The blend clearly shows a two-step intensity increase on cooling, corresponding first to the crystallization of LPE and then to the crystallization of HBD. Since the SALS pattern does not change in shape, but simply increases in intensity, this demonstrates that the HBD does not crystallize isotropically. Rather, the HBD chain axis must be lo- cally aligned with the LPE chain axis. This would be accomplished if the

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CRYSTALLIZATION-INDUCED PHASE SEPARATION

g1600-

33

b \

3 2 (d - 1200

h 4 -r( [I)

800 c)

k U

0 LPE D w O O HBO MMA blend

a Q) 400

+J

(d k bn Q) 4

A OiO ' 7b ' 40 t i 0 140 ' 140 ' 1 j O Temperature ("C)

FIG. 5. Integrated H, SALS intensities during cooling for the pure LPE (- , 0), pure HBD (--, O), and their 50/50 (w/w) blend (- -, A).

HBD crystallizes in elongated pockets, where the long axis of the pocket runs parallel to the radial direction of the spherulites.

It is interesting to interpret our findings in the context of the pioneering work of Keith et al. [30], which dealt with crystallization in linear PE (M, = 4500-726,000 g/mol) blended with n-C,,H,, at a 50/50 w/w ratio. This work showed that n-C32H66, which has a crystallization temperature lower than PE but essentially the same crystal structure, was segregated between lamellar stacks in spherulites of the PE, forming elongated pockets roughly 0.5 pm across, exactly as inferred above. Prior work from the same group [ 3 11 examined crystallization in isotactic polypropylene and polystyrene blended with their atactic analogs. There, it was suggested that the characteristic size of the pockets in which the noncrystallizable impurity is accu- mulated is given by the ratio DIG, where D is the diffusion coefficient of the impurity in the melt, and G is the radial growth rate of the spherulite.

In our dynamic cooling experiments (2OC/min), the LPE radial growth rate as observed by optical microscopy is quite fast, at least 1 pm/sec. The diffusion coefficient of linear polyethylene as a function of temperature and molecular weight was determined by Pearson et al. [18]. At 115OC, the LPE studied here has a diffusion coefficient of 4.5 pm2/sec. The HBD, which plays the role of impurity here, has a slightly higher molecular weight

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and may show a slightly reduced diffusion coefficient. These numbers indicate that the D/G ratio should be 5 pm or smaller. Such a size is consistent with both the SAXS results, which show a segregation size scale above 0.1 pm, and the SALS and optical microscopy results, which indicate a segregation size scale below 50 pm.

CONCLUSIONS

We have shown that cocrystallization does not occur in blends of low molecular weight, narrow-distribution LPE and HBD, for any attainable cooling conditions, despite their essentially identical crystal structures. It is well known that binary blends of LPE molecular weight fractions can segregate under certain cooling conditions [32]; the different fractions will have different melting points. Here, the two polymers have very similar molecular weights, but the ethyl branches in the HBD component provide the difference in melting point; in contrast to binary LPE blends with similar differences [32] in T,, no cooling conditions could be found which led to cocrystallization in the LPE/HBD blend. The LPE component forms essentially extended-chain crystals, as determined by SAXS. Rapid crystallization of the higher-melting LPE component traps the segregated HBD in pockets within the HDPE spherulites, but these pockets are large compared with the lamellar thickness. The mechanism of crystallization (primary LPE crystallization, secondary HBD crystallization) is reflected in the shape of the DSC cooling thermograms.

Block copolymers containing HBD blocks are currently of considerable interest as polymer blend compatibilizers [ 331; for example, using polystyrene-HBD diblocks to compatibilize blends of polystyrene and HDPE [ 341. The difficulty in achieving cocrystallization between HBD and LPE seg- ments, however, indicates that the improved mechanical properties [33, 341 of blends containing the diblock compatibilizer is due to the reduction of interfacial tension and the accompanying reduction in domain size and changes in domain continuity, rather than to cocrystallization between the LPE and HBD constituents.

ACKNOWLEDGMENTS

The authors thank Dr. L. J . Fetters for providing the HBD sample, and Dr. P. Rangarajan, Mr. D. J. Quiram, and Mr. D. M. Dean of Princeton University for assistance in sample preparation and characterization. Prin- cipal support for this research was provided by Mitsubishi Chemical Corpo- ration through a fellowship to M.U. Acquisition of the DSC was made possible by a grant from the AT&T Education Foundation and matching

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CRY STALLIZATION-INDUCED PHASE SEPARATION 35

funds from the Polymers Program of the National Science Foundation (DMR-9257565, to R.A.R.).

REFERENCES

1. 2.

3.

4.

5.

6.

7.

8.

9. 10. 11. 12.

13.

14.

15.

16.

17.

18.

19. 20 *

21 I 22. 23 I

K. Tashiro, R. S. Stein, and S. L. Hsu, Macromolecules, 25, 1801 (1992). K. Tashiro, M. M. Satkowski, R. S. Stein, Y. Li, B. Chu, and S. L. Hsu, Macromolecules, 25, 1809 (1992). K. Tashiro, M. Izuchi, M. Kobayashi, and R. S. Stein, Macromolecules, 27, 1221 (1994). K. Tashiro, M. Izuchi, M. Kobayashi, and R. S. Stein, Macromolecules, 27, 1228 (1994). K. Tashiro, M. Izuchi, M. Kobayashi, and R. S. Stein, Macromolecules, 27, 1234 (1994). K. Tashiro, M. Izuchi, M. Kobayashi, and R. S. Stein, Macromolecules, 27, 1240 (1994). M. T. Conde Braiia, J. I. Irragorri Sainz, B. Terselius, and U. W. Gedde, Polymer, 30,410 (1989). J. I. Irragorri, J . M. Rego, I. Katirne, M. T. Conde Braiia, and U. W. Gedde, Polymer, 33,461 (1992). M. T. Conde Braiia and U. W. Gedde, Polymer, 33, 3123 (1992). G . H. Edward, Br. Polym. J., 18,88 (1986). S. Hu, T. Kyu, and R. S. Stein, J. Polym. Sci. B: Polym. Phys., 25,71 (1987). F. M. Mirabella, Jr., and E. A. Ford, J. Polym. Sci. B: Polym. Phys., 25,777 (1987). H. Rachapudy, G. G. Smith, V. R. Raju, and W. W. Graessley, J. Polym. Sci., Polym. Phys. Ed., 17, 1211 (1979). T. M. Krigas, J. M. Carella, M. J. Struglinski, B. Crist, W. W. Graessley, and F. C. Schilling, J. Polym. Sci., Polyrn. Phys. Ed., 23, 509 (1985). R. G. Alamo, R. H. Glaser, and L. Manderkern, J . Polym. Sci. B: Polym. Phys., 26,2169 (1988). B. Crist and V. Yang, Polym. Prepr. (Am. Chem. SOC., Div. Polym. Chem.), 34(2), 470 (1993). W. W. Graessley, R. Krishnarnoorti, N. P. Balsara, R. J. Butera, L. J. Fetters, D. J. Lohse, D. N. Schulz, and J. A. Sissano, Macromolecules, 27, 3896 ( 1994). D. S. Pearson, G. Ver Strate, E. von Meerwal, and F. C. Schilling, Macromol- ecules, 20, 1133 (1987). M. Morton and L. J. Fetters, Rubber Chem. Technol., 48,359 (1975). P. Rangarajan, R. A. Register, and L. J. Fetters, Macromolecules, 26, 4640 (1993). F. A. Quinn, Jr., and L. Mandelkern, J. Phys. Chem., 88, 3548 (1963). R. A. Register and T. R. Bell, J. Polym. Sci. B: Polym. Phys., 30, 569 (1992). T. P. Russell, in Handbook on Synchrotron Radiation (G. S . Brown and D. E. Moncton, eds.), Vol. 3, North-Holland, New York, 1991.

Dow

nloa

ded

by [

Uni

vers

ity o

f L

eeds

] at

02:

28 2

0 A

ugus

t 201

4


24. 25.

26.

27. 28.

29. 30.

31. 32. 33.

34.

M. Ueda, M.S.E. Thesis, Princeton University (1995). N. P. Balsara, L. J. Fetters, N. Hadjichristidis, D. J. Lohse, C. C. Han, W. W. Graessley, and R. Krishnamoorti, Macromolecules, 25, 6137 (1992). T. L. Bluhm, G. K. Hamer, R. H. Marchessault, C. A. Fyfe, and R. P. Veregin, Macromolecules, 19,2871 (1986). P. R. Howard and B. Crist, J. Polym. Sci. B: Polym. Phys., 27,2269 (1989). H. Tadokoro, Structure of Crystalline Polymers, Wiley, New York, 1979, p. 375. R. S. Stein and M. B. Rhodes, J. Appl. Phys., 31, 1872 (1960). H. D. Keith, F. J. Padden, Jr., and R. G. Vadimsky, J. Polym. Sci. A-2, 4 , 267 (1966). H. D. Keith and F. J. Padden, Jr., J. Appl. Phys., 35, 1270 (1964). P. Smith and R. St. John Manley, Macromolecules, 12,483 (1979). D. R. Paul, in Thermoplastic Elastomers: A Comprehensive Review ( N . R. Legge, G. Holden, and H. E. Schroeder, eds.), Hanser, New York, 1987. R. Fayt, R. JerBme, and P. Teyssie, J. Polym. Sci. Polym. Phys. Ed., 19, 1269 (1981).

Received March 10, 1995 Revised April 3, 1995 Accepted April 6, 1995

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crystallization-induced phase separation in mixtures of model linear and short-chain branched...

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