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Chapter 14
Crystallization and Solid-State Structure of Model Poly(ethylene oxide) Blends James Runt Department of Materials Science and Engineering, The Pennsylvania State University, University Park, PA 16802
This paper presents an overview of several studies whose overall goal is to elucidate the solid state microstructure and crystallization kinetics of melt-miscible polymer blends. Model blends were prepared with poly(ethylene oxide) and four amorphous polymer diluents: two exhibiting relatively weak intermolecular interactions with PEO and two exhibiting strong interactions. Small-angle x-ray scattering experiments showed that the introduction of strong intermolecular interactions promoted diluent segregation over relatively large length scales, regardless of diluent mobility at the crystallization temperature. In addition, at a given T , spherulitic growth rates for blends with the strongly interacting polymers are considerably lower than those with weakly interacting polymers with comparable T s. To a first approximation, growth rates for PEO and the blends can be superposed when normalized by the degree of supercooling and Τ - T . Finally, the chapter concludes with a brief review of recent timeresolved small and wide-angle x-ray experiments conducted on PEO and the same blends. c
g
g
Polymer blends containing crystalline polymers are commonplace and a growing number of commercial materials are polymer mixtures in which at least one of the components is semi-crystalline. Crystallization and solid state microstructure of neat crystalline materials have generally been studied in some depth. However, these features of semi-crystalline blends are less well understood, due at least partly to the greater complexity of these systems. Important considerations in the case of meltmiscible blends are the influence of the polymeric diluent on crystallization kinetics and its ultimate locations) in the lamellar microstructure. Diluent molecules can reside between lamellae in lamellar stacks, between fibrils and/or in interspherulitic regions, yielding different microstructures, which in turn give rise to different material properties. It is important, therefore, to develop an understanding of the 218
© 2000 American Chemical Society
Cebe et al.; Scattering from Polymers ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
219 factors that influence the extent of diluent segregation. Strong intermolecular interactions are well known to play a critical role in determining melt-miscibility, yet their impact on crystallization and microstructure formation have been rarely examined in depth. In the past several years, we have undertaken a series of experimental studies focusing on the investigation of the factors responsible for controlling the nature of the solid state microstructure (and the development thereof) and the crystallization kinetics of a series of melt-miscible blends (1-4). Poly (ethylene oxide) [PEO] was chosen as the semi-crystalline component since it exhibits miscibility in the melt with a relatively broad range of amorphous polymers. The current paper presents an overview of several aspects this work: (1) initial small-angle x-ray scattering (SAXS) experiments to elucidate the final microstructure of PEO and the blends; (2) spherulitic growth rate experiments over a range of crystallization temperatures and concentrations and; (3) a real-time small- and wide-angle x-ray scattering study of microstructure development and crystallization of these systems.
Experimental Materials Throughout this work, we focused on four melt-miscible poly(ethylene oxide) blend systems. Two exhibit relatively weak intermolecular interactions [PEO with poly(methylmethacrylate) (PMMA) and poly(vinylacetate) (PVAc)] and two exhibit relatively strong interactions (PEO with ethylene - methacrylic acid (EMAA) and styrene - hydoxystyrene (SHS) copolymers). The PEO used in these studies has a viscosity-average molecular weight of 1.44 χ 10 . The EMAA copolymer contains 55% by weight methacrylic acid units while the SHS copolymer consists of 50 weight percent styrene and p-hydroxystyrene. The second, diluent polymer is non crystalline in all cases. In each series, one of the amorphous polymers was chosen to have a relatively low glass transition temperature (T ) [PVAc and EMAA] and the other a relatively high T [PMMA and SHS]. The two weakly interacting polymers, PMMA and PVAc, had measured T s of 113 °C and 31 °C, while the T s of EMAA and SHS were 36 °C and 150 °C, respectively. Molecular weights and polydispersities of the amorphous polymers can be found in refs. / and 5. Various compositions of each blend were prepared by solution casting as described in refs. / and 5 and solvent removal was accomplished at elevated temperature under vacuum. 5
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Small-angle x-ray scattering 'Static' SAXS experiments were performed on the Oak Ridge National Laboratory 10-m pinhole collimated SAXS camera using C u K radiation (λ = 0.154 nm) and a 20 χ 20 cm position-sensitive area detector. The scattering from each sample was determined at two sample-to-detector distances: 2.119 and 5.119 m. The data were azimuthally averaged and converted to an absolute differential cross-section by means of precalibrated secondary standards (6): a high density polyethylene, PES3, was used for the low q data and a vitreous carbon standard was employed for the a
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Cebe et al.; Scattering from Polymers ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
220 high q data. Here, q is the scattering vector defined as (4π/λ) sin(θ/2), where θ is the scattering angle and λ is the x-ray wavelength. Time-resolved SAXS experiments were conducted on two different beamlines at the National Synchrotron Light Source (NSLS) at Brookhaven National Laboratories. Initial experiments were conducted on Beamline X3A2 at a wavelength of 0.154 nm (2). The scattered intensity was collected by a linear position sensitive photodiode detector coupled to an optical multichannel analyzer. In these early experiments, real-time SAXS scattering during crystallization at 45 °C was analyzed for one composition of each of the model blends, as well as the scattering from neat PEO crystallized at two temperatures. A much more comprehensive series of experiments have been conducted more recently on the Advanced Polymer Beamline at NSLS, X27C (4). The beamline setup provided access to simultaneous small- and wide-angle x-ray (WAXS) scattering (7). The x-ray wavelength in this series of experiments was 0.1307 nm and pinhole collimation was used in conjunction with a sample to detector distance of 1.190 m. We utilized a specially designed sample holder to allow for a rapid temperature jump between the crystallization temperature, T , and the temperature in the melt (7). The scattering pattern from a duck tendon was used to calibrate the SAXS detector in the q range from 0 to 0.15 Å . A Lupolen standard was used to calibrate the WAXS detector in the 2Θ range of approximately 11° to 38°. Timedependent bulk crystallinities were determined by resolving the WAXS profiles into crystalline reflections and an amorphous halo, then calculating the areafractionof the crystalline reflections. Details of the data processing in the different studies can be found in the original publications (1,2,4). The one-dimensional correlation function was calculated for all experimental SAXS curves (8): c
-1
(1)
G(r)
where r is the correlation distance. The pseudo two-phase model was applied to the correlation functions of neat PEO and the blends to extract the experimental invariant, Q e x p t . Note that in the case of the 'static' SAXS experiments, Q t is in absolute units, whereas for the time-resolved experiments the Q tare relative values. The invariants in the time-resolved experiments are generally depicted as Q t / Q m a x , where Q is the maximum value obtained during the experimental run. Q is determined from the ordinate of the linearfitto the self-correlation portion of the correlation function: e x p
exp
e x p
m
a
x
e x p t
Qexpt
=
V W (l-Wc)Δη s
c
(2)
2 1
where w is the linear crystallinity (given by the ratio of the average lamellar thickness, l , to the average long period, L), v the volumefractionof lamellar stacks (given by the ratio of the bulk crystallinity, φ , to the linear crystallinity, w ) and Δη the linear electron density difference (defined as Δη1 = η - η , where η andηintare the electron densities of the crystal and interlamellar amorphous layers, respectively). In the case of complete incorporation of uncrystallized material in interlamellar regions, the linear crystallinity and electron density difference in eqn. 2 are replaced by their bulk values, φ and Δη, and the invariant becomes (9) c
c
s
ς
c
c
int
c
c
Cebe et al.; Scattering from Polymers ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
1
221 Qcaic
=
φο(1-φο)Δη
(3)
2
Δη = T|c - η , where η is the electron density of the amorphous layer when all of the amorphous diluent resides in the interlamellar regions. One can therefore calculate an 'all-interlamellar' invariant ( Q i ) and comparison with Q t indicates the extent of diluent polymer incorporation in interlamellar regions. The average long period can be written as: c a
c
e x p
2
L = ax\ i(dGi(k)
(4)
x
where (dGldx) is the slope of a linear fit to the self-correlation portion of G(r). For w > 0.5, the intersection of the linear fit with G(r) = 0 is: r
0
= lc(l-w ) = l w c
a
(5)
c
Q e x p t , (dGlaY) and ro are obtained directly from the fit to the self-correlation portion of G(r). Together with the long period (derived from the first correlation maximum), this permits quantitative analysis of the lamellar microstructure.
Spherulitic growth rates Dried films of PEO and the blends were heated to 100 °C (for 3 min.) in a hot stage to erase previous thermal history, then rapidly transferred to a second hot stage set at the desired crystallization temperature. The development of the spherulitic superstructure was then viewed with an Olympus BHSP-300 microscope and the crystallization event recorded with a video camera and VCR. In the case of non-linear growth, spherulite growth rates (G) were determined at the onset of growth. Average growth rates were determined on from 1 - 5 spherulites.
Results and Discussion 'Static' SAXS experiments (1) The two PEO blends that exhibit weak intermolecular interactions display quite different SAXS behavior after crystallization at T = 45 °C. For the PEO/PMMA blends, the scattering peaks move to significantly lower q (larger long periods), with increasing PMMA content. The correspondence between experimental and calculated 'all interlamellar invariants, as well as experimental and calculated 'all interlamellar' long periods, support complete incorporation of PMMA within lamellar stacks. This is in agreement with the results of other authors (10). However, for the PEO/PVAc blends, the scattering maxima shift only slightly to lower q with increasing PVAc content. In addition, the experimental and calculated invariants do not exhibit good correspondence. These findings, along with optical microscopy observations, suggest at least partial exclusion of the relatively mobile PVAc to interfibrillar regions during crystallization at 45 °C. Unfortunately, quantitative assessment of the extent of PVAc migration was not possible due to the relatively low electron density contrast between amorphous PEO and PVAc. c
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Cebe et al.; Scattering from Polymers ACS Symposium Series; American Chemical Society: Washington, DC, 1999.
c
222 Lorentz-corrected SAXS intensities vs. q for several PEO/EMAA blends are shown in Figure 1. A dramatic increase in long period (from -24 to 45 nm) is observed when EMAA content is increased to only 20%. That this is not primarily associated with incorporation of EMAA in lamellar stacks, as demonstrated in Figure 2. Figure 2 is a plot of the experimental invariant and the invariant calculated assuming all interlamellar incorporation of the EMAA chains. EMAA has an electron density that is considerably smaller than those of neat amorphous and crystalline PEO so that even modest inclusion of EMAA in lamellar stacks would be expected to lead to a significant increase in the invariant, as shown by the solid line in Figure 2. The data therefore indicate very significant displacement of EMAA to extralamellar regions. The microstructural parameters extracted from analysis of the correlation functions show clearly that the increase in long period with increasing EMAA content is due primarily to an increase in lamellar thickness, which is consistent with a lower degree of supercooling (ΔΤ). Similar conclusions were drawn for the SHS blends except that the observed increases in lamellar thickness are more modest, suggesting that the equilibrium melting point depression is not as pronounced for this blend. Equilibrium melting points of the EMAA and SHS blends were estimated by using the following approach. Using the Lauritzen, Hoffman, Miller theory of crystallization (//) and assuming that the thickening factor and end surface free energy are independent of blending then: l = ζ · 1/ΔΤ, where ζ is a constant for PEO and the blends. This then permits an estimation of the degree of supercooling at which the EMAA (and SHS) blends were crystallized by comparing the experimental crystal thicknesses with those determined by Arlie et al. for neat PEO (12). Further details of this analysis are described in references 1 and 3. For example, it was estimated that the addition of EMAA to PEO depresses the equilibrium melting point by as much as 5 -10 °C for as little as 10 - 20% EMAA. This seems plausible in light of the dramatic reduction in the crystallization rates of these blends with increasing EMAA content (see the next sections). The PEO/EMAA and PEO/SHS blends exhibit volume-filling spherulites for all compositions examined in this part of the study (i.e., < 20%). These observations, in concert with SAXS results, indicate at least partial exclusion of the diluent into regions between lamellar stacks at these compositions. The distribution of the second polymer between the interlamellar and interfibrillar regions can be estimated using measured bulk crystallinities and correlation function parameters as follows. The volumefractionof lamellar stacks is determined from the bulk and linear crystallinities (v = φο/Wc) and the average electron density difference between crystalline and interlamellar amorphous regions can be determined using eqn. 2. Since r| is known (0.676 mol e-/cc (13)% the average electron density of the interlamellar regions ( η 0 can be obtained, and the concentration of amorphous PEO in the interlamellar regions can be calculated as: c
s
c
ΐη
ΦΙΕ
=
( η ΐ ι « - η < ι ) / ( η » -η
