Work-Function and Surface Energy Tunable Cyanoacrylic Acid Small

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Work-function and Surface Energy Tunable Cyanoacrylic Acid Small-molecule Derivative Interlayer on Planar ZnO Nanorods for Improved Organic Photovoltaic Performance Swapnil B. Ambade, Rohan B. Ambade, Sushil S. Bagde, and Soo-Hyoung Lee ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b11865 • Publication Date (Web): 29 Nov 2016 Downloaded from http://pubs.acs.org on December 2, 2016

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ACS Applied Materials & Interfaces

Work-function and Surface Energy Tunable Cyanoacrylic Acid Smallmolecule Derivative Interlayer on Planar ZnO Nanorods for Improved Organic Photovoltaic Performance Swapnil B. Ambade, † Rohan B. Ambade, † Sushil S. Bagde† and Soo-Hyoung Lee*† †

School of Semiconductor and Chemical Engineering, Chonbuk National University, 664-14, 1ga Deokjin-dong, Deokjin-gu, Jeonju, Jeonbuk, 561-756, Republic of Korea.

KEYWORDS: planar ZnO Nanorods, inverted organic photovoltaic cell, hybrid organic photovoltaic cell, interfacial modifier, dipole moment, work function, surface energy

ACS Paragon Plus Environment

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ABSTRACT

The issue of work-function and surface energy is fundamental to ‘decode’ the critical inorganic/ organic interface in hybrid organic photovoltaics, which influences important photovoltaic events like exciton dissociation, charge transfer, photocurrent (Jsc), open circuit voltage (Voc), etc. We demonstrate that by incorporating an interlayer of cyanoacrylic acid small molecular (SML) on solution processed, spin-coated, planar ZnO nanorods (P-ZnO NRs), higher photovoltaic (PV) performances were achieved in both inverted organic photovoltaic (iOPV) and hybrid organic photovoltaic (HOPV) devices, where ZnO acts as an “electron transporting layer” and as an “electron acceptor”, respectively. For the tuned range of surface energy from 52.5 mN/m to 33 mN/m, the power conversion efficiency (PCE) in bulk heterojunction (BHJ) iOPVs based on poly(3-hexylthiophene) (P3HT) and phenyl-C60-butyric acid methyl ester (PC60BM) increases from 3.16% to 3.68% and that based on Poly[4,8-bis(5-(2-ethylhexyl)thiophen-2yl)benzo[1,2-b;4,5b']dithiophene-2,6-diyl-alt-(4-(2-ethylhexyl)-3-fluoro thieno [3,4-b] thiophene e-)-2-carboxylate-2-6-diyl)] (PTB7:Th): [6,6]-phenyl C71 butyric acid methyl ester (PC71BM) photoactive BHJ increases from 6.55% to 8.0%, respectively. The improved PV performance in iOPV devices is majorly attributed to enhanced photocurrents achieved as a result of reduced surface energy and greater electron affinity from the covalent attachment of strong electronwithdrawing cyano moiety; while that in HOPV devices, where PCE increases from 0.21% to 0.79% for SML-modified devices, is ascribed to large increase in Voc benefited due to reduced work-function effected from the presence of strong dipole moment in SML that points away from P-ZnO NRs.


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INTRODUCTION Organic photovoltaics (OPVs) have attracted scientiﬁc and economic interest due to the steep increase in power conversion efficiency (PCE).1-3 The use of solution-processed polymer organic semiconductors with inexpensive fabrication methods, such as ink-jet printing and roll-to-roll processing, can lead to low-cost photovoltaic devices. To date, the most effective OPV devices have been based on the bulk heterojunction (BHJ) concept that typically involves combining both the donor (D) and the acceptor (A) into a single composite layer where spontaneous phase separation occurs between the exciton generating polymer and a fullerene derivative [6, 6]phenyl C60 butyric acid methyl ester (PC60BM) or [6, 6]-phenyl C71 butyric acid methyl ester (PC71BM) that assist in the dissociation of the photogenerated excitons through the formation of localized nano-heterojunctions.4-6 The phase separation further progresses by extracting the charges to respective electrodes using percolation pathways formed by each phase. To achieve high PCE, it is important that charges be extracted with minimum recombinations. Organic semiconductors have intrinsically low free carrier densities, which can lead to serious injection barriers and hence a drastically reduced PCE.

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In this context, the intrinsic high electron

mobility of inorganics is a gifted advantage to overcome charge transfer limitation associated with organic materials.

[9]

Thus, the role of inorganic semiconducting nanomaterials in hybrid

organic photovoltaics (HOPVs) is getting more significant with time.11 More recently, n-type inorganics like TiO2,12-15 ZnO

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have effectively been explored as

promising electron tunneling layers in the inverted geometry of OPVs (iOPVs). From the technical perspective, the selection and use of inorganic materials in iOPVs typically necessitates matching their work function in the interface between photoactive BHJ and the electron collecting indium-doped tin oxide (ITO) transparent conducting electrode (TCE) since the barrier


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for charge transport across the organic semiconductor-electrode interface is proportional to the difference between the energy level of the highest occupied molecular orbital (HOMO) and the work function of the anode or the energy level of the lowest unoccupied molecular orbital (LUMO) and the cathode work function.24 Secondly, the difference in the surface energy of hydrophilic inorganics and hydrophobic photoactive layers possess quite a stiff challenge, resulting into trapping of charges and inadvertent backflow of charges leading to recombinations.25 Thus, engineering the interfaces between organic semiconductors and charge collecting electrodes becomes crucial in minimizing the interfacial resistances and backflow of charges in OPVs. One approach to engineer this crucial interface is by inserting an organic small molecular layer (SML) by adsorption on the inorganic oxide prior to the deposition of the photoactive layer. SMLs not only modify the nano-morphology of the photoactive layer by virtue of altered surface energy but are also responsible for lowering the cathode work function owing to the presence of net dipole moment. For example, Chen et al. demonstrated that oxetane-functionalized fullerene derivative anchored onto the electron transporting TiOx layer exerts improvement of exciton dissociation efficiency, reduction of charge recombination, a decrease of the interface contact resistance, and passivation of the surface electron traps at the interface of TiOx.26 In another study, Jen et al. demonstrated that SAMs of C60 inserted onto electron transporting ZnO can improve the electronic coupling of the inorganic/ organic interface leading to improved fill factor (ff) and photocurrent (Jsc).27 In another configuration, where ZnO plays a role of an acceptor with a polymer donor like P3HT, the interfacial modification can effectively lower contact resistance and reduce losses due to charge recombination.28 As an illustration, Wang et al. improved the P3HT/ TiOx interface by modifying with cyanoacrylic acid derivatives to achieve high charge


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dissociation of their HOPVs.29 Recently, we reported a study comprising similar SMLs as interfacial modifiers that were adsorbed (i) in between P3HT donor and sol-gel derived ZnO (ZnO-SG) electron acceptor in the bilayer HOPVs and (ii) on the top of electron transporting layer (ETL) of ZnO in iOPVs comprising of P3HT: PCBM BHJ.30 However, a detailed study of interfacial modification involving the use of solution processed planar ZnO NRs (P-ZnO NRs) is not reported yet. Compared with ZnO-SG/organic BHJ HOPVs, devices based on ZnO NRs can provide direct conduction pathways for electron transport. Importantly, the electron transfer in solar cells with ZnO NRs is known to be about two orders of magnitude faster than that with their colloidal counterparts.31 However, owing to a larger contact surface of ZnO NRs, the D/A interfacial area in ZnO-NRs/ organic BHJ HOPVs is extremely limited.32 In this work, we present a systematic study of interfacial modification of solution processed, spin-coated, P-ZnO NRs based inorganic/ organic BHJ HOPVs and iOPVs with a cyanoacrylic acid SML synthesized by us. Solvothermally synthesized P-ZnO NRs are advantageous as they can be easily spin-coated without any need of additional seed layer. It was observed that the interfacial contact in iOPVs comprising of SML-adsorbed P-ZnO NRs was significantly improved leading to enhancement in Jsc. Moreover, the conduction band edge of SML-adsorbed P-ZnO NRs was upshifted that resulted in lowering the cathode work function. In the case of HOPVs, the recombinations were found to be decreased for devices of SML- P-ZnO NRs owing to the increased interfacial area. Moreover, the enhancement in PCE in both iOPVs and HOPVs arise commonly by virtue of increased conducting pathways in devices of SML- P-ZnO NRs. RESULTS AND DISCUSSION Molecular structure and Device scheme


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Figure 1 depicts two OPV geometries studied in this work. The geometry presented in Figure 1b corresponds to “invert” structure of OPV and is termed as iOPV. In this architecture, the excitons generated in the donor polymer are dissociated into electrons and holes at the donor: acceptor interface. Through the fullerene acceptor (PCBM), the electrons are transported to ITO via ZnO and holes move towards the other electrode. Broadly, ZnO in iOPVs acts as an electron tunneling layer. In the second geometry (HOPV) presented in Figure 1c, the role of ZnO in the absence of PCBM is mainly that of an “electron acceptor”. Depending on the nature of organic interlayer, it can also be used to control the upper layer growth (BHJ in the case of iOPV and HJ in the case of HOPV) and distribution of phases, shift the interfacial energy offset between donor-acceptor materials and passivate inorganic surface trap states. The role of SML is typically to control the surface energy and tune the work function of P-ZnO NRs. To realize the excellent photovoltaic capability of OPVs, unperturbed charge transfer, and efficient exciton separation between the inorganic oxide (in this work, P-ZnO NRs) and the interfacial layer as well as between the interfacial layer and the photoactive polymers that form BHJ is necessary, which necessitates tuning the work function.33 The SML presented in Figure 1a was inserted between P-ZnO NRs layer and the photoactive layers using dipping (optimized to 90s). Briefly, P-ZnO NRs coated ETLs were dipped in a solution of SML in THF for a stipulated time. On loading SML, films were rinsed to remove non-bonded SML and dried using nitrogen (N2). SML typically consists of a carboxylic acid terminal functional group that forms covalent bonding with the hydroxyls of P-ZnO NRs thereby allowing anchoring of SML to P-ZnO NRs. The terminal functional group promotes electronic coupling between the lowest unoccupied molecular orbital (LUMO) of P3HT and the delocalized acceptor levels of the P-ZnO NRs’ conduction band and mediates charge injection. Depending


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on the packing, the SML would either stand up or lie flat on the surface of P-ZnO NRs. Such structural differences will certainly account for differing electron injection or recombinations at the site of ETL. The bonding of SML is most likely through the bridged bidentate coordination mode of the carboxylate ions (COO-) to the surface of P-ZnO NRs. Secondly, as it is known that the bridged bidentate coordination mode is the most stable one,

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strong coupling is thus

expected to result eventually in faster electron injection from SML to P-ZnO NRs. Next, on the side of the terminal functional group, an electron withdrawing moiety (cyano) is incorporated to direct the molecular dipole in the right direction, i.e., away from ZnO. Interfacial engineering is crucial in the sense that organic molecular layers sandwiched between inorganic/ organic interface provide interfacial dipole between the two layers of varying physical nature. The direction and magnitude of dipole moment are crucial in determining the work function of ZnO after interfacial modification with organic molecular layers. For example, Yip et al. proposed that devices with SAMs inducing negative dipoles (i.e. pointing away) show higher performance with enhancement in all the photovoltaic parameters than those with SAMs inducing positive dipole moment (i.e. pointing towards).35 Moreover, the protonation of P-ZnO NRs surfaces stemming from the dissociative adsorption of the carboxylic acid group to form a carboxylate bond, in which the positive charge of the proton on the surface and the negative charge on the carboxylic group also forms an interfacial dipole. Our DFT studies 30 revealed that the optimized geometry of SML has the thiophene-benzothiadiazole-thiophene (D-A-D) unit coplanar with respect to the cyanoacrylic acid group, indicating proper conjugation across the entire molecule. Due to the incorporation of benzothiadiazole between the thiophene units, the LUMO of SML has electron density spreading across the whole molecule. From the DFT


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calculations, the measured dipole moment was as high as 8.9D, ascribed to having additional electron withdrawing benzothiadiazole between two thiophene units in SML. Work Function and Surface Energy As mentioned earlier, the role of interfacial engineering is mainly to tune the work function (Φ); work functions of the ITO/ P-ZnO NRs/ SML were measured by photoelectron yield spectroscopy (AC-2; Riken-Keiki, Japan) in air.31 For better understanding, the loading of SML over P-ZnO NRs was varied by varying the dipping time in the ascending order of 30s interval each and used as electrodes further for photovoltaic evaluation. Thus the resulting electrodes with increasing SM dipping times were SML0 (corresponds to no dipping in SML, i.e. pristine P-ZnO NRs), SML30, SML60, SML90, and SML120. The results are presented in Figure 2. It is observed that the loading of SM had a significant influence on Φ that decreased from 4.4 eV (SML0) to 4.28 eV (SML30), 4.14 eV (SML60) 3.98 eV (SML90) and 4.22 eV (SML120). This indicates that thickness of SM has a certain influence in modulating Φ of P-ZnO NRs. In the interfacially engineered photovoltaic devices, Φ is tuned by many factors like direction of dipole, surface energy or the chemical composition at the site of modification. The reduction of workfunction is appreciable since SML has a strong dipole moment pointing away from ZnO. Kim et al. explained that work-function is influenced by the direction of dipole wherein for a dipole pointing towards ZnO, Φ is increased while for that pointing away, Φ is reduced.36 The reduction in Φ is ascribed to strong electron-withdrawing and donating power of the substituent in SML. From this result, it is observed that Φ continues to decrease with increase in SML loading, meaning more number of SML substituent to participate in effective strong electron-withdrawing and donating power of the substituent in SML. From this result, it is observed that Φ continues to


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decrease with increase in SML loading, meaning more number of SML substituent to participate in effective coupling with P-ZnO NRs. The surface energy calculated from water contact angle measurements (Figure 2, Table S1) indicate a linear decrease in surface energy (increase in water contact angle) with respect to the time of SML loading. The surface energy was found to be tuned from 51.03 mN m-1 (SM0) to 44.76 mN m-1 (SM30) 37.98 mN m-1 (SM60) 35.15 mN m-1 (SM90), 46.29 mN m-1 (SM120). The lowering of surface energy suggests specific adsorption of SML over P-ZnO NRs. The difference in surface energies is ascribed to the difference in binding constants of P-ZnO NRs and SML that affects surface coverage. Since the concentration of SML is same, it is obvious that the monolayer coverage and surface area occupied by SML is different for each time of loading. Typically, the surface coverage increases with increase in dipping time from 0s to 90s. However, further loading leads to additional co-adsorption of SML and the eventual intercalation of solvent molecules. We believe that higher loading of SML leads to surface aggregation of SML over P-ZnO NRs, which is in line with our previous findings. Surface aggregation typically causes etching or chemical heterogeneity at the interface. Thus, we ascribe the increase in surface energy for SML120 to overloading of organic SML. This result demonstrates that work function is tuned with variation in surface energy as resulted from different loading density of SML onto P-ZnO NRs. Photovoltaic Testing SML-Modified iOPV Devices To understand the effect of improvisation by SML at the P-ZnO NRs interface, iOPVs consisting of ITO/ P-ZnO NRs/ SML/ P3HT: PC60BM/ MoO3/ Ag were fabricated using SML


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loaded on P-ZnO NRs (Figure 3a) at same conditions (SML30, SML60, SML90, SML120). Device with no SML (SML0) as an ETL was also fabricated as a reference. The detailed fabrication process is described in the experimental section. The average performance of all iOPV devices is summarized in Table 1. The illuminated current density-applied voltage (J-V) curves are shown in Figure 4a. For the reference device without SML, the maximum PCE attained was 3.14% owing to Jsc of 9.38 mA/ cm2, Voc of 0.55 V and ff of 0.61. These were consistent with our previous device results, with slight variations in photovoltaic parameters. After adsorbing SML initially for just 30s, the magnitude of the photocurrent decreases while Voc increased by 0.03V. As the maximum current of 10.02 mA/ cm2 was obtained for iOPVs of SML loaded for the 90s (Device SML90), the Voc and ff of this device were 0.61V and of 0.6 respectively to result in maximum PCE of 3.68 %. For the thicker device (SML120), PCE drastically decreases to 2.64% arising from a weaker photocurrent of 8.27 mA/ cm2 and Voc of 0.58V. Strikingly, with the higher loading time of 120s, the ff drops sharply to 0.55. The enhancement of photocurrent is indicative of improved charge transfer characteristics of SML-modified iOPV devices. Figure 4b compares the external quantum efficiency (EQE) of all the iOPV devices. It is evident from Figure 4b that the improvement in EQE is in parallel with enhancement in Jsc except for the device SML30. It is very strange that despite lowering of Jsc in device SML30, the EQE is still higher. Lower Jsc in SML30 is attributed to incomplete coverage of SML owing to a very short time dipping. Under such circumstances, it is possible for an undesired interface to form, most likely pin-holes that would hamper the charge transfer across the interface. From the Hall electron mobility data (Table S2), it is evident that enhancement in charge transfer is majorly due to increase in electron mobility for SML-modified iOPV devices. The decreased electron mobility for SML120 is suggestive of the increased aggregation of SML


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at the interface. Surface aggregation is detrimental as it leads to etching of the oxide surface. Additionally, the carboxylic groups of excess SML typically act as charge traps that inhibit charge transfer across the interface. The universality of interfacial modification of P-ZnO NRs was verified by fabricating iOPVs based on Poly[4,8-bis(5-(2-ethylhexyl)thiophen-2-yl)benzo[1,2-b;4,5b'] dithiophe ne-2,6-diylalt-(4-(2-ethylhexyl)-3-fluorothieno[3,4-b]thiophen e-)-2-carboxylate-2-6-diyl)] (PTB7: Th): [6,6]-phenyl C71 butyric acid methyl ester (PC71BM) photoactive blends. As shown in Figure 4c and Table 1, the SML modified devices show obvious enhancements in their photovoltaic characteristics compared to that of iOPVs of only P-ZnO NRs ETL. The optimized device SML90 exhibits a striking PCE of 8.0 % owing to an enhanced Jsc of 14.50 mA/cm2, ff of 0.69 in contrast to a lower PCE of 6.55 % yielded by iOPVs of P-ZnO NRs ETL. The enhancement in Jsc upon interfacial modification of P-ZnO NRs ETL by SML is supported by the increase in EQE responses (Figure 4d) of all the SML-modified devices compared to iOPVs of only P-ZnO NRs as ETLs. Also, the EQE response onsets for all devices are consistent with the absorption onset of donor PTB7: Th. Photovoltaic characteristics of P3HT: PC60BM BHJ iOPVs as a function of surface energy in the range of 51.03 mN/m to 35.15 mN/m for SML loaded P-ZnO NRs are presented in Figure 5. It is evident from Figure 5 that the maximum PCE is obtained at the minimum surface energy (35.15 mN/m) that corresponds to device SML90. From Figure 5, it appears that the trend in Jsc and PCE is closely consistent. Voc also follows a similar pattern; however, the slopes for the trend in Voc are slightly different in comparison to the slopes of a trend in Jsc and PCE. It is thus obvious that the enhancement in PCE is majorly dominated by the increase in photocurrent, Jsc although the contributions by enhancements in Voc and ff cannot be ignored. The trend in Voc


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follows that Voc changes linearly with the change in work function. As explained earlier, the work-function has a strong influence on surface energy. Thus, as the surface energy is reduced, the interfacial contact of P-ZnO NRs with BHJ photoactive is enhanced that eventually helps in eliminating charge traps that typically arise from surface defects. This explanation is authentic since the dark current is considerably suppressed in SML-modified iOPV devices. SML-Modified HOPV Devices To obtain the evidence of exciton dissociation, hybrid organic photovoltaic (HOPV) devices with configuration ITO/ P-ZnO NRs/SML/ P3HT/ MoO3/ Ag were fabricated (Figure 6), where P-ZnO NRs typically act as an electron acceptor. The resulting J-V curves are presented in Figure 7 and performances summarized in Table 2. Reference device with SML (SM0) showed poor photovoltaic capability yielding PCE of 0.21% resulting from Jsc of 1.15 mA/ cm2, Voc of 0.4V and ff of 46.02%. The Voc and ff of SM0 are too poor. Since in the HOPV devices, Voc is determined by the difference in the donor highest occupied molecular orbital (HOMO) and the conduction band (CB) of n- type semiconductor. Considering that the work function of SML0 is higher (4.4 eV), it is thus obvious to expect a lower Voc. The low ff in the SM0 device is due to the poor interface between P-ZnO NRs and P3HT. As semiconducting oxides are known to possess surface hydroxyls, these act as recombination media to result in a poor interface. It is, however, noteworthy that this PCE is many orders higher than that observed for HOPV devices of sol-gel derived ZnO. In our experience, this PCE is perhaps the most competitive as far as HOPV devices comprising spincoated P-ZnO NRs is concerned. This is ascribed to the moderate surface energy of P-ZnO NRs that we synthesized. With a further decrease in surface energy on adsorbing SML, the interface is


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considerably improved that enhances the photovoltaic performance. The highest PCE of 0.79% was observed for an SML90 device that exhibited lowest surface energy amongst all the devices under consideration. When modified with the SML, the HOPV devices show an overall improvement in almost all photovoltaic parameters (Jsc, Voc and ff), depending on the SML loading time. Figure 8 shows the plots comparing the effect of the different SML and their surface energy on the Jsc, Voc, ff and PCE of HOPV devices. Comparing the slopes in photovoltaic parameters of iOPV (Figure 5) and HOPV devices (Figure 8), it follows that entirely different mechanism exists in both the device architectures.

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In the case of iOPV devices, SML- P-ZnO NRs

typically act as electron selective layers that help the smooth transfer of electrons to electron collector (ITO). In these OPVs, exciton dissociation takes place at the interface between donor and acceptor (P3HT: PCBM) that forms a photoactive BHJ. Whereas, in the case of HOPV devices, ZnO acts as an electron acceptor along with polymer (P3HT) to form a heterojunction, where a critical process of exciton dissociation takes place at the ZnO: P3HT interface. Secondly, after dissociation, ZnO provides continuous pathways for transporting charge carriers to electrodes. In HOPVs, obtaining high efficiencies is challenging because of the strong incompatibility between the polymer and inorganic semiconductor. The role of interfacial modifier (in this study, SML) is typically to improve the compatibility between the polymer and inorganic semiconductor and to prevent the back recombination by tuning the energy level alignment. From Figure 8, it is evident that all photovoltaic parameters improve on modifying P-ZnO NRs by SML. More precisely, the enhancement is linearly related to the surface energies. From Figure 8b, we see that the Voc rises abruptly on modifying the interface of P-ZnO NRs by SML.


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The values of Voc for all modified devices are in the range of 0.7 V, in sharp contrast to 0.4 V for the unmodified device (SM0). Voc in HOPV devices originates from the difference in electronic distance between the conduction band level of metal oxide and the HOMO level of the polymer. [37]

In our studies, we tune the work function by varying the loading time of SML on P-ZnO

NRs. Typically; we observed that the decrease in work function was in parallel to decrease in surface energy. Thus, for the lowest surface energy, the work-function obtained was lowest. i.e. the conduction band shifts closer to the LUMO of P3HT or farther from HOMO of P3HT. The lowering of work-function and the eventual increase in Voc is majorly attributed to a strong dipole induced by SML that points away from P-ZnO NRs to uplift the conduction band. In previous studies, Wu et al. explained that one of the causes for increased Voc is the elevated quasi-Fermi level of the HOPV devices that also contribute to enhancement in Jsc.39 In one of the devices, SML30, Voc increases irrespective of Jsc, suggesting that overall, the major contribution to Voc comes as a result of work function tuning. Figure 9a compares the EQE spectra of all HOPV devices. All devices show UV and visible bands in the regions 345-420 nm and 450-600 nm, relating to photocurrent from acceptor (ZnO) and donor (P3HT) respectively. Compared to the absorption spectrum of unmodified P-ZnO NRs (SML0), the EQE spectra of SML-modified P-ZnO NRs show considerable enhancements that are attributable to the photoelectrons from SML at the critical interface between P-ZnO NRs and P3HT donor. A close look at the absorbances in the EQE reveals that the UV absorbance of all SML modified devices red-shifts while the absorbance in the visible region blue shifts compared to those in SML0 HOPV device. We observe that the shifts in UV region are more prominent while the ones in the visible region are minor. Moreover, it is interesting to note that the redshifts in UV region are dependent on the loading of SML on P-ZnO NRs. Typically, as the


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loading time increases, the magnitude of red-shift of EQE peak in the UV region increases. This suggests that with an increase in the loading of SML on P-ZnO NRs, the density of conjugation around P-ZnO NRs increases, obviously as a result of the electronic coupling of SML and PZnO NRs. To prove if such electronic coupling is effective or not, we measured the ratio of the intensity of peaks in UV and visible region (IUV/ IVis) of the EQE spectra. It is revealed that as the time of loading increased from 0s to 90s, IUV/ IVis ratio increases from 0.65 (for SML0) to 0.85 (for SML90), suggesting that the contribution to photocurrent from P-ZnO NRs acceptor gets dominant rather than that from P3HT donor. For, thicker SML adsorption (SML120), IUV/ IVis decreases, probably as a result of overloading of organic SML that would possibly increase the rate of recombinations at the interface. In HOPVs, another event that is responsible for photocurrent is exciton-charge dissociation.40 To assess whether the change in Jsc is related to the exciton dissociation at the donor: acceptor (P-ZnO NRs/ P3HT) interface, photoluminescence (PL) quenching measurements were conducted. Figure 9b presents the PL emission spectra for the different SML modifications resulting from excitation of the P3HT at 500 nm. The maximum PL intensity is exhibited by the P3HT sample that is coated on glass. The PL spectrum of P3HT sample exhibited two known shoulders at 648 nm and 725 nm, corresponding to the typical PL emission of P3HT are in agreement with earlier reports.41 The second highest PL intensity is observed for P-ZnO NRsP3HT sample, which shows distinguished quenching in comparison to that of the only P3HT sample. Furthermore, the PL intensity of SML modified P-ZnO NRs is reduced remarkably compared to that of unmodified P-ZnO NRs sample. The most efficient quenching is exhibited by SML90 sample, which is attributed to the maximum exciton dissociation efficiency in SML90 amongst all the modified samples. Thus, the emergence of highest Jsc for ML90 is justified. The


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PL quenching analysis suggests that there are appreciably large charge transfer states at the SML90-P3HT interface compared to other loading conditions. This is attributed to a much favorable energy level alignment in case of SML90, which is ultimately tuned as a function of surface energy. The improved PCEs of HOPV devices are also attributed to enhancement in ff values in the SML-modified devices. The increase in ff is attributed to varying resistances in SML-modified devices. The series (Rs) and shunt (Rsh) resistances of HOPV devices calculated from the slopes of dark J-V characteristic curves (Figure 7b) presented in Table 2 reflects that device SML90 exhibited minimum Rs and maximum Rsh values. Effectively, SML90 also exhibited highest ff (60 %) that was about 30 % greater than the unmodified device (SML0). Reduced leakage as evident from dark J-V curves corroborates the increase in Rsh in the SML-modified devices. Moreover, a largely improved rectification ratio from 0.0034 to 0.112, 75.76, 222.83 and 118.28 for SML0, SML30, SML60, SML90, and SML120 respectively, further proves that SMLmodified devices had much-improved diode characteristics. A drastically reduced rectification ratio for SML120 indicates that higher SML loading is detrimental to yield a good photovoltaic effect in HOPVs. It is evident since the high loading increases the area at the interface which would only lead to greater charge recombination. It may be inferred from the best possible device (SML90) that the chemical coupling of SML and P-ZnO NRs is more optimum in case of SML loaded for the 90s on P-ZnO NRs. To confirm the formation of favorable coupling in SML90, the surface coverage was monitored using fieldemission scanning electron microscope (FE-SEM) analysis and the chemical composition was revealed from energy dispersive X-ray (EDS) microanalysis while the elemental depth profile was studied using X-ray photoelectron spectroscopy (XPS).


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Morphological Evaluation of SML90 FE-SEM image in Figure 10a reveals that P-ZnO NRs are ~80 nm long and ~25 nm in diameter. It is seen that solvothermally processed P-ZnO NRs are planar and not oriented vertically. The planar geometry of 1-D nanomaterials offers higher surface than that of vertical and hence is advantageous from the perspective of electronic coupling with immediate layers. On adsorbing SML on P-ZnO NRs by immersing in tetrahydrofuran (THF) solution of SML for an optimized time (the 90s) results into uniform coverage of SML over P-ZnO NRs (Figure 10b). The absence of any NRs in Figure 10b clearly indicates that P-ZnO NRs have been completely covered by SML. Such compact structure is certainly necessary to improve contact between the photoactive layers that would be coated over it. The EDS microanalysis of SML- P-ZnO NRs sample (Figure 10c) confirms the presence of all elements corresponding to ZnO (Zn, O) as well as those corresponding to SML (S, C). From the XRD pattern in Figure 10d, the intense (100), (002) and (101) Bragg peaks of ZnO are in good agreement with the typical wurtzite crystal structure of ZnO with Zn 2+ and O 2− terminated polar surfaces along c-axis (hexagonal phase, space group P63mc, JCPDS No. 36-1451). [40] 42 The intensity of peaks is seen to be reduced in SML- P-ZnO NRs sample suggesting that P-ZnO NRs have been covered. In the XRD, no signals related to SML were detected and nor were any pronounced peak shifts observed. Chemical Composition Quantitative chemical compositions of P-ZnO NRs and SML-P-ZnO NRs samples were investigated by XPS (Table 2). Figure 11a, b shows the comparison of Zn and O XPS peak positions of both the samples, respectively. The Zn2p spectrum of P-ZnO NRs shows two strong


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peaks at 1022.58 and 1045.5 eV, corresponding to Zn 2p3/2 and Zn 2p1/2, while those for SML- PZnO NRs are observed at relatively lower binding energies, 1020.593 and 1043.635 eV, respectively. The O1s XPS spectrum for P-ZnO NRs is deconvoluted into three sub-peaks43 namely at low binding energy (P1) corresponding to O2- ions surrounding Zn atoms of normal wurtzite structure of ZnO single crystal, the middle peak (P2), wurtzite attributed to O2- ions in the oxygen vacancies on the ZnO nanostructure and the peak positioned at higher binding energy (P3), attributed to chemisorbed oxygen. O1s spectra indicate that the ratio of P1: P2: P3 changes after interfacial modification of P-ZnO NRs with SML. Moreover, the shifts in peak positions of Zn as well as O indicate that SML has fully covered P-ZnO NRs, consistent with morphological investigations (FE-SEM, Figure 10b). The SML-ZnO NRs sample exhibited characteristic S2p peaks (arising from sulfur atoms of benzothiadiazole moiety of SML) that are assigned into S2p3/2 and S2p1/2 respectively at 163.977 eV and 165.391 eV binding energies (Figure 11c). An additional S2p shoulder appears at 167.266 eV. Kim et al. ascribed this S2p high binding energy emission lines to the possible formation of a doped state of the sulfur atom from a polymer comprising benzothiadiazole units.44 As presented in Figure 11d, N1s core energy peak is also observed at 399.06 eV for SML- P-ZnO NRs sample. A new peak appearing at higher binding energy (401.172 eV) is attributed to N-O bond formation, which is consistent with previous findings.45 Thus, from the evidence of strongly doped states of both S as well as N atoms, we believe that SML is evenly distributed over P-ZnO NRs. CONCLUSIONS We demonstrate that the surface energy and the work function of solution processed P-ZnO NRs can be tuned by incorporating a conjugated cyanoacrylic acid, small molecule derivative. The presence of strong dipole moment in SML that points away from ZnO lowers the work


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function and upshifts the conduction band of ZnO. More precise control of electronic coupling by optimizing the density of SML at the interface, by executing control on thickness can achieve a much needed suitable interface to result in higher performance in both iOPV and HOPV devices. SML-modified P-ZnO NRs have used as the electron transporting layers and as an electron acceptor in iOPV and HOPVs, respectively. The change in Voc is consistent with the change in work function, while the increase in Jsc is supported by lowering of surface energy at the inorganic-organic interface. Moreover, the presence of strong electron withdrawing cyano moiety enhances the electron affinity, which also helps in extracting photogenerated charge carriers. The results of both the OPV architectures suggest that the inorganic-organic interfacial engineering is crucial in realizing excellent charge transfer interfaces. While, limited loading is expected to have a lower coverage suggesting the occurrence of possible undesired pinholes, a much thicker SML simply increases the interfacial area that serves as recombination centers for charge carriers. AUTHOR INFORMATION Corresponding author, *E-mail: [email protected], Tel: +82 63 270 2435 Fax: +82 63 270 2306.

ACKNOWLEDGEMENTS This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future planning (2015R1A2A2A01004404). This research was also supported by a grant from the Program of the Korea Research Institute of Chemical Technology (KRICT), Republic of Korea. Supplementary Information


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Experimental details, General Instrumentation, Characterizations, and tables. "This information is available free of charge via the Internet at http://pubs.acs.org/." REFERENCES (1) Su, Y. W.; Lan, S. C.; Wei, K. H. Organic Photovoltaics. Mater Today, 2012, 15, 554-562. (2) Li, G.; Zhu, R.; Yang, Y. Polymer Solar Cells. Nat. Photonics, 2012, 6, 153-161. (3) Kumar, P.; Chand, S. Recent Progress and Future Aspects of Organic Solar Cells. Prog. Photovoltaics, 2012, 20, 377-415. (4) Vanlaeke, P.; Swinnen, A.; Haeldermans, I.; Vanhoyland, G.; Aernouts, T.; Cheyns, D.; Deibel, C.; D’Haen, J.; Heremans, P.; Poortmans, J.; Manca, J. V. P3HT/PCBM Bulk Heterojunction Solar Cells: Relation Between Morphology and Electro-optical Characteristics. Sol. Energy Mater. Sol. Cells, 2006, 90, 2150-2158. (5) Kim, Y.; Cook, S.; Tuladhar, S. M.; Choulis, S.; Nelson, J.; Durrant, J. R.; Bradley, D. D. C.; Giles, M.; McCulloch, I.; Ha, C. S.; Ree, M. A Strong Regioregularity Effect in SelfOrganizing Conjugated Polymer Films and High-Efficiency Polythiophene:Fullerene Solar Cells. Nat. Mater., 2006, 5, 197-203. (6) Guo, S.; Ning, J.; Körstgens, V.; Yao, Y.; Herzig, E. M.; Roth, S. V.; Müller-Buschbaum, P. The Effect of Fluorination in Manipulating the Nanomorphology in PTB7: PC71BM Bulk Heterojunction Systems, Adv. Energy Mater. 2015, 5, 1401315. (7) Tress, W.; Merten, A.; Furno, M.; Hein, M.; Leo, K.; Riede, M. Correlation of Absorption Profile and Fill Factor in Organic Solar Cells: The Role of Mobility Imbalance. Adv. Energy Mater., 2013, 3, 631-638. (8) Proctor, C. M.; Kim, C.; Neher, D.; Nguyen, T.-Q. Nongeminate Recombination and Charge Transport Limitations in Diketopyrrolopyrrole-Based Solution-Processed Small Molecule Solar Cells. Adv. Funct. Mater. 2013, 23, 3584-3594. (9) Savoie, B. M.; Movaghar, B.; Marks, T. J.; Ratner, M. A. Simple Analytic Description of Collection Efficiency in Organic Photovoltaics. J. Phys. Chem. Lett., 2013, 4, 704-709. (10) Sofos, M.; Goldberger, J.; Stone, D. A; Allen, J. E.; Ma, Q.; Herman, D. J.; Tsai, W.–W; Lauhon, L. J.; Stupp, S. I. A Synergistic Assembly of Nanoscale Lamellar Photoconductor Hybrids. Nat. Mater., 2009, 8, 68-75. (11) Dayal, S.; Nikos, K.; Olson, D. C.; Ginley D. S.; Rumbles, G.; Photovoltaic Devices with a Low Band Gap Polymer and CdSe Nanostructures Exceeding 3% Efficiency. Nano Lett., 2010, 10, 239-242. (12) Zhu, R.; Jiang, C. Y.; Liu, B.; Ramakrishna, S. Highly Efficient Nanoporous TiO2Polythiophene Hybrid Solar Cells Based on Interfacial Modification Using a Metal-Free Organic Dye. Adv. Mater., 2009, 21, 994-1000.


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Figure 1. (a) Chemical structure of the conjugated cyanoacrylic acid derivative small molecule (SML) incorporating a strong electron-withdrawing -CN moiety adjacent to the -COOH molecular anchoring group, (b) iOPV and (c) HOPV representative device structures used in this study.


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Figure 2. Work function and surface energy as a function of SML loading on P-ZnO NRs.


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Figure 3. (a) Schematic representation of iOPV BHJ device structure, (b) Tuned energy alignment on introducing an SML dipolar layer (pointing away).


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Figure 4. (a) J-V characteristics, (b) EQE spectra of ZnO/P3HT: PCBM iOPV devices with different SML loaded P-ZnO NRs, (c) J-V characteristics, (d) EQE spectra of ZnO/PTB7-Th: PCBM iOPV devices with different SML loaded P-ZnO NRs.


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Figure 5. Photovoltaic characteristics of P3HT: PC60BM BHJ iOPVs as a function of surface energy of SML loaded P-ZnO NRs: (a) short-circuit current density, Jsc, (b) open-circuit voltage, Voc, (c) fill factor, ff and (d) power conversion efficiency, PCE obtained under AM 1.5 solar illumination.


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Figure 6. (a) Schematic representation of HOPV device structure, (b) Band alignment and charge transfer processes at the Donor: acceptor interface after incorporating an SML dipolar layer (pointing away).


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Figure 7. Current density-voltage characteristics of ZnO: P3HT HOPV devices with different SML loaded P-ZnO NRs tested under (a) simulated AM 1.5G illumination at 100 mW/ cm2. (b) dark conditions.


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Figure 8. Photovoltaic characteristics of HOPVs as a function of surface energy of SML loaded P-ZnO NRs: (a) short-circuit current density, Jsc, (b) open-circuit voltage, Voc, (c) fill factor, FF and (d) power conversion efficiency, PCE obtained under AM 1.5 solar illumination.


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Figure 9. (a) EQE spectra of HOPV devices (b) Photoluminescence emission spectra of P3HT on various SML- P-ZnO NRs.


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Figure 10. FE-SEM images of (a) P-ZnO NRs, (b) SML- P-ZnO NRs. (c) EDS microanalysis spectra of SML- P-ZnO NRs (Inset: Elemental maps of Zn (green), O (Red), S (blue), C (white)), (d) corresponding XRD spectra.


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Figure 11. XPS spectra corresponding to (a) Zn 2p, (b) O 1s, (c) S 2p, (d) N 1s.


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Table 1. Photovoltaic Performance Data iOPV Devices with BHJ of P3HT: PC60BM and PTB7Th: PC71BM. Best (average) values from ~50 P3HT: PC60BM cells and ~40 PTB7-Th: PC71BM cells. a

Calculated

from best cells


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Table 2. Average Photovoltaic Performance Data of HOPV Devices with P3HT as a donor (averaged for ~40 cells). a

Calculated

from best cells


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Table of Content (TOC)


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Work-Function and Surface Energy Tunable Cyanoacrylic Acid Small

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