Block Copolymer-Tuned Fullerene Electron Transport Layer Enhances

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A block copolymer–tuned fullerene electron transport layer enhances the efficiency of perovskite photovoltaics His-Kuei Lin, Yu-Wei Su, Hsiu-Cheng Chen, Yi-Jiun Huang, and Kung-Hwa Wei ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b07690 • Publication Date (Web): 30 Aug 2016 Downloaded from http://pubs.acs.org on August 31, 2016

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ACS Applied Materials & Interfaces

A block copolymer–tuned fullerene electron transport layer enhances the efficiency of perovskite photovoltaics Hsi-Kuei Lin, Yu-Wei Su, Hsiu-Cheng Chen, Yi-Jiun Huang, Kung-Hwa Wei＊

Department of Materials Science and Engineering, National Chiao Tung University, 300 Hsinchu, Taiwan KEYWORDS Perovskite, photovoltaics, copolymer, electron transport layer, grazing-incidence small-angle Xray scattering

ACS Paragon Plus Environment

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ABSTRACT

In this study, we enhanced the power conversion efficiency (PCE) of perovskite solar cells by employing an electron transfer layer (ETL) comprising [6,6]phenyl-C61-butyric acid methyl ester (PC61BM) and, to optimize its morphology, a small amount of the block copolymer polystyrene-b-poly(ethylene oxide) (PS-b-PEO), positioned on the perovskite active layer. When incorporating 0.375 wt% PS-b-PEO into PC61BM, the PCE of the perovskite photovoltaic device increased from 9.4 to 13.4%—a relative increase of 43%—because of a large enhancement in the fill factor of the device. To decipher the intricate morphology of the ETL, we used synchrotron grazing-incidence small-angle X-ray scattering for determining the PC61BM cluster size, atomic force microscopy and scanning electron microscopy for probing the surface and transmission electron microscopy for observing the aggregation of PC61BM in the ETL. We found that the interaction between PS-b-PEO and PC61BM resulted in smaller PC61BM clusters that further aggregated into dendritic structures in some domains, a result of the similar polarities of the PS block and PC61BM; this behavior could be used to tune the morphology of the ETL. The optimal PS-b-PEO–mediated PC61BM cluster size in the ETL was 17 nm, a large reduction from 59 nm for the pristine PC61BM layer. This approach of incorporating a small amount of nanostructured block copolymer into a fullerene allowed us to effectively tune of the morphology of the ETL on the perovskite active layer and resulted in enhanced fill factors of the devices and thus their device efficiency.

INTRODUCTION


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Perovskite solar cells are promising devices for new energy technologies with lower materials and fabrication costs relative to traditional silicon solar cells, while retaining relatively high power conversion efficiencies (PCEs). Perovskite, a lead halide, has several attractive features for use in light-harvesting solar cells, including good optical, excitonic, and electrical conductivity properties. In recent years, the PCEs of perovskite solar cells have increased rapidly to greater than 19% 1-3, depending on the types of structures involved. Perovskite solar cells can be divided into conventional and inverted structures, defined in terms of the layered structures that determine the direction of flow of the electrons and holes in the devices. Conventional perovskite solar cells usually provide higher PCEs, but require high-temperature sintering of the mesoscopic metal oxide (e.g., Al2O3, TiO2, ZrO2) used as the electron transport layer (ETL) beneath the perovskite active layer. Therefore, it can be more appealing to use inverted-type perovskite solar cells because they can be manufactured more rapidly through low-temperature processes. Most studies reported so far have adopted an approach incorporating [6,6]phenyl-C61butyric acid methyl ester (PC61BM) as the ETL in an inverted perovskite solar cell structure5. Because the ETL facilitates the transport of electrons generated from the perovskite layer to the metal cathode, the morphology of the PC61BM layer should feature minimal defects to facilitate electron–hole recombination at the perovskite–cathode interface4-5. Recently, Xia et al. reported6 the incorporation of oleamide into the PC61BM layer on the perovskite active layer; the resultant enhancement in PCE arose from improved ETL coverage on the active layer and enhanced electron collection efficiency. Block copolymers comprise two (or more) chemically distinct polymer segments (blocks) that were connected through a covalent linkage7, and each block, depending on its molecular weight, has dimensions on the order of 10–100 nm8. The segments of a block copolymer will,


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therefore, have different polarities allowing self-assembly into various nanostructures, including spherical micelles, worm micelles, lamellae, cylinders, and vesicles. The self-organization of block copolymers and the potential applications of their phase transition behaviors have been subjects of strong focuses in previous studies9-12. Block copolymers have also been used as templates for producing, for example, TiO2 photo-anodes and three-dimensional structures in solid lithium batteries

13-14

. One particular block copolymer, polystyrene-b-poly(ethylene oxide)

(PS-b-PEO), has been used as a dispersant, stabilizer, and non-ionic surfactant 15-16. In a previous study, we incorporated a small amount of nanostructured PS-b-PEO into the active layer of small-molecule solar cells to optimize the morphology and, thus, enhance the devices’ PCEs17. The large polarity difference between the PEO (µ = 1.04 D) and PS (µ = 0.13 D) blocks in PS-bPEO resulted in their interacting differently with fullerenes that has a small polarity, thereby allowing effective tuning of the morphology of the ETL. In this study, we incorporated small amounts (0.25, 0.375, and 0.5 wt%) of low-molecularweight PS5k-b-PEO5k that comprises a semi-crystalline PEO block and an amorphous PS block, into PC61BM to form the ETL. We expected the low-polarity PS block of PS5k-b-PEO5k to interact with low-polarity PC61BM to form smaller clusters that would aggregate into dendritic nanostructures that provide more pathways for electron transport, thereby tuning the morphology of this ETL to facilitate transport of electrons. We probed the resulting structures’ absorption spectra, photoluminescence (PL) spectra, electron mobility, and photovoltaic performance. Morphological studies and grazing incidence X-ray scattering patterns provided insight into the different properties obtained when using different concentrations of PS-b-PEO.

EXPERIMENTAL


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Materials. indium tin oxide (ITO)-Coated glass substrates (5 Ω cm–2) was obtained from Merck. The hole transport layer was polyethylenedioxythiophene:polystyrenesulfonate (PEDOT:PSS, CleviosTM P VP AI 4083). The perovskite active layer was made from Ossila precursor I201; this ink contained methyl ammonium iodide (MAI), PbI2, and PbCl2 at 40 wt% in anhydrous N,N-dimethylformamide (DMF) at a stoichiometry of 4:1:1. PC61BM was obtained from FEM Technology. The PS-b-PEO diblock copolymer that was obtained from Polymer Source had PS and PEO blocks molecular weights of 5000 and 5000 g mol–1, respectively. PC61BM (20 mg mL–1) in chlorobenzene (CB) and doping with PS-b-PEO at certain ratios that were stirred continuously in a glove box for 12 h at 85 °C were prepared as the ETL precursor solutions. Device Fabrication. Different solvents such as detergent, water, acetone, and isopropyl alcohol (ultra sonication: 20 min each) were used stepwise to clean the patterned ITO glass substrates. The cleaned substrates were dried in an oven for 1 h; prior to use, the substrates were additionally treated with UV ozone for 15 min. Subsequently, PEDOT:PSS was spin-coated (4000 rpm, 40 s) onto the ITO substrates. The substrates were transferred to a N2-filled glove box after they were subject to baking at 150 °C for 15 min under N2. The perovskite ink was spincoated (4500 rpm, 30 s) onto the ITO/PEDOT:PSS surfaces. After drying for more than 15 min, the film was annealed at 85 °C for 50 min. The precursor solutions of the ETL were filtered through a PTFE filter (0.2 µm) and then spin-coated (1000 rpm, 30 s) onto the ITO/PEDOT:PSS/perovskite surfaces. Device fabrication was completed after thermal deposition of a 100-nm-thick film of Ag as the cathode under vacuum of ca. 10–7 torr. During the thermal evaporation process, a shadow mask was used to define a device area of 0.1 cm2.


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Device Characterization. Using a Keithley 2400 source meter, we measured the current density–voltage (J–V) characteristics of the devices and the photocurrent of the devices under simulated AM 1.5 G illumination (100 mW cm–2) that was generated with a Xe lamp–based 150W solar simulator (Newport 66902). We employed a calibrated Si photodiode with a KG-5 filter to verify the illumination intensity during the measurements and calculated the spectral mismatch factor, which can be obtained by comparing the solar simulator spectrum with the AM1.5G (ASTM G173) spectrum. We measured EQEs with an EQE-D-3011 system (Enlitech, Taiwan). A calibrated mono-silicon diode that exhibits a response at 300–800 nm was used as a reference. Sample films were prepared by spin-coating of PEDOT:PSS/perovskite/PC61BM:PS-b-PEO structures onto either 4-cm2 quartz (for UV–Vis spectroscopy) or a silicon wafer (for grazingincidence wide-angle X-ray scattering (GIWAXS) and grazing-incidence small-angle X-ray scattering (GISAXS)). The crystal structures of the samples were characterized by roomtemperature Cu-radiation powder X-ray diffraction (XRD) using a Bruker D8 diffractometer. We used a fluorescence spectrophotometer, Hitachi F-7000, to record the steady-state PL spectra under ambient conditions in air and adopted a spectrophotometer equipped with an integrating sphere, Hitachi U-4100, to acquire UV–Vs absorbance spectra. We determined film morphologies through atomic force microscopy (AFM) (Veeco Innova) with tapping mode and scanning electron microscopy (SEM) (Hitachi SU-8010) operated at 20 kV. We used Transmission Electron Microscopy (TEM), JEOL-2010, to record images at beam energy of 200 keV. We conducted synchrotron GIWAXS and GISAXS analysis [X-ray beam energy: 10 keV (λ = 1.24 Å); incident angle: 0.15°] at the BL23A SWAXS beam line in the National Synchrotron Radiation Research Center, Hsinchu, Taiwan.


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Single-carrier Space-charge Limited Conduction (SCLC) Mobility Measurement. We fabricated electron-only devices having the structure ITO/ZnO/PC61BM:PS-b-PEO/Ag (100 nm) and measured the dark J–V characteristics of these electron-only devices, using the Mott–Gurney

law described SCLC model. The Mott–Gurney law is given as J = , in which J is the

current density, Ɛ0 Ɛr is the dielectric permittivity of the layer (PC61BM:PS-b-PEO); µ is the zerofield mobility; L is the thickness of the layer; V is the internal voltage in the device, and V = Vappl – Vbi – Vrs Where Vappl is the applied voltage, Vbi is the built-in voltage resulted from the difference in work functions between the two electrodes, and Vrs is the voltage drop resulting from the difference in work functions. RESULTS AND DISCUSSION Figure 1(a) displays the molecular structures of PC61BM and PS-b-PEO as well as the planar-heterojunction structure of the photovoltaic devices having the configuration ITO/PEDOT:PSS/CH3NH3PbI3–xClx/PC61BM:PS-b-PEO/Ag.

We

chose

CH3NH3PbI3–xClx

perovskite as the active layer because it is more stable than CH3NH3PbI3 and exhibits a substantially longer carrier diffusion length (ca. 100 nm)18-19. We applied focus ion beam technique with two cuts to prepare a thin (100nm) slice of the device for observing the crosssectional morphologies and interfacial layer structures using TEM. Figure 1(b) presents a crosssectional TEM image of an actual device; the layer thicknesses of PEDOT:PSS, CH3NH3PbI3– xClx,

PC61BM:PS-b-PEO, and Ag were 20, 250, 75, and 100 nm, respectively. Figure 2(a) displays the PL spectra of films of CH3NH3PbI3–xClx, CH3NH3PbI3–

xClx/PC61BM,

and CH3NH3PbI3–xClx/PC61BM:PS-b-PEO. Hereafter, we use “PVK” to represent

CH3NH3PbI3–xClx perovskite. The ETL in this structure can extract free carriers and dissociate


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excitons into free charges. PL spectroscopy can provide a good indication of the extent of recombination of electrons and holes and, thus, the probability of having free carriers. We observed a strong emission peak at 765 nm for the neat PVK film, with a large decrease in the intensity of this signal after depositing the PC61BM layer on the perovskite layer, suggesting that a relatively large amount of free carriers—electrons—were extracted from the perovskite active layer rather than recombining with the holes. The incorporation of 0.25 wt% PS-b-PEO into the PC61BM layer on the PVK layer led to further quenching of the PL peak intensity, relative to that of the neat PC61BM layer on the PVK layer. Thus, incorporation of PS-b-PEO should have resulted in a distribution of PC61BM aggregates that might provide more effective charge transfer and less charge trapping in the ETL20-21. A reduction in the amount of traps in the ETL would presumably reduce the probability of interfacial recombination of carriers, resulting in an enhancement in the fill factor (FF) characterizing the performance of the solar cell device22-23. When the PS-b-PEO concentration increased to 0.375 wt%, the PL peak intensity reached a minimum; it increased slightly in the case when using 0.5 wt% PS-b-PEO. Although the incorporation of PS-b-PEO can induce a better morphology of PC61BM film for electron transport, PS-b-PEO itself is not a good conductor for charge carriers. When the amount of incorporated PS-b-PEO exceeding its optimum concentration in the PC61BM film, the mobility of the charge carriers was influenced more by the excess PS-b-PEO than by the tuned PC61BM clusters in the film; as a result the mobility of carriers decreases and in turn produces an increase in the trap-assisted recombination of electrons and holes—the cause of photoluminescence— in the film. Consequently, the PL intensity of the film decreased with the increasing concentration of incorporated PS-b-PEO up to 0.375wt% (the optimum value) but increased when the concentration of PS-b-PEO reached 0.5wt%. Table S1 shows that the electron mobility for the


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film with a maximum value for the case of 0.375wt% PS-b-PEO, providing strong support for our speculation. Figure 2(b) displays UV–Vis absorption spectra of the PVK film and PVK/PC61BM films incorporating various concentrations of PS-b-PEO. The PVK film absorbed light strongly over the whole visible range (300–550 nm)24, and an absence of adsorption peak at 500 nm for the PC61BM:PS-b-PEO films (see Figure S1). We have re-prepared all the PC61BM:PS-b-PEO (0, 0.25, 0.375, 0.5wt%) film samples with calibrated film thickness of about 300 nm for UV-Vis measurement and found that their UV-Vis absorption intensity is almost identical in the range from 350 to 800 nm. We used the SCLC model to determine the electron mobility (µ) with the devices having the structure ITO/ZnO/PC61BM:PS-b-PEO/Ca/Al

25-26

. The SCLC plots show that the out-of-plane

direction electron mobility increased slightly with the PS-b-PEO content: from 2.75 × 10–4 m2 V– s for the device incorporating a layer of neat PC61BM to 4.03 × 10–4 and 3.75 × 10–4 m2 V–1 s–

1 –1

1

for devices incorporating layers of PC61BM containing 0.375 and 0.5 wt% PS-b-PEO,

respectively (see Figure S1 and Table S1, Supporting Information). Figures 3(a) and 3(b) exhibit the planar perovskite solar cells’ current density–voltage (J–V) characteristics under 1 sun (simulated AM 1.5 G irradiation, 100 mW cm–2) as well as their external quantum efficiency (EQE) curves along with the integrated short-circuit current (Jsc), respectively. Table 1 presents statistical plots of the data of the perovskite devices (10 devices), including their open-circuit voltages (Voc), values of Jsc, FFs, and PCEs. These J–V curves reveal that the value of Voc increased from 0.98 V for the device incorporating the neat PCBM layer to 1.02, 1.01, and 1.00 V for the devices featuring the 0.25, 0.375, and 0.5 wt% PS-b-PEO– incorporated PC61BM layers, respectively. The device containing the ETL incorporating 0.25


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wt% PS-b-PEO had the lowest value 16.56 mA cm–2 of Jsc; it increased to 18.19 and 17.91 mA cm–2 for the devices incorporating 0.375 and 0.5 wt% PS-b-PEO, respectively, in their ETLs, from 16.56 mA cm–2 for the case of neat PC61BM as the ETL. These results are consistent with the trend of the hole mobilities of these materials. Thus, incorporating a small amount of PS-bPEO improved the electron transport ability, but it required an optimal concentration because PSb-PEO itself is not a good electron transport material (Figure S2, Supporting Information). In these devices incorporating different concentrations of PS-b-PEO, the trend in the PCEs followed that of the FFs; therefore, the main reason for the increased PCE was because the FF of the device increased, the result of PS-b-PEO tuning the PC61BM ETL. In Figure 3(b) we observe that devices featuring PC61BM:PS-b-PEO as the ETL achieved a maximum EQE of 79% at 540 nm; the control device (neat PC61BM) provided a slightly lower EQE of 76% at 540 nm. The control device exhibited a best PCE of 9.3%; the device featuring an ETL of 0.375 wt% PS-b-PEO–incorporated PC61BM exhibited the best PCE of 13.4%, an increase of 43% relative to that of the control device (see Figure S3 for histogram). The main contribution to the significantly improved PCE for the device containing 0.375 wt% PS-b-PEO– incorporated PC61BM was the increase in FF to 69.7%, up from 53.4% for the device based on neat PC61BM. We attribute this enhancement in FF to the improved PVK–ETL interfaces in the presence of PS-b-PEO in PC61BM that can enhance the contact interfacial area with electrodes2728

, thereby facilitating the transport of electrons and decreasing the accumulation of charges at

the electrode interface. Typically, the series resistance (Rs), shunt resistance (Rsh), and electron mobility are all related to the FF29-30. It was well established that the value of Rs of a device expresses the integral conductivity of the device directly related to its internal carrier mobility, whereas the value of Rsh refers to the loss of photocurrent through carrier recombination within


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the device, particularly at the interfaces of each layer

31-32

. The values of Rs and Rsh can be

measured from the slopes of the J–V curves near the values of Voc and Jsc, respectively33. The value of Rs of the ETL featuring 0.375 wt% PS-b-PEO incorporated in PC61BM was 4.45 Ω cm2, lower than the value of 9.68 Ω cm2 in the control device (neat PC61BM). The device featuring 0.375 wt% PS-b-PEO had the highest value of Rsh (34,600 Ω cm2), suggesting that the power loss in the solar cell through an alternate current path was very small, resulting in the high FF. Moreover, the electron mobility (Table S1) was highest for the ETL containing 0.375 wt% PS-bPEO, implying that the interface between this ETL and Ag facilitated electron transport. The interface between the PVK and the ETL plays a key role affecting device performance. Figures 4(a–d) present top-view SEM images of PC61BM containing PS-b-PEO at concentrations of 0, 0.25, 0.375, and 0.5 wt%, respectively. The PC61BM:PS-b-PEO films [Figures 4(b–d)] had more compact structures than the film prepared without PS-b-PEO [Figure 4(a)]. The large difference in polarity of the PS and PEO blocks influenced the specific aggregation of PC61BM, which aggregated along the PS blocks in solution during the process of drying the ETL

34-37

. When PS-b-PEO was incorporated, the PC61BM clusters were distributed

well and the aggregation size became more uniform. The phase separation of the PC61BM became clearer upon increasing the content of PS-b-PEO, resulting in the formation of large grains of PC61BM and a high aggregation size [Figures 4(c) and 4(d)]

38

. Excess phase

separation generated voids that destroyed the electron transport paths, thereby decreasing the value of Jsc and the FF. A nanoscale phase separation morphology for the interpenetrating network would be most suitable for optimal contact of the perovskite or Ag, with the best devices containing 0.375 and 0.5 wt% PS-b-PEO giving FFs of greater than 70% 40. In addition to decreased charge recombination in the perovskite layers, due to their highly crystalline nature


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and passivation by PC61BM, the compact and leakage-free perovskite active layers formed by the inter-diffusion process should also have contributed to the large FF

39

. Thus, PS-b-PEO played

an important role as an agent with properties intermediate between those of self-assembling surfactants and aggregating PC61BM. Nevertheless, an excessive amount of PS-b-PEO in the ETL made the film uneven because the spherical pores became too large 40. We used AFM and TEM to gather more details about the structures of the ETLs. Figures 5(a–d) display the surface textures of PC61BM layers prepared with and without PS-b-PEO, recorded in AFM tapping mode. The root-mean-square (RMS) roughness of the PC61BM layer decreased slightly, from 8.8 nm in the absence of PS-b-PEO to 8.5–8.6 nm in the presence of various amounts of PS-b-PEO; thus, the incorporation of PS-b-PEO resulted in smoother films. Figures 5(e–h) present bright-field TEM images of the films that were prepared using the same conditions as those that have been applied for AFM characterization. The TEM images reveal that in the presence of PS-b-PEO in some parts of the films, fullerenes were induced to form two nanostructured morphologies: spherical and dendritic. When doping initially with PS-b-PEO, the structure in the TEM image appeared more uniform and contained some dendritic (close to lamellar) structures and spherical cells. These dendritic structures caused PC61BM clusters to aggregate in that specific region, with the dispersion of PC61BM becoming more even than that of PC61BM in the absence of PS-b-PEO

41-42

. Nevertheless, PS-b-PEO formed more of the

spherical cells as we increased its content. These spherical cells did not favor the dispersion of PC61BM, causing it to aggregate around them, resulting in decreased electron mobility 43. Typically, it is beneficial for electron transfer across the interfacial layers to occur when there is a close contact (large interfacial area) between the layers. The topology of the ETL layer are smoother for the PS-b-PEO incorporated PC61BM film than that for the pristine PC61BM film,


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as indicated by the smaller value of RMS. As a result, the better contact between the ETL and Ag electrode facilitates the electron transport across the interface and thus has higher electron extraction efficiency. Based on the mass balance principle, larger PC61BM clusters leads to fewer of PC61BM clusters in the film, and the electron transport process that takes place along these clusters and depends on the number of the clusters will be less efficient because of more trappings. The micellar structure of PS-b-PEO in PC61BM can effectively provide a better dispersion of PC61BM, resulting in the enhancement of the electron transport. Figure S4 displays the X-ray diffraction (XRD) spectra of all of our samples. Signals were clearly evident at 5–10°; they did not change after depositing the ETL on top of the standard substrate. In the range 5–10°, however, a new peak appeared after we had added a small amount of PS-b-PEO into PC61BM; this peak characterized the crystallization of the PEO phase of PS-bPEO. Because PS-b-PEO can self-assemble into ordered structures, the peak of the PS-b-PEO– incorporated PC61BM layer was different from that of the pristine PC61BM layer 44. Figure 6 presents GIWAXS patterns of PC61BM on perovskite films in the presence of different amounts of PS-b-PEO. GIWAXS is sensitive to the crystalline parts of a structure, allowing determination of the crystal structure and orientation of the crystalline regions in thinfilm organic materials

45-46

. The orientation of a crystal affects the solar cell efficiency;

GIWAXS can be used to define the orientation of PS-b-PEO crystals

47

. We measured the

GIWAXS pattern of the pure perovskite film to identify its peaks from those of PC61BM:PS-bPEO (see Figure S4, Supporting Information). The PC61BM films containing PS-b-PEO featured peaks in low-q region that were not evident for the PC61BM film prepared without dopant; this result is consistent with Figure S4. We have observed that there is no actual PC61BM crystallites


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owing to the absence of diffraction peak in 2-D GIWAXS patterns, but PC61BM usually has a weak halo at the q=1.4 Å-1, as reported in our previous paper48. Figures 7(a–d) display the 2-D GISAXS images of PC61BM:x% PS-PEO films (x = 0, 0.25, 0.375, 0.50) spin-cast onto PVK films on a silicon wafer substrate. Figure 7(e) reveals that the 1D profiles of the PC61BM:x% PS-PEO films (x = 0, 0.25, 0.375, 0.50) decreased along the inplane direction (i.e., Yoneda peak), qx, of the 2-D GISAXS images. These 1-D GISAXS profiles were deducted from the background of PVK/wafer. The profile of the neat PC61BM film did not feature the shoulder that represents the PC61BM cluster size. Thus, the pristine PC61BM film most likely featured very large PC61BM clusters, beyond the detection limit of the q range of the GISAXS measurement. The GISAXS profiles of the PC61BM:PS-b-PEO blended films prepared at three different doping concentrations (0.25, 0.375, and 0.5 wt%) featured almost identical and significant shoulders in the q region from 0.01 to 0.04 Å–1. We determined the PC61BM cluster sizes in these PC61BM:PS-b-PEO blend films by fitting the GISAXS I(q) profiles using Eq. (1), comprising a power-law model (first term) and polydisperse spheres having the assumed Schulz size distribution (second term) 49, as given by 2

I (q) = Aq

−m

2 2 ∞  4π  +   N0 ( ∆ρ ) ∫ f ( R ) R6  F ( qR )  dR 0  3 

 z +1  f (r ) =    Ra 

z +1

  z +1  1 R z exp  −   R R Γ z + 1) (   a  

eq. (1)

eq. (2)

where m is the power law of a 1-D GISAXS profile in low-q region, Ra is the mean radius of clusters. The Schulz distribution f(r) is defined in Eq. (2) to describe the probability of the cluster’s radius size. The parameter z is related to the polydispersity (p = 0–1) by the = −


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1, where = √ / , and is the variance of distribution. The scattering amplitude for a sphere is defined in Eq. (3): 3  sin ( qR ) − qR cos ( qR )  F ( qR ) =  3 ( qR )

eq. (3)

The GISAXS profiles of the blended films containing 0, 0.25, 0.375, and 0.5 wt% of PSPEO displayed the behavior of power-law scattering [I(q) ∝ q–m; 1 ≤ m ≤ 3] in the q range from 0.002 to 0.01 Å–1. The model fitting is was, therefore, an approximation approach. The fitted values is 18.3, 17.0, and 20.4 nm for the mean PC61BM cluster diameters in the PC61BM: x% PSb-PEO (x = 0.25, 0.375, and 0.5) blend films with x of 0.25, 0.375, and 0.5 were 18.3, 17.0, and 20.4 nm, respectively. The average diameter for the pristine PC61BM (without PS-b-PEO) clusters is calculated to about 59 nm, which is much larger than the sizes (17-20 nm) for the PSb-PEO incorporated PC61BM films (see Supporting Information, Table S2). Correlating to the device data, we found that the ETL incorporating 0.375% PS-b-PEO featured the smallest PC61BM clusters size and the highest volume fraction (indicated by the remarkable shoulder of in the GISAXS profile) dispersed in the blend layer, which thereby provides providing more paths between clusters for electron transporting to the electrode and results in the highest PCE of 13.4%. Based on the mass balance principle, the morphology comprised of larger PC61BM clusters has fewer clusters and smaller and more PC61BM clusters that result in less trappings or more pathways for electron transporting to the electrode, leading to improved fill factor of the devices. CONCLUSION We have developed a new ETL layer comprising a small amount of a block copolymer PSb-PEO incorporated within PC61BM. By taking advantage of the large polarity disparity between


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the PEO and PS blocks of the PS-b-PEO copolymer, we are able to differentiate their interactions with PC61BM that led to the formation of clusters that further aggregated into large domains, tuning the phase separation. Through systematic characterization of these ETL films using GISAXS, SEM, AFM, and TEM, we found that the electronic, morphological, and inner structural properties of the PS-b-PEO–incorporated PC61BM layers were superior to those of the pristine PC61BM film, making them suitable for use as ETLs for inverted perovskite solar cells. The devices prepared with PS-b-PEO exhibited significant enhancements in FF, especially when the PC61BM layer incorporated 0.375 wt% PS-b-PEO; this device provided a PCE of 13.4%, significantly greater than that (9.4%) for the control devices prepared without PS-b-PEO. Consequently, incorporating a small amount of this large-polarity-disparity diblock copolymer with PC61BM in the ETL allowed us to effectively tune of the ETL morphology and resulted in high fill factor and thus enhanced device efficiency.

ASSOCIATED CONTENT Supporting Information. Current density–electric field log–log plots for electron-only diodes; electron mobility calculation table; XRD profiles. This information is available free of charge on the ACS Publication website. AUTHOR INFORMATION ＊E-mail: [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENT


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Figure 1. (a) Schematic representation of the structure of the perovskite cells; molecular structures of PC61BM and PS-b-PEO, the ETL materials. (b) Cross-sectional TEM image of a perovskite solar cell.


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Figure 2. (a) PL and (b) UV–Vis absorption spectra of films processed using perovskite and perovskite/PC61BM incorporating PS-b-PEO at various contents (0, 0.25, 0.375, 0.5 wt%).


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Figure 3. (a) J–V characteristics and (b) EQE spectra of perovskite devices incorporating PS-bPEO at various contents (0, 0.25, 0.375, 0.5 wt%) in the ETL.


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Figure 4. Top-view SEM images of (a) the perovskite/PC61BM film, (b) the perovskite/PC61BM:PS-b-PEO (0.25 wt%) film, (c) the perovskite/PC61BM:PS-b-PEO (0.375 wt%) film, and (d) the perovskite/PC61BM:PS-b-PEO (0.5 wt%) film.


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Figure 5. (a–d) AFM topographic and (e–h) TEM images of PC61BM films incorporating PS-bPEO at various contents on the perovskite layer; (a, e) 0; (b, f) 0.25; (c, g) 0.375; (d, h) 0.5 wt%.


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Figure 6. 2D GIWAXS patterns of perovskite/PC61BM films incorporating PS-b-PEO: (a) 0, (b) 0.25, (c) 0.375, and (d) 0.5 wt%.


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Figure 7. (a–d) 2D GISAXS patterns of PC61BM films containing (a) 0, (b) 0.25, (c) 0.375, and (d) 0.5 wt% PS-b-PEO. Green lines represent the integral areas. (e) GISAXS profiles of perovskite/PC61BM films containing various weight fractions of PS-b-PEO. All profiles were fitted using calculated model intensities (solid lines)


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Table 1. Device performance of perovskite solar cells. PS-b-PEO Concentration

Voc a) [V]

0 wt%

0.98±0.01

17.16±0.87 53.4±2.2 (58.0)

0.25 wt%

1.02±0.01

0.375 wt% 0.5 wt%

Jsc b) [mA cm–2]

FF c) (FFbd)) [%]

ηe) (ηbf)) [%]

Rshg) Rsh) (kΩ cm2) (Ω cm2) 0.39

9.68

16.56±0.63 65.1±1.9 (63.9) 11.03±0.41 (11.7)

0.92

8.77

1.01±0.01

18.19±0.48 69.7±3.3 (73.5) 12.74 ±0.42 (13.4)

3.46

4.45

1.00±0.01

17.91±0.79 67.4±3.2 (72.8) 12.08±0.25 (12.5)

2.18

4.20

9.05±0.3 (9.4)

a)

Voc: Open-circuit voltage; b) Jsc: short-circuit current density; c) FF: fill factor; d) FFb: best fill factor ; e) η: power conversion efficiency; f) ηb: best power conversion efficiency; g) Rsh: shunt resistance; h) Rs: series resistance.


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TOC Figure


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Block Copolymer-Tuned Fullerene Electron Transport Layer Enhances

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