Versatile Interpenetrating Polymer Network Approach to Robust

Aug 28, 2017 - The pursuit of intelligent optoelectronics could have profound implications on our future daily life. Simultaneous enhancement of the e...
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Versatile Interpenetrating Polymer Network Approach to Robust Stretchable Electronic Devices Guoyan Zhang,†,+ Michael McBride,† Nils Persson,† Savannah Lee,† Tim J. Dunn,⊥ Michael F. Toney,⊥ Zhibo Yuan,‡ Yo-Han Kwon,† Ping-Hsun Chu,† Bailey Risteen,† and Elsa Reichmanis*,†,‡,◊ †

School of Chemical and Biomolecular Engineering, Georgia Institute of Technology, 311 Ferst Drive NW, Atlanta, Georgia 30332, United States ‡ School of Chemistry and Biochemistry, Georgia Institute of Technology, 901 Atlantic Drive, Atlanta, Georgia 30332, United States ◊ School of Material Science and Engineering, Georgia Institute of Technology, 771 Derst Drive NW, Atlanta, Georgia 30332, United States ⊥ Stanford Synchrotron Radiation Light Source, Menlo Park, California 94025, United States + Peking University, No. 5 Yiheyuan Road Haidian District, Beijing 100871, P.R. China S Supporting Information *

ABSTRACT: The pursuit of intelligent optoelectronics could have profound implications on our future daily life. Simultaneous enhancement of the electrical performance, mechanical stretchability, and optical transparency of semiconducting polymers may significantly broaden the spectrum of realizable applications for these materials in future intelligent optoelectronics, i.e., wearable devices, electronic skin, stretchable displays, and a vast array of biomedical sensors. Here, semiconducting films with significantly improved mechanical elasticity and optical transparency, without affecting the film’s electronic conductivity even under 100% strain, were prepared by blending only a small amount (below 1 wt %) of either p-type or n-type commercial semiconductor polymers. We demonstrate that a self-organized versatile conjugated polymer film displaying an interpenetrating polymer network is formed in the semiconducting films and is crucial for the observed enhancement of elasticity, optical transparency, and charge-carrier mobility. On the basis of this versatile semiconducting film, we explored a new practical approach to directly integrate all the stretchable components for a large area transistor array through solution processing and a final single, mechanical peel-off step. We demonstrate robust transistor arrays exhibiting charge carrier mobilities above 1.0 cm2/ V s with excellent durability, even under 100% strain. We believe our achievements will have great impact on stretchable optoelectronic devices for practical applications and represent promising directions for industry-scale production of stretchable displays and wearable electronic devices. components.21−26 Among these strategies, most efforts so far have concentrated on the design and preparation of an intrinsically elastic active layer with high charge carrier mobility.15,21,22,25,26 Generally, the manufacturing processes used to fabricate such devices include vacuum evaporation,11,16 photolithographic patterning,18−20 or multistep peel-off transfer printing.15,22,25,27 Because of the elaborate synthesis process of intrinsically elastic semiconducting polymers and the multistep transfer manufacturing process for device fabrication, such stretchable transistors have been mostly achieved in small-area devices, and few studies have focused on the fabrication of large

1. INTRODUCTION Stretchable microelectronics has garnered significant attention from research and industry due to its essential role in the realization of large-area, wearable,1−3 epidermal, and biomedical electronics.4−9 Polymer transistors, as the fundamental building blocks for state-of-the-art stretchable electronics, are necessary for economically viable wearable optoelectronic devices.10−14 Simultaneous enhancement of the electrical performance, mechanical stretchability, and optical transparency of polymer transistors may significantly broaden the spectrum of realizable applications for these materials.15 Approaches for stretchable transistors have been intensively investigated: geometric structure design has been used to impart stretchability to brittle materials,16−20 and intrinsically elastic electrode or semiconductor materials have been used directly as device © 2017 American Chemical Society

Received: July 18, 2017 Revised: August 28, 2017 Published: August 28, 2017 7645

DOI: 10.1021/acs.chemmater.7b03019 Chem. Mater. 2017, 29, 7645−7652

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Figure 1. Versatile conjugated semiconducting films with enhanced stretchability, transparency, and conductivity. (A) Chemical structures of insulating materials: PDMS base and PDMS curing agent; commercial p-type semiconducting polymers P3HT and DPP-DTT and n-type semiconducting polymer DPPDPyBT. (B) Stress−strain behavior of freestanding films composed of PDMS and commercial semiconductors. (C) Conductivity as a function of tensile strain. Measurements were performed 10 times in ambient air.

We took advantage of the propensity of dissimilar polymers to undergo phase separation and form an IPN of semiconducting polymer within a rubbery matrix. The IPN ensured formation of efficient charge carrier transport pathways, even with only ∼0.5 wt % semiconductor, and surprisingly exhibited improved elasticity, transparency, and hole or electron mobility. Upon incorporation into organic field effect transistor (OFET) devices, the semiconducting films could be stretched by 100% with retention of their high charge carrier mobility. Such semiconducting films offer expanded flexibility for scaling to large area future intelligent optoelectronics with practically useful levels of functionality. The fabrication scheme we present offers significant opportunities to realize the vision of flexible and stretchable electronics for displays, wearable devices, devices suitable for epidermal applications (i.e., electronic skin), and a vast array of biomedical sensors.

area highly stretchable transistors. To date, none of these approaches allow for scalable fabrication of large area, highly stretchable polymer transistors with stable carrier mobility via a simple and robust, fully solution-processed methodology. Thus, the practical implementation of stretchable devices has not been realized, and alternative approaches are required. Mixtures of polymers are known to phase segregate due to their small entropy of mixing.28−31 A phase-separated interpenetrating polymer network (IPN)32,33 might also enhance mechanical compliance due to cocontinuity of the components.34 For semiconducting polymers, an IPN might afford a spatially distributed interface that would facilitate the formation of percolation networks needed to support charge transport in semiconducting polymers.35−38 In addition, a semiconductor− insulator IPN would reduce the quantity of costly semiconductor consumed and could impart visual transparency. Thus, precise control of the phase separation in a blend system to generate an IPN structure might be an efficient strategy for realizing high-performance transistors at drastically reduced materials cost, obviating the need to design all performance requirements into the active polymer itself. Here, we present a versatile approach to fabricate transparent, elastic semiconducting films using commercially available semiconducting and insulating polymer materials. We selected commercial polydimethylsiloxane (PDMS) as the soft material mainly because of specific characteristics such as good transparency, low cost, and ease of fabrication. Additionally, PDMS is biocompatible, an advantage for envisioned biomedical applications. The active materials were integrated with stretchable device layers to create stretchable polymer transistors exhibiting exceptional performance. Finally, a 10 × 10 cm2 (86 transistors) high performance polymer transistor array was fabricated via a robust and fully solution processed approach.

2. RESULTS AND DISCUSSION 2.1. Stretchable, Transparent, Interpenetrating Semiconducting Polymer Networks. Transparent, elastic semiconducting films were fabricated using PDMS as the transparent stretchable insulating component39 and three distinctly different commercially available semiconducting polymers. Poly(3-hexylthiophene-2,5-diyl) (P3HT) was selected as a readily accessible model semiconducting polymer, enabling investigation of fundamental mechanisms associated with process-induced organization and device performance. The versatility of our methodology was demonstrated using two commercially available high performance polymer semiconductors, namely the hole transport (p-channel) and electron transport (n-channel) polymers poly[2,5-(2-octyldodecyl)-3,6diketopyrrolopyrrole-alt-5,5-(2,5-di(thien-2-yl)thieno [3,2-b]thiophene)] (DPP-DTT) and poly(2,5-bis(2-octyldodecyl)3,6-di(pyridin-2-yl)-pyrrolo[3,4-c]pyrrole-1,4(2H,5H)-dionealt-2,2′-bithiophene) (DPPDPyBT), respectively (Figure 1A). 7646

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Table 1. Optical, Mechanical, and Electrical Properties of the Semiconducting Films mobility (cm2/V s)

semiconductors type p

material P3HT DPP-DTT

n

DPPDPyBT

fsemi (wt %) 0.49 100 0.83 100 0.62 100

original 0.17 0.02 1.53 0.81 1.25 0.43

(±0.03) (±0.01) (±0.22) (±0.11) (±0.09) (±0.14)

stretchability (strain %)

parallel 100% 0.15 (±0.01) 0.006 (±0.001) 1.35 (±0.04) 0.18 (±0.01) 1.08 (±0.02) 0.12 (±0.01)

Solutions of a given semiconductor in chloroform (CHCl3) were mixed with PDMS (wbase: wcuring agent = 10:1) to afford semiconductor/PDMS solutions in a range of semiconductor weight fractions ( fsemiconductor = [Wsemiconductor/(Wsemiconductor + WPDMS)] × 100%) (Figure S1, Supporting Information). The optimum semiconductor to PDMS ratio was determined through OFET performance evaluation. Semiconductor weight fractions as low as 0.49−0.83 wt % afforded the best performance: mobilities up to 0.23 cm2/Vs were obtained for UV irradiated40 P3HT (0.49 wt %). The high mobility, pchannel polymer DPP-DTT exhibited a mobility of 1.75 cm2/ Vs when blended with PDMS (0.83 wt % semiconductor), while the electron transport polymer DPPDPyBT, displayed a mobility as high as 1.34 cm2/Vs (0.62 wt % DPPDPyBT). Under optimized conditions, the results were reproducible, as demonstrated by analyzing 100 OFETs prepared in different batches (Figure S2, Supporting Information). The approach reported here afforded semiconducting films exhibiting ∼10fold, 2-fold, and 3-fold mobility enhancement compared to control samples (100 wt %) of P3HT, DPP-DTT, and DPPDPyBT, respectively (Table 1). Freestanding semiconductor/PDMS films were further characterized from a mechanical perspective (Figure 1B, Table 1, Figure S3, and Video S1). Surprisingly, incorporation of less than 1 wt % of conjugated polymer led to films that could be stretched to a greater degree than the 100% PDMS control (Table 1). No cracks were observed under 100% strain, which was in marked contrast to the neat semiconductors (Figure S4, Supporting Information). The increased stretchability is likely due to the semiconducting film structure, as discussed below. The low proportion of conjugated polymer also facilitated high film transmittance at 300−1000 nm (Figure S5, Supporting Information). Semiconductor/PDMS film direct-current conductivity (σDC) as calculated from the source-drain current (IDS) and source-drain voltage (VDS) relationship at low VDS, demonstrated that the average σDC of the blend films was significantly improved compared with that of the parent semiconductor controls (Figure 1C, Table 1). Furthermore, applying a 100% strain to the blend films led to only a slight decrease in observed σDC. 2.2. Electronic Performance of the Semiconducting Films under Strain. A delamination−stretching−relamination process27,41 was used to evaluate electrical performance of the elastic semiconducting films under strain (Figure S6, Supporting Information). Good ohmic contact42 was retained throughout, and thus, device performance was not hindered during the process (Figure S7, Supporting Information). As shown in Figure 2A, subjecting the semiconductor/PDMS to 100% strain led to a slight decrease in charge carrier mobility; however, the mobility under strain was still higher than that of the control films. Moreover, the semiconductor transfer curves under 100% strain remained unaffected (Figures 2B and C).

crack on-set

elongation break 182.1 (±8.6)

32.6 (±4.8) 140.1 (±10.2) 20.7 (±5.2) 133.5 (±7.6) 22.5 (±4.5)

visual transparency (% at 550 nm) 90.8 47.1 94.9 79.6 97.1 73.8

(±0.6) (±0.4) (±0.5) (±0.6) (±0.3) (±0.2)

Figure 2. (A) Mobility under 100% strain parallel (solid) and perpendicular (empty) to the charge transport direction. Left: normalized mobility performance of the semiconducting film with 0.49 wt % P3HT (red) and the neat P3HT film (black). μ0 denotes the original mobility. Right: mobility from the DPP-DTT/PDMS (0.83 wt %, blue) and DPPDPyBT/PDMS (0.62 wt %, green). (B and C) Transfer curves obtained from the semiconducting films in their original condition, under 100% strain parallel to the charge transport direction and its release condition, and under 100% strain perpendicular to the charge transport direction and its release condition.

The origin of the observed exceptional charge transport performance is discussed below. 2.3. Morphology and nanofibrillar crystallization of the “model” P3HT/PDMS through phase separation. The P3HT/PDMS blend was used as a model to determine the origins of the composite film performance characteristics. Specifically, the relationship between the film microstructure and semiconductor performance were of interest. It has been reported that vertical phase separation between two dissimilar polymers can promote the formation of an IPN microstructure.32 Contact angle experiments allowed calculation of the surface free energy (Υs) of P3HT (19.3 (±0.7) mJ/m2) and 7647

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Figure 3. (A) Hansen spheres of P3HT (black sphere) and PDMS (red sphere) constructed on Hansen 3D space. Each axis represents three different components consisting Hansen solubility parameters (HSP: dispersion, polar, and hydrogen-bonding related cohesion energy). (B) Normalized UV−visible absorption of neat P3HT (black color) and P3HT/PDMS with 0.49 wt % P3HT (red color): solution state (dot line) and film state (solid line). (C) Left: evolution of percentage of ordered aggregates and right: exciton bandwidth of the P3HT/PDMS film with 0.49 wt % P3HT.

PDMS (7.3 (±0.3) mJ/m2) (Figure S8 and Tables S1−S3). Additionally, Hansen solubility parameter studies confirmed a lower relative energy difference (RED) between PDMS and CHCl3 (0.74) compared to that of P3HT and CHCl3 (0.98), indicating that PDMS is more soluble in CHCl3 than P3HT (Figure 3A and Table S4).43 The higher surface energy of P3HT and its lower solubility are key driving forces that can induce vertical phase separation during film formation. UV−vis spectroscopy results suggest that in the presence of PDMS, P3HT undergoes more extensive aggregation than pristine P3HT: low energy P3HT bands at ∼555 nm (0−1) and ∼605 nm (0−0) suggest that the presence of PDMS led to P3HT aggregation in solution, likely due to the large repulsive interactions (χ: 0.55) between the two components (Figures 3B and S9 and Table S4). Furthermore, this unfavorable interaction is expected to continue to promote P3HT aggregation via favorable π−π interactions during CHCl3 evaporation. According to Spano’s model, the lower exciton bandwidth (W) calculated for 0.49 wt % P3HT/PDMS indicates the presence of a higher degree of intramolecular ordering and longer conjugation length within the aggregates; the extent of aggregation was significantly increased to ∼71% (∼30% in solution) vs the P3HT control (61% in film, 1.5% in solution) (Figure 3C).44 Atomic force microscopy (AFM) images of the top and bottom surfaces of the P3HT/PDMS films (Figures 4A and S10) showed that at 0.49 wt %, P3HT appears to be buried in the PDMS matrix, presenting an aligned nanofibrillar structure. The nanofibrillar alignment was quantified by the orientational order parameter (S2D) which yielded a value of 0.56. X-ray photoelectron spectroscopy (XPS) depth profiling facilitated visualization of the location of P3HT sulfur (S 2p, ∼164 eV) and PDMS silicon (Si 2p, ∼102 eV) within the film (Figures 4B and S11). In addition to top to bottom depth profiling, the composite semiconductor was delaminated from the substrate and then etched from the bottom side. The results demonstrated that, while the P3HT component was dispersed within the PDMS matrix, importantly, the polymer film− substrate interface (bottom) was highly enriched in P3HT, with PDMS appearing throughout the entire thickness. Grazingincidence wide-angle X-ray-scattering data of P3HT/PDMS films are presented in Figures 4C and S11. First, a halo similar to that of neat PDMS was observed for both surfaces of the film; the characteristic P3HT (100) diffractions were only observed at the bottom surface overlapping with PDMS. After subtracting PDMS, the (100) reflection was clearly apparent,

Figure 4. (A) AFM phase image (left), color orientation map (right), and orientation distribution (inset) of the bottom interface of the P3HT/PDMS film with 0.49 wt % P3HT. The color orientation map was extracted from the original AFM phase image using GTFiber software [gtfiber.github.io]. Orientation distribution was extracted from the orientation map, and the radial axis indicates the count of pixels of a given orientation. The diametrical black line segment indicates the average orientation. (B) XPS color images of the P3HT/ PDMS film with 0.49 wt % P3HT obtained from the etching depth profile via a Fourier transformation showing the S 2p and Si 2p elements distribution along the etching direction (bottom to top). From each sample, four different points were selected to collect the depth profile. (C) Two-dimensional GIWAXS maps show the crystalline quality of the P3HT/PDMS film with 0.49 wt % P3HT at bottom interface, and both P3HT (100) and PDMS diffractions were observed (left). After removal of the PDMS phase background, the characteristic P3HT (100) and (010) diffractions are well-resolved, indicating an edge-on orientation of P3HT nanofiber in the PDMS matrix.

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and (010) P3HT diffractions, attributed to interchain π−π stacking, were visible. The results indicate that in the PDMS blend, P3HT is more crystalline than pristine semiconductor. The increased P3HT (010) intensity indicates that the conjugated polymer mainly exhibited edge-on orientation, which facilitated charge transport in the devices.45 2.4. Mechanism and Evidence for Semiconductor/ PDMS Interpenetrating Polymer Network Formation. The data presented above suggest that film formation takes place in a multistage process as summarized in Figures 5A and

P3HT, PDMS interdiffusion is limited but sufficient to enhance P3HT aggregate localization. As a result, a film comprising an interpenetrating interfacial network is formed (Figures 5B−III). Selective removal of the PDMS phase confirmed the presence of an IPN structure (SEM and AFM, Figures 5C and S13). As presented in the three-dimensional and cross sectional SEM image (Figure 5C), the residual P3HT film appeared as a porous, networked structure. The phase separated IPN structure proved advantageous because the solubility and surface energy differences between P3HT and PDMS led to increased interactions between the conjugated polymer chains, leading to a highly networked fibrillar structure that extended into the cross-linked PDMS host. In turn, the networked semiconducting film might play a role similar to that of a plasticizer, helping increase the fluidity of PDMS chains in the biphasic region and thereby affording enhanced stretchability. The resultant IPN structure not only improved the effectiveness of charge transport pathways and elasticity but also provided for environmental stability (Figure S14, Supporting Information). 2.5. Solution-Processed Large Area Stretchable Polymer Transistor Arrays. Stretchable polymer transistor arrays comprising 86 transistors were fabricated using a poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate)-bis(trifluoromethane) sulfonimide lithium salt (PEDOT:PSSLiTFSI) composition for the stretchable electrodes and a composite of PDMS-Ecoflex (PDMS-Eco.) as the substrate and passivation layer. The solution processing steps used to fabricate the stretchable arrays are summarized in Figures 6A and S15. A major difficulty in this process is the patterning of polymer-based electrodes. Since evaporation methods were not feasible, we utilized a patterned inorganic substrate to obtain polymeric electrodes. The remaining layers were deposited on top of the electrodes, and then the entire device was peeled away from the inorganic substrate. The peeling step must be designed so that the layers with the weakest mutual adhesion occur between the bottom-most device components and the inorganic substrate; thus, a thin-layer of PDMS (10−20 nm) was added between the SiO2 substrate and the PEDOT:PSS

Figure 5. (A) Aggregated nanofibrillar structures in the solution state. (B) The film separates into two metastable layers, and finally, an IPN structure is formed. (C) SEM and AFM images of the IPN structure. Note: To illustrate the fiber aggregation and IPN structure clearly, the cartoons are not drawn to scale..

B. Large repulsive interactions between the two immiscible polymers may accelerate P3HT aggregation in solution. Then, during the coating process, differences in surface energy and solubility facilitate vertical phase separation, resulting in a selfstratified metastable two-layer structure (Figures 5B−I). During film formation, parent PDMS and curing agent chains have an opportunity to penetrate and diffuse into the evolving P3HT layer (Figures 5B−II). Because of the low weight fraction of

Figure 6. (A) Schematic illustration of the preparation of stretchable transistor arrays via solution processing. (B) Schematic illustration of the stretchable transistor configuration and high-resolution cross-sectional SEM image of a complete stretchable transistor fabricated by the solution processing methods (left). The integrated configuration of the stretchable transistor after 300 stretching−releasing cycles confirmed by the crosssectional SEM image (right). (C) Photographic images illustrate the stretchable transistor arrays. Optical image shows transistors with integrated electrodes and clearly visible channel with and without strain (100%). 7649

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Figure 7. Photograph showing the high visual transparency of the transistor arrays without (A) and with (B) strain. (C) Transfer curves obtained from P3HT/PDMS (0.49 wt % P3HT: left), DPP-DTT/PDMS (0.83 wt % DPP-DTT: middle), and DPPDPyBT/PDMS (0.62 wt %, DPPDPyBT: right), in its original condition, under 100% strain (parallel and perpendicular to the charge transport direction), and release condition. (D) Changes in mobility of the stretchable in their original condition, under 100% strain, and after release. (E) Changes in the mobility after multiple stretching− releasing cycles at 100% strain parallel to the charge transport direction.

transistors fabricated on Si/SiO2 is due to differences in the dielectric and the slight decrease in electrode conductivity (Figure S19, Supporting Information). Note that the under 100% strain, mobility decreased only slightly, primarily because of the associated changes in channel length, dielectric capacitance, and electrode conductivity (Figures 7C and D, S20, S21, and Table S6). Upon releasing strain, the mobility recovered its initial values. In addition, the fully stretchable transistors exhibited stable device performance over 300 stretching cycles (Figures 7E and S22).

electrodes. Work of adhesion (W) between the relevant layers was calculated using the surface free energy.46 WPDMS-Eco/ PEDOT:PSS-LiTFSI was found to have a higher value compared to that of WPEDOT:PSS-LiTFSI/PDMS but lower than that of WPEDOT:PSS-LiTFSI/(Semi.-PDMS). The different W values between the various layers ensured that the one-step peel-off process could be realized (Figure S16 and Table S5, Supporting Information). The multilayer device architecture was further confirmed through cross-sectional SEM imaging (Figures 6B, left). After 300 stretching−releasing cycles, the integrated layered structure was retained, confirming strong adhesion between device layers (Figure 6B, right). Figure 6C shows transistors with integrated electrodes and clearly visible channel with and without strain (100%). 2.6. Electronic Performance of the Fully Stretchable Polymer Transistors. The device yield was above 90%, and the devices had good transparency and stretchability (Figures 7A and B, Figure S17, and Video S2). The P3HT-based stretchable transistors exhibited the anticipated transfer characteristics with an average mobility of 0.12 (±0.02) cm2/ V s and low threshold voltage (Figure S18, Supporting Information). An average mobility above 1.0 cm2/V s was obtained for DPP-DTT (1.23 (±0.03) cm2/V s) and DPPDPyBT (1.02 (±0.05) cm2/V s), p- and n-channel devices, respectively (Figures S18B and C, Supporting Information). The lower mobility obtained for the stretchable devices vs

3. CONCLUSIONS In summary, we present a robust semiconducting film with enhanced conductivity, mechanical stretchability, and visual transparency achieved through an IPN structure. Commercial p-channel and n-channel conjugated polymers confirmed the generality of this method. Mechanistically, the enhanced performance derived from the formation of an IPN, confirmed through multiple imaging techniques. The self-organized IPN formed within semiconducting polymer−PDMS blend films is crucial for the observed enhancement of elasticity and chargecarrier mobility. Our approach to stretchable electronics requires only a single mechanical peel-off step to integrate all of the components, providing new opportunities for large-area production of stretchable devices. The IPN strategy and process protocols are expected to bring stretchable electronic 7650

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systems to a practical level, making possible the industrial-scale production of stretchable displays and wearable electronic devices.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b03019. Experimental details and additional characterization data (PDF) Video of freestanding semiconducting film (0.49 wt % P3HT) under sequential stretching, twisting, and poking (AVI) Video showing the stretchable high visual transparency of the transistor arrays (AVI) Video showing the fully stretchable transistor array fabrication process (AVI)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Nils Persson: 0000-0003-0750-5880 Michael F. Toney: 0000-0002-7513-1166 Elsa Reichmanis: 0000-0002-8205-8016 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Kenneth A. Gall and Andrew T. Miller for help in mechanical testing, Ho-Yee Hui for SEM data collection, Rui Chang for optical image collection, and Lu Jiang for the contact angle experiments. G.Z. gratefully acknowledges financial support from the International Postdoctoral Exchange Fellowship Program of China Postdoctoral Council-Peking University (Grant 20150008) and the authors also appreciate support from the National Science Foundation (Grant CBET1264555). N.P. gratefully acknowledges financial support from the NSF FLAMEL IGERT Traineeship program, NSF Grant 1258425, IGERT-CIF21. M.M. acknowledges an NSF IGERT NESAC Traineeship: NSF Grant 1069138. This work was performed in part at the Georgia Tech Institute for Electronics and Nanotechnology, a member of the National Nanotechnology Coordinated Infrastructure, which is supported by the National Science Foundation (Grant ECCS1542174). Support from the Georgia Institute of Technology, the Brook Byers Institute for Sustainable Systems, funds associated with the Pete Silas Chair in Chemical Engineering, the Renewable Biomaterials Institute, and the Georgia Tech Polymer Network is also acknowledged.

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