Well-Defined Nanostructured, SingleCrystalline TiO2 Electron Transport Layer for Efficient Planar Perovskite Solar Cells Jongmin Choi,†,§ Seulki Song,†,§ Maximilian T. Hörantner,‡ Henry J. Snaith,‡ and Taiho Park*,† †
Department of Chemical Engineering, Pohang University of Science and Technology (POSTECH), San 31, Nam-gu, Pohang, Kyungbuk 790-784, Korea ‡ Department of Physics, Clarendon Laboratory, University of Oxford, Parks Road, Oxford OX1 3PU, United Kingdom S Supporting Information *
ABSTRACT: An electron transporting layer (ETL) plays an important role in extracting electrons from a perovskite layer and blocking recombination between electrons in the fluorine-doped tin oxide (FTO) and holes in the perovskite layers, especially in planar perovskite solar cells. Dense TiO2 ETLs prepared by a solutionprocessed spin-coating method (S-TiO2) are mainly used in devices due to their ease of fabrication. Herein, we found that fatal morphological defects at the S-TiO2 interface due to a rough FTO surface, including an irregular film thickness, discontinuous areas, and poor physical contact between the S-TiO2 and the FTO layers, were inevitable and lowered the charge transport properties through the planar perovskite solar cells. The effects of the morphological defects were mitigated in this work using a TiO2 ETL produced from sputtering and anodization. This method produced a well-defined nanostructured TiO2 ETL with an excellent transmittance, single-crystalline properties, a uniform film thickness, a large effective area, and defect-free physical contact with a rough substrate that provided outstanding electron extraction and hole blocking in a planar perovskite solar cell. In planar perovskite devices, anodized TiO2 ETL (A-TiO2) increased the power conversion efficiency by 22% (from 12.5 to 15.2%), and the stabilized maximum power output efficiency increased by 44% (from 8.9 to 12.8%) compared with S-TiO2. This work highlights the importance of the ETL geometry for maximizing device performance and provides insights into achieving ideal ETL morphologies that remedy the drawbacks observed in conventional spin-coated ETLs. KEYWORDS: perovskite, solar cell, electron transport layer, TiO2, nanostructures, anodization
S
maximum power conversion efficiency that exceeded 20% in solar cells employing a TiO2 mesoscopic scaffold layer and the formamidinium lead iodide (FAPbI3) perovskites.8 Devices prepared without a mesoscopic scaffold layer (e.g., a planar perovskite cell) have attracted significant attention for their simple device architecture and fabrication procedures. The PCEs of these devices, however, are around 12%, as reported by Snaith and Yang, independently,9,10 apparently lower than the values obtained from devices with a TiO2 mesoscopic scaffold layer. Efforts have been applied toward improving the PCE and stability of the devices by developing advanced materials (perovskites11−18 and HTMs19−27) while also examining the roles of the components and their operating principles.28−32
olid-state organic−inorganic hybrid perovskite solar cells are one of the most important recent discoveries in the field of photovoltaics.1−4 The first descriptions of solar cells that employed methylammonium lead iodide (MAPbI3) perovskite nanocrystals as the light-absorbing material in a liquid electrolyte, such as an iodine/triiodide redox couple, were reported by Miyasaka et al.5 Since those studies, Snaith6 and Park7 independently employed 2,2′,7,7′-tetrakis(N,N′-di-pmethoxyphenylamine)-9,9′-spirobifluorene (spiro-OMeTAD) as an organic hole transporting material (HTM) in place of the liquid electrolyte, which improved the power conversion efficiency (PCE, with an increase from 3.8 to 9.7−10.9%) and stability. In general, perovskite solar cells are composed of a counter electrode/HTM/perovskite/electron transport layer (ETL) on a fluorine-doped tin oxide (FTO)-coated glass. The active surface area may be increased by forming a perovskite layer on a mesoscopic scaffold layer, such as a layer composed of Al2O3 or TiO2 nanoparticles. Recently, Soek reported a © 2016 American Chemical Society
Received: March 4, 2016 Accepted: May 16, 2016 Published: May 16, 2016 6029
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Figure 1. High-resolution scanning electron microscopy images of (a−c) top and the corresponding (d−f) cross sections of S-TiO2 (the number indicates the average thickness of the S-TiO2 layer on the nanometer scale) on FTO glass. The inset in (a) shows an SEM image of bare FTO glass. The red circle shows a pinhole in the S-TiO2 layer. The light red regions of the cross-sectional SEM images indicate the STiO2 layer.
the FTO glass, although numerous discontinuous S-TiO2 areas (marked as pinholes in Figure 1a) appeared. The corresponding vertical image (Figure 1d) revealed that the bare FTO surface was exposed and could form direct contact with the perovskite layer. Pinholes were observed in the 40 nm thick STiO2 layer (Figure 1b), and they were absent from the ca. 80 nm thick S-TiO2 layer (Figure 1c). The surface smoothness increased with the S-TiO2 layer thickness on the rough FTO glass, as determined by atomic force microscopy (AFM) imaging (Figure S1a,b). The root mean square (rms) roughness of the S-TiO2 sample was clearly lower than that of the bare FTO glass due to overfilling of the valleys on the FTO surface with S-TiO2, as is evident in the vertical images (Figure 1e,f). For instance, the 80 nm thick S-TiO2 layer displayed a vertically nonuniform morphology at thickness values of 45−160 nm. The irregular thick S-TiO2 layer (and the resulting smooth surface) or the presence of pinholes may have significantly impacted the device performance. For instance, if the S-TiO2 were too thin or the perovskite layer directly contacted the FTO layer, electrons in the FTO layer easily recombined with holes due to the lack of hole blocking (② and ③ in Figure 2a). On the other hand, the electron flow from the perovskite layer to the thick S-TiO2 region was ineffective due to the high series resistance, and the electrons recombined with the holes (④ in Figure 2a). The resulting smooth surface decreased the interfacial surface area between the perovskite and S-TiO2 layers, resulting in a low device performance, especially in the planar perovskite solar cells. The irregular ETL morphology significantly influenced the device performance, as observed in the photocurrent−photovoltage (J−V) curves of the devices prepared with various STiO2 ETL thicknesses (Figure 2 and Table 1). First, a device without an ETL was prepared such that the perovskite layer directly contacted the FTO surface. This structure was analogous to a structure with a gigantic pinhole. The PCE was 2.0%, and the photovoltaic parameters were poor (the short-circuit current density (JSC) was 14.0 mA/cm2; the opencircuit voltage (VOC) was 0.49 V, and the fill factor (FF) was 0.29). In addition to inefficient electron extraction, the low FF resulted mainly from recombination between the photoinduced electrons and holes (③ in Figure 2a), which again decreased
However, the importance of the ETL has not yet been extensively explored. The ETL plays an important role in extracting electrons from the perovskite layer and blocking recombination between electrons in the FTO and holes in the perovskite layer. An excessively thick ETL can minimize recombination, although the electron flow may be hampered due to the high series resistance. A few solar cell structures have been prepared without an ETL, although the resulting PCEs tend to be low compared to planar perovskite cells prepared with an ETL.33,34 Recently, several research groups have developed new ETL materials, including graphene/TiO 2 nanocomposites, 35 ZnO,36,37 and yttrium-doped TiO2 (Y-TiO2),38 that provide efficient electron transport and permit room temperature device fabrication. Dense TiO2 ETLs have not been extensively investigated, although they are the ETLs most widely used in the field. Titanium isopropoxide (TTIP) solutions spin-coated onto FTO (denoted to S-TiO2) provide stable characteristics under harsh chemical conditions and excellent electrical properties.39−42 Here, we describe the relationship between the morphology of a conventional S-TiO2 layer on FTO and the electron transport properties. The spin-coating process formed a highly irregular thick S-TiO2 layer on top of the rough FTO layer, resulting in the reduced interfacial area at the perovskite/STiO2/FTO layer interface and the poor electron extraction. These problems were addressed by implementing a TiO2 ETL based on a simple sputtering and anodization method (denoted A-TiO2). The resulting A-TiO2 ETL yielded well-defined nanostructured morphology with excellent physical properties, including single-crystalline properties, the absence of pinholes, a uniform film thickness, enhanced transmittance, and strong connections between the FTO and ETL.
RESULTS AND DISCUSSION Figure 1 shows high-resolution scanning electron microscopy (HR-SEM) top and cross-sectional images of the S-TiO2 layer prepared with different thicknesses (20−80 nm) on the FTO substrate. The rough FTO glass (see the inset of Figure 1a) became smooth as the thickness of the S-TiO2 layer increased. The 20 nm thick S-TiO2 layer retained the rough top surface of 6030
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12.5% for the 40 nm thick S-TiO2 layer. These values were equivalent to those reported independently by Snaith and Yang in studies of planar devices.9,10 The FF values increased to 0.45 and 0.67 for 20 and 40 nm thick S-TiO2 layers, respectively, indicating that electron extraction properly occurred, and direct contact between the perovskite layer and the FTO surface was minimized. Further increases in the S-TiO2 ETL thickness could increase the length of the electron transport pathway. The electrons were trapped in the thick S-TiO2 ETL and recombined with holes (④ in Figure 2a). The VOC and FF values in the device prepared with 60 nm thick S-TiO2 layers were 0.98 V and 0.57, smaller than those (1.03 V and 0.67) obtained from the 40 nm thick S-TiO2 layers. These values again decreased as the S-TiO2 thickness increased (80 nm), affording VOC and FF values of 0.92 V and 0.26. These results revealed the importance of the S-TiO2 thickness and the absence of direct contact (pinholes) between the perovskite layer and the rough FTO surface. These results also suggested that vertical uniformity in the TiO2 layer, along with complete coverage of the rough FTO surface, could maximize electron extraction and minimize recombination to enhance the photovoltaic performance in planar devices. A TiO2 layer with vertical uniformity and a high surface roughness was achieved on the FTO surface using a proposed method (Figure 3a−d). This process involved three simple steps: (i) sputtering a Ti layer onto the FTO surface, (ii) anodization, and (iii) thermal annealing. The thickness of the Ti layer was controlled by the sputtering time (ca. 10 min for 40 nm thickness). After the Ti deposition, a mild anodization process at 5 V in ethylene glycol solution containing 0.25 wt % NH4F and 0.3 vol % of distilled water converted the opaque Ti layer to a transparent and amorphous TiO2 layer at room temperature for 5 min (see Figure S2 for effects of the anodization time and voltage on the TiO2 surface morphology). The anodizing conditions are described in greater detail in the Supporting Information.44 The amorphous TiO2 layer was thermally annealed at 500 °C for 2 h to produce a crystalline (anatase) phase.45 The A-TiO2 layer completely retained the rough surface of the FTO glass. The thickness of the A-TiO2 layer on the rough surface of the FTO glass was highly uniform, as shown in the surface and vertical SEM images (Figure 3c,d). The AFM analysis confirmed the uniform surface morphologies in the films (Figure S1a,c). The rms roughness value of the A-TiO2
Figure 2. Schematic illustration of the predicted electron collection process in the S-TiO2 for various thicknesses (a), and Photocurrent−photovoltage (J−V) properties of the planar perovskite solar cells prepared using S-TiO2 ETLs with various thicknesses, measured under AM 1.5 solar illumination (b).
Table 1. Photovoltaic Parameters of the Planar Perovskite Solar Cells Prepared Using S-TiO2 ETLs with Various Thicknessesa average thickness (nm)
VOC (V)
JSC (mA/cm2)
FF
η (%)
0 20 40 60 80
0.49 1.04 1.03 0.98 0.92
14.0 18.6 18.1 17.9 16.8
0.29 0.45 0.67 0.57 0.27
2.0 8.7 12.5 10.0 4.2
a Measured under AM 1.5 solar illumination. Cell size: 0.09 cm2. Scan direction: from 1.4 to 0 V (reverse sweep).
VOC and JSC.43 Next, devices were fabricated with 20 or 40 nm thick S-TiO2 ETLs to improve the electron extraction efficiency and minimize direct contact between the perovskite layer and the FTO surface, thereby decreasing the number of pinholes, as shown in Figure 1a,b. The PCE values improved dramatically to
Figure 3. Surface HR-SEM images of bare FTO glass (a), Ti-deposited FTO glass (b), and anodized TiO2 (A-TiO2) on FTO glass (c). (d) Cross-sectional HR-SEM image of A-TiO2 on FTO glass. (e) Transmittance spectra of the bare FTO, A-TiO2, and spin-coated TiO2 glass (STiO2). The thickness of the deposited TiO2 layer was 40 nm. The inset shows photographs of the substrates after each treatment. 6031
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Figure 4. High-resolution transmission electron microscopy images: cross section (a), surface lattice (b), FFT pattern (c), and EELS composite (carbon (red) and Ti (green)) (d) and oxygen (yellow) (e) mapping images of the A-TiO2. Cross section (f), surface lattice (g), FFT pattern (h), and EELS composite (i) and carbon (white) (j) mapping images of the S-TiO2.
film was almost identical to that of the bare FTO. Therefore, high effective contact area between the A-TiO2 and perovskite and efficient electron transport properties are expected at the A-TiO2. The 40 nm thick A-TiO2 layer deposited on the FTO glass exhibited a transparency better than that of the corresponding S-TiO2 layer on the FTO glass and was comparable to the bare FTO glass, as shown in Figure 3e and its inset. Notice that the 60 nm thick A-TiO2 layer displayed a transmittance similar to that of the 40 nm thick STiO2 layer (Figure S3). The outstanding transparency of the ATiO2 layer was ascribed to the uniform film morphology and was expected to improve the light harvesting. X-ray diffraction studies were conducted in an attempt to investigate the crystal structure and the chemical property of the thin A-TiO2 layer; however, these studies failed due to the strong peak intensities of the FTO. High-resolution transmission electron microscopy (HR-TEM) was employed to characterize the crystal morphology and the chemical property of the thin A-TiO2 layer after thermal annealing. The vertical HR-TEM images of the A-TiO2 and S-TiO2 layers agreed with the results obtained from the SEM measurements (Figure S4). Interestingly, close examination of the TiO2 layers revealed that the nanostructures of the A-TiO2 and S-TiO2 layers differed significantly. Unlike the aggregation of the nanoparticles of various sizes, observed in the S-TiO2 layer (Figure 4f), the magnified image of the A-TiO2 layer clearly revealed the formation of vertically aligned and well-connected onedimensional nanorods 10−20 nm in width and 40 nm in length (Figure 4a). The surface HR-TEM images indicated the formation of highly ordered lattice structures (Figure 4b) and randomly oriented (Figure 4g) crystal alignment in the A-TiO2 and S-TiO2 layers, respectively. These properties were confirmed in the fast Fourier transform (FFT) patterns collected from the S-TiO2 and A-TiO2 layers. Single-crystalline anatase patterns were clearly observed in the A-TiO2 layer (Figure 4c; see Figure S5 for more lattice HR-TEM images of A-TiO2). The S-TiO2 layer displayed a polycrystalline anatase structure with randomly oriented crystal axes, as identified from the ring patterns (Figure 4h, inset). More interesting features observed during a deep analysis of the HR-TEM images of the A-TiO2 and S-TiO2 layers indicated the difference of physical contact between the TiO2 layer and the rough FTO surface. Figure 4f shows an empty area within the deep sharp valley
regions of the FTO, indicating the absence of physical contact between the S-TiO2 and the FTO surface. A liquid to solid TiO2 anatase crystal phase transformation and the concomitant volume shrinkage could be induced in the deep sharp valley regions. Composite (carbon and Ti in Figure 4i) and carbon (Figure 4j) mapping images of the S-TiO2 layer after being filled with viscous carbon liquid, obtained from electron energy loss spectroscopy (EELS) analysis, clearly revealed an empty area between the S-TiO2 and the FTO layers (see Figure S6 for more images). Meanwhile, the A-TiO2 layer displayed robustly constructed highly ordered single crystals at the FTO surface, as identified from the absence of carbon (C) atoms in the composite mapping image of the A-TiO2 layer (Figure 4d). These morphologies were confirmed using an oxygen (O) atom mapping image of the A-TiO2 layer on the surface of the FTO glass (Figure 4e). Oxygen atoms in this layer corresponded to TiO2 and FTO units. This image also confirmed that the Ti layer was completely converted to TiO2 during anodization. TiO2 conversion from Ti is also identified by secondary ion mass spectrometry (SIMS), as shown in Figure S7. Compared with Ti depth profile, A-TiO2 exhibits increased 16O signal, and the enhanced 16O intensity of A-TiO2 with increased film thickness (as shown in SIMS depth profiles of 40 and 80 nm ATiO2) supports TiO2 formation. We also confirmed that the surface morphology and the electrical property (2.0−2.1 ohm/ sq by the four-point probe resistivity measurements) of the FTO were not changed at all during the anodization process due to its high chemical resistance (Figure S8).46 Additionally, we also fabricated the electron-only device and confirmed the improved mobility of the A-TiO2 electron-only device (Figure S9, from 3.09 × 10−5 to 1.42 × 10−4 cm2/Vs). In addition to improved light absorption and an increase in the effective contact area between at the A-TiO2 and perovskite layer, as mentioned above, fast electron transport and a long electron lifetime in the single-crystalline anatase TiO247 and defect-free physical contact between the A-TiO2 and FTO surfaces could improve the photovoltaic performances of devices prepared with an A-TiO2 ETL. The advantages of the A-TiO2 ETL in comparison to the STiO2 ETL were made apparent in tests of a planar perovskite solar cell (Table 2). Figure 5 shows the J−V properties of planar perovskite solar cells prepared with A-TiO2 ETLs of 6032
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the defect sites generated due to the increased series resistance48,49 became dominant and increased recombination (④ in Figure 2a), as observed in the device prepared using an 80 nm thick S-TiO2 ETL (Figure 2). The 80 nm thick A-TiO2 exhibited the worst performance (PCE = 8.5%), with a JSC = 18.3 mA/cm2, FF = 0.49, and VOC = 0.95 V. The device prepared with a 20 nm thick A-TiO2 yielded a FF of 0.57, much lower than that obtained in the device prepared with a 40 nm thick A-TiO2 layer (0.70), although the thickness uniformity was excellent (Figure S10). This result indicated that a thickness of 20 nm was insufficient to prevent the recombination reaction (③ in Figure 2a). The photovoltaic performances and spectroscopic characteristics of two devices prepared using S-TiO2 or A-TiO2 were directly compared, as shown in Figure 6. The perovskite layer deposited on different ETLs exhibited similar film morphology (Figure S11). Overall, the A-TiO2 ETL performance was superior to that of the corresponding S-TiO2 ETL (see Figure S12 for a comparison of the 40 nm ETL device reproducibility characteristics). For instance, the device prepared using a 40 nm thick A-TiO2 (PCE = 15.2%) provided a PCE that was 122% of the value obtained from the 40 nm thick S-TiO2 layer (PCE = 12.5%) due to the uniform thickness, high transparency, and increased surface area. The incident photon-toelectricity efficiency values (IPCE) agreed well with the JSC values (Figure 6b). The better rectifying shape of the A-TiO2 device, compared to that of the S-TiO2 device, indicated improved prevention of recombination and better leakage current blocking (Figure 6c). In addition, the A-TiO2 also exhibited little hysteresis, seen in the smaller difference between the back and forward sweeps compared to that of the S-TiO2 (Figure 6a and Table S1), and it is ascribed to superior electron extraction ability of A-TiO2 reducing charge accumulation in the perovskite layer.50,51 A highly stabilized efficiency (12.8%)
Table 2. Photovoltaic Parameters of the Planar Perovskite Solar Cells Prepared Using A-TiO2 ETLs with Various Thicknessesa thickness (nm)
VOC (V)
JSC (mA/cm2)
FF
η (%)
80 60 40 20
0.95 1.08 1.06 1.04
18.3 20.4 20.5 21.1
0.49 0.62 0.70 0.57
8.5 13.7 15.2 12.5
a Measured under AM 1.5 solar illumination. Cell size: 0.09 cm2. Scan direction: from 1.4 to 0 V (reverse sweep).
Figure 5. J−V properties of the planar perovskite solar cells prepared using A-TiO2 ETLs with various thicknesses, measured under AM 1.5 solar illumination.
various thicknesses (20−80 nm). A device employing a 40 nm thick A-TiO2 layer exhibited the best photovoltaic performance, with a PCE of 15.2%, a JSC of 20.5 mA/cm2, a FF of 0.70, and a VOC of 1.06 V. The PCE of 15.2% was apparently greater than the best values (around 12%) in the planar perovskite solar cells, as reported by Snaith and Yang, independently.9,10 As the thickness of the A-TiO2 layer increased (40 → 60 → 80 nm),
Figure 6. Comparison of J−V properties under reverse and forward scans (scan range: from 0 to 1.4 V) with dark current measurement (a), IPCE (b), rectifying current density−voltage characteristics (FTO/ETL/spiro-OMeTAD/Au) (c), and stabilized maximum power output performances of the planar perovskite solar cells prepared using A-TiO2 or S-TiO2 ETLs (d). Comparison of PL (e) and PL decay (f) spectra obtained from perovskite layers on bare FTO, S-TiO2, and A-TiO2. 6033
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continuously stirred to prevent fluctuation of the current. The resulting sample was cleaned using ethanol to remove the acidic salts and debris, followed by thermal annealing at 450 °C for 2 h. Device Fabrication. The planar perovskite solar cells were fabricated on FTO glass. Some parts of FTO were treated with 2 M HCl solution containing zinc powder to make insulator regions. The resulting FTO was washed using detergent, ethanol, acetone, and isopropanol. The spin-coated TiO2 compact layer (S-TiO2) was prepared by the following method.9 A solution of 0.5 mM TTIP in ethanol and 40 mM HCl solution in ethanol were slowly mixed together. The mixed solution was added dropwise onto the substrate and spin-coated at 2000 rpm for 1 min. The S-TiO2-coated substrates were heated at 500 °C for 30 min. The perovskite precursor was prepared following the reported method.53 Perovskite precursor solution was spin-coated on the substrate at 2000 rpm for 60 s and then crystallized by heating at 100 °C for 5 min in the glovebox. The 2,2′,7,7′-tetrakis(N,N′-di-p-methoxyphenylamine)-9,9′-spirobifluorene (spiro-MeOTAD) (Merck KGaA) in chlorobenzene (180 mg/1 mL) solution with Li-bis(trifluoromethanesulfonyl)imide (Li-TFSI) and tBP (tert-butylpyridine) was spin-coated on the perovskite layer. Finally, a 100 nm silver electrode was vacuum-deposited on the spiroMeOTAD overlayer. Characterization. Field emission scanning electron microscopy (Hitachi S 4800) was utilized for morphological and structural characterization. High-resolution transmission electron microscopy analyses were made with a JEM-2200F electron microscope with an accelerating voltage of 200 kV. Photovoltaic performances of the devices were measured under air AM 1.5G illumination of 100 mW/ cm2 (Oriel 1 kW solar simulator), which was calibrated with a KG5 filter-covered silicon photovoltaic solar cell traceable to the national renewable energy laboratory. A mask was used to define the device illumination area of 0.09 cm2 to minimize photocurrent generation from the edge of the electrodes. Secondary ion mass spectrometry measurement was done using the SIMS system (IMS-6f, Physical Electronics USA, Chanhassen, MN). Beam energy was 15 keV using CS+ ions as the primary ions, and raster was 250 μm (depth) and 100 μm (image). Base pressure was always below 2 × 109 Torr. Spacecharge-limited current was also measured in the same condition but in dark conditions. The IPCE spectra were measured under a constant white light bias of approximately 5 mW/cm2 supplied by an array of white-light-emitting diodes using a power source with a monochromator (Zahner GmbH) and a multimeter chopped at approximately 4 Hz. AFM (Dimension 3100, VEECO) was operated in tapping mode to acquire images of the surfaces of films. Transmittance and PL spectra were obtained by employing an OPTIZEN POP UV− vis−near-IR single-beam spectrophotometer and Jasco FP-6500 spectrometer, respectively. Time-resolved PL decay measurements were obtained using a FluoTime 300, PicoQuant GmbH. The samples were photoexcited using a 507 nm laser head (LDH-P-C-510, PicoQuant GmbH) pulsed at 10 MHz substrate samples, with a pulse duration of 117 ps and a fluence of 30 nJ/cm2/pulse. The PL was collected from a high-resolution monochromator and a hybrid photomultiplier detector (PMA Hybrid 40, PicoQuant GmbH). The sheet resistance of the FTO surface was measured with the four-point probe method (Keithley 2400 sourcemeter).
at the A-TiO2 device was obtained (Figure 6d). This value (12.8%) exceeded that obtained from the S-TiO2 device (8.9%, consistent with values reported previously52), which explained the efficient charge collection in the A-TiO2 layer, 84.2% of its first scan value (Figure 6a). Figure 6e shows the photoluminescence (PL) quenching spectra of perovskite layers of identical thickness prepared on different substrates (Figure S13). These results clearly revealed that the PL intensity decreased due to quenching effects if the ETL were placed on the FTO glass. The emission of the perovskite layer was remarkably quenched (81% quenching) by the A-TiO2 layer, as compared to a 62% quenching by the S-TiO2 layer. The PL decays of both ETLs were measured (Figure 6f). The perovskite layers on the bare FTO displayed very long halflives (250.5 ns) that shortened upon use of the TiO2 ETLs. The half-life of the A-TiO2 was 83.2 ns, much shorter than that (146.5 ns) obtained from the S-TiO2 . These results demonstrated how efficiently electrons could be extracted from excited perovskites into the A-TiO2 ETL.
CONCLUSIONS Conventional S-TiO2 ETLs produce highly irregular film morphologies on rough FTO surfaces, and they exert a bad influence on the performances of planar perovskite devices. The morphological problems in the S-TiO2 ETLs were addressed here by an alternative ETL fabrication method based on sputtering and anodization. The resultant A-TiO2 displayed well-defined nanostructure morphology with a uniform film thickness on the rough FTO surface. The unique morphological characteristics of A-TiO2 enhanced the transmittance to improve light absorption and increase the contact area between the perovskite and TiO2 ETL. In addition, the A-TiO2 was single-crystalline in nature and provided defect-free physical contact between the A-TiO2 and FTO surface, which improved the effective electron extraction and hole-blocking properties. In a planar perovskite device, A-TiO2 increased the PCE by 22% (from 12.5 to 15.2%) and increased the stabilized maximum power output efficiency by 44% (from 8.9 to 12.8%) to provide 84.2% of the first scan efficiency. The highly stabilized maximum power output efficiency of A-TiO2 reveals that the A-TiO2 device functioned well under realistic device operation conditions. The good performance of the device prepared using A-TiO2 was attributed to the excellent electron extraction and hole-blocking properties of A-TiO2. Spectroscopy and rectifying analysis results clearly demonstrated the outstanding electron transport capabilities of A-TiO2. These results revealed that the performance of the ETL was predominantly affected by the morphological properties of the ETL. Our ETL fabrication method provides important insights into achieving ideal ETL morphologies that remedy the drawbacks observed in conventional spin-coated ETLs.
ASSOCIATED CONTENT S Supporting Information *
METHODS
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b01575. Atomic force microscopy, transmittance spectra, SIMS, SEM and TEM images of ETLs, supplemental I−V characteristics of devices (PDF)
Preparation of A-TiO2 on FTO Glass. The A-TiO2 was prepared by a potentiostatic anoidzation method. Anodization was carried out using a Ti-deposited fluorine-doped SnO2 (FTO) conducting glass (Pilkington TEC 7, 7 Ω/□) as a working electrode and a carbon plate as a counter electrode. Before Ti deposition, the FTO glass was rinsed with detergent, distilled water, ethanol, and acetone. The Ti layer was deposited on cleaned FTO glass by a radio frequency sputter (AMAT E5500 at National Institute for Nanomaterials Technology (NINT) in POSTECH). Anodization was conducted at 5 V for 5 min in ethylene glycol solution containing 0.25 wt % NH4F and 0.3 vol % of distilled water at room temperature. During anodization, the solvent was
AUTHOR INFORMATION Corresponding Author
*E-mail:
[email protected]. 6034
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J.C. and S.S. contributed equally.
Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MSIP) (Code No. 2015R1A2A1A10054230), Center for Advanced Soft Electronics under the Global Frontier Research Program (Code No. NRF-2012M3A6A5055225), and Nano Material Technology Development Program (2012M3A7B4049989). REFERENCES (1) Heo, J.-H.; Im, S. H.; Noh, J. H.; Mandal, T. N.; Lim, C.-S.; Chang, J. A.; Lee, Y. H.; Kim, H.-J.; Sarkar, A.; Nazeeruddin, M. K.; Grät zel, M.; Seok, S. I. Efficient Inorganic-Organic Hybrid Heterojunction Solar Cells Containing Perovskite Compound and Polymeric Hole Conductors. Nat. Photonics 2013, 7, 486−491. (2) Burschka, J.; Pellet, N.; Moon, S.-J.; Humphry-Baker, R.; Gao, P.; Nazeeruddin, M. K.; Grätzel, M. Sequential Deposition as a Route to High-Performance Perovskite-Sensitized Solar Cells. Nature 2013, 499, 316−319. (3) Cai, B.; Xing, Y. D.; Yang, Z.; Zhang, W.-H.; Qiu, J. S. High Performance Hybrid Solar Cells Sensitized by Organolead Halide Perovskites. Energy Environ. Sci. 2013, 6, 1480−1485. (4) Etgar, L.; Gao, P.; Xue, Z.; Peng, Q.; Chandiran, A. K.; Liu, B.; Nazeeruddin, M. K.; Grätzel, M. Mesoscopic CH3NH3PbI3/TiO2 Heterojunction Solar Cells. J. Am. Chem. Soc. 2012, 134, 17396− 17399. (5) Kojima, A.; Teshima, K.; Shirai, Y.; Miyasaka, T. Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. J. Am. Chem. Soc. 2009, 131, 6050−6051. (6) Lee, M. M.; Teuscher, J.; Miyasaka, T.; Murakami, T. N.; Snaith, H. J. Efficient Hybrid Solar Cells Based on Meso-Superstructured Organometal Halide Perovskites. Science 2012, 338, 643−647. (7) Kim, H.-S.; Lee, C.-R.; Im, J.-H.; Lee, K.-B.; Moehl, T.; Marchioro, A.; Moon, S.-J.; Humphry-Baker, R.; Yum, J.-H.; Moser, J. E.; Grätzel, M.; Park, N.-G. Lead Iodide Perovskite Sensitized AllSolid-State Submicron Thin Film Mesoscopic Solar Cell with Efficiency Exceeding 9%. Sci. Rep. 2012, 2, 591. (8) Yang, W. S.; Noh, J. H.; Jeon, N. J.; Kim, Y. C.; Ryu, S.; Seo, J.; Seok, S. I. High-Performance Photovoltaic Perovskite Layers Fabricated Through Intramolecular Exchange. Science 2015, 348, 1234−1237. (9) Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341− 344. (10) Chen, Q.; Zhou, H.; Hong, Z.; Luo, S.; Duan, H.-S.; Wang, H.H.; Liu, Y.; Li, G.; Yang, Y. Planar Heterojunction Perovskite Solar Cells via Vapor-Assisted Solution Process. J. Am. Chem. Soc. 2014, 136, 622−625. (11) Eperon, G. E.; Stranks, S. D.; Menelaou, C.; Johnston, M. B.; Herz, L. M.; Snaith, H. J. Formamidinium Lead Trihalide: a Broadly Tunable Perovskite for Efficient Planar Heterojunction Solar Cells. Energy Environ. Sci. 2014, 7, 982−988. (12) Noel, N. K.; Stranks, S. D.; Abate, A.; Wehrenfennig, C.; Guarnera, S.; Haghighirad, A.-A.; Sadhanala, A.; Eperon, G. E.; Pathak, S. K.; Johnston, M. B.; Petrozza, A.; Herz, L. M.; Snaith, H. J. LeadFree Organic-Inorganic Tin Halide Perovskites for Photovoltaic Applications. Energy Environ. Sci. 2014, 7, 3061−3068. (13) Hao, F.; Stoumpos, C. C.; Guo, P.; Zhou, N.; Marks, T. J.; Chang, R. P. H.; Kanatzidis, M. G. Solvent-Mediated Crystallization of 6035
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