Amorphous Tin Oxide as a Low-Temperature-Processed Electron

Subscriber access provided by University of Newcastle, Australia

Article

Amorphous Tin Oxide as Low Temperature-Processed Electron Transport Layer for Organic and Hybrid Perovskite Solar Cells Jeremy Barbe, Max L. Tietze, Marios Neophytou, Banavoth Murali, Erkki Alarousu, Abdulrahman El Labban, Mutalifu Abulikemu, Wan Yue, Omar F. Mohammed, Iain McCulloch, Aram Amassian, and Silvano del Gobbo ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b13675 • Publication Date (Web): 08 Feb 2017 Downloaded from http://pubs.acs.org on February 10, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

ACS Applied Materials & Interfaces is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Applied Materials & Interfaces

Amorphous Tin Oxide as Low TemperatureProcessed Electron Transport Layer for Organic and Hybrid Perovskite Solar Cells Jérémy Barbé,a Max L. Tietze,*a Marios Neophytou,a Banavoth Murali,a Erkki Alarousu,a Abdulrahman El Labban,a Mutalifu Abulikemu,a Wan Yue,a Omar F. Mohammed r,a Iain McCulloch,a Aram Amassian,a Silvano Del Gobbo*a a

King Abdullah University of Science and Technology, KAUST Solar Center,

Physical Science and Engineering Division. Thuwal 23955-6900, Saudi Arabia Email: [email protected]; [email protected]

KEYWORDS. Perovskite Solar Cells, Tin Oxide Electron Transport Layer, Chemical Bath Deposition, Ultraviolet Photoelectron Spectroscopy, Organic Solar cells.

ABSTRACT: The chemical bath deposition of tin oxide (SnO2) thin films as electron transporting layer (ETL) in planar heterojunction n-i-p organohalide lead perovskite and organic bulk heterojunction solar cells is reported. The amorphous SnO2 films are grown from a nontoxic aqueous bath of tin chloride at very low temperatures (55°C) and do not require postannealing treatment to work very effectively as ETL in planar heterojunction n-i-p organohalide

ACS Paragon Plus Environment

1


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 32

lead perovskite or organic bulk heterojunction solar cells, in lieu of the commonly used ETL materials TiO2 and ZnO, respectively. Ultraviolet photoemission spectroscopy measurements on the glass/ITO/SnO2/MAPbI3/Spiro-OMeTAD device stack indicate that extraction of photogenerated electrons is facilitated by a perfect alignment of the conduction bands at the SnO2/MAPbI3 interface, while the deep valence band of SnO2 ensures strong hole-blocking properties. Despite exhibiting very low electron mobility, the excellent interfacial energetics combined with high transparency (Egap,optical > 4 eV) and uniform substrate coverage make the amorphous SnO2 ETL prepared by chemical bath deposition an excellent candidate for potentially low cost and large scale fabrication of organohalide lead perovskite and organic photovoltaics.

INTRODUCTION Perovskite solar cells (PSCs) based on hybrid organic-inorganic absorbers have recently attracted tremendous attention in the research community due to their rapid progress in power conversion efficiency (PCE) while being amenable to solution processing. In just a few years, this technology reached more than 20% certified efficiency at laboratory scale.1-2 The combination of high PCE, scalability and low-cost solution processability of organohalide perovskite solar cells is likely to make these competitive with more mature thin-film photovoltaic technologies. The impressive photovoltaic performances of PSCs is mostly attributed to the intrinsic properties of organohalide perovskites, such as suitable and tunable bandgap,3 high optical absorption coefficient,4 high carrier mobility,5 very long carrier diffusion length,6 and the capability for ambipolar carrier transport.7 However, for PSC devices to work to


2

Page 3 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


their full potential, a careful choice of selective contacts must be made. Recent efforts in this field have therefore focused on synthesis and development of suitable electron transport layers (ETL) and hole transport layers (HTL).8-9 These selective layers are critical for achieving high efficiency solar cells because they prevent interfacial recombination which directly affects the open-circuit voltage (VOC) and fill factor (FF). To date, the highest efficiency and majority of research into PSCs uses the standard architecture with mesoporous or compact TiO2 as ETL directly deposited on ITO/glass or FTO/glass substrates, and Spiro-OMeTAD (2,2′,7,7′tetrakis(N,N-di-p-methoxyphenylamine)-9,9′-spirobifluorene) as HTL on top of the photo-active perovskite in the so-called n-i-p device architecture. Despite the high efficiency obtained with TiO2 ETL, this metal-oxide still exhibits some drawbacks that hinder the commercialization of stable and reliable PSCs. In particular, UV light exposure results in modification of the TiO2 contact angle,10 which makes the processing of perovskite films more challenging and less reliable since the solution wetting and perovskite crystallization properties are tightly related to the surface energetics. TiO2 surfaces also suffer from UV light instability due to photocatalytic activity which may also hinder reproducibility over time.11,12 Moreover, researchers have shown a significant misalignment between the conduction bands of high quality TiO2 prepared by atomic layer deposition and the photoactive perovskite layer, to which they have attributed the strong hysteretic behavior of the J-V characteristic.13 Considerable effort has been made in recent years to replace TiO2 by other metal oxides possessing similar or even better electrical and optical properties, while being processed in low temperature conditions.8,14,15 For example, Liu et al. achieved 15.7% efficient planar PSCs with room-temperature ZnO ETL.8 They took advantage of the fact that ZnO is substantially thinner


3


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 32

than TiO2 and does not require sintering and, thus, could prepare flexible organohalide perovskite solar cells with efficiencies above 10%. The drawback of ZnO is its chemical instability toward organohalide perovskite compounds under visible and UV light irradiation16. CdS deposited by chemical bath deposition (CBD) at low temperature (< 100°C) was also used as ETL in perovskite solar cells reaching 13.2% stabilized efficiency, however the devices suffered from external quantum efficiency (EQE) losses in the UV range due to parasitic absorption in the rather low bandgap CdS film (Egap,optical = 2.4 eV)17,18. Another concern is the high toxicity of cadmium, which makes the liquid waste management of cadmium-containing solutions critical. Hence, the migration towards high bandgap and low toxicity chemical bath deposited materials is pursued. Within the material class of metal oxides, SnO2 was shown to possess higher bandgap, electron mobility and better stability than TiO2.19,20 Despite these remarkable properties and the deep valence band maximum, which is supposed to effectively block photo-generated holes from the absorber, the use of SnO2 as ETL for PSCs has been reported only very recently. Among the first, Li et al. presented PSCs with mesoporous SnO2 single crystals as ETL.21 The rather poor performance of ca. 3.8% was mainly caused by strong charge recombination due to the large interfacial area between SnO2 and the perovskite. Ke et al. showed the best performing planar perovskite cells using SnO2 ETLs prepared from a solution of SnCl2:2H2O in ethanol spin-coated on FTO, demonstrating an average efficiency of 16.0% and higher JSC, VOC and FF than their TiO2-based counterparts.22 The authors attributed the performance improvement in the SnO2 cell to higher transparency of the FTO/SnO2 stack as well as to lower carrier recombination at the SnO2/MAPbI3 interface. Concurrently, Song et al. showed PSCs with 13% efficiency using sinter-less SnO2 nanoparticles ETL.23 These solar cells exhibited promising stability even after


4

Page 5 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


one month in ambient air, while TiO2 based PSCs lost 30% of the PCE in 20 days. Recently, the same group reported that MAPbI3 deposited on pure SnO2 exhibited better thermal stability as compared to ZnO, which was explained by the reduced number of hydroxyl groups or/and residual acetate ligands at the surface.24 These different studies show the potential of SnO2 as an alternative ETL to TiO2 or ZnO for high efficiency PSCs. Yet, much of the work previously performed on SnO2 utilizes extensive annealing at 180-200°C to achieve crystalline films with tetragonal rutile structure. This represents a critical obstacle to the potential application of solution-processed SnO2 to flexible transparent substrates such as PET or PEI which can withstand temperatures of only 100°C to 150°C. Here, we report on the low temperature (55°C) chemical bath deposition (CBD) of amorphous SnO2 thin films and their successful implementation as ETLs in efficient planar heterojunction ni-p PSCs. CBD is an inexpensive process which has widely been used for the deposition of CdS or Zn(S,O,OH) and is well-adapted for large-area deposition.25,26 In our work, the CBD-SnO2 is based on a one-step, simple non-toxic aqueous bath of tin (IV) chloride precursor deposited at very low temperature, and does not require post-deposition annealing. We investigate the optical and electronic properties of CBD-SnO2 and determine in detail the energetic level alignment of an ITO/SnO2/MAPbI3/Spiro-OMeTAD stack through the use of ultraviolet photoemission spectroscopy (UPS). The best performing perovskite devices yield a PCE of 14.8% thanks in part to the electronic properties, high bandgap (Egap,optical > 4 eV) and excellent transparency of the amorphous SnO2 ETL. The versatility of the CBD-SnO2 and its applicability across different areas of photovoltaics are further highlighted by fabricating so-called inverted polymer solar


5


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 6 of 32

cells and achieving performance parity with ZnO, the most commonly used inorganic ETL in organic solar cells.

EXPERIMENTAL Materials. All materials were used without purification. Anhydrous dimethylsulfoxide (DMSO), gamma-butyrolactone (GBL), ethanol, diethyl ether, methylamine, 57 wt% hydroiodic acid in water, 4-tert-butylpyridine (TBP), anhydrous acetonitrile, anhydrous chlorobenzene, and lithium bis(trifluoromethane-sulfonyl)imide (Li-TFSI) were purchased from Sigma–Aldrich, whereas 99.999% lead iodide (PbI2) was purchased from Alfa Aesar. Tin (IV) chloride pentahydrate (SnCl4:5H2O) 98% was purchased from Acros Organics and tris(2-(1H-pyrazol-1yl)-4-tert-butylpyridine)cobalt(III) bis(trifluoromethanesulphonyl)imide (FK 102 Co(III) TFSI Salt) was purchased from Dyesol. Spiro-OMeTAD 98% was purchased from Bodrum Chem. Methylammonium iodide (MAI) was synthesized in accordance with the literature.27 CH3NH3PbI3 was prepared by reacting CH3NH3I and PbI2 (1:1 molar ratio) in GBL/DMSO (3:2 vol. ratio) with a concentration of 800 mg.ml-1, at 60°C for 12 h. Spiro-OMeTAD solution was prepared by dissolving 80 mg of Spiro-OMeTAD in 1 ml of chlorobenzene alongside 30 µL Tert-butyl pyridine (TBP), 18 µL Li-TFSi stock solution (520 mg Li-TFSI in 1 ml acetonitrile) and 30 µL FK 102 Co(III) TFSI Salt stock solution (300 mg FK 102 Co(III) TFSI Salt in 1 ml acetonitrile). This solution was stirred for 1h at room temperature.

Deposition of SnO2 and TiO2 films. SnO2 films were synthesized by chemical bath deposition on ITO, glass and quartz substrates as described elsewhere.28 Briefly, 1 x 1 inch2 samples were placed face down in aqueous solution of tin chloride pentahydrate (0.028 M) at


6

Page 7 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


55°C with stirring at 400 rpm for 11.5 min. This duration was optimized to ensure deposition of a uniform film without pinholes before the solution becomes turbid. It was also attempted to increase the film thickness by carrying out multiple depositions but this process resulted in poor film quality and increased surface roughness. After deposition, they were finally rinsed thoroughly with deionized (DI) water and dried with N2. For deposition of compact TiO2, a solution of 730 ml titanium isopropoxide (98% Acros Organics) in 10 ml of anhydrous ethanol acidized by 73 ml of HCl (99.999% metal basis, Sigma Aldrich) was spin-coated on the FTO substrates at 2000 rpm for 30 s, dried at 125ºC for 10 min and then annealed in an oven for 1 h at 500ºC. Subsequently, the films were then rinsed by DI water and annealed once again at 500ºC for 30 min.

Perovskite device fabrication. Glass substrates with indium tin oxide patterns were first scrubbed with Extran 300 detergent before being sequentially cleaned with detergent in DI water, acetone and isopropanol in an ultrasonic bath. Subsequently, they were dried with N2 and exposed to UV-O3 plasma for 10 min. SnO2 CBD was carried out immediately after plasma cleaning. For passivation study of the SnO2/perovskite interface, PC70BM diluted in chlorobenzene at a concentration of 10 mg.ml-1 was spin-coated on SnO2 at 4000 rpm. The methylammonium lead iodide (MAPbI3) solution was then deposited inside a glovebox onto the SnO2 (or SnO2/PC70BM) films by a consecutive two-step spin-coating process at 1000 rpm for 10 s followed by 5000 rpm for 40 s. During the high-speed step, chlorobenzene was drop-casted on the rotating sample after 20 s of spinning to promote crystallization. The samples were annealed at 80°C for 10 min. After cooling down, the Spiro-OMeTAD solution was spincoated on the MAPbI3 layer at 2600 rpm for 30 s. The films were scratched with a razor blade to


7


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 32

reveal the ITO bottom contact on the edge of the sample. Finally, the samples were placed in a thermal evaporator for evaporation of 30 nm of gold and 80 nm of silver electrodes at a pressure below 10-7 Torr.

Polymer-based device fabrication. BHJ solar cells were fabricated with an inverted (ITO/aSnO2 or ZnO/TBTIT-h:PC71BM/MoOx/Ag) configuration. ZnO nanoparticles were spin-coated at 5000 rpm and used as received. Devices were prepared with a TBTIT-h:PC71BM blend ratio of 1:2 solution processed from o-dichlorobenzene.

Films and devices characterization. The structure of SnO2 films was analyzed by X-ray diffraction (XRD) using a Bruker Advance D8 diffractometer with Cu Kα radiation (λ = 1.5418 Å). Cross-section scanning electron microscopy (SEM) images were taken at 5 kV accelerating voltage with 30 µm beam aperture, using a Zeiss Auriga microscope equipped with an in-lens detector. Atomic force microscopy (AFM) scans were performed with Dimension Icon from Veeco on 10 x 10 µm2 surface areas in tapping mode. Specular transmittance measurements were carried out in the range from 175 to 800 nm using Cary 5000 double beam mode UV-Vis spectrophotometer. Absorbance measurements of the perovskite films were taken using a Cary 6000i spectrophotometer equipped with a diffuse reflectance accessory (integrating sphere). The sample was positioned at the center of the sphere and tilted by 20o with respect to the incident light beam in order to collect the entire specular and diffuse light, after reflection and transmission through the sample. The UPS measurements were performed with an Omicron SPHERA hemispherical analyzer under He-I excitation (21.22 eV) of an attenuated discharge lamp (Omicron/Focus HIS 13, 1/12 attenuation). The base pressure of the analysis chamber was


8

Page 9 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


8·10-10 mbar. As substrates, ITO covered glasses were used. The samples were prepared in glovebox according to the device procedures and transferred to the UPS analysis chamber via ambient conditions, estimating the total air exposure time to less than 2 min. The film work functions (Wf) are determined from the spectra high binding energy cut-offs via:

ܹ݂ = 21.22eV − ‫ܧ‬ு஻ா஼ ,

where EHBEC is the respective cut-off position. The film ionization energies (IE) are obtained by adding the photoemission onset of the valence edge EVB w.r.t. to the Fermi level EF:

‫ ݂ܹ = ܧܫ‬+ ‫ܧ‬௏஻ .

The multi-channeltron detector and analyzer were calibrated to the Fermi edge of a sputtercleaned Au foil, yielding an energy resolution of 100 meV at room temperature. J-V measurements of solar cells were performed in glovebox with a Keithley 2400 source measure unit and an ABET Technology Sun 3000 class AAA solar simulator calibrated to 1 sun (AM1.5 G). All devices possess active areas of 0.1 cm2. They were measured from 1.2 V to -0.2 V with a scan rate of 0.28 V.s-1. The external quantum efficiency (EQE) measurements were performed at zero bias by illuminating the device with monochromatic light supplied from a Xenon arc lamp in combination with a dual-grating monochromator. The number of incident photons on the sample was measured for each wavelength by using a NIST calibrated silicon photodiode.


9


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 32

The electron mobility of SnO2 was determined by means of space charge limited current (SCLC) measurements in a vertical glass/ITO/SnO2/Au device configuration. Respective I-V curves were acquired by a Keithley 2400 Source Measure Unit at 300 K under dark conditions. The Au metal contact (area of 0.0314 cm2) was thermally evaporated at 0.2 A.s-1 to form the ITO/SnO2/Au geometry on a 20 nm SnO2 thin film ensuring the accurate thickness parameter for the global fitting. The dielectric constant of the SnO2 was considered to be 9.86 as reported elsewhere.29 The space charge limited current was ensured by reversing the polarity of the electrodes, since dissimilar electrodes were used in the present case.

RESULTS AND DISCUSSION Microstructure and properties of SnO2 prepared by chemical bath deposition Low and high magnification top-view scanning electron micrographs (SEM) of a CBD-SnO2 layer formed on top of an ITO-coated glass substrate after only 11.5 min are shown in Fig. 1a and b. The SnO2 film forms a conformal, compact and pin-hole free layer with ~20 nm thickness, uniformly coating the surface of ITO. The film has moderate surface roughness as shown by the respective atomic force microscopy (AFM) scan of Fig. 1c (quadratic mean Rq = 5.5 nm). As a matter of comparison, the compact TiO2 film deposited on FTO possesses a surface roughness of Rq = 16.2 nm (cf. Fig. S1). Based on these coating characteristics, we expect minimal shunt pathways between the perovskite absorber and the ITO contact.30 An X-ray diffraction (XRD) spectrum of as-deposited SnO2 on glass is plotted in Fig. 1d. It shows that the low-temperature SnO2 is in the amorphous state per XRD and does not exhibit any diffraction features commonly attributed to the tetragonal rutile SnO2.23 CBD-prepared a-SnO2 films processed in the same way


10

Page 11 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


and subsequently annealed on a hot plate at 100°C or 150°C for 1h do not show evidence of further crystallization either (Fig. S2).

(a) (b)

(b)

10 µm

200 nm

(c)

50 nm

2 µm

0 nm

Figure 1. SEM top-views at (a) low and (b) high magnification of chemical bath deposited aSnO2 on ITO/glass substrate. (c) 10 x 10 µm2 AFM scan of a-SnO2 on ITO. (d) X-ray diffraction spectrum of a-SnO2 on glass. The electron mobility of the amorphous SnO2 (a-SnO2) is investigated by SCLC measurements as shown in Fig. S3. By fitting the Child’s regime with the Mott-Gurney equation, a mobility of µ = 1.5 x 10-8 cm2V-1s-1 is determined. The low mobility of the chemical bath deposited SnO2 obtained here is ascribed to its amorphous nature and could potentially be problematic for sufficient electron conduction in an ETL. However, as shown later, this is actually not the case, probably due to the sufficient thinness of the film. Furthermore, it should be noted that efficient inverted organic BHJ solar cells were recently obtained by using amorphous ZnO as ETL, which also exhibited much lower electron mobility (10-5-10-4 cm2V-1s-1) than its crystalline counterparts.31 The high optical transmittance of the ETL is important to maximize the passage of sunlight into the perovskite absorber in n-i-p solar cells. The absorption of UV light in the ETL must be


11


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 32

minimized by using a semiconductor of equal or greater bandgap than the ITO electrode and the glass substrate, so as to enhance the external quantum efficiency and photocurrent. In Fig. 2, the optical transmittance of different substrate-film combinations is shown in the range from 180 nm to 850 nm. Owing to the high bandgap of a-SnO2, CBD a-SnO2 films were deposited on high quality quartz substrates. Indeed, bare quartz is highly transparent and its absorption onset is outside the measurement range, whereas typical float glass substrates exhibit a sharp absorption onset at 320 nm. The absorption onset of a-SnO2 is observed at 280 nm, which corresponds to an optical gap of 4.4 eV (cf. corresponding Tauc plot in Fig. S4). This value is higher than the bandgap of crystalline SnO2 (3.6 eV), as determined by ellipsometry and transmittance measurements in another work13, which is reasonable since amorphous phase materials have wider bandgaps than their crystalline phases. In comparison, the absorption onset of compact TiO2 is observed at ∼380 nm (Egap,optical = 3.2 eV). As TiO2 is commonly deposited on FTO/glass, which can sustain high annealing temperatures required to form the rutile phase of TiO2, we also compare the transmittance of bare FTO-coated glass with TiO2/FTO-coated glass. It is clear that the transmittance is mostly limited by the TiO2 film that suffers from severe losses in the 300 – 450 nm region. Interference oscillations are also observed in the visible spectrum due to the high thickness (∼350 nm) of the FTO film, potentially causing more losses due to interference effects. In comparison, the transmittance spectra of ITO-coated glass (100 nm) and a-SnO2/ITO–coated glass are almost identical, including the low wavelength range, which indicates that the transparency is limited by ITO in this case. Hence, CBD a-SnO2 is optically most suitable as a highly transparent ETL since it does not prevent any UV light to reach the actual photo-active perovskite film. The absorbance of the MAPbI3 thin film on glass was also measured using an integrating sphere, as indicated on the right axis of Fig. 2. A sharp absorption onset can be


12

Page 13 of 32

observed at 780 nm owing to the bandgap of the perovskite absorber layer. Therefore, the range of useful solar harvesting spectrum available to single junction solar cells using a-SnO2 as ETL is only limited by the transmittance and absorbance onsets of ITO (300 nm) and MAPbI3 (780 nm), respectively. This is in contrast to other ETLs like TiO2, ZnO or CdS having lower bandgap values.17 100

100

80

80

60

60

bare quartz bare glass ITO on glass FTO on glass SnO2 on quartz

40

MAPbI3

40

Absorbance (%)

Transmittance (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


TiO2 on quartz

20

20

SnO2 on ITO TiO2 on FTO MAPbI3 on glass

0 200

300

400

500

600

700

800

0

Wavelength (nm)

Figure 2. Transmittance spectra of various substrates (quartz, glass, ITO, FTO) and thin films (aSnO2 on quartz or ITO, compact TiO2 on quartz or FTO). The absolute absorbance of MAPbI3 film on glass is shown on the right axis (black dashed line).

We also assess the sensitivity of a-SnO2 to UV light exposure by measuring the contact angle of water on a-SnO2 before and after exposure to UV-O3 for 10 minutes. As shown in Fig. S5, the contact angle of a-SnO2 remains nearly unchanged. This suggests that the a-SnO2 film is more resistant to UV radiation and does not undergo critical surface chemical and energetic modifications unlike TiO2 surfaces, which become more hydrophobic by irradiation with UV light.10 This will give rise to modifications at the TiO2/MAPbI3 interface and possibly initiate the MAPbI3 degradation.

Ultraviolet photoelectron spectroscopy and interfacial energetics


13


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 14 of 32

Next, we provide ultraviolet photoelectron spectroscopy (UPS) investigations on the a-SnO2 as well as the entire stack to gain insight into the energy level alignment at the various ITO/aSnO2/MAPbI3/Spiro-OMeTAD hetero-interfaces. Respective samples were transferred from the glovebox to the UPS analysis chamber via ambient conditions, ensuring air exposure times of less than 2 min for minimizing surface contamination/modification effects. In Fig. 3a, we show UPS spectra corresponding to the following incrementally deposited and independently measured stack sequences: ITO/a-SnO2, ITO/a-SnO2/MAPbI3 and ITO/a-SnO2/MAPbI3/SpiroOMeTAD. The deduced energy level diagram of the stack is presented in Fig. 3b. Determining the valence band onset of a-SnO2 w.r.t. the ITO Fermi level yields EVB = 4.03 eV, i.e., the conduction band onset is only 0.37 eV above EF when adding the optical gap of 4.4 eV, which strongly points to an n-type a-SnO2. The film ionization energy (IE) is 8.61 eV. In comparison, the IE and EVB of highly crystalline SnO2 grown on FTO by atomic layer deposition (ALD) were reportedly IE = 7.86eV and EVB = 2.21 eV.13 Spin-coating and thermally converting MAPbI3 on top of the CBD a-SnO2 results in a valence band onset of 1.31 eV w.r.t. EF, Wf = 4.35 eV, and an ionization energy of 5.66 eV which is fairly similar to the value reported by Correa Baena et al. for MAPbI3 deposited on ALD-SnO2 (IE = 5.65eV). The energetic offset of the conduction bands at the a-SnO2/MAPbI3 interface is practically negligible (0.02 eV), if considering an electronic gap for the perovskite of 1.7 eV, as determined previously by combined UPS/IPES.32 This favorable alignment provides efficient electron transfer from the MAPbI3 layer to the a-SnO2 ETL. In contrast, it has been shown that even a slight mismatch between the conduction bands of TiO2 and MAPbI3 can hinder the charge transfer from which, eventually, the photovoltaic performance suffers.13 Apart from that, the valence band offset at the CBD a-SnO2/perovskite interface of about 2.7 eV (cf. Fig. 3b) provides a large barrier for blocking photo-generated holes


14

Page 15 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


efficiently. Finally, UPS measurements on the sequence ITO/a-SnO2/MAPbI3/Spiro-OMeTAD yield EVB = 0.40 eV, Wf = 4.75 eV, and IE = 5.15 eV for the Spiro-OMeTAD HTL on top. Thus, we deduce a valence band offset of 0.91 eV at the MAPbI3/Spiro-OMeTAD interface, which corresponds to a hole injection barrier ca. 0.5 eV larger than reported previously for a TiO2based stack.32 However, Schulz et al. measured an IE of just 5.4 eV for MAPbI3, whereas more recent reports point towards much higher values in the range of 5.6 to 6.4 eV.33,34 Therefore, we presume that the larger barrier found here is not exclusively due to the usage of a-SnO2 but also to the higher IE of the MAPbI3, whose ionization energy values have been shown to strongly depend on the actual processing conditions, conversion degree, and thus, stoichiometry.33,34 We have estimated a Pb/N ratio of ∼1.05 by XPS (Fig. S6). Finally, it shall be noted that the above analysis must be regarded as an estimation since it does not account for possible level bending effects within the layers of the stack,35 in particular at the MAPbI3/Spiro-OMeTAD interface.32


15


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 32

Figure 3. (a) UPS spectra of incrementally stacked a-SnO2/perovskite/Spiro-OMeTAD films deposited on ITO/glass substrates. (b) Corresponding energy level diagram of the ITO/aSnO2/perovskite/Spiro-OMeTAD layer stack determined by UPS and UV-Vis spectrophotometry.

Methylammonium lead iodide perovskite solar cells Planar heterojunction n-i-p solar cells with the structure glass/ITO/ETL/MAPbI3/SpiroOMeTAD/Au are fabricated to test the effectiveness of CBD a-SnO2 as ETL in comparison to TiO2 in a similar architecture. A cross-sectional SEM image of a typical perovskite device with SnO2 ETL is given in the inset of Fig. 4a. The thickness of the perovskite and Spiro-OMeTAD layers are 250 nm and 170 nm, respectively. The CBD a-SnO2 film is too thin (∼ 20 nm) to yield a clear contrast, hence, being hardly visible in the SEM image. Current density-voltage (J-V) curves for the best performing solar cells incorporating either a-SnO2 or reference TiO2 as ETL


16

Page 17 of 32

are plotted in Fig. 4a. The corresponding PV figures of merit are summarized in Table 1. The aSnO2-based device achieved a PCE of 14.8% with VOC = 1.05 V, JSC = 21.3 mA.cm-2 and FF = 66.3%. Compared to our TiO2-based reference with the same perovskite thickness, a slight improvement in JSC is observed, which can be attributed to the higher transmittance of a-SnO2 in the UV and visible range (cf. Fig. 2). Furthermore, the FF and VOC are notably improved by using CBD a-SnO2, which is ascribed to the lower surface roughness of a-SnO2 on ITO as well as the favorable energy alignment at the a-SnO2/MAPbI3 interface. The charge extraction effectiveness of the CBD a-SnO2 is further confirmed by time-resolved photoluminescence (TRPL) measurements (cf. Fig. S7). For the a-SnO2/MAPbI3 interface, a two-component exponential PL decay with τ1=7.4 ns and τ2=20.4 ns is found, which is on par with application of SnO2 nanoparticles as ETL (τ=16.1 ns)24. Instead, pure perovskite on glass possesses much longer PL lifetimes, e.g., measured to τ1=64.1 ns and τ2=444.2 ns45. Interestingly, the PL decay of perovskite on ITO is slightly faster, which suggests efficient charge extraction at the ITO/MAPbI3 interface, probably due to the low surface roughness of ITO. However, all devices prepared directly on ITO, without SnO2 or TiO2 ETL, resulted in very poor PV performances.

(a)

20 15 10 a-SnO2 TiO2

5 0 0.0

0.2

(b)

80 EQE (%)

Current density (mA.cm-2)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


60 40

a-SnO2

20

0.4 0.6 0.8 Voltage (V)

1.0

0

TiO2 300 400 500 600 700 800 Wavelength (nm)


17


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 32

Figure 4. J – V curves (a) and respective EQE spectra (b) of the obtained best performing perovskite solar cells with CBD a-SnO2 or compact TiO2 as ETLs, respectively. Inset: crosssection SEM image of a typical perovskite solar cell with SnO2 ETL. Table 1. PV parameters of a-SnO2-based and TiO2-based perovskite solar cells with best performances. Electron transport layer

VOC (V)

JSC (mA.cm-2)

FF (%)

PCE (%)

a-SnO2

1.05

21.3

66.3

14.8

TiO2

0.99

20.8

61.6

12.8

Unfavorably, a strong hysteresis is observed between the forward and reverse scans of perovskite devices with a-SnO2 ETL (Fig. S8). In forward, the device suffers from strong decreases in VOC and fill factor. As described by several recent studies,36,37,38,39,40,41 such strong hysteresis effects can be assigned to three factors: the ferroelectric nature of the perovskite material, the migration of I- and CH3NH3+ ions within the MAPbI3 crystal, and by accumulation of charges in trap states at the a-SnO2/MAPbI3 interface. Such undesired hysteresis effects were reported in previous studies on devices comprising crystalline SnO2 as ETL as well.30 However, in the case of high quality SnO2 grown by atomic layer deposition (ALD), almost no hysteresis was observed, which was consistently explained by a favorable energy level alignment and trapfree configuration of the ALD-SnO2/perovskite interface.13 Therefore, passivation of the here reported a-SnO2 surface is implemented by inserting a thin PC60BM layer on top of the a-SnO2. The J-V curve of the respective device is given in Fig. S8, and indeed, a surface passivation effect is observed, i.e. the hysteresis is strongly reduced compared to bare a-SnO2-based devices without PC60BM. This result indicates that a significant contribution to the J-V hysteresis arises from trap states at the a-SnO2/MAPbI3 interface. Due to the amorphous nature of the CBD-deposited a-SnO2, it is reasonable to assume that the surface of the a-SnO2 itself is composed of a high density of trap states, which eventually induce polarization of the device and decrease the VOC and FF. In


18

Page 19 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


contrast, the VOC remains almost constant, around 1.02 V, if the interface is passivated by PC60BM. These surface defects are an intrinsic drawback of the low temperature chemical bath deposition, which does not allow to grow highly crystalline and defect-free SnO2. A similar kind of passivation was previously demonstrated by Wojciechowski et al.,42 where an effective decrease of the device J-V hysteresis was attributed to a covalent binding of the used fullerene derivative C60-SAM to the TiO2 ETL surface. Besides, Shao et al. showed that fullerenes can passivate defects that originate from grain boundaries in the perovskite film.40 Hence, PC60BM used here as interlayer at the ETL/perovskite interface could basically passivate both SnO2 surface trap states and perovskite grain boundaries. Migration of negative iodine ions in the perovskite film is another phenomenon that has to be taken into account to consistently explain the observed hysteresis behavior. A recent study showed that an important cause of the hysteric behavior comes from the formation of an internal electric field upon ions migration in the perovskite layer.41 Introducing a PCBM interfacial layer provided a way to decrease the hysteresis by trapping I- ions in the fullerene film, which eventually becomes negatively doped. This self-sustained internal field improves charge extraction at the anode side and sensitively decreases the hysteresis behavior. The values reported in Table 1 are further verified by plotting the evolution of the current density versus time at the firstly determined maximum power point (Fig. S9). For devices using a-SnO2 ETL, the current density undergoes a decrease of 10% after 50 s, and then stabilizes until the end of measurement. For TiO2-based devices, a more pronounced drop of the current is observed within the first 50 s, but then the current gradually increases to reach a 10% decrease after 300 s, which is on par with the SnO2-based device. This trend confirms the importance of ions migration in the operation of the device when an external field is applied. The time scale for


19


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 32

stabilization of the device under operation (50 s) is in good agreement with the dynamics of ions migration and cannot be attributed to electrons trapping, which is generally a much faster process. The stabilization of the photocurrent after 50 s operation confirms that the a-SnO2 is not light soaking sensitive and has good stability towards UV irradiation, which is consistent with the contact-angle measurements. In case of TiO2 ETL, the device current starts to increase after the ions migration stage, which evidences the instability of compact TiO2, probably due to photochemical reactions. The external quantum efficiency (EQE) of the best performing a-SnO2/MAPbI3 device is plotted in Fig. 4b and compared to PSCs using compact TiO2. For a-SnO2-based devices, it can be observed that the EQE cut-off at low and high wavelengths rigorously follows the absorption onsets of ITO/a-SnO2 and MAPbI3, respectively (cf. Fig. 2). The JSC obtained from the integral of the EQE spectrum matches the value obtained from J-V curves, which confirms the reliability of the current-voltage measurements. More than 80% EQE is reached from 350 nm to 700 nm, without any shoulder in the UV range. On the contrary, the TiO2-based device exhibits significant EQE losses in the 300 – 400 nm range due to parasitic absorption in the TiO2 film. A histogram of cell efficiencies for 20 separate a-SnO2-based perovskite devices is shown in Fig. S10, which yields an efficiency spread from 11.5% to 14.8%, having 11 devices within 12.5% and 13.5%, thus, indicating high processing reliability for CBD-deposition of the a-SnO2 ETL. As regard to the previous results, it can be concluded that a-SnO2 is a suitable material as an electron transport layer for overall light harvesting, but also benefits from attractive processing conditions like the low temperature. Besides, it was found that the a-SnO2 films could be stored for several days in air without any degradation being induced by moisture or UV irradiation, and hence, could be further used for device fabrication without additional treatments.


20

Page 21 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


The finding that planar heterojunction n-i-p solar cells based on a-SnO2 as ETL achieve better performances than devices based on TiO2 is surprising, especially when considering the low electron mobility in amorphous SnO2, which is in the order of 10-8 cm2V-1s-1. The solar cell results are therefore remarkable and indicate that this low mobility has not considerably increased the series resistance (Rs) and reduced the FF. The observation of reasonable solar cell performance despite very low mobility in the ETL is not new. We recently demonstrated efficient organic bulk heterojunction solar cells using an amorphous ZnO as ETL. The electron mobility in this instance (approx. 10-5 cm2V-1s-1) was several orders of magnitude lower than that in crystalline ZnO, yet devices performed as well if not better than ZnO processed at higher temperature,31 and were nearly on par with Al:ZnO nanocrystals,43 partly because the ETL was a thin but compact film with excellent coverage of the ITO and exhibited favorable energetics for electron extraction. Therefore, we believe that the combinations of favorable conduction band matching between a-SnO2 and the perovskite and excellent hole blocking capabilities of a-SnO2, along with a thin film with compact, uniform and pin-hole free morphology can overcome potential resistive losses and prevent shunting, making the amorphous SnO2 ETL developed here very effective. Trap states at the a-SnO2/MAPbI3 interface, which potentially could cause interfacial recombination, can be passivated by a thin PC60BM interlayer. Finally, we found that annealing the a-SnO2 films on a hot plate at 100°C or 150°C is neither significantly improving nor degrading the cell performances. Much higher annealing temperatures of at least 400°C would actually be necessary to crystallize the a-SnO2 thin films and improve its transport and interface properties, however making it unsuitable for future solar cell manufacturing on plastic substrates.28


21


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 32

Polymer solar cells Finally, we demonstrate the suitability of CBD a-SnO2 as ETL for organic BHJ solar cells on the example of so-called inverted cells comprising the thieno[3,2-b][1]benzothiophene isoindigobased donor polymer (TBTIT-h) blended with PC71BM acceptor.42 For this demonstration, we compare a-SnO2 head-to-head with ZnO, one of the most successful metal oxide ETLs used in the OPV community. We compare the performance of devices based on the architecture glass/ITO/ETL/TBTIT-h:PC71BM/MoOx/Ag

cells

using

either

the

CBD

a-SnO2

or

nanocrystalline ZnO as ETL. The ZnO film was spin-coated as received at room temperature and not annealed after deposition. The J-V characteristics and EQE spectra are shown in Fig. 5a and b, respectively. Both types of solar cells possess practically equal EQEs over the whole wavelength range, thus, similar short circuit currents of about JSC = 15.5 mA.cm-2, which explains their similar PCEs (cf. Table S1, 7.3% vs. 7.5%). Strikingly, the VOC of approx. 0.72 V is almost equal for both metal-oxide ETLs and comparable with devices reported previously44, i.e., no limitations, e.g., due to barrier effects, are found when applying the a-SnO2. These results strongly highlight the remarkable versatility of the low-temperature CBD-SnO2, which is shown to achieve performance parity with TiO2 and ZnO, the most commonly used and successful ETL materials in the perovskite and organic photovoltaics communities, respectively.


22

Page 23 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5. (a) J-V and (b) EQE curves of ITO/ETL/TBTIT-h:PC71BM/MoOx/Ag organic BHJ solar cells using either nc-ZnO or CBD-deposited a-SnO2 as ETL.

CONCLUSIONS In summary, preparation of amorphous SnO2 ETL was demonstrated by very low-temperature (55°C), scalable chemical bath deposition, using a simple non-toxic aqueous bath of tin chloride. In contrast to other approaches, the films did not require annealing, sintering or crystallization after deposition and could be used successfully in the amorphous state. The energetic level alignment within the solar cell stack was studied in detail by ultraviolet photoelectron spectroscopy, and shown to be very favorable to electron extraction and hole blocking in conjunction with the perovskite light absorber. Given the excellent optical properties of a-SnO2 and its high bandgap, the ETL results in very little to no parasitic absorption losses in comparison with other ETL materials, such as TiO2, ZnO or CdS. The best performing planar heterojunction n-i-p perovskite solar cells reached a PCE = 14.8%, outperforming our reference TiO2-based PSCs by 2%, despite being prepared slightly above room temperature and in conditions which yield low electron mobility. The J-V hysteresis caused by interfacial charge


23


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 24 of 32

build-up was then minimized by passivating the SnO2 surface with a thin PC60BM interlayer. The combination of excellent optical properties, interfacial energetics, complete and uniform coverage of the ITO surface despite being very thin, make the amorphous SnO2 ETL behave remarkably well for hybrid perovskite and organic solar cells.

SUPPORTING INFORMATION Atomic force microscopy scans of ITO, a-SnO2 and TiO2 surfaces, XRD spectrum of a-SnO2 thin-film after annealing, space-charge limited current (SCLC) of a-SnO2 thin-film, Tauc plot of as-deposited a-SnO2 thin-film, contact angles of water on a-SnO2 or TiO2 surfaces, survey XPS spectrum of MAPbI3 and a-SnO2 thin-film, current-voltage curves of a-SnO2-based perovskite solar cells with forward and reverse scans, evolution of the current density of a-SnO2 or TiO2based perovskite solar cell at working condition, histogram of power conversion efficiency for 20 perovskite solar cells using a-SnO2 ETL, PV parameters of the organic BHJ solar cells. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author * (M.T.) Email: [email protected] * (S.D.G.) Email: [email protected] Present Addresses †If an author’s address is different than the one given in the affiliation line, this information may be included here. Author Contributions


24

Page 25 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


J.B. prepared the SnO2 films and carried out SEM, AFM, XRD, UV-Vis and contact-angle characterizations; M.T. carried out UPS and XPS analysis; A.E.L., M.N. and M.A. fabricated and measured perovskite and polymer solar devices; M.B. carried out SCLC measurements and prepared TiO2 films; W.Y. synthesized TBTIT-h; O.F.M., E.A. carried out TRPL measurements, I.M.C. contributed to fruitful discussions; J.B. and M.T. wrote the manuscript; A.A. and S.D.G. reviewed the manuscript and supervised the project. ACKNOWLEDGMENT This research was supported by funding from King Abdullah University of Science and Technology. REFERENCES (1)

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 (7458), 316–319.

(2)

Yang, W. S.; Noh, J. H.; Jeon, N. J.; Kim, Y. C.; Ryu, S.; Seo, J.; Seok, S. I. HighPerformance Photovoltaic Perovskite Layers Fabricated through Intramolecular Exchange. Science 2015, 348 (6240), 1234–1237.

(3)

Noh, J. H.; Im, S. H.; Heo, J. H.; Mandal, T. N.; Seok, S. Il. Chemical Management for Colorful, Efficient, and Stable Inorganic-Organic Hybrid Nanostructured Solar Cells. Nano Lett. 2013, 13 (4), 1764–1769.

(4)

De Wolf, S.; Holovsky, J.; Moon, S. J.; Loper, P.; Niesen, B.; Ledinsky, M.; Haug, F. J.; Yum, J. H.; Ballif, C. Organometallic Halide Perovskites: Sharp Optical Absorption Edge and Its Relation to Photovoltaic Performance. J. Phys. Chem. Lett. 2014, 5 (6), 1035– 1039.


25


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(5)

Page 26 of 32

Stranks, S. D.; 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 2014, 341 (2013), 341–345.

(6)

Dong, Q.; Fang, Y.; Shao, Y.; Mulligan, P.; Qiu, J.; Cao, L.; Huang, J. Electron-Hole Diffusion Lengths > 175 Um in Solution-Grown CH3NH3PbI3 Single Crystals. Science 2015, 347 (6225), 967–970.

(7)

Snaith, H. J. Perovskites : The Emergence of a New Era for Low-Cost, High-Efficiency Solar Cells. J. Phys. Chem. Lett. 2013, 4, 3623–3630.

(8)

Liu, D.; Kelly, T. L. Perovskite Solar Cells with a Planar Heterojunction Structure Prepared Using Room-Temperature Solution Processing Techniques. Nat. Photonics 2013, 8 (2), 133–138.

(9)

Ye, S.; Sun, W.; Li, Y.; Yan, W.; Peng, H.; Bian, Z.; Liu, Z.; Huang, C. CuSCN-Based Inverted Planar Perovskite Solar Cell with an Average PCE of 15.6%. Nano Lett. 2015, 15 (6), 3723–3728.

(10)

Stevens, N.; Priest, C. I.; Sedev, R.; Ralston, J. Wettability of Photoresponsive Titanium Dioxide Surfaces. Langmuir 2003, 19 (8), 3272–3275.

(11)

Leijtens, T.; Eperon, G. E.; Pathak, S.; Abate, A.; Lee, M. M.; Snaith, H. J. Overcoming Ultraviolet Light Instability of Sensitized TiO2 with Meso-Superstructured Organometal Tri-Halide Perovskite Solar Cells. Nat. Commun. 2013, 4, 2885.

(12)

Barbé, J.; Thomson, A. F.; Wang, E. C.; Mcintosh, K.; Catchpole, K. Nanoimprinted TiO2 Sol–gel Passivating Diffraction Gratings for Solar Cell Applications. Prog. Photovoltaics Res. Appl. 2012, No. 20, 143–148.


26

Page 27 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(13)

Correa Baena, J. P.; Steier, L.; Tress, W.; Saliba, M.; Neutzner, S.; Matsui, T.; Giordano, F.; Jacobsson, T. J.; Srimath Kandada, A. R.; Zakeeruddin, S. M.; Petrozza, A.; Abate, A.; Nazeeruddin, M. K.; Grätzel, M.; Hagfeldt, A. Highly Efficient Planar Perovskite Solar Cells through Band Alignment Engineering. Energy Environ. Sci. 2015, 8, 2928–2934.

(14)

Bera, A.; Sheikh, A. D.; Haque, A.; Bose, R.; Alarousu, E.; Mohammed, O. F.; Wu, T. Fast Crystallization and Improved Stability of Perovskite Solar Cells with Zn2SnO4 Electron Transporting Layer : Interface Matters. ACS Appl. Mater. Interfaces 2015, 7 (51), 28404–28411.

(15)

Wang, K.; Shi, Y.; Dong, Q.; Li, Y.; Wang, S.; Yu, X.; Wu, M.; Ma, T. Low-Temperature and Solution-Processed Amorphous WOX as Electron-Selective Layer for Perovskite Solar Cells. J. Phys. Chem. Lett. 2015, 6 (5), 755–759.

(16)

Yang, J.; Siempelkamp, B. D.; Mosconi, E.; De Angelis, F.; Kelly, T. L. Origin of the Thermal Instability in CH3NH3PbI3 Thin Films Deposited on ZnO. Chem. Mater. 2015, 27 (12), 4229–4236.

(17)

Liu, J.; Gao, C.; Luo, L.; Ye, Y. Q.; He, X.; Ouyang, L.; Guo, X.; Zhuang, D.; Liao, C.; Mei, J.; Lau, L. W. M. Low-Temperature, Solution Processed Metal Sulfide as Electron Transport Layer for Efficient Planar Perovskite Solar Cell. J. Mater. Chem. A 2015, 3, 1– 6.

(18)

Dunlap-Shohl, W. A.; Younts, R.; Gautam, B.; Gundogdu, K.; Mitzi, D. B. Effects of Cd Diffusion and Doping in High-Performance Perovskite Solar Cells Using CdS as Electron Transport Layer. J. Phys. Chem. C 2016, 120 (30), 16437–16445.

(19)

Kılıç, Ç.; Zunger, A. Origins of Coexistence of Conductivity and Transparency in SnO2. Phys. Rev. Lett. 2002, 88 (9), 95501.


27


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(20)

Page 28 of 32

Snaith, H. J.; Ducati, C. SnO2-Based Dye-Sensitized Hybrid Solar Cells Exhibiting near Unity Absorbed Photon-to-Electron Conversion Efficiency. Nano Lett. 2010, 10 (4), 1259–1265.

(21)

Li, Y.; Zhu, J.; Huang, Y.; Liu, F.; Lv, M.; Chen, S.; Hu, L.; Tang, J.; Yao, J.; Dai, S. Mesoporous SnO2 Nanoparticle Films as Electron-Transporting Material in Perovskite Solar Cells. RSC Adv. 2015, 5, 28424–28429.

(22)

Ke, W.; Fang, G.; Liu, Q.; Xiong, L.; Qin, P.; Tao, H.; Wang, J.; Lei, H.; Li, B.; Wan, J.; Yang, G.; Yan, Y. Low-Temperature Solution Processed Tin Oxide as an Alternative Electron Transporting Layer for Efficient Perovskite Solar Cells. J. Am. Chem. Soc. 2015, 137, 6730–6733.

(23)

Song, J.; Zheng, E.; Bian, J.; Wang, X.-F.; Tian, W.; Sanehira, Y.; Miyasaka, T. LowTemperature SnO2-Based Electron Selective Contact for Efficient and Stable Perovskite Solar Cells. J. Mater. Chem. A 2015, 3, 10837–10844.

(24)

Song, J.; Zheng, E.; Wang, X. F.; Tian, W.; Miyasaka, T. Low-Temperature-Processed ZnO-SnO2 Nanocomposite for Efficient Planar Perovskite Solar Cells. Sol. Energy Mater. Sol. Cells 2016, 144, 623–630.

(25)

Jackson, P.; Hariskos, D.; Wuerz, R.; Kiowski, O.; Bauer, A.; Friedlmeier, T. M.; Powalla, M. Properties of Cu(In,Ga)Se2 Solar Cells with New Record Efficiencies up to 21.7%. Phys. status solidi - Rapid Res. Lett. 2015, 9 (1), 28–31.

(26)

Shin, D.; Kim, S.; Kim, J. Study of Band Structure at the Zn(S,O,OH)/Cu(In,Ga)Se2 Interface via Rapid Thermal Annealing and Their Effect on the Photovoltaic Properties. ACS Appl. Mater. Interfaces 2013, 5, 12921–12927.

(27)

Zhou, H.; Chen, Q.; Li, G.; Luo, S.; Song, T.; Duan, H.-S.; Hong, Z.; You, J.; Liu, Y.;


28

Page 29 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Yang, Y. Interface Engineering of Highly Efficient Perovskite Solar Cells. Science 2014, 345 (6196), 542–547. (28)

Khallaf, H.; Chen, C.-T.; Chang, L.-B.; Lupan, O.; Dutta, A.; Heinrich, H.; Haque, F.; del Barco, E.; Chow, L. Chemical Bath Deposition of SnO2 and Cd2SnO4 Thin Films. Appl. Surf. Sci. 2012, 258 (16), 6069–6074.

(29)

Roessler, D. M.; Albers, W. A. Infrared of. j. Phys. Chem. Solids 1972, 33, 293–296.

(30)

Dong, Q.; Shi, Y.; Wang, K.; Li, Y.; Wang, S.; Zhang, H.; Xing, Y.; Du, Y.; Bai, X.; Ma, T. Insight into Perovskite Solar Cells Based on SnO2 Compact Electron- Selective Layer. J. Phys. Chem. C 2015, 119, 10212–10217.

(31)

Jagadamma, L. K.; Abdelsamie, M.; El Labban, A.; Aresu, E.; Ngongang Ndjawa, G. O.; Anjum, D. H.; Cha, D.; Beaujuge, P. M.; Amassian, A. Efficient Inverted BulkHeterojunction Solar Cells from Low-Temperature Processing of Amorphous ZnO Buffer Layers. J. Mater. Chem. A 2014, 2 (33), 13321.

(32)

Schulz, P.; Edri, E.; Kirmayer, S.; Hodes, G.; Cahen, D.; Kahn, A. Interface Energetics in Organo-Metal Halide Perovskite-Based Photovoltaic Cells. Energy Environ. Sci. 2014, 7 (4), 1377.

(33)

Emara, J.; Schnier, T.; Pourdavoud, N.; Riedl, T.; Meerholz, K.; Olthof, S. Impact of Film Stoichiometry on the Ionization Energy and Electronic Structure of CH3NH3PbI3 Perovskites. Adv. Mater. 2016, 28, 553–559.

(34)

Kim, T. G.; Seo, S. W.; Kwon, H.; Hahn, J.-H.; Kim, J. W. Influence of Halide Precursor Type and Its Composition on the Electronic Properties of Vacuum Deposited Perovskite Films. Phys. Chem. Chem. Phys. 2015, 17, 1–6.

(35)

Guerrero, A.; Juarez-Perez, E. J.; Bisquert, J.; Mora-Sero, I.; Garcia-Belmonte, G.


29


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 30 of 32

Electrical Field Profile and Doping in Planar Lead Halide Perovskite Solar Cells. Appl. Phys. Lett. 2014, 105 (13), 133902. (36)

Snaith, H. J.; Abate, A.; Ball, J. M.; Eperon, G. E.; Leijtens, T.; Noel, N. K.; Stranks, S. D.; Wang, J. T.-W.; Wojciechowski, K.; Zhang, W. Anomalous Hysteresis in Perovskite Solar Cells Henry. J. Phys. Chem. Lett. 2014, 17 (5), 1511–1515.

(37)

Tress, W.; Marinova, N.; Moehl, T.; Zakeeruddin, S. M.; Mohammad K., N.; Grätzel, M.; Nazeeruddin, M. K.; Grätzel, M. Understanding the Rate-Dependent J–V Hysteresis, Slow Time Component, and Aging in CH3NH3PbI3 Perovskite Solar Cells: The Role of a Compensated Electric Field. Energy Environ. Sci. 2015, 8 (3), 995–1004.

(38)

Chen, B.; Yang, M.; Priya, S.; Zhu, K. Origin of J-V Hysteresis in Perovskite Solar Cells. J. Phys. Chem. Lett. 2016, 7 (5), 905–917.

(39)

Wu, Y.; Shen, H.; Walter, D.; Jacobs, D.; Duong, T.; Peng, J.; Jiang, L.; Cheng, Y. B.; Weber, K. On the Origin of Hysteresis in Perovskite Solar Cells. Adv. Funct. Mater. 2016, 6807–6813.

(40)

Shao, Y.; Xiao, Z.; Bi, C.; Yuan, Y.; Huang, J. Origin and Elimination of Photocurrent Hysteresis by Fullerene Passivation in CH3NH3PbI3 Planar Heterojunction Solar Cells. Nat. Commun. 2014, 5, 1–7.

(41)

De Bastiani, M.; Dell’Erba, G.; Gandini, M.; D’Innocenzo, V.; Neutzner, S.; Kandada, A. R. S.; Grancini, G.; Binda, M.; Prato, M.; Ball, J. M.; Caironi, M.; Petrozza, A. Ion Migration and the Role of Preconditioning Cycles in the Stabilization of the J-V Characteristics of Inverted Hybrid Perovskite Solar Cells. Adv. Energy Mater. 2016, 6 (2), 1–9.

(42)

Wojciechowski, K.; Stranks, S. D.; Abate, A.; Sadoughi, G.; Sadhanala, A.; Kopidakis,


30

Page 31 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


N.; Rumbles, G.; Li, C.; Friend, R. H.; Jen, A. K.; Snaith, H. J. Efficient OrganicInorganic Perovskite Solar Cells. ACS Nano 2014, 8 (12), 12701–12709. (43)

Jagadamma, L. K.; Al-Senani, M.; El-Labban, A.; Gereige, I.; Ngongang Ndjawa, G. O.; Faria, J. C. D.; Kim, T.; Zhao, K.; Cruciani, F.; Anjum, D. H.; McLachlan, M. A.; Beaujuge, P. M.; Amassian, A. Polymer Solar Cells with Efficiency >10% Enabled via a Facile Solution-Processed Al-Doped ZnO Electron Transporting Layer. Adv. Energy Mater. 2015, 5 (12), 1–12.

(44)

Yue, W.; Ashraf, R. S.; Nielsen, C. B.; Collado-Fregoso, E.; Niazi, M. R.; Yousaf, S. A.; Kirkus, M.; Chen, H. Y.; Amassian, A.; Durrant, J. R.; McCulloch, I. A Thieno[3,2b][1]benzothiophene Isoindigo Building Block for Additive- and Annealing-Free HighPerformance Polymer Solar Cells. Adv. Mater. 2015, 27 (32), 4702–4707.

(45) Liang, P.-W.; Chueh, C.-C.; Williams, S. T.; Jen, A. K.-Y. Roles of Fullerene-Based Interlayers in Enhancing the Performance of Organometal Perovskite Thin-Film Solar Cells. Adv. Energy Mater. 2015, 5 (10), 1402321.


31


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 32 of 32

ToC graphic:


32

Amorphous Tin Oxide as a Low-Temperature-Processed Electron

Recommend Documents