Observation of Nanoscale Morphological and Structural Degradation

Research Article www.acsami.org

Observation of Nanoscale Morphological and Structural Degradation in Perovskite Solar Cells by in Situ TEM Bin Yang,*,†,⊥ Ondrej Dyck,†,‡,⊥ Wenmei Ming,§,⊥ Mao-Hua Du,§ Sanjib Das,∥ Christopher M. Rouleau,† Gerd Duscher,‡ David B. Geohegan,† and Kai Xiao*,† †

Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States Department of Materials Science and Engineering, University of Tennessee, Knoxville, Tennessee 37996, United States § Materials Science & Technology Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States ∥ Department of Electrical Engineering and Computer Science, University of Tennessee, Knoxville, Tennessee 37996, United States ‡

S Supporting Information *

ABSTRACT: High-resolution in situ transmission electron microscopy (TEM) and electron energy loss spectroscopy were applied to systematically investigate morphological and structural degradation behaviors in perovskite films during different environmental exposure treatments. In situ TEM experiment indicates that vacuum itself is not likely to cause degradation in perovskites. In addition, these materials were found to degrade significantly when they were heated to ∼50− 60 °C (i.e., a solar cell’s field operating temperature) under illumination. This observation thus conveys a critically important message that the instability of perovskite solar cells at such a low temperature may limit their real field commercial applications. It was further unveiled that oxygen most likely attacks the CH3NH3+ organic moiety rather than the PbI6 component of perovskites during ambient air exposure at room temperature. This finding grants a deeper understanding of the perovskite degradation mechanism and suggests a way to prevent degradation of perovskites by tailoring the organic moiety component. KEYWORDS: perovskite solar cell, organometallic halide perovskite, degradation, transmission electron microscopy, electron energy loss spectroscopy

■

INTRODUCTION Solution-processed photovoltaic cells, based on organometallic trihalide perovskites (OTPs), have emerged as one of the most promising solar energy harvesters due to a power conversion efficiency (PCE) that is rapidly approaching 22%,1−7 which is comparable to that of existing thin film photovoltaic technologies based on polycrystalline silicon, cadmium telluride, or copper indium gallium selenide.8 Despite the significant progress in increasing device PCE, a major barrier facing the commercialization of this emerging technology is the chemical stability of OTP materials.8,9 It has been shown that these perovskites can degrade severely under environmental factors such as heat10−12 or moisture13,14 because of the weak van der Waals interactions between organic and inorganic components in the perovskite structure.15,16 To address this issue, much effort has been devoted to improving the material stability and associated device lifetime.17,18 For example, the charge transport layer has been modified to protect the perovskite photoactive layer against detrimental infiltration of moisture and oxygen.19 Specifically, it has been successfully demonstrated how device lifetimes can be improved by using charge transport layers with organic hydrophobic materials17 and dense metal oxides (ZnO and NiOx)18,19 that prevent the diffusion of moisture and oxygen molecules in ambient air into © 2016 American Chemical Society

the perovskite photoactive layer. Very recently, Grätzel and coworkers incorporated an additive into the perovskite precursors as a cross-link between neighboring grains in the perovskite structure to improve the surface functionality of perovskites and increase the stability of OTPs by utilizing hydrogen bonding between groups in the additives (PO(OH)2+, NH3+) and the iodide ions.20 However, further improvement in device lifetime urgently requires a fundamental understanding of the degradation mechanisms in OTP materials enabled by systematic investigations by multiple characterization techniques.13,21−23 For instance, in situ X-ray diffraction and UV−vis absorption spectroscopy revealed that the CH3NH 3PbI3 (MAPbI3) perovskites initially form a hydration product (CH3NH3)4PbI6·2H2O and ultimately decompose to PbI2 in ambient air.21 In addition, electroluminescence and photoluminescence imaging spectroscopy techniques were applied to spatially examine the degradation of optoelectronic properties of the perovskite film in ambient air, concluding that an applied bias could induce the degradation of the perovskite layer and the device in ambient air as suggested by the degraded Received: September 7, 2016 Accepted: November 4, 2016 Published: November 4, 2016 32333

DOI: 10.1021/acsami.6b11341 ACS Appl. Mater. Interfaces 2016, 8, 32333−32340

Research Article

ACS Applied Materials & Interfaces

maintained a PCE that was ∼86% of the starting efficiency (Figure 1a). The primary cause for the reduction in PCE is the decrease in fill factor (FF) of the device, which is most likely due to slight oxidation of the metal electrode in the N2-filled glovebox (O2 ≤ 20 ppm). However, the device exhibited a fast degradation when exposed to ambient air, and the PCE decreased to less than 0.1% in only 21 days. Figure S1 shows that the color of the device turned from dark brown to light yellow when stored in ambient air for ∼55 days, while similar devices remained dark brown when stored in the N2-filled glovebox. To understand morphological and structural degradation at the nanoscale in perovskite solar cells, high-resolution TEM imaging was used to examine a nonencapsulated device ex situ under air-exposure. As shown in Figure 2, high-angle annular

luminescence intensity.22 Although some inroads have been made on this topic, the degradation process of the microscopic structure and morphology of perovskite films is poorly understood. In this work, high-resolution transmission electron microscopy (TEM) imaging and electron energy loss spectroscopic (EELS) elemental mapping were applied to monitor the microstructural degradation process in MAPbI3 OTP films during different environmental exposures. Compared to energydispersive X-ray spectroscopy (EDS), EELS has better resolution to examine the spatial distribution of light elements such as carbon (C), nitrogen (N), and oxygen (O) in degraded perovskite films. In particular, both C and N are major components of organic moieties (e.g., CH3NH3+) of perovskite, while O is the key element that participates in the degradation process. It was found that O was present in the air-degraded perovskites, most likely attacking the CH3NH3+ organic moiety rather than the PbI6 inorganic part of perovskites during ambient air exposure. This finding presents a fundamental insight of the perovskite degradation mechanism and suggests a feasible route to prevent degradation of perovskites by tailoring the organic moiety component. Moreover, MAPbI3 perovskites were shown to be stable for an extended period of time (30 days) in the high vacuum TEM chamber, which enters a debate about the deleterious impact of vacuum conditions on the stability of the perovskites.

■

RESULTS AND DISCUSSION We first investigated the degradation of photovoltaic performance for nonencapsulated perovskite solar cells with a typical device architecture (ITO/TiO2/MAPbI3/Spiro-OMeTAD/Ag) that were fabricated as following our previous reported procedures.4 Figure 1 summarizes the degradation of photovoltaic parameters, which were obtained from devices stored in ambient air (∼30% humidity) and a nitrogen (N2)-filled glovebox (H2O ≤ 0.1 ppm and O2 ≤ 20 ppm). The devices stored in N2 degraded very slowly and, after 120 days,

Figure 2. (S)TEM images showing the variation of vertical phase morphology and material microstructure before and after air exposure in a typical perovskite solar cell with device structure of glass/ITO/ TiO2/CH3NH3PbI3/Spiro-OMeTAD/Ag. HAADF images acquired from the device before (a) and after (b) degradation in ambient air for 6 days to display the vertical phase morphology evolution. Highresolution TEM images (c) and (d) acquired from regions as indicated by red spots of the sample before (a) and after (b) air exposure for 6 days, respectively. Inset pictures of (c) and (d) show the corresponding FFT patterns.

dark-field (HAADF) cross-sectional scanning transmission electron microscope (STEM) images were acquired to display the changes of vertical phase morphology after air exposure. In the as-prepared device, the vertically aligned large grains and grain boundaries are seen (Figure 2a), which is consistent with our previous report.4 However, after the sample was exposed to ambient air for 6 days, the perovskite photoactive layer exhibited a severe change in morphology. Specifically, the film segregated into two major components as characterized by bright and dark areas in the perovskite layer (Figure 2b). In order to further examine structural changes of the perovskite material, corresponding high-resolution TEM images were acquired as shown in Figure 2c,d. In the as-prepared device, the perovskite layer exhibited high crystallinity and a well-ordered crystalline structure as evidenced by lattice fringes (Figure 2c) and a clear fast Fourier transform (FFT) pattern (inset of

Figure 1. Photovoltaic performance degradation of CH3NH3PbI3 perovskite solar cells in both nitrogen (N2)-filled glovebox (blue triangles) and ambient air (red squares): (a) normalized power conversion efficiency (PCE), (b) normalized short circuit current density (JSC), (c) normalized open circuit voltage (VOC), and (d) normalized fill factor (FF). 32334


Research Article


Figure 3. EELS mapping to reveal chemical elemental distribution along the vertical direction. A high-resolution HAADF image (a) and corresponding EELS mapping images to show the atomic areal densities of titanium (b), oxygen (c), and iodine (d) crossing the trilayer TiO2/ CH3NH3PbI3/Spiro-OMeTAD.

Figure 4. EELS mapping to reveal the formation of lead (Pb) particles after ambient air exposure. A high-resolution HAADF image (a) and corresponding EELS mapping images (acquired from green rectangular area in (a)) to show the relative concentration of lead (b), iodine (c), and carbon (d) crossing the trilayer TiO2/CH3NH3PbI3/Spiro-OMeTAD. (e, f) Example spectra showing the presence and absence of the Pb N45 EELS edge when the beam is located on and off a particle as marked in (b). (g) Complementary background subtracted Pb M edge acquired on a similar particle and from the surrounding perovskite. Note that the positions and shapes of the Pb particles in the recorded EELS spectra do not match well that of the HAADF image due to the extreme beam sensitivity of the Pb particles.

S2. However, sharp and well-defined rings associated with PbI2 were not observed, and likely what we observed was a mixture of polycrystalline PbI2, MAPbI3 perovskite, or other unidentified degradation products. To better elucidate the components of the bright region, EELS mapping was used to determine the elemental distribution in the degraded OTP layer. As shown in Figure 3a and Figure 3b−d, respectively, we acquired a HAADF image and corresponding EELS map while crossing the TiO2/ MAPbI3/Spiro-OMeTAD trilayer of the degraded device. Figure 3b−d shows the atomic areal densities of titanium (Ti), O, and iodine (I), respectively. As shown in Figure 3b,c, both Ti and O were mainly present in the TiO2 layer; however, interestingly, a small amount of O was also found in the MAPbI3 perovskite layer, demonstrating that O participated in the degradation process. The heavy element, I, mainly resided in the bright area (Figure 3d), which was most likely composed of PbI2,21,22,26 or other materials that contain I such as the hydrate product (CH3NH3)4PbI6·H2O,27 and result from the decomposition of MAPbI3 perovskite in ambient air.21 Since the dark region corresponds to light atoms, the presence of O

Figure 2c). We note that large and single crystalline grains were on the order of ∼50 nm in size. After air-exposure for 6 days, high-resolution TEM images of the similar area of the perovskite showed that the lattice fringes disappeared in the dark area (Figure 2d), and when combined with corresponding FFT patterns (inset of Figure 2d), it was confirmed that the dark area was mainly composed of amorphous material. In Zcontrast HAADF imaging, the brightness or contrast intensity (Y) increases with increasing the atomic number (Z) according to the relation Y∼Z1.7 in the imaging region,24,25 making it possible to correlate the bright and dark areas in Figure 2b with heavy and light elements, respectively. Thus, it is most likely that the hybrid perovskite decomposed into a relatively lighter amorphous material, such as the organic moiety CH3NH3+, within the dark area in Figure 2b. To identify the material structure within the brighter area, a selected area electron diffraction (SAED) pattern was taken as indicated by the red circle (Figure S2), and it showed lowquality polycrystalline material. The green curves are the expected location of the PbI2 polycrystalline ring pattern due to (002), (011), (012), and (013) reflections as shown in Figure 32335


Research Article


Figure 5. Real-time cross-sectional HAADF images of a typical perovskite solar cell with device structure of glass/ITO/TiO2/CH3NH3PbI3/SpiroOMeTAD/Ag. Day 1: bright particles at the surface of the perovskite are most likely due to FIB milling damage, which were beam-sensitive and disappeared on Day 2. Days 7−15: growth of an electron beam damaged area which occurred accidentally (yellow rectangular region). Days 20−30: greater care was taken to prevent the stationary beam from dwelling on the sample, and no further changes were observed. During the HAADF measurements, the sample remained in the TEM chamber with a high vacuum level of ∼10−7 Torr during the entire time (i.e., 30 days). After 4 h heating at ∼50−60 °C on the sample Day 30, degradation was observed from beam damage sites. After additional 10 h heating at ∼50−60 °C, the degradation became more severe.

sources. For example, it was recently proposed that vacuum conditions can facilitate the degradation of perovskites due to the pumping out of volatile species.29,30 To examine vacuum-induced degradation mechanisms, we specifically designed an in situ STEM imaging experiment so that we were able to watch the perovskite degradation and monitor structural and morphological changes in the material in high vacuum conditions (∼10−7 Torr). Figure 5 shows representative HAADF images acquired from the specimen, and a zoomed-out view of the place where in situ STEM images were taken is shown in Figure S4. Note that, during the HAADF measurements, the sample remained in the TEM chamber with a high vacuum level of ∼10−7 Torr during the entire time (i.e., 30 days). On Day 1, vertically aligned large grains and grain boundaries were observed in the perovskite layer. It should be noted that, near the top surface of the perovskite layer, we observed similar bright Pb particles to those mentioned previously, which, however, disappeared by Day 2. During the course of the first few days (up to Day 15) of observation, no attempt was made to prevent the focused electron beam from dwelling on the perovskite layer during microscope alignment. After observing that the perovskite was degrading in this area (and identifying that this was beaminduced), greater care was taken to prevent the stationary electron beam from dwelling on the perovskite layer, which reduced the beam exposure time on the perovskite from several seconds to microseconds (the pixel dwell time during imaging). Surprisingly, we did not observe any further change in the perovskite layer after ensuring limited beam exposure. This is illustrated in the images acquired from Day 15 to Day 30 of Figure 5. We did not observe the expansion of Pb particles or the aggregation of additional Pb particles, but instead, these particles disappeared during prolonged storage in high vacuum conditions. On the basis of these results, we were forced to conclude that vacuum itself is not likely to be responsible for

in the dark region of the perovskite layer indicates that O most likely attacks the organic species such as CH3NH3+. To the best of our knowledge, this represents the first microscopic observation of the element O in an air-degraded sample, as a further concrete evidence to support the hypothesis of O participation in the degradation process. In addition, we observed the formation of lead (Pb)-rich particles in an air-exposed sample as shown in Figure 4. The EELS maps (Figure 4b−d) were acquired to reveal the elemental composition of bright particles observed in the HAADF image in Figure 4a. The EELS image acquisition position is indicated by the green rectangle as shown in Figure 4a. The analyzed spectral maps shown in Figure 4b−d indicate the relative concentration of Pb, I, and C, respectively, suggesting that these bright particles are composed of Pb, rather than I and C. It should be noted that the positions and shapes of the particles in the recorded EELS spectra do not match those of the HAADF image very well due to extreme beam sensitivity of these particles. They moved easily under the beam and appeared to be reabsorbed into the perovskite as the electron beam added energy to the system as shown in Figure S3. To further identify the bright particles, two spectra were taken on and off a particle location as marked by squares in Figure 4b. It is clearly observed that the Pb N45 EELS edge shows up in the spectrum acquired on a bright particle (Figure 4f), whereas it disappears in the spectrum taken from the surrounding area (Figure 4e). As a second confirmation, the Pb M edge was also recorded and it was observed that there was a much higher concentration of Pb on the particle and a reduced concentration in the surrounding area (Figure 4g). The observation of Pb-rich particles as a degradation product is consistent with a recent photoelectron spectroscopy study,28 where a small Pb0 signal (137 eV) was observed after heating the perovskites. Note that it is still unclear whether the formation of Pb-rich particles is due to air exposure or other 32336


Research Article

ACS Applied Materials & Interfaces the degradation. At the current stage, it is still unclear about the origin of formation of Pb aggregates and we believe that further study is needed. In addition to moisture and oxygen, temperature is another non-negligible factor because solar cells normally operate at ∼50−60 °C31,32 in field applications. The structural and morphological degradation in perovskite films during heating at ∼50−60 °C would shed light on the understanding of device operation and lifetime under real-world conditions. Thus, we moved the stabilized specimen (Day 30 as shown above) to the air lock (moisture- and oxygen-free environment, filled with argon) in the same microscope where a built-in halogen lamp was used to provide heat. To avoid overheating the sample, we used a continuous flow of argon gas to vent the air lock and cool the sample to achieve a stable temperature of ∼50−60 °C as measured by an IR thermometer. After 4 h, the device exhibited a significant change in the observed intensity of HAADF images in the beam damaged area (indicated in the rectangular region). This observation suggests that degradation can indeed occur through a redistribution of the heavy and light atoms, which may be due to the elemental migration induced by heating as reported recently by Divitini and co-workers.26 Since the degraded area is primarily located in the beam damage region, this indicates that defects play a significant role in the degradation process. After heating for an additional 10 h, we observed more pronounced degradation around the beam damaged sites, as well as an obvious degradation of the entire perovskite layer. In order to rule out the possibility of excessive electron beam exposure on the particular imaging region shown in Figure 5 as the source of degradation, we also acquired images from a different area of the same specimen that had undergone significantly reduced beam exposure (Figure S5). Compared to the Day 2 image, the image that was acquired on Day 30 plus 14 h also exhibits an obvious degradation in the entire perovskite layer. Therefore, the degradation was not caused by excessive imaging of the same area, but instead resulted from a combined light/heat exposure at ∼50−60 °C. Since a halogen lamp was used to heat the specimen in the TEM air lock, a concern arises as to whether photoexcitation upon illumination could also trigger decomposition of MAPbI3 in a moisture- and oxygen-free environment. To gain more insight on the stability of MAPbI3 upon illumination, we calculated the enthalpy of formation for MAPbI3, ΔHf, with respect to PbI2 and CH3NH3I (MAI) based on density functional theory (DFT). ΔHf is given by ΔHf = EMAPbI3 − EMAI − E PbI2

Figure 6. Enthalpy of formation (ΔHf) for MAPbI3 with respect to MAI and PbI2 calculated using PBE and various vdW functionals. The black squares represent results obtained using experimental lattice parameters, while the red circles represent those based on theoretically relaxed lattice parameters. Lines are guides for the eye.

can be incentivized by lowering the energy of the excited-state electrons. The decomposition of MAPbI3 on the surface to PbI2 and MAI under an AM 1.5 global solar spectrum lowers the photoelectron energy by about 6 × 10−6 eV/f.u., which is too small to drive the decomposition of MAPbI3. Therefore, we can safely rule out the possibility of photoexcitation as a driving force for the decomposition of MAPbI3. Note that, during the experimental portion of our study, we did not separate the effects of light and heat. Our DFT calculations indicate that light-induced degradation should be unlikely; however, without strong experimental backing for purely heat-induced degradation, we cannot conclusively demonstrate that degradation is purely heat-induced. It is still possible that the combined effect of heat (50−60 °C) and light are driving such degradation.

■

CONCLUSIONS We have studied the microscopic degradation of MAPbI3 perovskites. We examined the effect of vacuum condition on the stability of perovskite. Pb particles were found to be formed initially as observed using in situ TEM; however, they disappeared by the second day. No further degradation was observed in the perovskite layer for 30 days at room temperature in high vacuum conditions aside from the initial electron beam damage, thus entering a debate about the deleterious impact of vacuum conditions on the stability of the perovskites. In addition, we showed the microscopic observation (by in situ TEM) of decomposition at 50−60 °C (under illumination), i.e., a solar cell’s field operating temperature. We also showed that the excited-state electrons and holes under solar irradiation are not the driving force for decomposition, while we are cautious to conclude that heat alone is responsible for the observed degradation without strong experimental backing to exclude the effect of light. Nevertheless, this observation conveys a critically important message that the instability of perovskite solar cells at such a low temperature under illumination may limit their real field commercial applications at the current stage. We revealed further that O most likely attacks the organic moiety CH3NH3+ of perovskite in ambient air, which is a new finding that may help to prevent degradation of perovskite by tailoring its molecular structure such as organic moiety. While these discoveries enhance the understanding of the degradation mechanisms of MAPbI3 perovskites, further study is required to understand the cause of Pb aggregations and additional in situ heating experiments

(1)

where EMAPbI3, EMAI, and EPbI2 are the total energy (per formula unit) of MAPbI3, MAI, and PbI2, respectively. As shown in Figure 6, our calculated ΔHf are generally positive regardless of the choice of lattice parameters and functionals (except the TSSCS functional) used in the calculations (see the Supporting Information for detailed calculations). These results support the proposition that MAPbI3 is unstable against decomposition into PbI2 and MAI.23,33 Although the effect of entropy is not considered in the present calculation, Zhang et al. showed that the contribution of the entropy should not reverse the instability predicted based on the calculated enthalpy of formation.23 We further compared the energies of photogenerated excited-state electrons in MAPbI3, MAI, and PbI2 to determine if the decomposition of MAPbI3 under illumination 32337


Research Article


Since it is still unclear what causes the formation of the Pb aggregates observed, precise details of sample handling are provided here. Particularly, we note the detailed experimental procedure for Figure 4. This particular sample was prepared via FIB milling as described previously and initially imaged in the Zeiss Libra 200MC TEM. During this initial TEM examination, Pb aggregates were not observed. The sample was then transferred from the Libra TEM to the VG HB501 STEM, exposing it to ambient air for a few minutes (∼5− 10 min) during the sample exchange. After introducing the sample back into vacuum, it was subsequently left in the STEM column overnight before being imaged again the following morning. The data presented in Figure 4 were acquired at this time, which shows Pb aggregates had formed within the perovskite photoactive layer. Electron Energy Loss Spectroscopy. The electron energy loss spectroscopy (EELS) images were acquired (Zeiss Libra 200MC) with a beam current of 43 pA, as measured with a calibrated CCD in vacuum, with an exposure time of 1 s per pixel, and a pixel size of 11.1 × 11.1 nm2, and sub-pixel scanning on. The convergence angle was 9 mrad, and the collection angle was 65 mrad. Spectra were backgroundsubtracted and fit with a Hartree−Slater model scattering cross section. The acquired and analyzed edges were the oxygen K, titanium L, iodine M, as well as lead M and N edges. With a known beam current and calibrated spectrometer, the fitting parameter could be interpreted as atomic areal density, and a quantitative value was obtained within ∼10%. Density Functional Theory Calculations. DFT calculations were conducted with the projector augmented-wave pseudopotential.34 Various functionals were adopted for approximating the exchange-correlation potential, including PBE35 and added-in vdW corrections/functions: D2 method of Grimme;36 D3 method of Grimme;37 D3 method with Becke−Jonson damping;38 the method of Tkatchenko and Scheffler;39 the method of Tkatchenko and Scheffler with self-consistent screening;40 vdW functional by Dion;41 vdWDF2;42 optPBE;43 optB88;43 optB86b.43 An energy cutoff of 500 eV and reciprocal space sampling mesh with a k-spacing of 0.3/Å were used for all calculations.

will be required to separate the degradation effects of light and heat.

■

EXPERIMENTAL SECTION

Device Fabrication and Characterization. Device fabrication followed previously reported procedures.4 A titanium dioxide (TiO2) precursor solution was spin-coated onto UV-ozone treated indium tin oxide (ITO) glass substrates at 2000 rpm for 60 s in ambient air. To increase the crystallinity of TiO2, the TiO2 films were sintered at 500 °C for 30 min in a furnace. A PbI2 solution, which was dissolved in dimethylformamide (DMF) with a concentration of 550 mg/mL, was then spin-coated onto the TiO2 layer at 6000 rpm for 30 s in an N2filled glovebox. When the PbI2 films were completely dry, methylammonium iodide (CH3NH3I) solution (70 mg/mL, isopropanol as solvent) was spin-coated onto the PbI2 layer at 6000 rpm for 30 s. After ambient air exposure for 60 min, the samples were annealed at 100 °C for 2 h in an N2-filled glovebox. Spiro-OMeTAD solution (90 mg/mL) was spin-coated at 2000 rpm for 40 s, and then left in a desiccator overnight. Finally, a silver layer with a thickness of 100 nm was thermally evaporated at a vacuum level of 10−6 mbar and a deposition rate of 1 Å/s. The device active area was ∼6.5 mm2, which was carefully measured using optical microscopy. The photovoltaic performance was characterized in an N2-filled glovebox. A Keithley 2400 source meter was used to acquire the J−V curves of the devices under illumination of 100 mW/cm2 (AM 1.5 G solar spectrum). The illumination was provided by a solar simulator (Radiant Source Technology, 300 W, Class A). The lamp light intensity was calibrated with a National Institute of Standards and Technology (NIST) certified Si-reference cell. The J−V curves were measured by sweeping from reverse bias (−0.2 V) to forward bias (1.2 V), and then forward bias (1.2 V) to reverse bias (−0.2 V), with a 35 mV step size, and 50 ms sweep time. Due to the observed J−V hysteresis from both forward and backward scan directions, the power conversion efficiencies that were used to assess the device performance degradation process were calculated using the J−V curves from the forward direction scan (scanning from forward bias to reverse bias). Note that we soaked the devices under 1 sun illumination for ∼60 s near the maximum power output point (to stabilize device performance) before each J−V curve scan. Transmission Electron Microscope. Transmission electron microscopy (TEM) samples were prepared via focused ion beam (FIB) polishing in a Zeiss Auriga CrossBeam dual FIB SEM microscope. The bulk sample was transferred from the fabrication N2-filled glovebox to the FIB station via a sealed stainless steel sample transfer tube immediately after fabrication to ensure minimal exposure to air. Once the transfer tube was opened, the sample was sputtercoated with gold and loaded into the FIB chamber in less than 5 min. Sample extraction was performed with an FIB accelerating voltage of 30 kV and current of 2 nA. Once the lamella was extracted and welded to a copper TEM half-grid, it was thinned to electron transparency with an FIB accelerating voltage of 5 kV and a current of 20 pA. The SEM accelerating voltage was kept at 3 kV for the entire process. The final thicknesses of TEM samples were typically in the 40−60 nm range. Once the TEM sample was finished, it was immediately loaded into the STEM with, again, less than 5 min of exposure to air. For the images shown in Figures 2, 3, and S2, a Zeiss Libra 200MC transmission electron microscope (TEM) was used. For Figures 4, 5, and S3−S5, the microscope used was a VG HB501 third-order aberration-corrected dedicated scanning transmission electron microscope (STEM). The accelerating voltage was 100 kV. Images were acquired each day on the same area of the specimen at the same magnification. The probe was switched on only for alignment and subsequent imaging (typically two images per day, though on Day 1 and Day 30, for example, more than two images were acquired), after which it was positioned off the sample and turned off. The sample was not removed from the microscope vacuum until after Day 30 and was not exposed to air at any time except during FIB preparation as mentioned previously.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b11341. Additional pictures of the real degraded devices, highresolution TEM images, and the selected area electron diffraction (SAED) patterns are included in Figures S1− S5 to show the degradation of MAPbI3. The detailed calculation of the enthalpy of formation of both groundstate and excited-state electronic energy of MAPbI3, Figure S6, and Table S1 are also provided (PDF)

■

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (B.Y.). *E-mail: [email protected] (K.X.). ORCID

Mao-Hua Du: 0000-0001-8796-167X Kai Xiao: 0000-0002-0402-8276 Author Contributions ⊥

These authors contributed equally to this work.

Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS Part of this research was conducted at the Center for Nanophase Materials Sciences (CNMS), which is a DOE Office of Science User Facility. O.D. and G.D. are thankful for 32338


Research Article

ACS Applied Materials & Interfaces the financial support from the Sustainable Energy and Education Research Center (SEERC), Tennessee Solar Conversion and Storage using Outreach, Research and Education (TN-SCORE), the Department of Energy, Basic Energy Sciences (DOE-BES). M.-H.D. and W.M. are supported by the Department of Energy, Office of Science, Basic Energy Sciences, Materials Sciences and Engineering Division.

■

conductivity Response of CH3NH3PbI3 Perovskites Suggesting Structural Changes under Working Conditions. J. Phys. Chem. Lett. 2014, 5, 2662−2669. (17) Hwang, I.; Jeong, I.; Lee, J.; Ko, M. J.; Yong, K. Enhancing Stability of Perovskite Solar Cells to Moisture by the Facile Hydrophobic Passivation. ACS Appl. Mater. Interfaces 2015, 7, 17330−17336. (18) You, J.; Meng, L.; Song, T.-B.; Guo, T.-F.; Yang, Y. M.; Chang, W.-H.; Hong, Z.; Chen, H.; Zhou, H.; Chen, Q.; Liu, Y.; De Marco, N.; Yang, Y. Improved Air Stability of Perovskite Solar Cells via Solution-Processed Metal Oxide Transport Layers. Nat. Nanotechnol. 2016, 11, 75−81. (19) Chen, W.; Wu, Y.; Yue, Y.; Liu, J.; Zhang, W.; Yang, X.; Chen, H.; Bi, E.; Ashraful, I.; Grätzel, M.; Han, L. Efficient and Stable LargeArea Perovskite Solar Cells with Inorganic Charge Extraction Layers. Science 2015, 350, 944−948. (20) Li, X.; Dar, M. I.; Yi, C.; Luo, J.; Tschumi, M.; Zakeeruddin, S. M.; Nazeeruddin, M. K.; Han, H.; Grätzel, M. Improved Performance and Stability of Perovskite Solar Cells by Crystal Crosslinking with Alkylphosphonic Acid ω-ammonium Chlorides. Nat. Chem. 2015, 7, 703−711. (21) Yang, J.; Siempelkamp, B. D.; Liu, D.; Kelly, T. L. Investigation of CH3NH3PbI3 Degradation Rates and Mechanisms in Controlled Humidity Environments Using In Situ Techniques. ACS Nano 2015, 9, 1955−1963. (22) Okano, M.; Endo, M.; Wakamiya, A.; Yoshita, M.; Akiyama, H.; Kanemitsu, Y. Degradation Mechanism of Perovskite CH3NH3PbI3 Diode Devices Studied by Electroluminescence and Photoluminescence Imaging Spectroscopy. Appl. Phys. Express 2015, 8, 102302. (23) Zhang, Y.-Y.; Chen, S.; Xu, P.; Xiang, H.; Gong, X.-G.; Walsh, A.; Wei, S.-H. Intrinsic Instability of the Hybrid Halide Perovskite Semiconductor CH3NH3PbI3. 2015, arXiv:1506.01301. arXiv preprint. https://arxiv.org/abs/1506.01301. (24) Krivanek, O. L.; Chisholm, M. F.; Nicolosi, V.; Pennycook, T. J.; Corbin, G. J.; Dellby, N.; Murfitt, M. F.; Szilagyi, Z. S.; Oxley, M. P.; Pantelides, S. T.; Own, C. S.; Pennycook, S. J. Atom-by-Atom Structural and Chemical Analysis by Annular Dark-Field Electron Microscopy. Nature 2010, 464, 571−574. (25) Hartel, P.; Rose, H.; Dinges, C. Conditions and Reasons for Incoherent Imaging in STEM. Ultramicroscopy 1996, 63, 93−114. (26) Divitini, G.; Cacovich, S.; Matteocci, F.; Cinà, L.; Di Carlo, A.; Ducati, C. In situ Observation of Heat-Induced Degradation of Perovskite Solar Cells. Nat. Energy. 2016, 1, 15012. (27) Christians, J. A.; Miranda Herrera, P. A.; Kamat, P. V. Transformation of the Excited State and Photovoltaic Efficiency of CH3NH3PbI3 Perovskite upon Controlled Exposure to Humidified Air. J. Am. Chem. Soc. 2015, 137, 1530−1538. (28) Philippe, B.; Park, B.-W.; Lindblad, R.; Oscarsson, J.; Ahmadi, S.; Johansson, E. M.; Rensmo, H. k. Chemical and Electronic Structure Characterization of Lead Halide Perovskites and Stability Behavior under Different Exposures: A Photoelectron Spectroscopy Investigation. Chem. Mater. 2015, 27, 1720−1731. (29) Alberti, A.; Deretzis, I.; Pellegrino, G.; Bongiorno, C.; Smecca, E.; Mannino, G.; Giannazzo, F.; Condorelli, G. G.; Sakai, N.; Miyasaka, T.; Spinella, C.; La Magna, A. Similar Structural Dynamics for the Degradation of CH3NH3PbI3 in Air and in Vacuum. ChemPhysChem 2015, 16, 3064−3071. (30) Smecca, E.; Numata, Y.; Deretzis, I.; Pellegrino, G.; Boninelli, S.; Miyasaka, T.; La Magna, A.; Alberti, A. Stability of Solution-Processed MAPbI3 and FAPbI3 layers. Phys. Chem. Chem. Phys. 2016, 18, 13413− 13422. (31) Raman, A. P.; Anoma, M. A.; Zhu, L.; Rephaeli, E.; Fan, S. Passive Radiative Cooling below Ambient Air Temperature under Direct Sunlight. Nature 2014, 515, 540−544. (32) Pathak, M.; Girotra, K.; Harrison, S.; Pearce, J. M. The Effect of Hybrid Photovoltaic Thermal Device Operating Conditions on Intrinsic Layer Thickness Optimization of Hydrogenated Amorphous Silicon Solar Cells. Sol. Energy 2012, 86, 2673−2677.

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, 316−319. (2) Im, J.-H.; Jang, I.-H.; Pellet, N.; Grätzel, M.; Park, N.-G. Growth of CH3NH3PbI3 Cuboids with Controlled Size for High-Efficiency Perovskite Solar Cells. Nat. Nanotechnol. 2014, 9, 927−932. (3) 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. (4) Yang, B.; Dyck, O.; Poplawsky, J.; Keum, J.; Puretzky, A.; Das, S.; Ivanov, I.; Rouleau, C.; Duscher, G.; Geohegan, D.; Xiao, K. Perovskite Solar Cells with Near 100% Internal Quantum Efficiency Based on Large Single Crystalline Grains and Vertical Bulk Heterojunctions. J. Am. Chem. Soc. 2015, 137, 9210−9213. (5) 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. (6) Das, S.; Yang, B.; Gu, G.; Joshi, P. C.; Ivanov, I. N.; Rouleau, C. M.; Aytug, T.; Geohegan, D. B.; Xiao, K. High-Performance Flexible Perovskite Solar Cells by Using a Combination of Ultrasonic SprayCoating and Low Thermal Budget Photonic Curing. ACS Photonics 2015, 2, 680−686. (7) Saliba, M.; Matsui, T.; Seo, J.-Y.; Domanski, K.; Correa-Baena, J.P.; Nazeeruddin, M. K.; Zakeeruddin, S. M.; Tress, W.; Abate, A.; Hagfeldt, A.; Gratzel, M. Cesium-Containing Triple Cation Perovskite Solar Cells: Improved Stability, Reproducibility and High Efficiency. Energy Environ. Sci. 2016, 9, 1989−1997. (8) McGehee, M. D. Materials Science: Fast-Track Solar Cells. Nature 2013, 501, 323−325. (9) Stranks, S. D.; Snaith, H. J. Metal-Halide Perovskites for Photovoltaic and Light-Emitting Devices. Nat. Nanotechnol. 2015, 10, 391−402. (10) Misra, R. K.; Aharon, S.; Li, B.; Mogilyansky, D.; Visoly-Fisher, I.; Etgar, L.; Katz, E. A. Temperature-and Component-Dependent Degradation of Perovskite Photovoltaic Materials under Concentrated Sunlight. J. Phys. Chem. Lett. 2015, 6, 326−330. (11) Yang, B.; Dyck, O.; Poplawsky, J.; Keum, J.; Das, S.; Puretzky, A.; Aytug, T.; Joshi, P. C.; Rouleau, C. M.; Duscher, G.; Geohegan, D. B.; Xiao, K. Controllable Growth of Perovskite Films by RoomTemperature Air Exposure for Efficient Planar Heterojunction Photovoltaic Cells. Angew. Chem., Int. Ed. 2015, 54, 14862−14865. (12) 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, 4229−4236. (13) Niu, G.; Li, W.; Meng, F.; Wang, L.; Dong, H.; Qiu, Y. Study on the Stability of CH3 NH3PbI3 Films and the Effect of Postmodification by Aluminum Oxide in All-Solid-State Hybrid Solar Cells. J. Mater. Chem. A 2014, 2, 705−710. (14) Mosconi, E.; Azpiroz, J. M.; De Angelis, F. Ab Initio Molecular Dynamics Simulations of Methylammonium Lead Iodide Perovskite Degradation by Water. Chem. Mater. 2015, 27, 4885−4892. (15) Yin, W.-J.; Shi, T.; Yan, Y. Unusual Defect Physics in CH3NH3PbI3 Perovskite Solar Cell Absorber. Appl. Phys. Lett. 2014, 104, 063903. (16) Gottesman, R.; Haltzi, E.; Gouda, L.; Tirosh, S.; Bouhadana, Y.; Zaban, A.; Mosconi, E.; De Angelis, F. Extremely Slow Photo32339


Research Article

ACS Applied Materials & Interfaces (33) Ganose, A. M.; Savory, C. N.; Scanlon, D. O. CH3NH3)2Pb(SCN)2I2: A More Stable Structural Motif for Hybrid Halide Photovoltaics? J. Phys. Chem. Lett. 2015, 6, 4594−4598. (34) Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector Augmented-Wave Method. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 59, 1758−1775. (35) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865−3868. (36) Grimme, S. Semiempirical GGA-Type Density Functional Constructed with a Long-Range Dispersion Correction. J. Comput. Chem. 2006, 27, 1787−1799. (37) Grimme, S.; Antony, J.; Ehrlich, S.; Krieg, H. A Consistent and Accurate ab initio Parametrization of Density Functional Dispersion Correction (DFT-D) for the 94 Elements H-Pu. J. Chem. Phys. 2010, 132, 154104. (38) Grimme, S.; Ehrlich, S.; Goerigk, L. Effect of the Damping Function in Dispersion Corrected Density Functional Theory. J. Comput. Chem. 2011, 32, 1456−1465. (39) Tkatchenko, A.; Scheffler, M. Accurate Molecular van der Waals Interactions from Ground-State Electron Density and Free-Atom Reference Data. Phys. Rev. Lett. 2009, 102, 073005. (40) Tkatchenko, A.; DiStasio, R. A., Jr; Car, R.; Scheffler, M. Accurate and Efficient Method for Many-Body van der Waals Interactions. Phys. Rev. Lett. 2012, 108, 236402. (41) Dion, M.; Rydberg, H.; Schröder, E.; Langreth, D. C.; Lundqvist, B. I. Van der Waals Density Functional for General Geometries. Phys. Rev. Lett. 2004, 92, 246401. (42) Lee, K.; Murray, É. D.; Kong, L.; Lundqvist, B. I.; Langreth, D. C. Higher-Accuracy van der Waals Density Functional. Phys. Rev. B: Condens. Matter Mater. Phys. 2010, 82, 081101. (43) Klimes, J.; Bowler, D. R.; Michaelides, A. Chemical Accuracy for the van der Waals Density Functional. J. Phys.: Condens. Matter 2010, 22, 022201.

32340


Observation of Nanoscale Morphological and Structural Degradation

Recommend Documents