Langmuir 1999, 15, 2911-2915
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Surface-Induced Ordering in Graft Copolymer Thin Films Shouren Ge, Lantao Guo, Miriam H. Rafailovich,* and Jonathan Sokolov Department of Materials Science and Engineering, State University of New York at Stony Brook, Stony Brook, New York 11747-2275
Dennis G. Peiffer Exxon Research and Engineering Company, Annandale, New Jersey 08801
Steven A. Schwarz Department of Physics, Queens College, Flushing, New York 11367
Ralph H. Colby Department of Materials Science and Engineering, The Pennsylvania State University, University Park, Pennsylvania 16802-5007
William D. Dozier Fannie Mae, Washington, D.C. 20016 Received May 11, 1998. In Final Form: January 29, 1999 The surface-induced ordering of deuterated poly(ethyl acrylate)-polystyrene (dPEA-PS) graft copolymers with different average number x of graft chains per copolymer chain (dPEA-g-xPS, x ) 1, 3, 5) was investigated using neutron reflection, secondary ion mass spectrometry, and atomic force microscopy. The ordering starts only at the vacuum surface and decays through the bulk of the film to the silicon surface. The dPEA-g-3PS thin films (f, fraction of styrene monomers ) 0.28) appear to have dPEA-PS lamellar structures and a nearly pure dPEA layer next to the vacuum surface. The dPEA-g-5PS thin films (f ) 0.48) are found to be cylindrical except for the top layer, adjacent to the vacuum surface, which is lamellar. On comparison with diblock copolymers, graft copolymers show a phase diagram shift and smaller micelle spacing. We find that long-range order in graft copolymers only occurs within ∼15 lamellar spacings of the vacuum surface. Linear viscoelasticity of bulk samples reveals that none of the shear-history and low-frequency terminal response complications observed for diblock copolymers are observed for graft copolymers, owing to their lack of long-range order.
Introduction The equilibrium domain structures of block copolymers have been studied extensively both theoretically and experimentally. The microphase-separated bulk morphology of diblock copolymers with chemically dissimilar blocks is determined by f, the fraction of monomers of type A in an A-B diblock copolymer.1,2 In a thin film, a strong orientational aspect is additionally introduced owing to the constraint imposed by the two interfaces (air/copolymer, copolymer/substrate), the surface energy difference between the blocks, and the chemical affinity of one of the blocks for the substrate.3-5 If the orienting potential is strong, the localized segregation of the chemically dissimilar block has the effect of orienting the microdomains parallel to the surface and only film thicknesses corre* To whom correspondence should be addressed. (1) Thomas, E. L.; Alward, D. B.; Kinning, D. J.; Handlin, D. L.; Fetters, L. J. Macromolecules 1986, 19, 2197. (2) Hasegawa, H.; Tanaka, H.; Yamasaki, K.; Hashimoto, T. Macromolecules 1987, 20, 1651. (3) Green, P. F.; Christensen, T. M.; Russell, T. P.; Jeˆro´me, J. J. Chem. Phys. 1990, 92, 1478. (4) Russell, T. P.; Menelle, A.; Anastasiadis, S. H.; Satija, S. K.; Majkrzak, C. F. Macromolecules 1991, 24, 6263. (5) Liu, Y.; Zhao, W.; Zheng, X.; King, A.; Singh, A.; Rafailovich, M. H.; Sokolov, J.; Dai, K. H.; Kramer, E. J.; Schwarz, S. A.; Gebizlioglu, O.; Sinha, S. K. Macromolecules 1994, 27, 4000.
sponding to a multiple of the domain spacing are stable.6-8 Any excess material is pushed out, forming islands on the surface approximately one lamella in height.9,10 Graft copolymers, A-g-B, are “macromolecular combs” composed of a linear backbone A onto which pendant side B block(s) are grafted. The two blocks A and B are generally incompatible. These graft copolymers are technologically important, particularly in their ability to compatibilize polymer blends11 and to form thermoplastic elastomers.12 Indeed, they constitute the primary compatibilizers in commercial applications. Compared with block copolymer films, the case of graft copolymer films is more complicated and has not been studied extensively. The aggregation structures depend on a number of parameters such as the graft level and their distribution along the backbone, the relative molecular weights of the backbone, and the (6) Fredrickson, G. H. Macromolecules 1987, 20, 2535. (7) Bates, F. S.; Fredrickson, G. H. Annu. Rev. Phys, Chem. 1990, 41, 525. (8) Anastasiadis, S. H.; Russell, T. P.; Satija, S. K.; Majkrzak, C. F. Phys. Rev. Lett. 1989, 62, 1852. (9) Coulon, G.; Collin, B.; Ausserre, D.; Chatenay, D.; Russell, T. P. J. Phys. (Paris) 1990, 51, 2801. (10) Coulon, G.; Ausserre, D.; Russell, T. P. J. Phys. (Paris) 1990, 51, 777. (11) Peiffer, D. G.; Rabeony, M. J. Appl. Polym. Sci. 1994, 51, 1283. (12) Legge, N. R.; Holden, G.; Schroeder, H. E. In Thermoplastic Elastomer; Macmillan: New York, 1987.
10.1021/la980556o CCC: $18.00 © 1999 American Chemical Society Published on Web 03/19/1999
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Table 1. Summary of Graft Copolymers Used samples
Mw,dPEA
Mw,PS
wt % PS
no. of PS grafts
dPEA-g-1PS dPEA-g-3PS dPEA-g-5PS
136 000 108 000 78 000
15 000 15 000 15 000
9.3 28.4 48.0
0.9 2.8 4.8
pendant grafts. The number of structural variables is large compared with conventional diblock and triblock copolymers. Furthermore, the syntheses of these materials (macromonomer techniques13 and graft onto11 and graft from 11 syntheses) lead to a statistical placement of the graft onto the chain backbone. This random placement of the grafts endows these copolymers with a somewhat less defined molecular architecture as compared with anionically polymerized block copolymers. It appears that microphase segregation occurs between the blocks, but no distinct morphologies are observed with long range order. Gido and co-workers14,15 were able to synthesize a series of graft copolymers of monodisperse molecular weight and uniform, well-defined positioning of grafts. In contrast to randomly grafted copolymers, long-range order was observed in these copolymers. The results showed that if the graft copolymer structure could be approximated by a series of Y-shaped structure, the agreement could be obtain with the phase diagram calculated by Milner for star-shaped polymers.16 Polydisperse graft copolymers with random grafting points are more commonly used in various commercial applications since they are easy to produce and inexpensive. Therefore, it is very important to understand the phase and rheological behavior of these polydisperse copolymers as well. In this paper, we focus on the surfaceinduced ordering of graft copolymers with different numbers of grafted chains (or f, fraction of monomers) synthesized by a macromonomer process since these copolymer do not order easily in the bulk. These graft copolymers have deuterated polydisperse poly(ethyl acrylate) (dPEA) backbones with pendant chains of monodisperse polystyrene (PS) attached randomly along their length. The equilibrium domain structure formed by these graft copolymers in thin films was studied using neutron reflection (NR) and secondary ion mass spectrometry (SIMS). The associated surface topology was studied with atomic force microscopy (AFM). The viscoelasticity of bulk samples was investigated with oscillatory shear. Experimental Section The graft copolymers used in this study are listed in Table 1. The dPEA backbone (synthesized by free-radical polymerization) had a polydispersity index Mw/Mn ) 2.5. This polydispersity and the randomness of the grafting points made these polymers somewhat ill-defined but considerably more predictable in structure than those produced by in-situ grafting. Thin films of graft copolymers of thickness ranging from 400 to 20 000 Å were spun cast from toluene solution onto HF-treated polished silicon substrates. These samples were then annealed at T ) 450 K for 24 h in a vacuum of 10-4 Torr. The samples used for SIMS17 were prepared by floating on top of the samples a “sacrificial” layer of PS, approximately 300 Å thick, prior to analysis. This extra layer allows the ion-beam-sputtering rate to come to a steady state before the sample surface is reached. The sputtering was (13) Meijs, G. F.; Rizzardo, E. Macromol. Chem. Phys. C 1990, 30, 305. (14) Lee, C.; Gido, S. P.; Poulos, Y.; Hadjichristidis, N.; Tan, N. B.; Trevino, S. F.; Mays, J. W. J. Chem. Phys. 1997, 107, 6460. (15) Lee, C.; Gido, S. P.; Poulos, Y.; Hadjichristidis, N.; Tan, N. B.; Trevino, S. F.; Mays, J. W. Polymer 1998, 39, 4631. (16) Milner, S. T. Macromolecules 1994, 27, 2333. (17) Schwarz, S. A.; Wilkens, B. J.; Pudensi, M. A. Mol. Phys. 1992, 76, 937.
Figure 1. SIMS depth profile of the dPEA volume fraction (proportional to the D-ion intensities, dotted line) and concentration profile of dPEA used to fit the NR data shown in Figure 2 (solid line) for a dPEA-g-3PS graft copolymer film, 877 Å thick, annealed at 450 K for 24 h. performed on an Atomika 3000-30 ion microprobe using a 2 keV, 20 nA beam of Ar+ ions at 30° off normal incidence, rastered over a 0.25 mm2 region. Negative ions of C, D, CN, CH, CD, and O were monitored. At a sputtering rate of approximately 500 Å/h, we obtain a depth resolution corresponding to a Gaussian with full width half-maximum of 100 Å. Neutron reflection data were obtained at the H-9A reflectometer at the Brookhaven National Laboratory High Flux Beam Reactor. The neutron wavelength was fixed by pyrolitic carbon monochromators to be 4.16 ( 0.02 Å, and slit settings were adjusted to fix the resolution at a constant value of ∆qZ/qZ ) 0.02, where qZ is the momentum transfer perpendicular to the sample surface. The details of the NR technique have been considered elsewhere.18 AFM was performed on the surfaces of the ordered films using a Nanoscope III instrument in the contact mode, with a Si3N4 tip and a force of approximately 18 nN. A Rheometrics RMS-800 rheometer was used for the viscoelastic measurements with 25 mm diameter parallel circular plates and roughly 2 mm sample heights. Measurements were made at two temperatures, 390 and 428 K. Time-temperature superposition19 was used to shift the data at 390 K to 428 K. This empirical superposition is known to apply to homopolymers19 and the isotropic phase of block copolymers.20
Results and Discussion The SIMS depth profile for a dPEA-g-3PS graft copolymer film (877 Å) annealed at 450 K for 24 h is shown in Figure 1 (dotted line). The dPEA volume fraction is assumed proportional to the D- ion yields. From the figure, it can be seen that the sample is well ordered in a direction perpendicular to the Si substrate surface. The lower height of the first half-lamella at the vacuum interface is just due to convolution with the SIMS instrumental resolution function. The much smaller amplitude of the last layer is indicative of a loss of alignment at the Si interface. To obtain a more accurate determination for the structure of the respective layers and the interfaces between them, neutron reflectivity measurement were done on the same sample. The data are shown in Figure 2, where we plot the reflectivity, R, and the ratio of reflected to incident neutron intensity, vs kZ, the component of the neutron wavevector normal to the surface. While NR data of high quality can be rather straightforward, the analysis of the results is often not. This is especially true with multilayered or quasiperiodic samples. The lack of phase information can make a unique interpretation of NR data (18) Russell, T. P. Mater. Sci. Rep. 1990, 5, 171. (19) Ferry, J. D. Viscoelastic Properties of Polymers, 3rd ed.; Wiley:, 1980. (20) Rosedale, J. H.; Bates, F. S. Macromolecules 1990, 23, 2329.
Surface-Induced Ordering in Graft Copolymer Thin Films
Figure 2. Neutron reflectivity vs qZ, corresponding to the sample in Figure 1.
nearly impossible in a complex sample. If the profile of the sample is not periodic or uniform in depth profile, it is impossible to determine with confidence how the periodicity or nonuniformity is distributed in the sample. By using a lower resolution real-space probe (i.e., SIMS), however, we can more correctly choose a model with which to describe the depth profile. We fitted the NR data by dividing the film profile into a histogram of slabs of varying thickness, neutron scattering length density, and roughness (taken to be Gaussian). We chose the starting parameters to approximate those obtained from the SIMS. We also required that fitting parameters be physically reasonable. For example, the profile must preserve the total integrated material, no slab could be more than 100% of one component, and roughness must be less than the slab thickness. The profile that provided the best fit to the data is shown in Figure 1 as a solid line. From the profile, dPEA-g-3PS appears to have dPEA-PS lamellar structure with alternating layers of PS and dPEA. A nearly pure dPEA layer is situated next to the vacuum surface, since dPEA is the component with lower surface tension. The following layers alternate between pure dPEA layers and PS-rich layers. The last lamella adjacent to the Si interface appears to be somewhat disordered with alternating layers containing 80% and 40% PS, respectively. The lamellar period is 250 Å. The PS layers have thicknesses of ≈100 Å, and the dPEA layers have thicknesses of ≈150 Å. The interfacial width between the layers, ≈30-40 Å, is also consistent with the interfaces previously reported between lamellar structures in diblock copolymers with similar χN parameters.4,18 If we take the radius of gyration (Rg) of a Gaussian PS polymer with a MW of 15k to be ≈30 Å, then we can see that the PS layers here are stretched to almost double this length. On the other hand, the total MW of dPEA in a chain of the 28 wt % PS material is 108k, leading to an Rg ≈ 80 Å. The dPEA layers are therefore only slightly thinner than what we would find for Gaussian chains. Figure 3 shows the SIMS depth profiles for a 17 350 Å thick dPEA-g-3PS graft copolymer film annealed at 450 K for 24 h. Comparing with Figure 1, we can see that the surface-induced ordering exists only at the polymer vacuum interface and decays after approximately 15 periods (∼0.5 µm) into the bulk of the film. As the film thickness increases, each lamellar plane must distort less to accommodate the extra material in the surface layer; consequently, the surface-induced orientation can decay after several lamellar periods. No orientation is observed at the Si interface, in contrast to the PS-PMMA, PS-
Langmuir, Vol. 15, No. 8, 1999 2913
Figure 3. SIMS depth profiles of the dPEA volume fraction for the dPEA-g-3PS graft copolymer film, 17 350 Å thick, annealed at 450 K for 24 h.
Figure 4. SIMS depth profile fraction of the dPEA volume (dotted line) and concentration profile of dPEA used to fit the NR data shown in Figure 5 (solid line) for a dPEA-g-5PS graft copolymer film, 828 Å thick, annealed at 450 K for 24 h.
Figure 5. Neutron reflectivity vs qZ, corresponding to the sample in Figure 4.
PVP diblock copolymer,5 where the non-PS block prefers the Si surface. Figure 4 shows the SIMS data (dotted) for a dPEA-g5PS graft copolymer film with thickness of 828 Å annealed at 450 K for 24 h. From the figure, it can be seen that the sample is also well ordered in a direction perpendicular to the surface. The neutron reflectivity data are shown in Figure 5, and the profile that provided the best fit to the data is shown in Figure 4 (solid line). In this case, the layers in the interior of the film alternate between 80% PS and about 65% dPEA, and the layer spacing is 190 Å. The roughness of these layers is about half the thickness
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Figure 6. AFM image and height profile along the line in the AFM image for the dPEA-g-3PS graft copolymer film annealed at 450 K for 24 h.
of the dPEA layers. The results show that the dPEA-g5PS appear to have cylindrical structures with a PS matrix containing dPEA cylinders. Both the dPEA and the PS layers in the interior are thinner than those in the lamellar profiles of dPEA-g-3PS, with the PS layers getting close to their Gaussian thickness. The dPEA layers are considerably thinner than Gaussian. If we assume that the dPEA-g-5PS is a completely phase-separated system and the dPEA cylinders have hexagonal structure as shown in Figure 4, we calculated the radius of the cylinders and maximum volume fraction of dPEA from the measured layer spacing to be 80 Å and 70%, respectively. Considering the partially phase mixing on the PS side of the dPEAPS interface, we can obtain the maximum volume fraction for dPEA to be close to 65%, in good agreement with the value measured with SIMS and neutron reflection. The layers adjacent to the vacuum and Si interface differ from the others. In addition to having a different thickness, the in-plane structure of these layers is consistent with pure and partial lamellar morphology at vacuum and Si interfaces, respectively. This phenomenon was also observed for PVP-dPS-PVP triblock copolymer films ordered on a native oxide covered silicon substrate.5 When the difference between the surface interaction energies of two components is very large, the lamellar configuration is most efficient in minimizing the surface tension. Figure 6 shows the AFM image for the dPEA-g-3PS graft copolymer film (828 Å) annealed at 450 K for 24 h. From the figure, we see that the thin films of graft copolymers exhibited island/hole structures similar to those previously reported in linear block copolymers.5 In Figure 6b, we show the profile of the dPEA-g-3PS. The contour of the islands is relatively sharp, indicating a high
Ge et al.
degree of ordering within the film. Similar features are also seen for dPEA-g-5PS. The height/depth of islands/ holes in films of PEA-g-1PS, PEA-g-3PS, and PEA-g-5PS measured by AFM are 300, 250, and 190 Å, respectively. The previous studies show that when a block copolymer system orients with the microdomains parallel to the surface, the excess material will be pushed out and form islands/holes on the surface approximately one lamella in height.9,10 The results of AFM observation match the SIMS and NR measurements well. The morphology of block copolymers with chemically dissimilar blocks is determined by f, the fraction of monomers of type A in an A-B block copolymer.1,2 For symmetric (f ) 1/2) block copolymer, the interfacial tension is minimized by arranging the copolymers in a lamellar structure. As f decreases, the morphologies with the minimum interfacial free energy change from bicontinuous double diamond (0.28 < f < 0.34) to cylindrical (0.13 < f < 0.28) and to spherical (f < 0.13). In the case of graft copolymers, the thin films of dPEA-g-1PS (f ) 0.09),21 dPEA-g-3PS (f ) 0.28), and dPEA-g-5PS (f ) 0.48) show spherical, lamellar, and cylindrical structures, respectively. On comparison with a known phase diagram of diblock copolymer, it is clear that a phase shift occurs because of the different chain architecture. The phase diagram shift is expected to be determined by the number of grafts16 and will be complicated by polydispersity.22 Through the use of the constituting block copolymers, Milner’s calculation predicts the morphology of dPEA-g3PS and dPEA-g-5PS would be cylindrical and lamellar structures, respectively.16 The disagreement between our experimental results and the theoretical calculation can be attributed to the polydispersity of the backbone molecular weight between grafts, which leads to a lessthan-expected curvature of the interface toward the PS graft component. In addition to this effect, owing to the difference in the reactivity ratios of the reactants (γ2 ≈ 0.7), there are probably many attachments where the molecular weight between grafts is very small, especially toward the end of the chain. This situation will lead to more phase mixing on the PS side of the dPEA-PS interface. For incompatible A-B block copolymers, the dependence of the micelle spacing and interfacial width on the properties of the polymers (such as molecular weight, segment length, and interaction parameter) are well understood in terms of mean field theory in the strong segregation limit. The principal results for the lamellar and cylindrical micelle spacings, h, are given by 23-25
hlamellar ) 1.10aN1/2(χN)1/6
(1)
hcylinder )
2.06 aN1/2(χN)1/6 1/3 (1.645 - ln f)
(2)
hsphere )
1.78f1/3 aN1/2(χN)1/6 1/3 1/3 (1 - 0.57f )
(3)
The derivation assumes that the two polymers have similar statistical segment lengths, a (which is the case for dPEA and PS, 6.8 Å), and the Flory-Huggins interac(21) Rabeony, M.; Peiffer, D. G.; Behal, S. K.; Disko, M.; Dozier, W. D.; Thiyagarajan, P.; Liu, Y. J. Chem. Soc., Faraday Trans. 1995, 91, 2855. (22) Erukhimovich, I.; Dobrynin, A. V. Macromol. Symp. 1994, 81, 253. (23) Semenov, A. N. Macromolecules 1992, 25, 4967. (24) Semenov, A. N. Sov. Phys. JETP (Engl. Transl.) 1985, 61, 733. (25) Rubinstein, M.; Obukhov, S. P. Macromolecules 1993, 26, 1740.
Surface-Induced Ordering in Graft Copolymer Thin Films
tion parameter, χ for PS-PEA blends is 0.04. N is the total number of monomers in the diblock, and f is the fraction of monomers of type A. If we substitute the values of f from Table 1 for the graft copolymer dPEA-g-1PS (spherical), dPEA-g-3PS (lamellar), and dPEA-g-5PS (cylindrical), we obtain hspherical ) 453 Å, hlamellar ) 547 Å, and hcylinder ) 771 Å, much larger than hspherical ) 295 Å,21 hlamellar ) 250 Å, and hcylinder ) 190 Å measured with SIMS and neutron reflection. The disagreement between the experimental results and the theoretical calculation increases with the number of grafts. Furthermore, if we approximate each dPEA-g-xPS (x, number of grafted chains ) 3, 5) graft copolymer chain as x diblock chains with the molecular weight of Mw,total/x, we obtain micelle spacing hlamellar for dPEA-g-3PS to be 263 Å, in good agreement with the value measured with SIMS and neutron reflection. For dPEA-g-5PS calculated hcylinder ) 264 Å is still larger than the measured value and indicates that the backbone of dPEA-g-5PS is more stretched than that of dPEA-g-3PS.26 Figure 7 shows the rheological data for the dPEA-gxPS graft copolymers at a reference temperature of 428 K. At both temperatures where measurements were made, strain amplitude variations at constant frequency indicated a wide range of strain amplitudes at low strains where the complex modulus is independent of strain amplitude. The range of this linear viscoelastic response was similar to that observed for the isotropic materials (i.e., homopolymers) and markedly different from the ordered phases of block copolymers, which usually do not exhibit any range of linear viscoelasticity.27 Samples that were presheared at 1 s-1 shear rate had identical linear viscoelastic response to samples that had been annealed in a compression mold. This insensitivity to shear history is also in sharp contrast to the behavior of ordered block copolymers28 and is again more akin to the response of isotropic materials. For each sample, the low-frequency response is that of a viscoelastic liquid, with the storage modulus G′ and loss modulus G′′ approaching their limiting terminal frequency dependence (G′ ∝ ω2 and G′′ ∝ ω) characteristic of liquids.19 All isotropic polymer liquids exhibit these characteristics, whereas ordered phases of block copolymers do not.29,30 It is clear from the results that this family of graft copolymer exhibits only weakly ordered microstructures with no long-range order in bulk samples. Even in thin films, where there is a highly selective surface to provide an orientational anchor, the order persists with an extinction depth of at best about ∼15 lamellar spacings. Rheological studies show that the viscoelastic response of these graft copolymers in their ordered phase is qualitatively similar to isotropic polymer liquids and markedly different from the ordered phases of block copolymers. Attempts to shear-align samples achieved no success. These graft copolymers simply do not appear to be capable of long-range order. We estimate the largest length scale for order from the modulus scale at which the terminal response starts, G(τ) = 3 × 104 dyn/cm2, l = (kT/G(τ))1/3 = 100 Å. This suggests that long-range order in bulk (26) The disagreement between the measured and calculated values for the f ) 0.09 single graft copolymer illustrates the effect of a junction point along the backbone in a Y-shaped molecule. Even though ref 16 calculates the phase diagram for this structure, no prediction is given as yet for micelle spacing h. (27) Larson, R. G.; Winey, K. I.; Patel, S. S.; Watanabe, H.; Bruinsma, R. Rheol. Acta 1993, 32, 245. (28) Colby, R. H. Curr. Opin. Colloid Interface Sci. 1996, 1, 454. (29) Ghijsels, A.; Raadsen, J. Pure Appl. Chem. 1980, 52, 1359. (30) Koppi, K. A.; Tirrell, M.; Bates, F. S.; K., A.; Colby, R. H. J. Phys. II 1992, 2, 1941.
Langmuir, Vol. 15, No. 8, 1999 2915
Figure 7. Viscoelastic response of the dPEA-g-xPS graft copolymer at 428 K (filled symbol) and 390 K shifted to 428 K using time-temperature superposition (open symbol). Storage modulus G′ are circles and loss modulus G′′ are squares.
samples is lost at the scale of the size of a single graft copolymer chain. This estimate is consistent with lack of long-range ordering observed in transmission electron microscopy cross-sectional scan of the graft copolymer and the lack of structure on the SANS spectra.21 Conclusion Phenomena of surface-induced ordering of PEA-PS graft copolymer were investigated by a variety of techniques. The study reveals information on properties such as whether the ordered structure is spherical, lamellar, or cylindrical for each composition of the graft copolymers. NR results are well correlated with SIMS profiles of the ordered graft copolymer thin films, indicating substantial long-range order near the vacuum surface. However, roughly 0.4 µm away from the surface (∼15 lamellar spacings) the long-range order is lost. There results are consistent with our rheological measurements, which show that the graft copolymers in their “ordered” phase respond like simple viscoelastic liquids that cannot be shearaligned, indicating that there is no long-range order. Acknowledgment. Support from the NSF (DMR9316157), NSF-MRSEC (DMR9632525), and DOE (DEFG02-93ER45481) is gratefully acknowledged. LA980556O
