Elastic Mechanical Response of Thin Supported Star-Shaped Polymer

Mar 15, 2016 - We show evidence of thickness-dependent elastic mechanical moduli that are associated largely with the effects of architecture (topolog...
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Elastic Mechanical Response of Thin Supported Star-Shaped Polymer Films Peter C. Chung,† Emmanouil Glynos,†,‡ Georgios Sakellariou,§ and Peter F. Green*,† †

Department of Materials Science and Engineering, Biointerfaces Institute, University of Michigan, Ann Arbor, Michigan 48109, United States § Department of Chemistry, University of Athens, Panepistimiopolis, Zografou, 15771, Athens Greece S Supporting Information *

ABSTRACT: We show evidence of thickness-dependent elastic mechanical moduli that are associated largely with the effects of architecture (topology) and the overall shape of the macromolecule. Atomic force microscopy (AFM) based nanoindentation experiments were performed on linear chain polystyrene (LPS) and star-shaped polystyrene (SPS) macromolecules of varying functionalities (number of arms, f) and molecular weights per arm Mwarm. The out-of-plane elastic moduli E(h) increased with decreasing film thickness, h, for h less than a threshold film thickness, hth. For SPS with f ≤ 64 and Mwarm > 9 kg/mol, the dependencies of E(h) on h were virtually identical for the linear chains. Notably, however, for SPS with f = 64 and Mwarm = 9 kg/ mol (SPS-9k-64), the hth was over 50% larger than that of the other polymers. These observations are rationalized in terms of the structure of the polymer for high f and sufficiently small Mwarm and not in terms of the influence of interfacial interactions.

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Long and co-workers, relying on free volume based dynamic heterogeneity models, correctly predict these trends. They moreover describe the general case of the linear elastic behavior of linear-chain polymer films confined between two substrates for varying polymer/substrate intermolecular interactions.7 Nanoindentation measurements of polymer thin films supported by stiff (noncompliant) substrates, performed using and atomic force microscope (AFM), reveal that the effective out-of-plane modulus E(h) increases with decreasing h, for film thickness h smaller than a threshold film thickness ht.8−13 This enhancement of E(h < hth) is due to the propagation of the indentation-induced stress field, from the free surface of the polymer in the direction normal to the polymer/substrate interface, throughout the entire polymer film and interacting with the underlying substrate.9,11 This is the so-called substrate ef fect. Nanoindentation stress fields, even for indentations of a few nanometers, extend hundreds of nanometers throughout a film. We recently showed that the length-scale of the propagation of the externally imposed stress field is correlated with the local vibrational force constants (i.e., local chain

he thickness-dependent behavior of the physical properties of polymers, with thicknesses that range from nanometers to as much as hundreds of nanometers, is of considerable scientific and technological interest. This behavior impacts diverse applications, from membrane separation to nonvolatile memory and plastic electronic applications. Generally, the thickness-dependent behavior stems from a combination of intermolecular interactions at the external interfaces and entropic effects, which have the effect of inducing subtle changes in local structure, with more significant implications on the physical properties at larger length-scales. Most notable examples include the glass transition temperature, phase separation, and dynamics.1,2 Here we are interested in understanding the mechanical properties of thin polymer films. Molecular dynamics (MD) simulations reveal that the elastic moduli of freely standing linear-chain polymer films with thicknesses less than 100 nm are lower than that of the modulus of the bulk.3,4 This behavior is due to the enhanced configurational freedom of polymer segments at the free surface compared to that of the bulk. Additionally, buckling experiments of the in-plane elastic moduli of polystyrene (PS) and poly(methacrylate) thin films, supported by soft (low moduli) substrates, indicate that the elastic moduli decrease with decreasing film thickness h, for thicknesses h < 40−80 nm.5,6 © XXXX American Chemical Society

Received: December 28, 2015 Accepted: March 8, 2016

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DOI: 10.1021/acsmacrolett.5b00944 ACS Macro Lett. 2016, 5, 439−443

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ACS Macro Letters stiffness) of the specific polymer.11,12 For this reason, ht is not a constant; it is a function of compliance of the polymer. Here we report nanoindentation measurements of thin film multiarm stars with a wide range of functionalities 2 < f < 64. We show for the first time the influence of molecular architecture on the out-of-plane elastic mechanical response of star-branched polystyrene (SPS) films supported by silicon substrates. We suggest that the behavior is based on differences in molecular architecture/packing of the stars, and not due to interfacial, free surface or substrate effects. These observations compliment our earlier studies showing how compliance of the polymer based on chemical composition could account for differences in the elastic mechanical response. Star-shaped macromolecules exhibit physical properties that differ appreciably from those of their linear-chain analogs; notable examples include the glass transition temperatures,14,15 surface wetting properties,16,17 and physical aging.2,18,19 Much of this behavior stems from entropic effects, inherently associated with the molecular architecture and overall molecular shape, responsible for subtle, yet consequential, changes in segmental “packing” and, hence, the density. A typical star-shaped macromolecule is composed of a central particle (branch point) onto which f (the functionality) chains (arms) are attached. Macromolecular segments close to the branch point constitute the core region, within which the segments are stretched, compared to segments at the ends of the arms; so segments that constitute the core region suffer larger conformational entropy losses compared to segments outside the core region, composing the corona. The size of the core increases as f1/2. The monomer density inside the core region is necessarily higher than that of the corona, and it increases with increasing f.20−24 The existence of the core region is responsible for entropic repulsion between the macromolecules, which increases with increasing f. This entropic repulsion is responsible for the tendency of the star-shaped macromolecules to exhibit an increasing degree of positional order, with increasing f and decreasing weight per arm, Mwarm.15,17 This was particularly evident in very high functionality stars, as revealed by atomic force microscopy of thin films, showing that the macromolecules self-assemble like spherical particles (or colloids).2,15 The interactions of star-shaped polymers at interfaces are very different from their linear chain analogs, leading to different wetting properties on surfaces and thickness dependent properties.2,15,17,24 A single, isolated, star-shaped molecule adsorbs more strongly to the same surface than its linear chain analog of the same degree of polymerization. The strength of the adsorption increases with increasing f.25 This is attributed to the increasing number of monomers in contact with the surface with increasing f and the decreasing entropic penalty to do so; both indicating a gain in the free energy compared to the linear analogues. In a melt, where multiple molecules are adsorbed, there exists a competition between the gain in free energy from surface adsorption of a large number of monomers/molecule and the unfavorable entropic repulsion between the macromolecules that are “packed” on the surface.2,15 The entropic repulsion is associated with the loss of compliance of the core region with increasing f. These effects are manifested in the behavior of the anisotropy in the radii of gyration of the macromolecules, normal to the substrate Rgz, the wetting properties (a minimum in the macroscopic contact angle14), and the glass transition.14,15 In light of the influence of these entropic forces, responsible for qualitative and quantitative differences between the physical

properties of star-shaped and linear chain polymers, it would be important to investigate and to understand the thin film elastic mechanical response of star-shaped macromolecules. To this end, we investigate the elastic mechanical response of thin starshaped PS (SPS) films supported by silicon oxide (SiOx) substrates. Nanoindentation measurements, using atomic force microscopy (AFM), were performed on a series of SPS films of varying f (2 < f < 64) and Mwarm (7 kg/mol < Mwarm < 140 kg/ mol) with film thicknesses: 200 nm < h < 900 nm. We investigated the following two situations: (1) The effect of functionality; this was performed by considering SPS where the Mwarm remained fixed at ∼10 kg/mol, while f varied from 2 to 64; (2) The effect of the Mwarm; This was performed by studying SPS with f fixed at 64, while Mwarm changed from 9 to 140 kg/mol. The terminology that will be used throughout is SPS-f-Mwarm, so SPS-64-9k denotes star shaped polystyrene with 64 arms and Mwarm = 9 kg/mol. Dokukin and Sokolov showed that nanoindentation measurements provide reliable results if the tip is not excessively sharp, thereby avoiding nonlinearities in the stress−strain behavior, and second, if adhesion between the tip of the indentation probe and the polymer surface are accounted for in the analysis, avoiding artifacts such as the “skin-effect”.26 The “skin-effect” artifact is the depth dependence of the modulus. For this reason, we used the following experimental protocols.11,12 A typical force−distance (FD) curve is shown in Figure 1; the

Figure 1. Typical FD curve obtained from an AFM nanoindentation measurement of h ∼ 870 nm thick 64-arm SPS (Mwarm = 9 kg/mol) film is shown: approach curve (open squares), retraction curve (open circles), and the JKR fitting (solid line).

approach (represented by open black squares) and retraction (represented by open blue circles) curves are virtually identical; the absence of a hysteresis effect is indicative of the fact that the deformation is elastic.27 The effective elastic moduli E(h) and indentation depths d were estimated from the FD curves by fitting the retraction curve (represented by solid line) using the prediction of the Johnson-Kendall-Roberts (JKR) model,28 which provides a relationship between a the contact radius of the tip, R the radius of the indenter, K the reduced modulus, P the applied force, and w the work of adhesion. a3 =

R [P + 3πRw + (6πRwP + (3πRw)2 )1/2 ] K ⎛ 8πwa ⎞ a2 ⎟ −⎜ ⎝ 3K ⎠ R

(1a)

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(1b) DOI: 10.1021/acsmacrolett.5b00944 ACS Macro Lett. 2016, 5, 439−443

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ACS Macro Letters The effective modulus, E, was estimated from K using the following equation: 1/K = (3/4)[(1 − νfilm2)/Efilm + (1 − νindenter2)/Eindenter], where νfilm is the Poisson’s ratio of the film, νindenter is the Poisson’s ratio of the indenter, Efilm is the elastic modulus of the film, and Eindenter is the elastic modulus of the indenter. Since Eindenter ≫ Efilm, 1/K ≈ (3/4)[(1 − νfilm2)/ Efilm].29 The detachment of the indenter from the surface of sample occurs when the pull-off force is larger than Ppull = −3/ 2(πRw).28 The work of adhesion w was estimated from the corresponding FD curve. A value for the Poisson’ ratio of ν = 0.33 was used for PS.30 A fixed maximum force of Pmax = 400 nN was used for all the nanoindentation measurements. The elastic moduli were independent of indentation rates (from 20 to 80 nm/s), confirming the absence of hysteresis effects and therefore the absence of viscoelastic effects. Therefore, all measurements were performed under elastic deformation conditions. The elastic mechanical moduli E(h) increased with decreasing h, for h < hth, for all films that were investigated, as expected. Notably, however, the dependence of E(h) on h, for the SPS-64-9k molecule was significantly different: hth was over 50% larger than that of the other polymers and E(h) exhibited a stronger dependence on h in the regime h < hth. This behavior, as discussed below, would be related to important differences between the structures of these macromolecules. The moduli E(h) are plotted as a function of h in Figure 2a, for all the SPS films. In all cases, E(h) increased with decreasing h for h < hth, as expected and is understood in terms of substrate effect: the indentation-induced stress field propagates throughout the entire polymer film and interacts with the noncompliant substrate, leading to an enhanced local stress field, and hence the increased magnitude of E.9,11 The E(h) for the SPS-64-9K SPS molecule increases at approximately 50% higher value for h < hth, compared to the SPS films with lower functionalities. More importantly, and noteworthy, is that the hth for this polymer is larger than that of the other polymers: hth(SPS-64-9K) ∼ 650 nm, whereas hth(SPS-4-7K) ∼ hth(SPS8-14K) ∼ hth(linear-PS-6K) ∼ 400 nm. In order to compare the extent of the substrate effect, the values of E(h) (Figure 2a) are normalized with average effective moduli for h > hth, E(h)/E(h > hth) and plotted in Figure 2b as a function of ratio of contact radius to film thickness, a/h. The a/h dependencies of moduli for SPS-4-7K and SPS-8-14K films are virtually identical, and the trends are similar to that of linear-PS-6K. However, the significantly deferent behavior of the SPS-64-9K molecule is even more apparent in this plot. This effect might initially appear to be exclusively due to the fact that the functionality of the molecule is comparatively very high f = 64; so it was valuable to consider the response of other SPS molecules of the same functionality f = 64, but with longer arms. Such data are plotted in Figure 3a,b for a series of 64-arm SPS films with different values of Mwarm: 9, 36, and 140 kg/mol. With the exception of the SPS-64-9K molecules, for which hth(SPS-64-9K) ∼ 650 nm, the threshold thicknesses for the other 64-arm SPS molecules are virtually the same as those of the other polymers: hth(SPS-64-36K) ∼ hth(SPS-64-140 K) ∼ hth(LPS-6K) ∼ 400 nm. This point is even more apparent in Figure 3b, where E(h)/E(h > hth) is plotted as a function of a/ h, consistent with the fact that the propagation of the indentation imposed stress field is quantitatively different in the SPS-64-9k films than the others. This effect is associated

Figure 2. (a) Effective moduli, E(h), for linear chain PS (LPS; Mw = 6 kg/mol), 4-arm SPS (Mwarm = 7 kg/mol), 8-arm SPS (Mwarm = 14 kg/ mol), and 64-arm SPS (Mwarm = 9 kg/mol) films are plotted as a function of film thickness, h. (b) E(h) normalized with average effective moduli for h > hth, E(h > hth) are plotted as a function of ratio of contact radius to film thickness, a/h. Note that maximum force, Pmax, was fixed at a constant value of 400 nN for all the nanoindentation measurements. Each data point is an average of 15 measurements, and dashed lines are guides for the eyes.

with a combination of high functionality and short arms, not one or the other. Generally, the elastic mechanical behavior of thin films is understood to be largely due to the influence of intermolecular interactions at the external interfaces. Recall that buckling experiments show that the elastic modulus is a function of thickness, for films in the thickness range of less than 40−80 nm.5,6 This behavior is fundamentally related to interfacial processes; such processes also influence the glass transition of thin films. The reduction of the modulus of PS thin films, with decreasing film thickness is associated with the increase in configurational freedom of molecules at the free surface, beyond that of the bulk, as shown by simulations and theory based on the notion of dynamic heterogeneity. This is the same reason that the average glass transition temperatures of PS films decrease with decreasing h, for h less than approximately 60 nm. In fact, the Tg at the free surface Tgsurface is lower than that of the bulk Tgbulk, that is, Tgsurface < Tgbulk. In the case of PS, the glass transition temperature at the substrate Tgsubstrate is comparable to Tgbulk, that is, Tgsubstrate ∼ Tgbulk,1,31 so the average Tg of this polymer decreases with decreasing h. The glass transition behavior of SPS molecules is appreciably different from linear chain PS. To begin with, the dependence of the bulk Tg on chain length, for low molecular weight polymers, is understood in terms of chain-end effects; chain ends contribute excess free volume. However, as shown by 441

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in our experiments. The length scales over which interfacial interactions would affect the modulus is on the order of nanometers,7,36 whereas in nanoindentation experiments (with a stress field normal to the interfaces of a supported film), the actual stress field extends hundreds of nanometers into the film, even for indentation depths of d < 3 nm nanometers.9,11 The indentation-induced stress field could extend over ∼450 nm for indentation depths of d < 3 nm in polymers.11,12 Therefore, these nanoindentation experiments are not sensitive to interfacial effects; they provide information about an out-ofplane modulus for the entire film. The buckling experiments, on the other hand, are sensitive to the effects of interfaces, which is significant for thinner films. It is evident from all the measurements for SPS molecules, 2 < f ≤ 64 and varying Mwarm, that the behavior of the SPS-64-9K molecule is exceptional. The glass transition of this molecule was also exceptional: Tgsurface and Tgsubstrate were comparable to Tgbulk, that is, Tgsurface ∼ Tgsubstrate ∼ Tgbulk.15 The average Tgs of films made of these polymers did not exhibit a dependence on h. The rationalization for the very different E(h)/E(h < hth) versus a/h response exhibited by the SPS-64-9K films, compared to the other samples, must be based on other reasons, particularly the structure. Simulations show that high functionality star-shaped macromolecules pack more efficiently, with higher packing densities, than their longer armed higher molecular weight counterparts.20,24 In fact, as mentioned previously, experiments and simulations showed that “starshaped” macromolecules with sufficiently large f and sufficiently small Mwarm self-assemble to form highly ordered layers across the film due to strong intermolecular inter-core repulsions.37−39 In fact, in a recent publication, AFM topographies of SPS-649K films revealed layered structures with each layer height corresponding to a length scale of ∼2Rg.15 Since a significant enhancement of E(h) and larger hth were exhibited by the SPS64-9k films, with the effect disappearing when Mwarm is increased above 36 kg/mol, we suggest that the strong degree of enhancement is associated with the higher packing densities of SPS-64-9K molecules, accommodating more efficient stress transfer throughout the film. It would not be related to the interactions of the SPS-64-9k polymer with the substrate, as these interactions are weaker than those of the other polymers. In conclusion, the elastic moduli of polymer films supported by SiOx substrates increased with decreasing film thickness h for h < hth due to the propagation of the indentation-induced stress field and its impingement with the noncompliant substrate. While this general behavior exhibited by all SPS molecules, the dependence was quantitatively different for the SPS-64-9K molecule; hth(SPS-64-k) was 50% larger than that of all the polymers examined, indicating that the indentationimposed stress field propagated more effectively through this polymer than the others. This behavior is consistent with the more efficiently “packed” structure of this macromolecule than the others, 20,24 enabling more efficient stress transfer. Molecular simulations would be very useful toward getting a deeper understanding of the structure property behavior of these systems.

Figure 3. (a) Effective moduli, E(h), for LPS and 64-arm SPS (Mwarm = 9, 36, and 140 kg/mol) films are plotted as a function of film thickness, h. (b) E(h) normalized with average effective moduli for h > hth, E (h > hth) are plotted as a function of ratio of contact radius to film thickness, a/h. Note that maximum force, Pmax, was fixed at a constant value of 400 nN for all the nanoindentation measurements. Each data point is an average of 15 measurements, and dashed lines are guides for the eyes.

Chremos et al., the Tg of star-shaped molecules may not be understood in terms of chain-end effects, but effects associated with the molecular topography (architecture) and shape, as these factors dictate packing effects and the density. It should be apparent from the discussion in the first paragraph of this Letter that the segmental density in the core region is higher than in the corona. Consequently, the mobilitiy of segments in the core are slower than near the chain ends (and the corona). Hence, as the arm lengths become smaller, the overall number density of SPS molecules increases; in contrast, for linear chains, the density decreases as the chain length become small. Star-shaped polymers, provided the arms are not too long and f > 4, exhibit enhanced glass transition temperatures, compared to the bulk, at the free surface: Tgsurface > Tgbulk and are functions of f and Mwarm.14,15 This behavior is associated with the fact that star-shaped molecules more strongly adsorb (adsorption of a larger number of monomers for a molecule of the same degree of polymerization) and exhibit spatial order at interfaces than their linear chain analogs, as discussed earlier.17,32−35 Our experiments reveal no connection between the trends in the mechanical response and trends in Tg and the interfacial processes. More specifically, the interactions at the substrate, which would lead to stiffening of the modulus, or the low compliance free surface (especially for linear chains) leading to a low surface modulus are not important factors, to first order,

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EXPERIMENTAL SECTION

LPS was purchased from Pressure Chemical and 4-, 8-, and 64-arm SPS were synthesized by means of anionic polymerization.40,41 The polymers used in this study are listed in Table S1 of the Supporting Information. Thin polymer films were prepared by spin coating 442

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(11) Chung, P. C.; Glynos, E.; Green, P. F. Langmuir 2014, 30 (50), 15200−15205. (12) Chung, P. C.; Green, P. F. Macromolecules 2015, 48 (12), 3991− 3996. (13) Silbernagl, D.; Cappella, B. Scanning 2010, 32 (5), 282−293. (14) Glynos, E.; et al. Phys. Rev. Lett. 2011, 106 (12), 128301. (15) Glynos, E.; et al. Macromolecules 2015, 48 (7), 2305−2312. (16) Glynos, E.; Frieberg, B.; Green, P. F. Phys. Rev. Lett. 2011, 107 (11), 118303. (17) Glynos, E.; et al. Macromolecules 2014, 47 (3), 1137−1143. (18) Frieberg, B.; et al. ACS Macro Lett. 2012, 1 (5), 636−640. (19) Frieberg, B.; Glynos, E.; Green, P. F. Phys. Rev. Lett. 2012, 108 (26), 268304. (20) Chremos, A.; Douglas, J. F. J. Chem. Phys. 2015, 143, 111104−1. (21) Vlassopoulos, D. J. Polym. Sci., Part B: Polym. Phys. 2004, 42 (16), 2931−2941. (22) Daoud, M.; Cotton, J. P. J. Phys. 1982, 43 (3), 531−538. (23) Likos, C. N. Phys. Rep. 2001, 348 (4−5), 267−439. (24) Chremos, A.; Glynos, E.; Green, P. F. J. Chem. Phys. 2015, 142 (4), 044901. (25) Chremos, A. C.; Camp, P. J.; Glynos, E.; Koutsos, V. Soft Matter 2010, 6, 1483. (26) Dokukin, M. E.; Sokolov, I. Macromolecules 2012, 45 (10), 4277−4288. (27) Butt, H. J.; Cappella, B.; Kappl, M. Surf. Sci. Rep. 2005, 59 (1− 6), 1−152. (28) Johnson, K. L.; Kendall, K.; Roberts, A. D. Proc. R. Soc. London, Ser. A 1971, 324 (1558), 301−313. (29) Passeri, D.; et al. Anal. Bioanal. Chem. 2013, 405 (5), 1463− 1478. (30) Nielsen, L. E. Mechanical Properties of Polymers; Van Nostrand Reinhold: New York, 1962. (31) Forrest, J. A.; Dalnoki-Veress, K. Adv. Colloid Interface Sci. 2001, 94 (1−3), 167−196. (32) Striolo, A.; Prausnitz, J. M. J. Chem. Phys. 2001, 114 (19), 8565−8572. (33) Qian, Z. Y.; et al. Macromolecules 2008, 41 (13), 5007−5013. (34) Minnikanti, V. S.; Archer, L. A. Macromolecules 2006, 39 (22), 7718−7728. (35) Kosmas, M. K. Macromolecules 1990, 23 (7), 2061−2065. (36) Xia, W.; Keten, S. Extreme Mechanics Letters 2015, 4, 89−95, http://dx.doi.org/10.1016/j.eml.2015.05.001,. (37) Vlassopoulos, D.; et al. Europhys. Lett. 1997, 39 (6), 617−622. (38) Pakula, T. Comput. Theor. Polym. Sci. 1998, 8 (1−2), 21−30. (39) Pakula, T.; et al. Macromolecules 1998, 31 (25), 8931−8940. (40) Hadjichristidis, N.; et al. J. Polym. Sci., Part A: Polym. Chem. 2000, 38 (18), 3211−3234. (41) Uhrig, D.; Mays, J. W. J. Polym. Sci., Part A: Polym. Chem. 2005, 43 (24), 6179−6222. (42) Hutter, J. L.; Bechhoefer, J. Rev. Sci. Instrum. 1993, 64 (7), 1868−1873.

solutions of LPS and SPS using toluene as the solvent onto oxidized (∼1.7 nm native oxide layer) silicon substrates (Wafer World). Thin films were subsequently annealed under vacuum at a temperature of ∼30 °C above the bulk Tg of the polymer for 2 h for LPS and for at least 24 h for SPS in order to remove residual solvent. Film thicknesses were measured using spectroscopic ellipsometry (JA Woolam, M2000). Nanoindentation measurements were performed using an AFM (Asylum Research, MFP-3D), equipped with a hemispherical AFM tip. All measurements were performed under closed loop mode conditions in order to ensure a constant indentation rate of 40 nm/s. The hemispherical AFM tip (radius, R ∼ 550 nm) was prepared by annealing the AFM probe (NanoWorld, NCH) in air for ∼3.5 h at 1200 °C (Supporting Information, Figure S1). The shape of the tip was imaged prior, and after nanoindentation measurements, with a scanning electron microscope (FEI, Nova Nanolab 200). The sensitivity of the AFM cantilever was calibrated on a mica substrate, and the spring constant was measured using the thermal tune method and found to be approximately 30 N/m.42 The average surface rootmean-square (RMS) roughness measured over the area of 5 × 5 μm2 was 0.31 nm. All measurements were performed at a constant temperature of T = 30 °C, appreciably below the glass transition temperatures of the polymers. Finally, please note that the maximum force Fmax was fixed at 400 nN.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsmacrolett.5b00944. Figure S1 and Table S1 (PDF).

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Present Address ‡

Institute of Electronic Structure and Laser, Foundation for Research and Technology- Hellas, 711 10 Heraklion Crete, Greece (E.G.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Support for this research from the National Science Foundation (NSF), Division of Material Research (DMR-1305749), is gratefully acknowledged.

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REFERENCES

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DOI: 10.1021/acsmacrolett.5b00944 ACS Macro Lett. 2016, 5, 439−443