Letter pubs.acs.org/JPCL
Ferroelectric Polarization in CH3NH3PbI3 Perovskite Hui-Seon Kim,†,⊥ Sung Kyun Kim,‡,⊥ Byeong Jo Kim,‡,⊥ Kyung-Sik Shin,‡ Manoj Kumar Gupta,‡ Hyun Suk Jung,‡ Sang-Woo Kim,*,‡,§ and Nam-Gyu Park*,† †
School of Chemical Engineering and Department of Energy Science, ‡School of Advanced Materials Science and Engineering, and SKKU Advanced Institute of Nanotechnology, Sungkyunkwan University, Suwon 440-746, Korea
J. Phys. Chem. Lett. 2015.6:1729-1735. Downloaded from pubs.acs.org by UNIV OF SUNDERLAND on 11/07/18. For personal use only.
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ABSTRACT: We report on ferroelectric polarization behavior in CH3NH3PbI3 perovskite in the dark and under illumination. Perovskite crystals with three different sizes of 700, 400, and 100 nm were prepared for piezoresponse force microscopy (PFM) measurements. PFM results confirmed the formation of spontaneous polarization in CH3NH3PbI3 in the absence of electric field, where the size dependency to polarization was not significant. Whereas the photoinduced stimulation was not significant without an external electric field, the stimulated polarization by poling was further enhanced under illumination. The retention of ferroelectric polarization was also observed after removal of the electric field, in which larger crystals showed longer retention behavior compared to the smaller sized one. Additionally, we suggest the effect of perovskite crystal size (morphology) on charge collection at the interface of the ferroelectric material even though insignificant size dependency in electric polarization was observed.
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ferroelectricity was measured from MAPbI3, but the ion migration was suggested to interpret the phenomena.10,19 This means that the argument on ferroelectricity in MAPbI3 has not been terminated with clear conclusion, which motivates us to investigate the ferroelectric behavior of MAPbI3. We report here on ferroelectric polarization behavior in MAPbI3 by a piezoresponse force microscopy (PFM) investigation. We measured the electric polarization change of the samples with different crystal size with and without an electric field in the dark and under illumination. In addition, polarization retention characteristics after removal of the external electric field were investigated. All experiments were designed similar to the practical solar cell operation including the morphology of MAPbI3 and the range of the external electric field in order to closely correlate with the solar cell characteristics and give insight into the working principle of perovskite solar cells. Photovoltaic performance was previously found to be significantly altered by the crystal size of MAPbI3,20 where a large crystal was beneficial to light harvesting and charge extraction compared to a small one. Moreover, a large crystal was less sensitive to current−voltage (I−V) hysteresis than the small one.12 Thus, perovskite layers with different crystal sizes were selected as samples for PFM study in this work in order to correlate the PFM results with the studied photovoltaic properties. The crystal size was controlled by varying the concentration of the CH3NH3I (MAI) solution in the two-step spin-coating deposition method, as described elsewhere,20
ince the reports on perovskite solar cells based on the mesoscopic sensitization concept1 and the mesosuperstructured extremely thin layer concept2 in 2012, organometal halide perovskite has been considered as a promising light harvester for high-efficiency solar cells. The power conversion efficiency (PCE) of perovskite solar cells surged to over 15% in 2013 by modifying a perovskite deposition method from a conventional one-step method to two-step sequential deposition3 and vapor evaporation process.4 So far, the highest PEC of 19.3% was reported using a planar structure,5 and 20.1% was certified by the National Renewable Energy Laboratory.6 The outstanding photovoltaic performance from perovskite solar cells has urged researchers to investigate the fundamental properties of perovskite. Consequently, balanced charge transport7,8 and charge accumulation9 properties of organolead iodide perovskite were unraveled. Besides light harvesting and balanced charge transport properties, unusual properties were observed with CH3NH3PbI3 (MAPbI3), such as the ambipolar self-doping property,10 high dielectric constant,11 current−voltage hysteresis,12 and slow dynamics.13 Polarization was suggested as one of the bases for the origin of these phenomena. In ABO3 oxide perovskite materials, ferroelectricity along with high dielectric constants was observed.14 In addition, a phase transition was accompanied in ferroelectric oxide perovskite materials. Thus, anomalous properties observed for MAPbI3 are likely to be related to the underlying ferroelectric property because of its high dielectric constant and cubic−tetragonal phase transition. Regarding the issue on ferroelectricity in MAPbI3, theoretical calculations suggested the possibility of ferroelectricity of MAPbI3.15−17 Ferroelectric domains were directly observed from the solution-processed MAPbI3.18 Interestingly, no © 2015 American Chemical Society
Received: April 3, 2015 Accepted: April 21, 2015 Published: April 21, 2015 1729
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Figure 1. Scanning electron microscopy (SEM) images for the MAPbI3 crystals with average cuboid sizes of (a) 700, (b) 400, and (c) 100 nm prepared using 0.038, 0.044, and 0.063 M MAI, respectively. Scale bars represent 500 nm. XRD patterns of (d) 700, (e) 400, and (f) 100 nm sized MAPbI3 cuboids directly formed on a fluorine-doped tin oxide (FTO) substrate.
atomic force microscopy (AFM) and PFM. It is well-known that PFM is a powerful tool that detects polarization states based on the detection of the local piezoelectric deformation induced by an external electric field in a ferroelectric sample. The topography image of the MAPbI3 was built by measuring the z-axis movement of the AFM tip in a contact mode. For poling the sample, we applied a DC bias voltage between the MAPbI3 cuboid and the conductive tip. During the PFM measurement, an AC voltage was applied from the conductive tip as the top electrode to the sample’s surface. The amplitude and the phase signals of the response of the MAPbI3 material to the applied electrical signal were detected using the vertical displacement of the AFM tip. To determine the ferroelectric polarization of the MAPbI3 perovskite material, positive and negative poling processes were sequentially applied, and the PFM phase was measured both in the dark and under illumination. Figure 3 shows the PFM phase along with the topography observed in the dark, where no significant deformation of the perovskite sample is assumed throughout the successive poling process due to the sequential distinctive image of the PFM phase. The topographic image of the 700 nm sized perovskite sample and the corresponding PFM phase images obtained without bias voltage (unpoled) and with bias voltages of +3 V (positively poled) and −3 V (negatively poled) are depicted in Figure 3a−d, respectively. Figure 3a shows some extent of the PFM phase contrast even in the absence of an electric field (unpoled), indicative of spontaneous polarization. Unlike the change in displacement of the B-site cation in oxide perovskites,21,22 the polarization behavior in organolead iodide perovskite is likely due to large freedom of rotation of the polar CH3NH3+ cation.15 When positive poling (+3 V) is applied, stimulated polarization in the positive direction is distinctively observed, as shown in Figure 3c. Figure 3d shows the PFM image under the negative poling process. Interestingly the opposite electric field (−3 V) is incapable of rotating dipoles 180°, which implies that the dipoles cannot easily rotate by the applied electric field due to the interaction with defects in crystals.23 The spontaneous and stimulated polarizations are also observed for the 400 nm and the 100 nm sized cuboids. As can be seen in Figure 3f−h and j−l, PFM phase contrast is enhanced upon positive poling and relieved with negative poling, but the dependence of polar-
where 0.038, 0.044, and 0.063 M MAI produce average MAPbI3 crystal sizes of about 700, 400, and 100 nm, respectively, as shown in Figure 1a−c. Compared to the compact nature of the 100 nm sized cuboids, the number of gaps appears and increases upon increasing size, which is consistent with the previous results.20 No structural change was observed depending on the crystal size, as confirmed by X-ray diffraction (XRD) patterns in Figure 1d−f, which shows that the chemical composition is not modified by crystal size. Because there is no significant difference in XRD peaks among three samples with different morphologies, we expect that the grain sizes in the samples with polycrystalline nature are comparable irrespective of the different crystal size. In addition, the prepared MAPbI3 shows phase purity without unreacted PbI2. The piezoresponse measurements were conducted using a PFM setup, as illustrated in Figure 2. The detailed measurement condition is described in the Methods section. To confirm the ferroelectric characteristics of the MAPbI3, we measured the topography and the piezoresponse of the MAPbI3 on the fluorine-doped tin oxide (FTO)/glass substrate using
Figure 2. Schematic illustration of the experimental setup for PFM study of the MAPbI3. 1730
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Figure 3. AFM topographic images and PFM phase images measured in the dark for the MAPbI3 cuboids with size of (a−d) 700, (e−h) 400, and (i−l) 100 nm. Scale bars in (a−h) represent 1 μm, and those in (i−l) represent 500 nm. Measured areas were 5 μm × 5 μm for 700 and 400 nm sized samples and 2 μm × 2 μm for the 100 nm sized sample.
Figure 4. AFM topographic images and PFM phase images measured under illumination (50k lux) for the MAPbI3 cuboids with a size of (a−d) 700, (e−h) 400, and (i−l) 100 nm. Scale bars in (a−h) represent 1 μm, and those in (i−l) represent 500 nm. Measured areas were 5 μm × 5 μm for 700 and 400 nm sized samples and 2 μm × 2 μm for the 100 nm sized sample.
ization on the crystal size seems to be not significant. It is suggested that the negligible size dependency originates from
the polycrystalline nature of MAPbI3 composed of multiple grains. Thus, we conclude that grain size is a critical factor to 1731
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Figure 5. PFM phase images under illumination for (a) 700, (b) 400, and (c) 100 nm MAPbI3 cuboids, measured at positive poling (0 min) and at 30 and 60 min after removal of poling (+3 V). Scale bars represent 1 μm for 700 and 400 nm sized MAPbI3 and 500 nm for 100 nm sized MAPbI3. Measured areas were 5 μm × 5 μm for 700 and 400 nm sized samples and 2 μm × 2 μm for the 100 nm sized sample.
PFM phase of positively poled crystals is measured at an interval of 30 min under illumination (Figure 5). Highly aligned dipoles in 700 nm sized crystals are maintained for 60 min, as shown in Figure 5a, while dipoles in the 400 and 100 nm sized MAPbI3 are retained only for 30 min as shown in Figure 5b and c, respectively. It is obvious from our observation that the retention of the light-induced polarization is dependent on MAPbI3 crystal size, in which the 700 nm sized large crystal shows better retention behavior than smaller ones. Generally, a large crystal is assumed to have a lower defect density compared to the smaller crystal where significant defects hinder the retention of polarization.24 In this study, our results reveal ferroelectric characteristics of MAPbI3 and its dependence on light. Last, we correlate photoferroic behavior with photovoltaic performance and suggest a carrier collection mechanism in the ferroelectric MAPbI3 material. A forward bias voltage more than 1 V is applied to the real device while measuring the photocurrent− voltage curve. Poling at 3 V in this study is acceptable for correlation with real device performance because the actual extent of applied external electric field between the tip and plane in the PFM measurement is two or three times weaker than the case between the plane and plane in solar cell measurements.18 Even though the ferroelectric property is comparable regardless of crystal size, excluding retention behavior, photovoltaic performance is highly dependent on crystal size.12,20 Here, we suggest a charge collection mechanism related to the effect of morphology of the ferroelectric MAPbI3 material. Figure 6 illustrates the carrier pathway depending on the perovskite crystal morphology. Photoinduced stimulated polar-
realize effective manipulation of ferroelectric polarization rather than crystal morphology. However, careful interpretation is still needed considering the fact that the perovskite crystals with different sizes of 700, 400, and 100 nm simultaneously generate different thicknesses of 360, 290, and 170 nm, respectively. Because few studies of the MAPbI3 ferroelectric behavior on thickness effect have been conducted, little information was given to explain the impact of thickness around hundreds of nanometers. Light-induced polarization of the perovskite cuboids was also investigated with respect to the crystal size and poling process using the PFM measurement system equipped with a white light source (50 000 lux). Figure 4 shows topographic images and the corresponding PFM phase images under illumination depending on the perovskite polycrystal size. For the case of unpoled samples (Figure 4b,f,j), it is found that the spontaneous polarization is not enhanced by light irradiation compared to the unpoled samples in the dark (Figure 3b,f,j). The 700 nm sized large sample shows notably attenuated polarization even under illumination (Figure 4b) compared to that in the dark (Figure 3b), which indicates that the ferroelectric polarization is efficiently screened by photogenerated conducting electrons.22 However, photoinduced stimulated polarization is clearly distinct in the presence of the electric field, especially positive poling (Figure 4c,g,k) compared to those in the dark (Figure 3c,g,k). Upon negative poling, similar tendency is observed showing weakened polarization (Figure 4d,h,i) compared to the positively poled samples (Figure 4c,g,k). Retention of ferroelectric polarization is investigated as a function of time after removal of the external electric field. The 1732
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presence of a mesoporous TiO2 layer where the ferroelectric domain could be limited by pore size. In terms of light effect, they also found increased polarization upon illumination, which was confirmed by the PFM amplitude. Xiao et al. recently reported that ion drift occurs during a poling process, and they could not find any trace of ferroelectric characteristics from their PFM measurements even though they used a sequential two-step deposition method (the same method used in this study) to prepare the MAPbI3 perovskite film.10 The difference might come from the crystal quality or/and significant crystal size difference because much smaller crystal size would be expected due to the high concentration of the 45 mg/mL MAI solution.10 It was mentioned that the ion drift by poling was negligible in the large grain (600−1000 nm) owing to the reduced vacancy concentration based on the fewer grain boundaries. In conclusion, MAPbI3 was found to exhibit ferroelectric characteristics according to the PFM study. Spontaneous polarization behavior was commonly found in MAPbI3 where the size dependency was not significant. The electric polarization was further stimulated by positive poling and released by negative poling, in which 180° dipole rotation by the opposite electric field is inhibited by defects in crystals. Photoinduced polarization is negligible in the absence of an electric field but evident in the presence of an electric field showing strong polarization. Photoinduced polarization with positive poling is basically beneficial to charge transport and collection especially for the large crystals by providing an efficient electron pathway. Polarization was remained after removal of the electric field, and retention of polarization was longer in the large crystal than smaller ones. The observed ferroelectric behavior in organolead iodide perovskite is expected to give crucial insight into the design and working mechanism of perovskite solar cells.
Figure 6. Expected charge collection mechanism with the carrier pathway depending on the MAPbI3 crystal morphology under illumination and forward bias condition.
ization is evident in MAPbI3 irrespective of crystal size. It is theoretically found that the well-aligned dipoles in MAPbI3 have a substantial effect on the PbI3 inorganic cage and lead to an enhanced current response.25 In addition, at positive poling (forward bias direction in an actual I−V measurement), the external electric field will be essentially weakened by a generated internal electric field due to strong polarization, which is good for separation and collection of photogenerated electrons toward an interface of the perovskite and selective contact. However, transported electrons are easy to be trapped at the interface and thus hardly injected into the selective contact.26 Whereas densely populated small crystals cause inefficient electron injection due to the higher density of trapped electrons (Figure 6b), sparsely populated large crystals sufficiently provide an efficient pathway for electrons along the edge (Figure 6a). Then, the sparse morphology with large crystals is beneficial for charge collection, leading to high photocurrent density. This result is well in accordance with the previous report, where the perovskite solar cells incorporating large crystals of about 700 and 400 nm exhibit higher amounts of the extracted charges compared to the smaller cuboids of around 100 nm.20 Furthermore, a similar interpretation can be applied to I−V hysteresis or perovskite solar cells depending on the crystal size. It is reported that the I−V hysteresis problem is intensified with decreasing crystal size.12 The higher density of trapped electrons in small crystals can afford to induce severe hysteresis. It was observed that there was a strong correlation between transient ferroelectric polarization induced by negative poling in the dark and I−V hysteresis enhancement in photovoltaic characteristics.27 Our results support this because the internal electric field formed by negative poling is adverse to carrier collection. Another important aspect is the retention of polarization, as observed in Figure 5. A large crystal of 700 nm sized MAPbI3 showed longer retention of polarization than smaller ones. Actually, we expected that the decay (relaxation) of the polarization would be around a few seconds according to our previous result,12 where the current response time of the capacitive current leading to I−V hysteresis was observed in the range of seconds (∼5 s) and its corresponding capacitance also responded in the low-frequency range (∼0.28 Hz). However, the observed retention time of the induced polarization by poling was much longer than expected. Therefore, we suggest that the transient current response is related to the traps for charge carriers at the interface. On the other hand, Coll et al. observed ferroelectric retention behavior in the one second regime from the similar morphology of MAPbI3.28 This definitely different retention behavior might result from the
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METHODS Synthesis of MAI. The MAI was synthesized according to a previous report.12,20 A volume of 28 mL of methylamine (40% in methanol, TCI) was added to 30 mL of hydroiodic acid (57 wt % in H2O, Aldrich), and the mixture was stirred in the ice bath for 2 h. The product was obtained by evaporating the solvent at 70 °C for 1 h and washed three times with diethyl ether and filtered. For higher purity, the product was dissolved in a small amount of ethanol (∼150 mL), and the solution containing the product was dropped into a large amount of diethyl ether (∼1000 mL) with stirring. The white precipitate was filtered and dried under vacuum overnight. Preparation of MAPbI3 Films for PFM Measurements. FTO glasses (Pilkington, TEC-8, 8Ω/sq) were cleaned with detergent, which was followed by sonication in an ethanol bath for 30 min and UV-ozone treatment in a UVO cleaner for 30 min. MAPbI3 was prepared by a sequential deposition method.20 A 0.5 M PbI2 (99%, Aldrich) solution dissolved in N,N-dimethylformamide (99.8%, Aldrich) was spin-coated on the cleaned FTO substrates. The spin-coated films were heated at 40 °C for 5 min and 100 °C for 5 min to form the PbI2 layer. A volume of 200 μL of 2-propanol solution of MAI (0.038, 0.044, and 0.063 M were used in this experiment) was dropped on the PbI2-deposited substrate with dimensions of 2.4 × 2.4 cm2, which was spun after a given loading time (30, 20, and 10 s for 0.038, 0.044, and 0.063 M, respectively). The spin-coated film was heated at 100 °C for 5 min to form MAPbI3. Surface perovskite cuboid morphologies were obtained using a field 1733
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(4) Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient Planar Heterojunction Perovskite Solar Cells by Vapour Deposition. Nature 2013, 501, 395−398. (5) Zhou, H.; Chen, Q.; Li, G.; Luo, S.; Song, T.-b.; Duan, H.-S.; Hong, Z.; You, J.; Liu, Y.; Yang, Y. Interface Engineering of Highly Efficient Perovskite Solar Cells. Science 2014, 345, 542−546. (6) National Renewable Energy Laboratory Best Research-Cell Efficiencies. http://www.nrel.gov/ncpv/images/efficiency_chart.jpg (2015). (7) Stranks, S. D.; Eperon, G. E.; Grancini, G.; Menelaou, C.; Alcocer, M. J. P.; Leijtens, T.; Herz, L. M.; Petrozza, A.; Snaith, H. J. Electron−Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341− 344. (8) Xing, G.; Mathews, N.; Sun, S.; Lim, S. S.; Lam, Y. M.; Grätzel, M.; Mhaisalkar, S.; Sum, T. C. Long-Range Balanced Electron- and Hole-Transport Lengths in Organic−Iorganic CH3NH3PbI3. Science 2014, 342, 344−347. (9) Kim, H.-S.; Mora-Sero, I.; Gonzalez-Pedro, V.; Fabregat-Santiago, F.; Juarez-Perez, E. J.; Park, N.-G.; Bisquert, J. Mechanism of Carrier Accumulation in Perovskite Thin-Absorber Solar Cells. Nat. Commun. 2013, 4, 2242. (10) Xiao, Z.; Yuan, Y.; Shao, Y.; Wang, Q.; Dong, Q.; Bi, C.; Sharma, P.; Gruverman, A.; Huang, J. Giant Switchable Photovoltaic Effect in Organometal Trihalide Perovskite Devices. Nat. Mater. 2015, 14, 193−198. (11) Juarez-Perez, E. J.; Sanchez, R. S.; Badia, L.; Garcia-Belmonte, G.; Kang, Y. S.; Mora-Sero, I.; Biszuert, J. Photoinduced Giant Dielectric Constant in Lead Halide Perovskite Solar Cells. J. Phys. Chem. Lett. 2014, 5, 2390−2394. (12) Kim, H.-S.; Park, N.-G. Parameters Affecting I−V Hysteresis of CH3NH3PbI3 Perovskite Solar Cells: Effects of Perovskite Crystal Size and Mesoporous TiO2 layer. J. Phys. Chem. Lett. 2014, 5, 2927−2934. (13) Bertoluzzi, L.; Sanchez, R. S.; Liu, L.; Lee, J.-W.; Mas-Marza, E.; Han, H.; Park, N.-G.; Mora-Sero, I.; Bisquert, J. Cooperative Kinetics of Depolarization in CH3NH3PbI3 Perovskite Solar Cells. Energy Environ. Sci. 2015, 8, 910−915. (14) Nuraje, N.; Su, K. Perovskite Ferroelectric Nanomaterials. Nanoscale 2013, 5, 8752−8780. (15) Frost, J. M.; Butler, K. T.; Brivio, F.; Hendon, C. H.; Van Schilfgaarde, M.; Walsh, A. Atomistic Origins of High-Performance in Hybrid Halide Perovskite Solar Cells. Nano Lett. 2014, 14, 2584− 2590. (16) Liu, S.; Zheng, F.; Koocher, N. Z.; Takenaka, H.; Wang, F.; Rappe, A. M. Ferroelectric Domain Wall Induced Band Gap Reduction and Charge Separation in Organometal Halide Perovskites. J. Phys. Chem. Lett. 2015, 6, 693−699. (17) Butler, K. T.; Frost, J. M.; Walsh, A. Ferroelectric Materials for Solar Energy Conversion: Photoferroics Revisited. Energy Environ. Sci. 2015, 8, 838−848. (18) Kutes, Y.; Ye, L.; Zhou, Y.; Pang, S.; Huey, B. D.; Padture, N. P. Direct Observation of Ferroelectric Domains in Solution-Processed CH3NH3PbI3 Perovskite Thin Films. J. Phys. Chem. Lett. 2014, 5, 3335−3339. (19) Gottesman, R.; Haltzi, E.; Gouda, L.; Tirosh, S.; Bouhadana, Y.; Zaban, A. Extremely Slow Photoconductivity Response of CH3NH3PbI3 Perovskites Suggesting Structural Changes under Working Conditions. J. Phys. Chem. Lett. 2014, 5, 2662−2669. (20) Im, J.-H.; Jang, I.-H.; Pellet, N.; Gratzel, M.; Park, N.-G. Growth of CH3NH3PbI3 Cuboids with Controlled Size for High-Efficiency Perovskite Solar Cells. Nat. Nanotechnol. 2014, 9, 927−932. (21) Dubourdieu, C.; Bruley, J.; Arruda, T. M.; Posadas, A.; JordanSweet, J.; Frank, M. M.; Cartier, E.; Frank, D. J.; Kalinin, S. V.; Demkov, A. A.; et al. Switching of Ferroelectric Polarization in Epitaxial BaTiO3 Films on Silicon without a Conducting Bottom Electrode. Nat. Nanotechnol. 2013, 8, 748−754. (22) Gupta, M. K.; Lee, J.-H.; Lee, K. Y.; Kim, S.-W. TwoDimensional Vanadium-Doped ZnO Nanosheet-Based Flexible Direct Current Nanogenerator. ACS Nano 2013, 7, 8932−8939.
emission scanning electron microscope (FE-SEM, JSM-7600F, JEOL). XRD patterns were measured by a HP-Thin Film XRD system (D8 Advance for high power, Bruker Corporation) using a turbo X-ray source of 18 kW with a scan rate of 2° min−1. PFM Measurement. The AFM-based investigations were carried out using a XE-100 AFM (Park Systems). The ferroelectric polarization of the halide perovskite samples was confirmed by PFM. A lock-in amplifier (Stanford Research Systems SR830) was used to detect the piezoresponse signal equipped with conductive Pt/Cr-coated silicon tips (spring constant of 3 N/m), operating in contact mode for imaging of topography and relative polarization by PFM. Samples were basically kept in the dark before measuring without any pretreatment. For ferroelectric-induced dipole orientation measurements, a poling voltage of +3 or −3 V was applied to the tip to align the dipoles in a single direction. Afterward, the PFM signal was measured. The deflection of the cantilever was set to 30 nN, and the scanning rate was 0.5 Hz. The PFM measurement was conducted at room temperature in ambient air with 20% humidity. The sample scanning area was 5 × 5 μm2 for 700 and 400 nm sized perovskite cuboids and 2 × 2 μm2 for the 100 nm sized perovskite cuboids sample. A whitelight point source (50000 lx) with a light guide GS6-1000F integrated with the AFM system was used to detect lightinduced polarization.
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected] (S.-W.K.). *E-mail:
[email protected] (N.-G.P.). Author Contributions ⊥
H.-S.K., S.K.K., and B.-J.K. contributed equally to this work.
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was supported by the National Research Foundation of Korea (NRF) grants funded by the Ministry of Science, ICT & Future Planning (MSIP) of Korea under Contracts NRF2010-0014992, NRF-2012M1A2A2671721, NRF2012M3A7B4049986 (Nano Material Technology Development Program), and NRF-2012M3A6A7054861 (Global Frontier R&D Program on Center for Multiscale Energy System). H.-S.K is grateful for the Global Ph.D. Fellow Grant funded by NRF. S.-W.K. acknowledges financial support by the Basic Science Research Program (2009-0083540) through the National Research Foundation (NRF) of Korea Grant funded by the Ministry of Science, ICT & Future Planning.
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