Effects of Molecular Orientation of Fullerene Derivative at Donor

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Effects of Molecular Orientation of Fullerene Derivative at Donor/ Acceptor Interface on Device Performance of Organic Photovoltaics Kouki Akaike, Takumi Kumai, Kyohei Nakano, Shed Abdullah, Shun Ouchi, Yuuki Uemura, Yuta Ito, Akira Onishi, Hiroyuki Yoshida, Keisuke Tajima, and Kaname Kanai Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b03659 • Publication Date (Web): 01 Nov 2018 Downloaded from http://pubs.acs.org on November 2, 2018

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Chemistry of Materials

Effects of Molecular Orientation of Fullerene Derivative at Donor/Acceptor Interface on Device Performance of Organic Photovoltaics Kouki Akaike*,†, Takumi Kumai†, Kyohei Nakano‡, Shed Abdullah¶, Shun Ouchi¶, Yuuki Uemura¶, Yuta Ito†, Akira Onishi†, Hiroyuki Yoshida¶, Keisuke Tajima*,‡, Kaname Kanai† †

Department of Physics, Faculty of Science and Technology, Tokyo University of Science, 2641 Yamazaki, Noda, Chiba 278-8510, Japan ‡ RIKEN Center for Emergent Matter Science (CEMS), 2-1 Hirosawa, Wako 351-0198, Japan ¶

Graduate School of Advanced Integrated Science, Chiba University, 1-33 Yayoi-cho, Inage-ku, Chiba 263-8522, Japan

ABSTRACT: Designing donor/acceptor (D/A) interfaces that can efficiently generate free carriers is an attractive research target for organic photovoltaics (OPVs). While many reports suggest that the molecular orientation of the donor at the D/A interface influences the free-charge generation and recombination, the effects of the acceptor orientation on these processes remain elusive. In this work, we demonstrate that [6,6]-phenyl-C61-butyric acid methyl ester (PC61BM) changes its molecular orientation at the film surface on crystallization, resulting in the preferential surface exposure of the side chains. Photoelectron spectra of amorphous- and crystallinePC61BM/sexithiophene (6T) interfaces and analysis of the external quantum efficiency and electroluminescence of bilayer OPVs in the charge-transfer absorption range reveal that the orientational change of PC61BM raises the energy of the charge-transfer state at the D/A interface. In addition, the PC61BM side chain at the crystalline-PC61BM/6T interface reduces the electronic coupling between the charge-transfer and ground states, suppressing carrier recombination without sacrificing photocurrent. These two factors lead to the higher open-circuit voltage of crystalline-PC61BM/6T OPV compared with its amorphous counterpart. This work directly links the interface and photovoltaic properties, highlighting the role of the acceptor’s orientation in determining the efficiency of OPVs.

heterojunction can be formed by direct contact of the π-conjugated backbone or by indirect contact through the side chains. In the indirect case, the physical distance between the hole at the donor and the electron at the acceptor is larger, which reduces the Coulomb interaction binding the charge pairs in the charge-transfer (CT) states. Holcombe et al. reported that this type of steric effect can increase short circuit current density (JSC) by a factor of two27. The other potential effect, inferred from research on the effect of the spacer at the D/A interface on photovoltaic performance28,29, is that the side chains at the interface may suppress non-geminate recombination, if the side chains are insulating. Consequently, the indirect contact between the donor and acceptor increase VOC. Despite the large number of experimental and theoretical studies that highlight the importance of the geometrical structure of the donor in the interfacial charge processes15–20, there are fewer studies on the effect of acceptor’s orientation. This is probably because there is no general methodology for manipulating the orientation of the acceptor. However, it has been speculated from the measurements of solid-state ionization energy and electron affinity that the orientational change at the film surface of a prototypical acceptor, [6,6]-phenyl-C61-butyric acid methyl ester (PC61BM) (molecular structure shown in the inset of Figure 1a) occurs through thermal treatment30. We could address the impact of PC61BM’s orientation on photovoltaic performance if the change in the molecular orientation were observed and if

1. INTRODUCTION The applications of organic photovoltaics (OPVs) in energyharvesting devices have been explored based on the unique potentials of OPVs, such as their excellent mechanical-flexibility1–3 and large-area production at low cost. The power conversion efficiency of single-junction OPVs has already exceeded 13%4–6, whereas, this year, 15%7 and 17%8 has been achieved for multijunction OPVs that were fabricated by combining vacuum and solution processes and solely utilizing solution processes, respectively. Since Tang’s report9, the steady increase in the performance of OPVs has been supported by the exploration of new molecular semiconductors10–13. In addition to the development of materials, clarifying the role of the local geometry at donor/acceptor (D/A) interfaces during the operation of OPVs is necessary to provide the ground rules for efficient charge generation and suppression of geminate and non-geminate recombination14. The molecular orientation of a donor relative to an acceptor (and vice versa) at the D/A interface is an essential geometrical property critical in these electronic processes15–20. The effects of the orientational change on the photovoltaic properties of state-of-the-art organic semiconductors appear to be complex. The semiconductors typically have various side chains to ensure solubility, carrier transport, and photoabsorption11–13,21, and to control the energies of the highest-occupied molecular orbital (HOMO) and lowest-unoccupied molecular orbital (LUMO)22–26. In principle, the interface in a bulk

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structurally well-defined D/A interfaces comprising PC61BM films with different orientations were prepared. In this study, we use metastable atom electron spectroscopy (MAES), which probes the electronic structure of the outermost surface, to demonstrate the orientational change of PC61BM at the surface on crystallization. The excellent surface sensitivity of MAES enabled us to determine the relative orientation of the side chain and C60-backbone of PC61BM because the photoemissions originating from the molecular orbitals (MOs) of these moieties are separate. The analysis of the MAES spectra for amorphous and crystalline PC61BM films clearly suggests that the orientational change is accompanied by the exposure of the side chain to the surface. Based on this finding, we fabricated the following sharp D/A interfaces: amorphous PC61BM (bottom)/sexithiophene (6T; top) and crystalline PC61BM/6T interfaces. At the amorphous PC61BM/6T interface, the C60 backbone of PC61BM is in direct contact with the 6T molecules, which are in an end-on orientation. In contrast, at the crystalline PC61BM/6T interface, the C60 backbone is in indirect contact with the end-on 6T molecules through the PC61BM side chain. Using these bilayer systems, we investigated the effect on the photovoltaic properties by the molecular orientation of PC61BM at the D/A interfaces. The measurements of current densityvoltage (J-V) characteristics suggest that the PC61BM side chain at the crystalline PC61BM/6T interface increases VOC from 0.41 to 0.59 V without decreasing JSC. Analyses of the D/A interfaces with ultraviolet photoelectron spectroscopy (UPS) and of the high-sensitivity external quantum efficiency (EQE) and transient photovoltage (TPV) of the bilayer OPVs suggest that the increase in VOC is attributed to two factors. The first is the increase in the energy of the CT state (ECT) and the second is the decrease in the electronic coupling between the CT and ground states due to the PC61BM side chain at the D/A interface. This study clarifies the multiple roles of PC61BM’s orientation in determining the photovoltaic performance.

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with an electron analyzer (PHIBOS-100, SPECS). In the MAES measurements, helium atoms were excited to the metastable states, 21S (20.62 eV) and 23S (19.82 eV), by a dc discharge. To obtain pure He* (23S) atoms, He* (21S) was quenched by a separate dc helium lamp. A sample bias of –1 V was applied during MAES spectra acquisition to increase spectral intensity. X-ray photoelectron spectroscopy (XPS). XPS spectra of the PC61BM films were recorded using PHI 5000 VersaProbe (Ulvac-Phi) with monochromatic Al Kα as an excitation source. Atomic force microscopy (AFM) and X-ray diffraction (XRD) measurements. Surface morphology was analyzed by AFM (SPA300, Seiko Instrument) in non-contact mode under ambient atmosphere. XRD was carried out using a diffractometer (SmartLab, Rigaku) with Cu Kα as a radiation source. Fabrication and characterization of OPV. PC61BM/6T heterostructures were constructed on PEIE-coated ITO substrates. PEIE was spin-coated from a 0.4 wt % 2-methoxyethanol solution of PEIE at 5000 rpm for 60 s after acceleration for 5 s. The PEIE-modified substrate was annealed at 100 °C for 10 min. The resulting thickness was 3.4 nm, which was determined by XRR. PC61BM films were prepared on PEIE, followed by a two-step evaporation of 6T. Up to nominal thickness of 2.4 nm, the evaporation rate was kept at 1 Å/min to grow sharp interfaces identical to the specimen used for the electronic and structural measurements. Above 2.4 nm, the rate was increased to approximately 0.2 Å/s and 30-nm-thick 6T layers were vacuum-deposited on top of PC61BM films. Subsequent deposition of MoO3 (~0.3 Å/s, 10 nm) and Al (1 Å/s, 50 nm) completed the OPV cells. For C60/6T OPVs, a 23-nm-thick C60 film was evaporated onto PEIE at approximately 0.2 Å/s. The J-V characteristics of the devices were measured under simulated solar illumination (AM 1.5, 100 mW cm–2) from a solar simulator based on a 150 W Xe lamp (PEC-L11, Peccell Technologies). The light intensity was calibrated with a standard silicon solar cell (BS520, Bunkoh-Keiki). The active area of the devices was defined by using a 0.12 cm2 photomask. The EQE of the devices was measured on a Hypermonolight system (SM-250F, Bunkoh-Keiki). Electroluminescence was measured on a spectrofluorometer equipped with an InGaAs detector (Nanolog, Horiba). A constant DC voltage was applied to the devices and current was measured by source measurement unit (2400, Keithley). TPV and transient photocurrent measurements. TPV and transient photocurrent (TPC) measurements were carried out at room temperature. The light source for measuring small perturbations of VOC (ΔVOC) was an N2-dye pulse laser (KEC-160, Usho) with an excitation wavelength, repetition rate, and pulse duration of 532 nm, 10 Hz, and 0.4 ns, respectively. The intensity of the laser pulse was controlled by a neutral density filter to keep ΔVOC below 3 mV. The bias light source was a white LED (3 W LED XM-L, Cree) with a neutral density filter. The electrical signal was detected with a digital oscilloscope (DS5632, Iwatsu). For the TPV measurements, a low-noise differential FET amplifier with 1 MΩ input impedance (SA-421F5, NF) was used, which kept the sample under open-circuit conditions. For TPC measurements, 50 Ω resistance was put parallel to the input of the oscilloscope, and the transient current was calculated using Ohm’s law. The carrier density and lifetime were calculated using the method reported by Durrant and coworkers31. To calculate the carrier density, the differential capacitance method was used. The small perturbation lifetime, τΔn,

2. EXPERIMENTAL SECTION Materials. PC61BM, 6T, polyethyleneimine ethoxylated (PEIE), and MoO3 were purchased from Sigma-Aldrich and used without further purification. Crucibles of 6T and MoO3 were thoroughly degassed in evaporation chambers before vacuum deposition. Prior to use, ITO substrates were ultrasonicated in water, acetone, and isopropanol for 10 min each, followed by UV/ozone treatment for 15 min. To construct PC61BM/6T heterojunctions, PC61BM films were prepared from spin-casting chloroform solution (5 mg/mL) at 1200 rpm for 60 s in a glove box filled with N2. Thermal annealing of PC61BM spin-coated film was carried out at 160 °C for 10 min under N2. The resulting thicknesses of a-PC61BM and c-PC61BM were 23 and 22 nm, respectively, which were measured by X-ray reflection (XRR). The heterojunctions were completed by vacuum depositing 6T onto the PC61BM films typically at an evaporation rate of 1 Å/min. UPS and MAES measurements. UPS measurements were performed under ultra-high vacuum (base pressure < 5 ´ 10–8 Pa) using a homemade apparatus with an electron analyzer (SES200, VG Scienta) and helium discharge lamp. The He Iα resonance line (21.22 eV) was used as an excitation source to acquire the UPS spectra. The energies of the vacuum levels were deduced by using the secondary electron cutoff (SECO) of the UPS spectra at normal emission with a sample bias of –5 V. MAES spectra were recorded also under ultra-high vacuum


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Figure 1. (a) UPS spectra of a- and c-PC61BM are shown with the calculated DOS of an isolated PC61BM. The intensities of the measured spectra are normalized at the center of the first band. The energy axis is scaled relative to the center of the first band. Vertical bars below the calculated DOS indicate the distribution of MOs. The signals in energy ranges A and B correspond to the photoemission from π orbitals of the C60 backbone and to π and non-bonding orbitals localized on the side chain, respectively. The molecular structure of PC61BM is also shown. (b) MAES spectra of a- and c-PC61BM are shown with the calculated DOS. Energy ranges A and B are the same as in (a). The magnified spectra highlight the decrease in the intensity of the first band for c-PC61BM. The energy axis is scaled relative to the center of the first band. (c) Direction of the permanent dipole of PC61BM, indicated by an orange arrow. Gray, red, and white balls in the molecular structure represent carbon, oxygen, and hydrogen atoms, respectively.

contains the π orbitals on the phenyl and carbonyl groups, and the non-bonding orbital of the carbonyl group on the side chain32. The valence spectrum of c-PC61BM films in energy range A appears to be the same as that of a-PC61BM films, whereas the spectrum in range B at the high binding energy side shows a slight increase in the spectrum of c-PC61BM films. Considering the surface sensitivity of UPS, the enhancement of the band in range B, which originates from the MOs distributed on the PC61BM side chain, may imply that the side chain is extended near the film surface on crystallization. To detect the enhanced band in this range explicitly, we employed MAES, which probes the outermost electronic structure and qualitatively determines molecular orientation by examining the relative intensities of the bands originating from MOs with different characters. The excellent surface sensitivity of MAES arises from the preferential interaction of the excited helium atom in the triplet state, which is the excitation source for acquiring MAES spectra, with the wave functions that spatially extend to the outermost surface33. The photoemission in energy range A, which is attributed to the exposure of π orbitals of the C60-backbone to the outermost surface, was detected in the MAES spectrum of a-PC61BM (Figure 1b). In contrast, the clear feature in range B is seen in the MAES spectrum of c-PC61BM. This result is consistent with the enhanced photoemission visible in the UPS spectrum of c-PC61BM (Figure 1a). The stronger peak in this range clearly suggests that the PC61BM side chain is exposed toward the surface of the crystalline film. In principle, the MAES spectra might change if the desorption of the solvent molecules and/or surface degradation during heat treatment occur under the presence of solvent vapor and remaining O2 with light illumination. To inspect these undesired contaminations at the film surfaces, we measured the XPS spectra of a- and c-PC61BM (Figure S2 in the Supporting Information). The survey spectra showed no Cl 2p peak (Figure S2a), suggesting that chloroform molecules should be desorbed during the spin-coating the films. Besides, the shapes of C 1s and O 1s core level spectra for c-PC61BM are the same as a-PC61BM

was determined by fitting the TPV signal with single-exponential decays. Molecular orbital (MO) calculations. MO calculations were performed using the Gaussian03 package. After the structural optimization, DFT calculations were performed using the B3LYP exchange-correlation function and 6–311G basis set. The simulated DOS was obtained by convoluting delta functions at MO energies with a Gaussian function with a full width at half maximum of 0.5 eV. Time-dependent DFT calculations of ECT were performed using Gaussian16 at the CAMB3LYP/6-31G+(d,p) level. To examine the dependence of ECT on the selection of the functional and basis set, calculations using B3LYP/6-31G(d) and B3LYP/6-31G+(d,p) were also carried out. 3. RESULTS AND DISCUSSION Exposure of PC61BM side chain toward the film surface on crystallization. Comparing the UPS spectra for the ascoated and heated PC61BM films hinted at the exposure of the side chain outside the film surface. Prior to characterizing the electronic structure with UPS, we confirmed the crystallization of PC61BM by annealing the spin-coated film at 160 °C for 10 min under N2 and using in-plane XRD (Figure S1 in the Supporting Information). Hereafter, based on the difference in crystallinity of PC61BM, the as-coated PC61BM film is referred to as amorphous PC61BM (a-PC61BM) and the heated film is referred to as crystalline PC61BM (c-PC61BM). Figure 1a shows the normalized UPS spectra of a- and cPC61BM in the valence region. The energy axis is scaled with respect to the center of the first band and the spectral intensity is normalized at the height of the first band. Both the specimens display similar spectral features and the density of states (DOS) calculated with density functional theory (DFT) reproduces the measured spectra well. According to the DFT calculations, the first band and the lower-energy part of the second band (range A) originate from π orbitals of the C60 backbone, whereas the second band at the higher binding energy side (range B)



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Figure 2. Structural properties of the PC61BM/6T heterostructures. (a) and (b) show the evolution of the MAES spectra for a-PC61BM/6T and c-PC61BM/6T heterostructures, respectively, as a function of 6T thickness (d). The top yellow curves in (a) and (b) correspond to the MAES spectra of the 4.8-nm-thick 6T film vacuum-deposited on SiOx. (c) and (d) show out-of-plane and in-plane XRD profiles of the heterostructures, respectively, with those of neat a- and c-PC61BM films. The numbers in parentheses correspond to diffraction planes of 6T. (e) and (f) show AFM images of 16-nm-thick 6T layers on top of a- and c-PC61BM, respectively. The image scale is 1 ´ 1 µm. (g) and (h) present cross-sectional profiles plotted along the white dashed lines in (e) and (f), respectively. The horizontal dashed lines are guides for visualizing the layered structures of the 6T top layer. The inset in (g) illustrates the molecular structure of 6T together with the length of the long molecular axis40.

interface is formed, owing to the outermost sensitivity of MAES, the photoemission from only the top layer material should be detected immediately after the deposition of the top layer. In contrast, the bands originating from the bottom layer are progressively attenuated as the coverage increases. The MAES results in Figure 2a and b for a-PC61BM/6T and c-PC61BM/6T heterostructures indeed demonstrate the predicted trend. In both heterostructures, the bands originating from PC61BM disappear upon the deposition of 6T at d = 2.4 nm, indicating that the surfaces of the a- and c-PC61BM films are fully covered with 6T at this nominal thickness and that sharp D/A interfaces are formed. The spectral shapes of the 6T overlayers on both a- and c-PC61BM appear to be unchanged with increasing d, which could suggest that there is no orientational evolution in the 6T overlayer. In addition, the spectral shape is consistent with that of a 4.8-nm-thick 6T film on SiOx. Because 6T molecules on SiOx exhibit an end-on orientation37, where the molecular long axis is aligned vertically to the film surface, the molecules vacuum-deposited on the PC61BM films should adopt the same orientation as 6T on SiOx. The detailed assignment of the MAES spectra and qualitative determination of the molecular orientation of 6T are described in the Supporting Information (see Figure S4 and related discussion). The standing orientation of 6T was confirmed by XRD measurements. Figure 2c shows the out-of-plane XRD profiles of the PC61BM/16-nm-thick 6T heterostructures. (00l)-diffraction resulting from the 6T overlayer37 was clearly observed. In addition, (110), (021), and (120) diffractions of the layer were detected in the in-plane profile (Figure 2d). These results suggest the end-on orientation of 6T on the PC61BM film. For the cPC61BM/6T heterostructure, the Bragg peaks of c-PC61BM remained unchanged, even after the 6T deposition on top of cPC61BM. The MAES and XRD analyses suggest the formation

(Figure S2b, c), indicating that the chemical modification by remaining O2 and H2O34 in the glove box is not significant. The ionization energy of the PC61BM film in the solid state (Is) was found to decrease from 6.16 to 5.89 eV on crystallization, owing to the reduction of the work function (from 4.65 to 4.47 eV) (Figure S3a, b in the Supporting Information) as well as the increase of the polarization energy30. The preferential exposure of the side chain at the surface of c-PC61BM explains the change in the work function. A dipole layer with the positive charges at the outermost surface is formed because the permanent dipole of an isolated PC61BM extends from the C60 backbone to the phenyl group of the side chain (see Figure 1c). As a consequence, the work function of c-PC61BM decreases and Is of cPC61BM becomes smaller than that of a-PC61BM. Further discussion on the electronic structure of a- and c-PC61BM is available in the Supporting Information. Structural investigation of PC61BM/6T interfaces. Based on the finding that the molecular orientation of PC61BM changes on crystallization, we fabricated the two types of D/A interfaces using 6T as a prototypical donor. That is, D/A interfaces with the C60-backbone in direct contact with 6T (aPC61BM/6T interface) and in indirect contact through the side chain (c-PC61BM/6T interface). To investigate how the difference in the interface geometry affects the photovoltaic properties, it is essential to study the structure of the D/A interface. This is because the direct comparison of the device performance is difficult if the molecules of the top layer diffuse into the bottom layer forming an intermixed amorphous layer during the interface growth17,19,35. MAES is also a good tool for investigating possible diffusion of deposited materials into a layer underneath36. If a sharp


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Figure 3. Schematics of the evolution of the molecular orientation of 6T and the morphology of 6T overlayers deposited on (a) a-PC61BM and (b) c-PC61BM.

of the lack of molecular diffusion39,40 and remain small. In contrast, 6T wets better on c-PC61BM owing to the higher surface energy, resulting in crystal grains with larger terraces. The cross-sectional profiles of the AFM images of the 6T overlayer also suggest a layered structure in each crystal grain (Figure 2g and h). Parallel dashed lines in Figure 2g and h indicate the height of one layer. The average height of one layer was determined to be 2.8 and 2.0 nm for the 6T layers on a- and cPC61BM, respectively, in close agreement with the length of the long axis of 6T41 (inset of Figure 2g). The analyses may indicate the end-on orientation of 6T on both the PC61BM films, in agreement with the XRD results (Figure 2c and d). We also employed MAES to investigate the molecular orientation of 6T at sub-monolayer coverage. More discussion of our results is available in the Supporting Information. The 6T layer on a-PC61BM may be disordered or slightly tilted from the surface normal, as suggested by the decrease in band (iii) from 6 to 7 eV (Figure S4a in the Supporting Information) compared with the c-PC61BM/6T heterostructure (Figure S4b in the Supporting Information). However, as d was increased to 1.2 nm, the relative intensity of the 6T-derived bands for both the heterostructures became similar to those for the 6T film on SiOx. That is, except for the initial stage of the interface growth, 6T on the PC61BM films generally adopted an end-on orientation. Figure 3 summarizes the evolution of the molecular orientation and the morphology of the 16-nm-thick 6T overlayer deposited on a- and c-PC61BM. Energetics at the D/A interfaces. Next, we evaluated the energetics at the PC61BM/6T interfaces, employing UPS to clarify the evolution in the energy of the frontier orbitals of both molecules, because the energy difference between the donor HOMO and acceptor LUMO is directly related to VOC in OPVs42,43. The left panel of Figure 4a shows the evolution of representative UPS spectra in the SECO region for the aPC61BM/6T interface. As d increases to 2.4 nm, where PC61BM surface is fully covered with 6T (see the MAES spectra in Figure 2a), the SECO onsets gradually shift toward higher binding energy, lowering the vacuum level as 6T is deposited on aPC61BM. The vacuum level shift is 0.3 eV. The HOMO peak of

of the sharp D/A interfaces with the least structural perturbation of the bottom PC61BM layer. The roughness of the bottom PC61BM layer may affect the sharpness of the D/A interface. Figure S5a and b in the Supporting Information show the surface morphology of a- and cPC61BM films, respectively, which were measured with atomic force microscopy (AFM). Both the films exhibit leaf-like domains, and the root mean square roughness of the films was not significantly different (1.82 and 1.68 nm for a- and c-PC61BM, respectively). This indicates that when 6T is deposited on the PC61BM films, the interface areas of a-PC61BM/6T and cPC61BM/6T are similar. Although the crystal structures of the 6T overlayer on a- and c-PC61BM are similar, a major difference in the surface morphology was found. Figure 2e and f show AFM images of the 16-nm-thick 6T layer on top of a- and c-PC61BM, respectively. Cone-like 6T domains grow on a-PC61BM with a lateral size of 200–300 nm (Figure 2e). On the other hand, on c-PC61BM, 6T islands spread two-dimensionally with larger terraces than those on a-PC61BM (Figure 2f). The difference in the morphology of the 6T overlayer is demonstrated by the cross-sectional profiles (Figure 2e and f for the 6T overlayer on a- and cPC61BM, respectively). The surface wettability of the bottom PC61BM films can be related to the morphological change. To study the surface properties, we performed contact angle measurements for a- and cPC61BM using water. The contact angle decreased from 94.6 ± 0.7° to 86.9 ± 0.8° after the crystallization (Figure S5c and d for a- and c-PC61BM, respectively, in the Supporting Information). This result implies an increase in the surface energy of PC61BM, which was attributed to the surface exposure of the PC61BM side chain. In general, surface energy comprises dispersion and dipole contributions38. As the MAES results for c-PC61BM indicated (Figure 1b), crystallization by thermal annealing is accompanied by the formation of the dipole layer at the film surface, increasing the dipole contribution to the surface energy of c-PC61BM. The contact angle measurements explain the morphological difference of the 6T overlayer. On a-PC61BM with the lower surface energy, 6T grain nuclei tend to form because



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Figure 4. Evolution of representative UPS spectra for (a) a-PC61BM/6T and (b) c-PC61BM/6T interfaces as a function of 6T thickness (d). Left and right panels show the spectra in the SECO and valence regions, respectively. The red-dashed curves in the right panels of (a) and (b) are the PC61BM contributions to the spectra at d = 1.2 nm, which were extracted by subtracting the 6T fraction from the spectra. To produce the 6T fraction, the bulk 6T spectrum was normalized at the first band of the interface spectra. (c) and (d) Energy diagrams near the a-PC61BM/6T and c-PC61BM/6T interfaces, respectively. The red (black) horizontal lines indicate the locations of the HOMO and LUMO for PC61BM (6T). The energies are given in units of electronvolts. The LUMO onsets were estimated from the HOMO onsets and bandgaps determined by UPS and inverse photoemission spectroscopy of neat PC61BM30 and 6T48 films.

6T, initially appearing at a binding energy of 0.8 eV, similarly shifts by 0.1 eV (right panel of Figure 4a). The photoemission from the underlying PC61BM, originally located at approximately 2.2 eV, also shifts toward a higher binding energy by 0.1 eV. Similar energy shifts were measured for the c-PC61BM/6T interface (Figure 4b). Again, the vacuum level was lowered by 0.3 eV upon the deposition of a 2.4-nm-thick 6T layer, and the first peaks of 6T and PC61BM shifted by 0.1 eV toward high binding energy. The energy level alignment at these heterostructures appears to be similar to poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate) (PEDOT:PSS)/C60/diindenoperylene44, PC61BM/dihexylsexithiophene45, and MoO3/C60/zinc phthalocyanine heterostructures46. We performed UPS analysis for both heterostructures three times, and the diagrams are shown in Figure 4c and d. The results have implications for OPV performance. The energy difference between the 6T HOMO and PC61BM LUMO (EDA), which is related to VOC, increases after the crystallization of PC61BM. The HOMO onsets of 6T for the 6T layers on top of the a- and c-PC61BM films are at 0.46 and 0.53 eV, respectively. The HOMO onsets of PC61BM are at 1.61 and 1.51 eV for the amorphous and crystalline films, respectively, after the film surfaces are fully covered with 6T. By using these values and the previously reported bandgaps of a-PC61BM and c-PC61BM30,

the LUMO onsets of both a-PC61BM and c-PC61BM were estimated to be 0.59 eV relative to EF. Using the energy gaps of neat molecular films for estimating the LUMO levels of PC61BM and 6T in the heterostructures is reasonable, although the effect of the interface structure on the bandgaps has been discussed previously35,47. In the present work, structural perturbation, such as intermixing and concurrent formation of disordered phase on the interface growth, was not confirmed by the MAES and XRD analyses (Figure 2a–d). Consequently, EDA was roughly estimated to be 1.0 and 1.1 eV for the a-PC61BM/6T and cPC61BM/6T heterostructures, respectively. The larger EDA for the c-PC61BM/6T interface implies that the VOC of the cPC61BM/6T bilayer OPV is higher42,43. OPV performance and carrier behavior at the D/A interface. Finally, we investigated the device performance of the OPVs that contain a-PC61BM/6T and c-PC61BM/6T heterostructures as the photoactive layer. The VOC for c-PC61BM/6T OPV was higher than that for its amorphous counterpart. Figure 5a shows the J-V curves of the bilayer OPVs with a structure of ITO/polyethylenimine ethoxylated (PEIE)/aor cPC61BM/6T/MoOx/Al. The crystallization of the PC61BM film on PEIE-covered ITO upon heat treatment was confirmed by the in-plane XRD measurement (Figure S1 in the Supporting Information). The OPV performances are summarized in Table 1 along with EDA determined by UPS (Figure 4c and d). The OPVs with a c-PC61BM layer had VOC values 0.18 V higher than


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Figure 5. (a) J-V curves under AM1.5 100 mW/cm2 light irradiation (closed symbols) and dark (open symbols) and (b) EQE spectra for OPVs with structures of ITO/PEIE/a- or c-PC61BM/6T/MoOx/Al. (c) EQE spectra of the OPVs in a logarithmic plot. (d) Charge carrier lifetime plotted against the charge carrier density.

those with a a-PC61BM layer, whereas JSC was not significantly affected and FF slightly increased on crystallization. The EQE below 400 nm for this OPV with the c-PC61BM layer was slightly improved (Figure 5b). On average, the OPVs with the c-PC61BM layer showed superior performance to those with the a-PC61BM layer. For comparison, we also tested a C60/6T bilayer OPV to examine the effect of the solubilizing side chain on photovoltaic performance. The J-V characteristics and EQE spectra on a linear scale for this device are shown in Figure S6a and b, respectively, in the Supporting Information, and the extracted device parameters are summarized in Table 1 together with those of the PC61BM/6T OPVs. VOC of the C60/6T device was lower than those of a- and c-PC61BM/6T OPVs, whereas JSC was almost doubled. The smaller VOC arose from the lower LUMO level of C60 compared with PC61BM32,49. The difference in FF among the three OPVs was independent of the electron mobility (µe) of the acceptor layers; µe of C60 and PC61BM were reported to be 8 ´ 10–2 and 2 ´ 10–3 cm2 V–1 s–1, respectively50.

Table 2. Summary of VOC and extracted fitting parameters of the CT absorption in the EQE response along with EDA determined by UPS and the transport gaps of PC61BM and 6T.

The difference in VOC between a-PC61BM/6T and cPC61BM/6T OPVs can be attributed to the change in the electronic structure of the PC61BM/6T interfaces. In Table 1, we compared VOC and EDA estimated from the interface energy diagrams of the D/A interfaces (Figure 4) because the energy difference between the electron affinity of the acceptor and the ionization energy of the donor are strongly correlated with VOC51,52. However, the difference in EDA between a-PC61BM/6T and c-PC61BM/6T OPVs was roughly 0.1 eV, which can only partly explain the observed difference in VOC. Thus, the modification in the local structure at the PC61BM/6T interfaces, caused by the change in the molecular orientation of PC61BM on crystallization, probably explains the remaining increase in VOC. To investigate the effect of the orientational change of PC61BM on the OPV performance further, we performed highsensitivity EQE measurements53. In Figure 5c, the EQE spectra show a Gaussian photoresponse below the optical bandgap of the materials, which was assigned to the absorption of the

Table 1. Device performance of OPVs with acceptor/6T heterostructures. The performances are the averages of ten PC61BM/6T devices and four C60/6T bilayer OPVs. PCE denotes power conversion efficiency. Numbers in parentheses are standard deviations. EDA estimated from the UPS is also shown.



interfacial CT state. Based on Marcus theory, CT absorption in the EQE spectra is described as43 𝑓 – (𝐸CT + 𝜆– 𝐸)0 𝐸𝑄𝐸 = exp + 1 (1) 4𝜆𝑘𝑇 𝐸√4𝜋𝜆𝑘𝑇

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The change in the electronic coupling at the D/A interfaces can also affect non-geminate recombination processes. Burke et al. established a model in which the free charges and the interfacial CT state are in equilibrium and the lifetime of the CT state can dominate the recombination rate59. The radiative and nonradiative recombination rates of the CT state are proportional to the square of the electronic coupling between the CT state and ground state60,61, and f can be directly related to the recombination rates62. In Figure 5d, the carrier lifetime determined by TPV measurements is plotted against the charge carrier density. Small perturbation lifetime τΔn is converted to charge lifetime τn by τn = (λ + 1)τΔn, where λ + 1 is the reaction order31,63. Comparing the three OPVs, we found that the carrier lifetime at the same carrier density had a negative correlation with f; the C60/6T heterostructure with the strongest electronic coupling showed the shortest charge lifetime, whereas the c-PC61BM/6T heterostructure with the weakest electronic coupling showed the longest lifetime due to the side chain between the C60 backbone and 6T. In general, the lifetime depends on the D/A interface area64, but this effect is negligible in PC61BM/6T heterostructures because the interface area is similar (see the AFM images of the PC61BM surfaces in Figure S5a and b in the Supporting Information). Therefore, the difference in the measured lifetime can be attributed to the recombination kinetics through the interfacial CT state. The PC61BM side chain between the fullerene backbone and 6T reduces the electronic coupling in the CT state, resulting in the longer lifetime of the CT state at the c-PC61BM/6T interface. In other words, weak electronic coupling between the donor and acceptor materials leads to low reverse saturation current density (J0) in the equivalent circuit model of solar cells because J0 is proportional to the square of the electronic coupling matrix element62. Indeed, a monotonic decrease of J0 has been observed as the distance between the donor and acceptor material was increased by inserting an insulating layer28. We estimated J0 of C60/6T, a-PC61BM/6T and cPC61BM/6T OPVs to be 5.8 ´ 10–4, 2.2 ´ 10–4, and 3.6 ´ 10–6 mA cm–2 respectively, from JSC-VOC plots with varying illumination intensity, following the methodology reported by Tvingstedt and Deibel65. c-PC61BM/6T OPV shows the lowest J0, in good agreement with the trend in the electronic coupling deduced from the high-sensitivity EQE measurements (Figure 5c). The reduction in the electronic coupling, determining f, should decrease the voltage loss resulting from radiative recombination. Moreover, the non-radiative recombination rate follows the energy gap law, in which high ECT reduces non-radiative recombination61,66–68. Hence the higher ECT of the c-PC61BM/6T device could have an additional positive effect on the suppression of the voltage loss resulting from non-radiative recombination. These multiple effects of PC61BM’s orientation on suppressing the carrier recombination explain the majority of the increase in VOC on the crystallization of PC61BM. Note that even if the sharp interface is formed using bilayer systems, as demonstrated by this work, minor molecular contacts with different interface geometry, for instance, the contact of C60 cage of PC61BM with the molecular plane of 6T, may exist. This situation is possible if the surface roughness of underlying PC61BM is large, although the effect of minor interfaces on device performance should not be significant. Reference experiments will be required to clarify the impact of the other geometry on non-geminate recombination and VOC.

where E is the photon energy, λ is the reorganization energy in the photoexcitation of the CT state, k is Boltzmann’s constant, T is the absolute temperature, ECT is the energy of the CT state, and f is the prefactor, which is proportional to the square of the electronic matrix element describing the electronic coupling of the ground state with the CT state54. The sub-band features in the EQE spectra were fitted with this equation and the extracted parameters are summarized in Table 2, along with EDA. For accurate fitting, electroluminescence from the CT states was also measured (Figure S7a in the Supporting Information) and used for the fitting (Figure S7b–d in the Supporting Information). The trend in ECT extracted from the EQE spectra agrees with that of VOC, as reported previously43. The peak shift of the electroluminescence from the CT state shown in Figure S7a in the Supporting Information is also clear evidence that ECT reflects the change in the electronic structure at the D/A interfaces. The time-dependent DFT calculations of the singlet CT states with a long-range corrected functional, CAM-B3LYP55, qualitatively supported the EQE observations (Table S1 and Figure S8 in the Supporting Information). CAM-B3LYP can be used to describe both excited and CT states56. We used the extreme geometries that correspond to a-PC61BM/6T and c-PC61BM/6T interfaces (Figure S8a and c in the Supporting Information, respectively). When the PC61BM side chain points to 6T (Figure S8a in the Supporting Information), the calculated ECT is larger than that of the molecular geometry where the C60 backbone is in direct contact with 6T (Figure S8c in the Supporting Information). Thus, the difference in the local structure at the PC61BM/6T interface may affect ECT. However, similar to the comparison in EDA, the measured difference in ECT between aand c-PC61BM (0.06 eV, see Table 2) was too small to explain the observed difference in VOC (0.18 V) fully. Therefore, in addition to the difference in the energetic structures, charge recombination factors should be considered43 to explain the higher VOC of the c-PC61BM/6T OPV with relating the charge process to the difference in the interface geometry. To address charge recombination, the intensities of the photoresponse from CT state excitation (f) were compared to discuss the electronic coupling between the donor and the acceptor at the interface. The C60/6T device has the highest EQE subband photoresponse and the largest f value, which reflects the direct contact of C60 with the standing 6T57. The OPVs with the a-PC61BM layer showed a decreased photoresponse with a 52% smaller f value, which may be attributed to the spacer effect of the substituents in PC61BM on the CT absorption. The reduction of the photoresponse from the CT band was more drastic when the c-PC61BM layer was used, with f decreased by 81% compared with the C60/6T OPVs. This was attributed to the exposure of the side chains at the surface of PC61BM layer at the cPC61BM/6T interfaces induced by thermal annealing (Figure 1b). The difference in the density of CT states should be trivial because the interface area was expected to be similar in the two bilayer OPVs based on the similar surface roughness of aPC61BM and c-PC61BM (Figure S5a and b in the Supporting Information), and because the densities of PC61BM (amorphous: 1.64 g/cm3, crystalline: 1.69 g/cm3)30 and C60 (1.72 g/cm3)58 are similar.


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4. CONCLUSION We demonstrated that the crystallization of PC61BM film exposes the side chains of the molecules toward the film surface. Based on this finding, we fabricated two types of D/A interfaces with the following interface geometries: direct contact of the C60 cage of PC61BM with 6T and indirect contact through the side-chain layer at the PC61BM/6T interface. Structural analyses revealed the formation of sharp D/A interfaces with 6T in an end-on orientation, although the growth mode of 6T was altered by the change in the surface energy of the PC61BM film. The formation of the sharp D/A interface, which was demonstrated by MAES and XRD, enabled us to link the local structure at the D/A interface directly to the photovoltaic performance. Analyses of the high-sensitivity EQE spectra and TPV measurements for the PC61BM/6T bilayer OPVs revealed that the PC61BM side chain at the D/A interface leads to a higher ECT and suppresses the carrier recombination at the interface due to the decreased electronic coupling. Consequently, the increased ECT and the decreased electronic coupling lead to the VOC of the c-PC61BM/6T bilayer OPV being higher than that of aPC61BM/6T device by a factor of 1.4 without loss of JSC. Our results show that there are two key ways to improve photovoltaic performance. (1) The orientation of PC61BM should be controlled so that the molecular side chains are exposed toward the donor, although the control of the relative orientation of PC61BM to the donor molecules will be challenging, particularly for bulk heterojunctions. (2) More generally, the impact of the molecular orientation on the photovoltaic performance should depend on which functional groups in the acceptors (or donors) are in contact with the counterpart. For instance, the lengthy linear and branched alkyl chains that are frequently used may increase the physical distance between the donor and acceptor, possibly suppressing the charge recombination at the D/A interface more. Further investigation of the correlation between the types of the side chains and carrier behavior, probably by using well-defined bilayer systems, will provide rational designs for superior OPVs.

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ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website. Supplementary discussion on UPS spectra, XPS, XRD profiles, AFM images, water contact-angle for PC61BM films, MAES spectra of PC61BM/6T heterostructures, J-V curve and EQE spectra of C60/6T OPV, EL spectra of OPVs, TD-DFT results for PC61BM/6T interfaces. (PDF)

AUTHOR INFORMATION Corresponding Authors *[email protected] (K.A.) *[email protected] (K.T.)

ACKNOWLEDGMENT K. A. and K. K. acknowledge JSPS for financial support (18K04944 and 16K05956, respectively). DFT and time-dependent DFT calculations for the CT states were performed using the HOKUSAI-Great Wave system at the Advanced Center for Computing and Communication, RIKEN.



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