Nanoscale Domain Imaging of All-Polymer Organic Solar Cells by

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Nanoscale Domain Imaging of All-Polymer Organic Solar Cells by Photo-Induced Force Microscopy Kevin L. Gu, Yan Zhou, William A. Morrison, Katherine Park, Sung Park, and Zhenan Bao ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.7b07865 • Publication Date (Web): 16 Jan 2018 Downloaded from http://pubs.acs.org on January 17, 2018

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Nanoscale Domain Imaging of All-Polymer Organic Solar Cells by Photo-Induced Force Microscopy Kevin L. Gua, Yan Zhoua, William A. Morrisonb, Katherine Parkb, Sung Parkb, and Zhenan Baoa* a

Department of Chemical Engineering, Stanford University, Stanford, CA 94305, USA

b

Molecular Vista, 6840 Via Del Oro, Suite 110, San Jose, CA 95119, USA

ABSTRACT

Rapid nanoscale imaging of the bulk heterojunction layer in organic solar cells is essential to the continued development of high-performance devices. Unfortunately, commonly used imaging techniques such as tunneling electron microscopy (TEM) and atomic force microscopy (AFM) suffer from significant drawbacks. For instance, assuming domain identity from phase contrast or topographical features can lead to inaccurate morphological conclusions. Here we demonstrate a technique known as Photo-induced Force Microscopy (PiFM) for imaging organic solar cell bulk heterojunctions with nanoscale chemical specificity. PiFM is a relatively recent scanning probe microscopy technique that combines an AFM tip with a tunable infrared laser to induce a dipole for chemical imaging. Coupling the nanometer resolution of AFM with the chemical specificity of a tuned IR laser, we are able to spatially map the donor and acceptor domains in a model all-

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polymer bulk heterojunction with resolution approaching 10 nm. Domain size from PiFM images is compared to bulk-averaged results from resonant soft X-ray scattering, indicating excellent quantitative agreement. Further, we demonstrate that in our all-polymer system, the AFM topography, AFM phase, and PiFM show poor correlation, highlighting the need to move beyond standard AFM for morphology characterization of bulk heterojunctions.

Keywords: all-polymer solar cell, atomic force microscopy, photo-induced force microscopy, Xray scattering, nanoscale chemical imaging

Accurate and rapid nanoscale imaging is vital to the research and development of nextgeneration organic electronic devices.1,2 In particular, polymer-based organic photovoltaics (OPVs) have emerged as a promising renewable energy candidate suitable for inexpensive and scalable production, being lightweight, flexible, and amenable to low-energy processing under ambient conditions.3 Recent efforts have achieved 13% power conversion efficiency (PCE),4 highlighting the potential of OPVs for large-scale commercial applications. Despite the impressive performance, in order to address the current challenges OPVs still face – stability,5 environmentally benign solvents,6 large-scale printing,7–9 limited absorption,10 etc. – much needs to be understood about the internal morphology of the light-absorbing layer. Thus, tools to structurally characterize the active layer morphology are as indispensible as the techniques used to fabricate them. Rapid morphological characterization would dramatically improve the processing-performance-morphology feedback loop.11 In OPVs, the light-absorbing layer is comprised of a network of self-assembled interdigitating electron donor and electron acceptor materials known as a bulk heterojunction (BHJ). Among the


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various morphological parameters affecting photovoltaic performance including domain purity, crystalline stacking, vertical phase separation, etc., phase separated domain size is held to be of the most principal and is therefore of particular interest to image or quantify.1,12 However, this phase separation sizescale between donor and acceptor domains occurs on the order of tens of nanometers for high-performance devices.13,14 The small sizescales involved, combined with the relative lack of electron contrast between organic materials, makes characterization and imaging of OPV BHJs uniquely challenging. Many common techniques such as optical microscopy and tunneling and scanning electron microscopy are either ineffective or unreliable for highperformance OPVs.1 This problem is exacerbated in non-fullerene-based BHJs, as there is even less contrast between donor and acceptor materials due to their similar molecular segments. As the field currently stands, comprehensive morphology characterization is a laborious and involved process requiring the use of synchrotron facilities to obtain nanoscale spatial information. In addition to the ever-growing issue of facility availability, synchrotron X-ray methods only provide bulk-averaged quantities, losing any direct and fine structural information present in the BHJ. The relative dearth of lab-scale characterization methods highlights the need for additional facile imaging techniques. Despite decades of active research, there are relatively few imaging techniques capable of achieving a spatial resolution of 10 nm in organic semiconductor BHJs, none of which directly provide chemical spatial information. Atomic force microscopy (AFM) is perhaps the most commonly utilized and most accessible technique in characterization of OPV BHJs, but it suffers from several shortcomings owing to its surface-sensitive nature. Phase separated domains are commonly inferred from either topographical features or phase contrast from AFM images. However, this serves as a proxy for true chemical imaging, with neither consistently accurately


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reflecting the spatial distribution of donor/acceptor phases.2 Again this is particularly a challenge for polymeric acceptors as both their chemical and mechanical properties are similar to those of the donor materials. Thus, AFM can lead to tenuous quantitative results or inaccurate conclusions, exacerbated in instances involving capping layers at the air interface,15 and still other systems show no correlation between phase and topographic features.16,17 Several AFMcoupled techniques have been developed to image contrast in various materials properties such as photoconductivity,18,19 thermal conductivity,20 and electronic properties.21–23 However, none of these techniques provide unambiguous chemical identification, instead relying on assumptions of chemical identity from a featured map. An excellent review of these various characterization techniques can be found elsewhere.2 The only method providing true chemical specificity is AFM-infrared, in which infrared-induced thermal expansion is measured using an AFM tip.20 However, this method has not found use in modern OPVs likely due to insufficient lateral resolution on organic materials, despite having demonstrated 20 nm in perovskite photovoltaics.24 Another common strategy for imaging BHJs is to utilize transmission techniques. In particular, transmission electron microscopy (TEM)1,25,26 and scanning transmission X-ray microscopy (STXM)27–29 have been successfully applied to OPVs, both capable of resolving nanometer-scale features. A shortcoming of both techniques is that signal integration over the thickness of the film results in an image which is a projection of the entire film depth. 3D tomography is possible by taking multiple images at different angles, but this method requires significant effort and specialized equipment that is not widely available.30 Due to full-film projection and weak contrast, 2D images are typically not relied upon to make strong quantitative morphological conclusions.31 Outside of a few cases utilizing beam defocusing or energy filtered TEM,25,26,30


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the issue remains that many OPV systems simply exhibit insufficient contrast between the two compositionally similar semiconductor materials. The lack of appropriate characterization methods that can achieve nanoscale spatial mapping highlights the need for a rapid imaging method that can penetrate the surface, yet avoid full-film projection.

Figure 1. Simplified schematic of the photo-induced force microscopy setup. PiFM measures the dipole force between the absorbing sample and the AFM tip. The sample is excited at a specified IR wavelength, inducing a dipole resulting from the vibrational motion of chemical bonds. A mirror dipole is induced in the tip leading to an attractive force between the sample molecules and the tip. Multiple channels can be acquired simultaneously (e.g. AFM topography and PiFM as shown).

Here we present Photo-induced Force Microscopy (PiFM) as a method to image OPV BHJs with nanoscale chemical specificity (Figure 1). PiFM is a relatively recent scanning probe


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microscopy technique that combines an AFM with a tunable infrared laser (IR) to produce contrast for chemical imaging. The IR laser is focused at the region where the tip and sample interact, and the gold-coated AFM tip acts as a nano-antenna, creating a strong near-field light focused underneath the tip apex.32 When sample molecules absorb this light, an electric dipole is induced from the vibrational motion of IR-resonant bonds, which leads to a mirror dipole being created in the metallic AFM tip and an attractive force between the sample molecules and the tip.33 Coupling the nanometer resolution of AFM with the chemical specificity of a tuned IR laser, PiFM has been successfully used to image several polymer systems including polystyrene, poly (methyl methacrylate), and poly (2-vinyl pyridine) and resolve the features arising from the self-assembly of block copolymers32. A detailed description of the operating principles of PiFM can be found elsewhere.33–35 We demonstrate the application of PiFM on an all-polymer OPV system, in which both donor and acceptor materials are conjugated donor-acceptor polymers. In comparison to the more common polymer donor-small molecule acceptor, all-polymer OPVs boast tunable spectral absorbance through molecular design, more amenable solution processing properties, and enhanced morphological and chemical stability.13,36,37 Due to these distinct advantages, a recent surge of rapid development has improved the performance from 3% PCE to 9% PCE within the past 5 years.13,38 In comparison to small molecule systems,39 all-polymer systems are particularly suited as a model to demonstrate the ability of PiFM to overcome the shortcomings of the aforementioned characterization techniques. For example, polymer-polymer electron density contrast is much lower than that between polymer-fullerene or other small molecules. Additionally, the surface penetrating nature of PiFM circumvents skin effects that limit the usefulness of surface characterization methods – PiFM directly images domains rather than


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relying on indirect features in AFM topography or phase correspond to donor or acceptor phases. Here we report the application of PiFM on all-polymer OPV BHJs, demonstrating chemical specificity at a morphologically relevant spatial resolution of ~10 nm. The results are compared to an established bulk measurement method to quantify domain size in all-polymer BHJs, demonstrating excellent quantitative agreement.

Scheme 1. All-polymer OPV model system.

RESULTS AND DISCUSSION Two semiconducting polymers exhibiting distinct infrared absorption spectra were selected as a model OPV system. The p-type electron donor chosen is a low-bandgap polymer, poly isoindigo-bithiophene with 10 mol% polystyrene sidechains (PII-2T-PS), for which our research group has previously developed expertise in synthesizing and processing (Scheme 1).36,40,41 A perylene diimide-based n-type polymer, PPDI-T, is used as the electron acceptor.36,40 Notably, this system has achieved 5% PCE using non-halogenated solvents and is one of the highest performing isoindigo-based all-polymer solar cells.41,42 It should be noted that this system has no


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unusual attributes making it particularly amenable to PiFM. In principle, any material pair with at least one set of non-overlapping vibrational peaks can be selectively mapped. First, bulk Fourier Transform Infrared (FTIR) spectra of neat PII-2T-PS and PPDI-T films were obtained. Solutions of PII-2T-PS or PPDI-T dissolved in toluene were spin coated on glass substrates to fabricate neat films. A PiFM point spectrum on each sample was then taken, revealing excellent agreement with bulk FTIR spectra (Figure 2). In order to obtain singlecomponent mapping, a non-overlapping spectral region is selected for chemical imaging. The donor polymer, PII-2T-PS, exhibits a modest peak at 1453 cm-1 corresponding to the polystyrene methylene bending modes where the signal from the acceptor polymer is quite weak. Similarly, the carbonyl peak at 1706 cm-1 is chosen to selectively image the acceptor polymer, PPDI-T. While both the donor and acceptor backbone units contain carbonyl groups, a sufficient wavenumber shift arises from their different chemical environments. It should be noted that the excitation laser used has a varying power profile, with two low power zones between 1330-1370 cm-1 and 1570-1660 cm-1 (Figure S1). This explains in particular the absence of a PiFM peak at 1605 cm-1 for PII-2T-PS, where there is a very large peak according to FTIR.


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Figure 2. FTIR and PiFM spectra of neat polymer films and PiFM images of bulk heterojunctions. (Left) Bulk FTIR spectra of PII-2T-PS and PPDI-T neat polymer thin films (a). The excitation wavenumber chosen for selective imaging of donor and acceptor polymers are 1453 cm-1 and 1706 cm-1, respectively. PiFM point spectra overlaid on FTIR spectra of PII-2TPS (b) and PPDI-T (c). We observe excellent agreement between bulk FTIR and PiFM spectra in the range of 1300 – 1800 cm-1. (Right) PiFM image of a PII-2T-PS/PPDI-T BHJ fabricated with 0.2 vol% CN additive at 1453 cm-1, highlighting the donor phase (d) and acceptor phase (e). We refer the reader to the Supporting Information media file for an overlay of these two images.

PII-2T-PS/PPDI-T dissolved in toluene with small amounts of the processing additive 1chloronaphthalene (CN) was then spin coated on glass to form bulk heterojunctions. The high-


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boiling point CN additive modifies the phase separation behavior during drying, influencing the domain size in the final dried film.43 PII-2T-PS/PPDI-T has been shown to demonstrate a decrease in domain size with increased amounts of 1-chloronaphthalene additive,41 which is useful for assessing the qualitative and quantitative accuracy of PiFM. The IR excitation wavelength was tuned to 1453 cm-1 to selectively image the donor polymer and the PiFM images of BHJs resulting from four different CN concentrations are presented in Figure 3. The images acquired are 1 µm (256 by 256) at total time of 5.5 minutes per image. The PiFM signal penetrates the top ~20 nm,17,32 revealing a clear nano-fibrillar network of donor polymer apparent in all four thin films. We observe a narrowing of domain widths as CN content is increased from 0.2 vol% to 2.0 vol% as can be seen in Figure 3. Many high-performance OPV systems are known to form fiber networks as it enables good domain interconnectivity while maintaining small domain widths.25 Smaller domains – closer to the exciton diffusion length of 10 nm – enhance charge separation and improve device photocurrent.14


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Figure 3. PiFM images of PII-2T-PS/PPDI-T bulk heterojunctions. (a-d) BHJs from four different concentrations of 1-chloronaphthalene additive, revealing that higher concentration of additive reduces the domain width of the donor polymer nanofibrils. The wavelength is tuned to 1453 cm-1, where the donor polymer is selectively imaged.

In order to quantify domain size, image analysis was performed on the PiFM images in Figure 3. Each image was first binarized using Otsu thresholding44, which bins pixels into two populations while minimizing the intraclass variance. Binarization was selected because it allows


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for quantitative comparison with the center-to-center domain periodicity from scattering techniques, and an Otsu threshold was chosen to reflect the 1:1 donor:acceptor weight ratio in the initial solution. Next, a medial axis (skeleton) transform was performed to yield the midline of the donor domains.45 By successively removing boundary pixels, only the midline remains, preserving the connectivity of the morphology. Loose branches shorter than 20 nm were pruned. In parallel, a Euclidean distance transform was calculated from the binary image, where each foreground pixel is assigned a value corresponding to its distance to the nearest background pixel.46 The resulting image is used as a weighting factor for the previous skeleton, thus representing half of the fiber width at each point. Weighted skeleton pixels were then averaged to obtain a measure of the average domain width. The computational steps on Figure 3 (a) (0.2 vol% CN) are presented in Figure 4, and we refer the reader to the Supporting Information for the remaining images (Figures S2 – S4). The calculated domain sizes are presented in Figure 5 (b) and (c), from which it is observed that increasing CN concentration results in a decrease in the donor phase domain width from 42 nm for 0.2 vol% CN to 31 nm for 2.0 vol% CN.


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Figure 4. Image analysis procedure on PII-2T-PS/PPDI-T with 0.2 vol% CN to obtain average domain width of donor polymer domains. (a) Raw PiFM image, where highlighted regions indicate presence of polymer donor; (b) binarized image; (c) medial axis transform overlaid on binarized image; (d) distance transformation of binarized image; (e) distance-weighted skeleton, where the intensity of each skeleton pixel represents the distance from a donor/acceptor interface; (f) distance-weighted skeleton overlaid raw PiFM image.


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Figure 5. (a) 1D RSoXS curves for BHJs fabricated with varying CN concentrations. The average domain spacing is represented by the peak in reciprocal space, where the center-tocenter periodicity is given by 2π/q. (b) Averaged domain widths as calculated from PiFM and RSoXS, compared to the inverse of the device short circuit current. We find that the PiFM domain size tracks the inverse current very well, whereas RSoXS is skewed for large domains. (c) Similarly, photoactive volume calculated from the distance transform plotted vs PiFM and RSoXS domain size.

To corroborate these results with an established method, resonant soft X-ray scattering (RSoXS) was used to characterize the bulk-averaged domain size in the series of BHJ films. RSoXS is often considered the method of choice for OPV BHJs, providing a robust measurement of domain periodicity.47 Despite being a powerful measurement technique providing robust quantitative results, it should be emphasized that measurements require access to a synchrotron facility and are thus infeasible for rapid morphological feedback. Peaks/shoulders in the intensity vs scattering vector plot – which is often presented in intensity times the square of the scattering vector vs scattering vector, I*q2, for better visualization – represent the period of scatterers, which may be used as a measure of twice the average phase separated domain size:


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݀=

1 2ߨ 2 ‫ݍ‬௣௘௔௞

where d is the domain size in Å, and qpeak is the peak/shoulder position in the scattering vector in inverse Å.47–49 In Figure 5 (a) we observe that higher CN content shifts the peak towards higher scattering vector indicating smaller domain spacing periodicity. Converting from reciprocal space to real space, we find a decrease in domain size with increasing amounts of CN in qualitative agreement with results from PiFM. Additionally, the excellent quantitative agreement between bulk measurement from RSoXS and surface measurement from PiFM suggests the bulk and surface morphologies in PII-2T-PS/PPDI-T are indeed similar. Furthermore on quantitative comparison, it should be noted that RSoXS data representation in I*q2 is commonly used to identify peaks from shoulders but is known to skew low-q data as q ~ 0 drags down the curve, reducing the peak position.50 Indeed we find while RSoXS quantitatively agrees quite well with our results from PiFM for small domain sizes, large domains represented by peaks at low-q diverge dramatically from the directly measured domain width (Figure 5 b, c). To address this issue, inverted solar cells were fabricated to correlate the observed morphological differences directly to device performance. Domain size is commonly reported as a useful metric for quantifying the probability of charge separation since excitons generated within ~10 nm of a donor/acceptor interface are considered within close enough proximity to diffuse to an interface and separate. Thus, a true measure of domain size should correlate inversely with photocurrent. The device performance improves significantly with increasing CN concentration (Table 1). Examining the individual performance parameters, we observe that the open circuit voltage (VOC) and fill factor (FF) remain constant while the short circuit current (JSC) increases. JSC depends on


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several factors including absorption, exciton diffusion, and dissociation.13 Absorption between samples is similar since the film thicknesses are identical as measured by profilometry (101 ± 1.6 nm), and dissociation is determined by the energy levels of the material. Necessarily, the improvement in JSC arises from the increased probability of excitons reaching the donor/acceptor interface where the energy level offset dissociates excitons into free charge carriers, i.e. the domain width.

Table 1. OPV device performance with varying amounts of CN additive.

CN content

PCE

JSC

VOC

FF

(v/v%)

(%)

(mA cm-2)

(V)

(%)

0.2

3.29

7.92

0.96

43.1

0.5

3.83

9.09

0.97

44.6

1.0

4.16

9.69

0.97

44.5

2.0

4.52

10.57

0.97

44.1

The domain sizes from PiFM and RSoXS are plotted against the inverse of the short circuit current (Figure 5 b). We see that the PiFM domain width tracks JSC-1 quite well, whereas RSoXS diverges for large domains (low-q region). From a photophysical perspective, it is more appropriate to consider all photoactive volume, where those further from the nearest interface are expected to contribute less to the total photocurrent. Therefore, we utilize the previously calculated distance transform from the PiFM images to obtain a more accurate measure of the photoactive volume. Non-zero pixels in Figure 4 (d) are averaged and the resulting quantity of pixel times distance from the nearest interface provides a more robust measure of photoactive


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volume contributing to photocurrent. We find that this value follows JSC-1 very well, which correlates well with the domain size from PiFM rather than RSoXS (Figure 5 c). This suggests that despite data representation in I*q2, RSoXS is quantitatively valid at domain sizes of 30 nm but becomes more inaccurate for larger domains. Another point of emphasis is that PiFM provides direct domain imaging rather than interpretation from topography or phase images. While in polymer/fullerene systems, AFM may provide a reasonable estimation of domain size due to the higher modulus of fullerene aggregates, in all-polymer systems there is often little correspondence between AFM topography and domains. Figure 6 displays PiFM at 1453 cm-1, AFM topography, and AFM phase of PII-2TPS/PPDI-T BHJs with 0.2 – 2.0% CN. The three signals were acquired simultaneously on different channels from the same scan, eliminating the possibility of drift. Indeed, in PII-2TPS/PPDI-T the spatial correlation between PiFM and topography or phase is minimal. It is clear that inference of domain size from topography or phase would lead to overestimation or underestimation, respectively. In PiFM, the polymer fibers are distinct and clearly visible, providing an accurate measurement of domain width.


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Figure 6. Donor phase PiFM (top), AFM topography (middle), and AFM phase (bottom) of PII2T-PS/PPDI-T BHJ films with 0.2 – 2.0% CN additive. It can be seen from images within each column that there is very poor correlation between PiFM, topography, and phase. All images are 1 x 1 µm.

CONCLUSIONS In summary, we have demonstrated the use of PiFM for imaging all-polymer BHJs providing true chemical imaging. By tuning the IR excitation laser to specific absorption peaks of different chemical species, we are able to spatially map the donor and acceptor domains with a lateral resolution approaching 10 nm. Results from image processing are corroborated with RSoXS, which show that domain size decreases with increasing CN content. Further, we have shown that


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in PII-2T-PS/PPDI-T AFM topography, phase, and PiFM show very poor correlation, highlighting the need to move beyond standard AFM for morphology characterization. We emphasize that PiFM is high-throughput, lab-scale, ambient, and requires no special sample preparation, filling an important role in non-projective surface imaging of nanoscale bulk heterojunctions. PiFM provides a powerful analytical method for morphological imaging of solar cell materials, enabling accurate and rapid morphological feedback towards the next generation of high-performance devices.

METHODS Materials. PII-2T-PS and PPDI-T were synthesized according to previously reported procedures36,40,41. Anhydrous toluene and 1-chloronapthalene were purchased from Sigma Aldrich and used as received. 15 wt% diethylzinc solution and tetrahydrofuran were used as received (Sigma-Aldrich). Silicon nitride (Si3N4) windows for resonant soft X-ray scattering (RSoXS) measurements were purchased from Norcada Inc. Solar Cell Fabrication and Testing. Inverted structure devices were fabricated on ITOpatterned glass substrates (13 Ω/sq, Xinyan Technology). The glass/ITO substrates were ultrasonicated sequentially in detergent, distilled water, acetone, and isopropanol. Following a 30 min UV-ozone treatment, ZnO was fabricated by spincoating a diethylzinc precursor solution (15 wt% diethylzinc solution diluted 1:7 with tetrahydrofuran) at 5000 rpm for 30 s in air. The substrates were annealed in air at 180 ˚C for 20 min and cooled to room temperature. 5 mg each of PII-2T-PS and PPDI-T was dissolved in toluene with varying amounts of 1-chloronaphthalene and stirred in a nitrogen glovebox at 80 ˚C overnight. Active layer solutions were spincoated at 700 rpm for 45 s and annealed in nitrogen at 180 ˚C for 5 min. 10 nm MoO3 and 100 nm Ag


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were thermally evaporated in high vacuum (6 x 10-6 Torr) to produce solar cells of 4 mm2 active area. For testing of performance metrics, devices were illuminated under a simulated AM 1.5G solar spectrum (Newport Solar Simulator 94021A) calibrated to a KG5 Si photodiode and measured by a Keithley 2400 analyzer. Resonant soft X-ray scattering. Films for RSoXS were fabricated on Si/polystyrene sulfonate and then float transferred in water to Si3N4 windows for transmission measurements. RSoXS was performed at Beamline 11.0.1.2 at the Advanced Light Source at Lawrence Berkeley National Laboratory. The X-ray energy was scanned from 270 to 290 eV to maximize the scattering intensity, which was determined to be at 285.3 eV. Photo-induced force microscopy. The microscope used is a VistaScope from Molecular Vista, Inc., operated in dynamic mode using commercial gold-coated silicon cantilevers (NCHAu) from Nanosensors. The excitation laser is a LaserTune IR Source from Block Engineering. We note that the low-power zones in Figure 2 (b) are a result of different laser systems available at certain points in time. For images and spectra, there is 0.5 - 2 mW of light focused by the parabolic mirror to a spot size of ~20 µm. The images are 256 by 256 pixels, recorded at 0.78 lines per second, for a per-pixel acquisition time of 1.5 ms, and 5.5 minutes per image. The PiFM spectra shown are acquired by sweeping the laser at 50 ms/cm-1, or 25 seconds per spectrum. Bulk FTIR spectra for comparison were obtained on a Nicolet iS50 FT/IR Spectrometer.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publications website.


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PiFM laser power profile, image analysis procedure, RSoXS 2D scattering profiles, OPV current-voltage curves, MATLAB and Mathematica image processing code (PDF). Fading overlay of Figures 2 (d) and (e) (.mp4).

The authors claim no competing financial interest.

AUTHOR INFORMATION Corresponding Author *

E-mail: [email protected]

Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

ACKNOWLEDGMENTS This work is supported by the Office of Naval Research (grant no. N00014-17-1-2214). K. G. was supported by the Department of Defense (DoD) through the National Defense Science & Engineering Graduate Fellowship (NDSEG) Program. Portions of this work were carried out at Beamline 11.0.1.2 at the Advanced Light Source, which is supported by the Director, Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231.


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Table of Contents Figure


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Nanoscale Domain Imaging of All-Polymer Organic Solar Cells by

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