Nanoscale characterization of squaraine-fullerene-based photovoltaic active layers by atomic force microscopy mechanical and electrical property mapping

Nanoscale characterization of squaraine-fullerene-based photovoltaic active layers by atomic force microscopy mechanical and electrical property mapping

Accepted Manuscript Nanoscale characterization of squaraine-fullerene-based photovoltaic active layers by atomic force microscopy mechanical and elect...

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Accepted Manuscript Nanoscale characterization of squaraine-fullerene-based photovoltaic active layers by atomic force microscopy mechanical and electrical property mapping

Tonya Coffey, Andrew Seredinski, Jake N. Poler, Crystal Patteson, William H. Watts, Kenny Baptiste, Chenyu Zheng, Jeremy Cody, Christopher J. Collison PII: DOI: Reference:

S0040-6090(18)30730-2 doi:10.1016/j.tsf.2018.10.046 TSF 36963

To appear in:

Thin Solid Films

Received date: Revised date: Accepted date:

17 July 2018 11 October 2018 26 October 2018

Please cite this article as: Tonya Coffey, Andrew Seredinski, Jake N. Poler, Crystal Patteson, William H. Watts, Kenny Baptiste, Chenyu Zheng, Jeremy Cody, Christopher J. Collison , Nanoscale characterization of squaraine-fullerene-based photovoltaic active layers by atomic force microscopy mechanical and electrical property mapping. Tsf (2018), doi:10.1016/j.tsf.2018.10.046

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ACCEPTED MANUSCRIPT Nanoscale Characterization of Squaraine-Fullerene-Based Photovoltaic active layers by Atomic Force Microscopy mechanical and electrical property mapping. Tonya Coffeya *, Andrew Seredinskib , Jake N. Polera , Crystal Pattesona , William H. Wattsa , Kenny Baptistec,

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Chenyu Zhengc,d , Jeremy Codyc, Christopher J. Collisonc,d,e . Department of Physics and Astronomy, Appalachian State University, Boone, NC 28608, USA.

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Department of Physics, Duke University, 120 Science Drive, Durham, NC 27705, USA.

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a

c

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School of Chemistry and Materials Science, NanoPower Research Laboratory, Rochester Institute of

Technology, 84 Lomb Memorial Drive, Rochester, NY 14623, USA. (585 475 6142) NanoPower Research Laboratory, Rochester Institute of Technology, 84 Lomb Memorial Drive,

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d

Microsystems Engineering, Rochester Institute of Technology, 84 Lomb Memorial Drive, Rochester, NY

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e

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Rochester, NY 14623, USA.

14623, USA.

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Abstract

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The efficiency of organic solar cells can be increased by careful control of the nanoscale morphology of a dispersed bulk heterojunction device. Atomic force microscopy (AFM) has often been used to

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characterize morphology but debate persists over the value of traditional AFM measurements since the technique only addresses the active layer topography, and provides insufficient contrast to differentiate between components in a well-mixed composite. Using newer Kelvin Probe Force Microscopy (KPFM) and Quantitative Nanomechanical Mapping (QNM) modes, we demonstrate contrast due to differing elastic modulus and surface potentials between donor and acceptor materials and highlight the value of these techniques to understand critical materials properties as part of a comprehensive nanomorphology study. We test the value of our approach using blends of each of two anilinic 1

ACCEPTED MANUSCRIPT squaraines with phenyl-C61-butyric acid methylester. These two squaraine materials differ in chemical compatibility with the standard fullerene acceptor. We vary annealing conditions for our blended films and use the described AFM approaches to demonstrate changing domain sizes, which are affected by chemical compatibility with the fullerene. We demonstrate how KPFM measurements go beyond QNM

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to provide contrast between materials with reproducibility at a higher image resolution. With the ability

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to measure contrast between donor and acceptor material, we make a strong case for non-destructive

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microscopy data to measure effects of variations in annealing temperature on squaraine film morphology, which we confirm influences device performance and efficiency. These conclusions are

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important for informing material selection for long-term use of associated commercial devices in the

Introduction

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1.

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field.

Organic photovoltaic devices (OPV) have many advantages, including the promise of manufacture onto

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flexible and low weight substrates.[1–3] However, high efficiency that is comparable to traditional solar

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cells is difficult to achieve due to the morphology concerns and reduced mean free paths of the excited states (the excitons’ diffusion lengths) in the electron donor and acceptor materials comprising the OPV

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active layer.[4,5] To overcome the typically short exciton diffusion length, the active layer materials are

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blended together to make a dispersed bulk heterojunction (BHJ),[6] rather than a bi-layer junction. It is well known that the efficiency of OPV devices can be increased by careful control of the structure of this dispersed bulk heterojunction. The BHJ morphology can vary, from a homogeneous mixture to well phase-separated domains that are enriched by one component over the other.[7] The optimal domain size radius is considered to be comparable to the exciton diffusion length. Careful control of the blend morphology is critical to optimize the solar cell performance[8] and this can be accomplished by varying the blend ratio of the organics, and by changing fabrication parameters such as annealing temperature.[9,10] Long term annealing also occurs during operation in the field at typical running 2

ACCEPTED MANUSCRIPT temperatures higher than the ambient temperature, which influences the drop-off in efficiency over a commercial device’s lifetime. In any case, morphology must be measured in order to fully confirm models and nuanced mechanisms of operation so that materials and devices can be more effectively designed. Specifically, morphology control can only be confirmed when the domain size and purity in

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OPV active layers can be measured.

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Squaraines (SQ) are near-infrared-absorbing small donor molecules, which have been used successfully

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in recent OPV studies.[11–13] They are push-pull molecules representative of a larger class of organics

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where HOMO and LUMO energy levels can be tailored based on the relative strengths of intramolecular electron donating or accepting groups.[1,14] The excitonic diffusion length of squaraines is likely

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correlated with the formation of aggregates[15–18] and their inter-molecular charge transfer interactions[10,19] within purified crystal-like domains,[20] based on the importance of both charge

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transfer and Coulombic interaction as described by Hestand et al.[21] Annealing of active layers may be

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necessary to enable formation of crystal-like domains but it is difficult to control phase separation between SQs and fullerenes during annealing. Zheng et al. have described how squaraine domain size

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seems critical but that dark states formed in H-aggregates may also contribute to poor OPV

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performance.[22] Therefore it is vital that the purity and size of bulk heterojunction domains can be measured so that the simple but representative squaraine molecules may continue to be used for

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mechanistic study in small molecule OPV‘s.

Atomic force microscopy (AFM) has often been used to characterize the morphology of the BHJ in organic photovoltaics[23–29] as it is a nondestructive, nanoscale technique. Yet, there may always be some debate over the value of AFM measurements in understanding nanoscale morphology of the bulk, since the technique only provides a perspective of the device (or active layer) topography. New techniques such as Conductive AFM or Photoconductive AFM are starting to play a role in understanding 3

ACCEPTED MANUSCRIPT the influence of the nanoscale morphology on the efficiencies of OPV devices[30] but there may be very little contrast between the different donor and acceptor materials and the resulting interpretation may be limited.

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To help differentiate between donor and acceptor materials, newer AFM modes such as Kelvin Probe

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Force Microscopy (KPFM) and Quantitative Nanomechanical Mapping (QNM) are sometimes

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utilized.[31] In this work we acquire topographical maps created in peak force tapping mode, maps of the effective elastic modulus from the QNM technique, and maps of the surface potential in KPFM mode

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to investigate the nanomorphology of bulk heterojunction films representative of the efficient SQ-based

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OPV devices. We use the difference in elastic modulus of the donor and acceptor materials to differentiate the domains in blended films. The QNM findings have been verified with Raman spectral

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maps of the blended films. We also demonstrate how KPFM measurements go beyond QNM to provide

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contrast between materials with efficient reproducibility at a higher image resolution. With the ability to measure contrast between donor and acceptor material, we make a strong case for scanning probe

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microscopy data to measure the effect of variations in the annealing temperature on squaraine film

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morphology, which we confirm influences device performance and efficiency. We wish to test the value of our approach using two similar aniline-squaraines (see figure 1), with the

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only difference being their solubilizing side chains, i.e. (2,4-bis[4-(N,N-dibutylamino)-2,6dihydroxyphenyl]squaraine), hereafter refered to as DBSQ(OH)2, vs. (2,4-bis[4-(N,N-dihexylamino)-2,6dihydroxyphenyl]squaraine), hereafter refered to as DHSQ(OH)2. These two materials have different chemical compatibilities with the standard fullerene compound, PC61BM ([6,6]-phenyl C61 butyric acid methyl ester (hereafter refered to as PCBM), used as a reference material in the vast majority of OPV research to date. DBSQ(OH)2 has been shown to mix more readily with PCBM in bulk heterojunction devices21 as evidenced by i) smaller feature sizes in TEM upon spin-coating, and ii) spectra features 4

ACCEPTED MANUSCRIPT assigned to squaraine monomers in dilute solid solutions. DHSQ(OH)2, on the other hand, phase separates from PCBM even during spin coating and the spectrum of blended films is representative of the neat DHSQ(OH)2 films, with spectral features being assigned to contributions from intermolecular charge transfer and H-aggregation. We use the described AFM approaches to confirm the different

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domain feature sizes in blended films made from these two squaraines. We also investigate the effect of

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a subtle difference in chemical compatibility between these two squaraines and PCBM on the domain

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properties through a range of different annealing temperatures, until a thermal equilibrium is

Materials and Methods

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2.

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approached analogous to the long-term use of associated commercial devices in the field.[32]

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2.1 Device Fabrication

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Two SQ molecules (DBSQ(OH)2, and DHSQ(OH)2) were synthesized according to a one-pot two-step procedure. 26 The corresponding amine was purchased and condensed at reflux with 1,3,5-

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trihydroxylbenzene in a toluene : n-butanol (3:1, v/v) mixed solution. The yielded aniline intermediates

were green solids.

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were directly mixed with half equivalent squaric acid for the second condensation. All final products

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PCBM (>99.5%) was purchased from Solenne BV; MoO3 (molybdenum trioxide, >99.5%) was purchased from Sigma Aldrich. All the materials are stored in a N 2-filled glovebox and used as received. SQ:PCBM solutions in chloroform were prepared in vials by weighing solids and adding chloroform. All solutions were sonicated and then heated on a hot plate at 55 0C for 5 min to ensure the materials were fully dissolved. Thin films were spin cast from chloroform solution at a spin speed of 1500 RPM with total solute concentration equal to 16 mg/mL. Thermal treatment, when indicated in the text, was

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ACCEPTED MANUSCRIPT performed by placing the film on a hot plate in a nitrogen glove box. The hot plate was calibrated against an infrared thermometer and a ± 5 ºC deviation from the displayed value is presented. For organic photovoltaic devices described in this work, patterned ITO substrates were consecutively cleaned in an acetone and an isopropanol ultrasonic bath. A thin layer (8 nm) of MoO 3 was evaporated

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onto the pre-cleaned ITO substrates at a rate of 0.5 Å/s. (Alternatively, PEDOT:PSS solution (Cleavios PH

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750, σ = 10−100 S cm−1, diluted 1:1 with deionized water) was spin-cast onto ITO at a spin speed of

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5000 rpm.) The ITO\MoO3 (or ITO\PEDOT:PSS) substrates were then transferred directly into a N 2-filled

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glovebox for active layer spin casting. Solutions of SQ and PCBM (varying concentration corresponding to blend ratios) were prepared in chloroform as described above. The active layer solutions were spin

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coated at 1500 RPM in the glovebox, and were placed in a dark vacuum chamber to allow further evaporation of residual solvent. Finally, a shadow mask was applied and a 100 nm aluminum layer was

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vacuum evaporated through the shadow mask under low pressure (< 10-6 torr) to form the cathode. J-V

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characterization was performed on a Newport 91192 solar simulator at a power density of 100 mW/cm -2 (calibrated against standard InGaP solar cells fabricated in NASA Glenn Research Center, Photovoltaics

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Branch 5410) and by using a Keithley 2400 sourcemeter.

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2.2 Characterization Techniques

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To characterize our samples, we acquired Polarized Light Microscopy (PLM) images, UV-Vis absorbance (Shimazu-2401PC spectrophotometer), Raman spectral maps, AFM images, Peakforce QNM images, and Peakforce KPFM images from a Bruker Icon system. The PLM images were acquired with an Olympus IX81 microscope with either a 10x or 20x objective. 2.2.1 Introduction to Peakforce QNM and the Calibration Procedure In all Peakforce modes, the cantilever oscillates well below its resonant frequency (1 kHz – 2 kHz), and obtains at least one force curve per pixel in the image acquired. In all Peakforce modes, the topographic 6

ACCEPTED MANUSCRIPT map is generated from the z position of the tip and cantilever, which is controlled in a feedback loop by the maximum (or peak) force applied to the sample in each force curve acquired. In Peakforce QNM mode, the force curve at each pixel is also used to obtain the values of several material properties, including tip and sample adhesion, sample deformation, and elastic modulus. These values are then

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plotted as maps of the material properties. Topographic maps and maps of the material properties are

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acquired simultaneously in different channels during each scan. This allows the user to relate material

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properties directly to sample surface morphology. QNM has been described in detail elsewhere [34].

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In this work we are primarily concerned with creating maps of the elastic modulus of an organic semiconductor blend, and comparing the modulus maps to the surface morphology to determine the

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structure of the bulk heterojunction. We were therefore comparing the height channel on each file. The

obtain the modulus: 4

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linear region of the force curve is fit using the Derjaguin-Muller-Toporov (DMT) model[35] equation to

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𝐹 − 𝐹𝑎𝑑ℎ = 𝐸 ∗ √𝑅(𝑑 − 𝑑𝑜 ) 3 = 𝑘∆𝑧 3

(1)

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𝐹 − 𝐹𝑎𝑑ℎ is the force on the cantilever, which is calculated using Hooke’s law (multiplying the change in height ∆𝑧 by the spring constant of the cantilever, 𝑘.) R is the tip radius, and 𝑑 − 𝑑𝑜 is the deformation

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of the sample. The result of the fit is the reduced elastic modulus 𝐸 ∗ . This is not the same as the true

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elastic modulus for a bulk material, but rather a modulus of the tip and sample interaction:

𝐸∗ = [

1−𝜈𝑠2 𝐸𝑠

+

1−𝜈2 𝑡𝑖𝑝 𝐸𝑡𝑖𝑝

−1

]

(2)

Here 𝐸𝑠 and 𝐸𝑡𝑖𝑝 are the moduli of the sample and tip, respectively, and 𝜈𝑠 and 𝜈𝑡𝑖𝑝 are the Poisson’s ratio of the sample and tip. Our samples were molecular blends with elastic moduli in the 1 – 10 GPa range. We therefore used Bruker’s RTESPA-525 probes, with nominal spring constants of 200 N/m and resonant frequencies of approximately 525 kHz, as recommended by Bruker. Careful calibration of each 7

ACCEPTED MANUSCRIPT probe was carried out for Peakforce QNM. Our goal here was to show the compositional structure of the bulk heterojunction and, hence, we were primarily concerned with creating maps showing the relative effective elastic modulus of the two component materials in the blend, rather than measuring absolute modulus values. We therefore performed a “relative” calibration for our cantilevers, as described in the

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literature provided by the manufacturer, [36] briefly described here. First, the spring constant of each

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cantilever was calibrated by performing a thermal tune. The deflection sensitivity of each cantilever was

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calibrated using the linear region of a force curve on a sapphire substrate. Combined with the spring constant determined in the thermal tune, calibration of the deflection sensitivity converts the voltage

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change from the deflection of the laser spot on the photodiode into the vertical deflection of the

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cantilever in nanometers. Using Hooke’s law, we can then measure the force exerted on the sample by the tip. Finally, each tip’s radius was calibrated by scanning a reference sample with known elastic

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modulus. The reference sample chosen was a Polystyrene:Polymethylmethacrylate (PS:PMMA) blend

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sample. The radius of the tip was adjusted until the measured modulus of the PS:PMMA blend was between 3-4 GPa. The reference sample was chosen because its modulus was close to the modulus of

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the SQ:PCBM blend samples.

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As the tip scanned the SQ:PCBM blends, material adhered to the tip. We therefore calibrated the tip radius between each sample to ensure consistency in the modulus measurements. Again, while these

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calibration steps are not sufficient to state the absolute values of the moduli of our SQ:PCBM blends, it is sufficient for making comparative statements about the modulus of our samples and creating compositional maps. 2.2.2 Introduction to Peakforce KPFM Bruker’s Peakforce KPFM combines Peakforce tapping mode with Frequency Modulated KPFM (FMKPFM). KPFM measures the potential, or work function, of the sample surface. In Peakforce KPFM mode,

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ACCEPTED MANUSCRIPT each line of the image is scanned first to determine the topography using Peakforce tapping mode as described above. The same line is scanned again in “lift mode”. In all lift modes, a second trace of each scan line is repeated with the tip completely out of contact with the sample, held a set height above the sample surface. We used lift heights of 30 nm for our KPFM imaging. FM-KPFM has been described

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elsewhere, [37] but briefly, the tip and sample form a capacitor and a changing distance between the tip

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and sample causes the capacitance (𝐶) to vary. An electrostatic force (𝐹𝑒𝑙 ) acts on the cantilever that is

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proportional to the spatial derivative of the capacitance and the difference in surface potential between the tip and sample (∆𝑉). 1 𝜕𝐶

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𝐹𝑒𝑙 = −

2 𝜕𝑧

(∆𝑉 )2

(3)

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By applying an equal and opposite potential (or backing potential) between the tip and sample, this

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electrostatic force will disappear. This is the primary concept that underlies all KPFM techniques. A conductive cantilever held in an electric field will experience a shift in spring constant (𝑘) proportional to 𝜕𝐹𝑒𝑙

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the gradient of the electrostatic force (

), causing the resonant frequency of the cantilever (𝜔) to

∆𝜔 ≈ −

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change.

𝜕𝑧

𝜔 2𝑘

(

𝜕𝐹𝑒𝑙 𝜕𝑧

)

(4)

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In FM-KPFM, the difference in surface potential between the tip and sample is measured by determining the gradient in the electrostatic force. The electrostatic force gradient is found by modulating its value by applying a low frequency (several kHz) AC signal (𝜔𝑚 ) while simultaneously driving the cantilever at its resonant frequency. This causes side bands to be present at 𝜔 ± 𝜔𝑚 and 𝜔 ± 2𝜔𝑚 when the amplitude response of the cantilever is monitored during a drive frequency sweep. The amplitude of the side bands is monitored while a DC backing potential is applied until the side bands disappear, determining the force gradient and hence the surface potential difference between the tip and sample.

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ACCEPTED MANUSCRIPT We used PFQNE-AL cantilevers for our measurements, as recommended by Bruker. These probes have nominal spring constants of 1.5 N/m and resonant frequencies of 300 kHz. In Peakforce KPFM mode, it is possible for many samples to use these probes to acquire mechanical properties and surface potential difference information simultaneously. However, our blends had moduli on the order of 3-4 GPa—too

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high for reliable modulus measurements for probes with such a low spring constant. We therefore

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measured the material properties and the electrical properties of our samples with different cantilevers

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on different days.

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As all of our measurements were performed in ambient conditions, we did not attempt to perform absolute measurements of the surface potential difference between the tip and sample, since the

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measured surface potential can change with varying humidity levels, among other complications. As with the QNM measurements, we are simply interested in compositional maps of the surface potential

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that can be compared with the sample’s topography. We measured the average surface potential of

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pure DHSQ(OH)2 and pure PCBM using KPFM mode for unannealed samples and for samples annealed at 150 o C and 200 o C for 5 minutes. When we compared the average surface potential difference between

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tip and sample determined from these maps, we found that the average surface potential difference

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was 0.3 eV smaller (versus PCBM) for the neat unannealed SQ, and 0.1 eV smaller (versus neat PCBM) for the neat SQ, annealed at 150 o C, and 200 o C. Since the surface potential of the neat squaraine films

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was consistently lower than the surface potential of the neat PCBM films at comparable annealing temperatures, we therefore concluded that regions in the potential maps of the SQ:PCBM blends with higher surface potential differences are hereafter considered to be PCBM-rich, and lower surface potential differences are hereafter associated with squaraine-rich regions of the sample. Raman Mapping

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ACCEPTED MANUSCRIPT Raman spectra were acquired with a LabRAM Aramis microRaman system using a 532nm laser source with a nominal spot size on the order of 0.5µm, a x100 objective lens, and a 600 line/mm grating. Measurements were taken in air and at room temperature, with a spectral range of 200 to 1800 cm -1. Laser light incident on the sample surface that is scattered back through the objective is segregated by

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the grating based on the shift in energy from its interaction with the material. The intensity of these

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different bins of photons is measured, resulting in a spectrum characteristic of the material under study. Mapping spectra were acquired with an automated stage moving in half micron steps in both the x and y

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directions over a 30 µm x30 µm field. By comparing control spectra of PCBM and DHSQ(OH)2 to the spectra acquired over the mapped regions, we can visualize the heterogeneity of the samples. Spectra

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were acquired and mapping visualizations were created using LabSpec version 5.78 (HORIBA Scientific).

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3. Experimental Results and Analysis

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3.1 PLM images

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PLM images for blended films of DHSQ(OH)2 and PCBM are shown in Figure 2. In the unannealed and 150 °C anneal samples, localized dark spots represent asperities but we otherwise focus on the changes

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to the overall surface, which here is featureless and amorphous. As the annealing temperature increases, crystallites begin to form and are clearly recognizable throughout the PLM images. These

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crystallites are present in all films annealed at 175 °C and above. This is consistent with previously described differential scanning calorimetry data with phase separation exotherms of films[10] and crystallization exotherms in powder,[22] all appearing below 175°C. In the sample annealed at 200 °C for 4 hours, the crystals are significantly longer in dimension. We also note that the color of the films changes. Associated absorption spectra for these samples are presented in Figure 3. The subtle color change observed as annealing temperature is raised through 175

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ACCEPTED MANUSCRIPT °C indicates that the extent of aggregation is impacted with the physical changes of increased crystallinity. The annealing at 200 °C has a profound effect on the sample and degradation is not ruled out. Spectral changes represent the decrease in population of monomers in amorphous regions and the increase in aggregate population based on assignments made by Zheng et al.[10] Spectral data (Figure 3)

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demonstrate that there is not a significant increase in aggregate population once a certain crystallinity

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has been achieved across the entire film. However, the domains will likely keep growing as annealing

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time or temperature increases (which cannot be inferred from spectral information) and here we focus

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now on the evolution of domain size once the monomer population has been depleted.

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3.2 AFM, Peakforce QNM and KPFM.

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The average effective elastic modulus data are shown in Table 1. After full calibration of the system (see

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experimental section for details), the maps of the effective moduli of pure PCBM and pure DHSQ(OH)2 samples with the same annealing treatment as the blended samples were measured to set up a

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reference to differentiate the DHSQ(OH)2 and PCBM components present in the blend. In Peakforce QNM mode, the mean value of the modulus from the modulus channel map was calculated from

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Bruker’s analysis software. All acquired images were 100 square micrometers in size, with 512 x 512

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pixels per image, acquired at a scan rate of approximately 0.5 Hz. The averaged elastic modulus values for neat and blended samples are shown in Table 1. The variation in modulus data, represented as a standard deviation over all temperatures is unsurprising because, for the QNM technique, 20-40% experimental uncertainty in the measured value of the modulus is expected even for a calibrated system due to compounding of uncertainty in the modulus from calibration of the tip radius, cantilever spring constant and deflection sensitivity. [38] The average modulus of the pure DHSQ(OH)2 film is approximately 1 GPa (low), and the average modulus of the pure 12

ACCEPTED MANUSCRIPT PCBM films is approximately 7 GPa (high). Although our measured effective elastic moduli are somewhat lower for PCBM than those reported in previous studies, the larger values of the PCBM acceptor material relative to the donor material are consistent with those previous studies of organic semiconducting blends using QNM techniques [31,39]. These values do not significantly change with

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increasing anneal times and temperatures (for neat films), outside of our expected experimental

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uncertainty. Since the relative modulus values of the pure PCBM films and pure DHSQ(OH)2 films differ

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by a factor of 7, we can distinguish between the component materials in the blend, despite the large experimental uncertainty. For blended films, the average modulus across each entire image was

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between 2-4 GPa, consistent with the blend being a mixture of high and low modulus materials. As with

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the pure films, the hardness of the blended films does not significantly change as the annealing temperature increases. Similar results were observed for DBSQ(OH)2:PCBM blends but are not reported.

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In figures 4 and 5, below, we present the AFM topographical images and Peakforce QNM maps of 100

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m2 regions of the blend films for the DHSQ(OH)2:PCBM blends and DBSQ(OH)2:PCBM blends. The topographical images were corrected for tilt, and the DMT Modulus images have been flattened so that

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the contrast in the images is clear to aid in visual interpretation. Again, we emphasize the difference in

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relative modulus across the sample for these images, rather than the absolute modulus value. For the images shown in Figures 4 and 5, we report the root mean square roughness of each topography image

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and the average grain sizes of 20 representative grains in Table 2, below. Previous AFM studies of OPV films have used surface roughness as a quantitative measure of the changing morphology of the film with variations in the processing techniques such as annealing temperature or blend ratios [28,29]. We therefore report this statistic here for comparison with previous studies. The grain sizes of the blends increase significantly with increasing annealing temperature, as can be seen qualitatively in the images in Figures 4 and 5, and quantitatively (after AFM image analysis) in Table 2. The root mean square (RMS) roughness in the height images of the blended materials in Table 2 13

ACCEPTED MANUSCRIPT increases as annealing temperature is increased. Previous studies have linked increasing selforganization and aggregate formation with increased surface roughness [29], and we believe we are seeing the same effect here. This is supported by the fact that the RMS roughness in DHSQ(OH)2 is consistently higher than the corresponding DBSQ(OH)2 sample with the same thermal treatment. This

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reflects the larger SQ crystals formed by DHSQ(OH)2 in the blends and the greater phase separation

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previously seen in the TEM images and associated with the greater chemical incompatibility of

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DHSQ(OH)2 and PCBM over that between DBSQ(OH)2 and PCBM. [10,22] The RMS roughness in as-cast DHSQ(OH)2:PCBM blends is already higher than that of DBSQ(OH)2:PCBM films annealed at 175 °C.

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Furthermore, the grain area quickly expands to 3.8 (± 1.8) × 105 nm2 and then to 4.9 (± 1.5) × 105 nm2 for

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DHSQ(OH)2 whereas the DBSQ(OH)2 films reach an apparent maximum size of 2.0 (± 0.7) × 105 nm2. One should note that all samples have fully crystallized in the 200 o C anneal samples shown in Figures 4 and

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5. The crystalline structure of the film can clearly be seen with careful examination of the height images.

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When we compare the AFM images of the DHSQ(OH)2:PCBM blend film (unannealed, and 150 °C, 5 minute anneal, Figure 4) to the PLM images of those same samples (Figure 2), we can see bright (raised)

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areas in the AFM height images that correspond to the dark spots (asperities) in the PLM images.

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Comparing the AFM height images to the QNM maps, we conclude that these asperities are primarily squaraine because they have a low effective elastic modulus. The asperity (crystallite) size increases

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substantially up to the order of 1-2 microns in diameter upon annealing at 150 °C. Data acquired on a region where a large crystallite has formed in a DHSQ(OH)2:PCBM blended film, annealed at 175 °C for 5 minutes, are shown in Figure 4. From the QNM map, we can see that the backbone of the crystal has a higher effective elastic modulus than the surrounding material, indicating that the squaraine and PCBM are phase separating near the crystals. This is consistent with the rapid unmixing from a quasi-stable mixture to form two coexisting phases described in more detail elsewhere.[9,10] Similar to the DHSQ(OH)2:PCBM film, the DBSQ(OH)2:PCBM film annealed at 175 °C for 5 minutes forms needle-like 14

ACCEPTED MANUSCRIPT squaraine crystals of about 1 μm in length with a lower modulus than the surrounding matrix, as can be seen in Figure 5. Yet, the phase separation is not total for the DBSQ(OH)2, unlike for DHSQ(OH)2. Focusing only on the blended DHSQ(OH)2:PCBM, extreme 200 °C 5 min annealed sample, first on the QNM maps in Figure 4, we can see well-defined regions of high and low effective elastic modulus. This

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indicates that within the matrix, the organics begin to completely phase segregate at such high

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annealing temperature, and that regions of primarily squaraine and primarily PCBM form. Small regions

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of lower elastic modulus against a very large backdrop of higher elastic modulus material can still be

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seen, which implies vertical phase separation. At 200°C we have approached the melting point for DHSQ(OH)2 at 203 °C.[22] Thus the very large extent of phase separation is expected. We can also see

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the large amount of crystal structure in the PLM images for all 200 °C annealed samples (Fig. 1). Furthermore, there is evidence of cracking and thinning in PLM and optical microscope images of the

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film, which supports the fact that material is moving as the crystals are growing, also consistent with

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vertical phase separation. With such significant phase separation, the total surface area of the bulk heterojunction interface is too low for effective device operation. For example, DHSQ(OH)2:PCBM

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devices with a 1:1 blend ratio (ITO/PEDOT:PSS/ DHSQ(OH)2:PCBM/Al) drop from a power conversion

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efficiency of 0.86% ±0.12% when unannealed to an efficiency of 0.02% with 175 °C annealing, effectively showing diode behavior.[10] Nevertheless, this microscopy data provides a link which allows

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a confirmation that the contrast in DMT Modulus is associated with the difference in materials properties. The confirmation of materials assignment in these images is also made with Raman spectral mapping of the DHSQ(OH)2:PCBM 200 °C annealed films, which is explored later. These images, the quantitative data in Table 2 and their combined interpretation highlight how advanced AFM measurement techniques can provide insight into materials properties to confirm inferences from device and spectroscopy data that may not be directly made with TEM based approaches.

15

ACCEPTED MANUSCRIPT In general, BHJ morphology is well studied with transmission electron microscopy (TEM).[40] However, film samples for OPV are typically ca. 100nm thick and so if a variance in measured height is observed to be 10 nm, then this represents at least a 10% thickness variation. Such variations in film thickness could also generate contrast in the TEM image that is not associated with chemical composition. Given that

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the electron beam passes through multiple alternating domains of both donor and acceptor material in

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a bulk heterojunction film, which reduces the ability to find contrast in a TEM image, coupled with

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typical thickness variance, we therefore expect some limitations to the resolution in TEM measurements compared to specialized tomography approaches,[41] for example. When we access the surface with

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AFM, we are able to only look at individual domains, so long as a contrast mechanism is in place to

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differentiate those domains. We therefore state the value of using AFM to complement TEM data sets when trying to understand morphology and phase separation in these films.

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While we make the claim that combination of complementary morphology data sets from a variety of

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different techniques is necessary to enhance materials and mechanistic understanding for OPV, we can also recognize the value of individual data sets to obtain a deeper insight into the nature of the

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interactions between materials and some qualitative assessment of an intrinsic property. For instance,

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the DBSQ(OH)2:PCBM films annealed at 200 °C shown in Figure 5 have fully crystallized, similar to the DHSQ(OH)2:PCBM films annealed at 200 °C. However, the fully expanded domains of the DBSQ(OH)2

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crystals are much smaller (see Table 2) than those of the DHSQ(OH)2. Variations in modulus between the crystals is present for both samples, but the contrast is better defined for the DHSQ(OH)2:PCBM blend film. These data support the idea that the equilibrium domain size will depend upon the intrinsic chemical compatibility between the two materials, which will influence selection of the ideal materials blend for commercial devices. The KPFM images of 1 m2 regions for the blend films annealed at various temperatures are shown in Figures 6 and 7. In general, the KPFM technique yielded higher resolution in chemical contrast for the 16

ACCEPTED MANUSCRIPT blended films than did the QNM technique. For the QNM technique the PeakForce setpoint was at a much higher force than for the KPFM technique due to the higher stiffness of the cantilvers as discussed earlier. Therefore our pure SQ and SQ:PCBM blend samples exhibited a high adhesion to the tip, especially for the softer unannealed, and 150 °C annealed samples most likely due to the larger tapping

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force during QNM measurements. This caused dulling of the tips and lower resolution images for QNM.

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(In some cases, calibration of the tip radius using the PS:PMMA blend samples showed a marked 10-25

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nm increase in tip size after only one acquired image on SQ:PCBM samples.)

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Another key observation has led to another advantage of KPFM over QNM for distinguishing chemical and topographical contrast in these blends. In analyzing our results, we noted that in QNM maps of our

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blends, topographical changes frequently caused artifacts in the measured effective elastic modulus of the sample. We explain by recognizing how the DMT model (equation 1) describes samples that are

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initially perfectly flat, and as mentioned above, our tip radii increased during QNM scans due to the

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sample adhering to the tip. As an illustration, in Figure 8 below, a pure film of DBSQ(OH)2 annealed at 200 °C is imaged. For a pure film, one does not expect to see any contrast in the map since the relative

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modulus would not change for an even coating of the same pure material. However, in the modulus

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map, it appears that regions of high elastic modulus surround each grain. To further analyze, values of the deformation, adhesion, and DMT modulus channels were acquired at the red line in Figure 8. On top

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of the grains in this image, the measured effective elastic modulus was approximately 4 GPa, but reached roughly 11 GPa in the crevice between the grains. This is most likely due to the increased contact area between the tip and sample in the crevice. This is supported by the information in the adhesion and deformation channels. Knowing there are no differences in material, we expect the adhesion to be proportional to the contact area between the tip and sample. On the red line, the adhesive force is roughly 70 nN on the top of the grain and approximately twice that in the crevice. Similarly, in the absence of material differences, deformation should be proportional to the pressure 17

ACCEPTED MANUSCRIPT between the tip and sample, or inversely proportional to the area of contact. On top of the grains on the red line, the sample is deformed by roughly 2 nm, and it is deformed by approximately 1 nm in the crevice. The information from the deformation and adhesion channels therefore support a factor of two increase in the contact area when the tip taps in the crevice between these grains. Using equation 1 and

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knowing the peak force applied to the sample by the tip is 160 nN, one can predict an increase in the

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output of the DMT modulus channel of a factor of ~3 in the crevice compared to the top of the grain,

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matching our observations. In summary, certain topographical changes are responsible for artifacts in

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the measured effective elastic modulus and these will be problematic to the untrained user. Previous studies have noted similar problems with AFM mechanical measurements and techniques.[42]

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We therefore recommend that those who use QNM techniques carefully consider any potential impact that topographical features may have on the interpretation of the maps of the material properties. Since

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the possibility of such artifacts is completely absent with our KPFM work, we focus on KPFM images at

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smaller length scales in Figures 6 and 7 as opposed to 100 μm2 images.

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With these higher resolution KPFM images of 1 μm2 regions for the blend films in Figures 6 and 7, we attempted to move away from larger asperities, when possible, to focus on the structure of the matrix

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of the films. Looking first at the KPFM high resolution images of unannealed films (Figure 6a), we can see

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that the domain size for DHSQ(OH)2 looks larger when compared to DBSQ(OH)2 and that the image may be described in terms of a collection of small but discrete islands, with each high-work-function island having a width of 10-50 nm. On the other hand, the unannealed DBSQ(OH)2 (Figure 7a) could be described as a small-feature bicontinuous network. When the DBSQ(OH)2 sample is annealed at 150 °C for 5 min, the result is a similar bicontinuous “lace-like” network but with an enlarged feature size. For the 150 °C annealed DHSQ(OH)2 sample, the feature size does not seem to increase much beyond the unannealed sample.

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ACCEPTED MANUSCRIPT Both the DBSQ(OH)2:PCBM and DHSQ(OH)2:PCBM blend films have crystallized significantly in the 175 °C and 200 °C anneal images. In fact, it is almost impossible to find regions without crystallites in the DBSQ(OH)2 175 °C sample; yet the crystallites always remain small, perhaps 100 nm wide across the small axis. With smaller crystals at 175 °C, the DBSQ(OH)2:PCBM matrix retains its lacey bicontinuous

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structure in the KPFM maps. The 175 °C annealed DHSQ(OH)2:PCBM and 200 °C annealed

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DBSQ(OH)2:PCBM KPFM potential maps are similar in that for each case, low surface potential crystals

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(assumed to be rich in squaraine) are interspersed with smaller high surface potential crystals (assumed to be rich in PCBM). These combined images demonstrate that a higher temperature is required for

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more significant domain growth in the DBSQ(OH)2 than in the DHSQ(OH)2. For 200 °C in DHSQ(OH)2 the

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entire 1 μm2 image sits within a large squaraine crystallite, where little contrast is observed resulting from larger crystals at thermal equilibrium and extensive vertical phase separation. SQ Device data

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generally show a decrease in power conversion efficiency upon annealing, as the domain size increases.

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We illustrate this here with results of optimized DBSQ(OH)2:PCBM Devices with a blend ratio of 0.8:1.2 (ITO/MoO3/ DHSQ(OH)2:PCBM/Al) measured in our lab, in parallel to this microscopy study (Table 3).

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There is a reduction of power conversion efficiency (PCE) from 3.12±0.17% (Unannealed) to 3.01±0.09%

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(90 °C , 5 min) to 2.92±0.10% (110 °C , 5 min) as temperature of annealing increases. These trends are fully consistent with data published elsewhere [18] and, notably, DHSQ(OH)2:PCBM devices with a 1:1

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blend ratio (ITO/PEDOT:PSS/ DHSQ(OH)2:PCBM/Al) annealed at 175 °C effectively fail, showing only diode behavior.[10] Domains are consistently larger in DHSQ(OH)2 than for DBSQ(OH)2, and therefore these data support the explanation that large domains restrict the diffusion of photogenerated excitons to the BHJ interface, thereby reducing the power conversion efficiency for DHSQ(OH)2:PCBM active layer devices versus DBSQ(OH)2:PCBM active layer devices.[6] In summary, we have observed that for both squaraine samples, the extent of crystallinity increases as annealing temperature increases. Yet for DBSQ(OH)2 the crystal domain growth seems capped to a 19

ACCEPTED MANUSCRIPT lower size; DHSQ(OH)2 phase separation and domain size increase appears to have no observable limit, within our ability to measure with high resolution imaging. From a materials perspective, one may infer that the chemical compatibility of the DBSQ(OH)2 and PCBM is such that the phase separated crystal morphology coexists with a mixed state through higher temperatures. One would naturally infer that a

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phase separation exotherm onset would occur at much higher temperatures than for DHSQ(OH)2,[10]

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and that even when thermal equilibrium is achieved close to the melting point of DBSQ(OH)2, the

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domain size remains relatively small.

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The study of phase separation and morphology in two squaraines suggests that, in general, small molecules can further be designed to maintain some optimal capped crystal size even in devices “burnt

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in” through standard operation.[32] Our results suggest that a maximum crystal size is achievable and, knowing how significant crystallinity and domain size are to OPV efficiency, these morphological caps

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may limit the loss of efficiency in working devices after some burn-in period.[43] These data, combined

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with observations of morphology-based efficiency changes, provide significant design strategy

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opportunities. 3.3 Raman Spectra and Mapping

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To ultimately confirm the assignment of contrast between the DHSQ(OH)2 and PCBM, Raman spectra, with associated sample mapping, were collected from similar samples to those measured using QNM

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and KPFM techniques. The surfaces of control samples were initially exposed to a range of different laser intensities to determine a threshold for any surface damage, observed as a visible darkening of points on the samples. Hence, to avoid damage to the surface, the laser source current was adjusted, the laser intensity was attenuated by 90% and exposure at each point was limited to 1s for 200°C 4 hr samples and 0.2 s for 150°C and 175 °C 5 min anneal samples.

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ACCEPTED MANUSCRIPT Raman spectra of the representative control samples for DHSQ(OH)2 and PCBM are shown in Figure 9. The measurements in Figure 9a were made with the same optimized acquisition parameters as for subsequent mapping to ensure fair comparison. There is a significant difference in overall signal intensity between the PCBM and DHSQ(OH)2 samples as can be seen by a direct comparison of the red

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and black curves in Fig. 9a. Consequently, each spectrum from the blended films is dominated by the

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DHSQ(OH)2 character (Figures 10a and 10b). However, there is a prominent peak in the spectrum of

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PCBM at 1462cm-1 which does not correspond to any peaks in the DHSQ(OH)2 spectrum. We select a range of 1450cm-1 to 1474cm-1 in order to isolate the PCBM peak. We baseline correct the data within

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that range, and the counts at each Raman shift value are then summed. These summations are plotted

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in a map of PCBM peak intensities in the mixed samples (Figures 10c and 10d). This gives a visualization of the distribution of PCBM spectral character throughout the sample.

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There is a large variation of summed 1462cm-1 peak intensities across the pixels in the maps in Figures

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10c and 10d. This variation leads us to infer the presence of large but separate domains of DHSQ(OH)2 and PCBM within the 200 °C annealing preparations. Domain sizes are estimated to be on the order of

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0.25 to 1 µm2, based on high intensity (red and above) and no intensity (black) regions, within the 30µm

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by 30µm area. However, mapping resolution limitations preclude further precision. These results are nevertheless consistent with the grain size analysis from Table 2, which indicates a grain area of 4.9±1.5

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× 105 nm2 (0.49 μm2) for these samples, annealed at high temperatures. Mapping of the unannealed, 150 °C annealed, and 175 °C annealed samples revealed no meaningful peaks at 1462cm-1 and so no heterogeneity could be measured. This is consistent with domains being significantly smaller than the step size of the mapping, which would give an overall impression of homogeneity. The Raman mapping technique is itself restrictive in that each pixel is 500 nm wide and therefore incorporates a large number of domains. The penetration is typically greater than 200nm and so the entire thickness of the film is being sampled in the measurement. 21

ACCEPTED MANUSCRIPT Nevertheless, in 200°C annealed films, there is always a significant difference in the intensity of the 1462 cm-1 peak in the regions of high intensity (red and above) and low intensity (black) regions as are clearly seen in Figures 10c and 10d. These data provide further support for significant phase separation in DHSQ(OH)2 blended films. They confirm that assignments made in QNM images are appropriate.

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The use of Raman spectral mapping confirms the assignment of different materials to the different

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modulus and surface potential data maps and provides complementary imaging data for samples that

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are heavily annealed. However, since such high temperature annealing is detrimental for device

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optimization and given that the Raman spectral mapping has a relatively low resolution, Raman spectral

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mapping is limited in use to providing support for the modulus assignment.

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4. Conclusion.

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In summary, AFM data collected in both PeakForce QNM mode and KPFM mode allow the study of BHJ phase separation in films targeted for OPV active layers because there is significant contrast observed

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between PCBM and SQ. With Peakforce QNM we have a different foundation for determining chemical

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composition. Harder compounds do not deform as much as softer compounds. Pure harder compounds do not deform as much as hard-soft compound composites. Therefore Peakforce QNM and KPFM AFM

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study is an excellent complementary approach (along with TEM) to the measurement of domain size and purity. It may provide more convincing proof for how well mixed and how well separated the donor and acceptor materials are.

With AFM, we are able to measure the film in situ upon a typical substrate, which is representative of the device itself. By contrast, with TEM measurements, the film must be removed from the surface of the substrate and deposited onto a TEM grid. The film is inevitably modified in the process, even if by a small extent. The electron beam itself will also have deleterious effects on the sample; use of AFM to 22

ACCEPTED MANUSCRIPT study the surface of an active layer film will have minimal effects on the sample. We therefore conclude that both QNM and KPFM provide a complementary new AFM approach to morphology measurement. Despite the improved contrast between materials that QNM imaging provides, force constant restrictions on the tip use may limit this technique given that reproducibility testing becomes expensive.

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However, KPFM studies show a similar contrast between materials in a blended sample along with

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considerably higher resolution.

Finally, we confirm that chemical compatibility between donor and acceptor materials will make a big

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difference in the thermally equilibrated morphology of an OPV device. We provide supporting evidence that the domains in bulk heterojunction films grow larger as the chemical incompatibility between a

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squaraine and PCBM increases. This is observed through a range of different annealing temperatures

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and times, and will likely play a big role during long-term use of associated commercial devices in the field.[32]

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Acknowledgements

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We would like to thank Dr. John Thornton at Bruker for his valuable advice, and we also thank Bruker for the loan of the QNM and KPFM modes. This work was carried out with the financial support from

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National Science Foundation (CBET- 1603372 and CBET-1236372). The purchase of the Bruker Icon

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system was possible thanks to the funding from the National Science Foundation, DMR-Award 0821124. We would like to thank Dr. Sergey Yarmolenko at North Carolina A&T for assistance with the Raman spectroscopy.

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[32] I.T. Sachs-Quintana, T. Heumüller, W.R. Mateker, D.E. Orozco, R. Cheacharoen, S. Sweetnam, C.J. Brabec, M.D. McGehee, Electron Barrier Formation at the Organic-Back Contact Interface is the First Step in Thermal Degradation of Polymer Solar Cells, Advanced Functional Materials, 24 (2014) 3978-3985. [33] C.W. Dirk, W.C. Herndon, F. Cervantes-Lee, H. Selnau, S. Martinez, P. Kalamegham, A. Tan, G. Campos, M. Velez, Squarylium Dyes: Structural Factors Pertaining to the Negative Third-Order Nonlinear Optical Response, J. Am. Chem. Soc. 117 (1995) 2214–2225. doi:10.1021/ja00113a011. [34] T. Young, M. Monclus, T. Burnett, W. Broughton, S. Ogin, P. Smith, The use of the PeakForce quantitative nanomechanical mapping AFM-based method for high-resolution Young’s modulus measurement of polymers, Meas. Sci. Technol. 22 (2011) 125703. [35] D. Maugis, Contact, Adhesion and Rupture of Elastic Solids, Springer-Verlag Berlin, Heidelberg, 2000. [36] B. Pittenger, N. Erina, C. Su, Quantitative Mechanical Property Mapping at the Nanoscale with PeakForce QNM, 2012. https://www.bruker.com/fileadmin/user_upload/8-PDFDocs/SurfaceAnalysis/AFM/ApplicationNotes/AN128-RevB0Quantitative_Mechanical_Property_Mapping_at_the_Nanoscale_with_PeakForceQNMAppNote.pdf (accessed January 3, 2018). [37] S. Sadewasser, Experimental Technique and Working Modes, in: S. Sadewasser, T. Glatzel, (Eds.) Kelvin Probe Force Microscopy, Springer-Verlag, Berlin, 2018, pp. 3–22. [38] K. Sweers, K. van der Werf, M. Bennink, V. Subramaniam, Nanomechanical properties of alphasynuclein amyloid fibrils: a comparative study by nanoindentation, harmonic force microscopy, and Peakforce QNM, Nanoscale Res. Lett. 6 (2011) 270. [39] D. Wang, K. Nakajima, F. Liu, S. Shi, T.P. Russell, Nanomechanical Imaging of the Diffusion of Fullerene into Conjugated Polymer, ACS Nano. 11 (2017) 8660–8667. doi:10.1021/acsnano.6b08456. [40] D.M. DeLongchamp, R.J. Kline, A. Herzing, Nanoscale structure measurements for polymerfullerene photovoltaics, Energy Environ. Sci. 5 (2012) 5980–5993. doi:10.1039/C2EE02725A. [41] S. van Bavel, E. Sourty, G. de With, S. Veenstra, J. Loos, Three-dimensional nanoscale organization of polymer solar cells, J. Mater. Chem. 19 (2009) 5388–5393. doi:10.1039/B900901A. [42] M.P. Nikiforov, S.B. Darling, Concurrent quantitative conductivity and mechanical properties measurements of organic photovoltaic materials using AFM, J. Vis. Exp. (2013). doi:10.3791/50293. [43] X. Tong, N. Wang, M. Slootsky, J. Yu, S.R. Forrest, Intrinsic burn-in efficiency loss of small-molecule organic photovoltaic cells due to exciton-induced trap formation, Sol. Energy Mater. Sol. Cells. 118 (2013) 116–123. doi:10.1016/j.solmat.2013.08.006.

List of figure and table captions

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ACCEPTED MANUSCRIPT Figure 1. Molecular structure of the two squaraines used in this work, referred to as DBSQ(OH)2, and DHSQ(OH)2. Figure 2. PLM images of the DHSQ(OH)2:PCBM blended films annealed for various temperatures and times.

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Figure 3: Absorbance spectra of DHSQ(OH)2 samples after annealing.

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Table 1. The average effective elastic modulus for the pure and blended DHSQ(OH) 2:PCBM films.

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Figure 4. The AFM topographical images and Peakforce QNM maps of 100 m2 regions of the blend films

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for the DHSQ(OH)2:PCBM blends for (a) no anneal; (b) 150 o C anneal for 5 minutes; (c) 175 o C anneal for 5 minutes and (d) 200 o C for 5 minutes.

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Figure 5. The AFM topographical images and Peakforce QNM maps of 100 m2 regions of the blend films

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for 5 minutes and (d) 200 o C for 5 minutes.

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for the DBSQ(OH)2:PCBM blends for for (a) no anneal; (b) 150 o C anneal for 5 minutes; (c) 175 o C anneal

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Table 2. Average size and their standard deviations of 20 representative grains and root mean square roughness of 100 μm2 images of select blend films.

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Figure 6. The AFM topographical images and Peakforce KPFM maps of 1 m2 regions of the blend films for the DHSQ(OH)2:PCBM blends for (a) no anneal; (b) 150 o C anneal for 5 minutes; (c) 175 o C anneal for

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5 minutes and (d) 200 o C for 5 minutes. Figure 7. The AFM topographical images and Peakforce KPFM maps of 1 m2 regions of the blend films for the DBSQ(OH)2:PCBM blends for (a) no anneal; (b) 150 o C anneal for 5 minutes; (c) 175 o C anneal for 5 minutes and (d) 200 o C for 5 minutes. Figure 8: Images of the height, deformation, adhesion, and DMT modulus channels for DBSQ(OH)2 annealed at 200 °C. Values along the bold white line for each image were analyzed and compared.

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ACCEPTED MANUSCRIPT Table 3. The device performance of DBSQ(OH)2:PC61BM at a 0.8:1.2 blend ratio with and without thermal treatment. a The thermal treatment was performed in a N2-filled glovebox before evaporation of cathode layer. b The device performance is characterized at AM1.5G illumination and the results are averaged over 6, 4 and 7 devices respectively for thermal treatments None, 90 °C, 5 min, and 110 °C, 5

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min. Standard Deviations are provided as the uncertainties.

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Figure 9. Raman spectra of pure PCBM and DHSQ(OH)2. a. Comparison of DHSQ(OH)2 spectra (black)

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with PCBM spectra acquired using the same parameters (red) and the PCBM signal magnified by a factor

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of 10 (blue). b. Detailed PCBM spectra taken with a longer exposure time and higher Raman shift

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density.

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Figure 10. a(b). Sample spectra from 200°C 5min (200°C 4 hr) annealed sample corresponding to contrasting bright (red) and dark (black) points of c(d), illustrating the presence and absence of a

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1462cm-1 peak in different sample regions. c(d). Intensity maps of the 1462cm-1 peak in the 200°C 5min

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(200°C 4 hr) annealed sample with intensity scaled internally and relatively.

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ACCEPTED MANUSCRIPT Table 1. The average effective elastic modulus for the pure and blended DHSQ(OH) 2:PCBM films.

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PCBM DHSQ(OH)2 PCBM:DHSQ(OH)2

Measured effective modulus averaged over sample (GPa) No anneal 150 °C, 5 200 °C, 5 200 °C 4 hr Average min min for all T 7 8.3 7 5 6.8 1.1 1.2 0.9 0.76 0.99 2.1 2.8 3.7 2.1 2.7

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St. Devn for all T 0.20 0.20 0.28

ACCEPTED MANUSCRIPT Table 2. Average size and their standard deviations of 20 representative grains and root mean square roughness of 100 μm2 images of select blend films. Grain area (nm2)

DHSQ(OH)2 :PCBM no anneal DHSQ(OH)2 :PCBM 150 °C anneal DHSQ(OH)2 :PCBM 175 °C anneal DHSQ(OH)2 :PCBM 200 °C anneal DBSQ(OH)2 :PCBM no anneal DBSQ(OH)2 :PCBM 150 °C anneal DBSQ(OH)2 :PCBM 175 °C anneal DBSQ(OH)2 :PCBM 200 °C anneal

4.0±1.7 × 104 5.8±1.3 × 104 3.8±1.8 × 105 4.9±1.5 × 105 5.3±1.3 × 104 3.7±1.2 × 104 1.6±0.7 × 105 2.0±0.7 × 105

Root mean square roughness (nm) 4.5 6.7 6.6 7.2 1.4 1.4 3.2 5.8

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ACCEPTED MANUSCRIPT Table 3: The device performance of DBSQ(OH)2:PC61BM at a 0.8:1.2 blend ratio with and without thermal treatment. a The thermal treatment was performed in a N2-filled glovebox before evaporation of cathode layer. b The device performance is characterized at AM1.5G illumination and the results are averaged over 6, 4 and 7 devices respectively for thermal treatments None, 90 °C, 5 min, and 110 °C, 5 min. Standard Deviations are provided as the uncertainties.

FF (%) 44.18 ± 1.58 43.98 ± 0.40 52.53 ± 0.93

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PCE(%) 3.12 ± 0.17 3.01 ± 0.09 2.92 ± 0.10

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VOC(V) 0.814 ± 0.007 0.782 ± 0.012 0.735 ± 0.013

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Device Parametersb JSC (mA/cm2) -8.663 ± 0.106 -8.748 ± 0.157 -7.555 ± 0.104

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Thermal Treatmenta None 90 °C, 5 min 110 °C, 5 min

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Two squaraine compounds are blended with a fullerene for organic solar cells Mechanical and electrical property maps characterize morphology of the films Domain size increases when chemical compatibility with the fullerene lessens Film annealing conditions are varied, also impacting domain size

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