CsSnBr3, A Lead-Free Halide Perovskite for Long ... - ACS Publications

Oct 18, 2016 - unprecedented high power conversion efficiencies in recent years. However ... (Sn) is a possible alternative (as is germanium, Ge). How...
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CsSnBr3, A Lead-Free Halide Perovskite for Long-Term Solar Cell Application: Insights on SnF2 Addition Satyajit Gupta,† Tatyana Bendikov,‡ Gary Hodes,*,† and David Cahen*,† †

Department of Materials & Interfaces and ‡Chemical Research Support Unit, Weizmann Institute of Science, Rehovot 76100, Israel S Supporting Information *

ABSTRACT: Solar cells based on “halide perovskites” (HaPs) have demonstrated unprecedented high power conversion efficiencies in recent years. However, the well-known toxicity of lead (Pb), which is used in the most studied cells, may affect its widespread use. We explored an all-inorganic lead-free perovskite option, cesium tin bromide (CsSnBr3), for optoelectronic applications. CsSnBr3-based solar cells exhibited photoconversion efficiencies (PCEs) of 2.1%, with a short-circuit current (JSC) of ∼9 mA cm−2, an open circuit potential (VOC) of 0.41 V, and a fill factor (FF) of 58% under 1 sun (100 mW cm−2) illumination, which, even though meager compared to the Pb analogue-based cells, are among the best reported until now. As reported earlier, addition of tin fluoride (SnF2) was found to be beneficial for obtaining good device performance, possibly due to reduction of the background carrier density by neutralizing traps, possibly via filling of cation vacancies. The roles of SnF2 on the properties of the CsSnBr3 were investigated using ultraviolet photoemission spectroscopy (UPS) and X-ray photoelectron spectroscopy (XPS) analysis. ∼1.65 and 0.95%, respectively, also achieved through SnF2 modification.17 With MASnBr3 (EG ≈ 2.15 eV that was sequentially deposited without SnF2), an efficiency up to ∼1.1% was obtained when P3HT was used as a hole conductor.18 Here, we use CsSnBr3 as an active absorber for the solar cells in the “n−i−p” configuration. The solar cells with CsSnBr3 (with SnF2 addition) were optimized and tested using several “hole-selective” and “electron-selective” contacts, and a best PCE of 2.1% was obtained using 2,2′,7,7′-tetrakis(N,N-di-4methoxyphenylamino)-9,9′-spirobifluorene (spiro-MeOTAD) as a hole-transport material (HTM). To understand the fundamental role of SnF2 on CsSnBr3 as a PV material, we analyzed the structure and energetic properties (work function, WF, and top of the valence band, EVB) using XPS and UPS. The time-dependent XPS analysis revealed that pristine CsSnBr3 is susceptible to beam damage (metallic Sn formation was observed within ∼0.5 h of constant X-ray irradiation), while the addition of ∼20 mol % SnF2 improves HaP stability to such an extent that no beam damage was observed after 2.5 h of constant X-ray irradiation. The structural and optical properties of CsSnBr3 were analyzed using X-ray diffraction (XRD) and UV−visible spectroscopy. An XRD pattern of CsSnBr3 deposited over mesoporous TiO2 is shown in Figure S1. The indexed peaks correspond to the pure cubic perovskite phase of CsSnBr3.17 The optical absorption edge of the as-prepared CsSnBr3 (with 20 mol % SnF2) was analyzed using UV−visible spectroscopy,

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alide perovskite (HaP)-based solar cells have shown efficient light to electrical energy conversion. HaPs have the general composition AMX3 (A: monovalent cation, primarily methylammonium (MA+), formamidinium (FA+) or cesium (Cs+); M: most commonly Pb2+; X: halide, Br−, Cl−, I−).1−10 Some specific compositions of HaP-based cells have demonstrated power conversion efficiencies (PCEs) over 21%.11 However, from the environmental point of view, Pb, a major structural component of HaPs, is globally considered as a harmful element. This explains the interest in replacing it, and tin (Sn) is a possible alternative (as is germanium, Ge). However, Sn perovskites suffer from instability caused by ready Sn2+→ Sn4+ oxidation and, likely related, small diffusion lengths. While several Sn-based HaP solar cells have been reported, finding ways to increase the stability is an important challenge to allow one to assess the future of Sn-based HaP cells. It was observed that MASnI3 (band gap EBG of ∼1.23 eV) can be used for solar cells with efficiencies as high as 5−6%.12,13 FASnI3 (EBG of 1.41 eV) based cells were observed to generate high photocurrents (23 mA cm−2), with an optimized PCE of 2.1%.14 Addition of SnF2 was found to improve the solar cell performance of Sn-based HaPs by reducing the excessive p-type doping of the resulting HaPs caused by air oxidation of Sn2+ to Sn4+.15,16 It was also reported that SnF2 improved the morphology and substrate coverage of the Sn perovskites.14 In addition to SnF2, pyrazine addition to the perovskite precursor solution was reported to be beneficial, yielding a PCE of 4.8% for FASnI3, a result that was ascribed to homogeneous dispersion of SnF2 in the HaP layer.15 The highest reported PCEs for CsSnI3 (EG of 1.27 eV) and CsSnBr3 (EG of 1.75 eV) cells are © 2016 American Chemical Society

Received: September 1, 2016 Accepted: October 18, 2016 Published: October 18, 2016 1028

DOI: 10.1021/acsenergylett.6b00402 ACS Energy Lett. 2016, 1, 1028−1033

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during the spin-coating process. This is most likely due to our as yet limited experience with Sn perovskites compared to the Pb ones, rather than any intrinsic property of the CsSnBr3. The absorption coefficient (α) values of CsSnBr3 as a function of wavelength are given in Figure S2B. The peak that we observe at 500−520 nm for CsPbBr3 and MAPbBr3 (actually a dip in the transmission mode) could be due to an excitonic contribution, as discussed in our earlier study.2 Devices were fabricated with the following architecture: FTO/d-TiO2/mp-TiO2 or mp-Al2O3; after perovskite deposition on the substrate, one of three HTMs (4,4′-bis(N-carbazolyl)1,1′-biphenyl (CBP), poly[bis(4-phenyl)(2,4,6 trimethylphenyl)amine] (PTAA), or spiro-MEOTAD) was deposited and the device structure was completed by depositing a Au contact onto the HTM. Using mesoporous TiO2 as the electron-transport material (ETM) and spiro-MEOTAD as the HTM gave the best photovoltaic (PV) performance. The J−V characteristics of the solar cells in the dark and under illumination are shown in Figure 2A (with SnF2 added during HaP preparation) and Figure 2B (without SnF2 addition). A dramatic, ∼200 times improvement in the PCE and a large improvement in all PV parameters resulted when 20 mol % SnF2 was added. The distributions in various device parameters are shown in Figure S3 (with SnF2) and Figure S4 (without SnF2). The device cross section observed using SEM analysis with various detectors is shown in Figure S5, indicating the various layers of the cell, including the presence of the Sn perovskite in the mesoporous layer. Note that there was no capping layer of the perovskite above the TiO2 because during spin-coating of the HTM the capping layer is essentially removed, possibly due to the complexation between 4-tert-butylpyridine (tBuP, used as a dopant for the HTM) with the CsSnBr3, leading to the formation of a coordination complex, similar to what was observed in a recent study with MASnI3.19 Such a capping layer is normally necessary for highefficiency Pb-based perovskite cells, presumably to prevent shorting between the HTM (or the contact made to it) and TiO2.

Figure 1. Reflection-corrected normalized transmittance of MAPbBr3, CsPbBr3, and CsSnBr3 (∼250−300 nm thick) on d-TiO2 (∼70 nm)/ mp-TiO2 (∼450 nm). The transmittance profiles for MAPbBr3 and CsPbBr3 are taken from previous work.3

and the reflection corrected transmittance (%) spectrum is shown in Figure 1. The Tauc plot analysis (Figure S2A) indicates a direct band gap of ∼1.75 eV for CsSnBr3. The optical absorption profile of the CsSnBr3 layer (∼300 nm) was also compared to that of Pb-based bromide HaPs (CsPbBr3 and MAPbBr3, both approximately 300 nm thick), as shown in Figure 1. Besides the expected red shift in the absorption onset of the Sn perovskite due to its lower band gap, the absorption is less strong compared to that of the Pb perovskites, something that has been observed for Sn perovskites in general.12 The lower transmission of the Sn perovskite in the sub-band-gap region is due to higher scattering of these films compared to that of the Pb perovskites, only part of which is measured and corrected for by the integrating sphere. The higher scattering (lower transmission in the nonabsorbing wavelength region) of the CsSnBr3 is probably due to the inhomogeneous nature of the film, formed

Figure 2. J−V characteristics (cf. experimental section) of the best-performing CsSnBr3 [(A) with and (B) without SnF2] based cells in the dark and under illumination. PV parameters of the cells are tabulated at the bottom. [PCE: power conversion efficiency; FWD: forward; REV: reverse.] 1029

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Figure 3. Energies of the VBM (from UPS measurements) and CBM (calculated from the UPS data and the optical band gap, i.e., neglecting the exciton binding energy) and the Fermi level (from UPS data) of dense TiO2, pristine CsSnBr3, and CsSnBr3 (with 20 mol % SnF2). The HOMO level of spiro was taken from the literature.20

concentration of SnF2 between 5 and 80 wt % revealed no significant change in the WF and EVBM energies of CsSnBr3 compared to those of CsSnBr3 with 20 wt % SnF2. Also, it was found that CsSnBr3 (with 20 mol % SnF2) gave the best device performances. In addition, the conduction band minimum (CBM, calculated from the EVBM and the optical band gap, neglecting exciton binding energy) for CsSnBr3 (with and without SnF2) can be estimated, assuming that the band energies do not change when the interfaces are formed. This shows a significant electron injection barrier between the perovskite and the CBM of the TiO2 layer (0.48 eV for the pristine CsSnBr3 and a smaller barrier of 0.25 eV for the material prepared with SnF2). This barrier may, if viewed simplistically, preclude electron injection from the perovskite into the TiO2. There are two rather obvious reasons why electron injection may still occur: (1) While holes can easily be injected into the HTM, electrons will accumulate in the perovskite. This accumulation of electrons will lead to negative charging of the perovskite (i.e., the perovskite bands will move toward the vacuum level, upward in the scheme, on the electron energy scale), and this shift will continue until electrons can be injected into the TiO2 CB and an equilibrium is reached. At the same time, the greater the accumulation of electrons in the perovskite, the greater electron−hole recombination is likely to be; this means that it is preferable to minimize the need for this charging. Thus, the CsSnBr3 with SnF2, with a smaller barrier (in the dark) to electron injection from the perovskite to TiO2, should exhibit less recombination than the pristine perovskite for this reason alone. (2) Upon formation of the actual interface, a dipole may form at the TiO2/CsPbBr3 interface that changes the band offsets at the interface even in the dark. Such a possible situation is not reflected in the band diagram in Figure 3, which shows the band positions either of the isolated phases or of the surface of the perovskite and not the interface with TiO2. It is also possible that the SnF2 creates such a dipole at the interface. This is an issue that can be explored in principle by in situ photoemission spectroscopy during interface formation, an approach that we are exploring in collaboration with a synchrotron research group.

The external quantum efficiency (EQE) of the CsSnBr3 (with 20 mol % SnF2) cell in the dark and under bias light (∼3% of full sun) is shown in Figure S6. The EQE profile is consistent with the transmittance profile (Figure 1), both in terms of the wavelength of absorption onset (∼710 nm) and the gradual decrease in the EQE going to longer wavelengths due to the suboptimal thickness. The EQE shows a maximum photoconversion of ∼55% at 375 nm. The JSC, calculated by integrating the area under the curve, is 6.8 mA cm−2 (no bias light) and 7.1 mA cm−2 (∼3% of full sun bias light), which is 2 mA/cm−2 less than the directly measured Jsc (cf. Figure 2, left). Part, and possibly all, of the reason for this difference in JSC is that the intensity of the white light bias used in the EQE measurements was only 3% of sunlight due to technical reasons (note that the EQE under light bias is higher than that in the absence of light bias). However, there are a number of reasons why the EQEs measured under low-intensity monochromatic radiation are less (or sometimes even more) than those measured from I−V curves. A further possible reason is that the EQE was measured after the I−V measurement and some degradation may have occurred. Devices without any HTM gave much poorer device performance than those with a HTM, that is, a Voc of 0.04 V, ISC of 2.6 mA/cm2, and PCE of ∼0.02% (Figure S7). With mesoporous alumina (instead of TiO2) as a scaffold, a reduction in the photocurrent is observed, decreasing the PCEs (Figure S8) of the devices, while the VOC remains essentially the same as the best cells with mp-TiO2 (0.42 V), the latter contrary to expectation. It may be that the increased recombination in this cell canceled out the extra voltage expected by replacing TiO2 with Al2O3. Use of other HTMs such as PTAA (Figure S9) or CBP (Figure S10) resulted in a decrease in VOC (0.2 V for PTAA and 0.16 V for CBP) and in overall device efficiency. To see if SnF2 affects the energy levels of CsSnBr3, we carried out UPS valence band analysis of CsSnBr3 with and without SnF2 (Figure 3). Addition of SnF2 decreases the work function (WF) and, to a greater extent, the ionization potential (valence band maximum, EVBM) of CsSnBr3 (WF = 4.52 eV and EVBM = 6.1 eV) compared to pristine CsSnBr3 (WF = 4.62 eV and EVBM = 6.33 eV). This brings the VBM of the perovskite with SnF2 somewhat closer to the HOMO level of the HTM, spiro (∼5.0 eV),20 slightly decreasing the large loss of voltage at the HTM/HaP interface compared to pristine CsSnBr3. Varying the 1030

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with increasing time for the sample with SnF2 is probably due to a combination of reduced photogenerated current and increased dark current (see Figure S11) in the cell (operating at a constant 0.2 V forward bias). When the same experiment was performed (CsSnBr3 with 20 mol % SnF2) with lower light intensity (∼0.1 sun) to give a comparable current density to the cell without SnF2, the decay in photocurrent (as a function of time) was observed to be much slower (Figure 4 (2)). This suggests that the photogenerated charges can cause instability, possibly by influencing the moisture and/or oxygen degradation rate. When the operational stability of the devices (with and without SnF2) was measured under an inert atmosphere in a glovebox (O2 < 0.1, H2O < 0.1 ppm illumination with a highintensity LED ≈ 2 sun intensity), the devices were observed to be nearly stable (Figure 5 (1) and (2)), maintaining most

However, the barrier can be one cause for the large current loss in cells made with the CsSnBr3/TiO2 interface, particularly for the pristine CsSnBr3 (Figure 3). This misalignment may result in high recombination of the photoexcited charge carriers with the photogenerated holes during illumination due to buildup of electron density in the perovskite. There is also a large voltage loss across the perovskite/HTM (spiro) interface (0.9 V). Hence, the calculated voltage loss (considering both ETM/perovskite and HTM/perovskite) is 1.15 V. This implies that CsSnBr3 with an EG of 1.75 eV should result in a VOC of 0.6 V (using TiO2 as an ETM with ECB ≈ 4.1 eV and spiro as a HTM, with a HOMO level of ∼5.0 eV), although this may be reduced due to negative charging of the perovskite by accumulated electrons upon illumination. We observe a VOC of 0.4 V (Figure 2A), indicating an additional 0.2 V loss. This loss is most likely due to relatively high electron−hole recombination within the perovskite, seen from the relatively low JSC of the cells relative to the band gap and relatively low EQE (Figure S6). The results indicate that HTMs with a deeper HOMO level (than spiro) may help to minimize the voltage loss at the perovskite/HTM interface. However, we did not observe an improved VOC using PTAA (with a HOMO level of −5.2 eV)21 and CBP (with a HOMO level of −5.7 eV)22 but rather a much lower one, a result that requires further investigation. The operational stability (near the maximum power voltage) of nonencapsulated devices (with and without SnF2) was evaluated under continuous illumination in ambient and in an inert atmosphere. For both samples, a strong reduction in the photocurrent density (Figure 4 (1) and (3)) was observed in the

Figure 5. Current density measured close to the maximum power point in an inert atmosphere with continuous illumination with white LED as a function of time for (1) CsSnBr3 with 20 mol % SnF2 with 0.2 V applied bias under ∼2 sun intensity, (2) CsSnBr3 measured at 0.08 V applied bias under ∼2 sun intensity, and (3) CsSnBr3 with 20 mol % SnF2 measured at 0.08 V applied bias under ∼0.1 sun intensity.

of their initial photocurrent (∼4% decrease compared to the initial photocurrent) over a 5 h period (the results of the same experiment with on−off cycles are shown in Figure S12). When the device with SnF2 was illuminated with low-intensity white light (∼0.1 sun LED light) to give a comparable current density to the cell without SnF2, no reduction in the photocurrent was observed (Figure 5 (3)). To investigate the role of SnF2 on the stability of the CsSnBr3, time-dependent XPS measurements were carried out on CsSnBr3, (with and without SnF2) by continuously irradiating the sample with the (1.4867 keV) X-ray beam of the XPS spectrometer. The results showed that, in contrast to CsSnBr3 with SnF2 (Figure 6B), pristine CsSnBr3 was highly susceptible to beam damage (Figure 6A). For the latter, within 30 min of continuous irradiation, in addition to the expected 3d5/2 peaks for divalent Sn (∼495.5 and ∼487.5 eV), additional peaks at ∼493 and ∼484.5 eV appear (marked with circles in the figure), which correspond to zero-valent Sn. However, no zero-valent Sn signal could be detected even after 3 h of continuous X-ray irradiation for CsSnBr3 with 20 mol % SnF2. This shows that addition of SnF2 prevents Sn reduction from X-ray irradiation of the perovskite. From XPS, it was also observed that the fluoride

Figure 4. Current density measured close to the maximum power point in an ambient atmosphere with continuous illumination as a function of time for (1) CsSnBr3 with 20 mol % SnF2 at 0.2 V applied bias under 1 sun intensity (∼AM 1.5, 100 mW cm−2), (2) CsSnBr3 with 20 mol % SnF2 measured at 0.08 V applied bias with ∼0.1 sun intensity, and (3) CsSnBr3, measured at 0.08 V applied bias with 1 sun intensity (∼AM 1.5, 100 mW cm−2).

first few minutes, when measured in ambient atmosphere (∼65% RH with 1 sun intensity). The sample without SnF2 appears more stable over the longer time than that with SnF2 (Figure 4 (3)), but it should also be noted that this sample also gave, and was tested at, a much smaller photocurrent density. This strong degradation is likely due to a coupled effect of moisturemediated decomposition and oxygen-mediated oxidation of Sn2+ to Sn4+ (possibly SnO2 formation), leading to degradation of the perovskite structure. The reversal of the “photocurrent” 1031

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (D.C.). *E-mail: [email protected] (G.H.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

This research work was supported by the Israeli Ministry of National Infrastructures, Energy and Water Resources as part of an ERA-net project and by the Israel Ministry of Science and Technology’s China−Israel program. D.C. holds the Sylvia and Rowland Schaefer Chair in Energy Research.

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Figure 6. Time dependence of XP spectra for (A) pristine CsSnBr3 and (B) CsSnBr3 prepared with 20% mol SnF2, showing severe beam damage only for CsSnBr3 prepared without added SnF2.

content at the film surface was 3−4 atom % (of the total halide content) in the CsSnBr3 prepared with 20 mol % SnF2. Our analyses suggest that an important cause for the low cell voltages is that the energy levels are badly mismatched, which would mean that for CsSnBr3-based devices to become useful for solar cell application, especially better-matching selective conductors are needed. Achieving 2% PCE was possible because of SnF2 addition, which, besides the known reduction in charge carrier concentration due to reduction in Sn vacancies, was found to (1) lead to somewhat less detrimental band alignment at the HaP interfaces and (2) increase stability to X-ray beam damage compared to pristine CsSnBr3 and an apparent increase in stability under operating conditions. The operational instability in ambient air and relative stability under inert conditions suggest that, with suitable encapsulation, Sn HaP-based solar cells may become useful, either/both as stand-alone modules or/and as part of larger cell systems.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.6b00402. Experimental section, XRD analysis, Tauc plot, absorption cross section as a function of wavelength, cross-sectional SEM, J−V plots of various devices, and current densities with on−off cycles (PDF) 1032

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ACS Energy Letters (15) Lee, S. J.; Shin, S. S.; Kim, Y. C.; Kim, D.; Ahn, T. K.; Noh, J. H.; Seo, J.; Seok, S. Fabrication of Efficient Formamidinium Tin Iodide Perovskite Solar Cells through SnF2−Pyrazine Complex. J. Am. Chem. Soc. 2016, 138, 3974−3977. (16) Kumar, M. H.; Dharani, S.; Leong, W. L.; Boix, P. P.; Prabhakar, R. R.; Baikie, T.; Shi, C.; Ding, H.; Ramesh, R.; Asta, M.; et al. LeadFree Halide Perovskite Solar Cells with High Photocurrents Realized Through Vacancy Modulation. Adv. Mater. 2014, 26, 7122−7127. (17) Sabba, D.; Mulmudi, H. K.; Prabhakar, R. R.; Krishnamoorthy, T.; Baikie, T.; Boix, P. P.; Mhaisalkar, S.; Mathews, N. Impact of Anionic Br− Substitution on Open Circuit Voltage in Lead Free Perovskite (CsSnI3‑xBrx) Solar Cells. J. Phys. Chem. C 2015, 119, 1763−1767. (18) Jung, M. C.; Raga, S. R.; Qi, Y. Properties and Solar Cell Applications of Pb-Free Perovskite Films Formed By Vapor Deposition. RSC Adv. 2016, 6, 2819−2825. (19) Yokoyama, T.; Cao, D. H.; Stoumpos, C. C.; Song, T. B.; Sato, Y.; Aramaki, S.; Kanatzidis, M. G. Overcoming Short-Circuit in LeadFree CH3NH3SnI3 Perovskite Solar Cells via Kinetically Controlled Gas−Solid Reaction Film Fabrication Process. J. Phys. Chem. Lett. 2016, 7, 776−782. (20) 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, 1377−1381. (21) Ryu, S.; Noh, J. H.; Jeon, N. J.; Kim, Y. C.; Yang, W. S.; Seo, J.; Seok, S. I. Voltage Output of Efficient Perovskite Solar Cells with High Open-Circuit Voltage and Fill Factor. Energy Environ. Sci. 2014, 7, 2614−2618. (22) Zhang, T.; Liang, Y.; Cheng, J.; Li, J. A CBP Derivative as Bipolar Host for Performance Enhancement in Phosphorescent Organic Light- Emitting Diodes. J. Mater. Chem. C 2013, 1, 757−764.

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